thermal and electrical properties of graphene …...ultra-low dielectric loss of polymer/carbon...

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Thermal and Electrical Properties of Graphene-Based Polymer Nanocomposite Foams

by

SeyedMahdi Hamidinejad

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

Graduate Department of Mechanical and Industrial Engineering University of Toronto

© Copyright 2019 by SeyedMahdi Hamidinejad

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Thermal and Electrical Properties of Graphene-Based Polymer Nanocomposite Foams

SeyedMahdi Hamidinejad

Doctor of Philosophy

Graduate Department of Mechanical and Industrial Engineering

University of Toronto

2019

Abstract

Recently, multifunctional polymer-graphene nanoplatelet (GnP) composites have demonstrated

great promise as next-generation materials for energy management and storage, electromagnetic

interference (EMI) shielding and heat dissipation components in electronic industries. However,

the practical underpinning needed to economically manufacture graphene-based polymer

composites is missing. Therefore, this dissertation aims to demonstrate how some of the

challenges for efficient manufacturing of functional polymer composites, can be strategically

tackled by using supercritical fluid (SCF)-treatment and physical foaming technologies.

In this PhD research, an industrial-scale technique for in situ exfoliation and dispersion of GnP

in polymer matrices was developed and invented. This thesis also developed an in-depth

understanding of the effects of cellular structures, GnPs’ orientation, arrangement, and

exfoliation on the thermal/electrical conductivity, percolation threshold, dielectric performance,

and EMI shielding effectiveness of the graphene-based polymer composites. In particular, it was

demonstrated how SCF−treatment and physical foaming can significantly enhance thermal

conductivity of polymer-GnP composites. The SCF-treatment and physical foaming exfoliated

the GnPs in situ and microscopically tailored the nanocomposites’ structure to enhance the

thermal conductivity. The research findings in this thesis have also demonstrated that the

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introduction of foaming and microcellular structure can substantially increase the electrical

conductivity, EMI shielding effectiveness and can decrease the percolation threshold of the

polymer-GnP composites. This research also presented a facile technique for manufacturing a

new class of ultralight polymer-GnP composite foams with excellent dielectric performance.

The generation of a microcellular structure provided a unique parallel-plate arrangement of

GnPs around the cell walls. This significantly increased the real permittivity and decreased the

dielectric loss.

This dissertation developed a fundamental understanding of structure-property relationships and

new routes to microscopically engineer the structures and properties of graphene-based polymer

composites for various application such as heat management (heat sink materials), EMI

shielding, energy storage and capacitors (dielectric materials).

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Acknowledgment

First and foremost, I want to thank my supervisors Professor Chul B. Park and Professor Tobin

Filleter. It has been a great honor to be mentored by them. I appreciate all their contributions of

time, ideas, and funding to make my PhD experience productive and stimulating. The joy and

enthusiasm they have for their research was contagious and motivational for me, even during

tough times in the PhD pursuit. I am also thankful for the excellent opportunity they provided

for me to grow as a research scientist.

I have been also very fortunate to have Professor Hani Naguib (University of Toronto),

Professor Sanjeev Chandra (University of Toronto), Professor Chandra Veer Singh (University

of Toronto) and Professor Aiping Yu (University of Waterloo) for serving as my PhD

committee members. They have provided me with inspiration, advice, and support to address the

challenges of my dissertation.

I Would like to give special thanks to my colleagues and my co-authors Dr. Raymond k.M. Chu,

Dr. Biao Zhao, Lun Howe Mark, Jung Hyub Lee, Chongxiang Zhao, and Azadeh Zandieh who

spent hours helping me with experiments. I gratefully acknowledge Mr. Doug Holmyard’s help

for preparing the TEM samples and Dr. Raiden Acosta for XRD analysis. I am highly grateful to

Mrs. Kara Kim (MPML’s Assistant Director) for all help and support.

I am also grateful to NanoXplore Inc., Montreal, QC for financial support and donation of

materials throughout my PhD studies. I would like to particularly thank Dr. Nima Moghimian

(NanoXplore Inc.) for his suggestions, guidance and his help for preparing polymer-graphene

masterbatches. I gratefully acknowledge funding from Natural Sciences and Engineering

Research Council of Canada’s (NSERC) Alexander Graham Bell Canada Graduate Scholarship

Program and the Ontario Graduate Scholarship (OGS).

Lastly, I would like to deeply thank my family for all their love and encouragement. Specially

for my deceased mother and my dear father who raised me with a love of science and supported

me in all my pursuits. And most of all for my loving, supportive, encouraging, and patient wife

Nooshin whose faithful support during my 2nd PhD is so appreciated.

This thesis is dedicated to My lovely, dear Nooshin, to the soul of my beloved Mother and to my

dear Father.

Mahdi Hamidi

University of Toronto

January 2019

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Contributions of Co-Authors

I am the principle or co-principle author of all the articles and patent for which this thesis is

partially based. I have performed all or the majority of the experiments, data collection, analysis,

and manuscript preparation for each article and patent. I would like to acknowledge all of my

co-authors for their valuable contribution. Also, I acknowledge my supervisors, Prof. Chul B.

Park and Prof. Tobin Filleter, who provided knowledge, ideas and directions that were pivotal in

my entire research and they both were involved in manuscripts’ preparation. The following

outlines the contributions of each co-author:

Patents and copyrights submitted

[1] Hamidinejad, S.M., Park, C.B., and Nazarpour, S., (2017) Method of Exfoliating and

Dispersing High Concentration Graphene Nanoplatelets (GnP) into Polymeric Matrices Using

Supercritical Fluid (SCF), applied for US Provisional Patent, Application Serial No. 62/512,790

▪ C.B. Park and S. Nazarpour were involved in patent preparation. Discussions regarding the

research concepts and presentation of data and claims were continuously conducted with all

contributing authors.

Articles Published or Accepted in Refereed Journals

[1] Hamidinejad, S.M., Chu, R.k.M., Zhao, B., Park, C.B., and Filleter, T. (2018) Enhanced

thermal conductivity of graphene nanoplatelet-polymer nanocomposites fabricated via

supercritical fluid assisted in-situ exfoliation, ACS Applied Materials and Interfaces, 10 (1):

1225−1236

▪ R.K.M. Chu and B. Zhao assisted in performing foam injection molding and sample

preparation. C.B. Park and T. Filleter were involved in manuscript preparation. Discussions

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

all contributing authors.

[2] Hamidinejad, S.M., Zhao, B, Zandieh, A., Moghimian, N., Filleter, T., and Park, C.B.

(2018) Enhanced electrical and electromagnetic interference shielding properties of polymer-

graphene nanoplatelet composites fabricated via supercritical-fluid treatment and physical

foaming, ACS Applied Materials and Interfaces, 10 (36): 30752–30761

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▪ B. Zhao performed electromagnetic shielding measurements. A. Zandieh assisted in sample

preparation through foam injection molding. N. Moghimian performed melt-mixing and

masterbatch preparation. T. Filleter and C.B. Park were involved in manuscript preparation.

Discussions regarding the research concepts and presentation of data were continuously

conducted with all contributing authors.

[3] Hamidinejad, S.M., Zhao, B., Chu, R.k.M., Moghimian, N., Naguib, H. E., Filleter, T., and

Park, C.B. (2018) Ultralight microcellular polymer-graphene nanoplatelet foams with enhanced

dielectric performance, ACS Applied Materials and Interfaces, 10 (23): 19987–19998

▪ B. Zhao and R.K.M. Chu assisted in performing extrusion foaming and sample preparation.

H.E. Naguib, T. Filleter and C.B. Park were involved in manuscript preparation. Discussions

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

all contributing authors.

[4] Zhao, B.†, Hamidinejad, S.M.†, Zhao, C., Li, R., Wang, S., Kazemi, Y., and Park, Chul B.

(2019) A versatile foaming platform to fabricate unprecedentedly high dielectric permittivity,

ultra-low dielectric loss of polymer/carbon composites, Journal of Materials Chemistry A, 7

(1), 133-140, DOI: 10.1039/C8TA05556D (†Equal Contribution)

▪ B. Zhao performed the electrical conductivity and dielectric performance measurements. C.

Zhao assisted in performed batch-foaming. R. Li, S. Wang and Y. Kazemi performed XRD

and FTIR analysis. B. Zhao and C.B. Park were involved in manuscript preparation.

Discussions regarding the research concepts and presentation of data were continuously

conducted with all contributing authors.

http://mpml.mie.utoronto.ca/info/personnel/yasamin-kazemi/

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Contents

Abstract .......................................................................................................................................... ii

Acknowledgment .......................................................................................................................... iv

Contributions of Co-Authors ......................................................................................................... v

List of Figures .............................................................................................................................. xii

List of Tables ............................................................................................................................. xvii

Nomenclature ............................................................................................................................ xviii

CHAPTER 1: Introduction ........................................................................................................... 1

1.1 Motivation of the Thesis ................................................................................................. 3

1.2 Scope of the Thesis ......................................................................................................... 5

Bibliography .............................................................................................................................. 6

CHAPTER 2: Background & Literature Review ....................................................................... 13

2.1 Summary ....................................................................................................................... 13

2.2 Introduction ................................................................................................................... 13

2.3 Bottom-Up Graphene .................................................................................................... 14

2.4 Top-Down Graphene .................................................................................................... 14

2.4.1 Direct Exfoliation of Graphite ............................................................................... 15

2.4.2 Graphite Oxide (GO) ............................................................................................. 16

2.4.3 Chemical Reduction of GO ................................................................................... 17

2.4.4 Thermal Exfoliation and Reduction....................................................................... 17

2.5 Preparation of Graphene-Based Polymer Nanocomposites .......................................... 18

2.5.1 Filler Dispersion Methods ..................................................................................... 19

2.6 Thermoplastic Polymer Foams ..................................................................................... 20

2.6.1 Physical Foaming................................................................................................... 22

2.7 Functional Properties of Polymer Composites ............................................................. 22

viii

2.7.1 Thermal Conductivity of Polymer Composites ..................................................... 23

2.7.2 Electrical Conductivity of Polymer Composites ................................................... 25

2.7.3 Electromagnetic Interference (EMI) Shielding ..................................................... 28

2.7.4 Dielectric Properties of Polymer Composites ....................................................... 31

Bibliography ............................................................................................................................ 34

CHAPTER 3: Development of a Facile Technique to in situ Exfoliate and Disperse Graphene

Nanoplatelets in Polymer Matrices .............................................................................................. 48

3.1 Summary ....................................................................................................................... 48

3.2 Introduction ................................................................................................................... 49

3.3 Experimental Section .................................................................................................... 50

3.3.1 Materials and sample preparation .......................................................................... 50

3.4 Characterization ............................................................................................................ 55

3.5 Results and discussion .................................................................................................. 55

3.5.1 Effect of SC-N2-treatment and physical foaming on GnP’s exfoliation and

dispersion in an injection molding process .......................................................................... 55

3.5.2 Effect of the SC-CO2-treatment physical foaming on the GnP’s exfoliation and

dispersion in extrusion foaming........................................................................................... 59

3.6 Conclusion .................................................................................................................... 61

Bibliography ............................................................................................................................ 62

CHAPTER 4: Enhancement of the thermal conductivity of polymer-GnP composites via facile

SCF-assisted manufacturing ........................................................................................................ 65

4.1 Summary ....................................................................................................................... 65

4.2 Introduction ................................................................................................................... 66



4.4 Characterization ............................................................................................................ 70

4.5 Results and discussion .................................................................................................. 72

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4.5.1 Defect density of GnPs .......................................................................................... 72

4.5.2 Microstructure and morphology of polymer-GnP composites .............................. 73

4.6 Thermal conductivity .................................................................................................... 77

4.6.1 Effect of the GnP content on the thermal conductivity ......................................... 77

4.6.2 Effect of GnP’s exfoliation and dispersion on the thermal conductivity............... 78

4.6.3 Effects of GnPs’ re-orientation on the thermal conductivity of HDPE-GnP

composites ........................................................................................................................... 80

4.6.4 Optimal degree of foaming on the thermal conductivity ....................................... 82

4.6.5 Solid phase thermal conductivity .......................................................................... 83

4.7 Conclusion .................................................................................................................... 86

Bibliography ............................................................................................................................ 87

CHAPTER 5: Enhancement of electrical and electromagnetic interference (EMI) shielding

properties of the polymer-GnP composites ................................................................................. 94

5.1 Summary ....................................................................................................................... 94

5.2 Introduction ................................................................................................................... 95



5.3.2 Characterization ................................................................................................... 100

5.4 Results and Discussion ............................................................................................... 101

5.4.1 Microstructure and morphology of the HDPE-GnP composites ......................... 101

5.4.2 The effect of physical foaming on the GnP’s exfoliation and dispersion............ 102

5.4.3 The electrical conductivity of the polymer-GnP composites .............................. 103

5.4.4 The dielectric properties of polymer-GnP composites ........................................ 107

5.4.5 The EMI shielding effectiveness (SE) of the polymer-GnP composites ............. 110

5.5 Summary & Conclusions ............................................................................................ 114

Bibliography .......................................................................................................................... 114

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CHAPTER 6: Enhancement of the dielectric performance of polymer-GnP composites using

SCF-treatment and physical foaming-Part I ............................................................................. 122

6.1 Summary ..................................................................................................................... 122

6.2 Introduction ................................................................................................................. 123

6.3 Experimental Section .................................................................................................. 126

6.3.1 Materials and sample preparation ........................................................................ 126



6.4.1 Microstructure and morphology of the polymer-GnP composites ...................... 128

6.4.2 Electrical conductivity of the polymer-GnP composites ..................................... 131

6.4.3 Dielectric properties of the polymer-GnP composites......................................... 133

6.5 Conclusion .................................................................................................................. 143

Bibliography .......................................................................................................................... 143

CHAPTER 7: Enhancement of the dielectric performance of polymer-GnP composites using

SCF-treatment and physical foaming-Part II ........................................................................... 150

7.1 Summary ..................................................................................................................... 150

7.2 Introduction ................................................................................................................. 151

7.3 Experimental Section .................................................................................................. 153

7.3.1 Materials .............................................................................................................. 153

7.3.2 Fabrication of PVDF-GnP Solid Composites ...................................................... 153

7.3.3 Fabrication of PVDF-GnP Composite Foams ..................................................... 154



7.5 Conclusion .................................................................................................................. 160

Bibliography .......................................................................................................................... 160

CHAPTER 8: Contributions and Future Work ........................................................................ 164

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8.1 Contributions ............................................................................................................... 164

8.2 Future Work ................................................................................................................ 166

8.2.1 Thermal and Electrical Conductivities of Graphene-Based Polymer Composites

with the Geometrical Characteristics of GnPs ................................................................... 166

8.2.2 The Development of the Thermally Conductive Graphene-Based Polymer

Composites with High Thermal Stability .......................................................................... 167

8.2.3 SCF-Assisted Manufacturing of Hexagonal Boron Nitride (hBN)-Polymer

Composites with Enhanced Thermal Conductivity ........................................................... 167

8.2.4 Generalizing the SCF-Assisted Exfoliation Method to Other 2D Materials ....... 168

8.2.5 Fatigue Behavior of Graphene-Based Nanocomposite........................................ 168

8.2.6 3D Nanostructured Graphene for Heat Management in Microelectronic Devices

169

8.2.7 Development of Lightweight Superthermal Insulation Graphene-Based

Nanocomposites ................................................................................................................. 170

8.2.8 Fabrication of 3D Architected Nanostructures of 2D Materials .......................... 170

xii

List of Figures

Figure 2.1. Different top-down methods for producing graphene or functionalized graphene

from graphite or GO. ................................................................................................................... 15

Figure 2.2. Surface chemistry of GO containing containing carboxyl, poxide and hydroxyl

groups and double bonds. ............................................................................................................ 16

Figure 2.3. SEM of thermally reduced GO adapted from Ref. [33]. .......................................... 18

Figure 2.4. Physical foaming of thermoplastic polymer consisting of these steps of: (i) the

dissolution of blowing agent into the polymer matrix and formation of single-phase gas/polymer

mixture; (ii) phase separation of gas due to thermodynamic instability; and (iii) curing when the

blowing agents are replaced with the ambient air. ...................................................................... 21

Figure 2.5. Percolation curve of compression-molded HDPE-GnP composite (A typical

percolation curve of conductive polymer composites). ............................................................... 26

Figure 2.6. Diagram of electron-transfer mechanisms between adjacent sites separated by a

potential energy barrier. Adapted from Ref. [84]. ....................................................................... 27

Figure 2.7. Schematic of shielding mechanisms of a plane wave by a shielding material.

Adapted from Ref. [94]................................................................................................................ 30

Figure 2.8. Polarization of a dielectric material by an applied electric field. ............................. 32

Figure 2.9. Real (ε') and imaginary (ε'') parts of permittivity as a function of frequency for a

material showing interfacial, orientational, ionic, and electronic polarization. Adapted from Ref.

[111]. ............................................................................................................................................ 33

Figure 3.1. The injection molding processes (i.e. IMS, HPIMF and IMF) and injection-molded

parts.............................................................................................................................................. 53

Figure 3.2. The schematic of the extrusion process and processing parameters ........................ 54

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Figure 3.3. (a) XRD spectra of neat HDPE, GnP powder, IMS samples (HDPE-9 vol.% GnP)

and their HPIMF and IMF counterparts with various degrees of foaming; (b) magnified XRD

pattern of Figure 3.3a over 2θ=40°-50° highlighted with light green, to examine (100)

diffraction peaks and illustration of the GnPs’ orientation and their effect on the (002) and (100)

diffraction peaks of the XRD pattern; (c) residual values (%) of I(002) (intensity of the (002)

diffraction at 2θ = 26.5°) before and after SCF-treatment and physical foaming; (d)

representative TEM micrographs of the IMS of HDPE-4.5vol.% GnP and; (e) IMF of HDPE-

4.5vol.% GnP; (f) ideal conceptualization of various phenomenon resulting in further exfoliation

and dispersion of GnPs in IMF samples. DF stands for degree of foaming. ............................... 57

Figure 3.4. (a) Representative TEM micrographs of the SCM of the HDPE-4.5vol.% GnP and;

(b) Foam-extruded 4.5 vol.% HDPE-GnP; (c) XRD spectra of neat HDPE, GnP powder, SCM

samples (4.5 vol.% HDPE-GnP), and their extruded-foam counterparts with different densities

..................................................................................................................................................... 60

Figure 3.5. Ideal conceptualization of various phenomenon resulting in further exfoliation and

and parallel-plates arrangement of the GnPs in the extruded foam samples ............................... 61

Figure 4.1. The schematic of the ISO/DIS 22007-2.2 setup for measuring the thermal

conductivity using TPS 2500 ....................................................................................................... 71

Figure 4.2. (a) Raman spectroscopy of the GnPs; (b) deconvoluted C 1s XPS spectra. Raman

spectra of GnP.............................................................................................................................. 72

Figure 4.3. (a) SEM micrographs of skin and core regions for IMS, HPIMF and IMF HDPE-9

vol% GnP nanocomposites. Scale bars are all 10 μm; (b) ideal 2-D conceptualization of the

evolution of GnPs interconnectivity, orientation and further exfoliation due to SCF-treatment

and physical foaming; (c) SEM micrographs of IMF HDPE-9 vol% GnP nanocomposites

showing different types of cells generated in the microstructure. FD stands for flow direction. 75

Figure 4.4. (a) SEM micrographs of the FIM -9 vol% HDPE-GnP composites with 7% degree

of foaming; (b) Zoomed-in SEM micrographs of Figure 4.4a. ................................................... 76

Figure 4.5. (a) The total thermal conductivity (λtotal) of IMS, HPIMF, and IMF HDPE-GnP

composites as a function of the GnP content and; (b) the thermal conductivity of IMS, HPIMF,

xiv

and IMF samples (HDPE-9 vol.% GnP) before (total) and after removing their skin (core); (c)

the total thermal conductivity (λtotal) of IMS, HPIMF, and IMF HDPE-GnP composites as a

function of the degree of foaming and the GnP content; (d) the total thermal conductivity (λtotal)

of the samples as a function of the degree of foaming (GnP vol. % has been reported with

respect to the polymer volume) ................................................................................................... 79

Figure 4.6. (a) Differential Scanning Calorimetry (DSC) of the IMS, HPIMF and IMF sample

(HDPE-4.5 vol.% GnP); and (b) High Pressure Differential Scanning Calorimetry (HPDSC) of

HDPE-4.5 vol.% GnP samples .................................................................................................... 82

Figure 4.7. Solid phase thermal conductivity (ksolid) of IMS, and IMF HDPE-GnP composites as

function of (a) the GnP content and; (b) the degree of foaming and the GnP content. DF stands

for degree of foaming .................................................................................................................. 85

Figure 5.1. The schematic of the injection molded parts and the location of cut samples ......... 99

Figure 5.2. (a) SEM micrographs of the skin and core regions of the solid and foamed (16 %

degree of foaming) HDPE-GnP composites at 9.8 vol % GnP content, and (b) Ideal

conceptualization of the GnPs’ arrangement in the solid and foamed samples. The arrow shows

the melt’s flow direction in the injection-molding process. ...................................................... 102

Figure 5.3. (a) XRD spectra of neat HDPE, GnP powder, solid, foamed samples with 4.5 vol.%

GnP. The inset figure (a) shows an ideal conceptualization of the SCF-assisted exfoliation of the

GnPs in the foamed samples. (b) Representative TEM micrographs of the foamed and (c) solid

samples of the HDPE-4.5vol.% GnP ......................................................................................... 103

Figure 5.4. (a) The AC conductivity of the solid, and foamed HDPE-GnP composite; and (b)

The DC conductivity of the solid, and foamed HDPE-GnP composite measured at 0.1 Hz

(degree of foaming of foamed samples is 16%) ........................................................................ 104

Figure 5.5. (a) Variations of the foaming degree on the electrical conductivity of the HDPE-

GnP composites; (b) The evolution of the percolation threshold with the foaming degree ...... 107

Figure 5.6. (a) Real dielectric permittivity (ε'); and (b) The dielectric loss (tan δ) of the solid

and foamed (16% degree of foaming) nanocomposites as a function of the GnP content

measured at 1×10+3 Hz. (GnP vol.% is reported in relation to the polymer volume) .............. 108

xv

Figure 5.7. Broadband dielectric permittivity of (a) The solid samples, and (b) The foamed 9.8

vol.% HDPE-GnP composites. Broadband dielectric loss of (c) The solid samples, and (d) The

foamed 9.8 vol.% HDPE-GnP composites ................................................................................ 110

Figure 5.8. K-band EMI SE of (a) the solid; and (b) the foamed HDPE-GnP composites with

various GnP content. .................................................................................................................. 111

Figure 5.9. (a) The K-band EMI SE of the solid and foamed HDPE-GnP composites as a

function of their GnP content; (b) The contributions of the reflection and absorption

mechanisms to the total K-band EMI SE of the solid and foamed HDPE-GnP composites as a

function of their GnP content; (c) schematic diagrams of the scattering and multiple reflections

of the electromagnetic waves..................................................................................................... 112

Figure 6.1. The SEM micrographs of the (a) as-received GnP powder; (b) SCM HDPE-9.8

vol.% GnP composites; (c)-(d) Foam-extruded nanocomposites counterparts; and (e) TEM of

extruded foam samples showing parallel-plates arrangements of GnPs within the cell walls .. 129

Figure 6.2. Representation of the density of HDPE-GnP composite foams vs the foaming

temperature, together with the related SEM micrograph. The scale bar is 300 m. (GnP vol.%

was reported with respect to the polymer volume) .................................................................... 130

Figure 6.3. (a) Broadband conductivity of the SCM and the extruded HDPE-GnP composite

foams. The extruded foam samples had 0.14±0.01 g.cm-3 (corresponding to ~8 times the foam

expansion ratio); (b) The DC conductivity of the SCM and extruded HDPE-GnP composite

foams as a function of the GnP content measured at 0.1 Hz (X-axis is logarithmic and scales

before and after break are not equidistant). Note that the extruded foams of the 12, 15 and 19

vol.% samples could not be obtained due to the excessive viscosity as discussed in Section

5.4.1. .......................................................................................................................................... 132

Figure 6.4. (a) Real dielectric permittivity (ε'); and (b) Dielectric loss (tan δ) of the extruded

foam (with the density of 0.14±0.01 g.cm-3 or ~8 times foam expansion ratio) and the SCM

HDPE-GnP composites as a function of GnP content measured at 1×10+3 Hz. Note: X-axis is

logarithmic and scales before and after break are not equidistant. ............................................ 135

xvi

Figure 6.5. (a) Broadband dielectric permittivity; (b) Broadband dielectric loss of the SCM

HDPE-9.8 vol.% GnP composites and their extruded foam (with a density of 0.15 g.cm-3 or ~8

times foam expansion ratio) counterparts. ................................................................................. 137

Figure 6.6. Variations in real permittivity and dielectric loss measured at 1×10+3 Hz as a

function of density in the extruded HDPE-GnP composite foams made from solid precursors

containing 4.5 vol% GnP ........................................................................................................... 138

Figure 6.7. SEM and TEM micrographs of the extruded HDPE-GnP composite foams made

from solid precursors containing 4.5 vol% GnP, which show the GnPs’ arrangement at different

densities including: (a) 0.13 g.cm-3; (b) 0.08 g.cm-3; and (c) 0.05 g.cm-3. (d) Ideal 2-D

conceptualization of GnP’s arrangement in cell walls as the density decreased. ...................... 140

Figure 7.1. A schematic illustration of the home-made batch-foaming device ........................ 154

Figure 7.2. A schematic diagram of the PVDF-GnP foam fabrication process ........................ 155

Figure 7.3. (a) Expansion ratio of PVDF/GnP composite foams; (b) SEM image of FG3 foam

sample, and the inset is the corresponding magnification SEM; the cell density of PVDF/GnP

................................................................................................................................................... 157

Figure 7.4. (a) Frequency-dependent electrical conductivity of the solid and foamed PVDF/GnP

composites, (b) Real permittivity, (c) Imaginary permittivity, and (d) Dielectric loss of the solid

and foamed PVDF/GnP composites as a function of applied frequencies ranging from 1 Hz to

300,000 Hz, (e) Real permittivity and dielectric loss of the solid and foamed PVDF/GnP

composites in the 100 Hz frequency, (f) The correlation amongst the real permittivity, the

dielectric loss and the expansion ratio of the foamed PVDF/GnP composites in the 100 Hz

frequency. .................................................................................................................................. 158

xvii

List of Tables

Table 3.1. Processing parameters used in injection molding of solid and foamed composites .. 52

Table 4.1. Thermal conductivity of various batch-type graphene/polymer nanocomposites ..... 69

Table 5.1. Processing parameters used in injection molding of solid and foamed composites .. 99

Table 6.1. Dielectric performance and density of different polymer nanocomposites ............. 141

Table 7.1. Expansion ratio of PVDF-2wt% GnP foams obtained at various saturation

temperatures ............................................................................................................................... 155

xviii

Nomenclature

1D One-Dimensional

2D Two-Dimensional

3D Three-Dimensional

BNNT Boron Nitride Nanotubes

CNT Carbon Nanotubes

dB Decibels

DF Degree of Foaming

DMF (N,N)-Dimethylformamide

DSC Differential Scanning Calorimetry

EMI Electromagnetic Interference

FD Flow Direction

FG Foamed PVDF-GnP composite

GnP Graphene Nanoplatelet*1

hBN Hexagonal Boron Nitride

HDPE high-density-polyethylene

HPDSC High-Pressure Differential Scanning Calorimetry

HPIMF High-Pressure-Injection-Molded Foam

IAA Infrared Attenuated Agent

IMF Injection-Molded Foam

IMS Injection-Molded Solid

MWCNT Multi-Walled Carbon Nanotubes

MWS Maxwell–Wagner–Sillars

PVDF Poly(Vinylidene Fluoride)

1 * It is notable that, based on the recommended nomenclature for 2D carbon materials by

Bianco et al. (Carbon, 65 (2013) 1–6) the filler used in this study is graphite nanoplates (GNP).

However, the commercial name of this filler (i.e. graphene nanoplatelets (GnP)) introduced by

the manufacturer (NanoXplore Inc.) was used.

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SC-CO2 Supercritical CO2

SCF Supercritical Fluid

SCM solid compression molded

SC-N2 Supercritical Nitrogen

SE Shielding Effectiveness

SEM Scanning Electron Microscopy

tan Dielectric Loss

TEM Transmission Electron Microscope

TPS Transient Plane Source

WAXD Wide Angle X-ray Diffraction

XPS X-ray Photoelectron Spectroscopy

XRD X-ray Diffraction

ε' Real Permittivity

ε'' Imaginary Permittivity

λgas Gas Thermal Conductivity

λsolid Solid Phase Thermal Conductivity

λtotal Total Thermal Conductivity

σAC Alternative Current Conductivity

σDC Direct Current Conductivity

1

CHAPTER 1

1 Introduction

Next-generation, multifunctional materials are considered to be the foundation for technological

innovations in the 21st century. By combining science with specialized engineering knowledge,

research on advanced functional materials will enable the design and development of cutting-

edge, multifunctional, lightweight, and high-performance materials for a wide variety of

applications that can be used in the automotive, aerospace, telecommunication, energy, and

microelectronics industries. Over the last decades, conductive polymer composites have shown

great potential as a highly-desirable class of advanced functional materials. They have an

attractive array of properties. These include their light weight, ease of processing, non-linear

voltage-current behaviour, and environmentally-sensitive resistivity [1–5]. These features have

led to their use in numerous energy storage applications, such as in capacitors and super-

capacitors [6,7], energy conversion (bipolar plates of fuel cells [8,9]), electromagnetic

interference shielding [10–14], and electrostatic discharge [15,16]. Further, the global

electromagnetic Interference (EMI) shielding market is expected to grow from $2.3 billion in

2009 to $3.8 billion in 2016 [17]. Compared with metallic and ceramic composites, polymer

composites have superior resistance to chemicals and corrosion. They also use inexpensive

materials and processing methods while providing higher specific toughness and ductility, and

lighter weight [1,2,12–16,3,4,6–11].

