RHEOLOGY MODIFICATION AND FOAMING OF
POLYPROPYLENE – CLAY NANOCOMPOSITES WITH
COUPLING AGENTS
By
Amit Kumar Chaudhary

A DISSERTATION
Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of
DOCTOR OF PHILOSOPHY
Chemical Engineering
2010

ABSTRACT
RHEOLOGY MODIFICATION AND FOAMING OF POLYPROPYLENE – CLAY
NANOCOMPOSITES WITH COUPLING AGENTS
By
Amit Kumar Chaudhary
This research was motivated by the need to develop closed cell thermoplastic foam based
on linear polypropylene to form the core of a wide variety of polypropylene components. This
would require rheology modifiers since linear polypropylene does not have adequate tensile melt
strength to prevent cell wall rupture during foaming or bubble expansion in the melt state. The
objectives of this research were (1) to understand the effects of composition on melt extensional
strain hardening in composites of linear polypropylene with nanoclay, compatibilizer and various
coupling agents and (2) to evaluate selected nanocomposite formulations in extrusion foaming
with chemical blowing agents.
Nanocomposites were prepared with organically modified montmorillonite, maleic
anhydride grafted polypropylene as compatibilizer and linear polypropylene as the bulk matrix.
The organoclays were also treated with amino silane coupling agents to increase the interaction
between nanoclays and the compatibilizer. The effects of different polymeric compatibilizers and
varying organoclay loading on the melt extensional viscosity behavior and the rate of
crystallization of the nanocomposites were examined and shown to be critical for the resulting
foam quality. Besides providing large numbers of nucleation sites for bubble formation during
the foaming process, addition of nanoclay with appropriate compatibilizer provided strain
hardening in melt extensional flow and slower crystallization.
In composites where the organoclay was used without additional treatment, hydrogen
bonding on the surfaces of the organoclay holds the compatibilizer chains. Comparison of

composites with compatibilizers having different chain lengths revealed that the length of
compatibilizing polymer chain that is anchored at one or more clay surfaces must be sufficient to
affect the dynamics of the entanglement network during extensional flow and produce strain
hardening in the melt.
Organically modified nanoclays and the compatibilizer have mainly electrostatic or
hydrogen bonding type interactions, which are weaker forces of attraction. Amino silane
coupling agents were grafted on the clay faces and/or edges to replace these weaker forces of
attraction with stronger chemical bonds. Chemical reactions between the silated clays and
appropriate compatibilizer resulted in nanocomposites with significantly greater strain hardening
compared to the unsilated clay nanocomposites. These nanocomposites were further evaluated
for extrusion foaming using a chemical blowing agent. The nanocomposites exhibiting uniform
strain hardening over a range of strain rates resulted in foams with small cell size and highest cell
densities within the scope of this research.
Since recyclability drives the substitution of thermoset polyurethane foams with closed
cell PP foams, the effect of re-processing on the melt strain hardening behavior was investigated.
Various nanocomposites were annealed and extruded through a uniform cross-section die to
study the effect of heating and shear flow on the structural morphology and strain hardening of
the nanocomposites. The results indicated that annealing had no negative impact on morphology
or strain hardening of the nanocomposites. However, extrusion of these nanocomposites
indicated that nanocomposites prepared using multiple hydrolysable groups resulted in formation
of layered aggregates that are difficult to orient during uniaxial extensional flow and thus
resulted in loss of the strain hardening behavior.

Copyright by
AMIT KUMAR CHAUDHARY
2010

DEDICATED
TO
MY FAMILY
AND
IN THE MEMORY OF MY LATE BROTHER

v

ACKNOWLEDGEMENTS

Achieving success in any endeavor required not only personal efforts but also the support
of coworkers, colleagues and family. I have been fortunate enough to have this support during
the entire course of my graduate studies here at Michigan State University. First and foremost, I
would like to express my deepest gratitude to my academic advisor, Dr K. Jayaraman, for his
advice and constant encouragement. Apart from being encouraging and supportive of my ideas, I
thank you for the endless discussions and guidance to help shape my hypothesis into realities. I
would also like to thank my committee members Dr R. Narayan, Dr A. Lee and Dr L. Matuana
for serving on my committee and providing valuable guidance during the course of my research.
I would also like to express my gratitude to my fellow colleagues, John, Tanmay, and
Katie, who helped me get started with my research during the early stages of my graduate
studies. I would also like to thank Mr. Mike Rich for training me on various equipments and
allowing me to use the facilities at the Composite Materials and Structure Center lab. I would
also like to thank Mr. Brian Rook for helping me with various instruments and equipments
during the course of my research. Further I would like to extend my sincere appreciation to Dr A.
Pastor and Ms E. Danielewicz for training me on various microscopy tools along with sample
preparations.
I would also take this opportunity to thank my parents and grand parents for their
immense faith and unconditional love, which helped me stay miles away from home and earn a
doctoral degree from Michigan State University. I couldn’t have done it with you – Mom and
Dad. Finally I would like to say “Thank you” to my wife Deboleena, for her constant support and

vi

understanding through difficult times. Her love and support turned the tortuous path of graduate
research work into a smooth one.
I want to take this opportunity to thank the chemical engineering staffs, Jennifer
Somerville, Jennifer Peterman, Donna Fernandez, Lauren Brown, and Nichole Shook for all the
help. I cannot thank Ms. JoAnn Peterson enough for her help and support during my initial tough
days here at MSU. Finally, to my friends and lab-mates, thanks for all the laughs and
unforgettable moments. You have all made my stay at Michigan State University a truly
enjoyable experience.

vii

TABLE OF CONTENTS

LIST OF TABLES..........................................................................................................................X
LIST OF FIGURES ...................................................................................................................... XI
1.

INTRODUCTION ...................................................................................................................1
1.1. Overview and Objectives................................................................................................ 1
1.2. Polymer Foams: General Introduction............................................................................ 4
1.3. Polymer Nanocomposites ............................................................................................... 8
1.4. Polymer Nanocomposite Foams ................................................................................... 12
1.5. Polypropylene Foams.................................................................................................... 15
1.6. Research Objectives...................................................................................................... 21

2. POLYPROPYLENE / CLAY NANOCOMPOSITES – PREPARATION AND
CHARACTERIZATION PROCEDURES ....................................................................................25
2.1. Materials ....................................................................................................................... 25
2.2. Nanocomposite Preparation .......................................................................................... 26
2.3. Silane Treatment of Nanoclays..................................................................................... 27
2.4. BET Surface Area Measurements................................................................................. 28
2.5. X-Ray Diffraction ......................................................................................................... 30
2.6. Transmission Electron Microscopy .............................................................................. 31
2.7. Differential Scanning Calorimetry................................................................................ 32
2.8. Thermo-gravimetric Analysis ....................................................................................... 34
2.9. Rheology ....................................................................................................................... 34
2.9.1.
Shear Rheology......................................................................................................35
2.9.2.
Extensional Rheology ............................................................................................37
2.10.
Polymer Foaming Procedure..................................................................................... 39
2.11.
Characterization of Foam Samples ........................................................................... 40
2.11.1. Density Measurement ............................................................................................40
2.11.2. Scanning Electron Microscopy ..............................................................................41
3.

EXTRUSION OF LINEAR POLYPROPYLENE/CLAY NANOCOMPOSITE FOAMS ..46
3.1. Introduction................................................................................................................... 46
3.2. Materials and Experimental Procedures ....................................................................... 49
3.3. Results and Discussion ................................................................................................. 50
3.3.1.
Morphology of Nanocomposites............................................................................50
3.3.2.
Uniaxial Extensional Rheology .............................................................................51
3.3.3.
Crystallization Behavior ........................................................................................55
3.3.4.
Extruded Foam Structure .......................................................................................56
3.4. Conclusion .................................................................................................................... 60

4. THE USE OF SILANE COUPLING AGENTS IN POLYPROPYLENE CLAY
NANOCOMPOSITES ...................................................................................................................90
viii

4.1. Introduction................................................................................................................... 90
4.2. Materials and Experimental Procedures ....................................................................... 93
4.3. Results and Discussion ................................................................................................. 94
4.3.1.
Characterization of Silated Clays...........................................................................94
4.3.1.1. Thermo-gravimetric analysis ............................................................................ 95
4.3.1.2. BET surface area measurements....................................................................... 95
4.3.1.3. X-ray diffraction study...................................................................................... 97
4.3.2.
Characterization of Polypropylene – Silated Clay Nanocomposites .....................98
4.3.2.1. Structural characterization ................................................................................ 98
4.3.2.2. Uniaxial Extensional Rheology ........................................................................ 99
4.3.2.3. Effect of solvent used during silation on final PP/clay nanocomposite property
102
4.4. Conclusions................................................................................................................. 103
5. CONTINUOUS FOAMING OF POLYPROPYLENE CLAY NANOCOMPOSITES
HAVING DIFFERENT STRAIN HARDENING CHARACTERISTICS .................................124
5.1. Introduction................................................................................................................. 124
5.2. Materials and Experimental Procedures ..................................................................... 126
5.3. Results and Discussion ............................................................................................... 127
5.4. Conclusions................................................................................................................. 130
6. EFFECT OF COMPATIBILIZER PROPERTIES ON STRAIN HARDENING
BEHAVIOR OF POLYPROPYLENE CLAY NANOCOMPOSITES.......................................140
6.1. Introduction................................................................................................................. 140
6.2. Materials and Experimental Procedures ..................................................................... 142
6.3. Results and Discussion ............................................................................................... 142
6.3.1.
Uniaxial extensional rheology .............................................................................142
6.4. Conclusions................................................................................................................. 147
7. EFFECT OF MELT PROCESSING ON STRAIN HARDENING BEHAVIOR OF
POLYPROPYLENE CLAY NANOCOMPOSITES - RECYCLABILITY ...............................157
7.1. Introduction................................................................................................................. 157
7.2. Materials and Experimental Procedures ..................................................................... 159
7.2.1.
Thermal Annealing ..............................................................................................159
7.2.2.
Extrusion ..............................................................................................................160
7.3. Results and Discussion ............................................................................................... 160
7.3.1.
Effect of Thermal Annealing ...............................................................................160
7.3.2.
Effect of Extrusion...............................................................................................161
7.4. Conclusions................................................................................................................. 164
8.

CONCLUSIONS AND RECOMMENDATIONS ..............................................................181
8.1. Conclusions................................................................................................................. 181
8.2. Recommendations....................................................................................................... 185

REFERENCES ............................................................................................................................188

ix

LIST OF TABLES

Table 2.1:

Details of various polymers used in this study. .....................................................43

Table 3.1:

Composition of various nanocomposites and polymer blends used in this study..62

Table 3.2:

Crystallization parameters of various nanocomposites and polymer blends
estimated from non-isothermal crystallization study in a DSC. ............................63

Table 3.3:

Comparison of various PP nanocomposites and polymer blends. .........................64

Table 3.4:

Characteristics of various foam extrudates of polymer blends and polymer
nanocomposites foamed in a single screw extruder using 3 wt% of
azodicarbonamide. .................................................................................................65

Table 4.1:

Formulation details of linear PP/ clay nanocomposites with 85 wt% Profax 6523,
12 wt% of PO 1015, and 3 wt% to treated clays. ................................................104

Table 4.2:

BET surface area measurements of silane treated nanoclays. .............................105

Table 4.3:

Basal spacing (d-spacing) of various treated nanoclays and their corresponding
nanocomposites prepared using 85 wt% linear PP and 12 wt% PO 1015. ..........106

Table 6.1:

Composition of linear PP/clay nanocomposites. (*The silated clay used here is
Diamino dimethoxy-silated I.44P, except for PPNC-S15 where it is diamino
monomethoxy-silated I.44P)................................................................................148

x

LIST OF FIGURES

Figure 1.1:

Classification of layered-silicate polymer nanocomposites based on the extent of
dispersion of the nanolayers...................................................................................23

Figure 1.2:

Effect of die processing temperature on evolution of cellular structure of semicrystalline polymer foam [46]................................................................................24

Figure 2.1:

A schematic of Extensional Viscosity Fixture (EVF) used for the measurement of
uniaxial elongational viscosity of polymer melts. .................................................44

Figure 2.2:

Thermal decomposition profile of Cell span–693k, as obtained from thermogravimetric analysis. At 200°C, 15.2 wt% of the chemical blowing agent is
decomposed to produce gas (primarily nitrogen) for foaming various polymers
used in this study....................................................................................................45

Figure 3.1:

Melt extensional viscosity of polymer blend of linear and branched PP mixed at
two different mixing conditions. The blend prepared at low mixing rpm show
greater strain hardening behavior...........................................................................66

Figure 3.2:

X-ray diffraction patterns of various linear polypropylene–clay nanocomposites
along the X-ray diffraction pattern of the nanoclay...............................................67

Figure 3.3:

Melt extensional viscosity of PP, PPNC-N1 and PPNC-N2 under uniaxial
extensional flow. None of the polymer melt show strain hardening behavior under
uniaxial extensional flow. ......................................................................................68

Figure 3.4:

Melt extensional viscosity of PPNC-S1, PPNC-S2 and PPNC-S6 nanocomposites
under uniaxial extensional flow. All polymer melts shows significant strain
hardening phenomena under uniaxial extensional flow.........................................69

Figure 3.5:

Strain hardening parameter of PPNC-S1, PPNC-S2 and PPNC-S6
nanocomposites......................................................................................................70

Figure 3.6:

Melt extensional viscosity of TPO and TPONC under uniaxial extensional flow.71

Figure 3.7:

Non-isothermal crystallization of PPNC-N1 and PPNC-N2 nanocomposites
having same wt% of different PP-g-MA and same organoclay.............................72

Figure 3.8:

Non-isothermal crystallization of PPNC-S1, PPNC-S2, PPNC-S6, PPNC-N2
nanocomposites and L80B20 polymer blend.........................................................73

Figure 3.9:

Scanning electron micrograph of linear polypropylene foam extrudate................74

Figure 3.10: Scanning electron micrograph of PPNC-N2 foam extrudate.................................75
xi

Figure 3.11: Scanning electron micrograph of PPNC-N1 foam extrudate.................................76
Figure 3.12: Scanning electron micrograph of PPNC-S1 foam extrudate. ................................77
Figure 3.13: Scanning electron micrograph of PPNC-S2 foam extrudate. ................................78
Figure 3.14: Scanning electron micrograph of PPNC-S6 foam extrudate. ................................79
Figure 3.15: Scanning electron micrograph of L80B20 foam extrudate....................................80
Figure 3.16: Scanning electron micrograph of TPO foam extrudate. ........................................81
Figure 3.17: Scanning electron micrograph of TPONC foam extrudate....................................82
Figure 3.18: Cell size distribution of PPNC-N2 foam. ..............................................................83
Figure 3.19: Cell size distribution of PPNC-N1 foam. ..............................................................84
Figure 3.20: Cell size distribution of PPNC-S1 foam. ...............................................................85
Figure 3.21: Cell size distribution of PPNC-S2 foam. ...............................................................86
Figure 3.22: Cell size distribution of PPNC-S6 foam. ...............................................................87
Figure 3.23: Cell size distribution of L80B20 foam...................................................................88
Figure 3.24: Cell size distribution of TPONC foam...................................................................89
Figure 4.1:

General structure of a silane coupling agent........................................................107

Figure 4.2:

Reaction scheme between silane treated clays and maleic anhydride grafted
polypropylene. .....................................................................................................108

Figure 4.3:

Chemical structure of silane coupling agents (a) Diamino monomethoxy-silane,
(b) diamino dimethoxy-silane, (c) diamino trimethoxy-silane, and (d) Phenyl
trimethoxy silane. The ring opening hydrolysis of diamino monomethoxy-silane is
also shown along with its chemical structure. .....................................................109

Figure 4.4:

TGA profile of various silane treated I.44P clays................................................110

Figure 4.5:

TGA profile of silane treated I30P clays. ............................................................111

Figure 4.6:

BET surface area data from He et al. [189] on reduction in surface area due to
surfactant adsorption............................................................................................112

Figure 4.7:

X-ray diffraction pattern of silane treated I.44P clays (top) and silane treated I.30P
clays (bottom). The treatment of clays is represented as (a): Diamino
monomethoxy-silated clay, (b) Diamino dimethoxy-silated clay, (c) diamino

xii

trimethoxy-silated clay, and (d) Phenyl trimethoxy silated clay. The d-spacing of
the corresponding organoclays are presented alongside the XRD curve.............113
Figure 4.8:

X-ray diffraction pattern of linear PP nanocomposites with silane treated I.44P
clays (top) and silane treated I.30P clays (bottom). The treatment of clays is
represented as (a): Diamino monomethoxy-silated clay, (b) Diamino dimethoxysilated clay, (c) diamino trimethoxy-silated clay, and (d) Phenyl trimethoxy
silated clay. The d-spacing of the corresponding organoclays are presented
alongside the XRD curve. ....................................................................................114

Figure 4.9:

TEM micrographs of linear PP nanocomposites with untreated I.44P clay (PPNCS6) (top) and with diamino dimethoxy silane treated I.44P clay (PPNC-S7)
(bottom)................................................................................................................115

Figure 4.10: TEM micrographs of linear PP nanocomposites with untreated I.30P clay (PPNCS11) (top) and with diamino dimethoxy silane treated I.30P clay (PPNC-S4)
(bottom)................................................................................................................116
Figure 4.11: A comparison of melt extensional viscosity of PPNC-S10 (phenyl trimethoxy
silated I.30P) and PPNC-S11 (unsilated I.30P) nanocomposites under uniaxial
extensional flow. ..................................................................................................117
Figure 4.12: Effect of amino silane treatment of I.30P clay on the melt extensional viscosity of
linear PP nanocomposites under uniaxial extensional flow. The number of
hydrolysable groups has negligible effect on the strain hardening behavior of the
nanocomposites....................................................................................................118
Figure 4.13: Effect of amino silane treatment of I.44P clay on the melt extensional viscosity of
linear PP nanocomposites under uniaxial extensional flow.................................119
Figure 4.14: Dynamic time sweep test at 180°C for various PP nanocomposites along with its
matrix (PP+compatibilizer)..................................................................................120
Figure 4.15: Variation of strain hardening parameter with strain rate for PPNCs prepared with
silated I.44P clays. ...............................................................................................121
Figure 4.16: Variation of strain hardening parameter with strain rate for PPNCs prepared with
silated I.30P clays. Open symbols are for Hencky strain = 2.9 and closed symbols
are for Hencky strain = 2.25. ...............................................................................122
Figure 4.17: Effect of solvent used for silation process on melt extensional viscosity of PPNCS7 nanocomposites...............................................................................................123
Figure 5.1:

Non-isothermal crystallization exotherm for linear PP/clay nanocomposites
prepared using 85 wt% linear PP, 12 wt% PO 1015 and 3 wt% of different treated
organoclays. .........................................................................................................131

xiii

Figure 5.2:

Evolution of cell size as a function of time demonstrated by various theoretical
and experimental work [196-199]. (θr represents the residence time in the die) .132

Figure 5.3:

Scanning electron micrograph of PPNC-S6 (prepared using 3 wt% of unsilated
I.44P) foam extrudate...........................................................................................133

Figure 5.4:

Scanning electron micrograph of PPNC-S7 (prepared using 3 wt% of diamino
dimethoxy silane treated I.44P organoclay) foam extrudate. ..............................134

Figure 5.5:

Scanning electron micrograph of PPNC-S4 (prepared using 3 wt% of diamino
dimethoxy silane treated I.30P organoclay) foam extrudate. ..............................135

Figure 5.6:

Cell size distribution of PPNC-S6 foam. .............................................................136

Figure 5.7:

Cell size distribution of PPNC-S7 foam. .............................................................137

Figure 5.8:

Cell size distribution of PPNC-S4 foam. .............................................................138

Figure 5.9:

Progressive decrease in cell size and increase in cell density as a function of
material property of linear PP/clay nanocomposite. ............................................139

Figure 6.1:

Melt extensional viscosity comparison for various linear PP/clay nanocomposites,
showing the effect of molecular weight of the compatibilizer on the strain
hardening characteristics of the derived nanocomposites....................................149

Figure 6.2:

Melt extensional viscosity comparison for two nanocomposites under uniaxial
extensional flow, showing the effect of molecular structure of the compatibilizer
on the strain hardening characteristics of the derived nanocomposites (PPNC-S7
is prepared using PO 1015 and PPNC-S14 is prepared using PO 1020). ............150

Figure 6.3:

Melt extensional viscosity comparison for two nanocomposites under uniaxial
extensional flow, showing the effect of molecular structure of the compatibilizer
on the strain hardening characteristics for nanocomposites having Bondyram 1001
(PPNC-S17) and Bondyram 1001CN (PPNC-S16) as the compatibilizer...........151

Figure 6.4:

Melt extensional viscosity data under uniaxial extensional flow, for
nanocomposites having Bondyram 1001CN as the compatibilizer. PPNC-N3 is
prepared from Diamino dimethoxy silane treated I.44P and has a compatibilizer to
clay ratio of 4 and PPNC-S15 is prepared from Diamino monomethoxy silane
treated I.44P and has a compatibilizer to clay ratio of 7......................................152

Figure 6.5:

Molecular structure of PPgMAs. (a) AC-950, (b) PO 1020, and (c) PO 1015....153

Figure 6.6:

TEM micrographs of a stretched section of a representative nanocomposite
(PPNC-S4) sample. The nanoclay platelets are oriented along the direction of
stretch (shown by the arrow). The cartoon shows the section of the stretched
samples used for TEM. ........................................................................................154

xiv

Figure 6.7:

Variation of storage modulus with frequency for various linear PP-clay
nanocomposites and the matrix polymer, prepared using Bondyram 1001CN as
the compatibilizer.................................................................................................155

Figure 6.8:

Van-Gurp Palmen plot for various linear PP-clay nanocomposites prepared using
Bondyram 1001CN and Bondyram 1001 as the compatibilizer along with the
matrix (linear PP + Bondyram 1001CN). ............................................................156

Figure 7.1:

X-ray diffraction pattern showing the effect of thermal annealing and extrusion
for PPNC-S6 nanocomposites. ............................................................................166

Figure 7.2:

X-ray diffraction pattern showing the effect of thermal annealing and extrusion
for PPNC-S7 nanocomposites. ............................................................................167

Figure 7.3:

X-ray diffraction pattern showing the effect of thermal annealing and extrusion
for PPNC-S3 nanocomposites. ............................................................................168

Figure 7.4:

X-ray diffraction pattern showing the effect of thermal annealing and extrusion
for PPNC-S4 nanocomposites. ............................................................................169

Figure 7.5:

Melt extensional viscosity of PPNC-S6 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow. .....................................................................................................................170

Figure 7.6:

Melt extensional viscosity of PPNC-S7 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow. .....................................................................................................................171

Figure 7.7:

Melt extensional viscosity of PPNC-S3 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow. .....................................................................................................................172

Figure 7.8:

Melt extensional viscosity of PPNC-S4 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow. .....................................................................................................................173

Figure 7.9:

TEM micrographs of linear PP/diamino dimethoxy silated I.44P (PPNC-S7)
nanocomposites showing the effect of extrusion on nanoclay dispersion (top:
before extrusion, bottom: after extrusion). ..........................................................174

Figure 7.10: TEM micrographs of linear PP/diamino monomethoxy silated I.30P (PPNC-S3)
nanocomposites showing the effect of extrusion on nanoclays dispersion (top:
before extrusion, bottom: after extrusion). ..........................................................175
Figure 7.11: TEM micrographs of linear PP/diamino dimethoxy silated I.30P (PPNC-S4)
nanocomposites showing the effect of extrusion on nanoclays dispersion (top:
before extrusion, bottom: after extrusion). ..........................................................176

xv

Figure 7.12: Schematic representing state of nanolayers dispersion/agglomeration after mixing
and after extrusion through a capillary die. .........................................................177
Figure 7.13: Schematic representing agglomeration of diamino dimethoxy silated I.30P layers
in PPNC-S4 due to extrusion through a capillary die. .........................................178
Figure 7.14: Effect of extrusion on storage modulus (G’) of PPNC-S3 and PPNC-S4
nanocomposites. (open symbols: before extrusion and closed symbol: after
extrusion) .............................................................................................................179
Figure 7.15: Effect of extrusion on complex viscosity of PPNC-S3 and PPNC-S4
nanocomposites. (open symbols: before extrusion and closed symbol: after
extrusion) .............................................................................................................180

xvi

1. INTRODUCTION
1.1. Overview and Objectives
During the last few decades, polymeric materials have evolved into a global industry and
have gained ever-increasing acceptance as an integral part of modern technology by steadily
replacing other materials in a wide variety of applications. In order to satisfy the growing needs
of the current market, nearly few 100 million tons of polymeric materials are produced
worldwide, as compared to 5 million tons produced around six decades ago.
Ease of processing, light weight, economy and versatility makes these polymeric
materials a preferred and competing commodity. The weight of these polymeric materials can
further be reduced by foaming them with suitable blowing agents. Polymer foams, also referred
to as sponge, expanded or cellular plastics are a class of polymeric materials that are extensively
used as a common consumer product due to their economic and performance advantages, starting
from styrofoam coffee cups to polyurethane foam seat cushions in automobiles. Polymer foams
were first introduced in early 1940s and since then its use in modern technology continues to
grow at a rapid pace across the globe due to their light weight, excellent strength/weight ratio,
superior insulating and energy absorbing ability, buoyancy, comfort features and many more,
most of which are absent in the unfoamed resin. The polymer foam industry is expected to grow
steadily due to the unique advantages offered by these materials compared to traditional
materials. In U.S. alone, the polymer foam consumption for the year 2010 is estimated to be
around 7.5 billion pounds and is expected to reach around 8.6 billion pounds by the year 2015 at
an annual growth rate of around 2.5% for all polymeric foams produced from commercial
polymers.

