INVESTIGATIONS ON THE FRACTURE BEHAVIOR AND NOTCHSENSITIVITY OF INJECTION MOLDINGS

By

BERNARD CHUKWUEMEKA OGAZI-ONYEMAECHI(ID: 07851502)

A DISSERTATION

Submitted in Fulfillment of the Requirements for the Degree of

DOCTOR OF ENGINEERINGAT

KYOTO INSTITUTE OF TECHNOLOGY
Kyoto, Japan, 606-8585

SEPTEMBER 2010

© (Ogazi-Onyemaechi Bernard Chukwuemeka, September 2010)

ACKNOWLEDGEMENTS
The task of accomplishing three years of PhD research work has been an uphill one. This isagainst the background that the path trailed in this great academic adventure was a darktunnel with the expectation of a great light at the end. It is said that behind every cloud, there
is a silver lining. Professor U. S. Ishiaku once reminded me that “where there is a will, there
is always a way”. Realizing this golden and seasoned saying, I willed it within me that I mustovercome all challenges in the pursuit of this academic goal. Today, it is a reality. The
requirements for the degree of Doctor of Engineering have been accomplished.
However, the will and consequent accomplishment would have been a mirage if
people builders were not available. To GOD is the glory that Professor Hiroyuki Hamada has
been available. Therefore, I acknowledge the great opportunity he has given to me toaccomplish this task under his guidance. I acknowledge also the academic contributions of
Associate Professor A. N. Nakai and Associate Professor M. Kotaki. Profoundly, I express
my thanks to Asst. Professor Leong Yew Wei for his effective and critical support in myresearch work, especially in the preparation of this dissertation. I also thank Dr. Kazushi
Yamada for his technical support.
In a special way, I acknowledge and thank the administration of KIT for extending
tuition waiver/reduction to me, and for the 2007 KIT Foundation Award, which were the only
sources for the payment of my tuition fees and running costs for the many research works
conducted outside Kyoto. I thank the staff members of the International Planning Division for
their kindness and friendship.
I owe much gratitude to Dr. Kiyotaka Tomari and Dr. Koji Yamada from Osaka
Municipal Research Institute not only for making the Institute’s laboratory equipments
available for my use but also for been guides and trainers.
I express my heartfelt thanks and love to my wife, Dr. (Mrs.) Chinyere G. Ogazi, andmy children, Divine E. I. Onyemaechi, Stephanie E. C. Onyemaechi, and Chelsea U. C.
Onyemaechi for their love, and for always been a source of inspiration during this period ofstudentship. My brothers and sisters are all well appreciated for been available for me.
To the evergreen memory of my parents, Mr. Longinus Sunday Onyemaechi and Mrs.
Elizabeth Alunma Onyemaechi, I say peace of GOD.
In all situations, GOD is SUPREME. I thank HIM for HIS WILL in my life. To HIM
be all the GLORY, HONOR and ADORATION, AMEN.

iDEDICATION
To the Glory and Majesty of GOD, I dedicate this Dissertation to my children, Divine
Ebubechukwu Ihechidere Onyemaechi, Stephanie Ezichidindu Chinelo Onyemaechi, Chelsea
Uchechi Chidiogo Onyemaechi, and other children. May this be one of the sources ofinspiration for them to achieve full accomplishments in their chosen careers.

“There is nothing constant in the universe, all ebb and flow; and every shape that is born
bears in its womb the seed of change.”*YOU CAN BE WHAT YOU CHOOSE TO BE.*ii

TABLE OF CONTENTS

Acknowledgementi

Dedicationii

Table of Contentsiii

List of Tables

xii

List of Figures

xiii

Abstract

xxvi

Chapter 1: General Introduction

1

1.1 Background of Studies on Fracture Behavior and Fracture Toughness of Materials

1

1.2 Failure of Structures and the Need for Fracture Mechanics

2

1.3 Characteristics and Recycling of PET

3

1.4 Fracture Behavior of PET

5

1.5 Effect of Annealing on the Notch Sensitivity of PET

7

1.6 Crack Tip Temperature and Essential Work of Fracture7

1.7 The Need for Investigation on the Effect of V-notches and Hairline Cracks

9

1.8 Composite Materials

10

1.9 Injection Molded Composites

11

1.10 Thesis Structure

12

References

13

Chapter 2: Review of Related Literatures17

2.1 Neat (V-PET) and Recycled (R-PET) Poly(ethylene terephthalate) Injection Moldings

17

2.2 Effect of Notches on the Fracture Behavior of PET

18

iii

2.3 Effect of Hairline Cracks and Annealing on the Fracture Behavior of PET Injection
Moldings

21

2.4 Essential Work of Fracture Analysis

24

2.5 Crack Tip Temperature27

2.6 Mechanical Properties and Fracture Toughness of Glass Fiber Reinforced
Polypropylene Composites (GFPP)

30

2.6.1 Composites

30

2.6.2 Calculation of Fracture Toughness under Linear Elastic Fracture Mechanics

30(LEFM)
2.6.3 Calculation of Interfacial Shear Strength of GFPP Composites

33

2.6.4 Mechanical and Fracture Properties of GFPP Composites

34
44

References

Chapter 3: Fracture Modes and Notch Sensitivity of V-notched V-PET and R-PET

54

Injection Moldings
Abstract

54

3.1 Introduction

55

3.2 Characterization methods

58

3.2.1 Material and Sample Preparation

58

3.2.2 Notching and Mechanical Property Characterization

58

3.2.3 Determination of Notch Sensitivity Factors

59

3.2.4 Measurement of Crystallinity in PET

60

3.2.5 Birefringence Observation of PET Injection Moldings

61iv

3.3 Results and Discussions

62

3.3.1 Cross-Sectional Observation of V-PET by Polarized Light Microscopy

62

3.3.2 Effect of V-notch depth on the fracture modes of V-PET

63

3.3.3 Correlation between Notch Sensitivity Factor for Strength (kS) and Skin-Core

Morphology of V-PET

64

3.3.4 Correlation between Notch Sensitivity Factor for Energy (kT) and Skin-Core

Morphology of V-PET

65

3.3.5 The role of the skin-core morphology on the notch sensitivity of V-PET atintermediate loading points during tensile test

66

3.3.6 SEM characterization of tensile fracture surfaces of V-PET

67

3.3.7 Effect of V-notches on the Izod Impact Fracture Behavior of V-PET

68

3.3.8 SEM Characterization of Izod Impact Fracture Surface of V-PET

69

3.3.9 Summary and Observations on the Crack Propagation Behavior and Toughnessof V-notched V-PET Injection Moldings

70

3.3.10 Cross-sectional observation of R-PET by polarized light microscopy

70

3.3.11 Effect of V-notches on the fracture modes of R-PET

71

3.3.12 Correlation between notch sensitivity factors for strength (kS) and the skin-coremorphology of R-PET

73

3.3.13 Correlation between Notch Sensitivity Factors for Energy (kT) and the Skin-

Core Morphology of R-PET

73

3.3.14 The role of Skin-Core Morphology on the Fracture Modes of R-PET at

Intermediate Loading Points during Tensile Test
3.3.15 SEM characterization of tensile fracture surfaces of R-PET

v

74
76

3.3.16 Effect of V-notches on the Izod Impact Fracture Behavior of R-PET

77

3.3.17 SEM Characterization Izod Impact Fracture Surfaces of R-PET

78

3.3.18 Summary and Observations on the Crack Propagation Behavior and Toughnessof V-notched R-PET Injection Moldings

79

3.4 Conclusion on the Crack Propagation Behavior and Toughness of V-notched Injection
Molded V-PET and R-PET Components

79

References

80

Chapter 4: The Fracture Modes and Notch Sensitivity of Hairline Cracked V-PET

119

and R-PET Injection-Moldings
Abstract

119

4.1 Introduction

120

4.2 Characterization Methods

121

4.2.1 Material and Sample Preparation

121

4.2.2 The Annealing of Dumbbell Specimens

121

4.2.3 Measurement of Crystallinity in PET

122

4.2.4 Laser-Raman Spectroscopy of the Materials

122

4.2.5 Birefringence Observation of PET Injection-Moldings

123

4.2.6 Hairline Cracking by Saicas DN-20S Micro Cutting Machine

124

4.2.7 Tensile and notch sensitivity characterization

124

4.2.8 Fracture Modes and Notch Sensitivity Characterization

125

4.3 Results and Discussions

126

4.3.1 Evaluation of the Crystallinity in V-Pet Materials before and after Annealing byvi

the use of DSC and Raman Spectroscopy
4.3.2 The Morphology of V-PET before and after Annealing

126
127

4.3.3 Effect of Hairline Cracks on the Fracture Modes of V-PET before and after

Annealing128

4.3.4 Correlation between Notch Sensitivity and Skin-Core Morphology of V-PET

before and after Annealing

129

4.3.5 SEM Characterization of Tensile Fracture Surfaces of V-PET

131

4.3.6 Evaluation of the Crystallinity in R-PET Materials before and after Annealing by

the use of DSC and Raman Spectroscopy
4.3.7 The Morphology of R-PET before and after Annealing

132
134

4.3.8 Effect of Hairline Cracks on the Fracture Modes of R-PET before and after

Annealing135

4.3.9 Correlation between Notch Sensitivity and Skin-Core Morphology of R-PET

before and after Annealing

136

4.3.10 SEM Characterization of Tensile Fracture Surfaces of V-PET

138

4.4 Conclusion

138

References

139

Chapter 5: Essential Work of Fracture Evaluation of the Fracture Modes of V-

177

Notched PET Components
Abstract

177

5.1 Introduction

177

5.2 Characterization Methods

180

vii

5.2.1 Material and Sample Preparation

180

5.2.2 Measurement of Crystallinity in PET

180

5.2.3 Birefringence Observation of Skin-Core Morphology of Injection- Molded PET

1805.2.4 Notching and mechanical property characterization

181

5.2.5 Monitoring of Temperature Distribution in Specimens during Tensile Test

181

5.2.6 Notch Sensitivity and Temperature Distribution Characterization

182

5.2.7 Determination of Essential Work of Fracture (EWF)

183

5.3 Results and Discussion

184

5.3.1 Essential Work of Fracture Evaluation and the Analysis of Crack Tip

Temperature Distribution of the Fracture Modes of V-Notched V-PET
Components

184

5.3.1.1 Measurement of Crystallinity in V-PET

184

5.3.1.2 Cross-sectional Observation by Polarized Light Microscopy of V-PET

185

5.3.1.3 Mechanical Properties of V-PET at Different Notch Depths

185

5.3.1.4 Distribution of Temperature at the Crack-tip During Deformation in V-

PET

186

5.3.1.5 The Effect of Notch Sensitivity and Fracture Modes on the Change in

Temperature and the Shape of Plastic Deformation Zones in V-PET

189

5.3.1.6 Effects of V-Notches on the Fracture Properties of V-PET and

Determination of Essential Work of Fracture (EWF)

190

5.3.1.7 Relationship between Change in Crack Tip Temperature and Ratio of

Energy Dissipated in the Process (We) and Plastic (Wp) Deformation
Zones

193

viii

5.3.1.8 Correlating the Notch Sensitivity Factors for Strength (kS) and Energy (kT)with the Ratios of Work done in both Deformation Zones of EWF

195

5.3.2 Essential Work of Fracture Evaluation and the Analysis of Crack Tip

Temperature Distribution of the Fracture Modes of V-Notched R-PET
Components

196

5.3.2.1 Measurement of Crystallinity in R-PET

196

5.3.2.2 Cross-sectional observation by polarized light microscopy of R-PET

196

5.3.2.3 Mechanical Properties of R-PET at Different Notch Depths

197

5.3.2.4 Distribution of Temperature at the Crack-Tip during Deformation in R-

PET

197

5.3.2.5 The Effect of Notch Sensitivity and Fracture Modes on the Change in

Temperature and the Shape of Plastic Deformation Zones in R-PET

199

5.3.2.6 Effects of V-Notches on the Fracture Properties of R-PET and

Characterization under Essential Work of Fracture (EWF)

200

5.3.2.7 Relationship between Change in Crack Tip Temperature and Ratio of

Energy Dissipated in the Process (We) and Plastic (Wp) Deformation
Zones in R-PET

202

5.3.2.8 Correlating the Notch Sensitivity Factors for Strength (k S) and Energy (kT)with the ratios of Work done in both Deformation Zones of EWF

203

5.4 Conclusions

205

References

206

Chapter 6: Investigations on the Fracture Toughness of GFPP Composites

239

ix

6.1 Introduction

239

6.2 Characterization Methods

240

6.2.1 Material and sample preparation

240

6.2.2 Tensile Tests and Determination of Mechanical Properties and Fracture

Toughness

240

6.2.3 Scanning Electron Microscopy (SEM)

242

6.2.4 Charpy and Izod Impact Tests

242

6.3 Results and Discussions242

6.3.1 Mechanical properties and fracture toughness of GFPP from tensile data242

6.3.2 Scanning Electron Microscopy of GFPP

244

6.3.3 Impact Properties of Un-notched and Notched GFPP Samples

244

6.4 Summary and Conclusions

245

References

245

Chapter 7: Characterization of the Fracture Toughness of GFPP Composites under264

Elastic Plastic Fracture Mechanics
Abstract

264

7.1 Introduction

265

7.1.1 Crack Tip Opening Displacement (CTOD)

266

7.1.2 J-Contour Integral

267

7.1.2.1 Nonlinear Energy Release Rate

268

7.1.2.2 J as a Stress Intensity Parameter

269

7.2 Characterization Methods

270x7.2.1 Material and Sample Preparation

270

7.2.2 Tensile Tests of GFPP Composites

271

7.2.3 CCD Camera Monitored Tensile Test and Further Characterization of the

Fracture Toughness of GFPP under EPFM

271

7.2.4 Charpy and Izod Impact Tests272

7.2.5 Scanning Electron Microscopy (SEM)

272

7.2.6 Determination of the Fracture Toughness of GFPP Composites272

7.2.7 Determination of Interfacial Shear Strength of GFPP Composites

273

7.2.8 Correlation of the Mechanical Properties/Impact Properties, Molecular Weightand Fracture Toughness Properties274

7.3 Results and Discussion

275

7.3.1 Tensile and Fracture Toughness

275

7.3.2 Micrographs and CCD Camera Images of GFPP Materials

277

7.3.3 Interfacial Shear Strength of GFPP Composites

279

7.3.4 Relationship between Fracture Toughness, PP Molecular Weight,

Mechanical/Impact Properties and Interfacial Shear Strength

280

7.3.5 Strength of Association between Fracture Toughness and other Properties of

GFPP

281

7.4 Summary and Conclusion

282

References

283

Chapter 8: General Conclusion

314

xi

LIST OF TABLES

Table 3.1 The Crystallinity of V-PET and R-PET Injection Moldings

82

Table 3.2 Gray Values from Gray Image Plots of the Cross-sectional View of Dumbbell
Samples Showing the Sizes of the Skin, Interface and Core Layers of V-PET
and R-PET

83

Table 4.1 Crystallinity of V-PET Injection Moldings before and after Annealing

141

Table 4.2 Crystallinity of R-PET Injection Moldings before and after Annealing

142

Table 4.3 Mechanical Properties of V-PET and R-PET Samples before and after
Annealing

143

Table 5.1 The Crystallinity of V-PET and R-PET Injection Moldings

208

Table 5.2 Gray Values from Gray Image Plots of the Cross-sectional View of Dumbbell
Samples Showing the Sizes of the Skin, Interface and Core Layers of V-PET
and R-PET

209

Table 5.3 Effects of Different Depths of V-notches on the Yield Stress (MPa) and Strain(%) of V-PET and R-PET

210

Table 6.1 Properties of Polypropylene Compounded with Glass Fiber

248

Table 7.1 Properties of Polypropylene Compounded with Glass Fiber

287

xii

LIST OF FIGURES

Figure 1.1

V-notch showing the radius (1030m) and angle of curvature (45°C) as
fitted with circle at tangent with the base of the notch

Figure 1.2

Hairline-crack showing the radius (90m) and angle of curvature (about
8°C) as fitted with circle at tangent with the base of the notch

Figure 1.3

14

15

Apparatus setup showing the annealing of specimens in hot edible oil andsubsequent quenching in ice-chilled oil (7 - 10°C)

16

Figure 2.1

Mode I Fracture48

Figure 2.2

Mode II Fracture

49

Figure 2.3

Mode III Fracture

50

Figure 2.4

Geometry conditions for fracture toughness

51

Figure 2.5

Dumbbell specimen for tensile test

52

Figure 2.6

Determination of σQ at 5° off tangent

53

Figure 3.1

Arrows indicate the position where samples were extracted for DSC
experiments

84

Figure 3.2

Preparation and polishing of Samples for birefringence observations

85

Figure 3.3

Birefringence observation of the cross-section of dumbbell Specimens
showing the skin and core structures of V-PET

Figure 3.4

Image contour plot showing the skin, interface and core regions of V-PET
indicated with arrows

Figure 3.5

8687

Gray image plot showing the gray values of the skin, Interface and coreregions of V-PET [W = graph of gray values on the width direction, X =xiii

graph of gray values on the thickness direction]
Figure 3.6

Stress - Strain curves of V-PET showing the effects of V-notch of different
depths [{ }* notch depths of 2.0 - 4.5mm]

Figure 3.7

8889

Crack initiation at notch root showing: a, the development of two lines at anangle of 45⁰ prior to shear banding; b, Shear banding and plastic flow due
to stress and consequent semi ductile failure [arrow shows micro-cracks andeventual coalescence]

Figure 3.8

90Notch sensitivity factors for strength kS for V-PET at the various notch
depthsFigure 3.9

91

Notch sensitivity factors for energy (kT) for V-PET at the various notch
depthsFigure 3.10

92

Load-displacement curve of 0.6mm deep notch sample showing threeintermediate loading points [1, 2 and 3 = 1.2, 1.4 and 1.6kN respectively]

Figure 3.11

93

Birefringence of V-PET showing the formation and extension of two
families of intersecting slip lines when a, 1.2kN; b, 1.4kN; and c, 1.6kN
load is applied as shown by points 1, 2 and 3 on the Load/Displacement
curve in Figure 3.10

Figure 3.12

94

Load-displacement curve of 1.5mm deep notch sample showing 1.2kN
loading at point 1

Figure 3.13

95

Birefringence of V-PET showing the formation of two straight crack lines
prior to brittle fracture when [1= 1.2kN] load is applied on a 1.5mm deep
notched sample as shown on the Load/Displacement curve in Figure 3.12

Figure 3.14:

Tensile fracture surface of V-PET showing ductile failure with high plastic

xiv

96

deformation at 0.6mm deep V-notch [white arrow indicates notch direction;
vertical line is the end of notch]
Figure 3.15

97

Tensile fracture surface of V-PET showing semi ductile failure with crazing
at 1.5mm deep V-notch [white arrow indicates notch direction; vertical lineis the end of notch]

Figure 3.16

98

Tensile fracture surface of V-PET brittle failure at 2.0mm deep V-notch[white arrow indicates notch direction; vertical line is the end of notch]

Figure 3.17:

Izod impact strength of V-PET showing gradual loss of strength at different
V-notch depths

Figure 3.18

99

100

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of V-PET at 0.6mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 3.19

101

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of V-PET at 1.5mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 3.20

102

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of V-PET at 2.0mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 3.21

Birefringence observation of the cross-section of dumbbell Specimens
showing the skin and core structures of R-PET

Figure 3.22

104

Image contour plot showing the skin, interface and core regions of R-PET
indicated with arrows

Figure 3.23

103

105

Gray image plot showing the gray values of the skin, interface and corexv

regions of R-PET [ W = graph of gray values on the width direction, X =graph of gray values on the thickness direction]
Figure 3.24

Stress - Strain curves of R-PET showing the effects of V-notch of different
depths [{ }* notch depths of 2.0-4.5mm].

Figure 3.25

108

Notch sensitivity factors for energy kT for the materials at the various notch
depthsFigure 3.27

107

Notch sensitivity factors for strength (kS) for V-PET and R-PET at thevarious notch depths

Figure 3.26

106

109

Birefringence of R-PET showing the formation and extension of two
families of intersecting slip lines when a, 1.2kN; b, 1.4kN and c, 1.6kN load
is applied as shown by points 1, 2 and 3 on the Load/Displacement curve in
Figure 3.10

Figure 3.28

110

Birefringence of R-PET showing the formation of two straight crack lines
when [1= 1.2kN] load is applied on 1.5mm deep notched sample as shownon the Load/Displacement curve in Figure 3.12

Figure 3.29

111

Tensile fracture surface of R-PET showing semi ductile failure with crazing
at 1.5mm deep V-notch [white arrow indicates notch direction; vertical lineis the end of notch]

Figure 3.30

112

Tensile fracture surface of R-PET showing semi ductile failure with crazing
at 1.5mm deep V-notch [white arrow indicates notch direction; vertical lineis the end of notch]

Figure 3.31

113

Tensile fracture surface of R-PET showing brittle failure at 2.0mm deep Vnotch [white arrow indicates notch direction; vertical line is the end ofxvinotch]Figure 3.32

114

Izod impact strength of V-PET and R-PET showing gradual loss of strength
at different V-notch depths

Figure 3.33

115

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of R-PET at 0.6mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 3.34

116

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of R-PET at 1.5mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 3.35

117

SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod
impact specimen of R-PET at 2.0mm deep V-notch [arrows indicate notch
direction; vertical lines are the end of notches]

Figure 4.1

118

Annealing apparatus showing dumbbell sample held in-between two
aluminum bars, oil-filled beaker held inside oil-filled electric pot, and a bath
containing Ice-chilled oil (7 - 10°C)

Figure 4.2

144

DSC sample slices at 50m thick each showing samples representing thedifferent regions of V-PET a = before annealing, b = after annealing [X1 =Pseudo-skin, X = skin, Y = interface, and Z = core]

Figure 4.3

145

Epoxy embedding and smooth-polishing of un-annealed and annealed VPET and R-PET samples for Raman spectroscopy scanning in the flow
direction [X1 = Pseudo-skin, X = Skin, Y = Interface, and Z = Core]

146

Figure 4.4

Apparatus for cutting hairline cracks on the dumbbell specimens

147

Figure 4.5

FWHM Raman intensity showing crystallinity along the skin-core region ofxvii

V-PET before (UA) and after (A) annealing [X1 = pseudo-skin, X = skin, Y= interface, and Z = core]
Figure 4.6

148

Crystallinity along the skin-core of V-PET before and after annealing [X1 =pseudo-skin, X = skin, Y = interface, and Z = core]

Figure 4.7

Birefringence observation of the cross-section of V-PET dumbbell samples
before annealing showing their skin-core structure

Figure 4.8

149

150

Gray image plots of V-PET before annealing showing the gray values of theskin, interface and core regions [W & X= graph of gray values in the width
and thickness directions, respectively].

Figure 4.9

Image contour plots of V-PET before annealing showing the skin, interface
and core regions

Figure 4.10

151

152

Birefringence of the cross-section of V-PET dumbbell samples after
annealing Showing the pseudo-skin, skin and core structure [pseudo-skin isindicated by the arrows]

Figure 4.11

153

Gray image plots of V-PET after annealing showing the gray values of thepseudo-skin, skin, interface and core regions [W & X= graph of gray values
in the width and thickness directions, respectively]

Figure 4.12

154

Image contour plots of V-PET after annealing showing the pseudo-skin,skin, interface and core regions

155

Figure 4.13

Crack initiation and brittle fracture of V-PET at 10m hairline crack depth

156

Figure 4.14

Crack sensitivity factors for strength (kS) of V-PET before and after
annealing

Figure 4.15

157

Crack sensitivity factors for energy (kT) of V-PET before and after annealing 158

xviii

Figure 4.16

Fracture surface of V-PET showing semi-ductile failure at 5m hairline
crack depth before annealing

Figure 4.17

159

Fracture surface of V-PET showing brittle failure at 10m hairline crack
depth before annealing

Figure 4.18

160

Fracture surface of V-PET showing semi-ductile failure at 10m hairline
crack depth after annealing

Figure 4.19

161

DSC sample slices at 50m thick each showing samples representing thedifferent regions of R-PET a = before annealing, b = after annealing [X1 =pseudo-skin, X = skin, Y = interface, and Z = core]

Figure 4.20

FWHM Raman intensity showing crystallinity along the skin-core of R-PET
before and after annealing

Figure 4.21

163

Crystallinity along the skin-core of R-PET before and after annealing [X1 =pseudo-skin, X = skin, Y = interface, and Z = core]

Figure 4.22

164

Birefringence observation of the cross-section of R-PET dumbbell samples
before annealing showing their skin-core structure

Figure 4.23

162

165

Gray image plots of R-PET before annealing showing the gray values of theskin, interface and core regions [W & X= graph of gray values in the width
and thickness directions, respectively]

Figure 4.24

Image contour plots of R-PET before annealing showing the skin, interface
and core regions

Figure 4.25

166

167

Birefringence of the cross-section of R-PET dumbbell samples after
annealing showing the pseudo-skin, skin and core structure [pseudo-skin isindicated by the arrows]

168

xix

Figure 4.26

Gray image plots of R-PET after annealing showing the gray values of thepseudo-skin, skin, interface and core regions [W & X= graph of gray values
in the width and thickness directions, respectively]

Figure 4.27

Image contour plots of R-PET after annealing showing no identifiableboundaries between the pseudo-skin, skin, interface and core regions

Figure 4.28

171

Crack sensitivity factors for strength (kS) of R-PET before and after
annealing

Figure 4.30

172

Crack sensitivity factors for energy (kT) of R-PET before and after
annealing

Figure 4.31

173

Fracture surface of R-PET showing semi-ductile failure at 5m hairline
crack depth before annealing

Figure 4.32

174

Fracture surface of R-PET showing semi-ductile failure at 10m hairline
crack depth before annealing

Figure 4.33

175

Fracture surface of R-PET showing semi-ductile failure at 10mm hairline
crack depth after annealing

Figure 5.1

170

Shear-banding and semi-ductile fracture of R-PET at 5 - 60m hairline
crack depth

Figure 4.29

169

176

DSC samples taken from slices of dumbbell samples at 50m per slice
which were peeled off from the thickness side showing the skin/core
appearance of the materials at different slice depths [a = V-PET and b = RPET]

Figure 5.2

211

Gray image plot showing the gray values of the skin, interface and coreregions of V-PET [W = graph of gray values on the width direction, X =xxgraph of gray values on the thickness direction]
Figure 5.3

212

Determination of maximum change in temperature (MΔT) at full ligament
yielding of the stress-strain curve for the different notch depths; [inset
infrared images (a) before yielding, (b) at yielding, and (c) full ligament
yielding]

Figure 5.4

213

Crack tip temperature distributions of V-PET samples showing degree ofchanges at the skin, interface and core deep regions [a = un-notched, b =0.5-0.9 mm, c = 1.0-1.5 mm and d = 1.7-3.0 mm deep notches; arrows
indicate the temperature at full ligament yielding]Figure 5.5

214Infrared images of V-PET showing the fracture modes at different notch
depths where maximum changes in temperature occurred at maximum atfull ligament yielding [a = Un-notched samples, (b and c) represent 0.5 0.9mm, d = 1.0mm, e = 1.3mm, f = 1.5mm, and (g – i) represent 2.0 –3.0mm]

Figure 5.6

215

Maximum change in notch root temperature (V-PET) in the process zone
showing drastic rise at the interface (1.0-1.5mm notch) between the skin andcore regions where mixed fracture occurred with crazing.

Figure 5.7

216

Stress/Strain curves of V-PET showing the essential (We) and non-essential(We) works of fracture deformation zones for the specimens at different
notches ranging from un-notched (Un) to 1.5mm]

217

Figure 5.8

Total work of fracture for V-PET at different ligament depths

218

Figure 5.9

Essential work of fracture for V-PET at different ligament depths

219

Figure 5.10

Non-essential work of fracture for V-PET at different ligament depths

220

xxi

Figure 5.11

Energy dissipation in the process (We) and plastic deformation zones (Wp)with increasing ligament length

Figure 5.12

221

Ratio of essential (We) and non-essential (Wp) works of fracture for thematerials at the various notch depths

Figure 5.13

222

Correlation between change in temperature and the ratio of energy
dissipated in both the process (We) and plastic (Wp) deformation zones

223

Figure 5.14

Correlation between We/Wp and notch sensitivity factors for strength

224

Figure 5.15

Notch sensitivity factors for energy kT and the ratio of essential and nonessential work of fracture (We/Wp) for the V-PET at the various notch depth

Figure 5.16

225

Gray image plot showing the gray values of the skin, interface and coreregions of R-PET [W = graph of gray values on the width direction, X =graph of gray values on the thickness direction]

Figure 5.17

226

Crack tip temperature distributions of X = V-PET and Y = R-PET samples
showing degree of changes at the skin, interface and core deep regions [a =un-notched, b = 0.5-0.9mm, c = 1.0-1.5mm and d = 1.7-3.0mm deep
notches; arrows indicate the temperature at full ligament yielding]Figure 5.18

227

Infrared images of RPET showing the fracture modes at different notch
depths where maximum changes in temperature occurred at full ligament
yielding [a = Un-notched samples, (b and c) represent 0.5-0.9mm, d =1.0mm, e = 1.3mm and f = 1.5mm]

Figure 5.19

Comparing maximum change in temperature between V-PET and R-PET in
the Process zone showing drastic rise at the interface (1.0-1.5mm notch)
between the skin and core regions where mixed fracture occurred with

xxii

228

crazing
Figure 5.20

229

Stress/Strain curves showing the essential (We) and non-essential (Wp)works of fracture deformation zones for the specimens at different notches
ranging from un-notched (Un) to 1.5mm [a = V-PET and b = R-PET]

230

Figure 5.21

Total work of fracture at different ligament depths

231

Figure 5.22

Essential work of fracture for V-PET and R-PET at different ligament depths 232

Figure 5.23

Essential work of fracture for V-PET and R-PET at different ligament depths 233

Figure 5.24

Energy dissipation in the process (We) and plastic deformation zones (Wp)

Figure 5.25

Ratio of essential (We) and non-essential (Wp) work of fracture for thematerials at the various notch depths

Figure 5.26

236

Correlation between We/Wp and notch sensitivity factors for strength for VPET and R-PET

Figure 5.28

235

Correlation between change in temperature and the ratio of energy
dissipated in both the process (We) and plastic (Wp) deformation zones

Figure 5.27

234

237

Notch sensitivity factors for energy kT and the ratio of essential and nonessential works of fracture (We/Wp) for the materials at the various notch
depths238

Figure 6.1

Mode I Fracture249

Figure 6.2

Mode II Fracture

250

Figure 6.3

Mode III Fracture

251

Figure 6.4

Dumbbell and rectangular specimens for (a) tensile and (b) impact tests

252

Figure 6.5

Determination of σQ at 5° off tangent

253

Figure 6.6

Tensile Strength and Modulus for Un-notched Specimens

254

xxiii

Figure 6.7

Elongation at Break (%) of the Un-notched Specimens

255

Figure 6.8

Notch Sensitivity of GFPP

256

Figure 6.9

Maximum Stress Intensity Factor (K1max/KQ) for GFPP

257

Figure 6.10

Maximum Strain Energy Release Rate (G1max/GQ)

258

Figure 6.11

Micrograph of GFPP (J-700GP)

259

Figure 6.12

Micrograph of GFPP (J-900GP)

260

Figure 6.13

Micrograph of GFPP (J-3000GP)

261

Figure 6.14

Charpy Impact Strength for Un-notched and Notched Specimens

262

Figure 6.15

Izod Impact Strength for Un-notched and Notched Specimens

263

Figure 7.1

Schematic comparison of the stress-strain behavior of elastic-plastic andnonlinear elastic materials

288

Figure 7.2

Dumbbell specimen for tensile test

289

Figure 7.3

Schematic of sample portions burnt off for fibers measurement in thedetermination of interfacial shear strength

Figure 7.4

Grayscale image of glass fibers used in the calculation of interfacial shear
strength

Figure 7.5

290

291

Correlation of tensile strength and modulus of PP/GFPP and molecularweight292

Figure 7.6

J1max, JQ, K1max and KQ for GFPP of different molecular weights

293

Figure 7.7

CTODmax and rymax for GFPP of different Molecular Weights

294

Figure 7.8

Fracture surface of A = GFPP(J-700GP), B = GFPP(J-900GP), C = GFPP(J3000GP) [x = notch, y = ductile region, z = brittle region]

Figure 7.9

Correlation between the lengths of ductile /brittle fracture regions and

xxiv

295

molecular weight of PP
Figure 7.10

296

Interfacial bonding between GF and PP of different molecular weight [a =GFPP(J-700GP), b = GFPP(J-900GP) and c = GFPP(J-3000GP)]

Figure 7.11

297

Crack Propagation of GFPP composites monitored by CCD Camera; a =GFPP (J-700GP), b = GFPP (J-900GP), c = GFPP (J-3000GP)

298

Figure 7.12

Correlation between interfacial shear strength and molecular weight of PP

299

Figure 7.13

Average and critical fiber lengths of the composites

300

Figure 7.14

Fiber lengths distribution of GFPP (J-700GP) in percentage frequency

301

Figure 7.15

Fiber lengths distribution of GFPP (J-900GP) in percentage frequency

302

Figure 7.16

Fiber lengths distribution of GFPP (J-3000GP) in percentage frequency

303

Figure 7.17

Fracture toughness and molecular weight of PP

304

Figure 7.18

Fracture toughness and tensile modulus of GFPP

305

Figure 7.19

Fracture toughness and tensile strength306

Figure 7.20

Fracture toughness and interfacial shear strength

307

Figure 7.21

Average/critical fiber length and molecular weight of PP

308

Figure 7.22

Fracture toughness and Izod impact strength (un-notched)

309

Figure 7.23

Fracture toughness and Izod impact strength (notched)

310

Figure 7.24

Fracture toughness and Charpy impact strength (un-notched)

311

Figure 7.25

Fracture toughness and Charpy impact strength (notched)

312

Figure 7.26

Degree of association between molecular weight, mechanical, impact,
interfacial properties and fracture toughness

xxv

313

ABSTRACTThe prevention of catastrophic failures in engineering components depends largely on the accurate predictionof their fracture mechanism. Much research efforts have been made to understand the fracture toughness ofpolymeric materials and composites but these are limited to investigations conducted on samples introducedwith single-depth standard V-notches. However, there is inadequate understanding on the notch sensitivity of
injection moldings especially when a hierarchical skin-core structure is present. Therefore, this thesis focusedon characterizing the fracture toughness of injection molded polyethylene terephthalate (PET) and glass fiber
reinforced polypropylene (GFPP) composites, both of which are known for being notch sensitive. Chapters 3
and 4 of this thesis were dedicated to understanding the effects of standard V-notch and sharper hairline
cracks, respectively, on the fracture mechanism of neat- (V-PET) and recycled-PET (R-PET) injection
moldings. The PET specimens exhibited very distinct skin-core structures and the depths of these regions
were well defined. Upon tensile test, the specimens would remain semi-ductile as long as the V-notch depth
was within the skin region. The skin region was tougher due to its lower crystallinity and would therefore,
yield extensively during tensile loading. However, brittle fracture was imminent when the notch penetrated
the skin into the core, which had higher crystallinity than the skin. Nevertheless, except upon annealing, thetoughness of the skin region was compromised when sharper hairline cracks were introduced. The
dependence of crack-tip temperature on stress intensity and notch sensitivity was investigated in Chapter 5.
High crack-tip temperatures would indicate full ligament yielding, which provided the necessary
requirements for the application of essential work of fracture principles in characterizing the fracture
toughness of the specimens. Chapters 6 and 7 of the thesis were dedicated to investigating the effect ofpolypropylene (PP) molecular weight on the fracture toughness of glass fiber reinforced PP (GFPP)
composites. Elastic plastic fracture mechanics have revealed that GFPP containing the lowest molecularweight PP (J-3000GP) exhibited the highest fracture toughness due to its high interfacial shear strength that
resulted in more fiber breakage and better resistance to crack propagation.

xxvi

Chapter 1
General Introduction

1.1 Background of Studies on Fracture Behavior and Fracture Toughness of Materials
The phenomenon of structural failure by catastrophic crack propagation at average stresses well
below the yield strength has been known for many years. Rashes of such brittle failures occur
with increasing frequency as the strength and size of engineering structures have increased. In
the past, each series of failures has given rise to a set of empirical tests and procedures thatsometimes provided a solution to the specific problem at hand but did not result in a generally
useful approach that would permit avoiding future failures. The military and aerospacerequirements for very-high-strength, lightweight hardware have given added importance to the
problem of brittle fracture and greatly emphasized the need for a quantitative approach to thegeneral problem of crack tolerance in structures. This need was dramatically highlighted several
years ago by the repeated failure of early Polaris rocket motor cases at stresses well below thedesign value. ASTM committee was formed with the goal of providing laboratory tests andAnalytical techniques, which will permit a quantitative measure of crack tolerance useful not
only in evaluating materials for a given application but also in the development of rational
procedure for design against fracture. To achieve this goal requires the development of an
essentially new branch of engineering science, and this, of course, is an evolutionary process,which will take considerable time to complete. However, with the Irwin linear elastic fracture
mechanics as a basis, considerable progress has been made in the desired direction, and today
there are available reliable procedures for avoiding failure by fracture in new structures.

1

1.2 Failure of Structures and the Need for Fracture Mechanics
Fracture is a problem that society has faced for as long as there have been fabricated structures.
The problem may actually be worse today than in previous centuries, because more can gowrong in our complex technological society. Major airline crashes, for instance, would not be
possible without modern aerospace technology. Fortunately, advances in the field of fracture
mechanics have helped to offset some of the potential dangers posed by increasing technologicalcomplexity. The understanding of how materials fail, and the ability to prevent such failures has
increased considerably since World War II. Much remains to be learned, however, and existingknowledge of fracture mechanics is not always applied when appropriate. Failures of structures
are classified into two common causes, namely negligence during design, construction oroperation of the structure, and introduction of a new design or material, which produces an
unexpected and undesirable result. In the first instance, existing procedures are sufficient toavoid failure, but are not followed by one or more of the parties involved, due to human error,
ignorance, or willful misconduct. Poor workmanship, inappropriate or substandard materials,
errors in stress analysis, and operator error are examples of where the appropriate technology andexperience are available, but not applied. The second type of failure is much more difficult toprevent. When an "improved" design is introduced, there are invariably factors that the designer
does not anticipate. New materials can offer tremendous advantages, but also potential problems.
Consequently, a new design or material should be placed into service only after extensive testing
and analysis. Such an approach will reduce the frequency of failures, but not eliminate them.

