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GREEN BIOCOMPOSITES FROM POLYHYDROXY—
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GREEN BIO-COMPOSITES FROM POLYHYDROXY-
BUTYRATE-CO-VALERATE (PHBV), WOOD FIBER
AND TALC

By

Sanjeev Singh

A DISSERTATION
Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of

DOCTOR OF PHILOSOPHY
Packaging

2009

ABSTRACT

GREEN BIO-COMPOSITES FROM POLYHYDROXY—BUTYRATE-
CO-VALERATE (PHBV), WOOD FIBER AND TALC

By
Sanjeev Singh

Wood/Natural fiber plastic composites (W/NPCs) have been widely accepted and
are an enormously growing segment of the commodity plastics market. These are
predominantly being used in decking, household interiors, etc., and more recently have
made inroads in automotive interior applications. These are short fiber composites
developed by using wood/natural fiber as reinforcement in the plastic matrix. Generally,
plastic constitutes from 30 to 70% of the entire product mass of the composite and the
most used plastics are conventional petroleum based i.e., polyethylene, polyvinyl
chloride, polypropylene and nylon. Polyhydroxyalkanoates (PHAs) are bacterially
produced plastics from commonly available biomass, and are projected to be sustainable
and viable alternatives to petroleum based plastics. This research focuses on a holistic
approach to shift W/NPCs from partially to fully green biocomposites, by replacing the
conventional plastic constituent in the W/NPCs with a renewable source based and
biodegradable bioplastic i.e., PHAs. This broadly addresses the evolving environmental
concerns related to conventional petro-plastics. In this process of transformation from
partially to fully green bio-composites, issues coupled to biopolymers based W/NPCs
such as, cost, processablity, performance, compatibility of reinforcing agent and matrix,
and raw material availability are addressed. These fully green biocomposites can find

potential applications in rigid packaging, distribution, household items and automotive

interiors.

This dissertation advances in four consecutive syllogistic stages:

1. In the first stage, the development of wood fiber reinforced bio-plastic i.e.,
polyhydroxybutyrate—co-valerate (PHBV) bio-composites is accomplished on the
performance evaluation, and comparison to conventional WPCs is addressed.

2. In the second stage, bio-composites from natural fiber, i.e., bamboo fiber, and
PHBV, at 30 & 40 wt% fiber content, were fabricated and their analogy to wood
fiber reinforced PHBV bio-composites is determined.

3. In the third stage, a comparative study of anisotropy, static and dynamic mechanical
evaluation of biocomposites fabricated using two different molding processes, i.e.
injection and compression molding, was accomplished.

4. In the final stage, “green hybrid bioplastic composites” containing synergistic
reinforcements of talc and wood fiber were designed and fabricated.

The biocomposites were extrusion processed followed by injection molding and
analyzed for static and dynamic mechanical, thermal & morphological aspects. The
tensile and flexural modulus of the bio-composites with 40 wt% of the wood fiber was
enhanced by ~167%, and the heat deflection temperature by 21%. Statistically, there was
no effect of the fiber type (i.e., bamboo and wood fiber) on the mechanical properties,
except for notch impact strength and heat deflection temperature. The squeeze flow test
revealed the anisotropy in the injection molded short fiber bio-composites. Synergetic
reinforcement of the tale platelets and cylindrical wood fibers increased Young’s and
flexural modulus, by ~200%, at 20 wt% of each in PHBV. The high surface energy of
talc platelets allows good talc-PHBV interfacial interactions as shown by the

morphological study and no additional crystallization of PHBV was observed.

ACKNOWLEDGMENTS

I would like to express my thanks to my major supervisor Dr. Mohanty for his
guidance and support during the course of this study. His perseverance to hard work
had been a great source of motivation during the progress of this research.

My special thanks to Dr. Sher Paul Singh, who has always been there for his
support during my stay in School of Packaging. His kind and giving attitude would
always be indebted.

I express my greatest gratitude to Dr. Pourboghrat, a person of great humility,
intelligence and mentorship. I would always carry his wisdom and guidance with me as
an example of great teacher and professionalism.

I also express my utmost thanks and best regards to Dr. Harte for all his support,
guidance and suggestions during the course of study as a graduate student in School of
Packaging.

This research was made possible through the funding from USDA-CSREES
under the McIntire-Stennis program of the Michigan Agricultural Experiment Station,

Michigan State University. This support is greatly appreciated.

iv

TABLE OF CONTENTS

LIST OF TABLE ....................................................................................................... viii
L IST OF FIGURES .................................................................................................. ix
INTRODUCTION ..................................................................................................... 1
References .................................................................................................................. 5
CHAPTER 1: Literature Review ............................................................................... 6
1.0: Introduction ........................................................................................................ 6
1.1 Polyhydroxyalkanoate (PHA) Biopolymers ........................................................ 16
1.1.1 Synthesis and Production ..................................................................... 16
1.1.2 Structure and Properties ....................................................................... 23
1.1.3 Methods of modification and processing .............................................. 32
1.1.3.1 Copolymerization .................................................................... 32
1.1.3.2 Plasticization ............................................................................ 34
1.1.3.3 Blending .................................................................................... 36
1.2 Wood Plastic Composites ................................................................................... 41
1.2.1 Conventional Wood Plastic Composites ......................................... 41
1.2.2 Wood as a natural reinforcing component ......................................... 46
1.2.3 Talc as a reinforcing inorganic filler/component ............................. 49
1.3 Reference ..................................................................................................... 52
CHAPTER 2: Wood fiber reinforced bacterial bioplastic composites: Fabrication
and Performance Evaluation .............................................................. 62
2.0 Introduction ........................................................................................................ 62
2.1 Experimental Section .......................................................................................... 64
2.1.1 Materials .............................................................................................. 64
2.1.2 Fabrication of composite employing micro-compounding
Technique .......................................................................................... 64
2.1.3 Testing and Characterization ............................................................... 66
2.1.3.1 Mechanical properties ........................................................... 66
2.1 .3 .2 Therrnomechanical properties:
2.1.3.2.1 Dynamic Mechanical Analysis ............................... 66
2.1.3.2.2 Heat deflection temperature ................................... 66
2.1.3.2.3 Coefficient of linear thermal expansion ................. 66
2.1.3.3 Therrnogravimetric analysis ................................................. 67
2.1.3.4 FTIR analysis ......................................................................... 67
2.1.3.5 .Morphological studies .......................................................... 67
2.2 Results and Discussion ........................................................................................ 68
2.2.1 Tensile Properties ................................................................................ 68
2.2.2 F lexural properties ................................................................................ 75
2.2.3 Dynamic Mechanical Properties .......................................................... 77
2.2.4 FTIR analysis ....................................................................................... 83
2.2.5 Heat Deflection Temperature ............................................................... 85

2.2.6 Coefficient of linear thermal expansion ............................................... 85

2.2.7 Thermogravimetric analysis ................................................................. 88
2.2.8 Impact Strength ................................................................................... 91
2.2.9 Morphology ......................................................................................... 93
2.3 Conclusion .......................................................................................................... 97
2.4 References ........................................................................................................ 98

CHAPTER 3: Renewable Resource based Biocomposites from Natural fiber
i.e., Bamboo fiber, and Polyhydroxybutyrate-co-valerate (PHBV)
Bioplastic and its analogy to Wood fiber reinforced bioplastic

composites ........................................................................................ 1 01
3.0 Introduction ........................................................................................................ 101
3.1 Experimental Section .......................................................................................... 103
3.1.1 Materials .............................................................................................. 103

3.1.2 Fabrication of composite employing micro-compounding
Technique .......................................................................................... 103
3.1.3 Testing and Characterization .................................... 105
3.1.3.1 Mechanical properties ............................................................ 105

3 .l .3 .2 Thermomechanical properties

3.1.3.2.1 Dynamic Mechanical Analysis ............................... 105
3.1.3.2.2 Heat deflection temperature .................................... 105
3.1.3.3 Thermal Properties ................................................................. 106
3.1.3.3.1 Thermogravimetric analysis ................................... 106
3.1.3.3.2 Differential Scanning Calorimetry ......................... 106
3.1.3.4 FTIR analysis ......................................................................... 106
3.1.3.5 Morphological studies ........................................................... 106
3.1.3.6 Statistical Analysis ................................................................. 107
3.2 Results and Discussion ........................................................................................ 107
3.2.1 Tensile Properties ................................................................................ 107
3.2.2 Flexural properties ................................................................................ 113
3.2.3 Dynamic Mechanical Properties .......................................................... 115
3.2.4 FTIR analysis ....................................................................................... 120
3.2.5 Thermogravimetric analysis ................................................................ 123
3.2.6 Heat Deflection Temperature ............................................................... 125
3.2.7 Impact Strength .................................................................................... 126
3.2.8 Morphology .......................................................................................... 127

3.2.9 Analogy of Bamboo fiber-PHBV and Wood fiber-PHBV
biocomposites ..................................................................................... l 30
3.3 Conclusion .......................................................................................................... 143
3.4 References ........................................................................................................... 144
CHAPTER 4: Anisotropy .......................................................................................... 146
4.0 Introduction .......................................................................................................... 146
4.1 Experimental Section ......................................................................................... 148
4.1.1 Materials ............................................................................................ 148
4.1.2 Fabrication of composite ...................................................................... 148

vi

4.1.2.1 Injection molding ................................................................... 148

4.1.2.2 Compression molding ............................................................ 149

4.1.3 Testing and Characterization ................................................................ 150

4.1.3.1 Mechanical properties ............................................................ 150

4.1.3.2 Dynamic mechanical analysis ................................................ 150

4.1.3.3 Anisotropy analysis .............................................................. 150

4.1.3.3.] Squeeze flow test .................................................... 150

4.1.3.4 Morphological studies ........................................................... 151

4.2 Results and Discussion ........................................................................................ 153

4.2.1 Tensile properties .................................................................................. 153

4.2.2 Flexural Properties ................................................................................ 155

4.2.3 Dynamic Mechanical Analysis ............................................................. 156

4.2.4 Squeeze flow testing ............................................................................. 160

4.2.5 Morphology ......................................................................................... 166

4.3 Conclusion ........................................................................................................... 166

4.4 Reference ............................................................................................................. 170
CHAPTER 5: Hybrid Green Composite from Talc, Wood fiber and Bioplastic:

Fabrication and Characterization ....................................................... 172

5.0 Introduction .......................................................................................................... 172

5.1 Experimental Section ......................................................................................... 174

5.1.1 Materials ............................................................................................... 174

5.1.2 Fabrication of composite ...................................................................... 174

5.1.3 Testing and Characterization ................................................................ 176

5.1.3.1 Mechanical properties ............................................................ 176

5.1.3.2 Dynamic Mechanical Analysis .............................................. 176

5.1.3.3 Coefficient of linear thermal expansion ................................ 176

5.1.3.4 Heat deflection temperature ............................................... 176

5.1.3.5 Differential scanning calorimetry .......................................... 177

5.1.3.6 Morphological studies .......................................................... 177

5.2 Results and Discussion ........................................................................................ 178

5.2.1 Tensile Properties ................................................................................ 178

5.2.2 Flexural properties ................................................................................ 181

5.2.3 Dynamic Mechanical Properties ........................................................... 187

5.2.4 Differential Scanning Calorimetry ....................................................... 193

5.2.5 Impact Strength ..................................................................................... 196

5.2.6 Morphology .......................................................................................... 197

5.2.7 Thermomechanical Properties .............................................................. 200

5.3 Conclusion .......................................................................................................... 202

5.4 References ............................................................................................................ 203

CHAPTER 6: Conclusion and Future Scope ............................................................. 206

vii

Table 1.1

Table 2.1

Table 2.2

Table 3.1

Table 3.2

Table 3.3

Table 3.4

Table 3.5

Table 4.1

Table 5.2

Table 5.2

Table 5.3

Table 5.4

Table 5.5

LIST OF TABLES

Listing 29 values corresponding to d-spacing of PHBV lattice.
Adapted from reference [24] ............................................................ 28

Processing parameters of various PHBV-wood fibers
Compositions ..................................................................................... 65

Value of coefficient “C” with different loading levels of
Wood fiber ...................................................................................... 78

Processing parameters of PHBV and fiber compositions ................. 104

Description of Bamboo and Wood‘ fiber based PHBV
biocomposites properties ................................................................... 131

Description of P and F values ........................................................... 132

Detailed Information Obtained from Differential Scanning
Calorimetry of PHBV and its Composites ......................................... 139

Lignin and Cellulose content of Wood and Bamboo fiber ................ 142

Processing parameters of the fabricated PHBV and it

wood fiber loaded composites ........................................................... 149
Processing parameters of the composites .......................................... 175
Density of the PHBV and its composites .......................................... 179
Surface energy parameters of PHBV Maple wood ............................ 185
Calculated Interfacial energy and Termodynamical Work of Adhesion

........................................................................................................... 187
Details of DSC of PHBV and its composites ................................... 194

viii

Figure 1.1
Figure 1.2

Figure 1.3

Figure 1.4

Figure 1.5

Figure 1.6

Figure 1.7
Figure 1.8
Figure 1.9

Figure 2.1

Figure 2.2

Figure 2.3

Figure 2.4

Figure 2.5

Figure 2.6

Figure 2.7

Figure 2.8

LIST OF FIGURES

Sketch of long and continuous fiber composite ................................ 7
Sketch of randomly oriented sort fiber composite ............................. 8

Stress Equilibrium on fiber aligned parallel to applied load.
(Redrawn from reference [6]) ........................................................... 14

Sketch of the tensile stress acting at the fiber length and
shear stress acting at the interface (Redrawn from reference [6]) ..... 15

Enzymatic Path of PHB synthesis (Redrawn from reference) ........ 17

PHA production process
(Redrawn from source: Bioinformation Associates, Boston MA) ....22

Structures of different PHAs ............................................................. 23
FTIR spectra of PHBV .................................................................... 24
Schematic of coupling in interface region ......................................... 44

Tensile properties of PBV-wood fiber composites with different
loading level of wood fiber .............................................................. 70

Comparison of Tensile modulus; Halpin-Tsai and Tsai-Pagano
relation data experimental data points ............................................... 73

Comparison of Tensile strength Nicolais and Nicodemo

relation data points and experimental data points .......................... 74
Flexural properties of PHBV-wood fiber composites with

different loading level of wood fiber ............................................... 76
Temperature dependency of storage modulus with varying

loading level of wood fiber in PHBV matrix. ................................... 80
Temperature dependence of Loss Modulus with varying loading

level of Wood fiber in PHBV matrix. ............................................... 81
Temperature dependency of Tan Delta with varying loading

level of Wood fiber in PHBV matrix. ............................................... 82
F TIR spectra of PHBV and its composites ...................................... 84

ix

Figure 2.9

Figure 2.10

Figure 2.11

Figure 2.12

Figure 2.13

Figure 2.14

Figure 3.1

Figure 3.2

Figure 3.3

Figure 3.4

Figure 3.5

Figure 3.6

Figure 3.7:

Figure 3.8

Figure 3.9

Figure 3.10

Effect of wood fiber content on heat deflection temperature and

coefficient of linear thermal expansion values of PHBV-wood fiber

composites ....................................................................................

Thermogravimetric analysis of PHBV, wood fiber and their

Composites ...................................................................................

..... 87

..... 89

Derivative weight loss of PHBV, wood fiber and their composites..90

Notch Impact strength of PHBV-wood fiber composites with

different loading levels of wood fiber ..........................................

Low magnified SEM photomicrographs of impact - fractured

samples of PHBV-wood composite ..............................................

High Magnified SEM photomicrographs of impact-fractured

samples of PHBV-wood ...............................................................

Tensile properties of Bamboo fiber-PHBV composites at

30 and 40 wt% fiber ......................................................................

Experimental and Theoritical data plots .......................................

Flexural properties of Bamboo fiber/PHBV composites

having 30 and 40 wt% fiber ..........................................................

Storage modulus vs temperature of PHBV with varying loading

levels of bamboo fiber ..................................................................

Tan Delta vs temperature of PHBV with varying loading

levels of bamboo fiber ..................................................................

Loss modulus vs temperature of PHBV with varying loading

levels of bamboo fiber ..................................................................

FTIR spectra of (a) PHBV (b) PHBV/Bamboo fiber Composite

Percentage weight loss of PHBV, PHBV/Bamboo fiber (70:30)

and bamboo fiber ..........................................................................

Derivative weight loss of PHBV, PHBV/Bamboo fiber (70:30)

and bamboo fiber ..........................................................................

HDT of Bamboo fiber/PHBV composites ....................................

..... 92

..... 95

..... 96

..... 111

..... 112

..... 114

..... 117

..... 118

..... 119

..... 122

..... 124

..... 124

..... 125

Figure 3.11

Figure 3.12

Figure 3.13

Figure 3.14

Figure 3.15

Figure 3.16

Figure 4.1

Figure 4.2

Figure 4.3

Figure 4.4

Figure 4.5

Figure 4.6

Figure 4.7

Figure 4.8

Figure 4.9

Notch impact strength of PHBV/ Bamboo fiber composites ........... 126

SEM photomicrographs of impact-fractured samples of
PHBV/Bamboo (70:30) composite (a) 90X, 200 um,
(b) 450X, 50 um ................................................................................. 128

SEM photomicrographs of impact-fractured samples of
PHBV/Bamboo (60:40) composite (a) 90><, 200 um,
(b) SOOX, 50 um ................................................................................. 129

SEM Photomicrographs at high magnification
(a) PHBV/ Wood fiber composite 1500X, 20 um ............................. 135

Crystallization and melting curves of
(a) PHBV (b) PHBV/Bamboo fiber (70:30) composite

(c) PHBV /Wood Fiber (70:30) composite ........................................ 138
FTIR spectra of wood and bamboo fiber ......................................... 141
Illustration of Squeeze flow .............................................................. 152

Tensile properties of PHBV and its composite loaded
with 30 wt% wood fibers ................................................................. 154

Flexural properties of PHBV and its composite loaded
with 30 wt% wood fibers ................................................................. 155

Storage modulus of PHBV and its 30 wt% wood fibers
loaded composite with varying temperature ..................................... 157

Loss modulus of PHBV and its 30 wt% wood fibers
loaded composite with varying temperature ...................................... 159

Tan 5 of PHBV and its 30 wt% wood fibers loaded
composite with varying temperature ................................................. 159

Squeeze flow test plots of 20 wt% fiber loaded
injection molded composites ............................................................. 162

Squeeze flow test plots of 40 wt% fiber loaded
injection molded composites ............................................................. 163

Squeeze flow test plots of 30 wt% fiber loaded
compression molded composites ....................................................... 165

xi

Figure 4.10

Figure 4.11

Figure 5.1
Figure 5.2
Figure 5.3
Figure 5.4
Figure 5.5
Figure 5.6
Figure 5.7

Figure 5.8

Figure 5.9

Figure 5.10

SEM micrographs of 30 wt% wood fiber loaded PHBV
compression molded composite ........................................................ 168

SEM micrographs of 30 wt% wood fiber loaded PHBV

injection molded composite ........................................................... 169
Tensile Properties of PHBV and its Composites ............................... 180
Flexural Properties of PHBV and its Composites .......................... 182
Storage Modulus of PHBV and its filled composite ......................... 190
Loss Modulus of PHBV and its filled composite ............................. 191
Tan 6 of PHBV and its composite ..................................................... 192
DSC curves of PHBV and its composites .......................................... 193
Impact strength of PHBV and its hybrid composite .......................... 196

SEM Photomicrographs of fractured hybrid composite surface
(a) Low magnification 270><, 100nm
(b) High magnification 3000X, 10pm. ............................................... 198

SEM Photomicrographs of Hybrid composite
(a) Wood fiber pulled out interface with PHBV matrix,2700><, 10pm
(b) High magnification fractured surface of composite,5500><, Sum.199

Coefficient of linear thermal expansion and
Heat distortion temperature of PHBV and its hybrid composite ....... 201

xii

INTRODUCTION

In the perspectives of the environmental concerns and fluctuating oil prices,
consumer markets all across the world have initiated a drive for bio based material as
an alternatives for holistically, sustainable economic growth. Automotive industry in
Europe and America have, by now, embraced the rationality and feasibility of replacing
the plastic components with the biomaterials, and are estimating an opportunity of more
than 50 billion US $ by 2015 [1,2,3]. Medical world have also ventured in to exploring
the use of biomaterial in medical application, and the projected market is more than 100
billion US $. The forecast to grow in the near future is by 14% [4,5]. Biomaterials for
other applications, such as, industrial and structural, are also being explored for the
sustainable economic growth, in conjunction with green environment. Biopolymers
from the biomass, like polylactic acid, starch based plastic, Polyhydroxyalkanoates, and
other bio-oil based resin are being envisaged as a prospective alternate to their
counterpart in the olefin plastics. PLA and starch plastic products are already in the
consumer market for textile, packaging, foam etc., applications [6,7]. PHAs, bioplastics
have yet to meet the challenges based on cost and inherent characteristics, and make
inroad as a considerable probable. Polyhydroxybutyrate-co-valerate (PHBV), is an
intracellular bacterially produced biodegradable bioplastic from a renewable source,
i.e., biomass like corn, switch grass etc. The current challenges it posses for

00mmacialization are cost, brittleness and processing. The objective of this research
was to develop an all green biomaterial from PHBV utilizing the conceptual
fundamentals of wood plastic composite (WPC) in processing/manufacturing and

applications. Besides, this research also takes into the consideration of the trio, cost-

 

characteristics-processing, co-relationship and optimizing the three, most appropriately,
during the development of the biomaterial. These fully green biocomposites from
various bio-sources can run the gamut of applications in rigid packaging, distribution,
household items and automotive interiors.

WPC are the fastest growing segment in the plastic industry with their ongoing
growth of 25% [8]. The current market of wood plastic composites (WPC) is 1.2 billion
pounds and is expected to grow with an average annual growth rate of 9.8%, giving a
market of 1.9 billion pounds by the year 2009 [9].The wide acceptance of W/NF PC by
the global market has been due to their reduced cost, weight, superior properties over
wood and plastic alone, recycling ability, adaptation to the existing plastic processing
techniques, and the desire for the eco—friendly products.

This thesis is divided in to six chapters, beginning from the literature survey in
the first chapter to the development of hybrid all green bio composite from
polyhydroxy-co-valerate (PHBV) in the fifth chapter. The subsequent second, third, and
fourth chapters proceed in a logical sequence based on the illations drawn from the
preceding chapter. The second chapter targets the development and characterization
of the biocomposite from wood fiber and PHBV. The third, describes the development
and characterization of biocomposites from a natural fiber i.e., bamboo fiber, and
PHBV, and comparison of the two, PHBV based, biocomposites from wood fiber and
natural fiber. The fourth chapter illustrates on the isotropy of the short wood fiber
biocompostes processed with two different processing methods. The final and sixth

chapter summarizes the research.

The first chapter gives the description from the literature regarding the
conventional short wood and natural fiber plastic composite, biopolymers and the other
two components of the hybrid composites, i.e., wood and tale. The literature collected
from various sources provides the current and the future scenario of the wood plastic
composites in the consumer market, in terms of demand and growth. The economic
viability discussion of the WPC is followed by the section on the mechanics of the short
fiber reinforced composites that details the understanding of the isotropic characteristic
of the short fiber composites. The interfacial shear stress based on the shear lag model
is described as well. The interface of wood and plastic plays a vital role in the structure
of wood plastic composites, and the literature also focuses on various kinds of
interfacial phenomenons. A subsection of the second chapter focuses on the synthesis,
manufacturing, structure, properties, limitations and modification of
polyhydroxyalkanoates. Wood fiber and talc are the two different kinds of
reinforcements that had been used to reinforce the thermoplastic, i.e.,
polyhydroxybutyrate-co-valerate, therefore it becomes imperative to discuss their
fundamental structures and properties. The last two sections of this chapter focus on
these.

The second chapter describes the fabrication of a wood fiber reinforced PHBV
composites and evaluate the static and dynamic mechanical and thermo-mechanical
properties as a function of wood fiber weight content in the PHBV matrix. The
experimental findings were correlated with theoretical models presented by Halpi-Tsai
and Tsai-Pagano. The morphological aspects such as fiber dispersion and fiber-matrix

interface were also investigated using scanning electron microscopy (SEM).

The third chapter illustrates the fabrication and evaluation of the bamboo fiber
based PHBV biocomposites, and also describes the rationality of bamboo fiber quid pro
quo to wood fiber. Characteristic study of the bamboo based biocomposites using
differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), Fourier
transform infrared spectroscopy (FTIR), and their morphology is also illustrated. The
experimental results of the tensile modulus were compared with the persisting
theoretical model and additionally, comparative analysis by performing a two-way
ANOVA on the mechanical properties of the bamboo fiber biocomposites and the wood
fiber-PHBV biocomposites has been discussed.

The fourth chapter discusses the comparison between the two different
processing techniques, i.e., injection molding and compression modeling of the wood
fiber composites from PHBV, in terms of the anisotropy associated with them. This
chapter also draws a direct comparison of the static and dynamic mechanical properties
between the compression molded specimens and injection molded specimens of the
same biocomposite.

The fifth chapter elaborates the development of a completely
biodegradable hybrid biocomposite consisting of talc and wood fibers in PHBV. It also
discusses the experimentally determined dynamic-mechanical and thermo-mechanical
properties of the hybrid biocomposites. A theoretical discussion, based on the surface
energy parameters, of the interacting components in the composite system is described
to explain the reinforcing effect of talc and wood fiber. The Morphological analysis of
the hybrid composite was carried out to study the interfacial interactions among the

different components in the hybrid biocomposite system.

References

1) Crawford, C. Ontario Bioauto Council, Presentation. 4/11/08. Presentation available
at; http://www.sobin.ca/files/BioAuto%20Council%20-%20Apri1%20l L%202008.pdf.
Date accessed; 12/8/08.

2) Khalil, H. Woodbridge Foam Corporation, “The use of bio material in automotive
parts manufacturing”, Presentation to Canadian conference on industrial bioproduct
innovation, Montreal. 1 1/6/08. Available at;
httJ)://www.biop.ca/PPT_pdf/KHAI.1L_Hantdef. Date accessed; 12/08/08.

3) Cascadden, A. M. “Bio—Materials — Why is the Auto Industry interested?” AUTOZI
NC E, Toronto. Ontario, July 12, 2006, available at;
http://wwwwagrifoodforum.com/2006/presentations/AFIFO6-05-Cascadden.pdf. Date
accessed; 12/08/08.

4) Amsden, B. CHEE 340 course lecture, Queens University. Available at;

hgpz/lwwwchemengqueensu.ca/courses/C HEE340/1ectures/documents/biomaterials1 2
006.ppt#264,LBIOMATERIALS. Date accessed; 12/08/08.

5) Grarnmenou, E. Medical Devices Research Analyst, “Technology development
drive the market in Europe”, published on 5/4/2006, available at;
http://www.frost.com/prod/servlet/market—insight-toppafldocid=6852598 1 , Date
accessed; 12/08/08.

6) Available at; http://wwvw.ides.com/generics/PLA/PLA_applicationshtm, Date
accessed; 12/8/08.
7) Available at; http://www.starch.dk/isi/applic/applic.htm, Date accessed; 12/8/08.

8) Principia partner’s presentation in, 7th international conference of wood fiber-
plastic composites 2003. Madison (WI), USA.