For these reasons, research and development on polymer composites and nanocomposites have

received intense attention from a variety of industries. Polymer composites, in particular

nanocomposites, can substantially enhance mechanical, thermal, and electrical performances.

And they do so at much lower filler concentrations than the conventional micro-size additives

2

such as graphite and carbon fibers. This eventually lowers the component weight and enhances

their processability. Moreover, tailorable mechanical and functional property enhancements

achieved by the incorporation of different fillers in polymer nanocomposites can help to address

the requirements of a broad range of cutting-edge applications.

In the last decade, the emergence of microcellular plastics (i.e., plastics with micron-size cells)

has expanded the market for plastic products to include high-value industrial and consumer

applications like thermal insulation, soundproofing, impact absorption and safety equipment,

and packaging products. The improved mechanical properties (e.g., impact strength, fatigue, and

toughness) achieved with microcellular foaming have also made plastics attractive to many

types of industry. Given this context, the multi-functional micro-porous structures of

microcellular materials would make polymer nanocomposites integral to the design of

innovative products (e.g., filtration membranes and fuel cells [8,9]). Industries can benefit from

the cost, material, and weight saving that accompanies microcellular polymer nanocomposite

technology. Thus, they would be able to develop lightweight multifunctional parts and

components with tailored electrical, thermal and mechanical properties that can be applied in

virtually every industrial sector.

With the recent advances in nano-materials and technologies, the types of fillers available for

polymer composites have been significantly developed and their functions have been increased.

One example of the filler candidates now emerging is graphene. It is an atomically thick layer

composed of sp2 carbon atoms that have formed a two-dimensional (2D), honeycomb-structured

lattice. In recent years, graphene has attracted great attention due to its exceptional mechanical,

electrical, and thermal properties. Notably, it is the strongest material ever measured with an

ultimate strength of 130 GPa and Young’s modulus of 1 TPa. The thermal and electrical

conductivities of single-layer graphene have been respectively reported 5,000 W/(m.K) and

6,000 S/cm [18]. However, the practical underpinning needed to economically manufacture

graphene-based polymer composites is missing. This has been due to the complexities that exist

in the exfoliation, dispersion, and control of the graphene nanoplatelets’ (GnP) orientation

within the composites [19]. Various strategies, such as in-situ polymerization [20,21], GnP

surface modification [22,23], GnP alignment by electrical field [24], and the use of hybrid

additives [22,25] have all been proposed to develop conductive polymer composites.

3

1.1 Motivation of the Thesis

Polymer nanocomposites are a new emerging class of advanced materials. They have unique

physical and mechanical properties tailored for a broad range of applications in diverse areas.

Novel characteristics (e.g., thermal, and electrical) of polymer nanocomposites can be derived

from suitable combinations of the properties of parent constituents. Hence, the research and

development of nanocomposites have been of great importance for application areas such as

electronics, sensors, computing, and biomedical materials, where miniaturized and lightweight

components play critical roles [26].

Unlike the batch-type synthesis methods [27–32] injection molding and extrusion processes are

economically viable and continuous methods to manufacture polymer composites. When these

methods are combined with supercritical fluid (SCF) treatment and physical foaming, another

layer of flexibility is added, which can which could further customize the functionalities of

conductive polymer composites for a broader range of structural, energy, and irradiation

shielding applications [33,34,43–46,35–42]. Thus, in-depth understanding of the structure-

function relationships is essential to facilitate the development of advanced graphene-based

polymer composites and foams.

Therefore, the objective of this thesis is (i) to develop a SCF-assisted manufacturing of the

graphene polymer nanocomposites with enhanced functionalities; and (ii) to develop a

fundamental understanding on how the graphene-based polymer composite foams can achieve

high functionalities (i.e. thermal/electrical conductivity, electric performances and EMI

shielding effectiveness). Achieving higher functionalities and electrical and thermal

conductivity at lower filler content has always been challenging in manufacturing of polymer

composites [22,39]. Therefore, the current research pointes towards the further development of

lightweight and functional graphene-based polymer nanocomposites with enhanced

functionalities at lower filler contents using facile and industrially-viable manufacturing

techniques. Thus, studies in this thesis has addressed the following objectives:

• Development of a facile technique to in situ exfoliate and disperse GnPs into polymeric

matrices

Exploiting the full potential of GnPs in polymer nanocomposites is highly challenging due to

the complexities that exist in the exfoliation and dispersion[19]. Supercritical Fluid (SCF)-

4

treatment and physical foaming show a great promise to further exfoliate and disperse the GnPs

in situ [36,41].

• Enhancement of the thermal conductivity of polymer-GnP nanocomposites via facile SCF-

assisted manufacturing

The development of new generations of smaller, lighter and more powerful electronic devices

requires more compact and lighter heat sinks. Heat dissipation functionality is extremely critical

in high-energy density systems such as next-generation miniaturized electronic devices [47].

Multifunctional, highly thermally conductive polymer composites show promise as candidates

to be employed as heat dissipation components in the electronic packaging technology. This

thesis presents a facile SCF-assisted manufacturing method for producing thermally conductive

polymer-GnP composites high heat dissipation functionality.

• Enhancement of electrical and electromagnetic interference (EMI) shielding properties of the

polymer-GnP nanocomposites

One major class of graphene-based polymer nanocomposites are those which take advantage of

the electron transport characteristics of graphene for applications such as EMI shielding, where

the focus has been on achieving a higher electrical conductivity at lower graphene

concentrations [48]. EMI shielding of radio frequency radiation is a serious concern in our

technological society and graphene has attracted great attention for the fabrication of efficient

EMI shields [32,49–51]. This thesis demonstrates that the introduction of a microcellular

structure can substantially increase the electrical conductivity, EMI shielding effectiveness (SE)

and can decrease the percolation threshold of the polymer-GnP composites.

• Enhancement of the dielectric performance of polymer-GnP composites using SCF-treatment

and physical foaming

High performance dielectric materials are vital to the development of next-generation

miniaturized electronic devices. Dielectric materials with high dielectric permittivity (ε') and

low dielectric loss (tan ) have been receiving increasing interest in modern electronics as the

capacitors and integrated capacitors [41,52–55]. This thesis presents an industrially-viable

technique for manufacturing a new class of ultralight polymer composite foams using

commercial GnPs with excellent dielectric performance is presented.

5

1.2 Scope of the Thesis

This thesis focuses on the SCF-assisted manufacturing of the graphene-based polymer

composites with functional properties. This includes thermal and electrical properties (e.g.

electrical conductivity, dielectric performance and EMI shielding effectiveness) as heat sinks for

electronic devices, dielectric materials for capacitors and lightweight electrically conductive

composites for EMI shielding.

Chapter 2 presents a background of the recent literatures on graphene, graphene-based polymer

composites, and methods of their fabrication. Fundamentals of polymer composites’

functionalities (i.e. thermal, electrical and dielectric) are also reviewed.

Chapter 3 presents a novel technique for in situ exfoliating and dispersing GnPs within polymer

matrix via SCF-treatment and physical foaming.

In Chapter 4, the thermal properties of polymer-GnP fabricated using SCF-treatment and foam

injection molding is discussed. The effects of GnPs’ exfoliation and dispersion, their orientation

and interconnectivity on thermal conductivity caused by foaming are studied and compared with

regular injection-molded counterparts.

In Chapter 5 it is demonstrated that how the introduction of a microcellular structure can

substantially increase the electrical conductivity and EMI shielding effectiveness (SE) and can

decrease the percolation threshold of the polymer-GnP composites.

In Chapters 6 and 7, the dielectric properties of polymer-GnP fabricated using extrusion

foaming and bath foaming are discussed. The effects of GnPs’ orientations and arrangements

developed via foaming on electrical conductivity, real permittivity and dielectric loss are

scientifically studied.

In Chapter 8, contributions and future work are discussed. In particular, the SCF-assisted

exfoliation of other 2D materials (e.g. hexagonal boron nitride (hBN) and transition metal

dichalcogenides (i.e. MoS2, WS2 and WSe2), the development of 3D nanostructured graphene,

6

and the creation of 3D architected nanostructures of 2D materials, are presented as worthy areas

of future investigation.

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13

CHAPTER 2

2 Background & Literature Review

2.1 Summary

In this chapter, a background of the recent literatures on graphene and graphene-based polymer

composites is presented. A variety of methods used to fabricate graphene-based polymer

composites are reviewed, along with methods for exfoliating and dispersing graphene

nanoplatelets (GnP) in various polymer matrices. The functional properties of the polymer

composites containing conductive fillers such as thermal, electrical, dielectric and EMI

shielding effectiveness are also reviewed.

2.2 Introduction

Graphene, single atomic layer of sp2-bonded carbon atoms, is an emerging nanomaterial with

excellent electrical, thermal and mechanical properties [1–3]. Graphene has a thermal

conductivity of 5,000 W/(m.K) and electrical conductivity of 6,000 S/cm [4]. Furthermore,

graphene is the strongest materials ever measured [4]; the Young’s modulus of monolayer

graphene has been reported 1 TPa with ultimate strength of 130 GPa, [5]. Therefore, it has

attracted great attention in last decade in both academia and industries to develop economically

viable methods for manufacturing of graphene-based polymer nanocomposites.

14

In this chapter, a survey on the different methods for fabrication of graphene and graphene-

based polymer nanocomposites is presented. The fundamentals of electrical and thermal

properties of conductive polymer composites are also discussed.

2.3 Bottom-Up Graphene

In bottom-up processes, a wide variety of methods for graphene synthesis including chemical

vapor deposition (CVD) [6,7], epitaxial growth on SiC [8], arc discharge [9], unzipping carbon

nanotubes [10], chemical conversion [11], and self- assembly of surfactants [12] have been

reported. Usually, epitaxial growth and CVD can produce large size and defect free graphene

layers. These methods are highly attractive for electronic applications and fundamental studies.

The fabrication of high-quality graphene films has also been scaled through roll-to-roll CVD

synthesis [13]. However, these methods are hardly scalable for bulk 3-dimensional production

of graphene and are not suitable source of graphene for fabrication of polymer nanocomposites.

2.4 Top-Down Graphene

In top-down methods, graphite or graphite derivatives (e.g. graphite oxide (GO)) are exfoliated

to produce graphene or functionalized graphene. Generally, the top-down methods can yield

large quantities of graphene which is suitable for polymer composites application. Starting from

graphite which is a commodity material offer a great economic opportunity over bottom-up

processes. The annual global production of graphite is over 1.19 million tons at $1,250/ton in

2015 [14]. Thus, the top-down methods, are discussed in this chapter. Figure 2.1 summarizes the

different top-down methods for producing graphene or functionalized graphene from graphite or

GO.

15

Figure 2.1. Different top-down methods for producing graphene or functionalized graphene

from graphite or GO.

Acid-intercalated graphite can be expanded due to thermal shock/treatment resulting in

expanded graphite (EG) with thickness of around 100 nm [15]. The EG is a common filler in

manufacturing conductive polymer composites [4]. Lee et al. [16] produced a thinner form of

EG by thermal expansion of fluorinated EG which is known as graphene nanoplatelets (GnP).

Due to the rigidity and large diameter of GnPs obtained in this method, the electrical

conductivity improved at significantly lower loadings as compared to graphite or EG [17,18].

2.4.1 Direct Exfoliation of Graphite

High-quality, large-size sheets but very limited quantities of graphene can be produced by

micromechanical cleavage of bulk graphite. Therefore, it is not suitable for manufacturing

polymer composites, however, suitable for fundamental studies [19]. Graphite can be directly

exfoliated to multi- and single-layer graphene through sonication in the presence of N-

methylpyrrolidone [20] or polyvinylpyrrolidone [21], via dissolution in superacids [22], and

through graphite electrochemical functionalization assisted with ionic liquids [23] (Figure 2.1).

The direct exfoliation of multi- and single-layer graphene from graphite can be very

challenging. For instance, the removal cost of the hazardous materials (e.g. hydrosulfonic)

16

resulted from dissolution of graphite in superacids (chlorosulfonic acid) [22], can significantly

limit the potential for large scale production.

2.4.2 Graphite Oxide (GO)

The most viable techniques for large-scale production of graphene are mainly developed on the

basis of exfoliation and reduction of GO [4]. For the first time, GO was prepared in the

university of Oxford, UK, over 160 years ago by Brodie [24]. Later, GO was also produced

using Hummers method [25]. In this method, graphite is oxidized using oxidants such as

NaNO2, KClO3 and KMnO4, in the presence of nitric acid mixed with sulfuric acid. The GO is

composed of stacks of graphene oxide sheets with expanded interlayer distances between 6-10

Å which can be affected by the water content [26]. The GO surface can contain the epoxide,

carboxyl and hydroxyl groups as shown in Figure 2.2. The GO approximately has the 2/1/0.8 of

C/O/H atomic ratio [24].

Figure 2.2. Surface chemistry of GO containing containing carboxyl, poxide and hydroxyl

groups and double bonds.

The exfoliation of GO can provide large-scale production of graphene and functionalized

graphene sheets. Meanwhile, GO can also be dispersed in organic solvent such as water and

Tetrahydrofuran (THF), however, the it is thermally unstable and electrically insulative [4,27].

Therefore, the reduction of GO is required to restore the electrical conductivity. There are

17

various methods to exfoliate and reduce the GO to obtain functionalized graphene. The term

“functionalized graphene” is used because the full reduction of GO to graphene has not been

reported yet [4]. In the following sections two methods of exfoliation and reduction of GO are

presented.

2.4.3 Chemical Reduction of GO

In these processes, a colloidal dispersion of GO is provided, and then chemical reduction of the

exfoliated GO is conducted. A stable colloidal dispersion of GO can be produced using organic

solvents such as water and alcohols followed by vigorous stirring or sonication. The stable

colloidal dispersion of GO, can be chemically reduced using hydrazine [28,29], or

dimethylhydrazine [30]. The GO reduction restores the electrical conductivity, however, a high

oxygen content of C/O 10/1 remains [27].

The chemical reduction is an efficient method to produce the reduced GO (rGO), however, the

hazardous nature of chemicals and cost required for chemical reduction, significantly limit the

use of these techniques for large-scale production of graphene. Alternatively, the chemical

reduction can be conducted via dehydration of the hydroxyl groups at high pressure and

temperature (120−200 °C) in water [31,32].

2.4.4 Thermal Exfoliation and Reduction

Thermal exfoliation and reduction of GO can be conducted by thermal shock (rapid heating) of

the GO at high temperature (1000 °C for 30 s) under inert gas [33,34]. Due to the rapid heating,

the hydroxyl and epoxide groups on the surface of GO are decomposed into CO2. When the

generated pressure as a result of CO2 expansion exceed the van der Waals and hydrogen

bonding holding the graphene layers together the exfoliation occurs. Usually, the thermal

reduction is accompanied with 30% weight reduction which is associated with decomposition of

oxygen sites and water evaporation [34]. The thermal exfoliation results in a volume expansion

of 100-300 times. As shown in Figure 2.3, Thermally reduced GO are highly wrinkled due to

structural defects, and pressures generated by the expansion CO2 [34].

18

Figure 2.3. SEM of thermally reduced GO adapted from Ref. [33].

The C/O ratio of thermally reduced GO is about 10/1 as compared to 2/1 of GO. This ratio can

be further enhanced up to 660/1 by increasing the reduction time or increasing the heat

treatment temperature (e.g. 1500 °C) [35].

2.5 Preparation of Graphene-Based Polymer Nanocomposites

Recently, various processing techniques have been reported for dispersing GnP and GnP-

derived fillers in polymeric matrices. Most of the preparation routs for graphene-based polymer

nanocomposites are similar to those used for manufacturing other polymer nanocomposites [36].

The bonding interaction between filler and polymer matrix is highly critical for the final

properties of the composite. Most of the processing methods produce composites with non-

covalent bonding interactions where the interfaces between the polymer matrix and filler are

relatively weak forces (e.g. van der Waals or hydrogen) [37]. On the other hand, covalent

bonding interaction between filler and polymer matrix can provide stronger interfaces. Due to

the relative sparseness of the functional groups on the surface of pristine GnPs, the formation of

covalent bonding between polymer matrix and pristine GnPs is quite challenging. However, GO

usually offers a functional group-rich surface which can be used to introduce covalent bonding

between polymer matrix and GO. In this thesis, a pristine grade of GnP is used for the

fabrication of polymer nanocomposites, thereby the likelihood of the covalent bonding

formation is negligible [4]. Therefore, this thesis focuses on the preparation of graphene-based

19

nanocomposites based on the non-covalent bonding interaction between filler and polymer

matrix.

2.5.1 Filler Dispersion Methods

2.5.1.1 Solution-Based Methods

The solution-based processing routs involve mixing the colloidal suspensions of graphene (e.g.

rGO of functionalized graphene) with the polymer in a solution. The polymer can also be

dissolved in the colloidal suspensions of graphene by shear mixing or stirring. The resulting

colloidal suspensions can be directly casted by molding and solvent removal. However, this

method can result in the agglomeration of filler which can deteriorate the final properties of the

composite [36]. Alternatively, a non-solvent for the polymer can be used to encapsulate the

fillers by polymer chain due to precipitation. The resulting precipitated composite can then be

dried and further processed [37].

Solution mixing is widely used for preparation of polymer/GO-derived composites, using

organic solvent such as water and alcohols. This approach has been used for variety of polymers

such as polystyrene [30], Poly(methyl methacrylate) [38], polycarbonate [39] and polyimides

[40]. Particularly, solution mixing via sonication of aqueous suspensions of polymer/GO-

derived composites is a facile and appealing technique for water-soluble polymers such as

poly(vinyl alcohol) [41,42]. In solution mixing routes, the level of exfoliation and dispersion of

graphene within the polymer matrix are mainly governed by exfoliation procedures conducted

prior to mixing [37].

2.5.1.2 Melt Mixing

In melt mixing the filler in powder form is compounded under high shear forces. Melt mixing is

an economical and industrially viable processing route as compared to solution mixing and in

situ polymerization [37]. However, solution mixing and in situ polymerization can provide

better dispersion of fillers as compared to melt mixing [43]. Specifically, the dispersion of

20

single- or multi-layer graphene has not previously been reported using melt mixing without

prior exfoliation [4,37].

The melt mixing of the polymer/GnP has been reported in several studies [44,45]. In this

method the GnP powder can fed directly into the extruder and compound with the molten

polymer without any solvent. However, the extremely low density of the GnP powder poses a

highly challenge process. In a different approach, Kalaitzidou et al. [17] uniformly coated the

surface of GnPs with polymer powder by sonication of the polymer powder with GnPs in a non-

solvent before melt mixing. The polymer/GnP composites fabricated in this study showed lower

percolation threshold.

2.6 Thermoplastic Polymer Foams

Polymer foams offer unique mechanical, thermal, and physical properties which are mainly

governed by its constituents (e.g. cellular structure and polymer matrix) [46]. Polymeric foams

are considered as a composite structure in which the gaseous phase is uniformly dispersed in the

polymer matrix [46]. The properties of the polymeric foams are characterized by the expansion

ratio, cell density, cell size distribution and open cell content. The cellular structure of polymer

foams can be controlled by foaming technology implemented in processing. The choice of

foaming technology greatly depends on the type of polymer. Therefore, different foaming

technologies such as batch (single-step and multi-step) foaming, semi-continuous and

continuous processing routes have been developed based on various polymeric foams [46].

A wide variety of gases can act as blowing agent, however, some gases such as N2 and CO2 are

more suitable in terms of solubility, volatility, and diffusivity. After the fabrication of polymeric

foam, most of blowing agents are eventually replace with the ambient air as time progresses

[46]. A foaming process is a succession of three steps. These steps occur in different foaming

techniques such as batch, extrusion, and injection foaming. These steps include: (i)

implementation, which is the dissolution of predetermined amount of blowing agent into the

polymer matrix to form a single-phase gas/polymer mixture; (ii) liberation, which is the phase

separation of gas due to thermodynamic instability (e.g. rapid depressurization); (iii) evacuation,

21

which is the replacement of the blowing agents within the cellular structure with the ambient air

[46]. The foam processing path is illustrated in Figure 2.4.

Figure 2.4. Physical foaming of thermoplastic polymer consisting of these steps of: (i) the

dissolution of blowing agent into the polymer matrix and formation of single-phase gas/polymer

mixture; (ii) phase separation of gas due to thermodynamic instability; and (iii) curing when the

blowing agents are replaced with the ambient air.

In order to generate cells within the single-phase gas/polymer mixture, a driving force is

required. This driving force is induced by rapid depressurization or heat which can provide

thermodynamic instability. Due to the thermodynamic instability the gas solubility in the

polymer rapidly drops leading to the nucleation of cells [46]. There are two types of nucleation

mechanism including: homogeneous and heterogeneous nucleation. Homogeneous cell

nucleation randomly occurs in the gas/polymer mixture and needs higher activation energy as

compared to heterogeneous nucleation. However, heterogeneous cell nucleation takes place at

the interfaces of polymer with another phase (e.g. polymeric crystals and/or fillers), due to the

lower activation energy [47].

22

Due to the high pressure in the nucleated cells with compared to the surrounding pressure, the

nucleated bubbles start growing to decrease the pressure. This results in the diffusion of gas

from polymer matrix into the nucleated cells [46,47]. Bubble growth is mainly governed by melt

temperature, viscosity, gas concentration, diffusion coefficient and the number of nucleated cells

[47,48]. Temperature can change the diffusion and viscosity which play important roles in

bubble growth. A decrease in temperature can limit cell growth due to the lower diffusivity and

viscosity. However, increasing the temperature, can decrease melt strength leading to pinhole

formations and the cell wall ruptures with opened-cell structure [46,47,49,50].

2.6.1 Physical Foaming

Physical foaming is a well-developed foaming technology which includes the dissolution of

physical blowing agent in polymer melt. Injection molding and extrusion combined with

physical fomaing are well-known methods for thermoplastic polymer foaming [46]. The high-

speed continuous foam processing is main advantage of these processing routes. Due to a high

pressure in the mixing section generated by a specially designed screw, gas can be dissolved in

the polymer. The high pressure in the mixing section is required to produce a homogenized

polymer melt/gas mixture [50–52]. The resulting melt/gas mixture is subsequently extruded into

a low-pressure environment (e.g. mold cavity) where foaming takes place [50–52]. Another lab-

scale foaming technology is batch foaming. The batch foaming technique is mainly used for

fundamental studies to investigate the effect of blowing agent, materials composites and

foaming parameters on the physical and microstructures of the foamed samples.

2.7 Functional Properties of Polymer Composites

Polymer composites have attracted significant attention in past few years for both various

technological applications and fundamental studies. With the broad diversity in the surface

chemistry of fillers and polymer matrices via different technological routes, a number of

functional polymer composites have been developed with unique thermal, electrical, mechanical

properties, chemical stability, and tailored functionalities. This has greatly expanded the

23

research horizon of the functional polymer composites. The functional properties of polymer

composites can be readily tailored and tuned by controlling the hierarchical structure and

intricate interaction between polymer and fillers within the polymer composites. In this section,

fundamentals of polymer composites’ functional properties including: thermal conductivity,

electrical conductivity, dielectric performance and electromagnetic interference (EMI) shielding

performance are discussed.

2.7.1 Thermal Conductivity of Polymer Composites

2.7.1.1 Basics of Heat Conduction

Thermal conductivity is defined by Fourier’s law as following[53]:

𝑞 = −𝐾∇𝑇 (2.1)

Where q, K, and ∇𝑇 are respectively are heat flux, thermal conductivity and temperature

gradient. In this equation, K is valid for small temperature variations, because thermal

conductivity is a function of temperature. In solid materials, heat is transferred by acoustic

phonon and electrons as following [53]:

𝐾 = 𝐾𝑝 + 𝐾𝑒 (2.2)

Where, 𝐾𝑝 are phonon and 𝐾𝑒 are electron contributions. In metallic materials the electron

contribution is dominant due to the large concentration of free charge carriers. For instance, in

pure copper with K 400 W m-1K-1 at room temperature, the electron contribution (𝐾𝑒) in heat

transport is 99–98% [53]. The Wiedemann–Franz law defines 𝐾𝑒 through the following equation

[53]:

𝐾𝑒

𝜎𝑇=

𝜋2𝑘𝐵2

3𝑒2 (2.3)

Where, 𝜎 is electrical conductivity, T is temperature, the 𝑘𝐵 and 𝑒 are respectively the

Boltezman constant and the charge of an electron. Heat conduction in graphene and graphene-

based polymer composites is mainly dominated by 𝐾𝑝 even though they can have metal like

24

properties [53,54]. This concept can be described by strong sp2 bonded carbon atoms leading to

efficient heat transport through lattice vibrations. However, the contribution of electron can be

significant in doped materials [53].

The 𝐾𝑝 can be defined as [53,54]:

𝐾𝑝 = ∑ ∫ 𝐶𝑗(𝜔)𝜗𝑗2(𝜔)𝜏𝑗(𝜔)𝑑𝜔 (2.4)

Where 𝑗 is the phonon polarization branch including one longitudinal acoustic branch and two

transverse acoustic branches, 𝐶 and 𝜔 are respectively the heat capacity and the phonon

frequency, the 𝜗 is the velocity of phonon group (which is approximated by sound velocity in

solids), the 𝜏 is relaxation time of phonon. The phonon mean-free path (Λ) is a function of

phonon relaxation time and velocity (Λ = 𝜏𝜗) [53,55]. The acoustic phonon in typical solids are

scattered by other phonons, impurities, defects, interfaces and conduction electrons [53,56].

Phonon transport has two regimes including diffusive and ballistic regimes, which is important

to distinguish them from each other. If the samples size (L) is much larger than phonon mean-

free path (Λ), the thermal transport is called diffusive, where phonons experience several

scattering events. In Fourier’s law thermal transport is assumed to be diffusive. On the other

hand, when the L is smaller than Λ the thermal transport is ballistic [53,55]. In nanostructures,

the thermal conductivity decreases due to phonon scattering from boundaries and interfaces.

2.7.1.2 Thermally Conductive Graphene-Based Polymer Nanocomposites

Among a wide array of thermally conductive materials, graphene a highly promising

nanomaterial with exceptional thermal conductivity [1], can be used to fabricate polymer

nanocomposites with high thermal conductivity for various applications such as heat sink

components in electronic packaging.

Thermal conductivity of monolayer graphene has been reported ~5000 Wm-1K-1 when it is

suspended [1] and ~600 Wm-1K-1 when supported on a SiO2 substrate [57]. Carbon nanotubes

exhibit similar thermal conductivities, however, the 2-dimentional (2D) geometry of graphene

platelets may offer lower interfacial thermal resistance and consequently provide higher thermal

conductivity enhancement in polymer composites [58,59]. The anisotropic thermal conductivity

25

of the graphene and graphite [60] very often results in higher in-plane thermal conductivity (as

much as ten times) with compared to the through-plane thermal conductivity [61][62]. The

closer interparticle distance of the dispersed fillers in the polymer matrix decreases the

interfacial thermal resistance which can be explained by percolation theory [37,63].

In polymers and carbon-based polymer composites, phonon transferring is the main mode of

thermal conduction, therefore, the stronger bonding (e.g. covalent bonding) between the

polymer and the conductive filler can reduce the interfacial thermal resistance. The stronger

interfacial bonding reduces phonon scattering at the interface of filler and polymer leading to

higher thermal conductivity [64,65]. However, the thermal conductivity of the polymer

nanocomposites can be compromised due to functionalization of graphene which is required for

the enhancement of interfacial bonding [66].

The majority of studies on the thermal conductivity of graphene-based polymer nanocomposites

have focused on epoxy-matrix composites [59,61,64,65,67–69]. The thermal conductivity

enhancement achieved in these systems range for 3 to 12.4 Wm-1K-1 while the thermal

conductivity of the neat epoxy is approximately 0.2 Wm-1K-1. However, the required additive

loading levels are rather high (50 wt% and higher). Very small thermal conductivity difference

between carbon-based nanomaterial and polymer matrices (with compared to electrical

conductivity) can justify the much lower increase in the thermal conductivity achieved versus

electrical conductivity at the same loading [36]. Various methods have been used to achieve

higher thermal conductivity at lower filler loadings. These include: in-situ polymerization

[70,71], GnP surface modification [64,72], GnP alignment by electrical field [73], and the use of

hybrid additives [64,74] have all been proposed to develop polymer composites with high

thermal conductivity.

2.7.2 Electrical Conductivity of Polymer Composites

2.7.2.1 Basics of Electrical Conductivity in Polymer Composites

The formation of a conductive network within conductive polymer composites can be defined

and predicted by various geometric-, thermodynamic-, and statistical-based models [75–77].

The most common and reputable model is the percolation theory, however, this model is only

26

valid at filler loadings above the percolation threshold [75–77]. The percolation threshold of the

conductive polymer composites can be defined by statistical percolation theory as [76]:

𝜎 = 𝜎𝑓 . (𝜑 − 𝜑𝑐)𝑡 (2.5)

Where, 𝜑𝑐 is the percolation threshold, 𝜑 is the filler volume faction, the 𝜎𝑓 is the filler

conductivity, 𝜎 is the composite conductivity, and 𝑡 is the scaling exponent. The fillers are not

necessarily required to be in direct contact for electron transport, rather the electron conduction

can occur through tunneling and hopping between fillers which are separated by a very thin

insulating polymer layer (1.8 nm) [78,79], and this tunneling resistance is the governing factor

in the electrical conductivity of polymer composites [80].