In this research effort, the fracture behavior and fracture toughness of PET and glass fiber
reinforced polypropylene composites were evaluated. The PET materials will be investigated to2determine the effect of the hierarchical skin-core structure on their notch sensitivity when Vnotches were introduced. The dependence of crack-tip temperature on stress intensity and notch
sensitivity of PET will also be studied. Essential work of fracture theory will be applied toanalyze the fracture toughness of the materials. Another investigation on PET will focus on theeffect of hairline crack and annealing on its fracture mode.

Investigations will equally be

conducted on the fracture toughness of glass fiber reinforced polypropylene (GFPP) composites.
These shall be investigated for characterization within the context of Linear Elastic Fracture
Mechanics (LEFM) and Elastic Plastic Fracture Mechanics (EPFM).

1.3 Characteristics and Recycling of PET
Poly(ethylene terephthalate) (PET) is a semi-crystalline thermoplastic that has wide end-use
applications because of its excellent tensile and impact strength, chemical resistance, clarity,
processability, color ability and reasonable thermal stability. PET is said to have a wide range of
intrinsic viscosity [ ] that varies from 0.45 to 1.2 dl g-1 with a polydispersity index generally
equal to 2.0. The PET chain is considered stiff above the glass transition temperature (T g) un-likemany other polymers. The low flexibility of PET chain is a result of the nature of the short
ethylene group and the presence of the p-phenylene group. The chain inflexibility significantly
affects PET structure-related properties such as thermal transitions. Commercial PET has amelting temperature (Tm) of between 255 and 265ºC, while a more crystallized PET has Tm of
265ºC. The Tg of virgin PET (V-PET) varies between 67 and 140ºC. This phenomenon is
attributable to morphological and structural re-organization. As the temperature increases, better
crystal structures are achieved because of the re-organization of the crystals that were not
perfect. V-PET is well known for having very slow crystallization rate. The highest

3

crystallization rate takes place at 170ºC, or 190ºC. Cooling PET rapidly from the melt to atemperature below Tg can produce an amorphous, transparent PET. Semi-crystalline PET can beobtained by heating the solid amorphous PET to a temperature above Tg where 30% crystallinity
can be achieved. The rate of crystallization of V-PET depends greatly on temperature, andreaches its maximum at a temperature of 150 – 180ºC. The rate of crystallization also depends onother factors such as MW, the presence of nucleating agents, the degree of chain orientation, thenature of the polymerization catalyst used in the original production of PET and the thermal
history. Pure PET has little value when injection molded. Injection molded PET has poor
mechanical properties due to limited crystallization, which is attributed to the high Tg of PET
that can occur during cooling after injection molding. Improvement has been made by many
researchers and companies through the addition of nucleating agents and/or increasing moldtemperatures to as high as 140ºC to promote crystallization.

The recycling of PET began out of the pressure to keep the environment free from postconsumer PET products. However, as new and cheaper technologies for recycling developed,
recycled PET (R-PET) becomes a source of cheap PET supply. The physical and chemical
properties of R-PET, nonetheless, are undermined by the presence of contaminants; therefore
minimizing the level of these contaminants improves the quality of R-PET. Such contaminants
include: (a) acid producing contaminants - acetic acid produced by poly(vinyl acetate) closuresdegradation; rosin acid and abietic acid produced by adhesives and hydrochloric acid produced
by PVC. The acids act as catalysts for chain scission reactions during POST Consumer-PET melt
processing; (b) Water - reduces MW during POSTC-PET recycling through a hydrolysis reaction;

4

(c) Coloring contaminants - fragments of colored bottles and printed ink labels cause undesirable
colors during processing.

1.4 Fracture Behavior of PET
Much research effort has been made towards improving the physical and chemical properties of
R-PET. These research efforts have yielded good and improved results. The improved propertiesof R-PET give researchers the opportunity to study the fracture behavior of R-PET incomparison with V-PET. PET products are highly ductile when there are no notches on them.
However, PET products are belied to be very notch-sensitive when notches are introduced onthem. There are sufficient data on the fracture behavior of injection moldings when there are no
cracks, which gives products designers the opportunity to predict their fracture modes.
Nevertheless, there are insufficient data on the prediction of the fracture behavior of
PET/polymeric products when there are different notch geometries, particularly in PET products,
which are semi crystalline thermoplastic material known to be sensitive to notches. Fractures
occur in materials when there are pre-existing cavities or nucleation of voids due to stressful
loading. Such defects in engineering thermoplastic components are stress concentrators that tend
to undermine the structural stability of such components during their life span. The resultant
effects of cavities, voids or defects in engineering structures so often can be catastrophic.
Catastrophic failure of structures claims heavily in economic and human cost. Therefore, thestudy of fracture mechanics is to investigate the materials’ response to fracture under a predetermined crack condition. Understanding the fracture behavior of materials with different sizesof defects helps components designers and engineers to balance between the choice of discarding
very expensive defective material and safety requirements of engineering construction. PET is a5semi-crystalline thermoplastic material known to be very notch sensitive. The notch sensitivityof PET products had led to many research investigations, which reveal that both V-PET and RPET products exhibit low fracture resistance in the presence of notches.

It is believed that the fracture modes of injection-molded materials vary differently with different
notch geometries. The effect of v-notches on the fracture behavior of polymer materials appears
to be different from the effect of hairline cracks. Because of the large angle and radius ofcurvature of v-notch as shown in Figure 1.1, there tends to be wall rounding at the crack tip
during crack propagation, which results in slow crack propagation and more resistance tofracture. In contrast, because of small angle and radius of curvature as shown in Figure 1.2,
hairline (microscopic) crack tends to result in fast crack propagation during fracture. However,the notch depth at which v-notched and hairline-cracked PET materials fail is not yet very clear.
The hierarchical skin-core structure of PET injection moldings is believed also to affect the notch
depth where the material fails under each of the notch geometries. This study will attempt toinvestigate the fracture modes of PET injection moldings at increasing notch depth, and the roleof the skin-core structure to determine the critical notch depths at failure for v-notch and hairline
cracks. The results will be evaluated in part with notch sensitivity factor for fracture strength (kS)and notch sensitivity factor for energy (kT).

1.5 Effect of Annealing on the Notch Sensitivity of PET
Very sharp and microscopic defects are known to affect the fracture modes of engineering
materials. Unlike v-notches, crack initiation and propagation in sharp microscopic cracks (alsoknown as hairline cracks) is very rapid and catastrophic. Some studies have shown that a mere6

10m deep hairline crack would cause a brittle fracture in PET injection moldings. The skin andcore structure of the injection moldings is believed to affect their fracture modes. The slow
crystallization nature of PET results in a distinct amorphous skin and more crystallized coremorphology when the materials were injection molded. By precisely controlling the notch depth,
the notch root can be situated within the skin, interfacial or core regions of the moldings. The
fracture behavior of the notched specimens during tensile and impact loading was found to be
dependent not only on the notch depth but also the region where the notch roots were situated(skin, interface or core). By annealing the specimens in 160°C oil for 35s and quenching it in 710°C ice-chilled oil as shown in Figure 1.3, the skin-core ratio was altered in order to evaluate its
effect on the fracture behavior of hairline-cracked specimens. The effect of annealing on theskin-core morphology and fracture behavior of hairline-cracked samples was correlated with thetoughness of PET products. This will help to draw a conclusion on whether the annealing of PET
injection moldings can reduce their notch sensitivity.

1.6 Crack Tip Temperature and Essential Work of Fracture
Fracture is a dissipative process during which plastic work is converted partly into heat, or heatis released during the fracture process itself. When the process is adiabatic, the released heatcauses a local temperature rise. In the formation of an adiabatic shear band there is a substantial
increase in the local temperature in the band. Therefore, heat dissipation due to the combined
effect of plastic deformation (plastic flow) and fracture during tensile loading of PET is believed
to affect the temperature around the stress concentration point, including the crack-tip. This
could equally reflect the magnitude of stress intensity on the stress field of the materials. The
presence of voids, cracks, cavities, or even impurities can influence the stress and temperature7

distribution upon tensile loading. The introduction of cracks on injection molded flat bars affects
their lines of stress. In an un-cracked uniform bar (dumbbell), the stress concentration point islocated at the center, and runs parallel to the direction of applied force.

In injection-molded semi-crystalline thermoplastics such as polyethylene terephthalate (PET),
the presence of skin and core layers becomes another factor that influences their fracture
behavior. It is known that some materials are more sensitive to cracks than others are. Therefore,
it was suggested that stress-strain tests should be conducted on both notched and un-notched test
specimens for comparison. Stress-strain tests enable design engineers to predict the fracture
behavior of perfectly smooth (uniform) objects, but these do not allow the prediction of thefracture behavior of objects with cuts. Deeply notched specimens, however, do not allow
sufficient consideration of the contributions from the skin region towards fracture behavior of thebulk.

An earlier study revealed three distinct fracture behaviors during tensile test, i.e. ductile, semiductile, and brittle, when notches ranging from 0.5 to 4.5 mm in depth were introduced to one
side of the specimen of 10.0 mm in width. This variation in fracture behavior appears to relate to
the skin-core structure of PET injection moldings. Semi-ductile failure was imminent when thenotch depth fell within the thickness of the skin region. Therefore, it would be premature to
assume that PET injection moldings are notch sensitive especially when a deep notch thatpenetrates through the skin into the core region is introduced. This study evaluates the fracture
toughness of PET injection moldings when the notch depths are within the thickness of the skinregion. The theories of essential work of fracture (EWF) were applied to elucidate the fracture8

toughness of the moldings. An essential requirement for the application of EWF is that fullligament yielding has occurred prior to crack displacement. Therefore, the temperature of thespecimens was monitored during tensile tests to confirm localized ligament yielding to satisfy
the requirements of EWF as well as to pinpoint the regions where the highest plastic flow has
occurred. Shallow notches of various depths were carefully machined into the injection moldings
and their influence on the toughness and notch sensitivity of the PET bulk was elucidated.

1.7 The Need for Investigation on the Effect of V-notches and Hairline Cracks
Most studies on the fracture behavior of polymer materials investigated only one single standard
notch depths (ASTM D256). Since injection-moldings are known to exhibit anisotropic skin-coremorphological structure, these single depth notches are usually deep enough to penetrate throughthe skin region. As such, the introduction of standard notches does not provide adequate
consideration for the influence of the hierarchical skin-core structure on the fracture
characteristics of the bulk material. Furthermore, in slow crystallizing materials such as PET, thedevelopment of a distinct skin-core structure during injection molding is inevitable. PET is
known also to be a very notch-sensitive material and would fracture in a very brittle manner
upon introduction of notches but remains very tough if un-notched. However, the effect of notch
depth on the fracture characteristics is still unclear. Therefore, it is necessary to introduce
notches of different geometries at various depths and examine their effects on the crack
propagation behavior of injection molded PET products. If a correlation between fracture
characteristics and morphology could be established, the information will be very useful to assist
molders in optimizing molding conditions as well as improve the safety and reliability of themoldings.

9

In order to understand these effects, V-notches ranging from 0.5-4.5mm and hairline cracks
ranging from 5 – 60m depths were introduced on single edges of dumbbell and Izod impact
samples. These depths were chosen so that the notch root was positioned along the skin throughthe core regions of the moldings. Skin and core regions were determined through polarized
optical microscopy. The crystallinity along the skin-core width was determined by differential
scanning calorimetery and Raman Spectroscopy. The fracture characteristics of the samples atvarious notch depths were monitored, during tensile and impact testing, by various opticaldevices including a CCD camera.

1.8 Composite Materials
Composites are compound derivatives of two different and heterogeneous materials with eachhaving distinct and homogeneous property characteristics. These two distinct materials arebrought into one form, otherwise known as composites, by means of chemical reaction. The
resultant composites have new property characteristics, which are more superior to individual
properties of the consisting materials. The possibility of combining favorable properties fromvarious materials has made composite materials the choice for many engineering applications
and environments. Composites can be made from neat or blended thermoplastics that are
reinforced with fibers. In whichever case, the compatibility of the fibers and resins are enhanced
by incorporating coupling agents. Composites can come into effect when they are subjected toindustrial processes such as extrusion, molding processes such as injection molding, hand lay-up,
resin transfer molding, etc. During the extrusion process, pre-mixed thermoplastic resin andreinforcements are fed into an extruder, where further blending occurs at temperatures above themelting point of the resin. The compounded materials are subsequently extruded and pelletized10

before being used for secondary processes such as injection molding. This research focuses oncomposites fabricated by the injection molding process.

1.9 Injection Molded Composites
Injection molding is one of the most important methods for processing plastics. To this end both
reinforced and un-reinforced thermoplastics materials are processed by injection molding inorder to obtain a variety of molded end-use products. In order to obtain a desired part by
injection molding process, consideration is given for the selection of a plastic material, which has
the required properties.

In this research effort, focus will be on the mechanical and fracture mechanism of injection
molded glass fiber reinforced polypropylene (GFPP) composites, with emphasis on the effects ofmolecular weight and compatibilizers on the fracture toughness properties.

1.10 Thesis StructureThis thesis is arranged into eight main chapters in order to deal with each specific subject area.
Chapter 1 provides some introductory descriptions and general aspects of PET and composite
materials that are related to the study. The overview of the thesis presentation is also outlined. In
chapter 2, literature materials related to the subject areas of investigation are reviewed. The
review encapsulates the various published works and findings related to the present study.
Fundamental principles that were applied during the course of this study are described in detail.
Chapter 3 deals with the fracture modes and notch sensitivity of V-notched PET injectionmoldings. In Chapter 4, the fracture modes of hairline-cracked PET injection-moldings and the11role of annealing on the notch sensitivity of PET will be evaluated. Chapter 5 studies theapplication of EWF theory in characterizing the fracture behavior of PET and dependence of
crack-tip temperature on notch sensitivity of PET injection-molded components. The crack-tip
temperature distribution was monitored with an infrared thermography camera. In Chapter 6, thefracture toughness properties and fracture behavior of three grades of polypropylene withdifferent molecular weights reinforced with glass fiber (GFPP) were investigated. This
investigation was conducted within the context of linear elastic fracture mechanics (LEFM) andthe fracture toughness was characterized by using the critical stress intensity factor and criticalstrain energy release rate approaches, which were designated as K1max and G1max, respectively. In
chapter 7, GFPP was characterized under elastic plastic fracture mechanics (EPFM) by using JIntegral and crack tip opening displacement parameters. Further study was conducted on thefracture behavior of single-edge-notched GFPP composites by using tensile test and this wasmonitored by a CCD camera equipped with high-resolution optical microscope. Investigation
was also conducted on the interfacial shear strength of the composites. Lastly, Chapter 8 gives ageneral summary of results and findings of all the investigations conducted on V-PET/R-PET
and GFPP composites.

References
1. Anderson, T. L., Fracture Mechanics, Fundamentals and Applications (2nd ed.), Boca
Raton, London, 3-7, (1995)
2. Garwood, S.J., Private Communication, (1990).
3. Jones, R.E. and Bradley, W.L., Failure Analysis of a Polyethylene Natural Gas Pipeline,
Forensic Engineering, 1, (47-59, 1987)

12

4. Jones, R.E. and Bradley, W.L., Fracture Toughness Testing of Polyethylene Pipe
Materials, ASTM STP 995, 1, 447-456, (1989), American Society for Testing andMaterials, Philadelphia
5. Duga, J.J., Fisher, W.H., Buxbaum, R.W., Rosenfield, A.R., Burh, A.R., Honton, E.J.,
and McMillan, S.C., The Economic Effects of Fracture in the United States, NBS Special
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6. Shank, M.E., A Critical Review of Brittle Failure in Carbon Plate Steel Structures Other
than Ships, Ship Structure Committee Report SSC-65, National Academy of Science,
National Research Council, Washington, DC, December, (1953)
7. Love A.E.H, A Treatise on the Mathematical Theory of Elasticity. Dover, New York.
(1944)
8. Griffith, A.A. The Phenomena of Rupture and Flow in Solids. PhilosophicalTransactions, Series A, 221, 163-198, (1920)
9. Inglis, C.E., Stresses in a Plate Due to the Presence of Cracks and Sharp Corners,
Transactions of the Institute of Naval Architects, 55, 219-241, (1913)
10. Bannerman, D.B. and Young, R.T., Some Improvements Resulting from Studies ofWelded Ship Failures, Welding Journal, 25, (1946)
11. Irwin, G.R., Fracture Dynamics, Fracturing of Metals, American Society for Metals,
Cleveland, 147-166, (1948)
12. Orowan, E., Fracture and Strength of Solids, Reports on Progress in Physics, XII, 185232, (1948)
13. Mott, N.F., Fracture of Metals: Theoretical Considerations, Engineering, 165, 16-18,(1948).
14. Irwin, G.R., Onset of Fast Crack Propagation in High Strength Steel and AluminumAlloys, Sagamore Research Conference Proceedings, 2, 289-305, (1956)
15. Westergaard, H.M., Bearing Pressures and Cracks, Journal of Applied Mechanics, 6, 4953, (1939).
16. Irwin, G.R., Analysis of Stresses and Strains near the End of a Crack Traversing a Plate,
Journal of Applied Mechanics, 24, 361-364, (1957).

13

Radius of curvature

Angle of curvature

Figure 1.1: V-notch showing the radius (1030mm) and angle of curvature(45°) as fitted with circle at tangent with the base of the notch
14

Radius of curvature

Angle of curvature

Figure 1.2: Hairline-crack showing the radius (90mm) and angle of curvature(about 8°) as fitted with circle at tangent with the base of the notch
15

Dumbbell in-between
two aluminum bars

Oil
160⁰C

Hot OilElectric Pot

Ice-blockFigure 1.3: Apparatus setup showing the annealing of specimens
in hot edible oil and subsequent quenching in ice-chilled oil (7-10°C)
16

Chapter 2
Review of Related Literatures

2.1 Neat (V-PET) and Recycled (R-PET) Poly(ethylene terephthalate) Injection Moldings
Poly(ethylene terephthalate) (PET) is a semi-crystalline thermoplastic that has wide end-use
applications because of its excellent tensile and impact strength, chemical resistance, clarity,
processability, color ability and reasonable thermal stability. PET is said to have a wide range of
intrinsic viscosity [ ] that varies from 0.45 to 1.2 dl g-1 with a polydispersity index generally
equal to 2.0. The PET chain is considered stiff above the glass transition temperature (Tg) un-likemany other polymers. The low flexibility of PET chain is a result of the nature of the short
ethylene group and the presence of the p-phenylene group. The chain inflexibility significantly
affects PET structure-related properties such as thermal transitions [1]. Commercial PET has amelting temperature (Tm) of between 255 and 265ºC, while a more crystallized PET has Tm of
265ºC. The Tg of virgin PET (V-PET) varies between 67 and 140ºC. This phenomenon is
attributable to morphological and structural re-organization. As the temperature increases, better
crystal structures are achieved because of the re-organization of the crystals that were less perfect[1]. V-PET is well known for having very slow crystallization rate. The highest crystallization
rate takes place at 170ºC, or 190ºC [1]. Cooling PET rapidly from the melt to a temperature
below Tg can produce an amorphous, transparent PET. Semi-crystalline PET can be obtained byheating the solid amorphous PET to a temperature above Tg where 30% crystallinity can beachieved [1]. The rate of crystallization of V-PET depends greatly on temperature, and reaches
its maximum at a temperature of 150 – 180ºC. The rate of crystallization also depends on otherfactors such as MW, the presence of nucleating agents, the degree of chain orientation, the nature17

of the polymerization catalyst used in the original production of PET and the thermal history [1].
Pure PET has little value when injection molded. Injection molded PET has poor mechanicalproperties due to the limited crystallization, which is attributed to the high Tg that PET has, thatcan occur during cooling after injection molding. Improvement has been made by many
researchers and companies through the addition of nucleating agents and/or increasing mould
temperatures to as high as 140ºC to promote crystallization [1]. Examples of nucleating agents
which improve crystallinity are 1,5-pentanediol and 1,8-octanediol [1].

The recycling of PET began out of the pressure to keep the environment free from postconsumer PET products. However, as new and cheaper technologies for recycling developed,
recycled PET (R-PET) becomes a source of cheap PET supply. The physical and chemical
properties of R-PET, nonetheless, are undermined by the presence of contaminants; therefore
minimizing the level of these contaminants improves the quality of R-PET [1]. Suchcontaminants include: (a) acid producing contaminants - acetic acid produced by poly(vinyl
acetate) closures degradation; rosin acid and abietic acid produced by adhesives and hydrochloric
acid produced by PVC. The acids act as catalysts for chain scission reactions during POSTCPET melt processing [1]; (b) Water - reduces MW during POSTC-PET recycling through ahydrolysis reaction [1]; (c) Coloring contaminants - fragments of colored bottles and printed ink
labels cause undesirable colors during processing [1].

2.2 Effect of Notches on the Fracture Behavior of PET
Much research effort has been made towards improving the physical and chemical properties of
R-PET. These research efforts have yielded good and improved results. The improved properties18

of R-PET give researchers the opportunity to study the fracture behavior of R-PET incomparison with V-PET. There are sufficient data on the fracture behavior of injection moldings
when there are no cracks, which gives products designers the opportunity to predict their fracturemodes [2]. However, there are insufficient data on the prediction of the fracture behavior of
PET/polymeric products when there are different notch geometries, particularly in PET products,
which are semi crystalline thermoplastic material known to be sensitive to notches [2].

Fractures occur in materials when there are pre-existing cavities or nucleation of voids due tostressful loading. Such defects in engineering thermoplastic components are stress concentrators[3] that tend to undermine the structural stability of such components during their life span. The
resultant effects of cavities, voids or defects in engineering structures so often can becatastrophic. Catastrophic failure of structures claims heavily in economic and human cost.
Therefore, the study of fracture mechanics is to investigate the materials’ response to fracture
under a pre-determined crack condition. Understanding the fracture behavior of materials withdifferent sizes of defects helps components designers and engineers to balance between the
choice of discarding very expensive defective material and safety requirements of engineering
construction. PET is a semi-crystalline thermoplastic material known to be very notch sensitive.
The notch sensitivity of PET products had led to many research investigations, which reveal thatboth V-PET and R-PET products exhibit low fracture resistance in the presence of notches. The
effect of standard V-notches on polymeric materials including PET was investigated byconducting tensile test on dumbbell test bars with single and double edge notches of 1.27 and3.175mm deep [4]. Equations (2.1) & (2.2) were used to calculate the notch sensitivity factors
for fracture strength (kS) and energy to fracture (kT), respectively. Results show that if a notch19has no detrimental effect on the toughness of a polymeric material, the notch sensitivity factor is
1.0 otherwise; it is greater than 1.0. The investigation shows that the values of kS for PET at 1.27
and 3.175mm notch depths were 1.02 and 1.26, respectively; while kT at the two notch depths
were 7.33 and 11.21 [4]. These results therefore, prove that notches ≥1.27mm are detrimental to
PET products as they fracture in a brittle manner.

 

 (YS ) [t (w  a )] i 0(YSi )  (tw) 




(2.1)

A
 t (w  ai )  
 0ssc T  A
tw


issc



(2.2)

S

Where kS = notch sensitivity factor for yield strength, kT = notch sensitivity factor for energy. YS0
and YSi = yield stress (strength) of un-notched and notched samples respectively [subscript i is
the successive notch depths], t = thickness, w = width, and a = notch depth. A0ssc and Aissc are theareas under the stress-strain curves for un-notched and notched specimens respectively [2, 4].

In recent years, there is much effort to improve the mechanical properties of PET products.
Investigation was conducted on the improvement in toughness of different blends of PET byconducting tensile test on standard 2.54mm deep-notched Izod bars at a temperature range of -20– 55oC [5]. The work showed that PET samples blended with un-functionalized styrenebutadiene-styrene (SEBS) yielded at the crack tip thereby fracturing in a brittle manner. On theother hand, blending PET with small amounts of functionalized SEBS elastomers was effective

20

in increasing the ductility (toughness) of PET products [5]. The fracture toughness of a material
is known to be dependent on its morphological structure and crystallinity [2, 6]. Therefore,
investigation was conducted on the effect of molecular weight and crystallinity on the notch
sensitivity of PET injection moldings. Two samples of varying degrees of crystallinity wereobtained by varying the sample cooling time in the mold and the mold temperature [6]. Tensile
test was conducted on the dumbbell samples cut with two notch conditions of 1.91 and 2.79mm
deep. The investigation revealed that crystallinity affects both the brittle fracture strength andyield behavior of PET injection moldings. It was shown that at ≤ 35% crystallinity there waslittle influence on the brittle stress, but the yield stress increased. However, the brittle stressgreatly decreased at 44% crystallinity. The results also indicate that notched samples had higher
yield stress in the cold-drawing region than the un-notched when tested at temperatures above
5⁰C [6].

2.3 Effect of Hairline Cracks and Annealing on the Fracture Behavior of PET Injection
Moldings
Very sharp and microscopic defects are known to affect the fracture modes of engineering
materials. Unlike v-notches, crack initiation and propagation in sharp microscopic cracks (alsoknown as hairline cracks) is very rapid and catastrophic. However, the fracture modes of hairline
cracked PET injection moldings are not yet understood. Investigations are lacking on the effectsof hairline cracks on PET injection moldings, therefore, there is urgent need for research efforts
into the fracture modes of PET injection moldings when there are hairline cracks.

Another element of interest on the fracture modes of hairline-cracked injection-molded PET
components is the possible effect of annealing. It is believed that annealing the components will21

create a crystallized outermost skin (referred to as the pseudo-skin). Annealing is equally done inorder to remove internal stresses, thus improving the fracture toughness of the materials. A study
on the effect of annealing revealed that the differences in the percentage crystallinity and residual
stress levels between the skin and core layers were decreased considerably after annealing [7].
However, in another study, injection molded PET was annealed at 120⁰C for 4hrs in order toincrease its crystallinity [8]. The study concluded that annealing was necessary to enable
subsequent selective dissolution of the elastomers phase. The effect of annealing on the fracture
behavior of PLA was investigated. Amorphous state was obtained by quenching, and annealing
was performed to the quenched plates to obtain highly crystallized solid structure. DSC
measured the thermal properties of the two different types. Mode I fracture toughness values, KIC
and GIC, were evaluated under quasi-static and impact loadings. Fracture micro mechanisms
were also investigated by polarizing microscopy and SEM. Thermal properties of PLA obtainedby DSC analysis showed that Tg and Tm slightly increased due to annealing as a result of
crystallization. The quenched PLA had low crystallinity of 2.7%, and annealing formed a
crystallized state with moderate crystallinity of 48.3%. Investigation revealed that annealing
lowered KIC and GIC under quasi-static loading; in contrast annealing increased the toughness
under impact loading. Multiple craze formation resulted in the high K IC and GIC of the quenchedstatic, and the suppression of the crazing lowered the toughness at impact loading rate. For theannealed condition, micro cracking at the spherulite interfaces might have been responsible forthe increase in the toughness [9]. Annealing of rectangular Polycarbonate specimens ofdimensions 150mm by 50 mm by 3 mm or 5 mm was in an airflow oven at 130°C for 5 to 250h
and then cooled at a slow rate (5°C h-1). Tensile tests on both Single-edge notch (SEN) and
Surface notch (SN) specimens were performed at ambient temperature. SEN was induced both22

by razor blades and by machining with a very sharp fly cutter. SN was induced by machining.
Results show that annealing increases the degree of brittleness. The 3 mm single-edge notch
specimens of un-annealed polycarbonate exhibited ductile failure with gross yielding andnecking over the range of the cross-head speed of the testing machine (0.005 to 50cm min -1)[10]. However, the 3mm SEN specimens became completely brittle when annealed below theirglass transition temperature. The 5mm thickness SEN specimens of un-annealed polycarbonate
exhibited a duplex behavior (both gross ductile and brittle fracture), while the 5 mm annealed
specimens were always brittle with a Kc value lower than that of un-annealed specimens(3.87MNm -3/2 compared to 4MNm -3/2 of un-annealed specimens). This transition has also been
observed in Izod and Charpy impact studies of notched samples of annealed polycarbonate.
Yield stress of annealed specimens are higher than that of the un-annealed ones, and this may
well have some influence on the change in behavior. Many other investigators have also reported
the increase in the yield stress. Annealing increases the density and produces an endothermic
peak at glass transition temperature, Tg, as measured on a differential scanning calorimeter. It
was suggested that the preheat treatment of polycarbonate in the temperature range 80 to 130°C
produces a greater degree of order within the amorphous region, resulting in an increase instrength. No effect of cooling rate on the fracture behavior of PC was observed. The specimensthat were cooled immediately by immersing in iced water showed the same fracture behavior as
those cooled at a slow rate (5°Ch-1). Increasing the annealing time from 5 h to 2 weeks has also
no effect on fracture behavior. Annealing temperature was found to be a governing factor sinceannealing below 100°C produced no change in fracture behavior [10].

232.4 Essential Work of Fracture Analysis
The essential work of fracture (EWF) theory splits the total energy required to fracture a precracked ductile material, with full ligament yielding into two components, namely the essential(We) and non-essential or plastic (WP) works of fracture. Therefore, the total work of fracture(Wf) is the area under the load-displacement curve and defined by equation (2.3)(2.3)

We is the energy needed to fracture the pre-cracked specimen in the process zone to create new
surfaces; Wp on the other hand is the actual work done in the outer plastic region where various
energy dissipation mechanisms take place. Because We is surface dependent and Wp is volume
dependent, equation (2.3) can be written in specific terms [11, 12]:

(2.4)
(2.5)

Where l = specimen ligament, t = specimen thickness, β = shape factor, wf, we and wp = specifictotal work of fracture, specific work of fracture and specific non-essential work of fracture
respectively.

However, the application of EWF on the characterization of materials is valid if the pre-cracked
material has full ligament yielding prior to fracture, and if there is a strong linear relationship(R2) between specific total work of fracture (wf) and ligament size [11, 13].

24

The fracture toughness of isotropic and flow-oriented linear low-density polyethylene (LLDPE)
was evaluated by the Essential Work of Fracture (EWF) concept. Special setup of CCD camera
was used to monitor the process of deformation. Allowing for molecular orientation, floworiented sample, prepared via melt extrusion drawing, is stretched parallel (oriented-08) andperpendicular (oriented-908) to its original melt extrusion drawing direction, respectively. The
obtained values of specific EFW we are 34.6, 10.2, and 4.2 N/mm for the oriented-08, isotropic
and oriented-908 sample, respectively. With knowledge of intrinsic deformation parameters
deduced from uniaxial tensile tests, moreover, a relationship between specific EFW we the ratioof true yield stress to strain hardening modulus

was established. This suggested that thefracture toughness of polyethylene is determined by both crystalline and amorphous parts, rather
than by one of them [12]. The essential work of fracture concept was extended to cover ductile
tearing of polymeric materials that neck before fracture. It is shown that the plane stress specificessential fracture work (we) can be obtained from deeply edge-notched tension specimens,
containing either single or double notches, by extrapolating the straight-line relationship between
the total specific fracture work (wf) and ligament length (l) to zero ligament [14, 17]. In this way,
specific essential fracture works have been obtained for nylon 66 and two polyethylenes. The
study suggests that we is a material property for a given sheet thickness being independent ofspecimen geometry. The straight-line relationship between wf and l breaks down when theligament length to sheet thickness ratio is less than about three, since the fracture data fall in theplane stress-plane strain transition region. However, a plane strain specific essential fracture
work can still be obtained by extrapolating the least squares curve of the data to zero ligament
provided the thickness satisfies plane strain condition. If this condition is not satisfied, a nearplane strain value is obtained which is dependent upon thickness. This method is also appropriate

25

for ductile polymers like the rubber-modified polystyrenes that craze rather than neck. J R curves
have also been obtained for nylon 66 and the polyethylenes. Under strictly J-controlled crack
growth conditions, it was shown that the intercept and slope of the JR curve, i.e. Jc and dJ/da, arerelated to the intercept and slope of the wf versus l plot [14].

Plane-stress fracture toughness of amorphous copolyester (COP) sheets of different composition
and molecular mass characteristics was determined by the essential work of fracture (EWF)
concept using tensile-loaded deeply double-edge notched (DDEN-T) specimens. It wasdetermined that these COPs meet the basic requirement of the EWF concept since their yielding
along the full ligament preceded crack growth. A drop in load in the corresponding loaddisplacement (F − x) curves indicated yielding and allowed us to split both the specific essential
and non-essential work of fracture (we and wp, respectively) into their contributing terms basedon yielding (wy) and necking including fracture (wn). Development and size of the plastic zone
were studied by light microscopy (LM) and infrared thermography (IT) [15].

The skin-core structure of injection molded poly(ethylene terephthalate) (PET)/polyethylene(PE) and polycarbonate (PC)/PE blends was investigated. The results indicate that both shapeand size of the PET and PC phases depended not only on the nature properties of PET/PE and
PC/PE blends, but also on the injection molding parameters such as injection speed and thepositions in the molded bars. The morphology in the section perpendicular to the melt flow
direction included four layers, surface, sub-skin, intermediate layers as well as core zone. The
surface layer was ignored in the present study. The sub-skin layer contained more or less fibrous
structure and its thickness gradually decreased along the molded bar from the gate toward the26non-gate end. At the same injection speed, the concentration of the injection-induced fibers in
PC/PE blend was much higher than that in PET/PE blend. In the core region, the dispersed phase
was mainly composed of spherical particles whose diameter increased along the melt flow
pathway. Between these two layers, there was an intermediate layer where the dispersed particles
mainly assumed the form of fibers, ellipsoids or spheres. Generally, no matter whether thedispersed particles were elongated or not during injection molding, the PET particles were larger
than PC ones [16].

2.5 Crack Tip Temperature
In engineering materials, when a crack propagates at very high velocity, it branches into two
cracks known as bifurcation. The focus of most investigation is to determine the minimum crack
propagation velocity, VB, for a particular homogeneous material at which the crack bifurcation
takes place, and the angle Ɵ0 that the branched crack makes with the original crack. A crack in amaterial is supposed to propagate at high velocity c, if it is greater than or equal to one third of
the shear wave velocity for the material, which is given by c2 equivalent toshear modulus of elasticity and, wheres theis the mass density of the material. When a crack propagates atsuch a high velocity, the energy balance at the crack tip can be written as:

G = ES + EP + EK

(2.6)

where G is the strain energy release rate; and ES , EP , EK are the specific surface energy, theplastic energy, and the kinetic energy, respectively. Part of the plastic energy E P contained in the27plastic deformation of the material ahead of the crack tip leads to the generation of heat Q F,
which increases the material temperature near the crack tip, which increases with the increase in
the crack speed. At low crack velocities, the heat generated by plastic work at the crack tip hasenough time to be conducted away from the newly formed crack surface and thus the rise of
temperature ahead of the crack tip will be negligible. The time for heat conduction reduces with
the increase in crack velocity and further eventually leading to the adiabatic temperature rise.
This causes local softening of the material and a consequent decrease in modulus of elasticity
and the critical stress intensity factor. Thus at high crack velocities the consideration of thethermal stresses becomes necessary. If the heat were to evolve adiabatically and uniformly about
the crack tip, the rise in temperature would be the plastic work per unit volume of the deformed
region [18].

The initiation and propagation of fracture cracks can result in a large local temperature rise near
the crack tip. This temperature rise strongly affects the nature of the near-tip deformation field,
and hence can lead to a decrease in the dynamic fracture toughness of the material. For thisreason, over the last two decades there has been a growing interest in investigating the thermal
behavior of polymer materials during crack initiation and propagation. The experimental
techniques have included calorimeters, contact thermocouples, and, more recently, non-contact
IR detectors. These efforts, though proven helpful for understanding the crack behavior, wererestricted to measuring this temperature rise at only a few points, and failed to get the full-field
information around the crack tip. It has been pointed out by various authors, that it is importantto get the full-field information on the temperature distribution around the crack tip in order tounderstand the crack propagation dynamics and fracture mechanisms. This was achieved by28

using advanced IR camera to investigate the full field, real time, thermal-wave imaging ofdynamically propagating cracks in notched polymer materials during tensile testing. The timedependent thermal behavior was studied, and results reveal that the behavior includes elastic
cooling, followed by heating caused by plastic deformation in the stressed region prior to crack
formation, and localized heating as the crack initiates and propagates, followed by diffusive
cooling after failure. The full field picture provides also information about the interaction of
cracks with pre-existing flaws, and the dynamics of the interaction between two cracks [19].

An analytical/numerical method has been developed to find the temperature rise near the crack
tip under fatigue loading. The cyclic plastic zone ahead of the crack tip is assumed the shape ofthe source of heat generation and some fraction of plastic work done in cyclic plastic zone asheat generation. Plastic work during fatigue load was found by obtaining stress/strain distribution
within the plastic zone by Hutchinson, Rice and Rosengren (HRR) crack-tip singularity fields
applied to small scale yielding on the cyclic stress strain curve. A two-dimensional conductionheat transfer equation, in moving coordinates, was used to obtain temperature distribution aroundthe crack tip. Temperature rise was found to be a function of frequency of loading, applied stressintensity factor and thermal properties of the material. A power–law relation was found between
the rise in temperature at a fixed point near the crack tip and range of stress intensity factor [20].

292.6 Mechanical Properties and Fracture Toughness of Glass Fiber Reinforced
Polypropylene Composites (GFPP)

2.6.1 Composites
The importance of composites as engineering materials has been realized in view of theiradvantageous properties, which include high fracture toughness, lightweight, thermal stability
and its ability to be processed into variety of designs. The mechanical and fracture properties ofcomposite materials are evaluated in order to predict their performance in their respective end
use application. The properties of the fiber-reinforced composites depend on many factors suchas fiber matrix adhesion, fiber orientation, stress transfer efficiency of the interface, andadditionally in the case of fracture toughness, on the geometry of the material. In fracture
mechanics, there are three modes of fracture denoted as ‘mode I’, ‘mode II’ and ‘mode III’
fractures. These different modes are illustrated in Figures 2.1 to 2.3 and the arrows show thedirection of applied force. In this work, investigation is on the mode I fracture of composite
materials.

2.6.2 Calculation of Fracture Toughness under Linear Elastic Fracture Mechanics (LEFM)
The application of a particular model to calculate the fracture toughness of a material depends on
the satisfaction of certain geometry conditions. A specimen to be tested can fall within any of thefollowing conditions as illustrated in Figure 2.4. A condition of plane stress applies when a test
specimen is very thin. Under this condition, the strain is not evenly distributed because of thegeometry effect, thus producing inconsistent fracture toughness results. Plane strain condition
applies when the thickness of a material is larger than a critical value, which enables a wide andeven distribution of strain at an applied stress. Therefore, under plane strain condition, the30fracture toughness of materials becomes constant. Hence, it can be applied as a material constant.
Fracture toughness was calculated from the tensile test data (see Figure 2.5 the geometries of atensile specimen) by using the stress intensity factor approach for the determination of K 1max
(maximum stress intensity factor) while the strain energy release rate approach was used for thedetermination of G1max (maximum strain energy release rate). Equations 2.7 and 2.8 are supposed
to be used to calculate K1max and G1max respectively under a plane-strain condition. However,since the dumbbell specimens used were rather thin a plane stress condition applies, thus
equations 2.9 and 2.10 were applied instead.

