Bio-nanocomposites for enhancing water vapor barrier of

Bio-nanocomposites for enhancing water vapor barrier of cellulose-

based packaging materials

by

Madjid Farmahini-Farahani

M.Sc., Polymer Engineering, University of Tehran, Iran, 2006

A Dissertation Submitted in Partial Fulfillment of the Requirements for the Degree of

Doctor of Philosophy

in the Graduate Academic Unit of Chemical Engineering

Supervisor: Huining Xiao, PhD, P.Eng., Chemical Engineering

Examining Board: Yonghao Ni, PhD, P.Eng., Chemical Engineering

Felipe Chibante, PhD, Chemical Engineering

Huining Xiao, PhD, P.Eng., Chemical Engineering

Ying Hei Chui, PhD, P.Eng., FIWSci., Forestry & Environmental

Management

External Examiner: Dr. Baiyu (Helen) Zhang, Ph.D., Faculty of Engineering & Applied

Science, Memorial University of Newfoundland

This dissertation is accepted by the Dean of Graduate Studies

THE UNIVERSITY OF NEW BRUNSWICK

January, 2016

©Madjid Farmahini-Farahani, 2016

ii

ABSTRACT

The use of renewable materials within the packaging industry has gained increasing

interest in recent years. Cellulose fiber is the most abundant biopolymer on planet earth

due to it is sustainable, biodegradable and environmental friendly nature, it is widely

utilized in many fields. However, the barrier properties of porous and hydrophilic

cellulose fiber network products are inadequate for most barrier applications.

In order to improve its barrier properties and widen its applications, extensive studies

have been carried out in the current thesis work.

In the first section, two biopolymers, poly lactic acid (PLA), poly (3-hydroxybutyrate-co-

4-hydroxybutyrate), (PHBV) and their nanocomposites with different nanoclays and with

various clay contents were coated on paper.

Moreover, various coating methods were also used to control the hydrophilic surface and

cover the porous structure of paper in an attempt to improve the water vapor resistance of

paper products.

It was found that the coating method and clay exfoliation were the most important factors

affecting water vapor permeability. The papers coated with exfoliated PHBV

nanocomposites drastically improved the water vapor barrier of the paper by lowering

the water vapor transmission rate (WVTR) to a level that is similar to those of polypropylene

(PP) and low density polyethylene (LDPE).

In the second section, a simple route was used to produce regenerated cellulose and

regenerated cellulose nanocomposites with sodium-montmorillonite (Na-MMT). A series

of novel modified montmorillonites were prepared via solution blending in aqueous

alkaline/urea solvent at low temperatures, followed by regeneration of films. The effect

iii

of nanoclays loading on the mechanical, crystallinity and water vapor transmission

properties of the nanocomposites was investigated. Generally, nanoclays loading

improved the mechanical and water vapor barrier properties of the cellulose films.

Although, the WVTR values of the resulted nanoclay-cellulose films was much lower than

regular paper, but the WVTR values remain much higher than most of the fossil-based

polymers.

iv

DEDICATION

I dedicate this dissertation to my beloved wife and parents. Thank you for all the love,

support and sacrifice throughout the years. And to whom I owe a lot.

v

ACKNOWLEDGEMENTS

I owe my deepest gratitude to my supervisor, Professor Huning Xiao, for his inspiration

and encouragement. This thesis would not have been possible without his full support and

trust throughout my PhD studies.

I would like to also thank my beloved, gorgeous, patient, and my only true companion in

short life journey, Azam Gholami.

I also wish to extend my sincere thanks to my beloved family, who have given me their

full heart whelming support during my entire life; my mother, Mari Fatemi, my father,

Ali Farmahini-Farahani, and my sisters, Masi, Mahnaz, my twins sisters Mahtab and

Mahnoosh who persuaded and supported me.

I cannot find words to express my gratitude to all my dear friends who are always

supporting me. Especially “Dr. Zainab Ziaee, Dr. Pedram Fatehi, Dr. Alemayehu H.

Bedane, Dr. Zhaoping Song, Dr. Hao Wang, Khosro Seddighi, Babak Shirani, Dr. Peng

Lu, Dr. Felipe Chibante, Dr. Yonghao Ni, Dr. Mladen Eic, Carl E. Murdock, Adon

Briggs, and Sylvia Demerson. I feel obliged to say thanks to so many people who have

been doing so many good things for me.

Thanks are also given to the colleagues and staffs in our research group, Department of

Chemical Engineering and Limerick Pulp & Paper Center, UNB, for their infinite help.

At last, but not the least, I am really thankful to the financial support from Green Fiber

Network, NSERC Canada and UNB President's Doctoral Tuition Awards.

vi

Table of Contents

ABSTRACT ........................................................................................................................ ii

DEDICATION ................................................................................................................... iv

ACKNOWLEDGEMENTS ................................................................................................ v

Table of Contents ............................................................................................................... vi

List of Tables .................................................................................................................... xv

List of Schemes ................................................................................................................ xvi

List of Figures ................................................................................................................. xvii

List of Symbols, Nomenclature or Abbreviations ......................................................... xxiii

Chapter 1 Introduction ........................................................................................................ 1 1.1. Background .............................................................................................................. 1

1.2. Thesis objectives ...................................................................................................... 6

1.3. Thesis layout ............................................................................................................ 7

References ..................................................................................................................... 10

Chapter 2 Literature Review ............................................................................................. 16 2.1. Introduction ............................................................................................................ 16

2.2. Water vapor transmission rates and permeability .................................................. 16

2.3. Factors affecting permeability of polymeric films ................................................. 20

2.3.1. Chemical structure .......................................................................................... 20

2.3.2. Crystallinity..................................................................................................... 22

2.3.3. Glass transition temperature ........................................................................... 23

2.3.4. Porosity and Voids .......................................................................................... 23

2.4. Barrier properties improvement of a polymer ....................................................... 24

vii

2.4.1. Barrier Improvement from Polymer Blends ................................................... 24

2.4.2 Barrier Improvement from Nanocomposites ................................................... 25

2.5. Biodegradable materials......................................................................................... 28

2.5.1. Poly lactic acid: ............................................................................................... 28

2.5.2. Poly(3-hydroxybutyrate) and (Poly(3-hydroxybutyrate-co-3-hydroxyvalerate)

................................................................................................................................... 31

2.5.3. Cellulose and Regenerated Cellulose .............................................................. 37

References ..................................................................................................................... 42

Chapter 3 Poly Lactic Acid Nanocomposites Containing Modified Nanoclay

with Synergistic Barrier to Water Vapor for Coated Paper ............................... 64 Abstract ......................................................................................................................... 64

3.1. Introduction ............................................................................................................ 65

3.2. Experimental .......................................................................................................... 67

3.2.1. Materials ..................................................................................................... 67

3.3. Characterization ..................................................................................................... 68

3.4. Preparation of the Surface Modified Clays ............................................................ 70

3.5. Preparation of PLA/MMT Nanocomposites .......................................................... 71

3.6. Paper Coating ......................................................................................................... 72

3.7. Results and discussion ........................................................................................... 72

3.7.1. Morphologies of the PLA Nanocomposites ........................................... 74

3.7.2. Thermal decomposition of the PLA/Caly NanoComposites ...................... 77

viii

3.7.3. The WVTR Measurement ........................................................................... 79

Conclusions ................................................................................................................... 81

Acknowledgment .......................................................................................................... 81

References ..................................................................................................................... 81

Chapter 4 Preparation and Characterization of Exfoliated PHBV Nanocomposites to

Enhance Water Vapor Barriers of Calendared Paper ....................................................... 87 Abstract ......................................................................................................................... 87

Keywords: ..................................................................................................................... 88

4.1. Introduction ............................................................................................................ 88

4.2. Experimental .......................................................................................................... 91

4.2.1. Materials ..................................................................................................... 91

4.3. Methods.................................................................................................................. 92

4.3.1 Preparation of the Surface Modified Clays (Scheme 4-1): .............................. 92

4.3.2 Paper Coating .............................................................................................. 94

4.4. Characterization ..................................................................................................... 95

4.5. Results and Discussion .......................................................................................... 97

4.5.1. Analysis of chemical structure of surface modified clay ............................ 97

4.5.2. XRD pattern of PHBV Nanocomposites .................................................... 99

4.5.3. Morphologies of PHBV nanocomposites revealed by TEM .................... 103

4.5.4. The water vapor transmission rate (WVTR) measurement ...................... 105

ix

4.6. Conclusion ........................................................................................................... 111

Acknowledgement ...................................................................................................... 112

References ................................................................................................................... 112

Chapter 5 Surface morphological analysis and synergetic water vapor barrier

properties of modified clay/Poly(3- hydroxybutyrate-co-3-hydroxyvalerate)

composites ..................................................................................................................... 119 Abstract ....................................................................................................................... 119

5.1. Introduction: ......................................................................................................... 119

5.2. Materials .............................................................................................................. 122

5.3. In situ polymerization of β-butyrolactone on Cloisite®30B ............................... 123

5.4. Bio-nanocomposites films preparation ................................................................ 123

5.5. Characterization ................................................................................................... 124

5.6. Results and discussion ......................................................................................... 126

5.6.1. Characterization of the ring opening polymerization on CS30B .................. 126

5.6.2. The surface characterization of the nanocomposites ................................ 127

5.6.3. Contact Angle and surface roughness correlation in nanocomposites. ..... 130

5.6.4. Morphologies of PHBV nanocomposites revealed by XRD and TEM .... 133

5.6.5. Moisture adsorption .................................................................................. 136

5.6.6. The water vapor transmission rate (WVTR) measurement ...................... 139

5.7. Conclusion ........................................................................................................ 141


x

References ................................................................................................................... 142

Chapter 6 Thermal analyses and rheological behavior of exfoliated

poly(hydroxybutyrate-co-hydroxyvalerate) bionanocomposite: Influence of pre-

intercalated nanoclay incorporation ................................................................................ 150 Abstract ....................................................................................................................... 150

6.1. Introduction .......................................................................................................... 151

6.2. Materials .............................................................................................................. 153

6.3. Bio-nanoomposite film preparation ..................................................................... 154

6.4. Characterization of the bio-nanocomposites ........................................................ 154


6.5.2. Morphology of Nanocomposites................................................................... 157

6.5.3. Thermal properties and Crystallization behavior of nanocomposites........... 158

6.5.4. Thermogravimetric analyses ......................................................................... 163

6.5.6. Viscoelastic behavior of PHBV and its nanocomposites .............................. 166

6.6. Conclusion ........................................................................................................ 170



References ................................................................................................................... 172

Chapter 7 Cellulose/nano-clay composite films with high water vapor resistance and

mechanical strength ........................................................................................................ 179 Abstract ....................................................................................................................... 179

7.1. Introduction .......................................................................................................... 180

xi

7.2. Experimental section ............................................................................................ 183

7.2.1. Materials ................................................................................................... 183

7.2.2. Preparation of cellulose nanocomposite films .......................................... 183



7.4.1. FTIR spectroscopy .................................................................................... 190

7.4.2. X-ray diffraction ....................................................................................... 192

7.4.3. Crystallinity of regenerated cellulose and nanocomposite based films .... 197

7.4.4. Water vapor transmission and oil resistance ............................................. 198

7.4.5. Mechanical Properties ............................................................................... 202

7.5. Conclusion ........................................................................................................ 206

Acknowledgment ........................................................................................................ 206

References ................................................................................................................... 207

Chapter 8 Preparation and Characterization of surface modified Exfoliated Na-

Montmorillonite nanoclay with quaternized cellulose .................................................... 213

Abstract ....................................................................................................................... 213

8.1. Introduction .......................................................................................................... 213


8.2.1. Materials & Methods ................................................................................ 216

8.2.2. Synthesis of quaternized celluloses in NaOH/Urea Aqueous Solutions.

(Scheme 8-1) ........................................................................................................... 216

xii

8.2.3. Synthesis of quaternized celluloses Intercalated Na-MMT Adsorbent

(Scheme 8-2) ........................................................................................................... 218



8.4.1. Chemical compositions and morphology.................................................. 222

8.4.2. Water vapor adsorption of quartenized cellulose and modified clays powder

226

8.4.3. Bulk morphology of polycation-clay ........................................................ 228

8.4.4. Surface topology of polycation-clay thin film .......................................... 231

8.4.5. Zeta potential analyses at various pH values ............................................ 234

8.5. Conclusion ........................................................................................................... 237

Acknowledgment ........................................................................................................ 237

References ................................................................................................................... 238

Chapter 9 WVTR, morphology and mechanical properties of regenerated cotton linter

nanocomposites reinforced with pre-exfoliated Polycation-montmorillonite ................ 244 Abstract: ...................................................................................................................... 244

9.1. Introduction: ......................................................................................................... 245


9.2.1. Materials ....................................................................................................... 248

9.2.1. Quaternization of α-cellulose and microcrystalline cellulose ....................... 249

xiii

9.2.2. Synthesis of quaternized celluloses polycation modified montmorillonites

(PMMT) .................................................................................................................. 249

9.2.3. Regenerated Cellulose based nanocomposites film preparation ................... 250



9.4.1. Characterization of quarterinzed cellulose and quarterinzed celluloses

modified Na-MMT.................................................................................................. 254

9.4.2. XRD analysis of polycation modified Na-MMT (Filler Structure in

Nanocomposites.).................................................................................................... 256

9.4.3. TEM morphology analysis of nanocomposites............................................. 257

9.4.4. Crystallinity of Cellulose II based films and its nanocpmposites ................. 259

9.4.5. Mechanical properties ................................................................................... 260

9.4.5. Water vapor mass transfer properties ........................................................... 263

9.5. Conclusion ........................................................................................................... 267

Acknowledgment ........................................................................................................ 268

References ................................................................................................................... 268

Chapter 10 Conclusions, and Future Perspectives .......................................................... 275

10.1. Summary of the preparation and characterization of biopolymers nanocomposites

coated paper to enhance water vapor resistance cellulose-based material in packaging

application ................................................................................................................... 275

xiv

10.1.1. Summary of the preparation and characterization of PLA nanocomposites

coated on paper ....................................................................................................... 276

10.1.2. Summary of preparation and characterization of PHBV nanocomposites and

water vapor barrier properties of nanocomposite-coated paper .............................. 277

10.2. Summary of the preparation and characterization of regenerated cellulose-based

film nanocomposites. .................................................................................................. 280

10.3. Recommendations for future research ............................................................... 282

Appendices .................................................................................................................. 284

Appendix-I: Effect of super-hydrophobicity on WVTR of hand-sheet .................. 284

Appendix-II: Summary of WVTR of various material have been measured using

IGA ......................................................................................................................... 286

Appendix-III: Primary study on effect of pore size on WVTR .............................. 286

Curriculum Vitae

xv

List of Tables

Table ‎3.1. WVTR results of W&P paper Bar-Coated and Compress Hot Melt Coated with

PLA, PLA/CS30B, and PLA/M-CS30B Nano-Composites ..................................... 80

Table ‎4.1. Modified clay with various molar ratio of catalyst to β-BL in Toluene and

bulk polymerization .................................................................................................. 98

Table ‎4.2. WVTR of samples in details .......................................................................... 111

Table ‎5.1. Surface roughness and water contact angle data of various PHBV and its

nanocomposites ....................................................................................................... 133

Table ‎6.1. Thermal properties of PHBV nanocomposites .............................................. 161

Table ‎6.2.Thermal stabilities of the molten PHBV, PHBV/CS and PHBV/PGCS nano

composites under air ............................................................................................... 164

Table ‎7.1. Experimental results of the regenerated cellulose from XRD paterns ........... 198

Table ‎7.2. Water vapor transmission rate of regenerated cellulose film at various

thickness, temperature and RH% ............................................................................ 205

Table ‎8.1. Surface charge density of quaternized celluloses .......................................... 217

Table ‎8.2. The recipes, remained ash of modified clays at 600 °C and coding of the poly-

cation Na-MMT samples ........................................................................................ 219

Table ‎8.3. Summarized results from AFM test ............................................................... 233

Table ‎9.1. Experimental results of the cellulose II, cellulose II nanocomposites from

XRD patterns .......................................................................................................... 260

xvi

List of Schemes

Scheme ‎3-1. Schematic illustration of IGA setup for WVTR testing ............................... 70

Scheme ‎3-2. Preparation of the Surface Modified Clays .................................................. 71

Scheme ‎4-1. Preparation of the Surface Modified Clays .................................................. 93

Scheme ‎5-1. Structures of the 1-Ethoxy-3-chlorotetrabutyldistannoxane ...................... 122

Scheme ‎5-2. Structure of PHB-g-CS30B........................................................................ 123

Scheme ‎6-1. Structure of CS30B and PHB-g-CS30B .................................................... 153

Scheme ‎7-1. IGA experimental set-up for WVTR ......................................................... 188

Scheme ‎7-2. Oil barrier properties set-up ....................................................................... 189

Scheme ‎8-1.Quaternization of Cellulose (QC) ............................................................... 218

Scheme ‎8-2. Poly-cation modified MMT (PMMT) with QC ......................................... 219

xvii

List of Figures

Figure ‎2-1. Micro-composite, intercalated, and exfoliated structure ................................ 26

Figure ‎2-2. 3D pattern of layered silicate of exfoliated nanocomposites ......................... 27

Figure ‎2-3.Chemical structure of PLA.............................................................................. 29

Figure ‎2-4. Chemical structure of PHB ............................................................................ 31

Figure ‎2-5. Chemical structure of PHBV ......................................................................... 32

Figure ‎2-6.‎ Chemical Structure of Cellulose .................................................................... 38

Figure ‎3-1. FTIR spectra of (a) CS30B and (b) M-CS30B............................................... 74

Figure ‎3-2. X-ray diffraction patterns of MMT and PLA nanocomposites. ..................... 76

Figure ‎3-3. TEM images of PLA nanocomposites: (a) PLA/M-CS30B 95/05, (b) PLA/M-

CS30B 90/10 and, (c) PLA/CS30B 90/10 at different magnification. ..................... 77

Figure ‎3-4. Weight loss for pure PLA and PLA nanocomposites .................................... 78

Figure ‎3-5. The WVTR of PLA/M-CS30B-coated paper with bar and compress hot melt

coating methods ........................................................................................................ 80

Figure ‎4-1. FTIR spectra of PHBV, CS30B, and HPGC .................................................. 98

Figure ‎4-2. X-ray diffraction patterns of Clays .............................................................. 101

Figure ‎4-3. X-ray diffraction patterns of a) PHBV b) PCS1 c) PCS3 d) PCS5 and e)

PCS10 ..................................................................................................................... 102

Figure ‎4-4. X-ray diffraction patterns of a) HPGC1, b) HPGC3, c) HPGC5, d) HPGC10,

e) MPGC1 f) MPGC5, and g) MPGC10 ................................................................. 103

Figure ‎4-5 TEM images of PHBV nanocomposites a) PCS5, b) PCS10, c) MPGC5, d)

MPGC10, e) HPGC5, and f) HPGC10 ................................................................... 105

Figure ‎4-6. The WVTR of PHBV coated paper with LC and SPLC coating methods... 108

xviii

Figure ‎4-7. SEM cross section of coated paper with a) LC coating and b) SPLC coating

method..................................................................................................................... 109

Figure ‎4-8. The WVTR values of PHBV nanocomposites coated paper with SPLC

coating methods a) at 23 °C and 50 RH% and b) 38 °C and 90 RH% ................... 110

Figure ‎5-1. FTIR spectra of PHB, PHB-g-CS30B, and CS30B ..................................... 127

Figure ‎5-2. AFM phase images of a) PHBV, b) PHBV/CS30B 99/1 w/wt%, c)

PHBV/CS30B 97/3 w/wt% ,d) PHBV/CS30B 95/5 w/wt% ,e) PHBV/CS30B 90/10

w/wt% ,f) PHBV/CS30B 80/20 w/wt% ,g) PHBV/CS30B 70/30 w/wt%,

h)PHBV/CS30B 60/40 w/wt%, i) PHBV/PHG-g-CS30B 99/1 w/wt%, j) PHBV/

PHG-g-CS30B 97/3 w/wt% ,k) PHBV/PHG-g-CS30B 95/5 w/wt% ,l) PHBV/PHG-

g-CS30B 90/10 w/wt% ,m) PHBV/PHG-g-CS30B 80/20 w/wt% ,n) PHBV/PHG-g-

CS30B 70/30 w/wt%, and o)PHBV/PHG-g-CS30B 60/40 w/wt%. ....................... 129

Figure ‎5-3. AFM 3D microroughness and water contact angel of a) PHBV, b)

PHBV/CS30B 99/1 w/wt%, c) PHBV/CS30B 97/3 w/wt% ,d) PHBV/CS30B 95/5

w/wt% ,e) PHBV/CS30B 90/10 w/wt% ,f) PHBV/CS30B 80/20 w/wt% ,g)

PHBV/CS30B 70/30 w/wt%, h)PHBV/CS30B 60/40 w/wt%, i) PHBV/PHG-g-

CS30B 99/1 w/wt%, j) PHBV/ PHG-g-CS30B 97/3 w/wt% ,k) PHBV/PHG-g-

CS30B 95/5 w/wt% ,l) PHBV/PHG-g-CS30B 90/10 w/wt% ,m) PHBV/PHG-g-

CS30B 80/20 w/wt% ,n) PHBV/PHG-g-CS30B 70/30 w/wt%, and o)PHBV/PHG-g-

CS30B 60/40 w/wt%............................................................................................... 132

Figure ‎5-4. X-ray diffraction patterns of a) PHBV, b) PHBV/PHG-g-CS30B 99/1 w/wt%,

c) PHBV/PHG-g-CS30B 97/03 w/wt% ,d) PHBV/PHG-g-CS30B 80/20 w/wt% ,e)

xix

PHBV/PHG-g-CS30B w/wt% ,f) PHBV/CS30B 97/03 w/wt% , and g)

PHBV/CS30B 90/10 w/wt% ................................................................................... 135

Figure ‎5-5. TEM images of PHBV nanocomposites a) PHBV/PHG-g-CS30B 90/10and b)

PHBV/CS30B 90/100 w/wt% ................................................................................. 136

Figure ‎5-6. Water vapor sorption isotherms of clays ...................................................... 138

Figure ‎5-7. Water vapor sorption isotherms of a) PHBV/CS30B b) PHBV/PHB-g-CS30B

................................................................................................................................. 138

Figure ‎5-8. The WVTR values of PHBV nanocomposites film a) at 23 °C and 50 RH%

and b) at 38 °C and 90 RH% ................................................................................... 140

Figure ‎6-1. FTIR spectra of PHBV, PHBV/CS 90/10 (wt/wt%), and PHBV/PGCS 90/10

(wt/wt%) ................................................................................................................. 157

Figure ‎6-2. TEM images of PHBV nanocomposites a)PHBV/PCS 95/05 wt/wt% , b)

PHBV/PCS 80/20 wt/wt%, c) PHBV/PGCS 95/05 wt/wt%, and d) PHBV/PGCS

80/20 wt/wt% .......................................................................................................... 158

Figure ‎6-3. DSC PHBV/CS nanocomposites ................................................................. 162

Figure ‎6-4. DSC of PHBV/PGCSs nanocomposites ...................................................... 162

Figure ‎6-5. a) TGA and b) DTGA thermograms of pristine PHBV, PHBV/CS and

PHBV/PGCS nanocomposites ................................................................................ 165

Figure ‎6-6. Storage modulus G’(ω) and loss modulus G”(ω) for PHBV, PHBV/CSs, and

PHBV/PGCSs nanocomposites .............................................................................. 169

Figure ‎6-7. Complex viscosity η*(ω) of PHBV, PHBV/CSs, and PHBV/PGCSs

nanocomposites ....................................................................................................... 170

xx

Figure ‎7-1. photographs of a) RCL10 and b) RMMC10 transparent film for wrapping and

packaging application ............................................................................................. 185

Figure ‎7-2. FTIR spectra of RMCC, RCL, and the nanocomposites with 1, 3, 5 , and 10

Na-MMT clay content in transmittance in the region of 1500-1800 cm−1 ............ 191


Na-MMT clay content in transmittance in the region of 800-1400 cm−1 .............. 192

Figure ‎7-4. XRD diffractograms of CL, MCC, RCL and RMCC .................................. 194

Figure ‎7-5. XRD patterns of a) RCL, b) RCL1, c) RCl3, d) RCL5, and e) RCL10 films