1

With the growing need for energy and cost efficient systems primarily in automotive,
aerospace and construction industries, polymer foams have evolved as another promising area of
intensive research. With the advancement of foaming technology over the past few decades, the
foaming technology has evolved from being purely empirical to more fundamental, relying on
science and technology. Foams can be prepared from any polymer, by the introduction of a
secondary gaseous phase within the polymer matrix. Different polymers are selected for different
industrial foam applications depending upon their properties, their ease of manufacture, and the
economics of the foaming system. This has led to foaming of nearly all types of polymeric resin
– thermoset [1-8], thermoplastic [9] and even biopolymers [10-18] and has thereby led to the
expansion of polymer foam application from automotive, construction, packaging, furnishing,
transportation, sports industries all the way up to its use in tissue engineering [19-21] and drug
delivery system [21-23]. Various polymers have been used for foaming applications, for
example, polyurethane (PU) [1-8], polystyrene (PS) [24-28], poly-vinyl chloride (PVC) [29-33],
polycarbonate (PC) [34-36], polyolefins (polyethylene (PE), polypropylene (PP)) [37-49], poly
(ethylene terephthalate) (PET) [50], poly (lactic acid) (PLA) [51, 52]. Amongst all polymer
foams, polyurethane (PU) foam continues to be the largest product segment in the global
polymeric foams market followed by polystyrene (PS) foam. Even though PU foams cannot be
recycled, thereby causing a strong impact on environmental issues, its market is expected to
grow at a rate of 2.6% annually. This is due to the fact that PU leads to the production of closed
cell foams with a wide range of densities and hardness, forming both rigid and flexible foams
depending on the final application. Also due to the high impact resistance properties of PU
foams, it is widely used in automotive industries for its use in door and instrument panels along
with cushioning of seats. The production of such automotive interiors involves multistep

2

fabrication and thus produces as lot of non-recyclable scrap. This demands a sooner replacement
of this material with more environmental friendly and recyclable polymer foam.
Thermoplastic polyolefins (TPO) has all the required properties necessary to replace PU
foams from automotive industry [53, 54]. TPO is a two phase blend of semi-crystalline
polypropylene and ethylene-octene elastomer. This material is widely used in the automotive
industry for the production of various automobile exterior and interior parts such as bumper
fascia, claddings, wire insulations etc. because of its weatherability, low density, low cost, high
impact strength and scratch resistance [53-56]. However, utilization of TPO foam has not been
implemented successfully due to various technological challenges involved in its production and
has not yet been fully developed. Nonetheless, there lies a significant amount of interest in
developing foamed TPO products due to its attractive economic and environmental properties.
Polypropylene is the major component of the TPOs and has various attractive properties, for
example, it can be used at higher service temperatures, has good stiffness, excellent chemical and
abrasion resistance and above all has low material cost and is easily processable. However, PP
has poor impact properties. Even though polyethylene has better impact properties than PP, PE
has failed to satisfy the market demands for high performance and high heat resistance foamed
polymers. Even though rheologically PE is more favorable for foam production than PP, PP
foams are preferred for applications where stiffness, chemical resistance, good heat and sound
insulation, and higher end use temperatures is required. The addition of elastomers to the PP
matrix improves the impact strength of the resulting material by improving its ductility and crack
resistance [55-57], thereby making it the most desirable polymer for foaming applications.
However, foamability of linear polypropylene is also largely restricted by its rheological [47, 58,
59] and thermal properties [60, 61]. This research primarily focuses on understanding

3

foamability and modifying the rheological properties of linear polypropylene by the use of
organically modified montmorillonite nanoclays that can lead to successful production of
polypropylene foam, thereby leading to TPO foams for various interior and exterior applications
in the automotive industry.

1.2. Polymer Foams: General Introduction
Thermoplastic foams are cellular plastics having dispersed, usually spherical or
hexagonal, voids in a continuous thermoplastic polymer matrix. These materials possess unique
physical, mechanical and thermal properties, usually governed by the matrix, the foam cellular
structure (for example, cell size, cell density and bulk density).
Thermoplastic polymer foams represent a class of light weight materials that find
immense potential in a variety of applications, such as, transportation, insulation, packaging,
sealing materials, sports, impact and sound absorption and many more [62]. Thermoplastic
foams have been classified by cell size, expansion ratio and cell wall integrity. These
classifications are independent of processing technology and the polymer used, and they reflect
different applications and different physical properties of the foams. Polymer foams with an
average cell size above 100 µm are classified as conventional (coarse) foams, between 10 and
100 µm as fine-celled foams, and foams below 10 µm are known as microcellular foam.
Microcellular foams were first introduced by Suh et al. [63] in 1980s. Microcellular foams have
large number of small cells, which reduces the stress concentration around the cells due to
reduction in cell size. Compared to conventional and fine-celled foams, microcellular foams are
known to exhibit superior impact strength, high toughness and stiffness to weight ratio, high

4

thermal stability, low thermal conductivity, high fatigue life, and reduced material weight and
overall cost.
Foams are also classified as either rigid or flexible depending on whether their glass
transition temperature is below or above ambient temperature and on their cell morphology.
Rigid foams are selected for applications where high mechanical strength is required, such as in
building and construction, transportation, packaging, furniture, decorative moldings, food and
drink containers, and floatation devices. Flexible or semi-flexible foams are used in gasketing,
sports applications, shock and sound attenuation, shoes, furniture cushioning, textile and primary
seals in automobiles.
Polymer foams are also classified as closed cell or open cell foams depending upon the
cell structure. Closed cell foams have cell cavities that are closed and isolated from each other
surrounded by cell walls. This leads to foams with better insulation properties and higher
mechanical strength. Open cell foams have cell cavities that are interconnected and the cell walls
are broken and not continuous. The open cell foams are superior materials in terms of absorptive
capability.
The foaming of polymeric materials can be achieved by various methods, such as use of
blowing agents, mechanical whipping of air and dispersion of hollow microspheres [62]. The
most widely used process is the expansion of molten polymer containing a blowing or foaming
agent. Blowing agents used in the foaming processes can be classified into two categories –
physical blowing agents (PBA) and chemical blowing agents (CBA) [64]. Chemical blowing
agents are chemicals that function by releasing a gas due to heat-induced decomposition above a
critical decomposition temperature. The gas thus generated (commonly nitrogen or carbon
dioxide) is dissolved in the hot melt under high pressure, which later nucleates and grows into

5

bubbles as the pressure is reduced. These can either be endothermic or exothermic, or a
combination of both. Proper selection of the CBA allows a good temperature control in the
foaming process. Most widely used CBAs are azodicarbonamide, benzene sulfonylhydrazide,
and sodium bicarbonate. PBA are substances that form gaseous products under the foaming
conditions (high temperature and/or reduced pressure). In this case a homogeneous mixture of
polymer melt and physical blowing agent is first obtained by directly metering the physical
blowing agent into the polymer melt at high pressure. Subsequent increase in temperature or
reduction in pressure results in bubble nucleation and growth. Commonly used physical blowing
agents are volatile liquids such as, hydrocarbons, alcohols, inert gases such as nitrogen, and
carbon dioxide. Amongst all, CO2 is the most favorable PBA, as it is inexpensive, nontoxic, and
environmentally benign. However, due to limitation in solubility in the polymer melt and the
high rate of diffusion out of the melt, significantly limits the development of CO2 as an industrial
scale foaming agent.
Polymer foams can be manufactured by a variety of continuous to batch processes, for
e.g., continuous extrusion, injection molding, rotational molding, compression molding, and
lamination of foams, or foam board production [62]. Irrespective of the type and nature of
blowing agent, the polymer matrix and the foaming technology, in general, the foaming process
involves three fundamental steps:


Bubble / cell nucleation



Cell growth



Cell stabilization

Nucleation is a classical phenomenon that exists in processes like condensation and
crystallization. Colton et al. [65] presented a theoretical model for cell nucleation in amorphous
6

thermoplastic microcellular foams based on the classical nucleation theory. The process of cell
nucleation within a polymer melt takes place when the polymer melt gets supersaturated with the
blowing agent and a sudden pressure drop results in phase separation. If the nucleant is generated
from a single homogeneous phase of dissolved gas in a polymer melt without impurity or dirt,
the process is called homogeneous nucleation. However, if cells are formed at solid/liquid
interface, e.g., at the surface of foreign particles, the process is called heterogeneous nucleation.
Fletcher [66], in his study presented the effect of nucleant size on heterogeneous cell nucleation.
In heterogeneous nucleation, high nucleation efficiency can only be achieved when nucleation is
energetically favored and the nucleant is well dispersed in the polymer matrix. According to the
theory developed [65, 66], heterogeneous nucleation rate is primarily a function of particle
concentration and dimensions, nature of polymer-gas-particle surface interactions, and pressure
of the system. Qualitatively, it was inferred that a small contact angle and a large surface
curvature results in an increase in heterogeneous nucleation rate.
Once the gas bubbles are nucleated, they will continue to grow by diffusion of gas from
polymer melt phase into the bubble, until eventually they are stabilized or ruptured. The growth
process is very complicated due to several factors governing the polymer melts rheological
response to the deformation introduced by the expansion of the gaseous phase. With the
continuation in cell growth, the probability of cells coming in contact with each other increases,
causing neighboring cells to share a common cell wall. This decreases the cell wall thickness, as
a result of which they become highly stretched and unstable. This could eventually lead to cell
coalescence because of which the average cell size increases and cell density decreases. One of
the important factors in controlling cell growth and thereby reducing cell coalescence is the
stability of the cell walls, which depends upon the tensile melt strength or melt strain hardening

7

of the surrounding polymer matrix. Strain hardening or tensile melt strength of a polymer melt is
defined as the degree of resistance to extensional flow. Therefore, polymers with higher tensile
melt strength are expected to increase the cell wall stability and thereby prevent cell coalescence
leading to polymer foams with higher cell density, lower cell size and narrow cell size
distribution. Higher melt viscosity would also reduce cell coalescence by slowing down the
drainage of polymer from the cell wall during the foam solidification period.
In order to obtain a controlled and uniform cellular structure of the polymer foam,
heterogeneous nucleating agents such as, talc, silicon oxide, titanium oxide, and zinc oxide are
commonly added to the polymer matrix. Work done by Baldwin et al. [67] indicates that
production of microcellular foams using an extrusion system requires stringent processing
conditions – high pressure drop or pressure drop rate and processing temperature close to the
glass transition temperature (for amorphous polymers), resulting in a narrow operating window.
In order to ameliorate these strict processing conditions, work was done to incorporate nucleating
agents in the polymer matrix to produce microcellular foams. Work done by Park et al. [68]
using talc (microfillers) as the heterogeneous nucleating agent in the polymer matrix, showed an
increase in cell density and better control over cell size distribution compared to microcellular
polymer foam without any nucleating agents. This encouraged researchers to explore the use of
nanoparticles as heterogeneous foaming nucleating agent in polymer matrix for the production of
microcellular foams.

1.3. Polymer Nanocomposites
Traditionally, polymeric materials have been filled with various natural or synthetic
micron sized inorganic fillers forming polymer composites, thereby improving their material
8

properties and reducing cost. Commonly used micro fillers are in the form particles (calcium
carbonate), fibers (glass or cellulose fibers), or plate-like (glass flake, mica, talc). Even though
these materials impart significant improvement in various material properties, they are often
associated with few drawbacks, for example, increase in density, brittleness and opacity [69-74].
Polymer nanocomposites, on the other hand, are a different class of composite in which the
fillers incorporated have at least one dimension in the nanometer scale.
Polymer nanocomposite is a field of research that has been studied in great detail over the
past two decades. Interest in polymer nanocomposites has been results of improvements in their
various physical and mechanical properties due to the addition of a small amount of anisotropic
nanoparticles in the polymer matrix. These small additions are effective at low concentrations
due to the large surface area to volume ratio of the nanofillers. This surface to volume ratio can
significantly alter the macroscopic properties of polymers, imparting high thermal and
dimensional stability, enhanced mechanical properties, high heat deflection temperature,
improved flame retardance, reduced gas permeability and increased barrier properties without a
significant increase in density and opacity [73, 74]. Polymer nanocomposite was first introduced
by Toyota Motor Co. by developing nylon–6/montmorillonite nanocomposites [75, 76] in late
1980s. In the time since their introduction, this field has expanded to work on many different
polymer systems such as polyolefins, polyamide, polystyrene, poly(methyl methacrylate),
poly(ethylene terephthalate), polycarbonate, polyvinyl alcohol, polyacrylic acid, and
polyethylene oxide [73, 74]. The derived nanocomposites have been extended for various
applications such as solar calls, electronic devices, and structural applications. Commonly used
nanoparticles used in polymer nanocomposites are layered silicates (nanoclay), fibers (carbon

9

nano fibers, and rod-shaped (carbon nano tubes). Amongst all the nanofillers, nanoclays are most
widely studied probably because of its availability and cost.
Polymer/layered silicate (PLS) nanocomposites are synthesized by dispersing nano
silicates in the polymer matrix either by in situ polymerization, solution blending or melt mixing.
Amongst these, melt mixing is the most economical process and is widely used for commercial
purposes. Nanoclays are high aspect ratio particle (100 – 300), that provide considerable
enhancement in various material properties at low weight fractions (0.5 – 5 wt%) compared to
traditional fillers (glass, mica, talc), which are incorporated at a higher weight fractions (10 – 40
wt%). This result in significant increase in density of traditional composites compared to
nanocomposites. Most common nanoclays used in the synthesis of these PLS nanocomposites is
montmorillonite that are obtained naturally in the form of tactoids comprised of several layers.
These nanolayers are comprised of aluminum-oxygen-hydroxyl octahedral layers sandwiched
between two silicon-oxygen tetrahedral sheets and are generally called 2:1 smectite clays. These
nanolayers are held together by Van-der Walls forces of attraction and are thus difficult to
separate from each other. Since most engineering polymers are hydrophobic in nature, dispersion
of these tactoids into discrete nanolayers is further hindered by the hydrophilic nature of these
layered silicates. The extent of dispersion of these clay platelets defines whether the final
nanocomposite is intercalated or exfoliated (Figure 1.1). In intercalated structure, the polymer
chains penetrate between layers of clay galleries and increase the inter-platelet spacing but retain
a well-ordered multilayer morphology of alternating polymer and clay layers. In case of
exfoliated nanocomposites, the nanoplatelets are completely separated from one another and are
well dispersed within polymer matrix. There is no repeating pattern of polymer and clay layers in
exfoliated nanocomposites as seen in intercalated composites. Due to better dispersion of clay

10

platelets, exfoliated nanocomposites will have higher effective particle concentration at the same
weight percentage of nanoclay loading.
In order to achieve better dispersion of these nanolayers, the naturally occurring
montmorillonite has to be organically modified to synthesize polymer compatible nanoclay. In
order to facilitate the interaction between the hydrophilic clay layers and hydrophobic polymers,
the galleries of the clay tactoids are intercalated with long chain alkyl ammonium surfactants
[77-81]. The ammonium ion is attached to the clay face and the dangling hydrocarbon tails
swells the clay galleries to facilitate polymer chain intercalation and thus improves dispersion of
these nanolayers in the polymer matrix. The functionality present in the polymer chain along
with its extent of hydrophobicity defines the selection criteria of the surfactant to be intercalated
in the naturally occurring montmorillonite. Reichert et al. [82] studied the effectiveness of alkyl
chain length of the surfactant towards dispersion of the organically modified nanosilicates in the
polypropylene matrix and reported that the extent of dispersion of organosilicates improves with
the increase in alkyl chain length of the surfactant. Manias et al. [83] and Ton That et al [84]
reported that multiple alkyl tails of the surfactant results in better dispersion of nanoclay in the
polypropylene matrix. However, the selection criteria are different for polymers having
functional groups (nylon, polystyrene) as opposed to polyolefins [85].
Even with the organic functionalization of montmorillonites, the intercalation of
hydrophobic polyolefins in the organically modified montmorillonites is difficult and is
generally facilitated by the use of less hydrophobic compatibilizers. These compatibilizers are, in
general, polyolefins functionalized by grafting hydroxyl groups or more commonly maleic
anhydride groups at the ends of the polyolefin chains [78, 86, 87]. Okada et al. [88] were the first
to demonstrate the use of polypropylene-grafted maleic anhydride (PP-g-MA) as a

11

compatibilizer to synthesize PP-clay nanocomposites. The importance of such polar
compatibilizers for the use in non-polar clay nanocomposite preparation is well demonstrated by
various researchers [86, 89-94].

1.4. Polymer Nanocomposite Foams
In recent years, the integration of the polymer nanocomposite technology with the
traditional polymer foaming process, has led to synthesis of polymer nanocomposite foams with
greatly improved characteristics in terms of reduced average cell size and increased average cell
density. The presence of nanoparticles provides additional surface to facilitate heterogeneous
nucleation of the large number of cells during the foaming process. The heterogeneous
nucleation effect of nanoparticles has been reported for various polymers, for example,
polystyrene, nylon, poly (lactic acid), polycarbonate, polyurethane, poly (vinyl chloride), poly
(ethylene terephthalate), and polypropylene.
The use of nanoclay in foaming polystyrene to obtain closed cell foams with narrower
cell size distributions and high cell densities was demonstrated by Lee and coworkers [9, 26, 28,
95-99]. Han et al. [28] foamed polystyrene (PS) / clay nanocomposites in a continuous extrusion
foaming process using carbon dioxide as foaming agent and studied the influence of
nanoparticles on cell size, cell density and mechanical properties of the foam. They observed an
eight fold increase in cell density and the average cell size decreased by approximately 40%
upon addition of 7.5 wt% organically modified nanoclay. This is attributed to the much higher
surface area made available for nucleation of bubbles by well dispersed nanolayers.
Shen et al. [99] presented a direct impact of aspect ratio of nanoparticles on cell size and
cell density of PS / carbon nanofibers (CNFs) nanocomposite foam. They studied the foaming of
12

PS / CNFs nanocomposites via continuous extrusion process (melt compounding) using carbon
dioxide as foaming agent. In order to overcome Van der Waals forces of attraction amongst the
CNFs and to improve mixing and dispersion of CNFs uniformly in the polymer matrix, higher
level of shear was applied during the extrusion process. The intensive shear exerted by the twin
screw extruder resulted in significant break-up of CNFs (average CNF length reduced from ~60
µm to ~0.28 µm). In order to study the effect of change in aspect ratio of the CNFs, they made
PS / CNFs nanocomposites via solution blending (minimum shear on the CNFs) with the help of
ultrasonication, instead of high shear mixing, to achieve complete dispersion of CNFs in the
polymer matrix. Solution blending resulted in nanocomposites with average fiber length of
several dozens of microns, indicating minimum damage to the CNFs compared to that in the
extrusion process. PS / CNFs composite made via these two processes (melt compounding and
solution blending) were subsequently foamed in a batch process. The average cell density of the
8

3

solution blended sample (5.7x10 cells/cm ) was found to be higher than that of the melt
8

3

compounded one (1.85x10 cells/cm ) and the average cell size was also smaller (10.2 µm for
solution blended sample and 14.9 µm for melt compounded one).
Work done by Zeng et al. [27, 95] also looked into the importance of aspect ratio of
nanoparticles along with its surface properties in foaming of PS nanocomposites. Two types of
nanoparticles were selected for preparing PS nanocomposites – plate-like nanoparticles
(nanoclay) [95], nanofibers (carbon nanofibres (CNFs) and single wall carbon nanotubes
(SWCNTs)) [27]. The synthesized nanocomposites were subsequently foamed in a batch process
using carbon dioxide as the foaming agent. During the synthesis of PS / SWCNTs
nanocomposites, it was found that the dispersion of SWCNTs in the PS matrix was poor and
resulted in large amount of ball shaped aggregates at higher levels. However, the results indicate
13

that even with poor dispersion of SWCNTs in the PS matrix, the resultant nanocomposite foam
with 0.1wt% SWCNTs displays a much higher cell density and a smaller cell size compared to
exfoliated PS / clay nanocomposites with 5wt% nanoclay. PS / CNFs nanocomposites with well
dispersed CNFs (1wt%) also resulted in foams with higher cell density and a smaller cell size
compared to the exfoliated PS / clay nanocomposites. However all three nanocomposites
significantly increased the cell density and reduced the cell size compared to pure PS foam. One
important factor leading to lower cell density and higher cell size of PS / clay nanocomposite
foams compared to CNFs and SWCNTs system is lower aspect ratio. However, in order to
explain the comparatively poor performance of exfoliated PS / clay nanocomposites towards
heterogeneous nucleation even at such higher clay levels (5 wt%), Zeng et al. found that
polymer-gas-particle contact angle (θ) was 105.5°compared to 20° for PS /CNFs system. This
indicated that poor wetting property of nanoclays is responsible for lower cell density and higher
cell size compared to that of CNFs system. These results indicate that low particle loading of
high aspect ratio nanoparticles having favorable wettability (low contact angle) serve well as
heterogeneous nucleating agents during the foaming process for producing microcellular foams
under mild processing conditions.
Apart from increasing the nucleation efficiency, high aspect ratio of nanoparticles offers
the potential for cell wall reinforcement. Since the thickness of foam cell walls is in the micron
and sub-micron regime [100], the nanoscale dimension of the nanoparticles make them desirable
elements to reinforce the cell walls, resulting in foams with enhanced mechanical properties [99].
The success behind microcellular foams can be attributed to the mechanism of cell structure
evolution, which can be due to increased cell nucleation, controlled cell growth, reduced cell
coalescence, and enhanced cell stability. Any one of the above mentioned factors or a

14

combination of these factors can result into foams with superior physical and mechanical
properties.