K1max   max  a  2 Or (  max  a )
1

(2.7)

where K1max is the fracture toughness in MPa m corresponding to a mode I crack in which
the crack is normal to the direction of largest applied force.

' max ' is the maximum applied stress at failure in MPa.
' a ' is the initial crack length in meters.
'  ' Is equal to 22 ; and7

G1max 2
 a max

E



(2.8)

where G1max is also the fracture toughness in kJ/m2 equally corresponding to the mode I crackand ‘E’ is the tensile modulus in GPa.

Equation (2.8) is the maximum strain-energy release rate measured in kJ / m2 for a plane-strain
situation in a mode I fracture. It measures the energy per unit area of the crack surface at the31maximum applied stress; i.e. it is the strain-energy release rate at the point of largest crack
opening.

The dumbbell specimens used in this investigation do not measure up to the plane-strain size
requirement; therefore, the plane-stress situation is applied. The following equations were
employed for the determination of maximum fracture toughness:

K1max   max  a  2 1.12  0.23  Wa   10.6  Wa   21.7  Wa   30.4  Wa 
1G1max 

2

3

4



(2.9)

2
K1max
E

(2.10)

where E = Young’s modulus measured in GPa.
The values of KQ and GQ were also calculated following a similar adaptation from ASTM D
5045–99. The ASTM model is actually for compact tension test in a plane-strain situation;
however, it is adapted only for the determination of σQ (which is the stress at 5° off the tangent ofa stress-strain curve, refer to Figure 2.6 for the determination of

KQ(which is aconditional K1max ). The value  Q is imputed into equation (2.9) above for the calculation of

KQ  K1max , thus:

2
3
4
1 


a
a
a
a 
KQ   Q   a  2  1.12  0.23    10.6    21.7    30.4   
 w
 w
 w
 w 


32(2.11)

According to the standard if σQ > σmax then σmax is to be used to calculate KQ otherwise σQ isused.

2.6.3 Calculation of Interfacial Shear Strength of GFPP Composites
In this investigation, the determination of the strength of the interface between glass fiber andmatrix resin is considered very important in order to have a good understanding of the fracture
mechanism of GFPP composites vis-à-vis the de-bonding and fiber pullout phenomenon. It will
equally help to form the knowledge on how the interface contributes to the overall mechanicaland fracture properties of the composites through an effective and efficient stress transfer
between the matrix and the fibers. It is known that a high interfacial shear strength (andshorter critical fiber length (lc) characterize a superior interface in a fiber reinforced composite.
The theoretical strength of a short glass fiber reinforced composite is given by the following
equation: τ lVlσ cu = C   i2ri i +  σ fuV j 1 - 2lc  + σ m 1 -V f
j
 i

j





(2.12)

Where  cu = the ultimate strength of the composite, σ fu = the ultimate strength of the fiber.
 = the matrix stress at failure,  i = interfacial shear strength, C = the fiber orientationmefficiency. lc = the critical fiber length, li and l j are the sub-critical and super-critical fiber
lengths, respectively. Vi and V j are the fiber volume fractions of the sub-critical and super-33

critical fiber lengths, respectively. V f = the overall fiber volume fraction, and r is the fiber
radius.

The critical fiber length ( lc ) and the interfacial shear strength (  ), respectively, are calculated
from equations (2.13) and (2.14) below:lc  4l
3(2.13)

r σ
f fu
τ=
lc

(2.14)The interfacial shear strength was calculated with software which module was based onequations (2.12) – (2.14).

2.6.4 Mechanical and Fracture Properties of GFPP Composites
Some related works have been done to investigate the properties of GFPP composites and many
findings were made. A study was conducted on the fracture behavior of glass fiber reinforced
composites with single-edge crack under tensile loading by means of the traditional stressmeasurement method combined with digital speckle correlation method [22]. By this system, thehorizontal and vertical displacement fields of the crack tip in different loading processes were
measured. A three-element strain gauge was also used in the investigation to measure the strain
responses near the crack during the whole damage and fracture process of the single-edge crack
specimen. The digital speckle correlation method is confirmed to be an effective way to study the34micro deformation of the composite materials. The effect of macromolecular coupling agent on
the property of GF/PP composites was studied [23]. Silane was employed as a coupling agent
between GF and PP in order to improve their interfacial adhesion, which is essential for thetransfer of stresses from matrix to fibers thereby improving mechanical properties of thecomposite. Results agree that the coupling agent actually improved the mechanical properties of
the composites so treated as against the untreated composites. According to the study, there wasan increase in the tensile modulus of the GF/PP composites, which resulted from the higher
tensile modulus of GF. Hence, the mechanical properties of the composites were improved.
Investigation was conducted on the failure behavior and the damage zone growth in a glass fiber
mat-reinforced thermoplastic polypropylene (GMT-PP) on a static tensile loaded single-edge
notched (SEN-T) and compact-tension (CT) specimens [24]. Different techniques, such as lightmicroscopy (LM), acoustic emission (AE) and infrared thermography (IT) were used. It wasestablished by AE that the failure mode, and especially the onset of fiber fracture, strongly
depends on the fiber content (Vf) of the GMT-PP. The size of the damage zone was considerably
underestimated when deduced from the stress-whitened zone in the LM pictures taken during theloading of the specimens (diameter = 5 mm). Location of the AE by a four sensors array resulted
in much higher values (diameter ~ 30mm) being practically unaffected by Vf in the studied range(Vf ~ 20 to 30 vol. %). The damage zone size derived from IT lay between those of LM and AE(diameter ~ 11-16 mm). It was assumed that the whole damage zone (encompassing a process
and a dissipation part) is assessed by AE during the whole fracture process, whereas LM and IT
reflect mostly its process zone constituent in the crack initiation and propagation stages,
respectively.

35

Effects of reinforcement on the mechanical and fracture properties of injection molded shortglass fiber (SGF)/short carbon fiber (SCF) reinforced polyamide 6 (PA 6) hybrid composites
were investigated [25]. The short fiber composites of PA 6 glass fiber, carbon fiber, and the
hybrid blend were injection molded using a conventional machine whereas the two types ofsandwich skin–core hybrids were co-injection molded. The fiber volume fraction for all
formulations was fixed at 0.07. The carbon fiber/PA 6 (CF/PA 6) formulation showed the highest
values for most tests. The sandwich skin-core hybrid composites exhibited values lower than the
CF/PA 6 and hybrid composite blends for the mechanical and fracture tests. CF/PA 6 had thehighest values for each mechanical and fracture test due to the higher strength and stiffness ofcarbon fibers compared to glass fibers. There was better adhesion between carbon fibers and PA
6 compared to glass fibers due to surface roughness. DSC results confirmed that carbon fibers
also had better nucleating abilities compared to glass fibers. The skin–core sandwich composites
exhibited only moderate values for each mechanical and fracture test. However, the CFs/GFc/PA
6 composites exhibit higher mechanical properties compared to GFs/CFc/PA 6 composites.
Understanding the parameters that regulate the mechanical and interfacial properties ofcomposites was the focus of another study [26]. Grafted PP and silane coupling agent were
employed as adhesion promoters. The study showed that the addition of grafted PP increased the
crystallization temperatures; and that the increase was more important in the presence of glass
fibers. Further, it was revealed that PP-g-AA bearing long AA branching showed an increased
secondary crystallization temperature interval, smaller crystalline morphology, and non-welldefined crystalline entities compared to the net spherulitic structures of the pure PP composite.
Therefore, the work showed that the interface between the matrix and fibers of composites was
one of the main parameters controlling the mechanical properties of composites. The results

36

showed that if the addition of an adhesion promoter might be necessary, the level of grafting andto a lesser extent the migration of the additive play important roles in the enhancement of themechanical properties of the composite. Finally, appropriate sizing can promote the adhesion of
the matrix to the fibers in two ways: by physical interactions and by chemical reactions between
the substrates. In understanding the role of glass fibers and fiber-matrix adhesion on deformation
mechanism, Composites of different fiber contents, with or without adhesion promoter were
investigated [27]. Results of the study showed that the introduction of glass fibers shifts thedeformation mechanism from shear yielding to crazing in contrast to pure polypropylene, whichshows shear yielding under a static loading condition. It was also shown that at a weak interface,
the craze deviated towards the interface, and at strong interface, the craze grew in its original
direction, causing breaking of fibers. Finally, the results revealed that craze thickening occurredbecause of craze-fiber interaction and concluded that there was no influence of interfacialadhesion on this phenomenon.

Investigation was conducted on the fracture toughness of Polymers in the ductile-to-brittle
transition region with a view to treating the scatter and nonlinearity effects in fracture toughnessof polymers in the region [28]. A statistical analysis was performed on large toughness data setof PPH and PPH/POes that, under the given test conditions, presented large scatter in the fracture
toughness values. The experimental point’s distribution tended toward a toughness thresholdvalue, J0 > 0. 3P-Weibull model proved adequate to describe the PPH toughness behavior over
the whole scatter range, whereas for PPH/POes blend, the experimental points larger than J =16.0 KJ/m2 did not fit because these values were influenced by excessive crack growth and crack
tip plasticity. This limitation resulted in adequate fracture toughness threshold-value, J0, both for37the fracture toughness of polymers 3683 analyzed materials. In a study on the experimental
Analysis and Computational Modeling of Damage and Fracture, the failure analysis of shortglass fiber reinforced polypropylene (SGFPP) composite was investigated by using opticalparticle tracking method from Hentschel [29]. The Compact-Tension (CT) test used for thisanalysis was characterized by the presence of high deformation gradients in front of the notch of
the specimen. The experiments were carried out on specimens of different sizes with different
notch depths. The study revealed that the rate-sensitivity of the SGFPP was negligible addingthat the dominant failure mechanisms for this material were fiber-matrix de-bonding, matrix
cracking and fiber pullout. The study reasoned that the non-local nature of these failure modes
was obvious, which provided the necessary experimental basis for the use of an internal lengthparameter in regularized numerical models. The local CT-results permit assessment of non-local
numerical models in order to fit the physical reality on the local and on the global level. In an
investigation on the microscopic damage mechanism in GFPP, the damage mechanisms in two
structurally different glass mat reinforced polypropylene materials were studied [30].
Micrographs of the damage processes in the two materials were presented and the study revealed
that the major points of damage initiation were transversely oriented fibers and fiber bundles. In
the swirled mat material, cracks grew along the fiber bundles; crack formation and growth was
relatively unaffected by macroscopical stress concentration. In the short fiber material, crack
growth occurred at the notch. Damage initiation occurred at fiber bundles or single fibers
oriented transverse to the direction of loading. For the short fiber material, the absence of fiber
bundles provided no inherent initiation points for crack formation. During crack growth, asmaller material volume was active in energy dissipation in the short fiber material, partly
because of the smaller resistance to crack growth from single fibers. The effect of pretreatment38

of fiber with NaOH and coupling with maleated polypropylene (MAPP) on the physical andmechanical properties of polypropylene reinforced with radiate pine fiber was investigated [31].
The study revealed that there was an increase of 106% in tensile strength and 302% in Young’s
modulus obtained by using 60 wt% fiber and 2 wt% MAPP when compared with un-reinforced
polypropylene. However, for the NaOH pretreated composite there was a reduced strength andan improved modulus by 370%. In a related study, focus was on the effects of various surface
treatments and different compounding and processing methods on the mechanical properties ofshort glass fiber (GF) reinforced Polypropylene (PP) injection molded parts; and on the fracture
toughness and fatigue behavior of such composite parts [32]. The study showed thatpolypropylene-based binder forms a superior interfacial bonding with the matrix polypropylene
and thus have a superior tensile strength with greater ductility, which leads to higher fracture
toughness. It also has a better long-term fatigue behavior under pre-cracked conditions. The
compounding method was shown to have strong influence on the composites. Among othermethods, their results showed that the Buss Kneader extruder was a superior method of extrusionin that it caused less damage to fibers, thus leading to better interfacial bonding than otherextrusion methods. This results to improved superior tensile and fracture values of composites
extruded from the Buss Kneader.

The thermal behavior and dynamic mechanical properties of isotactic polypropylene (PP) andreactor blend PP/ethylene–propylene copolymer (EPM) which was reinforced with different
amounts of short glass fibers (GF) and/or polyester fibers (PETF) were examined by differential
scanning calorimetry (DSC) and dynamic mechanical thermo analysis (DMTA) of imposed
tensile load on rectangular film specimens [33]. DMTA spectra revealed an increase in the39

stiffness and a decrease of the damping with increasing GF content. The variation of frequency
affected the secondary relaxations considerably and, in the presence of GF, the glass transitions.
With increasing GF content the E′ modulus increases and tan θ decreases in a regular manner.
The positions of the relaxations in the amorphous regions do not change but the relaxation in the
crystalline phase is strongly affected. The presence of PETF enhances the damping values of PP
at low temperatures. An evaluation of the effect of molecular weight and fiber diameter on theinterfacial behavior in GFPP examined the influences of the surface state of the fiber and theheat-treatment condition on the interfacial morphology and the spherulitic formation process inthe matrix [34]. Consequently, it was found that both the fiber diameter and molecular weight ofthe polymer significantly influence the thickness of the transcrystalline layer. In addition, as theinterfiber spacing becomes smaller, the spherulites in the matrix polymer were not seen to form
between the transcrystalline layers developed on the glass-fiber surface. In addition, the radius of
the largest spherulites in the matrix polymer was found to be about the same as the thickness oftranscrystalline region and largely depend on the holding time at the crystallization temperatureand cooling condition (or rate). The effect of processing temperature on the interfacial adhesion,mechanical properties and thermal stability of bio-flour-filled, polypropylene (PP) composites asa function of five different maleic anhydride-grafted PP (MAPP) types were evaluated [35]. In
order to investigate the effect on the interfacial adhesion of the composites, the five MAPP types
were subjected to characterization tests. The MAPP-treated composites with sufficient molecularweight and maleic anhydride (MA) graft (%) showed improved mechanical and thermal stability.
The enhanced interfacial adhesion and mechanical and thermal stability of the MAPP-treatedcomposites was strongly dependent on the amount of MA graft (%) and the MAPP molecularweight. The morphological properties of the MAPP-treated composites showed strong bonding40and a paucity of pulled-out traces from the matrix in the two phases. In addition, the improved
interfacial adhesion of the MAPP-treated composites was confirmed by spectral analysis of thechemical structure using attenuated total reflectance (FTIR-ATR).

Bio-fiber reinforced (lignocellulosic) polypropylene composites, which were generated by a
combination of fiber modification and matrix modification methods and prepared by reactive
extrusion processing, were examined [36]. In the investigation, PP matrix was modified byreacting with maleic anhydride while the fiber surface was modified by reacting it with a silane
in aqueous reaction system. The modified fibers and modified matrix were extruded andinjection-molded. The results of the study revealed that the resulting composite had improved
mechanical properties in toughness, tensile and impact strength due to improved adhesion
resulting from reactions and enhanced polar interactions at phase boundaries. The study
concluded that it was possible to enhance the properties of bio-fiber reinforced composites
through functionalization of the polymer matrix with MA and fiber surface modification. Addingthat composites based on the modified matrix had, in general, superior mechanical properties tothose containing the unmodified matrix, primarily because of improved adhesion and enhanced
polar interactions at the fiber/matrix interface. Effect of silane coupling agent on the morphology
and fracture performance of particulates-filled polypropylene was investigated [37]. The onset ofthe crystallization at the filler-matrix interface was observed by optical microscopy. This
analysis showed that the silane alters the matrix morphology by inducing nucleation at the fillermatrix interface. The fracture performance of the composites was analyzed by the post-yieldfracture mechanics approach. The effect of coupling agent on the fracture performance wasfound to occur at a very low concentration, whereas subsequent increase in silane concentration

41

had no effect. Furthermore, the effect of coupling agent depended on concentration of filler.
Silane treatment had a detrimental effect at low concentration of filler. However, at highconcentration of filler, an improvement in the filler-matrix adhesion had a beneficial effect on
the fracture performance. Two phenomena, the morphological effect and the competition
between the contribution of the matrix-filler interface and the contribution of the matrix, were
found to be possible explanations for the variation of KRmax with filler concentration andadhesion at the matrix-filler interface. The combined effect of fiber content and microstructureon the fracture toughness of SGF and SCF reinforced polypropylene composites were studied,focusing on the combined effect of fiber volume fraction and microstructure (fiber length andalignment) on the fracture toughness of short fiber composites [38]. It was revealed that thefracture toughness of the composite kept almost unchanged in the transverse direction anddecreased in the longitudinal direction with increasing fiber volume fraction; concluding that the
major factor influencing fiber breakage is fiber-fiber interaction – especially where there arefiber bundles. With high volume fiber fraction, the interaction between fibers leads to moredamage to fiber length. Hence, fiber length decreases with increasing fiber content. SGFPP
composites show a three-layer structure. Failure mechanisms, which contribute to the overallcomposite toughness, include fiber fracture, interfacial de-bonding, fiber pullout and matrix
fracture. Note also that fibers that are parallel to the crack propagation direction do not produce
any fiber pull-out energy, therefore, it contributes much less than transverse and oblique fibers to
the composite fracture toughness. The skin layers contribute more to the composite fracture
toughness than the core layers. Compression and injection molding processes were carried out inorder to evaluate the better mixer method for fiber (sugarcane bagasse, bagasse cellulose andbenzylated bagasse) and matrix (polypropylene) [39]. The samples (composites and42polypropylene plates) were cut and submitted to mechanical tests in order to measure flexuraland tensile properties. The morphological and micro-structural analyses of fracture surface andspecimens from composites were evaluated by microscopic techniques. The fracture surface was
examined by SEM and selected specimens from composites were analyzed by reflected light in
OM. The better tested method for composites obtainment was the injection molding undervacuum process, by which composites were obtained with homogeneous distribution of fibers
and without blisters. The mechanical properties show that the composites did not have goodadhesion between fiber and matrix; on the other hand, the fiber insertion improved the flexural
modulus and the material rigidity.

The fabrication of kenaf fiber reinforced polypropylene sheets, which could be thermoformed fora wide variety of applications with properties that are comparable to existing synthetic
composites, was studied [40]. This study has proven the ability to fabricate kenaf-polypropylene
natural fiber composites into sheet form. The optimal fabrication method for these materials wasdetermined to be a compression molding process utilizing a layered sifting of a microfine
polypropylene powder and chopped kenaf fibers. The use of a coupling agent, 3% Epolene
enabled successful fiber-matrix adhesion. The kenaf-PP composites compression molding in thisstudy proved to have superior tensile and flexural strength when compared to other compression
molded natural fiber composites such as other kenaf, sisal, and coir-reinforced thermoplastics.
The kenaf–maleated polypropylene composites manufactured in this study have a higher
Modulus/Cost and a higher specific modulus than sisal, coir, and even E-glass thereby providing
an opportunity for replacing existing materials with a higher strength, lower cost alternative thatis environmentally friendly.

43

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Figure 2.1: Mode I Fracture48

Figure 2.2: Mode II Fracture

Figure 2: Mode II Fracture49

Figure 2.3: Mode III Fracture

Figure 3: Mode III Fracture

50

in

Plane
Stress

um

M

ic
k

Mixed Mode

i
ed

Th

Fracture Toughness, KC

Th

K1C

Plane Strain

Thickness, B
Figure2.4:4:Geometry
conditions
for fracture
toughness
FigureGeometry Conditions
for Fracture
Toughness
51

WidthSpan length (115mm)
‘a’
Notch

‘B’ Thickness

Figure 2.5:
Dumbbell specimen for tensile test
Figure 3.2: Dumbbell specimen for tensile test
52

σQ

25

σmax

Stress (MPa)20

15

Failure Stress
10

5

0
0

0.002

0.004

0.006
Strain0.008

Figure 2.6: Determination of σQ at 5O off tangent

0.01

0.012

53

Chapter 3
Fracture Modes and Notch Sensitivity of V-notched V-PET and R-PET
Injection-Moldings

Abstract
The development of an anisotropic skin-core structure in injection moldings is wellacknowledged. The role of the skin and core regions in controlling the fracture behavior of Vnotched PET injection moldings during tensile and impact loadings was investigated. Results
revealed that with increasing notch depth there were three distinct fracture behaviors namely
ductile, semi ductile and brittle fracture transitions. The notch sensitivity factor for strength(kS) in the ductile and semi ductile transitions indicates that the fracture strength was notsensitive to ≤ 1.5mm deep notches, which is considered the skin region. The introduction ofcore-deep notches (˃ 1.5mm) resulted in a rapid increase in kS. On the other hand, the notch
sensitivity factor for energy (kT) shows that the fracture energy was not sensitive at ≤ 0.5mm
deep notches. However, kT increased drastically when notches ˃ 0.5mm deep wereintroduced. Notably, there was a drastic change in the fracture pattern from ductile to semi
ductile at a critical 0.6mm-deep notch. A constant semi ductile fracture behavior between 0.61.0mm-deep notches was then noted. The specimens experienced a mixed fracture behavior
at 1.5mm-deep notch, which marks a transitional fracture pattern at the interface between theskin and core regions. Lastly, brittle fracture behavior was observed at ≥ 1.5mm-deep notches.
Results show that crack opening, in the samples that had semi ductile fracture, was a post
necking phenomenon. Prior to shear yielding, two shear lines that intersected at an angle of
45⁰ were seen to originate from the crack root when a 1.2kN load was applied. Conversely,
crack opening and failure occurred simultaneously in brittle fractures. It is obvious that Vnotches provided a gradual transition in fracture behavior from the skin to the core regions,54

which confirms that the fracture behavior of PET injection moldings can be dependent on
their morphological characteristics.

3.1 Introduction
Fractures occur in materials due to propagation of pre-existing cavities or nucleation of voids
with the presence of mechanical loading. Such defects in engineering thermoplastic
components are stress concentrators [1, 2] that tend to undermine the structural stability ofsuch components during their life span. The resultant effects of cavities, voids or defects inengineering structures so often can be catastrophic. Catastrophic failure of structures claims
heavily in economic and human cost. The study of fracture mechanics therefore is meant toinvestigate the materials’ response to fracture under a pre-determined crack condition.
Understanding the fracture behavior of materials with different sizes and depths of cracks
helps components designers and engineers to balance between the choice of discarding very
expensive defective material and safety requirements of engineering construction.
Polyethylene terephthalate (PET) is a semi-crystalline thermoplastic material known to bevery notch sensitive. The notch sensitivity of PET products had led to many research
investigation which reveals that both virgin (V-PET) and recycled PET (R-PET) products
exhibit low fracture resistance in the presence of notch. Takano and Nielson [3] examined theeffect of standard V-notches on polymeric materials including PET by conducting tensile teston dumbbell test bars with single and double edge notches of 1.27 and 3.175mm deep.
Results show that if a notch has no effect on the toughness of a polymeric material, the notch
sensitivity factor is 1.0, otherwise it will be greater than 1.0. Their investigation shows thatthe values of kS for PET at 1.27mm and 3.175mm notch depths were 1.02 and 1.26,
respectively; while kT at the two notch depths were 7.33 and 11.21. These results therefore,55

prove that notches ≥1.27mm are detrimental to PET products as they fracture in a brittle
manner.

In recent years, there is much effort to improve the mechanical properties of PET products asdemonstrated in the work of Tanrattanakul, et al [4]. They investigated the improvement in
the toughness of different blends of PET by conducting tensile test on standard 2.54mm deepnotched Izod bars at a temperature range of -20 – 55oC. Their work showed that the PET
samples blended with un-functionalized styrene-butadiene-styrene (SEBS) yielded at thecrack root thereby fracturing in a brittle manner. On the other hand, blending PET with small
amounts of functionalized SEBS elastomers was effective in increasing the ductility(toughness) of PET products. The fracture toughness of a material is known to be dependent
on its morphological structure and crystallinity [5, 6]. Therefore, Stearne and Ward [6]investigated the effect of molecular weight and crystallinity on the notch sensitivity of PET
injection moldings. They obtained two samples with varying degrees of crystallinity bycontrolling the cooling time and mold temperature during sample preparation. Tensile tests
were conducted on the dumbbell samples that were notched at depths of either 1.91 or
2.79mm. Their investigation revealed that crystallinity affects both the fracture strength andyield behavior of PET injection moldings. At crystallinity of less than 35%, there was little
influence on the fracture strength, but the yield stress increased. However, at 44%
crystallinity the fracture strength was greatly reduced. Their results also indicate that notchedsamples had higher yield stress in the cold-drawing region than the un-notched when tested attemperatures above 5⁰C.

The mentioned studies would typically include notch depths of 1.91, 2.54 or 2.79mm, which
are standard notch depths stipulated in ASTM D256. Since injection moldings are known to56exhibit anisotropic skin-core morphological behavior, these notches are usually deep enoughto penetrate through the skin region (which is typically 1.5mm deep). As such, theintroduction of standard notches does not provide adequate consideration for the influence ofthe skin region on the fracture characteristics of the bulk material. Furthermore, in slow
crystallizing materials such as poly(ethylene terephthalate) (PET), the development of a
distinct skin-core structure during injection molding is inevitable. PET is known also to be avery notch-sensitive material and would fracture in a very brittle manner upon introduction ofnotches but remains very tough if un-notched. However, the effect of notch depth on thefracture characteristics is still unclear. Therefore, it is necessary to introduce notches atvarious depths and examine their effects on the crack propagation behavior of injection
molded PET products. If a correlation between fracture characteristics and morphology could
be established, the information will be very useful to assist molders in optimizing molding
conditions as well as improve the safety and reliability of the moldings.

In order to understand these effects, V-notches ranging from 0.5 - 4.5 mm depths wereintroduced on a single edge of dumbbell and Izod impact samples. These depths were chosen
so that the notch root would be positioned at regions ranging from the skin to the core of themoldings. The depths of the skin and core regions were determined through polarized optical
microscopy. The fracture characteristics of the samples at various notch depths weremonitored, during tensile and impact tests, by various optical devices including a CCD
camera.

57

3.2 Characterization methods

3.2.1 Material and Sample Preparation
Amorphous grade neat PET (V-PET) pellets (MA2103 LOT 601K; Mw = 23,000) wereobtained from UNITIKA, whereas recycled PET (R-PET) flakes were used as received from
Yasuda-Sangyo, Co. Ltd. R-PET flakes were dried at 120⁰C for at least 2 hrs and extruded
by a set of single-screw extruder (SR-Ruder Bambi SRV-P70/62 from Nihon Yuki, Japan) to
produce pellets. Barrel temperature of the extruder was set between 255 and 290⁰C, andscrew rotation speed was 50 rpm. Before injection molding, the V-PET and R-PET pellets
were dried for at least 4 hrs at 130⁰C, in a PICCOLO hopper-dryer from Itswa Co. Ltd. The
fabrication of dumbbell samples was performed with a TOYO PSS TI-30F6 injectionmolding machine at a barrel temperature range of 250–270⁰C. Mold temperature was set at
30⁰C, while the injection and holding pressures were set at 60 kgf.

3.2.2 Notching and Mechanical Property Characterization
V-notches of 0.5–4.5 mm depths were introduced on dumbbell and Izod impact test pieces
with a Tool A-3 Digital Notching Machine (from Toyo Seiki Manufacturing). The notching
was performed in stages of 0.5 mm depth to minimize internal deformation of the samples.
The angle of the V-notches was 45o in accordance with ASTM D256.

Tensile testing of notched and un-notched dumbbell specimens was performed with an
Instron 4466 universal testing machine mounted with 10kN load cell and set at a crosshead
speed of 10mm/min. At least seven sample pieces were tested for each material and notchingcondition. New sets of samples bearing the notch depths of 0.6, 1.5 and 2.0mm respectively
were loaded to intermediate points on the stress-strain curves and stopped. These points
58

include 1.2, 1.4 and 1.6kN. The 1.2kN was determined based on 75% of the yield stress of
un-notched samples. Subsequent intermediate load points were increased by 12.5% from thepreceding load. The samples were carefully unloaded, polished and observed in a polarized
optical microscopy to determine the crack behavior at these points.

As the test progresses, CCD camera was used to monitor the notch tip during crack opening
and propagation. The crack propagation behavior of V-PET and R-PET during fracture could
provide an account of the toughness of the bulk material. This investigation was conductedby observing the crack propagation of the materials under tensile test with a CCD Camera
mounted with a high-resolution optical microscope.

Izod impact testing was conducted on samples introduced with 0.5 - 4.5mm deep notcheswith a TOYOSEIKIIDigital Impact Tester machine fitted with a standardized JIS K7110
hammer bearing impact energy of 2.75J.

Scanning electron microscopy (SEM) was used to observe the fracture surfaces of tensile andIzod impact specimens. The fracture surfaces were cut to size, mounted on sample stands,
and spotted with gold coatings. Electron scanning was conducted at a magnification of 200
with Model S-3000N Scanning Electron Microscope, from Hitachi Co. Ltd. The component
images were joined together by the use of Adobe Photoshop CS image processing software.

3.2.3 Determination of Notch Sensitivity Factors
Investigations were conducted on the effects of V-notches by analyzing the notch sensitivity
factors kS and kT at the corresponding notch depths. This gives an understanding of theeffects of notches on the fracture behavior of the materials. If notches have no detrimental

59effects on the tensile strength and total energy absorption upon deformation, the notch
sensitivity factors for the strength (kS) and energy (kT) are equal to 1.0. Conversely, if they
have detrimental effects, the notch sensitivity factors are greater than 1.0 [3]. In this study,
the notch sensitivity factors are determined by the following equations:

 (YS ) [t (w  a )] i 0(YSi )  (tw) 




 
S

A

T   0





(3.1)



ssc  t (w  ai )  
Aissc  tw 

(3.2)



Where kS = notch sensitivity factor for yield strength, kT = notch sensitivity factor for energy.
YS0 and YSi = yield strength of un-notched and notched samples respectively [subscript i is
the successive notch depths], t = thickness, w = width, and a = notch depth. A0ssc and Aissc are
the areas under the tensile stress-strain curves for un-notched and notched specimens
respectively [3].

The notch sensitivity factor for yield strength gives an understanding of how notches affectthe strength of materials. On the other hand, the notch sensitivity factor for energy measuresthe effect of notches on the toughness of the materials.

3.2.4 Measurement of Crystallinity in PET
Sample measuring about 5.6mg was extracted from across the middle portion of the crosssection of dumbbell samples as shown in Figure 3.1. DSC was conducted with a TA
Instrument model 2920 modulated differential scanning analyzer. The sample was heated60

from 30 to 300⁰C at 10⁰C/min under nitrogen atmosphere. The enthalpies due to cold
crystallization [EC] and melting [EM] were measured from the resulting DSC enthalpy curves.
The crystallinity of the materials due to processing [EP] was then calculated by [EP = EM –EC] as shown in Table 3.1.

3.2.5 Birefringence Observation of PET Injection Moldings
Birefringence observation was performed on the cross-section of dumbbell samples to gauge
the depths of the skin and core regions. Prior to polarized-optical-microscopy, a short sectionof dumbbell specimen was cut and embedded in epoxy. One side of the cross-section waspolished with a rotational polishing wheel mounted with abrasives of different grain sizes toachieve smooth surfaces. The surface was further made as smooth and clean as possible by
the use of alumina particles. Thereafter, the polished side was glued to a clear and clean slide
glass with araldite glue to hold it in place; and ensure that liquid does not penetrate inbetween the glass and sample surface during the polishing of the opposite end of thespecimen. The polishing steps described above were repeated for the opposite side of thespecimen until a thin and translucent film of about 30m thickness was obtained, as shown in
Figure 3.2. Polarized optical micrographs were taken from the polished specimen through a
3.34 mega pixel Nikon digital camera mounted on a Nikon ECLIPSE E600 Polarizer. The
micrographs were combined by the use of Adobe Photoshop CS software. The resulting
composite image was converted to gray mode with Origin 7J software by OriginLab
Corporation; the plots of the gray values over the entire image were taken and analyzed.

The tensile test samples that were loaded midway to 1.2, 1.4, and 1.6kN, equally, wereobserved for fracture initiation. The width view of the portion containing the notches on thetest pieces were cut and embedded in epoxy. One side of the width view was polished with61

rotational polishing wheel as described previously in the preceding paragraph. The polishedview was glued to a clean and clear slide glass and kept to cure. The polishing steps
described above were repeated for the opposite width view of the samples until a thin smoothand translucent view of about 30m thickness is achieved. Polarized optical micrographs
were taken from the polished specimen and subsequently observed for evidence of possible
crack opening at the three intermediate load points.

3.3 Results and Discussions

3.3.1 Cross-Sectional Observation of V-PET by Polarized Light Microscopy
The polarized optical micrograph of the cross-section of V-PET specimen is shown in Figure
3.3. Distinct skin and core regions could be recognized from the birefringence and contrast of
the micrograph. The figure was then converted to grayscale and subsequently contour plots
defining the intensity of the gray areas are depicted in Figures 3.4 and 3.5, respectively. The
contour plots revealed a more complex morphology where the existence of an interface
between the skin and the core is clearly visible in the material.

Derivations from the intensity of the gray areas were used to determine precisely the
thickness of the skin and core regions, as shown in Figure 3.5. The plots marked with W andX represent the intensity values of the gray scale in the width and thickness directions of thesamples, respectively. The horizontal and vertical cursors on the gray scale image were
adjusted along the width and thickness directions until well-defined frequency peaks wereobtained. The center region in between the two maximum peak intensities (marked by arrows
in Figure 3.5) obtained from W and X plots were taken to be the core width and thickness,
respectively. The remaining areas surrounding the core will be regarded as the skin. The

62

thickness values of the skin and core were obtained from the W and X plots with data points
and screen readers provided in the Origin 7J processing software, which are presented inTable 3.2. Since notches were introduced in the width direction, the thickness of the skinregion in this direction is noted to be about 1.5 mm on each side of the specimen.

3.3.2 Effect of V-notch depth on the fracture modes of V-PET
Figure 3.6 shows the typical stress/strain curves during tensile tests for V-PET specimens.
The curves depict the transitions in fracture behavior corresponding to various notch depths.
Results indicate that the un-notched samples maintained high ductility and did not fracture
below 300% strain. The introduction of a 0.6 mm deep notch onto the materials, however,resulted in a drastic reduction in the ductility of both materials. This marked a change in thefracture behavior of the materials from ductile to semi-ductile where shear yielding andtearing at the notch tip was evident with increasing strain as shown in Figure 3.7. Prior toshear yielding, two lines originated from the crack root at 45o angles [7]. Crack propagation
eventually occurred along one of these lines through which the samples ultimately failed.
V-PET exhibited semi-ductile behavior when notches between 0.6 - 1.5 mm deep wereintroduced while brittle failure was imminent when deeper notches were present (2.0 - 4.5mm) as shown in Figure 3.6. It is important to note that the transitions in fracture behavior
from ductile to semi-ductile to brittle were not gradual. A very shallow V-notch, i.e. less than
0.6 mm, would not effectively cause stress concentration, thus the material remained ductile
as in un-notched specimens. When the notch is deep enough to act as a stress concentration
point (0.6 mm ≤ a ≤ 1.5 mm), yielding would occur at the root of the notch. However, theonset of crack propagation would be slow since the notch root was still located within thehighly amorphous skin region that is 1.5 mm in thickness, as indicated in Table 2. At thisnotch level there would be sufficient mobility of the polymeric chain segments for plastic

63

flow to occur from the skin region towards the crack tip [7]. With the introduction of notches
deeper than 1.5 mm, which would have already penetrated the more crystallized core region,
brittle failure was imminent, as plastic flow from the skin region to the crack tip was notpossible.

3.3.3 Correlation between Notch Sensitivity Factor for Strength (kS) and Skin-Core
Morphology of V-PET
The presence of cracks in materials builds up stress concentration around the defect. It meansthat the stress acting on the defect is higher than any other part of the material. An increase instress results in increased intensity until a critical value where the material yields and crack
propagates. Therefore, notch sensitivity factor for strength (kS) gives a measure of thesensitivity of materials to cracks. Figure 3.8 shows the notch sensitivity factors for strength(kS) for V-PET at the various notch depths determined by equation (3.1). The notch
sensitivity factor for strength (kS) for the un-notched sample is 1.0; and the notch ranges 0.5 1.0 mm were slightly lower than 1.0. This indicates that the notches had no detrimental
effects on the strength of the materials if the notch depths remained ≤ 1.5 mm, which isconsidered to be within the skin region. However, the lower kS values in samples notchedbetween 0.5 and 1.0 mm was attributed to their higher yield stress as compared to the unnotched specimens. This could be the effect of the stress field at the notch root as it changes
from biaxial stress to tri-axial [3] thereby increasing the materials’ resistance to fracture.
ASTM D5045 provides that in the fracture toughness study of a material, the notch depth towidth ratio (a/w) would have to be 0.45 ≤ a/w ≤ 0.55 so that the material will have minimum
resistance to crack propagation. It is obvious that the notches between 0.6 - 1.5 mm fall shortof this standard provision since the width of the specimens was 10.0mm. Studies show thatthe strain rate at the crack tip is greater than the strain rate at any other region on the test64

piece [3]. This brings about the twisting of the test sample resulting in a tearing fracture. In
addition, results show that the yield stress of V-PET is approximately 60 MPa. This could be
due to higher level of crystallinity, which results in high modulus of resilience, as shown in
Table 3.1. High molecular weight could be another factor that results in high yield stress in
V-PET. It is evident that the notch depths of 0.6 - 1.0 mm, did not affect the yield stress ofthe material.