................................................................................................................................. 195

Figure ‎7-6. XRD patterns of a) RMCC, b) RMCC3, c) RMCC1, d) RMCC5, and e)

RMCC10 films ........................................................................................................ 196

Figure ‎7-7. WVTR of RCL and RMCC at various thicknesses and temperature at 50

RH% ........................................................................................................................ 199

Figure ‎7-8. Water vapor permeability of RCL, RMCC and regenerated cellulose-Na-

MMT nanocomposite films with various Na-MMT contents measured at different

thickness at 50 RH% and 23 °C .............................................................................. 201

Figure ‎7-9. Tensile strength, modulus and elongation at break of RCL and RMCC/Na-

MMT nanocomposite films at various clay content ............................................... 204

Figure ‎8-1. FT-IR spectra of modified nanoclays from 450 to 1950 cm-1

...................... 223

Figure ‎8-2. XRD of modified nanoclays at various pH, a) pH 6 and b) pH 10, ............. 225

Figure ‎8-3. The isothermal moisture adsorption of quaternized celluloses and modified

nanoclay .................................................................................................................. 227

xxi

Figure ‎8-4. TEM micrograph of diluted C2NC1 and C2NM1 at pH 6 and 8 a) C2NC1 at

pH8, b) C2NC1 at pH6 c) C2NC1 at pH6 (lower magnification), d) C2NM1 at pH8,

e) C2NM1 at pH 6, and f) C2NM1 at pH 6 (lower magnification) ........................ 229

Figure ‎8-5. Higher magnification TEM micrograph of diluted C2NC1 and C2NM1 at pH

6 and 8 a) C2NC1 at pH8, b) C2NC1 at pH6 c) C2NM1 at pH8, c) C2NM1 at pH 6

................................................................................................................................. 230

Figure ‎8-6. TEM of CNC3 tiny film at various pH; a) pH 6, b) pH 8, and c) pH 10 ..... 231

Figure ‎8-7. AFM 3D microroughness of a) C2NC1, b) CNC1 , c) CNC2,d) CNC3 ,e)

C2NM1 ,f) CNM1,g) CNM2, and h)CNM3 tiny film ............................................ 232

Figure ‎8-8. AFM phase images of a) C2NM1, b) CNM2, c) CNM3,d) C2NC1, e) CNC2

and, f) CNC3 ........................................................................................................... 234

Figure ‎8-9. ξ potential analyses at various pH for a) Quaternized Cellulose samples b)

Modified-MMT, and c) Na-MMT .......................................................................... 236

Figure ‎9-1. FT-IR spectra of QC from 1000 to 1750 cm-1 wavenumber ....................... 254

Figure ‎9-2. FT-IR spectra of modified clays from .......................................................... 255

Figure ‎9-3. XRD of Cellulose II, Na-MMT, and Cellulose II nanocomposites ............. 257

Figure ‎9-4. TEM micrograph of a) Cellulose II/CNC3, b)Cellulose II/CNC2, c) Cellulose

II/CNM3, and d) Cellulose II/CNM1 nanocomposites ........................................... 258

Figure ‎9-5. Tensile strength, modulus and elongation at break of Cellulose II, A)

Cellulose II/CNC1, B) Cellulose II/CNC2, C) Cellulose II/CNC3, D) Cellulose

II/CNM1, Cellulose II.CNM2, and Cellulose II/CNM3 at various thickness ........ 263

xxii

Figure ‎9-6. WVTR (at RH 50% and 23 °C) of Cellulose II, A) Cellulose II/CNC1, B)

Cellulose II/CNC2, C) Cellulose II/CNC3, D) Cellulose II/CNM1, Cellulose

II.CNM2, and Cellulose II/CNM3 at various thickness ......................................... 265

Figure ‎9-7.Transparency of Cellulose II and its nanocomposites film at visible

wavelength light with a thickness of 20 µm ........................................................... 266

xxiii

List of Symbols, Nomenclature or Abbreviations

A Effective area

AFM Atomic Force Microscope

ASTM American Society for Testing and Materials

C2NC1 Modified Na-MMT with NC1 (CEC ratio1/2)

C2NM1 Modified Na-MMT with NM1 (CEC ratio1/2)

CEC Cation exchange capacity

Cellulose I Native cellulose

Cellulose II Regenerated cellulose

CHPTAC 3-Chloro-2-hydroxypropyltrimethylammonium chloride

CL Cotton linter

CNC1 Modified Na-MMT with NC1 (CEC ratio1/1)



CNM1 Modified Na-MMT with NM1 (CEC ratio1/1)



CS30B Montmorillonite Cloisite®30B

Dm Pore diameter

DSC Differential scanning calorimetry

DTGA Differential thermogravimetric

FT-IR Fourier transform infrared spectroscopy

xxiv

G’ Storage modulus

G’’ Loss modulus

GPa Gigapascal

GPC Gel permeation chromatography

HCl Hydrochloric acid

HPGC1 PHBV/PHB-g-CS30B (99/01 w/w%)

HPGC3 PHBV/ PHB-g-CS30B (97/03 w/w%)






ISO International Organization for Standardization

KBr Potassium bromide

LDPE low density polyethylene

LiOH Lithium hydroxide

MCC Microcrystalline cellulose

MFC Microfibrillated cellulose

Mn Number average molecular weight

MPa Megapascal

Mv Viscosity average molar weight

Mw Weight average molecular weight

xxv

MMT Montmorillonite

Na-MMT Sodium-montmorillonite

NaOH Sodium hydroxide

NC1 Quaternized -α-cellulose with molar ratio of CHPTAC to

anhydroglucose units (12/1)





NCC Nano-crystalline cellulose

NCF Nano-cellulose fibers

NM1 Quaternized -MMC with molar ratio of CHPTAC to






PCS1 PHBV/CS30B (99/01 w/w%)





xxvi



PCL Poly(ε-caprolactone)

PE Polyethylene

PHAs Polyhydroxyalkanoates

PHB Poly (3-hydroxybutyrate)

PHB-g-CS30B PHB grafted on CS30B via ring opening polymerization

of β-Butyrolactone

PHBV poly (3-hydroxybutyrate-co-4-hydroxylvaleriate)

PLA Poly (lactic acid)

PMMT Poly-cationic modified sodium montmorillonite

Poly-DADMAC poly diallyldimethylammonium chloride

Poly-sorbate 80 Poly-oxyethylene (20) sorbitan monooleate

POSS polyhedral oligomeric silsesquioxanes

PVSK Potassium poly(vinyl sulfate)

Q Volume of the permeate water per unit time

RCL Regenerated cotton linter film

RCL1 Regenerated cotton linter/ Sodium-montmorillonite

99/01 (w/w) film


97/03 (w/w) film


xxvii

95/05 (w/w) film


90/10 (w/w) film

RMCC Regenerated microcrystalline cellulose film

RMCC1 Regenerated microcrystalline cellulose / Sodium-

montmorillonite 99/01 (w/w) film

RMCC3 Regenerated microcrystalline cellulose / Sodium-


RMCC5 Regenerated microcrystalline cellulose/Sodium-


RMCC10 Regenerated microcrystalline cellulose/Sodium-


ROP Ring opening polymerization

SEM Scanning electric microscope

Tc Cold crystallization temperature

TEM Transmission electron microscopy

Tg Glass transition temperature

TGA Thermal gravimetric analysis

THF Tetrahydrofuran

Tm Melting temperature

χc Relative crystallinity

XRD X-ray Diffraction

xxviii

WVP water vapor permeability

WVTR Water vapour transmission rate

Wwet Wet membranes

Wdry Dry membranes

ΔH*m Perfect enthalpy of 100% crystal

ΔHm Melting enthalpy

η* Complex viscosity

ω Angular frequencies

RH% Relative humidity

.ƞ Absolute viscosity

.l Thickness of the membrane

∆P Pressure difference

ρw Density of cellulose film

1

Chapter 1 Introduction

1.1. Background

The multibillion-dollar packaging industry seeks to meet the needs of fresh

produce, frozen baby food and the beverage sector in North America. With over 12

million tons of production every year, the packaging industry is the single largest market

for plastics. In the U.S. alone, the packaging industry comprises roughly a quarter of the

country’s overall plastic production [1]. In effort to enhace the quality of packaging to

meet product safety requirements, shelf-life extension, cost-efficiency, consumer con-

venience, and minimizing environmental damage, several innovative-modified and active

packaging materials are being developed[2].

Green fiber network (Paper), which has been widely used in packaging

applications, has a biodegradable advantage in that it is both environmentally friendly

and sustainable. Paper consists of a porous cellulose fiber network structure made up of

micro-fibrils which are comprised of long-chain cellulose molecules. In addition to being

an abundant natural polymer resource and easily renewable, cellulose also possesses

unique characteristics that include hydrophilicity, biodegradability, and a remarkable

chemical modifying capacity[3]. The hydrophilic nature of cellulose reduces the water-

vapor-barrier properties of paper mostly due to the presence of pendant hydroxyl group in

the repeating unit of cellulose (C6H10O5) and fiber network porosity. With a

predisposition to adsorbing water from its surroundings or food, paper packaging has a

tendancy to lose its physical and mechanical strength. Paper can undergo moisture

depletion following a process where the diffusion of water vapor molecules takes place

2

through pore spaces as well as in condensed form through the fiber cell walls

(Bandyopadthay et al., 2002). It has been established that large pores are primarily

responsible for controlling the water vapor transmission rate (WVTR) in regular paper.

These large pores create direct pathways for water vapor to penetrate through the regular

paper substrate, thus decreasing water vapor barrier properties as compared to a

thermoplastic polymer and even water soluble polymer. To prevent the growth of

microorganisms in foods, the permeation of water vapor through the packaging material

is essential[4, 5].

Yeasts, mold, and bacteria require an relative humidity (RH) level of more than

60% to start growing consequently, creating a vapor hydration atmosphere where high

water vapor permeability packaging has the potential to accelerate the process of food

spoilage[6]. It is imperative to accurately determine the moisture content of the food itself

when deciding on the most suitable material for food packaging[6]. Water vapor uptake

in dry products can increase moisture absorption and lead to a loss of firmness and

freshness while water vapor loss from fruits and vegetables can accelerate color change,

and as the products dry, leave them shrunken and visibly wrinkled. The quality of

sensitive frozen or chilled food also deteriorates rapidly when facing similar conditions[7,

8]. In this context that decreasing water vapor permeability becomes a central concern to

the development of paper as sustainable packaging material. Effective mechanisms that

contribute to the enhancement of water vapor resistance in hand sheet and cellulose-based

products include the simultaneous covering of the porous structure and hydrophilic

surface and/or decreasing pore size to reduce the ratio of direct pathways to tortuous

3

pathways for water vapor in order to penetrate the cellulose fiber network substrate[9-

11].

All these can be achieved by wet-end sizing[12], extrusion coating[13, 14], hot

melt lamination of polymer[15], chemical modification of cellulose like cellulose acetate

film[16], dissolving cellulose to produce regenerated cellulose[17], the mechanical and

chemical treatment of cellulose fiber to reduce fiber size in micro or nanometer in order

to make microfibril cellulose (MFC)[18], Nanofibrillar Cellulose (NFC)[19] and

nanocrystalline cellulose (NCC) paper[20]

Fossil-based synthetic thermos plastics have consistently been in high demand for

various reasons including being low cost and available in large quantities as well as

possessing exceptional mechanical and barrier properties, thermal stability, and

processability. These materials are extensively used in paper coating applications to

overcome paper’s porosity and hydrophilicity via extrusion coating and hot melt

laminating. In addition to it’s superior barrier properties against oxygen, grease, and

water vapor, fossil-based synthetic thermos plastics exhibit advanced heat sealability. Its

major downfall, however, is that fossil-based materials are non-biodegradable, and the

existence of such polymers in coated paper makes it all the more difficult to separate,

recycle, or composts the elements. The use of fossil-based materials can also lead to

environmental pollution and serious ecological and environmental issues.

As consumers are becoming increasingly more environmentally conscious, the

packaging industry has expressed interest in producing biodegradable and

environmentally friendly biopolymers and green methods.

4

Water soluble biopolymers like wheat-gluten[7], whey protein isolates (WPI)[8,

21], soy protein isolates (SPI)[22, 23], carboxylated methylcellulose, cellulose acetate,

and starches [24] have already been considered as coating and wet-ends sizing agents for

paper and paperboard in an effort to enhance the barrier properties of the cellulosic

substrates. Despite the fact that water vapor transmission rate of the resulted paper treated

with the aforementioned materials has been lowered to some extent, some key challenges

continue to remain.

Bio-thermoplastic hydrophobic such as Polyglycolide (PGA), Poly(lactide-co-

glycolide) (PLGA), Polycaprolactone (PCL), Poly(butylene succinate) (PBS), Poly(p-

dioxanone) (PPDO), Poly(trimethylene carbonate) (PTMC), poly(butylene adipate-co-

terephthalate) (PBAT), Polybutylene succinate (PBS), Polylactide (PLA), and

Polyhydroxyalkanoates (PHAs) are considered as qualified substitutes for general non-

degradable fossil-based polymers[25-27].

Among the different classifications of biodegradable polymers, polylactic acid

(PLA) and polyhydroxyalkanoates (PHAs) are generally used as a food-packaging

polymers for short shelf life products such as drinking cups, salad cups, containers,

overwrap, paper coating, and lamination films[28]. Obtained from 100% renewable

resources such as corn, wheat, and wood residue [29, 30], the biodegradable PLA and

PHAs are made without any difficulty. From the standpoint of biomedical applications,

PLA and PHAs have garnered great interest not only because of the necessity to

ultimately replace synthetic polymers, but also because of their useful physical and

mechanical characteristics[31, 32]. However, the use of both PLA and PHAs in food

5

packaging applications is quite limited due to relatively low gas barrier and water vapor

permeation in comparison to the conventional fossil-based polymers[33, 34].

One of the most important goals for the application of native cellulose in

packaging applications is the regeneration of cellulose based film. Transparent cellulose

films, regenerated from dissolved cellulose fiber, offer promising alternatives in terms of

reducing the pore size of cellulose fiber network and enhancing the production of

polymeric cellulose films[35]. To this end, alkaline/urea aqueous solution has proven to

be relatively low-cost, efficient, and non-toxic[36]. This solvent system has considerable

potential for the development of environmentally friendly and cost-efficient regenerated

transparent cellulose films. Regenerated cellulose film has substantially higher oxygen

barrier than fossil-based polymers such as poly(vinylidene chloride) and poly(vinyl

alcohol), even at relatively low humidity[17]. Regenerated cellulose based films have

exhibited far greater resilience to cellophane films in terms of their mechanical properties

as well as water vapor and oxygen barrier resistance owing to higher crystallinity [17, 37,

38]. However regenerated cellulose based films are notorious for their high water vapor

permeability and low thermo-mechanical properties compared to conventional polymers

and bio-thermoplastics [39-41].

To combat such debilitating problems, layered silicate nanocomposites have been

successfully used to enhance the properties of polymers and are being effectively utilized

by the paper coating and food packaging industries. Research has shown that the

properties of polymer matrices can be significantly altered with relatively low nanoclay

loadings (2−5 wt %)[42, 43]

6

1.2. Thesis objectives

This dissertation seeks to expand on the existing literature that addresses the

characterization of two bio-thermoplastics and their nanocomposites in an effort to

enhance the water vapor barrier properties of regular paper. Moreover, this study tackles

issues related to the characterization of regenerated cellulose nanocomposites intent on

developing “green-based” food packaging materials.

The overall objectives and goals of this dissertation are as follows:

To improve the water vapor barrier properties of cellulose fiber network through

the development of PLA and hydrophobically modified montmorillonite/PLA

nanocomposites for coating;

To synthesize a novel and compatible montmorillonite (MMT) nanoclay that can

be used in a well-dispersed phase in PHBV matrix system in order to improve

paper’s water vapor resistance through the application of various coating

methods;

To explore the extent of MMT clay exfoliation, enhance water vapor resistance

and surface adsorption, thermal stability, and rheology properties of PHBV;

To investigate the impact of different types of cellulose on WVTR of regenerated

cellulose film as a new-generation of green food packaging;

To assess the effects of sodium montmorillonite (Na-MMT) on mechanical,

crystallinity, water vapor adsorption, and water vapor barrier properties of

regenerated cellulose film;

7

To analyze the impact of various MMT modified with poly-cation cellulose on

mechanical, crystallinity, water vapor absorption, and water vapor barrier

properties of regenerated cellulose film;

To explore the extent of MMT clay exfoliation and clay content in order to

enhance water vapor resistance, surface adsorption, and the mechanical properties

of regenerated cotton linter film.

1.3. Thesis layout

This dissertation is comprised of ten chapters and the content; and the contributions of

each chapter are briefly described.

Chapter 1 describes the background of the project and the objectives of the

research.

Chapter 2, which consists of two sections, provides an overview of the literature

most salient to the current study. This Chapter initially delves into existing literature that

sheds light on the current study and highlights its significance. The second section of

Chapter 2 offers an overview of the various methods applied in the literature and the

methods utilized in the current research to assess the results.

Chapter 3 focuses on the grafted hydrophobic chemical, polyhedral oligomeric

silsesquioxanes (POSS) on Cloisite®30B montmorillonite (CS30B) surface applied for

the purposes of enhancing hydrophobicity clay and making it more compatible with PLA.

This chapter also explores the effects of various clay contents on PLA. PLA/CS30B and

PLA/POSS-CS30B were prepared via solution blending and characterized using FTIR,

XRD, and TEM. The superior dispersion of POSS-CS30B in PLA was observed using

8

TEM and XRD images. Moreover, PLA nanocomposites coated on copy paper were

carried out through two coating methods. The WVTR value of the PLA/POSS-CS30B

coated paper all clay content was much lower than PLA/CS30B coated paper.

In Chapter 4, CS30B was modified with grafting poly (3-hydroxybutyrate) (PHB)

via ring opening polymerization β-Butyrolactone monomers onto the clay surface. The

modification intended to enhance the compatibility of nanoclay with the PHBV matrix in

the process of achieving the exfoliated structure. After modification, the PHBV/CS30B

and exfoliated PHBV/PHB-g-CS30B nanocomposites were made using a solvent casting

method. Consequently, the properties of composites, as well as the WVTR of paper

coated with the composites, were determined. The results indicate that the water vapor

barrier of the coated papers was improved significantly.

Chapter 5 demonstrates the impact of modified CS30B (PHB-g-CS30B) and

CS30B nanoclay on surface morphological analysis, water adsorption and water vapor

barrier of PHBV nanocomposites films at various content up to 40% using AFM, XRD,

Contact Angel, and WVTR.

Chapter 6 illustrates the influence of modified CS30B (PHB-g-CS30B) and how it

is synthesized through the ring opening polymerization of β-Butyrolactone into the

layered structure of CS30B and CS30B nanoclay on melt processability, viscoelastic

behaviour, thermal stability and thermal properties of PHBV nanocomposites films at

various content up to 20% using melt rheometer, TGA and DSC.

Chapter 7 focuses on the preparation of regenerated microcrystalline cellulose

(RMMC) and cotton linter (RCL)/Na-MMT nanocomposites film through solution

blending in alkaline/urea/water solvent at low temperature. This process demonstrates the

9

impact of Na-MMT on the properties of regenerated cellulose films. The Na-MMT

dispersion played a pivotal role in improving the strength of the regenerated cellulose and

water vapor resistance. It was observed that the regenerated cellulose film with higher

molecular weight presented superior mechanical properties for all clay contents, while

RMMC and RMMC/Na-MMT nanocomposites showed lower WVTR value at all clay

content.

Chapter 8 investigates the preparation and characterization of surface modified

pre-exfoliated-montmorillonite (polycation-MMT), carried out by a series of ion

exchange reactions with quaternized microcrystalline cellulose (MCC) and Cotton linter

(CL). The celluloses were prepared and quaternized through grafting with 3-Chloro-2-

hydroxypropyltrimethylammonium chloride (CHPTAC). Quaternary celluloses with

various charge densities (380 to 1080 µeq/g) were used to modify Na-MMT. The effects

of quaternized celluloses on the particle size, surface characteristics, and moisture

adsorption properties of modified Na-MMT have also been examined at different pH

levels.

Chapter 9 concentrates on exfoliated and/or intercalated transparent

nanocomposites of regenerated cellulose and series of polycation-MMT (PMMT)

modified with six quaternized celluloses (QCs) with various charge densities from two

types of cellulose that were blended with cotton linter (CL) using environmental friendly

aqueous LiOH/urea solvent. The experiments conducted in this study affirm that the

mechanical barrier properties and even crystallinity of regenerated cotton linter film

improved substantially regardless of clay type. The UV results indicated that regenerated

cotton linter remains transparent, irrespective, of clay type.

10

Chapter 10 offers a detailed description of the effect of two types of polycation-

MMT, which were modified with quaternized-MCC and quaternized-α-Cellulose on

regenerated cotton linter at different clay content from 4 to 20 wt%. The UV results

reveal that regenerated cotton linter remains transparent regardless of clay content. Both

clays, which were modified with quaternized-MCC and quaternized-α-Cellulose, were

able to improve the cellulose strength and water vapor barrier of paper. While both

succeeded in increasing clay content significantly, the latter proved to be considerably

more efficient.

Chapter 11 summarizes the results stemming from this dissertation and discusses

recommendations for future research.

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16

Chapter 2 Literature Review

2.1. Introduction

This chapter presents an overview of the relevant literature and addresses

problems in line with the current research. This section starts off by contextualizing water

vapor transmission mechanism and water vapor barrier improvement methods in

polymer. This is followed by a summary description of PLA, PHBV, and their usage in

food packaging and paper coating. More importantly, the literature as an integral part of

the thesis outlines the need for PHBV and PLA nanocomposites for improving the

properties of biodegradable polymers.

The last part of the review offers a general overview of cellulose based fiber

network (paper) and regenerated cellulose film. The application of biodegradable and

sustainable materials used in packaging, as well as those with water vapor and/or grease

barrier properties relevant to the papermaking process will also be reviewed.

2.2. Water vapor transmission rates and permeability

Water Vapor Permeability (WVP) for different packaging material is considered a

key element in determining the most suitable application fields and content material

quality, particularly in the domain of food packaging. Since water vapor permeation in

packaging applications exerts considerable influence on food product quality, it is

essential to select material with the proper permeability characteristics while designing

packages for particular products.[1, 2].

17

WVP processes through polymer materials include moisture contact with the

polymer surface, adsorption on the side with higher moisture concentration, diffusion

through the polymer, and desorption on side with the lowest moisture concentration [3].

Based on an analogy of heat conduction through solid media, Fick’s first law can

propose a general law for diffusion. Essentially, Fick’s law establishes that the diffusive

flux causes the transfer of a substance from regions of high concentration to low

concentration with a magnitude that is proportional to the concentration gradient. One

way to describe Fick’s law is as follows Equation (2.1):

𝐽 = −𝐷𝑒𝑓𝑓

𝑑𝑐

𝑑𝑧 [2.1]

In this formula, J is the flux or the amount of material crossing a plane of unit area per

unit time (g/m2.s or ml/m

2.s); 𝐷𝑒𝑓𝑓 is the diffusion coefficient and has the dimensions of

(m2/s); dC is the diffusing substance’s differential of concentration in the film (g/m

3 or

ml/m3), and dZ is the differential thickness of the film (m)[4].

The negative sign in Fick’s law denotes a downward trend in the direction of

diffusion indicating that flow is directed toward decreased concentration. The diffusion

coefficient measures the amount of material that would diffuse across a unit area under a

unit concentration gradient in unit time. The application of Fick’s first law of diffusion

requires determining the flux J and the concentration gradient[5]. The fact that the law

specifies conditions of a constant concentration gradient is indicative of a steady state

Equation (2.2).

𝐽 = −𝐷𝑒𝑓𝑓

𝑐2 − 𝑐1

𝑧=

𝑄

𝐴 × 𝑡 [2.2]

18

The application of Henry’s Law, e.g., low concentration region, allows expression

of the driving force in terms of the partial pressure differential of gas[6].

A modified form of Henrys laws in terms of permeability is illustrated in the following

equation (2.3).