Figure 2.5 shows a typical percolation curve of conductive polymer composites (high-density-

polyethylene (HDPE)-GnP). In general, a percolation curve for conductive polymer composites

has three distinct zones including: (i) insulative zone which is below the percolation threshold;

(ii) percolation zone where the percolation takes place; and (iii) conductive zone which is above

the percolation threshold.

Figure 2.5. Percolation curve of compression-molded HDPE-GnP composite (A typical

percolation curve of conductive polymer composites).

27

In the insulating zone, the concentration of conductive fillers is low, and the conductive fillers

are far from each other. Therefore, the charge transport is governed by the polymer matrix and

the polymer composites exhibit an electrical conductivity close to neat polymers’ conductivity

(10-13 −10-11 S/cm) [2,3,81]. In this zone due to the low filler loading, the insulating gap is too

large and the chance of charge transport between the conductive fillers is very small.

By increasing the filler loading, the insulating gaps between the fillers decrease. When the

interparticle distance drops to below 1.8 nm, electron tunneling and hopping govern the electron

transport mechanism [78,79]. Basically, a much higher electric field can be developed over a

narrower insulating gap as compared to the macroscopic applied electric field by a factor of M

(i.e. the ratio of average size of conductive fillers’ to the average interarticular distance) [82,83].

This higher electric field provides enough energy for the free electrons to hop or tunnel over the

insulating gaps (see Figure 2.6). In the tunneling mechanism, the insulating gap should be

narrow enough for the tail of electron wavefunction to pass through a barrier. However, in

hopping mechanism, an electron needs sufficient energy to reach an energy barrier for changing

its lattice site [82–84].

Figure 2.6. Diagram of electron-transfer mechanisms between adjacent sites separated by a

potential energy barrier. Adapted from Ref. [84].

28

At higher concentrations, and in the percolation zone, the conductive fillers come in direct

contact with each other letting the free charges pass through the polymer composite. Therefore,

in the percolation region, the conductivity of composites dramatically increases by several

orders of magnitude. By incorporating more conductive filler, a 3-dimentioanl conductive

network is formed, however, the electrical conductivity marginally increases. This marginal

increase can be attributed to high contact resistance between conductive fillers which results in a

plateau region in the percolation curve [85].

2.7.2.2 Electrically Conductive Graphene-Based Polymer Nanocomposites

One major aspect of graphene-based polymer nanocomposites is for applications such as EMI

shielding which rely on the electrical conductivity and electron transport characteristics of the

composites. Graphene with exceptional electrical conductivity (∼6,000 S/cm) [4], is an ideal

nanomaterial for manufacturing conductive polymer composites.

Different methods such as surface modification of graphene platelets [86,87], synergism of the

hybrid nanomaterials [88,89] and in-situ polymerization [87,90] have been implemented to

develop more efficient graphene-based polymer composites with enhanced electrical properties.

Yousefi et al. [86] fabricated self-aligned reduced graphene oxide (rGO)-polymer

nanocomposites by dispersing monolayer graphene in epoxy using an aqueous casting method

through the in situ reduction of graphene oxide (GO) [86]. They achieved very low percolation

threshold of 0.12 vol% [86]. Kim et al. [89] fabricated a hybrid polymer nanocomposite by

chemical vapor deposition of carbon nanotubes onto rGO oxide platelets followed by solution

mixing. They reported a dielectric constant of 32 with a dielectric loss of 0.051 at 0.062 wt%

loading of hybrid fillers and 1×102 Hz [89].

2.7.3 Electromagnetic Interference (EMI) Shielding

Electromagnetic interference shielding of radio frequency radiation is a serious concern in our

technological society. The EMI shielding is needed to protect our environment and workplaces

from the radiation that is emitted by the new ultra-high (GHz) frequency products. These

29

include mobile/communication devices, high-frequency microwave radio relay transmitters,

microwave remote sensors, wireless Local Area Network (LAN), communications satellites,

cable and satellite television broadcasting, direct-broadcast satellite (DBS), and radio

astronomy. Consequently, there is a great demand in industry for lightweight and easily

processable materials with effective EMI shielding properties.

To efficiently shield EMI, the CISRP organization (Comité International Spécial des

Perturbations Radioélectriques) has established electromagnetic compatibility (EMC)

regulations for electronic enclosures. Based on these regulations, a minimum of 30 decibels

(dB) EMI shielding effectiveness (SE) is commercially needed to shield 99.9% of the incident

electromagnetic wave [91].

2.7.3.1 Shielding Mechanism

In a plane EMI wave (in this thesis EMI shielding measurements have been conducted using

plane wave), the electric and magnetic fields are normal to each other and are in a plane which

is perpendicular to the direction of wave propagation.

Electrons’ response to electric and magnetic fields in a shielding material can be defined based

on the Lorentz’s force law [91]:

�⃗� = 𝑞�⃗⃗� + 𝑞�⃗� × 𝜇0�⃗⃗⃗� (2.6)

Where, �⃗⃗� and �⃗⃗⃗� are respectively the electric and magnetic fields’ strength; 𝑞 is the charge of

particles traveling with the velocity of �⃗�, and 𝜇0 is the free space magnetic permeability (4π×10-

7 H·m-1 H.m-1 (henries/meter)).

The EMI shielding consists of three mechanisms including: (i) reflection; (ii) absorption; and

(iii) multiple reflection [3,91–93] as it is shown in Figure 2.7.

30

Figure 2.7. Schematic of shielding mechanisms of a plane wave by a shielding material.

Adapted from Ref. [94].

A fraction of an incident electromagnetic wave is reflected mainly due to the presence of the

charge carriers (that is, the electrons and holes) and/or the surface charges [95–97]. A fraction of

the electromagnetic wave is also transmitted within the shield and its energy is dissipated though

absorption. The absorption mechanism originates from the Ohmic and polarization losses [98].

The Ohmic loss results in energy attenuation via the current flow through the conduction and

tunneling mechanisms. The polarization loss is correlated to the interfacial polarization’s density

and is thereby related to the absorber’s real permittivity [97,99]. The higher amount of mobile

charges increases the ability of shielding material to attenuate electromagnetic wave through

reflection and absorption mechanisms [100,101]. A fraction of the electromagnetic wave is

transmitted through the interface of the shielding material, while a fraction of it is reflected from

the interface [100,101]. Multiple reflection mechanism is the re-reflection of electromagnetic

waves which have been already reflected inside the shielding material [92].

31

The ability of materials for attenuating of electromagnetic waves is defined in though EMI

shielding effectiveness (SE). The shielding performance for a given electromagnetic radiation is

defined as [92,102]:

𝐸𝑀𝐼 𝑆𝐸 = 10log10(𝑃𝑖

𝑃𝑡) (2.7)

Where, Pi is the incident power and Pt is the transmitted power in dB. For instance, a material

with a SE of 40 dB can block 99.99% of the incident wave. This equation can be utilized to

obtain the contributions of reflection and absorption mechanisms (more details are discussed in

Chapter 5).

2.7.4 Dielectric Properties of Polymer Composites

High performance dielectric materials are extremely critical in the development of the next-

generation miniaturized electronic devices. Dielectric materials with high dielectric permittivity

(ε') and low dielectric loss (tan ) are receiving an ever-increasing interest in different cutting-

edge industries such as energy storage devices [103], optoelectronics [104], piezoelectric

generators [105], inverters and transistors [106]. Lightweight, multifunctional, low cost,

polymer nanocomposites are very promising for use as dielectric materials. They have large

tunability of dielectric permittivity and dielectric loss, superior resistance to chemicals, ease of

processing, and tailorable thermal and mechanical properties [107,108].

Polymers have extremely low dielectric loss and high dielectric breakdown strength; however,

they suffer from low dielectric constants (3< ε' <10). The incorporation of fillers of different

types and shapes can further improve the dielectric properties of polymer nanocomposites. In

general, two different types of fillers, including nonconductive and conductive, are used to

fabricate dielectric polymer nanocomposites [107].

32

2.7.4.1 Dielectric Mechanism

A dielectric material can be polarized by an applied field and is able to store energy. The

response of a dielectric material to an applied field can be defined by the concept of parallel-

plate capacitors. When a direct current (DC) of voltage (V) is applied, the electric field (E)

generated between the plates (with d distance) is equal to E=V/d. As schematically shown in

Figure 2.8, the applied electric field polarizes the dielectric material by separation of positive

and negative charges.

Figure 2.8. Polarization of a dielectric material by an applied electric field.

When the dielectric material is free space in a parallel-plate capacitor, the charges stored on the

plates is equal to Q = ε0E (ε0: free space permittivity). When a free space is substituted with

dielectric material, an extra charge (P) can be stored by capacitor due to higher polarizability of

the dielectric material as compared to the free space. The extra charge can be expressed as:

𝑃 = 𝑄(𝜀𝑟 − 1) (2.8)

Where, 𝜀𝑟 is the permittivity of the dielectric material with respect to the free space

permittivity. A dielectric material with higher 𝜀𝑟 can store more energy on the surfaces of the

parallel-plate capacitor. Generally, permittivity is defined as a complex function [108]:

33

ε (ω) = ε' (ω) - iε'' (ω) (2.9)

where ω is the frequency, ε'(ω) is the real part and ε''(ω) is the imaginary part of dielectric

permittivity. The real part of equationε (ω) = ε' (ω) - iε'' (ω) (2.9) is related to

charge displacement which is affected by different types of polarization within the material

[108]. On the other hand, the imaginary part of equation ε (ω) = ε' (ω) - iε'' (ω) (2.9)

indicates the energy dissipation or dielectric loss which is quantified by the ratio of ε'' to ε' (tan

). Dielectric loss of polymer composites are generally governed by, polarization loss of space

charges, ohmic loss, and the molecular dipole movement (dipole loss) [109,110].

In general, there are several polarization mechanisms contributing in the overall real

permittivity. These include: interfacial, dipolar (orientational), ionic (atomic), and electronic

polarization [108,111]. Higher polarization density and thereby higher real permittivity can be

obtained using more polarization types. However, more complicated frequency dependencies

can be induced by various structures within the material [108]. As shown in Figure 2.9 each type

of polarization at various frequency ranges is associated with a peak in dielectric loss (ε'').

Figure 2.9. Real (ε') and imaginary (ε'') parts of permittivity as a function of frequency for a

material showing interfacial, orientational, ionic, and electronic polarization. Adapted from Ref.

[111].

34

Electronic polarization originates from the delocalization of electrons surrounding nucleus

inside atoms in response to an applied electric field. In the electronic polarization, the dielectric

loss is only induced over optical and infrared frequencies [108]. Ionic polarization is induced by

the displacement of charged ions and separation of positive and negative ions in a crystal lattice.

Ionic polarization is the main type of polarization in glass, ceramics and inorganic crystals.

However, the contribution of ionic polarization in the total polarization for organic materials due

to the absence of ions is limited [112].

In polymer composites, the dipolar polarization can be induced due to the orientation of

permanent molecular dipole moments [111]. Depending on the nature of dipoles, frequency and

temperature, the dipolar relaxation can vary between 0.1 to 107 Hz [111]. By controlling the

density and structure of dipoles, polymers with different levels of real permittivity and dielectric

loss can be synthesized (e.g. ferroelectric, relaxor ferroelectric and dipolar glass, polymers)

[113].

However, over the frequency range (<1MHz) the governing polarization is mainly interfacial

polarization (e.g. polarization of the matrix/filler interface) [101,108,114]. According to the

Maxwell–Wagner–Sillars (MWS) effect [110], charges can be accumulated at interface of filler

and polymer matrix due to a considerable contrast between the electrical conductivity and/or

permittivity of the fillers and polymer matrix. This can be explained by the concept of relaxation

time expressed as following [110]:

𝜏 = 𝜀𝜎⁄ (2.10)

Where, 𝜀 is the dielectric constant and 𝜎 is the conductivity. When an electric field is applied

across a composite material of two constituents, charge can be accumulated at the interface of

two materials with different relaxation times (𝜏). The large difference in conductivity (for

conductive fillers) or dielectric constant (for ceramic fillers) is highly desirable for higher

interfacial polarization.

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48

CHAPTER 3

3 Development of a Facile Technique to in situ Exfoliate and Disperse Graphene Nanoplatelets in Polymer Matrices

The following section is based on text from

Hamidinejad, S.M., Park, C.B., and Nazarpour, S., "Method of Exfoliating and Dispersing High

Concentration Graphene Nanoplatelets (GnP) into Polymeric Matrices Using Supercritical Fluid (SCF)”,

US Provisional Patent, applied for, Application Serial No. 62/512,790, May 31, 2017s

Hamidinejad, S.M., Chu, R.k.M., Zhao, B., Park, C.B., and Filleter, T., “Enhanced Thermal

Conductivity of Graphene Nanoplatelet-Polymer Nanocomposites Fabricated via Supercritical Fluid

Assisted In-Situ Exfoliation”, ACS Applied Materials and Interfaces, 2018, 10 (1), 1225−1236

Hamidinejad, S.M., Zhao, B., Chu, R.k.M., Moghimian, N., Naguib, H., Filleter, T., and Park, C.B.,

“Ultralight Microcellular Polymer-Graphene Nanoplatelet Foams with Enhanced Dielectric

Performance”, ACS Applied Materials and Interfaces, 2018, 10 (23), 19987–19998

3.1 Summary

This chapter presents a SCF-assisted

method for in situ exfoliation and

dispersion of GnPs in polymer

matrices. In this method, the molten

polymer/GnP mixture is treated with

SCF (CO2, N2), while the act of shearing of the mixture is constantly applied throughout the

process (i.e. extrusion, and injection processes). This technique is then followed by rapid

49

depressurization of the SCF-treated polymer-GnP mixture to further exfoliate and disperse the

GnPs in the polymer matrix.

3.2 Introduction

Next-generation, multifunctional materials are considered to be the foundation for technological

innovations in the 21st century. By combining science with specialized engineering knowledge,

research on advanced functional materials will enable the design and development of cutting-

edge, multifunctional, lightweight, and high-performance materials for a wide variety of

applications that can be used in the automotive, aerospace, telecommunication, energy, and

microelectronics industries [1,2]. Recently polymer nanocomposites have shown enormous

potential as a highly-desirable class of advanced functional materials. They have an attractive

array of properties. Tunable functionality of the polymer composites achieved by the

incorporation of different fillers can help to address the requirements of a broad range of

cutting-edge applications.

One example of the filler candidates now emerging is graphene. Graphene is an atomically thick

layer composed of sp2 carbon atoms has exceptional mechanical, electrical, and thermal

properties. However, the underpinning for economically-viable manufacturing of graphene-

based polymer composites is missing. It is greatly challenging and expensive to exploit the full

potential of graphene due to the complexities in the exfoliation and dispersion of graphene

layers in the polymer matrix. Physical and mechanical properties of GnP-based polymer

composites not only depend on the chemical composition, defects concentration, aspect ratio

and level of exfoliation of GnPs; but also, highly depend on how well the GnPs are dispersed in

polymer matrix [3,4].

There are several approaches to produce and exfoliate graphene nanoplatelets (GnP). (i):

Formation of graphene oxide platelets followed by reduction. In this approach, the natural

graphite materials are treated with intercalant and an oxidant to produce graphite intercalated

compound (GIC) also called graphite oxide (GO). The resulting product will be then subjected

to exfoliation procedure which can be either by solution-based separation approach assisted by

sonication or by thermal shock exposure. (ii): Formation of pristine GnPs directly from natural

graphite without going through a chemical intercalation routes such as direct sonication of

50

graphite flakes to separate and exfoliate GnP. (iii) Small scale production of GnPs using

chemical vapor deposition and epitaxial growth which is called bottom-up method [3,4].

In the most of abovementioned graphite intercalation and exfoliation methods, considerable

amount of chemicals is required which leads to tedious washing steps. The GO, prepared in

approach (i) even after chemical and thermal reductions still shows much lower thermal and

electrical conductivities as compared to the pristine GnPs, due to the complexities in complete

reduction of highly oxidized graphite in this method. The GnPs produced through the

approaches (ii) and (iii) can be highly conducting and defect-free, however, these approaches

are cost prohibitive for large-scale production of GnPs [4].

One of the most economically viable and scalable technique for dispersing GnPs into

thermoplastic polymers is melt mixing due to more compatibility with current industrial

practice. However, the knowledge about melt mixing of GnP-based polymer nanocomposites is

still insufficient. Restacking and agglomeration of GnPs during melt blending can significantly

reduce their effectiveness in the functionality enhancement of the products. Thus, it is crucial to

develop an economically-viable and scalable method of melt blending not only to assure that the

exfoliated GnPs remain exfoliated but also to further exfoliated and dispersed them during

blending.

Thus, the thesis develops an innovative method of in situ exfoliation and dispersion of GnPs in

polymer matrices for various functionalities. In this technique, the GnPs, are melt-blended with

the polymer to form a treatable polymer/GnP mixture. Then the molten polymer-GnP mixture is

subjected to SCF in an extrusion or injection process. The process is followed by rapid

depressurizing to transform the SCF penetrated between the layers of GnPs to the gaseous state.

During the phase transition, the expanding SCF can exfoliate graphene layers.

3.3 Experimental Section

3.3.1 Materials and sample preparation

A commercially available grade of high-density-polyethylene (HDPE), HHM 5502BN

Marlex®, (MFI:0.35 dg/min.-1 at 230ºC/2.16 kg, with a specific gravity of 0.955 g.cm-3,

51

Chevron Phillips Chemical) was used as the polymer matrix. A commercial GnP grade heXo-g-

V20 with a specific gravity of 2.2 g.cm-3, a surface area of 30 m2/g, average lateral dimensions

of 50 μm and an average thickness of 20 nm, was provided by NanoXplore Inc. Montreal, QC,

Canada, and was used to fill the polymer matrix. It is notable that, based on the recommended

nomenclature for 2D carbon materials [5] the filler used in this study is graphite nanoplates

(GNP). However, the commercial name of this filler (i.e. graphene nanoplatelets (GnP))

introduced by the manufacturer (NanoXplore Inc.) was used. Commercial Nitrogen (N2) and

carbon dioxide (CO2), supplied by Linde Gas, Canada, were used as the physical blowing agent

and supercritical fluid.

A 35 wt.% HDPE-GnP masterbatch was produced by melt compounding using a TDS-20 twin-

screw extruder with a screw diameter of 22 mm and L/D: 40. The temperature profile was set to

180-220°C and a rotational speed of 45 rpm and a throughput of 5 kg.hr-1 were used. HDPE-

GnP composites with different GnP loading content were then obtained by diluting the HDPE-

35 wt.% GnP masterbatch with neat HDPE and mixing them in a twin-screw extruder (with

diameter of 27 mm and L/D: 40).

3.3.1.1 Foam injection molding

A 50-ton Arburg Allrounder 270/320C injection molding machine (Lossburg, Germany), with a

30-mm diameter screw equipped with MuCell® technology (Trexel, Inc., Woburn,

Massachusetts) was used to fabricate the GnP-HDPE nanocomposite samples. The mold

contained a rectangular cavity with a fan gate after the sprue. The mold cavity dimensions were

132 × 108 × 3 mm. More details about the implemented mold in this study were reported by Lee

et al. [6].

Three different types of HDPE-GnP nanocomposites, namely injection-molded solid (IMS),

injection-molded foam (IMF) and high-pressure-injection-molded foam (HPIMF), were

prepared.

The IMS samples were fabricated using the conventional injection molding process without the

SCF-treatment and physical foaming. For the HPIMF and IMF samples, 0.4 wt% N2 (as the

SCF), was injected into the barrel in its supercritical form using the MuCell module. The

52

MuCell module is a built-on, commercially available, system for an injection molding machine

to facilitate injecting the physical blowing agent into the barrel.

In the IMF samples, after the GnP-polymer mixture was treated with the SCF, the mold cavity

was partially filled with a gas-GnP-polymer mixture. In the HPIMF, the mold cavity was fully

filled with the single-phase gas-GnP-polymer mixture. Then the filling step was followed by a

composite melt packing step to re-dissolve the nucleated cells back into the melt. The nominal

degrees of foaming in the IMF samples were controlled by partially filling the mold cavity. The

processing parameters used in the injection molding of the IMS and IMF nanocomposites were

optimized based on their microstructure integrity and thermal conductivity. Table 3.1

summarizes these processing parameters.

Table 3.1. Processing parameters used in injection molding of solid and foamed composites

Parameter IMS HPIMF IMF

Melt temperature (°C) 210 210 210

Barrel pressure (MPa) 16 16 16

Screw speed (rpm) 300 300 300

Metering time (s) 12 12 12

Injection flow rate (cm3s-1) 90 90 90

Mold temperature (°C) 75 75 75

Pack/hold pressure (MPa) 30 30 N/Aa

Pack/hold time (s) 15 60 N/A

Gas injection pressure (MPa) N/A 24 24

N2 content (wt.%) N/A 0.4 0.4

Degree of foaming (%) N/A N/A 7, 16, 26

aN/A: not applicable

53

The schematic of the injection molding processes (i.e. IMS, HPIMF and IMF) and injection-

molded parts have been presented in Figure 3.1. The IMF’s actual degrees of foaming were

measured using samples via the water-displacement method (the ASTM D792-00) after

fabrication.

Figure 3.1. The injection molding processes (i.e. IMS, HPIMF and IMF) and injection-molded

parts

3.3.1.2 Extrusion foaming

A tandem foam extrusion system was used to fabricate the foam samples with different

densities. The tandem extruder used in this study could extrude foam filaments with a 4 mm

diameter at a rate of 2 kg/h. The tandem foam system had two single-screw extruder barrels. The

first extruder (Brabender 05−25−000) had a 5-hp extruder drive with a mixing screw of 19 mm

and L/D:30. The second extruder consisted of a 15-hp extruder drive (Killion KN-150) with a

mixing screw of 38.1 mm and L/D:18.

54

In the first extruder, the HDPE-GnP experienced a 215°C temperature. The HDPE-GnP

composites were completely melted due to the temperature and the screw motion, which

generated shear heating. Then, using a syringe pump, 4 wt. % CO2 at a constant flow rate was

injected into the melt through the first extruder barrel. The rotating screw inside the first

extruder facilitated the CO2’s dissolution in the HDPE-GnP mixture. In the second extruder, the

homogeneous GnP/HDPE/CO2 mixture experienced temperatures around the melting points of

127°C -145°C). Foaming occurred at the die exit due to the rapid depressurization, during which

the GnP/HDPE/CO2 mixture was extruded down to ambient conditions. This resulted in the

gas’s phase separation [7]. A stainless-steel capillary die with a circular pinhole with a diameter

of 1.2 mm and a channel length of 10 mm was used. The temperature of die and the second

extruder were brought down, and when the system’s temperature had reached an equilibrium,

then the foamed samples were collected. The flow rate was kept constant at 5.5 g/min. The

schematic of the extrusion process and processing parameters have been presented in Figure 3.2.

Figure 3.2. The schematic of the extrusion process and processing parameters

The disk-shape samples (1 mm thickness and 10 mm diameter) solid compression molded

(SCM) nanocomposites were hot-pressed at a temperature of 215°C for 7 minutes under a 6 kN

55

pressing force. Next, the samples with a mold assembly were quenched by using compressed

air.

3.4 Characterization

To examine the exfoliation and dispersion of GnPs in the polymer matrix, Wide Angle X-ray

Diffraction (WAXD) analyses were conducted on the injection-molded nanocomposites using a

Rigaku MiniFlex 600 X-ray diffractometer (Cu Kα radiation, λ = 1.5405 Å). To further evaluate

the level of exfoliation and dispersion of different samples, transmission electron microscope

(TEM; FEI Tecnai 20) were conducted. The TEM samples were prepared by cryo-

ultramicrotomy (Leica EM FCS). The microstructure and morphology of the fabricated samples

were investigated using scanning electron microscopy (SEM; Quanta EFG250). The samples

were frozen in liquid nitrogen, cryofractured, and sputter-coated prior to electron microscopy.

3.5 Results and discussion

3.5.1 Effect of SC-N2-treatment and physical foaming on GnP’s exfoliation

and dispersion in an injection molding process

To quantify the GnPs’ exfoliation level after the SCF-treatment, WAXD analyses were

conducted. Figure 3.3 (a and b) shows the WAXD patterns for the neat HDPE, GnP powder, the

IMS samples (HDPE-9 vol.% GnP) and their HPIMF and IMF counterparts. The diffraction

peak at 2θ = 26.6° is characteristic of the (002) reflection of the graphite (I002), associated to the

d-spacing between the monolayer graphene sheets. By monitoring the (002) diffraction peak of

the XRD pattern the stacking nature of GnP’s can be identified. As the ratio of exfoliated GnPs

to stacked (unexfoliated) GnPs increases, the intensity of (002) diffraction decreases [8–14].

While a low-angle shift of the (002) diffraction peak indicates GnP d-spacing expansion and

intercalation [9,10], the decrease in the intensity of (002) diffraction has been frequently used as

the evidence of exfoliation in literature [8–14].

The SCF-treatment and physical foaming of the HDPE-GnP composites produced a 94%

decrease in the intensity of the I(002) initial value’s diffraction peak, which corresponded to the

untreated nanocomposites (IMS) (Figure 3.3c). This suggested very efficient exfoliation of

56

GnPs which is in good agreement with the literature [8–14]. However, there was still a layered-

GnP structure retained in each flake [12], as evidenced by the presence of a small diffraction

peak (I002), even after the SCF-treatment and physical foaming of the HDPE-GnP composites.

Moreover, the SCF-treated HDPE-GnP composites’ (002) diffraction peaks shifted to somewhat

lower angles, indicating a slight d-spacing expansion in the layered GnPs’ structure, based on

Bragg’s law.

To further support the GnPs exfoliation in the SCF-treatment technique we have also conducted

WAXD on all of the samples over 2θ angles of up to 50° to examine the effect of GnPs

orientation on the intensity of the (002) and (100) peaks. The intensity of the (002) and (100)

peaks of layered structures such as GnP and hBN can be used to identify the orientation of these

fillers within polymer composites [13,15,16]. Vertically and horizontally oriented flakes are

responsible for magnifying the (100) and (002) peaks respectively [15,16] as schematically

shown in Figure 3.3b. The (100) peaks of all the samples are found to be very weak and they are

not evident in Figure 3.3a. In a magnified XRD pattern over 2θ=40° to 50°, presented in Figure

3.3b, it was notable that the IMS and IMF samples had very small (100) diffraction peaks with

similar intensity. However, the intensity of the (002) peak of IMS samples was more intense as

compared to those of the IMF counterparts. This suggests that (i) the GnPs were horizontally

oriented on the surfaces of IMS and IMF samples, and (ii) the decrease in the intensity of (002)

peaks of the IMF samples is caused solely by exfoliation of GnPs and not by the orientation of

GnPs [13,15,16].

Figure 3.3c also shows that, once the HDPE-GnP composites had been treated with the SCF, the

I002’s intensity considerably decreased. However, the degree of foaming (that is, the void

fraction in percentage) did not significantly reduce the I002’s intensity in the range of 7-26 %. It

is also noteworthy that the SCF-treatment provided almost the same level of exfoliation in the

HPIMF samples as it had in the IMF samples. This was even after the nucleated bubbles had re-

dissolved back into the composite melt under high pressure.

57

Figure 3.3. (a) XRD spectra of neat HDPE, GnP powder, IMS samples (HDPE-9 vol.% GnP)

and their HPIMF and IMF counterparts with various degrees of foaming; (b) magnified XRD

pattern of Figure 3.3a over 2θ=40°-50° highlighted with light green, to examine (100)

diffraction peaks and illustration of the GnPs’ orientation and their effect on the (002) and (100)

diffraction peaks of the XRD pattern; (c) residual values (%) of I(002) (intensity of the (002)

diffraction at 2θ = 26.5°) before and after SCF-treatment and physical foaming; (d)

representative TEM micrographs of the IMS of HDPE-4.5vol.% GnP and; (e) IMF of HDPE-

4.5vol.% GnP; (f) ideal conceptualization of various phenomenon resulting in further exfoliation

and dispersion of GnPs in IMF samples. DF stands for degree of foaming.

58

In the IMF samples, the HDPE-GnP mixture is subjected to the SCF before being injected into

the mold cavity. It is well known that SCF can help to enhance dissolution behavior [17]. Over a

sufficient duration, the SCF is capable of intercalating the graphitic layered structures. This

weakens the nanoplatelets’ bonding force and makes their exfoliation easier. Moreover, the

dissolution of the SCF in the HDPE melt creates a favorable interaction between the GnPs’

surfaces and the polymer melt which reduces system energy. This decreases interfacial tension

between the GnPs and the polymer matrix, and it allows for a better GnP dispersion in the

polymer melt [18]. Furthermore, the SCF’s plasticizing effect enhances the polymer molecules’

diffusivity. Furthermore, the SCF’s plasticizing effect enhances the polymer molecules’

diffusivity. This would likely increase the likelihood for polymer chains to penetrate the GnP

nanoplatelets’ interlayer regions because of (i) the higher mobility of the chains by the

plasticizing SCF dissolved in the polymer matrix; and (ii) the increase in the GnP nanoplatelets’

interlayer distances of the SCF-GnP intercalated structure. As a result, the GnPs’ exfoliation and

layer separation are more effectively induced. To completely dissolve the SCF in the HDPE-

GnP composite, it is necessary to maintain the HDPE-GnP/gas mixture’s single-phase

throughout the injection molding process.