3.3.4 Correlation between Notch Sensitivity Factor for Energy (kT) and Skin-Core
Morphology of V-PET
The notch sensitivity factors for energy (kT) at the various notch depths calculated fromequation (3.2) are shown in Figure 3.9 for V-PET. It is seen clearly that when the notch was ≤0.5 mm deep, the kT value remained at 1.0, which indicates that the material still has highplastic deformation within these notches. However, when the notch depth was extended
further to 0.6 - 1.00 mm, which is still within the skin region, the kT values promptly
increased to about 6.0. This marks a transition from ductile to semi-ductile fracture behavior.
This shows that while kS remained stable at ≤ 1.5 mm deep notches, kT experienced lower
plastic deformation. It is clear that V-PET experienced a transition in fracture behavior fromductile to semi-ductile at notch depths of between 0.6 - 1.5 mm. However, brittle failure wasimminent when deeper notches of 2.0 mm and above were introduced. This is evident in the
successively higher kT recorded between 2.0 - 4.5 mm deep notches. As mentioned earlier,
three distinct transitions in fracture behavior exist, namely ductile, semi-ductile, and brittle
fracture behaviors. The values of kS has shown that the material will remain tough andmaintain their structural stability if the notch depth is ≤ 0.5 mm. On the other hand, notches
between 0.6 - 1.0 mm deep will be structurally unstable, while deeper notches ≥ 1.5 mm willcause catastrophic failure in the material.

65

3.3.5 The Role of the Skin-core Morphology on the Notch Sensitivity of V-PET at
Intermediate Loading Points during Tensile Test
Further investigations were conducted on the three distinct fracture transitions that wereidentified earlier in Figure 3.6 for samples tested at various notch conditions. This was toinvestigate if crack opening occurred prior to the necking or failure of specimens introducedwith 0.6, 1.5, and 2.0 mm deep notches. Figure 3.10 shows the three loading points 1, 2, and3 [1.2, 1.4, and 1.6kN, respectively], on the load-displacement curve for 0.6 mm deep
notched samples. The birefringence observation of the samples that were loaded up to points
1, 2 and 3 are shown in Figure 3.11[a, b, and c] respectively. It is seen from Figure 3.11(a)that two shear lines at angle of 45⁰ had already appeared at the crack root when a 1.2kN load
was applied as indicated by point 1. The loading of a new sample to 1.4kN, at point 2,
resulted in the extension of the two lines towards the opposite end of the sample width asshown in Figure 3.11(b). The lines continued to extend with higher load (1.6kN - point 3)
until they reach the width end of the samples as shown in Figure 3.11(c) where shear yielding
occurred. Figure 3.11[a, b, and c] show that there was no crack opening in the 0.6 mm deep
notched samples prior to shear yielding; this means that crack opening did not occur in thematerial until after shear banding. It is worthy to note that the sample was able to shear band
because the 0.6 mm deep notch was still located in the amorphous skin region.

However, the load-displacement curve of 1.5 mm deep notched specimens shown in Figure
3.12 shows that the material could sustain only 1.2kN load before necking and subsequent
failure. This means that the material failed below 1.4kN load when a 1.5 mm deep notch wasintroduced. Figure 3.13 shows the birefringence of the material. In the figure, what would
have been the two lines originating from the crack root at angle of 45⁰ became two straightlines in the direction of the pre-crack. It is evident that the material yielded at this notch
66depth, nonetheless brittle failure was imminent. Figure 3.13 clearly shows that the notch was
already at the interface between the skin and the core; hence, there was no further plastic flowof the skin towards the crack tip resulting in brittle failure as the crack penetrates the more
crystallized core region. This fracture transition marks the onset of brittle failure in thematerial when deeper notches were introduced, i.e. the 2.0 mm deep notched samples did not
attain 1.2kN load before brittle fracture.

3.3.6 SEM characterization of tensile fracture surfaces of V-PET
Figure 3.14 shows the fracture surface of tensile test samples of V-PET. It is evident that highplastic deformation occurred in the material at a notch depth of 0.6 mm. This is indicative ofductile failure. It is observed that during crack propagation there were micro voids nucleation
around the crack tip (indicated by arrows in Figure 3.7), which eventually coalesced along thecrack path thereby leaving lobe-like fibrils [indicated by broken line arrows in Figure 3.14].
The micro voids are clearly seen from the images of high magnification camera monitored
movie recorded during tensile testing [refer to Figure 3.7]. It is evident also that there wastapering in thickness towards the notch direction as the materials fractured in a semi-ductile
manner from the notch end of the specimen’s width to the opposite side as shown in Figure
3.14. This tapering in thickness is caused by a continuous plastic flow [9] from theamorphous skin layer with lower stress concentration towards areas of high stressconcentration, such as the notch root, which can only occur when the notch does not
penetrate through the skin, i.e. notch depth is less than 1.5 mm. In Figure 3.15, there was amixed failure (semi ductile to brittle fracture behavior) [9]. Similar tapering of specimen as
seen in Figure 3.14 also occurred in the 1.5 mm deep-notched samples. Nevertheless, it iseasy to observe two fracture phases as the notch approaches the interface between the skinand core regions. At the notch root, which contains the terminal part of the skin region, there67

is significant degree of plasticity. However, as the crack progresses through this interface into
the core region, crazing was observed within the core region, as indicated by broken line
arrows in Figure 3.15, which lead to semi brittle failure.

Figure 3.16 shows the brittle fracture surface of the materials, when notch depths ≥ 2.0 mm
were introduced. Cavities or voids were seen as circular voids (pointed to by broken line
arrows) at the roots of the notches but these were unable to coalesce and resist fracture
thereby forming craze that leads to eventual brittle fracture. It is obvious from Figure 3.16
that there was no tapering of thickness in the specimens when notches have alreadypenetrated the core region. When the notch was core-deep the plastic flow of the amorphous
skin region towards the notch tip was not possible, hence brittle fracture was imminent.

3.3.7 Effect of V-notches on the Izod Impact Fracture Behavior of V-PET
Figure 3.17 shows that V-PET had a gradual loss of strength when impacted upon as thenotches deepen. The un-notched and 0.5 mm-notched samples did not fracture upon impact.
However, V-PET experienced brittle fracture and recorded impact strengths of about 25kJ/m2
when a critical notch depth of 0.6 mm was introduced. When subsequent deeper notches ofbetween 0.6 - 1.0 mm were introduced, the specimens experienced a gradual but consistentdecline in impact resistance. This could be due to the gradual thinning of the amorphous skinregion with increasing notch depth, which reduces the effectiveness of the skin in suppressing
crack propagation during impact loading. When the notch approaches the boundary between
the skin and the core, i.e. the interface 1.5 ± 0.2 mm, there was a transitional change in theloss of impact strength. Minimum impact resistance in the material was observed when a 2.0mm deep notch was introduced. Further increments in notch depth resulted in a consistently
low impact strength.

68

The fracture behavior of the material under impact testing is seen to be different from its
behavior under tensile loading. Under impact testing, the impact is very sudden that plastic
flow from the skin to the crack tip was not possible; hence, the material exhibited brittle
failure, which was characterized with a gradual and consistent loss of impact strength.
However, under tensile loading three distinct fracture behaviors were revealed at different
notch depths. This is because while the notches remain at the amorphous skin regions, themolecular chains can slide past one another and realign thereby causing plastic flow [10]towards the crack tip.

3.3.8 SEM characterization of Izod impact fracture surfaces of V-PET
Figures 3.18[a, b] show the two opposite impact fracture surfaces of V-PET at the 0.6 mm
notch depth. The circular-shaped craze at the crack root indicates sudden disruption of voids
nucleation and coalescence due to impact. There is a similar fracture behavior in Figures 3.19
and 3.20 [a, b] as explained in Figure 3.18. The micrographs show similar brittle fractureeven as the notches extend deeper into the core region. However, the wavy fracture surface of
Figure 3.18[c] suggests that the 0.6 mm deep notched sample absorbed more impact energy
than the samples with deeper notches. Comparing Figure 3.18[c] with Figure 3.19[c] reveals
that the sample that was introduced with 1.5 mm deep notch has a less wavy fracture surface.
This means a gradual decline in the impact energy absorption resulting from the deepening ofnotches. This is consistent with the change in impact strength at 1.5 mm deep notch as shownin Figure 3.17. Further deepening of the notch to 2.0 mm creates a more flat fracture surface,which indicates little or no resistance to fracture by the samples. The impact fracture surfaces
do not show outstanding difference in the fracture behavior of V-PET at the different notch
depths as was observed in the tensile test samples. It can be concluded that due to sudden

69

impact on the material no plastic flow from the skin region towards the stress concentration
region was possible.

3.3.9 Summary and Observations on the Crack Propagation Behavior and Toughness of
V-notched V-PET Injection Moldings
This investigation has revealed that the fracture behavior of V-PET injection moldings
responded in transitional phases from the skin through the interphase to the core regions of
the samples. Thus, three distinct fracture transitions were clearly observed from the results
indicating that the skin and core morphology plays significant role in the fracture behavior of
V-PET injection molded components. The notch sensitivity factor (kS) shows that the fracture
strength of the material will drop drastically if the notch is deeper than 1.0 mm. On the other
hand, kT indicates that notches deeper than 0.5 mm have detrimental effects on the toughnessof the material. Tensile loading provided a transitional fracture behavior from ductile, semiductile to brittle failure of V-PET while impact testing resulted in their brittle failure andgradual loss of impact strength. By loading the specimens to intermediate points on the loaddisplacement curves, it has been shown that crack opening did not occur in the specimens
prior to yielding. Therefore, it could be concluded that crack opening was a post-necking
phenomenon in the specimens that failed in a semi-ductile manner.

3.3.10 Cross-sectional observation of R-PET by polarized light microscopy
Figure 3.21 shows the birefringence of R-PET. The process that defines the derivation of theimage was previously shown in Figure 3.1, and discussed in section 3.3.1 for V-PET. The
difference in the intensity between the birefringence of V-PET [Figure 3.3] and R-PET shows
not only different levels of skin-core boundary, but suggests different levels of crystallinity inthe materials. It could suggest also different levels of internal stress in the materials resulting70

from injection molding. Figure 3.21 suggests that there is no sharp boundary between the skinand core layers of R-PET. This could suggest also that crystallinity [refer to table 3.1] andinternal stresses are lower in R-PET than V-PET. The difference in birefringencecharacteristics equally can mean lower molecular chain orientation at the skin layers resulting
in a more amorphous skin region in R-PET than in V-PET. These suggested characteristics
tend to be favorable conditions for high modulus of toughness, which on the other handresults in low modulus of resilience. The contour image of R-PET in Figure 3.22 shows thenon-distinct boundaries between the skin, interface and core regions. Although the skin andcore regions still could be identified, the image shows a more dispersed and homogenous
amorphous bulk in R-PET than in V-PET. As previously described for V-PET in Figure 3.5
section 3.3.1, Figure 3.23 depicts the gray image plot of R-PET showing the determination of
the sizes of the skin-core regions. Plot W measures the skin-core sizes in the width direction,
while plot X measures the sizes in the thickness direction. The values of the skin-core sizes
are shown in Table 3.2. It can be seen that R-PET has the same skin-core sizes with V-PET
irrespective of the more boundaries between the regions as seen in Figure 3.5 for V-PET.

3.3.11 Effect of V-notches on the fracture modes of R-PET
Figure 3.24 is a collection of the typical stress-strain curves of R-PET at different notch
depths. It is obvious that R-PET has similar transitional fracture behavior with V-PET(section 3.3.2). Similar to V-PET, three distinct fracture transitions could be identified. These
include ductile fracture, which occurred in samples that has ≤ 0.5 mm-deep notches
introduced in them; semi ductile fracture occurred in samples that were cut with 0.5 ≥ a ≤ 1.5
mm deep notches; and brittle fracture where the depths of the notches are > 1.5 mm. Figure
3.7 represents the deformation behavior of V-PET and R-PET as monitored by a CCD
camera. As discussed previously for V-PET in section 3.3.2, R-PET equally had two slip

71lines (shear lines) originating from the notch root at an angle of 45⁰. Shear banding occurredon the two lines. Crack propagation and failure eventually occurred along one of the shear
lines. It could be concluded that R-PET and V-PET have similar fracture behavior; however,the cluster of the stress-strain curves for R-PET particularly at full yielding for ≤ 0.5, and 0.5≥ a ≤ 1.0 mm deep notches suggests a more homogenous fracture behavior in the bulk. This
could have resulted from the more dispersed and homogenous bulk with less distinct
boundaries between the skin and core regions as was observed previously.

Figure 3.6 reveals that V-PET has a gradual fracture transition as the notches deepen.
However, it is obvious that the transition from semi ductile to brittle failure appears to bevery sudden as suggested by the distance between stress-strain curves of ≤ 1.0 and 1.5 mm
deep notches on one hand; and the distance between 1.5 and > 1.5 mm deep notches on theother hand. This sudden change from semi ductile to brittle fracture could not have provided
enough warning that is necessary to prevent catastrophe. In contrast, R-PET provided
adequate transitional warning from ductile to semi ductile fracture and particularly from semi
ductile to brittle failure. This is evident in Figure 3.24 where there was a significant drop instress when 1.5 mm deep notch was introduced in R-PET. Equally important to observe isthe isolation of the stress-strain curve of the 1.5 mm interface-deep notched sample from thesemi ductile and brittle fracture sets of samples. Figure 3.24 show that R-PET has lower
stress values than V-PET. This could have resulted principally from lower crystallinity in RPET (Table 3.1) which does not favor high modulus of resilience [10, 11, 12], and could
result in higher notch sensitivity for strength (kS) in R-PET.

72

3.3.12 Correlation between notch sensitivity factors for strength (kS) and the skin-coremorphology of R-PET
Figure 3.25 shows the kS for R-PET compared with V-PET. It is clear from the figure that VPET and R-PET behaved in a similar manner at the different notch depths. The figure reveals
that kS for the two materials remained 1.0 when the notches are ≤ 1.00 mm deep. This meansthat at these notch depths the materials retained high yield stress. Nevertheless, there is aslight difference in the notch sensitivity of the V-PET and R-PET when the notches became
deeper than 1.0 mm considered the interface between the skin and core regions. The higher kS
in R-PET could be attributed to lower crystallinity [Table 3.1], otherwise more amorphousnature of the bulk of the material. It is known that low crystallinity/high-amorphous nature inpolymeric materials results in high modulus of toughness and low modulus of resilience;
conversely, high crystallinity results in high modulus of resilience and low modulus of
toughness [10, 11, 12,].

3.3.13 Correlation between Notch Sensitivity Factors for Energy (kT) and the Skin-Core
Morphology of R-PET
Figure 3.26 compares the notch sensitivity factors for energy (kT) of V-PET and R-PET. It
shows great similarity in kT for both materials. As previously discussed in section 3.3.4 for
V-PET, it is clear that the toughness of R-PET is equally not affected when the V-notches are≤ 0.5 mm deep. This is evident in the value of kT = 1.0 [1, 3]. However, when the notches
deepen to 0.6 ≥ a ≤ 1.0 mm there was semi ductile failure. This marks the first transition in
the fracture behavior of the materials within the skin region; and signifies the first major
effect of notches, where kT increased significantly to about 6.0 [1, 3]. The materials were able
to fracture in a semi ductile manner because there was significant portion of less stressed
amorphous skin region that permitted plastic flow to the crack tip, which was the stress73

concentration point in the materials. At 1.5 mm deep notch, a slight but visible transition isobserved as suggested by the value of kT. The 1.5 mm deep notch marks the interface
between the skin and the core regions, which were observed previously in Figures 3.22 –3.24. Further deepening of the notch (≥ 1.5 mm) into the core region resulted in a brittle
fracture with phenomenal high notch sensitivity factors for energy (kT) (above 12.0). The
different transitions in the fracture behaviors of the materials suggest that the fracture modes
depend on the skin-core structure of the materials.

3.3.14 The role of Skin-Core Morphology on the Fracture Modes of R-PET at
Intermediate Loading Points during Tensile Test
The fracture modes of V-PET and R-PET at intermediate loading points during tensile test
were investigated. The fracture modes were discussed previously in section 3.3.5. Figures
3.10 and 3.12 represent the load-displacement curves of the materials when they wereintroduced with 0.6 and 1.5 mm deep notches respectively. It is important to remark at thispoint that the samples, which were introduced with 2.0 mm deep notch fractured below
1.2kN load. Figure 3.10 shows that the 0.6 mm deep notched samples fractured above 1.6 kN
load, and this occurred after full ligament yielding. The three load points under investigation
include 1(1.2 kN), 2 (1.4 kN) and 3 (1.6 kN). Further deepening of the notch to 1.5 mm(interface depth) resulted in the fracture of the material below 1.4 kN load. Although thematerials had full ligament yielding prior to fracture, it is important to observe that there wasa drop in the yield stress of R-PET signifying a transition into the interfacial region. In
contrast to V-PET, the drop in yield stress at this transitional region suggests the lastsignificant warning prior to catastrophic brittle fracture in R-PET. Figure 3.27 [a – c] shows
the birefringence of R-PET at the three intermediate loading points of the 1.5 mm deep
notched specimens. Similar to Figure 3.11[a – c] for V-PET, two slip lines (shear lines)
74

intersecting at an angle of 45⁰ are seen to originate from the notch root. Subsequent increasein the load from 1.2 to 1.4, and 1.6 kN resulted in an extension of the shear lines towards the
opposite end of the specimen until shear banding and full ligament yielding occurred. The
two slip lines originating at an angle of 45⁰ from the notch root, at the different load points is
consistent with Figure 3.7 as monitored by the CCD camera. Figures 3.11 and 3.27 show
similarity in the fracture modes of both materials, nonetheless it is clear from Figures 3.11[b,c] and 3.27 [b, c] that V-PET did not have much plastic flow prior to fracture. This could be
because of its level of crystallinity as shown in Table 3.1, and long molecular chain
configuration. Conversely, R-PET had more plastic deformation than V-PET. This indicatesthat there is a high plastic flow from the areas of less stress concentration of the amorphous
skin region to the crack tip, which is considered the stress concentration point. The highplastic deformation suggests high modulus of toughness. This means that R-PET had moreresistance to fracture and absorbed more energy, particularly in the actual crack opening andfracture (i.e. the plastic deformation zone).

Figure 3.28 show that R-PET has similar fracture behavior as V-PET as was shown in Figure
3.13 when a 1.5 mm deep notch was introduced. Figures 3.13 and 3.28 reveal that the two
slip lines that originated from the notch root at 45⁰ angles at ≤ 1.5 mm deep notches became
two straight lines. This marks a transition from semi ductile to brittle fracture behavior.
Further increase in the depth of the notch resulted in brittle failure. This observation isconsistent with the transition in fracture behavior recorded earlier as the notches deepened tothe interface between the skin and the core regions. Results have shown similarities in thefracture behavior of V-PET and R-PET, particularly in the notch sensitivity factors for energy(kT). However, results suggest that R-PET had higher plastic deformation than V-PET,
suggesting higher modulus of toughness.
75

3.3.15 SEM characterization of tensile fracture surfaces of R-PET
Figure 3.29 shows the tensile fracture surface of samples of R-PET that were introduced with
0.6 mm deep notch. Similar to V-PET in Figure 3.14, and discussed in section 3.3.6 R-PET
equally had high plastic deformation at this notch depth. The fracture behavior of thematerials at the 0.6 mm deep notch is characterized by the nucleation and coalescence of
micro-voids; and tearing, which leaves fibrillations on the fracture surface. The suggested
coalescence of micro-voids was observed earlier in Figure 3.7b. The tapering of sample width
towards the opposite end of the crack is also prominent in R-PET as it was observed in VPET. This tapering is possible only if the notch depth was within the skin region (i.e. ≤ 1.5
mm deep notches); where there would be continuous plastic flow [9] from amorphous skinlayer with lower stress concentration to the crack tip, considered the stress concentration
point. These characteristics indicate significant ductility at the 0.6 mm notch depth.

In Figure 3.30, R-PET had similar mixed fracture (semi ductile to brittle fracture behavior)[9] with V-PET, when the notch penetrates the interface between the skin and core regions at
1.5 mm deep notch. Similar tapering of specimen as seen in Figure 3.29 also occurred in the
1.5 mm deep-notched samples. Nevertheless, it is easy to observe two fracture phases as thenotch approaches the interface between the skin and core regions. At the notch root, whichcontains the terminal part of the skin region, there is significant degree of plasticity.
However, as the crack progresses through this interface into the core region, crazing wasobserved within the core region, as indicated by broken line arrows in Figure 3.30, whichlead to semi brittle failure.

Figure 3.31 shows the brittle fracture surface of R-PET, when notches ≥ 2.0 mm deep wereintroduced. Similar to V-PET, cavities or voids were seen as circular voids (pointed to by76

broken line arrows) at the root of the notch but these were unable to coalesce and resist
fracture thereby forming craze that leads to eventual brittle fracture. It is obvious from Figure
3.31 that there was no tapering of thickness in the specimens when notches have alreadypenetrated the core region. When the notch was core-deep the plastic flow of the amorphous
skin region towards the notch tip was not possible, hence brittle fracture was imminent.
Figure 3.16 and 3.23 for V-PET and R-PET, respectively, indicate brittle fracture and show
micro-voids but it is important to observe that there was greater number of micro-voids in RPET than in V-PET. This greater number of micro-voids in an attempt to coalesce became amechanism for energy absorption prior to brittle fracture. This suggests that even at the core
deep notches, R-PET appeared to have more resistance to fracture than V-PET. Thisobservation is evident in Figures 3.6 and 3.24. The percentage values of strain for V-PET
appears to be clustered for the core-deep notched samples, while the percentage strain for RPET samples within the same region seemed separated and decreased gradually.

3.3.16 Effect of V-notches on the Izod Impact Fracture Behavior of R-PET
The effects of V-notches on the Izod impact properties of V-PET, as depicted in Figure 3.17,
was discussed previously in section 3.3.7. Figure 3.32 shows the impact behavior of R-PET
when notches of different depths were introduced. It is obvious that R-PET behaved in asimilar manner as V-PET when impacted upon at different notch depths. However, V-PET isseen to have higher impact strength than R-PET when 0.6 mm deep notch was introduced.
This result is consistent with the higher yield stress in the tensile samples, which resulted
from higher level of crystallinity in V-PET [Table 3.1]. Similar to V-PET, there was a
gradual and consistent loss of impact strength as the notch deepens to ≤ 1.0 mm. The
materials fractured in a brittle manner under impact at different notch depths, nevertheless,the transitional fracture behavior at 0.6, 1.4 – 1.6, and 2.0 mm deep notches was obvious.

77

This is consistent also with the tensile fracture behaviors of the materials at the same notch
depths, considered the interface between the skin and the core regions. When the notches
become core-deep, the materials fractured with the same impact strength irrespective of thenotch depth.

3.3.17 SEM Characterization Izod Impact Fracture Surfaces of R-PET
The two opposite sides of Izod impact fracture surfaces of R-PET representing the fracturetransitions are shown in Figures 3.33 – 3.35. The fracture surfaces are consistent with thebrittle fracture behaviors discussed previously, but the wavy fracture surface of Figure3.33(c) (impact test specimen) shows that the material absorbed more energy when it was cut
with 0.6 mm deep notch. The wavy pattern can be seen in Figure 3.33 [a, b] as solid chunks(lumps) indicated by broken line arrows. Similar to V-PET, the brittle fracture in R-PET waspreceded by crazing at the notch root, which is seen as circular/oval discs at the crack tip.
Figure 3.34 shows a more plane fracture surface at 1.5 mm deep notch. Figure 3.34 [a, b] hasan evidence of undeveloped lumps, which rather appeared as cavitations as indicated byarrows. The two opposite sides of the fracture surfaces equally show small grains spread
evenly across the entire fracture area, which indicates evenly distributed impact energy on thematerial with little or no resistance. The near zero impact resistance of R-PET is revealed in
the apparent smooth fracture surface shown in Figure 3.35. Unlike V-PET, there is noevidence of solid chunks or significant grains on the Izod impact fracture surface of R-PET.
Nonetheless, the brittle fracture was preceded equally by crazing.

78

3.3.18 Summary and Observations on the Crack Propagation Behavior and Toughnessof V-notched R-PET Injection Moldings
Results have shown that low level of crystallinity in R-PET resulted in resulted in highmodulus of toughness. This means that R-PET absorbed more energy in the plastic
deformation zone, particularly in the crack opening and fracture process. While V-PET
provided distinct morphological boundaries, the more dispersed phase (not well-definedboundaries) and near homogenous morphology of R-PET provided a good warning signal
prior to the catastrophic (brittle) failure of the material when deeper notches were introduced.
This characteristic was evident by the significant decrease in yield stress when 1.5 mm deep
notch was introduced. The interface-deep notch marked a transition from semi ductile tobrittle fracture. It is obvious that V-PET did not show such decrease in yield stress prior tosudden brittle fracture. Loading the notched materials to intermediate load points have shown
the there is higher plastic flow in R-PET than was observed in V-PET, which could result instrain-induced crystallinity on the skin layer that will eventually toughen the material andresist fracture.

3.4 Conclusion on the Crack Propagation Behavior and Toughness of V-notched
Injection Molded V-PET and R-PET Components
The role of the skin and core regions in controlling the effects of V-notches, on the fracture
behavior of PET injection moldings, was correlated with their tensile and impact properties.
Investigations revealed that there were three distinct fracture behaviors: ductile, semi ductile
and brittle fracture transitions. The notch sensitivity factor for strength (kS) in the ductile andsemi ductile transitions indicates that the fracture strength was not sensitive to ≤ 1.5mm deep
notches, which is considered the skin region. The introduction of core-deep notches (˃1.5mm) resulted in a rapid increase in kS. On the other hand, the notch sensitivity factor for79

energy (kT) shows that the fracture energy was not sensitive at ≤ 0.5mm deep notches.
However, kT increased drastically when notches ˃ 0.5mm deep were introduced. The
development of an anisotropic skin-core structure in injection moldings is well acknowledged.
This is revealed in a constant fracture behavior between 0.6-1.0mm deep notches. Notably,
there was a drastic change in the fracture pattern from ductile to semi ductile at a critical
0.6mm deep notch. The specimens experienced a mixed fracture behavior at 1.5mm deep
notch, which marks a transitional fracture pattern at the interface between the skin and coreregions. Lastly, a constant fracture behavior was observed at notch depths ≥ 1.5mm. Results
show that crack opening, in the samples that had semi ductile fracture, was a post necking
phenomenon. Prior to shear yielding, two shear lines that intersected at an angle of 45⁰ wereseen to originate from the crack root when a 1.2kN load was applied. Conversely, crack
opening and failure occurred simultaneously in brittle fractures. It is obvious that V-notches
provided a gradual transition in fracture behavior from the skin to the core regions, which
confirms that the fracture behavior of

PET injection moldings can be dependent on the skinand core structure.

References
1.

Ogazi-Onyemaechi, B. C., Leong, Y. W., & Hamada, H., Crack Propagation Behavior
and Toughness of V-Notched Polyethylene Terephthalate Injection Moldings, Journalof Applied Polymer Science, 116, 132-141 (2010)

2.

Broek, D., The Practical use of Fracture mechanics, pg. 23, Kluwer Academic
Publishers, Boston, 1989

3.

Takano, M., and Nielsen, L. E., Journal of Applied Polymer Science, 20, 2193 (1976)

4.

Tanrattanakul, V., Perkins, W. G., Massey, F. L., Moet, A., Hiltner, A., and Baer, E.,
Journal of Materials Science, 32, 4749 (1997)
80

5.

Kuriyama, T., Narisawa, I., Shiina, F., and Kotaki, M., Journal of Vinyl and Additive
Technology, 4, 164(1998)

6.

Stearne, J. M., and Ward, I. M., Journal of Materials Science, 4, 1088(1969)7.

MA, M., Vijayan, K., Hiltner, A., and Baer, E., Journal of Materials Science, 24,2687 (1989)

8.

Perkins, W.G., Polymer Engineering and Science, 39, 12, (1999)

9.

Mills, N. J., Walker, N., Journal of Materials Science, 15, 1832 (1980)

10.

Askeland, D. R., Phule, P. P., The Science and Engineering of Materials, 4th ed., pg.
669-719, Brooks/Cole, USA. 2003.

11.

Park, B.D. and Balatinecz, J. J., Polymer Composites, 18, 1, 79 (1997)

12.

Brown, N., and Ward, I. M., Journal of Materials Science, 18, 1405 (1983)

81

Table 3.1: The Crystallinity of V-PET and R-PET Injection moldings

Cold-crystallization (EC)EP = (EM – EC)

Melting (EM)Materials
Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

Temp.
(⁰C)

V-PET

19.62

133.7241.68

257.82

22.06

124.10

R-PET

24.30

129.6938.09

255.2713.79

125.5882

Table 3.2: Gray values from gray image plots of the cross-sectional view of dumbbell samples showing the sizes of the skin, interface
and core layers of V-PET and R-PET

Width direction (mm)Thickness direction (mm)Region of Specimen
V-PET

R-PET

V-PET

R-PET

Skin region

1.47 (x2)1.47 (x2)0.65 (x2)0.65 (x2)Interphase

0.53 x 2

0.53 x 2

0.11 (x2)0.11 (x2)Core region

5.98

5.98

1.46

1.46

Total

9.98

83

2.98

Figure 3.1: Arrows indicate the position where samples were extracted for DSC
experiments

84

Figure 3.2: Preparation and polishing of samples for birefringence observations

85

Figure 3.3: Birefringence observation of the cross-section of dumbbell Specimens showing the skin and core structures of
V-PET.

86

CoreInterface
Skin

Figure 3.4: Image contour plot showing the skin, interface and core regions of V-PET indicated witharrows.

87

W

X
Figure 3.5: Gray image plot showing the gray values of the skin, Interface and core regions of V-PET[ W = graph of gray values on the width direction, X = graph of gray values on the thickness direction].

88

Figure 3.6: Stress - Strain curves of V-PET showing the effects of V-notch of different
depths [{ }* notch depths of 2.0 - 4.5 mm].

89

Figure 3.7: Crack initiation at notch root showing : a, the development of two lines at an angle of 45⁰ prior to shear
banding; b, Shear banding and plastic flow due to stress and consequent semi ductile failure [arrow shows microcracks and eventual coalescence].

90

Figure 3.8: Notch sensitivity factors for strength kS for V-PET at the various notch depths.

91

Notch sensitivity factors for energy (kT)Figure 3.9: Notch sensitivity factors for energy (kT) for V-PET at the various notch depths.

92

Figure 3.10: Load-displacement curve of 0.6mm deep notch sample showing threeintermediate loading points [1, 2 and 3 = 1.2, 1.4 and 1.6kN respectively].

93

Figure 3.11: Birefringence of V-PET showing the formation and extension of two families of intersecting slip lines when a,
1.2kN; b, 1.4kN; and c, 1.6kN load is applied as shown by points 1, 2 and 3 on the Load/Displacement curve in Figure 3.10.

94Figure 3.12: Load-displacement curve of 1.5mm deep notch sample showing 1.2kN
loading at point 1.

95

Skin

Core

Figure 3.13: Birefringence of V-PET showing the formation of two straight crack lines prior to brittle fracture when [1=1.2kN] load is applied on a 1.5 mm deep notched sample as shown on the Load/Displacement curve in Figure 3.12.

96Figure 3.14: Tensile fracture surface of V-PET showing ductile failure with high plastic deformation
at 0.6 mm deep V-notch [white arrow indicates notch direction; vertical line is the end of notch].

97

Figure 3.15: Tensile fracture surface of V-PET showing semi ductile failure with crazing at
1.5 mm deep V-notch [white arrow indicates notch direction; vertical line is the end of notch].

98

Figure 3.16: Tensile fracture surface of V-PET brittle failure at 2.0 mm deep V-notch[white arrow indicates notch direction; vertical line is the end of notch].

99

Impact strength (kJ/m2)

25
V-PET

20
15

10
5
0
0

1

2
3
V-notch depth (mm)45Figure 3.17: Izod impact strength of V-PET showing gradual loss of strength atdifferent V-notch depths.

100

Figure 3.18: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof V-PET at 0.6 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

101

Figure 3.19: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof V-PET at 1.5 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

102

Figure 3.20: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof V-PET at 2.0 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

103

Figure 3.21: Birefringence observation of the cross-section of dumbbell Specimens showing the skin and core structuresof R-PET.

104

Core
Interface
Skin

Figure 3.22: Image contour plot showing the skin, interface and core regions of R-PET indicated witharrows.

105

W

X
Figure 3.23: Gray image plot showing the gray values of the skin, interface and core regions of R-PET[
W = graph of gray values on the width direction, X = graph of gray values on the thickness direction].

106

Figure 3.24: Stress - Strain curves of R-PET showing the effects of V-notchof different depths [{ }* notch depths of 2.0-4.5mm].

107

Figure 3.25: Notch sensitivity factors for strength (kS) for V-PET and R-PET
at the various notch depths.

108

Figure 3.26: Notch sensitivity factors for energy kT for the materials at thevarious notch depths.

109

Figure 3.27: Birefringence of R-PET showing the formation and extension of two families of intersecting slip lines when a,
1.2kN; b, 1.4kN and c, 1.6kN load is applied as shown by points 1, 2 and 3 on the Load/Displacement curve in Figure 3.10.

110Figure 3.28: Birefringence of R-PET showing the formation of two straight crack lines when [1= 1.2kN] load
is applied on 1.5mm deep notched sample as shown on the Load/Displacement curve in Figure 3.12.

111

Figure 3.29: Tensile fracture surface of R-PET showing semi ductile failure with crazing at
1.5 mm deep V-notch [white arrow indicates notch direction; vertical line is the end of notch].

112

Figure 3.30: Tensile fracture surface of R-PET showing semi ductile failure with crazing at
1.5 mm deep V-notch [white arrow indicates notch direction; vertical line is the end of notch].

113

Figure 3.31: Tensile fracture surface of R-PET showing brittle failure at 2.0 mm deep
V-notch [white arrow indicates notch direction; vertical line is the end of notch].

114

Figure 3.32: Izod impact strength of V-PET and R-PET showing gradual loss of strength atdifferent V-notch depths.

115

Figure 3.33: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof R-PET at 0.6 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

116

Figure 3.34: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof R-PET at 1.5 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

117

Figure 3.35: SEM of two opposite fracture surfaces [a, b] and fracture sample [c] of Izod impact specimenof R-PET at 2.0 mm deep V-notch [arrows indicate notch direction; vertical lines are the end of notches].

118

Chapter 4
The Fracture Modes and Notch Sensitivity of Hairline Cracked V-PET and
R-PET Injection-Moldings

Abstract
The effect of the skin and core structure of injection-moldings on their fracture modes when
standard v-notches were introduced was discussed in my previous investigations. However,in most engineering components in daily applications, the cracks on them are very
microscopic that they are hardly visible without the use of microscope. The sharp nature ofsuch microscopic cracks (also known as hairline cracks), their small angle and radius ofcurvature, is believed to have more detrimental effect on the fracture behavior of skin-core
structured injection moldings than the standard v-notches. Therefore, the focus of this study
is to investigate the effect of different hairline crack depths on the fracture behavior of neat
(V-PET) and recycled (R-PET) polyethylene terephthalate injection moldings. The slow
crystallization nature of V-PET and R-PET results in a distinct amorphous skin and more
crystallized core morphology when the materials are injection molded. Annealing the moldedcomponents creates an outer pseudo-skin and minimizes internal stresses in the samples. By
precisely controlling the hairline crack depth, the crack-tip can be situated within the pseudoskin/skin regions of the moldings. The fracture behavior of hairline-cracked specimens, in atensile loading, was found to depend not only on the crack depth but also on the region where
the crack-tips were situated (pseudo-skin, or skin). The crystallinity at the various layers
within the skin and core regions also was found to affect the fracture modes of the materials.
Specimens were found to exhibit ductile failure when the notch depth was within theamorphous skin region. By annealing the specimens, the morphological characteristics aswell as residual stresses of the specimens, especially at the skin region, could be altered. This119

was intended to elucidate the dependence of notch sensitivity of PET on the toughness of theskin region. Results revealed that annealing created a pseudo-skin layer and improved thedeformability of the materials, especially for V-PET, which fractured in a semi ductile
manner at all crack depths in contrast to its brittle failure even at shallow crack depths when
the specimens were not annealed.

4.1 Introduction
Very sharp and microscopic defects are known to affect the fracture modes of engineering
materials. Unlike standard 45o v-notches [1], crack initiation and propagation in sharp
microscopic cracks (also known as hairline cracks) is very rapid and catastrophic. Meanwhile,the skin and core structure of injection moldings is also believed to affect their fracturemodes, as have been proven in previous chapters. Another element of interest on the fracturemodes of hairline-cracked injection-molded components is the possible effect of residual
stresses. Annealing has been known to induce stress relaxation while enhancing the
crystallinity in injection-molded components. Therefore, in this chapter, the effect of
annealing on the skin-core morphology and fracture behavior of hairline-cracked samples
was correlated with the toughness and crack sensitivity of V-PET and R-PET injection
moldings. The notch sensitivity factors kS and kT for the un-annealed and annealed samples
were determined by introducing hairline cracks ranging from 5 – 60m and conducting
tensile tests on the specimens.

120

4.2 Characterization Methods

4.2.1 Material and Sample Preparation
Amorphous grade V-PET pellets (MA2103 and viscosity average molecular weight = 41499
g/mol), and R-PET (viscosity average molecular weight = 23131 g/mol) were used in thisinvestigation. R-PET flakes dried at 120⁰C for at least two hours were extruded with a set of
single screw extruder SR-Ruder Bambi SRV-P70/62. Barrel temperature was set between 255– 290⁰C, and screw rotation speed was 50rpm. Production of the R-PET pellets was done atroom temperature.

Prior to the fabrication of dumbbell samples with a TOYO PSS TI-30F6 injection-molding
machine, both V-PET and R-PET pellets were dried for four hours at 130⁰C (in a PICCOLO
Dryer). The barrel temperature was set between 250 – 270⁰C and the mold temperature wasset at 30⁰C. Injection pressure was set at 60kgf with a holding pressure of 60kgf. The coolingtime was maintained at 30s.

4.2.2 The Annealing of Dumbbell Specimens
Annealing of the specimens was performed by immersion into hot oil. An oil-filled beaker
was placed inside an equally oil-filled electric pot with temperature regulating capability as
shown in Figure 4.1. The set-up was heated and stirred regularly with a rod to maintainhomogeneous temperature within the oil in the beaker. The temperature of the oil in the
beaker was regulated at 160⁰C. In order to maintain their geometrical form, the samples were
sandwiched in-between two flat aluminum plates prior to dipping into the hot oil for 35
seconds. Subsequently, the annealed specimen was cooled in oil that was ice-chilled to lessthan 10 oC (Figure 4.1). It is important to note that the annealing was conducted via oil
121dipping rather than oven heating since it is intended to heat the samples at a fast rate andthereby modify only the outer skin region rather than the bulk of the specimens.