1.5. Polypropylene Foams
Linear polypropylene, in today’s market, is one of the most widely used polymeric
materials due to its outstanding combination of cost performance, excellent physical properties,
strong and continuous expansion of process versatility, and environmental friendly processes and
materials during manufacturing, use, and recycling stages. Due to its outstanding functional
characteristics and low material cost, polypropylene foams have been considered as a substitute
for other thermoplastic foams in industrial applications. With the growing attention towards total
recyclability and to take advantage of the benefits of polypropylene, these materials are
considered to be the best replacement for thermoset (polyurethane) foams. Even though
polypropylene has higher rigidity compared to other polyolefins and it offers better impact
strength than polystyrene along with a higher service temperature range and good temperature
stability, continuous foaming of polypropylene foams is largely restricted by the semi-crystalline
nature of the polymer [60, 61, 101] and its rheological properties [47, 58, 59].
The evolution and control of the cellular structure in semi-crystalline polymer foams is
quite different from that of amorphous polymer foams. The semi-crystalline nature of the
polymers makes the foaming process more complicated because of simultaneous nucleation and
growth of gas bubbles and crystals. Reignier et al. [61] have demonstrated with batch foaming of
poly-(e-caprolactone) that crystallization affects both bubble nucleation and growth. Park and
coworkers [46, 60, 102] have investigated foaming of polypropylene with supercritical CO2.
Work done by Doroudiani et al. [60] indicated that the cellular structure of semi-crystalline
15

polymer foams is difficult to control because of the fact that the gas would not be able to diffuse
into the crystallites leading to non-uniform bubble nucleation. Hence, early onset of
crystallization, before the gas bubbles are nucleated, could lead to reduction in bubble nucleation
efficiency resulting in non-uniform foaming with regions of sparse and dense bubbles. Naguib et
al. [46] reported that both early and delayed onset of crystallization is detrimental to foam
cellular structure (Figure 1.2). In their work done on foaming of branched PP using N-butane as
a blowing agent in a tandem foam extrusion system, they report an optimum melt temperature for
maximum expansion ratio of the foam samples. They observed a decrease in foam extrudate
diameter as a function of the distance from the die when the melt is at a higher temperature, due
to delayed onset of crystallization causing the gas to diffuse out of the sample. On the contrary, a
lower melt temperature resulted in early onset of crystallization leading to insufficient foaming.
Park et al. [45, 103, 104] have reported effects of other processing parameters – melt temperature
and pressure, on continuous foaming of polymers with a physical blowing agent.
The most important rheological property for producing superior quality foam is the
ability of the polymer melt to show strain hardening under uniaxial extensional flow. Linear PP
does not have adequate tensile melt strength or strain hardening under uniaxial extensional flow,
a necessary property for producing superior quality foam, particularly in crystallizing polymers.
Strain hardening, is defined as the transient extensional viscosity rise above the linear
viscoelastic curve, at a constant strain rate. Polymers that do not demonstrate strain hardening
under uniaxial extensional flow, develop thinner cell walls during bubble expansion in the
foaming process, which results in coalescence of bubbles and even open-cell foams [105, 106].
An increase in extensional viscosity during the expansion process provides stability to the cell
walls and has been shown to be important for producing good quality foams [58, 100]. Ghijsels

16

et al. [107] studied the temperature dependency of tensile melt strength for polyolefins and
showed that the tensile melt strength of polypropylene decreases with increase in temperature.
As mentioned earlier, melt strain hardening behavior of a polymer is an important parameter for
foaming applications. Due to weak tensile melt strength at higher temperature, the processing
window for foaming polypropylene is extremely narrow, the lower limit being determined by the
melting temperature of polypropylene (160°C - 165°C) and the upper temperature limit being
determined by minimum accepted tensile melt strength that can result in microcellular foams.
Burt [108] reported in his work that the operating temperature range for producing acceptable
polypropylene foam, using linear polypropylene, is merely 4°C as opposed to 50°C for
polystyrene.
The mechanism of polymer deformation during bubble growth in a polymer foaming
process is biaxial stretching. Munstedt et al. [48, 109, 110] has measured the biaxial extensional
viscosity for linear PP as well as branched PP using lubricated squeezing flow and have observed
that the biaxial extensional viscosity shows similar trends to the uniaxial extensional viscosity. In
particular, they concluded that if strain hardening is observed in uniaxial extensional flow, it will
be observed in biaxial extensional flow as well. This is useful because more reliable equipment
and procedures are available for measurement of uniaxial extensional viscosity of melts than for
the measurement of biaxial extensional viscosity of polymer melts. Hence uniaxial elongational
flow measurements may be used to identify relevant trends in the biaxial extensional viscosity.
Various approaches have been adopted to produce melt strain hardening of linear PP and
eventually improve its foamability. Chemical or radiation crosslinking is a common method used
for introducing strain hardening property in the resulting polymer melts [44, 111, 112].
Crosslinking the polymer results in melt strain hardening of PP and results in better foam;

17

however, the crosslinking process leads to loss of recyclability, chain scission and oxidative
degradation of the PP, which affects its physical properties [113-115]. Another approach to
impart strain hardening behavior in polypropylene is to introduce tree type branching network. It
is well established that branched polypropylene show strong strain hardening character [110,
116, 117]. Blends of linear and branched PP have also resulted in polymer melt showing strain
hardening in uniaxial extensional flow [116, 118-120]. Due to the presence of strain hardening
property in blends of linear and branched PP, it has been reported [48, 58] that such blends foam
better than the linear resin, leading to smaller cell sizes and higher cell densities compared to
linear PP. These blends typically have around 20 wt% of the branched polymer that increases the
overall cost of the polymer material significantly.
Successful application of nanoparticles to serve as heterogeneous nucleating agent to
improve foamability of amorphous polymers led to its use to improve the foamability of linear
polypropylene foam. However, the extent of success in producing good quality polypropylene
foams with the assistance of nanoparticles was limited due to the absence of strain hardening
property of the linear polypropylene melt in uniaxial extensional flow. Numerous work has been
done on foaming linear polypropylene in a wide array of processes, starting from continuous
extrusion process [48, 49, 121-123], injection foaming process [42], and batch foaming process
[41, 124, 125]. Gendron et al. [121] have reported extrusion foaming of nanocomposites based
on linear PP and branched PP having 2 wt% of organically modified nanoclay, with direct carbon
dioxide injection as the blowing agent to produce low density foams (0.2 – 0.3 gm/cc). In
addition to the nanoclay, they had also added 0.5 wt% of talc to serve as nucleating agent. They
observed severe coalescence in linear PP foam samples that showed mild improvement with the
addition of nanoclay at low concentrations (0.3 wt%) of the blowing agent. However, they

18

observed severe collapse in foam cellular structure of the PP/clay nanocomposite foam with
increase in the blowing agent concentration from 0.3 wt% to 0.5 wt%. However, foam samples
produced from branched PP system resulted in better cellular structure. But no significant
improvement in cellular structure was observed upon the addition of the nanoclay on top of 0.5
wt% of talc. Branched PP with talc as the nucleating agent resulted in foams with cell sizes in the
range of 250 – 600 µm. Similar results were obtained with both talc and nanoclay present in the
branched PP matrix. However, in the absence of talc, branched PP/clay nanocomposite resulted
in poor foam with cell sizes of the order of 600 – 1000 µm. This was probably due to poor
dispersion of the nanoclays, hence suggesting the reduced capability of nanoclay serving as
efficient nucleating agent in polypropylene matrix. In the absence of strain hardening in linear
PP matrix, even in the presence of both talc and nanoclay, they observed foams with large
coalesced zones, thereby indicating the importance of strain hardening in the polymer melt for
foaming applications.
Carreau et al [42] reported foaming studies of both linear and branched PP
nanocomposites containing 2 wt% of organically modified montmorillonite nanoclay in an
injection foaming process, using 1 wt% of a chemical blowing agent (azodicarbonamide). With
controlled injection foaming conditions, they were able to produce high density linear PP foams
with good cellular structure. They observed a small drop in average cell size from 60µm to
45µm, due to the addition of clay in linear PP matrix. An even better cellular structure was
obtained for branched PP and its nanocomposite foams. The addition of nanoclay in the branched
PP system resulted in a small reduction in the average cell size from 40µm to 35 µm. Zheng et
al. [49] also reported extrusion foaming of linear PP/clay nanocomposites with supercritical
carbon dioxide as the foaming agent. They reported improvement in cellular structure and
19

reduction in average cell size and increase in average cell density with the addition of nanoclay.
The linear PP foam had cell sizes in the range of 150 – 500 µm, which reduced to around 50 – 80
µm upon addition of 1 wt% of nanoclay. This improvement in foam quality can be attributed
partly to the use of supercritical carbon dioxide as the foaming agent that provides greater
pressure drop during the foaming process, usually not achievable by liquid carbon dioxide and
other chemical blowing agents. However, its use in industrial application demands stringent
safety regulations and cost. It has also been shown that production of polypropylene foam with
smaller cell size (20 – 60 µm) is possible by the use of a non-conventional batch foaming process
[41, 124, 125]. This limited improvement in foam quality by the use of nanoclays in linear
polypropylene is primarily because of absence of strain hardening behavior in PP melts and
crystalline nature of the polymer. This is not the case with amorphous polymers like polystyrene
that show strain hardening behavior under uniaxial extensional flow [126, 127] and do not have
any crystallization issues, leading to the successful application of nanoparticles in improving
foamability of amorphous polymers. This suggests that there is clearly a need for preparing and
foaming linear PP/clay nanocomposites that display significant strain hardening under uniaxial
extensional flow.
Takahashi et al. [128] studied the extensional flow behavior of polyethylene composites
with microscale fillers (glass beads, glass fiber, talc) and found that strain hardening goes down
with filler addition. Similar observations were made while studying extensional flow behavior of
ethyl-vinyl acetate / nanoclay nanocomposites [129]. Similar to the effect of fillers on strain
hardening behavior in microcomposites, Gendron et al. [121] and Carreau et al. [42] also
observed no improvement in melt strain hardening under uniaxial extensional flow upon the
addition of nanoclays to linear PP. However, Seong et al. [130] reported in their work done on

20

polyamide-6/clay nanocomposite that melt strain hardening behavior can be obtained in a
nanocomposite if the organically modified nanoclay tactoids are completely exfoliated. Recently,
nanoclays were used by Okamoto et al. [131] to introduce strain hardening property in maleated
polypropylene melt. Pathak and Jayaraman [132] have found that selection of high molecular
weight maleated polypropylene and the composition of the nanocomposite were critical to obtain
significant strain hardening in uniaxial extensional flow of the nanocomposite melt. In their
subsequent work, Okamoto et al. [100] evaluated the performance of maleated polypropylene /
clay nanocomposite for foam applications. They reported reduction in cell size and an increase in
cell density upon addition of nanoclay to the maleated polypropylene foam. However, it is well
established that functionalizing with maleic anhydride results in polypropylene with
unacceptable mechanical properties [133]. Hence there lies a strong need for preparing and
foaming linear PP/clay nanocomposites that display significant strain hardening under uniaxial
extensional flow.

1.6. Research Objectives
The objective of this research is to develop novel and controlled approach to impart strain
hardening behavior to linear polypropylene by the addition of organically modified nanoclays
and maleic anhydride grafted polypropylene without the loss of recyclability and economy and
investigate foamability of the resulting nanocomposites in a continuous extrusion foaming
process using chemical blowing agents for the manufacture of low-density polypropylene foams.
While formulating these approaches, one needs to develop understanding of the role of
compatibilizer and organically modified nanoclay, towards imparting strain hardening property
to the linear PP/clay nanocomposite melt.

21

To address this bigger objective, this research is categorized into three specific research
problems:
a) To synthesize linear polypropylene/clay nanocomposites showing strain hardening
behavior under uniaxial extensional flow and investigate the specific material properties
affecting foamability of linear polypropylene/clay nanocomposites. This was then used to
demonstrate superior foamability of TPO nanocomposite that is prepared from strain
hardening PP/clay nanocomposite, compared to TPO without any nanoclays.
b) To study the effect of additional nanoclay treatment by carefully chosen silane coupling
agents on the strain hardening property of linear polypropylene-clay nanocomposites and
their foamability.
c) To develop basic understanding on the effect of compatibilizer molecular structure on the
strain hardening behavior of the linear polypropylene/clay nanocomposites under uniaxial
extensional flow.
d) To investigate the effect of thermal annealing and melt processing by extrusion through a
uniform die on strain hardening behavior of polypropylene-clay nanocomposites, in the
context of recyclability.

22

Figure 1.1: Classification of layered-silicate polymer nanocomposites based on the extent of
dispersion of the nanolayers.

23

Figure 1.2: Effect of die processing temperature on evolution of cellular structure of semicrystalline polymer foam [46].

24

2. POLYPROPYLENE / CLAY NANOCOMPOSITES –
PREPARATION AND CHARACTERIZATION PROCEDURES
2.1. Materials
Linear polypropylene, Profax 6523, was used as the matrix polymer for most of the
study. In order to prepare PP/clay nanocomposite, maleic anhydride grafted polypropylene was
used as a compatibilizer. A tree type branched polymer (PF-814) was used to prepare a blend of
linear and branched PP. This was done so as to compare the foaming results of linear PP/clay
nanocomposites with the blend of linear and branched PP. Since the ultimate objective of this
study is to replace polyurethane foam with TPO foam, a TPO blends was made by blending an
ethylene–octane copolymer (Engage 8480) with linear PP. The details of all the polymers (melt
flow rate (MFR), molecular weight (Mw), and supplier) used in this study are presented in Table
2.1.
The nanoclays used in this study were organically modified montmorillonite clay
provided by Nanocor. Two different organoclays were used for the synthesis of the PP/clay
nanocomposites: (a) I.44P: with 35 wt% of dimethyl-di-octadecyl ammonium cation and (b)
Nanocor I.30P: with 24 wt% octadecyl ammonium cation. Both these organoclays have same
molecular equivalent of organic surfactants in the clay galleries.
The chemical blowing agent used in this study was Cell-span 693K (gas volume
70cc/gm), provided by Phoenix Plastics with the active ingredient being azodicarbonamide. Four
different silane coupling agents were used for the treatment of nanoclays:
a) N-aminoethyl-Aza-2,2,4-trimethylsilacyclopentane (Diamino monomethoxy-Silane)
b) N-(2-aminoethyl)-3-aminopropylmethyldimethoxysilane (Diamino dimethoxy-Silane)
c) N-(2-aminoethyl)-3-aminopropyltrimethoxysilane (Diamino Trimethoxy-Silane)
25

d) Phenyl trimethoxy silane.

2.2. Nanocomposite Preparation
The nanocomposites for the scope of this research were prepared by melt mixing
technique. For the preparation of polymer/clay nanocomposites, both continuous and batch
mixing process have been used for research purpose as well as on commercial scale. Twin screw
extruder has been used as a continuous mixing process for the preparation of polymer/clay
nanocomposite. Batch mixing process has been increasing used, primarily for research purpose,
for the preparation of polymer nanocomposites at small scale. Common batch mixers used for
this purpose include Banbury mixers [134-137], DACA mixers [138] and Gelimat mixers [139].
The residence time and shear rate in all these mixing techniques were optimized so as to result in
nanocomposites with best possible dispersion of the nanoplatelets with minimal degradation of
the polymer and the organically modified clay.
In this study a 75 cc Banbury batch mixer, manufactured by CW Brabender Instruments
Inc., with two counter rotating Banbury screws was used for the preparation of
polypropylene/clay nanocomposites. Prior to compounding, maleated polypropylene pellets and
the nanoclay powder were dried in a vacuum oven at 80ºC and 20” Hg vacuum, for 12 hours.
The dried polymers were then premixed with the nanoclay power in a plastic beaker, according
to the proportions required for a 47gm batch nanocomposite formulation. This homogenized
polymer – clay mixture was then fed to the preheated Banbury mixer bowl. For majority of the
research, the nanocomposite was prepared by melt mixing at 180ºC, 150rpm for 10 minutes. The
mixing temperature of 180ºC was chosen based on the fact that most polypropylene has a
melting temperature of around 165 ± 2ºC. This mixing protocol was chosen so as to maximize
26

shearing time and minimize polymer and/or clay thermal and/or shear degradation and has been
optimized in our lab through previous work [135, 140]. In order to minimize oxidative thermal
degradation, a constant nitrogen purge was maintained inside the mixing bowl. For part of the
study additional mixing conditions used were: (a) 200ºC, 80 rpm, 10 minutes mixing time and
(b) 200ºC, 40 rpm, 10 minutes mixing time. This was done so as to study the effect of shearing
on nanocomposite structure and its effect on strain hardening under uniaxial extensional flow.
The reason for increasing the temperature was to match the actual melt temperature of 180ºC
mixing case, which was measured to be 200ºC due to additional heat generated due to shear at
150 rpm. The final nanocomposite was then collected from the mixing bowl as clumps with the
help of a brass scraper. These clumps were used directly for the preparation of samples for
rheological studies and structural characterization of the final nanocomposite, while for foaming
studies these clumps were granulated in a mini granulator.

2.3. Silane Treatment of Nanoclays
In order to carry out silane treatment of the nanoclays, 15g of an organoclay was
dispersed in 500ml of a solvent (80 wt% methanol + 20 wt% de-ionized water) using a magnetic
stirrer. In a separate flask, 4.5gm of silane was diluted with 200ml of the same solvent. This
diluted silane solution was slowly added to the clay dispersion and left under constant stirring
condition for 6 hrs under ambient conditions. After 6 hrs, the clay suspension was filtered and
washed at least 3 times using the same solvent to remove any unreacted silane. The resulting clay
cake was then left for 24 hrs in a vacuum oven at 80ºC, 20” Hg pressure for curing. The cured
clay was then powdered in a mortar-pestle and sieved with a No. 200 sieve. The powdered clay

27

was used for various characterizations and for the preparation of the nanocomposites with linear
polypropylene, using PP-g-MA as a compatibilizer.
Three different solvent systems were chosen for the study:
(a) A mixture of 80 wt% methanol and 20 wt% de-ionized water,
(b) A mixture of 80 wt% methanol and 20 wt% de-ionized water, and
(c) De-ionized water.
The organoclays were easily dispersed in the first solvent, however in order to disperse
the clay in the second solvent, a lab sonicator was used to prepare the clay dispersion. Since the
nanoclays used for silane treatment is already pretreated with organic surfactants, it was next to
impossible to disperse the clay in de-ionized water. Hence an aqueous cake of the organoclay
was obtained from Nanocor, an unfinished product obtained from an intermediate step during the
production of the organoclay, which was easily dispersed in de-ionized water.

2.4. BET Surface Area Measurements
Brunauer, Emmett, and Teller in 1938 proposed a theory [141] with an objective to
explain the physical adsorption of gas molecules on a solid surface that can further be used as the
basis for an important analytical technique for the measurement of the specific surface area of
the surface. This theory was invariably known by the initials of the authors, BET theory. This
theory is an extension of Langmuir theory [142] for monolayer molecular adsorption and is
applicable to multilayer adsorption. The fundamental hypotheses of the BET theory are:
1. The surface is homogeneous
2. Gas molecules physically adsorb on a solid surface in layers infinitely
3. There is not interaction between each adsorption layers
28

4. Langmuir theory can be applied to each layer.
Based on these hypotheses, the BET equation is given as:
(2.1)
where, P and P0 are the saturation and equilibrium pressure of the adsorbates, v and vm are the
total and monolayer adsorbed gas volume and c is the BET constant, given by equation (2.2):
(2.2)
where, EA is the heat of adsorption of the first layer and EL is the heat of liquefaction of the
second and higher layers.
This theory is widely used to calculate the total surface area (Stotal) of solids by physical
adsorption of gas molecules, using equation (2.3)
(2.3)
where, ‘N’ is Avogadro’s number, ‘V’ is the molar volume of adsorbent gas and ‘s’ is the cross
section area of the adsorbing species.
In the present work, Micrometrics (Gemini V) Surface area and pore size analyzer is used
to determine the surface area of the various organically modified montmorillonite nanoclays,
using nitrogen as the adsorbent gas. Prior to subjecting the nanoclays for BET surface area
analysis, around 1.5gms of clay samples were degassed at 120ºC with nitrogen as the purge gas
for 24 hours. The sample was then weighed accurately and subjected to the BET surface area
analysis following the adsorption procedure. Helium gas was used to estimate the free volume in
the sample tube followed by the adsorption of nitrogen gas, under liquid nitrogen temperature, to

29

calculate the total volume. These informations were then used to estimate the specific surface
area (area/gm) of the clay samples.

2.5. X-Ray Diffraction
X-ray diffraction (XRD) technique was used for structural characterization of various
organically modified montmorillonite and its resulting polypropylene/clay nanocomposites. This
technique is based on the fact that if a crystal was an assembly of regularly spaced atoms, then
they can act as scattering sites for X-ray beam that has a wavelength approximately equal to the
inter atomic distance in a crystal. Thus it can provide insight to the microsctructure of the crystal
by diffracting the X-ray beam. A mathematical expression for such diffraction phenomenon was
proposed by Bragg and is given by equation 2.1, known as Bragg’s Law of diffraction.
(2.4)
Where ‘n’ is the order of reflection, ‘λ’ id the wavelength of the X-ray beam, ‘d’ is the spacing
between atoms or planes, and ‘θ’ is the angle of diffraction. The first order reflection is only
considered in this work [143].
The X-ray diffractometer used in this study was a Rigaku Rotaflex Ru-200BH X-ray
diffractometer, which is equipped with a Ni-filtered Cu Kα radiation source and operated at
45kV and 100mA. This instrument has four diffraction slits; one between the X-ray source and
the sample (1/6º); two between the sample and the goniometer (1/6º and 0.3º) and the last slit is
mounted on the detector (0.3º). The selection of the optimal diffraction slits combination was on
the basis of previous work done in our lab. To obtain the X-ray diffraction pattern of the clay
samples, the powder clay samples were compacted on a 1”x1” cavity on a glass slide. For the
polymer nanocomposite specimens, a 20x40x1 mm3 sample was compression molded at 200ºC
30

and 10 tonnes pressure in a Wabash compression molder. All measurements are recorded for a
2θ scan range of 0.5º to 10º, at a rate of 0.5º/min and at equal increments of 0.01º.
Natural montmorillonite clay has a mean interlayer basal spacing of the (001) plane (dspacing) of about 0.95nm. The organic treatment by various octadecyl ammonium surfactant
leads to intercalation of the surfactant molecules inside the clay galleries resulting in an increase
in d-spacing. For the nanoclay treated with octadecyl ammonium cation (single tail alkyl chain),
the mean d-spacing increase to about 2.34 nm and the clay treatment with dimethyl-di-octadecyl
ammonium cation results in a mean d-spacing of about 2.54 nm.

2.6. Transmission Electron Microscopy
Transmission electron microscopy (TEM) is a valuable analytical tool, where a beam of
electrons is transmitted through an ultra thin sample. The high energy electron beam interacts
with the sample during transmission and produces an image containing finest details of the
internal structure of the specimen. The electron beam produced from an electron gun is
controlled using a series of lenses (condenser and objective lens) before it passes through the
sample. Even though TEM is a powerful tool to produce high resolution images, up to the nanodimensions of specimens, it comes with various limitations. The biggest drawback is that TEM
can scan very small area of a given sample (usually around 0.3 mm2). Another limitation is the
sample preparation. Since the technique relies on transmission of electron beam through the
sample, the thickness of the sample has to be less than about 90nm. Moreover, the thickness
should be uniform or else it will produce images with varying degree of brightness due to
variation in sample thickness. Hence, this technique is difficult to be used with a lot of materials.
A sectioning device called an ultra-microtome is generally used to prepare samples. For polymer
31

samples, the sectioning has to be done below its glass transition temperature in order to preserve
the structure of the specimen under investigation. Since the glass transition temperature of
polypropylene samples below 0ºC, cryo-sectioning is required to provide enough rigidity so as to
facilitate uniform sectioning of the sample and to maintain its structure. Cryo-sectioning is a
technique where thin, uniform sections of specimen are produced by cutting the material with a
diamond tip knife, below room temperature. The low temperature is maintained using liquid
nitrogen as a coolant.
The ultra microtome used for this study was PTXL ultra-microtome (RMC, Boeckeler
Instruments) connected to a CRX cryo unit. In order to prepare the sample, a small rectangular
piece (1-2 mm2) was cut out of a compression molded sample. The cut sample was then mounted
on the ultra-microtome maintained at -120ºC using concentrated sucrose solution. This section
close to 90nm was then cut and mounted on 200 mesh size formvar coated copper grids. These
sections were then used to study the dispersion and structure of nanoclay layers in various
PP/clay nanocomposites using JEOL 100CX TEM at 120KeV acceleration voltage.

2.7. Differential Scanning Calorimetry
Differential scanning calorimetry (DSC) is a thermoanalytical technique that measures
the change of the difference in heat flow rate to the experimental sample and to a reference
sample while both are subjected to a controlled temperature program. Throughout the experiment
both the experimental and the reference sample are maintained at a constant temperature. The
reference is generally an empty sample holder (hermetically sealed aluminum pan). This
technique was used to determine the melting temperature, crystallization temperature, fractional
crystallinity, and rate of crystallization of the pure polymer or polymer/clay nanocomposites. The
32

principle underlying this technique is that, when the polymer sample undergoes a physical
transformation (phase change), which is associated with a change in enthalpy, the instrument
either increases or decreases the amount of heat flow to the sample so as to maintain both the
experimental and the reference sample at the same temperature.
The differential scanning calorimeter used in the present study was a TA Q-Series
instrument – DSC Q10TA. The sample (8-10mg) was heated to 200 °C (5°C /min ramp), then
cooled to 40 °C (-5°C /min ramp) under nitrogen atmosphere and the entire cycle was repeated.
Results are reported from the second run as the first run is primary to remove any thermal history
present in the sample. The peak obtained from the heating cycle denotes the melting temperature
(Tm) and the peak from the cooling cycle represents the crystallization temperature (Tc). The
fractional crystallinity (XC) and the average rate of crystallization (k) in each sample were
obtained by the following relations.

where, w is the weight fraction of polymer in the sample, ΔHc is the enthalpy of crystallization of
polymer in the sample obtained by computing the area under the crystallization peak, the
quantity 207 J/g refers to the enthalpy of crystallization of 100% crystalline PP and tC refers to
the time span over which crystallization is completed.