This process was followed by a rapid depressurization to transform the dissolved and

intercalated-SCF state into a gaseous state. During the phase transition, the expanding SCF can

further separate and exfoliate graphene layers. Moreover, during the phase transformation, many

small cells were generated between the platelets within the intercalated polymer/gas mixture.

This led to further delamination and separation of individual platelets in the polymer matrix.

Meanwhile, during the SCF’s depressurization and phase transformation, an additional driving

force for the delamination and dispersion of the GnPs was generated. Nucleated cells growing

near the GnPs acted like nucleating agents, and this further delaminated and uniformly dispersed

the GnPs in the polymer matrix. Figure 3.3d and Figure 3.3e, respectively show representative

TEM micrographs of the IMS and IMF (containing 7% degree of foaming) samples. It is notable

that agglomerated and thick GnPs in the IMS samples (Figure 3.3d) were further exfoliated to

thinner layers after SCF-treatment in the IMF samples (Figure 3.3e). This result is in a good

agreement with the WAXD results and provides further evidence of the higher level of

exfoliation and better dispersion after SCF-treatment and physical foaming. Figure 3.3f shows

the ideal conceptualization of the various phenomenon that resulted in further GnP exfoliation

and dispersion in the IMF samples. Moreover, the dissolved gas in the composite melt reduced

59

the melt’s viscosity and, therefore, lowered the shear stresses that were applied to the fillers.

This helped to reduce the GnP’s mechanical breakdown.

3.5.2 Effect of the SC-CO2-treatment physical foaming on the GnP’s

exfoliation and dispersion in extrusion foaming

Figure 3.4a-b show the TEM micrographs of the SCM HDPE-4.5 vol.% GnP composite and its

extruded-foam counterparts with a density of 0.05 g.cm-3. It is notable that after foaming, the

agglomerated and thick GnPs in the SCM samples (Figure 3.4a) were further exfoliated into

thinner layers. This was in good agreement with WAXD results as shown in Figure 3.4c. The

intensity of the diffraction peak of the (002) plane at 2θ = 26.5°, which is a characteristic of the

graphite, decreased 86% compared with the SCM samples (inset Figure 3.4a). The decreased

intensity of the (002) diffraction is the indication of GnPs’ exfoliation [8–14,19]. It is notable

that, the SCF-assisted exfoliation resulted in thinner GnPs in foam-extruded samples as

compared to those in the SCM counterparts. However, there were still layered GnP structures

retained in the foam-extruded nanocomposites. This is shown by the small diffraction peak of

the (002) plane corresponding to the foam-extruded nanocomposites.

In the extruded foam samples, the HDPE-GnP mixture was subjected to supercritical-CO2 (sc-

CO2) before being extruded down to ambient conditions. The sc-CO2 was able to diffuse in

between the layers of GnPs due to its small molecular size and high diffusivity [20]. This

weakened the nanoplatelets’ interlayer Van der Waals forces and made their exfoliation easier.

Moreover, the dissolution of the sc-CO2 in the polymer matrix provided a favorable interaction

between the polymer melt and the surface of GnPs, leading to lower interfacial tension between

them [10,18]. At the same time, it enhanced the polymer molecules’ diffusivity, which increased

the chance of polymer chains to penetrate in between the layers of GnPs [10,18].

60

Figure 3.4. (a) Representative TEM micrographs of the SCM of the HDPE-4.5vol.% GnP and;

(b) Foam-extruded 4.5 vol.% HDPE-GnP; (c) XRD spectra of neat HDPE, GnP powder, SCM

samples (4.5 vol.% HDPE-GnP), and their extruded-foam counterparts with different densities

After the sc-CO2-treatment of the HDPE-GnP composite, a rapid depressurization followed,

where the GnP exfoliation driving forces were effectively generated. The expansion of the sc-

CO2 which had been diffused between the layers of the GnPs delaminated and exfoliated the

graphene layers. Hamidinejad et al. [19], Zhamu, and Jang [21], Kaschak et al. [22], and Pu et

al. [20] also reported similar phenomena and observed that when the GnP was first subjected to

a supercritical fluid, and then followed by a rapid depressurization, the SCF-intercalated GnPs

were forced to exfoliate due to the sc-CO2’s expansion after the rapid depressurization [20–22].

Eventually the growth of the nucleated cells near the GnPs induced another driving force for the

GnPs’ delamination and dispersion in the polymer matrix. This is shown by the differences in

the I(002)’s intensities in the foamed samples. A decrease in the density of the foamed samples

from 0.27 g.cm-3 to 0.05 g.cm-3 reduced the I(002)’s intensity. This can be attributed to the biaxial

18 20 22 24 26 28 30 32

GnP

powder

Solid Foam Foam0

20

40

60

80

100

% R

esid

ua

l In

ten

sity

(I002)

(0.05 g.cm-3)(0.27 g.cm

-3)

Solid-HDPE-GnP

Foam-HDPE-GnP

0.05 g.cm-1 density

Neat HDPE

Foam-HDPE-GnP

0.27 g.cm-1 density

GnP Powder

Inte

nsity (

a.u

.)

2 (°)

(c) (002)

61

stretching of the polymer matrix during foaming which created an additional driving force for

the GnP’s exfoliation. A similar phenomenon has also been reported for PP-nanoclay by Zhao et

al. [18] and for PP-GnP by Ellingham et al. [10]. Figure 3.5 shows an ideal 2D

conceptualization of the phenomena attributed to further exfoliation and parallel-plate

arrangement of the GnPs in the extruded foam samples.

Figure 3.5. Ideal conceptualization of various phenomenon resulting in further exfoliation and

and parallel-plates arrangement of the GnPs in the extruded foam samples

3.6 Conclusion

This research work offers a scalable method of manufacturing GnP-based nanocomposites

having uniformly dispersed and highly exfoliated GnPs within the polymer. The SCF-assisted

method consolidates the subsequent steps of (i): exfoliation of GnPs and (ii): their dispersion

within the polymer, both in one step. This is greatly in favor of large-scale production of

multifunctional GnP-based polymer nanocomposites with a reasonable cost. The findings of this

research are also critically important to the advancement and optimization of industrial-scale

processing of GnP-based polymer nanocomposites with tailored properties for various

applications.

62

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CHAPTER 4

4 Enhancement of the thermal conductivity of polymer-GnP composites via facile SCF-assisted manufacturing


Hamidinejad, S.M., Chu, R.k.M., Zhao, B., Park, C.B., and Filleter, T., “Enhanced Thermal

Conductivity of Graphene Nanoplatelet-Polymer Nanocomposites Fabricated via Supercritical Fluid

Assisted In-Situ Exfoliation”, ACS Applied Materials and Interfaces, 2018, 10 (1), 1225−1236

4.1 Summary

As electronic devices become increasingly

miniaturized, their thermal management becomes

critical. Efficient heat dissipation guarantees their

optimal performance and service life. Graphene

nanoplatelets (GnPs) have excellent thermal properties

that show promise for use in fabricating commercial

polymer nanocomposites with high thermal

conductivity. Herein an industrially-viable technique

for manufacturing a new class of lightweight polymer-GnP composites with high thermal

conductivity is presented. Using this method, high-density-polyethylene (HDPE)-GnP

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nanocomposites with a microcellular structure are fabricated by melt mixing, which is followed

by supercritical fluid (SCF)-treatment and injection molding foaming which adds an extra layer

of design flexibility. Thus, the microstructure is tailored within the nanocomposites, and this

improves their thermal conductivity. Therefore, the SCF-treated HDPE-17.6 vol.% GnP

microcellular nanocomposites have a solid-phase thermal conductivity of 4.13±0.12 Wm-1K-1.

This value far exceeds that of their regular injection-molded counterparts (2.09±0.03 Wm-1K-1)

and those reported in the literature. This dramatic improvement results from an in-situ GnPs’

exfoliation and dispersion, and from an elevated level of random orientation and

interconnectivity. Thus, this technique provides a novel approach to the development of

microscopically-tailored structures for the production of lighter and more thermally conductive

heat sinks for the next generations of miniaturized electronic devices.

4.2 Introduction

Heat dissipation functionality is extremely critical in high-energy density systems such as next-

generation miniaturized electronic devices [1]. The continuous development of the smaller,

lighter, and faster electronic components of such devices means that the heat they generate

needs to be efficiently dissipated by more compact and lightweight heat sinks. Lightweight,

multifunctional, low cost, and highly thermally conductive polymer composites show promise

for use as heat dissipation components [2]. When compared with metallic and ceramic

composites, polymer composites have an attractive array of properties, including ease of

processing, superior resistance to chemicals and corrosion, and tailorable physical/mechanical

properties[3–5]. The thermal conductivity of polymer composites is intensely affected by their

interfacial thermal resistance and interfacial phonon scattering [6,7], by their dispersion and

orientation, and by the type of fillers used [8].

Conventionally, thermally conductive polymer composites are filled with a high loading (50−80

vol.%) of micro-size fillers to achieve target thermal conductivity values (>1 Wm-1K-1) [9].

With such a high filler loading level, however, the amount of polymer matrix left to support the

fillers and the composite’s structural integrity is insufficient. This leads to expensive and

heavyweight composites, which are difficult to process. One promising way to address this

67

drawback is to incorporate nanomaterials with extraordinary thermal conductivity, higher aspect

ratios, and mechanical properties during the creation of these composites.

With the recent advances in nanomaterials and their growing availability, the types and

functions available for polymer composites have been significantly increased. This has

increased the opportunities to develop polymer nanocomposites with superior thermal

conductivity. Extraordinary heat transport properties in such nanomaterials as graphene, carbon

nanotubes (CNT), and boron nitride nanotubes (BNNT) have driven the research of polymer

nanocomposites. However, the expected dramatic enhancement of thermal conductivity by the

incorporation of CNTs [8] and BNNTs [10] has not yet materialized in polymer

nanocomposites, even at very high additive loading levels. Despite the excellent thermal

conductivity reported for individual nanotubes[11], CNTs and BNNTs have not been shown to

substantially improve the thermal transport properties of polymer nanocomposites [8,10]. This

has been attributed to the nanotubes’ one-dimensional nature, which leads to their having

anisotropic thermal conductivity in the axial direction [11–13]. However, it has been suggested

[8,14,15] that 2D nanomaterials such as graphene can be a more effective nanomaterial for

polymer nanocomposites with a high thermal conductivity.

In recent years, graphene has attracted a great deal of attention due to its exceptional

mechanical, electrical, and thermal properties. Notably, the thermal conductivity of single-layer

graphene has been reported, as ∼5000 W/ (m.K) [16–18]. However, the practical underpinning

needed to economically manufacture graphene-based polymer composites is missing. It has been

extremely challenging to exploit graphene’s full potential. This has been due to the complexities

that exist in the exfoliation, dispersion, and control of the GnPs’ orientation within the

composites [19].

Various strategies, such as in-situ polymerization [20,21], GnP surface modification [9,22], GnP

alignment by electrical field [23], and the use of hybrid additives [2,9] have all been proposed to

develop polymer composites with high thermal conductivity. Table 4.1 summarizes some of the

recent advances made in the development of thermally conductive polymer nanocomposites.

Nevertheless, all of these fabrication techniques have been batch-type processes. This makes

them expensive, time-consuming, and not easily scalable. Furthermore, in most cases, the

required additive loading levels remain rather high.

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On the other hand, SCF-treatment and physical foaming have shown promise in enhancing the

electrically conductive polymer composites’ functionalities in different applications [3,24–31].

Incorporating the optimum microcellular foaming structures into the conductive polymer

composites can significantly reduce the product’s weight. At the same time, it can also add

another degree of design flexibility to help control the polymer composites’ functional

properties. During foaming, cell growth can change both the alignment and orientation fillers

around the growing bubbles through the biaxial stretching of the polymer matrix

[3,24,28,29,32,33]. Furthermore, applying the SCF-treatment and physical foaming to the

polymer composites can enhance the dispersion [31,32,34] and distribution [28,29] of the

additives in the polymer matrix. It also can lower the fillers’ mechanical breakdown [29,30]

during processing. In this way, the optimized SCF-treatment and microcellular foaming can

introduce tailored structures that support the various functionalities such as electromagnetic

interference shielding effectiveness [25,27,29,30,35,36], electrical conductivity,[3,24,28–30]

and the dielectric properties of conductive polymer composites [3,24]. However, to the best of

our knowledge, no attention has been paid to the role of SCF-treatment in promoting heat

dissipation in thermally conductive polymer composites.

In contrast to the batch-type methods (Table 4.1), injection molding is a common and

economically-viable industrial technology used to manufacture polymer parts. Therefore,

injection molding combined with a SCF-treatment of polymer composites can be an easy

solution to generate tailored microstructures that improve the heat dissipation properties in

graphene-based polymer nanocomposites. However, to the best of our knowledge, no effort has

yet been reported on the heat dissipation performance of injection-molded microcellular

nanocomposites.

Our study demonstrates a SCF-assisted manufacturing method for producing thermally

conductive HDPE-GnP composites by using injection molding to create heat dissipation

components. The SCF-treated microcellular HDPE-GnP composites exhibited heat transport

properties that were remarkably superior to those of the regular injection-molded

nanocomposites [41,42]. Furthermore, they are comparable to the overall heat transport

performances of bath-type methods reported in the literature [2,9,20–23,37–39,41]. This was

due to the tailored microstructure of the HDPE-GnP composites created by the proposed

technique.

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Table 4.1. Thermal conductivity of various batch-type graphene/polymer nanocomposites

materials filler content thermal conductivity

[Wm-1k-1] fabrication method ref

Epoxy/GnP 25 vol% 12.4 Surface treatment + planetary centrifugal

mixing

[9]

Epoxy/GnP 10 wt.% 1.53 Functionalization, solution mixing and a

curing process

[22]

CBT/GnP 20 wt.% 7.1a Solvent-free melting process followed by

in-situ polymerization

[37]

PLA/hBN/GnP 16.65/16.65 vol% 2.77 Melt mixing and compression molding [2]

PVDF/GnP 20 wt.% 0.562 High-shear solution mixing followed by

bath-sonication

[23]

PA-6/GnP 10 wt.% 0.416 In situ polymerization with simultaneous

thermal reduction

[20]

PA-6/GnP 12 wt.% 2.49 One-step in situ intercalation

polymerization

[21]

PC/GnP 20 wt.% 1.76 Melt mixing and compression molding [38]

SBR/GnP 24 vol.% 0.48 Solution mixing and a sonication [39]

PA6/hBN/GnP GnP /1.5/20 wt.% 1.76 Liquid exfoliation, solution blending and

hot-pressing

[40]

a In-plane Thermal conductivity



A commercially available HDPE (Marlex® HHM 5502BN with a melt flow index 0.35 dg/min.-

1 at 230 ºC/2.16 kg) with a specific gravity of 0.955 g cm-3 was used as the polymer matrix. The

HDPE was filled with GnP grade heXo-g-V20, with average lateral dimensions of 50 μm, a

surface area of 30 m2/g, and a specific gravity of 2.2 g.cm-3 (Group NanoXplore Inc. Montreal,

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QC, Canada). Commercial nitrogen (N2), supplied by Linde Gas, Canada, was used as the

environmentally friendly SCF.

The HDPE nanocomposites with a different GnP content were made by diluting HDPE-35 wt.%

GnP masterbatch with the as-received neat HDPE through mixing in a twin-screw extruder (27

mm, L/D:40). A 50-ton Arburg Allrounder 270/320C injection molding machine (Lossburg,

Germany), with a 30-mm diameter screw equipped with MuCell® technology (Trexel, Inc.,

Woburn, Massachusetts) was used to fabricate the HDPE-GnP composite samples. Three

different types of HDPE-GnP nanocomposites, namely injection-molded solid (IMS), injection-

molded foam (IMF) and high-pressure-injection-molded foam (HPIMF), were prepared. Details

of preparing samples were discussed in Section 3.3.1.1 and Table 3.1. The schematic of the

injection molding processes (i.e. IMS, HPIMF and IMF) and injection-molded parts were

presented in Figure 3.1.

A die cutter was used to cut disk-shape samples with a 20 mm diameter × 3 mm thickness from

the injection-molded nanocomposites at a distance of 100 mm from the cavity gate. The

schematic of the injection molding processes (i.e. IMS, HPIMF and IMF) and injection-molded

parts have been presented in Figure 3.1. The actual degrees of foaming in the IMF samples were

measured using samples via the water-displacement method (the ASTM D792-00) after

fabrication.

4.4 Characterization

The relative GnP powder’s defects, and an estimation of the number of layers it had, were

determined using Raman spectroscopy (Renishaw, 532 nm laser excitation). X-ray

Photoelectron Spectroscopy (XPS) was also conducted on the GnP powder to identify its surface

chemistry and functional groups, and also to measure the C/O ratio. The results of Raman

spectroscopy and XPS on the GnP powder are discussed in Supporting Information. An X-ray

photoelectron spectrometer (ThermoFisher Scientific K-Alpha) equipped with an AlKα X-ray

source was used to collect XPS data to analyze the qualitative defect density of the GnPs. To

examine the exfoliation and dispersion of GnPs in the polymer matrix, Wide Angle X-ray

Diffraction (WAXD) analyses were conducted on the injection-molded nanocomposites using a

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Rigaku MiniFlex 600 X-ray diffractometer (Cu Kα radiation, λ = 1.5405 Å). To further evaluate

the level of exfoliation and dispersion of different samples, transmission electron microscope

(TEM; FEI Tecnai 20) were conducted. The TEM samples were prepared by cryo-

ultramicrotomy (Leica EM FCS). The microstructure and morphology of the fabricated samples


were frozen in liquid nitrogen, cryofractured, and sputter-coated prior to electron microscopy.

The thermal conductivities of the polymer-GnP composites were measured using the transient

hot disk method. A transient plane source (TPS) hot disk thermal constants analyzer (Therm

Test Inc., TPS 2500, Sweden) was used to measure the samples’ thermal conductivity under

ambient conditions with a Kapton (C7577) sensor. Measurements were taken based on the

ISO/DIS 22007-2.2 standard. In this method, an electrically conductive double spiral disk-shape

sensor made of nickel foil works as both a heater, to increase the temperature, and a dynamic

thermometer to record the change in samples’ temperature as a function of time. The sensor is

placed between two pieces of the sample and the increase in the samples’ temperature is

evaluated by the analyzer to calculate the thermal conductivity (See Figure 4.1). Therefore, the

generated heat will be dissipated in any direction (e.g. through-plane and in-plane) and the

measured thermal conductivity is the overall (total) thermal conductivity [43,44].

Figure 4.1. The schematic of the ISO/DIS 22007-2.2 setup for measuring the thermal

conductivity using TPS 2500

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4.5 Results and discussion

4.5.1 Defect density of GnPs

The level of defect density greatly affects GnPs’ inherent thermal conductivity [4]. The GnPs’

D-to-G peak intensity ratios of Raman spectroscopy below 0.3, suggest that the GnPs have low

defect densities [45]. For the GnPs we tested, this value was about 0.1, as shown in Figure 4.2a.

This result was in good agreement with the XPS measurement. The deconvoluted C 1s XPS

spectra (Figure 4.2b) showed that various carbon bonds had different energy values: C=C bonds

(∼284.6 eV), C-O (∼286.1 eV) and O-C=O bonds (∼290.1 eV). The C 1s spectra also showed

that the sp2 C=C bonds’ intensity, which is critical for the GnPs’ superior thermal conductivity,

was much higher than both the C-O and the O-C=O bonds. Based on the XPS results, the

qualitative defect density of the GnPs was 23% (qualitative defect density =100%-sp2% [9]).

Figure 4.2. (a) Raman spectroscopy of the GnPs; (b) deconvoluted C 1s XPS spectra. Raman

spectra of GnP

73

The most pronounced peaks in the Raman spectra, which were D, G and 2D, were observed at

wavelengths of 1350, 1580, and 2700 cm-1, respectively. It has been reported that the 2D peak

clearly evolves along with the number of graphene layers (up to 10 layers) in the GnPs.

However, the 2D peaks for graphene nanoplatelets with more than 10 layers are quite similar

[46]. The 2D peak for bulk graphite consists of two peaks (roughly ½ and ¼ of the G peak

intensity). However, the 2D peak for a single-layer graphene is single-peak and roughly 4 times

more intense than the G peak [46]. As shown in Figure 4.2a, it is noteworthy that the 2D peak

(deconvoluted using the Gaussian model) consisted of two distinct peaks. This suggested that

the GnPs used in this study have more than 10 layers of graphene.

4.5.2 Microstructure and morphology of polymer-GnP composites

The IMS samples were fabricated without SCF-treatment and physical foaming. In the IMF

samples, we first obtained a single-phase gas-Gn P-polymer mixture, by SCF-treatment. When

the mold cavity was partially filled with this mixture, physical foaming occurred due to the

depressurization process. However, in the HPIMF samples the mold cavity was fully filled with

the same mixture. Yet, as had happened with the IMF samples, the physical foaming occurred

when the mixture had entered the mold cavity. However, the next step occurred under high

pressure, and the nucleated cells re-dissolved back completely into the polymer-GnP mixture.

The SCF-treatment and physical foaming produced a tailored microcellular structure, which

increased the GnPs’ exfoliation and random orientation. A thinner skin layer also resulted.

4.5.2.1 Effects of SCF-treatment on the cellular microstructure, GnPs’

orientation, and skin layer

Figure 4.3a shows the skin and core microstructure of the IMS HDPE-9 vol.% of GnP and that

of its counterparts: HPIMF and IMF (7% degree of foaming). As expected, the IMS samples’

core and skin layers had completely solid structures. The IMF samples had a microcellular

structure with a random cell morphology in both the skin and core layers, with an average cell

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size of 3 μm. The HPIMF samples’ structure was almost solid. And its cellular structure was

barely visible due to the nucleated cells’ redissolution under high pressure.

In IMF samples, cell growth caused different degrees of GnP rotation and displacement, which

led to the GnP’s random orientation and further dispersion. To be specific, the GnPs get oriented

more perpendicular to the radial direction with bubble growth and, consequently, the GnPs

come to meet each other along the bubble surface. In other words, there was a greater chance of

interconnectivity and direct GnP-GnP contact. This led to a particular morphology in which the

IMF samples were greatly differentiated from the flow-induced structure found in the IMS

samples.

In the IMS samples’ skin layers (about 500 μm on each side), the GnPs were aligned in the

machine direction (Figure 4.3a). This was due to the rapid cooling and the applied shear stresses

in the direction of flow during the melt injection. This preferred filler alignment in the

composites fabricated via injection molding has been well covered in the literature

[24,28,29,47]. In the IMS samples’ core layer, the GnPs followed the fountain flow orientation

and were relatively randomly oriented. The HPIMF samples had the same skin-core

morphology; however, the skin layer was thinner compared to their IMS counterparts’ (about

350 μm on each side). The composite melt’s lower viscosity, which was due to the SCF-

treatment, had reduced the GnPs’ flow-induced orientation in the skin layer. This resulted in a

lower skin layer thickness with oriented GnPs. Similar phenomenon has also been found in

polymer/fiber composites [3,28,29,47]. Conversely, in the IMF samples, the skin-core

morphology and the oriented skin layer were hardly identified. This can be attributed to not only

the composite melt’s lower viscosity but also to physical foaming. Figure 4.3b shows the ideal

2-D conceptualization of the evolution of the GnPs’ interconnectivity, orientation, and their

further exfoliation due to SCF-treatment and physical foaming.

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Figure 4.3. (a) SEM micrographs of skin and core regions for IMS, HPIMF and IMF HDPE-9

vol% GnP nanocomposites. Scale bars are all 10 μm; (b) ideal 2-D conceptualization of the

evolution of GnPs interconnectivity, orientation and further exfoliation due to SCF-treatment

and physical foaming; (c) SEM micrographs of IMF HDPE-9 vol% GnP nanocomposites

showing different types of cells generated in the microstructure. FD stands for flow direction.

In the IMF samples, three different types of cells were found: (i) small cells that had nucleated

in the polymer matrix, which led to cells with polymeric walls (shown in Figure 4.3c by green

76

circles); (ii) cells that had nucleated at the edge, or on the surfaces, of the GnPs, which acted as

nucleating agents, and which led to the cells being surrounded by a combination of polymeric

walls and GnPs (shown in Figure 4.3c by yellow arrows); and (iii) cells formed by the phase

transition of the SCF to a gaseous state, which led to cells encompassed only by GnPs (shown in

Figure 4.3c by red arrows). To more clearly elucidate the microcellular structure, we have

conducted additional electron microscopy imaging for IMF samples with lower magnifications

which are presented in Figure 4.4.

Figure 4.4. (a) SEM micrographs of the FIM -9 vol% HDPE-GnP composites with 7% degree of

foaming; (b) Zoomed-in SEM micrographs of Figure 4.4a.

Because of the non-homogeneity of the structure with the dispersed and distributed GnP

particles, the observed cells were quite non-homogeneous as shown in Figure 4.3c. This non-

77

homogeneity cannot be explained with the non-homogeneous growth alone because some cell

walls have a clean GnP surface. The bubbles must have been nucleated very non-

homogeneously and non-uniformly. It is quite well accepted that the heterogeneous cell

nucleation scheme will be preferred at the interface because of the lower activation energy for

cell nucleation[48]. The clean surfaces of the cavities observed from Figure 4.3c indicate that

those cells were nucleated at a surface of the GnP, based on the heterogeneous cell nucleation

mechanism. It is clear from Figure 4.3c that most cells were nucleated this way. The size of this

type of cells approximately ranges from 3μm to 20μm. But we could also observe the smaller

cells (1μm) nucleated inside the polymer matrix alone because these cells were completely

encapsulated by the polymer melt (see the green color circles in Figure 4.3c). On the other hand,

Figure 4.3c also shows the other category cavities that were neither formed at the polymer-GnP

interface, nor inside the polymer matrix. In fact, there are so many of these types of cavities that

are observed from the SEM images. Since these cavities’ boundaries are GnP particles alone,

not a polymer melt, these must have been formed by the expanding action of the SCF that

diffused into the GnP layers before expansion (see cavities shown by red arrows in Figure 4.3c).

The size of these bubbles is approximately 1-20μm.

4.6 Thermal conductivity

4.6.1 Effect of the GnP content on the thermal conductivity

Figure 4.5a shows the total thermal conductivity of the IMS samples as a function of their GnP

content. Their thermal conductivity is reported as a function of the final GnP content. The GnPs’

volume percent was calculated with respect to the total volume of foamed HDPE-GnP

composites, including both gaseous and solid phases. As we had expected, in all of the samples,

including the IMS, HPIMF and the IMF, the thermal conductivities of the HDPE-GnP

composites were highly dependent on the GnP content.

In the IMS samples, the total thermal conductivity corresponded to a 404% increase over the

neat HDPE samples at an 18 vol.% of GnP and an increase of 21.3% per 1 vol.% GnP loading.

This accorded with the enhancement efficiency of such traditional fillers as graphitic

78

microparticles, which typically show an increase of ∼20% per 1 vol.% filler loading [9,49].

However, this value for HPIMF and IMF samples was 31% and 46% respectively.

On the other hand, for the IMF samples, introducing a 26% degree of foaming into the neat

HDPE (i.e., at zero GnP loading), would decrease the thermal conductivity by 50% because of

the creation of the voids with low thermal conductivity (0.026 Wm-1K-1 for ambient air) [50].

But, interestingly, as the GnP loading increased, the detrimental effect of physical foaming on

the IMF nanocomposites’ thermal conductivity became insignificant at around 7 vol.% GnP

loading. The thermal conductivities of the IMF nanocomposites started to outpace those of the

IMS and HPIMF composites at a GnP loading of more than 7 vol.%. We attributed this to a

sufficiency of GnPs in the polymer nanocomposites to form thermally conductive paths. In other

words, below a 7 vol.% GnP loading, the polymer matrix mediated between the GnPs. The

result was a polymeric gap that broke the direct GnP-GnP contact. This caused the phonon

scattering and high interfacial thermal resistance [9,51].

4.6.2 Effect of GnP’s exfoliation and dispersion on the thermal conductivity

It is interesting to note that the SCF-treated HPIMF counterparts’ total thermal conductivity was

563% greater than the neat HDPE samples at an 18 vol.% of GnP, and there was an increase of

31% per 1 vol.% of GnP loading. The increase in the HPIMF samples’ thermal conductivity

over that of the IMS samples can be attributed largely to their higher level of GnP exfoliation,

when compared to their IMS counterparts (Figure 3.3). In other words, at the same GnP loading

level, the number of effective GnPs in the HPIMF samples was greater than the number of GnPs

in the IMS samples due to a higher level of exfoliation with the SCF treatment. This increased

the chance for direct GnP-GnP contact, which has a much lower interfacial thermal resistance

than a polymer mediated structure would have (GnP-polymer contact) due to a lower amount of

phonon scattering [9,51]. Consequently, thermally conductive paths were formed more likely.

It should be emphasized that the degree of foaming of HPIMF samples was almost negligible

because of the high packing pressure (~300 MPa) used in the process. So, there would be

negligible effect of the foaming on the thermal conductivity for the HPIMF samples.

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Figure 4.5. (a) The total thermal conductivity (λtotal) of IMS, HPIMF, and IMF HDPE-GnP

composites as a function of the GnP content and; (b) the thermal conductivity of IMS, HPIMF,

and IMF samples (HDPE-9 vol.% GnP) before (total) and after removing their skin (core); (c)

the total thermal conductivity (λtotal) of IMS, HPIMF, and IMF HDPE-GnP composites as a

function of the degree of foaming and the GnP content; (d) the total thermal conductivity (λtotal)

of the samples as a function of the degree of foaming (GnP vol. % has been reported with

respect to the polymer volume)

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4.6.3 Effects of GnPs’ re-orientation on the thermal conductivity of HDPE-

GnP composites

It is also interesting to note that the IMF samples exhibited much higher thermal conductivities

than their IMS and HPIMF counterparts, indicating that the local interconnectivity and the

amount of direct GnP-GnP contact became much higher with foaming. This was due to the

reduced orientation of the GnPs in the flow direction as well as the re-orientation of the GnPs

surrounding the bubbles, caused by the foaming action that occurred in the IMF samples. This

resulted in a lower interfacial thermal resistance than what would be found in a polymer

mediated structure.