9) Available at; Hltp.'//Www.Principiuconsu/Iing.(Torn/"(7027.9tilting/News.Cfin. Date
accessed 05/ 1 5/2006,

Chapter 1

Literature Review

1.0 Introduction

Wood or natural fiber plastic composites predominantly exist as short fiber
composites in most of the applications. Time and again, since their inception in 1990’s,
these have occupied ever-increasing share of market year by year. In the future
prospects, W/NPC’s have the potential to grow up to $ 60 billion market, from the
existing just over $ one billion US market [1].Their modern-day principal application
fields in North America are building structural products such as, decking, fencing,
railings, window and door profiles, and shingles, where they constitute over 85% of the
total demand. Other application arenas like, infrastructure, automotives and
industrial/consumer segment are also catching up with good pace. In Europe, the
automotive application have an edge over North America, where the W/NPC’s account
for over half of the total consumption [2]. A boom in the W/NPC’S market has also lead
to 3 grth explosion in the injection molding industry, as the extensively used process
for W/NPC’s manufacturing has been extrusion, and this is supplemented by injection
molding. Anticipated demand for the injection molded wood-plastic composites is 70%
average annual growth, which was supposed to reach a market of over $350 million in
2008 [3, 4]. The driving impulse for this industry is that W/NPC’s lend a hand in
reducing the dependence on the petroleum based plastics by limiting the use of plastics
in them. The fiber constitutes from 30 to 70 wt% in the composite, and the major
cellulose fiber used in these composites are wood fiber (pine or maple saw dust),

recycled newsprint, hemp, sisal, kenaf, jute, flax or rice hulls. The scarce of the natural

or wood fiber, as a long and continuous, but in abundance as a short length and low-
price, presented an opportunity to use these as short fiber reinforcement in the matrix,
which has been plausible with the available plastic processing technology. The leading
thermoplastics used are PE, PVC, PP, and nylon. Unlike, long fiber composites, which
are termed as orthotropic, the short fiber composites are assumed to be isotropic in
nature, due to their randomly oriented dispersion in the matrix. The mechanics of the
isotropic materials can be well appreciated by assuming these to be a simplified case of
orthotropic composites, where the stress-strain response of the material can be assumed
to equal/similar along the x, y and z axis. Figure 1.1 and 1.2 represents the schematic

sketches of long and short fiber reinforced thermoplastic composite.

 

 

Figure 1.]: Sketch of long and continuous fiber composite

 

Figure 1.2: Sketch of randomly oriented sort fiber composite

Mechanical properties of an anisotropic material are defined by the following

stiffness and compliance matrix.

 

 

 

 

 

 

'q‘ "a. q. a. 0 0 0"«2‘
05 Czr Cu (:23 0 O 0 82
<05}: (3161.36.30 0 0 3.,
2;, O O 0 C44 0 O 723
G3 0 0 0 0 C55 0 7,3
5,, _0 0 0 0 0 C664 [91,)

’g,‘ ’5“ 512 5,, 0 0 o‘ia,‘
52 S21 S22 S23 0 O 0 0'2
<5, >= 5,1 5,, 5,, 0 0 0 <03}
y,, 0 0 0 S44 0 O 7,,
7,, 0 0 0 0 S55 0‘ 2'],
.712) _ 0 O O 0 0 S66_ [7121

 

 

 

 

 

 

Where 8i , 0'1- , Tij , yij , are strain, stress, shear stress and shear strain, respectively,

along i,j=1,2,3 directions. [Sij] and [C ij] are compliance and stiffness matrices defined

as.
Sn = 1/E1, 512 = -l'12/EI. 513 = - 1'13 /E1
522: I/EI' 523 =-l'12/E2, S33 = I/E]
S44 = 1/G23. 555 = 1/G13. 566 = U012
Ei and 0,} are elastic and shear modulus, ii]- is poison’s ratio in their respective
directions.
C11 = (522 S33 - S22 S33)/5 C12 = (513 523 - 512533)/S
C22 = (533511 — 513 513)/5 C 13 = (512523 — S13522)/S

C33 = (511522 - 512512)/S C23 = (512513 - 523511}/S

C44 = 1/544 C55 = 1/ 555 C66 = 1/566

For Isotropic material, imposing the condition that it has same properties in all

the directions, therefore
EI=E2=E3=E, i13=i23=i12=13
G12: 023: 013: G =E/2(I+ i)

The compliance matrix become

 

 

5.. 5.. 5.. 0 0 0
512 SH 812 0 0 0
S12 S,2 SH 0 0 0
0 0 0 S44 0 0
0 0 0 0 5,, 0
_ 0 0 0 0 0 5,,_
where511=I/E, S12Z'i/E, S44=I/G=2(I+ 1')/E
and the stiffness matrix becomes
"C” C,,_ C,, 0 0 0 ‘
Cl2 Cll Cl2 0 0 0
Cl2 C,2 CH 0 0 0
0 0 0 C,, 0 0
0 0 0 0 C4, 0
_ 0 0 0 0 0 C4,

 

 

10

where C11 =E/((1+ i)(1+2 i)), C12 = ((1- i)E)/((I+ 1')(1+2 1'»,

C44=G=E/(2(1+ 1?)

As elaborated by Bhagwan and Lawrence [5], a good estimate of average
mechanical properties of short fiber composites can also be obtained by assuming them
as a quasi-isotropic laminate made from unidirectional laminae. Quasi-isotropic
laminate has an elastic coefficient that is independent of continuous fiber orientation in
the plane. Other models for the prediction of properties are described in the subsequent
chapters.

The interface properties between fiber and matrix plays a vital role while
describing the micro- structure of the composite and therefore, it becomes necessary to
discuss the stress transfer mechanism from matrix to fiber along its length. The shear
lag model is the most widely used model, which is based on the transfer of tensile stress
from matrix to fiber by means of interfacial shear stress. This model is based on
considering the interfacial bonding between the matrix and fiber. According to the
shear lag model, the tensile stress experienced by the fiber, aligned along the external
axial loading, is given by the following equation, neglecting the stress transfer at the

fiber ends.

of = E f8, [1 — cosh(nx / r) sec h(n5)
The modified equation that takes in to account of the stress transferred at the

fiber end is as follows.

11

of = e,[Ef — (Ef — E;)cosh(nx/r)sec h(n5)]

 

E' = [Ef(1— sec h(n5) + Em]

Where, "7 2

l/2

2E

m

Ef(1+vm)ln(1/f)

 

‘ 0- f ’is the tensile stress in the fiber along the length ‘x’, ‘r’ is the radius of

the fiber, ‘5’ is the aspect ratio i.e., L/r, ’L ’is the length of the fiber. ‘Ef’ and ‘Em ’ are

the elastic modulus of fiber and matrix, respectively. ‘f’ is the fiber volume fraction in

matrix and ‘v m’ is the matrix poison ratio. ‘8 1’ is the over all composite strain.

The interfacial shear stress, “C i, along the fiber length is given by

118
_ l '
r, — 2 E f srnh(nx / r) sec h(n5)

The tensile stress or the axial stress in the fiber along the fiber length attains
maximum in the mid section of the fiber. It decreases gradually towards the fiber ends,
as depicted in the figure 1.3 and 1.4. For higher aspect ratio fibers, the tensile stress
forms a plateau over a considerable length of the fiber before plummeting to its

minimum values at the fiber ends. The interfacial shear stress reaches maxima near the

12

fiber ends due to the shear deformation of the matrix under an externally applied, axial
load, along the fiber length, to the composite. The interfacial shear stress in the mid
section of the fiber has minima due to low shear deformation of the matrix in the
middle section of the fiber along its length. For higher aspect ratio fibers, the minimum
interfacial stress values continuous to a considerable length of the fiber before attaining
the maxima at the fiber ends. Therefore, higher aspect ratio fibers are more efficient in
stress transfer than the short fibers because these have sufficient length to attain the

maximum tensile stress in the fiber or in other words to transfer the stress from matrix

to fiber. Stress transfer aspect ratio ‘5,’ is defined as the minimum aspect ratio of the

fiber, when the stress reaches its maximum possible value. It is given as

Stz 3/71

Stiffer the matrix shorter is the St, and stiffer the fiber longer is the St,

As the composite strain increases, the stress in the fiber also increases with the
increased interfacial shear stress, and subsequently the fiber undergoes fracture. A

reduced fiber length does not allow the stress to reach it maximum value. Therefore, a

critical aspect ratio ‘Sc’ of a fiber can be defined as the minimum length below which

the fiber does not undergo fracture, i.e., the value of tensile stress in fiber when it just

at:
reaches its ultimate strength ‘O'f ’ .

13

Of f

i |

 

 

 

 

 

 

 

of: i l

Gf+ de

 

 

 

 

 

 

Figure 1.3: Stress Equilibrium on fiber aligned parallel to applied
load. afis the tensile stress on fiber, 1' is the shear stress, I is the length
of fiber, ac is the stress acting on composite ( Redrawn from reference

[6])

5, = (Sf/2 of

*
G; is the critical shear stress at which the onset of inelastic strain/behavior

occurs in the composite.
Generally, incase of composites having a fibers well above the critical aspect
ratio, as the composite load increases, sliding at the interface or yielding phenomena is

predominant before the fiber fractures [5, 6].

l4

 

i W

 

 

‘l

 

 

h J

 

Figure 1.4: Sketch of the tensile stress acting at thefiber length
and shear stress acting at the interface (Redrawn from
reference [6])

15

1.1 Polyhydroxyalkanoate (PHA) Biopolymers
1.1.1 Synthesis and Production

These belong to the family of microbial energy reserve materials that
are accumulated as granules within the cytoplasm of cells [7-10]. PHA accumulation in
microorganism gets triggered when it faces the metabolic stress such as limitation of
essential nutrients (e. g. nitrogen, oxygen and phosphorus) in excess of carbon
source[11]. PHAs are produced in various species of bacteria as spherical bodies. The
diameter of these inclusions in bacteria is roughly 0.5 pm [8]. Under regulated
conditions PHAs may constitute up to 90% of the biomass of some bacteria species.
The PHB synthesis mechanism in the bacteria Alcaligenes eutrophus is carried out as
follows. The first step constitutes the enzymatic self condensation of Acetyl CoA
(available from the fermentation of various feedstocks such as glucose and other simple
sugars, methanol, and acetic acid) to form acetoacetyl-CoA. Subsequently, in the
second step this compound is further enzymatically reduced to R-A-hydroxybutyryl-
CoA which is finally added to the growing PHB chain by PHB synthase in the third and
final step. Other PHAs can be produced by inducing bacteria to the varying growth
media. One such other PHA is polyhydroxyvalerate, PHV. The appropriate
concentration of propionic acid are added to the glucose feedstock in the fermenter to
induce bacteria to make a copolymer of PHB and PHV, polyhydroxybutyrate valerate
(PHBV) with a controlled ratio of hydroxyl- butyrate (HB) to hydroxyl valerate (HV)
monomers varying from 0 to 47 mole%. Higher HV fractions, up to 90 mole%, in the
copolyester can be attained by using pentanoic acid in the feedstock with glucose [9].

Although commercially not available, PHAs can be chemically synthesized.

l6

p yruva te 0

Acetyl-CoA

 

,B - ketothiolase l
0 0
II II
CoASH + CH, —C—CH2 —C—C-—SCoA
Acetoacetyl-CoA

AcetoacetyI-CoA
reductase

‘1’” if
CH3—C—CH2—C—C—SC0A

(R)—B-hydroxybutyryl-CoA

PHB

synthase

f”: it
[—CH—CH,—C—0—]n + CoASH

PHB

Figure 1.5: Enzymatic Path of PHB synthesis (Redrawn from reference)

This approach provides one to more easily and predictably accommodating variations

17

in molecular weight, co-monomer type, and co-monomer content with the availability
of pure monomers.
PHA copolymerization is a ring-opening polymerization of a-lactones as shown

following.

I
I,"

 

(o

R can be, C3H7, C5HH,C15H31

Initiator used to prepare the isotactic PHA copolymers is difunctional zinc compound,
ethylzinc isopropoxide (EZDP). Reaction temperature is reported to be at 60-65 °C
because the 3-alkyl-a-lactones are unstable at higher temperatures. Lactones become
more thermally labile with the longer the alkyl group, thus decomposing to carbon
dioxide and terminal alkene. Therefore, there is always a need of balance among the
polymerization temperature and reaction time with the thermal stability of the o-lactone
monomers. Incorporation of the co-monomers is well accommodated in the molecular
chain, and even at lower conversions of 30-50% the co-monomer incorporation is
greater than 50% of the feed ratio. Other initiators reported in literature for this type of
synthesis are aluminoxanes, distannoxanes and alkylzinc alkoxide. These initiator
system give PHA with molecular weights >30 000. The appreciable mechanical
properties of isotactic PHA polymers are only realized at the weight-average molecular
weights (Mw) of 500 000 or greater. The use of highly purified and highly

enantiomerically (enantiomeric excess > 90%) enriched monomers give isotactic PHA

l8

polymers and copolymers with Mw up to several hundred thousand. The differences in
the synthetic PHA are because these synthetic copolymers are not perfectly isotactic as
the starting monomer 3-alkyl-a-propiolactones had only 92-95% stereo-purity.
Stereochemistry plays a vital role in chemical synthesis as low levels of
stereoimperfections causes 10 °C drop in melting point for synthetic PHB compared to
the melting point of bacterial PHB. As reported in the literature this initiator system
performs well in bulk polymerizations as well as in solution polymerization with
toluene and tetrahydrofuran (THF) as a reaction media (in case of THF if the
copolymer produced is soluble in THF). The polydispersities (PDI) of the polymer can
be seen to be around 1.5 [12]. Utilizing the genetic engineering technique, genes in R.
eutropha responsible for all the three enzymes in synthesis of PHB were transfer to
other forms of bacteria Escherichia coli that have tenfold larger cell thus producing ten
times as much polymer. In another, isolated genes were added to an E. coli strain that
bursts at 42 0C thereby facilitating separation. Strains containing only the synthase gene
can express enzyme in sufficiently large quantities for isolation and purification. The
purified enzyme is stable in aqueous solution and was used for in vitro polymerization
reactions of a wide variety of 3- and 4-hydroxyalkanoate-COA monomers. In vitro
polymerization reactions forms “living polymers”, thus giving a polymerization process
with no chain termination reaction therefore, propagating end group remains active
indefinitely and very high molecular weight polymers can be prepared in vitro[13].

The genome sequence of the PHA producing bacteria, Ralstonia eutropha H16
had been reported by Pohlmann et al [14]. This opens new ventures for bioplastic

producer improving PHA biosynthesis. The decoding of genome sequence provides an

19

opportunities for optimizing PHA biosynthesis and for developing metabolic-
engineering strategies to generate bioplastics with improved properties[15].

In another the most exciting genetic engineering synthesis involves the transfer
of PHA genes in plants so that they will convert acetyl CoA to PHAs as they grow.
This accomplishment in the plant Arabidopsis thaliana came by a group of scientists
from our university (Michigan State University) and James Madison University in
1992. Monsanto in 1994 took initiative to produce PHAs in the leaves and stem of corn.
Thus, plastic production could, theoretically, be produced from plants. However, the
challenge of efficient energy and cost remains in separating the PHAs fi'om the plant
material and purifying it [7-13].

Production

The first prototype of PHAs, polyhydroxybutyrate (PHB), was discovered in
1927 at the Pasteur Institute in Paris and the first commercialization was done by W.R.
Grace Co. in the 19505. Recently, a series of PHA copolymers are being commercially
produced under the trade name of BIOPOL by Zeneca Bio Products (formerly ICI Bio
Products),[7]. More recently, Metabolix has announced new grades of PHBV under the
trade name of Mire] Natural Plastics for injection molding and paper coating. General
mass production of PHAs is done by fermentation process as illustrated in figure 1.6.
Extraction of PHAs from the bacteria is carried out in the final stage. This can be
accomplished by the following three techniques 1) Solvent extraction, 2) Sodium
hypochlorite digestion, and 3) Enzymatic digestion.

Solvent extraction is expensive and requires large volumes of solvent but

produces high purity and high molecular weight PHA. Benefit of sodium hypochlorite

2O

digestion is that it degrades and dissolves all cellular components of biomass except
PHAs, but it can also damage the polymer product and is difficult to separate from the
final product. The extraction process of enzymatic digestion involves the use of several
enzymes to maintain an appropriate temperature to dissolve DNA that is released along

with PHAs otherwise it becomes difficult to separate it from the viscous suspension [9].

21

Growth of bacteria and

Carbon Source

A

 

 

RAW MATERIAL

 

I

\

 

a; f ,

E. / CELL DISRUPTION

8‘ ll

”5

C'.

%{ FERMENTATION WASHING

(U

E

E 3 >

(U

( MEDIA CENTRIFUGATION
PREPARATION

fl

 

DRYING
Q J
PHA PRODUCT

Jaunfiod jo uorieogrmd

Figure 1.6: PHA production Process (Redrawn from source: Bioinformation
Associates, Boston MA)

22

1.1.2 Structure and Properties

«as O
CZ<—I/H\
O O

P(3 HB-co—HV)

 

CH3 n= 3-7

(CH2 )n

O

'5

CZ<}I/IJ\
O O

P(3HB-co—mcl—3I—IA)

 

CH3

«an. O
CZ<FI/U\
O O

P(3 HB-co-3 H H)

 

CH3

   
   
 

O

P(3HB-co—6HH)

 

P(3HB—co—4HB)

Figure 1. 7: Structures of different PHAs, (HB) hydroxy butyrate, (HA) hydroxy
valerate, (HH) hydroxy hexanoate.

23

The general structures of PHAs are shown in figure 1.7. By altering the
feedstocks and the microorganisms PHAs copolymers of different compositions can be
framed with properties ranging from highly crystalline, rigid plastics to rubbery

elastomers. The side chains can be of different functional groups or with different

number of methylene groups at a-position (3rd carbon from carbonyl group) or the

polymer’s main chain itself can have different number of carbon atoms. There are more
than hundreds of monomers reported as PHA constituents, but only few have been
commercialized. PHA’s thus produced by microbial synthesis bear a highly
stereoregularity in their molecular structure and all the chiral carbon atoms in the
backbone chain are in the R (-) configuration [10, I6]. PHB is isotactic with methyl

groups on carbon chain in a single conformation [17].

phbv

 

 

1750 1550 1350 1150 950 750
Figure 1.8: F TIR spectra of PHB V

24

The figure 1.8 above is the FTIR spectrum of PHBV and describes various

absorption bands by functional groups of the polymer. In the spectrum, band at 1720

ch was the C=O stretch of the ester group present in the molecular chain of highly

ordered crystalline structure[18]. A small hump at 1740 cm-1 was of C=O stretch in the
amorphous region of a semicrystalline PHBV. Characteristic peaks of symmetric

—C—O—C— stretching vibration were present from 800 cmil to 975 cm.1 and the
antisymmetric —C—O—C— stretching between 1060 cm"l and 1150 cm". Band around
1378 cm-I corresponds to the symmetrical wagging of CH3 groups and at 1453 cm'1

was due to asymmetric deformation of methylene groups, the intensity of both the

peaks is independent of crystallinity or in other words insensitive to cystallinity. The

absorption bands at 1720, 1276, 1228, and 980 cm-1 arose from the crystalline regions(

designated with arrows) and bands at 1740, 1453, 1176 cm.1 from the amorphous

regions (designated with double arrows, pointed both ways in opposite directions) of

PHBV. Peak at 1228 cm-1 was purely of the conformational band of crystalline helical

molecular chains and therefore are absent in amorphous region [19]. Crystallinity Index

(CI ) give the relative measure of degree of crystallinity in PHBV. It is given as

CI = Iatl378cm"

all 176cm—l

 

(ratio of intensity of the band insensitive to

crystallization to the band sensitive to crystallization), where I is intensity of the peak

at a given wave number.

25

13C and 1H NMR of PHB reveals the details of its structure by elucidating the

environment of carbon and hydrogen atoms, respectively in the molecule. The peak at
about 69.6 ppm is assigned to the a carbons that are at branching of methyl or ethyl
groups. Similarly, the peak at 42.8 ppm is attributed to a carbon neighboring the
carbonyl groups and peak of carbonyl carbons is observed at 171 ppm. Branched

methyl group carbons of PHB show peak at 21.3 ppm and in case of PHBV branched

ethyl has two carbon atoms (CH3CH2), carbons corresponding to CH3 show peak at

about 10 ppm and that corresponding to CH2 shows at about 37 ppm (intensity of

peaks very small). Intensity of peaks corresponds to the % of valerate (HV) content in
the PHBV copolymer [20]. The sharpness of peaks or signals may arise from the

orderliness of the structure therefore, an estimation of the crystalline regions can be

made using NMR. For further details reference can be consulted [21]. In 1H NMR

spectra, the methine proton attached to the asymmetric carbon or or carbon shows

multiplet at 5.25 ppm. Proton or hydrogen atoms on the methylene (CH2) shows

multiplet at 2.53 ppm, which consists of eight peaks due to the conformational

inequivalence of the two hydrogens on it. Protons on methyl group (CH3) have a
simple doublet at 1.27 ppm. In case of PHBV methyl protons shows a triplet at around
0.89 ppm. Due to various stereo isomers of PHBV or PHB, different conformations of

branched groups on the main chain are possible. Thus the different tacticity of the

stereomers may lead to further splitting of multiplets. Further discussions regarding the

l 13
use of H and C NMR to study stereoregulation of PHB and PHV can be read in

26

reference [22]. Using the 1H NMR spectra of PHBV the compositions of copolyester

PHBV samples can be determined from the areas of the methyl resonances in the side

groups of HB and HV repeating units. In another option, the peak area of the proton

from methylene resonance due to the HV (CH3CH2) ethyl side group can also be

utilized to determine copolymer compositions. However, presence of water in the

sample leads to faulty inference of high copolymer due to impurity peak at -1.67 ppm,

which overlaps the methylene resonance of HV in CDC13 therefore this is not a

recommended analysis [23].

PHB or PHV molecule forms a left handed 21 helix with two molecular chains

passing through the unit cell. The unit cells of the lattice system for both PHB and PHV
are orthorhombic (or = B = y = 900) with P 212121 (D2 4) index and the axis are a =

0.576 nm, b = 1.320 nm, and c = 0.596 nm (fiber repeat) for PHB and a = 0.952 nm, b
= 1.008 nm, and c = 0.556 nm (fiber repeat) for PHV). PHBV or P(HB-co-HV) is a
copolymer of HB and HV monomer units and is an isodimorphous system (i.e. it can
crystallizes with two different structures depending on the HV content in the
copolymer). The PHB or P(3HB) crystal lattice (fiber repeat 0.596 nm) is observed
when the compositions of HV ranges from 0 to 37 mol % while the PHV or P(3HV)
crystal lattice (fiber repeat 0.556) is observed when the compositions of HV ranges
from 53 to 95 mol %. Lattice transformation takes place at about 40 mol % of HV
content and it is assumed that both crystal structures coexist in an intermediate

pseudoeutectic composition ranging from 36 to 56 mol% HV. The d-spacing of PHBV

27

or P(3HB-co-3HV) are listed in Table 1.1, borrowed from the reference [24]. For the
copolyester PHBV with HV content ranging from 0 to 37 mol % which has the P(3HB)
lattice pattern. Among the calculated (020), (110), and (002) d spacing, the (001) d
spacing was found to be increasing and (020) and (002) d remained unchanged with the
increasing HV mol %. The author suggests that the increasing (110) d spacing which is

associated with the “a” parameter of the unit cell was due to the ethyl group (which is

Table 1.]: Listing 26 values corresponding to d—spacing of PHBV lattice
Planes 020 1 10 1 1 1 031 040 002

d'si’zgmg 6.58 5.254 4.048 3.493 3.300 2.988
20(0) 13.75 17 22 25.3 27 30

 

 

 

 

 

 

 

 

 

 

 

bigger in size than methyl) of HV units which expand the (110) plane of the P(3HB)
lattice because of the steric hindrances. Whereas for the compositions of PHBV with
HV content ranging from 53 to 95 mol % which have P(3HV) lattice pattern, there was
no change in all the observed {(020), (110), and (211)} d spacings, thus leading to the
conclusion that the P(3HV) unit cell parameters were not influenced by the methyl side
chains of 3HB units (as the methyl units as smaller than ethyl units thus fits well in the

units of PHV lattice) [24, 25]. The degree of crystallinity can be obtained from the X-

ray diffracted intensity data in the range of 2 O = 10-38.50. The peaks at 28 =17O

corresponds to (110) diffraction patterns and at 13.70 corresponds to (020) diffractions

[26]. The approximate 29 values corresponding to the d spacing of the lattice plane are

listed in the table 1.1.

28

PHB produces bacterially has a low nucleation density and thus have a
spherulites of very large dimensions which may grow up to several millimeters in
diameter at high crystallization temperatures during the crystallization from the melt
[27].

As suggested by the author in reference[28] there are some intermolecular

interactions between the C: O and CH3 groups in the lattice, and these interactions

decrease along the 0 axis of the crystal lattice with temperature. (Ozaki) The author in
the same reference discusses the conformational rearrangements of PHB during the
melt crystallization using 2D-IR- correlation spectroscopy and infers that there is a
cooperative structural change of various functionalities within a polymer chains. During
the melt-crystallization of PHB, a different rate of changes was observed for the C-O-C
and C-C backbones. The author further suggests that a sequential change of chains
occurs during their structural adjustment and these sequential changes occur in the
order stated. During the crystallization process, Firstly, the ester groups are the

localized and flexible moieties and instantly adjust their conformation. Secondly, long

and regular 21 helix chain sequences starts developing through annealing and in the

final stage these regular 2] helix sequences packs themselves into the crystal unit cell,

which takes the longest annealing time [28].
Mechanical properties

Generally, the properties of PHAs corresponds to the type, composition, and
distribution of co-monomer units in the polymer chains, and also to the weight average

molecular weight (MW) and molecular weight distribution (poly dispersity index) of the

29

polymer. Their mechanical characteristics can be tailored to resemble elastic rubber or
hard crystalline plastic. The minimum Mw required for these polymer to exhibit
appreciable mechanical properties is 500,000 Da [12]. Often its properties have been
compared to polypropylene thermoplastic as there elastic modulus and strength are of
the similar order. PHB and PHBV produced in bulk quantities have elastic modulus of
1.7 and 1.2 GPa, with the strength (or yield stress) of 35 and 25 MPa. respectively.
These are brittle materials with low impact strength of the order of 3 kJ/mm and low
elongating (> 7-8%)[29]. In fact PHBV copolymer was produced in order to improve
upon the brittleness of PHB by inducing the valerate (HV) units in to the PHB
molecular chain. It was assumed that the ethyl group may interfere with the ordered
unit cell of PHB and disrupt the crystallinity, which was responsible for its brittleness.
Because of the bio origin, PHB has a high order of stereo-regularity of the isotactic
chain thus leading to high crystallinity and its subsequent brittleness. The primary
reason associated with its brittleness is the secondary crystallization. During the aging
or the storage time at the room temperature post crystallization phenomena of the
amorphous phase occurs. which further constrains it within the crystallites. Due to high
purity of bacterial synthesis, PHB has low nucleating density leading to large spherulite
size which may run up to several millimeters in size and exhibit interspherulitic cracks.
Both these factor leads to rapid decrease of elongation at break or embrittlement,
increase in its elastic modulus, strength and density. The elastic modulus increases from
~1.5 GPa of the up to ~35 GPa, whereas the impact strength decreases accompanying

the reduction in elongation to > 7-10 % [12, 30, 31]. Brittle failure of PHB due to

30

presence of radial and circumferential cracks in spherulites which grows on straining
has also been reported in literature [32].