33

2.8. Thermo-gravimetric Analysis
Thermo-gravimetric analysis or thermal gravimetric analysis (TGA) is a technique used
to determine weight loss (or weight gain) of a sample as a function of temperature under
controlled atmosphere. As a sample is heated, it can undergo weight loss due to drying, due to
liberation of gases from chemical reactions or oxidative decomposition.
The Thermo gravimetric analyzer, TA instrument TGA Q500 series, was used in this
study to study the percentage weight change with respect to change in temperature for different
organically modified montmorillonite clays and the chemical blowing agent. A high precision
platinum pan was used to and the experiment was carried out under an inert nitrogen atmosphere.
About 6 – 10mg of sample was used in this study. A high resolution temperature ramp procedure
was selected to increase the temperature from room temperature to 800°C at a ramp rate
10°C/min. This procedure allows maintaining an almost constant temperature when the sample is
undergoing weight change at a particular temperature, thereby resulting in greater accuracy.

2.9. Rheology
Rheology is the science of flow and deformation [144] that gives valuable information
about the microstructure of a material. In this study, we have used rheological studies to
investigate the role of various organically modified montmorillonite nanoclays on the melt
rheological properties of polypropylene, under both shear and uniaxial extensional flow.

34

2.9.1. Shear Rheology
Rheology involving simple shear flow is the most common type of rheological study,
generally, used to investigate the linear viscoelastic response of a material (in this case molten
polymer). Linear viscoelastic studies of polymer melt produces valuable information about its
molecular structure. For this study, simple shear flow is produced by sandwiching the test
material between a set of smooth parallel plates, where one of the plates is stationary while the
other one is mobile. This moving plate produces a shear deformation that is analyzed using
relationships derived from equation of motion using the following assumptions [144]:
1. Steady, laminar, isothermal flow
2. υθ (r, z) only, υr = υz = 0
3. Negligible body forces
4. Cylindrical edge
Based on these assumptions, the equation of motion for parallel plate geometry transforms into:

Further, using these relationships and assuming a “no-slip” condition at the plate
surfaces, the relationship between the shear stress and the torque generated is given as:

35

In the present study, we have used strain controlled small amplitude oscillatory shear type
flow to obtain various linear viscoelastic properties of the polymer melt. The oscillatory strain
and stress in such flow is given as:
(2.11)
(2.12)
The oscillatory stress can further be decomposed into an in-phase plastic or elastic
component and an out-of-phase viscous component, which are used to obtain an in-phase plastic
or elastic or storage modulus (G’) and an out-of-phase viscous or loss modulus (G”) respectively.
These are defined as:

In order to generate the mentioned and many more such linear viscoelastic polymer melt
rheological properties, strain controlled oscillatory shear measurements were performed on an
AR2000 rheometer from TA Instruments with 40mm diameter steel parallel plates. The linear
viscoelastic regime for various PP and PP/clay nanocomposites was first determined by running
a strain sweep tests at a fixed frequency. The strain beyond which the storage modulus starts
changing as a function of strain is defined as the non-linear viscoelastic range and before that is
the linear viscoelastic range. A strain of 1% was selected for testing as this value of strain was
well within the linear viscoelastic limit for various polymeric materials used in this study. In
order to generate linear viscoelastic data as a function of angular frequency, an oscillatory
frequency sweep test was run with angular frequency ranging from 0.01 to 100 rad/s at 180ºC
36

and 1% strain. An oscillatory time sweep was also conducted at a frequency of 1 rad/s for 1 hour
to observe the stability of the nanocomposite melt with respect to time. In this test, a constant
value of the storage modulus over time within ± 5% was considered as a stable nanocomposite
with minimal changes in the microstructure. To ensure minimal oxidative thermal degradation, a
continuous purge of nitrogen was maintained during the tests.

2.9.2. Extensional Rheology
As opposed to shear rheology, extensional rheology is a study of material properties
under shear free elongational flow. Extensional flow is of particular importance in various
commercial processing, for example, fiber spinning, foaming, blow molding, thermoforming,
film blowing etc, however, material properties determined from shear flow alone cannot be
applied for the prediction of such processes. This makes elongational or extensional flow
behavior of polymer melts a very important property that needs attention for our research. Due to
the presence of shear stresses at stationary boundaries, homogeneous extensional flows are more
difficult to be produced compared to simple shear flows [145]. In this work, polymer melts are
subjected to uniaxial extensional flow in order to study its extensional rheological behavior.
Uniaxial extension is an axi-symmetric deformation, where a tensile stress is applied
along the longest dimension and the other free surfaces of the specimen are under compression.
In order to study the uniaxial flow behavior of polymer melts used in this research, TA ARES
rheometer mounted with “Extensional Viscosity Fixture” (EVF) has been used. The design of the
fixture is based on the original concept by Meissner [146]. This fixture consists of two cylinders,
one stationary connected to the transducer to measure the force and the other rotating around the
stationary one to wind up the sample (Figure 2.1). The extensional force (FS(t)) measured by the

37

transducer and the sample area (AS(t)) , both a function of time, is converted to the extensional
stress (σE(t)) as follows:
(2.15)
This is then used to calculate the transient uniaxial viscosity (ηE(t)) at a particular strain rate as
follows:
(2.16)

In order to measure the transient extensional viscosity of the various polymer melts used
in this study, a constant Hencky stain (εH) of 3 and four different strain rates (0.1, 0.5, 1.0, 2.0 s1

) have been used. The lowest test temperature has been maintained at 180ºC, well above the

melting point of various polymer melts used in this study (165 ± 3ºC). A higher temperature of
200ºC has also been used for few tests to observe the behavior of the transient extensional
viscosity at higher temperatures. Since these tests are very sensitive to the sample condition, care
has been taken so that the test specimens are free of any voids or inhomogeneities. The test
samples were prepared by compression molding the polymer chunks into 18 x 10 x 0.75 mm3
rectangular specimens. The polymer chunks were allowed to melt in the compression molder at
200ºC and then the melt was subjected to a pressure of 1 tonnes for 5mins followed by 5 tonnes
for 5 mins. The melt was then cooled to room temperature using cooling water maintaining the
pressure of 5 tonnes. The samples were then carefully removed from the mold and examined for
any defects (voids, cracks, grain boundaries, foreign particles and uniform shape). The state of
the stretched specimen was also inspected immediately after the test, so as to observe any

38

sagging of the polymer melt due to gravity effect and any non-uniform stretching in order to
ensure that there was uniform stretching in the central portion of the specimen during the test.

2.10. Polymer Foaming Procedure
Continuous extrusion foaming process was used to prepare PP nanocomposite foams. A
single screw extruder was used for this purpose with a 2mm diameter die (L/D=15) and a tapered
transition from the extruder barrel. The optimum temperature for extrusion of foam from each
compound was selected based on the crystallization temperature of the compound. Park et al.
[45, 103, 104] have reported effects of die temperature and pressure, on continuous extrusion
foaming of polymers with a physical blowing agent. Naguib et al. [46] reported that both early
and delayed onset of crystallization is detrimental to foam cellular structure. In their work done
on extrusion foaming of branched PP using n-butane as a blowing agent, they report an optimum
melt temperature that provides maximum expansion ratio in the foam samples. They observed
foam collapse away from the die when the melt is at a higher temperature, due to delayed onset
of crystallization causing the gas to diffuse out of the sample. On the contrary, a lower than
optimum melt temperature resulted in early onset of crystallization leading to insufficient
foaming. The conclusions of these studies were used to maintain the optimum temperature in the
extruder for the foaming experiment. The extruder has four different temperature zones: three in
the barrel and one in the die. The first three temperature zones were maintained at 180°C, 200°C
and 170°C respectively for all foaming experiments done using cell span 693K as the blowing
agent. The decomposition profile of the blowing agent with increase in temperature was obtained
from TGA (Figure 2.2). According to the TGA result, majority of the blowing agent is
decomposed by 215°C. However, in order to avoid thermal degradation of the material in use
39

and to take advantage of maximum volume of gas produced from the blowing agent, the second
zone was maintained at 200ºC. The first temperature zone was maintained at a lower temperature
compared to the second zone so as to prevent any loss of gas from the hopper due to early
decomposition. The second temperature zone is the region where maximum decomposition of
the blowing agent is expected. The fourth temperature (die temperature) zone was maintained at
approximately 35°C above the crystallization temperature; this was found to be optimal in our
foam extrusion work, as it helped to maintain a constant pressure drop of 4000psi with various
polymer samples during the foaming experiments. The third temperature zone was maintained in
between the second zone and the die temperature. The chemical blowing agent (CBA) was used
at 3 wt% in this study. The CBA was premixed with the granulated polymer in a bag and was fed
through the hopper. The extruder was operated at a constant speed of 20 rpm in this study. While
flowing through the single screw extruder the CBA and the PP nanocomposite fuse to form a
homogeneous melt. Near the die exit, due to steep pressure drop through the die, the gas laden
polymer melt started to nucleate bubbles and the extruded polymer foam strand continued to
expand until it solidified under ambient conditions. The polymer strands were collected for
further foam characterization.

2.11. Characterization of Foam Samples
2.11.1.

Density Measurement

Archimede’s principle is used for determining the density of the foam extrudates using a
Sartorius YDK 01 density measurement kit. According to Archimede’s principle when a solid is
immersed in a liquid, it is subjected to the force of buoyancy. The value of this force is same as

40

that of the weight of the liquid displaced by the volume of the liquid. This kit is equipped with a
hydrostatic balance that enables the measurement of weight of the foam samples in air (W(a)) as
well as in water (W(l)). For the present study, water is used to do these measurements and
equation 2.17 is used to obtain the density of foam samples. A minimum of three measurements
are done with different specimen for a particular foam sample in order to get an average value of
density.
(2.17)

2.11.2.

Scanning Electron Microscopy

Scanning electron microscope (SEM) is an electron microscope that provides
topographical information of the surface by scanning it with a high energy electron beam. The
beam of electrons interacts with the surface of the specimen, thereby providing with important
information about its surface topography and chemical composition along with various other
properties.
Scanning electron microscopy is based upon “electron in” – “electron out” process in
ultra-high vacuum (UHV) environment. As the primary electron beam is scanned across the
surface, electrons of a wide range of energies are emitted from the surface in the region where
the beam is incident. These electrons include backscattered electrons and Auger electrons, but
the vast majority of these are the secondary electrons formed in multiple inelastic scattering
processes. These secondary electrons are collected by the detector, converted to a voltage and
amplified. The amplified voltage is then applied to the grid of the cathode-ray tube (CRT) and
modulates or changes the intensity of the spot of light on the surface. For example, if the beam is
on a projection on the surface, a large number of secondary electrons will be detected, causing a

41

large voltage in the detector and results in a bright spot on the surface of the CRT. If the beam of
electron then moves to a depression on the sample, fewer electrons will be detected, resulting in
a darker spot on the surface of the CRT. The contrast in the micrograph arises from variations in
the surface topography. Consequently, the secondary electron micrograph is a direct image of the
real surface structure.
Since an electron beam is bombarded on the specimen, care needs to be taken to avoid
any accumulation of electron on the sample, called charging of the sample, leading to hot spots
in the image. For samples that are conductive this is not an issue, however for non-conductive
polymer samples, the specimen need to be coated with conductive gold or osmium tetraoxide
particles. SEM technique is used in this research to study the cellular structure of the foam
extrudates of various PP and PP/clay nanocomposite samples using JEOL 6400 scanning
electron microscope. Foamed extrudates were cryogenically fractured to maintain the structural
integrity of the foam cells and the fractured surface was coated with osmium tetroxide particles.
SEM micrographs were taken from these fractured surfaces at a magnification of 60X. The cell
size, cell size distribution and cell density were evaluated using ImageJ software. A minimum
number of five SEM images were used to get an average value these parameters. The cell density
(N) was calculated using the following equation:

(2.18)
where n is the number of cells in the defined area A, derived from SEM micrographs.

42

Maleic
anhydride
content
(wt%)

Name

Polymer
Type

Supplier

Mol. Wt

Melt Flow
Rate
(g/10 min)

Profax 6523

Linear PP

Lyondell-Basell

390,000

4.0

--

Profax PF 814

Branched PP

Lyondell-Basell

--

3.0

--

Exxelor PO 1015

PP-g-MA,
copolymer

ExxonMobil
Chemical Co.

180,000

150.0

0.25-0.5

Exxelor PO 1020

PP-g-MA,
homopolymer

ExxonMobil
Chemical Co.

80,000

430.0

0.5-1.0

Bondyram
1001CN

PP-g-MA,
copolymer

Polyram Ind.

--

100.0

0.75-1.0

Bondyram 1001

PP-g-MA,
homopolymer

Polyram Ind.

--

100.0

0.75-1.0

Honeywell Ind.

22,000

--

4.0

Dow Chemical
Co.

--

1.0

--

AC 950
Engage 8480

PP-g-MA,
homopolymer
Ethylene
Octene
elastomer

Table 2.1: Details of various polymers used in this study.

43

Transducer shaft
Rotating drum

Figure 2.1: A schematic of Extensional Viscosity Fixture (EVF) used for the measurement of
uniaxial elongational viscosity of polymer melts.

44

Figure 2.2: Thermal decomposition profile of Cell span–693k, as obtained from thermogravimetric analysis. At 200°C, 15.2 wt% of the chemical blowing agent is
decomposed to produce gas (primarily nitrogen) for foaming various polymers used
in this study.

45

3. EXTRUSION OF LINEAR POLYPROPYLENE/CLAY
NANOCOMPOSITE FOAMS
3.1. Introduction
A variety of thermoplastic polymer foams have been produced with polystyrene [27, 95]
poly (lactic acid) [51], poly (ethylene terephthalate) [50] and polycarbonate [36]. Amongst the
various foamed polymers, polyurethane and polystyrene continues to dominate the foam market
today. Polypropylene foams are expected to replace polystyrene and polyurethane foams in
various applications, because polypropylene can be used at higher service temperatures, and has
greater stiffness, excellent chemical, abrasion and impact resistance and low material cost along
with total recyclability. However, continuous foaming of linear PP has been largely restricted by
its rheological properties [47, 58, 108] and the semi-crystalline nature of the polymer [60, 61].
Linear PP does not show strain hardening under extensional flow, a necessary property for
producing good foam particularly in crystallizing polymers [58, 100]. Polymers without strain
hardening property develop thinner cell walls during bubble expansion in the foaming process,
which results in coalescence of bubbles and even open-cell foams [47, 105, 106]. The semicrystalline nature of the polymer also makes the foaming process more complicated because of
simultaneous nucleation and growth of gas bubbles and crystals. Reignier et al. [61] have
demonstrated with batch foaming of poly–(ε-caprolactone) that crystallization affects both
bubble nucleation and growth.
Blends of linear PP and branched PP foam better than the linear resin [48, 58, 59],
leading to smaller cell sizes and higher cell densities than with the linear resin. This is because an
increase in the extensional viscosity of the blends during the expansion process provides stability
to the cell walls [116, 118-120]. These blends typically have 20 wt% of the branched polymer
46

and add to the cost significantly. Partial crosslinking also increases the tensile melt strength of
PP [44, 111, 112] and results in better foam; however, the crosslinking process leads to loss in
recyclability, chain scission and oxidative degradation of the PP, which affects its physical
properties [113-115]. This section of the research describes an alternative and novel method of
introducing strain hardening property in linear PP melts that has the potential for lower cost and
easier recyclability. This method incorporates small amounts of organoclay with an appropriate
maleated PP as compatibilizer.
The use of nanoclay in foaming polymers to obtain closed cell foams with narrower cell
size distributions and high cell densities was demonstrated for the amorphous polystyrene by Lee
et al. [9, 95]. They reported over 50% reduction in average cell size and around six to ten fold
increase in cell densities of the polystyrene nanocomposite foam by the addition of 5wt%
organically modified nanoclay. This is attributed to the much higher surface area made available
for nucleation of bubbles by well dispersed nanolayers [9, 42, 147]. The enhanced nucleation
efficiency has also been reported for foaming of various other polymer nanocomposites [95, 147150].
Other workers [41, 42, 49, 121] have tried this approach with the semi-crystalline
polypropylene with limited improvement in foam quality produced in continuous extrusion or
injection molding process. For example, Carreau et al. [42] reported foaming studies of both
linear and branched polypropylene nanocomposites containing 2 wt% of nanoclay in an injection
foaming process, using 1wt% of a chemical blowing agent (azodicarbonamide). They observed a
25% drop in average cell size, due to the addition of clay in linear PP. An even smaller reduction
in cell size (12.5%) was observed for branched PP foam samples due to the addition of clay.
Gendron et al. [121] have reported extrusion foaming of nanocomposites based on linear PP and

47

branched PP having 2 wt% nanoclay, with direct carbon dioxide injection as the blowing agent.
They observed large coalesced zones in linear PP foam samples when foamed with lower
concentrations of CO2. Addition of 2wt% of nanoclay to the linear PP matrix slightly improved
the cellular structure of the foam samples, however, when foamed with higher concentrations of
CO2, the PP-clay nanocomposite again led to formation of large coalesced regions. With the
addition of nanoclay to the branched PP matrix, they obtained foams with cell sizes in the range
of 600–1000 µm. However, when they added 0.5wt% of talc serving as nucleating agent, they
obtained foams with cell sizes in the range of 250 – 600 µm. However, no further reduction in
cell size was observed when both talc and nanoclay was added to the branched PP matrix. With
their research they showed that the efficiency of nanoclay serving as nucleating agent is greatly
dependent on the level of nanoclay dispersion obtained in PP nanocomposites and that its
nucleating ability is inferior to talc. Carreau et al. [42] and Gendron et al. [121] also reported
uniaxial extensional flow behavior of linear PP/clay nanocomposite melt and observed no
improvement in melt strain hardening upon addition of nanoclay to linear PP. There is clearly a
need for preparing and foaming polypropylene clay nanocomposites that display significant
strain hardening in extensional flow.
The use of nanoclay to introduce strain hardening in maleated polypropylene melts alone
was demonstrated first by Okamoto et al. [100] and this was shown with a high molecular weight
polypropylene. Pathak and Jayaraman [132] have found that selection of a high molecular weight
maleated polypropylene and the composition of the nanocomposite were critical to obtain
significant strain hardening in uniaxial extensional flow of the nanocomposite melt.
The objective of the present research is to investigate differences in foam quality
produced from two sets of polypropylene clay nanocomposites, all prepared with good
48

dispersion of clay – one set that does not display melt strain hardening and another that has much
improved strain hardening behavior under uniaxial extensional flow. These nanocomposites were
prepared by using different ratios of compatibilizer to clay and two different grades of the
compatibilizer. The foams were produced in a single screw extruder with the same proportion of
a chemical blowing agent in all cases. It is worth noting here that the literature on foaming
polymer nanocomposites has been focused largely on physical blowing. The results showed that
besides providing a greater concentration of nucleating sites for foaming the polymer matrix,
addition of nanoclay with appropriate compatibilizer has the potential for producing melt strain
hardening and altering the crystallization kinetics favorably in extrusion of closed cell
polypropylene foams. In order to compare the foaming results of linear PP/clay nanocomposite
foams with a blend of linear and branched PP in the same process, a blend of linear and branched
PP was prepared and foamed using the same process. Since the final objective of this research is
to foam TPOs for automotive application, this technology was extended to foam TPOs.

3.2. Materials and Experimental Procedures
Linear polypropylene (Profax 6523) was used in this part of the research along with two
grades of polypropylene-grafted-maleic anhydride (PP-g-MA): Exxelor PO-1015, labeled as PPg-MA1 is a random copolymer of propylene and ethylene, and AC-950, labeled as PP-g-MA2 is
a homopolymer. A tree type branched polypropylene, PF 814, was used at 20 wt% to prepare an
80:20 blend of linear and branched PP so as to compare the foaming results with linear PP/clay
nanocomposite system. The nanoclay used in this study was organically modified
montmorillonite clay, I.44P. The chemical blowing agent used in this study was Cell-span 693K.

49

In order to prepare a TPO material, Engage 8480 was used as the elastomer. The details of all
materials are provided in section 2.1.
Five different PP-clay nanocomposites were compounded for this study. A TPO blend
and TPO nanocomposite, derived from a strain hardening PP/clay nanocomposite, was also
prepared for foaming studies. The details of the composition are tabulated in Table 3.1. A 47g
Banbury batch mixer at 180°C and at a rotation speed of 150 rpm for 10 minutes was used for
the preparation of all the nanocomposites. However, to prepare a blend of pure polymers the
same mixer was used at 200ºC and at a rotation speed of 30 rpm for 10 minutes. This was done
because at higher rotation speed the entanglements formed amongst the branches of the branched
PP are lost [116], affecting the strain hardening behavior of the polymer melt (Figure 3.1). The
compounded materials were then foamed in a single screw extruder using 3 wt% of the chemical
blowing agent. XRD, DSC, extensional rheology was used for detailed characterization of the
polymer materials, and SEM and its image analysis along with density measurements was used
to characterization the foam samples. The details of all the preparation and characterization
procedures are provided in Chapter 2.

3.3. Results and Discussion
3.3.1. Morphology of Nanocomposites
X-ray diffraction (XRD) patterns obtained for all the PP nanocomposites along with the
nanoclay (I.44P) in the range of 2θ = 1–10º are presented in Figure 3.2. Lack of distinct features
in the XRD curve implies the existence of random-well dispersed clay platelets, however,
presence of distinct peaks indicates presence of ordered clay stacks and the corresponding 2θ

50

value can be used in equation 2.4 to calculate the spacing of the stacked platelets. The mean
interlayer spacing of the (001) plane (d-spacing) for the nanoclay used in this study was found to
be 2.54nm (2θ = 3.48º). The d-spacing of the (001) plane was found to be higher in all the PP
nanocomposites here, confirming intercalation by polymer. A comparison between PPNC-N1
and PPNC-N2 indicates a larger d-spacing of 3.53 nm in PPNC-N2, than in PPNC-N1 (2.85nm).
The higher maleic anhydride content in PP-g-MA2 (2.6 wt% bound maleic anhydride) compared
to PP-g-MA1 (0.5 wt% bound MA) allows greater hydrophilic interaction between the polar
montmorillonite surface and the maleic anhydride groups of the polymer chains and results in
greater intercalation of the polymer chains within the clay galleries. A comparison between
PPNC-N1 and PPNC-S1 nanocomposites, which have similar clay loadings and different
proportions of the same compatibilizer (PP-g-MA1) shows an increase in d-spacing with an
increase in the proportion of compatibilizer, thus indicating that similar extents of nanolayer
intercalation can be achieved by increasing the weight fraction of the lower maleic anhydride
content compatibilizer. Further increases in the proportion of this compatibilizer provided very
little change in the extent of intercalation as seen by comparing PPNC-S1, PPNC-S2 and PPNCS6.

3.3.2. Uniaxial Extensional Rheology
Although gas expansion within foaming polymers creates biaxial extensional flow in the
melt, uniaxial extensional flow measurements are relevant because measurements of biaxial
extensional viscosity in lubricated squeezing tests display trends that are similar to trends in
measurements of uniaxial extensional viscosity – cf. Munstedt et al. [48, 109]. In particular, they
conclude that if strain hardening is observed in uniaxial extensional flow, it will be observed in
51

biaxial extensional flow as well. This is useful because more reliable equipment and procedures
are available for measurement of uniaxial extensional viscosity of melts than for the
measurement of biaxial extensional viscosity of polymer melts.
The melt extensional viscosity transients for the linear PP and its nanocomposites at 180
°C and several strain rates are presented in Figure 3.3 and 3.4. For all specimens, as expected, the
transient extensional viscosity curves measured at different strain rates superimpose at low
strains to give the linear viscoelastic limit. Strain hardening of the melt is defined as an upward
deviation of the extensional viscosity from this linear viscoelastic envelope usually at strains
greater than one.
For a quantitative estimation of strain hardening effect in the nanocomposites, the strain
hardening parameter (χ) is defined as follows:

(3.1)

where,

is the transient extensional viscosity as a function of time and strain rate and
is the transient extensional viscosity in the linear viscoelastic regime. The linear

viscoelastic

may be determined in one of two ways: (a) as 3 times the transient shear

viscosity growth curve at very low strain rates or (b) by extrapolating the superimposed portion
of the curves for different strain rates. The two methods give the same result for homogeneous
melts but may give different results for filled polymer systems and polymer clay nanocomposites
in particular because of differences in orientation of anisotropic filler particles in elongational
and shear flows [128, 131, 151]; hence method (b) was used for this paper.