For example, the total thermal conductivity of the IMF samples (with 7% degree of foaming)

exhibited a higher increase of 46% per 1 vol.% GnP loading over the IMS samples with an

increase of 20% per 1 vol.% GnP loading. Likewise, the total thermal conductivity of the IMF

samples was also significantly higher than that of the skinless HPIMF samples. This outstanding

improvement in the IMF samples’ total thermal conductivity over the IMS and HPIMF samples

was attributed to the re-oriented GnPs’ microstructure in which the IMF samples were greatly

differentiated from their IMS and HPIMF counterparts.

Moreover, reduction of the GnPs’ orientation in the skin layer provided more isotropic heat

transport functionality. The skin-core morphology, with highly oriented GnPs in the skin, was

much more pronounced in the unfoamed IMS and HPIMF than in the foamed IMF samples

(Figure 4.3). This resulted in a highly anisotropic heat dissipation property which deteriorated

the product’s total thermal conductivity. It is worthy of noting Gong et al.’s claim [52] that too

high orientation of the conductive fibers will increase the percolation threshold even in the

oriented direction. In fact, we observed an increased thermal conductivity from 1.20±0.01 to

1.40±0.04 Wm-1K-1 for the IMS samples with a 9 vol.% of GnP, after we removed, by

machining, their skins with highly oriented GnPs (see

Figure 4.5b).

On the other hand, the higher thermal conductivity of the HPIMF samples over the IMS samples

discussed in Section 4.6.2 may also have been affected by the lower thickness of the skin layer

with highly oriented GnPs. Because of the reduced viscosity with the SCF treatment, the skin

81

layer of the HPIMF would be reduced and, therefore, the thinner skin layer with highly oriented

GnPs will increase the conductivity. As shown in

Figure 4.5b, after removing the skins of the HPIMF samples (with 9 vol.% GnP), the thermal

conductivity of the parts increased from 1.47±0.04 to 1.62±0.06 Wm-1K-1. However, the thermal

conductivity of the IMF counterparts remained approximately constant after removing the skin

layer (2.09±0.01 to 2.12±0.02Wm-1K-1). This is caused by the re-orientation of GnPs in the IMF

samples due to the physical foaming leading to a more isotropic structure as compared to the

IMS and HPIMF counterparts.

In a nutshell, the enhanced thermal conductivity of the polymer-GnP composites with foaming

was attributed to: (i) the reduced orientation of the GnPs in the flow direction; (ii) an increased

local interconnectivity among the GnPs surrounding each bubble; and (iii) reduced orientation

of the GnPs in a thinner skin-layer.

It is notable that the crystallinities of the IMS, HPIMF and IMF samples are very similar (see

Figure 4.6 showing Differential Scanning Calorimetry (DSC) and High-Pressure Differential

Scanning Calorimetry (HPDSC)) on the HDPE-4.5 vol.% GnP). We also investigated the effects

of the dissolved gas on the crystallinity of the HDPE-4.5 vol.% GnP using HPDSC. To

investigate the non-isothermal crystallization in HPDSC, the HDPE-4.5 vol.% samples were

heated and equilibrated at 200 °C for 30 min. The heating and thermal history removals were

implemented under the N2 pressures of 1 and 48 bars. Then, the samples were cooled to 30 °C at

a cooling rate of 10 °C/min, under N2 pressures of 1 and 48 bars in the HPDSC. We observed

that the crystallinities and crystallization temperatures at different N2 pressures were very

similar. This result was in good agreement with the DSC results. This can be attributed to the

very fast crystallization kinetics of HDPE which may not have been significantly affected by the

parameters studied in this work. Therefore, we believe that the effect of crystallinity on the

thermal conductivity in this study is negligible.

82

Figure 4.6. (a) Differential Scanning Calorimetry (DSC) of the IMS, HPIMF and IMF sample

(HDPE-4.5 vol.% GnP); and (b) High Pressure Differential Scanning Calorimetry (HPDSC) of

HDPE-4.5 vol.% GnP samples

4.6.4 Optimal degree of foaming on the thermal conductivity

Although foaming can enhance the thermal conductivity of the polymer-GnP composites, too

high degree of foaming would be undesirable because of the non-conductive nature of the voids.

This indicates that there exists an optimal degree of foaming to maximize the thermal

conductivity.

Figure 4.5 (c and d) shows the thermal conductivity variations with the GnP loading and the

degree of foaming for the HDPE-GnP composites. When a 7% degree of foaming was

introduced to the HDPE-GnP composites in the IMF samples, the total thermal conductivity was

increased from 2.09±0.03 to 3.75±0.12 Wm-1K-1 at 17.6 vol.% of GnP. However, increasing the

degree of foaming to beyond 7%, decreased the total thermal conductivity. This optimal

behavior is attributed to the competing relationship between the favorable GnPs’ re-orientation

effects (as discussed in Section 4.6.3) and the voids’ ultralow thermal conductivity. An

excessive degree of foaming (that is, 16% and 26% in the current case) resulted in higher voids

in the structure, which led to a lower thermal conductivity. Thus, we lean to conclude that the

two competing mechanisms mentioned above govern the total thermal conductivity and that 7%

was the optimal degree of foaming.

83

4.6.5 Solid phase thermal conductivity

The total thermal conductivities of IMF samples were affected by two antagonistic parameters

which included: (i) a constructive tailored morphological structure in the solid phase; and (ii) the

negative impacts of insulating voids in the structure. To further analyze the net effects of the

SCF-treatment and the foaming actions on the HDPE-GnP’s intrinsic thermal conductivity

improvement, a theoretical model was employed in order to exclude the effect of voids and

determine the thermal conductivity of the solid phase alone. This theoretical work was

undertaken to quantitatively clarify the positive and negative effects of the voids on the total

thermal conductivity.

The convection that results from gas movement within the cells is negligible if the cell sizes are

less than 4-5 mm [53]. The contribution of radiation to the thermal conductivity of cellular

plastics is less than 5% if their relative density are greater than 0.3 [53]. It has also been reported

that carbonaceous materials as the infrared attenuated agents (IAA) can block the radiation [54].

The IAAs reported on in different studies can include surface-modified nano-graphite

particulates [55], carbon nanotubes [56,57], and dispersed graphene fillers [58]. Therefore, the

contribution of the radiative heat transfer does not apply to this study.

In the IMF nanocomposite samples, the heat flux must pass through either the solid phase

(HDPE-GnP phase) or through the gaseous phase. Then, the total thermal conductivity (λtotal) of

the IMF nanocomposites includes the solid conductivity (λsolid), and the gas conductivity (λgas)

and is expressed as follows [50]:

λ𝑡𝑜𝑡𝑎𝑙 = λ𝑠𝑜𝑙𝑖𝑑 + λ𝑔𝑎𝑠 4.1)

However, in confined spaces, the gas molecule collisions become lower. Thus, the gas

conduction is governed by the energy transfer between the cell walls and the gas molecules. The

Knudsen number (Kn) is defined to relate the dependency of the gas conductivity to the cell sizes

as follows [59]:

d

lK mean

n = (4.2)

(2)

84

where d is approximated by the cell size, and lmean is the mean free path of gas molecules, which

is 68 nm in the ambient condition. The gas conductivity in polymeric foams is as follows [59]:

0

21

1gas

n

gas kBK

k+

= (4.3)

where B is the energy transfer efficiency between the cell walls and the gas molecules,1.94. The

0

gask is the bulk gas’s conductivity, which is 26 mWm-1K-1 for air [56].

We used the Maxwell-Eucken I model [60] in this study. This model is suitable for materials in

which the thermal conductivity of the dispersed phase is lower than the continuous phase (i.e.,

ksolid > kgas) such as polymeric foams [61]. The Maxwell-Eucken I model is expressed as

follows:

ggassolidgassolid

ggassolidgassolid

solidtotalkkkk

kkkkk

)(2

)(22

−++

−−+= (4.4)

where, ksolid and kgas respectively represent the thermal conductivity of the solid phase (HDPE-

GnP) and the gaseous phase. The ʋg is the degree of foaming (that is, the void fraction) of the

IMF samples.

Based on the cell sizes, the gas conductivities (kgas) of the IMF HDPE-GnP composites are

calculated via Equation (3.3). The calculated kgas, the measured values of the ʋg, and the λtotal

(that is, the total thermal conductivity of the IMF HDPE-GnP composites shown in

Figure 4.5) are substituted in Equation (3.4). The thermal conductivity of the solid phase (ksolid)

was then calculated and is plotted in Figure 4.7.

85

Figure 4.7. Solid phase thermal conductivity (ksolid) of IMS, and IMF HDPE-GnP composites as

function of (a) the GnP content and; (b) the degree of foaming and the GnP content. DF stands

for degree of foaming

Figure 4.7a presents only the thermal conductivities of the solid phase, which were extracted

from the IMF samples’ thermal conductivity using the Maxwell-Eucken I model. We note that

the thermal conductivities of the solid phase in all of the IMF samples with various degrees of

foaming (i.e., 7%, 16% and 26%) coincided at approximately the same values. Thus, it was

possible to evaluate the actual efficiency of the SCF-treatment and the physical foaming in the

thermal conductivity enhancement of the HDPE-GnP composites.

86

The thermal conductivity of IMF nanocomposites’ solid phase showed up to a 1,000%

enhancement over the neat HDPE samples and an enhancement of 56% per 1 vol.% loading of

GnP. In Figure 4.7b, the thermal conductivity of the IMF samples’ solid phase remained almost

intact with a change in the degree of foaming. This occurred when the HDPE-GnP composites

were first treated with the SCF, and then underwent the physical foaming. It is also noteworthy

that the thermal conductivities of the IMF HDPE-GnP composites in the solid phase were higher

than those of their IMS counterparts, even at GnP loadings of less than 7%. However, the

difference between the thermal conductivities of the IMF’s solid phase and IMS samples was

more pronounced with a higher GnP content.

4.7 Conclusion

In our study, we introduced a new class of highly thermally conductive microcellular polymer-

GnP composites. Microcellular nanocomposites containing highly exfoliated GnPs were

developed by an industrially-viable technique of melt mixing followed by SCF-treatment and

physical foaming in an injection molding process. This process provided a tailored structure that

effectively supported the improved thermal conductivity of polymer-GnP composites. For

example, the SCF-treated HDPE-17.6 vol.% GnP nanocomposites had a solid thermal

conductivity of 4.13±0.12 Wm-1K-1 which was vastly superior to the values of their regular

injection-molded counterparts (2.09±0.03 Wm-1K-1) as well as to those reported in the literature

[41,42]. The reasons for this dramatic improvement include the following: (i) a higher level of

GnPs’ exfoliation and dispersion in the polymer matrix; (ii) a decreased degree of GnP

orientation from the reduced viscosity and the foaming action; (iii) an increased local

interconnectivity among the GnPs surrounding each bubble; and (iv) a reduced skin-layer

thickness. Our research shows that SCF-treatment of HDPE-GnP composites can add an extra

layer of design flexibility in the manufacture of polymer-GnP composites with tailored

morphologies and thermal conductivity. This design can be readily scaled up to an industrial

level to make efficient and lightweight thermally conductive products for heat dissipation

components in various miniaturized electronic devices.

87

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94

CHAPTER 5

5 Enhancement of electrical and electromagnetic interference (EMI) shielding properties of the polymer-GnP composites


Hamidinejad, S. M., Zhao, B, Zandieh, A., Moghimian, N., Filleter, T., and Park, C.B., “Enhanced

Electrical and Electromagnetic Interference Shielding Properties of Polymer-Graphene Nanoplatelet

Composites Fabricated via Supercritical-fluid Treatment and Physical Foaming”, ACS Applied

Materials and Interfaces, 2018, DOI: 10.1021/acsami.8b10745

5.1 Summary

Lightweight high-density

polyethylene (HDPE)-

graphene nanoplatelet

(GnP) composite foams

were fabricated via a

supercritical-fluid (SCF)

treatment and physical

foaming in an injection-molding process. We demonstrated that the introduction of a

95

microcellular structure can substantially increase the electrical conductivity and can decrease the

percolation threshold of the polymer-GnP composites. The nanocomposite foams had a

significantly higher electrical conductivity, a higher dielectric constant and a higher

electromagnetic interference (EMI) shielding effectiveness (SE) and a lower percolation

threshold compared to their regular injection-molded counterparts. The SCF treatment and

foaming exfoliated the GnPs in situ the fabrication process. This process also changed the GnP’s

flow-induced arrangement by reducing the melt viscosity and cellular growth. Moreover, the

generation of a cellular structure rearranged the GnPs to be mainly perpendicular to the radial

direction of the bubble growth. This enhanced the GnP’s interconnectivity and produced a

unique GnP arrangement around the cells. Therefore, the through-plane conductivity increased

up to a maximum of nine orders of magnitude and the percolation threshold decreased by up to

62%. The lightweight injection-molded nanocomposite foams of 9.8 vol.% GnP exhibited a real

permittivity of ε'=106.4, which was superior to that of their regular injection-molded (ε'=6.2). A

maximum K-band EMI SE of 31.6 dB was achieved in HDPE−19 vol. % GnP composite foams,

which was 45% higher than that of the solid counterpart. In addition, the physical foaming

reduced the density of the HDPE-GnP foams by up to 26%. Therefore, the fabricated polymer-

GnP nanocomposite foams in this study pointed towards the further development of lightweight

and conductive polymer-GnP composites with tailored properties.

5.2 Introduction

Polymer composites have shown impressive potential as a highly desirable class of advanced

functional materials for use in various applications such as capacitors (dielectric materials [1]),

electromagnetic interference (EMI) shielding [2,3], electro-static dissipation [4], and energy

conversion (bipolar plates of fuel cells [5,6]). Polymer composites offer tailorable electrical,

thermal, and mechanical properties. They are also low cost, offer ease of processing, and their

chemical resistance is superior to their metallic and ceramic counterparts [6–8]. The recent

advances in conducive nanofillers such as graphene have significantly increased the

opportunities to develop polymer nanocomposites with tailored functionalities [1,2,9].

Graphene provides a unique combination of exceptional electrical, thermal, and mechanical

properties. Notably, the electrical conductivity of monolayer graphene has been reported as

96

∼6,000 S/cm [8]. One major class of graphene-based polymer nanocomposites are those which

take advantage of the electron transport characteristics of graphene for applications such as EMI

shielding, where the focus has been on achieving a higher electrical conductivity at lower

graphene concentrations [10]. EMI shielding of radio frequency radiation is a serious concern in

our technological society and graphene has attracted great attention for the fabrication of

efficient EMI shields [11–14].

Polymer-graphene nanocomposites also exhibit promise for use as dielectric materials with high

dielectric permittivity (ε') and low dielectric loss (tan ) for high-performance capacitors [1].

The high electrical conductivity of graphene, when compared to that of the polymer matrix,

results in interfacial polarization and, consequently, improved the dielectric permittivity [15,16].

However, the dielectric properties of percolative polymer nanocomposites change significantly

near the percolation threshold. The dielectric loss abruptly increases due to the formation of

conductive paths throughout the composite system. Therefore, the dielectric properties of the

percolative polymer composites need to be optimized within an “adjustable window” near the

percolation threshold, where the dielectric constant can be enhanced while the dielectric loss is

still limited [17]. This is, however, extremely challenging [18]. In addition, reaching graphene’s

potential to improve the polymer-graphene nanocomposites’ electrical conductivity, EMI

shielding performance, and dielectric properties involves highly complex processes. There are

challenges associated with exfoliation, homogeneous dispersion, and the microscopic

arrangement of the graphene platelets within the polymer [19].

Different methods have been used to develop more efficient graphene-based polymer

composites with enhanced electrical and EMI shielding properties. These have included

modifying the graphene platelets’ surfaces [1,20], exploiting the synergistic behavior of the

hybrid nanomaterials [21,22], and in-situ polymerization [1,23]. Zhao et al. [2] fabricated hybrid

poly-(vinylidene fluoride)-5 wt.% carbon nanotube/10 wt.% GnP thin films of 0.1 mm

thickness, using solution casting followed by hot pressing with the EMI SE of 27.58 dB. Wu et

al. [14] developed EMI shielding graphene foam (GF)/poly(3,4-

ethylenedioxythiophene):poly(styrenesulfonate) (PEDOT:PSS) composites by drop coating of

the PEDOT:PSS on the cellular structure of the freestanding GFs. The fabricated composites

exhibited EMI SE of 91.9 dB. Yousefi et al.[20] fabricated self-aligned reduced graphene oxide

(rGO)-polymer nanocomposites by dispersing monolayer graphene in epoxy using an aqueous

97

casting method through the in-situ reduction of graphene oxide (GO) [20]. They achieved a very

low percolation threshold of 0.12 vol% [20]. Kim et al. [22] fabricated a hybrid polymer

nanocomposite through the chemical vapor deposition of carbon nanotubes onto rGO oxide

platelets, followed by solution mixing. They reported a dielectric constant of 32 with a dielectric

loss of 0.051 at 0.062 wt% loading of hybrid fillers and 1×102 Hz [22]. Soliman et al. [24,25]

developed porous-organic polymers (POPs)-GnP with enhanced electrical conductivity. They

utilized the POP-GnP interactions and homogeneous in-situ coating of the POP atop GnP

through a bottom-up assembly on the dispersed GnPs.

Unlike the batch-type synthesis methods [1,14,20–23] injection molding is an economically

viable and continuous method to manufacture polymer composites. When it is combined with

physical foaming, another layer of flexibility is added, which can tailor the polymer composites’

functional properties. In addition to weight reductions, supercritical fluid SCF treatment and

physical foaming can enhance the fillers’ dispersion [26] and exfoliation [27–29], and can re-

arrange their orientation within the polymer matrix [18,29,30]. Foaming can also enhance

various polymer composite functionalities, including their electrical conductivity [7,31], their

dielectric performance [18,32], their thermal conductivity [29], and their electromagnetic

interference shielding effectiveness [33–37]. However, to the best of our knowledge, no

research has been published on the electrical properties of injection-molded graphene-polymer

nanocomposite foams.

In this study, we have presented a facile manufacturing platform to decrease the percolation

threshold and to enhance the electrical properties and the EMI SE of high-density-polyethylene

(HDPE)-graphene nanoplatelet (GnP) composites. Herein, we have demonstrated that the

generation of a microcellular structure can substantially enhance the electrical conductivity and

reduce the percolation threshold of the GnP based polymer composites. The microcellular

HDPE-GnP composite foams were fabricated using melt mixing, SCF-treatment and, finally,

foaming in an injection-molding process. The generated microcellular structure re-orientated

and changed the arrangement of well exfoliated GnPs within the polymer matrix. The HDPE-

GnP nanocomposites foams had a lower percolation threshold, enhanced the electrical

conductivity, the EMI SE, and the dielectric constant, which made them superior to the regular

injection-molded and compression-molded nanocomposites.

98



An HHM 5502BN Marlex® grade HDPE (MFI:0.35 dg min−1 230 °C/2.16 kg) with a density of

0.955 g cm−3 was loaded with GnP powder provided by NanoXplore Inc. (heXo-g-V20 with a

density of 2.2 g.cm-3, a surface area per unit mass of 30 m2/g) to make a HDPE-35 wt.% GnP

masterbatch. The HDPE-35 wt.% GnP masterbatch was produced by melt compounding using a

TDS-20 twin-screw extruder with a 22 mm screw diameter and a 40 L/D ratio. The temperature

profile was set to 180°C - 220°C. A rotational speed of 45 rpm and a throughput of 5 kg.hr-1

were used. HDPE-GnP composites with a different GnP loading content were then obtained by

diluting the HDPE-35 wt.% GnP masterbatch with neat HDPE and mixing them in a twin-screw

extruder (with a diameter of 27 mm and L/D: 40). Nitrogen (N2), supplied by Linde Gas,

Canada, was used as the SCF.

A 50-ton Arburg Allrounder 270/320C injection-molding machine (Lossburg, Germany), with a

30-mm diameter screw equipped with MuCell® technology (Trexel, Inc., Woburn,

Massachusetts) was used to fabricate the HDPE-GnP composites. The following two types of

HDPE-GnP composite samples were fabricated: injection-molded solid (Solid), injection-

molded foam (Foam). The degrees of foaming in the foamed samples were controlled by

partially filling the mold volume. The degree of foaming is indicative of the void fraction in the

injection-molded foam samples.

The processing parameters used to fabricate injection-molded-solid and injection-molded-foam

nanocomposites were obtained based on the nanocomposites’ microstructure integrity. Table 5.1

summarizes the employed processing parameters in this study.

The solid nanocomposites were manufactured using the conventional injection molding process

without the SCF-treatment (N2 was used as the SCF) and physical foaming. For the foamed

samples, the MuCell module was used to inject 0.4 wt.% N2 at its supercritical form into the

barrel. After treating the polymer-GnP with the SCF, the mold cavity was partially filled with

the polymer/GnP/gas mixture. The degrees of foaming in the foamed samples were controlled

by partially filling the mold volume. The degree of foaming is indicative of the void fraction in

the injection-molded foam samples.

99

Table 5.1. Processing parameters used in injection molding of solid and foamed composites

Parameter

Solid Foam

Melt temperature (°C) 210 210

Barrel pressure (MPa) 16 16

Screw speed (rpm) 300 300

Metering time (s) 12 12

Injection flow rate (cm3s-1) 90 90

Mold temperature (°C) 75 75

Pack/hold pressure (MPa) 30 N/Aa

Pack/hold time (s) 15 N/A

Cooling time (s) 60 60

Gas injection pressure (MPa) N/A 24

N2 content (wt.%) N/A 0.4

Degree of foaming (%) N/A 7, 16, 26

The solid and foamed samples were cut from the injection-molded parts at a distance of 100 mm

from the cavity gate. The disk-shape samples with a 20 mm diameter × 3 mm thickness were

used to measure electrical conductivity, dielectric constant, and dielectric loss and samples with

the dimensions of 10.6mm×4.3mm×3.0mm were used to measure the EMI SE values of the

HDPE-GnP composites. The schematic of the injection-molded parts has been presented in

Figure 5.1.

Figure 5.1. The schematic of the injection molded parts and the location of cut samples

100

5.3.2 Characterization

Scanning electron microscopy (SEM) imaging was performed using a Quanta EFG250. The

SEM samples were prepared through cryofracture and subsequently sputter-coating with gold.

Transmission electron microscopy (TEM) imaging was performed using a FEI Tecnai-20 TEM

to investigate the level of GnP’s exfoliation within the polymer matrix. The TEM samples were

prepared by cryo-ultramicrotomy (Leica EM FCS).

The through-plane electrical conductivity, the dielectric constant, and the dielectric loss of the

samples with a 20 mm diameter × 3 mm thickness, were measured using an Alpha-A high

performance dielectric impedance analyzer (Novocontrol Technologies GmbH & Co. KG). The

broadband electrical properties of the HDPE-GnP composites were analyzed at frequencies that

ranged from 1×10-1 Hz to 3×10+5 Hz. The electrical conductivity was measured at a frequency

of 0.1 Hz and was reported as the direct current (DC) conductivity (σDC) [7,31,35]. The

comparative analyses of the dielectric properties were conducted at a frequency of 1×10+3 Hz

[17,38].

The EMI SE values of the HDPE-GnP composites with dimensions of 10.6 mm×4.3 mm×3.0

mm were measured over a frequency range of 18−26.5 GHz (K-band) using the waveguide

method via the Agilent N5234A vector network analyzer. The power coefficient of the

reflection (R), transmission (T), and absorption (A) were calculated from the S-parameters (that

is, the S11 and S21) based on the following Equations [39–42]:

R = |S11|2 (5.1)

T = |S21|2 (5.2)

A = 1 ˗ R ˗ T (5.3)

Thus, the total EMI shielding (SET), including the shielding by absorption (SEA) and the

reflection (SER), can be described by the following Equations [40,41,43,44]:

101

SET = SER + SEA (5.4)

SER = −10log10(1 − 𝑅) (5.5)

SEA = −10log10(𝑇

1−𝑅) (5.6)

5.4 Results and Discussion

5.4.1 Microstructure and morphology of the HDPE-GnP composites

Figure 5.2 shows the microstructure of the core and skin regions of the solid and foamed HDPE-

9.8 vol.% GnP composites. As was expected, the solid samples’ structure was completely solid.

The GnPs were highly oriented in the flow direction in the skin (about 500 μm on each side)

region of the solid samples. This was due to the high shear stresses caused during injection

molding [45]. However, in the core region of the solid samples, the GnPs had a relatively more

random orientation.

The foamed nanocomposites had a microcellular structure with a non-homogeneous cell

morphology. The average cell size of the HDPE-GnP composites foams with a 16% degree of

foaming was 20±11μm. This non-homogeneous microcellular structure was a result of the

structure’s heterogeneities, which were caused by the dispersed GnPs, where the lower

activation energy for cell nucleation is required [46–48]. Moreover, in the foamed samples, the

GnPs’ orientation in both of these regions was random. This was mainly attributed to (i) the

nanocomposites’ lower melt viscosity due the SCF’s dissolution and (ii) the growth of cells

during the physical foaming. The growth of bubbles caused the rotation and displacement of the

GnPs and oriented c. This re-arranged the GnPs’ flow-induced orientation, and thus increased

the opportunities for their interconnectivity [31,51]. In Figure 5.2b, the schematic diagram

shows the GnPs’ arrangement and interconnectivity in the solid and foamed HDPE-GnP

composites.

102

Figure 5.2. (a) SEM micrographs of the skin and core regions of the solid and foamed (16 %

degree of foaming) HDPE-GnP composites at 9.8 vol % GnP content, and (b) Ideal

conceptualization of the GnPs’ arrangement in the solid and foamed samples. The arrow shows

the melt’s flow direction in the injection-molding process.

5.4.2 The effect of physical foaming on the GnP’s exfoliation and

dispersion

Following the SCF-treatment and physical foaming the thick and agglomerated GnPs in the

solid samples were further exfoliated into thinner layers. This process was discussed in detail in

Section 3.5.1. Figure 5.3 shows more analysis of the HDPE-4.5 vol.% GnP composites using

wide-angle X-ray diffraction (WAXD) (Figure 5.3a) and transmission electron microscopy

(TEM) (Figure 5.3b-c). The intensity reduction at the diffraction peak of the (002) plane

indicated the GnPs’ exfoliation [28,29,52,53].

Once the HDPE-GnP melt had received the SCF-treatment, the SCF was dissolved within the

polymer matrix, and then it was diffused between the GnPs’ layers. Due to the rapid

103

depressurization in the mold cavity, the SCF experienced phase transformation. The SCF’s

expansion during its transformation into a gaseous state exfoliated the graphene layers [29,54].

Moreover, the nucleated bubble growth near the GnPs further dispersed the GnPs within the

polymer matrix [27,29].

Figure 5.3. (a) XRD spectra of neat HDPE, GnP powder, solid, foamed samples with 4.5 vol.%

GnP. The inset figure (a) shows an ideal conceptualization of the SCF-assisted exfoliation of the

GnPs in the foamed samples. (b) Representative TEM micrographs of the foamed and (c) solid

samples of the HDPE-4.5vol.% GnP

5.4.3 The electrical conductivity of the polymer-GnP composites

Figure 5.4a shows the broadband conductivity of the nanocomposites across a frequency range

of 1×10-1 Hz to 1×10+5 Hz. The solid samples had a 7 to 12.6 vol.% GnP content. The foamed

samples were fabricated using the corresponding solid precursor, which contained 7, 9.8 and

12.6 vol.% of the GnP.

The broadband electrical conductivity of all the solid samples (containing 7, 9.8 and 12.6 vol.%

GnP) followed a frequency-dependent behavior across the whole frequency range. The

frequency-dependency of the electrical conductivity is one of the typical characteristics of

104

insulating polymer composites [17,55]. This indicates that the GnPs were distributed within the

polymer matrix without forming conductive channels. And this behavior is defined by σ = σDC +

σAC, where the σDC is the frequency-independent part and the σAC (alternative current (AC)

conductivity) is the frequency-dependent part of the total electrical conductivity. The frequency

below which the electrical conductivity shows a frequency-independent behavior is known as

the critical frequency [17,55]. The frequency-dependent conductivity of the solid samples

containing 9.8 vol.% GnP was decreased from 1.1×10-8 S.cm-1 to 2.0×10-14 when the frequency

was decreased from 1×10+5 to 1×10-1 Hz.

Figure 5.4. (a) The AC conductivity of the solid, and foamed HDPE-GnP composite; and (b)

The DC conductivity of the solid, and foamed HDPE-GnP composite measured at 0.1 Hz

(degree of foaming of foamed samples is 16%)

However, the physical foaming transformed the frequency-dependent behavior of the solid

samples (containing 9.8 vol.% GnP) into the frequency-independent behavior at frequency

ranges of below 2×10+3. By increasing the GnP content to 12.6 vol.%, the foamed samples

exhibited a frequency-independent behavior across the entire frequency range from 1×10-1 Hz to

1×10+5 Hz. Moreover, foaming enhanced the electrical conductivity of the solid HDPE-12.6

vol.% GnP composites by eight orders of magnitude at frequency ranges of below 1×100.