PHAs have poor thermal stability and thermal degradation begins ~200°C. The

melt temperatures is at ~ 1700C and even at this temperature with prolong processing

under conventional techniques molecular weight reduction has been observed. Thus,
this gives a very narrow processing window of these polymers [12, 30, 31].

The crystals of PHB bears no center of symmetry and therefore change in the
direction of average dipole moment occurs if the deformation of crystals are carried in a
particular way i.e., the material is a piezoelectric.

PHB has a lower solvent resistance but is more resistance to ultraviolet weathering than

polypropylene.

Being a crystalline material its film has an excellent gas barrier properties, five

times more effective to C02 as a barrier than PET and as strong as polypropylene but

not as tough as PET and PP [33].

In an attempt to improve upon mechanical properties of PHB and PHBV
copolymer nucleation, plasticization and lubrication have been carried through in
processing [29].

PHAs biodegrade in microbial active environments. Since, PHAs function as an
intracellular energy and carbon source and use them as reserve materials, bacteria can
degrade PHAs. Microorganisms attack PHBV by secreting enzymes (depolymerases)
that break down the polymer into its basic hydroxybutyrate (HB) and hydroxyvalerate
(HV) constituents. The HB and HV fragments are then consumed by the cells to sustain

growth. Under aerobic conditions, the final biodegradation products are water and

carbon dioxide; under anaerobic conditions, methane is produced as well. The
degradation of PHBV can be quite rapid in biologically active systems. A range of soil
micro-organisms, both bacterial and fungal, can utilize PHAs as a source of carbon and
energy [7, 9].

1.1.3 Methods of modification and processing

1.1.3.1 Copolymerization

A special approach to design the copolymer molecular chains was investigated
by McChalicher et al., in which the differences in mechanical properties between the
random copolymers and block copolymers of the same copolymer were investigated as
the properties change over time. It was shown that the block copolymers maintain more
elasticity over a period of time than the similar random copolymers of analogous
composition [34].

To improve upon the material properties or act as a useful functionalized
polymer, a copolymer containing 3-hydroxy butyrate (3HB) and 3-hydroxy-4-pentenoic
acid (3HPE) was reported. This was done to obtain a copolymer with low Tg, low
degree of crystallinity, and a functionality that could be cross-linked during processing
[35].

With an objective to reduce the melting temperature of the PHB, and assuming
this would be an effective method to prevent its thermal degradation, and allow it for
lower processing temperatures in the hot-drawing process. Tsuge et al, synthesized a
copolymer P(3HB-co-3HA) with a small (3HA)fraction. where 3HA were the medium
chain length hydroxy alkanoates with number of carbon atom. ranging from 6 to 12

[36]. In an another similar study to improve the flexibility of the PHB by incorporating

32

medium-chain length monomers (PHAmcl, C6—C12) in the molecular chain, the
synthesis of a copolymer, poly(3-hydroxybuty—rate-co-3-hydroxyoctanoate) [P (HB-
HO)] was reported [3 7].

To obtain a set of useful properties not achieved by more traditional PHA
polymers, like PHB homopolymer or PHBV copolymers another simplest form of mcl
3-hydroxyhexanoate (3HHx) units were introduced to obtain a copolymer P(3-HB-co-
3HHx), which is branded under the name of “Nodax”[38]. The industrial scale
production, and economical production of this copolymer from palm oil are also
listed[39, 40] .

Glycogen accmnulating organisms (GAOs) were used, in a study by Dia e tal, to
produce PHA copolymers from a simple and cheap carbon source. The production of
poly(3-hydroxybutyrate (3HB)-co-3-hydroxyvalerate (3HV)-co-3-hydroxy-2-
methylvalerate (3HMV)), with 7—35 C-mol% of 3HV fractions from acetate as the only
carbon source with the use of GAOs was described [41, 42].

To further widen the scope of PHAs, a unique biosynthesis of the copolymer of
3-hydroxybutyrate and 3-mercaptopropionate, poly(3HB-co-3MP), with a molar
fraction of 3MP of up to 43% was described [43]. In addition, PHAs with the side chain
substitution of phenyl, cyno and halogens were also reported [44].

PHA was attempted to copolymerize with Polyethylene glycol (PEG), thus
producing PHA-PEG terminated copolymers[45, 46]. The introduced PEG group may
react with a range of molecules, thus, setting up the opportunity of new conjugates
between PHAs and other natural or synthetic polymers. Poly(ester—urethane) (PU)

multiblock copolymers from hydroxyl-terminated poly(ethylene glycol) (PEG) and

33

hydroxylated poly[(R)-3-hydroxyalkanoate] (PHA-diol) were synthesized for blood
contact applications, using 1,6-hexamethylene diisocyanate as a coupling reagent [47].
For medical and pharmaceutical applications, such as drug delivery systems, an
investigation to observe the influence of P(3HB-co-4HB) composition ratio and drug
loading level on the biocompatibility of P(3HB-co-4HB) was reported [48].

In an another attempt to reduce the melting point and crystallization of PHB the
biosynthesis and characterization of HB-rich co- polymers containing 2-4mol% of
hydroxycaproate (HC) units, as well as a terpolymer containing HC and
hydroxyoctanoate (HO) units was reported [49].

To exploit the advantages of natural and synthetic polymers, methyl metha
acrylic acid (MMA) was grafted on PHA—soybean to obtain new composite materials
[50].

Investigation concerning the optimization of the cost effective production of
PHA copolymers had been carried out by many researchers by altering the various
parameters like, carbon feed source, bacteria, genetic modification, enzymes [38, 40,
51-58].

Other than copolymerizing different PHA monomers among themselves, a
copolymer of PHB and PLA monomeric units was carried out by Dahlia et al.. This was
achieved via ring-opening polymerization of L-lactide using PHA as a macroinitiator
with stannous octoate as catalyst, thus introducing PHA units, up to 20 wt %, into PLA
[59].

1.1.3.2 Plasticization

Plasticization of the PHB was another attempt to overcome its brittleness for

commercialization purpose. Plasticizers interfere with the intermolecular interactions

34

within the polymer molecular chains and therefore, ease their conformational degree of
freedom. This alters the glass transition and flexibility/deformability of the polymer.
Change in the degree or rate of crystallization can also be observed due to the
plasticization effect in a polymer. Preferably, plasticizer used for PHB must be
“biodegradable” so that it retains the integrity of biodegradability. Numerous
plasticizers irrespective of their biodegradability have been used to improve upon its
brittleness. Citric acid based esters like, triethyl citrate, acetyl triethyl citrate,
acetyltributyl citrate, butyryltrihexyl citrate were reported to be used in PHBV as a
biocompatible plasticizers[60]. Another type of plasticizers for biodegradable polymers
mentioned are glycerine based esters. Kunze et a1. give a description of salicylic esters,
acetyl salicylic esters, aryl propionic esters and phthalic acid esters as the plasticizers
in PHB films [60]. Among these, salicylic acid decyl ester, acetyl salicylic acid hexyl
ester and ketoprofen ethyl ester were stated to give increased elongation at break of
PHB [60].

Biber eta al., used a series of plasticizer, i. e., dioctyl sebacate, dibutyl sebacate,
polyethylene glycol, Lapr01503, and Lapr015003, These plasticizers were reported to
be compatible and forming monophase with the PHB up to 15 — 20 wt% in
composition. The relative breaking elongation of PHB with Laprol 503, at room
temperature was mentioned to increase up to 250-300%[61].

Acylglycerols were used as plasticizers by Iskawawa et al., and among these
glycerol triacetate (GTA) was reported to be the best for PHB. At 30 wt% the

elongation at break increased from 5% to 130% (as mentioned) [62]. An increase in the

35

elongation at break of pure PHB by addition poly (ethylene glycol) (PEG) in varying
proportions was reported as well [63].

Soya bean oil (SO) and epoxydized Soya bean oil (ESO) were used as the
biodegradable plasticizers in PHB and PHBV [64, 65]. Epoxydized Soya bean oil was
stated to be more effective as a plasticizer and in increasing the elongation at break and
impact strength of the films. The amount of these plasticizer was mentioned to be
varied from 5 to 30 % in the blend composition [64, 65]. In a comparative study,
plasticization effect of dibutyl phthalate (DBP) and triethyl citrate (TEC) was weighed
against the SO and E80 by the same author. It was affirrned that TEC or DBP were
better plasticizers than SO and ESO for PHBV [64, 65].

PHB was plasticized with another set of plasticizers dioctyl (o-)phthalate,
dioctyl sebacate, and acetyl tributyl citrate (ATBC) by Wang et al.[66]. Among these.
ATBC was stated to be a competent one.
1.1.3.3 Blending

Blends can be defined as physical mixtures of structurally different polymers
which may exist as a single phase (miscible blends) or distinct phases (immiscible
blends). The physical properties of blends are strongly dependent on this miscibility
and hence most work on PHB based blends is concentrated in achieving miscibility.
Thermodynamically miscible blends of PHB have been reported with poly (vinyl
alcohol), [67] poly (ethylene oxide), [68], [69] poly (lactide), [70, 71] poly (butylene
succinate-eo-butylene adipate), [72] poly (e-caprolactone-co-lactide), [73] poly
(butylene succinate-co—e-caprolactone), [72] cellulose acetate butyrate, [74] poly

(ethylene-co-vinyl acetate, [75] poly (vinylidene chloride-co-acrylonitrile) [76].

36

Compatible blends of PHB have been obtained with poly (e-caprolactone), [77, 78] and
polysaccharides [79]. Miscible blends of PHB have been formed with non-
biodegradable polymers such as poly (epichlorohydrin), [80] and partly miscible blends
are formed with poly (methylmethacrylate), [81] ethylene-propylene rubber, [82] and
poly (butylacrylate) [83].

The miscibility of the PHB with poly(caprolactone-eo-lactide) (P(CL-co-LA)) was
studied by Naoyuki Koyama et a1 [71]. The varying ratios of caprolactone to lactic acid
in P(CL-co-LA) describes the miscibility as a function of the ratio of two components

in P(CL-co-LA). He further provides the details of the persisting relationship among

the number average molecular weight (Mn) and the lactic acid fraction in P(CL-co-LA)

on the miscibility of the, PHB/(P(CL-co—LA), blend and discussed it in context of Flory
Huggin’s theory. The transition of immiscibility to the miscibility of the blends was
favorable at low Mn and over 21% LA content in (P(CL-co-LA) copolymer, therefore,
depicting the decrease in Flory Huggin parameter for the P[(R)-3HB]/P(CL-co-LA)
blend with the increasing LA fraction in the P(CL-co-LA) copolymer component.
These blends were, additionally, investigated for their spherullite morphology, further,
suggesting that the uncrystallizable P(CL-co-LA) component exists within the
spherulites of P[(R)-3HB]. The spherulites of P[(R)-3HB] exhibited banding in the
miscible blends, and in case of immiscible blends, phase separation of the P[(R)-3HB]
and P(CL-co-18 mol % LA) components in the interfibrillar regions within the P[(R)-
3HB] spherulites. In respect of the biodegradability of the blends, the author concludes

the decrease in the rates of enzymatic hydrolysis with the increasing P(CL-co-LA)

37

content, and the biodegradability was independent of the miscibility of the blending
components.

Wang et al.[74] discusses the PHB and cellulose acetate butyrate (CAB) blends
in different blending ratios. The author states the miscibility of both the components
over the entire scale of mixing ratios in a melt state. The decrease in the melting
temperature and the crystallinity of the PHB with the increasing CAB percentage in the
blends was reported, and the CAB content, more than 50%, lead to an improvement in
the PHB toughness with the development of homogeneous amorphous phase. CAB
content in the blends was a controlling parameter of the degradation rate of the blends.

In a study by Yoon et a1. [75], PHB and poly(ethylene-co-vinyl acetate) (EVA)
blends, with the 70% and 85% vinyl acetate(VA) content in EVA, confirms the
miscibility of both at a higher VA fraction. The blends were solution casted. The author

establishes the miscibility of PHB-EVA (at 85% VA) on the basis of Flory interaction

parameter (x12), which come out to be negative from the depression in melting

temperature with increasing EVA fraction in the blends. The growth rate of the PHB
spherulite was also shown to be reducing with the increasing EVA content at 85%.
Gonzalez et al.[76] reports the solution precipitated blends of PHB and
poly(vinylidene chloride-co-acrylonitrile) copolymer (VDC-eo—AN) with the improved
C02 permeability at 50 wt% of poly(VDC—co-AN). Blends, containing 32 mole % of
AN units in poly(VDC-co-AN) copolymer, were miscible over the complete set of
blended compositions. The appearance of distinct glass transition temperature, and the
depression in melting point of PHB, reflected the miscibility of the blends, and the

interaction parameter, obtained from the melting point depression analysis, was

38

reported to be -0.08. The miscibility of PHB and poly (VDC-co-AN) was contemplated
due to the conditioned dipole—dipole interaction between the carbonyl group of PHB
and chloride atoms of VDC-eo-AN copolymer.

Rosa et al.[84] studied the effects of oxidized polyethylene wax (OPW)
on the tensile properties of PHB/LDPE blends. The blends were stated to be immiscible
with or without the presence of OPW and no significant deviation in any of the
properties were demonstrated.

Binary blends of PHB with poly(-caprolactone) (PCL), poly(l,4-butylene
adipate) (PBA) and poly(vinyl acetate) (PVAc) were studied by Kumagai and Doi in
respect of their miscibility, morphology and biodegradability [77] PHB/PCL and
PHB/PBA blends were stated to be immiscible, and phase separation occurring at
macro and micro scale, respectively, whereas, the PHV/PVAc blend was miscible.
Enzymatic degradation of PHB was affected by the weight fraction of second
component in the blends. In case of PHB/PBA or PHB/PVAc blends the rate of
degradation plummeted with the increased weight fraction of PBA and PVAc. Though,
PHB and PCL blends were immiscible, Gassner and Owen[85] mentions that these
blends formed mechanically compatible films containing phase-separated partially

crystalline domains. In order to be temperature-resistant, the blended composition, due

to the low melting point of PCL (60°C), was restricted to have the PHB fraction no lees
than 0.6 in the entire composition.
Paglia et al [80] details that atactic poly(epichlorohydrin) (PECH), when above

the equilibrium melting temperature of the two components, has tendency to form a

thermodynamically stable miscible blends with PHB, and this was inferred from the

39

negative value of the 702 The miscibility of blends was also confirmed from the

appearance of distinct single glass transition temperature depending on the varying
composition of the components and their agreement with the Fox equation.

Tertiary blends of PHB with poly(ethylene oxide) (PEO) and poly(methyl
methacrylate) (PMMA) were attempted by Yoon et al [81] subsequently, after
establishing the miscibility of PEO individually, with PHB and PMMA, but PEO was
ineffective to stimulate the compatibility among the two. Miscibility of PHB with PEO

and immiscibility of PHB with PMMA was reasserted with their respective negative

and positive values of XI 2. The substantial effect of compositions on the X12 value was

emphasized.

Immiscibility of PHB and (ethylene-propylene rubber (EPR) and its
compatibility with poly(vinyl acetate) (PVAc)) was reported by Greco and Martuscelli
[82]. Miscibility of PHB and PVAc in perspective of distinct single glass transition
temperature of blends and depression of equilibrium melting temperature of PHB was
highlighted. Phase segregation of EPR and occlusion in intraspherulitic regions of PHB
spherulites during the growth from the melt was elucidated using optical microscopy.

A unique study by Avella et al.[83], with an objective to improve upon the
impact strength of PHB, reactively blended polymeric PHB with butyl acrylate
monomer to form inclusive poly(butyl acrylate) and PHB blends. The paper discusses
the enhanced impact strength with the poly(butyl acrylate) inclusions and details the
preparation of the reactive blending of the two. It also discusses the morphology of the

prepared blends in different proportion of the two components.

40

1.2 Wood Plastic Composites
1.2.1 Conventional Wood Plastic Composites

In the past few years wood and natural fiber plastic composite (W/NF PC) are
well-known in the consumer market of plastic composites, and there dynamic
emergence have made them the fastest growing segment in the plastic industry with
their ongoing growth of 25% and projecting a strong market demand in the near future
[86]. The current market of wood plastic composites (WPC) is 1.2 billion pounds and is
expected to grow with an average annual growth rate of 9.8% thus providing a market
of 1.9 billion pounds by the year 2009 [87]. Not only, do they have a proven market in
outdoor and indoor applications like decking, railing, fencing, furniture, etc. they also
hold a large share of the European market in automotive applications[88]. The wide
acceptance of W/NF PC by the global market has been due to their reduced cost,
weight, superior properties over wood and plastic alone, recycling ability, adaptation to
the existing plastic processing techniques, and above all the motivation of industry to
ahead for the coo-friendly products. Most of the thermoplastics being used in W/N F PC
are recyclable, and are being looked as a profitable opportunity for these materials.
Contemporary W/NF PC are based on conventional thermoplastics and naturally
available cellulose fiber, the fiber content varying from 30 to 70 wt% in the
composites. Polypropylene has the major market share of W/NF PC followed by
polyethylene. polyvinylchloride and other resins. Various kinds of lignocellulosic
natural and wood fiber have been used to achieve improved mechanical properties of

W/NF PC. These lignocellulosic filler includes various wood fibers/flour (e.g maple,

41

pine, birch etc.), flax, jute, bagasse, comcobs, cereal straw. kenaf, jute, sisal, coir, flax,
banana, rice hulls, newsprint, pulp, and even waste wood.

The enhancement in the mechanical properties of the W/NF PC are governed by
properties of the matrix and the reinforcement, phase state of matrix after
reinforcement, interface developed due to the interaction between the matrix and
reinforcement. Amount, shape (aspect ratio). distrihution/dispcrsion and alignment/
orientation of the fibers within the composite matrix, and fabrication process also plays
a crucial role in goveming the properties of the W/N F PC [89].

While, all afore listed factors affect the composite characteristics, among all, the
interphase between the matrix and fiber typically, plays an important role. It determines
the extent of the stress developed in the matrix that is transferred to the reinforcement.
A good interface gives an efficient stress transfer between the two due to the better
adhesion, and therefore, gives the best of mechanical properties (Strength and
toughness).

The possible adhesion mechanisms occurring between the matrix and the filler
reinforcement are adsorption, chemisorptions, electrostatic interactions, mechanical
interlocking. Adsorption happens due to the intimate intermolecular interactions
between the molecules of substrate (fiber) and resin (thermoplastic), which results from
Van derWaals forces or Hydrogen bonding. Chemical bonds created between the
substance and the surface of substrate as a result of chemical reactions is called
chemisorption. Electrostatic mechanism of adhesion is attributed to the creation of
opposite charges on the interacting surfaces of fibers and thermoplastic thus, an

interface consisting of two layers of opposite charges is formed, which accounts for the

42

adhesion of the two. The mechanical interlocking phenomena accounts for the adhesion
when a matrix penetrates into the pores, holes and crevices or other irregularities of the
substrate (fiber), and locks mechanically to it. A good wetting of the resin
(thermoplastic) on the substrate (fiber) is essential for all the mentioned adhesion
phenomena’s to exist in an appreciable amount.

The lignocellulosic fibers are hydrophilic, and apolar thermoplastics (PP, PE)
are hydrophobic therefore, there is a poor wetability and lack of adhesion between the
two. This results in an incompatible interface between the lignocellulosic (wood
/natural) fiber and thermoplastics that causes the filler debonding from the matrix and
the poor dispersion of fiber in the matrix. The stress transferred from the matrix to the
fiber due to poor interface is least efficient and subsequently, results in the poor
mechanical properties of the composite [90, 91]. This situation requires developing the
strategies for the surface modification of the lignocellulosic fiber surfaces that could
render an effective control over the fiber/polymer interface. One approach to this
situation is using a compatibilizer or a coupling agent that bridges the two different
surfaces of thermoplastic and fiber.

Compatiblizer has two segments in a block copolymer as depicted in figure 1.9.
One segment has a similar chemical nature to the thermoplastic matrix and the other is
able to adhere to the wood fibers by any of the adhesion mechanism. The segment
having similar chemical characteristics to the thermoplastic matrix is miscible in it.
This generates an interface, of appropriate strength, between the fiber and matrix.

The wood adhering segment of the block copolymer generally has functional

groups able to form covalent bonds with hydroxyl groups on wood fiber surface. A

43

covalent bond forms a stronger interface than a hydrogen bond and also reduces the

Polymer Chains Polymer compatible
Segment of compatiblizer

 

Fiber binding segment of
compatiblizer / coupling
agent

Figure 1.9: Schematic of coupling in interface region

hydroxyl groups ( hydroxyl groups accounts for the water absorbing tendency of
wood) on the wood fiber surface by reacting thus decreasing the hydrophilicity of wood
fiber and making the composite better water-resistant than wood.

Maleated polypropylene has been the most common compatiblizer used as a
coupling agent in polypropylene based WPC’s that performs with the same strategy as
mentioned and increases the strength of composites to a reasonable extent [92]. Surface
modification of fibers using enzymes is also being looked into by various researchers
and other chemical approaches to the compatiblization are carried out using silanes,
isocynates, zirconates, organotitanate, grafting some polymeric chain on fiber surface
(e.g MMA), impregnating fibers with the polymer compatible with the polymer matrix
and alkali treatment of fibers to change morphology and enhance mechanical
interlocking. Some fiber surface modifications using physical sources like plasma

treatment, sputtering and corona discharge have also been mentioned [93] .

44

Strength of fiber-matrix interface in composites is often related with the
mechanical dampening (Tan 6) in the dynamical mechanical analysis (DMA). Strong

fiber-matrix adhesion lowers the mechanical dampening (Tan 8) value due to the

reduced friction at fiber-matrix interface. Mechanical damping (Tan 50) within the

composite system is given as Tan 5c = 8p Tan 5p + 91 Tan 5i + 9f Tan 5f

where O with subscript f, p, i are the respective volume fractions of the fiber,
matrix and interface. Tan 8 with subscript f, p, i are the mechanical dampening of the
respective, fiber, matrix and interface. The adhesion factor (A) of the interface strength

can be related to Tan 8 as following.
A= {(Tan BC/ Tan 5p)/(1- Of)}-1

The molecular mobility surrounding the fiber is reduced; when the fiber-matrix
interface adhesion is increased. Low values of the adhesion factor (A) suggest
improved interactions at the fiber-matrix interface.

It has been assumed that adhesion factor is reasonable only if the development
of a transcrystallinity layer surrounding the fiber at the interface is ignored [94].

A significant amount of research. focusing to improve the compatibility
between the hydrophilic natural fiber and the hydrophobic thermoplastics, is referred
[95-98]. Nam-reinforcement using inorganic material like clay in wood-PP composites
has also been reported elsewhere [99]. The only major limitation these composites have
is their limited dimensional stability when exposed to moisture [100]. When comparing
their moisture absorbance with the wood and plastic alone, these have a better moisture

resistance than wood but lesser than plastics.

45

1.2.2 Wood as a natural reinforcing component

Wood has always been an essential part of human economy, since the beginning
of any civilization. It has been used in several ways as a structural material, fumishing,
tools etc., until recently, when it has found a new application as reinforcement in plastic
composite. Wood fibers are used most abundantly in thermoplastic based composite,
and in low severity applications for therrnoset composites. There use as has been
increasing with the growth of WPC market. This biomass industry is projected to
become a $1.5 billion fiber market by 2012 [101, 102]. These fibers invite attention of
the market due to their low cost, low density, easier to handle in production and
abundance, but their renewabilty and eco-friendliness also allow them to be used as a
sustainable material for future growth.

Wood fiber, also known as wood cells, forms the basic structural elements of
the wood. These fibrous cells are hollow, elongated, spindle shaped cells that are
arranged parallel to each other along the trunk of a tree. Their characteristics and their
arrangements essentially, affect the properties such as strength and shrinkage of the
wood. Woods are by and large classified in to hardwood (deciduous trees) and
softwood (coniferous trees), which is based on the cellular structure of the wood.
Hardwoods have four types of cells. Bulk of the wood is made from small wood fibers,
and large cells function’s to conduct sap up the tree through the mass of fibers , which
on the end grain appear as holes or pores,. These are also called tracheids, hardwoods
are also referred to as porous woods in contrast to nonporous softwoods. Other cells are
largely used to store food and are the strength-giving elements of hardwoods. They are

spindle shaped cells usually having small cavities and relatively thick walls. Softwoods

46

do not have vessels and have a simpler fibrous structure based on only two cell types.
The length of wood fibers is highly variable within a tree and among species.
Hardwood fibers average about 1 mm (1/25 in.) in length; softwood fibers range from 3
to 8 m (1/8 to 1/3 in.) in length [103, 104].

These fibrous cells have cell wall of the thickness. which may vary up to 100
um, and this makes the cell rigid. The supportive framework of the cell wall surrounds
the cell. This constitutes a layers cellulose microfibrils network that are embedded in a
matrix of lignin and hemicellulose. Dry wood is primarily composed of approximately
50% cellulose, which is the major component, 23% to 33% lignin in soft wood and
16% to 25% in hardwoods, hemicelluloses, and 5% to 10% of extraneous materials
[104].

The structure of cell wall is layered; a primary wall is the outer most layer with
a thickness of ~0.1um, which constitutes mainly, a non cellulosic substances like waxes
and pectin, but some cellulosic microfibrils that form an interwoven network; a
secondary wall is the middle layer, which is the main strengthening unit of the cell wall

that constitutes almost all of the cellulose in form of cellulose fibrils. To toughen the

structure, the fibrils are aligned at 10 to 300 to the tree trunk axis of the cell wall.

Lumen is the central canal, which chiefly holds the proteinaceous materials [104-106].
The secondary wall is actually an entity on three sub layers. First sub layer is
adjacent to primary layer, which act as a transitioning sub layer of thickness ~ 0.2 pm.

It consists of microfibrils, forming a lamina, crisscrossing each other and making an

angle of 40 to 500 to a fiber axis. Second sub layer is the central one with a thickness of

~ 1 to 5 pm, which has micriobrils arranged parallel or making small angle with the

47

fiber axis. The third and the innermost sub layer adjacent to the lumen is also called
tertiary wall. It has no fibrillar structure at all, but only a series of granules.

Cellulose, which forms a 50 to 55 % by wt. of the wood constituents, is a
polysaccharide. The monomeric unit of this polymer is an unbranched poly-B-(I,4)-D-
glucopyranose. In the cell walls these polymeric chains arrange to form ordered
crystalline structures and cellulose is up to 90% crystalline. These crystalline regions
form microfibrils, but less ordered or amorphous domains also exist. This is Important
to mention that in these crystalline domains, the hydroxyl groups of the glucose units
engage in ordered hydrogen bond systems, both within and between the cellulose
chains. Intramolecular hydrogen bonding holds the accountability of the individual
chain conformation, while intermolecular hydrogen bonding is responsible for
microfibrils formation. The metastable cellulose-l is the most common allomorphic
form of crystalline cellulose. It is a parallel arrangement of polymeric chains in a form
of flatribbon conformation with a strong intramolecular hydrogen bonding. The two
sub-allomorphs, I-a and 1-0 coexist in a cell wall and their ratio of prevalence varies in
different types of wood [105-107].