52

The transient extensional viscosity data for the linear PP melt did not show strain
hardening. The extensional viscosity data for molten PP-g-MA compatibilizer did not show any
strain hardening either. Amongst the PP nanocomposite melts, PPNC-N1 and PPNC-N2 showed
no strain hardening behavior. However, the transient extensional viscosity curves for PPNC-S1,
PPNC-S2 and PPNC-S6 show significant strain hardening of the polymer melt at all strain rates
tested. The strain hardening parameter, estimated at a Hencky strain of 2.25, is plotted against
strain rate for three strain hardening melts in Figure 3.5; these curves shows that strain hardening
in PPNC-S1, PPNC-S2 and PPNC-S6 melts decreased with increasing strain rate similar to the
trend shown by blends of linear PP with up to 20 wt% of branched PP [116] as well as blends of
linear PP with small amounts of crosslinked polymer [152, 153]. The melts of PPNC-S1, PPNCS2 and PPNC-S6 displayed similar strain hardening behavior indicating that 3 wt% of
organoclay and a 4:1 weight ratio of PP-g-MA to clay were adequate for producing this effect in
linear polypropylene/organoclay nanocomposites. The latter amounted to 0.20 g-mol bound
maleic anhydride on the compatibilizer chains per kg organoclay as presented in Table 3.1 which
represents one requirement for forming sufficient physical junctions (presumably by hydrogen
bonding) between the hydroxyl groups at the organoclay edges and the compatibilizer chains
which may also form bridges between clay particles.
The other requirement arises from the density of trapped entanglements along the
bridging compatibilizer chains. This may be seen from the fact that when the lower molecular
weight compatibilizer was used with more bound maleic anhydride amounting to 0.80 g-mol per
kg organoclay, the resulting nanocomposite melt did not display strain hardening. Hence, strain
hardening in the nanocomposite melts is analogous to that reported for blends containing small
amounts of crosslinked polymer. Just as blends of linear polypropylene with small amounts of

53

chemically crosslinked polymers contain trapped entanglements on chain segments between
crosslinks or junctions which respond to stretching differently from entanglements and chain
segments in the bulk [152], interactions between the surface or edge of the nanolayer and the
maleated polymer can form a network of physical junctions or anchor points between which
bridging compatibilizer chains trap entanglements to produce a similar effect. Thus, in addition
to good dispersion and sufficient interaction between the compatibilizer and the clay surface or
edge, the molecular weight of the compatibilizer should be high enough so as to form a sufficient
number of trapped entanglements along the compatibilizer chain between clay surfaces or edges.
The melt extensional viscosity transients for blend of linear (80 wt%) and branched PP
(20 wt%) at 180 °C is presented in Figure 3.1. The blend prepared at 30 rpm at 200ºC is used for
the foaming experiments in this study. As expected, the blend show significant deviation from
the linear viscoelastic regime at higher Hencky strain for all strain rates tested, indicating
significant strain hardening of the polymer melt. This blend was used for foaming along with
other linear PP/clay nanocomposites and a comparison of foam qualities was made later in this
study.
The uniaxial extensional viscosity behavior of TPO and TPONC melt at 180ºC is
presented in Figure 3.6. The TPO resulting from a blend of linear PP (70 wt%) and an elastomer
(30 wt%), shows absence of strain hardening behavior at all strain rates. However, the TPONC
resulting from the blend of PPNC-S2 (70 wt%) and the same elastomer (30 wt%), shows
significant improvement in the strain hardening behavior at all strain rates. This clearly
demonstrates that if the bulk linear PP component is a strain hardening one then the resulting
TPO nanocomposite would also show strain hardening behavior under uniaxial extensional flow.

54

Hence, if PPNC-S2 results in a good quality closed cell foam then TPONC should also lead to
similar foam properties.

3.3.3. Crystallization Behavior
Nonisothermal crystallization exotherms obtained from cooling ramps in DSC tests are
presented for all PP nanocomposites in Figure 3.7 and 3.8. The cooling rate for all the DSC tests
was 5°C /min. The fractional crystallinity (XC) and the average rate of crystallization (k) in each
sample were obtained by using equation 2.5 and 2.6 respectively. These results are tabulated in
Table 3.2.
It is clear from Table 3.2 that both compatibilizers crystallize more slowly than the linear
PP [42]. PP-g-MA1 being a random copolymer displayed a lower rate of crystallization from the
melt than PP-g-MA2 which is a homopolymer. Amongst the four nanocomposites, studied in this
work, PPNC-N2 has the highest rate of crystallization. The average rates of crystallization for
PPNC-N1, PPNC-S1, PPNC-S2 and PPNC-S6 (each containing PP-g-MA1) are all about half of
the rate for PPNC-N2 (which contains PP-g-MA2). This may be traced to the fact that PP-gMA1 crystallized at about half the rate of PP-g-MA2.
The nonisothermal crystallization exotherms for the blend of linear and branched PP is
presented in Figure 3.8. The fractional crystallinity and average rate of crystallization for
L80B20 is presented in Table 3.2. The average rate of crystallization for L80B20 is similar to the
PP/clay nanocomposites prepared using PP-g-MA2. This is probably because there is nothing
that would either interfere or slow down the folding of the polymer chains during crystallization,
thereby resulting into a blend that shows higher rate of crystallization.

55

3.3.4. Extruded Foam Structure
It is evident from the preceding sections that addition of nanoclay and a suitable
compatibilizer to linear polypropylene has the potential to produce three different effects: (a)
greater surface area for nucleation of voids, (b) changes in crystallization rate of the
polypropylene and (c) strain hardening in the melt state. PPNCs resulting in first two effects are
commonly observed due to the addition of nanoclays and PP-g-MA respectively, however a
combination of nanoclay and compatibilizer resulting in strain hardening behavior in linear PP
nanocomposites has not been reported previously. The five different polypropylene
nanocomposites of this work exhibit combinations of these effects to different extents as shown
in Table 3.3. Extruded foam samples were inspected visually to determine whether they were
open cell or closed cell foams. The foamed linear PP without any nanoclay particles turned out to
be open cell foam. Some open cell structure was seen in PPNC-N2 foam as well. The rest of the
nanocomposite foams were all closed cell foams. The scanning electron micrographs for all the
extruded foam samples are presented in Figures 3.9 – 3.17. The scanning electron micrograph of
linear PP foam sample, (Figure 3.9), shows evidence of severe cell coalescence and large scale
inhomogeneity, with regions of very low cell density. The micrograph of PPNC-N2 foam
presented in Figure 3.10 shows some improvement in quality with homogeneous cellular
structure; however the cells are still larger in size and the cell size distribution shows that most of
the cells are greater than 200 µm. The homogeneous nature of the cellular structure may be
attributed to the increased surface area for heterogeneous nucleation in PPNC-N2 over the neat
linear PP, but the absence of any strain hardening characteristics led to cell coalescence leading
to large size cells in the foam. The cell size distributions obtained from image analysis of the
micrographs are plotted for the various extruded foams in Figures 3.18 – 3.24. The bulk

56

densities, average cell sizes and cell number densities of the extruded polypropylene
nanocomposite foam samples are listed in Table 3.4. The bulk density of all the foam samples
are about 0.3g/cc indicating similar level of expansion was achieved during the foaming process
irrespective of the nature of the polymer melt. Thus any difference in cell size and cell size
distribution of the foam samples is not due to different levels of expansion achieved during
foaming but due to the additional surface area, slower rate of crystallization and, most important,
the strain hardening behavior of the polymer melt. Recall that neither PPNC-N1 nor PPNC-N2
showed extensional melt strain hardening and both have the same level of nanoclay, but PPNCN1 crystallizes slower than PPNC-N2. Hence a comparison between PPNC-N1 foam and PPNCN2 foam would establish the effect of slower crystallization on foam characteristics. It is evident
from the SEM micrograph of PPNC-N1 foam in Figure 3.11 that it has better cellular structure
compared to PPNC-N2 foam. The cell density for PPNC-N1 foam is six fold greater than that of
PPNC-N2 foam and the average cell size in PPNC-N1 foam is significantly lower as well. A
slower rate of crystallization of the polymer matrix in PPNC-N1 allows a greater rate of bubble
nucleation in the melt and also allows more uniform expansion of the bubbles. This is also
evident from the cell size distributions of the two nanocomposite foams presented in Figures
3.18 and 3.19. The cell size distribution of PPNC-N2 foam has 30% of cells greater than 200µm
in size with a small fraction of smaller cells whereas the cell size distribution for PPNC-N1 is
more uniform, indicating that greater surface for nucleation and slower rate of crystallization of
the polymer melt can definitely help in producing better foams. Even though the cellular
structure of PPNC-N1 is better than PPNC-N2 foam due to slower crystallization kinetics, the
final foam property of PPNC-N1 foam is not acceptable with a lot of larger size cells and a broad
cell size distribution.

57

Next, the effect of strain hardening in melt extensional flow may be seen by comparing
the cell size distribution of the foam produced with PPNC-S1 against that of the foam produced
with PPNC-N1. It is clear from Table 3.4 that PPNC-S1 resulted in foam with smaller average
6

3

cell size (86.5µm) and higher cell density (1.3x10 cells/cm ). The cell size distribution presented
in Figure 3.20 for PPNC-S1 also shows that the cells are mostly present in a narrow cell size
range of 40-100µm as opposed to the broad distribution in case of PPNC-N1. Since the clay
loading in PPNC-S1, the average rate of crystallization, amount of blowing agent used and the
pressure drop during foaming is very similar to that in PPNC-N1, the improvement in foam
quality with PPNC-S1 is solely due to the strain hardening of the polymer melt under uniaxial
extensional flow. The enhanced strain hardening behavior of the PPNC-S1 melt stabilizes the
cell walls and limits coalescence. The absence of strain hardening in PPNC-N1 resulted in cell
coalescence and led to larger cells.
Finally, the cell structures of foams produced with the three strain hardening
nanocomposites PPNC-S1, PPNC-S2 and PPNC-S6 may be compared. The three
nanocomposites differ only in clay and compatibilizer loading, however maintaining the strain
hardening property. All the three PPNCs show significant strain hardening behavior under
uniaxial melt extensional flow and have similar crystallization rates. From the SEM micrographs
in Figures 3.12 – 3.14 and the results in Table 3.4, it is evident that the three extruded foams
have comparable average cell sizes and average cell densities. The cell size distributions for
PPNC-S1, PPNC-S2 and PPNC-S6 nanocomposite foams (Figures 3.20 – 3.22) are also
comparable. This indicates that high quality foams can be produced with only 3 wt% of nanoclay
and using compatibilizer to clay ratio of 4, if the polymer nanocomposite melt has appropriate
rheological and thermal characteristics.
58

Further, foaming studies were done on a blend of linear and branched PP (L80B20) to
compare its results with those of linear PP/clay nanocomposites, in the same system. The
scanning electron micrograph of L80B20 extruded foam samples is presented in Figure 3.15 and
its cell size distribution is presented in Figure 3.21. The average cell size and cell size density is
also presented in Table 3.4. A comparison between PPNC-S6 and L80B20 foam samples clearly
shows that PPNC-S6 resulted in superior quality foams with smaller average cell size and higher
cell density. This is probably because of absence of any nanoparticles for heterogeneous
nucleation, despite the fact that the blend shows significant strain hardening characteristics. As
already established in our research along with other studies, presence of nanoparticles, serving as
additional nucleating sites for the cells, can significantly improve the foam characteristics.
Further, the faster rate of crystallization could also be a reason for poor foam characteristics of
the L80B20 foam. Hence, judicious selection of compatibilizer along with nanoclays can result
in an affordable approach for the production of superior quality linear PP foams compared to that
produced from blend of linear and branched PP.
Since the final objective of this research is to produce good quality TPO foams, foaming
studies were carried out on TPONC along with a control TPO. The SEM images of the two
foams are presented in Figures 3.16 and 3.17. From the micrographs it is clear that a strain
hardening PP component of the TPO can significantly improve the foam characteristics of the
TPO. Hence, the presence of strain hardening PP component and an elastomer can produce good
quality TPO foam that has necessary physical and mechanical properties to replace PU foams
from automotive applications.
Results presented in this section on foams extruded from polypropylene-clay
nanocomposites with a chemical blowing agent, show progressive improvement in the quality of

59

foam samples as one or more of the following features were present in the nanocomposites: high
surface area, slower crystallization and melt strain hardening in extensional flow. The maleated
polypropylene having low levels of maleic anhydride content and high molecular weight along
with slow crystallization kinetics led to the nanocomposites displaying slower crystallization
kinetics and strain hardening in melt extensional flow that are critical factors in foaming linear
semicrystalline polymers. This knowledge was further implemented to produce TPOs, showing
strain hardening behavior that further results in good quality closed cell foams. These results
demonstrate the effect of compatibilizer type on the thermal and rheological properties of the
polypropylene – clay nanocomposites. Further increases in melt strain hardening can be achieved
by silane treatment of the nanoclays used for making polypropylene nanocomposites, leading to
further improvement in the quality of extruded foams. This work is presented in the subsequent
chapters.

3.4. Conclusion
When several polypropylene clay nanocomposites all with good dispersion of the dialkyl
amine treated organoclay were foamed by extrusion with a chemical blowing agent, systematic
variations in foam quality were obtained based on differences among them in melt strain
hardening and in crystallization behavior. Even among nanocomposites without any melt strain
hardening, decreasing the rate of crystallization led to closed cell foams with smaller cell size
and higher cell density. Among nanocomposites where significant strain hardening was observed
in extensional flow along with the slower rates of crystallization, the extruded polypropylene
nanocomposite foams displayed the smallest cell sizes and the greatest cell density by reducing
cell coalescence. These enhancements may be attributed to attachment of high molecular weight
60

PP-g-MA chains to the surfaces of nanolayers due to hydrogen bonding, providing greater
constraints on chain stretch relaxation and also slowing down the crystallization.

61

Specimen

PP
6523
(wt%)

PPNC-N1

PP-g-MA (wt%)

Nanoclay
(wt%)

g-mol bound
MA / kg
organoclay

PF
814
(wt%)

Engage
8480
(wt%)

PP-gMA1

PP-gMA2

68.0

24

--

8.0

0.80

--

--

PPNC-N2

68.0

--

24

8.0

0.15

--

--

PPNC-S1

40.0

52.8

--

7.2

0.37

--

--

PPNC-S2

76.0

21.0

--

3.0

0.36

--

--

PPNC-S6

85.0

12.0

--

3.0

0.2

--

--

L80B20

80.0

--

--

--

--

20.0

--

TPO

70.0

--

--

--

--

--

30.0

TPONC

53.2

14.7

--

2.1

0.36

--

30.0

Table 3.1: Composition of various nanocomposites and polymer blends used in this study.

62

Specimen

Fractional
crystallinity

Time span tc
(min)

Average rate of
-1
crystallization (min )

PP

0.44

2.9

0.15

PP-g-MA1

0.34

3.9

0.09

PPNC-N1

0.35

5.0

0.07

PPNC-S1

0.36

5.1

0.07

PPNC-S2

0.38

5.0

0.07

PPNC-S8

0.39

5.1

0.08

PP-g-MA2

0.36

2.7

0.13

PPNC-N2

0.44

3.4

0.13

L80B20

0.42

3.9

0.11

Table 3.2: Crystallization parameters of various nanocomposites and polymer blends estimated
from non-isothermal crystallization study in a DSC.

63

Polymer
PPNC-N2

Surface
area
B

A

Average rate of
crystallization
S

E

C

Melt strain
hardening
A

S

E

PPNC-N1

=

_

= (None)

PPNC-S1

=

_

+

PPNC-S2

_

_

+

PPNC-S6

_

_

+

L80B20

x

=

+

(-): lower than base case, (+): higher than base case, (=): same as base case, (x): absent
Table 3.3: Comparison of various PP nanocomposites and polymer blends.

64

Polymer

Density
3
(g/cm )

Average Cell
Size (µm)

Average Cell
3
Density (no./cm )

Linear PP

0.4

--

--

PPNC-N2

0.4

170.9

0.1 x10

PPNC-N1

0.37

132.7

0.6 x 10

PPNC-S1

0.3

86.5

1.3 x 10

PPNC-S2

0.3

87.0

1.2 x 10

PPNC-S6

0.3

87.9

1.2 x 10

L80B20

0.35

112.1

0.8 x 10

TPO

0.43

--

--

TPONC

0.4

83.1

1.0 x 10

6
6
6
6
6
6

6

Table 3.4: Characteristics of various foam extrudates of polymer blends and polymer
nanocomposites foamed in a single screw extruder using 3 wt% of
azodicarbonamide.

65

Figure 3.1: Melt extensional viscosity of polymer blend of linear and branched PP mixed at two
different mixing conditions. The blend prepared at low mixing rpm show greater
strain hardening behavior.

66

Figure 3.2: X-ray diffraction patterns of various linear polypropylene–clay nanocomposites
along the X-ray diffraction pattern of the nanoclay.

67

Figure 3.3: Melt extensional viscosity of PP, PPNC-N1 and PPNC-N2 under uniaxial
extensional flow. None of the polymer melt show strain hardening behavior under
uniaxial extensional flow.

68

Figure 3.4: Melt extensional viscosity of PPNC-S1, PPNC-S2 and PPNC-S6 nanocomposites
under uniaxial extensional flow. All polymer melts shows significant strain
hardening phenomena under uniaxial extensional flow.

69

Figure 3.5: Strain hardening parameter of PPNC-S1, PPNC-S2 and PPNC-S6 nanocomposites.

70

Figure 3.6: Melt extensional viscosity of TPO and TPONC under uniaxial extensional flow.

71

Figure 3.7: Non-isothermal crystallization of PPNC-N1 and PPNC-N2 nanocomposites having
same wt% of different PP-g-MA and same organoclay.

72

Figure 3.8: Non-isothermal crystallization of PPNC-S1, PPNC-S2, PPNC-S6, PPNC-N2
nanocomposites and L80B20 polymer blend.

73

Figure 3.9: Scanning electron micrograph of linear polypropylene foam extrudate.

74

Figure 3.10: Scanning electron micrograph of PPNC-N2 foam extrudate.

75

Figure 3.11: Scanning electron micrograph of PPNC-N1 foam extrudate.

76

Figure 3.12: Scanning electron micrograph of PPNC-S1 foam extrudate.

77

Figure 3.13: Scanning electron micrograph of PPNC-S2 foam extrudate.

78

Figure 3.14: Scanning electron micrograph of PPNC-S6 foam extrudate.

79

Figure 3.15: Scanning electron micrograph of L80B20 foam extrudate.

80

Figure 3.16: Scanning electron micrograph of TPO foam extrudate.

81

500 µm
Figure 3.17: Scanning electron micrograph of TPONC foam extrudate.

82

Figure 3.18: Cell size distribution of PPNC-N2 foam.

83

Figure 3.19: Cell size distribution of PPNC-N1 foam.

84

Figure 3.20: Cell size distribution of PPNC-S1 foam.

85

Figure 3.21: Cell size distribution of PPNC-S2 foam.

86

Figure 3.22: Cell size distribution of PPNC-S6 foam.

87

Figure 3.23: Cell size distribution of L80B20 foam.

88

Figure 3.24: Cell size distribution of TPONC foam.

89

4. THE USE OF SILANE COUPLING AGENTS IN
POLYPROPYLENE CLAY NANOCOMPOSITES

4.1. Introduction
Strain hardening, as established in Chapter 3, is a very important rheological property of
polymers for foaming applications. However, absence of this property from linear PP has largely
restricted the use of such a potential polymer in the field of foaming technology. We have
successfully shown, in Chapter 3, that addition of suitable nanoclay and compatibilizer to linear
PP can result in significant strain hardening characteristics under uniaxial extensional flow,
which in turn led to significant improvement in foam quality.
Strain hardening in the nanocomposite melts is analogous to that reported for blends
containing small amounts of crosslinked polymer. Blends of linear polypropylene with small
amounts of chemically crosslinked polymers contain trapped entanglements on chain segments
between crosslinks or junctions which respond to stretching differently from entanglements on
chain segments in the bulk [152, 153]. Linkages or interactions between the clay surface or edge
and the maleated polymer can form a network of “physical crosslinks” or anchor points between
which entanglements may be trapped to produce similar effect. However, organically modified
nanoclays and the compatibilizer has mainly electrostatic or hydrogen bonding type interactions
[154], which are weaker forces of attraction. Hence, if these weaker forces of attraction can be
replaced with covalent bonds then one can expect a significant increase in the strain hardening
behavior of these PP nanocomposites. With the objective of producing PP nanocomposite foams
with even smaller cell sizes compared to PPNC-S6 foams, by an improvement in the strain
hardening characteristics, additional treatment of the nanoclays was investigated.
90

Over the past few decades, extensive research has been done on the use of
montmorillonite clay as nanofiller to prepare polymer nanocomposites. The potential for
significant improvements in various mechanical, barrier and rheological properties [69, 73-75,
88, 155, 156] upon the used of layered silicates in various polymer matrices has stimulated the
growth of polymer nanocomposites. The extent of increment in these properties was attributed
not only to dispersion of these nanolayers but also to the strong polymer-clay interactions [157160]. However attempts to disperse the clay in non-polar systems like polypropylene still remain
a challenge. This is primarily due to absence of favorable interfacial interactions between the
nanoclay and non-polar polymers like polypropylene. In order to increase interfacial interactions
between the two, functionalized polyolefin oligomers (compatibilizers) are commonly used [78,
86, 87, 89-91, 161, 162], where the polar functionality can favorably interact with the clay
surface and the oligomers can interact with the bulk matrix. In order to achieve high levels of
interaction between the clay and the polymer, a high level of maleic anhydride grafting is
required, which causes chain scission in polypropylene [133] and results in poor physical and
mechanical properties of the resulting nanocomposites.
Since montmorillonites has exchangeable interlayer cations and shows Bronsted and
Lewis acidity, the most commonly accepted way to improve the polymer-clay interaction was to
treat the nanoclays with various cationic surfactants [77-81]. Alternatively, due to the presence
of hydroxyl groups on the edges of the clay platelets and due to active sites on the surface of the
clay platelets, these nanoclays have also been treated with various silane-coupling agents to
improve polymer-clay interaction [160, 163-169]. With the use of silated clays, researchers have
shown significant improvements in clay dispersion in the polymer matrix [164, 170, 171],
resulting in improvements in various mechanical [166, 172, 173] and rheological properties [174,

91

175]. However, none of the work reported till date discusses changes in extensional rheology
characteristics by the use of silane treated nanoclays.
The properties and effects of silanes are determined by their molecular structures,
generally represented as shown in Figure 4.1. The silicon at the center has an organofunctional
group (Y) at one end and hydrolysable groups (most commonly alkoxy groups) on the other end.
A judicious selection of the silane-coupling agent can act as a bridge between the polymer and
the filler, where the alkoxy groups (after hydrolyzing) can link with the hydroxyl groups at the
edges of the clay or with the active sites on the clay surface and the organofunctional group, Y,
can interact with the polymer. The latter interaction can be in the form of covalent bonding or the
formation of interpenetrating networks [176] for non-reactive systems like polypropylene.
Depending upon the molecular structure of the silane coupling agents, it can end up either being
inside the layered clay galleries [83, 165] resulting in an increase in clay basal spacing or react
mainly with the edge hydroxyl groups [135, 165, 177, 178]. Shanmugharaj et al. [165] studied
the influence of dispersing medium, used during the silation procedure, on the grafting of
montmorillonite clays and reported that with the increase in solvent surface energy, the
possibility of intercalation of silane molecules within the clay galleries increases. The presence
of silane molecules within the clay galleries could provide greater interaction between the
polymer and the clay, resulting in better dispersion and improved mechanical properties.
In this work, we have investigated the effects of increased interaction using carefully
chosen silane coupling agents on the dispersion of nanoclays in linear polypropylene
nanocomposites and the associated improvements in strain hardening in uniaxial extensional
rheology of the polymer melt. In order to replace the weaker force of interaction (hydrogen
bonding) between the compatibilizer and the nanoclay with a stronger covalent bond, the

92

organoclays were additionally treated with carefully chosen amino silanes. This selection was
made based on the fact that amines present on the silane treated clays can react with the
anhydride group of the PP-g-MA [179-183], forming an amide or imide linkage as per the
reaction scheme shown in Figure 4.2. Lu et al. [181] investigated the reaction between PP-g-MA
and amines by melt mixing them in various proportions in an extruder. The study suggested that
the reaction between the two was very fast and was completed within 90s which was supported
by FTIR analysis. Song et al. [183] studied the reaction between different amines (primary,
secondary and tertiary amines) and anhydride group in polymeric systems and reported rapid
formation of the imide linkage between the polymer and the amines. Increased interaction
between the compatibilizer and the silated clay would further slow the relaxation of stretch in the
polymer chains trapped between particles along the direction of flow, under uniaxial extensional
flow at melt state, compared to the unsilated polypropylene nanocomposite system.