Figure 5.4b shows the variation of the DC conductivity of the solid and foamed HDPE-GnP

composite as a function of the GnP loading. The HDPE-GnP composite’s electrical conductivity

was significantly affected by the physical foaming. This occurred through two different

12 vol.% Solid Foam

9 vol.% Solid Foam

7 vol.% Solid Foam

10-1

100

101

102

103

104

105

10-14

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

Co

ndu

ctivity,

AC (S

.cm

-1)

Solid 7.0vol.% G

nP

Foam 12.6vol.% GnP

Frequency (Hz)

Foam 9.8vol.% GnP

(a)

0 2 4 6 8 10 12 14 16 18

10-15

10-14

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

Foam

Foam

Foam (GnP content

with respect to total volume)

Foam (GnP content

with respect to polymer volume)

Solid

Co

ndu

ctivity,

DC (S

.cm

-1)

GnP content (vol.%)

Solid

(b)

105

mechanisms, which included the following: (i) The foaming actions, such as bubble growth

which affected the GnPs’ arrangement and interconnectivity [29]; and (ii) The volume exclusion

effect of foaming which resulted in the GnPs’ localization within the struts and cell walls

[7,29,31]. To focus solely on how the foaming actions affected the electrical conductivity, the

GnP content was considered in relation to the polymer volume. In other words, the GnP content

in the foamed samples was reported the same as their solid precursors. To include how density

reduction in the foaming affected the electrical conductivity, the GnP content was calculated in

relation to the total volume of the nanocomposite foams.

The conductivity of all the HDPE-GnP composites showed a clear insulation-conduction

transition behavior. The abrupt insulation–conduction transition of the foamed samples began at

a much lower GnP content than that of their solid counterparts. Thus, the percolation threshold

of the foamed samples was found to be around 9.8 vol.% GnP (that is, in relation to the polymer

volume). This outcome was far superior to the 19 vol.% GnP that was found in the solid

samples. Moreover, by taking a 16 vol.% degree of foaming into account, the percolation

threshold of the foamed samples was further decreased from 9.8 vol.% to 8.2 vol.% GnP. In

other words, the generation of a microcellular structure within the injection-molded samples

decreased the percolation threshold for the nanocomposites by more than 2.3-fold. Meanwhile,

to achieve the same level of electrical conductivity in the given volume of the samples, the

required GnP content (in relation to the total volume) for the foamed nanocomposites was much

lower than it had been for the solid ones. For example, the foamed samples with a GnP content

of 8.2 vol.% had the same electrical conductivity, which had been achieved with 19 vol.% GnP,

in the solid nanocomposites.

The GnPs’ flow-induced orientation in the solid nanocomposites (discussed in Section 5.4.1)

significantly deteriorated their interconnectivity and the formation of a conductive network.

And, consequently, the through-plane electrical conductivity was inferior. This resulted in a

high percolation threshold and in the very slow increase of the electrical conductivity in the

solid samples with an increased GnP content.

The higher through-plane electrical conductivity and the lower percolation threshold of the

foamed samples, as compared to the solid counterparts, were mainly attributed to the changes in

the microstructures. This had been induced by the introduction of foaming, which operated in

several ways and included the following actions: (a) a higher level of GnPs’ exfoliation and

106

dispersion in the polymer; (b) a decreased flow-induced orientation of GnPs due to the foaming

actions and reduced viscosity; (c) enhanced local interconnectivity of GnPs due to the cell

growth during foaming; and (d) reduction in the skin layer’s thickness. It is also believed that

the GnPs’ aspect ratio is higher with foaming, due to the lower melt viscosity and the lower

fillers’ mechanical breakdown [31,35].

5.4.3.1 The effect of the foaming degree on the electrical conductivity

Figure 5.5a shows the variations of the σDC with the foaming degree in the foamed

nanocomposites with various GnP contents (in relation to the polymer volume). Below the

percolation threshold of the solid nanocomposites (that is, 9.8, 12.6 and 15.6 vol.% GnP in

Figure 5.4b), the generation of a 7% foaming degree caused the formation of conductive

percolative networks and resulted in a sharp increase in the σDC from 6 to 9 orders of magnitude.

Around the solid samples’ (19 vol.% GnP) percolation threshold, the conductivity enhancement

due to the foaming was less pronounced and increased only by 3 orders of magnitude. This was

attributed to the percolative networks that had already formed within the solid nanocomposites

at 19 vol.% GnP.

To further investigate how the foaming degree affected the electrical conductivity, the σDC of the

solid and foamed nanocomposites were plotted as a function of the GnP content in Figure 5.5b.

Notably, the percolation threshold was decreased by the increased foaming degree. The

percolation threshold sharply dropped from 19 to 9.1 vol.% GnP when a 7% degree of foaming

was generated. The percolation threshold was further decreased from 9.1 to 7.2 vol.%, when the

degree of foaming was increased to 26%. Therefore, the generation of the microcellular

structure decreased the percolation threshold by up to 62%. The decrease in the percolation

threshold that was obtained by the increase in the foaming degree from 7% to 26% was mainly

attributed to the volume exclusion effect induced in the gaseous phase.

107

Figure 5.5. (a) Variations of the foaming degree on the electrical conductivity of the HDPE-GnP

composites; (b) The evolution of the percolation threshold with the foaming degree

5.4.4 The dielectric properties of polymer-GnP composites

The dielectric permittivity presents in a complex function, which is composed of a real part ε'

and an imaginary part ε''. The real part is related to the charge displacement, which is governed

by the polarization within the material. Interfacial polarization is the most common type of

polarization that occurs across frequency ranges of less than 1 MHz [15]. Based on the

Maxwell–Wagner–Sillars (MWS) effect [56], charges are accumulated at the interface of the

polymer and filler. The imaginary part of the dielectric permittivity (ε'') is used to quantify the

dielectric loss (tan ), which is defined as the ratio of the imaginary part to the real part of the

dielectric permittivity.

Figure 5.6a-b exhibits the dialectic constant and loss of the solid and foamed (16% degree of

foaming) nanocomposites as a function of the GnP content. The dielectric constant (ε') in all of

the samples was enhanced by increasing the GnP content. The higher GnP content increased the

polymer-GnP interface area, which resulted in a higher interfacial polarization. Moreover, the

polymer-GnP nanocomposites can be considered as nanoscale parallel-plate capacitors, where

the GnPs act like electrodes, and the polymer matrix is considered to be dielectric [17,18].

Therefore, increasing the GnP content increased the number of nanocapacitors and decreased

the interspatial distances between the adjacent GnPs, thus leading to a higher real permittivity.

0 5 10 15 20 25

10-14

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

Co

ndu

ctivity,

DC (S

.cm

-1)

Degree of foaming (%)

19 vol.%

GnP

15.6 vol.% GnP

12.6 vol.% GnP

9.8 vol.% GnP

7.0 vol.% GnP

4.5 vol.% GnP

(a)

0 2 4 6 8 10 12 14 16 18

10-15

10-14

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

Foam (7% DF)

Foam (16% DF)

Foam (26% DF)

Solid

Conductivity,

DC (S

.cm

-1)

GnP content (vol.%)

Solid

(b)

108

Figure 5.6. (a) Real dielectric permittivity (ε'); and (b) The dielectric loss (tan δ) of the solid and

foamed (16% degree of foaming) nanocomposites as a function of the GnP content measured at

1×10+3 Hz. (GnP vol.% is reported in relation to the polymer volume)

However, with the same GnP content, the dielectric constant of the foamed samples was

considerably higher than that of their solid counterparts. For instance, at 9.8 vol.% GnP, the real

permittivity of the solid nanocomposites was 6.2. However, the introduction of the microcellular

structure substantially increased the real permittivity of the foamed nanocomposites to 106.4

with a 9.8 vol.% GnP content (Figure 5.6a). In other words, the real permittivity of the foamed

samples with a 9.8 vol.% GnP content was more than one order of magnitude higher than that of

their solid counterparts.

Figure 5.6b shows that the dielectric loss was increased by the increased GnP content in both

solid and foamed samples. The increased GnP content enlarged the number of the charge

carriers and the nanocapacitors which, respectively, resulted in a higher Ohmic and polarization

loss [18,32]. The foamed nanocomposites had a higher dielectric loss than the solid samples,

mainly due to the more random distribution of the fillers in the polymer matrix [18,56]. And this

led to the formation of GnP conductive networks and, thereby, a higher Ohmic loss [18,32]. On

the other hand, the introduction of foaming increased both the dielectric permittivity and the

dielectric loss of the nanocomposites[18]. However, it is interesting to note that the dielectric

loss of the foamed samples, around the percolation threshold, was still relatively low. For

instance, the real permittivity and the dielectric loss of the foamed samples with a 9.8 vol.%

GnP was 106.4 and 0.4, respectively.

0 2 4 6 8 10 12 14 16 18

0

50

100

150

200

250

300

0 2 4 6 8 10 12 14 16 18

0

50

100

150

200

250

300

Foam

Real perm

ittivity, (

' )

GnP content (vol.%)

Foam (GnP content


Foam (GnP content


Solid

Foam

Sol

id

(a)

0 2 4 6 8 10 12 14 16 18

10-3

10-2

10-1

100

101

102

103

Foam(GnP content


Foam (GnP content


Solid

Solid

Die

lectr

ic lo

ss, (t

an

)

GnP content (vol.%)

(b)

109

The increased real permittivity of the foamed samples, when compared with that of the solid

nanocomposites, was mainly attributed to the unique GnP parallel-plates arrangement in the cell

walls due to the cellular growth that occurred [18,32]. This led to a highly effective interface

area between the adjacent GnPs [32]. Moreover, a higher level of GnP exfoliation: (a) increased

the number of nanoscale capacitors; (b) raised the polymer-GnP interfaces; and (c) decreased

the interspatial distances between the adjacent GnPs, which enhanced the real permittivity [32].

Figure 5.7 shows the broadband real permittivity (ε') and the dielectric loss (tan ) of the solid

and foamed (with 16 % degree of foaming) nanocomposites with different GnP contents. The

broadband real permittivity of all the solid samples followed a relatively frequency-independent

behavior across the whole frequency range (Figure 5.7a). On the other hand, the generation of

the microcellular structure not only substantially increased the real permittivity, but it also

changed the frequency-independent behavior of the real permittivity in the solid

nanocomposites, which contained 9.8 and 12.6 vol.% GnP, into the frequency-dependent

behavior found in their foamed counterparts (Figure 5.7b). This frequency-dependent behavior

of the dielectric constant is a characteristic of the conductive composites [7,56]. It indicated that

conductive paths had formed within the foamed samples [31].

Figure 5.7c-d shows that, beyond the percolation threshold, the broadband dielectric loss of the

foamed nanocomposites was higher than in the solid counterparts. The higher dielectric loss of

the foamed samples was attributed mainly to their higher Ohmic loss, which was related to the

σDC. The total dielectric loss consisted of the Ohmic loss and polarization loss of the space

charges [57,58]. However, in the current polymer-GnP system, the Ohmic loss was the major

contributor to the total dielectric loss [57,58]. The higher σDC and, consequently, the higher

Ohmic loss, caused the frequency-dependency of the dielectric loss in the foamed

nanocomposites with a 9.8 and 12.6 vol.% GnP content.

110

Figure 5.7. Broadband dielectric permittivity of (a) The solid samples, and (b) The foamed 9.8

vol.% HDPE-GnP composites. Broadband dielectric loss of (c) The solid samples, and (d) The

foamed 9.8 vol.% HDPE-GnP composites

5.4.5 The EMI shielding effectiveness (SE) of the polymer-GnP

composites

The EMI’s shielding effectiveness represented the material’s ability to reduce the

electromagnetic waves’ intensity. The shielding performance for a given electromagnetic

radiation is defined as SE=10log (Pi/Pt), where Pi is the incident power and Pt is the transmitted

power in decibels (dB) [35,39]. For instance, a material with a SE of 40 dB can block 99.99% of

the incident wave. Figure 5.8 shows the EMI SE of the solid and foamed HDPE-GnP

100

101

102

103

104

105

101

102

103

12.6 vol% Gr 7.0 vol% Gr

9.8 vol% Gr 4.5 vol% Gr

Solid

Bro

ad

ba

nd

re

al p

erm

ittivity (

' )

Frequency (Hz)

(a)

100

101

102

103

104

105

101

102

103

Foam

12.6 vol.% GnP9.8 vol.% GnP

4.5 vol.% GnPBro

adband r

eal perm

ittivity (

' )

Frequency (Hz)

7.0 vol.% GnP

(b)

100

101

102

103

104

105

10-3

10-2

10-1

100

101

102

103

12 vol.% GnP

4.5 vol.% GnP

12.6 vol% GnP 7.0 vol% GnP

9.8 vol% GnP 4.5 vol% GnP

Solid

Bro

adband d

iele

ctr

ic loss (

tan

)

Frequency (Hz)

(c)

100

101

102

103

104

105

10-3

10-2

10-1

100

101

102

103

Bro

adband d

iele

ctr

ic loss (

tan

)

Frequency (Hz)

12 vol.% G

nP9 vol.% GnP

7 vol.% GnP

4.5 vol.% GnP

Foam(d)

111

composites over the K-band frequency range (between 18 GHz and 26.5 GHz). The EMI SE

values were greater at a higher GnP content in both the foamed and solid samples.

Figure 5.8. K-band EMI SE of (a) the solid; and (b) the foamed HDPE-GnP composites with

various GnP content.

As shown in Figure 5.9a, at a given GnP content, the foamed samples had higher SE values than

their solid counterparts. The grand average of the three sample replications’ measured values

over the K-band frequency range were plotted as the EMI SE shown in Figure 5.9. At a 19 vol%

GnP, the EMI SE of the foamed samples reached 31.6 dB, which corresponded to a 99.93%

blockage of the incident EMI wave. With the same GnP content, the solid samples had an EMI

SE of 21.8 dB.

Figure 5.9a also presents the foamed samples’ EMI SE as a function of the GnP content, which

was calculated in relation to the nanocomposite foams’ total volume. It is notable that to attain a

certain EMI SE value in a given nanocomposite volume, the GnP content required for the

foamed nanocomposites was considerably lower than it was for their solid counterparts. For

instance, to reach an EMI SE of about 21 dB, the final GnP vol.% was, respectively, 19 and 14

for the solid and foamed nanocomposites. This corresponded to a 26% reduction in the GnP

usage when foaming was done.

18 20 22 24 260

5

10

15

20

25

30

35

19vol.% GnP

15.6vol.% GnP

4.5vol.% GnP

7.0vol.% GnP

12.6vol.% GnP

4.5 vol.% GnP 15.6 vol.% GnP

7.0 vol.% GnP 19.0 vol.% GnP

12.6 vol.% GnP

EM

I S

E (

dB

)

Frequency (GHz)

(a)

Solid

18 20 22 24 260

5

10

15

20

25

30

35

EM

I S

E (

dB

)

Frequency (GHz)

4.5 vol.%GnP 15.6 vol.%GnP

7.0 vol.%GnP 19.0 vol.% GnP

12.6 vol.%GnP

(b)

4.5vol.% GnP

7.0vol.% GnP

15.6vol.% GnP12.6vol.% GnP

19vol.% GnP

Foam

112

Figure 5.9. (a) The K-band EMI SE of the solid and foamed HDPE-GnP composites as a

function of their GnP content; (b) The contributions of the reflection and absorption

mechanisms to the total K-band EMI SE of the solid and foamed HDPE-GnP composites as a

function of their GnP content; (c) schematic diagrams of the scattering and multiple reflections

of the electromagnetic waves

The wave reflection (SER) and the absorption (SEA) are the main electromagnetic attenuation

mechanisms [39–42]. To further demonstrate the shielding mechanisms in both the solid and

foamed nanocomposites, Figure 5.9b shows the contributions of the wave reflection and the

absorption to the total EMI SE (SET).

0 2 4 6 8 10 12 14 16 18

0

5

10

15

20

25

30

35(a)

Foam

Foam

Sol

id

EM

I S

E (

dB

)

GnP content (vol.%)

Foam(GnP content with respect to total volume)

Foam(GnP content with respect to polymer volume)

Solid

0 2 4 6 8 10 12 14 16 18

0

5

10

15

20

25

30

Absorption

EM

I S

E (

dB

)

GnP content (vol.%)

Foam-Absorption

Solid-Absorption

Foam-Reflection

Solid-Reflection Foam

Sol

id

Reflection

(b)

113

The contribution of the reflection to the total shielding in both the solid and foamed

nanocomposites was similar, and it reached 3.5 dB around the percolation threshold region.

However, the absorption mechanism clearly dominated the shielding mechanism, and it was

continuously increased by the addition of GnP in both the foamed and solid nanocomposites.

For example, the absorption mechanism contributed, respectively, 84% and 88% of the total

shielding in the solid and foamed HDPE-19 vol.% GnP composites. It was also notable that the

foamed samples’ SEA was higher than the solid counterparts’ with the same GnP content. This

gave the foamed nanocomposites a higher SET.

The reflection mechanism is related to the impedance mismatch between the shielding

composite and the air. The presence of the charge carriers (that is, the electrons and holes)

and/or the surface charge are mainly assumed to govern the reflection mechanism [2,59,60].

However, the absorption mechanism originates from the Ohmic and polarization losses [61].

The Ohmic loss results in energy attenuation via the current flow through the conduction and

tunneling mechanisms. The polarization loss is correlated to the interfacial polarization’s density

and is thereby transferred to the absorber’s real permittivity [2,4].

The foamed samples’ enhanced SE was mainly attributed to three factors. The first of these is

the electromagnetic wave’s multiple reflections on various surfaces (that is, of the cell-

composite matrix surface area), which created another shielding mechanism [13,31,35,36]. The

electromagnetic waves entering the nanocomposites foams were reflected and scattered in the

microcellular structure numerous times. Therefore, the adequate wave absorption capability of

the composite matrix combined with the multiple reflections inside the cells to further enhance

the shielding properties of the electromagnetic waves. Thus, the foamed nanocomposites’ SET

was improved. Figure 5.9c shows schematic diagrams of the scattering and multiple reflections

of the electromagnetic waves in both the solid and foamed nanocomposites. The second factor

was the GnPs’ increased interconnectivity and, hence, the samples’ resultant higher conductivity

and permittivity. It has been reported that higher conductivity and permittivity (ε') result in a

higher SE [2,31,62]. The third factor was a higher level of GnP exfoliation caused by the SCF

treatment and foaming processes. The higher level of GnP exfoliation would contribute to the

enhancement of the electrical conductivity and the dielectric permittivity of the foamed samples

(as discussed in Sections 5.4.4 and 5.4.5) and, thereby, would result in a higher EMI SE in the

foamed nanocomposites [2,31,62].

114

5.5 Summary & Conclusions

Herein, we have demonstrated that SCF-treatment and physical foaming can substantially

increase the electrical conductivity and reduce the percolation threshold of the polymer-GnP

composites. This facile technique at once enhanced the electrical conductivity, the dielectric

constant and the EMI shielding performance of the HDPE-GnP composites and decreased their

percolation thresholds. The lightweight HDPE-GnP composite foams were prepared by melt

compounding followed by foaming in an injection molding process. The SCF-treatment and

physical foaming were found to exfoliate the GnPs and change their flow-induced orientation by

reducing the viscosity and bubble growth. The generation of a microcellular structure re-

arranged the GnPs so that they were mainly perpendicular to the radial direction of the cellular

growth within the cell walls. This enhanced the GnPs’ interconnectivity which resulted in a

significantly higher conductivity and a lower percolation threshold. For example, in addition to

26% density reduction, the percolation threshold of 19 vol.% GnP in the solid samples was

sharply decreased to 7.2 vol.% GnP with the introduction of a 26% degree of foaming. Foaming

substantially enhanced the real permittivity of the foamed samples. The real permittivity of the

foamed samples with a 9.8 vol.% GnP was 106.4 while that of their solid counterparts was 6.2.

Moreover, the introduction of a microcellular structure enhanced the EMI shielding performance

of the HDPE-GnP composites. A maximum EMI SE of 31.6 dB was achieved in HDPE−19 vol.

% GnP composite foams, which was superior to 21.8 dB of the solid counterparts.

These research results show that SCF-treatment and physical foaming in an injection-molding

process offers a facile, cost-effective, and industrially viable method by which to develop

lightweight conductive polymer-GnP nanocomposites.

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122

CHAPTER 6

6 Enhancement of the dielectric performance of polymer-GnP composites using SCF-treatment and physical foaming-Part I


Hamidinejad, S.M., Zhao, B., Chu, R.k.M., Moghimian, N., Naguib, H., Filleter, T., and Park, C.B.,

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6.1 Summary

Dielectric polymer

nanocomposites with high

dielectric constant (ε'), and

low dielectric loss (tan δ) are

extremely desirable in the

electronics industry.

Percolative polymer-graphene

nanoplatelet (GnP) composites have shown great promise as dielectric materials for high-

performance capacitors. Herein an industrially-viable technique for manufacturing a new class

of ultralight polymer composite foams using commercial GnPs with excellent dielectric

performance is presented. Using this method, the high-density polyethylene (HDPE)-GnPs

123

composites with a microcellular structure were fabricated by melt mixing. This was followed by

supercritical fluid (SCF) treatment and physical foaming in an extrusion process, which added

an extra layer of design flexibility. The SCF treatment effectively in-situ exfoliated the GnPs in

the polymer matrix. Moreover, the generation of a microcellular structure produced numerous

parallel-plate nanocapacitors consisting of GnP pairs as electrodes with insulating polymer as

nanodielectrics. This significantly increased the real permittivity and decreased the dielectric

loss. The ultralight extruded HDPE-1.08 vol.% GnP composite foams, with a 0.15 g.cm-3

density, had an excellent combination of dielectric properties (ε'=77.5, tan δ=0.003 at 1×105

Hz) which were superior to their compression-molded counterparts (ε'=19.9, tan δ =0.15 and

density of =1.2 g.cm-3) and to those reported in the literature. This dramatic improvement

resulted from in-situ GnP’s exfoliation and dispersion, as well as a unique GnP parallel-plates

arrangement around the cells. Thus, this facile method provides a scalable method to produce

ultralight dielectric polymer nanocomposites, with a microscopically-tailored microstructure for

use in electronic devices.

6.2 Introduction

High performance dielectric materials are vital to the development of next-generation

miniaturized electronic devices. Dielectric materials with high dielectric permittivity (ε') and

low dielectric loss (tan ) have been receiving increasing interest in modern electronics as the

capacitors and integrated capacitors [1–4]. Multifunctional, lightweight, and low-cost polymer

nanocomposites show much promise for use as dielectric materials. Their dielectric permittivity

and dielectric loss tunability is large; their resistance to chemicals is outstanding; they are easily

processed, and they have tailorable thermal and mechanical properties [5,6].

Polymers have an extremely low dielectric loss and a high dielectric breakdown strength;

however, they suffer from a low dielectric constant (ε'<10). Incorporating different types and

shapes of fillers have been demonstrated to improve the dielectric properties of polymer

nanocomposites [5]. In general, there are two different types of fillers: nonconductive and

conductive that have been used to enhance the polymer nanocomposites’ dielectric properties

[5]. In order to exhibit an outstanding dielectric permittivity of the polymer nanocomposites, the

dielectric constant of the fillers should be significantly higher than that of the polymer matrix if

124

the fillers are nonconductive. Alternatively, if the fillers are conductive, the conductivity of the

fillers should be significantly higher than that of the polymer matrix. This difference, either the

dielectric constant or the conductivity, is the main cause of the interfacial polarization and,

consequently, the improved dielectric permittivity [6,7]. Nonconductive fillers can include

ceramics, with a high dielectric constant such as barium titanate (BaTiO3) [8], strontium titanate

(SrTiO3) [9], calcium titanate (CaTiO3) [10]. Such fillers can enhance the dielectric permittivity,

however, ceramic-based dielectric polymer composites are usually filled with a high ceramic

loading (<50 vol.%) to achieve a dielectric permittivity of approximately 50 [11]. With such a

high filler loading level, the amount of the polymer matrix left to support the fillers and the

composite’s structural integrity is insufficient. This leads to expensive and heavyweight

composites, which are difficult to process.

Percolative polymer nanocomposites, which contain highly electrically conductive additives

such as carbon black [12], carbon nanotubes [13,14], graphite nanoplates (GNP) [15], and

graphene [16,17], have the potential to overcome the limitations of ceramic-based polymer

nanocomposites. By using conductive nanofillers, higher dielectric permittivity values can be

achieved at much lower filler concentrations [5,11,13]. However, the dielectric properties

change substantially near the percolation threshold. The electrical conductivity of the polymer

nanocomposites abruptly increases when the filler concentration exceeds the percolation

threshold value. In addition, the dielectric loss is sharply increased near the percolation

threshold due to the high leakage current from the conductive channels that form across the

entire system [18]. Thus, the dielectric properties need to be optimized within the so-called

“adjustable window” near the percolation threshold. This is where the dielectric permittivity can

be maximized while the dielectric loss can be restrained [11,19]. However, the adjustable

window in percolative polymer nanocomposites is very narrow, and it is extremely challenging

to control the polymer nanocomposites’ dielectric performance within it. To address this

dilemma, a uniform and homogeneous dispersion of the conducting fillers in the polymer matrix

is critical.

The types and shapes of conductive fillers are also key factors that can contribute to improving

the dielectric performance of polymer nanocomposites [20]. It has been suggested [5] that

graphene nanoplatelets with an exceptional electrical conductivity (∼6000 S/cm [21]) can be a

more effective nanomaterial for dielectric polymer nanocomposites as compared to carbon black

125

and carbon nanotubes. This can be attributed to the higher interfacial polarization and, hence,

the higher dielectric permittivity that are related to the GnP’s higher specific surface area in

contrast to carbon black and carbon nanotubes [5]. However, achieving the GnP’s full potential

in dielectric polymer nanocomposites is extremely challenging. This is because of the

complexities associated with homogeneous dispersion, efficient exfoliation, and the GnPs’

microscopic arrangement within the polymer matrix [22]. It is notable that the GnPs are the 2-

dimensional carbon materials of up to 100 nm-thick consisting of hundreds of stacking graphene

layers [23]. The GnPs are resulted from exfoliating the graphite stacks into the thinner and

nanoscale graphite platelets [15]. Further exfoliation of the GnPs results in larger specific

surface area and higher aspect ratio of the filler which can benefit dielectric properties of the

polymer nanocomposites.

Different approaches have been undertaken to synthesize more efficient dielectric polymer-GnP

composites. These have included the surface modification [2,16] and coating [2,4] of the GnPs,

the use of hybrid fillers [2,4,24], and in-situ polymerization [16,25]. However, the previous lab-

scale techniques have not taken the complexities of the synthesis procedures into account. These

have a high material cost, and they are not easily scalable for the fabrication of dielectric

materials. To address these limitations of the existing technologies, herein a new method to

fabricate dielectric materials using an inexpensive, scalable, and high throughput process is

presented.

Unlike the lab-scale synthesis techniques [3,4,11,16,17,24–29], extrusion combined with

foaming is a continuous and industrially-viable technique for manufacturing polymer

composites. Physical foaming has shown promise in improving the polymer composites’

performance in different applications [13,30–36]. Physical foaming can add another degree of

design flexibility, which makes it possible to tune their functional properties. It also

significantly reduces the product’s weight. Moreover, physical foaming of polymer composites

can enhance the fillers’ exfoliation [37], dispersion [38,39] and distribution [34,35], and can

also change their orientation [34,39,40] in the polymer matrix. Optimized microcellular foaming

can create tailored microstructures that enhance the electrical conductivity [13,30,35], the

thermal conductivity [37] and the electromagnetic interference shielding effectiveness

[33,36,41] of the conductive polymer composites. However, to the best of our knowledge, there

has been no report on how physical foaming promotes the dielectric properties of the polymer-

126

GnP composites. Furthermore, physical foaming, combined with the extrusion process, can

easily produce tailored microstructures that improve the dielectric properties of GnP-based

polymer composites.

Herein a facile method of manufacturing ultralight, high-density-polyethylene (HDPE)-GnP

composite foams with high dielectric constant and low dielectric loss is reported. The dielectric

microcellular nanocomposites, which contained highly exfoliated GnPs, were developed by melt

mixing. This was followed by supercritical fluid (SCF)-treatment and physical foaming in an

extrusion process. This SCF-assisted technique effectively exfoliated the GnPs in the polymer

matrix. Moreover, a tailored parallel-plates arrangement of the exfoliated GnPs within the cell

walls was created by the microcellular structure. This unique tailored microstructure provided

an excellent combination of high dielectric permittivity and extremely low dielectric loss. This

made it superior to solid compression-molded (SCM) HDPE-GnP composites as well as to the

dielectric performances of batch-type methods [3,4,11,16,17,24,26–30,42,43].



An extrusion grade, commercially available HDPE, HHM 5502BN Marlex®, (MFI:0.35

dg/min.-1 at 230ºC/2.16 kg, with a specific gravity of 0.955 g.cm-3, Chevron Phillips Chemical)

was used as the polymer matrix. Commercial carbon dioxide (CO2), supplied by Linde Gas,

Canada, was used as the physical blowing agent and supercritical fluid.

HDPE-GnP composites with different GnP loading content were then obtained by diluting the

HDPE-35 wt.% GnP masterbatch (see the details in Section 3.3.1) with neat HDPE and mixing

them in a twin-screw extruder (with diameter of 27 mm and L/D: 40).

Two different types of HDPE-GnP nanocomposites, namely solid compression molded (SCM),

and foam-extrude (foam) were prepared. In the solid compression-molded (SCM)

nanocomposites, the HDPE-GnP with different GnP concentrations were hot-pressed into the

disk-shape samples (1 mm thickness and 10 mm diameter) at a temperature of 215°C for 7

127

minutes under a 6 kN pressing force. Next, the samples with a mold assembly were quenched by

using compressed air.