4.2.3 Measurement of Crystallinity in PET
The middle region of the V-PET and R-PET dumbbell specimens was sliced in the thickness
direction by using a microtome, JEOL TM-0015 from JEOL Ltd. Each slice measures about
50m in thickness. The slices were taken in sequence to correspond with the pseudo-skin,skin, interface and the mid-core regions of the specimens as shown in Figure 4.2 [a, b]. DSC
samples were extracted from one of the slices at each region. DSC was conducted with TA
Instrument model 2920 modulated differential scanning analyzer. The sample was heatedfrom 30⁰C to 300⁰C at 10⁰C/min under nitrogen atmosphere. Since crystallization occurredas the samples were heated, the enthalpies due to cold crystallization [EC] was subtracted
from the melting enthalpy [EM] to give the intrinsic enthalpy of the materials [EP]. Anapparent Crystallinity degree, %CDSC, was determined by equation (4.1). The resulting
thermal properties for the un-annealed and annealed specimens are shown in Tables 4.1 and4.2 respectively.

 H exp %CDSC  100  
0  H 

(4.1)

4.2.4 Laser-Raman Spectroscopy of the Materials
Short lengths of V-PET and R-PET samples before and after annealing were cut andembedded in epoxy solution with the width of the samples facing up and left to cure. The
samples were then polished core-deep by the use of rotational polishing wheel mounted withdifferent grain-sized abrasives and alumina to achieve very smooth and clean surface as122

shown in Figure 4.3. Laser-Raman spectroscopy was conducted in a scanning mode along
the melt flow direction of the polished surface of dumbbell samples. The Raman spectrum
was measured at room temperature with a light source of green laser (= 532 nm). The
diameter of the focused laser beam was 2m. All spectra were recorded by 1-3 time
integrations after 3-second irradiation to eliminate laser noise. The measured area was 30m
x 5000m with spatial intervals of 10m and 20m respectively. The average full width at
half-maximum (FWHM) Raman intensity was later determined from the scanned spectra
measurements to measure the crystallinity of the materials.

4.2.5 Birefringence Observation of PET Injection-moldings
Birefringence observation was performed on the cross-section of dumbbell samples to gauge
the sizes of the skin and core regions. Prior to observation by using a polarized-opticalmicroscope, the middle parallel section of the dumbbell specimen was cut and embedded in
epoxy with the cross section facing upward. One side of the cross section was polished with arotational polishing wheel mounted with abrasives of different grain sizes to achieve asmooth surface. The surface was further made as smooth and clean as possible by the use ofalumina particles. The polished side was mounted onto a clear and clean glass slide witharaldite glue to hold it in place and to ensure that liquid does not penetrate in-between theglass and sample during the polishing of the opposite side of the cross section. The polishing
was done until the specimen became thin and translucent with a thickness of about 30m.
Polarized optical micrographs were taken from the polished specimen through a 3.34 megapixel Nikon digital camera attached to a 10 X/0.25P magnification optical lens that ismounted on a Nikon ECLIPSE E600 Polarizer. The micrographs were digitally combined by
the use of Adobe Photoshop CS software. The resulting composite image was converted togray mode with Origin 7J [Rel v7.0265 (B265)] software by OriginLab Corporation. The
123

plots of the frequency and amplitude of the gray values were taken and analyzed over theentire cross-section [1]. The contour plots were equally derived from the birefringence toreveal a more complex hierarchical skin-core structure defining distinct boundaries between
the regions of PET injection moldings.

4.2.6 Hairline Crack Cutting by Micro Cutting Machine
Hairline cracks, a, ranging from 5 - 60m were introduced in the width direction at one edgeof the dumbbell samples. SAICAS DN-20S micro cutting machine, shown in Figure 4.4, that
was attached with a fresh and sharp cutter blade and mounted onto a 100 Newton load cell
was used for notching. The horizontal speed of the cutter was set at 25m/s. The angle andradius of curvature of the hairline crack tip were measured to be around 8º and 90m,
respectively, which were much smaller than that of the standard V-notch described in ASTM
D 256.

4.2.7 Testing and notch sensitivity characterization
Tensile testing was conducted on both un-annealed and annealed dumbbell specimens byusing an Instron 4466 universal tensile testing machine equipped with a 10kN load cell. The
tests were conducted with a crosshead speed of 10mm/min for both un-cracked and cracked
dumbbell specimens. The span length of the specimens measures at 115mm. At least sevensample pieces were tested for each material and notching conditions. All tests were conductedat room temperature.

124

4.2.8 Fracture Modes and Notch Sensitivity Characterization
Investigations were conducted on the effects of hairline cracks by analyzing the notch
sensitivity factors kS and kT at the successive crack depths for the un-annealed and annealed
samples. This gives an understanding of the effects of morphological changes on the notch
sensitivity of the material. If cracks have no detrimental effects on the tensile strength andtotal energy that is absorbed upon deformation, the notch sensitivity factors for strength (kS)and energy (kT) are equal to 1.0. Conversely, if they have detrimental effects, the notch
sensitivity factors are greater than 1.0 [1, 3,]. Therefore, the notch sensitivity factor of a given
material is defined in terms of the notch depth, which causes significant difference between
the mechanical properties of the notched and un-notched specimens. The effect of annealingon the toughness of the materials is determined by the difference between the notch
sensitivity factors for un-annealed and annealed samples. In this study, the notch sensitivity
factors are determined by the following equations [1, 3]: 

 (YS ) [t (w  a )] i 0(YSi )  (tw) 




(4.2)

A
 t (w  ai )  
 0ssc T  
Aissc  tw 




(4.3)

S

Where kS = notch sensitivity factor for yield strength, kT = notch sensitivity factor for energy.
YS0 and YSi = yield strength of un-notched and notched samples, respectively [subscript i
denotes the corresponding notch depth], t = thickness, w = width, and a = notch depth. A0ssc
and Aissc are the areas under the stress-strain curves for un-notched and notched specimens
respectively [1, 3].

125

The notch sensitivity factor for yield/fracture strength gives an understanding of how hairline
cracks affect the strength of materials. On the other hand, the notch sensitivity factor forenergy measures the effect of hairline cracks on the toughness of the materials.

4.3 Results and Discussions

4.3.1 Evaluation of Crystallinity in V-PET before and after Annealing
Table 4.1 shows the crystallinity at various depths of the V-PET materials before and after
annealing. The crystallinity at the topmost skin layer (<50m depth) of V-PET before
annealing was recorded at 30.1, which was slightly higher than deeper layers. This could be
the effect of the high shear rate at this region, which induced crystallization [4, 5]. Chainorientation and size of crystals could be the factors that increase crystallinity at the topmost
skin layer. The crystallinity, however, would decline at the depth of 550m, and then slightly
increased at 1400m, which is considered to be approaching the skin/core interface. Withfurther increment in depth of up to 2500m, a remarkable increase in crystallinity wasobserved. This region, which is the farthest from the mold surface, would cool at asignificantly lower rate and therefore develop into a more crystallized core region [4, 5].
These crystallinity values are consistent with the FWHM Raman intensity in Figure 4.5.
Similar to DSC results, it is clear from Figure 4.5 that the Raman intensity decreased from theskin region labeled X through the interface Y to the core Z, signifying increasing crystallinity
as the material crystallizes slowly from the fluid state through the rubbery to the crystalline
state [4, 5].

After the material was annealed, the crystallinity increased at all slice depths as shown in
Table 4.1. It is obvious that the annealed V-PET has similar hierarchical structure as was126

observed in the samples before annealing. The increase in crystallinity at the 50m slice
depth in the annealed sample when compared with the un-annealed sample reveals aremarkable effect of annealing on V-PET. This is evident in the FWHM Raman intensity in
Figure 4.5, which were spectra from Figure 4.3, where the introduction of the pseudo-skin isvisible at X1. Consistent with the samples before annealing, the Raman intensity gradually
decreased from the skin X through the interface Y to the core Z. This is evident that a more
crystallized pseudo-skin was created in V-PET after it was annealed.

The Full Width Half Maximum (FWHM) spectral intensity [Figure 4.5] of V-PET, before
and after annealing, shows strong consistence with the crystallinity derived from DSC in
Table 4.1. This consistence is depicted in Figures 4.6, which revealed significant differencein the crystallinity of the un-annealed and annealed specimens at different crack depths of thehierarchical skin-core structure of V-PET. Figure 4.6 shows that the crystallinity of V-PET
when it was not annealed was recorded at 30.1 in the 50m depth (skin layer, X). The
crystallinity however, decreased to 20.8 at the 550m depth (X) and increased through theinterface (Y) to the core (Z) regions. There was a remarkable increase in crystallinity to 37.9
at 50m depth when V-PET was annealed, revealing the introduction of a more crystallized
outermost skin (X1). Similar to the un-annealed specimens, the crystallinity decreased to 21.5
at 550m depth (X) and increased through the interface (Y) to the core (Z) regions.

4.3.2 The Morphology of V-PET before and after Annealing
Figure 4.7 shows the birefringence of the cross-section of V-PET before it was annealed [1].
The figure reveals distinct skin-core regions with well-defined boundaries, which enclose aninterfacial region. The image clearly revealed two distinct regions (skin and core) with theinterfacial layer in-between as revealed earlier in Figure 4.5. This establishes a relationship127

between the hierarchical structure of the material, crystallinity and FWHM Raman intensity.
The gray and contour plots of the matrix before annealing were shown in Figures 4.8 and 4.9,
which show that V-PET had well defined skin-core morphology with the presence of aninterface between the two regions. The well-defined morphology of V-PET is clearly shownby the intensity of the birefringence, gray and contour images in Figures 5.7 – 5.9
respectively, which is consistent with FWHM Raman spectra in Figure 4.5. The distinct
boundaries and high intensity of the contours of V-PET could have resulted from the size ofthe crystals and high molecular orientation, especially at the skin layer during injection
molding. The slow crystallizing nature of PET also made the intensity of the contours
stronger at the core. The degree of intensity of the contours depends on the length of themolecular chains, as well as the large size of the crystals. Figure 4.10 on the other handshows the birefringence of V-PET after it was annealed. The introduction of the pseudo skin
could be seen as indicated by arrows. The gray plots of the annealed specimens in Figure 4.11
show a more dispersed and diffused skin-core boundaries, which is depicted more clearly in
Figure 4.12. This suggests that annealing and quenching the samples did not only create amore crystallized outermost skin but resulted in the relaxation of internal shear stress.

4.3.3 Effect of Hairline Cracks on the Fracture Modes of V-PET before and after
Annealing
Table 4.3 partly shows that before annealing, V-PET had yield stress of 55MPa and did not
fracture below 300% strain when it was un-notched. However, when a 5m deep hairline
crack was introduced, the yield stress increased to 57.6MPa and ultimately failed in a semi
ductile manner. At 10m hairline crack, V-PET retained high yield stress even though thematerial fractured in a brittle manner as shown in Figure 4.13. Subsequent deepening of thecracks resulted in a corresponding increase in the yield stress, which could be attributed to the128tri-axial stress field generated from the crack as well as strain hardening. As stated earlier,
hairline cracks ≥ 10m caused brittle fracture in V-PET, and this could be attributed to highcrystallinity that resulted in high notch sensitivity resulting in a brittle fracture manner as
revealed in Figure 4.13.

In the case of the strain at break, the un-annealed and un-notched samples did not fracture
below 300% strain. However, the introduction of a mere 5µm deep hairline crack into thesample resulted in a drastic reduction in the fracture strain to about 14%. This marked atransition from ductile to semi-ductile fracture behavior. Further deepening of the crack to 10
m resulted in a brittle failure (3-4% strain) in V-PET. The brittle fracture of V-PET at 10m
crack depth could be attributed to higher crystallinity [1] as shown by the crystallinity in
Table 4.1 and the FWHM values of Raman spectroscopy in Figure 4.5.

Results reveal that in contrast to the un-annealed V-PET specimens, the annealed samples
experienced semi ductile fracture at all crack depths. This could be the direct effect of
annealing; and consequent quenching of the samples, which could result in the relaxation ofinternal stress within the bulk. It could be concluded that annealing V-PET created aprotective pseudo-skin, which did not only increase the yield stress but retarded the initialeffect of stress on the natural skin.

4.3.4 Correlation between Notch Sensitivity and Skin-core Morphology of V-PET
before and after Annealing
Figure 4.14 show the notch sensitivity factors for strength (kS) for V-PET as were calculated
by equation 4.2. The figure reveals that kS remained relatively the same (between 0.94 –0.96MPa) at different notch depths before annealing. However, after annealing kS increased,129

suggesting that V-PET became less sensitive to hairline crack because of increased modulusof resilience arising from the creation of more crystallized outermost skin. The value of kS is
known to be equal to 1.0 if notches have no detrimental effects on the yield stress of thematerials, otherwise it is greater than 1.0 [1, 3]. Figure 4.14 shows that kS values (0.9-1.0) for
V-PET before annealing at all notch depths are slightly lower than the values after annealing.
This suggests that V-PET is less affected by hairline cracks after annealing than before
annealing because of reduction in triaxial stress effect. It has to be noted that the increase inyield stress of V-PET did not translate into lower notch sensitivity factor for V-PET.
However, kS for V-PET increased slightly after annealing, which signified a decrease in thenotch sensitivity of V-PET in respect of yield stress. This could be attributed to the presenceof the more crystallized pseudo-skin, which would have increased the modulus of resilienceof the specimen. It important to mention that there could have been an increase in themodulus of resilience at the edges sides of the specimen that caused an increase in yield stress.
It minimized also the tri-axial effect of the hairline crack, which would have occurred at thenatural skin because of its high amorphous structure. It is important to observe that all thehairline cracks (≤ 60m) were within the pseudo-skin region (≈ 400m).

The notch sensitivity factors for energy (kT) are calculated with equation 4.3, and shown in
Figure 4.15. It was earlier defined that a material becomes notch sensitive at a notch depththat causes significant difference in the mechanical properties between the notched and unnotched specimens. The values of kT > 5.0 show that a mere 5m deep hairline crack had
detrimental effect on the fracture behavior of the un-annealed samples of V-PET. This meansthat the materials were sensitive to a mere 5m hairline crack, and had drastic change fromhighly ductile to semi-ductile fracture behavior. When the hairline crack was deepened to ≥10m, the kT for V-PET increased to more than 8.0 and remained constant, which signifies
130brittle fracture. The values of kT have shown that prior to annealing V-PET is more sensitive
to hairline cracks and fractures in a brittle manner when the crack is ≥ 10m deep. The major
reason for the high crack sensitivity of V-PET is its higher level of crystallinity (Table 4.1
and Figures 4.5 and 4.6).

Figure 4.15 shows a remarkable reduction in the notch sensitivity factors for energy (kT) of
V-PET. It is obvious that annealing greatly improved the toughness of V-PET, whichtranslates into a decrease in kT before annealing from between 5.0 and 9.0 to < 3.0. It is also
remarkable to see that V-PET fractured in a semi-ductile manner at all crack depths. Results
also reveal that after annealing, V-PET became less sensitive to hairline cracks in terms of
toughness.

4.3.5 SEM Characterization of Tensile Fracture Surfaces of V-PET
The fracture surfaces of V-PET at 5m and 10m hairline cracks are shown in Figures 4.16
and 4.17. Although V-PET has very high kT ≥ 8.0, it is evident in Figure 4.16 that the fracturesurface at 5m deep crack showed significant degree of plastic deformation, which ischaracteristic of fracture resistance. This suggests that there was plastic flow from adjacentareas towards the stress concentration point, which is at the crack-tip. The fracture surface
reveals evidence of fracture by tearing as the amorphous skin migrates to the crack-tip. It is
important to observe that the thin nature of the fibrils on the fracture surface is a result of lessdeformability arising from fast crack propagation because of the sharp hairline crack. The
fracture surface also shows that deformation in the geometry of the sample was symmetricaland tapered towards the opposite end of the sample. Figure 4.17 depicts the brittle fracture of
V-PET when the hairline crack was further deepened to 10m. The figure shows no evidenceof resistance to fracture as the material failed in a brittle manner. This indicates that the very131

sharp and microscopic nature of the hairline crack at 10m deep resulted in very quick crack
propagation without much plastic deformation. This is evident in the flat appearance of thefracture surface. Figure 4.18 shows that upon annealing, V-PET fractured in a semi-ductile
manner even when a 10m deep crack was introduced. The fracture surface resembles that of
the un-annealed specimen at 5m crack depth. The semi ductile fracture suggests that theannealing process resulted in relaxation of residual stresses, which subsequently improved thetoughness of the material and made it less sensitive to hairline crack. Evidence suggestingthat there was stress relaxation in V-PET after annealing could be seen in Figure 4.12 where
the intensity of the contour plots appears more dispersed than prior to annealing [Figure 4.9].
It is obvious that the well-defined skin-core boundaries that were seen in the samples before
annealing became diffused after the material was annealed, which indicates a reduction inmolecular orientation.

4.3.6 Evaluation of the Crystallinity in R-PET Materials before and after Annealingby the use of DSC and Raman Spectroscopy
Table 4.2 shows the thermal properties of R-PET materials before and after annealing. The
crystallinity and temperatures resulting from cold crystallization and melting, as recorded
from DSC experiment are measured from sample slices at different hairline crack depths. The
crystallinity varied at different samples slice depths taken from Figure 4.19 [a, b]. The
crystallinity of R-PET from the surface to about 550m depth was consistently recorded atabout 9.2, which is considered the skin region. A significant increment in crystallinity wasthen observed at a depth of about 1400m, which is considered the terminal end of the skinand the beginning of the core region. The crystallinity was consistent thereafter even until thedepth of 2500m. Because of the slow crystallization process of semi-crystalline PET, thecrystals in the core region had enough time to develop, resulting in higher crystallinity [4, 5].
132

Nonetheless, the changes in crystallinity from the skin to core regions were more gradual in
R-PET as compared to V-PET, which suggests a more homogenous morphological structurein R-PET. These values are consistent with the FWHM Raman intensity shown in Figure 4.20
where no significant difference in peak intensity could be observed within the skin-core subregions.

After the material was annealed, the crystallinity of the R-PET samples increased
significantly at all depths, especially at the 50m depth as shown in Table 4.2. It is obviousthat the highest increment in crystallinity was recorded at the outermost skin layer of R-PET
after annealing, followed by the deeper regions. The core regions at the depth of 1400m
onwards did not experience any significant increment in crystallinity. The availability ofmore amorphous regions at the skin layer allowed further crystallization upon annealing
whereas the core, which was saturated with crystalline regions, would allow very little freevolume for the formation of new crystals.

Figure 4.20 shows similar crystallinity structure as depicted in the FWHM spectral intensityof R-PET before and after annealing, which corroborates the dispersed skin-core structure of
the material as shown later in Figures 4.25 – 4.27. The crystallinity of R-PET measured at
different depths from DSC is also depicted in Figure 4.21. The figure reveals that annealing
resulted in a significant increase in the crystallinity at the 50m depth. It is obvious thatsimilar to FWHM, the crystallinity show a homogeneous structure at different depths, whichsuggests a more homogenous hierarchical skin-core structure as depicted in Figure 4.22 –4.27. It could be said that the crystallinity of R-PET before and after annealing appear similarfrom the 550m depth to the 2500m depth. These homogenous hierarchical transitionalchanges in Figures 4.20 and 4.21 is consistent with earlier observation in DSC values, and133

could be the result of thermal changes of the injected resin from the fluid state to crystallizedstate [4, 5]. It is important to note that the observation of skin-core structural homogeneity in
R-PET sharply contrasts with the well-defined and segmented skin-core structure of V-PET.

4.3.7 The Morphology of R-PET before and after Annealing
Figure 4.22 shows the birefringence of the cross-section of R-PET samples before they wereannealed. Although the intensity of birefringence is less than V-PET, there is significant
degree of clear boundaries between the skin-core regions. The conversion of Figures 4.22 tograyscale and subsequent plots defining the intensity of the gray areas for the un-annealed
sample is depicted in Figure 4.23 [1]. Figure 4.24 reveals a more complex skin-coremorphology, which shows the existence of the skin, interface and core regions in the samples.
Figure 4.25 shows the birefringence of R-PET after it was annealed, which revealed thecreation of an outer skin that is referred to as the pseudo-skin (indicated by the arrows at thefour sides). Hence, after the material was annealed, three regions, namely the pseudo-skin,
skin and core regions could be seen. The plots for the annealed samples shown in Figure 4.26
depict the introduction of the pseudo-skin resulting from annealing, in addition to theexistence of the skin, interface and core regions. It is clear that after annealing, the distinction
between the regions became less prominent, which appears to be the effect of stressrelaxation. The consequent quenching of the samples in ice-chilled oil could have disrupted
the development of crystals within the interface and core regions.

Derivations from the intensity of the gray areas were used to determine precisely the
thickness of the pseudo-skin, skin, interface and core regions, as shown in Figures 4.26. The
plots marked with W and X represent the intensity values of the gray scale in the width andthickness directions of the samples, respectively [1]. The horizontal and vertical cursors on134

the matrix were adjusted along the width and thickness directions until well-defined
frequency peaks were obtained at each region. The center region in between the two
maximum peak intensities obtained from W and X plots in each region were taken to be thecore width and thickness, respectively [1]. Before annealing, the skin region of R-PET
measures about 1500m. On the other hand, the pseudo-skin of the annealed R-PET
measures about 400m.

The contour plots of R-PET after annealing is shown in Figure 4.27, respectively. The
contour of the un-annealed sample in Figure 4.24 shows that R-PET had noticeable skin-coremorphology with the presence of an interface between the two regions. In contrast, theannealed R-PET in Figure 4.27 depicts a more dispersed and diffused skin-core morphology,
which shows that the intensity of the contour decreased, and the boundaries between theregions are non-existent. This observation corroborates the earlier results on FWHM in
Figure 4.20, suggesting that there was high hierarchical homogeneity in the bulk of thematerial. The degree of intensity of the contours of R-PET before annealing depended on theslow crystallizing nature of PET, which permitted the formation crystals. Conversely,
quenching the annealed R-PET from 160ºC to 7 - 10ºC in ice-chilled oil disrupted the slow
crystallization process, which would have resulted in crystals formation. Consequently, there
was a more relaxed and dispersed contour appearance in R-PET [Figure 4.27].

4.3.8 Effect of Hairline Cracks on the Fracture Modes of R-PET before and after
Annealing
Table 4.3 shows that before annealing, R-PET exhibited a yield stress of about 54MPa when
it was un-notched. The introduction of 5 – 60 m deep hairline cracks resulted in an increasein yield stress to 59MPa. However, when the samples were annealed the yield stress remained135

at an average of 57.7MPa for both the un-notched and notched samples. The slight decreasein the yield stress of R-PET could have resulted from a reduction in the size of the crystals
after annealing and subsequent quenching. Annealing is also believed to have removed
internal stresses from the materials thereby increasing the materials’ deformability.

The un-notched R-PET samples did not fracture below 300% strain regardless of whether
they were un-annealed and annealed, as shown in Table 4.3. However, the introduction of amere 5µm deep hairline crack resulted in a drastic reduction in strain from about 300% toabout 14%. This marked a transition from ductile to semi-ductile fracture behavior. The semi
ductility was retained in R-PET even as the crack deepened to 10 – 60m. Figure 4.28
depicts the semi ductile fracture behavior of R-PET when hairline cracks deeper than 5mwere introduced. Prior to shear banding, two shear lines were seen to originate from thecrack-tip at a 45º angle as shown in Figure 4.28[a]. Crack opening would then continue on
one of the shear lines until tearing at the crack tip occurred [Figure 4.28(b)]. The similar semi
ductile fracture behavior exhibited by R-PET before and after annealing could have been
influenced by the homogenous morphological structure as was previously discussed. It could
be concluded that annealing did not significantly change the morphological hierarchy of RPET; hence, the fracture behavior was comparable.

4.3.9 Correlation between Notch Sensitivity and Skin-core Morphology of R-PET
before and after Annealing
Figure 4.29 shows the notch sensitivity factors for strength (kS) for R-PET as were calculated
by equation 4.2. It is important to observe that all the hairline cracks (≤ 60m) were within
the pseudo-skin region (≈ 400m). The figure reveals that the kS value of un-annealed R-PET
specimens remained at about 0.9 regardless of crack depth, which indicates that the crack
136

slightly increased the yield stress of the material [1, 3]. This could be attributed to theformation of tri-axial stress fields that caused the specimens to twist during tensile loading,thus exerting higher resistance to deformation. Upon annealing, the kS values for R-PET wasapproximately equal to 1.0 at all notch depths, suggesting that the relaxation of internal
stresses and molecular orientation were able to minimize the tri-axial stress field effect andtherefore provided more flexibility to the specimens. The presence of the pseudo-skin on thesides of the samples increased the modulus of resilience at the edges and minimized the effectof hairline cracks on the natural skin. It could therefore be concluded that annealing was able
to minimize tri-axial stresses during tensile loading due to the relaxation of internal stresses
as well as oriented molecules.

The notch sensitivity factors for energy (kT) calculated from equation 4.3 is shown in Figure
4.30. The values of kT ≥ 5.0 show that a mere 5m deep hairline crack had effect on thefracture behavior of the un-annealed samples of R-PET, indicating that there was a drasticchange from highly ductile to semi-ductile fracture behavior. However, the kT for R-PET
remained < 6.0 as the material continued to fracture in a semi-ductile manner with notches ≤20m deep. The values of kT slightly increased above 6.0 but less than 7.0, which is still in asemi-ductile fracture. The values of kT have shown that before annealing R-PET is sensitive
to hairline cracks. Figure 4.30 shows a remarkable reduction in the notch sensitivity factors
for energy (kT) after R-PET was annealed. It is obvious that after annealing kT for 5 – 60m
hairline crack depths decreased from between 5.0 – 8.0 to 3.0 suggesting that R-PET showeda remarkable reduction in its sensitivity to all the hairline crack depths.

137

4.3.10 SEM Characterization of Tensile Fracture Surfaces of V-PET
Figures 4.31 and 4.32 represent the fracture surfaces of R-PET before and after annealing at
5m and 10m hairline cracks, respectively. Although R-PET has very high kT ≥ 6.0, it isevident in Figure 4.31 that the fracture surface at 5m deep crack shows significant degree ofplastic deformation, which is characteristic of ductile failure. The tapered geometry from thecrack initiation to the fracture points suggests that there was plastic flow from surrounding
areas towards the crack-tip. The layered structure on the fracture surface also indicatestearing as the amorphous skin migrates to the crack-tip in resistance to crack propagation.
Figure 4.32 depicts a similar semi ductile fracture even when the hairline crack was further
deepened to 10m. The figure reveals that at 10m deep crack, R-PET had experienced
significant plastic deformation that was shortly disrupted by craze midway into the crack
propagation process.

It can be concluded that the semi ductile fracture at the 10m deep crack suggests thatannealing R-PET in 160˚C hot oil and quenching in 7 – 10˚C ice-chilled oil resulted in stressrelaxation and more homogenous skin-core structure, which subsequently made the material
less sensitive to hairline crack.

4.4 Conclusion
The current study has shown that the annealed V-PET and R-PET experienced semi ductile
fracture behavior under tensile loading at notch depth of 5-60m. This semi ductile transition
suggests that annealing is effective in reducing brittle fracture in PET. Results have shownthat annealing created a more crystallized pseudo-skin, which measures ≈ 400m in thematerials. The pseudo-skin has provided a higher modulus of resilience to the materials,
especially V-PET, which delayed the deformation of the natural skin. Furthermore, annealing138

was able to cause substantial relaxation of the internal stresses as well as molecularorientation, thus minimizing tri-axial stresses and allowing the specimens to deform moreduring tensile loading.

References
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Differential

Scanning Calorimetry and Confocal

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Differential

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140

Table 4.1: Crystallinity of V-PET Injection moldings before and after annealing
Cold crystallization (EC)

Melting (EM)Slice depth

EP = (EM – EC)%C = 100(Materials
(m)

V-PET

Crystallinity

ΔH0 = 117.7J/g

Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

50

34.1

125.9

69.5

257.4

35.4

30.1

550

20.3

126.0

44.8

256.8

24.5

20.8

1400

22.2

124.6

46.2

256.7

24.0

20.4

2500

10.1

125.2

44.5

256.9

34.4

29.2

50

4.5

123.2

49.1

257.2

44.6

37.9

550

23.1

124.7

48.4

256.9

25.3

21.5

1400

24.2

123.8

52.8

257.4

28.6

24.3

2500

12.7

124.4

54.2

257.4

41.5

35.3

(28.1C/g)

Beforeannealing

V-PET
After
annealing

Note: Δ°: heat of fusion of an ideal PET crystal for which value of 117.7J/g (28.1Cal/g) was considered, Δexp.: measured heat of fusion

141

Table 4.2: Crystallinity of R-PET Injection moldings before and after annealing
Cold crystallization (EC)

Melting (EM)Slice depth

EP = (EM – EC)

Crystallinity%C = 100(Materials
Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy(J/g)

ΔH0 = 117.7J/g(28.1C/g)

50

26.8

130.7

37.3

253.6

10.8

9.2

R-PET

550

26.1

130.5

37.2

253.7

10.9

9.3

Beforeannealing

1400

35.0

130.7

49.9

232.9

16.1

13.7

2500

23.1

130.4

39.4

253.4

15.8

13.4

50

27.9

124.4

65.8

254.2

37.9

32.2

R-PET

550

19.3

128.5

37.7

254.2

18.4

15.6

After
annealing

1400

21.5

129.8

38.6

254.0

17.1

14.5

2500

20.9

129.9

42.4

253.6

21.5

18.3

Note: Δ°: heat of fusion of an ideal PET crystal for which value of 117.7J/g (28.1Cal/g) was considered, Δexp.: measured heat of fusion

142

Table 4.3: Mechanical Properties of VPET and RPET samples before and after annealing

Materials

V-PET

R-PET

Before Annealing

After Annealing

Hairline crack
depth (m)

Yield stress% Strain

Yield stress% Strain

Un-notched

55.2

> 300

59.9

< 100

5

57.6

14.8

61.7

13.2

10

58.0

4.2

61.7

13.3

20

58.1

3.8

61.6

13.3

40

58.3

3.7

62.8

13.4

60

57.9

3.5

61.1

13.1

Un-notched

54.19

> 300

57.6

> 300

5

59.1

15.3

57.7

14.1

10

59.3

14.7

57.7

14.1

20

59.3

15.1

57.5

14.2

40

60.0

14.0

57.8

14.2

60

58.7

14.7

57.8

14.3

143

Figure 4.1: Annealing apparatus showing dumbbell sample held in-between two aluminum bars,
oil-filled beaker held inside oil-filled electric pot, and a bath containing Ice-chilled oil (7 - 10°C)
144

X* = Pseudo-skin

Figure 4.2: DSC sample slices at 50mm thick each showing samples representing the different regions of
V-PET a = before annealing, b = after annealing [X* = Pseudo-skin, X = skin, Y = interface, and Z = core]145

V-PET after
annealing

R-PET after
annealing

V-PET before
annealing

R-PET before
annealing

X1

X1

X

X

X

X

Y

Y

Y

Y

Z

Z

Z

Z

Figure 4.3: Epoxy embedding and smooth-polishing of un-annealed and annealed V-PET and R-PET samples
for Raman spectroscopy scanning in the flow direction [X* = Pseudo-skin, X = Skin, Y = Interface, and Z = Core]
146

Figure 4.4: Apparatus for cutting hairline cracks on the dumbbell specimens

147

FWHM Raman intensity

30

X
Y

25

Z

X

20

X1

Y

Z

15
10
Un-annealed
Annealed

5
0

0

1000
2000
3000
4000
Hairline crack depth (mm)5000

Figure 4.5: FWHM Raman intensity showing crystallinity along the skin-core region of V-PET
before (UA) and after (A) annealing [X1: pseudo-skin, X: skin, Y: interface, and Z: core]148

Crystallinity (%)

40

X1
Z

30

Y

20
X

10

Un-annealed
Annealed

0

0

500

1000

1500

2000

2500

3000

Hairline crack depth (mm)Figure 4.6: Crystallinity along the skin-core of V-PET before and after annealing[X1: Pseudo-skin, X: skin, Y: interface, and Z: core]
149

Skin

Skin

Skin

Core

Skin
Figure 4.7: Birefringence observation of the cross-section of V-PET
dumbbell samples before annealing showing their skin-core structure

150

W

X
Figure 4.8: Gray image plots of V-PET before annealing showing the gray values of the skin, interface
and core regions [W & X= graph of gray values in the width and thickness directions, respectively].

151

Core
Interface
Skin

Figure 4.9: Image contour plots of V-PET before annealing showing the skin, interface and core regions.

152

Skin
Skin

Skin

Core
Skin
Figure 4.10: Birefringence of the cross-section of V-PET dumbbell samples after annealing
Showing the pseudo-skin, skin and core structure [pseudo-skin is indicated by the arrows]

153

W

X
Figure 4.11: Gray image plots of V-PET after annealing showing the gray values of the pseudo-skin, skin,interface and core regions [W & X= graph of gray values in the width and thickness directions, respectively]

154

Pseudo skin

Skin

InterfaceCore

Figure 4.12: Image contour plots of V-PET after annealing showing the pseudo-skin, skin, interface and core regions

155

a

b

Figure 4.13: Crack initiation and brittle fracture of V-PET at 10mm hairline crack depth

156

Crack sensitivity for strength (KS)

1.04
1

0.96
0.92

V-PET UA
V-PET A

0.88
0

10

20

30

40

50

60

70

Hairline crack (mm)Figure 4.14: Crack sensitivity factors for strength (kS) of V-PET before and after annealing157

Crack sensitivity for energy (K T)

108
6

V-PET UA
V-PET A

4

2
0
0

10

20

30

40

50

60

70

Hairline crack (mm)Figure 4.15: Crack sensitivity factors for energy (kT) of V-PET before and after annealing158

Figure 4.16: Fracture surface of V-PET showing semi-ductile
failure at 5mm hairline crack depth before annealing
159

Figure 4.17: Fracture surface of V-PET showing brittle
failure at 10mm hairline crack depth before annealing160

Figure 4.18: Fracture surface of V-PET showing semi-ductile
failure at 10mm hairline crack depth after annealing161

X1 = Pseudo-skin

Figure 4.19: DSC sample slices at 50mm thick each showing samples representing the different regions of
R-PET a = before annealing, b = after annealing [X1 = pseudo-skin, X = skin, Y = interface, and Z = core]162

FWHM Raman intensity

30
25
20

15
10
Un-annealed
Annealed

5
0
0

1000

2000

3000

4000

5000

Hairline crack depth (mm)Figure 4.20: FWHM Raman intensity showing crystallinity
along the skin-core of R-PET before and after annealing163

Un-annealed
Annealed

Crystallinity (%)

40
X1

30

20

Y

Z

10
X

0
0

500

1000

1500

2000

2500

3000

Hairline crack depth (mm)Figure 4.21: Crystallinity along the skin-core of V-PET before andafter annealing [X1: Pseudo-skin, X: skin, Y: interface, and Z: core]
164

Skin

Skin

Skin

Core
Skin
Figure 4.22: Birefringence observation of the cross-section of R-PET
dumbbell samples before annealing showing their skin-core structure

165

W

X
Figure 4.23: Gray image plots of R-PET before annealing showing the gray values of the skin, interface andcore regions [W & X= graph of gray values in the width and thickness directions, respectively].

166

Core
Interface
Skin

Figure 4.24: Image contour plots of R-PET before annealing showing the skin, interface and core regions.

167

Skin
Skin

Skin

Core
Skin
Figure 4.25: Birefringence of the cross-section of R-PET dumbbell samples after annealing
showing the pseudo-skin, skin and core structure [pseudo-skin is indicated by the arrows]

168

W

X
Figure 4.26: Gray image plots of R-PET after annealing showing the gray values of the pseudo-skin, skin,interface and core regions [W & X= graph of gray values in the width and thickness directions, respectively]

169

Figure 4.27: Image contour plots of R-PET after annealing showing no identifiableboundaries between the pseudo-skin, skin, interface and core regions.

170

a

b

Figure 4.28 Shear-banding and semi-ductile fracture of R-PET at 5 - 60mm hairline crack depth

171

Crack sensitivity for strength (KS)

1.04

R-PET UA
R-PET A

1

0.96
0.92

0.88
0

10

20

30

40

50

60

70

Hairline crack (mm)Figure 4.29: Crack sensitivity factors for strength (kS) of R-PET before and after annealing172

Crack sensitivity for energy (K T)

10
8
6
4

2

R-PET UA
R-PET A

0
0

10

20

30

40

50

60

70

Hairline crack (mm)Figure 4.30: Crack sensitivity factors for energy (kT) of R-PET before and after annealing

173

Figure 4.31: Fracture surface of R-PET showing semi-ductile
failure at 5mm hairline crack depth before annealing
174

Figure 4.32: Fracture surface of R-PET showing semi-ductile
failure at 10mm hairline crack depth before annealing
175

Figure 4.33: Fracture surface of R-PET showing semi-ductile
failure at 10mm hairline crack depth after annealing176

Chapter 5
Essential Work of Fracture Evaluation of the Fracture Modes of VNotched PET Components

Abstract
It is often thought that poly(ethylene terephthalate) (PET), like other amorphous or slow
crystallizing polymers, is extremely notch sensitive. However, in our previous studies, three
distinct fracture modes, i.e. ductile, semi-ductile and brittle, were observed during tensileloading of PET when the notch depth was gradually increased. Therefore, it is thought thatthe notch sensitivity of PET is dependent on its morphology and notch depth, while linear
elastic fracture mechanics would be inadequate to characterize the fracture toughness of PET,
especially when the material is injection molded and possesses the characteristic skin-core
structure. This study revealed that the temperature at the crack-tip would significantly
increase due to local plastic deformation when the crack-tip was situated within the skinregion. As such, full ligament yielding could be obtained, which allowed the application of
the principles of essential work of fracture (EWF) to characterize the fracture toughness of
PET. However, no significant temperature rise would occur when the specimens fractured ina brittle manner as the notch penetrated through the skin and into the core. At this point, thematerial is said to be notch sensitive and would no longer be able to undergo any plastic
deformation prior to failure.

5.1 Introduction
Fracture is a dissipative process during which either plastic work is converted partly into
heat; or heat is released during the fracture process itself [1]. Furthermore, when the process177is adiabatic, the released heat causes a local temperature rise, suggesting that in the formationof an adiabatic shear band there is a substantial increase in the local temperature in the band
[1].