4.2. Materials and Experimental Procedures
Three different amino silanes were used for this study:
a) N-aminoethyl-Aza-2,2,4-trimethylsilacyclopentane (Diamino monomethoxy-Silane)
b) N-(2-aminoethyl)-3-aminopropylmethyldimethoxysilane (Diamino dimethoxy-Silane)
c) N-(2-aminoethyl)-3-aminopropyltrimethoxysilane (Diamino trimethoxy-Silane)
d) Phenyl trimethoxy silane.
The chemical structures of all the silanes are presented in Figure 4.3. The main
differences among the three amino silanes are in the number of hydroxyl groups, which could be
significant if poly-condensation of silanes were more likely to occur with the large number of
hydroxyl groups. However, since the silation reaction is carried out under constant stirring and
93

also due to small amounts of silane used for the clay treatment chances of poly-condensation of
silanes is less likely. Phenyl trimethoxy silane is used as a control silane coupling agent, as it
does not have amine functionality. The methoxy groups of all the silanes get hydrolyzed in
presence of moisture and can then follow the reaction scheme as shown in Figure 4.2 (a similar
reaction scheme has also been proposed by Zheng et al. [184] and Nygard et al. [185]). However,
the monomethoxy-silane follows a ring opening mechanism during the hydrolysis step as shown
in Figure 4.3 and then follows the usual silane reaction scheme. The detail of the silation
procedure is laid out in Chapter 2.
Two different organically modified montmorillonite clays (I.44P and I.30P) were used
for the silane treatment. The silated clays were then used to prepare nanocomposites with 85
wt% linear PP, 12 wt% PO 1015 as the compatibilizer and 3 wt% of nanoclay, the details of
which are presented in Table 4.1. The various techniques used in this section for characterization
of nanoclays and nanocomposites are: XRD, TEM, TGA, BET, and uniaxial extensional
rheology. The details of all the experimental procedures are presented in section 2.

4.3. Results and Discussion
4.3.1. Characterization of Silated Clays
Thermo-gravimetric analysis, X-ray diffraction studies, and BET surface area
measurements were used to characterize the various silane treated organoclays. The information
from these techniques was then used to predict whether silation of the organoclays resulted in
edge grafting or face grafting.

94

4.3.1.1. Thermo-gravimetric analysis
Thermo-gravimetric analysis was used to estimate the organic content in the organoclays,
silated and unsilated. The TGA weight loss profile for silated and unsilated I.44P is presented in
Figure 4.4 and the same for I.30P clays are presented in Figure 4.5. To begin with, unsilated
I.44P has 35 wt% of organic content while I.30P has 24 wt% of organic content, however, they
both have same molecular equivalent of surfactant inside the clay galleries. The weight loss
profile (up to 500°C) for silated I.44P organoclay shows negligible change in the amount of
organic content after silation irrespective of the silane used, whereas the same for I.30P
organoclays, there is an increase in the organic content after silation (3.5 wt% for mono-silation
case and 7.3 wt% for diamino dimethoxy-silation case). This indicates that the silane molecules
have most probably intercalated inside the clay galleries for I.30P system whereas insignificant
change in the organic content of silated I.44P system indicated that the silane molecules are most
grafted on the edges. In order to further substantiate the grafting sites of the silane molecules in
the two organoclay system, BET surface area measurements and XRD was carried out.

4.3.1.2. BET surface area measurements
The nanoclays before and after silation were analyzed for their total surface area using
BET surface area measurements. The total BET surface area of sodium montmorillonite clays is
in the range of 43 m2/g to 50 m2/g [186-188]. It has been reported by He et al. [189] that with
progressive increase in adsorption of hexadecyltrimethyl ammonium surfactant inside the clay
galleries, the BET specific surface area, measured by nitrogen adsorption, of the resulting
organoclay decreases progressively (Figure 4.6). The BET specific surface area experimental
results for all the silated and unsilated organoclays are tabulated in Table 4.2. The BET specific

95

surface area of both the organoclays, I.30P and I.44P, is around 10 m2/g after surfactant
treatment, thus indicating that both these nanoclays are cationically exchanged with similar
molar equivalent of the cationic surfactant. A drop in BET surface area, from sodium
montmorillonite to organically modified montmorillonite, also indicates that a fraction of the
clay surface is occupied by bulkier ions. When I.44P was further treated with different silane
coupling agents, there was no significant drop in the BET surface area irrespective of the silane
used, indicating that silation of the clay has mainly occurred at the edges of the nanoclay
platelets. However, there is a significant drop in the BET surface area of the silane treated I.30P
nanoclays, indicating that silane coupling has occurred both at the edges and the faces of the
nanoclay platelets. This is probably due to different molecular structure and subsequently
different molecular conformations of the surfactants used in the two nanoclays, even though both
I.30P and I.44P are intercalated with same molar equivalents of the organic cation. Further, I.30P
has 24 wt% of surfactant while I.44P has 35 wt% of surfactant. Hence due to densely packed
conformation of I.44P, the silane coupling reaction happens predominantly at the edges while
that for I.30P is both at the edges and the faces.
The difference between surface area of diamino monomethoxy-silane treated I.30P and
diamino dimethoxy-silane treated I.30P could be explain by the fact that in case of diamino
dimethoxy-silane there are two methoxy groups that can graft to neighboring sites thereby
reducing the surface area available for nitrogen adsorption. Also due to the cyclic structure of the
diamino monomethoxy silane, the intercalation of the silane molecule might be sterically
hindered. However, there is no further drop in the surface area between diamino dimethoxysilane treated and diamino trimethoxy-silane treated I.30P. This is probably because due to the
tetrahedral conformation of the three methoxy groups on the silane atom making it sterically

96

difficult to adsorb at three neighboring sites. A very small reduction in the surface area of the
phenyl trimethoxy silane could be explained by the fact that the bulky nature of the phenyl group
makes it difficult for the silane molecule to enter the galleries [135].

4.3.1.3. X-ray diffraction study
X-Ray diffraction studies were conducted on these silane treated clays for the
characterization of their morphology. The analysis of TGA and BET surface area results
indicated that silation of I.44P clay resulted in edge grafting while that of I.30P clay resulted in
face grafting. This should result in an increase in d-spacing of I.30P nanoclay post silation and
the same for I.44P clay system should remain unchanged. The XRD patterns of silated I.30P clay
and silated I.44P clay are presented in Figure 4.7 and the corresponding d-spacings are presented
in Table 4.3. From the XRD pattern of the treated I.30P clay system, it is evident that the dspacing for all the amino silated clays have increased compared to the unsilated I.30P except
phenyl trimethoxy silane. This confirms that the amino silane molecules have resulted in face
grafting while phenyl trimethoxy silane is predominantly face grafted. However, the d-spacing
for I.44P clays remain unchanged after silation, indicating no possible intercalation of the silane
molecules inside the clay galleries and the silane molecules have been grafted to the edges of the
clay.
If the proposed hypothesis of stronger interaction between nanoclay and the
compatibilizer resulting in higher strain hardening characteristics is true, then nanocomposites
derived from amino silane treated nanoclays would result in greater strain hardening behavior
compared to untreated or phenyl trimethoxy treated nanoclay. Furthermore, nanocomposites
prepared using amino silated I.30P nanoclays should show even higher strain hardening

97

characteristics compared to those prepared from amino silated I.44P nanoclays, due to greater
number of sites for interaction with PP-g-MA.

4.3.2. Characterization of Polypropylene – Silated Clay Nanocomposites
4.3.2.1. Structural characterization
Various linear PP nanocomposites were prepared using 3 wt% of silated I.30P and I.44P
organoclays using PP-g-MA1 as the compatibilizer. The objective here is to study the effect of
different silation on the morphology and strain hardening behavior of the nanocomposites.
The X-ray diffraction pattern for various PP nanocomposites derived from silated I.30P
and I.44P nanoclays are presented in Figure 4.8 and corresponding d-spacing in Table 4.3. The
XRD pattern of silated I.44P nanocomposites shows negligible change in d-spacing and the
extent of dispersion (evident from the distinct peak in the XRD pattern) due to silation of the
organoclays. The TEM micrographs presented in Figure 4.9 also shows that nanocomposite
prepared from untreated I.44P as well as diamino dimethoxy silane treated I.44P have
intercalated structure with comparable thickness of clay stacks. This was expected because the
silane coupling has primarily occurred at the edges of the I.44P nanoclays. However the same is
not true for nanocomposites prepared form silated I.30P clay system. All the amino silane treated
I.30P clay nanocomposites showed an increase in d-spacing, indicating greater polymer
intercalation compared to untreated I.30P. The d-spacing for phenyl trimethoxy silated I.30P clay
nanocomposites is comparable to the unsilated I.30P clay. Also the XRD patterns of the
nanocomposites prepared from amino silane treated I.30P lacks distinct peaks, an indication of
random well dispersed clay platelets as compared to ordered clay stacks in case of untreated
I.30P and phenyl trimethoxy silane treated I.30P system. This is also evident from the TEM
98

micrographs for untreated I.30P and diamino dimethoxy silane treated I.30P nanocomposites
presented in Figure 4.10. Hence, the face grafting of the amino silanes in I.30P clay system
resulted in greater intercalation of the polymer chains leading to better dispersion, as opposed to
the edge grafting of amino silanes in I.44P clay system, which showed no appreciable change in
the extent of polymer intercalation or dispersion of nanoplatelets.

4.3.2.2. Uniaxial Extensional Rheology
The melt extensional viscosity transients for the linear PP nanocomposites prepared from
various I.30P nanoclays and I.44P nanoclays are presented in Figures 4.11-4.13. The melt
extensional viscosity transients for phenyl trimethoxy silated I.30P nanocomposites along with
unsilated I.30P nanocomposites are presented in Figure 4.11. As expected, the transient melt
extensional viscosity data of PP nanocomposites prepared with phenyl trimethoxy silane treated
I.30P nanoclay is similar to that of untreated I.30P nanoclay, indicating no improvement in the
strain hardening behavior. This is because presence of an additional phenyl group attached to the
clay surface does not change the nature of interaction between the compatibilizer and the clay.
However, as shown in Figures 4.12 and 4.13, all the nanocomposites prepared using amino silane
treated I.30P and I.44P shows a significant increase in the strain hardening behavior compared to
the unsilated system clay system. This is because the weaker hydrogen bonds between the clay
and the compatibilizer are now replaced with much stronger covalent bonds, which can help
orient the polymer chains to a greater extent and thus resulting in significant increase in strain
hardening characteristics. Moreover, there is no significant difference between the strain
hardening characteristics of clays treated with amino silanes with one or more hydrolysable
groups. This proves that the increase in polymer clay interaction is primarily due to the presence

99

of amine functionality on the silated clays. This also indicates that poly-condensation of the
silanes does not appear to be significant.
A comparison between amino silane treated I.30P and I.44P nanocomposites indicates
that I.30P system shows greater increase in strain hardening behavior. This is because, in case of
silated I.30P clays the silane molecules are present both on the edges as well as on the faces of
the nanoclay, which provides greater number of interaction sites between the polymer and the
clay surface as opposed to only edge functionalization in case of silated I.44P system. In order to
verify that the samples were stable during these tests and that the strain hardening data is not an
artifact of the continuously changing structure of the nanocomposite, dynamic time sweep tests
were conducted on these samples. The time sweep data (Figure 4.14) shows that the storage
modulus of these nanocomposite melts remained unchanged over a period of 1 hour, confirming
that the samples were stable over an hour.
For a quantitative evaluation of the strain hardening behavior of these nanocomposites,
the strain hardening parameters are calculated using equation 3.1, for a Hencky strain of 2.25,
and are presented in Figures 4.15-4.16. The choice of Hencky strain of 2.25 was done so as to
compare the strain hardening behavior of all the nanocomposites at the same strain and for few
nanocomposites, maximum achievable Hencky strain was 2.25. From the strain hardening data
for I.44P clay system (Figure 4.15), the significant change in strain hardening behavior of
nanocomposites prepared using amino silane treated I.44P is clearly reflected. However, it seems
as if there is not much of an improvement in strain hardening behavior of the I.30P system after
silation (Figure-4.16), even though the melt elongational viscosity transients presented in Figure
4.12 says otherwise. A closer look at the strain hardening data indicates that the deviation from
the linear viscoelastic curve starts at a higher strain for amino silane treated I.30P system and

100

continue to show an increase in the extensional viscosity under uniaxial extensional flow up to
the maximum value of Hencky strain used during the test, compared to unsilated I.30P
nanocomposites where the extensional viscosity starts to drop after a Hencky strain of 2.25.
Figure 4.16 also shows the strain hardening parameter of nanocomposites containing amino
silane treated I.30P clays at higher values of Hencky strain. It is evident from Figure 4.16 that the
strain hardening parameter (at εH = 2.9) of amino silated I30P nanocomposites is greater than
that measured at εH = 2.25 at all strain rates tested. While all other nanocomposites (including
silated I.44P system) shows a maximum value of extensional viscosity close to a strain of 2.25,
amino silane treated I.30P nanocomposites continues to show an increase in extensional viscosity
up to the maximum Hencky strain tested. This clearly demonstrates the superiority of the amino
silane treated I.30P nanocomposites over silated I.44P clay nanocomposites in withstanding the
uniaxial stretching force up to a larger Hencky strain at all rates tested. Another important
observation of these nanocomposites is the nature of change in the strain hardening parameter
with increase in strain rates. The I.44P system even after silation shows a drop in strain
hardening parameter with increase in strain rates (Type – 2 strain hardening behavior) while the
amino silated I.30P system is able to maintain its strain hardening characteristics with increase in
strain rates (Type – 1 strain hardening behavior) (Figures 4.15-4.16). Since the instrument used
to measure the extensional viscosity transients under uniaxial extensional flow, has a testing
limitation at higher strain rates, one can assume that amino silated I.30P nanocomposites will
continue to show strain hardening properties even at higher strain rates as opposed to amino
silated I.44P nanocomposites. This allows us to study the effect of type of strain hardening
behavior on foamability of polymer system.

101

4.3.2.3. Effect of solvent used during silation on final PP/clay nanocomposite
property
In order to study the influence of solvent surface energy on the coupling of the silane
molecule on the organoclay, the three different solvent systems were used to silate I.44P clay
using diamino dimethoxy-silane. They are:
(a) A mixture of 80 wt% methanol and 20 wt% de-ionized water,
(b) A mixture of 80 wt% methanol and 20 wt% de-ionized water, and
(c) De-ionized water.
The resulting diamino dimethoxy-silated I.44P clays were then used to prepare PP/clay
nanocomposites with 3 wt% clay and 12 wt% PP-g-MA1. These nanocomposites were then
tested for their strain hardening behavior under uniaxial extensional flow. The melt extensional
viscosity data for the three cases are presented in Figure 4.17. It can be seen from the figure that
silated clays synthesized using different solvent system has negligible effect on the strain
hardening behavior of the resulting I.44P clay nanocomposites, indicating that the silane grafting
in still predominantly on the edges of the clay. Hence, the choice of solvents, depending on
solvent surface energy, for silation does not influence the grafting location of silane molecules
for our system. This primary objective of this experiment was based on industrial interests. Since
aqueous suspension of sodium montmorillonite is used for surfactant system, hence the moist
clay cake obtained after surfactant treatment can directly be used for further silane treatment.

102

4.4. Conclusions
Results presented in this section reveals that carefully chosen silane-coupling agents can
significantly improve the interfacial interactions between the nanolayer and the compatibilizer.
Depending on the precursor nanoclays, the silane treatment can predominantly result in either
face or edge grafting. Since the available volume within the clay galleries in greater in I.30P
organoclays due to lower organic content (wt%) compared to I.44P clay, former system results in
face grafted silated clay while the later results in edge grafting of the silane molecules. This is
been convincingly proven by various characterization techniques. The linear PP nanocomposites
prepared using amino silane treated nanoclays resulted in significant increase in strain hardening
behavior under uniaxial extensional flow compared to the unsilated clay nanocomposite,
irrespective of the organoclays used. Phenyl trimethoxy silated I.30P, used as control, showed no
improvement in strain hardening behavior compared to the unsilated I.30P. Further, since I.30P
resulted in face grafted silated clay, the resulting nanocomposites showed greater strain
hardening, up to higher Hencky strains, compared to silated I.44P clays. This is due to greater
number of interaction points available on the silated I.30P clay because of face grafting as
opposed to edge grafting of the silated I.44P clay system. Another critical observation was the
nature of change in strain hardening behavior with the change in strain rate. Silate I.30P
nanocomposites resulted in more LDPE type strain hardening behavior that remain unchanged
with strain rates while silated I.44P system behaved more like blends of linear and branched PP.
These results demonstrate the effect of polymer-clay interaction on strain hardening behavior of
nanocomposites under uniaxial extensional flow. It will further be shown that these
nanocomposites with superior strain hardening property can result in closed cell foams with even
small size and greater cell densities compared to those obtained by using the unsilated clays.
103

Specimen

Clay treatment

PPNC-S3

Diamino monomethoxy-silane treated I.30P

PPNC-S4

Diamino dimethoxy-silane treated I.30P

PPNC-S5

Diamino trimethoxy-silane treated I.30P

PPNC-S10

Phenyl trimethoxy silane treated I.30P

PPNC-S11

Unsilated I.30P

PPNC-S7

Diamino dimethoxy-silane treated I.44P

PPNC-S12

Diamino monomethoxy-silane treated I.44P

Table 4.1: Formulation details of linear PP/ clay nanocomposites with 85 wt% Profax 6523, 12
wt% of PO 1015, and 3 wt% to treated clays.

104

2

Surface Area (m /g)

Treatment

I.30P

I.44P

No silane treatment

9.51

9.8

Diamino monomethoxy-silane

6.63

10.13

Diamino dimethoxy-silane

3.82

9.12

Diamino trimethoxy-silane

3.67

9.31

Phenyl trimethoxy silane

7.71

9.25

Table 4.2: BET surface area measurements of silane treated nanoclays.

105

I.30P

Treatment

I.44P

Clays

PPNCs

Clays

PPNCs

No silane treatment

2.34

2.95

2.58

3.16

Diamino monomethoxy-silane

2.76

3.33

2.54

2.99

Diamino dimethoxy-silane

3.21

3.33

2.54

3.33

Diamino trimethoxy-silane

3.1

3.16

--

--

Phenyl trimethoxy silane

3.21

2.27

--

--

Table 4.3: Basal spacing (d-spacing) of various treated nanoclays and their corresponding
nanocomposites prepared using 85 wt% linear PP and 12 wt% PO 1015.

106

Y – (CH2)n – Si – X3
Organofunctional Linker Silicon Hydrolysable
Group
atom
Groups

Figure 4.1: General structure of a silane coupling agent.

107

Figure 4.2: Reaction scheme between silane treated clays and maleic anhydride grafted
polypropylene.

108

(a)

H2O
Methanol

(b)

OCH3

(c)

(d)

Figure 4.3: Chemical structure of silane coupling agents (a) Diamino monomethoxy-silane, (b)
diamino dimethoxy-silane, (c) diamino trimethoxy-silane, and (d) Phenyl
trimethoxy silane. The ring opening hydrolysis of diamino monomethoxy-silane is
also shown along with its chemical structure.

109

Figure 4.4: TGA profile of various silane treated I.44P clays.

110

Figure 4.5: TGA profile of silane treated I30P clays.

111

Figure 4.6: BET surface area data from He et al. [189] on reduction in surface area due to
surfactant adsorption.

112

Figure 4.7: X-ray diffraction pattern of silane treated I.44P clays (top) and silane treated I.30P
clays (bottom). The treatment of clays is represented as (a): Diamino
monomethoxy-silated clay, (b) Diamino dimethoxy-silated clay, (c) diamino
trimethoxy-silated clay, and (d) Phenyl trimethoxy silated clay. The d-spacing of
the corresponding organoclays are presented alongside the XRD curve.

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Figure 4.8: X-ray diffraction pattern of linear PP nanocomposites with silane treated I.44P
clays (top) and silane treated I.30P clays (bottom). The treatment of clays is
represented as (a): Diamino monomethoxy-silated clay, (b) Diamino dimethoxysilated clay, (c) diamino trimethoxy-silated clay, and (d) Phenyl trimethoxy silated
clay. The d-spacing of the corresponding organoclays are presented alongside the
XRD curve.

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200 nm

200 nm
Figure 4.9: TEM micrographs of linear PP nanocomposites with untreated I.44P clay (PPNCS6) (top) and with diamino dimethoxy silane treated I.44P clay (PPNC-S7)
(bottom).

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500 nm

200 nm
Figure 4.10: TEM micrographs of linear PP nanocomposites with untreated I.30P clay (PPNCS11) (top) and with diamino dimethoxy silane treated I.30P clay (PPNC-S4)
(bottom).

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Figure 4.11: A comparison of melt extensional viscosity of PPNC-S10 (phenyl trimethoxy
silated I.30P) and PPNC-S11 (unsilated I.30P) nanocomposites under uniaxial
extensional flow.

117

Figure 4.12: Effect of amino silane treatment of I.30P clay on the melt extensional viscosity of
linear PP nanocomposites under uniaxial extensional flow. The number of
hydrolysable groups has negligible effect on the strain hardening behavior of the
nanocomposites.

118

Figure 4.13: Effect of amino silane treatment of I.44P clay on the melt extensional viscosity of
linear PP nanocomposites under uniaxial extensional flow.

119

Figure 4.14: Dynamic time sweep test at 180°C for various PP nanocomposites along with its
matrix (PP+compatibilizer).

120

Figure 4.15: Variation of strain hardening parameter with strain rate for PPNCs prepared with
silated I.44P clays.

121

Figure 4.16: Variation of strain hardening parameter with strain rate for PPNCs prepared with
silated I.30P clays. Open symbols are for Hencky strain = 2.9 and closed symbols
are for Hencky strain = 2.25.

122

Figure 4.17: Effect of solvent used for silation process on melt extensional viscosity of PPNCS7 nanocomposites.

123

5. CONTINUOUS FOAMING OF POLYPROPYLENE CLAY
NANOCOMPOSITES HAVING DIFFERENT STRAIN
HARDENING CHARACTERISTICS
5.1. Introduction
It is well established, as well as demonstrated in Chapter 3 that small cell and high cell
density closed cell polymer foams can be produced only with melts showing strain hardening
under uniaxial extensional flow [58, 100, 105, 106]. Linear PP displays lack of strain hardening
under uniaxial extensional flow, hence foaming of linear PP, even in the presence of
nanoparticles, have led to inferior quality foams [41, 42, 48, 121, 125]. Owing to its unique
physical and mechanical properties along with its recyclability and affordability, linear PP foams
can serve as an immediate replacement to market dominating polyurethane or polystyrene foams.
TPOs, a blend of linear PP and an elastomer, have all the required properties necessary to replace
the non-recycling polyurethane foams from automotive industry. Since linear PP is a major
component of these TPOs, foaming of such materials has also been largely restricted due to its
weak tensile melt strength. Various approaches have been adopted to increase the tensile melt
strength of linear PP, for example, long chain branching [48, 58, 59, 116], cross-linking [44, 111,
112, 190] to eventually improve its foamability, but these approaches have strong impact on
economy of foam production. As presented in Chapter 3, the novel approach of using organically
modified montmorillonite nanoclays and high molecular weight PP-g-MA as the compatibilizer
resulted in significant increase in the melt strain hardening behavior of linear PP
nanocomposites. These strain hardening nanocomposites were tested for foaming performance
and was compared with a blend of linear and branched PP. The results indicated that, even
though both showed significant melt strain hardening under uniaxial elongational flow, due to
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the presence of nanoparticles, serving as both strain hardening promoter and heterogeneous
nucleating agent, linear PP nanocomposites resulted in foams with much smaller cell size and
higher cell densities compared to the blend. The objective of this work is to study the extent and
nature of melt strain hardening on the final foam properties.
As presented in Chapter 4, linear PP/clay nanocomposites prepared using an amino silane
treated organoclays led to significant improvements in melt strain hardening under uniaxial
extensional flow. This novel approach for improvements in strain hardening provided an
opportunity to evaluate the effects and nature of strain hardening on foamability of linear PP.
The nature of strain hardening of polymer melts can be classified into two types. In one
type (labeled Type – 1 here), the strain hardening increases or remains constant with increasing
strain rate, e.g. branched polypropylene and low density polyethylene [191, 192]. In the other
type (labeled Type - 2 here) strain hardening decreases with increasing strain rate, e.g. blends of
linear and branched polypropylene [58, 116]. Stange et. al. [116] found that blends containing up
to 50 wt% of long chain branched PP with linear PP showed a drop in strain hardening with
increase in strain rates. This type of dependence of strain hardening on strain rates is also
observed for metallocene-catalyzed ethylene–α
 olefin random copolymers with “sparse longchain branching’ [193] and with sparse long-chain branched PP [194].
Spitael and Macosko [59] have reported that larger cell sizes were obtained upon
extrusion foaming tree-type branched PP (using carbon dioxide as the physical blowing agent)
than with blends that displayed strain hardening decreasing with increasing strain rates. They
attributed this to poorer crystallizability of the highly branched polypropylene. Results presented
in Chapter 2, also show that crystallization kinetics indeed plays a significant role in the
continuous foaming process of polypropylene, a semi-crystalline polymer. A slowly crystallizing
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polymer allows the cells to develop before crystals starts growing compared to the polymer melt
that crystallizes at a faster rate. Hence, in order to study the effect of strain hardening type or
extent of strain hardening, it must be ensured that the crystallization kinetics of the polymers
studied are comparable. Since the strain rate at which the foam cells are growing during foaming
process is not well known, an attempt is made here to study the effect of change in strain
hardening with strain rates on foamability of polymers.
In Chapter 2, the best PP nanocomposite foams produced in a continuous extrusion
process using a CBA had an average cell size of 6.7 mm with an average cell density of 1.3x106
cells/cc. The objectives of this work are to compare foams produced by extrusion with a
chemical blowing agent from two different polypropylene nanocomposites with different types
of strain hardening characteristic.