A tandem foam extrusion system was used to fabricate the foam samples with different

densities. Details of preparing samples were discussed in Section 3.3.1.2. The schematic of the

extrusion process and processing parameters have been presented in Figure 3.2.


The microstructures and morphologies of the SCM samples and the extruded foam samples


were frozen in liquid nitrogen, cryofractured, and sputter-coated prior to the SEM. A

transmission electron microscope (TEM; FEI Tecnai 20) were used to investigate the GnP’s

dispersion and exfoliation within the polymer matrix. The TEM samples were prepared by cryo-

ultramicrotomy (Leica EM FCS). The TEM samples were cut from the core region,

perpendicular to the flow direction of foam-extruded filaments, and from the core region in the

thickness direction of the SCM nanocomposites. Wide angle x-ray diffraction (WAXD) analyses

were conducted on the SCM and extruded foam samples using a Rigaku MiniFlex 600 x-ray

diffractometer (Cu Kα radiation, λ = 1.5405 Å) to further examine the GnP’s exfoliation in the

polymer matrix. In this study, WAXD analysis was conducted on the surfaces of both SCM and

foam-extruded nanocomposites. The foam-extruded filaments were degassed by compression

(perpendicular to the flow direction) at room temperature prior to the WAXD.

The through-plane electrical conductivity, the dielectric permittivity (both real and imaginary),

and the dielectric loss of the samples were measured using an Alpha-A high performance

dielectric impedance analyzer (Novocontrol Technologies GmbH & Co. KG) at a voltage of 1.0

V. The HDPE-GnP’s broadband electrical properties were analyzed across frequencies ranging

from 1×10-1 Hz to 3×10+5 Hz. The electrical conductivity at a frequency of 0.1 Hz was reported

as the direct current (DC) conductivity (σDC) [30,35,36]. The comparative analyses of the

dielectric properties were conducted at a frequency of 1×10+3 Hz [11,28].

128


6.4.1 Microstructure and morphology of the polymer-GnP composites

Figure 6.1 shows the microstructure of the as-received GnP powder, the SCM HDPE-9.8 vol.%

GnP composite and that of its extruded-foam counterpart. As expected, the SCM samples had a

completely solid structure (Figure 6.1b). The extruded foam samples, however, had a

microcellular structure with a random cell morphology. The average cell size of the

microcellular HDPE-GnP with a density of 0.15 g.cm-3 was 57±7μm. This non-homogeneous

cellular structure stemmed from the structure’s heterogeneities with the dispersed and

distributed GnP particles. This type of heterogeneity structure can lead to very non-

homogeneous cell nucleation. It has been shown that non-homogenous cell nucleation is

preferable at the interface, due to the lower activation energy for cell nucleation that is required

at that location [44]. The heterogeneous-melt structure (that is, the dispersed GnPs and the

polymer crystals) [44,45] and the low stiffness of the polymer melt [46,47] were the two main

reasons for the observed pinhole formations and the cell wall ruptures shown in the

microstructures in Figure 6.1c.

During foaming, cellular growth displaced the GnPs and changed their orientation [48].

Specifically, they were primarily oriented perpendicular to the radial direction of the cellular

growth. This generated a unique GnP parallel-plates arrangement within the cell walls (see

Figure 6.1d-e).

129

Figure 6.1. The SEM micrographs of the (a) as-received GnP powder; (b) SCM HDPE-9.8

vol.% GnP composites; (c)-(d) Foam-extruded nanocomposites counterparts; and (e) TEM of

extruded foam samples showing parallel-plates arrangements of GnPs within the cell walls

The foamed sample densities were a function of the gas they retained. A low-density foam was

indicative of a significant gas escape inhibition in the sample. To study the effect of the density

on the final properties of the extruded HDPE-GnP composite foams, the HDPE containing 4.5

vol.% of GnP was foamed in extrusion with different densities by changing the die temperature.

However, fabricating different densities of HDPE-GnP composites with a higher GnP content,

such as 7 and 9.8 vol.%, was technically complex due to the excessive viscosity. Figure 6.2

shows the microstructure of HDPE-GnP composite foams with different densities that were

fabricated at various foaming temperatures.

The densities of the HDPE-GnP composite foams at different die temperatures ranged from 0.05

to 0.95 g.cm-3. The minimum density for these foams was obtained at 131°C. This optimum

behavior, which has previously been reported in the literature for various polymers [47,49,50],

was mainly attributed to the competition between the gas escape at high temperatures and the

130

excessively stiff polymer melt at low temperatures with a reduced foam expansion ratio

[47,49,50].

Figure 6.2. Representation of the density of HDPE-GnP composite foams vs the foaming

temperature, together with the related SEM micrograph. The scale bar is 300 m. (GnP vol.%

was reported with respect to the polymer volume)

At higher foaming temperatures, when the melt strength is weak, the cell walls may not have

enough strength to withstand the biaxial stretching that occurs during cellular growth, and they

may easily rupture. This causes both cell-opening and coalescence. The cell morphology of the

extruded foam samples at 139°C and 145°C showed this. And this led to gas loss, lower foam

expansion, and thicker cell walls. On the other hand, at lower foaming temperatures, such as

131°C and 127°C, the polymer melt’s stiffness was increased due to the crystallization of the

HDPE [51]. Therefore, the cell walls were able to bear the biaxial stretching during foaming.

This minimized the cell wall rupture and gas escape. Thus, a more uniform cellular morphology

with thinner cell walls was generated.

131

6.4.2 Electrical conductivity of the polymer-GnP composites

Figure 6.3a shows the broadband electrical conductivity of the HDPE-GnP composites across a

range of frequencies from 1×10-1 Hz to 1×10+5 Hz. The SCM samples had 4.5 to 12 vol.% of the

GnP content. Their extruded-foam counterparts (with a density of 0.14±0.01 g.cm-3 or ~8 times

the foam expansion ratio) were fabricated using the corresponding solid precursor with 4.5, 7

and 9.8 vol.% of the GnP content.

For all the extruded foam samples and the SCM samples that contained 4.5 and 7 vol.% GnP,

the electrical conductivity followed a frequency-dependent behavior across the entire frequency

range which is a main characteristic of insulating polymer composites. It indicates that the

electrical conductivity in this case is mainly governed by the polymer matrix, and that the GnPs

are distributed without forming conductive percolative networks [11,52]. However, the

electrical conductivity in the SCM HDPE-9.8 vol.%-GnP composite exhibited a frequency-

independent behavior over a frequency range of less than 2×10+3. This indicated that conductive

paths were beginning to form in the solid composite. Furthermore, by increasing the GnP

content to 12 vol.%, the HDPE-GnP composites showed a frequency-independent behavior

across a much wider frequency range. This behavior was represented by σ = σDC + σAC, where

the alternative current’s (AC) conductivity (σAC) is the frequency-dependent part of the total

conductivity. The frequency below which the conductivity becomes frequency-independent,

which is known as the σDC, is also called the critical frequency [11,52].

The physical foaming, however, changed the frequency-independent conductivity behavior of

the SCM HDPE-9.8 vol.% GnP composites (below 2×10+3 Hz) into a fully frequency-dependent

behavior across the entire frequency range (see Figure 6.3a). This is a characteristic of highly

insulating materials. Moreover, the physical foaming decreased the electrical conductivity of the

HDPE-9.8 vol.% GnP composite by more than four orders of magnitude over a frequency range

of less than 1×100.

132

Figure 6.3. (a) Broadband conductivity of the SCM and the extruded HDPE-GnP composite

foams. The extruded foam samples had 0.14±0.01 g.cm-3 (corresponding to ~8 times the foam

expansion ratio); (b) The DC conductivity of the SCM and extruded HDPE-GnP composite

foams as a function of the GnP content measured at 0.1 Hz (X-axis is logarithmic and scales

before and after break are not equidistant). Note that the extruded foams of the 12, 15 and 19

vol.% samples could not be obtained due to the excessive viscosity as discussed in Section

5.4.1.

Figure 6.3b shows the HDPE-GnP composite’s direct electrical conductivity (σDC) at 0.1 Hz.

The GnP content of extruded-foam samples was given with respect to both the polymer volume

and the total volume of the foams [30,35]. To exclusively study the effect of foaming actions on

the σDC, the GnP content was considered with respect to the polymer volume. To include the

effect of density reduction due to foaming on the σDC, the GnP content was calculated with

respect to the total volume of the extruded nanocomposite foams.

The insulation-conduction transition behavior of the SCM HDPE-GnP composites was shown

by an abrupt increase in the σDC of about 5 orders of magnitude, when the GnP content was

increased from 4.5 to 9.8 vol.%. However, the extruded nanocomposite foam’s σDC increased

very gradually by less than 1 order of magnitude when the GnP content was increased from 4.5

to 9.8 vol.%. The percolation threshold of the SCM HDPE-GnP composite was found to be ~

9.8 vol.% GnP. Meanwhile, the extruded-foam counterpart (containing 9.8 vol.% GnP) was

highly insulating (9.8×10-12 S.cm-1). The lower conductivity of the extruded HDPE-GnP

composite foams was mainly attributed to the polymer’s high biaxial stretching during foaming.

100

101

102

103

104

105

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

10-3 4.5 vol.% GnP Solid Foam

7 vol.% GnP Solid Foam

9.8 vol.% GnP Solid FoamB

roa

db

and

co

nd

uctivity,

(

S.c

m-1)

Frequency (Hz)

Neat HDPE

12 vol.% GnP-solid

(a) Solid compression molded

Foam extruded(GnP content with respect to polymer volume)

Foam extruded(GnP content with respect to total volume)

0.5 1 5 10 15 20

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

10-3

Conductivity,

D

C (

S.c

m-1)

GnP content (vol.%)

(b)

4

133

This resulted in the separation of GnPs, and hence the loss of their interconnectivity. And, even

though the sc-CO2 treatment and physical foaming further exfoliated the GnPs, as described in

Section 5.4.2, the polymer’s excessive stretching reduced the conductive paths [32]. The

dilution of the GnPs, is shown in Figure 6.3b where the density reduction due to foaming was

considered. 11 vol.% of the extruded-foam samples was HDPE-GnP, and the gaseous phase

(cells) made up the remaining 89 vol.% (corresponding to 8 times expansion ratio). Therefore,

in the extruded-foam samples, the GnP content in the given volume of the extruded-foams was

only 11% of the GnP content in the corresponding SCM ones. This significantly decreased the

formation of conductive paths and reduced the σDC. For instance, the σDC in the SCM HDPE-9.8

vol.% GnP sample was more than four orders of magnitude higher than that of extruded

nanocomposite foam which was made using a solid precursor containing 9.8 vol% GnP. The

GnP content in this extruded nanocomposite foam was reduced to 1.08 vol.% due to foaming

(i.e. 9.8×0.11).

The observed gradual increase in the electrical conductivity with an increased GnP content in

the foam-extruded samples is favorable in the development of high-performance dielectric

materials. This led to a wider adjustable window near the percolation threshold, where a large

amount of well-dispersed GnPs maximized the dielectric constant. Meanwhile, the dielectric

loss was limited due to the low conductivity that existed [11].

6.4.3 Dielectric properties of the polymer-GnP composites

Generally, the dielectric permittivity is expressed in terms of the following complex function

[6]:

ε (ω) = ε' (ω) - iε'' (ω) (6.1)

where ω is the frequency, ε'(ω) is the real part, and ε''(ω) is the imaginary part of the dielectric

permittivity. The real part of equation (6.1) relates to the charge displacement, which is affected

by different types of polarization within the material [6]. However, in a frequency range below

1MHz, the governing polarization is mainly an interfacial one such as the polarization of the

matrix/filler interface [6,13,53]. Based on the Maxwell–Wagner–Sillars (MWS) effect [54], in

polymer composites with conductive fillers, charges are accumulated at the interface of the

134

fillers and the polymer matrix. This is due to the considerable contrast that exists between the

electrical conductivity of the fillers and the polymer matrix.

On the other hand, the imaginary part of Equation (6.1) indicates that the energy dissipation or

dielectric loss is quantified by a ratio of ε'' to ε' (tan ). The dielectric loss of polymer

composites is generally governed by the following conditions: polarization loss of space

charges, ohmic loss, and the molecular dipole movement (dipole loss) [54,55].

Figure 6.4a-b shows the real dielectric permittivity and the dielectric loss of the extruded foam

(with a density of 0.14±0.01 g.cm-3 or ~8 times foam expansion ratio) and SCM HDPE-GnP

composites as a function of the GnP content measured at 1×10+3 Hz. The GnP content of the

extruded nanocomposite foam was calculated with respect to either the total volume of the foam

or the polymer volume only [30,35]. To solely assess the effects of physical foaming on the

dielectric properties, the GnP content was taken into account with respect to the polymer

volume alone. To consider the effect of the density reduction due to physical foaming as well,

the GnP content was calculated with respect to the total volume of the extruded foam.

An increased GnP content enhanced the real dielectric permittivity of both the SCM and

extruded foam samples. The increased GnP content increased both the GnP-polymer interface

area and the interfacial polarization’s density [24,56]. Moreover, the higher GnP content likely

increased the number of nanocapacitors (consisting of GnP pairs as electrodes with mediated

insulating polymer as nanodielectrics) which enhanced the real permittivity. A higher GnP

content can also decrease the interspatial distances between adjacent GnPs (that is, thinner

nanodielectrics) and can further enhance the real dielectric permittivity [11]. The real

permittivity of the SCM samples at 4.5, 7 and 9.8 vol.% GnP was 9.6, 11.7 and 33.5,

respectively.

On the other hand, the generation of a microcellular structure in the HDPE-GnP composites

significantly improved the real permittivity. In the extruded foam samples, the real permittivity

was enhanced to 18, 39.7 and 77.5 at 4.5, 7 and 9.8 vol.% GnP content, respectively. It is also

notable (Figure 6.4a) that to achieve the same level of the real permittivity, the required GnP

content in the given volume of the extruded nanocomposite foam was significantly lower than

that of the corresponding SCM ones. For instance, to obtain the real permittivity of 80, 11 vol.%

135

GnP was required from the SCM samples, whereas only 1.08 vol.% GnP was required from the

extruded foam samples.

The increased real permittivity of the extruded foam samples, when compared with their SCM

counterparts, was attributed to several factors: (i) The sc-CO2-treatment and physical foaming

provided a higher level of GnP exfoliation than was found in their SCM counterparts (Section

5.4.2). In other words, because the GnPs had a higher level of exfoliation, the GnP-polymer

interfaces and the number of nanocapacitors significantly increased. This also decreased the

interspace distances between the adjacent GnPs. All these factors led to a higher interface

polarization density and dielectric permittivity. (ii) Moreover, the local stretching of the polymer

matrix caused by the cellular growth during the physical foaming generated a unique parallel-

plates arrangement of the GnPs in the cell walls [48]. Specifically, the GnPs were more

perpendicularly oriented in the radial direction with the cell growth. This could enhance the

effective interfaces between the adjacent GnPs. (iii) On the other hand, the polymer’s

compression in between the cells further reduced the interspacial distances between the adjacent

GnPs due to the cellular growth [48]. And this led to a higher real permittivity.

Figure 6.4. (a) Real dielectric permittivity (ε'); and (b) Dielectric loss (tan δ) of the extruded

foam (with the density of 0.14±0.01 g.cm-3 or ~8 times foam expansion ratio) and the SCM

HDPE-GnP composites as a function of GnP content measured at 1×10+3 Hz. Note: X-axis is

logarithmic and scales before and after break are not equidistant.

As Figure 6.4b shows, the SCM samples’ dielectric loss was sharply increased from

0.013±0.002 to 0.7±0.1, when the GnP content was increased from 4.5 to 9.8 vol.% GnP. The

0.5 1 5 10 15 20

0

20

40

60

80

100

120

140 Foam extruded

(GnP content with respect to total volume)

Foam extruded(GnP content with respect to polymer volume)

Solid compression molded

Re

al p

erm

ittivity

GnP content (vol.%)

(a)

4 0.5 1 5 10 15 20

10-2

10-1

100

101

102

103 (b)

Die

lectr

ic loss

GnP content (vol.%)

Foam extruded (GnP content with respect to total volume)

Foam extruded (GnP content with respect to polymer volume)

Solid compression molded

4

136

increased GnP content enhanced the number of (i) mobile charge carriers and (ii)

nanocapacitors, which contributed to both the increased Ohmic loss and the polarization loss.

The sharp increase in the dielectric loss of the SCM nanocomposites around the percolation

threshold (9.8 vol.% GnP) was mainly attributed to the formation of conductive paths, which led

to a significant Ohmic loss.

However, the introduction of a microcellular structure significantly decreased the dielectric loss

of the HDPE 9.8 vol.% GnP composites by two orders of magnitude. This was contrary to what

occurred with the SCM samples, because the dielectric loss of the extruded foam samples

increased very slightly from 0.010±0.005 to 0.014±0.005 when the GnP content was increased

from 4.5 to 9.8 vol.%. The lower dielectric loss of the extruded nanocomposite foams can be

largely attributed to the lower leakage current and consequently to the lower Ohmic loss, which

was proportional to the σDC. In other words, the polymer matrix’s high biaxial stretching during

the cellular growth resulted in the GnP’s dilution and, hence, the loss of their connections and a

lower σDC. The differences between the dielectric loss of the SCM samples and that of the

extruded foam samples were much more pronounced when the GnP content was near the

percolation threshold (9.8 vol.% GnP). This was where the introduction of the cellular structure

broke the conductive networks due to local stretching of the polymer matrix. Specifically, at a

9.8 vol.% GnP content, the σDC of the extruded foam samples was more than four orders of

magnitude lower than the SCM samples’.

Figure 6.5a-b show the broadband real dielectric permittivity and dielectric loss of the SCM and

extruded-foam samples for the HDPE-9.8 vol.% GnP composites. The real permittivity of the

SCM nanocomposites was highly frequency-dependent and increased as the frequency

decreased. This is the characteristic of conductive composites [30,54]. However, the

incorporation of the microcellular structure not only enhanced the real permittivity by more than

two-fold, but also changed the highly frequency-dependent behavior of the real permittivity in

the SCM samples to the relatively frequency-independent permittivity present in the extruded

foam samples, which is one of the requirements for reliable dielectric materials [13].

As seen in Figure 6.5b, the broadband dielectric loss of the extruded foam samples was around

two orders of magnitude lower than in the SCM samples across the entire frequency range. The

lower dielectric loss of the extruded foam samples was attributed to their lower Ohmic loss,

which was related to the σDC. In the current percolative system, the dielectric loss mainly

137

consists of the Ohmic loss and the polarization loss of space charges. And, generally, the

contribution of the Ohmic loss to the total dielectric loss is greater than the polarization loss

[56,57].

It was notable that in Figure 6.5b the dielectric loss of the extruded foam samples was relatively

frequency-dependent, with far lower values at higher frequencies (>103 Hz). The lower

dielectric loss at higher frequencies was proportional to the polarization loss [56,58], This was

further restricted after foaming was introduced. Thus, the dielectric loss dropped down to 0.003

at 1×105 Hz in the extruded foam samples. However, the dielectric loss of their SCM

counterpart samples was 0.18 at 1×105 Hz, and it sharply increased as the frequency decreased.

The reduction in the polarization loss in the extruded foam samples was attributed to a better

dispersion of the GnPs and to their better interfacial interaction with the polymer matrix. This

was due to the dissolution of the sc-CO2 in the polymer. Thus insulating layers formed among

the GnPs to prevent the migration of the space charge within the nanocomposites [55].

Figure 6.5. (a) Broadband dielectric permittivity; (b) Broadband dielectric loss of the SCM

HDPE-9.8 vol.% GnP composites and their extruded foam (with a density of 0.15 g.cm-3 or ~8

times foam expansion ratio) counterparts.

101

102

103

104

105

0

20

40

60

80

Neat HDPE

Bro

adband r

eal perm

ittivity (

' )

Frequency (Hz)

(a)

Solid

Foam

Foam

101

102

103

104

105

10-4

10-3

10-2

10-1

100

101

102

Solid

(b)

Neat HDPE

Bro

ad

ba

nd

die

lectr

ic lo

ss (

tan

)

Frequency (Hz)

138

6.4.3.1 Effect of the density on the dielectric properties

Figure 6.6 shows the variation of the real permittivity and the dielectric loss as a function of the

density of the extruded foam samples, which were made from solid precursors containing 4.5

vol.% GnP. When the density of the HDPE-GnP composites decreased from 1.07 to 0.08 g.cm-3,

the ε' increased from 9.6 to 22.3, and the tan decreased from 0.04 to 0.006. A further decrease

in the density to 0.05 g.cm-3, however, slightly decreased the ε' to 18.8. Meanwhile, the

dielectric loss continued to decrease to as low as 0.004. The optimal behavior of the ε' can be

mainly attributed to the changes in the microcellular structures, when their densities (that is, the

foam expansion ratios) were varied as was discussed in Section 5.4.1.

Figure 6.6. Variations in real permittivity and dielectric loss measured at 1×10+3 Hz as a

function of density in the extruded HDPE-GnP composite foams made from solid precursors

containing 4.5 vol% GnP

The parallel-plates arrangement of the GnPs within the cell walls enhanced the effective

interfaces between the adjacent GnPs (Figure 6.7a), and thus increased the real permittivity. As

the density continued to decrease, the cell walls were compressed in the thickness direction due

to bubble growth (Figure 6.7b), which can decrease the interspace distances between the

adjacent GnPs. This further enhanced the ε'. However, when the density was decreased to 0.05,

the cell wall thickness dropped to approximately 1 μm, and the number of adjacent GnPs in the

0.1 1

5

10

15

20

25 Real permittivity

Dielectric loss

Density (g.cm-3)

Real perm

ittivity (

' )

10-3

10-2

10-1

100

Die

lectr

ic loss (

tan

)

139

cell walls decreased dramatically, due to polymer matrix’s excessive compression. Interestingly,

almost no GnP can be found in the cell wall for the lower density foams (Figure 6.7c). This

resulted in lower effective interfaces between the adjacent GnPs and in lower real permittivity.

During the fast cellular growth process, the solid fillers (e.g. GnPs) could barely flow together

with the polymer melt [32]. Therefore, the number of GnPs in a unit area of the cell walls

decreased with the reduction in the density of foam-extruded samples. The decrease in the tan

when the density dropped can mainly be attributed to the higher foam expansion (up to 21-fold),

which further separated the GnPs and, thereby, resulted in a loss of conductive networks. This,

in turn, led to a lower Ohmic loss. Figure 6.7d shows the ideal 2-D conceptualization of the

change in the GnP’s alignment with the density.

Interestingly, the tailored microcellular structure further enhanced the HDPE-GnP composites’

dielectric performance. The facile sc-CO2-treatment and physical foaming of the HDPE-GnP

composites significantly increased their real permittivity (ε') and greatly decreased both their

dielectric loss (tan ) and their density. As a result, an excellent combination of the dielectric

properties, together with an ultra-low density, resulted. In the extruded 1.08 vol.% HDPE-GnP

composite foam with a density of 0.15 g.cm-3, the tan dropped down as low as 0.003 at 1×105

Hz while the ε' reached 77.5. These results were greatly superior to those of the SCM

nanocomposites (ε'~19.9, tan ~0.15 and density of ~1.2 g.cm-3).

140

Figure 6.7. SEM and TEM micrographs of the extruded HDPE-GnP composite foams made

from solid precursors containing 4.5 vol% GnP, which show the GnPs’ arrangement at different

densities including: (a) 0.13 g.cm-3; (b) 0.08 g.cm-3; and (c) 0.05 g.cm-3. (d) Ideal 2-D

conceptualization of GnP’s arrangement in cell walls as the density decreased.

Table 6.1 shows some of the recent advances made in the development of polymer

nanocomposites as dielectric materials. And they are compared with the dielectric performance

of the HDPE-GnP reported in our study. Most of the presented batch-type studies

[3,4,29,11,16,17,24–28] have undertaken complex synthesis procedures which are challenging

to be scaled up and/or have high material cost. For instance, Wen et al. achieved a very good

combination of the real permittivity (ε'=74) and the dielectric loss (tan =0.08) [16], however,

their fabrication method was a tedious multiple-step synthesis process (three-step process for

GnP preparation + synthesis of poly- (vinylidene fluoride-trifluorethylene) copolymer with

141

internal double-bonds through a dehydrochlorination process). Jin et al. also reported another

great combination of dielectric properties (ε'=71.7, and tan =0.045) [11] for the poly

(vinylidene fluoride) hybrid nanocomposites (containing MWNT/BaTiO3) which were

fabricated through a complex miscible-immiscible coagulation method. However, the required

nanomaterials loading was rather high which resulted in heavy (density of ~4.6 g.cm-3) and

expensive dielectric materials.

Table 6.1. Dielectric performance and density of different polymer nanocomposites

Materials Filler

content

Frequency

(Hz)

Dielectric

permittivity

Dielectric

loss

Fabrication

method

Density

(g.cm-3) Ref.

poly(vinylidenefluoride-

trifluorethylene)/graphene

nanosheets

4.0 vol% 103 74 0.08 solution casting,

functionalizatio

n and

crosslinking

~1.8 [16]

poly (vinylidene fluoride)/

BaTiO3 /BaTiO3 nanofibers

30 vol%/3

vol%

102 27 0.06 emulsion

polymerization

~1.1 [29]

poly (vinylidene fluoride-co-

hexafluoropropylene)/titaniu

m dioxide-modified reduced

graphene oxide (rGO)

20 wt.% 102 24.5 0.22 in-situ

assembling TiO2

on graphene

oxide (GO)+

solution mixing

and drop casting

~1.86 [4]

poly(p-phenylene

benzobisoxazole)/

Functionalized Graphene

Nanosheets

2.0 wt.% 103 66.27 0.045 polymer chains

grafting in

GnPs, in-situ

polymerization

followed by

further thermal

treatment

~1.57 [26]

poly (vinylidene

fluoride)/functionalized

graphene–BaTiO3

1.25/30

vol.%

106 65 0.35 GO synthesis

followed by

two-step

solution mixing

and hot-pressing

~3 [3]

142

polydimethylsiloxane/therma

lly expanded graphene

nanoplates

2.0 wt% 103 89 1.5 thermal

exfoliation of

graphene

followed by

solution mixing

and

vulcanization

~0.97 [17]

polyimide /graphene/BaTiO3 1.0/16vol.% 102 31 0.03 Multiple step

solution mixing

~3.5 [27]

cyanoethyl pullulan

polymer/carbon

nanotubes/rGO

0.062 wt.% 102 32 0.051 carbon

nanotubes/rGO

fabricated by

thermal CVD

followed by

solution mixing

~1.1 [24]

poly (vinylidene fluoride) /

multiwall carbon nanotubes

/BaTiO3

3.0/37.1vol.

%

103 71.7 0.045 miscible-

immiscible

coagulation

method

followed by hot

pressing

~4.6 [11]

diglycidyl ether of bisphenol-

A/rGO

1.0 wt. 103 32 0.08 covalent

functionalizatio

n and solution

mixing and

curing

~1.7 [28]

HDPE/GnP 1.08 vol%* 103 77.5 0.014 melt mixing and

foam extrusion

0.15 this work

HDPE/GnP 1.08 vol%* 105 77.1 0.003 melt mixing and

foam extrusion

0.15 this work

HDPE/GnP 0.8 vol.%* 103 39.7 0.012 melt mixing and

foam extrusion

0.14 this work

HDPE/GnP 0.5 vol.%* 103 18 0.010 melt mixing and

foam extrusion

0.13 this work

HDPE/GnP 0.5 vol.%* 103 22.3 0.006 melt mixing and

foam extrusion

0.08 this work

* the GnP vol.% is reported with respect to the total volume

143

6.5 Conclusion

In this study, a new class of ultralight, microcellular dielectric HDPE-GnP composites was

introduced. The composite foams of HDPE with highly exfoliated GnPs were developed using a

melt mixing method, followed by a sc-CO2-treatment and physical foaming via an extrusion

process. The generation of a microcellular structure provided a unique parallel-plate

arrangement of GnPs around the cell walls. This significantly enhanced the real permittivity and

greatly decreased the dielectric loss of the HDPE-GnP composites. For example, ultralight

extruded HDPE-1.08 vol.% GnP foam with a density of 0.15 g.cm-3 had a high dielectric

permittivity of ε'=77.5 and an extremely low dielectric loss of tan δ=0.003 at 1×105 Hz, which

made it superior to solid compression-molded samples (ε'=19.9 and the tan δ=0.15 with a

density of 1.2 g.cm-3). The extremely low dielectric loss, together with the enhanced real

permittivity, of the extruded foam samples provided an excellent combination of dielectric

properties. Our study showed that the tailored morphologies existing in the microcellular

structure within the HDPE-GnP composites offer a novel, industrially viable and cost-effective

method to develop ultralight dielectric materials with high permittivity and low dielectric loss.