Therefore, heat dissipation due to the combined effect of plastic deformation and crack
propagation when PET materials are subjected to tensile stress is believed to affecttemperature rise around the stress concentration point, including the crack tip. This could
equally reflect the magnitude of stress intensity on the stress field of the materials. The
uniformity of the specimen cross-section, size, and the presence of impurities, voids orcavities could influence the heat emitted by the specimens during the application of stress.
The introduction of cracks or notches on injection molded flat bars affects their lines of stress[2, 3]. The structural stability of the flat bar under tension depends not only on the depth ofthe crack but also on the axis of the stress field, as well as the location of the stressconcentration point. In an un-cracked uniform bar (dumbbell), the stress concentration point
is located at the center, and runs parallel to the direction of applied force. In injection-molded
semi-crystalline thermoplastics such as polyethylene terephthalate (PET), the presence ofskin and core layers becomes another factor that influences their fracture behavior [4]. It is
known that some materials are more sensitive to notches than others are. Therefore, it wassuggested that stress-strain tests should be conducted on both notched and un-notched test
specimens for comparison. They agreed that the stress-strain tests enable design engineers to
predict the fracture behavior of perfectly smooth (uniform) objects, but these do not allow theprediction of fracture behavior of objects with cuts [5]. Mouzakis et al noted that local
temperature rise accompanies full ligament yielding both in notched and un-notched
specimens [6]. The introduction of deeper notches in skin-core structured PET materials
gradually transits the fracture mode from semi ductile to brittle failure. This means that as the178notches become core-deep, brittle failure is imminent. However, there is a mixed fracturemode [7] in the materials as the terminal-skin portion fractures and crazing precedes brittle
failure. The sudden retraction of the fractured terminal-skin portion and the resultant craze ofthe core region generate heat in the process deformation zone, which increases thetemperature at the notch root. Bjerke et al observed that the heating that occurs in the process
deformation zone of an amorphous brittle polymer during dynamic fracture is the result ofpolymer crazing [8]. They added that craze formation and breakdown process includes manymechanisms that likely generate heat [8, 9]. They cited (Kramer, 1983; Kramer and Berger,
1990; Wool, 1995; and Kambour, 1964) that as crazes form and lengthen significant
molecular chain motion occurs within the active layer [8, 10, 11]. Such movement of thechains, which include sliding and rotation, would lead to heat generation [8, 11]. They
concluded that the dissipative processes, which involve heat that contributes to temperaturerise include, the failure of craze fibrils, which involves molecular chain scission [8, 11];molecular chain pullout from the surrounding material [8, 12] and the recoil of the fibril stubs[8, 12, 13].

This investigation involves introducing notches of various depths into injection moldings andmonitoring the temperature distribution at the notch root during tensile loading. The
temperature profiles corresponding to the various notch depths would then be analyzed basedon the fracture behavior and morphological characteristics of the specimens. The effects ofthe skin-core structure on the notch sensitivity and essential work of fracture will be
discussed. Equally to be discussed is the correlation between Full Width Half Maximum(FWHM) of Raman intensity, yield stress and notch sensitivity factors.

179

5.2 Characterization Methods

5.2.1 Material and Sample Preparation
Materials and samples preparation are the same as were described previously in Chapter 4.

5.2.2 Measurement of Crystallinity in PET
Fifty-six (56) slices from dumbbell specimens were obtained from the thickness side of adumbbell piece by using a JEOL TM-0015 microtome from Leica Microsystems, Germany.
Each slice measures 50 m in thickness. The slices were taken in sequence corresponding
with the skin, interface and the mid-core regions of the specimens as shown in Figure 5.1.
Samples measuring between 5.6 – 6.5 mg were cut from slices of (50, 550), 1400 and2500m deep, representing the skin, interface and core regions of the specimen as indicatedin Table 5.1. DSC was conducted as described in Chapter 3.

5.2.3 Birefringence Observation of Skin-Core Morphology of Injection- Molded PET
Birefringence observation was performed on the cross-section of dumbbell samples to gauge
the sizes of the skin and core regions. Prior to polarized-optical-microscopy, a short sectionof dumbbell specimen was cut and placed in epoxy. One side of the cross-section waspolished with a rotational polishing wheel mounted with abrasives of different grain sizes toachieve a smooth surface. The polished side was then mounted to a clean slide glass withepoxy glue to prevent moisture from penetrating in between the glass and sample surface
during the polishing of the opposite end of the specimen. The second polishing step involves
grinding the specimen down to a thickness of approximately 30 m. This results in a smoothcross section that is almost transparent. Polarized optical micrographs were taken from thepolished specimen through a 3.34 mega pixel Nikon digital camera attached to a 10 X/0.25P
180

magnification optical lens, which is mounted on Nikon ECLIPSE E600 Polarizer. The
micrographs were joined together by the use of Adobe Photoshop CS software. The resulting
composite image panorama [2] [refer to Figures 3.3 and 3.21 in Chapter 3] was converted togray mode [Figure 5.2] with Origin 7J [Rel v7.0265 (B265)] software by OriginLab
Corporation; the plots of the frequency and amplitude of their gray values over the entire
cross-section were taken and analyzed.

5.2.4 Notching and mechanical property characterization
V-notches of 0.5 – 3.0 mm depths were introduced on dumbbell test pieces with a Tool A-3
Digital Notching Machine (from Toyo Seiki Manufacturing Co., Ltd). The notching was
performed in stages of 0.5 mm depth in order to minimize internal deformation of thesamples. The angle of the V-notches is in accordance with ASTM D256.

Tensile testing of notched and un-notched dumbbell specimens was performed with an
Instron 4466 universal testing machine mounted with 10 kN load cell and set at a crosshead
speed of 10 mm/min. At least seven sample pieces were tested for each material and notchingcondition.

5.2.5 Monitoring of Temperature Distribution in Specimens during Tensile Test
During tensile test, an infrared video camera (TVS-200EX from NEC Avio ThermoTechnology, Japan) was used to monitor the temperature distribution on the entire surface ofthe test piece, including the crack tip. The camera speed was set at 5 frames per second where
each frame contains temperature readings taken from the start of tensile test through crack
initiation and propagation stages. The infrared camera was set at an emissivity coefficient of
1.0 to match that of polymer materials (plastic opacity) of 0.95. Tensile tests were conducted181

at a room temperature of 28oC. An infrared image analyzing software (Thermography Studioversion 4.8 from Nippon Avionics Co., Ltd) was used to export the temperature distributions
and graph plots for analysis. The tensile test and infrared camera were started simultaneously;
therefore, the temperature distribution in the specimens could be correlated to the stress-strain
curves as indicated in Figure 5.3.

The maximum change in temperature (ΔT) for each of the notch conditions was taken at fullligament yielding (as indicated in Figure 3c) and determined by equation (5.1) for notchedsamples that exhibited full ligament yielding [15]. Conversely, ΔT for samples with brittle
failures were taken at fracture. Corresponding infrared images, which show pre-yield, yieldand full ligament yieding modes of the specimens, were extracted from the thermo video
recordings taken during tensile test as indicated by insets (a), (b) and (c) in Figure 5.3.

T  Tmax  Tmin

(5.1)

Where Tmax = maximum temperature at full ligament yielding, Tmin = minimum temperatureof specimen at full ligament yielding.

5.2.6 Notch Sensitivity and Temperature Distribution Characterization
Investigations were previously conducted in Chapter 3 on the effects of V-notches on thefracture behaviors of V-PET and R-PET by analyzing the notch sensitivity factors for yieldstrength (kS) and energy (kT) at the different notch depths. In the present Chapter, the resultsof kS and kT are correlated with the essential work of fracture parameters, in order to gain anunderstanding of the effects of notches on the fracture behavior of the materials.

1825.2.7 Determination of Essential Work of Fracture (EWF)
The essential work of fracture (EWF) is used to characterize highly ductile materials, which
exhibited full ligament yielding even upon notching. The total works of fracture (W f) in
ductile materials under uniaxial tensile stress occur in two distinct deformation zones. They
are represented in the following equation [6, 7, 15]:

Wf  We  Wp

(5.2)

Essential work of fracture (We) is the work done in the process deformation zone; and nonessential work of fracture (Wp) is the work done in the plastic deformation zone for actual
crack opening and propagation. The essential work of fracture method requires that materials
considered for analysis should have full ligament yielding [15, 16]. It is also required that thecorrelation coefficient (R2) of the line of best fit between total specific work of fracture (W f)and length of ligament be very strong. Hence, the following equation applies [6, 7, 15]:

wf 

Wf
lt

 we   wp l

(5.3)

Where wf = specific work of fracture, l = ligament length and t = thickness of the samples, we= specific essential work of fracture = the failure work per unit surface area, and wp = specificnon-essential work of fracture = plastic work per unit volume of the plastic deformation zone.
β = shape factor of the plastic deformation zone.
Equation (5.3) indicates that wf is a linear function of l, therefore we is the positive intercept
at l = 0 of the linear regression line of the data; and βwp is the slope of the regression line. It
is important to note that one of the objectives of EWF is to show that we is a toughness
parameter [16].
183

5.3 Results and Discussion

5.3.1 Essential Work of Fracture Evaluation and the Analysis of Crack Tip
Temperature Distribution of the Fracture Modes of V-Notched V-PET Components

5.3.1.1 Measurement of Crystallinity in V-PET
Crystallinity is believed to affect the fracture behavior and energy absorption of notchedinjection-molded PET. It results in high modulus of resilience. The level of crystallinity insemi crystalline materials such as PET appears to influence its notch sensitivity. The
crystallinity of injection-molded specimens is believed to vary along different depths of theskin-core regions; therefore, knowledge of the level of crystallinity along different depths ofthe samples will help to understand its effect on the notch sensitivity of the material.

Figure 5.1[a] shows slices of V-PET samples peeled off along the specimen width at 50 mthickness per slice. The sections labeled 0 – 1.2, 1.201 – 1.740, and 1.741 – 5.00mm
respectively represent the skin, interface and core regions of the specimens. The slices give avisual understanding of the morphology of the materials. It is evident that slices from the skinregion (0 – 1.2mm) have clear color, which indicates the amorphous phase of the materials.
Deeper slices ≥ 1.2mm considered the onset of the interface show remarkable presence ofthin but growing white stripe in the middle of the clear colored layer. This signifies atransition from the amorphous skin to the crystalline phase. At ≥1.741mm there is greater
intensity and well defined white area that defines the core region. It is obvious from visualobservation that V-PET has a distinct core than R-PET where the intensity of the coreappears diminished. These visual segments of the skin and core structure of the materials areclearly depicted in Figure 5.2. The thermal properties obtained in a DSC experiment from the184slices are shown in Table 5.1. It indicates that V-PET has high enthalpy (EP) at 0.05mm slice
depth because of high molecular orientation resulting from injection molding. However, EP
decreased significantly in V-PET at the deeper skin region of 0.55mm; but increased throughthe interface (1.2mm) to the core (5.0mm) region.

5.3.1.2 Cross-sectional Observation by Polarized Light Microscopy of V-PET
The plots obtained from the gray scale images of the cross sectional area of the samples
revealed a more complex morphology where the existence of an interface between the skinand the core is clearly visible in V-PET injection moldings. Derivations from the intensity of
the gray areas were used to determine precisely the thickness of the skin and core regions, as
shown in Figure 5.2. The plots marked with W and X represent the intensity values of thegray scale in the width and thickness directions of the samples, respectively. The horizontal
and vertical cursors on the matrix (gray scale image) were adjusted along the width andthickness directions until well-defined frequency peaks were obtained. The center region in
between the two maximum peak intensities (marked with arrows) obtained from w and x
plots were taken to be the width and thickness of the core, respectively. The remaining areas
surrounding the core will be regarded as the skin. Since notches were introduced in the width
direction, the size of the skin region in this direction was noted to be about 1.5 mm on eachside of the specimen.

5.3.1.3 Mechanical Properties of V-PET at Different Notch Depths
Table 5.3 shows the yield stress of V-PET was about 55 MPa when the sample was unnotched. The introduction of notches ≤ 1.5mm deep resulted in a continuous increase in theyield stress. The increase in yield stress resulting from the deepening of the notches could be
the effect of triaxial stress resulting from the notches, which were still located within the185amorphous skin region. However, when the notches become core-deep at ≥ 2.0mm, the yieldstress decreased significantly as the specimens fractured in a brittle manner. Table 5.3
equally shows that V-PET did not fracture below 300% strain when ≤ 0.5mm deep notch wasintroduced. However, there was a drastic reduction in the percentage strain when the material
fractured in a semi-ductile manner with the introduction of 0.6 – 1.0mm deep notches. The
introduction of notches deeper than 1.5mm resulted in a brittle fracture. It is obvious that thefracture behavior of V-PET changes as the notches penetrate the skin-core regions of thematerial.

5.3.1.4 Distribution of Temperature at the Crack-tip During Deformation in V-PET
The area under the stress-strain curve of a material is the strain energy per unit volume
absorbed by the material during tensile loading. Conversely, the area under the unloading
curve is the energy released by the material during unloading. In the elastic range, the areas
under the stress-strain curve and unloading curve are equal; and no net energy is absorbed.
However, if the material is loaded into the plastic range, the energy absorbed exceeds theenergy released, and the difference is dissipated as heat. Therefore, the temperature recordedat the crack-tip is the effect the excessive energy that is dissipated as heat. The quantity ofthis heat could be influenced by crystallinity, crack depth or the hierarchical structure of thematerial.

The determination of maximum change in temperature (ΔT) was previously shown inEquation 5.1. Figure 5.4 shows the temperature distribution in V-PET specimens during
tensile test. The figure shows three distinct temperature distributions at three different notchranges as indicated by plots a, b-c, and d respectively. This covers the entire skin-core
regions of the samples. The arrow on plot a in Figure 5.4, indicates the level of change in186

temperature at full ligament yielding when there was no notch in the sample. It can be seenthat the point indicated by the arrow does not show any remarkable change from the normal
temperature distribution. Similar temperature distributions for the material occurred at semi
ductile failure, (0.5 - 0.9 mm deep notches, which is still within the skin region), as indicatedby plot b [Figure 5.4]. This marks a transition to semi ductile fracture behavior of thematerial; however, the temperature change is similar to the un-notched specimens as shownby the arrow on plot b. There was a phenomenal change in temperature and perhaps fracture
behavior when 1.0 - 1.5 mm deep notches (represented by plot c) were introduced. It is
obvious that at these notch depths there was a sudden increase in crack-tip temperature at fullligament yielding. This is evident from the sharp point indicated by the arrow on plot c. The
third temperature phase was in brittle fracture at 1.7 – 3.0 mm deep notches (depicted by plotd). The points indicated by arrows are the temperature readings corresponding with thestresses at full ligament yielding in the respective stress-strain curves at different notch
depths as shown in Figure 5.3c. The highest change in temperature occurred at these points
for the notches that attained full ligament yielding. For the brittle failures, the change intemperature was read off at the fracture point.

It is important to observe that there is no sharp point at the point indicated by the arrow onplot a. This indicates that the un-notched samples are uniform in geometry and the stress is
uniformly distributed and concentrated at the center of the material. This is evident in Figure
5.5[a] where the maximum temperatures at full ligament are located at the center of the test
pieces.

The introduction of deeper notches (0.5 – 0.9 mm) did not create significant sharp points(change in temperature) at the arrow point on plot b, which is the semi ductile phase. The
187

reason is that this range of notches was still within the skin region, and did not become
effective stress concentration points at the start of tensile test [2]. Hence, as the tensile stressincreases and crack begins to open, the stress concentration point moves from the center
towards the crack-tip as shown in Figure 5.5[b, c]. When the notch is deep enough and nearto the core region (1.0 – 1.5 mm), it becomes an effective stress concentration point rightfrom the start of tensile stress [Figure 5.5[d – f]. Therefore, at the point of shear and necking,semi ductile fracture proceeded by crazing occurred, resulting in a phenomenal change intemperature as indicated by a very sharp point indicated by arrow on plot c, and depicted in
Figure 5.6. This could be due to higher molecular chain movement, and strain-induced
crystallinity that results in semi ductile fracture preceded by high degree crazing; which is amajor mechanism for temperature change [8]. This is evident in Figure 5.5[d - f] for 1.0 – 1.5mm deep notches. Notches deeper than 1.5 mm resulted in brittle failure with zero (0.2⁰C)change in temperature, as shown in Figure 5.4 [plot d] and Figure 5.6. The corresponding
infrared images are shown in Figure 5.5[g-i].

It is important to note that the changes in temperature remained between 1.4 and 3.0⁰C as
shown in Figure 5.6 when the notch tip was not yet an effective stress concentration point.
This was revealed in the cases of un-notched to 0.9 mm deep notches depicted in Figure 5.5(a– c). This reveals an almost uniform distribution of stress on the bulk of the material, whichequally suggests some degree of structural stability in the materials. However, when thenotches became deep enough to act as an effective stress concentration point, the change intemperature increased significantly to 8 ≤ 10.2⁰C. This result is consistent with the changes inthe enthalpy and mechanical properties along the skin-core region as shown in Tables 5.1 and
5.3, respectively. It is also consistent with the skin-core structure of the material as depicted
in Figure 5.2 with the measured values in Table 5.2.
188

5.3.1.5 The Effect of Notch Sensitivity and Fracture Modes on the Change in
Temperature and the Shape of Plastic Deformation Zones in V-PET
Figure 5.6 shows the distribution of maximum change in temperature gained due todeformation processes in V-PET. The maximum change in temperature was taken at fullligament yielding of the stress-strain curves of the samples as was previously shown in Figure5.3. Figure 5.6 clearly shows that the changes in temperature from the un-notched to 0.9 mm
deep notch samples increased slightly from 1.3 to ≤ 3.0⁰C. This suggests there was a minimal
effect of notches on the material’s structure. It suggests also that longer ligament size resultsin lower change in temperature within the notch depths ≤ 0.9 mm. The structural stability ofthe material at these notch depths is demonstrated in Figure 5.5[a - c], the corresponding
infrared images taken full ligament yielding of the stress-strain curves. It shows that the stressconcentration point, with the highest temperature, is located at the center of the material. This
indicates that the lines of stress [3] are parallel to the direction of applied stress. However, as
the notches deepen and crack opening begins, the stress concentration point gradually moves
towards the crack tip, shown in Figure 5.5[b, c] representing 0.5 – 0.9 mm deep notches. It is
important to note the shape of the deformation zones. The un-notched V-PET samples[Figure 5.5(a)] have an intersecting oval-shaped deformation zones with circular disc at thecenter. On the other hand, the 0.5 - 0.9 mm deep notch samples [Figure 5.5(b, c)] have justoval-shaped deformation zones marking a transition to semi ductile fracture modes in thesamples.

Figure 5.6 shows that as the notch deepens to 1.0 mm, there is a phenomenal increase in thecrack tip temperature to about 9.0⁰C, when the notch root becomes an effective stressconcentration point. The increase in temperature continued at 1.3 and 1.5 mm deep notches
where the change in temperature was at maximum (10.1⁰C) and decreased to 9.0⁰C. The
189

phenomenal increase in temperature at 1.0 – 1.5mm deep notches could be explained by thefracture modes depicted in Figure 5.5[d - f]. The infrared images indicate a mixed mode offracture characterized by crazing at these notch depths, which are considered the terminal endof the skin to the interface between the skin and core regions. The regular shapes of thedeformation zones also point to a uniform mode of fracture that suggests semi-ductile
fracture without a clear warning about imminent brittle fracture. The major fracture
mechanism that activates the rise in temperature at these notch depths include crazing;
molecular chain alignment ahead of craze, molecular chain motion that occurs within theactive layer, and the failure of craze fibrils which involves molecular chain scission [9].
Figure 5.5[d–f] clearly shows that the fracture modes of V-PET at 1.0 – 1.5 mm deep notches
were reasonably consistent. At notches deeper than 1.7 mm, the material fractured in a brittle
manner with a zero change in temperature (0.2⁰C) as shown in Figures 5.5[g - i] and 5.6.

The degree of change in temperature in V-PET could be influenced by high crystallinity inthe material (see table 5.1), which absorbs more energy with less heat generation. Highcrystallinity equally appears to minimize the movement of molecular chains thereby reducing
the quantity of heat generated. The maximum ΔT has shown that the highest energy
dissipation in the fracture process occurred at full ligament yielding. It has also shown thatthe highest ΔT occurred at the interface between the skin and core regions when the notch
depth was ≥ 1.0 and ≤ 1.5mm.

5.3.1.6 Effects of V-Notches on the Fracture Properties of V-PET and Determination of
Essential Work of Fracture (EWF)
The area under the stress-strain curve up to a given value of strain is the total mechanicalenergy per unit volume consumed (total work of fracture) by a material in straining it to that190

value. In the absence of molecular slip and other mechanisms of energy dissipation, thismechanical energy is stored reversibly within the material as strain energy. When the stresses
are low enough that the material remains in the elastic range, the strain energy is just thetriangular area within the elastic range. If the stresses is high enough that the material extends
into the plastic range with high deformation, the strain energy is the area under the stressstrain curve.

Figure 5.7 depicts the typical stress-strain curves of specimens notched at various depths
where the deformation of the material extends into the plastic range. The un-notched and
0.5mm deep notch samples of V-PET maintained high ductility and did not fracture below
300 % strain. The introduction of a 0.6 mm deep notch, however, resulted in a drastic
reduction in the ductility of the material. This marks a change in the fracture behavior of thematerial from ductile to semi-ductile where shear yielding and tearing at the crack-tip was
evident with increasing strain.

Figure 5.7 indicates there was full ligament yielding in the notched samples, which satisfy animportant requirement for performing the essential work of fracture (EWF) analysis [16]. The
two distinct deformation zones are indicated as the process deformation zone (essential zone We), and the plastic deformation zone (non-essential zone - Wp). Therefore, the total work of
fracture (Wf) for a given notch depth can be determined by equation (5.2). The values in
Figure 5.8 show a strong linear correlation between Wf and the length of the specimens’
ligaments, l suggesting that an increase in l results in an increase in the Wf. Therefore, thespecific work of fracture (wf) is given by equation (5.3). Consequently, we and βwp are thepositive intercept (at l = 0) and slope of the linear regression line, respectively. This indicatesthat we can be used as a measure of the toughness of the material [16]. Figure 5.9 equally191

reveals a strong linear relationship between sample ligaments at different notch depths andthe process deformation zone (We). Similar to Wf, Figure 5.9 shows that We increases with anincreasing length of l. A strong linear relationship is also seen in Figure 5.10 between thelength of l of the samples and Wp, which shows that Wp increases with an increase in l.
Figures 5.8 – 5.10 have shown that increase in the length of l results in an increase in Wf, We
and Wp. This suggests a strong linear relationship between We and Wp as depicted in Figure
5.11. The figure shows that the energy dissipated in both deformation zones increases in a
linear manner as the sample ligament increases.

The modulus of resilience is the quantity of energy a material can absorb (within the process
deformation zone, We) without suffering damage. Resilience refers to the concept that up to
the point of yielding, the material is unaffected by the applied stress and upon unloading will
return to its original shape. On the other hand, modulus of toughness is the energy needed(within the plastic deformation zone, Wp) to completely fracture a material. It was explained
previously that when a material is loaded within the elastic limit, the energy absorbed is equal
to the energy released during unloading. However, if the material is loaded into the plastic
range, the energy absorbed exceeds the energy released and the difference is dissipated asheat. The slope of V-PET in Figure 5.11 indicates a decreasing difference between energies
absorbed and released as the ligament length decreases. Therefore, it could be reasonable tosuggest that the ratio of We and Wp can explain the effect of deepening notches on the
deformation zones, and their resultant effect on the total work of fracture in the material.

In trying to understand the degree of difference between the energy absorbed and the energy
released within the deformation zones, the ratio of We and Wp for the samples are plottedagainst their respective notch depths in Figure 5.12. The figure tends to suggest that the lower
192the ratio is, the more the energy consumed by the material in Wp than in We, suggesting thatthe material retains its ductility prior to a transition to mixed fracture behavior/brittle fracture.
At this transitional fracture behavior, the energy consumed in Wp decreases drastically until
Wp becomes zero. The figure shows that there is a gradual drop in We/WP ratio because ofdeepening notches (0.5 ≤ 1.0 mm). This indicates a more gradual decline in the work done in
the process (We) than the plastic (Wp) deformation zones, resulting from the increasing effectof notches on the modulus of resilience of the material. As the notches deepen, V-PET
appears to retain high modulus of resilience until sudden brittle failure. This could be thereason for the relatively high ratio of We/Wp = 0.69 at 1.0mm deep notch as shown in Figure
5.12. However, notches ≥ 1.0 mm deep resulted in a more decrease in Wp than in We at thetransition to mixed fracture mode (transition within semi ductile fracture) prior to brittle
fracture. The drastic drop in Wp results in a higher value of We/Wp. At brittle failure, when
Wp = 0, the ratio We/Wp ≈ ∞; therefore, the application of EWF becomes invalid. It is
important to note the proximity of the ratios We/Wp for V-PET, which fall between 0.74 –0.68 in decreasing order. These high ratios of We/Wp could suggest that V-PET has highmodulus of resilience resulting from high crystallinity (see table 5.1). The high ratios could
equally suggest that V-PET had a sudden transition from semi-ductile to brittle fracture atnotches deeper than 1.5mm.

5.3.1.7 Relationship between Change in Crack Tip Temperature and Ratio of Energy
Dissipated in the Process (We) and Plastic (Wp) Deformation Zones
Figure 5.13 reveals the relationship between changes in crack-tip temperature (temperature at
the stress concentration point) and the ratios of energy dissipated in both the process andplastic deformation zones (We/Wp) for samples at different notch depths. The plot of theratios and temperature indicates that the temperature changes remained constant (< 3.0⁰C)
193

between 0.5 - 0.9 mm deep notches. The minimal increase in values suggests similar fracture
behaviors in the material, and indicates that the lines of stress [3] in this range of notches
were relatively parallel to the direction of applied stress. Nevertheless, there was a drasticchange in temperature (8.96⁰C) when the notch deepens to 1.0 mm, which appears to be theborder between the skin and interface. As discussed previously there was a transition to
mixed mode of fracture, which was characterized by crazing that activated molecular motion
resulting in the generation of heat. The mixed fracture behavior continued at 1.3 and 1.5 mm
deep notches, with resultant increase in temperature to 10.1⁰C and 9⁰C respectively. Notches
deeper than 1.5 mm resulted in brittle fracture with little or no change in temperature. It isnoteworthy that 1.0 - 1.5 mm deep notches became effective stress concentration points in thespecimens right from the start of tensile test was shown in Figure 5.5(d – f).

Figure 5.13 equally appears to suggest that as the notches deepen, the ligament size decreases
resulting in lower energy absorption in We. This could result in more heat release, which
could cause an increase in the crack-tip temperature. Therefore, lower ratio We/Wp could be
related to an increase in temperature. Lower ratio of We/Wp equally indicates more work
done in the plastic deformation zone for crack propagation, suggesting retention of modulusof toughness prior to brittle fracture as the notches deepen. It also shows that We << Wp until
brittle fracture occurred. The level of change in temperature in V-PET could have resulted
from its level of crystallinity (refer to table 1); and perhaps more restricted molecular chain
motion.

194

5.3.1.8 Correlating the Notch Sensitivity Factors for Strength (kS) and Energy (kT) with
the Ratios of Work done in both Deformation Zones of EWF
Figure 5.14 depicts the correlation between notch sensitivity factors for strength (ks) and theratios (We/Wp) of work done in We and Wp deformation zones. The correlation indicates thatwhen 0.5 mm deep notch was introduced, the values of kS (0.96) and We/Wp (0.75) wereapproaching 1.0. Deeper notches resulted in simultaneous decrease in the both values,
suggesting that notches have similar effect on both parameters (kS and We/Wp). This meansthat as the material becomes more sensitive to deeper notches, the ratio of work done in bothdeformation zone decreases until brittle failure. This suggests that a decrease in yield stressdirectly affects the work done in the process deformation, with lesser effect on the work donein the plastic deformation zone. This fracture behavior subsists until mixed mode fracture
occurs at 1.0 – 1.5 mm deep notches, when the values [kS (0.93) and We/Wp (0.7)] began toincrease because of drastic decrease in Wp. When Wp = 0 at brittle failure, the values of
We/Wp tend to ∞; and EWF becomes invalid for further characterization of the material.

In Figure 5.15, the relationship between notch sensitivity factors for energy (kT) and the ratios
of work done (We/Wp) show that the 0.5 mm deep notch has kT = 1.0 and (We/Wp) = 0.75.
This means lower notch sensitivity and suggests slightly more work done in the plastic
deformation zone. The deepening of the notches to 0.6 - 1.0 mm resulted in higher but
apparent constant kT (2.75 J/m2) with corresponding decline in the ratios We/Wp (0.74) ofworks done in the two deformation zones. Decline in the ratios reveals that an increase innotch sensitivity results in a decrease in the work done in the deformation zones. Notches thatare deeper than 1.0 but ≤ 1.5 mm resulted in an increase in the ratios of works done while kT
appeared to remain constant. Therefore, the higher the notch sensitivity, the more brittle thematerial becomes. V-PET has shown to expend proportionately more energy in the process
195deformation zone than in the plastic deformation zone, where the actual crack opening andpropagation occurred.

5.3.2 Essential Work of Fracture Evaluation and the Analysis of Crack Tip
Temperature Distribution of the Fracture Modes of V-Notched R-PET Components

5.3.2.1 Measurement of Crystallinity in R-PET
The previous discussion in section 5.3.1.1 for crystallinity in V-PET is equally applicable to
R-PET. However, it is important to point out that the intensity of the white stripe in R-PET is
not as visible as in V-PET, particularly within the core region where there was a distinct skincore structure as shown in Figure 5.1[a, b]. Table 5.1 reveals also that V-PET is more
crystallized than R-PET as revealed by their enthalpy due to processing (Ep).

5.3.2.2 Cross-sectional observation by polarized light microscopy of R-PET
The methodology for cross-sectional observation used for V-PET, which was previously
discussed in section 5.3.1.2, was equally applied for R-PET. The polarized light microscopic
image of the cross section of R-PET shown in Figure 5.16 indicates that the skin and coredimensions are similar to V-PET. A major observation that is worthy of mention is thedifference in intensity of the gray appearances of the skin and core regions between V-PET
and R-PET specimens. The gray image of V-PET, which was shown previously in Figure 5.2
revealed a more distinct skin-core boundaries that contrasts with the dispersed and morediffused structure of R-PET. The gray images of V-PET and R-PET in Figures 5.2 and 5.16
reveal what appear to be meniscus shapes, which are more defined in V-PET than R-PET.
These contrasting observations are consistent with the materials’ optical appearances in
Figure 5.1[a, b].
196

5.3.2.3 Mechanical Properties of R-PET at Different Notch Depths
Table 5.3 shows that R-PET has a yield stress of about 54 MPa when the sample was unnotched. The introduction of notches ≤ 1.5mm deep resulted in a gradual increase in the yieldstresses, which could be the effect of triaxial stress resulting from the notches, which werestill located within the amorphous skin region. However, unlike V-PET the yield stress of RPET began to decrease when notches ≥ 0.9mm deep were introduced. When the notches
became core-deep at ≥ 2.0mm, the yield stress decreased significantly as the specimens
fractured in a brittle manner. Table 5.3 equally shows that R-PET did not fracture below
300% strain when ≤ 0.5mm deep notch was introduced. However, there was a drastic
reduction in the percentage strain when the material fractured in a semi-ductile manner with
the introduction of 0.6 – 1.5mm deep notches. The introduction of notches deeper than
2.0mm resulted in a brittle brittle fracture. It is obvious that the fracture behavior of R-PET
changes as the notches penetrate the hierarchical skin-core regions of the material.

5.3.2.4 Distribution of Temperature at the Crack-Tip during Deformation in R-PET
Figure 5.17[a, b] depicts the crack-tip temperature distribution of the materials during
deformation. As discussed previously in section 5.3.1.4, three different temperature changes
were observed. These indicate three different fracture modes, namely ductile, semi ductile,and brittle failure, which are represented by plots a, b-c, and d respectively. Plot a in the
figure show that V-PET and R-PET have similar temperature profile in the un-notchedgeometry (refer also to section 5.3.1.4). Nonetheless, they have different deformation shapes
as shown in Figures 5.5 and 5.18 respectively. In Figure 5.5[a] V-PET has a deformation
shape consisting of two oval-shaped deformation structures that crossed each other; incontrast, Figure 5.18[a] shows R-PET with one oval-shaped deformation structure. The shapeof the deformation zone could be a direct effect of the level of crystallinity [Table 5.1]. High197

crystallinity could have resulted in high modulus of resilience with a tendency to have more
centralized stress concentration point resulting in a double oval-shaped deformation
structures crossing each other. Conversely, low crystallinity could result in a more plastic
flow to create a single oval-shaped deformation structure that shears at 45° angle. Similar
temperature distributions for the materials occurred at semi ductile failure, (0.5 - 0.9 mm
deep notches), as indicated by plots b-c in Figure 5.17[X, Y]. At these notch depths, theshapes of the deformation zones are similar indicating a shift in the stress concentration
points towards the crack-tip as depicted in Figures 5.5 and 5.18, [b-c]. It is important toobserve that there is no remarkable sharp point in the un-notched samples as indicated by thearrows on plots a, while a mild sharp point occurred as indicated by arrows on plots b as
shown in Figure 5.17. This indicates that the un-notched samples are uniform in geometry
and the stress is uniformly distributed and concentrated at the center of the materials asshown in Figure 5.18[a]. On the other hand, similar to V-PET the samples introduced with ≤0.9 mm deep notches show evidence of structural deformation as depicted in Figure 5.18[b,c]. Figure 5.19 shows that ΔT for the un-notched samples is higher in R-PET (3.03⁰C) than
V-PET (1.38⁰C). This is equally true for the specimens introduced with ≤ 0.9 mm deep
notches, which shows that the ΔT in R-PET (3.4⁰C) is also higher than V-PET (2.2⁰C).

A phenomenal increase in temperature characterized by very sharp points when 1.0 - 1.5 mm
deep notches were introduced is represented by plot c, Figure 5.17[Y]; the infrared images
are shown in Figure 5.18[d – f]. The higher change in temperature in R-PET could be due tolow crystallinity [Table 5.1] that permits higher molecular chain movement, and straininduced crystallinity. This results in a semi ductile fracture proceeded by high degree crazing,
which is a major mechanism for temperature change [8]. Thus, higher molecular movement
results in the activation of more heat. Brittle fractures at 1.7 – 3.0 mm deep notches are198

depicted by plots d with 0.2⁰C change in temperature. The corresponding infrared images areshown in Figure 5.18[g – i].

5.3.2.5 The Effect of Notch Sensitivity and Fracture Modes on the Change in
Temperature and the Shape of Plastic Deformation Zones in R-PET
Figure 5.19 shows the distribution of maximum ΔT gained due to stress (work done) during
deformation process in R-PET. The ΔT in temperature, which was determined by equation
5.1, was taken at full ligament yielding as shown in Figure 5.3(c). The difference intemperature changes shown in Figure 5.19 clearly reveals that R-PET has higher changes intemperature than V-PET at all notch depths. The shape of the deformation zones was
discussed previously in the preceding section. The shapes of the deformation zones of the unnotched samples were seen to differ between V-PET and R-PET. The shape of thedeformation zone of R-PET appears to have resulted from molecular slip and plastic flow
arising from low crystallinity. A similarity in the shapes was observed in V-PET and R-PET
at the notch ranges of 0.5 – 0.9 mm deep. However, it is important to observe significant
differences in the deformation shapes when the notches penetrated the interface at notches ≥1.0 mm deep. Figure 5.5[d – f] shows that V-PET had a consistent deformation shape at 1.0 –1.5 mm deep notches prior to brittle fracture. In contrast, R-PET reveals changing
deformation shapes suggesting different degrees of mixed mode of fracture prior to brittle
failure, which could have provided enough warning mechanism about an imminent
catastrophic failure of the material. This warning mechanism is evident in the gradualdecrease in the yield stress of R-PET as shown in Table 5.3. Unlike V-PET, it is obvious thatthe difference in change in temperature between 1.0 – 1.5 mm deep notches in R-PET issignificant. Therefore, R-PET appeared to have provided enough warning mechanism against
imminent catastrophe with a significant change in temperature when the notch depth transited
199

from 1.3 – 1.5 mm. This observation is consistent with the deformation shapes in Figure 5.18
[d – f] discussed earlier; as well as the significant shift (isolation) of the stress-strain curvesof 1.5 mm deep notched samples at full ligament yielding [shown in Figure 5.18]. When thenotches were core-deep, brittle and catastrophic failure was imminent.

Without repetitions, sections 5.3.1.4 and 5.3.2.4 have shown that the changes in crack-tip
temperature were significantly influenced by the skin-core structure, the depth of notches,
and the materials’ notch sensitivity. Another factor that greatly influenced the changes intemperature is mixed fracture, which occurred at the terminal end of the skin region and theinterface between the skin and the core. It was observed previously that fracture at thesenotch depths were preceded with craze resulting in the generation of more heat. Higher
molecular chain movement resulting from lower level of crystallinity has also been shown to
be a major reason for higher changes in temperature in R-PET than V-PET.

5.3.2.6 Effects of V-Notches on the Fracture Properties of R-PET and Characterization
under Essential Work of Fracture (EWF)
Figure 5.20[a, b] depicts the typical stress-strain curves of V-PET and R-PET specimens withdifferent notch depths. Similar to V-PET (section 5.3.1.6), the un-notched samples of R-PET
maintained high ductility and did not fracture below 300 % strain. The fracture transitions of
V-PET and R-PET are essentially similar with full ligament yielding in the samples that wereintroduced with notches ≤ 1.5 mm, which satisfy an important requirement for performing theessential work of fracture (EWF) analysis [16]. The deformation zones were previously
discussed in section 5.3.1.6 where it was stated that the total work of fracture (Wf) for a given
notch depth can be determined by equation (5.2). It was equally stated that the values in
Figure 5.21 show a strong linear correlation between the total work of fracture and the length200

of the specimens’ ligaments. Therefore, the specific work of fracture (wf) is given byequation (5.3). Consequently, (we) and (βwp) are the positive intercept (at l = 0) and slope ofthe linear regression lines, respectively; and (we) is a measure of the toughness of the material[16]. Figure 5.22 equally reveals a strong linear relationship between different lengths ofsample ligaments and the process deformation zone (We). Similar to Wf, Figure 5.22 showsthat We increases with an increasing length of l. A strong linear relationship is also seen in
Figure 5.23 between the length of l of the samples and Wp. However, the slope of the plots
shows that the increase in Wp for R-PET when l increases is not as much as it is in V-PET.
This suggests that even at lower ligament sizes, R-PET retains high deformation. Figures 5.21
– 5.23 have shown that increase in the length of l results in an increase in Wf, We and Wp.
This suggests a linear relationship between We and Wp as depicted in Figure 5.24. The figure
shows that the energy dissipated in both deformation zones increases in a linear manner as
the sample ligament increases. However, unlike in V-PET as described in 5.3.1.6, Wp
remained almost the same in R-PET even as We increases with increasing length of ligament.
Therefore, it could be reasonable to suggest that the ratio of We and Wp can explain the effectof deepening notches on the deformation zones, and their resultant effect on the total work of
fracture in the material.