𝑄

𝐴 × 𝑡= 𝐷 × 𝑆

𝑝2 − 𝑝1

𝑍= 𝑃

∆𝑝

𝑍 [2.3]

Where Q is the amount of gas diffusing through the film (g or ml), c2 and c1 are

concentrations in the film (g/m3 or ml/m

3). In this equation, Z indicates thickness of the

film; A is the area of the solid (film) (m2), and t represents time (s). In addition, S is

Henrys law constant (solubility) given in mol/m3.atm or g/m

3.pa, Δp signifies the partial

pressure difference of the gas across the film (Pa), and P is the permeability coefficient

((ml or g). m/ m2.s. Pa).

According to the “solution-diffusion” model[7], a permeant (i.e. diffusing gas or

vapor) dissolves in the membrane material through diffusion in response to a

concentration gradient. Solubility is the process of adsorption of the sorbate in the

polymer and depends on the affinity (interaction energy) of the polymer with the

adsorbing molecule, the volume available for adsorption in the polymer, and the external

sorbate concentration. Diffusion is the concentration gradient driven process whereby the

adsorbed molecules are transported within the polymer, and diffusion properties are

characterized via diffusion coefficients.

Accordingly, the permeability constant, P, may be calculated by the diffusion

coefficient or diffusivity (D) and the solubility coefficients (S), which reflects the amount

of permeation in the polymer. The dimensions of D are area per unit time; S has

19

dimensions of the volume of gas (at stated conditions) per volume of the polymer film.

With a proper choice of units we may write:

𝑃 = 𝑆 × 𝐷 [2.4]

This correlation has been verified for permanent gasses (The molecules that make

up the air we breathe) and water vapor in a number of polymeric substances[8].

It is important to note that P, D, and S vary exponentially with reciprocal absolute

temperature for permanent gasses. Thus:

𝐷 = −𝐷𝑜𝑒𝑥𝑝 (−𝐸

𝑅 × 𝑇) [2.5]

𝑃 = −𝑃𝑜𝑒𝑥𝑝 (−𝐸𝑝

𝑅 × 𝑇) [2.6]

𝐷𝑜 and 𝑃𝑜 are a pre-exponential factors, R is the universal gas constant, and E and 𝐸𝑝 are

the activation energy and activation energy associated with the permeation, respectively.

Since P and D are essentially rate constants, E and Ep may be interpreted as the

energy of activation for permeation and diffusion, correspondingly. Moreover, 𝐷𝑜 and 𝑃𝑜

correspond to frequency factors in the Arrhenius equation of chemical kinetics. Barrel

has interpreted the measurements of these quantities in terms of transition state theory to

show that the activated state for diffusion involves many degrees of freedom, indicative

of the formation of a "hole" for the diffusion molecule to move into.

There are various mechanisms involved in diffusion processes in solids that

depend on the structure of the solid and the nature of the process. Diffusion through

porous materials could be molecular or bulk, Knudsen, and surface diffusion. Bulk

diffusion occurs when the pore diameter of the material is large in comparison to the

mean free path of the gas molecules. Knudsen diffusion is the transport mechanism

20

through pores that are small in comparison to the mean free path of the gas, e.g.,

mesoporous. The third type of mechanism for molecule transport in porous materials is

surface diffusion. In the process of surface diffusion, the adsorbed molecules on the

surface of the solid transport from one site to another towards decreasing concentration.

Owing to strong sorbate-adsorbent interactions, this becomes an activated process,

dissimilar to the other mechanisms involved in the diffusion process. However, in most

cases surface diffusion is anticipated to contribute minimally to the overall transport in

non-hydrophilic porous materials [9, 10].

2.3. Factors affecting permeability of polymeric films

The permeability of polymer materials is influenced by the internal structure of

the polymer including the degree of crystallinity, nature of polymer, chemical properties,

glass transition temperature and pores as well as pore size and pore size distribution [11,

12].

The two principal factors that affect macromolecular chain segmental motion are

temperature and the concentration level of the absorbed penetrant. An increase in

temperature intensifies the segmental motion which then allows the polymer to penetrate

structural transitions such as glass and melting transitions that impact the solution and

diffusion processes[13, 14].

2.3.1. Chemical structure

The temperature and the presence of the absorbed penetrant builds up excess free

volume or void in the system[14]. Noted for their balanced barrier properties and

characterized by their high strength, polymers containing polar groups such as ester

21

nitrile, Cl, and F, Poly(vinyl alcohol) - (PVA)- and uncoated cellophane are considered

excellent barriers for gases. They are however, inadequate in terms of moisture and water

vapor resistance and experience rapid deterioration in their gas barrier properties when

wet or swollen through water absorption[15, 16]. While polyolefin like high-density

polyethylene (HDPE) possess outstanding barriers against moisture, they tend to be

highly permeable to hydrocarbons[17]. Aliphatic polyamide (PA 6, PA 66), on the other

hand, is known for its excellent resistance to hydrocarbons while simultaneously being

highly permeable to water vapor[18].

Non-polar polymers such as polyethylene and polypropylene with flexible chains

have very low cohesive energy density (CED) and can influence the mass transport

properties. As a result, polyethylene possesses high oxygen and carbon dioxide

permeability. The low water vapor permeability of polyethylene has a much lower

affinity toward highly polar molecules and non-polar PE structural units, which cause

insignificant water solubility-diffusion of small water molecules through the substrate.

Large non-polar aromatic group on polymer backbone have the propensity to disrupt

chain packing and lead to an increased free volume, despite the rise in chain rigidity[19].

The methyl acrylate group-COOCH3 on PMMA (polymethylmethacrylate) causes higher

CED than PE and PP with corresponding lower gas permeability. Conversely,

functionalized hydrocarbon polymers with highly polar side groups such as PVOH

cellulose and nylon have a strong affinity for polar molecules including water vapor

molecules due to the presence of hydroxyl –OH groups [19].

These types of polymer customarily exhibit high CO2 and oxygen resistance when

dry. At the same time, the high water solubility of such polar polymers results in polymer

22

swelling and plasticization by dissolved water molecules. In a plasticized state, polymer

molecules become quite flexible. Plasticization occurs when the water molecule sorption

in the polymer matrix is considerably high. Sorbed water facilitates the diffusion and

permeation of oxygen or other gases into the polar polymer matrix[20].

Regenerated cellulose (i.e. cellophane) is also believed to have a high

concentration of hydroxyl groups per repeat unit and exhibits excellent gas barrier

properties if kept dry[21, 22]. While cellophane have been a popular choice for crisp and

cigar bags[23], their permeability to most gasses increases rapidly with increasing

humidity similar to that of EVOH. Cellophane is also unsuitable for coextrusion in

multilayer packaging.

2.3.2. Crystallinity

Polymers with oriented macromolecules tend to exhibit superior barrier properties

as opposed to polymers with un-oriented macromolecules. In this context, polymer

crystallinity plays a more effective role considering that permeability depends on the rate

of penetration of gas, liquid, or vapor molecules through the polymer matrix that also

depends on the physical structure of the polymer. Crystalline regions put up stiffer

resistance to penetrating molecules in comparison to amorphous domains. The inevitable

outcome is a decrease in permeability as the polymer crystallinity increases [24, 25].

As outlined in previous sections, polymer crystallinity plays a critical role on

barrier properties. While crystalline regions exhibit stronger resistance to a penetrating

molecule than amorphous domains, these impermeable crystal regions create a tortuous

path for any solute crossing the barrier. Consequently, the orientation of crystallites leads

23

to an increase in the diffusion path tortuosity. Additionally, the orientation of crystallites

reduces the segmental mobility of more flexible polymer chains in the amorphous phase

surrounding crystallites. It can thus be concluded that the higher the crystallinity fraction

of a semicrystalline polymer, the lower the water solubility in the polymer[26, 27].

2.3.3. Glass transition temperature

Glass transition temperature (𝑇𝑔) is an index commonly used to account the

rigidity of polymer chains. Polymers with high 𝑇𝑔 are renowned for their rigidity and

densely packed structure with lower diffusivities. To calculate the free volume in a

polymer, the sum of excess or unrelaxed free volume and the transient free volume

arising from segmental and sub-segmental rotational and vibrational motions need to be

calculated[28, 29]. In the glassy state (Below 𝑇𝑔), a polymer becomes hard and brittle as a

result of restricted mobility on the part of polymer chains consequently, the penetrant

diffusivity in this state is low while polymers in the rubbery state are characteristically

tough and flexible owing to free chain motion [30-32].

2.3.4. Porosity and Voids

The pore size distribution is also capable of impacting the water vapor

transmission rate[33]. For instance, narrow pore size distribution exhibits enhanced

barrier properties with increased resistance to the flow of water vapor across the

thickness of the film at STP[34, 35]. Similar to free volumes, pores also provide sites

where liquids and vapors can condense in a milieu that affords greater space for transport

24

than solid polymer A high level of porosity will amplify permeability by increasing the

solubility as well as the effective diffusion coefficient[36, 37].

An increase in film thickness leads to a decrease in the OTR values, supporting

the pore blocking theory, i.e., less connected pores throughout the film. Hence, density

measurements are indicative of a dense film structure. Therefore, density is deemed to be

higher in the bulk solution than at the surface of the film[37, 38].

In cases where the pores are linked (open pores) the diffusion rates through these

channels result in considerably higher permeation, in comparison to instances where the

pores are isolated (closed pores). Coating along with filler adding on a solid porous

surface has the potency to reduce surface porosity in addition to decreasing air and vapor

permeability[34].

2.4. Barrier properties improvement of a polymer

2.4.1. Barrier Improvement from Polymer Blends

The relatively low cost of the technology utilized to modify properties allows for

the widespread use of polymer blending[39, 40]. Some of the more conventional

techniques employed in polymer blending in an attempt to improve polymer barrier

properties include the addition of higher barrier plastics with a multilayer structure and

high barrier surface coatings. While effective in some aspects, these approaches have not

proven to be completely effectual on barrier properties[41]. The central premise of

creating a novel blend of two or more polymers is to extract the maximum potential from

the mixture instead of making siginificant alterations to the barrier properties of the

25

components. The past few decades have witnessed the development of an assortment of

blends of various polymers [42-45].

2.4.2 Barrier Improvement from Nanocomposites

Polymer nanocomposites are mixtures of a polymer matrix and reinforcing clay

that have at least one dimension in the nanometer range. Clay minerals such as

montmorillonite, hectorite, saponite, vermiculite and others are used as nanofillers. A

layer of silicate clay mineral is about 1 nm in thickness and consists of platelets of around

100 nm in width, so it represents filler with a significantly large aspect ratio. Nylon 6 –

clay nanocomposites developed by Usuki et al.[46, 47] was the first polymer

nanocomposite to be used practically. Since 1990 when it was first applied, many studies

and analyses have been reported.

Layered silicate nanocomposites have been successfully used to enhance the

properties of oil-based polymers and are being put in use by the paper coating and food

packaging industries. The properties of polymer matrices can be significantly altered at

very low nanoclay loadings (2−5 wt %)[48, 49]. These alterations are explained by the

high aspect ratio of nanoclay platelets and their nanolayered structure in the composites.

However, nanolayered structure creation in a polymer matrix is a major challenge, due to

the incompatibility of hydrophilic structure clay platelets with hydrophobic polymer

matrices.[50-52]

It is well known that the morphology and dispersion of silica nanoparticles and

other nanoparticles in polymers is one of the key factors affecting their properties. Besid

a good dispersion of the nano-filler, the modification of its surface with coupling agents

26

for a good compatibility with the polymer is also necessary[53, 54]. A high specific area

along with the respective loading determines the effective contact with polymer.

Depending on the interaction between clay and polymer and the morphology of

the clay layers in the polymer matrix, three different types of nanocomposite morphology

are obtained: conventional, intercalated, and exfoliated (Figure 2-1) .[55]

- Intercalated nanocomposites (Figure 2-1) obtained by insertion of polymer

macromolecules between clay platelets; in this case the clay particles are dispersed in an

ordered lamellar structure with large gallery height via insertion of macromolecules into

layers. In this case the insertion into the layered silicate structure occurs in a

crystallographically regular fashion regardless of organic modified clay.

- Exfoliated nanocomposites; (Figure 2-2) in this morphology the individual

layers of nano-clay are totally delaminated and dispersed in a continuous polymer matrix.

In an exfoliated nanocomposite the individual silicate layers are separated in a continuous

polymer matrix by an average distance that depends on the organic modified layered

Intercalation Exfoliation Micro-composite

Figure ‎2-1. Micro-composite, intercalated, and exfoliated structure

27

silicate loading. Usually the organic modified layered silicate of exfoliated

nanocomposites is much lower than that on intercalated one.

Figure ‎2-2. 3D pattern of layered silicate of exfoliated nanocomposites

The fully exfoliated nanocomposite morphology involves immense polymer

penetration with the individual delaminated platelets and randomly dispersed in the

polymer matrix. It is considered to be the most desirable morphology to improve the

water vapor barrier properties. The exfoliated morphology induces extremely large

surfaces and interfaces between the polymer and disperse phase which leads to an

increase in the tortuosity effect in the polymer matrix.[56, 57] A well-ordered staggered

polymer-silicate-layered nanostructure forces water vapour molecules to travel through

the film and to follow a tortuous path through the polymer matrix surrounding the silicate

particles. Therefore, this process will increase the effective path length for the diffusion.

28

That is why enhancing tortuosity is promising for the improvement of barrier properties

of biopolymer nanocomposites in food packaging and paper coating.[58, 59]

The chemical compatibility between clay and polymer matrices plays an

important role in the final nanocomposites morphology. Improving the compatibility of

clays with biopolymers, to generate homogeneous dispersion, is important for reducing

the WVTR for packaging purposes. One strategy to gain better interaction between clay–

polymer interfaces is to modify the clay surface.[60] Ring opening polymerization of

biopolymer monomers into the clay surface is an effective method to enhance the

compatibility of nano clay with biopolymer.[61-63] This kind of modified clay can be

used as a master-batch in an industrial extrusion process or extrusion paper coating

process to improve the water vapor resistance [61, 64, 65].

It is known that a low amount of organic modified clay dispersed in a polymer

matrix can lead to very high areas of interactions [61-63].

2.5. Biodegradable materials

2.5.1. Poly lactic acid:

Poly (lactic acid), PLA (Figure 2-3) is considered a typical biodegradable

synthetic semicrystalline polyester. Commercially obtained by ring opening

polymerization of lactones or by polycondensation of hydroxy-carbonic acids, PLA is

biodegradable material with exceptional mechanical and optical properties. Lactic acid,

the monomer of PLA, may easily be attained through the fermentation of renewable

resources such as corn, maize, wheat, wood residue, or other biomass[66, 67]. Production

29

of PLA from renewable resources and as a readily biodegradable material are the main

reasons behind a mounting commercial interest for this product.

Despite exhibiting low water vapor resistance, PLA displays excellent oxygen

barrier properties, even more so than low-density polyethylene (LDPE). It is these very

qualities that have earned PLA polymers the reputation of being one of the more cost-

effective alternatives to commodity petrochemical-based materials. In spite of its

advantages, the relatively high cost and low water vapor barrier properties of PLA in

comparison to fossil-based polymers has impeded its extensive use in areas such as

packaging application[68, 69]. The most important limitation facing PLA application,

particularly in food packaging is its low water vapor resistance. The biocompatible

properties inherent in PLA have nonetheless led to its widespread usage in the medical

field[70-72].

A study conducted by Shogren [24] reveals that crystalline PLA demonstrates

WVTR less than half of the amorphous PLA. The glass transition temperature (𝑇𝑔) leads

to distinct changes in the dynamic properties of amorphous PLA at and above the 𝑇𝑔, as a

consequence of chain mobility[73].

Figure ‎2-3.Chemical structure of PLA

30

The eco-friendly characteristics of PLA, has led to a recent upsurge in its

application as a polymer used to coat paper products. The relatively high resistance of

PLA films to aliphatic molecules is among a variety of factors attributing to its pervasive

use to coat paper. In line with this, paper coated with PLA was shown to resist grease

penetration after 120 h of testing at 55 °C, while paper coated with LDPE was incapable

of withstanding grease penetration in just 10 h[74].

Rhim and Kim[75] maintain that paperboard coated with PLA could be used as a

substitute for PE-coated paperboard to manufacture 1-way paper cups or containers for high

moisture foods such as beverage cartons and ice cream containers. It should be noted that

PLA film does not exhibit improved adhesion to paper in direct coating extruding

process. To achieve optimum results and sufficient adhesion, polylactic acid must be

applied in the highest possible temperature. The inherent brittleness of PLA causes leaks

and cracks whereby the coating does not endure the creasing or bending and extending

which are essential to producing plate or mold form products from PLA.

The main difficulty in the coating of PLA on paper surface is the rapid cooling of

the melt PLA after it emerges from the nozzle and before it ends up on the paper[76].

A number of researches have addressed aforementioned issue and investigated

various methods to overcome these weaknesses of PLA blends [77-80] and

nanocomposites [81, 82] in different applications. In one instance, Arrieta et.al[83]

prepared ternary PLA-PHB blends films for food packaging applications and plasticized

with a natural terpene D-limonene (LIM) with the dual objective of increasing PLA

crystallinity and obtaining flexible films.

31

Nanocomposites technology has the potential to improve polymer properties and

expand the applications of PLA. Raquel et al. [84] offer an overview of various PLA

based nanocomposites with different applications in packaging. In lieu of this, the effects

of fillers such as calcium sulfate and montmorillonite on PLA composites have also been

examined [85, 86].

2.5.2. Poly(3-hydroxybutyrate) and (Poly(3-hydroxybutyrate-co-3-hydroxyvalerate)

Poly(3-hydroxybutyrate) (PHB) (Figure 2-4) and its copolymers (Poly(3-

hydroxybutyrate-co-3-hydroxyvalerate), (PHBV), (Figure 2-5) is a biodegradable

semicrystalline polymer that can be produced by bacteria from biomass through natural

biosynthesis. This group of polymers is believed to have unique properties including

excellent chemical resistance, heat resistance, and rigidity. These are comparable to

isotactic polypropylene (PP), which posssess the highest water vapor resistance of

Polyhydroxybutyrate (PHB) compared to the existing biopolymers in the market. In light

of these qualities, these biodegradable polymers may very well come to be utilizied in

various applications in a not so distant future, particularly in packaging application[87-

89].

Figure ‎2-4. Chemical structure of PHB

32

Nevertheless, water vapor resistance of PHAs is much higher than those

conventional thermoplastics. Once they become relatively more cost-effective, the

application of PHB and PHBV in different industrial fields becomes economically viable

[90, 91].

Thus far, the high costs associated with PHB production have curtailed its

widespread application in various industries. Other factors that have attributed to PHB’s

lack of industrial development include its high crystallinity, mechanical brittleness, poor

processability[92], high melting temperature too close to its thermal degradation, low

thermal stability, and viscosity drop more than half for 5 min at melting temperature.

These considerations also prohibit a smooth transition from the use of conventional

plastics and their application to the biological denitrification process[93].

It is necessary to carry out further research to modify the mechanical, thermal,

rheological, and moisture barrier properties of PHAs. Furthermore, it should be explicitly

pointed out that the melting temperature close to the thermal degradation temperature

leads to a rather narrow window of processability. Therefore, the film blowing process of

PHB and PHBV are currently limited due to their slow crystallization and low melt

strength, i.e., low melt viscosity[94, 95].

In order to minimize these shortcomings, alternatives have been recommended.

These alternatives include the synthesis of copolymers by modification which leads to

Figure ‎2-5. Chemical structure of PHBV

33

lower melting point with hydroxyvalerate (HV). This copolymer is named PHBV.

However, copolymerization process is fairly important in avoiding thermal degradation

during melting process. This is while hydroxyvalerate (HV) increases the toughness of

pure PHB by reducing its crystallization, PHBV remains slightly thermally stable above

its decomposition temperatures and exhibits better processability. Nevertheless, the

commercially available PHBV biopolymer continues to be brittle with high elastic

modulus. As compared to traditional thermoplastic[96, 97] PHBV possesses a number of

key disadvantages including lower tensile strength, elongation at break, and lower water

vapor resistance.

There are several approaches that can be adopted to improve the thermal and

mechanical barrier properties, as well as melt processability of PHBV materials such as

blending[98-104] and nanocomposites[105-110].

Several researches have been carried out analyzing different polymers as potential

components to be blended with PHBV. These polymers, whose application is meant to

improve mechanical, thermal, and barrier properties, as well as the processability of

PHBV, include but are not limited to lignin,[111, 112] poly(D,L-lactic acid)[113-115],

starch[116, 117], Poly(p-vinyl phenol),[118] and cellulose[119].

Yet, research that focuses on PHBV rheological or viscoelastic properties has

been few and far between. For instance, Zhao et al. [120] studied the rheological

properties of PHBV as dispersed blended with PLA illustrated a Newtonian fluid

behavior. The study found that at 30/70 PHBV/PLA, blending ratio by weight exhibits

higher values of shear viscosity.

34

Modi [121] further examined the rheological properties of PHBV with PLA

grades 3051D, 4042D, and 6202D at 175 ºC using a micro-compounder. The study

revealed that the addition of PLA led to no significant improvement in the PHBV

viscosity.

One of the few attempts to prepare PHBV nanocomposites was carried out by

Wang et al. prepared poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV)/modified

montmorillonite (Cloisite 30B) nanocomposites with low clay content via melt blending.

Their investigatin revealed that intercalated modified montmorillonite acted as a

nucleating agent, increasing the temperature and rate of PHBV crystallization, while the

overall crystallinity, crystallite size, and melting temperature were diminished with the

increase of organo-montmorillonite (OMMT) in the nanocomposites [122].

Carli et al.[109] examined the effect of various nano-clays, Cloisite Na+, Cloisite

30B, and modified montmorillonite with trimethylene glycol mono-n-decyl ether with

three wt% clay concentration on the PHBV/nanoclay via melt mixing. They confirmed

that enhanced interfacial interactions between PHBV and modified montmorillonite with

trimethylene glycol mono-n-decyl ether leads to the production of nanocomposites with

increased toughness and higher modulus than those of Cloisite 30B and Cloisite Na+ the

pristine PHBV with significantly higher increase in the interlayer spacing and

intercalated morphology. Additionally, the study established that 𝑇𝑔 of nanocomposites

reduced slightly as a result of melt mixing.

The barrier properties of PHBV/ Cloisite®30B have also been investigated by

Wang et al. [123] with different clay contents using a twin-screw extruder. Their study

revealed that CO2 resistance diminished with increasing nanoclay content due to poor

35

PHBV – nanoclay interaction and a high affinity between CO2 molecules and quaternary

ammonium cations of modified clay. A decrease in the N2 permeability caused by the

tortuosity effect of the filler associated with a decline in the solubility within the matrix

was also measured. Moreover, the water vapor resistance of nanocomposites decreased

gradually for clay content of more than 2.5 wt% due to the increase in water solubility

and agglomeration of hydrophilic clay in PHBV matrix. The gradual decrease of water

vapor resistance could also be attributed to water vapor permeability process in

nanocomposites where there is a competition between the kinetic (diffusivity from

tortuosity effect) and thermodynamic (solubility).

Wang et al. [124] also noted that nanoclay has negative effect on biodegradability

of PHBV. This negative effect can be attributed to the intercalated morphology of clays

which deactivating the hydrolytic degradation process with reduction water and oxygen

permeability into polymer matrix. However, another study was conducted by these

researchers reported an increament in biodegradation tendency with PHB

nanocomposite[125]. A study by Carli et al.[126] revealed that the addition of 3% various

modified halloysites (HNTs) with organosilanes were prepared by melt processing in a

twin-screw co-rotating extruder. The addition of amino-silane-modified HNT with PHBV

reduced melt viscosity and had a negative impact on the thermal stability of PHBV. In the

remainder of organosilanes-PHBV nanocomposites, a rather insignificant enhancement of

the thermal stability was observed. The study also found out the extent of dispersion is

dependent on the clay organosilanes modifier with partially intercalated and

agglomerated morphology.

36

In another inquiry, Wolcott’s research group investigated the preparation and

properties of PHBV/cellulose nanowhiskers composites [127, 128]. The researchers

investigated the influence of homogeneously dispersed cellulose nanowhiskers through

solution blending on tensile strength, Young’s modulus, toughness, dynamic modulus

[128], and crystallization[129] in various clay content. The study demonstrated that well-

dispersed cellulose nanowhiskers enhance the mechanical and thermal stability of PHBV.