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150

CHAPTER 7

7 Enhancement of the dielectric performance of polymer-GnP composites using SCF-treatment and physical foaming-Part II

The following section is based on text from (†equal contribution)

Zhao, B.†, Hamidinejad, S. M.†, Zhao, C., Li, R., Wang, S., Kazemi, Y., and Park, Chul B., “A Versatile

Foaming Platform to Fabricate Unprecedentedly High Dielectric Permittivity, Ultra-Low Dielectric Loss

of Polymer/Carbon Composites”, Journal of Materials Chemistry A, 2019, 7 (1), 133-140, DOI:

10.1039/C8TA05556D

7.1 Summary

There is an urgent need for dielectric-

based capacitors to manage the increase

in storage systems related to renewable

energy production. Such capacitors must

have superior qualities that would include

light weight, a high dielectric constant,

and an ultra-low dielectric loss. The

poly(vinylidene fluoride) (PVDF)-

graphene nanoplatelet (GnP))

http://mpml.mie.utoronto.ca/info/personnel/yasamin-kazemi/

151

nanocomposite foams are considered promising alternatives to solid PVDF-GnP

nanocomposites. This is because they have excellent dielectric properties, which are due to the

preferred orientation of their GnPs occurring in the foaming process. In the PVDF-GnP foams,

the microcellular structure significantly influenced their electrical conductivity and dielectric

properties. The presence of a microcellular structure induced parallel arrangement of GnPs

isolated by insulating polymer or the air, as a medium between themselves. An unprecedentedly

high dielectric constant of 112.1 and an ultra-low dielectric loss of 0.032 at 100 Hz were

obtained from the PVDF-GnP composite foam with high expansion ratio of 4.4 due to the

charge accumulation at the aligned conductive filler/insulating polymer (or bubble air) interface.

7.2 Introduction

Due to population growth, global warming, and the energy crisis, the development of renewable,

cost effective and green energy techniques to supply future generations with renewable energy

is both challenging and urgent [1]. Of the various energy storage systems that exist, dielectric

capacitors with an ultrafast charging−discharging ability have become an important technology.

These could greatly benefit the high-performance power electronics used in military power

systems, hybrid electric vehicles, and in some portable electronics [2,3]. Ceramic-based

dielectric materials, such as SrTiO3 [4], SiC [5], and BaTiO3 [6], have high dielectric constant

values, which is why they play a major role in current practical applications. However, their

numerous serious defects, which include an insurmountable brittleness and a low electrical

breakdown strength, have hindered the development of dielectric materials [7–9]. Compared to

conventional ceramic-based dielectric materials, polymer-based dielectric materials have several

advantages: large-scale processability, mechanical flexibility, light weight, low cost, and high

electrical breakdown strength. However, most polymers’ dielectric constant is low compared to

the dielectric constant in inorganic ceramics. For example, polypropylene, polystyrene,

polyacrylates, and polymethacrylates usually have dielectric constant values between 2 and 5

[10].

To address these issues, significant efforts have been made to create polymer-based dielectric

materials with high permittivity. One effective strategy has been to introduce high-dielectric-

constant (high-k) ceramics into the polymer matrix, and this strategy has been extensively

152

investigated in the past few decades [11–14]. However, it was reported that these composites

always possess a high concentration of ceramic particles, which severely damaged the polymer

composites’ processability and mechanical flexibility. In addition, the poor compatibility

between the inorganic fillers and the organic polymer matrix resulted in weak interfacial

adhesion and aggregation. This led to a high dielectric loss and a further decrease in the

breakdown strength [3].

When a conductive filler is added to a polymer matrix, under an alternating electromagnetic field,

charge accumulation between the conductive filler and the polymer matrix would take place. In this

situation, numerous nanocapacitors are formed in which the conductive fillers and the polymer matrix

are considered as electrodes and a dielectric, respectively. When the content of the conductive filler is

less than the percolation threshold, the dielectric constant is enhanced dramatically with an increased

filler content, while the dielectric loss is moderately increased. Finally, it is possible to obtain a

supercapacitor. However, when the content of the conductive filler is very close to or above the

percolation threshold, conductive networks are generated and, thus, significant dielectric loss caused

by leakage conduction loss from the capacitors would occur. Consequently, the breakdown of

capacitors would occur, and the leakage conduction loss would be beneficial for the EMI shielding

properties. Thus, selecting a reasonable content of the conductive filler is the crucial step to determine

the application (i.e., capacitor application or EMI-shielding application) of the polymer–conductive

filler composites.

In recent years, one-dimensional (1D) and two-dimensional (2D) carbon nanomaterials with

large aspect ratios have become potential candidates to prepare high- k nanocomposites. Carbon

nanotubes (CNTs) and graphene nanoplatelets (GnPs) have been used most frequently for this.

In light of the percolation phenomenon, it is well accepted that a sudden increase in the

composites’ permittivity about one or even several orders of magnitude occurs when the loading

of conducting nanomaterials reaches a critical value, i.e., the percolation threshold [13,14]. We

noted that a large number of conductive networks would form in the composites if the percolation

threshold was exceeded. Therefore, to obtain a high permittivity, the amount of conducting

nanomaterials in the polymer matrix composites should be near the percolation threshold without

exceeding it. This means that the polymer composites with conducting fillers should still act as

insulators under this condition. Due to the excellent conductivity of carbon nanomaterials (CNTs or

GnPs), a high dielectric constant is easily obtained near the percolation threshold in carbon

153

nanomaterial/polymer composites. However, the formation of conductive channels also increases the

leakage current, which causes a high dielectric loss and a low breakdown strength [15,16]. This

dilemma seriously hampers the development of carbon polymer dielectric composites. Thus, the core

question is how to separate the adjacent carbon materials (CNTs or GnPs).

Recently, Ameli et al. investigated the dielectric properties of polypropylene (PP)–MWCNT

nanocomposite foams [17,18]. They found that adding a microcellular structure effectively improved

the dielectric constant and decreased the dielectric loss. They reported a real permittivity of 30 and a

dielectric loss of 0.07 for the PP–0.34 vol% MWCNT foams [18]. In addition, a microcellular PP–

1.25 vol% MWCNT had a dielectric permittivity of 57.2 and a dielectric loss of 0.05 [17]. In this

study, PVDF was used as the polymer matrix. This was because of the strong electric dipole moment

of its molecular chains and its high dielectric constants in the range of 8–10 [19]. Also, the GnPs were

selected as conductive fillers. Also, the correlation between a microcellular structure and dielectric

properties was investigated. Amazingly, an unprecedentedly high real permittivity with an extremely

low loss-tangent value was obtained from the foamed PVDF/GnP composites.


7.3.1 Materials

The PVDF (molecular weight 300,000–330,000 g/mol) was supplied by Solvay. The GnPs were

supplied by Group Nanoxplore, Inc. (N,N)-Dimethylformamide (DMF) was provided by

Caledon Laboratories Ltd. The raw materials were used as is, without further purification.

7.3.2 Fabrication of PVDF-GnP Solid Composites

The PVDF-GnPs solid composites were prepared by solvent casting. In this study, 2.0 wt%

GnPs were uniformly dispersed in the DMF solution using the ultrasonication process. These

content were chosen near the percolation threshold (5.0 wt% GnPs obtained for the same

materials in our previous study [20]) without exceeding it, as mentioned above. Then, the PVDF

particles were dissolved by magnetic stirring in the DMF mixture. Finally, the PVDF-GnPs

solid composites were obtained through the evaporation and compression-molding processes.

154

7.3.3 Fabrication of PVDF-GnP Composite Foams

We used a homemade batch foaming device, which is shown in Figure 7.1, to prepare the

PVDF-GnP samples’ foaming behavior. The foaming system consisted of a syringe pump filled

with CO2 as the physical blowing agent, and a foaming chamber with a thermal couple to detect

the temperature, along with a heater and a depressurizing valve. At the beginning of the

experiment, the chamber was heated to the desired temperature, and the sample (25 mm× 15

mm× 12 mm) was laid inside it. Subsequently, the CO2 was quickly released in and out of the

chamber so as to eliminate the air. Then, the CO2 was pressurized into 2,000 psi (13.8 MPa) at

an experimental temperature and was held there for 1 hour. The pressure was quickly released,

and the chamber was quenched in cold water. Finally, the sample was removed from the

chamber. Various foaming temperatures of the PVDF-GnP composites in the ranges of 167°C to

169°C were investigated. The foams prepared at various saturation temperatures (ranging from

167°C to 169°C) were conveniently denoted as FG1-FG5, as shown in Table 7.1.

Figure 7.1. A schematic illustration of the home-made batch-foaming device

The preparation of PVDF-GnP composite foams is a two-step process containing fabrication of

PVDF-GnP solid composites and foaming of PVDF-GnP solid composites. The PVDF-GnP

solid composites were prepared by solvent casting. This was followed by compression-molding,

Then, a homemade batch foaming device was used to prepare the PVDF-GnP composite foams.

155

Table 7.1. Expansion ratio of PVDF-2wt% GnP foams obtained at various saturation

temperatures

Foam FG1 FG2 FG3 FG4 FG5

Temperature (°C) 167 167.5 168 168.5 169

Expansion ratio 2.1 2.5 4.4 4.0 2.7

Figure 7.2 shows the PVDF-GnPs foams’ fabrication process. First, both the PVDF particles

and the GnPs were added to the DMF solvents and evenly blended. Second, the PVDF-GnP

solid composites were fabricated through the casting and compression processes. Third, a

homemade batch foaming device was used to foam the PVDF-GnP composites. The PVDF-GnP

foams were successfully fabricated.

Figure 7.2. A schematic diagram of the PVDF-GnP foam fabrication process


The densities of the solid (ρs) and foam (ρf) composites were measured using the water

displacement method (the ASTM D792-00). The cell density was calculated based on scanning

electron microscopy (SEM) images using the following formula [21]:

156

𝐶𝑒𝑙𝑙 𝑑𝑒𝑛𝑠𝑖𝑡𝑦 = (𝑛𝐴⁄ )

32⁄ × (

𝜌𝑠𝜌𝑓

⁄ ) (7.1)

where n is the number of cells in the designated area (A) in the SEM micrograph, respectively.

The volume expansion ratio was determined as 𝜌𝑠

𝜌𝑓⁄ .

The morphologies of the PVDF-based foam samples were examined using the SEM (JSM-6060)

The electrical conductivities and broadband dielectric spectroscopy measurements of PVDF-

GnP composite foams were carried out using an Alpha-N analyzer from Novocontrol

Technologies GmbH & Co. KG. The detailed experiment process was discussed in Section

6.3.2.


Figure 7.3 shows the cellular properties (the expansion ratio, the cell morphology and cell

density) of various PVDF-GnP composite foams. As shown in Figure 7.3a, the expansion ratios

PVDF-GnP composite foams firstly increased and then decreased with increasing the saturation

temperature. It is noteworthy that PVDF-GnP composite foams display a similar cell density

(Figure 7.3c). We believe that these changes were the result of variations in the crystal structure of

the PVDF matrix, which had been treated at different saturation temperatures [22–24]. This variation

of expansion ratio and morphological foam properties provided a complete platform by which

the dielectric properties of the PVDF-carbon composite foams can be tuned.

157

Figure 7.3. (a) Expansion ratio of PVDF/GnP composite foams; (b) SEM image of FG3 foam

sample, and the inset is the corresponding magnification SEM; the cell density of PVDF/GnP

Figure 7.4a-d show the frequency-dependent electrical conductivity, real permittivity (ε′), the

imaginary permittivity (ε′′), and the dielectric loss (tan δε) for the solid and foamed PVDF-GnP

composites. Astonishingly, all of the PVDF-GnP composite foams displayed a much higher dielectric

constant than that of their solid counterpart (Figure 7.4b). The Maxwell–Wagner–Sillars (MWS)

polarization, which is also called interfacial polarization, for heterogeneous systems plays a very

important role in improving the dielectric constant [25,26]. The MWS effect is associated with the

entrapment of the free charges between the insulator/conductor interfaces. Moreover, the large

number of nanocapacitors that form between two parallel GnPs would also enhance the dielectric

constant. Furthermore, the FG3 foam had the highest dielectric constant value of the five composite

foams, which was due to its relatively high electrical conductivity, and more parallel nanocapacitors.

Correspondingly, the imaginary permittivity (Figure 7.4c) and the dielectric loss (Figure 7.4d) of

both the solid and foamed PVDF-GnP composites had low values because of the insulation

_Ref523827054

_Ref523827054

_Ref523827054

_Ref523827054

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_Ref523827054

158

properties. Moreover, we compared both the dielectric constant and dielectric loss of the solid and

PVDF-GnP composite foams at 100 Hz (Figure 7.4e).

Figure 7.4. (a) Frequency-dependent electrical conductivity of the solid and foamed PVDF/GnP

composites, (b) Real permittivity, (c) Imaginary permittivity, and (d) Dielectric loss of the solid

and foamed PVDF/GnP composites as a function of applied frequencies ranging from 1 Hz to

300,000 Hz, (e) Real permittivity and dielectric loss of the solid and foamed PVDF/GnP

composites in the 100 Hz frequency, (f) The correlation amongst the real permittivity, the

dielectric loss and the expansion ratio of the foamed PVDF/GnP composites in the 100 Hz

frequency.

159

It is noteworthy that the dielectric constant values of all the foam samples had been significantly

increased when compared with the solid sample. Specifically, for the FG3 with the highest expansion

ratio, the dielectric constant improved nearly one order of magnitude, and a value of 112.1 was

obtained. In addition, it is inspiring that the dielectric loss had ultra-low values of less than 0.05. This

meets the requirements for the miniaturization of dielectric capacitors [2]. The GnP particles cannot

easily establish a good contact surface area with each other and, therefore, dielectric loss was not

high. So, the maximum dielectric tangent loss was still very low (0.032) at the maximum expansion

ratio, which is a unique feature of the GnP case.

The conductivity, the real permittivity, and the dielectric tangent loss became relatively larger with an

increased expansion ratio. Because of the limited expanding ability of the GnP samples (Figure

7.4e), we could not achieve a very high expansion ratio from the GnP samples. Furthermore, as the

expansion ratio was increased, the GnP particles would become more perpendicular to the radial

direction [27]. Consequently, the GnPs became more parallel with respect to each other regardless of

the initial orientation [28]. One can note that a large expansion ratio brings the platelets closer in

general as well. Consequently, the real permittivity increased. For FG3, the generation of a

microcellular structure produced numerous parallel-plate nanocapacitors consisting of GnP pairs as

electrodes with the insulating polymer as nanodielectrics. In addition, the increased dielectric constant

of the FG3 foam sample also resulted from several factors. FG3 had the highest expansion ratio

amongst all the samples. As the degree of foaming increases, the GnPs will be oriented more

perpendicular to the radial direction, regardless of their initial orientation. So, initially non-parallel

GnPs will become more parallel with respect to each other. Because of the highest expansion ratio of

the FG3 foam sample, the GnP–polymer interfaces and the number of nanocapacitors significantly

increased. This also decreased the interspace distances between the adjacent GnPs because the

polymer melt was compressed between the two growing cells. All these factors led to a higher

interface polarization density and dielectric permittivity. Thus, the FG3 foam sample possessed the

highest dielectric constant. Figure 7.4f presents the correlation between the real permittivity, the

dielectric loss and the expansion ratio of the foamed PVDF-GnP composites. Intriguingly, based on

this unique two-dimensional contour of the real permittivity and dielectric loss as a function of the

expansion ratio, the PVDF-GnP composite foams with a high real permittivity and a low dielectric

loss can be designed.

The PVDF-GnP composite foams showed that adding a blowing gas and thereby a microcellular

structure to the composites separated the adjacent GnPs, and this decreased the interconnectivity of

_Ref523827054

_Ref523827054

_Ref523827054

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GnPs leading to lower electrical conductivity of the composites (Figure 7.4a). But as the degree of

foaming increased, the GnPs oriented more perpendicular to the radial direction [27] and the

interconnectivity was improved while maintaining the insulating nature. The FG3 foam possessed a

high dielectric constant and an ultra-low dielectric loss (Figure 7.4c and d). In addition to the

Maxwell–Wagner–Sillars (MWS) polarization contribution to the dielectric constant, the formation

of a larger number of nanocapacitors in the nanocomposite foams (conductive GnPs as the electrodes;

an insulating bubble and polymer as the dielectrics) played a major role in improving the dielectric

constant [26]. More precisely, the graphene nanoplatelets were parallel to each other and were

isolated by the polymer layer, or air, as a medium between them. We believe that the excellent GnP

dispersion and the existence of many nanocapacitors in the PVDF polymer foams aid in achieving a

high dielectric constant, but with a low GnP loading and an ultra-low dielectric loss [26].

7.5 Conclusion

In summary, we have reported, for the first time, the fabrication of PVDF-GnP nanocomposite foams

with an enhanced dielectric constant. We effectively tuned the electrical conductivity and dielectric

constant by controlling the microcellular structure. In the PVDF-GnP composites, foaming

significantly enhanced the dielectric constant and decreased the dielectric loss, while the insulating

nature of the samples was maintained. This resulted in an excellent dielectric property combination.

For instance, the FG3 foam sample had the highest expansion ratio, the highest dielectric constant

(112.1) and a low dielectric loss of 0.032 at 100 Hz, which resulted from the interfacial polarization,

and a larger number of nanocapacitors. This novel methodology promises to lead the way to a new

design for lightweight energy storage capacitors, which are made from polymer-GnP foams, and

which have a high dielectric constant and an ultra-low dielectric loss.

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CHAPTER 8

8 Contributions and Future Work

8.1 Contributions

Polymer composites have shown impressive potential as a highly desirable class of functional

materials for use in various applications such as heat sinks in the miniaturized electronic

devices, electromagnetic interference (EMI) shielding, and capacitors (dielectric materials).

Moreover, the recent advances in conducive nanofillers such as graphene with exceptional

thermal and electrical conductivity, have significantly increased the opportunities to develop

functional polymer nanocomposites. However, it is extremely challenging to exploit graphene’s

full potential due to the complexities in the exfoliation, dispersion, and control of the graphene

nanoplatelets’ (GnP) orientation within the composites.

Therefore, this PhD research has aimed to strategically address these challenges by using

supercritical fluid (SCF)-treatment and microcellular foaming to manufacture graphene-based

polymer composites with enhanced functional properties. As a result, this thesis has achieved

the followings:

• The research in this thesis contributed in the invention and development of a novel and

industrial-scale technique for in situ exfoliation and dispersion of GnP in polymer matrices.

This method consolidated two subsequent steps for manufacturing graphene-based polymer

nanocomposites including (i) exfoliation of GnPs and (ii) their dispersion and compounding

within the polymer. The findings of this research are critically important to the large-scale

production of high quality and low-cost polymer-GnP composites with enhanced functional

properties. The result of this research has been filed as a patent and licensed by NanoXplore

Inc., Montreal, QC.

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• This study also developed a fundamental understanding of exfoliation process induced by

SCF-treatment and physical foaming. This in-depth understanding presents valuable

opportunities for generalizing this method to exfoliation of other 2D materials such as hBN.

• This research presented a scientific understanding of structure-thermal conductivity

relationships for graphene-based polymer composite foams. In particular, the effects of

GnPs’ orientation, exfoliation, and degree of foaming were fundamentally studied. Given this

developed knowledge, this PhD thesis presented a new class of lightweight microcellular

polymer-GnP composite foams with high thermal conductivity. In this study, the composites

were fabricated using melt mixing, followed by SCF-treatment and foam injection molding.

The SCF-treated nanocomposites offered much higher thermal conductivity as compared to

their regular injection-molded counterparts due to higher level of exfoliation, dispersion,

interconnectivity, and random orientation of the GnPs.

• This thesis also developed an in-depth understanding of the effects of cellular structures,

GnPs’ orientation, interconnectivity, and exfoliation on the electrical conductivity,

percolation threshold and EMI shielding effectiveness of the graphene-based polymer

composites. With this developed structure-electrical property relationships, the current

research demonstrated that the introduction of foaming and a microcellular structure can

substantially increase the electrical conductivity and can decrease the percolation threshold of

the polymer-GnP composites. The nanocomposite foams had a significantly higher electrical

conductivity, a higher dielectric constant and a higher EMI shielding effectiveness and a

lower percolation threshold compared to their regular injection-molded counterparts.

Foaming also changed the GnP’s flow-induced arrangement by reducing the melt viscosity

and cellular growth. Moreover, foaming rearranged the GnPs to be mainly perpendicular to

the radial direction of the bubble growth. This enhanced the GnP’s interconnectivity and

produced a unique GnP arrangement around the cells to efficiently shield the EMI.

• The development of the dielectric graphene-based polymer composite foams with high real

permittivity and low dielectric loss is greatly promising yet quite challenging. This PhD

research developed a fundamental understanding on the structure-dielectric performance

relationships in graphene-based polymer composites which is critical for the advancement of

these materials. In particular, the effects of microcellular structure and GnPs’ arrangement on

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the interfacial polarization, broadband real permittivity, broadband dielectric loss

(polarization loss and ohmic loss), DC conductivity and broadband conductivity were

scientifically studied. This developed knowledge provided a guide line to optimize the

structures of graphene-based polymer composites and to achieve high real permittivity and

low dielectric loss. For instance, this PhD research presented an industrially-viable technique

(i.e extrusion foaming process) for manufacturing a new class of ultralight (0.05-0.15 g.cm-3)

polymer-GnP composite foams with excellent dielectric. The introduction of microcellular

structure produced numerous parallel-plate nanocapacitors consisting of GnP pairs as

electrodes with insulating polymer as nanodielectrics. This significantly increased the real

permittivity and decreased the dielectric loss. The ultralight extruded HDPE-GnP composite

foams, with a 0.15 g.cm-3 density, had an excellent combination of dielectric properties

(ε'=77.5, tan δ=0.014 at 1000 Hz) which were superior to their compression-molded

counterparts (ε'=19.9, tan δ =0.15 and density of =1.2 g.cm-3) and to those reported in the

literature. The developed understanding of structure-dielectric performance relationships was also

implemented in another system with PVDF as the polymer matrix. This was because of the strong

electric dipole moment of the PVDF molecular chains and its high dielectric constants which is in

the range of 8–10. In the fabricated PVDF-GnP composites, the incorporation of microcellular

structure significantly enhanced the dielectric constant and decreased the dielectric loss. The

foamed PVDF-GnP sample exhibited the highest dielectric constant (112.1) and a very low

dielectric loss of 0.032 at 100 Hz.

This thesis pointed towards the further development of lightweight and functional graphene-

based polymer nanocomposites with tailored properties for various applications.

8.2 Future Work

8.2.1 Thermal and Electrical Conductivities of Graphene-Based Polymer

Composites with the Geometrical Characteristics of GnPs

The lateral size and thickness are important physical parameters of GnPs which can greatly

affect the thermal and electrical conductivities of the graphene-based polymer composites.

Therefore, it is necessary to clearly investigated the relationships between the thermal and

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electrical conductivities of graphene-based polymer composites by measuring the realistic size

(lateral and thickness) of GnPs within the polymer composites (using a nondestructive method

such as X-ray Computed Tomography). Prospective findings of this study will provide valuable

information to optimize the geometrical characteristics of GnPs to achieve desired thermal and

electrical conductivities.

8.2.2 The Development of the Thermally Conductive Graphene-Based

Polymer Composites with High Thermal Stability

Currently, electronic devices are getting smaller and smaller, leading to an increased heat

generation from the device. In order to develop efficient polymer composites as the heat sink

components in electronic devices, the thermal stability of the composites need to be improved.

Therefore, it is crucial to develop graphene-based polymer composites with high performance

engineering polymers with high glass transition temperature (Tg) or melting point (Tm) such as

polycarbonate (PC), polyether ether ketone (PEEK), polysulfone (PSU), polyethersulfone

(PESU) and polyphenylsulfone (PPSU).

8.2.3 SCF-Assisted Manufacturing of Hexagonal Boron Nitride (hBN)-

Polymer Composites with Enhanced Thermal Conductivity

The crystal structure of hBN and GnP platelets are the same and both are highly thermally

conductive. However, hBN are electrically insulative. Since, we have already demonstrated that

SCF-treatment and physical foaming can greatly enhance level of GnPs’ exfoliation and provide

a tailored structure that effectively improve the thermal conductivity of polymer-GnP

nanocomposites, it is highly promising to investigate the SCF-treatment and physical foaming

on hBN-polymer composites. The successful completion of the proposed research will result in

the fabrication of the light-weight thermally conductive electrically insulative polymer

composites.

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8.2.4 Generalizing the SCF-Assisted Exfoliation Method to Other 2D

Materials

2D nanomaterials such as hexagonal boron nitride (hBN), tungsten dichalcogenides (WS2 and

WSe2), and molybdenum disulfide (MoS2) have intriguing properties which can be useful for

various applications such as composites, thermal energy harvesting, electronics, batteries,

nanoelectromechanical systems and sensing. Therefore, efficient exfoliation of these 2D

materials is essential for their advancement. The SCF-assisted exfoliation of these 2D materials

is as worthy areas of future investigation. Prospective findings of this study will present

valuable opportunities for understanding of mechanical and functional properties of these 2D

materials at different size scale. Meanwhile, the efficiency of SCF-assisted exfoliation

implemented on various 2D materials can provide fundamental understanding on how different

interlayer interactions are in these materials.

8.2.5 Fatigue Behavior of Graphene-Based Nanocomposite

Graphene has emerged as attractive building blocks in next-generation carbon-based

nanocomposite materials with high strength/toughness-to-weight ratio. Composites

incorporating graphene have the potential to dramatically improve energy efficiency in the

aerospace industry which has recently made a dramatic transition towards the use of carbon

fiber composite in an effort to significantly reduce fuel consumption. Despite their extraordinary

intrinsic materials properties, efforts to scale up graphene-based nanocomposites and effectively

utilize them in developing lightweight composites have been limited. This is mainly due to poor

interfacial interactions and hierarchical structures of adjacent graphene and polymer matrix

elements. Furthermore, scientific reports on the fracture toughness, fracture energy, and fatigue

behavior of graphene-polymer nanocomposites are scarce and are highly required by the

automotive and aerospace industries. The adhesion and interfacial shear strength between

graphene of varying compositions and the polymer matrix as a function of functionalization

composition and type of the polymer matrix need to be fundamentally investigated to further

elucidate the effects of structural composition parameters graphene on fatigue and fracture

toughness graphene-based polymer composites. It is also possible to apply the nanocellular

169

structure into the composites. The nano/micro cellular foaming has the potential to further

improve the fatigue endurance of graphene-based polymer nanocomposites.

8.2.6 3D Nanostructured Graphene for Heat Management in

Microelectronic Devices

For decades, continuous miniaturization and high-power densification have been the hallmark of

microelectronic devices. However, overheating is often the greatest barrier in further

miniaturization due to high functional and power density requirements. Thus, efficient heat

dissipation is critical to guarantee optimal performance and extend the service life of

microelectronic devices. To overcome this challenge, thermally conductive porous structures

with significant surface areas can be utilized to substantially increase heat dissipation [36].

Among a wide array of thermally conductive materials, monolayer graphene with exceptional

thermal conductivity (~4000-5000 Wm-1K-1) [36], is a highly promising nanomaterials to

fabricate three-dimensional (3D) porous structures for heat management. Therefore, it is

suggested (1) to develop a novel method to fabricate 3D nanostructured porous graphene films

consisting of a network of hollow struts. The aim is to experimentally investigate the heat

dissipation efficiency of the fabricated 3D nanostructures with different morphologies in

electronic nanostructures and packaging. (2) In parallel with the experimental study,

fundamental physics and theoretical models of nano/micro-scale heat and phonon transport in

3D nanostructures can be investigated. The aim will be to bridge between nano- and micro-scale

heat transport mechanisms. Prospective findings of this study will present valuable opportunities

for future investigations. For instance, the synthesized porous 3D nanostructures can be

impregnated with phase change materials to store and release thermal energy during thermal

cycles.

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8.2.7 Development of Lightweight Superthermal Insulation Graphene-

Based Nanocomposites

Micro- and nano-cellular foams are of great interest to a broad range of industries. This is

because they possess lightweight, high chemical corrosion resistance and low thermal

conductivity. Thus, they have many applications. One of the most viable and promising

properties of micro- and nano-cellular composite foams is their thermal insulation behavior.

Polymeric foams with closed-cell morphologies have the lowest thermal conductivity of any

conventional non-vacuum insulation foams. However, thermal radiation is still significant in

these low-density polymeric foams. The thermal radiation contributes as much as 20% – 40% to

the total thermal conductivity. One promising approach to further decreasing polymeric foams’

thermal conductivity is to use carbonaceous materials as the infrared attenuated agents (IAAs) to

block the radiation. Among the allotropes of noted carbonaceous materials, graphene and CNTs

have received great attention due to their superior electrical and mechanical properties.

Therefore, it is suggested to develop superthermal insulation polymeric foams using physical

foaming technology to fabricate carbon-based polymeric foams with a reduced cell size to

further reduce these composites’ thermal conductivity to below 20 mW/mK. Until now, studies

on this novel application of graphene have been strikingly limited. An in-depth study of the

foam processing of a carbon-based polymeric composite and their nanoscale heat transfer

analysis are essential.

8.2.8 Fabrication of 3D Architected Nanostructures of 2D Materials

Nano-Architected structures merge structural and material properties into a single material.

Fabrication of ultralight and extremely strong materials can be achieved by applying

architecture into material design. The hieratical 3D nano-architectured graphene whose

constituents size ranges from several nanometers to microns to millimeters and centimeters, are

expected to exhibit superior mechanical, thermal and electrical properties and at ultralow

densities (lighter than aerogels). To address this research proposal, it is suggested to design and

create 3D polymeric scaffolds using two different methods including: (i) nano- and micro-

additive manufacturing (i.e. two-photon lithography direct laser writing) which is accurate yet

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hardly scalable; and (ii) scalable nano/microcellular open-cell foaming technique. The

fabricated polymeric scaffolds can be used for synthesizing graphene layers. The main objective

includes: fabrication and characterization of hierarchical 3D nanoarchitected and nanostructured

3D graphene as the next generation of ultralight and immensely robust structural components

for applications in biological and chemical devices, and ultralight energy storage materials.

thermal and electrical properties of graphene …...ultra-low dielectric loss of polymer/carbon...

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