As mentioned in 5.3.1.6, the plots of the ratios of We and Wp against notch depths for V-PET
and R-PET are compared in Figure 5.25 in trying to understand the degree of difference
between the energy absorbed and energy released (as heat) in both deformation zones. Unlike
V-PET, the lower ratio of We/Wp in R-PET in Figure 5.25 shows that R-PET retained its
plastic deformability even with short ligament lengths. It also reveals that even as themodulus of resilience decreased in R-PET, its modulus of toughness did not decrease in a
drastic manner. This is consistent with previous observation that R-PET had adequate
201

warning mechanism prior to catastrophic (brittle) fracture. It can be concluded that ratios
closer to 1.0 indicates high modulus of resilience; conversely, ratios further away from 1.0
indicates higher ductility (modulus of toughness). The retention of higher deformability
within the notches in R-PET could have resulted from morphological changes, which was theeffect of plastic flow from the more amorphous skin region to the crack tip.

5.3.2.7 Relationship between Change in Crack Tip Temperature and Ratio of Energy
Dissipated in the Process (We) and Plastic (Wp) Deformation Zones in R-PET
It was described earlier that when a material is loaded into the plastic region, the energy
absorbed exceeds the energy released, and the excessive energy is released as heat. The
released heat energy causes temperature rise in the material, especially at the crack-tip. The
quantity of the excessive energy released as heat affects the rise in crack-tip temperature.
Similar to the discussion on V-PET in section 5.3.1.7, Figure 5.26 reveals that when thenotches were still within the skin region (0.5 – 0.9mm deep), the energy released as heat waslow, arising from the small difference between energy absorbed and released. This resulted in
low crack-tip temperature as shown in Figure 5.26. The plots of the ratios of energy
dissipated in the process and plastic deformation zones (We/Wp) and temperature indicate anapparent constant temperature (3.55 - 3.85⁰C) between 0.5 - 0.9 mm deep notches. The
minimal increase in values suggests similar fracture behaviors in the material, and indicatesthat the lines of stress [3] in this range of notches were relatively parallel to the direction of
applied stress. However, when the notch deepens to the interface (1.0mm) between the skinand core, the difference between energy absorbed and released increases, resulting in morequantity of energy that is released as heat. Consequently, there was a drastic increase intemperature (12.78⁰C), which was greater than recorded in V-PET (8.96⁰C). As discussed
previously there was a transition to mixed mode of fracture, which was characterized by202

crazing (refer to Figure 3.15 in Chapter 3) that activated molecular motion resulting in thegeneration of heat. The mixed fracture behavior continued at 1.3 and 1.5 mm deep notches,with resultant increase in temperature to 18.07⁰C and 13.88⁰C respectively, which is muchhigher than the increase in V-PET (10.13⁰C and 9.0⁰C). This suggests that the difference
between energy absorption and energy release in R-PET is greater than in V-PET at thesenotch depths. Notches deeper than 1.5 mm resulted in brittle fracture with 0.2⁰C change intemperature. It is noteworthy that 1.0 - 1.5 mm deep notches became effective stressconcentration points in the specimens right from the start of tensile test [Figure 5.18(d – f)].

Figure 5.27 is consistent with previous observation, which revealed that the closer the ratio
We/Wp is to 1.0, the higher the modulus of resilience in the material, particularly in samples
that failed in a semi ductile manner. However, a ratio of We/Wp further away from 1.0
indicates more work done in the plastic deformation zone for crack propagation, suggesting
good retention modulus of toughness, and less effect of notches. It also suggests that the moreductile a material is, the higher the change in temperature, which confirms that there is highplastic flow from the amorphous skin to the crack tip during deformation process. The levelof change in temperature in R-PET could have resulted from its low crystallinity (refer totable 5.1); and perhaps more unrestricted molecular motion.

5.3.2.8 Correlating the Notch Sensitivity Factors for Strength (kS) and Energy (kT) withthe ratios of Work done in both Deformation Zones of EWF
Figure 5.27 compares the correlation between notch sensitivity factors for strength (ks) andthe ratios (We/Wp) of work done in We and Wp deformation zones. Similar to V-PET, thecorrelation indicates that when 0.5 mm deep notch was introduced, the values of kS (0.96)
and We/Wp (0.75) were approaching 1.0. Note that if notches do not have detrimental effect203on the materials, kS = 1.0, otherwise it kS > 1.0 [2, 14]. The values of kS < 1.0 recorded at 0.5
mm indicates the effects of triaxial stress resulting from the introduction of notch, whichraises the yield stress above un-notched samples. Unlike V-PET in section 5.3.1.8, deeper
notches resulted in a decrease in the values of We/Wp while kS increased in R-PET,
suggesting that notches decreased both the yield stress and the work done in the process
deformation zone (We), while the work done in the plastic deformation zone remained
apparently constant. This means that as the material becomes more sensitive to deeper
notches, the ratio of work done in both deformation zone decreases until brittle failure,suggesting that a decrease in yield stress directly affects the work done in the process
deformation zone, with lesser effect on the work done in the plastic deformation zone. This
fracture behavior subsists until mixed fracture occurred at 1.0 – 1.5 mm deep notches, whenthe values [kS (0.99) and We/Wp (0.59)] began to increase because of drastic decrease in Wp.
When Wp = 0 at brittle failure, the values of We/Wp tend to ∞; and EWF becomes invalid forfurther characterization of the material. It is obvious from Figure 5.27 that the curves of RPET and V-PET moved in opposite directions prior to brittle fracture, which indicates that theyield stress of R-PET is more sensitive to notches than V-PET.

Figure 5.28 shows the relationship between notch sensitivity factor for energy (kT) and theratios of work done (We/Wp). It shows that at the 0.5 mm deep notch R-PET has kT = 0.95
and (We/Wp) = 0.75, which suggests that more work is done in the plastic deformation zone.
The deepening of the notches to 0.6 - 1.0 mm resulted in higher but constant kT (2.73 J/m2)
with corresponding decline in the ratios We/Wp (0.59). Decline in the ratios reveals that anincrease in notch sensitivity results in a decrease in the work done in the deformation zones.
Notches that are deeper than 1.0 but ≤ 1.5 mm resulted in an increase in the ratios of works
done while kT appeared to remain constant. Therefore, the higher the notch sensitivity, the204more brittle the material becomes. In contrast with V-PET, R-PET has shown to dissipate
more energy in the plastic deformation zone (where the actual crack opening and propagation
occurred) than in the process deformation zone. The less notch-sensitivity of R-PET over VPET is made clear by the values of kT and We/Wp at brittle fracture. It is obvious that while
V-PET had mixed fracture prior brittle failure immediately after kT = 2.8 and We/Wp = 0.69,
R-PET had mixed fracture at kT = 2.73 and We/Wp = 0.59. Therefore, it can be concluded that
R-PET is less notch-sensitive in respect of toughness than V-PET.

5.4 Conclusions
Results revealed that the skin-core structure of the materials greatly influenced the fracturemodes and crack tip temperature of PET. Notch depth is another factor that had tremendous
effect on the temperature changes, especially in the process deformation zones of thematerials. It was observed that a phenomenal change in crack tip temperature occurred innotches that are as deep as the terminal end of the skin region, as well as the interface
between the skin and the core. Results reveal that the major fracture mechanism that activatedhigh change in temperature at these notch depths is crazing. Good correlations betweenchange in temperature, notch depth, notch sensitivity factors for energy, and the parametersof essential works of fracture were established. It can be concluded that the crack tip
temperature depends on the notch sensitivity of PET injection moldings. The results of thisinvestigation give valuable information on the thermoplastic and thermodynamic fracture
behavior of PET injection-moldings when notches of different depths were introduced.

205

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14. Takano, M., & Nielsen, L. E., The Notch Sensitivity of Polymeric Materials. J. App.
Pol. Sci. 20, 2193-2207 (1976).
15. Karger-Kocsis, J., For what kind of polymer is the toughness assessment by the
essential work concept straightforward?. Polymer Bulletin, 37, 119-126 (1996).
16. Chena, H., Karger-Kocsis, J., & Wu, J., Effects of molecular structure on the essential
work of fracture of amorphous copolyesters at various deformation rates. Pol. 45,
6375-6382 (2004).
17. Kuno, T., Yamagishi, Y., Kawamura, T., & Nitta, K., Deformation mechanism underessential work of fracture process in polycyclo-olefin materials. eXPRESS Pol.
Letters, 2, 404-412 (2008).

207

Table 5.1: The Crystallinity of V-PET and R-PET Injection Moldings
Cold crystallization (EC)Materials

Melting (EM)EP = (EM – EC)

Slice depth (m)Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

Temp.
(⁰C)

Enthalpy
(J/g)

50

34.1

125.9

69.5

257.4

35.4

550

20.3

126.0

44.8

256.8

24.5

1400

22.2

124.6

46.2

256.7

24.0

2500

10.1

125.2

44.5

256.9

34.4

50

26.8

130.7

37.3

253.6

10.5

550

26.1

130.5

37.2

253.7

11.1

1400

35.0

130.7

49.9

232.9

14.9

2500

23.1

130.4

39.4

253.4

16.3

V-PET

R-PET

208

Table 5.2: Gray values from gray image plots of the cross-sectional view of dumbbell samples showing the sizes of the skin,interface and core layers of V-PET and R-PET

Width direction (mm)Thickness direction (mm)Region of Specimen
V-PET

R-PET

V-PET

R-PET

Skin region

1.47 (x2)1.47 (x2)0.65 (x2)0.65 (x2)Interphase

0.53 x 2

0.53 x 2

0.11 (x2)0.11 (x2)Core region

5.98

5.98

1.46

1.46

Total

9.98

209

2.98

Table 5.3: Effects of different depths of V-notches on the yield stress (MPa) and strain (%) of V-PET and R-PET
Materials
Notch depth(mm)Materials’ Properties

Un-notched
0.5
0.6
0.7
0.8
0.9
1.0
1.5
2.0
2.5
3.0
3.5
4.0
4.5

V-PET

R-PET

Yield stress

Strain (%)

Yield stress

Strain (%)

55.2
58.6
59.0
59.7
59.9
60.3
60.7
61.1
55.3
51.5
48.9
45.9
43.5
41.8

> 300
> 300
14.9
14.3
14.1
14.0
14.3
11.3
2.6
2.1
1.9
1.7
1.5
1.2

54.1
56.2
57.6
58.1
58.2
57.9
57.7
54.1
51.3
47.7
44.2
41.8
39.6
37.0

> 300
> 300
14.9
14.5
14.3
14.7
15.0
12.7
3.7
2.0
1.7
1.4
1.3
1.2

210

a

0 – 1200mm

1201– 1740mm

1741 – 5000mmb
Figure 5.1: DSC samples taken from slices of dumbbell samples at 50mm per slice which were peeled off from thethickness side showing the skin/core appearance of the materials at different slice depths [a = V-PET and b = R-PET]
211

W

X
Figure 5.2: Gray image plot showing the gray values of the skin, interface and core regions of V-PET[W = graph of gray values on the width direction, X = graph of gray values on the thickness direction].

212

70
75%
60

b

87.50%

Stress (MPa)50

Yielding

c

40

FLY
Tearing

30
20

a

10
0
0

3

6

Strain (%)

9

12

15

Figure 5.3: Determination of maximum change in temperature (MΔT) at full ligament yielding of the stress-strain curvefor the different notch depths; [inset infrared images (a) before yielding, (b) at yielding, and (c) full ligament yielding]213

Figure 5.4: Crack tip temperature distributions of V-PET samples showing degree of changes
at the skin, interface and core deep regions [a = un-notched, b = 0.5-0.9 mm, c = 1.0-1.5 mm
and d = 1.7-3.0 mm deep notches; arrows indicate the temperature at full ligament yielding]214

Figure 5.5: Infrared images of VPET showing the fracture modes at different notch depths where maximum
changes in temperature occurred at maximum at full ligament yielding [a = Un-notched samples, (b and c)
represent 0.5 - 0.9mm, d = 1.0mm, e = 1.3mm, f = 1.5mm, and (g – i) represent 2.0 – 3.0mm].
215

Maximum change in temperature
at full ligament yielding (⁰C)

20
V-PET
16
12
8

4
0
0

0.5

1
1.5
2
2.5
V-notch depth (mm)3Figure 5.6: Maximum change in notch root temperature (V-PET) in the process zone showing drastic rise at theinterface (1.0-1.5 mm notch) between the skin and core regions where mixed fracture occurred with crazing.
216

Figure 5.7: Stress/Strain curves of V-PET showing the essential (We) and non-essential (Wp) works of fracture
deformation zones for the specimens at different notches ranging from un-notched (Un) to 1.5 mm].
217

Total work of fracture (wf)20
16
y = 2.1217x - 8.2114
R² = 0.909

12
8

4

VPET

0
8.2

8.45

8.7
8.95
Ligament (mm)

9.2

9.45

Figure 5.8: Total work of fracture for V-PET at different ligament depths

218

Essential work of fracture (We)

160
120

80
40
0
V-PET

8

8.5
9
9.5
Ligament width (mm)10Figure 5.9: Essential work of fracture for V-PET at different ligament depths

219

Non essential work of fracture
(WP)

200
150

100
50
0

V-PET

8

8.5
9
9.5
Ligament width (mm)10Figure 5.10: Non-essential work of fracture for V-PET at different ligament depths

220Non-essential work of fracture
(Wp)

240
200
160

120
80
40

VPET

0
90

100 110 120 130 140
Essential work of fracture (We)150

Figure 5.11: Energy dissipation in the process (We) and plastic
deformation zones (Wp) with increasing ligament length221

Ratio of essential (we) and nonessential (wp) works of fracture(we/wp)

0.8

VPET

0.55
0.3

0.6

0.9
1.2
1.5
V-notch depth (mm)1.8

Figure 5.12: Ratio of essential (We) and non-essential (Wp) works
of fracture for the materials at the various notch depths.
222

Change in temperature (⁰C)

20
16

VPET

12
8
4
0
0.55
0.6
0.65
0.7
0.75
0.8
Ratio of essential (we) and non-essential (wp)works of fracture (we/wp)

Figure 5.13: Correlation between change in temperature and the ratio of energy
dissipated in both the process (We) and plastic (Wp) deformation zones.
223

Ratio of essential (We) and non
essential (Wp) work of fracture(We/Wp)

0.8

VPET

0.55
0.92
0.94
0.96
0.98
1
1.02
Notch sensitivity factor for strength (kS)

Figure 5.14: Correlation between We/Wp and notch sensitivity factors for strength224

Notch sensitivity factor forenergy (J/m2)

3

2
1
VPET
0

0.55
0.6
0.65
0.7
0.75
0.8
Ratio of essential (we) and non-essential (wp)works of fracture (we/wp)
Figure 5.15: Notch sensitivity factors for energy kT and the ratio of essential andnon-essential work of fracture (We/Wp) for the V-PET at the various notch depths.
225

W

X
Figure 5.16: Gray image plot showing the gray values of the skin, interface and core regions of R-PET[W = graph of gray values on the width direction, X = graph of gray values on the thickness direction].

226

X

Y

Figure 5.17: Crack tip temperature distributions of X = V-PET and Y = R-PET samples showing degree ofchanges at the skin, interface and core deep regions [a = un-notched, b = 0.5-0.9mm, c = 1.0-1.5mm
and d = 1.7-3.0mm deep notches; arrows indicate the temperature at full ligament yielding]227

Figure 5.18: Infrared images of R-PET showing the fracture modes at different notch depths where maximum changes in temperature
occurred at full ligament yielding [a = Un-notched samples, (b and c) represent 0.5-0.9mm, d = 1.0mm, e = 1.3mm and f = 1.5mm].
228

Maximum change in temperature
at full ligament yielding (⁰C)

20
V-PET
R-PET

16
12
8

4
0
0

0.5

1
1.5
2
2.5
V-notch depth (mm)3Figure 5.19: Comparing maximum change in temperature between V-PET and R-PET in the Process zone showing drastic
rise at the interface (1.0-1.5 mm notch) between the skin and core regions where mixed fracture occurred with crazing.
229

a

b

Figure 5.20: Stress/Strain curves showing the essential (We) and non-essential (Wp) works of fracture deformation
zones for the specimens at different notches ranging from un-notched (Un) to 1.5mm [a = V-PET and b = R-PET].

230

Total work of fracture (wf)20
16
y = 2.1217x - 8.2114
R² = 0.909

128

y = 1.1862x - 0.9113
R² = 0.7389

4

VPET
RPET

0
8.2

8.45

8.7
8.95
Ligament (mm)

9.2

9.45

Figure 5.21: Total work of fracture at different ligament depths

231Essential work of fracture (We)

160
120

80
40

V-PET
R-PET

0
8

8.5

9

9.5

10

Ligament width (mm)Figure 5.22: Essential work of fracture for V-PET and R-PET at different ligament depths

232Non essential work of fracture
(WP)

200
150

100
50

Wp VP
Wp RP

0
8

8.5
9
9.5
Ligament width (mm)10Figure 5.23: Essential work of fracture for V-PET and R-PET at different ligament depths

233

Figure 5.24: Energy dissipation in the process (We) and plastic deformation zones (Wp)234

Ratio of essential (we) and nonessential (wp) work of fracture(we/wp)

0.8

VPET
RPET

0.55
0.3

0.6

0.9
1.2
1.5
V-notch depth (mm)1.8

Figure 5.25: Ratio of essential (We) and non-essential (Wp) work of
fracture for the materials at the various notch depths.
235

Change in temperature (⁰C)

20
VPET
RPET

16

12
8
4
0
0.55

0.6
0.65
0.7
0.75
0.8
Ratio of essential (we) and non essential
works of fracture (wp) [we/wp]

Figure 5.26: Correlation between change in temperature and the ratio of energy
dissipated in both the process (We) and plastic (Wp) deformation zones.
236

Ratio of essential (We) and non
essential (Wp) work of fracture(We/Wp)

0.8

RPET
VPET

0.55
0.92
0.94
0.96
0.98
1
1.02
Notch sensitivity factor for strength (kS)

Figure 5.27: Correlation between We/Wp and notch sensitivity factors for strength for V-PET and R-PET
237

Notch sensitivity factor for energy(KT) (J/m2)

3

2

1
VPET
RPET

0
0.55
0.6
0.65
0.7
0.75
0.8
Ratio of essential (we) and non essential (wp)works of fracture (we/wp)

Figure 5.28: Notch sensitivity factors for energy kT and the ratio of essential andnon-essential works of fracture (We/Wp) for the materials at the various notch depths.

238

Chapter 6
Investigations on the Fracture Toughness of GFPP Composites

6. 1

Introduction

Over the years, the world has lost greatly, in human and material resources, to accidents
caused by sudden failure in engineering components [1]. In many cases, flaws, cracks, voids,
cause the failures in engineering components, as there tend to be stress concentration aroundsuch flaws. Poor workmanship and non-strict adherence to engineering principles also causesome failures. It is obvious that engineers cannot be very sure that a material is flaw free andthis makes it a common practice to assume that flaws of some chosen size will be present incomponents. Therefore, the design of critical components are based on linear elastic fracture
mechanics, (LEFM) or elastic plastic fracture mechanics (EPFM), depending on the fracture
behavior of the composite materials used, in order to prevent sudden and catastrophic failuresof structures. In fracture mechanics, there are three modes of fracture denoted as ‘mode I’,‘mode II’ and ‘mode III’ fractures [2 - 4]. These different modes are illustrated in Figures 6.1
to 6.3 and the arrows show the direction of applied force. In this work, investigation is on themode I fracture of composite materials which is shown in Figure 6.1.

In this investigation, focus is on the fracture toughness of composite materials, which is very
important for fabricating components with a predictable life span. Fracture toughness is anindication of the amount of stress and/or strain energy (K1C and G1C) required in propagatinga pre-existing crack at fracture initiation. It is a very important material property since theoccurrence of flaws is not completely avoidable in the processing, fabrication, or service lifeof a material. The fracture toughness of three grades of polypropylene (PP) with different
molecular weights, which were reinforced with glass fiber (GF) was evaluated. Fracture239

toughness is used to characterize the toughness of these composites in terms of the maximum
Stress Intensity Factor, K1max, and the maximum Strain Energy Release Rate, G1max, atmaterial failure. By extension, in the case of composite materials that do not conform to theprinciples of LEFM, EPFM applies by using J-Integral (J1max) or crack tip opening
displacement (CTODmax). Such materials would exhibit significant levels of plastic
deformation and show non-linearity in its stress/strain relationship. The fracture toughness of
GFPP composites was determined by using single-edge-notched tensile (SENT) test. Izod andCharpy impact tests were also conducted on the notched and un-notched composites.

6. 2

Characterization Methods

6.2.1 Material and sample preparation
Idemitsu Co. Ltd., Japan, compounded and injection molded the dumbbell and rectangularshaped specimens that were used for the subsequent hot water immersion tests. The
composites consist of three grades of PP of different molecular weights namely GFPP (J700GP), GFPP (900GP) and GFPP (J-3000GP), (refer to Table 6.1 for details) which were
reinforced with GF and added with MAPP. The material compositions are noted below Table
6.1, while the specimen geometries are indicated in Figure 6.4 (a) and (b).

6.2.2 Tensile Tests and Determination of Mechanical Properties and Fracture
Toughness
Tensile test was conducted on both un-notched and single-edge-notch of GFPP samples. At
least five sample pieces were tested. Single-edge-notch was introduced with Ueshima N0.
820261 notching machine from Hitachi, Ltd. Tokyo, Japan, mounted with a V-shaped blade.
A fresh razor was used to introduce a sharp crack. The size ‘a’ of the notch was determined240

so that the ratio of ‘a’ to the sample width ‘W’ falls between 0.45≤ a/w ≤ 0.55. All tests wereconducted at room temperature of 25°C. Tensile test was performed on sample span length of
115mm at a crosshead speed of 10mm/min with an Instron Universal Tester 4466 that was
attached with a 10kN load cell. Tensile test was carried out on the un-notched test pieces inaccordance with ASTM D 3039. ASTM D 5045 was adopted for use in tensile testing of thenotched specimens.

The fracture toughness of the composites investigated under the LEFM was calculated byusing the following equations:

K1max   max  a  2 1.12  0.23  Wa   10.6  Wa   21.7  Wa   30.4  Wa 
1G1max 

2

3

4



2
K1max
E

(6.1)

(6.2)

where E = Young’s modulus measured in GPa.
The values of KQ and GQ were also calculated following a similar adaptation from ASTM D
5045–99. The ASTM model is actually for compact tension test in a plane-strain situation;
however, it is adapted only for the determination of σQ (which is the stress at 5ᴼ off thetangent of a stress-strain curve, refer to Figure 6.5 for the determination of K Q (which is aconditional K1max ). The value  Q is imputed into equation (7.1) above for the calculationof KQ  K1max , thus:

2
3
4
1 


a
a
a
a 
KQ   Q   a  2  1.12  0.23    10.6    21.7    30.4   
 w
 w
 w
 w 




241

(6.3)

According to the standard if σQ > σmax then σmax is to be used to calculate KQ otherwise σQ isused.

6.2.3 Scanning Electron Microscopy (SEM)
Micrographs were taken of the fracture surfaces of the tensile samples by the use of JEOL
JSM5200 Scanning Electron Microscope equipped with Olympus Digital Camera. The
micrographs are shown in subsequent chapters relating to each composite material.

6.2.4 Charpy and Izod Impact Tests
Charpy and Izod impact tests were conducted on both un-notched and notched test pieces inorder to determine the toughness of the materials. A 2 mm size notch was introduced with aType A-3 Notching Machine manufactured by Toyo Seiki manufacturing Co., Ltd. Ten
sample pieces were tested for each material composition. Both tests were conducted at room
temperature of 25ᴼC, with an Impact Tester model DG-UB by Toyo Seiki Manufacturing,Co., Ltd. The Impact Tester was mounted with standardized JIS K7110 hammer bearing theenergy of 5.5J and 2.75J for Charpy and Izod impact tests respectively.

6.3 Results and Discussions

6.3.1 Mechanical properties and fracture toughness of GFPP from tensile data
Fracture toughness of the composites was investigated under linear elastic fracture
mechanics. The maximum stress intensity factor (K1max) and maximum strain energy release
rate (G1max) under mode I fracture was calculated from the tensile data of the notched
specimens by the use of equations (6.1) and (6.2). KQ and GQ are calculated from equation(6.3) by applying the stress values determined at 5ᴼ off tangent of the stress-strain curves as242

shown in Figure 6.5 [5]. The results are shown in Figures 6.6 to 6.10. It can be seen that
GFPP (J3000GP) achieved the highest tensile strength and stiffness as compared to GFPP (J900GP) and GFPP (J-700GP). The materials had similar percentage of elongation at break, asevident from Figure 6.7. The notch sensitivity of GFPP is depicted in Figure 6.8. The value ofnotched stress are calculated based on the cross-sectional area of the remaining ligaments ofnotched samples and show that notches had significant effect on the tensile properties of
GFPP when compared with the tensile strength in Figure 6.6. It can be seen that when the
stress of the notched specimens were calculated based on the cross sectional area of theligament, GFPP (J-700GP) had the highest stress ratio of 58.7 between the notched and unnotched specimens. It also had the highest ratio of 30.4 between the tensile strength and the
maximum stress determined by treating the cross sectional area of the notched specimens as
if it were not notched. GFPP (J-900GP) had ratios of 54.6 and 28.4 respectively, while GFPP
(J-3000GP) had ratios of 54.6 and 28.4 for both properties. This implies that GFPP (J-700GP)
is less resistant to the effect of notch than the others are. The values of fracture toughness areshown in Figures 6.9 and 6.10. The values of K1max and G1max further show that GFPP (J700GP) had the lowest resistance to fracture. On the other hand, GFPP (J-3000GP) showed
the highest resistance to fracture with the highest fracture toughness value. Similar tendency
could be observed when the KQ and GQ values were calculated and represented in Figures 6.9
and 6.10. There is higher ductility in GFPP (J-3000GP) than the other materials. This isevident in the micrographs of the fracture surfaces of the materials shown in Figures 6.11 to
6.13. The high melt flow rate and less molecular chain entanglement of GFPP (J-3000GP)
incorporated with Maleic anhydride grafted PP might have caused a better wetting of GF by
the PP resin thus resulting in better interfacial bonding and higher mechanical properties.

2436.3.2 Scanning Electron Microscopy of GFPP
The region indicated as ‘X’ in Figures 6.11 to 6.13 shows that there is a good interfacialadhesion between GF and matrix PP. The zone labeled ‘a’ is the notch, ‘b’ is the damage(plastic deformation) zone and ‘c’ is the zone of brittle failure. It can be seen from the area ofthe micrographs marked ‘b’ and ‘c’ that there were two distinct zones in the fracture processof notched samples. It follows that when stress is applied on a pre-cracked material the stresswould concentrate at the crack tip. As the applied stress grows in magnitude, the matrixreinforcement interfacial strength resists the fracture until a critical stress where theinterfacial strength is overcome and then fracture occurs. The region marked ‘b’ in the
micrographs shows the level of damage in the specimens as the fracture propagates from thecrack tip towards the opposite end where there is a brittle failure. It can be seen from themicrographs that the fracture surfaces of the ductile region ‘b’ increases while the brittle
region ‘c’ decreases with decreasing molecular weight of PP. This means that the ductility
and fracture toughness of GFPP (J-3000GP) is higher than the rest of the specimens. Thisconfirms the values of the mechanical properties and fracture toughness that were discussed
in subsection 0. The zone labeled ‘b’ (damage zone) shows that there is significant plastic
deformation in the fracture surface of the composites. This implies that further
characterization of the materials is necessary under Elastic Plastic Fracture Mechanics(EPFM) using J-integral and Crack Tip Opening Displacement (CTOD).

6.3.3 Impact Properties of Un-notched and Notched GFPP Samples
The results of both Charpy and Izod impact tests conducted on GFPP are depicted in Figures
6.14 and 6.15. In the results, it was evident that the three materials responded in a similarmanner to both Charpy and Izod impact tests, with GFPP (J-3000GP) recording the highest
impact strength while GFPP (J-700GP) had the lowest impact strength for the un-notched and244notched samples. The impact strength ratio of notched and un-notched samples show that
GFPP (J-700GP) was more sensitive to notch with a ratio of 25 for Charpy impact test and 20
for Izod impact test. GFPP (J-3000GP) showed greater resistance to impact fracture with aratio of 23.6 and 19.7 respectively, for Charpy and Izod impact tests.

6.4 Summary and Conclusions
In this Chapter, the mechanical properties and fracture toughness of three grades of PP withdifferent molecular weights that were reinforced with GF and incorporated with Maleic
Anhydride Grafted Polypropylene (MAPP) were investigated. The study revealed thefollowing:
1. GFPP (J-3000GP) consistently showed the highest mechanical properties and fracture
toughness than GFPP (J-900GP) and GFPP (J-700GP) in tensile, Charpy and Izod
impact tests.
2. The results show that MAPP as a compatibilizer had greater interfacial bonding effects
on GFPP of lower molecular weights and high melt flow rates than those of higher
molecular weights and low melt flow rates.
It can be concluded from this study that incorporating MAPP into GFPP composites
achieved better properties with Polypropylene of lower molecular weight and higher melt
flow rate. Hence the mechanical and fracture properties of GFPP depend largely on themolecular weight of polymer matrix and the compatibilisation between GF and PP.

References
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2. David R., Introduction to Fracture Mechanics, Massachusetts Institute ofTechnology, Cambridge, 7, (2001).
245

3. Jireh J. Y., Energy Concepts for Fracture, Virginia Tech. Materials Science andEngineering, (1997).
4. Fracture Toughness, Education Resources, NDT Course Material, Materials andProcesses, (2006).
5. ASTM D 5045 – 99 Standard Test Methods for Plane-Strain Fracture Toughness
and Strain Energy Release Rate of Plastic Materials, ASTM International, WestConshohocken, United States, (2005).
6. Anderson, T. L., Fracture Mechanics: Fundamentals and Applications, CRC Press,
Boca Raton, (1991).
7. Barsom, J. M., ed., Fracture Mechanics Retrospective, American Society for
Testing and Materials, Philadelphia, (1987).
8. Collins, J. A., Failure of Materials in Mechanical Design, Wiley, (1981).
9. Courtney, T. H., Mechanical Behavior of Materials, McGraw-Hill, New York,(1990).
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Through the Floor, Princeton University Press, (1976).
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New York, (1973)
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with Applications, Duxbury Press, Boston, (1986).
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133, (1965).

246

15. Tetelman, A. S., and McEvily, A. J. Jr., Fracture of Structural Materials, Wiley,
New York, (1967).
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17. Portela A., Dual Element Analysis of Crack Growth, Computational Mechanics
Publications, (1993).
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Science Publishers, (1992).
19. Rolfe, Stanley, Barson J., Fracture and Fatigue Control in Structures, Prentice
Hall, Inc., (1977).
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21. Shackelford, J., Introduction to Materials Science for Engineers, Macmillan
Publishing Company, (1985).

247

Table 6.1: Properties of Polypropylene Compounded with Glass Fiber
Materials Designation

MW of PP

MW/Mn of PP

MFR of PP (g/10min)

GFPP (J-700GP)

310,000

4.3

8

GFPP (J-900GP)

230,000

4.1

13

GFPP (J-3000GP)

190,000

4.7

30[Note: Each composite contains 20wt % of GF, 79.2wt % of PP and 0.8wt % of Maleic Anhydride Polypropylene]

248

249
Figure 6.1: Mode I Fracture

Figure 2.1-1: Mode I Fracture

250
Figure 6.2: Mode II Fracture

Figure 2: Mode II Fracture

251
Figure 6.3: Mode III Fracture

Figure 3: Mode III Fracture

WidthWidth252

Thickness, B

(a)

(b)

‘B’ Thickness

Figure6.4: Dumbbell and rectangular specimens for (a) tensile and (b) impact tests
Figure 3.1-1: Dumbbell and rectangular specimens for (a) tensile and (b) impact tests

σQ

25

σmax

253

Stress (MPa)20

15

Failure Stress
10

5

0
0

0.002

0.004

0.006

0.008

StrainFigure 6.5: Determination of σQ at 5ᴼ off tangent

0.01

0.012

90
4.2

Modulus

75

3.6

60

3.0
2.4

45
1.8
30
1.2
15

0.6

0

0.0
GFPP (J700GP)GFPP (J900GP)GFPP (J3000GP)

Figure 6.6: Tensile Strength and Modulus for Un-notched Specimens

Modulus (GPa)254

Strength (MPa)Strength

4.0
3.5

255

Elongation at Break (%)

3.0
2.5
2.0
1.5
1.0
0.5
0.0
GFPP (J700GP)GFPP (J900GP)

Figure 6.7: Elongation at Break (%) of the Un-notched Specimens

GFPP (J3000GP)

256

Tensile strength, Notch and Max. stresses

90
80
70
60
50
40
30
20
10
0
GFPP (J700GP)GFPP (J900GP)

Tensile strength

Notch stress

GFPP (J3000GP)
Max. stress

Figure 6.8: Notch Sensitivity of GFPP

9

K1max

KQ

8

257

Kmax/KQ (MPa m1/2)

7
6
5
4
3
2
1
0
GFPP (J700GP)GFPP (J900GP)GFPP (J3000GP)

Figure 6.9: Maximum Stress Intensity Factor (K1max/KQ) for GFPP

16

Gmax

GQ

14

258

G1max /GQ(kJ/m2 )

12

10

8

6

4

2

0

GFPP (J700GP)GFPP (J900GP)GFPP (J3000GP)

Figure 6.10: Maximum Strain Energy Release Rate (G1max/GQ)

GFPP (J700GP)

X

259
Figure 6.11: Micrograph of GFPP (J-700GP) [a = notch region, b = plastic/semi ductile region,c = brittle region, x = fiber-matrix adhesion]

GFPP (J-900GP)

X

260
Figure 6.12: Micrograph of GFPP (J-900GP) [a = notch region, b = plastic/semi ductile region,c = brittle region, x = fiber-matrix adhesion]

GFPP (J3000GP)
X

261
Figure 6.13: Micrograph of GFPP (J-3000GP) [a = notch region, b = plastic/semi ductile region,c = brittle region, x = fiber-matrix adhesion]

14

50

12
10

40

8
30
6
20

4

10

2

0

0

Strength (KJ/m 2) - Notched

262

Strength (KJ/m 2) - Un-notched

60

GFPP
GFPP
GFPP
(J700GP) (J900GP) (J3000GP)
Un-notched

NotchedFigure 6.14: Charpy Impact Strength for Un-notched and Notched Specimens

9

40

8

35

7

30

6

25

5

20

4

15

3

10

2

5

1

0

0
GFPP
(J700GP)Strength (kJ/m 2) - (Notched)

Strength (kJ/m 2) - Un-notched

263

45

GFPP
GFPP
(J900GP) (J3000GP)

Un-notched

NotchedFigure 6.15: Izod Impact Strength for Un-notched and Notched Specimens

Chapter 7
Characterization of the Fracture Toughness of GFPP Composites under
Elastic Plastic Fracture Mechanics

Abstract
The accurate prediction of scientific and statistical parameters depends, largely, on thefunctional correlation between the variables involved, as well as their definite and predictablebehaviors around the line of best fit (the regression line). Hence, an accurate and reliableprediction of the fracture toughness and mechanical properties of glass fiber reinforced
polypropylene (GFPP) composites is very important in order to prevent sudden failure ofcomponents fabricated from them. It also gives the opportunity to save cost on
experimentation. Three grades of polypropylene (PP) of different molecular weights (MW)
had glass fiber reinforcements each. The incorporation of Maleic anhydride grafted PP (MAg-PP) into the composites was for compatibilisation between the matrix and reinforcements.
There was tensile testing on both the un-notched and notched dumbbell samples. The use of
J-Integral analysis was to investigate the fracture toughness of the composites. There was anevaluation of the effect of different MW of PP on the fracture toughness and impact propertiesof the composites. The examination of the morphology of the materials was by (SEM), whilecharge coupled device (CCD) camera monitored crack propagation behavior under tensileloading. There was further morphology study to analyze the fiber length distribution andinterfacial shear strength of the composites. Investigations revealed that the composite withthe lowest MW PP consistently showed the highest mechanical, impact and fracture toughness
properties than those with higher MW PP. It was also observed that incorporating MA-g-PP as
compatibilizers had greater interfacial bonding effects, resulting in higher interfacial shear

264

strength in GFPP of lower MW and high melt flow rates than those of higher M W and low
melt flow rate. The investigation further revealed that as the MW of PP decreased, the fracture
toughness and other mechanical properties increased in a linear correlation. Crack
propagation trailed the interface between PP matrix and GF because of weak bonding in thecomposite with higher MW PP. This study revealed that incorporating MA-g-PP into GFPP
composites achieved better properties with PP of lower MW and higher melt flow rate. Hencethe mechanical and fracture properties of GFPP depend largely on the MW of polymer PP
matrix. There was a strong linear correlation observed between the M W of PP, fracture
toughness and other impact properties thereby making it possible to make a reliableprediction of any property of the composites from another.

7.1 Introduction
In Chapter 6, investigation was conducted into the fracture mechanism of GFPP compositesof different molecular weights incorporated with maleic anhydride grafted polypropylene(MAPP). Effort was made to characterize the fracture toughness of these materials under
LEFM, which is valid only as long as nonlinear material deformation is limited to a small
region around the crack tip. However, results in Chapter 6 showed that the composites
exhibited high plastic deformation around the crack tip. Therefore, further characterization ofthese materials was necessary under Elastic Plastic Fracture Mechanics (EPFM). ElasticPlastic Fracture Mechanics is used to characterize the fracture behaviors of materials thatexhibit time-independent, nonlinear behavior or otherwise plastic deformation. Its parametersof measurement include the crack tip opening displacement (CTOD) and the J-contour
integral, and both parameters describe crack tip conditions in such elastic-plastic materials.
Critical values of CTOD or J-contour Integral give size-independent measures of fracture
toughness, even when there is a large amount of crack tip plasticity [1].