The networks of cellulose microfibrils in a cell wall are embedded in a matrix of
lignin and hemicellulose. Although, some fraction of lignin is distributed in the cell
wall, but major portion of it is concentrated around and between the fibrous cell, and
act as a binder or matrix for fibers. It is quit unclear that whether lignin is chemically or
physically bonded to the cellulose, but it is tightly bonded to one of the constituent of

the wood. Lignin is a three-dimensional phenylpropanol polymer and softens at the

temperature of about 1700C / 3400F. Wood lignins are predominantly an aromatic

48

mixture, and not fully soluble in the known solvent. Maple wood used in this research
has 23 % lignin in it [104, 108].

Hemicellulose is another prime constituent of the wood, which is associated
with the cellulose and lignin in a cell wall. It is a mixture of polysaccharides from
different kinds of sugar monomer. In hardwoods glucose and xylose sugar monomers
are more prevalent.

Other than three above mentioned major constituents wood has extraneous
compounds and mixtures like, waxes, essential oils, fats, resins, coloring matter, gums,
starch, and simple metabolic intermediates. These may vary from 5 to 30 % depending
upon the species and growth of a tree.

The density of the dry cell wall is in the vicinity of 1.53 gm/cc, and may vary
from specie to specie and other factors like growth conditions etc. The Young’s
modulus of wood cell wall has measured values ranging from 10 to 60 GPa, which
again varies fiom species to species.

1.2.3 Talc as a reinforcing inorganic filler/component

Tale is a significant industrial mineral, and has an application in material for
paints, rubber, plastics, insecticides, etc. It is a soft, platy, water repellent and
chemically inert mineral, and practically no two talc from two different sources are
same. It is available in various colors depending on the source ranging from white,

brown and green to blue. Tale is a 2:1 layer phyllosilicate mineral, and chemically, an

hydrated magnesium sheet silicate with the unit structure Mg3Si4O|0(OH)2 and

molecular weight, 379.27 gm. Crystal structure of tale is monoclinic or triclinic with

the space group C-1 and unit cell dimension of a=5.2900 A, b=9.1730 A, c=9.4600 A

49

a=90.460, B=98.680, y=90.090

=2. The elementary sheets are composed of magnesium ion in an oxygen and
hydroxyl octahedra that is sandwiched between the two layers of silicon and oxygen
tetrahedral. The elementary sheet thickness is ~ 6.6 A and the inter-sheet distance is
~2.8 A. The hydroxyl groups are not located on the basal faces of the sheets. rather
these are located within the sheets. i.e. on the lateral faces. which accounts for the talc's
hydrophobicity and inertness. The elementary sheets are stacked, and the binding
forces among these elementary sheets are weak Vander Wall forces. A few thousand
elementary sheets constitute to form a platelet and. these platelets slides apart at the
slightest touch, which imparts the characteristic of softness to the tale. The dimensions
of the individual platelet can range from 10 to 100 um depending on the deposit source.
The individual platelet dimensions determine the talc’s platyness or lamellarity. Large
individual platelet size will give a lamellar morphology and a much smaller platelet
dimensions would result in microcrystalline talc. Generally, it is classified on the basis
of mineralogy, morphology, and geographic source[109].

It is relatively chemically inert and thermally stable. Its melting point is at

1500°C, and re-crystallises into different forms of enstatite (i.e., anhydrous magnesium

silicate) above 10500C, after loosing the hydroxyl groups at 9000C.

Tale is used as an effective reinforcing agent in plastics because of its lamellar
structure. which can disperse uniformly in the matrix. although it requires an advance
and efficient milling technology to produce finest talc without compromising the aspect
ratio of the lamellar structure. Talc also act as a nucleating agent in the semi-crystalline

polymers. It is very commonly used in polypropylene to improve upon stiffness and

50

dimensional stability for food packaging, automotive and house hold appliance. In
LDPE it is used as an anti blocking agent. Other uses of talc are predominant in rubber,
paper, personal care, ceramic and coating industry. In rubber it is used as a processing
aid and a viscosity modifier, and in decorative coatings as extenders to improve upon

the hiding power and. titanium dioxide efficiency.

51

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61

Chapter 2

Wood Fiber Reinforced Bacterial Bioplastic Composites:

Fabrication and Performance Evaluation

2.0 Introduction

Wood reinforced plastics (WPC) are the fastest growing sector of the plastic
industry, ratcheting up with the quarter percent growth rate and reflecting a strong
market demand in the near future [1, 2]. The wide, global acceptance of these has been
due to the reduced cost, weight, superior properties over wood and plastic alone,
recycling ability, assimilation to the existing plastic processing techniques and a
mounting awareness regarding the sustainable and coo-friendly products. WPC, as a
structural material, predominantly exist in decking, railing, automotive and housing
interiors. Several kinds of WPC, using wood fiber and conventional polymers like PE,
PVC and PP, had been developed and researched [3-6]. Emerging issues with the
conventional petroleum based polymers such as, non-renewability, environmental
sustainability are the major challenges that needs to be addressed in the near future.
Renewable, bio-based, coo-friendly and fully green WPC are the potential alternative to
the conventional WPC.

Polyhydroxyalkanoates (PHAs) can be a new found renewable resource based
substitute to conventional polymers as a matrix in WPC. Polyhydroxybutyrate (PHB)
and polyhydroxybutyrate-co-valerate (PHBV) are prominent examples of PHAs
bioplastics. The potential benefit of these polymers as matrix in WPC is because of
their comparative properties to conventional polymers such as polypropylene coupled

with their eco- friendly properties like biodegradability, biocompatibility and

62

environmental sustainability. Among PHAs, PHBV was developed to compensate for
poor mechanical properties and process ability of PHB by introducing hydroxy valerate
(HV) units in PHB segments thereby, producing a PHBV copolymer. [7-9]. PHBV
produced through fermentation gives a good command on stereoregularity of the
polymer therefore, giving a highly isotactic and crystalline thermoplastic [10-11].
Increase of HV content in PHBV decreases its brittleness by limiting secondary
crystallization to some extent. Low HV content restricts segmental dynamics in the
amorphous region of PHBV thereby, showing reduced flexibility, elongation and
impact strength [12].

Natural fibers are the emerging reinforcing agent in the polymer matrix due to
their abundant availability, low cost, low density and compatibility with the nature [13].
Wood fibers, which fall in category of natural fiber, have been used to reinforce
thermoplastics [3-6]. The mechanical properties of PHBV were found to improve when
reinforced with cellulosic natural fibers [14-16]. PHBV can be reinforced with high
content of cellulosic wood fiber in order to enhance its mechanical properties and to
reduce the cost of molded product. Although, there are a few hydroxyl and carboxylie
groups on PHBV molecular chain ends, but some degree of dipolar interactions among
moderately polar (>C=O) and polar groups (-OH and —COOH) of PHBV and hydroxyl
group (-OH) of cellulosic wood fiber can provide an opportunity for an interfacial
interaction between the fibers and the matrix. [17,18]

The objective of this study was to fabricate maple wood fiber reinforced PHBV
composites and evaluate their mechanical and thennomechanical properties as a

function of wood fiber weight content in PHBV. The composites were fabricated by

63

extrusion-injection molding process. Efforts have been made to develop a theoretical
understanding of the experimental findings by correlating them with theoretical models
presented by Halpi-Tsai and Tsai-Pagano. The morphological aspects of the composites
such as fiber dispersion and fiber-matrix interface were investigated. using scanning
electron microscopy (SEM).
2.1 Experimental Section
2.1.1 Materials:

Polyhydroxybutyrate-co-valerate (PHBV) with the trade name Biopol (Zeneca
Bio Products) was obtained from Biomer, Germany. Maple wood fiber, trade name
2010 MAPLE wood flour was supplied by American Wood Fibers, Schofield, WI.
Length and diameter of the wood fibers were between 1600-1650 um and 300-400 pm,
respectively. The aspect ratio (l/d) of the wood fiber was ~ 4 [19].

2.1.2 Fabrication of composite employing micro-compounding technique:

Maple wood fiber was dried at a 110 0C under the vacuum for 2 hrs. PHBV was

dried at 80 0C for 4 hrs under similar vacuum conditions. Both, PHBV and wood fiber

were dried in order to avoid any moisture content during processing. The dried wood
fiber and PHBV were kept in a zip lock plastic bags until processing. Processing was
carried out in a micro extruder (manufactured and designed by DSM Research,
Netherlands) with a barrel volume of 15 cc. It was equipped with a co-rotating twin-
screw system with screw length of 150 mm and L/D (length/diameter) of 18.

Temperature, in the three respective zones of micro extruder, was set at 165,160 and

155 0C, for all the different compositions of PHBV and wood fiber. Processing

64

temperature was optimized between 155 and 1650C, to achieve low melt flow, and to

avoid thermal degradation of PHBV. The screw rotating speed was 100 rpm. Different
mix ratios of PHBV and wood fiber along with some other processing parameters are
mentioned in Table 2.1. Wood fiber and PHBV pellets were physically mixed, and fed
to the DSM extruder. After the fixed extrusion time, the molten mix was transferred to
a preheated mini-injection molding machine and the testing specimens were molded.

The optimized injection pressure for all PHBV-wood fiber composites was 551.5 KPa

and the temperature was 155 OC. Residenee/ Extrusion time, in the DSM extruder, of

compositions with 30 and 40 wt% was increased to 10 minutes in order to provide

better mixing and dispersion of the wood fiber in the matrix.

Table 2.1: Processing parameters of various PHBV-wood fiber compositions

 

PHBV (wt %) Wood fiber (wt %) Processing temp. Retention time

 

(° C) (min)
100 0 160 3
90 10 160 5
80 20 160 5
70 30 160 10
60 40 160 10

 

65

2.1.2 Testing and Characterization

2.1.2.1 Mechanical properties: Tensile and flexural properties of the specimens were
measured using a United Calibration Corp. SFM 20 testing machine as per ASTM
D638 and ASTM D790 standards, respectively. The tensile testing was carried out at a
rate of 2.54 mm/min with a pre load value of 0.023 kg. The system controls and data
analysis were performed using Datum software. The notch Izod impact testing was
carried out on a Testing Machine Inc. (TMI) 43-02-01 as per ASTM D256.using
pendulum of impact energy 1.35 J. The number of specimens used for testing were five.
2. 1. 2. 2 Thermomechanical properties:

2.1.2.2.] Dynamic Mechanical Analysis (DMA): The storage modulus, loss modulus

and tan delta of the specimens were evaluated using DMAQ800 supplied by TA

Instruments. The specimens were heated from -30 to 125 0C at a heating rate of 2

oC/min. A single cantilever mode was used to test the specimens at the oscillating

amplitude of 15 um and frequency of 1 Hz.
2.1.2.2.2 Heat deflection temperature (HDT): DMA Q800, TA instruments was used to
find the HDT of various specimens as per ASTM D648 standards. The rectangular bars,

1.99 x12 X 58 mm in three-point bending mode with an applying load of 66 psi were

used for testing. Samples were heated at the rate of 2°C/ min from the room

. o .
temperature, i.e., 20 C to the required temperature.

2.1.2.2.3 Coefficient of linear thermal expansion (CL TE): CLTE of the composites was

determined using TMA 2940, TMA was equipped with an expansion probe. Heating

66

rate of specimens was 3 oC/min under a constant applied load of 0.15 N. Testing was

performed in the temperature range from 30 to 90 oC.

2.1.2.3 T hermogravimetric analysis (T GA): TGA was tested, using TGA 2950

. . . . o
Instrument, wrth the heatrng rate of 20 OC/mrn from the room temperature to 600 C,

i.e., the temperature of complete thermal degradation. All TGA runs were performed in

N2 atmosphere.

Data analysis of all the therrno mechanical properties was performed using Universal
analysis software and the number of specimens used for testing were five.

2.1.2.4 FT IR analysis: Perkin Elmer, system 2000 FTIR Spectrometer, was used to
carry out F TIR spectra, using an attenuated total reflectance (ATR) accessory of all the
specimens.

2.1.2.5 .Morphological studies: Morphological studies were conducted using a JEOL
scanning electron microscope (model J SM-6400). Specimens were sputter-coated with

gold to a thickness of ~10 nm before the surface characterization. The SEM instrument

has a lanthanum hexaboride (LaB6) crystal as electron emitter. An accelerating voltage

of 13 KV was used to collect SEM photomicrographs.

67

2.2 Results and Discussion

2.2.] Tensile Properties: Tensile properties of neat PHBV and its wood fiber loaded
composites are depicted in Figure 1.1. A gradual increase in the Young’s modulus of
PHBV-wood fiber composites was observed with the uniform increment of wood fiber
loading in the total composition of composite. Young’s modulus ratcheted up by ~ 28
% with every 10 wt% increase of wood fiber loading in the composite. The tensile
modulus of PHBV based composite was increased by 167 %, when reinforced with 40
wt% of wood fiber as compared to neat PHBV. The modulus of neat PHBV was found
to be 1.02 i 0.09 GPa, while for the PHBV-wood composite with 40 wt% for wood
fiber was 2.73 :t 0.22 GPa. The increasing trend of the modulus could be attributed to
the compatibility between the wood fiber and PHBV. Wood fiber has polar groups in
the cellulose structure, and PHBV also has abundant moderate polar groups in its
chemical makeup, this may provide some opportunity of having a reasonable
interaction at the interface of the two. A unifonn dispersion of the filler in the matrix is
also a crucial parameter in determining the final properties of a composite. Based on the
visual observation of the processed specimens, the uniform dispersion of the fibers in
matrix allowed the fibers to share an evenly distributed load. Figure 1, shows a slender
decreasing trend in the tensile strength, with the increase in wood fiber content. Tensile
strength of neat PHBV decreased to 16.75 i 0.51 MPa from 21.42 i 1.5 MPa, when
loaded with 40 wt% of the wood fiber. There was an average decrease of 6.3 % tensile
strength with each subsequent increment of 10 wt% wood fibers in the composite. The
decreasing trend could be explained by taking the dewetting effect in to account. The

theory pertaining to the composite behavior under external loading suggested that the

68

interface region of fiber and matrix experiences the stress concentration around the
filler particle as a result; the interaction between the filler and the matrix weakens up
thus, gradually leading to debonding at the interface [20]. The phenomenon becomes
more predominant with the increase of fiber to matrix volume ratio. Comparing the
tensile modulus and strength of the PP based wood fiber composites with PHBV—wood
fiber composites, PP based composites have a higher order of tensile properties. The
modulus of the 30 and 40 wt% wood fiber-PP composites, as reported in literature,
have values of 3.33 and 4.72 GPa, respectively, and the strength of the same wood
fiber content composites are 27.1 and 25.6 MPa , respectively [6]. The experimental
data of tensile properties were compared with the most commonly, applied theoretical
models. Halpin -Tsai and Tsai-Pagano equations and Nicolais - Nicodemo model were

used to evaluate the modulus and tensile strength of the composites [21-23, 14].

69

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Halpin-Tsai and T sai-Pagano equations:
Halpin —Tsai and Tsai-Pagano proposed equations for calculation of tensile

modulus of composites. Elastic modulus of the maple wood is 11.3 GPa [24]. The

density of the PHBV and wood fiber are 1.25 and 1.4 gm/cm3 respectively. Aspect ratio

(l/d) of the wood fiber loaded in the composite is 4.02. Halpin-Tsai equation predicts
the longitudinal (EL) and transverse (E r) modulus of the aligned short fiber composite.
Longitudinal modulus (EL) and Transverse modulus (Er) are given by following

equations.

1+§77LVf

 

 

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1+277TVf
T =

m l—nTVf

4’ is the measure of geometry of reinforced fiber and. if the fiber has rectangular

77,, and 77., are given by the following equations.

_ (Ef/Em)-—l
_(Ef/Em)+§

 

(Ef/Em)—I
”T: .
(Ef/Em)+2

 

71

The following Tsai-Pagano equation predicts the moduli of the composite for a

randomly oriented and uniformly distributed fiber by taking in to account 5/8th of the

transverse moduli and 3/8th of the longitudinal moduli of composites with fibers

oriented in a single specified direction.

Emndom = (3 / 8)E, + (5 / 8)E,

random , is the tensile modulus of the randomly oriented fibers. Figure 2.2 depicts

the experimental and theoretical values of the tensile modulus. The experimental tensile
modulus values were, significantly, close to the theoretical data points. This suggested
that the wood fibers were randomly oriented in the PHBV matrix. The embedding of
wood fiber has stiffened the PHBV under the given processing conditions, thereby
increasing the tensile modulus. This could be attributed to an appropriate assumption of
model in respect to the fabricated composites i.e. short fiber length and random
distribution of fibers in matrix.
Nicolais and Nicodemo model:

Nicolais and Nicodemo proposed a model that proposes an equation for tensile

stress as a function of filler fraction, assuming no adhesion between filler and matrix

b = 2 / 3. ¢ is the Volume fraction of filler while, a is a constant. (Figure 2.3) The

tensile strength data fit very well with the proposed model assuming the value of a =

0.45. “a” is assumed to be a function of aspect ratio (l/d) of the fiber which decreases

72

was reported by Maiti and Singh

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2.2.2 Flexural properties: Flexural properties of PHBV-wood fiber composite with
varying wood fiber content from 10 to 40 wt% are shown in Figure 2.4. Flexural
modulus of the composite exhibited an increasing trend with increasing wood fiber
content in the PHBV-wood composites. The flexural modulus showed an average
increase of 28 % with every 10 wt% rise of wood fiber in the PHBV matrix. PHBV
based composite having 40 wt% of wood fiber content showed a notable increase of
168 % when compared with neat PHBV. The flexural modulus of neat PHBV was 1.28
:t 0.02 GPa while, for PHBV-wood composite with 40 wt% of wood fiber it was 3.44 d:
0.17 GPa. During the flexural testing, stress is more localized in the region of applied
load, which is in contrast to the tensile testing, where stress is distributed through out
the specimen. Therefore, the increase in the flexural modulus of the PHBV-wood
composite contribute to the assumption that a reasonable amount of interfacial
interaction existed between fiber and matrix, which allowed the transmission of stress
from matrix to fiber thereby, increasing the stiffness of the specimen. F lexural strength
of the PHBV-wood composite did not show any significant improvement with the
increase of wood fiber in PHBV-wood composition (Figure 2.4). Flexural strength of
composites remained at 30 MPa with a little variation of 3:1 MPa on adding wood fiber
to PHBV, which could be accounted in the standard deviation. Fiber alignment also
plays an important role in determining the flexural strength. The uniaxially aligned
fibers can give a better strength property than the randomly oriented fibers [18].
Flexural modulus of PP based wood fiber composite at 30 and 40 wt% was 3.4 and 4.6

GPa and strength at the same wt% was 51.4 and 55.1 MPa published by Huda [19].

75

Both the flexural properties have a higher order of values than the PHBV-wood fiber

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2. 2.3 Dynamic Mechanical Pr0pertie5:

Storage Modulus: The effect of temperature on the storage modulus of PHBV-wood
composites having differing wood fiber content is depicted in Figure 2.5. A general
trend of increase in storage modulus with the increasing fiber content in the composites
was observed. An average increase of 21 % in storage modulus was observed with
every 10 wt% increment of wood fiber in the PHBV. The storage modulus of PHBV-
wood composite having 40 wt% of wood fiber was improved by around 168 % higher
than that of neat PHBV at 25 0C. The storage modulus of all PHBV-wood composite
decreased with an increase in temperature and converged to a narrow range at high
temperature (Figure 2.5). The reduction in E ' with increasing temperature was due to
the softening of matrix and initiation of relaxation process. At temperatures above Tg,
the molecular mobility and thermal expansion increases thus reducing the interfacial
interaction as a result E' decreased sharply. Storage modulus is associated with the
elastic response of the composite and indicates the stiffness of the material. Being an
inherent property of a material, it is more associated with the molecular response and
therefore, can give a better estimation of the fiber to matrix interaction. The increasing
fiber concentration had increased the storage modulus because of the mechanical
limitation posed by the fibers embedded in the Viscoelastic matrix thereby, reducing the
mobility and the deformation of the matrix with the increasing temperature [26].
Interfacial interaction due to the dipolar nature of hydroxyl group on cellulose (wood
fiber) and carbonyl group on PHBV can not be ruled out [15]. A uniform dispersion of

fiber in matrix had also synergistically acted in the increase of storage modulus of

77

PHBV-wood composite. The effectiveness of the fibers on the modulus of the
composite can be represented by a coefficient factor “C"
I r
C _ (15./E».
r r
(E, /E,)m

 

r r
where Eg and Er are the storage modulus in the glassy and rubbery state[26]. The

subscripts ‘c ' represent composite and “m ' matrix. Lower the value of 'C’, higher is the

effectiveness of the reinforcement. The measured storage modulus at -20 and 20°C at a

r r
frequency of 1 Hz were taken as. E g and Er . respectively. The results are listed in

Table 2.2. The ‘C ’ value was found to be least in case of PHBV-wood composite
having 40 wt% of the wood fiber content. This showed that wood fiber had a good

reinforcing ability in the PHBV based composites.

Table 2.2: Value of coefficient “C” with different loading levels of wood fiber

 

Wood fiber Content Coefficient factor “C”
in PHBV (wt %)

 

10 1.07
20 0.87
30 0.82
40 0.75

 

78

Loss modulus: The Loss modulus of the material is associated with the viscous
response or the dampening effect ofthe material. It is the out of phase component of the
complex modulus and is associated with the dissipating energy. which is unrecoverable
and lost to the system. Figure 2.6 shows the effect on loss modulus (E") of the
composite with the varying wood fiber content and temperature. E " increased with the

increasing fiber concentration and gave a peak in the transition region. This peak can be

assigned as a glass transition temperature (Tg,) of the semi-crystalline material. The

peak shifted slightly towards the higher temperature thus indicating a small rise in Tg

of PHBV-wood composites with the increasing wood fiber content.

Tan§ .‘ The Tan6 is a ratio of the loss modulus to the storage modulus. It gives a peak
in a region, which has rate of decrease in storage modulus higher than that of loss
modulus, and after attaining the maxima. vice-vcrsa happens. thus an emergence of
peak indicates a transition of material from one region to the other. The height of tan
6peak decreased with an increase in the fiber content in the composite i.e., the
dampening effect had reduced with the increased fiber content in the matrix (Figure
2.7). This suggested that some degree of interfacial bonding existed between the fibers
and matrix in PHBV-wood composites. which increased with the addition of fibers to
the PHBV matrix. The higher wood fiber level induced a better fibers packing in the
matrix, and resulted in efficient stress transfer form the PHBV matrix to wood fibers.
thereby decreasing the dampening effect. Widening of the Tan§ peaks with the
increasing fiber content in the PHBV—wood composite suggested the increased fiber to

matrix interfacial regions in the fiber rich matrix as well as an increased molecular

relaxation in the composite system as compared to the neat polymer [26]. The

79

broadening of the peaks at higher level of fiber loading in the system reflects the

heterogeneity generated in it.

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g 4000 l (C)
8 3000 .4
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Figure 2.5: Temperature dependency of storage modulus with varying
loading level of wood fiber in PHBV matrix. (a) PHBV-wood fiber
(60:40), (b) PHBV-wood fiber (70:30), (0) PHBV-wood fiber (80:20), ((1)
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2.2.4 FTIR analysis: Figure 2.8 (a) shows FTIR spectrum of the specimen from

processed neat PHBV, (b) and (c) are of 30 and 40 wt% wood fiber-PHBV composites,
respectively. Bands appearing at 1275, 1226, and 1180 cm-] (marked with arrows) in

FTIR spectra of PHBV are sensitive to degree of crystallinity. [27]. In addition, bands

at 1130,1100,and 1060 cm-1 are also sensitive to degree of crystallization, but to a

lesser extent. Comparing the intensities of the bands at 1226 and 1180 cm-1 of neat

PHBV and 30 wt% and 40 wt% wood fiber loaded PHBV, it is observed that the

intensity of the band at 1226 cm-1 increases and that of l 180 cm-1 decreases thus giving
an inference of the rise in degree of crystallinity in PHBV with the addition of the wood
fiber. Crystallinity index (CI) is a relative measure of degree of crystallization and is

defined as a ratio of the intensity at band 1378 cm-1 (marked with asterisk) to the band

at 1176 cm (i.e., ratio of the ll‘ltCl‘lSlty of the band insensmve to crystallizatlon to the

band sensitive to crystallization) [26]. CI is not be confused with absolute crystallinity.
CI of neat processed PHBV was calculated to be 0.66 and that of PHBV with 30 wt%
and 40 wt% wood fiber was 0.95 and 0.9. Higher CI value of PHBV-wood fiber than
neat PHBV indicates the higher order in degree of cyrstallinity of the PHBV matrix in

presence of wood fiber.

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2.2.5 Heat Deflection Temperature (HDT): The effect of wood fiber reinforcement on
the HDT of the composite is of significant impedance for the reason, that HDT is
considered as an upper temperature limit of a material. HDT of crystalline polymers
depends on thermal history of polymer, crystallite sizes and presence of impurities or
other nucleating agents, which enhance crystallinity [28]. The HDT values of PHBV-
wood fiber composites with different wood fiber loading are depicted in Figure 2.9. The

HDT of PHBV-wood composite. loaded with 40 wt% of wood fiber. was improved by

24 0C as compared to neat PHBV. The major contribution to the improved HDT with

enhanced fiber content is the fiber reinforcement. which has higher HDT than the
matrix. The increased HDT can also be attributed to an increase in the degree of
crystallinity in PHBV due to addition of wood fibers, the fiber surface acted as
nucleating sites therefore, increasing the crystallinity of the wood fiber filled PHBV
composite [29,30]. This is supported. by the FTIR analysis of the PHBV- wood fiber
composite. HDT is directly proportional to the modulus, and inversely proportional to
the applied load at which deformation occurs therefore. an increased storage modulus
of PHBV based composites, at higher temperature. was also the contributing factor in
the increase of HDT with the increasing fiber content in matrix [14. 16. 31].

2.2.6 Coefficient of linear thermal expansion (CLTE): Coefficient of linear thermal
expansion (C LTE) is a vital property for composites in structural applications. A low
CLTE value is a desirable property in order to achieve dimensional stability. CLTE of
crystalline polymers depends on various parameters like processing conditions. themial
history, and degree of crystallinity and crystallite size. Filler aspect ratio. volume

fraction, orientation and distribution in matrix also play a deciding role in determining

85

the CLTE of composites [16, 30]. Figure 2.9 give the details of the CLTE with the

varying wood fiber level in matrix. The wood fiber lowered the CLTE of PHBV. The

CLTE value of PHBV-wood composite was reduced from 182 x 10.6 / OC (i.e., neat

PHBV) to 173 and 154 x 10'6/ 0c. when reinforced with 30 and 40 wt% ofwood fiber,

respectively. C LTE of a composite is affected by the mismatch of the high CLTE of the
matrix and low CLTE of the fibers. The fibers are tightly squeezed in the matrix during
mold cooling because of which the fiber poses a mechanical restrain on the opening of
the polymer chain during heating, and thus reduce the overall C LTE of the composite.
As the ratio of fiber to matrix volume increases the decrease in C LTE of the composite
is mainly due to the lower C LTE of the fiber. Surface of the wood fiber acted as an
additional nucleating site for the formation of crystallite of PHBV matrix therefore.
enhancing the crystallinity of the matrix as revealed from the FTIR spectra of the
PHBV-wood fiber composite [l l]. The increased crystallinity would pack the matrix
more densely, thus resulting in reduction of CLTE. This can be another identifiable
reason for the reduction of CLTE of a semi-crystalline polymer like PHBV when filled

with wood fiber.