The reduction in properties occurred over the 2.3 wt % cellulose nanowhiskers content

because of cellulose nanowhiskers agglomeration[127]. The content of cellulose

nanowhiskers strongly impacted the degree of alignment under the electric field and

viscosity of suspension increased significantly with cellulose nanowhiskers concentration

more than four wt%. The well dispersion of cellulose nanowhiskers were obtained using

solvent blending which led to improvement in mechanical properties and thermal stability

at higher cellulose nanowhiskers concentration [129].

Sridhar et al.[130] utilized the solution casting method to prepare nanocomposite

PHBV reinforced with graphene. Their research demonstrated that variation in graphene

content increases the storage modulus and crystallinity of biopolymer. Sridhar et al.

maintained that improvement in the thermal degradation of PHBV could potentially be

linked to the high thermal stability of graphene. The researchers asserted that graphene

oxide dispersed well in PHBV matrixes. However, these findings are contradictory to

SEM results that indicate an altogether different outcome.

Montanheiro et al.[106] investigated the incorporation of functionalized

multiwalled carbon nanotubes (MWCNTs) and pristine MWCNT into PHBV. Their

findings indicate that MWCNTs sharply increased the thermal stability of PHBV to up to

37

30 °C, decreased its surface roughness, and facilitated the PHBV nanocomposites

crystallinity. The study confirmed that the electrical conductivity performance of PHBV

improved with the addition of a small quantity of MWCNT. However, functionalized

MWCNT was found to reduce conductivity values. Studies conducted by Sanchez-Garcia

et al.[131] and Lai et al.[132] also revealed that the addition of carbon nanotubes into

PHBV using solution casting enhanced its thermal stability, electrical conductivity, and

barrier with one wt% carbon nanotubes. The thermal properties and the crystallization of

the composites were investigated using differential scanning calorimetry and X-ray

diffraction. Accordingly, the nucleating effect of carbon nanotubes on the crystallization

of PHBV was confirmed.

Yu et al. [133] investigated the effect of PHBV-grafted multi-walled carbon

nanotubes (functionalized carbon nanotubes) on the mechanical properties of transparent

PHBV through a solution casting method. Functionalized carbon nanotubes were

uniformly dispersed in the PHBV matrix, which in turn facilitated homogeneous

crystallization of PHBV. The seven wt.% PHBV-g-MWCNTs improved the tensile

strength and Young’s modulus of the nanocomposite film by 88% and 172%,

respectively. In addition, well disperse nano-carbon created a tortuous path that caused

significant reduction in water uptake and water vapor permeability in all nanocomposites,

which differed substantially from neat PHBV.

2.5.3. Cellulose and Regenerated Cellulose

The biodegradable backbone of polymers from renewable resources has led to

their growing importance as potential substitutes for petrochemicals in different fields.

http://www.sciencedirect.com/science/article/pii/S0266353814000293

38

Plant cellulose is predominantly believed to be among the richest natural

polymers on earth. While cellulose (Figure 2-6) is the most abundant, sustainable,

compostable, biodegradable and reusable organic material on earth, its application in

various fields has been rather limited. For the most part, this is due to the fact native

cellulose faces serious application limitations. Despite its shortcomings, the application

of cellulose to produce various products possesses major advantages that include the

protection of the environment from pollution and keeping the world’s limited petroleum

resources intact. In this vein, it is important to understand how the molecular structure of

cellulose, beta-(1-4)-D-glucan allows chain-packing by strong inter- and intramolecular

hydrogen bonding, material [134-136].

Figure ‎2-6.‎ Chemical Structure of Cellulose

Since there is no melting temperature (in other words, the melting temperature of

cellulose is higher than the degradation temperature due to its hydrogen-bonded and

partially crystalline cellulose structure ), it appears that dissolving cellulose in appropriate

solvents is the most cost effective and constructive way to assemble it into desirable

shapes such as fibers, films, food casings, membranes, and sponges among other

things[137, 138].

More than a century has passed since regenerated cellulose (RC) such as fibers

(viscose, rayon) and films (cellophane) were produced largely through a metastable

O

OH

OH

HOH

H

O

O

H

H

H

H

H

OO

C

C

H

OH

OR

H

H

OH

HOHOH

OH

OH

39

soluble cellulose derivative[139]. However, this process is extremely time consuming and

tends to be environmentally damaging. The carbon disulfide used in this process and its

hazardous by-products are notorious for their harmful effects on human health[140]. In

response to this problem, less polluting aqueous and organic solvent systems such as

dimethylacetamide/lithium chloride[141, 142], zinc chloride[143, 144], N-

methylmorpholine-N-oxide[145], and sodium hydroxide aqueous solution[146-149] have

been explored and developed in recent years. However, most of these organic solvent

systems are deemed as ineffective from an industrial standpoint. This is primarily due to

the high toxicity of the solvent itself, the emission of explosive by-products, and the

difficulty experienced in solvent recovery.

The use of renewable materials in the packaging industry has earned many

accolades particulary from those concerned with the future of the environemntLindstrand

[150]. Cellulose fiber is the most abundant biopolymer on earth and due to it is

sustainable, biodegradable, and environmental friendly nature, it is extensively utilized in

numerous fields. However, the barrier properties of porous and hydrophilic cellulose

fiber network products are inadequate for the majority of barrier applications[151]. In

particular, the porous structure of cellulose fiber-based network (paper) continues to be

the main problem associated with these fibers which restricts their usage in moisture and

greasy barrier applications[152]. The smaller pore size of the regenerated cellulose based

films, compared to paper makes them a promising and feasible alternaive in the

preparation of cellulose-based food packaging [153-156]. In the past two decades, several

organic systems such as dimethyl sulfoxide (DMSO)/triethylamine/SO2, 3-dimethyl-2-

imidazolidinone (DMI), ammonia/ammonium salt (NH3/NH4SCN), N,N-

40

dimethylacetamide (DMAc)/LiCl, and alkaline/urea aqueous solutions have been

examined with regards to their ability to contribute to the production of regenerated

cellulose based films [157]. In this context, Alkaline/urea aqueous solution has proven to

be relatively low-cost and non-toxic. This solvent system exhibits high potential for the

development of cost effective and environmentally friendly regenerated transparent

cellulose films [157].

Non-toxic regenerated cellulose films formed by the aforementioned solvent

systems possess excellent mechanical properties, transparency, and barrier resistance to

oxygen, carbon dioxide, and water vapor. This is mostly due to the presence of more

extensive hydrogen bonding among cellulose chains, compared to native cellulose I films

[158-160]. Regenerated cellulose based films have substantially higher oxygen barrier

than practical oxygen barrier fossil-base polymers such as poly(vinylidene chloride) and

poly(vinyl alcohol), even at low relative humidity[159]. In contrast, the higher

crystallinity of regenerated cellulose films allows for remarkably superior qualities than

commercial Cellophane films in terms of the mechanical properties and water vapor and

oxygen barrier resistance [135, 159, 161].

It has been clearly established that regenerated cellulose films are notorious for

their high water vapor permeability and low thermo-mechanical properties compared to

conventional polymers and bio-thermoplastics [69, 162]. The high brittleness, moisture

sensitivity, and high barrier permeability of regenerated cellulose films have increased

demands for the development of new strategies in an effort to enhance the barrier

properties of these films [161, 163-165]. Extensive research has been conducted to

fabricate cellulose films from a variety of cellulose fibers utilizing various methods such

41

as the modification of regeneration process [159, 163], film surface modification [166,

167], layer-by-layer coating [168-170], polymer blending, and regenerated cellulose

nanocomposites [171-173].

Sodium montmorillonite (Na-MMT) is known as a nontoxic, inexpensive,

chemically and thermally stable, and commercially available nanoclay while

simultaneously being one of the most widely accepted reinforcing materials in the

polymer nanocomposite industry. The distinctive geometrical dimensions and high aspect

ratios of MMT nanoclay gives it some very unique properties that allows it to behave

quite differently from conventional clays or micro-clays. In line with this, regenerated

cellulose/nanolayer clay nanocomposites have the potency to be an innovative solution

that can work to improve regenerated cellulose based film’s weakness even at low nano-

clay content.

Cerruti et al. investigated the morphologies and thermal and thermal oxidative properties

of cellulose/organo-modified montmorillonite (3% MMT) nanocomposites prepared, by

dissolving cellulose in N-methylmorpholine-N-oxide water solutions [174]. The

homogeneous dispersion of clay in a cellulose matrix led to improvement in thermal,

oxidative, and mechanical properties of cellulose. The nanocomposites also demonstrated

increased degradation temperatures compared to plain cellulose. Mahmoudian et al.

(2012) investigated the permeability of oxygen, water absorption, morphology, and

mechanical properties of regenerated cellulose/montmorillonite nanocomposites prepared

in 1-butyl-3-methylimidazolium chloride using a casting method. All the properties of

nanocomposites with six wt.% MMT loading were drastically enhanced compared to the

unfilled regenerated cellulose. Mahmoudian et al. also scrutinized the circumstances

42

where the decrease in permeability and thermal stability and tensile strength were much

more pronounced in the nanocomposites as compared to regenerated cellulose despite the

presence of higher nano-clay content. Yang et al. (2014) recorded the mechanical

strength and gas barrier water absorption curves of cellulose film (with a viscosity

average molecular weight of 8.6×104 gmol

−1) and the corresponding intercalated

montmorillonite nanocomposites prepared from LiOH/urea solutions. They found that an

increase in clay content had an inverse effect on the oxygen permeability of the

nanocomposite. The oxygen permeability of nanocomosite contained 15 (wt%) clay at 50

and 75 RH%.decreased 42 and 33%, respectively, as compared to that neat cellulose film.

The tensile strength of the nanocomposite films was 161 Mpa and the Young’s modulus

was 180% higher than that of neat cellulose film. However, the increase in water

absorption and oxygen permeability was apparent in above 15 (wt%) clay content , most

likely due to the aggregation of silicate layers.

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64

Chapter 3 Poly Lactic Acid Nanocomposites Containing

Modified Nanoclay with Synergistic Barrier to Water Vapor

for Coated Paper

Abstract

The binary nanocomposites of poly lactic acid (PLA) with the montmorillonite modified

with trisilanol polyhedral oligomeric silsesquioxanes (Trisilanolisooctyl POSS®) were

prepared via a solution-blending process and coated on paper by bar coating and

compress hot melt coating methods. The resulting components were characterized with

Fourier transform infrared spectroscopy, and X-ray diffraction (XRD) techniques.

Moreover, the water vapor transmission rates (WVTR) for the coated writing paper were

determined using a gravimetric analyzer IGA-003. The results indicated that the modified

clay PLA nanocomposites enhanced the water vapor barrier properties of coated paper

significantly. The permeability of PLA nanocomposites to water vapor decreased by 74%

[26.0 g/(m2 day)], as compared to those of the paper coated with pure PLA. The

dispersion and phase behavior of the modified montmorillonite in PLA matrix was

revealed by Transmission electron microscope. The intercalation of montmorillonite with

PLA was further demonstrated using XRD. WVTR results indicated that the compress

hot melt coating of the nanocomposites is an effective method to improve the water vapor

resistance of coated paper.

65

3.1. Introduction

Paper and paperboard are the only renewable materials widely used in packaging

applications. [1-4]However, the hygroscopic and porous nature of paper limits its

potential when shelf life concerns are taken into account. [5, 6]Paper and paperboard

intended for packaging materials are often coated with fossil-based and non-

biodegradable polymers, such as PE, [7, 8]HDPE [9] and, PP[10] with high barrier

properties whose impact on the environment is under debate. Food packaging represents

a high volume commodity with the use of paperboard-based products for shipping and

handling purposes. Therefore, with the growing awareness of sustainability the focus on

developing coated paper for food packaging has shifted from conventional plastic

materials to more environmentally friendly alternatives, biodegradable coatings in

particular.

Poly lactic acid (PLA) synthesized from renewable resources has attracted much

attention, as it can achieve excellent mechanical properties at a competitive cost as a

sustainable replacement for traditional petrochemical-derived polymers in packaging and

paper coating applications[11]. PLA shows a limitation for the application of gas barrier

to be used for food packaging, and it also has relatively low resistance to water vapor

permeation compared with conventional fossil-based polymers[12, 13]. Therefore, much

work has been performed to improve the gas and water vapor permeation resistance of

PLA[13, 14]. Blending polymers with additives such as nano-clays, nano-particles, and

other polymers is a common and cost-effective approach to render materials with desired

properties[15, 16].

66

Biopolymer/clay nanocomposites have a wide range of potential applications[17,

18]. Recently the research on improving vapor barrier properties by using a small amount

of clay with a high aspect ratio and nanoparticles in polymer composites has drawn

tremendous interest for packaging and paper coating applications[19, 20]. Clay

nanoparticles, montmorillonite (MMT) nano-clay in particular, can reduce oxygen, CO2,

and water vapor diffusion by creating a complex network in the biopolymer matrix[20].

The presence of fully exfoliated and dispersed silicate clay layers and nanoparticles

creates a complex tortuous path for vapor molecules moving through the biopolymer

matrix[20]. Exfoliated and intercalated silicate clay layers creates a large diffusion length

and lowers the permeability, allowing biopolymer/clay nanocomposites to use in the

packaging of foods and beverages, thus prolonging the shelf time as well as the freshness

of the foods[21]. MMT/biopolymer nanocomposites display significant improvements in

mechanical, thermal properties, thermal stability, and enhances the retardation of oxygen

and CO2 diffusion[20]. MMT nanoclays also improve the water vapor transimisson

rate.18–21 AS MMT is hydrophilic naturally, it is necessary to modify MMT with

hydrophobic chemicals in order to increase its compatibility with most hydrophobic

polymer matrices and expand the interlayer space. This process allows large polymer

molecules to enter between the clay plates for greater dispersion of organoclay to achieve

a homogeneous nanocomposite. Polyhedral oligomeric silsesquioxane (POSS), one of the

smallest silica particles (1–3 nm), is well-known commercially [22, 23]. POSS consists of

a silica cage in the core with organic substituents R attached at the edges of the cage. By

the variation of R group it could be a novel amphiphilic chemical structure and reveal

extraordinary chemical, thermal, and mechanical properties[23-25]. Thus, POSS has not

67

only been used for practical applications such as moisture sensitive shape memory,[26]

compatibilizer,[27] clay modification and promoting melt-crystallization,[28] barrier

properties improvement,[29-31] and coatings and membrane technology,[31, 32] but also

as a powerful tool for the characterization of hydrophobic and hydrophilic polymer nano-

based composites[33, 34].

The objective of this work was to explore a novel modified MMT as a disperse

phase for PLA nanocomposites, which were applied to the paper surface via two different

coating processes in an attempt to lower the water vapor transfer resistance of the coated

paper by a biodegradable polymer.

3.2. Experimental

3.2.1. Materials

The PLA used in experiments was Ingeo 2003D, purchased from NatureWorks,

which has a M̅w = 193000 g mol−1

and M̅n = 114000 g mol−1

as determined by gel

permeation chromatography (GPC). The sample was purified by dissolution in

Chloroform, filtered to remove any insoluble matter, and precipitated in methanol and

dried under vacuum overnight at 40C before GPC measurement. Paper was selected

from a commercial calendar paper (C-paper) with 75 grammage, manufactured by

Domtar. Organo-modified montmorillonite, Cloisite® 30B, with bis-(2-hydroxyethyl)

methyl (hydrogenated tallowalkyl) ammonium captions, was kindly donated by Southern

Clay Products (Gonzalez, TX). The organic content of the organo-modified

montmorillonite, determined by TGA, was 22 wt%. Prior to use, the clay was dried under

vacuum at 110 oC for 1 h. TriSilanollsooctyl POSS

® (98.2%) were purchased from

68

Hybrid Plastics Ltd. Isophoron diisocyanate (IPDI 98%), anhydrous Tetrahydrofuran

(THF), Chloroform and dibutyltin dilaurate (DBTDL) catalyst were obtained from Sigma

Aldrich and used without further purification. All powder materials were dried under

vacuum overnight at 40 oC before use.

3.3. Characterization

FTIR spectra were recorded to identify the chemical structure of the Cloisite 30B

modified with triSilanollsooctyl POSS using a Perkin Elmer Spectrum 100 FT-IR

Spectrometer.

X-ray Diffraction The powder diffraction patterns of samples were collected on

Bruker AXS D8 Advance solid-state powder diffraction X-ray diffraction (XRD) system

with CuKα radiation. The d-spacing of the intercalated clays and nanocomposites were

determined by fitting Bragg’s equation (𝑛𝜆 = 2𝑑𝑠𝑖𝑛𝜃) in the range of 2–12 degrees (2θ)

at a scan rate of 10 min-1

.

Thermogravimetric Analysis (TGA) The tests were carried out using a thermo-

gravimetric analysis instrument (TA Instruments SDT Q600) in a temperature range of

ambient to 600 °C at a rate of 15 K/min-1

under the flow of N2 with a flow rate of 100

ml/min-1

. The data obtained from the TGA instrument were in the form of weight

percentage versus temperature.

Transmission Electron Microscope (TEM) analysis was examined using a

JEOL 2011 STEM operated at an accelerating voltage of 200 keV. Molded

nanocomposites by hot press were first trimmed with iron knives. Subsequently, the

69

ultrathin sections were microtomed from these faces with a diamond knife. The

microtomed sections were collected in a water-filled boat.

Water Vapor and Gas Transmission Rates (WVTR) of all coated paper

samples were performed on IGA-003 (Hiden-Isochema, Warrington, UK) which consists

of a high sensitivity microbalance (0.1 µg) and a turbomolecular high vacuum pumping

system (see Scheme 3-1), in accordance to the methods described in TAPPI standard

T464 om-12 (2012) and ASTM E96/E96M-05 (2005). In this experiment, coated paper

samples (coating thickness 25 lm) was pre-conditioned on WVTR test measurement

conditions for 72 hours to reach water adsorption equilibrium. The round coated paper

samples were clamped in a permeation cell which was tightened by six screws. The

measurements were carried out at 50% of the relative humidity (RH%), which was

achieved by saturated magnesium nitrate salt and flowing dry nitrogen gas at a flow rate

of 20 mL/min. After the permeation cell was placed in a chamber, the data was collected

after 1 h to allow the transmission to reach a steady state. The chamber temperature was

controlled at 23 °C. The weight reduction of the container, because of the moisture

transfer, is proportional to the testing time. At constant temperature and RH (%), WVTR

can be calculated from the change in the weight of the container, at a specified time

interval, and the area of exposed coated paper, as described by the following equation

(3.1):

𝑊𝑉𝑇𝑅 =𝑊𝑒𝑖𝑔ℎ𝑡 𝐶ℎ𝑎𝑛𝑔𝑒

𝐴𝑟𝑒𝑎 × 𝑡𝑖𝑚𝑒 [3.1]

70

Scheme ‎3-1. Schematic illustration of IGA setup for WVTR testing

3.4. Preparation of the Surface Modified Clays

Six grams of Cloisite®30B (CS30B), (Scheme 3-2) which contains approximately

7.1 mmol/g hydroxyl groups, and 1.2 g of trisilanolisooctyl POSS® were added in THF

(100 mL) in a 250 mL of three-necked, round-bottom flask equipped with a magnetic

stirrer, nitrogen inlet, thermometer, and condenser with a drying tube. The suspension

was then heated up to 50 °C, and two drops of dibutyltin dilaurate (DBTDL) were added.

Then IPDI (0.2 mL) drop wise was added and the reaction continued for 5 h. Excess

ethylene glycol (0.2 g) was added to the mixture to react with the remaining IPDI for 1 h

and cooled down to 0 °C. The mixture was filtered, washed three times with chloroform,

subsequently washed three times with deionized water, and then dried overnight in a

vacuum oven at room temperature.

71

Scheme ‎3-2. Preparation of the Surface Modified Clays

3.5. Preparation of PLA/MMT Nanocomposites

The PLA nanocomposites were prepared through a solution method with the

addition of 5 and 10 wt % of Cloisite®30B (CS30B) and 1, 3, 5, and, 10 wt % surface

modified Cloisite®30B (M-CS30B). For the fabrication of the nanocomposites, the

following process was employed: First, appropriate amounts of PLA and clay were

dissolved separately in chloroform. Then, the PLA solution and the clay solution were

mixed together and stirred with sonication for 1 h. Nanocomposites powder was prepared

by precipitation of the solution into an excess of methanol, and then the filtering process

was followed. Finally, the end products were dried at 70 °C under vacuum for 3 days to

remove the solvent completely.

72

3.6. Paper Coating

The solution bar coating and compress hot melt coating methods were used to

coat PLA and PLA nanocomposites onto paper. The solution bar coating paper were

prepared at room temperature, under magnetic stirring, by dissolving 0.306 g powder of

each sample in 6 mL of chloroform sonicated for 1 h. Each coating solution was then

spread onto paper surface (10 × 10 cm, thickness 90 µm) mounted on a Teflon sheet

using a glass spoon-shape bar and dried under ambient conditions (23 ± °C) for 24 h. The

compress hot melt coated paper was prepared in the following steps: 0.306 g powder of

each sample was dispersed in 30 mL methanol and homogenized at 9000 rpm for 10 min

using a homogenizer (Nissei AM-9, Japan). Then the nanocomposites polymer were

spread onto paper (10 3 10 cm, thickness ca 90 lm) mounted on a Teflon sheet (10 × 10

cm) and dried overnight in an oven at 70 °C. Finally, covered nanocomposite powder

papers were placed between two Teflon sheets and inserted into a hydraulic hot press

(Carver laboratory Press model 3925; Carver, Wabash, IN), heated to 190 °C and

subjected to a pressure of 1 bar for 2 min.

3.7. Results and discussion

Analysis of the chemical structure of Surface Modified clay by Fourier

Transform Infrared (FTIR) spectra confirmed a reaction between the silanol hydroxyl

groups of trisilanolisooctyl POSS® and hydroxyl groups of CS30B. Figure 3-1 shows the

FTIR spectra of CS30B, and M-CS30B. Figure 3-1(a) (CS30B) exhibits a very broad

band centered at 3629 cm-1

corresponding to the structural –OH groups in the clay

sample. Assigned to the stretching vibration of the hydrogen-bond group, the absorptions

73

at 2924 and 2852 cm-1

can be attributed to asymmetric and symmetric stretching

vibrations of C–H bonds, whereas the absorption at 1469 cm-1

can be attributed to

methylene bending vibrations, 1032 cm-1

. Also an intense broad band in the range of

1100– 1000 cm-1

, can be readily assigned to the Si–O–R because of asymmetric Si–O–C

stretching; these peaks are also noted in surface modified CS30B. In addition, Figure 3-

1(b) exhibits the presence of a Si–CH2–R group indicated in 1231 cm-1

. A peak at 952

cm-1

confirms the presence of incompletely condensed Si-OH, 1028 cm-1

can be assigned

to Si-O-Si originating from cage-structured POSS. The N-H bending vibrations observed

at 1526 cm-1

and the strong peak at 1724 cm-1

confirmed the presence of cyclopentyl

carbonyl-amine. The peak at 1473 cm-1

corresponded to the –CH2- scissoring and

rocking modes. All these peaks provide strong evidence for the successful modification

of CS30B with trisilanolisooctyl POSS.

74

Figure ‎3-1. FTIR spectra of (a) CS30B and (b) M-CS30B

3.7.1. Morphologies of the PLA Nanocomposites

The clay minerals and the corresponding nanocomposites were analyzed by

XRD in order to obtain information on the basal spacing and, any changes

resulting from the processing. The interlayer distance (d-spacing) of clays inside

the nanocomposites is obtained from Bragg’s law:

𝑛𝜆 = 2𝑑𝑠𝑖𝑛𝜃 [3.2]

Where d is the interlayer distance of clay, kw is the wavelength used, n is the order,

which is equal to 1 for the first order, and h is the measured angle.

XRD analyses of the CS30B and M-CS30B clay PLA/CS30B and PLA/M-CS30B

nanocomposites are illustrated in Figure 3-2. After the grafting of trisilanollsooctyl

75

POSS® onto CS30B, the d spacing for organo-modified clay initially presents at

2θ=4.8 (corresponding to a basal space of 1.82 nm), decreases to 2θ=3.8 which

means d spacing is increased to 2.34 nm, thus indicating swelling of the nanoclay.