265

7.1.1 Crack Tip Opening Displacement (CTOD)
The CTOD describes the stress ratios (stress intensity factor) required for the moving apart ofpre-cracked faces prior to fracture; and the consequent blunting of an initially sharp crack as
the applied stress increases until fracture initiation and consequent failure. In effect, theCTOD is a component of the stress-intensity factor approach of fracture toughness
determination. CTOD is calculated from the following equations (note that in this thesis yieldstress will be substituted with maximum stress):(7.1)
Where δ = CTOD,  max = maximum stress, E = Young’s modulus, K1max = maximum stressintensity factor,(7.2)
Uy is the vertical distance between the center line of the original sharp crack and the blunted
crack wall. (a + ry) = effective crack length: where a is pre-crack and ry is the displacement
behind the effective crack tip,
(7.3)
is the extension of crack beyond the initial crack tip.
CTOD can also be solved in relation with the strain-energy release rate using the following
equation:(7.4)
Where G is the strain energy release rate and σmax is the maximum stress.
CTOD can also be estimated by the use of strip yield model as given below:266

(7.5)

However, it has to be noted that the strip yield model assumes plane stress conditions and a
non-hardening material. Thus, the actual relationship between CTOD and K1 and G1 depends
on stress state and strain hardening. Therefore, a more general form of this relationship isgiven by the following:(7.6)
Where m is a constant, and it is approximately 1.0 for plane stress and 2.0 for plane strain
conditions [1].

7.1.2 J-Contour Integral
The J-Contour integral is used to characterize nonlinear elastic materials. Such materials
show nonlinear elastic-plastic behavior and they are contrasted in behavior with elasticplastic materials. It is important to note here that the uniaxial stress-strain loading behavior ofboth elastic-plastic materials and nonlinear elastic materials are identical. However, they havedifferent unloading behavior. The elastic-plastic material follows a linear unloading patternwhere the gradient is equal to Young’s modulus. On the other hand, nonlinear elastic material
unloads along the same path as it was loaded (refer to Figure 7.1) [1]. It has to be noted that
there is a unique relationship between stress, and strain in an elastic material. On the contrary,
in an elastic-plastic material, a given strain can correspond to more than one stress value if

267

the material is unloaded or cyclically loaded. In effect, it is easier to analyze an elastic
material than a material that shows a plasticity that cannot be reversed.

J-contour integral is applied in nonlinear elasticity to analyze a crack in a nonlinear material.
Therefore, the nonlinear energy release rate, J, is a path-independent line integral thatuniquely characterizes crack tip stresses and strains in nonlinear materials. Hence, J-contour
integral (equals the energy release rate in a nonlinear elastic material that contains a crack)satisfies both the strain energy and stress intensity parameters.

The J-Contour Integral can be viewed as the following parameters: (a) nonlinear energy
release rate, (b) J as a path-independent line integral, (c) J as a stress intensity factor.
However, at present (a) and (c) will be considered.

7.1.2.1 Nonlinear Energy Release Rate
Here J is a path-independent contour integral used to analyze cracks. This integral J is equal
to the energy release rate in a nonlinear elastic material that has a crack. It follows that:(7.7)
Where П is the potential energy (П = U – F, U is the strain energy stored in the material
while F is work done by external forces) and A is the crack area. And for a fixed load andfixed grip conditions(7.8)
This translates to(7.9)

However, considering geometry factors J1 and JQ are computed from the following equations:268

(7.10)

Where  is Poisson Ratio,  max is maximum stress, and  is the CTOD. E is Young’s
modulus of elasticity while K1 is stress intensity factor.

7.1.2.2 J as a Stress Intensity Parameter
J is also used to characterize the crack tip conditions in a nonlinear elastic material. It is
assumed that there is a power law relationship between plastic strain and stress. Considering,also, the elastic strains, the uniaxial deformation is modeled as follows:(7.11)

Where  and  are the measured yield stress, and strain,  0 = yield strength (reference
stress value),(reference strain value),  is a unit less constant and n is the strain

hardening exponent. The following equations therefore define the variation of stress, andstrain condition ahead of the crack tip:

(7.12)

Where k1 and k2 are proportionality constants.

This study examined the effects of Maleic anhydride grafted polypropylene (MA-g-PP) anddifferent

molecular

weightsofpolypropylene

on

the

fracture

toughnessandmechanical/impact properties of glass fiber reinforced polypropylene (GFPP) composites.
269

There was investigation also on the interfacial shear strength. Tensile and Charpy impact tests
were conducted on the single-edge-notch samples of the composites. Characterization of thefracture toughness was by maximum J-integral (J1max) determined with the maximum stressof the samples; and the determination of JQ was by the stress at 5° off tangent of thestress/strain curves of the composites. There was analysis of the correlation between fracture
toughness and other properties in order to establish a functional relationship. Determinationof line of best fit and evaluation of degree of association between fracture toughness andother properties of the composites was by regression analysis and analysis of variancetechniques. The study of the material’s morphology was by SEM; and the monitoring of
crack propagation behavior of the composites during testing was by (CCD) camera.

7.2 Characterization Methods

7.2.1 Material and Sample Preparation
Glass fiber (GF) (diameter = 13µm) was used to reinforce polypropylene (PP) of different
molecular weights and melt flow rates (MFR); Maleic anhydride grafted PP (MA-g-PP),[POLYTACK grade H-1000P of molecular weight 72,000], was incorporated into thecomposites for compatibilization between the matrix and reinforcements (refer to Table 7.1).
The material compositions are noted below Table 7.1. The individual constituent materials
were dry-mixed in a mixer prior to compounding. Compounding was by a Technovel
Corporation KZW15-60MG-KIK extruder with a screw aspect ratio of 44.8 at a cylinder
temperature of 200ºC and screw rotation speed of 220rpm. The extrudate was pelletized andfabricated into dumbbell specimens for the various tests with an injection molding machinemodel TM30H by Toyo Machinery and Metal Co., Ltd. Barrel temperature was set at 200ᴼC
while the mold temperature was maintained at 30ᴼC; holding pressure of 40MPa was applied
for 15 seconds after the injection of resin.
2707.2.2 Tensile Tests of GFPP Composites
Tensile test was conducted on both the un-notched and single-edge-notch specimens of thecomposites. At least five sample pieces were tested for each condition. Introduction of singleedge-notch was by Ueshima N0. 820261 notching machine, [from Hitachi, Ltd. Tokyo,
Japan], that was mounted with a V-shaped blade. A fresh razor was used to introduce a sharp
crack into the pre-notch. The size ‘a’ of the notch was determined so that the ratio of ‘a’ tothe sample width ‘W’ falls between 0.45 ≤ a/w ≤ 0.55 [Figure 7.2]. All tests were conducted
at a room temperature of 25°C. Tensile test was performed at a crosshead speed of 10mm/minwith an Instron Universal Tester 4466 that was attached with a 10kN load cell. The unnotched test pieces were tested in accordance with ASTM D 3039. ASTM D 5045 wasadopted for use in the tensile testing of the notched specimens.

7.2.3 CCD Camera Monitored Tensile Test and Further Characterization of the
Fracture Toughness of GFPP under EPFM
Tensile tests were already conducted on the materials in subsection 7.2.2. The crack
propagation behavior of these composites during fracture toughness investigation could
provide an account of the toughness of the bulk material as well as the fiber-matrix interfacial
strength. This investigation was conducted by observing the crack propagation of GFPP
under tensile test with a CCD Camera mounted with a high-resolution optical microscope.
Test was done at room temperature of 25ᴼC on single-edge-notch specimens of gauge lengthof 115mm at a crosshead speed of 1mm/s. Instron Universal Tester 4466 was used to conduct
the test. The microscopic images obtained from this test are presented later in this chapter.

271

7.2.4 Charpy and Izod Impact Tests
Charpy and Izod impact tests were conducted on both un-notched and notched test pieces inorder to determine the toughness of the materials. A 2 mm size notch was introduced with aType A-3 Notching Machine manufactured by Toyo Seiki manufacturing Co., Ltd. Ten
sample pieces were tested for each material composition. Both tests were conducted at room
temperature of 25ᴼC, with an Impact Tester model DG-UB by Toyo Seiki Manufacturing,Co., Ltd. The Impact Tester was mounted with standardized JIS K7110 hammer bearing theenergy of 5.5J and 2.75J for Charpy and Izod impact tests respectively.

7.2.5 Scanning Electron Microscopy (SEM)
The morphology of the post-fracture materials was examined by the use of scanning electron
microscopy (SEM). The fracture surfaces of five samples of each composite were scanned inorder to determine the measured average lengths of the ductile and brittle regions of thefracture propagation. Micrographs were taken of the fracture surfaces of the tensile samples
by the use of JEOL JSM5200 Scanning Electron Microscope equipped with Olympus Digital
Camera. The micrographs are shown later in the chapter.

7.2.6 Determination of the Fracture Toughness of GFPP Composites
The fracture toughness (J1max or JQ) of the composites was investigated under the Elasticplastic fracture mechanics (EPFM) and was calculated by using equations (7.13) [3 and 4]:(7.13)

Where272

7.2.7 Determination of Interfacial Shear Strength of GFPP Composites
There was investigation also on the strength of the interface between glass fiber and matrix
resin. High interfacial shear strength ( ) and shorter critical fiber length (lc) arecharacteristics of superior interface in a fiber-reinforced composite [5]. The following KellyTyson equation (7.14) gives the theory-based strength of a short glass fiber reinforced
composite [5, 6, and 7].

(7.14)

Where σcu is the ultimate strength of the composite; and σfu is the ultimate strength of thefiber. σm is the matrix stress at failure; and τi is the interfacial shear strength. C is the fiber
orientation efficiency; and lc is the critical fiber length. li and lj are the sub-critical and supercritical fiber lengths, respectively, while Vi and Vj are their fiber volume fractions. Vf is theoverall fiber volume fraction, and r is the fiber radius.

The critical fiber length (lc) and the interfacial shear strength (τ), respectively, were calculated
from equations (7.15) and (7.16) below [5, 6, and 8]:(7.15)(7.16)

Fractured and un-fractured samples of GFPP materials were used to determine the fiber
length distribution and interfacial shear strengths of the respective composites. The fractured273

samples were chosen from the un-notched samples, which were initially tested. A 2mm
length was cut off from one end of the fractured surfaces of each specimen while a 4mm
length encompassing the fracture point was cut from a corresponding portion of an unfractured (un-stressed) specimen as shown in Figure 7.3. The cut portions, respectively, were
heated in an oven for 6 hours at a temperature of 600ᴼC in order to burn off the matrix resin
thereby leaving a residue of glass fibers. The glass fibers were carefully spread in-between
two transparent slide glasses in order to avoid overlapping and crossing. Optical micrographsof the samples were taken with a 1.5mm magnification microscope that was mounted on ahigh-resolution digital camera. The images were converted to grayscale and thresholddetermined in order to get a clear contrast between glass fibers and the background as shownin Figure 7.4. Image analyzing software was used to enumerate the glass fibers, measure thelengths of each fiber and the average length contained in the images. The analyzed fiber
lengths and other parameters were used for the calculation of interfacial shear strength with
software which module was based on equations (7.14) – (7.16) [5 and 6].

7.2.8 Correlation of the Mechanical Properties/Impact Properties, Molecular Weightand Fracture Toughness Properties
Regression analysis is used in the prediction of statistical parameters (variables). This
technique determines the line of best fit of a set of data [9]. The goodness of line of best fit
depends on the product moment correlation coefficient (r), which measures the degree of
linear association between the dependent and independent variables, or more precisely theamount of reduction in error in predicting values of one variable from the values of the other.
The various correlation coefficients between fracture toughness and other parameters were
determined from the following equations [9]:274(7.17)Or(7.18)
Or(7.19)
Where x = the independent variable (e.g. molecular weight),
variable, y = dependent variable (fracture toughness),
number of samples,= covariance of x and y;= mean of the independent= mean of the dependent variable, n =andare the standard deviations of thevariables x and y.

7.3 Results and Discussion

7.3.1 Tensile and Fracture Toughness
Figure 7.5 depicts the tensile properties of neat polypropylene (PP) and Glass fiber (GF)
reinforced PP composites (GFPP). Investigation revealed that reinforcing the lower molecularweight PP with GF, [GFPP (J-3000GP)] and incorporating maleic anhydride grafted
polypropylene (MA-g-PP) as compatibilizers achieved higher tensile strength and stiffness[5, 10-14] as compared to GFPP (J-900GP) and GFPP (J-700GP) [Figure 7.5]. The
superiority of GFPP (J-3000GP) in fracture toughness over the other composites is depicted
in Figure 7.6 for K1max, KQmax, J1max, and JQmax. It can be seen from the values of theseparameters that GFPP (J-3000GP) recorded higher values than GFPP (J-900GP) and GFPP
(J-700GP). The values of J1max and JQ show that GFPP (J-3000GP) expended more energy inthe resistance of fracture than the rest of the materials. The CTODmax and rymax values
depicted in Figure 7.7 further indicate that GFPP (J-3000GP) show higher fracture toughness.
CTOD max showed greater opening of the crack tip in GFPP (J-3000GP) which means that275

there was more blunting of the crack tip before brittle failure; suggesting higher ductility thanin the rest of the composites. It is also evident from the value of rymax that there was a longer
extension of the crack beyond the initial crack tip before the material’s brittle failure.
Therefore, it follows that GFPP (J-3000GP) had the longest effective crack length than theothers and this is supported by the measured values of the ductile region. It can be concluded
that this material resisted fracture until a very minimum size of ligament was attained before
brittle fracture occurred.

The micrographs in Figure 7.8 show that the lower molecular weight PP composite [C],showed greater resistance to fracture for both J1max and JQ as revealed by higher plastic
deformation. The resistance resulted from a better bonding between GF and PP, which
minimizes fiber pullout and de-bonding thereby leading to an efficient transfer of stress from
PP matrix to GF. The addition of functionalized maleic anhydride grafted polypropylene has
been shown to generate improved interface [10 and 14]; therefore, it is used to enhanceinterfacial adhesion between GF and matrix PP [5]. The average measured lengths of both theductile and brittle regions of the composites, derived from the fracture surfaces of eachcomposite, shown in Figure 7.8 are depicted in Figure 7.9. It is seen that the lowest molecularweight PP composite [C] had longer ductile region than others did. This indicates the degreeof resistance to fracture by each composite. Figure 7.9 equally shows that as the size of theductile region increases, the corresponding brittle region decreases. This means that there is aslow-down in the rate of crack progression due to higher resistance caused by either toughbulk of the matrix or bridging from GF due to better interface between GF and PP [Figure7.10(c)] [11]. There is an evidence of cement-like (glue-like) bonding between GF and PP in
the lowest molecular PP composite [Figure 7.10(c)]. The medium molecular weight PP
composite [Figure 7.10(b)] showed weak interfacial bonding. In contrast, there is PP matrix276

de-bonding (peel-off) from GF in the composite of high molecular weight PP as depicted by
the smooth surface of GF and the presence of holes between it and the matrix [Figure7.10(a)].

The superior strength, modulus and fracture toughness can be attributed to better wetting of
the fiber, which resulted from lower viscosity of the lower molecular weight matrix [5].
Consequently, there was better adhesion between GF and PP. This results in the efficienttransfer of stress from matrix to fiber during crack opening [5]. Therefore, if the fibers havehigh modulus and do not slide or pull out, during crack propagation, the GF will sustain
higher stress and both GF and matrix will fracture together. Thus, the increase in tensile
properties and fracture toughness will come from the toughness of the glass fibers [15].

7.3.2 Micrographs and CCD Camera Images of GFPP Materials
The fracture surfaces of notched specimens under tensile loading are indicated in Figure 7.8.
The notch, ductile and brittle zones are indicated as x, y and z respectively while A, B and C
represent the samples of different molecular weights, GFPP(J-700GP), GFPP(J-900GP) andGFPP(J-3000GP) respectively. The ductile (fracture) zone ‘y’ is indicative of the fracture
toughness or levels of resistance to fracture by the interfacial (surface) energy [15] of thematerials as stress builds up and concentrates at the crack-tip (i.e. the boundary between
zones ‘x’ and ‘y’). The moment the crack is initiated at the critical stress-intensity factor
value, if there is no bridging or toughness, it propagates rapidly as the resistance to fracture of
the material reduces until brittle failure [16]. A measurement of the damage zone
(ductile/plastic deformation), depicted in Figure 7.9, shows that GFPP (J-3000GP) had thelongest damage zone measuring 3.99mm and shortest brittle zone of 1.3mm of the total
length of both the ductile and brittle zones. The measurements also show that GFPP (J-277

900GP) had 3.85mm of ductile and 1.38mm of brittle zones, while GFPP (J-700GP) had
3.14mm and 1.89mm in the same order. This means that GFPP (J-3000GP) exerted much
more energy and took a longer time in resisting fracture thus indicating greater ductility andtoughness in the material than the others. This is revealed in the cement-like fiber-matrix
surface shown in Figure 7.10(c).

The fracture progression was captured by using a CCD Camera. The images are shown in
Figures 7.11[a – c]. Snap shots of the images were taken prior to failure of the materials. The
crack propagation in GFPP (J-3000GP) specimens falls along the pre-crack axis as depictedin schematic C. This could mean that the entire bulk of the material was tough and the precrack was deemed as the weakest region for the propagation of the crack. In the case of GFPP
(J-900GP) and GFPP (J-700GP) specimens, crack propagation did not follow the pre-crack
axis, instead deviated towards other regions of the specimens to yield a wavy fracture surface
as shown in schematics B and A. This indicates that there was weaker fiber-matrix interface,which created many sub-cracks and caused the main crack to diverge away from its originalpath. It is clear from the results and the images that GFPP (J-3000GP) had higher fracture
toughness than GFPP (J-900GP) and GFPP (J-700GP) specimens, owing to the stronger
fiber-matrix interface. The weak bonding in GFPP (J-700GP) and GFPP (J-900GP) could becaused by long molecular chains in the high molecular weight PP that made it difficult for thematrix to penetrate the GF treatment for better interaction. On the other hand, it was easier forthe low molecular weight PP to penetrate the GF treatment and achieve better adhesion due toits short molecular chains.

278

7.3.3 Interfacial Shear Strength of GFPP Composites
Procedure for the calculation of interfacial shear strength was described in subsection 7.2.7.
Sample extraction and the analyzed fiber lengths are shown in Figures 7.3 and 7.4. Theinterfacial shear strength of the composites was calculated with equations (7.14) – (7.16). Theinterfacial shear strengths and fiber distributions of the three GFPP composites are shown in
Figures 7.12 – 7.16. The high value of interfacial shear strength attained by GFPP (J3000GP) indicates a more superior interfacial bonding between glass fiber and matrix in thecomposite than the other specimens. As a result there is a more effective and efficient stresstransfer from matrix to glass fiber resulting in higher mechanical, impact and fracture
toughness properties [5, 17]. The respective values of interfacial shear strength attained by
the three materials could be one of the factors influencing the fracture mechanism of thecomposites [refer to Figure 7.11(a-c)]. It is evident from Figure 7.12 that as the molecularweight of the matrix PP decreases, the interfacial shear strength increases.

Figure 7.13 shows a corresponding decrease in both the average and critical fiber lengths as
the molecular weight decreased. The shorter average fiber length in specimens with low
molecular weight PP matrix could have resulted from extensive fiber breakage during
processing. The low molecular weight PP in GFPP (J-3000GP) leads to low melt viscosity,
which resulted in high shear stresses due to better matrix wetting on the fiber surface. The
shorter critical fiber length of the fractured GFPP (J-3000GP) indicates that the fibers wereable to provide better reinforcement and as such sustain more loads. This leads to more fiber
breakages, which gives shorter fibers before fiber-matrix de-bonding finally occur. However,with increasing PP molecular weight in GFPP (J-700GP) and GFPP (J-900GP) specimens,
de-bonding occurred at longer critical fiber lengths, signifying weak fiber-matrix interfacialinteraction.

279

The high interfacial shear strength between glass fiber and matrix in GFPP (J-3000GP) could
explain the linear crack propagation along the pre-crack axis in the composite [see Figure7.11(c)]. Conversely, the lower interfacial shear strength in GFPP (J-900GP) and GFPP (J700GP) has caused preferential crack propagation through weaker fiber-matrix interface
thereby creating a wavy crack movement within the material [refer to Figure 7.11(b and c)].

The percentage fiber-length distributions of both fractured and un-fractured samples of thecomposites are shown in Figures 7.14 – 7.16. These figures reveal that due to fiber breakages
during fracture there was movement of fibers from the fiber lengths, which measured above
the average fiber length over to those, which were below the average fiber length.

7.3.4 Relationship between Fracture Toughness, PP Molecular Weight,
Mechanical/Impact Properties and Interfacial Shear Strength
It can be seen from Figure 7.17 that the fracture toughness of GFPP composites incorporated
with maleic anhydride-grafted PP (MAPP) decreased as the molecular weight of PP
increased. This could be the result of a high degree of chain entanglement as the polymer
chain-length increases, which could significantly affect the wetting of the matrix resin on thefiber surfaces. In Figures 7.18 and 7.19, there was an increase in the fracture toughness andtensile modulus/strength of the composites as the molecular weight of PP decreases. This
could also be due to the better wetting of GF by the matrix resin resulting in a betterinterfacial adhesion between PP and GF. This would provide a more efficient stress transfer
mechanism from the matrix resin to the glass fiber, which is known to provide higher
stiffness and strength to the composite [3]. The high interfacial shear strength attained by
GFPP (J-3000GP) as shown in Figure 7.20 corroborates the preceding explanation and shows280

that as the molecular weight of PP decreases, both fracture toughness and interfacial shear
strength of GFPP composite increases.

Figure 7.21 equally shows that as the molecular weight of PP decreases, the average andcritical glass fiber lengths decrease. This indicates that due to high shear during processing
there were more fiber breakages in the GF reinforced low molecular weight PP than in thehigh molecular weight PP. The high interfacial shear strength also could lead to more fiber
breakages during fracture in GFPP (J-3000GP) than in GFPP (J-900GP) and GFPP (J700GP). The relationship between fracture toughness and Charpy/Izod impact strengths,
previously discussed in Chapter 6, for both un-notched and notched samples are shown in
Figures 7.22 to 7.25. The figures show a similar tendency as explained hitherto, i.e. thefracture toughness and impact strengths increased as the molecular weights of PP decreased.
All the figures under discussion show a similar trend in the relationship between fracture
toughness, interfacial shear strength, average/critical glass fiber lengths and other tensile andimpact properties of the composites. The J1max appears to be more linearly correlated to theinterfacial shear strength and mechanical/impact properties while JQ does not provide adefinite correlation function. This implies that J1max can be useful to predict the mechanicalproperties of GFPP and vice versa, irrespective of testing conditions or methods.

7.3.5 Strength of Association between Fracture Toughness and other Properties of
GFPP
The fracture toughness of the composites was determined under the elastic plastic fracture
mechanics by the use of J-Integral. The parameter is applied when there is significant plastic
deformation in thermoplastic composites. The correlation coefficient (r) measures the degreeof linear association between the dependent and independent variables of a set of data. In the281

present investigation, the correlation coefficients [Equations 7.17 – 7.19] of the dependentand independent parameters showed a very strong association between fracture toughness(J1max) and other individual GFPP composites properties [Figure 7.26] which indicates a very
good line of best fit. The value of -1 or +1 of the linear correlation coefficient (r) indicates aperfect linear association between the variables. This means that as (r) tends to -1 or +1 thedegree of association between the variables, along the line of best fit, becomes stronger withminimal error. It can be seen from Figure 7.26 that all the values of (r) between fracture
toughness (J1max) and other properties of GFPP composites are approaching or equal to ±1.
This is an indication that the parameters can be predicted from each other with high (above90%) degree of confidence.

7.4 Summary and Conclusion
In this Chapter, the mechanical properties and fracture toughness of three grades of PP withdifferent molecular weights that were reinforced with GF and incorporated with Maleic
Anhydride Grafted Polypropylene (MAPP) were investigated. Investigation using elastic
plastic fracture mechanics has revealed that GFPP (J-3000GP) with the lowest molecularweight exhibited the highest fracture toughness properties as compared to GFPP (J-900GP)
and GFPP (J-700GP) specimens. It was also observed that crack propagation in GFPP (J3000GP) was more aligned along the pre-crack axis while a deviation of the crack
propagation path was observed when higher molecular weight PP was incorporated in thecomposites. This suggests that GFPP (J-3000GP) had better interfacial bonding and higher
toughness in the bulk, as can be seen from its high interfacial shear strength; hence, theweakest region in the material was the pre-crack path. It can be concluded from these resultsthat the molecular weight of PP is one of the major factors influencing crack propagation
behavior and fracture toughness of GFPP composites. Calculations using J-integral have282

yielded interesting correlations between JQ, J1max and the various mechanical properties andPP molecular weights. J1max was found to exhibit a linear correlation to the static as well as
impact properties of the composites. The high values of correlation coefficient (r) obtained byregressing fracture toughness provides the possibility of predicting the mechanical properties
through fracture toughness measurements and vice versa.

References
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Plane Stress,Actn Metallurgica, 13, 293-306 (1965)
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Wiley and Sons, New York, (1996)
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Fracture Toughness: Proceedings of the 1971

National Symposium on Fracture Mechanics Part II, ASTM Special Publication 514,
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Composites and Specialty Applications, Marcel Dekker, New York, 1 – 54, (1989)
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35. Olagoke, O., (ed.), Handbook of Thermoplastics, Marcel Dekker, New York, 1017 –
1034, (1997)

286

Table 7.1: Properties of Polypropylene Compounded with Glass Fiber
Materials Designation

MW of PP

MW/Mn of PP

MFR of PP (g/10min)

GFPP (J-700GP)

310,000

4.3

8

GFPP (J-900GP)

230,000

4.1

13

GFPP (J-3000GP)

190,000

4.7

30[Note: Each composite contains 20wt % of GF, 79.2wt % of PP and 0.8wt % of Maleic Anhydride Polypropylene]

287

Stress
Nonlinear Elastic Material

288

Elastic-Plastic Material(Adapted from Figure
3.7, Anderson, T. L., pg.
123)

StrainFigure 7.1: Schematic comparison of the stress-strain behavior of elastic-plastic and nonlinear elastic materials

Width289

Span length (115mm)
‘a’
Notch

‘B’ Thickness

Figure Figure
7.2: Dumbbell
specimentensile3.2: Dumbbell
specimen forfor tensile
testtest

Figure 1

Fractured sample

Tensile fracture

290

2mm

4mm

Un-fractured sample

Figure 7.3: Schematic of sample portions burnt off for fibers measurement in the determination of interfacial shear strength291
Figure 7.4: Grayscale image of glass fibers used in the calculation of interfacial shear strength

90

292

Tensile Strength (MPa)

70

3.50
Strength (PP)Strength (GFPP)
Modulus (PP)Modulus (GFPP)

60
50

2.50

40
1.50
30
20
1800002100002400002700003000000.50
330000

Molecular weight of PP
Figure 7.5: Correlation of tensile strength and modulus of PP/GFPP and molecular weight

Modulus (GPa)

4.50

80

J1max

12

JQmax

KQ

K1max

11

8
7
)

8

5

7

4

6
3

5
4

2
GFPP (J-700GP)GFPP (J-900GP)GFPP (J-3000GP)

Figure 7.6: J1max, JQ, K1max and KQ for GFPP of different molecular weights

1/2

6

9

K1max, KQ (MPa m

293

J1max, JQ (kJ/m2)

10

1.2

CTODmax

rymax

294

CTODmax, rymax (mm)

1
0.8
0.6
0.4
0.2
0
-0.2
GFPP (J-700GP)GFPP (J-900GP)GFPP (J-3000GP)

Figure 7.7: CTODmax and rymax for GFPP of different Molecular Weights

x

y

z

x

y

z

x

y

z

A

295

B

C

500mm
Figure 7.8: Fracture surface of A = GFPP(J-700GP), B = GFPP(J-900GP),
C = GFPP(J-3000GP) [x = notch, y = ductile region, z = brittle region]296

Ductile and Brittle regions (mm)

5

4
3

2

Ductile fracture length
Brittle fracture length10180000210000240000270000
Molecular weight of PP

300000Figure 7.9: Correlation between the lengths of ductile /brittle fracture regions and molecular weight of PP

a

297

b

c
Figure 7.10: Interfacial bonding between GF and PP of different molecular weight[a = GFPP(J-700GP), b = GFPP(J-900GP) and c = GFPP(J-3000GP)]

a

b

2mm

2mm

1mm

298

c

1mm

Notch

A

Crack movement
B

Notch

Crack movement

2mm

C

1mm

Figure 7.11: Crack Propagation of GFPP composites monitored by CCD Camera;a = GFPP (J-700GP), b = GFPP (J-900GP), c = GFPP (J-3000GP)

Notch

299

Interfacial shear strength (MPa)

30

25
20
15
10

5
0
180000210000240000270000300000

PP molecular weight
Figure 7.12: Correlation between interfacial shear strength and molecular weight of PP

800
700

300

Fiber length (mm)

600
500
400
300
200

Average fiber length of un-fractured sample
Average fiber length of fractured sample

100
0

Critical fiber length (lc)0

1

GFPP (J-700GP)

2

3

GFPP (J-900GP) GFPP (J-3000GP)

Figure 7.13: Average and critical fiber lengths of the composites

4

301

Frequency of fiber lengths (%)

20
Range containing the critical fiber length (569mm)
Range containing the average fiber lengthof the un-fractured sample (669mm)

16
12

Un-fractured sample
Fractured sample

8
4
0
0

200

400
600
800
1000 1200
Range of fiber lengths (mm)Figure 7.14: Fiber lengths distribution of GFPP (J-700GP) in percentage frequency

1400 1600+

302

Frequency of fiber lengths (%)

20
Range containing the critical fiber length (470mm)
Range containing the average fiber lengthof the un-fractured sample (625mm)

16
12

Un-fractured sample
Fractured sample

8
4
0
0

200

400

600
800
1000 1200
Range of fiber lengths (mm)Figure 7.15: Fiber lengths distribution of GFPP (J-900GP) in percentage frequency

1400

1600+

303

Frequency of fiber lengths (%)

20
Range containing the critical fiber length (384mm)
Range containing the average fiber lengthof the un-fractured sample (581mm)

16
12
8

Un-fractured sample
Fractured sample

4
0
0

200

400

600

800

1000

1200

Range of fiber lengths (mm)Figure 7.16: Fiber lengths distribution of GFPP (J-3000GP) in percentage frequency

1400

1600+

11.4
11.2

304

J1max & JQ (kJ/m2)

J1max

3

JQ

2

1

11
3

10.8
10.6

2

10.4

1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

10.2

10
1800002100002400001270000

PP Molecular Weight
Figure 7.17: Fracture toughness and molecular weight of PP

30000033000011.4

305

J1max & JQ (kJ/m2)

11.2
11.0

J1max

JQ

1

3
3

2

10.8
2

10.6
10.4
10.2
10.0
3.85

1 = GFPP (J-100GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

1
3.9

3.95

4

4.05

4.1

Modulus (GPa)Figure 7.18: Fracture toughness and tensile modulus of GFPP

4.15

4.2

4.25

11.4
J1max

306

J1max & JQ (kJ/m2)

11.2
11

JQ

1

3
3

2

10.8
10.6

2

10.4

1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

10.2

1

10
72

74

76

78

80

Tensile strength (MPa)Figure 7.19: Fracture toughness and tensile strength

82

84

11.4
11.2

307

J1max & JQ (kJ/m2)

JQ

J1max

3

11

3

2

1

10.8

2

10.6
10.4
10.2

1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

1

10
14

16

18

20

22

Interfacial shear strength (MPa)Figure 7.20: Fracture toughness and interfacial shear strength

24

26

800
700

308

Fiber lengths (mm)

600
500
400
300
200

Average fiber length (Un-fractured)
Average fiber length (Fractured)

100

Critical fiber length0150000200000Linear (Average fiber length (Unfractured))fiber length)
250000 Linear (Critical
300000350000

Molecular weight of Polypropylene
Figure 7.21: Average/critical fiber length and molecular weight of PP

11.4
J1max

309

J1max, JQ (kJ/m2)

11.2

JQ

3

11

3

2
1

10.8
2

10.6
10.4
1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

10.2
1

10
38

38.5

39

39.5

40

40.5

Izod Impact Strength - Un-notched (kJ/m2)
Figure 7.22: Fracture toughness and Izod impact strength (un-notched)

41

11.4
J1max

310

J1max / JQ (kJ/m2)

11.2

JQ
3

11

3

2
1

10.8
2

10.6
10.4
1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

10.2
1

10
7.7

7.75

7.8

7.85

7.9

7.95
2

Izod Impact Strength - Notched (kJ/m )Figure 7.23: Fracture toughness and Izod impact strength (notched)

8

11.4

J1max

311

J1max & JQ (kJ/m2)

11.2
11

JQ
3
3

2
1

10.8

2

10.6
10.4
10.2

1

10
45

1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

46
47
48
49
Charpy Impact Strength - Un-notched (kJ/m2)
Figure 7.24: Fracture toughness and Charpy impact strength (un-notched)

50

11.4

312

J1max & JQ (kJ/m2)

11.2

J1max

JQ

11

3
3

2

1
10.8

2

10.6
10.4

1 = GFPP (J-700GP)
2 = GFPP (J-900GP)
3 = GFPP (J-3000GP)

10.2
10
11.35

1
11.4

11.45

11.5

11.55

11.6

11.65

Charpy Impact Strength - Notched (kJ/m2)
Figure 7.25: Fracture toughness and Charpy impact strength (notched)

11.7

313

Correlation coefficient ( r )

1.10
0.73

J1max
JQ

0.36
MW
-0.01
TS

TM

IFSS

UNC

NC

-0.38
-0.75

MW = Molecular weight, TS = Tensile strength, TM =Tensile modulus, IFSS = Interfacial shear strength, UNC &NC = Charpy impact strength (un-notched & notched).

-1.12
Materials properties
Figure 7.26: Degree of association between molecular weight,
mechanical/impact/interfacial properties and fracture toughness

Chapter 8General conclusion
Developing an understanding of the fracture behavior of polymeric materials and composites that
could help to prevent catastrophic failures in engineering components was the focus of this thesis. In
Chapters 3 and 4, the effects of different depths of v-notches and hairline cracks on the fracture modesof polyethylene terephthalate injection-moldings were investigated with emphasis on the influence of
hierarchical skin-core structure. Birefringence observations as well as various crystallinity
measurements showed that V-PET had a more distinct skin-core structure with well-definedboundaries than R-PET. On the other hand, R-PET had a more dispersed and diffused skin-core
structure. The fracture behavior of v-notched V-PET and R-PET consists of three distinct fracturemodes, namely ductile, semi ductile and brittle, which corresponds to the skin-core structure andcrystallinity distribution in the materials. Results showed that there would be full ligament yielding in
PET materials when v-notches remained in the skin region. When the v-notch penetrated the skinlayer into the core region, brittle fracture was imminent. The notch sensitivity factor for yield strength(kS) of the materials would increase drastically if the v-notch were deeper than 1.0 mm whereas thenotch sensitivity factor for energy (kT) would significantly increase when notches deeper than 0.5 mm
were present. These factors indicate the extent of change in mechanical performance in the materials
as the notches became deeper. By loading the specimens to intermediate points under the yieldstrength, it has been shown that crack opening did not occur in the specimens before yielding,indicating that crack opening was a post necking phenomenon in the specimens that failed in a semi
ductile manner. Although tensile loading provided a transitional fracture behavior from ductile, semi
ductile to brittle failure of the materials, all specimens that were subjected to impact testing resulted inbrittle failure. However, a gradual loss in impact strength could be observed especially in R-PET,
which was attributed to the less crystallized skin region and the more diffused skin-core structure.
However, unlike v-notches, sharper hairline cracks have a more drastic effect on PET injection
moldings even at low depths. The introduction of a mere 10m-deep hairline crack within the skinregion significantly affected the kS and kT values. Results also indicate that V-PET was more sensitive
to hairline cracks as compared to R-PET. In order to investigate further the effects of morphology
development on fracture behavior, annealing was conducted on the PET specimens with hot oil. The
annealing created a pseudo-skin layer and reduced the distinction between the skin and core regions.
Therefore, the moldings became more ductile especially V-PET since the skin-core structure appeared
to be more diffused after annealing.

314

Investigations also showed that full ligament yielding occurred in the PET materials when the notches
were situated within the skin region, which allowed the application of essential work of fracture(EWF) principles to characterize the fracture toughness of V-PET and R-PET. The skin-core structure
as well as notch depth greatly influenced the fracture modes and crack-tip temperature during tensileloading of the specimens. A phenomenal increase in crack-tip temperature occurred when notches
were deep enough to cause stress concentration but remained within the skin region. However, when
the notch penetrated the skin into the core, brittle fracture occurred with no noticeable change incrack-tip temperature, which indicates that no ligament yielding has occurred. The zero change incrack-tip temperature and brittle fracture marked the limit of notch depth where the EWF principles
would no longer be valid. The ductile to brittle transition was found to be more drastic in V-PET as
compared to R-PET. A more subtle ductile to brittle transition could be seen in R-PET due to its morediffused skin-core structure.
The fracture toughness of GFPP composites was initially characterized by using linear elastic fracture
mechanics (LEFM) to obtain maximum stress intensity factor (K1max) and maximum strain energy
release rate (G1max). The composites, however, exhibited significant plastic deformation, which
negated the assumptions of LEFM principles. Therefore, further characterization of the fracture
toughness of the composites was conducted under elastic plastic fracture mechanics (EPFM) by using
J-integral (J1max) and crack tip opening displacement (CTODmax) methods. The interfacial properties of
the composites were also evaluated. Investigations using EPFM have confirmed that GFPP (J3000GP) with the lowest molecular weight matrix exhibited the highest fracture toughness properties
as compared to GFPP (J-900GP) and GFPP (J-700GP) specimens. It was also observed that crack
propagation in GFPP (J-3000GP) was more aligned along the pre-crack axis while a deviation of thecrack propagation path was observed when higher molecular weight PP was incorporated in thecomposites. This suggests that GFPP (J-3000GP) had better interfacial bonding and higher toughness
in the bulk; hence, the weakest region in the material was the pre-crack path. It can be concluded fromthese results that the molecular weight of PP is one of the major factors influencing crack propagation
behavior and fracture toughness of GFPP composites. Calculations using J-integral have yieldedinteresting correlations between JQ, J1max and the various mechanical properties and PP molecularweights. J1max was found to exhibit a more linear correlation to the static as well as impact properties
and molecular weight of GFPP composites, thereby providing the possibility of predicting any
properties of the composites from another. From the observations in the various investigations, it canbe concluded that the fracture toughness of glass fiber reinforced polypropylene composites isdependent on the molecular weights of the matrix polymer. The linear correlation between fracture
toughness and molecular weight provided a mathematical means to predict the mechanicalperformance of the composites based on material composition.
315