5.2. Materials and Experimental Procedures
Three different nanocomposites were used for this study, PPNC-S4, PPNC-S6 and
PPNC-S7, all having 3 wt% of different organically treated montmorillonite clays and 12 wt% of
the same compatibilizer, PP-g-MA (Exxelor PO 1015). PPNC-S6 is prepared with I.44P
organoclay while PPNC-S7 is prepared with diamino dimethoxy silane treated I.44P and PPNCS4 was prepared with diamino dimethoxy silane treated I.30P. These nanocomposites were
characterized in detail and the results are presented in Chapter 3 and 4. These nanocomposites
were foamed in a single screw extruder using 3 wt% of the CBA. PPNC-S6 was foamed using a
2mm diameter die with L/D = 15 while the other two nanocomposites were foamed using a
higher L/D = 20. The first three temperature zones were maintained at the same temperature for
all the three nanocomposites (180 ºC, 200 ºC and 175 ºC respectively), however the die exit
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temperature was maintained at 15 ºC above the crystallization temperature for PPNC-S4 and
PPNC-S7 while the same for extruding PPNC-S6 foam was maintained at 35 ºC above the
crystallization temperature. This was done so as to maintain an overall constant pressure drop (~
4000psi) across the die. The resulting foam extrudates were characterized for the foam cellular
structure using SEM. The details of all the procedures are presented in Chapter 2.

5.3. Results and Discussion
Nonisothermal crystallization exotherms obtained from cooling ramps in DSC tests have
been compared for PPNC-S4, PPNC-S6 and PPNC-S7 in Figure 5.1. The DSC exotherms show
that the nature of crystallization (crystallization temperature and the rate of crystallization) is
similar for all the three PP/clay nanocomposites; hence any changes in the foaming behavior of
these nanocomposites cannot be attributed to differences in crystallization behavior.
The nanocomposites studied in this section had two different type of strain hardening
behavior. PPNC-S4 showed more like Type – 1 strain hardening, while PPNC-S6 and PPNC-S7
had Type – 2 strain hardening characteristics, with PPNC-S7 having greater strain hardening
parameter than PPNC-S6. These are presented in details in Chapter 4. The thermal characteristics
of these nanocomposites are comparable which is in contrast to different blends of linear and
branched PP. In blends of linear and branched PP, as the proportion of linear PP is increased, the
strain hardening characteristics changes from Type – 1 to Type – 2. Naguib and coworkers [195]
have studied the effect of blends of linear and branched PP on foamability. The crystallization
temperature of the blend increases with the increase in branched resin and they reported that the
optimum temperature for producing foams with maximum volume expansion ratio decreases
with increase in branched resin. In this work, we have successfully prepared nanocomposites
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with different strain hardening characteristics while keeping the thermal characteristics
comparable, in order to study the effect of different type and extent of strain hardening
characteristics on foaming.
Various theoretical and experimental studies [196-199] have shown the evolution of foam
cell size with time. Figure 5.2 shows the nature of the foam cell evolution with time. Cell size
growth with time profile indicates that the strain rates (slope of the curve) experienced by
smaller cells are much higher than the rates at which the strain hardening measurements are
conducted. For an extrusion foaming set-up, the highest strain rates are generally experienced
within the diverging section of the die, represented as Zone I. Immediately after the foam
extrudate exits the die region, the polymer starts to crystallize and arrest further growth of cell
size. This is represented as Zone II in Figure 5.1. Hence, in order to produce foams with cell
sizes in the range of microcellular regime, the polymer melt should show strain hardening
behavior at higher strain rates. With the nature of Type – 1 strain hardening behavior, it is
possible that the polymer melt would show higher strain hardening behavior inside the die region
(zone I) compared to Type – 2 polymer melt. This would result in more controlled and smaller
cell size at the end of zone I for Type – 1 strain hardening melt compared to Type – II melt.
Since the nanocomposites studied in this section are based on same compatibilizer and linear PP,
time taken for the melt to crystallize after exiting zone I would be similar which is also evident
from the DSC exotherms presented in Figure 5.1. Based on this hypothesis, it should be expected
that PPNC-S4 should result in foams with smallest possible cell size within the limitation of the
experimental conditions.
The SEM images for the cellular structures of polypropylene-clay nanocomposite foam
samples are shown in Figures 5.3 – 5.5, and the cell size distributions are presented in Figures

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5.6 – 5.8. A comparison between the SEM micrographs of PPNC-S6 and PPNC-S7 shows that,
PPNC-S7 results in foams with smaller cells and greater cell density as compared to PPNC-S6.
6

3

The average cell size and cell density of PPNC-S6 foam is 87.9 µm and 1.2x10 cell/cm
6

3

respectively while that for PPNC-S7 foam is 60.98 µm and 2.7x10 cells/cm . Even though both
these nanocomposites have same amount of nanoclays and Type – 2 strain hardening
characteristics, PPNC-S7 shows higher strain hardening behavior, thereby resulting in foams
with smaller cell size and higher cell density. This is presumably due to greater strain hardening
parameter of PPNC-S7 compared to PPNC-S6, in the range of strain rates encountered during
cell growth, arrested the growth of foam cells and restricted the cell size to lower values.
In order to compare the effect of strain hardening type, SEM images of PPNC-S4 and
PPNC-S7 foams are compared. Both have similar strain hardening parameter in the range of
strain rates tested, however PPNC-S7 shows a gradual drop in the strain hardening parameter
with increase in strain rates, while PPNC-S4 shows an almost constant strain hardening
parameter. The average cell size and cell density of PPNC-S4 nanocomposite foam is 37.09µm
6

3

and 6.2x10 cells/cm respectively. This shows that polymer melts with higher strain hardening
parameter along with Type – 1 nature of the strain hardening behavior can lead to closed cell
foams with smaller cell sizes and higher cell densities compared to Type – 2 melts. Since, the
strain rate of a growing bubble or cell is inversely proportional to its size, hence in order to arrest
the growth of cell size, strain hardening of polymer melts under uniaxial extensional flow at
higher rates is more important for foaming applications. This behavior is also reflected from the
cell size distributions presented in Figures 5.4 – 5.6. PPNC-S4 has maximum fraction of cells in
the range of 20 – 40 µm while the same for PPNC-S7 is in the range of 40 – 60 µm. A summary

129

of all the foam results is presented in Figure 5.8, which clearly shows superior foam
characteristics of Type – 1 strain hardening melt.

5.4. Conclusions
The importance of effect of extent and nature of strain hardening behavior on foamability
of polypropylene/clay nanocomposites was studied in this section. Nanocomposites with
different types of strain hardening behavior were tested for foam evaluation. A comparison of
foam quality (average cell size and cell density) produced by nanocomposites prepared by silated
and un-silated organoclays was presented in this section. The polymer melt that shows
significant strain hardening at the lowest strain rate but has very little strain hardening at the
higher strain rates (Type – 2) resulted in foams with cell sizes that are larger than polymer melts
displaying a steady strain hardening at all rates (Type – 1), indicating that the cell growth is
arrested at a smaller cell size in the presence of Type – 1 strain hardening than in the presence of
Type – 2 strain hardening.

130

Figure 5.1: Non-isothermal crystallization exotherm for linear PP/clay nanocomposites
prepared using 85 wt% linear PP, 12 wt% PO 1015 and 3 wt% of different treated
organoclays.

131

Zone I

Final cell size (Rmax)

Current cell size (R)

1.0

Zone II

0
θr

Growth Time (s)

Figure 5.2: Evolution of cell size as a function of time demonstrated by various theoretical and
experimental work [196-199]. (θr represents the residence time in the die)

132

Figure 5.3: Scanning electron micrograph of PPNC-S6 (prepared using 3 wt% of unsilated
I.44P) foam extrudate.

133

Figure 5.4: Scanning electron micrograph of PPNC-S7 (prepared using 3 wt% of diamino
dimethoxy silane treated I.44P organoclay) foam extrudate.

134

Figure 5.5: Scanning electron micrograph of PPNC-S4 (prepared using 3 wt% of diamino
dimethoxy silane treated I.30P organoclay) foam extrudate.

135

Figure 5.6: Cell size distribution of PPNC-S6 foam.

136

Figure 5.7: Cell size distribution of PPNC-S7 foam.

137

Figure 5.8: Cell size distribution of PPNC-S4 foam.

138

Figure 5.9: Progressive decrease in cell size and increase in cell density as a function of
material property of linear PP/clay nanocomposite.

139

6. EFFECT OF COMPATIBILIZER PROPERTIES ON STRAIN
HARDENING BEHAVIOR OF POLYPROPYLENE CLAY
NANOCOMPOSITES
6.1. Introduction
Polymer nanocomposite is a field of research that has been studied in great detail over the
past two decades. Interest in polymer nanocomposites has been results of improvements in their
various physical and mechanical properties due to the addition of a small amount of anisotropic
nanoparticles in the polymer matrix [73, 74]. Amongst all nanofillers, nanoclays are most widely
studied probably because of its availability and cost.
Polymer/layered silicate (PLS) nanocomposites are synthesized by dispersing nano
silicate layers in the polymer matrix. In order to achieve better dispersion of these nanolayers,
the naturally occurring montmorillonite has to be organically modified to synthesize polymer
compatible nanoclay. In order to facilitate the interaction between the hydrophilic clay layers and
hydrophobic polymers, the galleries of the clay tactoids are intercalated with long chain alkyl
ammonium surfactants [77-81].
Even with the organic functionalization of montmorillonites, the intercalation of
hydrophobic polypropylene in the organically modified montmorillonites is difficult and is
generally facilitated by the use of less hydrophobic compatibilizers. These compatibilizers are, in
general, polypropylene functionalized by grafting hydroxyl groups or more commonly maleic
anhydride groups at the ends of the polyolefin chains [78, 86, 87]. Okada et al. [88] were the first
to demonstrate the use of polypropylene-grafted maleic anhydride (PP-g-MA) as a
compatibilizer to synthesize PP-clay nanocomposites. Since then the importance of such polar
compatibilizers for the use in non-polar clay nanocomposite preparation is well demonstrated by
140

various researchers [86, 89-94]. The effect of maleic anhydride content and molecular weight of
the PP-g-MA on the dispersion of nanoclays in polypropylene nanocomposites have been studied
thoroughly by various researchers [136, 200-203]. Both these factors along with the loading of
the compatibilizer significantly affect the extent of polymer intercalation and subsequent
improvement in various physical and mechanical properties. Marchant and Jayaraman [136]
reported that optimum loading of the compatibilizer and higher maleic anhydride content
resulted in exfoliated structure of PP nanocomposites. The effect of molecular weight, maleic
anhydride content and loading of PP-g-MA was investigated by Wang et al. [161, 204]. They
reported that compatibilizers with lower melt flow index (higher molecular weight) leads to
better dispersion of the clay layers in the PP matrix. It is well established that state of dispersion
of these nanoparticles significantly affect the rheological properties in shear flow [69-71, 74, 78],
however very limited research has been done on the effect of PP-g-MA material properties on
extensional rheology of polypropylene nanocomposites.
It has been well established that maleation of linear polypropylene leads to reduction in
the molecular weight of the resulting PP-g-MA due to chain scission [205]. The grafting process
of maleic anhydride (MA) groups onto polypropylene follows a free radical mechanism, in
presence of organic peroxide [206, 207]. Due to the complexity of the grafting mechanism, two
different grafting mechanisms have been proposed in the literature; one where the MA groups
are distributed along the length of the PP chain as succinic anhydride units [208, 209], and other
where the MA groups are grafted only at the polypropylene chain-ends [210-212]. Sclavons et al.
[212] have recently validated the later mechanism of MA grafting at the chain-ends of
polypropylene by thorough FTIR studies. Further in their study they have mentioned that apart
from being grafted at the chain-ends, MA groups are present in a cluster of more than one unit.

141

The objective of this research is to understand or even propose a plausible explanation
that can help understand the role of compatibilizer in the strain hardening mechanism of linear
polypropylene/clay nanocomposites.

6.2. Materials and Experimental Procedures
Several compatibilizers were selected in order to understand the effect of molecular
details of PP-g-MA on the strain hardening of the resulting PP/clay nanocomposites. The
variables that led to the selection of these compatibilizers were molecular weight, maleic
anhydride content and the molecular structure of the resulting PP-g-MA. Maleic anhydride can
be grafted either on PP homopolymer or a copolymer. Based on these factors the various
compatibilizers used in this study were AC950, Exxelor PO 1015, Exxelor PO 1020, Bondyram
1001CN, and Bondyram 1001; the details of which are listed in Table 2.1 and the details of
corresponding nanocomposites are tabulated in Table 6.1.
The nanocomposites were characterized for their structural morphology and for their
rheological behavior in both shear and uniaxial extensional flow.

6.3. Results and Discussion
6.3.1. Uniaxial extensional rheology
The melt extensional viscosity curves presented in Figures 6.1-6.3 compares the effect of
molecular weight and molecular structure of the PP-g-MA on strain hardening behavior of linear
PP/clay nanocomposites. The strain hardening data suggests that certain minimum molecular
weight is required below which strain hardening is not observed in PP/clay nanocomposites.

142

Strain hardening in branched polymer melts undergoing extensional flow has been
attributed to slowly relaxing chain stretch by Larson and McLeish [213]. The dynamics of
entanglements involving the branches is accounted for by a chain relaxation equation in addition
to the dynamics of the main chain or backbone. Alternatively, strain hardening in the melt has
been attributed to compression of the tube around the main chain which amounts to an increase
in the entanglement density along the chain by Wagner et al. [214-218].
In cases where strain hardening in polymer melts has been achieved by partial
crosslinking of the polymer chains, the phenomenon has been attributed to the trapped
entanglements between crosslink junctions. Blends of linear polypropylene with small amounts
of chemically crosslinked polymers contain trapped entanglements on chain segments between
crosslinks or junctions which respond to stretching differently from entanglements and chain
segments in the bulk – cf. Yamaguchi and Suzuki [153] and Yamaguchi [152]. Ogura et al. [190]
have also indicated that this effect will be lost upon excessive crosslinking.
The mechanism of strain hardening in linear PP – clay nanocomposite melts appear to be
analogous to that reported for blends containing small amounts of crosslinked polymers.
Linkages or interactions between the clay surface or edge and the maleated polymer can form a
network of “physical crosslinks” or anchor points between which entanglements may be trapped
to produce a similar effect. In nanocomposites where the organoclay was used without additional
treatment, hydrogen bonding between the hydroxyl groups of the organoclays and the maleic
anhydride group of the compatibilizer holds the compatibilizer chains and the clay together.
Since MA groups are primarily grafted as multiple clusters at the chain-ends, a maximum of four
to six MA groups are available per chain for a maleated homopolymer [212]. However, in case

143

of a maleated PP-PE copolymer, there could be several MA groups all across the chain length, as
shown in Figure 6.5.
Comparison of strain hardening data (Figure 6.1) for nanocomposites with
compatibilizers having different molecular weight and therefore different chain lengths (PPNCN2, PPNC-S6 and PPNC-S13) revealed that the length of compatibilizing polymer chain that is
anchored at one or more clay surfaces is critical because it determines the number of trapped
entanglements along the chains that bridge between clay surfaces or edges. Trapped
entanglements and the chains which bridge between the clay edges or surfaces respond to stretch
differently from chains and entanglements in the bulk polymer and produce strain hardening in
the melt. However, a comparison between PPNC-S6 and PPNC-S13 suggests that beyond a
sufficient molecular weight, a further increase in molecular weight do not translate to progressive
increase in the strain hardening property of the nanocomposite prepared with organoclays
without further treatment. This is despite the fact that PPNC-S6 has maleated copolymer that has
greater number of MA groups along the chain that can interact with the organoclays. This is
because, the strength of such interaction is still dictated by weaker hydrogen bonding, which
limits further improvements in strain hardening.
A comparison between PPNC-S7 and PPNC-S14 (Figure 6.2) and PPNC-S16 and PPNCS17 (Figure 6.3) is made to understand the effect of molecular structure of the PP-g-MA on the
strain hardening of the nanocomposites. In this case, diamino dimethoxy silane grafted I.44P was
used to prepare the nanocomposites. As mentioned in Chapter 4, the amino silanes grafted on the
clay faces and/or edges can react with the maleic anhydride groups to replace the weaker forces
of attraction with stronger chemical bonds. Chemical reactions between the silated clays and
maleated PP-PE copolymer (PO 1015) resulted in nanocomposites with significant increase in

144

melt strain hardening compared to the unsilated clay nanocomposites. However, the same was
not true with nanocomposites prepared with maleated PP homopolymer (PO1020). This could be
due to the additional grafting of maleic anhydride on copolymer backbone or it could also be due
to the difference in the molecular weight between the two compatibilizers. Strain hardening
comparison between PPNC-S16 and PPNC-S17 shows that molecular weight or the chain length
of the compatibilizer is more important than the molecular structure of the compatibilizer.
PPNC-S16 is prepared from Bondyram 1001CN which is a propylene-ethylene copolymer while
PPNC-S17 is prepared from Bondyram 1001 which is a homopolymer and both the Bondyram
compatibilizers have comparable molecular weight (as MFR is same for both of them). This
confirms that longer chain length results in greater number of trapped entanglements along the
chains that bridge between clay surfaces and edges. The entangled polymer-clay structure when
stretched under uniaxial extensional flow, forces the clay and the polymer chains to orient along
the direction of the stretch. The orientation of the clay platelets has been confirmed by
conducting TEM on a stretched specimen (Figure 6.6). Since the polymer chains are anchored to
the clay edges and/or faces, the relaxation mechanism of the chain stretch would slow down
increasing the chain stretch relaxation time.
To further understand the strain hardening mechanism, nanocomposites prepared using
Bondyram 1001CN and Bondyram 1001 were studied. Strain hardening curves for PPNC-N3,
PPNC-S15, PPNC-S16 and PPNC-S17 are presented in Figure 6.3 and 6.4. A comparison
between PPNC-S16 and PPNC-N3 shows the effect of compatibilizer to nanoclays ratio on strain
hardening. With the same compatibilizer and the clay being used in these two nanocomposites,
as the ratio of compatibilizer to clay is increased from 4 to 7, strain hardening increases. This
suggests that some minimum polymer-clay interaction is required for strain hardening to appear.

145

Shear rheology has been used to get a qualitative idea of extent of entanglement formed
in various polypropylene clay nanocomposite systems. It has been well reported that with the
increase in polymer-particle interaction, the low frequency plateau modulus also increases [74,
92, 155]. Figure 6.7 shows the variation of storage modulus (G’) with frequency for various
nanocomposites prepared using Bondyram 1001CN and Bondyram 1001 as the compatibilizer. A
comparison between the matrix and the nanocomposites shows that with the addition of 3 wt% of
nanoclays the low frequency storage modulus increases confirming the there indeed exists
significant polymer-particle interaction. However, the low frequency plateau modulus cannot be
seen from such plots. An alternative way to estimate the low frequency plateau modulus is by the
use of van-Gurp Palmen (vGP) plots, a plot of phase angle (δ) against the complex shear
modulus (G*). A minimum at low frequency region in the vGP plot is an estimate of low
frequency plateau modulus of a polymer melt. For linear polymer melts or melts with high polydispersity index, the vGP plot generally shows a monotonic decrease in the phase angle [219,
220]. Figure 6.8 shows the vGP plot for linear PP/clay nanocomposites with Bondyram 1001CN
and Bondyram 1001 as the compatibilizer. A prominent minimum is observed for PPNC-S15,
PPNC-S16 and PPNC-S17 nanocomposites; however the vGP plot shows a monotonic drop in
the phase angle for PPNC-N3 nanocomposite and the polymer matrix (linear PP and
compatibilizer in the appropriate ratio). This suggests that the polymer-particle interaction is
greatest in case of PPNC-S15 and PPNC-S17 followed by PPNC-S16 and it is outside the test
range for PPNC-N3 and the matrix. The melt extensional viscosity data for PPNC-S15, PPNCS16 and PPNC-S17 shows that these nanocomposites melts show significant strain hardening
behavior while PPNC-N3 and the polymer matrix do not show such phenomenon. This confirms
that the presence of strain hardening in these nanocomposites is due to greater polymer-particle

146

interaction. The absence of strain hardening in PPNC-N3 could thus be due to insufficient
interactions due to low amount of PP-g-MA. Further a comparison between PPNC-S15, PPNCS16 and PPNC-S17 shows that the low frequency plateau modulus is greater for PPNC-S15 and
PPNC-S17 melts, which explains why PPNC-S16 shows lower extent of strain hardening
behavior compared to PPNC-S15 and PPNC-S17.

6.4. Conclusions
In composites where the organoclay was used without additional treatment, hydrogen
bonding on the surfaces of the organoclay holds the compatibilizer chains. Comparison of
composites with compatibilizers having different chain lengths revealed that the length of
compatibilizing polymer chain that is anchored at one or more clay surfaces must be sufficient to
affect the dynamics of the entanglement network during extensional flow and produce strain
hardening in the melt. Further van-Gurp Palmen plots were used to obtain a qualitative estimate
of low frequency plateau modulus, which was used to explain the strain hardening phenomenon
in linear PP/clay nanocomposite melts. The extent of strain hardening increased with an increase
in the low frequency plateau modulus of the polymer nanocomposite melt.

147

Specimen

PP 6523
(wt%)

Nanoclay (wt%)

PP-g-MA (wt%)
AC
950

PO
1020

PO
1015

Bond.
1001CN

Bond.
1001

I.44P

Silated
I.44P

PPNC-N2

68.0

24.0

--

--

--

--

8.0

--

PPNC-N3

85.0

--

--

--

12.0

--

--

3.0

PPNC-S6

85.0

--

--

12.0

--

--

3.0

--

PPNC-S7

85.0

--

--

12.0

--

--

--

3.0

PPNC-S13

85.0

--

12.0

--

--

--

3.0

--

PPNC-S14

85.0

--

12.0

--

--

--

--

3.0

PPNC-S15*

76.0

--

--

--

21.0

--

--

3.0

PPNC-S16

76.0

--

--

--

21.0

--

--

3.0

PPNC-S17

76.0

--

--

--

--

21.0

--

3.0

Table 6.1: Composition of linear PP/clay nanocomposites. (*The silated clay used here is
Diamino dimethoxy-silated I.44P, except for PPNC-S15 where it is diamino
monomethoxy-silated I.44P).

148

Figure 6.1: Melt extensional viscosity comparison for various linear PP/clay nanocomposites,
showing the effect of molecular weight of the compatibilizer on the strain hardening
characteristics of the derived nanocomposites.

149

Figure 6.2: Melt extensional viscosity comparison for two nanocomposites under uniaxial
extensional flow, showing the effect of molecular structure of the compatibilizer on
the strain hardening characteristics of the derived nanocomposites (PPNC-S7 is
prepared using PO 1015 and PPNC-S14 is prepared using PO 1020).