The presence of second-order signals in the M-CS30B indicates a high degree of

order in the modified clay mineral. Diffraction peaks were observed in the small

angle region of the XRD patterns of PLA/CS30B and PLA/M-CS30B. At high

nanoclay content (10 wt % clay), the shoulder centered on 2θ=3.0 and 2.6,

respectively. For PLA/CS30B with 5 wt % clay content, the shoulder centered on

2θ=2.8, which indicating that the dispersion is less accomplished in PLA/CS30B

that PLA/ M-CS30B nano-composite and partially intercalated structure with more

aggregations of nanoclay may take place in the PLA matrix. This aggregation may

impact the water vapor transmission rate because the aspect ratio of unmodified

clay, which is a crucial factor for nanocomposites permeation properties in

aggregated nanocomposite, is much lower than that of the intercalated or

exfoliated nanocomposite. The aggregated clay may cause the formation of pores

in the polymeric matrix, which potentially create preferential diffusion pathways

for water vapor transport within the nanocomposite[35-37]. The XRD pattern of

the PLA/M-CS30B 95/05 wt/wt % sample shows no peak, which might be an

indication of an exfoliated morphology or a disordered aggregate structure in the

matrix. In order to interpret the XRD observations, the phase behavior of the

blends was studied using a TEM.

76

Figure ‎3-2. X-ray diffraction patterns of MMT and PLA nanocomposites.

The nanostructures revealed by the TEM (Figure 3-3) corresponded to the

inferences drawn from XRD analysis and for confirming the formation of

intercalated structures in polymer nanocomposites. The PLA/MCS30B 95/05,

PLA/M-CS30B 90/10, and PLA/CS30B 90/10 Nano-composites micrograph

indicating the differences in the extent of intercalation and distribution of the clay

layers. In Figure 3-3(a), the TEM image shows that the clay is intercalated in the

polymer matrix with the samples at 5 wt % MCS30B loading, which creates a

tortuous path for vapor molecules moving through the biopolymer matrix. The

larger diffusion length leads to lower water vapor permeability [38, 39]. On the

77

other hand, as shown in Figure 3-3(c) with an increasing clay level, an increased

degree of aggregation has also been seen for PLA/M-CS30B 90/10. The large

particle aggregates were observed in TEM micrographs [Figure 3-3(c)], indicating

that the CS30B particles are neither intercalated nor dispersed uniformly in the

polymer matrix.

Figure ‎3-3. TEM images of PLA nanocomposites: (a) PLA/M-CS30B 95/05, (b) PLA/M-

CS30B 90/10 and, (c) PLA/CS30B 90/10 at different magnification.

3.7.2. Thermal decomposition of the PLA/Caly NanoComposites

TGA is a standard method of studying thermal decomposition of polymers. Figure

3-4 shows TGA thermogram PLA and PLA nanocomposites under N2 atmosphere. All

the samples show a large drop in the TGA curve. The TGA curves show that the onset of

78

decomposition temperature of each nanocomposite sample is higher than that of pure

PLA. Onset and peak of degradation temperatures of the nanocomposites shift to higher

temperatures. PLA/M-CS30B nanocomposites are more thermally stable than

PLA/CS30B nanocomposites. The intercalated and homogeneous nanocomposite is the

most thermally stable one among the samples. POSS itself is very thermally stable[22],

therefore, the MMT modified with POSS provides better thermal stability, compared to

that of alkyl-ammonium modified MMTs. As a result, the MMT modified with POSS

might diminish the application restriction of the organoclay in the engineering plastics,

especially for those processed at high temperatures.

Figure ‎3-4. Weight loss for pure PLA and PLA nanocomposites

79

3.7.3. The WVTR Measurement

Evaluated WVTR is defined as the steady water vapor flow over time through area of a

body, normal to specific parallel surfaces, and under 50% relative humidity (RH) at 23

°C.

Water vapor transmission rate is a measure of the passage of water vapor through a

substance. There are many industries where moisture control is critical. Moisture

sensitive foods and pharmaceuticals are put in packaging with controlled WVTR to

achieve the required quality, safety, and shelf life. Table I reveals variation in WVTR of

PLA nanocomposites coated on c-paper with two different coating methods. As shown in

Table 3-1 and Figure 3-5 by increasing the coated nanocomposites clay content from 0%

to 5%, the WVTR value was reduced from 72 to 26 g/(m2 day), because of the increase of

the effective diffusion path length for the water vapor[19-21]. By further increasing the

clay content to 10%, the aggregated structure of MMT, as observed in TEM and XRD

results, possibly provided channels or microvoids at the interface of polymer and

clay[35]. As shown in Table I, the greater reduction in WVTR associated with M-CS30B

might be attributed to the hydrophobic nature of trisilanolisooctyl POSS® and the higher

d spacing of M-CS30B, which is also crucial for good dispersion of MMTs in the

polymer matrix. Furthermore, Figure 3-5 shows that the WVTR of PLA/M-CS30B

nanocomposite coated paper by compress hot melt method was significantly lower than

that of the paper coated by bar coating method, regardless of M-CS30 content. This may

be explained by the fact that more compact structure and smoother surface were formed

through the compress hot melt coating method, thus creating the coating layer with less

80

cracks and pores, particularly compared with those formed via dispersion (bar) coating

method[14].

Figure ‎3-5. The WVTR of PLA/M-CS30B-coated paper with bar and compress hot melt

coating methods

Table 3.1. WVTR results of W&P paper Bar-Coated and Compress Hot Melt Coated with

PLA, PLA/CS30B, and PLA/M-CS30B Nano-Composites

Sample: wt/wt %

WVTR bar coated

WVTR Hot melt coated

C-Paper 700 ± 12 700 ± 12

PLA 72 ± 3.0 55 ± 2.1

PLA/CS30B 95/05 50 ± 3.5 35 ± 2.7

PLA/CS30B 90/10 67 ± 3.7 48 ± 2.9

PLA/M-CS30B 99/01 59 ± 3.1 51 ± 2.5

PLA/M-CS30B 97/03 34 ± 2.8 26 ± 3.2

PLA/M-CS30B 95/05 26 ± 2.6 19 ± 1.4

PLA/M-CS30B 90/10 41 ± 3.1 24 ± 1.2

a Test conditions: RH 50% @ 23o C coat weight 25±2 g/m2 and coat thickness 20 ± 2 µm.

81

Conclusions

This study clearly demonstrated that the incorporation of CloisiteVR 30B MMT clay and

the Cloisite® 30B clay modified with POSS, either via a solution blending process or

paper coating using bar coating and hot press coating methods, improved the water vapor

barrier properties of coated paper significantly. The paper coated with binary and ternary

nanocomposites diminishes the water vapor transmission rate by 74% [26 g/(m2 day)], as

compared to the paper coated with pure PLA. The dispersion or phase behavior of the M-

CS30B in PLA matrix was revealed by XRD. The interclation of clay with PLA was

further demonstrated using TEM. Overall, the hot press coating method is an effective

approach in improving the water vapor resistance of the coated papers.

Acknowledgment

Authors are grateful to NSERC Strategic Network—Innovative Green Wood Fibre

Product (Canada) for funding and NSF China (No. 51379077).

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87

Chapter 4 Preparation and Characterization of Exfoliated PHBV

Nanocomposites to Enhance Water Vapor Barriers of Calendared

Paper

Abstract

The water vapor transmission rates (WVTRs) of poly(3-hydroxybutyrate-co-4-

hydroxybutyrate) (PHBV) and PHBV/nanocomposite-coated papers were measured at

various levels of relative humidity and temperature. The coating of PHBV on base paper

was performed with two different methods, and the more effective one for lowering

WVTR values was utilized for coating the PHBV/nanocomposites on the paper.

Nanocomposites of PHBV were prepared with a commercial montmorillonite

Cloisite30B (CS30B), and a novel modified clay was obtained via a solution-blending

process. To prepare the series of modified clays, which were coined here as HPGC,

MPGC, and LPGC, we intercalated β-butyrolactone in the Cloisite30B clay by a ring-

opening polymerization. The morphology of the clay and nanocomposites was revealed

by X-ray diffraction. Additionally, the dispersion and phase behavior of the clay in the

PHBV matrix was observed using a transmission electron microscope. It was found that

the coating method and clay exfoliation were the most important factors affecting water

vapor permeability. The water vapor barrier of the coated papers was improved

significantly if the surface of the base substrate was prespray-coated with a PHBV

suspension prior to the laminate coating of the PHBV film with a hot press. The papers

coated with exfoliated PHBV/nanocomposites exhibited even lower WVTR values.

88

Overall, PHBV, PHBV/CS30B, and PHBV/HPGC coating treatments lowered the

WVTR values by 46, 56, and 118 times, respectively. The resulting coated paper is

promising as a green-based packaging material due to an improved moisture barrier.

Keywords: Nanocomposites, Surface treatment, Coatings; Clay,

polyhydroxybutyrate, Exfoliation, and WVTR

4.1. Introduction

Food packaging represents a high volume commodity with the use of paperboard

based products for shipping and handling purposes[1]. Paper is considered as the most

economical and environmental friendly material for food packaging. However; the use of

papers in food packaging applications largely depends on the control of the hydrophilic

nature and enhancement of the porous structure of paper [2-4].

Oil based polymers, due to their hydrophobic and film-forming characteristics, are

in most cases chosen as paper coating materials to overcome the inadequacies of the

paper[5-7]. However; due to coated synthetic polymer layers, the coated paper loses

recyclability and is non-biodegradable [8, 9]. Research on sustainable alternative

materials for paper coating has been a hot topic for over a decade. In an effort to produce

more environmental friendly and renewable materials, biopolymers produced from

renewable sources, are very promising as sustainable paper coating material[10, 11].

Biopolymers, for being environmental friendly, have the upper hand as opposed to oil

based synthetic paper coatings. However; in practice, biopolymer have many limitations

such as high production costs, limited oxygen barrier, physical and mechanical

89

properties, and a low water vapor barrier. These limitations do not allow ordinary

synthetic plastics to be fully substituted by the biopolymers [12, 13].

High WVTR of biopolymers is a crucial element that needs to be controlled in the

food packaging and paper coating industry. The overall water vapor barrier property of a

polymer is influenced by its chemical structure, polymer–gas interactions,[14] glass

transmission temperature, crystallinity, and hydrophobicity[15, 16]. Additionally, water

vapor diffusion rate is directly dependent on the polymer film thickness[17, 18].

The Poly (3-hydroxybutyrate) (PHB) is commercially produced biodegradable,

semicrystalline and hydrophobic aliphatic polyester that is synthesized by

microorganisms. It exhibits somewhat better performance in many situations than

synthetic oil base polymers and can be used to reduce the demand for landfill sites [19,

20]. PHB and its copolymers, especially poly (3-hydroxybutyrate-co-4-

hydroxylvaleriate) (PHBV), are becoming potential alternatives as green food packaging

material since they have very effective water vapor resistance and significantly higher

crystallinity than most of the biodegradable aliphatic polyesters like crystalline

polylactide (PLA) and Polycaprolactone (PCL)[21-23]. In the last decade, PHB coated

paperboard has been successfully used for the packaging of ready-to-eat (RTE) and

frozen meals, while PHBV coated boards have been used for dry products, dairy

products, and beverages[24, 25]. However, the use of both PHB and PHBV is still limited

in food packaging applications due relatively low gas barrier and water vapor permeation

compared with conventional fossil-based polymers. Much work has been done to

improve the gas and water vapor permeation resistance of PHB and PHBV, such as:

90

polymer blending[26], surface laser-treatment[27], carbon nanotubes composites[28], and

layered silicate nanocomposites[29].

Recently, layered silicate nanocomposites have been successfully used to enhance the

properties of oil-based polymers and are being put in use by the paper coating and food

packaging industries. The properties of polymer matrices can be significantly altered at

very low nanoclay loadings (2−5 wt %). These alterations are explained by the high

aspect ratio of nanoclay platelets and their nanolayered structure in the composites.

However, nanolayered structure creation in a polymer matrix is a major challenge, due to

the incompatibility of hydrophilic structure clay platelets with hydrophobic polymer

matrices [30, 31]. Depending on the interaction between clay and polymer and the

morphology of the clay layers in the polymer matrix, three different types of

nanocomposite morphology are obtained: conventional, intercalated, and exfoliated [32,

33]. The fully exfoliated nanocomposite morphology involves immense polymer

penetration with the individual delaminated platelets and randomly dispersed in the

polymer matrix. It is considered to be the most desirable morphology to improve the

water vapor barrier properties. The exfoliated morphology induces extremely large

surfaces and interfaces between the polymer and disperse phase which leads to an

increase in the tortuosity effect in the polymer matrix [34, 35]. A well-ordered staggered

polymer-silicate-layered nanostructure forces water vapor molecules to travel through the

film and to follow a tortuous path through the polymer matrix surrounding the silicate

particles. Therefore, this process will increase the effective path length for the diffusion.

That is why enhancing tortuosity is promising for the improvement of barrier properties

of biopolymer nanocomposites in food packaging and paper coating [36, 37].

91

The chemical compatibility between clay and polymer matrices plays an

important role in the final nanocomposites morphology. Improving the compatibility of

clays with biopolymers, to generate homogeneous dispersion, is important for reducing

the WVTR for packaging purposes. One strategy to gain better interaction between clay–

polymer interfaces is to modify the clay surface[38]. Ring opening polymerization of

biopolymer monomers into the clay surface is an effective approach to enhance the

compatibility of nanoclay with biopolymer [39, 40]. This kind of modified clay can be

used as a master-batch in an industrial extrusion process or extrusion paper coating

process to improve the water vapor resistance [41, 42].

In this work, a surface-modified organo-montmorillonite was prepared through

the ring opening polymerization of β-Butyrolactone in the presence of Closite®30B

montmorillonite and catalyst. Additionally, we investigated the water vapour

transmission rate of PHBV, PHBV/ Closite®

30B nanocomposites, and PHBV. The

surface modified organo-montmorillonite coated paper through two different coating

methods. In addition, the amount of incorporated nanoclays into the PHBV matrix, its

dispersion, spacing distance and morphology of the obtained nanocomposites were

assessed by means of XRD and transmission electron microscopy (TEM) techniques.

4.2. Experimental

4.2.1. Materials

The PHBV (grade ENMAT Y1000P) used in the experiments was purchased from

Tianan Biological Material Co., Ltd. (China, Ningbo), which has a M̅w = 387000 g

mol−1

, polydispersity of 2.70 as determined by gel permeation chromatography (GPC)

92

and ENMAT Y1000P content of 3% hydroxyvalerate. Calendared paper with 75 g/m2,

were prepared in Limerick Pulp & Paper Centre of University of new Brunswick.

Organo-modified montmorillonite, Cloisite® 30B, was kindly donated by Southern Clay

Products (Gonzalez, TX). The organic content of the organo-modified montmorillonite,

determined by TGA, was 22 wt%. Prior to use, the nanoclay was dried under a vacuum at

110 oC for 4 hours. The β-Butyrolactone (98 %) was purchased from Sigma Aldrich and

dried with Calcium hydride for 2 days by stirring, and distilled under reduced pressure

prior to use. Polysorbate 80 (a nonionic surfactant) was purchased from Alfa Aesar.

Dibutyltin(IV) oxide (98%), Dibutyltin dichloride (97%), anhydrous Ethanol (EtOH),

anhydrous Toluene, Chloroform, and hexane were obtained from Sigma Aldrich and used

without further purification. All powder materials were dried before use under vacuum

overnight at 50 °C.

4.3. Methods

4.3.1 Preparation of the Surface Modified Clays (Scheme 4-1):

A 4 g of Cloisite®30B (CS30B), (which contained 7.1 mmol hydroxyl group) and

138 mmol of β-Butyrolactone, were added in toluene (100 mL) in a 250 mL of three-

necked, round-bottom flask equipped with a magnetic stirrer, nitrogen inlet, thermometer,

and condenser with a drying tube. In order to ensure proper dispersion, the suspension

was homogenized by vigorous agitation stir bar overnight at room temperature.

Afterward, the desired amount of catalyst[43] was added to the mixture and slowly

heated up to 120 °C (Table 1). Polymerization was carried out at 120 °C for 12 hours.

The solvent was removed under reduced pressure. The solid mixture was washed three

http://en.wikipedia.org/wiki/Surfactant

93

times with chloroform and centrifuge subsequently at 10000 rpm. Then, it was washed

three times with ethanol and deionized water and then dried 72 h in a vacuum oven at

room temperature. The same modification procure was used for clay modification

without toluene through a bulk ring polymerization. As shown in Table 1, the

modification yield was not significant.

The nanocomposites were prepared with HPGC (40% reaction yield), MPGC

(22%), and CS30B (Table 1).

Scheme ‎4-1. Preparation of the Surface Modified Clays

4.3.2. Preparation of PHBV Nanocomposites and PHBV laminated films

PHBV laminated films were prepared by dissolving appropriate amounts of

PHBV in chloroform at 60 °C. Each coating solution was then casted onto a silicon

wafer, with weight of 10, 13, 17, 20, 23 and 25 g/m2 base on solid chemical, under a

fume hood at room temperature for 3 days to remove the solvent completely. Afterwards,

the films were removed from the silicon wafers.

For the fabrication of the nanocomposites, the following process was employed. At first,

appropriate amounts of nanoclay (CS30B, MPGC and HPGC) were dispersed in

chloroform for 4 hours followed by additional ultrasonic for 1 hour. The nanoclay

94

suspension was transferred to a 250 mL of round-bottom flask equipped with a magnetic

stirrer and a condenser, which was then placed in a water bath. Appropriate amounts of

PHBV were added to the mixture and the mixture temperature was gently increased to 60

°C by vigorous agitation stir to dissolve PHBV for more than an hour. Afterwards,

nanocomposites laminated films with 17 and 20 g/m2 weight were prepared by casting

each solution on a silicon wafer, and then dried at room temperature in a fume hood for 3

days to remove the solvent completely. In preparation for the next step, films were

removed from the silicon wafers. According to the various nanoclay content (1, 3, 5, and

10 wt%) the different nanocomposite formulations were labeled as, PCS1, PCS3, PCS5,

and PCS10 for PHBV/CS30B nanocomposites; MPGC1, MPGC3, MPGC5, and

MPGC10 for PHBV/MPGC nanocomposites; and HPGC1, HPGC3, HPGC5, and

HPGC10 for PHBV/HPGC nanocomposites.

4.3.2 Paper Coating

Paper laminated coating (LC coating) was used to coat PHBV nanocomposites

onto paper; whereas the combination of pre-spray coated with PHBV suspensions and

laminated coating (SPLC coating) methods were used to coat both PHBV and PHBV

nanocomposites onto paper.

In the LC coating method, the pre-prepared PHBV films with weight of 13, 20,

and, 25 g/m2 were mounted on a paper surface (10*10 cm) and then placed between two

Teflon sheets and inserted into a hydraulic hot press (Carver laboratory Press model

3925; Carver, Inc., Wabash, IN, USA). It was then heated to 190 °C and subjected to a

95

pressure of 2 bars for 2 minutes, and then the coated paper was cooled down to ambient

temperature.

In the SPLC coating process, the PHBV/water slurry was suspended with a non-

ionic biosurfactant (polysorbate 80 at 1% w/w PHBV) at 95 °C under continuous

agitation for 1 hour, and cooled down to room temperature. It was then treated with a

double-cylinder type homogenizer for 5 minutes. The stabilized PHBV suspension was

sprayed using an airbrush with a nozzle of 0.3 mm in diameter connected to a compressed

nitrogen tank (pressure 150 kPa) at a distance of 15 cm onto the paper substrate mounted

on Teflon sheet. The coating weight of PHBV was controlled at approximately 3 g/m2.

The coated paper was then dried at 70 oC in fume hood for 3 days to remove the water

completely. Afterward, the pre-prepared PHBV film with weight of 10, 17, and, 22 g/m2

and PHBV nanocomposites film (17 g/m2), were laminated on the pre-spray coated with

PHBV suspension paper according to LC coating method.


Fourier transform infrared spectrum (FTIR): FTIR was used to observe the

modified CS30B with ring opening polymerization of β-Butyrolactone using a Perkin

Elmer Spectrum 100 FT-IR Spectrometer. Clays were ground into KBR powder and

pressed into disks prior to being placed in the FTIR for scanning.

X-ray diffraction (XRD): The powder diffraction patterns of samples were

collected on a Bruker AXS D8 Advance solid-state powder diffraction XRD system with

Cu Kα radiation (λ= 0.1542 nm) at 45 kV and 100 mA. The diffraction patterns were

collected between 2θ of 2° and 10° at a scanning rate of 0.02°/min. The d-spacing of the

http://en.wikipedia.org/wiki/Surfactant

96

intercalated clays and nanocomposites were determined by fitting Bragg’s equation (nλ =

2dsinθ).

Scanning electron microscope (SEM): The coated layer thickness uniformity

was analyzed with a JEOL JSM6400 digital scanning electron microscope (SEM). The

coated paper was brittle fractured in liquid nitrogen and gold coated, and the fractured

surfaces were then observed.

Transmission electron microscope (TEM): Transmission electron microscopy

(TEM) analysis was examined using a JEOL 2011 STEM with an accelerating voltage of

200 keV. Nanocomposites that were molded by hot press were first trimmed with iron

knives. Subsequently, the ultrathin sections were microtomed from these faces with a

diamond knife. The microtomed sections were collected in a water-filled boat.

Water vapor and gas transmission rates (WVTR): Water vapor permeability

of all coated paper samples were performed on IGA-003 (Hiden-Isochema, Warrington,

UK) in accordance to the methods described in TAPPI standard T464 om-12 (2012) and

ASTM E96/E96M-05 (2005) as described in detail in our previous wor.[44]. In this

experiment, coated paper samples were pre-conditioned on WVTR test measurement

conditions for three days to reach water adsorption equilibrium. The measurements were

carried out at relative humidity 50 and 90 (ΔRH %) and at temperatures of 23 °C and 38

°C, respectively. All samples were pre-conditioned on WVTR test measurement

conditions for 72 hours to reach water adsorption equilibrium.

At constant temperature and ΔRH (%), a steady state is reached where WVTR can

be calculated from the weight change of the testing chambers described by Equation

(4.1).

97

𝑊𝑉𝑇𝑅 =𝑊𝑒𝑖𝑔ℎ𝑡 𝐶ℎ𝑎𝑛𝑔𝑒

𝐴𝑟𝑒𝑎 × 𝑡𝑖𝑚𝑒 [4.1]

4.5. Results and Discussion

4.5.1. Analysis of chemical structure of surface modified clay

Figure 4-1 shows the FTIR spectra for PHBV, CS30B, and HPGC. The CS30B

spectra exhibits a very broad band centered at 3629 cm-1

corresponding to the structural –

OH groups in the clay sample. Assigned to the stretching vibration of the hydrogen-

bonds, the absorptions at 2924 cm−1

and 2852 cm−1

can be attributed to asymmetric and

symmetric stretching vibrations of C–H bonds, whereas the absorption at 1469 cm−1

can

be attributed to methylene bending vibrations. Also, an intense broad band in the range

1100–1000 cm−1

can be readily assigned to the Si–O–R due to asymmetric Si–O–C

stretching These peaks are also noted in HPGC spectra. In addition, HPGC spectrum

demonstrated the presence of functional group of PHB, asymmetrical deformation of C-H

bond in CH2 groups and CH3 groups, C=O bond stretching and C-O ester bond, which are

represented by wave numbers 1391, 1300, 1737 and 1080 cm-1

, respectively, and

aliphatic C-H stretch at 2935 cm-1

, and a O-H stretch due to free PHB hydroxyl group at

3432 cm-1

. These FT-IR spectral observations are taken as an evidence for a successful

ring opening polymerization of β-Butyrolactone.

98

Figure ‎4-1. FTIR spectra of PHBV, CS30B, and HPGC

Table ‎4.1. Modified clay with various molar ratio of catalyst to β-BL in Toluene and

bulk polymerization

Samples Molar ratio of

catalyst to β-BL

Solvent Main

peak

[d(001)]

D-spacing Yield of reaction

base on β-BL (wt%)

MPGC

1/4000 Toluene 2.48

and

4.78

3.5 and1.87 ~22

LPGC

1/4000 Bulk

polymerization

3.54

and

4.77

2.49 and

1.88

<10

HPGC

1/7000 Toluene 2.4 3.71 ~40

CS30B

- - 4.78 1.86 -

99

4.5.2. XRD pattern of PHBV Nanocomposites

The crystal structure of nanocomposites was first determined by X-ray diffraction

analysis. The interlayer distance (d-spacing) of clays inside the nanocomposites is

obtained from Bragg’s law:

2𝑑 𝑠𝑖𝑛𝜃 = 𝑛𝜆𝑤 [4.2]

Where d is the interlayer distance of clay, λw is the wavelength used, n is the

order, which is equal to 1 for the first order, and θ is the measured angle.