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2.2.7 T hermogravimetric analysis (T GA): Figure 2.10 shows the wt% loss of PHBV,

wood fiber and their composites with increasing temperature. The TGA curve of wood

fiber indicated that the actual thermal degradation of wood initiates at 250°C, and its

maximum loss percent occurred at 390°C, completely degrading at 500°C. The TGA

of neat PHBV showed a gradual loss of weight with the increase in temperature. It

could be inferred from the TGA curves that the thermal degradation of PHBV was

uniform with maximum degradation occurring at 310 OC and onsets at 250°C. It has

been reported that PHBV is thermally unstable above 250°C. The degradation process

involves chain scission and hydrolysis which leads to reduction in molecular weight
and formation of crotonic acid [32], the presence of moisture can enhance hydrolytic
degradation of PHBV. The PHBV-wood fiber composites of PHBV having 30 and 40
wt% of the wood fiber showed similar trends of thermal degradation giving a little
information of any effect of increased wood fiber content on the thermal stability of
PHBV—Wood fiber composites. TGA curves of composites indicated that thermal

degradation of composite is a collective thermal degradation phenomenon of PHBV

and wood fiber. Degradation of the PHBV-wood composites onsets around 250 0C,

which is also an onset point of both wood fiber and PHBV, and completes at 390 C’C.

The TGA of PHBV-wood composites also showed that wood fiber was the final content
to degrade in composite. Maximum degradation of composites occurred at the
temperature, which corresponds, to the maximum degradation temperatures of PHBV

and wood fiber. Figure 2.11 shows the derivative weight loss of neat PHBV, wood fiber

88

and PHBV-wood fiber composite. TGA curve peak of the composites shifted towards
the lower temperature from that of neat PHBV, which indicated that the presence of
wood fiber made PHBV a little more thermally unstable than neat PHBV. This could be
due to presence of some other constituent in wood fiber or degradation effect of

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2.2.8 Impact Strength: The Impact test measures the energy required to completely
break the specimen. This energy is a combination of crack initiation and crack
propagation phenomena, which depends on various factors like fiber to matrix
adhesion, toughness of the matrix and fiber alone, defects in the packing of
fiber/matrix, crystalline morphology etc [34]. The notch izod impact strength test
basically measures the energy to propagate an existing crack. 2.12 represents the results
of the notched izod impact strength of PHBV and its wood fiber filled composites with
various concentration levels. The impact strength of PHBV decreased with the increase
of wood fiber content in the PHBV matrix. The impact strength of PHBV-wood
composite decreased from 34 i 2.7 J/m i.e. neat PHBV to 30 :1: 0.5 and 24.4 i 1.2 J/m
when reinforced with 30 and 40 wt% of the wood fiber, respectively. The crystalline
polymers are brittle, and the increase in degree of crystallinity further decreases the
impact strength [18].The crystallinity of PHBV increases with the addition of wood
fiber because the surface of cellulosic wood fiber acted as the additional nucleating
sites for the PHBV matrix. This increased the spherulite prominence in the matrix
around the fibers and constraining it thus, resulting in the reduced impact strength. The
reduced impact strength at a higher fiber loading of 40 wt% due to some fiber
clustering in specimens can not be ruled out. The molecular weight reduction due to
higher processing time (10 min.) of 30 and 40 wt% wood fiber composite as compared
to that of neat PHBV (3 min.) may also be a contributing factor to the reduced impact

strength of wood fiber-PHBV composite. [18, 32, 35].

91

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2.2.9 Morphology: Figure 2.13 and 2.14 shows the photomicrographs of the PHBV-
wood fiber composites with 30 and 40 wt% of wood fiber. Scanning electron
microscopy (SEM) was used to characterize the morphology of the composites. The
surface characterized under SEM was the fractured surface obtained as a result of the
notched impact test. Fiber pull out fiber breakage were the failure mode of the fractured
surface. Figure 2.13 (A) of 30 wt% wood fiber shows a good dispersion of wood fiber
in the PHBV matrix and relatively lesser number of pullouts to the fractured fibers.
Fiber breakage contributes much lesser energy than the fiber pullouts in the net
fractured energy. As can be observed from the SEM micrographs, fiber fracture was
predominant in the fractured surface of the specimen. This is consistent with the low
impact strength at 30 and 40 wt% wood fibers. A lesser amount of fiber dispersion is
observed in Figure 2.14 (A) of 40 wt% fiber than that of 30 wt% fiber specimen. More
number of fiber pullouts were observed on the fractured surface of a specimen with 40
wt% wood fiber content than the specimen with 30 wt% wood fiber content. More fiber
pullouts at 40 wt% of wood fiber should have had given more impact energy than at 30
wt% however, this did not happen because of higher fiber content at 40 wt% wood
fiber there was not sufficient fiber to matrix contact therefore, fiber to fiber contact
dominated in the matrix. Thus insufficient adherence of fiber to matrix lead easy fiber
pullout during the impact. Fiber matrix adhesion remained below the critical adhesion
limit due to which impact energy decreased with increased fiber loading [36]. Figure
2.13 (B) and 2.14 (B) are the photomicrographs at higher magnification which shows

that a significant amount of interfacial interaction exist between fiber and matrix. An

93

interface adhesion allowed better stress transfer from matrix to fiber, which accounted

for superior tensile and flexural modulus.

94

’1

Interface

 

Figure 2.13: SEM photomicrographs of impact - fractured samples
of PHBV-wood (70:30) composite;(A): 60x, 200nm, (B): 200x,
100nm.

95

a ,

. i ' 5‘ i

 

Figure 2.14: SEM photomicrographs of impact-fractured samples
of PHBV-wood (60:40) composite; A: 60x, 200nm, B: 200x,
100m.

96

2.3 Conclusion

Biodegradable composite from polyhydroxybutyrate-co-hydroxyvalerate
(PHBV) and wood fiber were fabricated using extrusion followed by injection molding.
The fabricated PHBV based biocomposites contained 10 to 40 wt% of the maple wood
fiber. The effects of increasing wood fiber weight contents on mechanical,
therrnomechanical and morphological properties of the PHBV based biocomposites
were evaluated. The tensile and flexural modulus of PHBV based biocomposites
reinforced with 40 wt% of wood fiber was improved by ~167% when compared with
neat PHBV. The theoretical tensile modulus values of PHBV based biocomposites
obtained from Halpin-Tsai and Tsai-Pagano equations were consistent with the
experimental tensile modulus values. The storage modulus values of the PHBV based
biocomposites also exhibited the increasing trend with the enrichment of wood fiber in
PHBV matrix. The effectiveness of fibers on the storage modulus of the composite was
evaluated by a coefficient factor, C. The heat deflection temperature was increased by
21%, and the coefficient of linear thermal expansion reduced by 18 % of the
biocomposite at 40 wt% of wood fiber content in PHBV. Scanning electron microscopy
(SEM) photomicrographs exhibited the existence of interfacial interaction between

wood fiber and PHBV thus, providing the compatibility between the two.

97

2.4 References

1.

10.

11.

12.

Principia Partners presentation at 7th International Conference of Wood fiber-
Plastic composites May 19-20, 2003; Madison, WI.

 

htt izt’lxx-‘vvvy'. rinci iiaconsultinueomr’consulting-‘nen'sctinfartit:lee—.78 accessed

date 05/15/2006).

Selke S E, Wichman I .Wood fiber/polyolefin composites. Composites Part A;
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Harper D, Wolcott M. Interaction between coupling agent and lubricants in wood-
polypropylene composites. Composites Part A; 2004; 35; 385—395.

Li T Q, Wolcott M P. Rheology of HDPE-wood composites. 1. Steady state shear
and extensional flow. Composites Part A; 2004; 35; 303-31 1.

Stark N M, Berger M J. Effect of Species and Particle Size on Properties of
Wood-Flour-Filled Polypropylene Composites; 1997.

Lenz R W, Marchessault R H. Bacterial Polyesters: Biosynthesis, Biodegradable
Plastics and Biotechnology. Biomacromolecules; 2005; 6(1); 1-8.

futilz’iisuji:.1191Luis:_.L‘Lm 1,1'T11Lixx sil,9???):‘ff-.’..SJ?UF l30~‘f1.il.l -i’_‘,;?_".20_‘:?30:31:11 ml (accessed date
5/15/2006)
ht t.ii::’:’.’.\:~;\5;\5;._1_uh11s..9Lim. :"biL. 259i {ha 0 Chitin}; ., f l 1013;: 03:01.0.uni!irtiL‘l L75-.:7;:iui L l 9.2 8.1.1.101 l.

(accessed date 5/15/2006).

Hao Y, Zhu M, Zhang Y, Chen Y. Mechanism of the Formation of Concentric

ring like Patterns on PHBV Spherulites. Polymer-Plastics Technology and Engg.;
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Reinsch V E, Kelley S S. Crystallization of Poly (hydroxybutyrate-co-
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Sics I, Ezquerra T A, Nogales A, BaltaC F J, Kalnins M, Tupureina V. On the
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during Isothermal Crystallization of Bacterial Poly(3-hydroxybutyrate-c0-3-
hydroxyvalerate) Copolymers. Biomacromolecules; 2001; 2(2); 581-587.

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Mohanty, A K, Misra M, Drzal LT. Sustainable Biocomposites from Renewable
Resources: Opportunities and Challenges in Green Materials World. Journal of
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Bhardwaj R, Mohanty A K, Drzal L T, Pourboghrat F, Misra M. Renewable
Resource-Based Green Composites from Recycled Cellulose Fiber and Poly(3-

hydroxybutyrate-co-3-hydroxyvalerate).Biomacromolecules; 2006;(Published on
web 04/26/2006).

Dufresne A, Dupeyre D, Paillet M. Lignocellulosic Flour-Reinforced
Poly(hydroxybutyrate-covalerate) Composites. J Appl Polym Sci ; 2003;
87;]302—1315.

Luo S, Netravali A N. Interfacial and mechanical properties of environment-
friendly “green” composites made from pineapple fibers and
poly(hydroxybutyrate-co- valerate) resin. Journal of Material 80.; 1999; 34; 3709
- 3719.

Jotier. J.J, Limieux. E, Prud’homme. ER. Miscibility of polyester/Nitrocellulose
blends: A DSC and FTIR study. Journal of Polymer Science: Part B: Polymer
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Fie.B, Chen.C, Wu.Hang, Peng.S, Wang. X, Dong.L. Comparative study of
PHBV/TBP and PHBV/BPA blends. Polymer International; 2004; 53; 903- 910.

Huda M S, Drazal L T, Mohanty A K, Misra M. Wood fiber reinforced Poly
(Lactic acid) composites. 5th Annual SPE Automotive Composition conference,
Troy, MI, USA, 2005, Sept 12-14.

Nielsen E L. Mechanical properties of Polymers and Composites. In: Marcel
Dekker, editor.vol.2. New York, 1976.

Tsai S W, Pagano N J, Composite Material Workshop, Technology Publishing
Co., New York; 1968.

Karrad S, Cuesta J, Crespy M L. Influence of a fine talc on the properties of
composites with high density polyethylene and polyethylene/polystyrene blends,
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23. Nicolais.L, Nicodemo.L. The Effect of particle shape on tensile properties of

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glassy thermoplastic composite; Intem.J.Polyeric Mater.;1974; 4; 229-243.

Wood Handbook: Wood as an Engineering Material; Forest Products Laboratory
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Maiti.S.N, Singh.K. Influence of wood flour on polypropylene: J. Applied
Polymer Sc.;1986; 32;4285-4289.

Pothana L A, Oommenb Z, Thomas S. Dynamic mechanical analysis of banana
fiber reinforced polyester composites. Composites Science and Technology;
2003; 63; 283—293.

Bloembergen. S, Holden.D.A, Hamer.G.K, Bluhm.T.L, Marchessault.R.H;
Studies of Composition and Crystallinity of Bacterial Poly(B-hydroxybutyrate-co-
B-hydroxyvalerate; Macromolecules; 1986; 19;2865-2871

Takemori M T. Towards an Understanding of the Heat Deflection Temperature of
Thermoplastics; Polymer Sc. and Engg.;1979;19(15);1 104-09.

Avella M, Martuscelli E, Raimo M. Properties of blends and composites based on
poly(3-hydroxy)butyrate (PHB) and poly(3-hydroxybutyrate -hydroxyvalerate)
(PHBV) copolymers; J. of Material Sc.; 2000; 35;523-545.

Tan .1 K, Kitano T, Hatakeyama T J. Crystallization of carbon reinforced
polypropylene. Mater Sci; 1990; 25; 3380-84.

Kardos J L, Raisoni J, Piccarolo S. Prediction and Measurement of the Thermal
Expansion Coefficient of Crystalline Polymers. Polymer Sc. and Engg.; 1979; 19;
5; 1004-09.

Gatenholm P, Mathiasson A. Biodegradable Natural Composites. II. Synergistic
effect of processing cellulose with PHB. Journal of Applied Polymer Sc.,1994;
51; 1231-1237.

Femandes E G, Pietrini M, Chiellini E. Bio-Based Polymeric Composites
Comprising Wood Flour as Filler. Biomacromolecules; 2004; 5(4); 1200-1205.

Kamdem.P, Jiang H C, Freed J W, Matuana M L. Properties of wood plastic
composite made of recycled HDPE and wood flour from CCA-treated wood
removed from service. Composites: Part A 35; 2004; 35; 347-355.

Zhang J, McCarthy S, Whitehouse R. Reverse temperature Injection Molding of
BiopolTM and effect on its properties. Journal of Applied Polymer Science; 2004;
94; 483-491.

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In. Jhon Wiley and Sons. Editor; New York, 1980.

100

Chapter 3

Renewable Resource Based Biocomposites from Natural Fiber
i.e., Bamboo Fiber, and Polyhydroxybutyrate-Co-Valerate
(PHBV) Bioplastic and its Analogy to Wood Fiber Reinforced

Bioplastic Composites

3.0 Introduction

In the established market of plastic composites there has been a dynamic
emergence of wood and natural fiber plastic composite (W/NF PC). These have a
proven market in outdoor and indoor applications like decking, railing, fencing,
furniture, etc., and they also hold a large share of the European market in automotive
applications [1]. Most of the thermoplastic used in W/NF PC are recyclable and is
being look to as a profitable opportunity for these materials. Contemporary, W/NF PC
are based on conventional themioplastics and naturally available cellulose fiber. with
the fiber content varying from 30 to 70 wt%. Various kinds of naturally available
fibers and wood fiber have been used to achieve good mechanical properties of WW F
PC. A significant amount of research is focused on improving the compatibility
between the hydrophilic natural fiber and the hydrophobic thermoplastics [2—5]. Nano-
reinforcement using inorganic material like clay in wood-PP composites has also been
reported elsewhere [6]. The major limitation these composites encompass is their
limited dimensional stability, when exposed to moisture [7].

The environmental concerns due to the accumulated plastic wastes over the last
few decades have obligated the biodegradable green plastics to surrogate petroleum

based polyolefins. Polyhydroxyalkanoates (l’HAs). a family of natural. biodegradable

101

polyesters, differ from petroleum-based synthetic polymers such as polypropylene and
polyethylene by their synthesis from renewable resources, and by their unproblematic
biodegradability. The similarity of some of its mechanical properties to the polyolefins
indicates that it may be a substitute for polyolefins. The biodegradability and
biocompatibility of PHBV has given it a remarkable breakthrough and application in
biomedical applications [8, 9]. Its composites with natural fiber have been studied
extensively [10-12]. Wood and other naturally available fibers are generally used to
reinforce thermoplastics due to their low cost, abundant availability, high performance
and low density [12]. Among the existing WPC products, maple wood fiber is the most
widely used fiber commercially; whereas among the naturally available fibers in Asia,
bamboo fiber has the most abundant applications [13, 14]. Therefore, these two fibers,
obtained from two different categories of lignocellulose-based sources are the most
used fibers in the commercial sector. In the previous chapter wood fiber was evaluated
for its reinforcing ability in PHBV.

The prime focus of the study in this chapter was to fabricate and evaluate
bamboo fiber based PHBV biocomposites. Bamboo is a more sustainable source of
natural fiber than wood fiber due to the short growth period of 3 to 5 years, regrowth
from its shoot therefore, avoid deforestation, and higher strength. In fact, bamboo is the
commonly used structural material in housing construction in south-east Asia.
Fundamental aspects of these bamboo based biocomposites were measured using
differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), Fourier
transform infrared spectroscopy (FTIR), and their morphology was also explored.

Furthermore, the experimental results of tensile modulus were compared with the

102

existing theoretical model. Comparative analysis of the processed bamboo fiber
biocomposites with the earlier reported wood fiber-PHBV biocomposites [15] in the
previous chapter has been discussed; along with performing a two-way ANOVA on
their mechanical properties.
3.1 Experimental Section
3.1.] Materials

Polyhydroxybutyrate-co-valerate (PHBV) under the trade name Biopol (Zeneca
Bio Products) was obtained from Biomer, Germany. Madake bamboo fiber with lengths
of ~ 500 um and diameters between 10- 100 um was supplied by the Kyoto Institute of
Technology, Japan.
3. .I .2. Fabrication of composites employing micro-compounding technique:

In order to avoid any moisture induced degradation during processing, the

bamboo fiber and PHBV were dried at a 110 oC and 80 0C under vacuum for two and

four hours, respectively. The dried fibers and PHBV were kept in air-tight plastic bags
until the time of processing. The composites were prepared using a micro-extruder
(DSM Research, Netherlands) with a barrel volume of 15 cc. The extruder was
equipped with co—rotating twin-screws having a length of 150 mm and L/D of 18. The

temperatures in three heating zones of the micro-extruder (top, center and bottom) were

165, 160 and 155 0C. The processing temperature was kept between 155 and 165 0C in
order to achieve low melt flow and avoid thermal degradation at higher temperature.
The screw rotation speed was 100 rpm. After fixed processing time in the micro—

extruder, the molten mix was transferred to a preheated mini-injection molding

machine for specimen fabrication. The optimized injection pressure for all PHBV-

103

bamboo fiber composites was 551.5 KPa and the temperature was 155 C)C. The

different mix ratios of PHBV and fibers along with some other processing parameters

are given in Table 3.1.

Table 3.1: Processing parameters of PHBV and fiber compositions

 

PHBV (wt%) Fiber (wt%) Processing Retention time

 

 

 

temperature (0C) (min)
100 0 160 3
70 30 160 10
60 40 160 10

 

104

3.1.3 Testing and Characterization
3.1.3.1 Mechanical Properties

Tensile and flexural properties of specimens were measured using a United
Calibration Corp. SFM 20 testing machine as per ASTM standards D638 and ASTM
D790, respectively. Tensile testing was carried out at a rate of 2.54 mm/min with a pre-
load of 0.023 kg. The system controls and data analysis were performed using the
supplied Datum software. The notched Izod impact testing was carried out on a Testing
Machine Inc. (TMI) model 43-02-01 as per ASTM D256 with a pendulum of impact
energy 1.35 J.
3.1.3.2 T hermo-mechanical Properties
3.1.3. 2.1 Dynamic mechanical analysis (DMA)

The storage modulus, loss modulus and tan delta of PHBV and its natural fiber

composites were measured using a TA Instruments model Q800 DMA. The specimens

were heated from -30 to 125 0C at a heating rate of 2 oC/min. The single cantilever

mode was used to test the specimens at amplitude of 15 pm and frequency of 1 Hz.
3.1.3.2.2 Heat deflection temperature (HDT)
The DMA was used to determine the HDT as per ASTM D648. Rectangular

bars 2 x 12 X 58 mm in three-point bending mode with an applying load of 66 psi were

used for testing. Samples were heated at the rate of 2 OC/min from room temperature to

the temperature required to produce 0.195% strain.

105

3.1.3.3 Thermal Properties

3.1.3.3.] Thermogravimetric analysis (T GA )

TGA was tested using TGA 2950 instrument with a heating rate of 20 OC/min

from room temperature to 600 0C, the temperature of complete thermal degradation.

All TGA runs were performed under a nitrogen purge.
3.1.3.3.2 Differential Scanning Calorimetry (DSC)
DSC of the samples was performed using a model Q100 DSC, TA Instruments.

The heat/cool method was adopted for all specimens testing. Specimens were heated to

190 c)C at a ramp rate of 10 oC/min and cooled with the same ramp. The data for

melting points and heat of crystallization were collected during the heating cycle. The
crystallization temperature and heat of fusion were obtained from the cooling cycle.
Data analyses of DMA, HDT, TGA and DSC results were performed using the
supplied software Universal Analysis.
3.1.3.4 F T IR analysis
A Perkin-Elmer system 2000 FTIR Spectrometer was used with an attenuated
total reflectance (ATR) accessory to gather the IR spectra of neat PHBV and its natural
fiber composites.
3.1.3.5 Morphological studies
Morphological studies were conducted using a JEOL scanning electron
microscope (model J SM-6400). Specimens were sputter-coated with gold to a thickness

of ~10 nm before surface characterization in order to make minimize charging. The

106

SEM was equipped with a lanthanum hexaboride (LaB6) crystal as electron emitter

source. An accelerating voltage of 13 KV was used to collect the SEM images.

3. I . 3.6 Statistical Analysis

Statistical analysis of tensile and flexural modulus results was carried out by

performing two-way ANOVA on the PHBV and its composites with wood and

bamboo fiber at 30 and 40 wt% using MIN ITAB® software.

3.2 Results and Discussions

3.2.1 Tensile Properties: Figure 3.1 shows the Young’s modulus of PHBV and its
composites with bamboo fiber at 30 and 40 wt.% loadings. The modulus of the PHBV
composites increased steadily with the fiber loading. At 30 wt% fiber loading, it
increased by 67%, and at 40 wt.% the increase was 175%. The tensile properties of
short fiber composites depend on fiber length, distribution, orientation and their
interfacial bond strength with matrix. The fiber orientations have been assumed to be
random; as observed visually, and the composites behavior isotropic. The increased
modulus of the composites can be ascribed to uniform dispersion of fibers in the PHBV
matrix, which leads to the even distribution and transfer of stress from matrix to fiber.
Uniform dispersion of fibers was physically visible in the specimens. The tensile
strength of PHBV decreased with addition of the bamboo fibers. This can be attributed
to the lack of sufficient interfacial interaction between the fiber and matrix. The
presence of polar group on the lignocellulosic natural fibers and the abundant moderate
polar groups in the chemical makeup of PHBV lead to an expectation of some
interfacial interaction between the two; although from the FTIR analysis, no chemical

or hydrogen bond interactions between the moderately polar groups of the two

107

components were observed. Reader may refer figure 3.7 for this purpose. Therefore, the
only possible interactions responsible for the load transfer from matrix to fiber were
frictional and Van der Waal forces. To increase the interfacial bonding between the
fiber and matrix use of compatiblizer lies under the scope of future work.
Christensen ’s Model

The theoretical Young’s modulus of the composites has been estimated using
Christensen’s model of equations [16, 17]. It provides a solution for the elastic

properties of composites with randomly oriented short fibers. The following expression

predicts the elastic modulus of composites below and above the critical length ([6) of

the fiber.

1 1,
Ecz—gva [l—El—j+vam at l>l

C

1 l

Ez E

c Evf +vE [<10

f E m m at

EC denotes the modulus of the composite, E f and Em that of fiber and matrix (in this
case bamboo fiber and PHBV), Vf and V”, are the volume fractions of fiber and
matrix, respectively, l is the fiber length and [c the critical fiber length which is given

by an expression of the Rosen model. For a bamboo fiber lc< l.

 

 

l/2
[c 115 l—vy2 Ef
df N ° v1,” Gm df is the diameter ofthe fiber

108

Gm has been estimated from a general relation

_ E
_ 2(1 + V) v is Poisson’s ratio

 

Figure 3.2 depicts the theoretical and experimentally determined values of the modulus.
Theoretical values of lignocellulosic natural fibers (bamboo) and the PHBV matrix
have been approximated to be 25 GPa [18] and 1.02 GPa, respectively. Volume fraction

of the fiber and the matrix at 30 and 40 wt.% were calculated using general equations

,0 f
m f m m and vm : (1 _ vf ) assuming density

Pf Pm

 

 

of natural fiber ( ,0 f ) and matrix (pm) to be ~ 0.84 gm/cm3 [19] for bamboo fiber and

1.25 gm/cm3 for PHBV. The modulus of bamboo fiber-PHBV composites as predicted

by Christensen’s model are in close approximation to the experimental results. This was
due to the uniform and narrow size distribution of bamboo fibers, although not
quantified but observed visually. The difference between the theoretical and
experimental data can be due to the agglomeration of the fillers at high volume
fractions, which results in the higher modulus than the expected. The agglomeration
was visible in SEM microphotographs. Although the composites have been assumed to

be isotropic; anisotropy of the fibers cannot be ruled out. The model does not account

109

for the effect of filler and processing conditions on the matrix (PHBV) e.g. change in

the degree of crystallinity, which had been noted from the F TIR and DSC analysis.

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3.2.2 Flexural Properties: Flexural modulus and strength of PHBV and its bamboo
fiber composites with different fiber loadings are shown in Figure 3.3. In general, the
modulus increased with incorporation of fibers. At 30 wt.% of fiber content the PHBV
modulus improved by 100% and the increase was greater at 40 wt.% fiber with the
modulus leap of 154%. The flexural strength of PHBV remained unchanged upon the
addition of fibers. The improved flexural modulus and retained strength of the
composites validated the reasonable amount of compatibility and interfacial interaction
of cellulose fiber to the natural polyester PHBV. Fiber orientation in the matrix plays
vital role in determining the flexural properties. Unlike tensile testing, where induced
tensile stress plays crucial role, the flexural testing incorporates both compression and
tensile stresses on the fiber, matrix and their interface across the specimen cross
section, therefore, the strength response to the stress during flexural testing would be
higher than during tensile testing. The tensile strength is chiefly govern by the
dewetting of matrix from the fiber under tensile load and also is more sensitive to the

defects originated at the interfaces of matrix and fiber.

113

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3. 2.3 Dynamic Mechanical Analysis (DMA)

The Viscoelastic characteristics of the materials were measured using DMA. It is
a complex quantity defined with elastic (storage) and viscous (loss) components;
denoted E' and E", respectively. In DMA, the effect of temperature on the Viscoelastic
properties of the composites was studied. This gives more distinct insight to the
interactions at the molecular level of the materials.

Figure 3.4, depicts the change in the E' of PHBV, and its composites with
temperature. Like a generic tendency of thermoplastic composites, E' of PHBV and its
composites decreases with rising temperature. E' of PHBV increased with the fiber
loading up to 40 wt.%. The difference between the modulus values of bamboo fiber

composites at 30 and 40 wt.% was not significant. E' of bamboo fiber composites

improved by 157 and 173% at 30 and 40 wt.% fiber content at 25 0C, respectively. The

increase in E' of the fiber filled PHBV with respect to neat PHBV was greater below Tg

than above it. This can be attributed to the temperature sensitivity of the natural fiber.