150

Figure 6.3: Melt extensional viscosity comparison for two nanocomposites under uniaxial
extensional flow, showing the effect of molecular structure of the compatibilizer on
the strain hardening characteristics for nanocomposites having Bondyram 1001
(PPNC-S17) and Bondyram 1001CN (PPNC-S16) as the compatibilizer.

151

Figure 6.4: Melt extensional viscosity data under uniaxial extensional flow, for nanocomposites
having Bondyram 1001CN as the compatibilizer. PPNC-N3 is prepared from
Diamino dimethoxy silane treated I.44P and has a compatibilizer to clay ratio of 4
and PPNC-S15 is prepared from Diamino monomethoxy silane treated I.44P and
has a compatibilizer to clay ratio of 7.

152

(a)

(b)

(c)

Figure 6.5: Molecular structure of PPgMAs. (a) AC-950, (b) PO 1020, and (c) PO 1015

153

200 nm
Figure 6.6: TEM micrographs of a stretched section of a representative nanocomposite (PPNCS4) sample. The nanoclay platelets are oriented along the direction of stretch
(shown by the arrow). The cartoon shows the section of the stretched samples used
for TEM.

154

Figure 6.7: Variation of storage modulus with frequency for various linear PP-clay
nanocomposites and the matrix polymer, prepared using Bondyram 1001CN as the
compatibilizer.

155

Figure 6.8: Van-Gurp Palmen plot for various linear PP-clay nanocomposites prepared using
Bondyram 1001CN and Bondyram 1001 as the compatibilizer along with the matrix
(linear PP + Bondyram 1001CN).

156

7. EFFECT OF MELT PROCESSING ON STRAIN HARDENING
BEHAVIOR OF POLYPROPYLENE CLAY
NANOCOMPOSITES - RECYCLABILITY

7.1. Introduction
With the growing interest towards sustainability, material recycling is getting serious
attention from industries, both from environmental point of view as well as resource
management and economics. One of the main objectives of this research is to replace currently
used thermosetting foams with completely recyclable thermoplastic foams. Since our objective is
to use the polypropylene – clay nanocomposites for foaming applications, it is important to study
the effect of reprocessing on nanoclays morphological stability and strain hardening of these
nanocomposites. In general, two main factors affecting the polymer properties during any
processing operations are temperature and shear. In this study we have studied the effect of
temperature and shear on strain hardening separately.
The effect of reprocessing of polypropylene was studied by Incarnato et al [221]. They
observed that due to reprocessing, the weight average molecular weight and shear viscosity of
the polypropylene decreases due to thermal degradation along with drop in elongation at break
with increase in recycling operations. Moreover, in case of polymer clay nanocomposites there is
a high likelihood that the dispersed nanoplatelets can agglomerate forming network like structure
due to reprocessing [222]. Hence, it is expected that after few recycling operations the virgin
polymer might show a drop in mechanical properties and rheological properties [52, 222-227].
In general, melt mixing is used to disperse the nanoclay in the polymer matrix, followed
by additional molding or processing steps for the intended applications. Work done by Reichert
157

et al. [228] and Manias et al. [83] on the effect of additional melt blending on nanolayer
morphology resulted in shear induced flocculation of nanolayers, also observed by Okamoto et
al. [229]. Reichert et al. [228] also conducted annealing studies of injection molded
polypropylene nanocomposites and found a decrease in spacing between the nanolayers due to
annealing. Manias et al. [83] reported the effect of thermal annealing on the morphology of
PP/clay nanocomposites and Pp-g-MA/clay nanocomposites. They observed severe collapse of
clay galleries in case of PP nanocomposite while the PP-g-MA nanocomposite remained stable
indicating enhanced stability of the nanoclays morphology due to better coupling between the
clay and the matrix.
Another important aspect of nanoclay composites is the anisotropy of the nanoparticles.
This allows them to be oriented differently along a particular direction resulting in significant
improvement in various material properties [230-232]. Various studies have investigated the
effect of shear [233-237] and elongational flow [131] on the orientation of nanoplatelets.
Okamoto et al. [131] observed alignment of the nanoplatelets perpendicular to the stretch
direction in PP-g-MA/clay nanocomposites under uniaxial extensional flow. as already
mentioned in Chapter 6, the reasons for strain hardening in polymer melts is the increase in the
chain stretch relaxation time. In organoclay nanocomposites the orientation of nanoparticles
could significantly alter this characteristic time. Hence any change in morphology of the
nanoclays that can prevent or reduce the orientability of the platelets should result in drop in
strain hardening characteristics.
Since recyclability drives the substitution of thermoset polyurethane foams with closed
cell PP foams, the effect of re-processing on the melt strain hardening under uniaxial extensional
flow was investigated in this chapter. Various nanocomposites were annealed and extruded

158

though a uniform cross-section die to study the effect of heating and shear flow on the structural
morphology and strain hardening of the nanocomposites.

7.2. Materials and Experimental Procedures
Various linear PP/clay nanocomposites (PPNC-S3, PPNC-S4, PPNC-S6 and PPNC-S7)
prepared using silated and unsilated I.30P and I.44P nanoclays, characterized in Chapters 3 and
4, were used to study the effect of additional melt processing on the morphology of the
nanoclays along with their strain hardening behavior. In order to carry out this study, the effect
of processing is divided into two steps – thermal annealing to evaluate thermal stability and shear
flow in an extruder to evaluate stability of the nanocomposites under shear flow. This was done
so as to understand the individual effect of temperature and shear on the stability and strain
hardening of the nanocomposites.

7.2.1. Thermal Annealing
In order to study the effect of thermal annealing alone on strain hardening behavior, the
PP nanocomposites were thermally annealed at a temperature of 180ºC for 5 hrs in an oven under
a blanket of nitrogen atmosphere so as to prevent any oxidative degradation of the nanoclays and
the compatibilizer. The annealed samples were then re-molded and tested for strain hardening
characteristics. The annealed samples were also characterized for their structural morphology
using XRD and shear rheology. A frequency sweep was run on the annealed samples in order to
see the effect annealing on the complex viscosity and shear modulus (G’ and G”) of the samples.

159

7.2.2. Extrusion
To study the effect of shear on strain hardening, various PP/organoclay nanocomposites
were extruded through a capillary die (d = 1mm, L/D = 30) at a shear rate of 500s-1 and a
temperature of 180ºC in a Dynisco LCR6000 Capillary rheometer. The PPNC granules were
carefully fed into the capillary rheometer with intermittent packing of the melt. The extruded
samples were then molded and characterized for various structural and rheological properties.
XRD and TEM were used to study the effect of extrusion on the structural morphology of the
nanocomposites. Both shear and uniaxial extensional rheology were used to study the rheological
properties of the extruded materials. The details of the characterization techniques are presented
in Chapter 2.

7.3. Results and Discussion
7.3.1. Effect of Thermal Annealing
The X-ray diffraction patterns of various linear PP/clay nanocomposites before and after
annealing are presented Figures 7.1-7.4. A comparison of the d-spacing before and after
annealing reveals that the spacing between the platelets have either remained unchanged or have
increased and have not resulted in coarsening of the nanoclays structure. This type of
phenomenon has also been reported in literature [238-240] for nanocomposites. Tang et al. [238]
and Fordiani et al. [239] reported the structural changes of maleated PP/organoclays due to
thermal annealing. FTIR studies conducted on these sample suggested that prolonged annealing
of the polymer nanocomposite melt, in presence of oxygen, led to a chemical reaction between
the hydroxyl group present at the nanoclays edges and the maleic anhydride group. This led to

160

greater intercalation of the polymer chains by so called “peeling-off” mechanism. The presence
of a small fraction of oxygen promotes this kind of mechanism. Even though we carried out the
annealing experiment under nitrogen; presence of traces of oxygen cannot be eliminated. The
effect of such reactions should also be reflected in the strain hardening behavior of the
nanocomposites, as the weaker hydrogen bonds are now replaced with stronger chemical bonds.
The effect of thermal annealing on the melt extensional viscosity data is presented in Figures 7.57.8. Figure 7.5 shows a significant increase in the strain hardening behavior for PPNC-S6 after
annealing. This is due to the replacement of the weaker H-bonds with chemical bonds between
the compatibilizer and the clay. However, thermal annealing resulted in no change in the strain
hardening behavior for the silated clay nanocomposites (figure 7.6-7.8). This is because there
were chemical bonds between the compatibilizer and the clay even before annealing. Hence, it
can be safely concluded that thermal annealing does not result in deterioration of the strain
hardening behavior of the PP/clay nanocomposites.

7.3.2. Effect of Extrusion
The X-ray diffraction patterns for extruded nanocomposite samples are presented in
Figure 7.1-7.4. From the XRD, it is evident that extrusion of the nanocomposites, through a
uniform die, resulted in a decrease in d-spacing of the nanocomposites along with a more
prominent peak. This indicates that extrusion of the nanocomposites through a uniform die
resulted in restacking of the nanolayers. The change in the morphology of the nanocomposites is
solely due to shear and not due to temperature as thermal annealing did not result in coarsening
of the nanocomposite structure.

161

From the strain hardening data presented in Figure 7.5-7.8, it is quite evident that thermal
annealing alone does not result in decrease in strain hardening characteristics. All four PP
nanocomposites subjected to thermal annealing at 180ºC for 5hrs were able to maintain their
strain hardening behavior. However, the effect of shear during extrusion has a completely
different story. In case of PPNC-S6 and PPNC-S7 nanocomposites, extruded samples resulted in
drop in the strain hardening behavior, especially at the higher strain rates. This could possibly be
due to re-stacking of the nanoplatelets during extrusion. In general, extrusion of nanocomposites
leads to agglomeration of nanoclay platelets (Figure 7.12). This could be either edge-to-edge
agglomeration or face-to-edge agglomeration (commonly known as House-of-cards formation).
The structure of the nanoclays after extrusion has also been studied by transmission electron
microscopy. The TEM micrographs for PPNC-S7 before and after extrusion, presented in Figure
7.9, demonstrate that extrusion led to more edge-to-edge re-stacking of the platelets, thereby
affecting the strain hardening characteristics of the nanocomposites. This is because, grafting of
the reactive silane coupling agents has occurred primarily at the edges of I.44P clay, resulting in
predominantly edge-to-edge aggregation of the nanoclays. The formation of such aggregates
increases the resistance towards orientation of the nanoplatelets during the extensional flow,
thereby reducing the strain hardening characteristics.
On the other hand, extrusion of PPNC-S4 nanocomposites resulted in complete loss in the
strain hardening characteristics while PPNC-S3 nanocomposites retains most of its strain
hardening characteristics even after extrusion. This could be due to severe agglomeration of the
clay platelets in case of PPNC-S4. The TEM micrographs for PPNC-S3 and PPNC-S4
nanocomposites are presented in Figures 7.10 and 7.11. In case of PPNC-S3, extrusion through a
uniform die did result in some re-stacking of the nanoclays that has been reflected in a small

162

drop in the strain hardening characteristics. However, extrusion of PPNC-S4 resulted in severe
face-to-edge kind of agglomeration. This results in decrease in anisotropy of the nanoparticles
thereby making it difficult to orient these agglomerates along the direction of stretch, resulting in
complete loss of strain hardening behavior in extruded PPNC-S4 nanocomposite. The formation
of such aggregates for PPNC-S4 could be explained using a simple cartoon shown in Figure
7.13.
The diamino dimethoxy-silane treated clay platelet have face grafted silanes with two
hydrolysable groups; one of them is involved in the condensation reaction with the active sites of
the nanoclay, while the other one is unreacted. These unreacted hydroxyl groups during the
extrusion process could interact with those of neighboring clay platelets forming a house of cards
structure. These structures might no longer behave as a nanoplatelet and orientating these
agglomerated structures during extensional flow to further assist orientation of polymer chains is
highly improbable. While in case of diamino monomethoxy-silane treated I.30P nanocomposites
(PPNC-S3) due to the absence of any free hydrolysable groups on the silane molecule attached to
the clay surface, this kind of re-agglomeration is not possible, thereby being able to retain its
strain hardening characteristics. Hence, diamino monomethoxy-silane treated I.30P-PP
nanocomposites is a potential candidate for producing foams that can be reused, as these
nanocomposites can maintain their strain hardening behavior under uniaxial extensional flow, an
indispensable property for producing high quality polymer foams, even after reprocessing.
To further verify the formation of such agglomerates, shear modulus and viscosity of
these samples were measured. A significant drop in the storage modulus and/or complex
viscosity at low frequency is an indication of loss of polymer-particle network. It has been well
established in literature that exfoliation of nanoclays leads to significant increase in both storage

163

modulus and complex viscosities at low frequency region as compared to nanocomposites with
poor dispersion of clay platelets [69, 71, 74, 92]. Figure 7.14 shows the effect of extrusion on
storage modulus and Figure 7.15 shows the effect of extrusion on complex viscosity for PPNCS3 and PPNC-S4 nanocomposites. It is evident from the figures that there is a significant drop in
both storage modulus and complex viscosity for PPNC-S4 compared to PPNC-S3. This is due to
the formation of face-to-face aggregates for PPNC-S4 that significantly lowers the reinforcement
effect of nanoclays on the polymer matrix, thereby affecting both its shear and extensional
rheology.

7.4. Conclusions
Recyclability of linear PP/clay nanocomposites drives the substitution of thermoset
polyurethane foams with closed cell PP foams. To study the effect of re-processing on the melt
strain hardening, various nanocomposites were annealing and extruded though a uniform crosssection die to study the effect of heating and shear flow on the structural morphology and strain
hardening of the nanocomposites. The results indicated that annealing had no negative impact on
morphology or strain hardening of the nanocomposites. However, extrusion of these
nanocomposites indicated that nanocomposites prepared using multiple hydrolysable groups
resulted in formation of face-to-face nanolayer aggregates that are difficult to orient during
uniaxial extensional flow and thus resulted in loss of the strain hardening behavior. The
formation of such agglomerates has been verified by transmission electron microscopy. Shear
rheology results showed a drop in storage modulus and complex viscosity for samples
confirming the formation of agglomerates during extrusion. However, the formation of
agglomerates can be avoided by the use of amino silanes having only one hydrolysable group.

164

Such nanocomposites showed retention of strain hardening behavior after extrusion and hence
can be used as a potential candidate for the replacement of thermoset polymers for applications
where strain hardening in polymer melt is essential.

165

Figure 7.1: X-ray diffraction pattern showing the effect of thermal annealing and extrusion for
PPNC-S6 nanocomposites.

166

Figure 7.2: X-ray diffraction pattern showing the effect of thermal annealing and extrusion for
PPNC-S7 nanocomposites.

167

Figure 7.3: X-ray diffraction pattern showing the effect of thermal annealing and extrusion for
PPNC-S3 nanocomposites.

168

Figure 7.4: X-ray diffraction pattern showing the effect of thermal annealing and extrusion for
PPNC-S4 nanocomposites.

169

Figure 7.5: Melt extensional viscosity of PPNC-S6 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow.

170

Figure 7.6: Melt extensional viscosity of PPNC-S7 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow.

171

Figure 7.7: Melt extensional viscosity of PPNC-S3 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow.

172

Figure 7.8: Melt extensional viscosity of PPNC-S4 nanocomposites, showing the effect of
thermal annealing and extrusion on strain hardening under uniaxial extensional
flow.

173

200 nm

200 nm
Figure 7.9: TEM micrographs of linear PP/diamino dimethoxy silated I.44P (PPNC-S7)
nanocomposites showing the effect of extrusion on nanoclay dispersion (top: before
extrusion, bottom: after extrusion).

174

200 nm

200 nm
Figure 7.10: TEM micrographs of linear PP/diamino monomethoxy silated I.30P (PPNC-S3)
nanocomposites showing the effect of extrusion on nanoclays dispersion (top:
before extrusion, bottom: after extrusion).

175

200 nm

200 nm
Figure 7.11: TEM micrographs of linear PP/diamino dimethoxy silated I.30P (PPNC-S4)
nanocomposites showing the effect of extrusion on nanoclays dispersion (top:
before extrusion, bottom: after extrusion).

176

Figure 7.12: Schematic representing state of nanolayers dispersion/agglomeration after mixing
and after extrusion through a capillary die.

177

Extrusion

After

Figure 7.13: Schematic representing agglomeration of diamino dimethoxy silated I.30P layers
in PPNC-S4 due to extrusion through a capillary die.

178

Figure 7.14: Effect of extrusion on storage modulus (G’) of PPNC-S3 and PPNC-S4
nanocomposites. (open symbols: before extrusion and closed symbol: after
extrusion)

179

Figure 7.15: Effect of extrusion on complex viscosity of PPNC-S3 and PPNC-S4
nanocomposites. (open symbols: before extrusion and closed symbol: after
extrusion)

180

8. CONCLUSIONS AND RECOMMENDATIONS
8.1. Conclusions
This study was focused towards improving the strain hardening property of linear
polypropylene using a suitable compatibilizer and organically modified montmorillonite as
nanoclays for foaming applications. This was achieved by improving the interactions between
the different components of a polypropylene nanocomposite system ― bulk polypropylene,
maleated polypropylene and the nanoclay intercalated with long chain alkyl ammonium
surfactants. The objectives of this research were (1) to understand the effects of composition on
melt extensional strain hardening in nanocomposites of linear polypropylene with nanoclay,
compatibilizer and various coupling agents and (2) to evaluate selected nanocomposite
formulations in extrusion foaming with chemical blowing agents.
Nanocomposites were prepared with organically modified montmorillonite, maleic
anhydride grafted polypropylene as compatibilizer and linear polypropylene as the bulk matrix.
Compatibilizers with different molecular weights and molecular structure (polypropylene
homopolymer or ethylene-propylene copolymer) were used. The organoclays were also treated
with amino silane coupling agents to increase the interaction between nanoclays and the
compatibilizer. The nanocomposites were characterized by X-Ray diffraction, TEM, melt
extensional rheometry and DSC. Selected nanocomposites were foamed in a single screw
extruder using a chemical blowing agent. The extrusion foaming process was optimized for each
nanocomposites based on their melting crystallization temperature. The cellular structure of the
extruded foam samples was characterized by scanning electron microscopy and its image
analysis was used to obtain average cell size and cell density of the foam samples.

181

When several polypropylene clay nanocomposites all with good dispersion of the dialkyl
amine treated organoclay were foamed by extrusion with a chemical blowing agent, systematic
variations in foam quality were obtained based on differences among them in tensile melt
strength and in crystallization behavior. Even among nanocomposites without any melt strain
hardening, decreasing the rate of crystallization led to closed cell foams with smaller cell size
and higher cell density. Among nanocomposites where significant strain hardening was observed
in extensional flow along with the slower rates of crystallization, the extruded polypropylene
nanocomposite foams displayed the smallest cell sizes and the greatest cell density by reducing
cell coalescence. These enhancements may be attributed to attachment of high molecular weight
PP-g-MA chains to the surfaces of nanolayers due to hydrogen bonding, providing greater
constraints on chain stretch relaxation and also slowing down the crystallization. The foam
extrudates were compared with those prepared from a blend of linear and branched
polypropylene. The cell size and cell density of foam prepared from the blend of linear and
branched polypropylene was significantly higher than that prepared from the linear
polypropylene/clay nanocomposites. This was due to absence of any nucleating agent present in
the blend. Hence nanoclays acts as both rheology modifier as well as nucleating agent and thus
improves the quality of foam produced form the linear PP/clay nanocomposites.
The organoclays and the compatibilizer present in the nanocomposites have mainly
electrostatic or hydrogen bonding type interactions, which are weaker forces of attraction. In
order to replace these weaker forces of attraction, the organoclays were further treated with
amino silane coupling agents. The amino silane treated nanoclays resulted in either edge grafting
or face and edge grafting depending upon the amount of intercalant already present in the clay
galleries. Characterization of the silated clays using BET, TGA, and XRD confirmed that

182

organoclays with low surfactant content (I.30P) resulted in face and edge grafting of the silanes
while the high surfactant containing I.44P organoclays led to edge grafting of the clays. The
amino groups of the face and/or edge grafted silanes is believed to react with the maleic
anhydride group of the selected compatibilizer to replace the weaker forces of attraction with
stronger chemical bonds. This resulted in nanocomposites with significant increase in melt strain
hardening compared to the unsilated clay nanocomposites. These nanocomposites were further
evaluated for extrusion foaming using a chemical blowing agent. It was found that the
nanocomposites showing melt strain hardening behavior that remained unchanged with strain
rates resulted in foams with lowest possible cell size and highest cell densities within the scope
of this research.
To further understand the effect of compatibilizer molecular structure on the strain
hardening behavior of the nanocomposites, compatibilizers having different chain lengths and
different chain structure were selected. It is known that grafting of maleic anhydride on
polypropylene occur predominantly at the two ends of the polypropylene chan. Hence if a
ethylene – propylene copolymer is selected for maleation, then along with the two ends of the
polymer backbone, grafting of maleic anhydride groups will also take place at additional sites
along the ethylene-propylene copolymer chain. A comparison of composites with
compatibilizers having different chain lengths revealed that the length of compatibilizing
polymer chain that is anchored at one or more clay surfaces must be sufficient to affect the
dynamics of the entanglement network during extensional flow and produce strain hardening in
the melt. Further, it was also found that maleated copolymer of propylene and ethylene with high
molecular weight provided greater improvements in strain hardening with the use of silated
clays. This is because the entangled network formed between polymer-clay systems would

183

involve greater number of clay layers/galleries and thereby would significantly increase the
characteristic chain stretch relaxation time. With the help of van-Gurp Palmen plots, it was
shown that nanocomposites having higher low frequency plateau modulus showed stronger strain
hardening behavior.
Since recyclability of linear PP/nanocomposites drives the substitution of thermoset
polyurethane for various foaming applications, the effect of re-processing on melt strain
hardening was also investigated. Various nanocomposites prepared different silane treated clay
were annealed and extruded though a uniform cross-section die to study the effect of heating and
shear flow on the structural morphology and strain hardening of the nanocomposites. The results
showed that annealing had no negative impact on morphology or strain hardening of the
nanocomposites. However, extrusion of these nanocomposites indicated that nanocomposites
prepared using multiple hydrolysable groups resulted in formation of layered aggregates that are
difficult to orient during uniaxial extensional flow and thus resulted in loss of the strain
hardening behavior. This was verified both by transmission electron microscopy and by shear
rheology. The nanocomposites prepared using diamino dimethoxy silated I.30P showed a
significant drop in both storage modulus and complex viscosity after extrusion. This can happen
if the polymer-particle network is destroyed due to extrusion. Hence proper selection of clay
treatment and compatibilizer can lead to linear polypropylene nanocomposites that can show
significant improvement in strain hardening characteristics, an indispensible property for
foaming polymers.

184

8.2. Recommendations
Based on the results of this research, the following recommendations can be made for
future research:
a) It was found in this study that the use amino silane coupling agents in synthesizing
linear PP/clay resulted insignificant increase in strain hardening behavior of the
polymer melt under uniaxial extensional flow. Since, the amino groups predominantly
react with the maleic anhydride group of the compatibilizer, the use of surfactant can
be eliminated. Hence, nanoclays treated with just amino silanes could provide enough
interaction between the clay and the compatibilizer to produce strain hardening effect
in linear PP/clay nanocomposite melt.
b) The use of amino silanes (amino propyl trimethoxy silanes) has been well reported in
literature for improving the interaction between the clay and the compatibilizer.
However, there has not been a single report that mentions about the improvement in
the strain hardening characteristics of the polymer melt. This could be due to short
length of the linker (three carbon) between the amine and the silicon groups. In our
research, we have used silanes with five carbon linker length and observed significant
improvement in strain hardening characteristics of the resulting nanocomposites.
Hence, the use of longer linker length in the silane coupling agents needs to be
evaluated in order to study the effects of chain length, between the amine and the
silicon group, on the coupling between the clay and the bulk polypropylene and
thereby on its melt strain hardening property.
c) The dispersion of clay layers has always been challenging in polypropylenes. In this
study we used batch mixing process to achieve intercalation of the polymer chains
185

within the clay galleries. For successful implementation of this research for industrial
purposes, continuous extrusion process needs to be evaluated for dispersion of
nanoclays and melt strain hardening of the resulting nanocomposites.
d) This research presented a qualitative understanding of the formation of polymer/clay
entangles network using van-Gurp Palmen plots. A more quantitative approach could
be taken to understand the formation of such networks and its effect on strain
hardening behavior of the nanocomposites.

186

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