In Figure 4-2, the curves represent XRD spectra of CS30B and ROP modified

clay with β-butyrolactone, LPGC, MPGC, and HPGC, respectively. The d spacing in

CS30B was initially present at 2θ=4.7o and corresponds to a basal space of 1.85 nm. The

2θ values were decreased by increasing the ring opening polymerization yield of β-

Butyrolactone in CS30B gallery to 2θ =3.54o, 2θ =2.5

o, and 2θ =2.4

o which correspond to

d spacing of 2.49 nm, 3.5 nm, and 3.7 nm, respectively. As listed in Table 1 the HPGC

nanclay with highest reaction yield indicates more intercalation in clay gallery. It is also

evident (From Figure 4-2) that all modified nanoclay samples exhibited the second

peak at around 4.7° (2θ) which is related to the amount of the original unmodified

CS30B. The intensity of peak was reduced with increasing the yield of reaction (Table 1).

As the second peak almost disappeared in HPGC, which means a smaller amount of

unmodified CS30B remained in a new clay gallery. Furthermore, two more weak peaks at

2θ=7.4o and 2θ=10.4

o are related to crystal structure of PHB.[45, 46]

The XRD patterns of PHBV/CS30B, PHBV/HPGC and PHBV/ MPGC

nanocomposites are illustrated in Figures 4-3, and 4-4, respectively. The diffraction peaks

of PHBV nanocomposites containing 3, 5 and 10 wt% of CS30B clays are presented in

100

Figure 4-3. XRD pattern of the nanocomposite with 3 wt% CS30B shows diffraction

peak shoulder centered on 2θ = 3.14°, as the amount of the CS30B increased to 5 and 10

wt% diffraction peaks shifted to higher angles, 2θ = 3.78° and 2θ = 4.22° ,respectively,

and the strong plane peak corresponding to d001 was clearly observed with increased

peak intensity.

The d001-spacings of the nanocomposites with CS30B were 2.8, 2.3 and 2.1 nm

for compositions 3, 5 and 10 wt%, respectively. The indexing of the d001 peak in the XRD

pattern suggests that the morphology of the PHBV/CS30B nanocomposites is

intercalated, which indicated that the dispersion is less pronounced in PHBV matrix. No

significant peak was observed for PHBV/ HPGC nanocomposites with 1, 3, 5 and, 10

wt% of HPGC in XRD pattern (Figure 4-4.). The non-existence of peaks in the XRD

pattern suggests that the PHBV/ HPGC1 and PHBV/ HPGC3 nanocomposites have an

exfoliated morphology. Further increase in the nanoclay content, (in PHBV/HPGC

nanocomposites) resulted in the appearance and shifting of a diffraction peak to higher

angles. Furthermore PHBV/MPGC nanocomposites XRD pattern illustrate less clay-

biopolymer interactions with weaker dispersion of clay compare to PHBV/HPGC

nanocomposites

A general conclusion from the XRD patterns is that the nanocomposites with

HPGC and MPGC (Figure 4-4) are rather exfoliated/Partial exfoliated while those with

CS30B could be considered as intercalated

101

Figure ‎4-2. X-ray diffraction patterns of Clays

102

Figure ‎4-3. X-ray diffraction patterns of a) PHBV b) PCS1 c) PCS3 d) PCS5 and e)

PCS10

103

Figure ‎4-4. X-ray diffraction patterns of a) HPGC1, b) HPGC3, c) HPGC5, d) HPGC10,

e) MPGC1 f) MPGC5, and g) MPGC10

4.5.3. Morphologies of PHBV nanocomposites revealed by TEM

Additional support for interpreting the phase behavior of nano-clays into the

PHBV matrix is based on the TEM observations. The TEM images were taken from a

representative region of the nanocomposite films. The nanostructures shown in Figures 4-

5 correspond to PCS, MPGC and HPGC. The PCS5 nanocomposite presents an

intercalated structure with a random dispersion of small tactoids (Figure 4-5a). However,

at the higher nanoclay content (10 wt% unmodified nanoclay), an aggregated structure is

obtained. The large particle aggregates observed by TEM (Figure 4-5b) indicated that the

CS30B particles were neither intercalated nor dispersed uniformly in the polymer matrix.

104

This aggregation may impact the water vapor transmission rate because some aggregates

may create preferential diffusion pathways.[47, 48]

Figures 4-5c and 4-5d illustrate the morphology pattern of MPGC5 and MPGC10

nanocomposites, respectively. It seems the nanocomposites with MPGC (with moderate

reaction yield) shows partially exfoliated structure and better dispersion of clay than

CS30B (unmodified clay) in PHBV matrix. Moreover, at 5 wt% of HPGC (the highest

reaction yield) nano-clay content, fully exfoliated structure with a random dispersion

reflecting to excellent homogeneity in the nanoclay dispersion is observed in Figure 4-5e.

In terms of barrier properties, the fully exfoliated structure creates tortuous and

complicated diffusion pathways for vapor molecules moving through the biopolymer

matrix, which can complicate the diffusion pathways taken by the water vapor molecules.

Furthermore, at the higher HPGC nanoclay content (10 wt%), the exfoliated nanoclay and

random structure were still present in PHBV matrix (Figure 4-5f).

105

Figure ‎4-5 TEM images of PHBV nanocomposites a) PCS5, b) PCS10, c) MPGC5, d)

MPGC10, e) HPGC5, and f) HPGC10

4.5.4. The water vapor transmission rate (WVTR) measurement

The WVTR measurement was aimed to investigate the influence of the both (LC

and SPLC) coating methods on the WVTR of PHBV coated paper at various coating

weights, thus identifying the optimum coats weight and the affective coating method.

Also the influence of exfoliated nanocomposites on the WVTR of coated paper was

investigated with the SPLC coating method.

Evaluated WVTR is defined as the steady water vapor flow over time through the

area of a body and under 50 and 90 (RH %) and at temperatures of 23 °C and 38 °C,

respectively. Water vapor transmission rate is a measure of the passage of water vapor

106

through a substance. Moisture control is critical in many moisture sensitive industries

such as, food and pharmaceuticals, as their products are put in packaging with controlled

WVTR to achieve the required quality, safety, and shelf life.

Figure 4-6 reveals the variation in WVTR of PHBV coated paper with both

coating methods at various coating weights. The results indicated that the SPLC coating

method led to a high water vapor barrier resistance, regardless of coating weight.

Especially at the lowest coating weight (13 g/m2) SPLC coating system showed

significantly better performance than the LC coating system. Obviously, the coating with

an increased thickness improved water vapor barrier. The WVTR values of the

investigated samples were also found to be strongly dependent on the temperature and the

relative humidity difference. SEM images of PHVB coated (25 g/m2 coat weight) paper

cross section illustrates that SPLC coated layer (Figure 4-7b) is more uniform and smooth

with lower penetration of biopolymer into the paper substrate than LC coated layer into

the Calendared paper substrate (Figure 4-7b).

Figures 4-8a and 4-8b show the water vapor barrier properties of

PHBV/nanocalys coated paper with SPLC coating method. PHBV/HPGC

nanocomposites coated papers had lower WVTR than PHBV/MPGC and PHBV/CS30B

nanocomposites coated papers at all clays content.

As listed in Table 2, the WVTR of PHBV/HPGC nancomposites reduced sharply

with just 1% HPGC nano-clay from 67 g/m2/day to 39.6 g/m

2/day (at 90 RH % and 38

°C) and 16.7 g/m2/day to 11.1 g/m

2/day (at 50 RH % and 23 °C) while by further

increasing the coated nanocomposites HPGC clay content up to 10 %, the WVTR values

were reduced gradually from 39.6 g/ m2/day (at 90 RH % and 38 °C) and 11.1 g/ m

2/day

107

(at 50 RH % and 23 °C) to 26.4 g/m2/day and 6.4 g/m

2/day, respectively, due to the

increase in the effective path length for diffusion as revealed by the XRD and TEM

results. Such low WVTR values enable the coated paper as green-based packaging

materials with extremely high barrier toward water vapor transfer (the control sample

without coating has the WVTR as high as 3120 g/m2/day (at 90 RH % and 38 °C) and

630 g/m2/day (at 50 RH % and 23 °C)).

Also water vapor permeability of PHBV/MPGC nanocomposites gradually

reduced from 67 (g/ m2/day) for neat PHBV to 41.5 (g/ m

2/day) (at 90 RH % and 38 °C)

up to 5 wt% MPGC loading and by further increasing the nano-clay content to 10% water

vapor permeability raised up to 54 (g/m2/day). Meanwhile, water vapor barrier of

PHBV/CS30B nanocomposites slightly improved up to 3% loaded CS30B. The higher

clay content was sharply diminished the water vapor resistance of coated paper even less

than the neat PHBV coated one.

This phenomenon, high WVTR values with increasing the clay content, is

probably because of aggregation of silicate layers. The aggregated structure of nano-

clays not only provides channels or micro-voids at the interface of polymer and clay, but

also may create direct pathways for water vapor through coated substrate in some cases

like PCS5 and PCS10. As a result, the nanocomposite coated paper could have even

worse water vapor barrier properties than the bare PHBV coated one.[32, 35]

However, the sharply reduction in WVTR associated with HPGC1 coated sample might

be attributed to fully exfoliated structure in this sample. The lower WVTR of

PHBV/HPGC samples in all modified nano-clay content in comparison with

PHBV/CS30B and PHBV/MPGC nanocomposite coated paper is related to exfoliated

108

structure of PHBV/HPGC nanocomposites, which is in agreement with the XRD and

TEM results.

Figure ‎4-6. The WVTR of PHBV coated paper with LC and SPLC coating methods

109

Figure ‎4-7. SEM cross section of coated paper with a) LC coating and b) SPLC coating

method

110

Figure 4-8. The WVTR values of PHBV nanocomposites coated paper with SPLC

coating methods a) at 23 °C and 50 RH% and b) 38 °C and 90 RH%

111

Table 4.2. WVTR of samples in details

Samples Coat weight (g/m2)

WVTR (g/m2/day)

at 23 °C and 50 RH%

WVTR (g/m2/day)

at 38 °C and 90 RH%

Coating method

Calendared paper

75 (g/m2)

630±15 3120±17 -

PHBV 13±2 33.2±2.3 149.7±12.1 LC

PHBV 20±2 20.4±1.9 95.4±7.6 LC

PHBV 25±2 15.9±0.7 70.8±3.5 LC

PHBV 13±2 22.1±1.2 84.5±3.6 SPLC

PHBV 20±2 16.7±0.3 67±1.6 SPLC

PHBV 25±2 13.6±0.1 56.9±1.2 SPLC

PCS1 20±2 14.2±0.4 59.8±1.4 SPLC

PCS3 20±2 13.3±0.3 56.3±1.1 SPLC

PCS5 20±2 14.6±0.5 61.6±2.7 SPLC

PCS10 20±2 18±1.2 76.4±3.5 SPLC

MPGC1 20±2 12.6±0.3 49.3±1.7 SPLC

MPGC3 20±2 11.4±0.1 44.8±1.4 SPLC

MPGC5 20±2 10.2±0.3 41.5±2.4 SPLC

MPGC10 20±2 13.2±0.6 54.0±3.2 SPLC

HPGC1 20±2 11.1±0.1 39.6±1.6 SPLC

HPGC3 20±2 9.3±0.1 34.8±1.4 SPLC

HPGC5 20±2 7.3±0.2 31.2±1.0 SPLC

HPGC10 20±2 6.1±0.4 26.4±1.6 SPLC

4.6. Conclusion

The water vapor barrier of the coated papers was improved significantly, if the

surface of the base substrate was pre-spray coated with PHBV suspension prior to the

coating of PHBV film via hot press. On the other hand, the coating formulation had more

profound effect on water vapor transmission rate of the coated paper than coating weight.

The dispersion of clay in PHBV matrix was verified by both XRD and TEM

characterization. The results indicated that the ring opening polymerization of β-

Butyrolactone into the layered structure of clay is an effective approach to enhance the

112

compatibility of clay with biodegradable polymer matrix. The superior water vapor

resistance obtained from PHBV/HPGC nanocomposites coated paper is attributed to the

fully exfoliated nano-clay structure. The nanocomposite coated paper reached the WVTR

value as low as 26.4 g/m2/day, thus leading to a very promising green-based packaging

material with high water vapor barrier.

Acknowledgement

Authors are grateful to NSERC Strategic Network – Innovative Green Wood Fibre

Product (Canada) and NSF China (No. 51379077) for funding.

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Chem., 59 (2011) 12384-12395.

[48] X. Tang, S. Alavi, Structure and physical properties of starch/poly vinyl

alcohol/laponite RD nanocomposite films, J. Agric. Food. Chem., 60 (2012) 1954-1962.

119

Chapter 5 Surface morphological analysis and synergetic water

vapor barrier properties of modified clay/Poly(3-

hydroxybutyrate-co-3-hydroxyvalerate) composites

Abstract

In this work, biodegradable bacterial copolyester, Poly(3-hydroxybutyrate-co-3-

hydroxyvalerate) (PHBV) based two different types of bio-nanocomposites were

prepared via solution intercalation process using two types of organo-modified clays (0 to

40 wt%). The effect of Cloisite®30B (CS30B) and an intercalated clay, which was

synthesized via ring opening polymerization of Polyhydroxybutyrate onto Cloisite®30B

(PHB-g-CS30B), on the wettability of the assemblies has been investigated in details.

Moreover, Phase morphology of the nanocomposite films was investigated by XRD and

atomic force microscopy (AFM). The moisture adsorption of PHBV and its

nanocomposite films was determined using Belsorp-Max. The nanocopmosite films

derived from PHB-g-CS30B exhibited higher wettability than that of the samples derived

from CS30B. An explanation, based on the surfaces topography, has been proposed.

Phase morphology, filler dispersion and RMS roughness were also analyzed; and the

results revealed the impact of aggregates on WVTR and contact angle of the composites.

5.1. Introduction:

Polymer surfaces and thin films play an important role in a wide range of

applications, for example as coatings, membranes, sensors and in medical purposes [1-6].

The properties of the polymer on a nanoscopic scale can be drastically different from

120

those observed on the macroscopic scale [7, 8]. This is due to the fact that both local and

global dynamics of the polymer, within the host galleries of nanoparticles, are strongly

influenced by the chain confinement of the polymer and polymer-surface interaction that

are not observed on the bulk scale. On a local scale, intercalated nanocomposites exhibit

simultaneously a fast and a slow mode of relaxation whereas on the global scale

relaxation of polymer chains are in close proximity to the host surface. In many

applications there is a need to know the surface characteristics and surface properties of

Polymer/nano-clay composites [1, 9].

Poly (3-hydroxybutyrate-co-3-hydroxyvalerate (PHBV) is a bio-thermoplastic

aliphatic polyester synthesized naturally by bacteria. It is a copolymer composed of

polyhydroxybutyrate (PHB) and polyhydroxyvalerate (PHV). PHBV has widespread

industrial applications in medicine packaging [10], bone tissue engineering [11-13],

construction, automotive industry, electronics, food packaging [14], and cellulose fiber

network coating [15, 16]. Moreover, PHBV exhibits properties that are comparable to

conventional fossil-base plastics especially polypropylene [17] and has hydrolysable

backbones that can easily biodegrade upon microbial attack. However, commercialization

of PHBV and most other bio-thermoplastics face certain challenges such as high

production cost, narrow processing temperatures, brittleness, and especially low water

vapor resistance [18, 19]. Several researchers believe that bio-thermoplastics with

improved properties can provide an effective role in the development of environmentally

friendly, disposable consumer packaging and cellulose fiber network coatings [6, 20].

The use of layered silicate nano-clays, as an additive in polymer industry, has

received great attention for decades not only because of its low cost but also because it

https://en.wikipedia.org/wiki/Bacteria

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enhances both the surface and bulk properties such as, wettability, adsorption, desorption

[21, 22], surface roughness, mechanical [23, 24], processability [25], electrical, thermal

stability [26, 27], and barrier [28] properties of a polymer. The PHBV reinforced with

montmorillonite clay showed better thermal, and mechanical properties because of the

nanoscale dispersion of the filler in the polymeric matrix[29-33]. Previous results also

confirmed that the addition of nano-clay significantly improved oxygen and water vapor

permeability of PHBV films while keeping their transparency [34].

It is reported that in an exfoliated nanocomposite the clay platelets are arranged in

a non-parallel manner and the random dispersion of nano-clays create a tortuous path for

the transmission of small molecules such as oxygen and water vapor through the

biopolymer matrix. This tortuous path is to be responsible for the enhancement of the

water vapor resistance of the nanocomposites. However, there remain the compatibility

issues with the mixing of hydrophilic [35], partially hydrophobic and organophilic[36,

37] nano-clays with a hydrophobic polymer such as, PHBV. As a result, one of the

significant challenges remained in PHBV/silica layer nanocomposite development is to

induce the hydrophobicity to nano-clay. Proper surface modification of nano-clay can

enhance the interfacial adhesion between the polymer matrix and the nano-clay, which in

turns help to reach complete dispersion and exfoliation of the nano-clay in the matrix[38-

42].

To the best of our knowledge, there is no previous report on investigating the

surface morphology of PHBV nanocomposites films. Furthermore, the present work was

focused on the detailed study of PHBV nanocomposites films prepared with a novel

organic-modified Cloisite®30B clay prepared through in situ polymerization of

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Polyhydroxybutyrate (PHB). Special attention has been paid to the exfoliated/intercalated

structure and surface characterizations of the nanocomposites, The bulk morphology of

the PHBV-nanocomposites was investigated by XRD, AFM. In addition, the impacts of

the level of exfoliation of the clay platelets and clay content on the water vapor

permeability, contact angel and water vapor adsorption in PHBV film were revealed.

5.2. Materials



mol−1


and ENMAT Y1000P content of 2% hydroxyvalerate. The β-Butyrolactone (98 %) was

purchased from Sigma Aldrich and dried with Calcium hydride for 2 days by stirring, and

distilled under reduced pressure prior to use. The 1-Ethoxy-3-

chlorotetrabutyldistannoxane catalyst (Scheme 5-1) was synthesized according to Hori Y.

et. al.[43], PHB-g-montmrollonite was synthesized in our lab (Scheme 5-2). The

Organo-modified montmorillonite, Cloisite® 30B, was kindly donated by Southern Clay

Products (Gonzalez, TX). All powder materials were dried before use under vacuum

overnight at 50 °C.

Scheme ‎5-1. Structures of the 1-Ethoxy-3-chlorotetrabutyldistannoxane

https://www.google.ca/url?sa=t&rct=j&q=&esrc=s&source=web&cd=2&cad=rja&uact=8&ved=0CCMQFjAB&url=http%3A%2F%2Fwww.linguee.fr%2Fanglais-francais%2Ftraduction%2Fa%2Bspecial%2Battention%2Bhas%2Bbeen%2Bpaid.html&ei=8eYaVbSYJ8OdyAS42oKYCw&usg=AFQjCNH6dRK8TIe2bf9qe7d1eUZlhLPUVg&sig2=MG6e2CfInICUMn-ZfKlLAA&bvm=bv.89381419,d.aWw

123

Scheme ‎5-2. Structure of PHB-g-CS30B

5.3.‎In‎situ‎polymerization‎of‎β-butyrolactone on Cloisite®30B

To obtain Pre-intercalated Clay, Cloisite®30B (4 g) was homogenized by

vigorous agitation stir bar overnight at room temperature in 100 ml dry toluene. The β-

butyrolactone was dissolved in toluene and added in 2:1 mass, to the Cloisite®30B

suspension (4 g/100ml), afterward, the desired amount of 1-Ethoxy-3-

chlorotetrabutyldistannoxane (as a catalyst) was added to the mixture [43] and slowly

heated up to 110 °C and stirred overnight at 110 °C temperature to carry out the

polymerization (Scheme 5-2 and Figure 5-1). Then, the suspension was filtered under

reduced pressure then washed three times with chloroform and centrifuge subsequently at

10000 rpm the organomodified clay (PHB-g-CS30) was dried under vacuum at 40 °C.

5.4. Bio-nanocomposites films preparation

The solvent-casted composites were prepared by dissolving PHBV powder in

chloroform (5%, w/v) at 50 °C. The overnight pre-suspended nano-clay was added into

the PHVB solution after the PHBV powder was dissolved completely, and the mixture

124

was vigorously stirred for 5 h. For reference, control PHBV films (no nano-clay) were

made by dissolving PHBV powder in chloroform and those solutions were casted onto

silicon wafer. The mixture was left in a fume hood overnight to evaporate the solvent.

The films were peeled off easily from the glass and stored in a vacuum desiccator until

use. Similarly, the composite films with 0, 1, 3, 5, 10, 20, 30 and 40% (w/w) of nano-

clays in PHBV matrix were also prepared.


Fourier transform infrared spectrum (FTIR)

FTIR was used to observe the CS30B, modified CS30B and PHBV using a Perkin

Elmer Spectrum 100 FT-IR Spectrometer. Samples were ground into KBR powder and

pressed into disks prior to being placed in the FTIR for scanning.

5.1.Characteristics of the bio-nanocomposite films surface

For the measurement of static water contact angle each dried sample was at first

mounted on a glass. Then, 3 µl of distilled water was dropped on each sample. Water

contact angle of samples was measured by using Theta Optical Tensiometer (Finland)

and Attension Theta software. Ten measurements were made on each sample. Atomic

Force Microscopic (AFM, Nanoscope IIIa, Veeco Instruments Inc., Santa Barbara, CA)

analysis was carried out to determine the surface morphology of PHBV and PHBV-

nanocomposites. The images were scanned in Multimode mode in air using a commercial

silicon tapping probe (NP-S20, Veeco Instruments) with a spring constant of 0.32 N/m, a

resonance frequency of 273 kHz and settings of 512 pixels/line and Hz scan rate. The

average roughness (Ra), and root mean square roughness (RMS), which indirectly

125

indicate the extent of clay dispersion, were measured from the AFM images by the

arithmetic average of absolute values of the surface height deviations.

Characteristics of the bio-nanocomposite films bulk morphology

Diffraction X-ray patterns (XRD) of pristine PHBV and nanocomposites

measurements were performed on a Bruker D8 Advance X-ray diffractometer using Cu-

Kα radiation (k = 1.5418 Å), in which the X-ray tube was operated at 40 kV and 30 mA

at room temperature. All the experiments were carried out with the wide angle varying

from 1 to 30°.

Moisture adsorption

Moisture sorption isotherms of the PHBV and its nanocomposites films was

determined using Belsorp-Max (BEL JAPAN, INC., Osaka, Japan) at 25 °C over 10 to

100 RH% the relative humidity range automatic moisture adsorption apparatus with the

previously reported method [44]. The samples were dried for 72 h at 60 °C under vacuum

before the isotherm moisture adsorption measurement. Generally the average value of the

equilibrium moisture content (EMC), which is per mg of adsorbed moisture to per g of

sample, was calculated according to Equation (5.1), which also refers to moisture content

of the product in each RH% condition.

𝐸𝑀𝐶 =𝑊𝑤𝑒𝑡 − 𝑊𝑑𝑟𝑦

𝑊𝑑𝑟𝑦∗ 1000 [5.1]

Where EMC is in percent dry basis; 𝑊𝑤𝑒𝑡 is sample equilibrium weight at a certain RH%

in and 𝑊𝑑𝑟𝑦 is sample initial dry weight.

Water vapor transmission rate (WVTR)

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Water vapor transmission rate (WVTR, g/m2/day) was measured by a microbalance

of gravimetric apparatus (IGA-003, Hiden Analytical Ltd., UK) as reported in our

previous works [45, 46] .All samples (thickness 30±1 μm) were pre-conditioned on

WVTR test measurement conditions, for 72 hours to reach water adsorption equilibrium.