The slope of the E' vs. temperature curve around the glass transition (Tg) temperature

(~ 1 0C) of PHBV increases with the fiber loading. This gave an indication of increased

brittleness of these composites. A small bump in the curve of PHBV at 52 °C is
indicative of the cold crystallization region. During the heating cycle as the specimen
enters the cold crystallization region, the increasing crystallinity leads to decrease in the

reduction rate of E', as a result a small bump in the curve emerges in this region. The

cold crystallization temperature (TCC) decreases with fiber loading. The ch decreases

115

 

to 46 0C in case of bamboo fiber composites. This could be due to very low diameter or

high aspect ratio of the bamboo fiber versus the wood fibers which gave it a better
opportunity to act as nucleating agent in PHBV matrix. Among the dynamic
mechanical properties, dampening is most sensitive to the structural heterogeneity,
morphological and transitions changes of system, and fiber-matrix interfacial
interactions. Figure 3.5 shows the dampening of the PHBV and its fiber composites.

Tan 6 expressed as the ratio of E" to E’ gives account of the energy dissipated as heat

during the dynamic testing. Tan 6 peak occurs approximately 10 to 20 c)C above the Tg

of the PHBV and its fiber composites. Increased stiffness of the PHBV with the fiber
loading decreases the tan 8 peak value thus reflecting the reduced energy losses in the
interfacial regions. Low dampening of the composites in the transition region is
attributed to the fineness of the bamboo fibers, which allows their uniform distribution
and crystallinity effect on the matrix. The widening of the tan 5 curve also indicates the
increased heterogeneity of the specimens which was observed from the SEM
photomicrographs. The maxima of the loss modulus of PHBV and its composites
(Figure 3.6) occur in the glass transition temperature region. The enhanced response of
molecular chains to the temperature in the amorphous region increases the viscous

characteristic of the material in the glass transition domain. The fiber incorporation into

the PHBV slightly increases its Tg towards higher temperature, thus delaying the

increasing viscosity component of the composites. This can be due to the restriction
imposed by the fibers on the movement of the expanding molecular chains in the

amorphous region.

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3. 2.4 F T IR analysis

In the PHBV IR spectra. Figure 3.7(a), the band at 1720 cm.l was the C=O

stretch of the ester group present in the molecular chain of highly ordered crystalline

structure [20]. A slight shoulder emerging at 1740 cm.1 was of the C=O stretch in the

amorphous region of semicrystalline PHBV. Characteristic peaks of symmetric —C—O—

C— stretching vibration were present from 800 cm.I to 975 cm.I and the
antisymmetric —C—O—C—stretching between 1060 cm-1 and 1150 cm-1 [21]. The band
around 1378 cm.1 corresponds to the symmetrical wagging of CH3 groups and at 1453

-l . . . .
cm was due to asymmetric deformation of methylene groups, the 1ntensrty of both
the peaks is independent of crystallinity or in other words insensitive to crystallinity.

The absorption bands at 1720, 1276, 1228, and 980 cm.1 crop up from the crystalline

regions (designated with arrows) and bands at 1740, 1453, 1176 cm-1 from the

amorphous regions (designated with two face arrows) of PHBV. Peak at 1228 cm-1

was purely of the conformational band of crystalline helical molecular chains and
therefore, are absent in amorphous region [22]. The crystallinity index (Cl) gives the

relative measure of degree of crystallinity in PHBV. It is given as

CI : Iatl378cm-I

atll76cm‘l

 

(ratio of intensity of the band insensitive to crystallization to

the band sensitive to crystallization), where I is intensity of the peak at a given wave

number. Figure 3.7 (b), (bamboo and PHBV composites) depicts the IR spectrum of

120

 

fibers and PHBV composites. On the comparison of the intensity of peaks at 1228

cm.1 and 1176 cm-1 of PHBV and its fiber composites, it is observed that these vary

inversely to one another with the changing degree of crystallization. The CI of PHBV
0.66 increased to 0.74 with the addition of 30 wt.% bamboo fiber indicating the
increase in the crystallinity of PHBV more, when the bamboo fiber was induced. CI

value of similar order was obtained even though when a 40 wt.% bamboo fiber was

added to the PHBV.

121

 

 

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3. 2.5 Thermogravimetric Analysis (T GA )

The TGA curves in Figure 3.8 depict the weight loss of fibers, PHBV and its

composites with temperature. Bamboo fibers began to thermally degrade at 250 OC,

reaching the maximum at 380 OC and entirely degrading at around 500 C’C. The PHBV

lost weight gradually with the increasing temperature and its thermal degradation

occurred maximum at 316 0C while insetting at 250 C)C. Its thermal instability above

250 c)C has earlier been reported stating that the degradation process involves chain

scission and hydrolysis which leads to reduction in molecular weight and formation of
crotonic acid [23] and the presence of moisture further promotes hydrolytic
degradation. The thermal degradation of biocomposites appeared to be cumulative

phenomena of thermal degradation of matrix (PHBV) and fiber alone. Degradation of

composites initiated around 250 C)C in case of both, fiber and PHBV and ceased around

400 oC. The TGA of composites also showed the fibers to be final constituent to

undergo degradation in composites. The composite degradation maxima occurred at the
temperature, a little higher than that of PHBV. Figure 3.9 illustrates the derivative
weight loss of PHBV, fibers and their composites at 30 wt.% fiber. Maximum
degradation temperature peak of the composite shifted to the higher temperature from
that of PHBV indicating it to be a slightly more thermally stable than PHBV. Thus the

presence of bamboo fiber (a lignocellulosic fiber ) did not had any degradation effect

on PHBV.

123

120

 

 

 

g
2 i (C)
.9 l
3.} l
a 40 -
g l (b)
& 20 l (a)
l
O 100 200 300 400 500
Temperature (°C)

Figure 3.8: Percentage weight loss of (a) PHBV, (b) PHBV/Bamboo fiber (70:30)

 

 

 

 

 

(c) Bamboo fiber.

4.5 1
A 4 ‘
o .
2— ; a
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E 2 a
2 15 .‘
‘3 ' <— b
e 1 i ( >
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0: m, -L . , 1
0 100 200 300 400 500 600
Temperature (°C)

Figure 3.9: Derivative weight loss of (a) PHBV, (b) PHBV/Bamboo fiber (70:30)
(c) Bamboo fiber.

124

3. 2.6 Heat Distortion Temperature (HDT)

The HDT is an upper use temperature limit of a material, therefore, the effect of
fiber type and loading in composites has a considerable importance in their structural
applications. The HDT of semicrystalline polymers like PHBV depends on thermal
history, crystallite sizes and distribution along with the presence of impurities and other
nucleating agents which affects their crystallinity [24]. Figure 3.10 has a description of
the HDT values of PHBV and its composites with bamboo fibers at 30 and 40 wt.%

fiber content. The HDT of PHBV was enhanced by addition of fibers at 30 and 40 wt.%

from 114 0C to 120 OC and 123 0C, respectively. The major contributing factor to the

improved HDT with higher degree of fiber content was the fiber reinforcement which
had better HDT than the matrix. The increased storage modulus of composites at higher
temperatures also contributed to the increased HDT. Similar types of observations have

been earlier reported [11, 15].

130 i
06‘ 120 i
110 ~-
l
100 - , ~ ~ . , ,
Neat PHBV 30/70 (BF/PHBV) 40/60 (BF/PHBV)
Fiber/PHBV (wt%)

Figure 3.10 : HDT of Bamboo fiber/PHBV composites having 30 and 40 wt% fiber.

125

 

3. 2. 7. Impact Strength

Impact strength depends on various factors like fiber-matrix adhesion,
toughness of the matrix and fiber, defects in the packing of fiber/matrix, crystalline
morphology, etc. [25]. The notch Izod impact strength test measures the energy to
propagate an existing crack. Figure 3.11 depicts the notched Izod impact strength of
bamboo fiber composites with PHBV. The impact strength of PHBV (34.3 i 2.7 J/mt)
decreased with the addition of fiber but there was no significant reduction when the
fiber content was increased from 30 to 40 wt.%. The higher fiber volume in matrix gave
manifold fiber to fiber contact (fiber clustering) with respect to fiber to matrix contact,
which did not allow the sufficient fiber matrix interfacial interactions to develop. As
seen from the micrograph discussed in the following section majority of the fi'acture
occurred through the fiber rather than through the matrix or the interface of matrix and

fiber.

 

Izod Impact Strengthn (Jlmt)
M
O

      

Neat PHBV 30/70 (BF/PHBV) 40/60 (BF/PHBV)
Fiber/PHBV (wt%)

Figure 3.11: Notch impact strength of PHBV/ Bamboo fiber composites at 30 and
40 wt% loading levels.

126

3. 2. 8 Morphology

Figure 3.12a and 3.12b shows the micro photographs characterized under SEM
at different magnifications of the impact fractured surfaces at 30 wt.% of bamboo fiber
composites. The failure modes observed during the fracture were fiber pull-out and
fiber breakage. Due to the higher volume fraction of bamboo fiber the majority of the
surface could be seen dispersed with the fibers along with the clustering. A number of
voids were also visible on the fractured surface that had contributed to the low impact
and tensile strength with the fiber loading. Fiber separation in the clusters contributes
much lesser to the energy than the fiber pull-outs from the matrix in the total fractured
energy during the impact. This was consistent with the observation of low impact
strength of the bamboo-PHBV composites, (Figure 3.13). At 40 wt.% fiber content
fiber clustering dominated over the fiber matrix adhesion thus giving insufficient
adherence of fiber to matrix which lead easy fiber pull-out and separation during
impact. F iber-matrix adhesion remained below the critical adhesion limit, due to which,

the impact energy decreased [26].

127

 

 

Figure 3.12 SEM photomicrographs of impact-fractured samples of
PHBV/Bamboo (70:30) composite: (a) 90X. 200 um, (b) 4SOX, 50 pm

128

 

fractured samples of

Figure 3.13 SEM photomicrographs of impact

200 mm, (b) 500x, 50 pm

9

PHBV/Bamboo (60:40) composite: (a) 90><

129

3. 2. 9 Analogy of Bamboo fiber-PHBV and Wood fiber-PHBV biocomposites

There wasn’t any significant difference among the tensile properties of bamboo
and wood fiber based composites at the same loading level even though the aspect
ratios of the two fibers were different (length and diameter of the wood fibers were
between 1600-1650 um and 300-400 um [15] and that of bamboo fiber were ~500 um
and 10-100 um, respectively). Table 3.2 lists various mechanical properties of the two
biocomposites at 30 and 40 wt.% fiber. The flexural strength of both types of
biocomposites remained somewhat unaltered with the addition of fibers up to 40 wt.%,
and the modulus in case of wood fiber composites was greater than the bamboo fiber
composites by 14 % .This may be due to the higher order stiffiiess of maple wood fiber
versus bamboo fiber. Furthermore, the effect on tensile and flexural modulus of fiber
type (wood and bamboo) and loading level (30 and 40 wt.%) of natural fibers in PHBV
was statistically analyzed by performing a two-way ANOVA. Table 3.3 represents the
F , F critic and P values on the effect of fiber type, their loading level and an
interaction of the type and loading levels on the afore mentioned characteristics.
ANOVA was performed at a significance level (a) of 0.05. From the above mentioned
values, it can be inferred that the type of fiber does not have any significant effect on
the elastic and flexural modulus of composites whereas, the fiber loading at 30 and 40
vvt% have significant effect on both the moduli. The interaction effect of both
parameters on tensile and flexural modulus varies. For the tensile modulus the
interaction effect of fiber class and loading was significant but for the flexural modulus

it was insignificant. This is consistent with the experimental results because the tensile

130

properties are more sensitive to the defects (stress concentration) originated during

processing or fabrication of the specimens.

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The HDT of the wood fiber biocomposites was greater than that of the bamboo

fiber biocomposites by 15 0C at both the loadings levels, whereas the increase from 30

wt% to 40 wt% was a marginal of 2 C>C in both the natural fibers biocomposites. The

difference between the HDT of bamboo and wood fiber composites can be attributed to
the more uniform distribution of low density and small fiber size of bamboo fiber
which increased the volume of fiber with respect to matrix in the composite. This did
not allow the sufficient matrix packing and crystallite size growth as it happened in
case of wood fiber based composites. This was supported by the FTIR and DSC
analysis depicting the increased crystallinity of wood fiber based composite versus the
bamboo fiber based PHBV composite.

From the comparative TGA of both kinds of composites, the thermal stability of
PHBV-Bamboo fiber composites was a little more than that of PHBV-Wood fiber
composites [15]. The reduced thermal stability of PHBV in presence of wood fiber has
earlier been reported in publication [15]. This can be due to the lesser amount of lignin
in bamboo fiber than maple wood fiber (table 3.5). The effect of delignification on the
mechanical properties of PP based WPC has been reported by Alinaghi et-al [27].

The impact strength of PHBV composites with wood fiber was observed to be
higher than that with bamboo fiber. This can be reasoned due to the higher volume
fraction of bamboo fiber than wood fiber at same wt.%. As observed from FTIR and
DSC results, low crystallinity in the bamboo fiber biocomposites may had resulted the
lower impact strength than the wood fibers biocomposites. Wood fiber surface gave a

nucleation growth effect thus giving a better embedment of fiber in the matrix.

133

 

Figure 3.14 are the photomicrographs at higher magnification of composites
depicting the fibers embedded in the matrix and its wettability on the fiber surface
indicating the similar interfacial interactions existing between the two different

lignocellulosic fiber and PHBV matrix.

134

 
   

Wood fiber

 

Figure 3.14: SEM Photomicrographs at high magnification (a) PHBV/ Wood
fiber composite 1500X , 20 um(b) PHBV- Bamboo fiber composite. 2300X, 10

135

Differential Scanning Calorimetry analysis of Bamboo fiber-PHBV and Wood fiber-
PHB V biocomposites.

Non-isothermal DSC study of PHBV and its fiber composites was carried out in
order to have an understanding of the effect of fibers on the nucleation and crystallinity
of PHBV. Figure 3.15 shows the thermograms of PHBV and its 30 wt.% wood and
bamboo fiber composites. The details are given in Table 3.4 where Tm is the melting
temperature and Hc is the cold crystallization enthalpy. A bimodal endotherrn of PHBV
during the heating cycle is due to the crystallization phenomena of the heterogeneous
crystal morphology formed during processing. Renstad et. al [28] have carried out a

detailed study of this. Embedding the fibers in PHBV does not vary its melting point

which remains at 154 0C for both kinds of fiber composites. The enthalpy of melting or

crystallization (Hf) in both the cases of fibers composites remains same, thus depicting
the degree of crystallization was unaffected by the type of fiber in the PHBV. Mergence
of the second peak with the primary peak of PHBV was an observable phenomenon in
30 wt.% bamboo fiber composite which was absent in 30 wt.% wood fiber composite.
The fine bamboo fiber interferes in the crystal morphology of PHBV during the
processing as a result the nucleation growth of PHBV became more homogenized in
presence of fine and uniformly blended bamboo fibers in the matrix. Spherulite
homogenization of PBS in presence bamboo fiber had been reported by Kori et. al [29]
.High volume fraction of bamboo than wood fiber at same wt.% may have increased the
nucleating density of the matrix but no appreciable difference between the net

crystallinity of both types of composites was observed. It is to be noted that no such

effect was noticed at 40 wt.%. The melt crystallization temperature (TC) of PHBV was

136

observed to be 101°C. The addition of 30 wt.% wood fiber increased this temperature

to 104 OC, whereas the bamboo fiber decreased it to 99 OC; indicating an increase in

the rate of crystallization in the presence of wood fiber and decrease in the case of
bamboo fiber. This may be due to slower diffusion of molecular chains though the well
dispersed bamboo fiber to the nucleus thus decreasing the crystallinity rate in case of
bamboo fiber composite. (Figure 3.15) Similarity in the melt crystallization peaks of
PHBV and its fibrous composites during cooling provided the evidence of having any

divergence in their crystalline morphologies.

137

Heat flow Exo up —>

 

(a)

 

(b)

 

60 80 100
temperature(°c )

0 50 100

(a)

(b)

4— heat flow down

120

150

140

160

200

Figure 3.15: Crystallization and melting curves of (a) PHBV (b) PHBV/Bamboo fiber

(70:30) composite (c) PHBV /Wood Fiber (70:30) composite

138

 

 

 

 

 

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FTIR of the bamboo and wood fiber were carried out to determine an estimate of their

contents and composition. Figure 3.16 shows F TIR spectra of wood and bamboo fibers

(described peaks marked with arrows). The region around 3400 cm.1 depicted
hydrogen bonded O-H stretching and at 2900 cm-1 C-H stretching. The peak at 1740
cm'1 was the C=O stretch in hemicelluloses present in both lignocellulosic fibers. 1505
and 1600 cm-l are characteristic peaks of lignin which arises due to the aromatic

skeletal vibration of benzene ring in lignin. The peak at 1600 cm-1 has a contributions

from the conjugated C=O group along with aromatic skeletal stretch [30]. The ratio of

the two (1505 cm-1 to 1600 cm-1) gives some qualitative aspects of the two different

types of lignins present in different kinds of natural fibers. Similar spectrums of wood
and bamboo fiber suggested the presence of same proportion of cellulose,
hemicellulose and lignin in both types of natural fibers, as supported in the literature
[30]. Composition of lignin and cellulose in both fibers are listed in Table 3.5 [31].
Wood fiber composites had a higher CI values (0.95) than bamboo fiber composites at
30 wt.% fiber content indicating greater crystallinity of PHBV in presence of wood
fiber than bamboo Similar results were noticed at 40 wt.% fiber content in both type of

composites [15].

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3.3 Conclusions

Biocomposites from a bamboo fiber and poly(hydroxybutyrate-co-
hydroxyvalerate) (PHBV) were fabricated using extrusion followed by injection
molding. Tensile and flexural modulus of this biocomposite increased with fiber
loading. The storage modulus also improved with fiber addition and no appreciable
difference among the two compositions (30 wt.% and 40 wt.% fiber) was noticed.
Comparative analysis of two different type of lignocellulosic fiber (wood and bamboo)
at 30 and 40 wt.% fiber loaded biocomposites were studied with respect to their
mechanical, therrnomechanical, thermal, FTIR and morphological aspects. Both types
of lignocellulosic fibers embedded in the PHBV matrix gave an appreciable rise in
tensile and flexural modulus. Statistically, there was no effect of the fiber type on
mechanical properties of composites. A difference between the net crystallinity of both
types of composites with different fibers was observed in differential scanning
calorimetry (DSC) and infrared spectroscopy (FTIR). Notch impact strength of PHBV
decreased with the fiber addition and the reduction was greater in case of bamboo fiber
composites. No effect of the fiber type on the thermal stability of composites was
present, but incorporation of wood fiber gave a higher heat deflection temperature
(HDT) versus the bamboo fiber. Scanning electron microscopy (SEM)
photomicrographs demonstrated the existence of similar kinds of interfacial interaction

between the two types of fibers and PHBV

143

 

3.4 References
1) Karus M, Ortmann S. Vogt D. Use of natural fibers in composites in the German
automotive production 1996 till 2003 (2004-10); http://\\-'ww.nachwachsende-

rohstol‘feinfo (date accessed: 12/4x’06).

2) Ferrer FM, Vilaplan F. Greus AR. Borras AB. Box CS. Journal of Polymer Science
2006; 99 : 1823-31.

3) Huda MS, Mohanty AK, Drzal LT, Shut E, Misra M. Journal of Materials Science
2005; 40: 4221 — 4229.

4) Rana AK, Mitra BC, Banerjee AN. Journal of Applied Polymer Science 1999; 71:
531—539.

5) Lu JZ, Negulescu 1 1. Wu Q. Composite Interfaces 2005; 1221-22125—140.

 

6) Yeh, Shu-Kai, Casuccio, Mary E, Al-Mulla, Adam. Annual Technical Conference -
Society of Plastics Engineers 2006; 64: 209-213.

7) Khalil HPSA, Shahnaz SBS. Journal of Reinforced Plastics and Composites 2006;
25:12: 1291.

8) Luzier WD. Proc. Nati. Acad. Sci. USA. Feb.1992; 89: 839-842. (Colloquium
Paper).

9) Williams SF, Martin DP, Horowitz DM, Oliver P, Peoples. International Journal of
Biological Macromolecules 1999; 25: 1 1 1-121.

10) Dufresne A, Dupeyre D, Paillet M. Journal of Applied Polymer Science 2003; 87:
1302-1315.

11)Bhardwaj R, Mohanty AK, Drzal LT, Pourboghrat F, Misra M. Biomacromolecules
2006; 6: 7: 2044-2051.

12) Femandes EG, Pietrini M, Chiellini E. Biomacromolecules 2004; 5: 4: 1200-1205.

13) Kitagawa K, Ishiaku SU, Mizoguchi M, Hamada H, in Natural fibers, Biopolymers.
and Biocomposites, Edited by Mohanty AK, Misra M, Drzal LT. Taylor and
Francis 2006.

14) Yongli M, Xiaoya C, Qipeng G, Journal of Applied Polymer Science 1997; 64:
1267—1273.

144

15) Singh S, Mohanty AK, Wood fiber reinforced bacterial bioplastic composites:
Fabrication and performance evaluation. Composites Science and Technology
2007; 67 (9): 1753-1763.

16) Berlin AA, Volfson SA, Enikolopian NS, Negmatov SS. Principles of Polymer
Composites. Springer-Verlag. Berlin Heidelberg; 1986.

17) Christensen R. Mechanics of Composite Material. John Wiley & Sons. New York;
1979.

18) Deshpande A P, Rao MB, Rao CL. Journal of Applied Polymer Science 2000; 76;
83—92.

19) Rajulu AV, Chary K, N, Reddy GR, Meng YZ. Journal of Reinforced Plastics and
Composites 2004; 23: 127.

20) Padermshoke A, Katsumoto Y, Sato H, Ekgasit S, Noda l, Ozaki Y. Polymer 2004;
45: 6547—6554.

21) Suthar V, Pratap A, Rawal H, Bulletin of Material Science 2000; 23: 3: 215—219.

22)Hong SG, Chen WM. e-Polymers 2006; 024: (www.epolymers.org , ISSN 1618-
7229)

 

23) Gatenholm P, Mathiasson A. Journal of Applied. Polymer Science 1994; 51: 1231—
7.

24) Takemori MT, Polymer Science Engineering 1979; 19: 15: 1104—9.

25) Kamdem P, Jiang HC, Freed JW, Matuana ML. Composite: Part A 2004; 35: 347—
55.

26)Agarwal BD, Broutman LJ, Analysis and performance of fiber composites. John
Wiley and Sons. New York.1980.

27) Alinaghi K, Saleh N, Ismaeil G, Mehdi T, Ghanbar E. Journal of Applied Polymer
Science 2006; 102: 4759—4763.

28) Renstad R, Karlsson S, Albertsson AC, Werner PE, Westdahl M. Polymer
International 1997; 43: 201-209.

29) Kori Y, Kitagawa K, Hamada H. Journal of Applied Polymer Science 2005; 98:
603—612.

30) Pandey K K. Journal of Applied Polymer Science 1999; 71: 1969—1975.

31) http://www.PaperListeningStudvorg (date accessed 03/03/07)

145

 

Chapter 4

Anisotropy and Characteristic Comparison of Short fiber

Bioplastic Composites .

4.0 Introduction

As discussed in previous chapters, bioplastics are coming up as an effective
alternatives to the petro-based plastics in various commercial sectors like packaging,
automotive interiors, structural, biomedical applications etc.. Among these,
polyhydroxyalkanoates (PHAs), a bacterial sourced bioplastic has found some foot
holds in the commercial sector. Significant serious attempts have been made to address
the processing, brittleness and cost relating issues of Polyhrdroxybutyrate-co-valerate
(PHBV) in past. These issues collectively have been approached either by blending,
plasticizing, filling with available cheap fillers, or by altering branching of the
molecular chain in order to alter its crystallinity. Fabricating PHBV based composites,
using naturally available fiber or wood fiber in order to enhance its performance as a
commodity material has been attempted by various researchers. The processing of the
composite has been carried out using either of the techniques, compression molding,
injection molding or solvent casting. In the short fiber composites, random and uniform
distribution of the fibers in a matrix leads to the theoretical assumption of an isotropy in
a composite material, with no preferred directional orientation of fibers. But, the
anisotropy associated firstly, with the fibers itself [1] and secondly, the alignment or
orientations of the fiber during the directional flow of the melt during the melt mixing
and injection molding engender an intrinsic anisotropy in the specimens. Therefore,

development of the anisotropy, during processing, is unavoidable in short fiber

146

,7

composites. Anisotropy in short fiber composites can become a governing factor in
their performance when the applied stress is not in a favorable direction during the
application.

Chao-Ming Lin et al. reported the anisotropic behavior of sheet molding
compound from polyester and chopped glass fiber, theromosets during processing with
compression molding [2]. In another paper, Ming Tian et a1. discuss the influence of
two roll mill processing on an anisotropic behavior in mechanical properties of fibrillar
silicate /rubber nanocomposites [3]. Bernd Lauke describes a modified rule of mixture
to analyze the anisotropic strength of the short fiber composites with respect to the
loading direction [4]. Different theoretical and computational analyses of short fiber
composites have been made by various authors to relate anisotropy of the composites
with their mechanical, thermo-viscoelastic behavior, the fiber ellipticity, and fraction
[5-10]. References [1 1-15] discuss the various aspects of anisotropic behavior of short
fiber reinforced polymeric composites associated with their mechanical characteristics.

Wanjun Liu discusses the effect of different processing methods on the kenaf
natural fiber reinforced soy based biocomposites [11] and in another kenaf reinforced
polypropylene composite Zampaloni elaborates the manufacturing difficulties and
associated solutions [12]. None of the published papers until now compare the two
different processing techniques, i.e., injection molding and compression modeling of
the natural short fiber composites from PHBV in terms of the anisotropy associated
With them. This chapter attempts to draw a direct comparison between the compression

molded specimens and injection molded specimens of the same composite in terms of

147

 

the anisotropy associated with the specimens and their static and dynamic mechanical
properties.

4.1 Experimental section

4.1.] Materials

Polyhydroxybutyrate-co-valerate (PHBV) was obtained from Biomer, Germany with
the trade name Biopol (Zeneca Bio Products). Wood fiber used was the maple wood
fiber from the American Wood fiber Schofield, WI with the trade name 2010 MAPLE.
4.1.2 Fabrication of composites

Two different processes were used to fabricate PHBV and wood fiber composite with
different compositions. Table 1, list the compositions and the fabrication method along

with other processing parameters associated with the specific process. The wood fiber

and PHBV pellets were dried at a 110 OC and 80 0C under vacuum for 2 and 4 hours

respectively, to devoid of any moisture during processing.

4.1.2.1 Injection molding

Specimens were fabricated using micro-compounder, (manufactured and designed by
DSM Research, Netherlands with a barrel volume of 15 cc) and a mini injection
molder. This equipment was a co-rotating twin-screw extruder with the length of 150
mm, L/D (length to diameter ratio) 18 and has three consecutive heating zones. PHBV
pellets and wood fiber in a fixed composition as listed in Table 1 were physically mixed
and fed to the micro-extruder, allowing the two to melt- mix for a fixed time period in

the extruder. The processing temperatures of the three consecutive zones were 165, 160

and 155 OC. The melt from the extruder was collected in a pre heated mini injection (at

148

 

O . . . . .
a temperature of 155 C), and was transferred to a m1n1-mold1ng machine for speC1men

fabrication. The optimized injection pressure listed in Table 1 was used to extricate the
melt-mix from the injection cylinder to the mold located in the molding machine.
Finally, the specimen was removed from the mold and conditioned for 48 hrs before
testing.