The average thickness was recorded using a micrometer caliper prior to the test. The

WVTR determination was according to TAPPI T464 om-12 at 37.8 °C and 90 RH% &

TAPPI T 448om-09 at 23 °C and 50 RH%. The WVTR values were calculated from the

slope of weight change versus time. The experiments were done in five times for each

sample. The WVTR values were calculated from the slope of the weight change against

time using the exposed surface area (120 mm2) of the sample, and the equation shown in

(5.2):

𝑊𝑉𝑇𝑅 = 𝑊𝑒𝑖𝑔ℎ𝑡 𝑐ℎ𝑎𝑛𝑔𝑒

𝐴𝑟𝑒𝑎×𝑡𝑖𝑚𝑒 (

𝑔(𝑚2 ∗ 𝑑𝑎𝑦)⁄ ) [5.2]


5.6.1. Characterization of the ring opening polymerization on CS30B

FTIR (Figure 5-1) and XRD (Figure 5-4) test was used to confirm the ring

opening polymerization of β-Butyrolactone in the presence of CS30B. Figure 5-1 shows

the FTIR spectra for PHB, CS30B, and PHB-g-CS30B. The PHB-g-CS30B spectra

exhibits a peak centered at 3432 cm-1

corresponding to O-H of stretch of free hydroxyl

groups in the modified clay. Furthermore, PHB-g-CS30B spectrum demonstrated the

presence of functional groups of PHB. The absorptions at 1391 cm−1

and 1300 cm−1

can

be attributed to asymmetrical deformation stretching of C-H bond in CH2 groups and CH3

127

groups respectively, whereas the absorption at 1737 and 1080 cm-1

can be attributed to

C=O bond stretching and C-O ester bond, respectively, and the peak at 2935 cm-1

assigned to the aliphatic C-H stretch.

Figure ‎5-1. FTIR spectra of PHB, PHB-g-CS30B, and CS30B

5.6.2. The surface characterization of the nanocomposites

The surface characterization of the nanocomposites was undertaken using AFM, which

was also used to study the topology of the two types of nanocomposite films surface at

different nano-clay content. Figure 5-2 and Figure 5-3 depicts the 2-dimensional and 3-

dimesional AFM images of PHBV/CS30B and PHBV/PHB-g-CS30B nanocomposite

films, respectively, in a scan range of 5 × 5 μm2. The average roughness (Ra), root mean

square roughness (RMS) and maximum height (Rmax) values are given Table 1. The

128

images recorded for the native PHBV (Figure 5-2a) indicate that the polymeric film has

smooth surface with the maximum height of 114 nm. In contrast, the microscopic images

of PHBV/CS30B nanocomposites (Figure 5-2b to 5-2h) indicate presence of some

disarrangement on the biopolymer film surface, demonstrating that the incorporation of

inorganic clay strongly influenced the films surface morphology. The AFM analysis also

showed that the surface of the PHBV/CS30B nanocomposites became rougher (with

respect to PHBV) when CS30B was incorporated, and the surface roughness values

increased consistently with the increase in nano-clay content of (Table 1) up to 10 wt%.

Thereafter, the values increased sharply with the addition of more nano-clay. This

phenomenon may be attributed to with the weak interactions between compounds and

poor homogeneity of the mixture components, in another word formation of

agglomerated nano-clay structure in PHBV matrix. As seen from Figure 5-3(b) to Figure

5-3(h), the micrograph of the PHBV/CS30B is comprised of hills and cavities. According

to literature [47], the lighter color area is ascribed to the hard segments of agglomerated

clay sheets, whereas the darker color area represent the polymer matrix. It seems that the

increase in nano-clay content results in an increase in the lighter color spots in

micrograph, which can be attributed to agglomerated morphology of CS30B at high clay

loading.

129

Figure ‎5-2. AFM phase images of a) PHBV, b) PHBV/CS30B 99/1 w/wt%, c)

PHBV/CS30B 97/3 w/wt% ,d) PHBV/CS30B 95/5 w/wt% ,e) PHBV/CS30B 90/10

w/wt% ,f) PHBV/CS30B 80/20 w/wt% ,g) PHBV/CS30B 70/30 w/wt%, h)PHBV/CS30B

60/40 w/wt%, i) PHBV/PHG-g-CS30B 99/1 w/wt%, j) PHBV/ PHG-g-CS30B 97/3

w/wt% ,k) PHBV/PHG-g-CS30B 95/5 w/wt% ,l) PHBV/PHG-g-CS30B 90/10 w/wt%

,m) PHBV/PHG-g-CS30B 80/20 w/wt% ,n) PHBV/PHG-g-CS30B 70/30 w/wt%, and

o)PHBV/PHG-g-CS30B 60/40 w/wt%.

Results from PHBV nanocomposites containing different modified clays also

revealed a significantly different distribution and contribution of clay in the polymer

matrix. The surface topography of PHBV/PHB-g-CS30B nanocomposites is different

from that of PHBV/CS30B nanocomposites based on AFM images (2i to 2o). The

average height and RMS of PHBV/PHB-g-CS30B surface were less than that of

130

PHBV/CS30B with smoother surface (Table 1). For example, the average surface

roughness (Ra) of the PHBV/CS30B and PHBV/PHB-g-CS30B composites at the same

clay content (10 wt%) were 95.1 and 85.7, respectively. The surface roughness of

PHBV/PHB-g-CS30B nanocomposites (up to 20 wt% clay content) were less or almost

same as the native PHBV, which indicates that the clay was dispersed uniformly and

randomly in PHBV matrix. As Illustrated in Figure 5-3(i) to 5-3(o), the dispersion of

lighter color spots related to clay seems homogeneous. This homogeneous dispersion of

lighter color spots (even with high content of modified clay) can be attributed to

intercalated or exfoliated morphology of PHBV/PHB-g-CS30B nanocomposites resulting

from stronger interaction of modified clay with and PHBV matrix compared to CS30B.

5.6.3. Contact Angle and surface roughness correlation in nanocomposites.

The static contact angle was measured on PHBV and the nanocomposites films to

study about the surface wettability. It is believed that the contact angle of the surface

depends on two main factors hydrophobicity and roughness or real area of contact[48].

Figure 5-3 and Table 5-1 show the water droplet contact angle of the neat PHBV,

PHBV/CS30B and PHBV/PHB-g-CS30B composite films surfaces that faced the silicon

wafer during the film preparation. The contact angel value of PHBV/PHB-g-CS30B

composites was lower than that of PHBV/CS30B composites regardless of the clay

content. For example, the water contact angle of the PHBV/CS30B and PHBV/PHB-g-

CS30B composites at the same clay content (10 wt%) were 89.9 and 78.2, respectively.

The water contact angle of the PHBV/PHB-g-CS30B composites remained almost

constant (up to 10 wt% clay). With the addition of more PHB-g-CS30B clay the water

131

contact angel increased gradually from 78.2 to 86.3 for the increase in clay content from

10 to 40%. However, an increment (from 86.3 to 101.8) can be also seen in the water

contact angle of PHBV/CS30B samples with increasing the clay content from 10 to 40

wt%.

Lower water contact angle of PHBV/PHB-g-CS30B (compared to CS30B)

samples despite being more hydrophobic and having less polarity, might be attributed to

topology of the composite films surface as seen in Figure 5-3 and table 1. As shown in

Table 5-1, the contact angle values of the nanocomposites (both PHB-g-CS30B and

CS30B) seem to correlate to the roughness values of the nanocomposites. The contact

angle increased gradually with the increase in roughness up to 20 wt% and 10 wt% clay

content for PHB-g-CS30B and CS30B, respectively. However, at high clay content,

while the roughness of the films increased sharply, the contact angel demonstrated only a

moderate increase, especially for PHBV/CS30B samples. As a result, the high contact

angle of the PHBV/CS30B composites might be attributed to the high surface roughness

found in this study. PHB-g-CS30B composites have smoother surface and less roughness

compared to PHBV/CS30B samples, hence lower contact angle despite being more

hydrophobic[48].

132

Figure ‎5-3. AFM 3D microroughness and water contact angel of a) PHBV, b)

PHBV/CS30B 99/1 w/wt%, c) PHBV/CS30B 97/3 w/wt% ,d) PHBV/CS30B 95/5 w/wt%

,e) PHBV/CS30B 90/10 w/wt% ,f) PHBV/CS30B 80/20 w/wt% ,g) PHBV/CS30B 70/30

w/wt%, h)PHBV/CS30B 60/40 w/wt%, i) PHBV/PHG-g-CS30B 99/1 w/wt%, j) PHBV/

PHG-g-CS30B 97/3 w/wt% ,k) PHBV/PHG-g-CS30B 95/5 w/wt% ,l) PHBV/PHG-g-

CS30B 90/10 w/wt% ,m) PHBV/PHG-g-CS30B 80/20 w/wt% ,n) PHBV/PHG-g-CS30B

70/30 w/wt%, and o)PHBV/PHG-g-CS30B 60/40 w/wt%.

133

Table ‎5.1. Surface roughness and water contact angle data of various PHBV and its

nanocomposites

Rrms (root mean squared)

(nm)

Ra (average surface

roughness) (nm)

Contact

angle

PHBV 112.3 87.0 77.2± PHBV/CS30B 99/1 w/wt% 113.4 90.5 77.0± PHBV/CS30B 97/3 w/wt% 114.3 91.6 84.5± PHBV/CS30B 95/5 w/wt% 116.4 91.1 86.6± PHBV/CS30B 90/10 w/wt% 118.0 95.1 89.9± PHBV/CS30B 80/20 w/wt% 173.5 138.8 95.8± PHBV/CS30B 70/30 w/wt% 231.0 188.7 101.8± PHBV/CS30B 60/40 w/wt% 224.0 183.7 99.0± PHBV/PHG-g-CS30B 99/1 w/wt% 112.3 94.5 78.4± PHBV/ PHG-g-CS30B 97/3 w/wt% 88.13 69.4 76.0± PHBV/PHG-g-CS30B 95/5 w/wt% 87.4 69.5 75.8± PHBV/PHG-g-CS30B 90/10 w/wt% 105.5 85.7 78.2± PHBV/PHG-g-CS30B 80/20 w/wt% 114.3 91.5 81.9± PHBV/PHG-g-CS30B 70/30 w/wt% 145.3 114.8 83.2± PHBV/PHG-g-CS30B 60/40 w/wt% 180.0 144.8 86.3±

5.6.4. Morphologies of PHBV nanocomposites revealed by XRD and TEM

Figure 5-4 represents the XRD patterns of CS30B and PHB-g-CS30B (modified

clay). The CS30B and PHB-g-CS30B exhibited characteristic peaks (Figure 5-4) of clay

at 2θ=4.7° and 2.4°, corresponding to the gallery height (d001 spacing) of 1.85 and 3.7

nm respectively. The diffraction peaks of the clay in PHBV/PHB-g-CS30B

nanocomposites (up to 10 wt% clay contents) nearly vanished. Diffraction peaks in the

lower 2θ region respond to higher d-spacing in clay gallery which suggests an

exfoliation/intercalated morphology of nanocomposites. However in PHBV/CS30B

134

nanocomposites CS30B had the fewer shifts to lower 2θ angles of the diffraction

maximum, which suggests a partial intercalation or aggregation of the clay.

TEM images were also obtained in order to confirm and support the XRD pattern.

The TEM image observations (Figure 5-5) indicated that the PHBV/PHB-g-CS30B

composites showed better dispersion compared to PHBV/CS30B. This can be attributed

to the presence of the PHB structure on CS30B surface, which lead to a well-dispersed

systems, due to the semi-identical structure of PHB with PHBV. The TEM image of

PHBV nanocomposites using PHB-g-CS30B at 10 wt% clay content exhibited (Figure 5-

5a) an exfoliated structure for the most part with some clay stacks could also exist, which

might correspond to the slight shoulder in XRD. However, the TEM image of PHBV

with 10 wt% CS30B (Figure 5-5b) showed a partial intercalated structure, and the

aggregated clay tactoids were also observed. This result implied that the CS30B particles

were not dispersed uniformly in the polymer matrix.

135

Figure ‎5-4. X-ray diffraction patterns of a) PHBV, b) PHBV/PHG-g-CS30B 99/1 w/wt%,

c) PHBV/PHG-g-CS30B 97/03 w/wt% ,d) PHBV/PHG-g-CS30B 80/20 w/wt% ,e)

PHBV/PHG-g-CS30B w/wt% ,f) PHBV/CS30B 97/03 w/wt% , and g) PHBV/CS30B

90/10 w/wt%

136

5.6.5. Moisture adsorption

Figure 5-6 illustrates the isothermal moisture adsorption of CS30B and PHB-g-

CS30B at 25 °C in range of 10 to 100 % RH. Regardless of the RH, PHB-g-CS30B

adsorbed less water vapor compared to CS30B. Much lower affinity of PHB-g-CS30B

clay to water molecules, suggesting that the grafted PHB significantly increased the

hydrophobicity of the clay, in contrast to characteristic of a hydrophilic CS30B.

Figure 5-7(a) and 5-7(b) show the moisture absorption study of pure PHBV and

its nanocomposites (unmodified and modified clay) in range of 10 to 100 % RH. It is

believed that the sorption values in semi-crystalline material, especially polymers are

controlled by the amorphous regions [49]. However it seems that morphology and

microstructure of nanocomposites also play a critical role in determining the adsorption

and transport phenomena[50]. As seen in Figure 5-7a and 5-7b at all clay content

Figure ‎5-5. TEM images of PHBV nanocomposites a) PHBV/PHG-g-CS30B 90/10and b)

PHBV/CS30B 90/100 w/wt%

137

partially and fully exfoliated nanocomposites (PHB-g-CS30B/PHBV system) show much

lower water vapor adsorption values, even at low clay content, compared to the

intercalated (CS30B/PHBV system) nanocomposites. The Moisture sorption increases

with increase in the CS30B content, particularly for the agglomerated nanocomposites

containing more than 10 wt% CS30B . Generally, the aggregation of nanoclay facilitates

the entry of water molecules inside the polymer nanocomposite films. At low CS30B

content, the water vapor value of the nanocomposite is less than pure PHBV in spite of

the theoretically high surface area, cation exchange capacity, and strong adsorption

capacity. Moreover, the addition of hydrophilic clay in a polymer system should lead to

an increase in hydrophilicity, leading to an increment in water vapor adsorption. This

abnormal behavior could be explained through the homogenous morphology of

nanocomposites at low clay content, in which the hydrophilic silica layers are shielded or

emerged and less available to the sorption of the water vapor molecules. On the other

hand the exfoliated structure of PHBV/PHB-g-CS30B nanocomposites (even up to 20

wt% clay) and hydrophobic structure of modified clay could explain the lower sorption of

nanocomposites. Furthermore, the synergetic reduction in water vapor absorption of

exfoliated structure of PHBV/PHB-g-CS30B (at less than 10wt% clay content) compared

to neat PHBV could be attributed to strong integration of clay and biopolymer which

restricted the access of water vapor molecules with free -OH group on nanocomposite

films. This intramolecular interaction does not allow the establishment of tremendous

hydrogen bonding with water.

138

Figure ‎5-6. Water vapor sorption isotherms of clays

Figure ‎5-7. Water vapor sorption isotherms of a) PHBV/CS30B b) PHBV/PHB-g-CS30B

139

5.6.6. The water vapor transmission rate (WVTR) measurement

The WVTR measurement was aimed to investigate the influence of clay type and

content on the WVTR of nanocomposites. WVTR is a measure of the passage of water

vapor through a substance. WVTR study was carried out at 50 and 90 (ΔRH %) and at 23

and 38 °C temperature, respectively. Controlling the water vapor permeability is critical

in many moisture sensitive industries such as food and pharmaceuticals, as their products

are put in packaging with controlled WVTR to achieve the required quality, safety, and

shelf life[51].

Figure 5-8 illustrates the WVTR values of PHBV film and PHVB nanocomposite

films with thickness of 30±2 µm. It is clearly observed that nano-clays had a great impact

on the reduction of WVTR values of PHBV films. The reduction of WVTR of PHBV

films can be explained by the creation of a tortuous path for the water molecule diffusion

in the PHBV matrix. Nanocomposite films containing 10 wt% PHB-g-CS30B exhibited

the lowest WVTR values and regardless of the PHB-g-CS30B content the WVTR values

were much lower that neat PHBV and PHBV/CS30B nanocomposites films.

As can be seen from Figure 5-8b, the WVTR of PHBV/PHB-g-CS30B

nanocomposite films reduced gradually from 50 g/ m2/day (at 90 ΔRH % and 38 °C) and

12.1 g/ m2/day (at 50 ΔRH % and 23 °C) to 18.4 g/m

2/day and 4 g/m

2/day, respectively,

with the increase in clay content up to 10 wt%. While further increasing the PHB-g-

CS30B content up to 40 w% results in a sharp increase of the WVTR values.

Interestingly, the WVTR value of the nanocomposites, even at 40 w% clay content, was

less than that of neat PHBV.

140

Figure ‎5-8. The WVTR values of PHBV nanocomposites film a) at 23 °C and 50 RH%

and b) at 38 °C and 90 RH%

This significant reduction in WVTR associated with PHBV/PHB-g-CS30B

nanocomposite samples might be attributed to the formation of fully exfoliated

morphology at less than 10 wt% clay content. Nanocomposites with higher clay content

exhibit a combination of interrelated and agglomerated structure, which causes an

141

increase in effective path length for the diffusion of water vapor through the polymer

matrix. This is in agreement with the XRD and TEM results.

Meanwhile, water vapor barrier of PHBV/CS30B nanocomposites film slightly

improved up to 3% loaded CS30B. The higher clay content sharply diminished the water

vapor resistance of nanocomposites. The high WVTR values with increasing the CS30B

content, is probably due to aggregation of clay. The aggregated structure of nano-clays

not only provides channels or micro-voids at the interface of polymer and clay, but also

may create direct pathways for water vapor that facilitate the migration of water vapor

molecules through PHBV matrix. Such low WVTR values enable the PHBV/PHB-g-

CS30B films as green-based packaging materials with extremely high barrier toward

water vapour transfer comparable with fossil-base polymer[52].

5.7. Conclusion

Biodegradable, environmental friendly, exfoliated and/or intercalated

nanocomposites were successfully prepared by incorporating modified clay in

biodegradable PHBV matrix. TEM and XRD analysis revealed homogeneous dispersion

of pre-intercalated silicate layers (PHB-g-CS30B) throughout the PHBV matrix even at

high clay content. It also showed that exfoliated structures were formed during the

preparation of PHBV/PHB-g-CS30B composites. AFM images showed that with the

addition of PHB-g-CS30B, the degree of agglomerated clay reduced. Also according to

the results obtained from AFM, presence of CS30B, at the surface of PHBV, increases

surface roughness due to its larger particle size as well as increased aggregation at the

surface, which caused higher water contact angle and much higher WVTR. The more

142

hydrophobic PHB-g-CS30B clay nanocomposites at all clay content showed lower

WVTR, water adsorption, and lower water contact angel.

Acknowledgement

Authors are grateful to NSERC Strategic Network – Innovative Green Wood

Fibre Product (Canada)

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150

Chapter 6 Thermal analyses and rheological behavior of exfoliated

poly(hydroxybutyrate-co-hydroxyvalerate) bionanocomposite:

Influence of pre-intercalated nanoclay incorporation

Abstract

A pre-intercalated nanoclay (PGCS) was synthesized via ring opening

polymerization of Polyhydroxybutyrate onto Cloisite®30B (CS). The

Poly(hydroxybutyrate-co-hydroxyvalerate) (PHBV) based bionanocomposites were

prepared by dispersing different amounts (1 to 20 parts per hundred resin, PHR) of the

nanoclay (CS and PGCS) into PHBV via solution intercalation process. The influence of

CS and PGCS on the rheology, morphology, thermal behavior of PHBV based

bionanocomposites has been investigated in details. The bionanocomposites drastically

improved the thermal stability of the pristine polymer matrix counterpart. The differential

scanning calorimetry (DSC) analysis revealed that the glass transition temperature and

melting point of the bionanocompositc decreased compared to its pristine counterpart.

The PHBV/PGCS bionanocomposites exhibited better thermal stability in comparison to

the PHBV/CS bionanocomposites. The transmission electron microgram (TEM) revealed

the exfoliated PGCS structure in the PHBV matrix. There was a great improvement in

viscoelastic behavior like complex viscosity and storage modulus, on incorporation of

PGCS in PHBV while CS had no influence on complex viscosity of PHBV.

151

6.1. Introduction

In last two decades, poly(hydroxybutyrate-co-hydroxyvalerate) (PHBV) has

attracted significant academic and industrial attention in the field of food packaging,

paper coating, pharmaceutical, toys , personal hygiene and automobile industry.

However, the industrial use of PHBV is still restricted due a number of factors such as, to

its high cost, high crystallinity, mechanical brittleness, poor process ability[1], high

melting temperature close to its thermal degradation[2], low thermal stability[3, 4], and

abrupt viscosity drop at melting temperature[5]. The close proximity of the melting

temperature to the thermal degradation temperature leads to a rather narrow window for

the process ability of PHBV. Moreover, slow crystallization and low melt strength i.e low

melt viscosity makes the melt processing of PHBV into films very difficult [6-8].

To overcome these weaknesses, several methods such as polymer blending[9-12],

grafting[13, 14], copolymerization[15, 16], and nanocomposites[17-19], have been

investigated,. Accordingly, it seems that the addition of nanoparticles to

poly(hydroxyalkanotes) (PHA) family is an effective ways to improve the physico-

mechanical properties of PHBV and make it more competitive with petroleum polyesters.

According to literature many researchers have made several attempts to prepare

bionanocomposites base on PHBV/layered nanoclay[20-23]. As mentioned before,

PHBV possesses low thermal stability and low melt flow viscosity, which occurs around

at 180 °C, i.e., around melting point of PHBV, is one of the most serious problems and

responsible for the poor process ability of PHBV. So far only few studies have been

carried out with a view to improve the rheological or viscoelastic properties of PHBV.

Zhao et al. [24] studied the rheological properties of PHBV/PLA blend, which showed a

152

Newtonian fluid behavior and found that an 30/70 PHBV/PLA blending ratio by weight

exhibited higher values of shear viscosity. While Modi[25] further examined the

rheological properties of PHBV with PLA grades 3051D, 4042D, and 6202D at 175 ºC

using a micro-compounder and no significant improvements has been reported in the

PHBV viscosity with the addition of PLA. Whereas, Zhao et al [26] reported that

PLA/PHBV silica layer nanocomposite showed a strong shear-thinning behavior, over

the full frequency range. As a crystalline nucleating agent, nanoclay significantly

improved the crystallinity of PHBV in the blend, thus leading to a relatively high

modulus for both solid and microcellular specimens. However, the addition of nanoclay

had less of an effect on the tensile strength and strain-at-break.

Moigne et. al. [27] found out that addition of organic motrmorillonite

(Cloisite®30B) to PHBV actually harmed the rheological properties of PHBV, which

were prepared through extrusion process at high Cloisite® 30B concentration.. Even

dilution of the Cloisite®30B masterbatch from 20 wt% to 2.5%did not show any

significant effect on viscoelastic behavior of PHBV. They claimed this phenomenon is

attributed to thermal degradation of PHBV via single-screw extrusion process. On the

other hand, Zembouai and co-workers[28]

showed that 3 wt% Cloisite®30B enhanced the

viscoelastic properties of PHBV/PLA blends just at low frequency significantly. This

effect became more substantial when 5 wt% of compatibilizer were used to more

homogeneously distribute clay in the blend matrix.

Up to now, many researchers worked on PHBV/clay nanocomposites, but none of

them could achieve much improvement in the rheological behavior and thermal stability,

while obtaining highly exfoliated morphology simultaneously. In the present work, for

153

the first time, the viscoelastic behavior, thermal properties and thermal decomposition of

exfoliated PHBV nanocomposites with modified nano-layer montmorillonite were

studied in detail and compared with PHBV-Cloisite® 30B (commercial available organo-

modified montmorillonite) systems. Furthermore nonlinear viscoelasticity was applied to

observe the Payne effect (decrease in storage modulus with strain amplitude) and

mathematical modeling is used to derive the filler–polymer interactions.

6.2. Materials



mol−1


and ENMAT Y1000P content of 3% hydroxyvalerate. Cloisite® 30B (CS), was kindly

donated by Southern Clay Products (Gonzalez, TX) and a Modified Cloisite 30B (PGCS)

were synthesized through ring opening polymerization of β-Butyrolactone on Cloisite30B

surface according to our previous article (Scheme 6-1) [1]. . All powder materials were

dried before use under vacuum overnight at 50 °C.