4.1.2.2 Compression molding

Both, PHBV pellets and wood fiber were physically mixed in a given proportion and

spread uniformly within the periphery of 3 mm thick window frame sandwiched

between the two heating platens of a Carver Press SP-F 6030 held a 160 oC. The

physical mixture was pre heated for 10 minutes and compressed to a pressure of 800
KPa initially and gradually increased up to 1245 KPa. The sheets were maintained at
this pressure for 15 minutes and eventually cooled to room temperature in order to
remove from the press. The resulting sheets were 2.7 mm thick and 11X 11 square
inches in area. The testing specimens for the appropriate test were cut from the sheets

as per the dimensions listed in the next section.

Table 4.1: Processing parameters of PHBV and PHBV-Wood fiber composite

 

 

 

Composition Injection Molding Compression molding
Injection Processing Processing Compression Processing
Pressure time temperature pressure time
(KPa) (minute) (0C) (KPa) (minute)
Neat PHBV 552 3 160 - -
Wood fiber-PHBV 825 6 160 1245 15
(30-70 wt %)

 

 

 

 

 

 

149

 

4.1.3 Testing and characterization

4. 1. 3. 1 Mechanical properties

Tensile and flexural properties of specimens were tested using a United Calibration
Corp. SFM 20 testing machine as per ASTM D638 and ASTM D790 standards,
respectively. The system controls and data analysis were performed using Datum
software.

4.1.3.2 Dynamic mechanical analysis (DMA)

The storage modulus, loss modulus and tan delta of the test specimens were evaluated

using DMAQ800 supplied by TA Instruments. The specimens were heated from -30 to

125 0C with a heating rate of 2 OC/min. Single cantilever mode was used to test the

specimens at an oscillating amplitude of 15 um and a frequency of 1 Hz.

4. 1.3.3 Anisotropy analysis

4.1.3. 3.1 Squeeze flow test:

This test was done to determine the developed orientations of wood fibers in the
specimens during injection or compression molding process. In case of injection
molding, three circular discs of 13 mm diameter were incised from three different
sections of the 4 mm thick specimen, i.e. gate area, middle portion and the far off end
section of the specimen. In case of compression molded samples, 25 mm diameter discs
were incised from the different areas of the compression molded sheet of thickness 3
mm. These discs were marked with an axis corresponding to zero degree direction and
positioned at the center between two teflon sheets. Teflon sheets and disc was

sandwiched between two smooth surface steel plates and this whole unit was placed

between the pre heated platens at 160 OC, i.e., the thermoforrning temperature. After

150

 

attaining an equilibrium temperature for 1 minute platens were compressed with a load
of 500 lbs for 60 seconds. The part was unloaded and allowed to cool immediately at
room temperature within the mold ensuring that it retains the deformed shape. An
illustration of the test along with the pre and post shapes of specimens is given in
Figure 4.1 on next page.

4. 1.3.4 Morphological studies

Morphological studies were conducted using a JEOL scanning electron microscope
(model ISM-6400). Specimens were sputter-coated with gold to a thickness of ~10 nm

in order to prevent charging during the examination. The SEM used was equipped with

a lanthanum hexaboride (LaB6) crystal as a source of electrons. An accelerating voltage

of 13 KV was used to collect SEM images.

151

 

 

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4.2 Results and Discussion

4. 2.1 Tensile properties:

Figure 4.2 illustrates the Young‘s modulus and tensile strength of PHBV and its 30
wt% loaded wood fiber composite with two different processing methods i.e.,
compression molding and injection molding. The modulus and the strength of the
composite were significantly affected by the two different processing methods. The
Injection molded specimens show better characteristics than the compression molded
specimens. The enhancement in the modulus of injection molding specimens was 104%
with respect to the neat PHBV, and practically, no loss in the strength. Where as, in
case of compression molded specimens the modulus remained same as of neat PHBV
with a significant loss of 50% in the strength. The wood fiber acted as filler in
compression molded specimens rather than a reinforcing agent. Lower tensile
characteristic of the compression molding composite than the injection molding can be
reasoned due to the poor interfacial adhesion of the matrix to the fiber. The mixing of
fibers in the matrix was expected to be more unifomted and percolated during the
extrusion than the physically mixed and spread fiber during compression molding.
More precise reason for the difference among the properties of the two methods as
discernible from the SEM micrographs was the entrapment of air in the matrix during
the compression molding. The entrapped air did not allow an intricate adhesion of the
matrix to the fiber as a result little stress transfer occurred from matrix to fiber during
loading conditions. The entrapped air lead to the formation of non-uniform and uneven
cell wall foamed like structure of the PHBV matrix, thereby reducing its specific

strength, and additionally, contributing to the reduction in strength.

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4. 2. 2 F lexural Properties:

Flexural properties of composites reflected no different trend than tensile
properties with both the different kind of molding processes. In case of injection
molding modulus doubled to 4.5 GPa from 2.2 GPa of PHBV with no deflated strength
of the composite but precipitously decreased strength and modulus of the compression
molded samples. Figure 4.3 show the details of the flexural strength and modulus of 30
wt% wood fiber filled PHBV composite fabricated using both the processes i.e.,

compression and injection molding process.

._ El Flexural modulus ._ 50

8
7 -- O Flexural strength
6

Flexural Modulus (GPa)
.5
'1+
Flexural Strength (MPa)

 

 

 

 

 

0 l l 0
Neat PHBV Compression Injection molded
molded

Figure 4.3: F lexural properties of PHBV and its composite loaded with 30 wt%
wood fibers

155

 

 

4.2.3 Dynamic Mechanical Analysis

Dynamic mechanical analysis was carried out to characterize the fiber matrix
interaction in composites. DMA allows a better estimate of interactions occurring at the
molecular level of a material. In visco-elastic materials modulus is a complex quantity,
which is the sum of elastic and viscous component of a material, and inherently
changes with the temperature of the material. The elastic component of the material is
described by the storage modulus, and the viscous component by the loss modulus.

Figure 4.4 shows the change in storage modulus of composites with the increasing

temperature at the rate of 2 OC/minute; The figure depicts the typical trends as of

thermoplastics and their filled composites in these cases. The storage modulus of
composites is higher than the neat PHBV because the high rigidity of the fiber induces
an additional stiffness to the matrix. The storage modulus of injection molded
specimens was higher than that of the compression molded at all the temperature. The

increase with respect to neat PHBV of injection molded composite was higher than that

of compression molded by 17 % before (at 0 OC) and 27 % after (at 25 0C) the glass

transition temperature (Tg ). This was due to the better compaction of fibers in the

matrix of the injection molded composites than the compression molded ones. At a

temperature above Tg in rubbery region fibers poses restrictions to the extending

molecular chains in the region of influence i.e., matrix around fibers, as a result, the
increase in the storage modulus in rubbery region is higher than in the glassy region.

Where the molecular chains are virtually immobile. In case of injection molded

156

 

composites these restriction becomes more predominant as fibers are imbedded more

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In case of compression molded composites the trapped air in the matrix and the fiber
interface provides an ample movement to molecular chain without much restriction
imposed by the loosely imbedded fibers thus contributing in low degree to the stiffness.
The trapped air in the matrix and loosely held fibers are visible in the SEM micrographs
discussed in the forth written section.

Figure 4.5 & 4.6 shows the loss modulus and tan 5 variation, respectively with
respect to the temperature on x- axis. Loss modulus is the outcome of the loss in
irrecoverable energy due to dampening effect of the molecular movements and
interactions in the material. A peak of loss modulus vs temperature represents the
transition region of the polymeric material from glassy to rubbery state. Dampening
effect in the fibrous composite can result from the molecular chain motion in the
amorphous phase, fiber-matrix interfacial regions and fiber-fiber frictional interactions
in the fiber conglomerated regions. Loss modulus of compression molded composite
was higher than the injection molded composite. This can be due to the more
amorphous region in the compression molded composite due to morphological
heterogeneity in the PHBV matrix because of air entrapment during processing as
visible in the SEM micrographs. Another source of dampening was the energy loss due
to friction among fibers in the fiber conglomerated regions. Compression molded
composites were processed just by physically mixing the fiber and PHBV pellets.

Tan 8 is the dampening factor generally decreases with the reinforcement in the
composite system. The theoretical decrease in the tan 5 value of the composite with

respect to neat matrix is by the weight percentage of the reinforcement filler in the

158

 

i

composition i.e., tan 5c = tan 5m (l-(D) where tan 5c , tan 5m are the values of

composite

Compression

 
 
   
  

Injection molded

. . "\R.‘ :-_- 35’

     

/’

 

Loss Modulus (MPa)
3
O

 

 

 

50‘ Neat PHBV
‘Efi ' 4
o I I I I I I I
-30 -10 10 30 50 70 90 110
Temperature (°C)

Figure 4.5: Loss Modulus of PHBV and its 30 wt% wood fiber loaded composite

    
   

 

 

0.1 -

0-08 ' Neat pHB 1 Compression
S
7, 0.06 -
1:
C
.2 0.041

0'02 ‘ Injection molded

0 V I j I I I fi
-30 -1o 10 30 50 7o 90 110

Temperature (°C)

Figure 4.6: Tan 5 of PHBV and its 30 wt% wood fiber loaded composite

159

and matrix respectively, and (I) is the weight fraction of fibers in the composite (30
wt%, (1) =03). The experimental decrease in the dampening factor of the injection
molded composite was 32% and that in compression molded composites was just 11%.
The broadening of the tan 5 peak accounts for the widening of the transition region.
Figure 6 depicts the broadening of the compression molded tan 5 peak more than that of
the injection molded composite. High tan 5 peak value and wide transition region
further confirms the added heterogeneity due to the trapped air and non uniform mixing
of fibers in the compression molded composite.
4. 2.4 Squeeze flow testing

Short fiber thermoplastic composites are assumed to have random fiber
orientation and therefore, behave isotropic. During the fabrication, either by injection
molding or by compression molding, short fibers do develop some orientation and
impart some directionality to the specimens. In case of injection molding process the
melt mix of thermoplastic and fibers flows through the channels of mold and narrow
gate opening. Fibers confluence in the thermoplastic melt during the flow tends to
develop orientation in the flow direction and adopt that orientation in the cooling stage
thus imparting a specific directionality to the specific regions of the same specimen. In
case of compression molding, since the manufacturing involve the physical mixing of
thermoplastic pellets and fibers, the random orientation distribution of fibers is not
guaranteed and fibers may sustain some directionality during fabrication. The goal was
to determine these inherently developed orientations during both fabrication processes.
A similar kind of anisotropy study of natural fiber filled thermoplastic composites had

been done by Zampaloni et al [17,18]

160

A squeeze flow test as illustrated in the previous section was utilized to
determine the fiber orientation of the wood fiber reinforced PHBV for injection and
compression molding composites. An ideal isotropic short fiber composite disc
deformed or compressed would result in a circular shape after the completion of
compression process. Otherwise, a non circular shape or elliptical shape if obtained
confirms the assumption of the directionality associated with the fiber orientations in
the thermoplastic matrix. This is due to the flow velocity and strain rate within the
matrix not being constant with respect to the fiber orientations described by angle “8”.

A polar coordinate system was used to assign the marking on the compressed discs with

. O . . O . . . .
an increment of 20 startmg Wlth 0 marked 1n1t1ally, before the compressmn. The

length of the lines from center to periphery of the compressed samples were recorded
and plotted against the marked angles as shown in figure 4.1. For an ideal circularly
deformed sample this plot must be a straight line. The squeeze flow test was performed
on 20 wt% and 40 wt% wood fiber loaded injection molded composites and 30 wt%
wood fiber loaded compression molding composite. Two different compositions were
taken in case of injection molded process in order to show the significant effect of fiber
loading on the inherently developed anisotropy of the composite.

Figures 4.7 and 4.8 shows the squeeze flow tests of 20 and 40 wt% wood fiber
loaded PHBV injection molded composites. These plots are utilized to determine the
preferred orientations of fibers by observing the trends of maxima and minima along
the X axis i.e., angle. The test was performed on the three different sections of

specimens as specified before. From the comparison of both the compositions it was

161

3'3:

 

observed that the length change with respect to the angle in 20 wt% wood fiber samples

is less variant

Elongation from center (mm)

Elongation from center (mm)

Elongation from center (mm)

40 ;

35 W
30 .

25

20

15

1O

5

o . . - -_ .
O 40 80 120 160 200 240 280 320 360

Angle along part (dog)

(a)

40
35

30 M
25 .

20 '

15

1O

0 4O 80 120 160 200 240 280 320 360
Angle along part (deg)

35 W
30

0 40 so 120 160 200 240 280 320 360
Angle along part (deg)
(C)
Figure 4.7 Squeeze flow test plots of 20 wt%
fiber loaded injection molded composites (a)

Gate section, (b) Center section, (0) Far off
end section

162

 

Elongation from center (mm) Elongation from center (mm)

Elongation from center (mm)

45
4O
35
3O
25
20
1 5
1O

5

O

45
4O
35

W

0 4O 80 120 160 200 240 280 320 360
Angle along part (deg)

(a)

M

30 .

25
20
15
1O

4O
35

0 4O 80 120 160 200 240 280 320 360
Angle along part (deg)

(b)

m

30 -
25 -
20 .

15

10-

0 40 80 120 160 200 240 280 320 360

Angle along part (deg)

(C)

Figure 4.8 Squeeze flow test plots of 40 wt% fiber loaded
injection molded composites (a) Gate section, (b) Center
section, (0) Far off end section.

163

 

than in the 40 wt% at all the three sections of the specimen i.e., gate, middle and far off
end. This was due to uneven resistance offered by fibers at higher concentration to the
thermoplastic melt flow and this gives a quantified estimate of the directionally
orientated distributions of fibers at certain angles. In other words, the certainty to obtain
a complete randomly oriented distribution of fiber at higher concentration in a
reinforced PHBV during injection molding was low. The squeeze flow tests at the three
sections of the specimen of each composition independently, (Figure4.7 (a, b, c) and
4.8(a, b, c)) were performed in order to examine the uniform random distribution of
fibers in the bulk of the specimen. The channeled flow of melt during the injection
molding causes a set of fibers to directionally align differently in different sections of
the same specimen. Investigating all the three sections (i.e., gate, center, far off) of the
specimen, it was evident from the plots that the change in lengths with respect to angles
differs from section to section. The trends of maxima and minima in any of the section
were not similar to the either of the other two sections. Had the fiber orientation been
similar through out the bulk of the specimen the plots of all the sections would be
similar to one another or a straight line in case of the randomly distributed fibers. A
close observation reveal an analogous underpinning trends in the maxima and minima’s
of the corresponding sections of the two compositions i.e., 20 wt% and 40 wt% fiber
loading. This signifies that fibers take up similar orientations in the corresponding
sections of both compositions irrespective of the concentration.

Figure 4.9 shows the squeeze flow test of compression molded specimens at 30
wt% fiber loading in PHBV. The plot of length change versus angle is linear to some

extent, which concludes the absolute random distribution of fiber orientations in the

164

compression molded samples although the specimens were prepared by the physical
mix-ups of fibers and pellets. From the analysis it was observed that the compression
molded PHBV-wood fiber composite exhibit more isotropic characteristic than the

injection molded.

 

E

gAeo:
LE!
§§40'
a; 20~
E! 0-
ll]

0 40 80 120 160 200 240 280 320 360
Angle along part (deg)

Figure 4.9 Squeeze flow test plots of 30 wt% fiber loaded
compression molded composites

 

165

4. 2.5 Morphology

Figure 4.10 and 4.11 are the SEM micrographs of 30 wt% wood fiber filled
PHBV composites processed by compression and injection molding, respectively. Air
channels developed due to air entrapment in the matrix during the processing were
observed in the compression molded composite which were not evident in the injection
molded specimens. At a higher magnification more inadequate interfacial contact
between the fiber and the matrix was observed in case of compression molded than
injection molded composite. Air entrapment and poor interfacial contact were primarily
responsible for the lower tensile and fiexural properties of compression molded

composite.

4.3 Conclusion

Short wood fiber—PHBV composites were fabricated using two different
molding processes i.e., injection molding and compression molding. These were
characterized for the static and dynamic mechanical prOperties and isotropic
characteristics. Injection molded composite had a better reinforcing effect from the
wood fibers than compression molding. This was evident from the better tensile and
flexural properties of injection molded composite which gave an improvement in
young’s modulus and fiexural modulus by 104%. High elastic modulus and low
dampening effect apparently suggested the same during dynamic mechanical analysis.
Morphological observations carried out using SEM to study the fiber matrix interaction
revealed poor adhesion and air entrapment in compression molded specimens.

Anisotropy study carried out using squeeze flow test of both the molding processes

166

 

 

indicated the anisotropic behavior of injection molded specimens. Anisotropy varied

with the sections of the specimens and. composition of wood fiber in the composites.

 

167

   

Figure 4.10: SEM micrographs of 30 wt% wood fiber loaded PHBV
compression molded composite (a) lOOX 20 um (b) 1700X 20 um

168

   

Figure 4 11: SEM micrographs of 30 wt% wood fiber loaded PHBV
injection molded composite (a) 60X 500 um (b) 500 X 50 um

169

4.4 Reference

1.

10.

11

F .R. Cichocki Jr., J .L. Thomason, Thermoelastic anisotropy of a natural fiber,
Composites Science and Technology 62 (2002) 669—678

Chao-Ming lin and Cheng-I Weng, Anisotropy in Sheet Molding Compounds
During Compression Molding polymer composites, October 1997, 18, 5.

Ming Tian, Lijun Cheng, Wenli Liang. Liqun Zhang The Anisotropy of
Fibrillar Silicate/Rubber Nanocomposites Macromol. Mater. Eng. 2005, 290.
681-687.

Bemd Lauke, Shao-Yun Fu, Strength anisotropy of misaligned short-fibre-
reinforced polymers, , Composites Science and Technology 59 (1999) 699i708.

A. Tounsi, E. A. Adda Bedia and A. Benachour, A New Computational Method
for Prediction of Transient Hygroscopicwith Different Degrees of Anisotropy
Stresses during Moisture Desorption in Laminated Composite Plates, Journal of
Thermoplastic Composite Materials 2005; 18; 37.

Ning Pan, Analytical characterization of the anisotropy and local hetrogenity of
short fiber composites: Fiber Fraction as variable, Journal of composite
Materials, 28,16, 1994.

P. S. Theocari & C. B. Demakod’ The effective anisotropy of short-fiber
composites due to fiber ellipticity, Composites Science and Technology 57
(1997), 677-685.

M. H. Santarecomposites, Anisotropic effective moduli of cracked short-fiber
reinforced R. S. Feltman, Journal of Applied Mechanics, 1999, Vol. 66, 709.

E. Kontou *, A. Kallimanis Thermo-visco-plastic behaviour of fibre-reinforced
polymer composites, Composites Science and Technology 66 (2006) 1588—
1596.

J.Beten, Creep damage and life analysis of anisotropic materials, Archives of
applied mechanics 71, 2001, 78-88

. G. V. Kozlov,l A. I. Burya,2 G. E. Zaikovl ,Anisotropy of 21 Fiber Structure and

the Frictional Wear of Composites on the Basis of Polyarylate, Journal of
Applied Polymer Science, Vol. 100, 2821—2823 (2006)

. Yuanxin Zhou, P.K. Mallick Fatigue Performance of an Injection-Molded Short

E-Glass Fiber-Reinforced Polyamide 6,6. I. Effects of Orientation, Holes, and
Weld Line, POLYMER COMPOSITES—2006

170

I

 

13.

14.

15.

16.

17.

18.

Xin-Xing Zhang, Can-Hui Lu, Mei Liang, Preparation of Rubber Composites
from Ground Tire Rubber Reinforced with Waste-Tire Fiber Through
Mechanical Milling Journal of Applied Polymer Science, Vol. 103, 4087—4094
(2007)

A. Godara, D. Raabe Influence of fiber orientation on global mechanical
behavior and mesoscale strain localization in a short glass-fiber-reinforced
epoxy polymer composite during tensile deformation investigated using digital
image correlation Composites Science and Technology 67 (2007) 2417—2427

Marie-Laure Dano, Guy Gendron, Francois Maillette and Benoit Bissonnette
Experimental Characterization of Damage in Random Short Glass Fiber
Reinforced Composites, Journal of Thermoplastic Composite Materials 2006;
19;79

Wanjun Liu, L. T. Drzal, A. K. Mohanty, M. Misra, Influence of processing
methods and fiber length on physicalproperties of kenaf fiber reinforced soy
based biocomposites, Composites: Part B 38 (2007) 352—359

M. Zampaloni, F.Pourboghart, Kenaf natural fiber reinforced polypropylene
composite: A discussion 0 manufacturing problems and solutions, Composites
:Part A 38, 2007, 1569-1580.

M. Zampaloni, A multiple fiber orientation constitutive model for fiber mat
reinforced thermoplastics with a random orientation applied to the stamp
thermo-hydroforrning prcess. PhD Dissertation , Michigan State University,
2003.

171

 

 

 

Chapter 5

Hybrid Green Composite from Talc, Wood Fiber and Bioplastic:

Fabrication and Characterization

5.0 Introduction

As discussed in previous chapters, certain limitations of PHBV like, its narrow
processing temperature range, thermal degradability, low toughness, and post
crystallization phenomena does not allow to render its capability to perform as a
material for the desired applications in structural, packaging, automotive and other
sectors. Natural or wood fiber reinforcement of PHBV matrix was a step forward in the
direction to improve its performance, while retaining biodegradability of the material.
The abundance, economical, renewable, environment friendliness, and easy

assimilation to the processing techniques make these as a sustainable reinforcing agent

in the plastics. Inorganic particulate (clay, talc, CaCO3) reinforcement in plastics is a

common practice to improve the mechanical properties of plastic in a composite

industry. Tale is a 2:1 layer phyllosilicate mineral with Mg3(Si4Olo)-(OH)2 as the unit

structure. Talc alone or in combination with other particulate filler has been used as
reinforcement in certain thermoplastics like polypropylene, nylon to improve upon their
stiffness, tensile strength, creep resistance, and deformation temperature. In addition to
these, talc reduces shrinkage, enhances scratch hardness and surface quality of the final
product [1-7]. Its low cost and ease of use during processing makes it a profitable

material to the processing industry. The major drawback of talc is that it limits the

172

 

 

ductility and eventually reduces the toughness of the thermoplastic. Its high density (~
2.5-2.7 gm/cm3) increases the weight to volume ratio of its composite, when used in a
significant amount as a second constituent. In composites, surface wettability is the

crucial parameter responsible for the interaction between talc particles and the polymer

. . . . O
matrlx. Gllllan et al. reports the water contact angle of talc as 90 , and some other

. . O . . . .
researchers Within a range of 60 to 90 , thus, charactenzmg 1t as a hydrophoblc

material [8].

The objectivity of this research focuses to develop a completely biodegradable
hybrid biocomposite consisting of talc and wood fibers in PHBV. Talc and PHBV have
similar water contact angles; therefore both have similar hydrophobic surface nature.
High surface energy of talc surface allows the wettability of PHBV and its miscibility
leading to subsequent reinforcement of talc platelets in PHBV matrix. The talc
reinforcement provides an additional benefit of cost, as the low cost talc replaces
PHBV in the system. The increased density of the tale and PHBV system can be
Counter balanced by surrogating the certain proportion of talc with wood fiber, thereby,
leveling the density of the composite to certain extent. This hybrid system of mineral
and wood fiber, when reinforced in a bioplastic, i.e. PHBV, is expected to be
cOmpletely biodegradable as all the components of the composite system are sourced

naturally and are independently, biodegradable.

173

 

5.1 Experimental section
5.1.] Materials

Polyhydroxybutyrate-co-valerate (PHBV) was obtained from Biomer, Germany
with the trade name Biopol (Zeneca Bio Products). Micro sized talc was supplied by
Luzenac America under the trade name of Jetfill 700C. The particulate dimensions of
the talc range from 2 to 10 um. Wood fibers used were the maple wood fiber from the
American Wood fiber Schofield, WI with the trade name, 2010 MAPLE.
5.1.2 Fabrication of composites
Hybrid biocomposite specimens were fabricated using micro-compounder
(manufactured and designed by DSM Research, Netherlands with a barrel volume of 15
cc) and a mini injection molder. It was equipped with co-rotating twin-screws of the

length 150 mm and L/D (length to diameter ratio) 18, and has three consecutive heating

zones. The wood fiber and PHBV were dried at a 1 10 OC and 80 0C under vacuum for

2 and 4 hours, respectively, to devoid of any moisture during processing. The hybrid
composite fabrication was carried out as two step process, and that of the PHBV and
wood fiber composite in a single step. The different compositions and processing
parameters used during composite fabrication are listed in the Table 5.1. For the hybrid
composite, PHBV pellets and wood fiber in fixed compositions as listed in Table 5.1
were physically mixed and fed to the micro-extruder. PHBV pellets and wood fibers
Were allowed to melt mix in the extruder for 4 minutes, and latter talc, as per the
Composition, was added to the melt mix of PHBV and wood fiber. Finally, all the three

Components were mixed for additional 4 minutes in the micro-extruder to attain an

174

appropriate dispersion of the talc and wood fibers in the melt. The processing

temperatures of the three consecutive zones were 165. I60 and 155 OC. The melt from

. . . . . . o
the extruder was collected 1n a pre heated m1n1 injecuon (at a temperature of 155 C),

and was transferred to molding machine for specimen fabrication. The optimized

injection pressure of 551.5 KPa was used to extricate the melt from the injection

cylinder and thrust to the mold located in the mini molding machine. Finally, the

specimen was removed from the mold and conditioned for 48 hours before testing.

Table 5.]: Processing parameters of the composites

 

 

 

PHBV Wood fiber Talc Processing Temperature Retention time
(wt%) (wt%) (wt%) (°C) (min.)

100 - - 160 3

60 40 160 7

70 30 10 160 8 (4+4)

80 20 20 160 8 (4+4)

90 10 30 160 8(4+4)

 

175

 

 

5.1.3 Testing and characterization
5. 1. 3. 1 Mechanical properties

Tensile and flexural properties of specimens were tested using a United
Calibration Corp. SFM 20 testing machine as per ASTM D638 and ASTM D790
standards, respectively. The system controls and data analysis were performed using
Datum software. The notch Izod impact testing was carried out on a Testing Machine
Inc. (TMI) 43-02-01 as per ASTM D256 with the pendulum of impact energy 1.35 J.
5.1.3.2 Dynamic mechanical analysis (DMA).

The storage modulus, loss modulus and tan delta of the test specimens were

evaluated using DMAQ800 supplied by TA Instruments. The specimens were heated

from -30 to 125 0C with a heating rate of 2 0C/min. Single cantilever mode was used to

test the specimens at an oscillating amplitude of 15 um, and a frequency of 1 Hz.
5.1.3.3 Coefficient of linear thermal expansion (CLTE)

CLTE of the test specimens was determined using TMA 2940, equipped with an

expansion probe. Heating rate was 3 OC/min with a constant applied load of 0.15 N.

Temperature range of the reported results is from 30 to 90 OC.

5.1.3.4 Heat deflection temperature (HDT):

DMA Q800, TA instruments was used to find the HDT of various specimens as per
ASTM D648 standards. The rectangular bars, 1.99 x12 X 58 mm in three-point bending
mode with an applying load of 66 psi were used for testing. Samples were heated at the

rate of 2°C/ min from room temperature to the required temperature.

176

 

 

5.1.3.5 Differential scanning calorimetry (DSC)
DSC of specimens was performed using a model Q100 DSC, TA Instruments.

The data was collected using heat and cool method. Specimens were heated up to 190

0C at a ramp rate of 10 OC/min and cooled under the similar ramp rate. The melting

points and heat of crystallization reported are from the heating cycle. The
crystallization temperature and heat of fusion are from the cooling cycle.

Data analyses of DMA, CLTE and DSC results were carried out on software
“Universal Analysis” and three replications of each examination were performed.
5. 1.3.5 Morphological studies

Morphological studies were conducted using a JEOL scanning electron
microscope (model J SM-6400). Specimens were sputter-coated with gold to a thickness

of ~10 11m in order to prevent charging during the examination. The SEM used was

equipped with a lanthanum hexaboride (LaB6) crystal as a source of electrons. An

accelerating voltage of 13 KV was used to collect the SEM images.

177

 

5.2 Results and Discussion
5. 2.1 Tensile Properties

Wood fiber and talc each added either alone or both in certain proportion
enhances the Young’s modulus of PHBV. Figure 5.1, provides the details of the tensile
strength and modulus of PHBV and its composites with various compositions of wood
fiber and talc. Though, there is increase in the Young’s modulus of the PHBV with the
loading of wood fiber at 40 wt% but the strength of the composites reduces, thereby not
giving significant reinforcing effect of wood fibers in the matrices. Where as, the
addition of 40 wt% talc not only increases the Young’s modulus by ~ 260 %, but the
strength as well by ~ 9 %, thus realizing the full potential of talc as reinforcement in
PHBV. Substituting the wood fiber with the talc in subsequent stages lowers the
strength of the hybrid composite a little but ratchet up the modulus by 9 and 5 % at 10
and 20 wt % wood fibers, respectively. Plummeted strength of the composite with the
wood fibers is due to the insufficient interfacial interaction of the PHBV matrix and the
fibers, which is visible in the SEM images of the impact fractured surfaces. Enhanced
strength of the PHBV due to the tale reinforcement provides an edge over the PHBV-
wood fiber composite, which has a lower strength than the hybrid composite. Hybrid
composite has a higher modulus than PHBV by 280% at 20 wt% of talc and wood fiber
each. Similar enhancement of 290% in the modulus of PHBV was observed at 30 wt%
talc and 10 wt% wood fibers loadings. High density of talc enhances the density of the
subsequent hybrid composite, therefore, at 20 wt% of talc and wood fiber, each, in

PHBV gives the optimum weight to strength ratio.

178

 

The density of hybrid composite at 20 wt% talc and wood fiber each, as

measured was 1.39 d: 0.03 grn/cm3, and table 5.2 lists the density of PHBV and its

composites with wood fiber and talc alone. The density was determined by using the
buoyancy method i.e., weighing the specimens, when fully immersed in water, and

compared to the determination form X-ray densometer. The results were comparable.

Table 5.2 : Density of the PHBV and its composites

 

 

 

Hybrid composite
2:21 Wits/tr Tartar
Density (gmlcm3) 1.25:0.01 1.31:0.02 1.54:0.02 1.391033

 

 

 

 

 

 

 

179

 

 

 

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5. 2. 2 Fl exural Properties

Figure 5.2 enumerate the flexural properties of PHBV and it composites in
various proportion of talc and wood fiber. Trends of flexural strength and modulus are
moderately identical to those of tensile. A leap of 270 % in modulus and 22 % in
strength was observed with the 40 wt% tale in PHBV. Modulus of the hybrid
composites plummeted while replacing the talc with wood fiber in each subsequent step
of 10 wt% substitution. There was no difference in strength and modulus between the
final two compositions (10/30/60 wt% talc/ wood fiber/ PHBV and 40/60 wt% wood
fiber/PHBV) of composites. The modulus and strength of the composition at 20 wt% of
talc and 20 wt% wood fiber was higher by 200% and 8%, respectively, from that of
neat PHBV and by 17 % and 8 % from that of 40 wt% wood fiber in PHBV. Meager
difference in the trends of flexural properties to that of tensile was due to different
directional stresses induced to the specimen during the testing. During the tensile
testing the specimen was induced to the uniaxial tensile stress while in the flexural
testing the specimen was induced to uniaxial compression as well as tensile stress

acting simultaneously in the opposite directions.

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182

Theory

The surface energy of material is a crucial property in determining the wetting
of the surface, which, in fact, governs the interface characteristics of the two adhering
surfaces. In composites, the strength of the interface between reinforcing agent and the
matrix plays an important role in deciding the efficiency of stress transferred from
matrix to the reinforcing agent. The strength of the interface in turn governs the overall
strength and modulus of the composite. Complete wetting of the matrix to the fiber or
particulate surface provides an opportunity to have intimate contact and thus adhesion
between the two. Wetting is an interfacial phenomenon that is governed by the surface
energy or tension of the two interacting phases, in this case matrix (PHBV) and fiber
(wood fiber) or inorganic particulate (talc).

Surface energy of the material is defined as sum of the two main components

that are based on the molecular interactions. Hitherto, these were based on polar (y p )
. d . .
and nonpolar/disperse (y ) interaction energy components. More lately, these are
. . LW . . AB . .
referred as Liftshitz-van der Waals (y ) and Lew1s ac1d-base (y ) interactions

refined by Good, Van 055 and Chaudhary (GVOC). 7A8 is further refined in to the acid

6 9
(y ) and base (y ) components and calculated from their geometric mean as. [9-11]

Y AB:2\/ Ye ><\/ 1'9

183

 

“fr

Total/net surface energy (y) is the sum of the yLW and 7A8, mathematically calculated

as «Ewen

The surface energy of PHBV, maple wood fiber and talc are listed in table 5.3. These
values have been taken from the various reference sources [10, 12-19].

Interfacial strength of the reinforcement and matrix can be well correlated with
the interfacial energy/interfacial tension [20]. A generalized expression for an
estimation of the interfacial tension between the two condensed interfaces is given by

Giese and Van Oss [11, 21, 22].

yxy = W_ W): +
217,573+ rfrf- 7:97?- 7,637.?

Here, LW are the Liftshitz—van der Waals interactions ®,® are the polar interactions

and x, y are the interacting phases.

184

 

The interfacial strength can also be quantified from the thermodynamical work
of adhesion (WA) between the two interacting phases. W A was related to surface tension
components of the two interacting phases by Owens and Wendt [8, 11, 23]. WA is taken
as the geometric mean of the polar and the non polar interactions given by the

following equation.

 

 

W4 =2\/}/fW><}/yLW+2\/}/:BX}/y43

Table 5.3: Surface energy parameters of PHBV Maple wood

 

 

)l 7LW ye) 2,6) yAB

(mJ/mz) (ml/m2) (mJ/mz) (mJ/mz) (mJ/mz)

 

PHBV 40.9@ ~32.0@ _ ~8.9@

Maple Wood fiber 47.7-53.3'3 45.5'3 0.02” 57.01” 2.14”

K K K

Talc 140440 " 61 159 11 83.64 "

 

Data taken from references @ [3 K

185

 

Table 5.4, provides the interfacial tension and thermodynamical work of

adhesion between the PHBV-talc and PHBV-wood fiber interfaces calculated using the

above equations. The theoretical values of the interfacial surface energy and W A of

PHBV-talc interface are higher than that of the PHBV-Wood fiber interface. This
suggests positive and better interfacial interactions of PHBV with talc than with the
wood fiber. Higher the interfacial energy or work of adhesion stronger the interface
between the two phases. This was reflected in the improved tensile and flexural
strength of the PHBV with the tale addition, and is verifiable from the SEM
micrographs of the fractured specimens of the hybrid composite. Relatively, higher 7 of
the talc provides an opportunity for the lower 7 PHBV matrix to completely wet its
surface and have intimate interactions of the two surfaces. At higher temperature (melt
temperature) 7 of the polymeric materials further reduces [24] thus enhancing the

wettability of PHBV in the melt state. The interfacial interaction between the PHBV

and talc surface are possible due to the contributions of high acidic surface energy (7 69)

component of talc which descend from the hydroxyl groups present on the lateral faces
of talc lamella and the basic tendency of carbonyl group in the PHBV polyester. For the
purpose of theoretical evaluation, the surface energy values of all the materials have
been adapted from the different sources in literature. Though, there are numerous
variation in the surface energy values of tale in various different available literature

sources, but most of the values are at the higher end, and therefore, we adapted the

. 2 .
theoretical value of 144 mJ/m for our calculations. The marked variation in the y of

186

 

 

talc is also due to its dependence on the type of processing (grinding/milling),

geographical location of resource, size of particle, ratio of exposed faces (basal/lateral,

. AB LW .
lateral faces contribute to the y and basal to y ) and methods of evaluation (

Table 5.4: Calculated Interfacial energy
and Termodynamical Work of Adhesion

Interface type 7x)» WA

 

(mJ/mz) (mJ/mz)

 

PHB V- Wood fiber -20.2 85.0

PHB V- Talc 29.9 142.9

 

experimental as well as theoretical). Another approximation taken for the yxy and WA

calculations was the assumption of equal contribution of 7 ® and 76)

(7 (‘9 =76) )totheyofPHBV.

5. 2.3 Dynamic Mechanical Analysis
Storage Modulus Figure 5.3, illustrates the storage modulus of the PHBV and its
reinforced composites. As observed E', the elastic component of the complex quantity,

increased by ~ 75% at 20 wt.% talc loading, and further increased by 120% with the

additional 20 wt% wood fiber loading, before the glass transition temperature (Tg), and

the increase was 103 and 208%, respectively after the Tg i.e., in the rubbery stage. This

187

 

suggested that the fillers were posing restriction to the free rotation of the molecular

chains above Tg. The relative increment of E' in rubbery stage was higher than that in

the glassy stage for the reason that Poisson’s ratio in glassy polymer is less than 0.5 and
the thermal stress induced due to the mismatch of the coefficient of linear thermal
expansion of filler and matrix reduces the modulus of polymer near the filler
particles[25]. The platelet shape of the talc also played a role in increasing the modulus,
when compared to the increase in the rod like shape of wood fiber filler[25].

The loss Modulus E" accounts for the viscous component of the complex

modulus or the out of phase component with the applied strain, which generally

increases with the addition of filler. 20 wt% of talc in PHBV increase its Tg by ~ 80C

as observed from the shift of the E" peak (Figure 5.4). Appearance of two distinct peaks
in case of hybrid composite at 20 wt% of talc and wood fiber, each loaded in PHBV,
have resulted due to two distinct phases of wood fiber and PHBV. The weak interface
of wood fiber and PHBV which is not observed in case of the interface between talc
and PHBV may have contributed to these.

The tan 5 which is associated with the dampening effect or energy lost as a
heat due interfacial interactions. The theoretical decrease in the Tan 5 of talc filled and
hybrid composite should be the same as percentage of loading in the composite as the
energy loss is mainly associated to the dampening from the viscous component of the
thermoplastic and not from the rigid filler particles. (Figure 5.5) Tan 5 decreases, in

case of a hybrid composite, by ~ 44%. but in case of only talc filled PHBV no

decrement in its value was observed. The increase in the Tg of talc filled composite was

188

 

confirmed by the shift of the Tan 5 peak towards the higher temperature. This specifies
the restrictions imposed by the surface of talc particles to the freedom of rotation of the
chain segment in the amorphous phase of matrix as a result of interfacial interaction
between the talc and PHBV. The additional dampening effect can be due to the
agglomerated talc particle—particle slippages or friction. The interfacial
interactionbetween the PHBV molecules and the talc surface may have lead to higher
energy losses due to the increased polymer particle friction, thus resulting in increased
dampening effect. Lower dampening in case of hybrid particle was due to the additional

stiffer wood fibers than the matrix.lt has been reported that talc in polypropylene

reduces its Tg due to the fact that talc act as nucleating agent which speeds up the

crystallization causing an amorphous phase with higher mobility in PP-talc composites.
This was observed in this case of these composites as well, suggesting that talc did
effect the crystallization of PHBV which speeded up the cold crystallization as
observed from the DSC results discussed next. Carboxylic groups (present at the chain
ends of PHBV) reaction with the hydroxyl group on the lateral faces of the tale cleaved
platelets can not be ruled out and those sites may have acted as additional nucleating
sites for the crystallization [26].The another possibility can be of the hydrogen bonding
between the carbonyl groups (present in the molecular chain of PHBV) and the
hydroxyl groups of talc. Although, this can not be the predominant phenomena because

the ratio of the lateral to basal cleavages is of the low order[8].

189

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5. 2. 4 Differential Scanning Calorimetry
A non isothermal DSC of PHBV and its composites with talc and wood fiber
was carried out. Figure 5.6 illustrate the thermograms of PHBV, PHBV/talc (60/40

wt%) composite and its hybrid composites with 20 wt% talc and wood fiber each.

Further details of the melting temperature (T,,,) , melt crystallization temperature (Tc),

    
 

PHBV/Talc/wood fiber
PHBV/Talc <—— (60/20/20)

(60/40)

1--

4 - A. 2 .

 

 

 

Heat flow (ng
O A N 0) 5 0| 0’ N 00

80 100 120 140 160 180 200

O
N
O
vh .
O
O
0

Temperature (°C)

 

   
 
 

o f I I I I I I l l I I
Q 20 40 60 80 100 120 140 160 180 200

-1 : PHBV/Talc
a 1 (60/40)
§ .2
5 .3 f PHBV/Talc/wood fiber /'
E ‘ (60/20/20)
'5
a “'

PHBV
.5 4

Figure 5.6: DSC curves of PHBV and its composites

193

 

cold crystallization enthalpy (AHC ), enthalpy of melting or crystallization (AHf) and
degree crystallinity (Xe) are provided in the Table 5.5. The melting point of PHBV
decreases a little by 2 0C in presence of talc, but there is no significant change in the

enthalpy of crystallization (AHf) or degree of crystallinity (X6) value.

_ - h“ . .;_gy~.;"n..!
l

 

Table 5.5: Details of DSC ol PHBV and its composites 1 4'
T).)”C)lAH)J/g) 11041414) (Xi)
PHBV 152.6 39.5 103.8 36.5 36.2
PHBVllalc (6040) 149.8 22.0 107.6 20.9 33.6
PHBVllalc/Wood liber(60/20/20) 150.1 21.9 107.5 22.2 33.5

 

 

 

 

 

The biomodal endothermic peaks observed in case of PHBV melting region due

to heterogeneous crystal morphology headed towards the homogenization of the

spherullite growth in the presence of talc, which resulted in the slight drop of Tm. AHC
and AHf of PHBV did not change in presence of either of the filler, thus not giving an

indication of any change in net crystallization. Tc increased to 107 0C from 103 0C in

194

presence of talc, thereby indicating the increase in crystallization rate of PHBV during
the cooling stage. The surface of talc acted as nucleating sites that allowed the PHBV
nucleation growth at a higher temperature than the neat PHBV. The increased

crystallinity rate is another indication of the interfacial interaction between the talc

surface and PHBV molecular chains. Xc was not effected by any significant factor on

the addition of either talc or both wood fiber and talc. XC was calculated

 

AH
XC(%)= f x100
AH X W '

f0

using, AHf was enthalpy of crystallization of

PHBV and AH!” was enthalpy of crystallization at 100 % crystallinity of PHBV

(109 J/g) [27, 28]. W , the weight fraction of PHBV in composite.

195

 

 

 

5. 2.5 Impact Strength

Figure 5.7, illustrates the impact strength of PHBV and its hybrid composite.
The impact strength of PHBV decreased from 25 J/mt to 19 J/mt on addition of 20 wt%
talc and wood fiber each. In the previous chapter, it is described that the impact
strength of PHBV decreases by 29% on addition of 40wt% wood or bamboo fiber. The
decrease in this case was 24%, i.e. 5% lesser than only wood fiber reinforced
composites, suggesting that talc does not lead to any additional drop in the impact
strength of the hybrid composite, but rather it is a common phenomenon due to the
addition of fillers. More over the decrease was due to the reduced interfacial interaction
between the wood fiber and PHBV in presence of talc, which was visible in SEM
micrographs.
Due to the low fiber-matrix interfacial interaction, fibers pull out used less energy that

resulted in lower impact strength.

30.

20j

 

H

10

 

 

 

 

lm pact Strength (Jlmt)

 

PHBV/Wood fibe rltalc(60l20l20)
(wt %)

 

Figure 5.7: Impact strength of PHBV and its hybrid composite.

196

 

 

5. 2. 6 Morphology

Figure 5.8 and 5.9, shows the SEM photomicrographs of the impact fractured
surface of hybrid composite Figure 5.8(a) and (b) are low and high magnification
micrographs, respectively, showing the uniform distribution of talc particles in PHBV
matrix. From the SEM images, it was observed that the major failure occurred due the
fiber pull out from the matrix. The interfacial interactions between the wood fiber and
matrix were not strong enough to resist the fiber pull out during the impact as a result
the energy consumed was small in magnitude. A poor interface between the wood fiber
and matrix is visible in the figure 5.8(b) at larger magnification. A poor interface
between wood fiber and PHBV was also due to the presence of talc particles within the
two surfaces that did not allow much of the interaction among the two and made the
fiber slippage easy during the pull out. This is visible in the image in figure 5.9(a) that
was obtained from the interface, where the fiber was pulled out. Figure 5.9(b) is a high
magnification image to reveal the talc and PHBV interface. It shows a good dispersion

and embedment of talc in PHBV matrix, suggesting a good interface between the two.

197

 

   

 

fiber .11.; * .. 4 :1:

 

:«e » ”u’w 4,904.: ‘1’ f 1' 4...; ,. loum

Figure 5.8: SEM Photomicrographs of fractured hybrid composite
surface (a) Low magnification 270x, 100um (b) High magnification
3000X, l0um.

198

 

Figure 5.9: SEM Photomicrographs of Hybrid composite (a) Wood fiber
pulled out interface with PHBV matrix, 2700><, lOum (b) High
magnification fractured surface of composite, 5500 X, Sum.

199

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5. 2. 7 Thermomechanical Properties
Figures 5.10, illustrate the CLTE and HDT of the neat PHBV in comparison to

the hybrid biocomposite, and reflected the beneficial improvement in both the

properties of the hybrid biocomposite. CLTE was reduced by 36% and HDT by 90C.

The reduction in CLTE was more due to the effect of talc than wood fiber, as reported
previously [29], the wood fiber loading in PHBV affected the CLTE of PHBV only a
little. Bhardwaj et al.[30] reported the reduction in CLTE of PHBV, approximately in
equal proportion, with the addition of talc. The decrement on CLTE of hybrid
biocomposite was due to the rigid and low CLTE filler in the polymeric matrix, which
was more pronounced due to the presence of tale. The HDT of the hybrid biocomposite
was more affected by the wood fiber than the tale. The HDT reduced, when the wood
fiber was replaced by the 20 wt% of talc. The increase in HDT was due to the increase

of stiffness at higher temperature of the fiber filled polymers.

200

 

 

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201

5.3 Conclusion

The hybrid composites from wood fiber, talc and a bioplastic PHBV using the
extrusion-injection molding technique were successfiilly fabricated. The composition at
20 wt% of wood fiber and talc in PHBV gave a leap of ~200% improvement in the
Young’s and flexural modulus. The dynamic-mechanical analysis supports the tensile
and flexural results as well as explains the talc-PHBV and wood fiber-PHBV interfacial
interactions. The high surface energy of talc particulates due to its high acidic character
provides an opportunity for talc—PHBV interfacial interactions. The DSC analysis does
not depict any additional crystallization of PHBV with the addition of talc in it. The
impact strength of the PHBV reduced with addition of fillers. The Morphological
analysis of the hybrid composite using SEM reveals the better interfacial interaction

and dispersion of talc with PHBV than wood fiber.

202

 

5.4 References

l.

10.

11.

Unal H., Mimaroglu A., and Alkan M., "Mechanical properties and morphology
of nylon-6 hybrid composites". Polym Int., 2004.53: p. 56-60.

Huneault M. A., Godfroy P.G., and Lafleur P.G., "Performance of
TalcEthylene-Octene Copolymer/Polypropylene Blends". Polymer Engineering
and Science, JUNE 1999.39(6): p. 1130-1138.

Ersoy O. G. and Nugay N., "Effect of Inorganic Filler Phase on Mechanical and
Morphological Properties of Binary Immiscible Polymer Blends". Polymer
Bulletin Springer- Verlag 2003(49): p. 465-472.

Browning R., Lim G. T., MoyseA., Sun L., and Hung-Jue Sue, "Effects of Slip
Agent and Talc Surface-Treatment on the Scratch Behavior of Thermoplastic
Olefins". Polymer Engineering and Science @2006 Society of Plastics
Engineers, 2006.46: p. 601-608.

Di'ez-Gutie'rreza S., Rodn'guez-Pe'reza M.A., De Sajaa J.A., and Velascob
J .I., "Dynamic mechanical analysis of injection-moulded discs of polypropylene
and untreated and silane-treated talc-filled polypropylene composites". Polymer
1999.40: p. 5345-5353.

McInerney L. F., Kao N., and Bhaitacharya N., "Melt Strength and Extensibility
of Talc-Filled Polypropylene". Polymer Engineering and Science, 2003.43(12):
p. 1821-8129.

H. Unal, "Morphology and mechanical properties of composites based on
polyamide 6 and mineral additives". Materials and Design 2004.25: p. 483-487.

Kaggwa G. B., Huynh L., Ralston J., and Bremmell K., "The Influence of
Polymer Structure and Morphology on Talc Wettability". Langmuir 2006.22: p.
3221-3227.

Leeden M.C.V. and Gert Frens, "Surface Properties of Plastic Materials in

Relation to their Adhering Performance (Review)". Advance Engineering
materials, 2002.4(5): p. 280-289.

Meijer M. de, Haemers S., Cobben W., and Militz I-l., "Surface Energy
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9

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Chapter 6

Conclusion

The objective, in entirety, of this research was to shift W/NPCs from
conventional olefin plastic to a fully green biocomposites from renewable source based
and biodegradable bioplastic, i.e., PHBV. During the course of changeover from
partially to fully green biocomposites, the issues pertaining to biopolymers are
addressed. This thesis also discusses the other emerging issues, during the processing,
of biopolymer based W/NPCs like cost, processablity, performance in comparison to
the conventional WPCs, compatibility of reinforcing agent and matrix, availability, and
anisotropy. The development of all green hybrid biocomposite from PHBV, wood fiber
and talc was attained in four sequential phases as follows.

In the first phase, biodegradable biocomposite from PHBV and maple wood
fiber were fabricated using extrusion followed by injection molding. The loading level
of wood fiber in biocomposites was varied from 10 to 40 wt%. It was noted that the
tensile and flexural modulus at 40 wt% of wood fiber were enhanced by ~167%, when
compared with neat PHBV. The theoretical Young’s modulus values of biocomposites
obtained from Halpin-Tsai and Tsai-Pagano equations were consistent with the
experimental values. The effectiveness of fibers on the storage modulus of the
composite was evaluated by the coefficient factor. C. The heat deflection temperature
was increased by 21%, and the coefficient of linear thermal expansion reduced by 18 %
at 40 wt% of wood fiber. The toughness of the biocomposites decreased with the fiber
loading. SEM photomicrographs exhibited the good dispersion at 30 wt% wood fiber

but fiber aggregation at 40 wt%.

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In the second phase, the biocomposites from a bamboo fiber and PHBV were
fabricated with the similar process, i.e., extrusion followed by injection molding, as
PHBV-Wood fiber biocomposite in stage one. The need for the bamboo fiber in
substitution to wood fiber was emphasized. Comparative analysis of the two different
kind of lignocellulosic fiber (wood and bamboo) at 30 and 40 wt.% fiber loading were
carried out with respect to their mechanical, therrnomechanical, thermal, FTIR and
morphological aspects. Both the fibers, embedded in the PHBV matrix, gave similar
tensile and flexural modulus, though they had different aspect ratios. Statistically, no
appreciable difference among the mechanical properties of two biocomposites, obtained
from two different types of lignocellulosic fibers, was noted. A difference between the
net crystallinity of the biocomposites was observed in DSC and FTIR. Notch impact
strength of PHBV decreased with the fiber addition and the reduction was greater in
case of bamboo fiber composites. The effect of the fiber type on the thermal stability of
composites was not observed, but incorporation of wood fiber gave a higher heat
deflection temperature (HDT) versus the bamboo fiber. Scanning electron microscopy
(SEM) photomicrographs demonstrated the existence of similar kinds of interfacial
interaction between the two types of fibers and PHBV

In the third phase, the research was focused on the processing of the PHBV-
wood fiber biocomposite. These were fabricated using two different molding processes
i.e., injection molding and compression molding, and characterized for the static and
dynamic mechanical properties, and isotropic characteristics. Injection molded
composite performed better than compression molding, which was apparent from the

better tensile and flexural properties of injection molded composite that gave an

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improvement in Young’s modulus and flexural modulus by 104%. High elastic
modulus and low dampening effect apparently suggested the same during dynamic
mechanical analysis. Morphological observations revealed poor adhesion and air
entrapment in compression molded specimens. Study of anisotropy was performed
using squeeze flow test of both the processes, and it indicated the anisotropic behavior
of injection molded specimens. Anisotropy varied;(a) along the sections of injection
molded specimens, and (b) composition of wood fiber in the biocomposites.

In the final stage “all green hybrid bioplastic composites” with the dual
reinforcements of talc and wood fiber were designed, fabricated and characterized. The
composition at 20 wt% of wood fiber and talc, each in PHBV, gave a leap of ~200%
improvement in the Young’s and flexural modulus. The analysis of the dynamic-
mechanical is in unison with the tensile and flexural analysis as well as explains the
talc-PHBV and wood fiber-PHBV interfacial interactions. The high surface energy of
talc particulates due to their high acidic nature provides better talc-PHBV interfacial
interactions than wood fiber-PHBV interface. The calorimetric analysis was carried out,
and no additional crystallization of PHBV in the presence of talc was observed. The
impact strength of the PHBV plummeted with addition of fillers. The Morphological
analysis reveals the better interfacial interaction and dispersion of talc, in PHBV, than
wood fiber.

The future scope of this research can be protracted with the realization of the
reinforcing effect of nano sized talc in the PHBV matrix. The nano size talc may reduce
the amount of talc with a substantial increase in mechanical properties of the matrix,

thus reducing the specific density of the biocomposite and providing the opportunity to

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load the matrix with higher fiber content. This would also reduce cost of the
biocomposites. The toughness improvement of the biocomposite still remains a
challenge that lays in the future scope of research. Toughness of these biocomposite
can be attained by modifying the surface functionality on surface of the wood fiber.
This would lead to either a chemically bonded interface between the wood fiber and
PHBV or high surface energy wood fiber that allows the complete wetting of the matrix
on the fiber surface. The second approach to enhance upon the toughness is to enhance

the toughness of the matrix by blending or nanostructured reinforcement.

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