Scheme ‎6-1. Structure of CS30B and PHB-g-CS30B

154

6.3. Bio-nanoomposite film preparation

The solvent-casted composites were prepared by dissolving PHBV powder in

chloroform (5%, w/v) at 50 °C. The pre-suspended nanoclay over night was added into

the PHVB solution after the PHBV powder was dissolved completely, and the mixture

was vigorously stirred for 5 h. For reference, pure PHBV films were made by dissolving

PHBV powder in chloroform and those solutions were cast onto silicon wafer. Similarly,

the composite films with 0, 3, 5, 10, 20% (w/w) of nanoclays in PHBV matrix were also

prepared subsequently. The mixture left in a fume hood overnight to evaporate the

solvent. The films were peeled off easily from the glass and stored in a vacuum

desiccator until use.

6.4. Characterization of the bio-nanocomposites

To determine the intermolecular interaction of nanocomposites component the

infrared spectra of the pre-dried PHBV, PHBV/CS and PHBV/PGCS nanocomposite

films were recorded with a Perkin Elmer Spectrum 100 FTIR spectrophotometer in an

ATR mode at room temperature (25 °C). The samples were scanned from 4000 cm- 1

to

600 cm-1

with a resolution of 4 cm- 1

All the spectra were taken after an average of

32 scan s for each films.

Morphology of nanocomposites were studied using Transmission electron

microscopic TEM (JEOL 2011 STEM) on very thin films (100 nm) of PHBV

nanocomposite, cast directly over the copper grids of 300 mesh size. Accelerate on

voltage was 200 kV. Prior to analysis, the samples were cut with ultramicrotome in water

bath. The inorganic components appear black/grey coloured on the micrographs.

155

A differential scanning calorimetry (DSC) (model DSC Q20, TA instruments,

USA) was used to determine the melting behavior and the crystallinity of film samples. A

set of heating/cooling/heating ramps was carried out following a three-step procedure. A

sample of approximately 5 mg was first heated from room temperature to 200 °C at 10

°C/min and held in the melting state for 2 min to erase the thermal history. Then, a

subsequent cooling was performed from 190 °C to -50 °C at 10 °C/min. In the second

run, the sample was heated from -50 °C to 190 °C at 10 °C/min again to investigate the

melting behavior and to measure the crystallinity of bionanocomposites. The DSC cell

was constantly purged with nitrogen at a flow rate of 50 mL/min. The relative

crystallinity (𝜒𝑐) was calculated by equation (6.1)

𝜒𝑐 =∆𝐻𝑚

∆𝐻𝑓 ∗ 𝜑∗ 100 [6.1]

where ΔHm is the melting enthalpy detected in the experiment, while 𝜑 is the

weight fraction of PHBV in the bionanocomposites and ∆𝐻𝑓 represents the perfect

enthalpy of 100% crystalline PHB, which is reported as 146 J/g[29, 30].

Thermo-gravimetric analysis (TGA) was performed using a thermogravimetric

analyzer (model SDT Q200, TA instruments, USA). A sample of approximately 15 mg

was heated from room temperature to 650 ˚C at 10 ˚C/min under air atmosphere. Both

thermogravimetric (TG) and differential thermogravimetric (DTG) thermograms versus

temperature were recorded to investigate the thermal stability of samples.

A rotational rheometer (Rheostress600, Haake Co) was used to study the

viscoelastic behavior of the samples. The parallel plate with a diameter of 25 mm and a

gap height of 2 mm was used for frequencies sweeps. The shear dynamic measurement

156

was carried out at temperature of 175 ˚C and 185 ˚C with the frequency sweeps ranged

from 0.001 to 1000 rad/s, while the maximum strain was used at 1% under dry air during

the measurement. The values of storage modulus (G’) and loss modulus (G’’) against the

frequencies were obtained in the results and complex viscosity calculated from G’ and

G”.


6.5.1. FT-IR analysis to investigate intermolecular interaction

The FTIR spectra of pristine PHBV and the nanocomposites are compared in

Figure 6-1. The peak associated with the stretching vibration of carbonyl group of PHBV

was sharper than it was for the nanocomposites. In addition, intensity of peaks that

belong to PHBV/PGCS, diminished extremely. Regarding the nanocomposite with PGCS

loading, the shift and broadening of the bands related to the O–H and C=O stretching are

considerably more pronounced compared to PHBV/CS nanocomoposites. Such behavior

can be attributed to the strong interaction between PHBV and PGCS in the interface of

nanocomposite. Moreover, a new peak appeared in spectra of PHBV/PGCS and

PHBV/CS at 460 cm-1

, which was more intense in the former one. This peak is attributed

to Si-O-Si and Si-O bending stretching vibration due to the changes in the intermolecular

interaction between PHBV and nanoclay.

157

Figure ‎6-1. FTIR spectra of PHBV, PHBV/CS 90/10 (wt/wt%), and PHBV/PGCS 90/10

(wt/wt%)

6.5.2. Morphology of Nanocomposites

In Figure 6-2, one can see TEM images from ultra-thin cuts of two

nanocomposites containing 5% and 20% of CS (Figure 6-2a and 6-2b, respectively) and

containing 5% and 20% of PGCS (Figure 6-2c and 6-2d, respectively). Differences

between the two nanocomposite and the nanoclay contents are clearly visible. For the 5%

CS,, separate layers and partly still connected layered (intercalated) structures were

observed. Whereas, with the increase in CS content to 20%, almost fully connected

layered (agglomerated) structures can be seen. Conversely for PHBV/PGCS 95/05 as

shown in Figure 6-2c nanoclay layered platelets are aligned in a complete non-parallel

manner (fully exfoliated network) with homogeneous dispersion. Furthermore, with

addition of 20wt% (Figure 6-2d), a uniform distribution of the PGCS can still be

observed with a complex exfoliated and intercalated morphology. This exfoliated

158

morphology can greatly improve the thermal properties, rheological behavior and melt

processing of PHBV for various industrial applications.

Figure ‎6-2. TEM images of PHBV nanocomposites a)PHBV/PCS 95/05 wt/wt% , b)

PHBV/PCS 80/20 wt/wt%, c) PHBV/PGCS 95/05 wt/wt%, and d) PHBV/PGCS 80/20

wt/wt%

6.5.3. Thermal properties and Crystallization behavior of nanocomposites

The characteristic parameters and thermal properties such as, glass transition

temperature (𝑇𝑔), melting temperature (𝑇𝑚), degree of crystallization (𝑋𝑐%), and heat of

159

fusion (∆𝐻𝑚) of the PHBV and as well as its nanocomposites, obtained from the

endotherm DSC patterns are listed in Table 6-1.

It can be seen in table 6-1, unlike the 𝑇𝑚, ∆𝐻𝑚 incorporation of nanoclay did not

show any significant change in the 𝑇𝑔 of PHBV. Nanoclay, irrespective of the type,

reduced the glass transition temperature around 3 ˚C while nanoclay loading hardly had

any effect in the 𝑇𝑔. This is in agreement with several other studies [31, 32], where it was

reported that 𝑇𝑔 does not depend upon the nanoclay concentration. It can be seen in

Figure 6-3 that all samples at second heating cycle show double melting peak. Bimodal

melting behavior of Polyhydroxyalkanoates (PHAs) and theirs copolymers has been

reported in literature [33-35]. It is well known that complex bimodal melting of PHBV at

various valerate contents is attributed to the melting–recrystallization during subsequent

second heating cycle [35]. As reported by Ziaee and Supaphol [36] melting peak with

lower temperature relates to the melting of principal crystals formed in cooling cycle, and

the latter peak with higher temperature is attributed to melting of the crystals melted and

recrystallized during second heating cycle. Moreover, in agreement with others report. It

can be seen in Figure 6-3 and 6-4 that there is a reasonable correlation between the lower

melting peak (heat of fusion) during second heating cycle (Figure 6-3a and 6-4a) and the

crystallization on cooling (heat of crystallization) (Figure 6-3b and 6-4b) at 10 C/min-1

in

all samples.

Comparatively, with increasing nanoclay loading, not only the double peaks still exist but

also the spacing between the peaks extended from 6 to 10 ˚C and the lower melting peak

of the bimodal endotherms is the predominant one. Moreover, with the content of

nanoclay increasing, double melting temperature shifted toward lower temperatures,

160

regardless of the nanoclay types. It seems in the presence of nanoclay, the crystallization

rate, size and/or size distribution of the crystal spherulites was effective throughout the

melt-crystallization process of the nanocomposites. This phenomenon has been reported

in various nanocomposites systems by several authors [37, 38]. Furthermore, with the

increase in nanoclay content, both the melting temperature values gradually shifted to

lower temperature but had no dependence on the nanoclay type. However, the addition of

the PGCS nanoclay to the PHBV did significantly affect the aforementioned thermal

parameters. As listed in table 1-6 melting temperature from 168.4 and 174.5 for PHBV

reduced to 157.8 and 168.05 for PHBV/PGCS 90/10 w/wt%.

Besides the melting temperature, the degree of crystallization decreased by the

addition of more than 1% of CS and 3% of PGCS to the system. It is believed that high

concentrations of nanaoclays (in a polymer matrix) prolong the crystal growth during the

cooling scan due to the intermolecular interaction between polymer matrix and nano-

layered clay galleries, which, partially immobilizes polymer chains. Resulting in the

reduction in the polymer chain mobility and increase in the energy required during

crystallization for the chain folding process, as a consequence the overall crystallization

of the system will be diminished. It can be observed that the ∆𝐻𝑚 supported the above-

mentioned phenomena and nanocomposites exhibited a lower ∆𝐻𝑚 than that of the

polymer matrix PHBV. However, as listed in Table 6-1, the ∆𝐻𝑚 of the exfoliated

nanocomposite (PHBV/PGCS) was lower than that of intercalated and agglomerated ones

(PHBV/CS). These results can be attributed to homogeneous dispersion and strong

intermolecular interaction with PGCS clay next to polymer chains, which prevent the

overall crystal growth severely.

161

On the other hand, as shown in Figure 6-4 addition nanoclay leads to a higher

cold crystallization temperature (𝑇𝐶), It means that the addition of nanoclay elevates the

nucleation stage, which leads to a faster crystallization grows with smaller crystals.

However, according to heat of diffusion results and lower overall crystallinity of the

nanocomposites, high clay concentration leads to the extension of imperfections in

crystals and the size and shape of crystals that lower the melting temperature.

Table ‎6.1. Thermal properties of PHBV nanocomposites

Sample Tg (°C) Tcc (

°C) Tm (°C) ∆𝐻𝑚 Xc

PHBV 5 77.62 168.4 174.5 87 59.6

PHBV/CS 99/01 2.8 78.24 167.8 172.9 90 61.6

PHBV/CS 97/03 2.3 86.32 168 172.6 73 50.0

PHBV/CS 90/10 1.8 84.36 166 171.6 70 47.9

PHBV/CS 80/20 -2.8 73.89 152 164.2 64 43.8

PHBV/PGCS 97/01 2 88.10 163 170.0 95 65.1

PHBV/PGCS 97/03 2.18 91.88 166 171.4 92 63.0

PHBV/PGCS 90/10 1.76 94.08 161 166.4 85 58.2

PHBV/PGCS 80/20 2.3 93.59 157.1 166.7 84 57.5

Conversely, the addition of 1% CS and 1% and 3% PGCS further increased the

degree of crystallinity of the nancomposites, This abnormal results probably due to the

fact that the nanoclays acted as strong nucleating and plasticizing agents both which

increased the polymer chain mobility with lower surface free energy barrier, as

consequence, the higher overall melt crystallization[39-41].

162

Figure ‎6-3. DSC PHBV/CS nanocomposites

Figure ‎6-4. DSC of PHBV/PGCSs nanocomposites

163

6.5.4. Thermogravimetric analyses

Basically the montmorilonite clays have greater thermal stability than the polymer

matrix and can physically hinder the polymer chains within and close to the silicate

layers, leading to greater resistance to thermal decomposition and has been assigned in

most of nanocomposite systems.

The TGA and DTGA results of the representative samples are shown in Figure 6-5

and 7, respectivley. All the degradation temperatures in details, calculated from the

curves are arranged in Table 6-2. Figure.6-5 displays the TGA curves for the pristine

PHBV, PHBV/CSs and PHBV/PGCS at various clay contents. It is observed that the

thermal stability of PHBV nanocpmposites improved compared to pristine PHBV for all

nanoclay contents and regardless of the type of nanoclay used.

The PHBV/CS did not show (Figure 6-5) any change in the onset and peak temperatures

of decomposition, irrespective of nanoclay content loaded. As can be seen from Figure

6-6a and Table 6-2 thermal resistant under air atmosphere increased gradually up to 3

wt% CS nanoclay content and afterward no more enhancement or negative effect on

thermal stability was observed with adding more CS. This phenomena may be attributed

to the agglomeration of CS at higher loading level may due to localized microscale hot

spots and lower oxygen barrier[42, 43], which is in agreement to TEM results.

On the other hand PGCS nanoclay exhibited superior thermal stability compared to the

corresponding pristine biopolymer and PHBV/CS nanocomposites at all nanoclay

loading. According to Table 6-2 of PHBV/PGCS 90/10 nanocomposites significantly

enhanced and the𝑇𝑑5%, 𝑇𝑑10%, 𝑇𝑑(𝑜𝑛𝑠𝑒𝑡), 𝑇𝑑(𝑝𝑒𝑎𝑘), and 𝑇𝑑(complection) increased 11.3,

164

11.6, 11.1, 16.4 and 20.8 ˚C compared to pristine PHBV.

Table ‎6.2.Thermal stabilities of the molten PHBV, PHBV/CS and PHBV/PGCS nano

composites under air

Sample Td5% Td10% Td(onset) Td(peak) Td(complection)

˚C ˚C ˚C ˚C ˚C

PHBV 254.1 259.4 263.1 269.9 274.5

PHBV/PGCS 97/03 262.6 267.8 270.3 280.6 288.2

PHBV/PGCS 95/05 265.4 271.1 274.0 284.3 292.3

PHBV/PGCS 90/10 265.7 272.0 274.2 286.3 295.3

PHBV/CS 97/03 262.0 267.3 271.7 278.2 283.2

PHBV/CS 95/05 261.7 266.8 269.1 278.3 283.9

PHBV/CS 90/10 260.6 265.9 268.0 276.1 281.2

Substantial differences between PHBV/PGCS and PHBV/CS in thermal stability,

thermal and rheological behavior indicate strong interactions and homogenous

dispersion of PGCS with PHBV. The enhanced thermal stability also points to the

highly exfoliated morphological structure of the bionanocomposite. The exfoliated or

intercalated nanoclay not only play critical role to increase mass transport barrier and

insulator for the gaseous molecules products of decomposition process but also the

silicate platelets act as an additional barrier, which most likely works against oxygen

diffusion towards the bulk of the PHBV[43].

In general, it can be concluded from TGA and DSC results that the more

compatible clay with exfoliated gallery played a very critical role on thermal properties

and thermal decomposition of PHBV. The modified nanoclays lowered the crystallization

165

temperature, the melting point and exhibited much higher thermal stability in comparison

with pristine PHBV.

Figure ‎6-5. a) TGA and b) DTGA thermograms of pristine PHBV, PHBV/CS and

PHBV/PGCS nanocomposites

166

6.5.6. Viscoelastic behavior of PHBV and its nanocomposites

In order to better understand the critical role of the exfoliated clay and their effect

on crystallinity, thermal stability, the inherent viscosity and melt rheological behavior of

PHBV nanocomposites were investigated. It is reasonable to assume that the exfoliated

morphology has a significant influence on other polymer nanocomposite properties. We

investigated the rheological behaviour of PHBV/nanoclay depending on composition and

clay type. Figures 6-6 and 7 illustrate the storage modulus (G’) and the complex viscosity

(η∗), respectively, of the PHBV and it nanocomposites with PGCS and CS at various clay

content.

Figure 6-6 illustrate G’ and G”, respectively, of the PHBV and it nanocomposites with

PGCS and CS at various clay content. As illustrated in from Figure 6-6a and Figure 6-6b,

it can be considered that G’ and G” of PHBV increase with increasing nanoclay content

regardless of clay type. PHBV/PGCS nanocomposites show much higher storage

modulus and the loss modulus, even for low nanoclay content compared to PHBV/CS.

The G’ and G” of the PHBV/HPGC nanocomposites improved sharply. Meanwhile, G’

and G” in PHBV/CS nanocomposites slightly improved with loading CS nanoclay. It

was observed that during the test regardless of temperature and full frequency ranges

PHBV shows a viscoelastic liquid like behavior, G”>G’ (G” almost 10 time more that G’

in logarithmic scale) with linear viscoelasticity over the all frequency (ω) test range.

These rheological behaviours implied that PHBV possesses a very narrow processing

window. Because of the crystallinity, laminate consolidation to be harmed at melting

point and higher. Same behavior can be observed for PHBV/CS nancomposites with

increasing CS nanoclay content up to 5 wt% but that behavior change to a viscoelastic

167

solid-like (where G”<G’) with further increasing in the the CS content. The PHBV/CS10

showed a viscoelastic solid-like behavior at frequency lower than 15 (S-1

). Conversely,

the presence of PGCS in PHBV matrix even at low nanoclay content leads to significant

change in viscoelastic behavior of PHBV from liquid-like to solid-like. It seems the

fractional interactions between nanoclay and polymer is the most parameter affecting the

viscoelastic behavior of nanocomposites. Figure 6-7 shows the complex viscosity (ƞ*) of

PHBV and its nanocomposites as a function of angular frequency at 175 °C. It can be

observed from Figure 6-6 the pristine PHBV and its nanocomposites were found to

exhibit non-newtonian flow behavior. However, as observed in Figure 6-7, all the PHBV

and its nanocomposites regardless of nanoclay content and clay type, exhibited shear

thinning melt viscosity behavior at 175 °C i.e. complex viscosity (ƞ

*) of nanocomposites

diminished with increasing angular frequency. However, no significant improvement in

ƞ*

of PHBV/CS nanocomposites was observed while, the ƞ*

of nanocomposites gradually

increased with increasing PGCS nanoclay content. This phenomenon in PHBV/PGCS

nanocomposites is consistent with the other observation [55] reporting that the intensive

interaction between the exfoliated nanoclay and matrix chains increases the complex

viscosity.

Furthermore, it has been shown Figure 6-6 that when the nanoclays was added to PHBV

for PGCS (at all nanoclay content) and for CS (more than 5 wt%) a plateau is appeared in

G’ graph. It can be seen clearly that G’ and G” were almost independent of the low strain

amplitude. According to literatures[44, 45] this phenomena can be attributed to strong

intermolecular interaction between nanoclay and polymer matrix and large agglomeration

of clay at high clay content as seen in PHBV/CS5 and PHBV/CS10. The linear

168

viscoelastic range reduces from 108 ω for the PHBV/PGCS1 to 9.3 ω, 3.8 ω, and 3.3 ω

for PHBV with 3, 5, and 10 wt% PGSC1 content, respectively, and from 140 for

PHBV/CS3 to 136 ω, 2.9 for PHBV with 5, and 10 wt% CS content, respectively.

However, all samples exhibited a non-linear viscoelastic behavior, with diminish in G’

when the strain amplitude increments, the phenomena known as Payne effect. The

reduction in G’ value exposes how amount of energy is disintegrated in the system.

These observations demonstrate the important role that compatible nanoclay can play in

PHBV market to be considered for various applications such as coating, blow molding,

extrusion process and, blow film extrusion.

169

Figure ‎6-6. Storage modulus G’(ω) and loss modulus G”(ω) for PHBV, PHBV/CSs, and

PHBV/PGCSs nanocomposites

170

Figure ‎6-7. Complex viscosity η*(ω) of PHBV, PHBV/CSs, and PHBV/PGCSs

nanocomposites

6.6.Conclusion

We have successfully prepared b io thermoplastic nanocomposites. It has been

demonstrated from TEM that the PGCS is finely and uniformly distributed within

exfoliated morphology in PHBV matrix, while Cloisite 30B dispersed within the

aggregates and within the particles that makeup the dual phase aggregates. The TGA

curves indicate improved thermal stability for the nanocomposites compared to the

pristine counterpart. Nanocomposites also exhibit lower melting temperature over the

entire concentration range regardless of nanoclay types than its pristine counterpart.

Moreover, the glass transition temperature of the PHBV of the nanocomposite decreases

171

independent of nanoclay concentration and crystallinity of system and reduced at high

nanoclay laoding. However nanocomposite contained PGCS illustrated better thermal

stability and thermal properties in general at all concentration. Cloisite 30B showed no

effect on complex viscosity and storage modulus of PHBV, while PHBV/PGCS

nanocomposites presented the remarkable viscoelastic behavior that can be implantable

in most of thermoplastic applications. In short, our study elucidates the importance of

nanoclay-polymer interfacial interaction and exfoliated nanoclay structure. It can be

concluded that surface modification can provide excellent dispersion of nanoclay in the

PHBV matrix and provide good interfacial interaction with desired properties. PGCS

not only provided the lower melting point and much higher thermal stability for PHBV

but also enhanced narrow melt processing window of PHBV with significant increment

in the storage modulus and complex viscosity.

Acknowledgement



172

Acknowledgement



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179

Chapter 7 Cellulose/nano-clay composite films with high water vapor

resistance and mechanical strength

Abstract

Cellulose based films were fabricated by dissolving the micro-crystalline

cellulose (MCC) and cotton linter (CL) in a LiOH/urea aqueous system, followed by

regeneration in acetone. The water vapor transmission rate (WVTR) values of the films

were measured and the results indicated that regenerated micro-crystalline cellulose

(RMCC) showed lower WTVR than regenerated cotton linter (RCL). WVTR values were

proportional to the thickness of samples. The films also exhibited high oil resistance as

no fat oil could pass through the RMCC and RCL films even after a week. The RMCC

and RCL/natural-montmorillonite nanocomposites (Na-MMT) films were prepared and

characterized. The effect of Na-MMT loading on the mechanical, crystalinity and water

vapor transmission properties of the nanocomposites was further investigated. Na-MMT

loading improved the mechanical properties and decreased the crystallinity of the

nanocomposite films. The results indicated partially intercalated nano-layered structures

with WVTR as low as 43 g/m2/day.

KEYWORDS: Regenerated-cellulose based film, water vapor barrier,

nanocomposites, and layered silicate

180

7.1. Introduction

The use of renewable materials within the packaging industry has gained

increasing interest in recent years Lindstrand [1]. Cellulose fiber is the most abundant

biopolymer on planet earth and due to its sustainable, biodegradable and environmental

friendly nature, it is widely utilized in many fields. However, the barrier properties of

porous and hydrophilic cellulose fiber network products are inadequate for most barrier

applications[2]. Particularly, the porous structure of cellulose fiber-based network

(paper) remains the main challenge and limits the use of these fibers in moisture and

greasy barrier applications[3]. The smaller pore size of the regenerated cellulose based

films, compared to paper, makes them promising and feasible candidates for the

preparation of cellulose-based food packaging [4-7]. In last two decades, several organic

systems such as, dimethyl sulfoxide (DMSO)/triethylamine/SO2, 3-dimethyl-2-

imidazolidinone (DMI), ammonia/ammonium salt (NH3/NH4SCN), N,N-

dimethylacetamide (DMAc)/LiCl and, alkaline/urea aqueous solutions, have been

investigated for producing regenerated cellulose films [8]. Alkaline/urea aqueous solution

is relatively low cost and non-toxic. This solvent system has a great potential for the

development of environmentally friendly and economical regenerated transparent

cellulose films [8].

Non-toxic regenerated cellulose films formed by all mentioned solvent system possess

excellent mechanical properties, transparency, and barrier resistance to oxygen, carbon

dioxide and water vapor due to more extensive hydrogen bonding among cellulose

chains, compared to native cellulose I films [9-11]. Regenerated cellulose based films

have substantially higher oxygen barrier than biodegradable synthetic polymers such as,

181

poly(vinylidene chloride) and poly(vinyl alcohol), even at low relative humidity [9]. On

the other hand, the regenerated cellulose films due to their higher crystallinity have

superior mechanical, water vapor and oxygen barrier properties [9, 12, 13], compared to

the commercial Cellophane films. However, the water vapor and oxygen barrier

properties of the regenerated cellulose, in comparison to the conventional bio-

thermoplastics, are still quite high [14, 15]. These apparent shortcomings have raised the

high demand on developing new strategies to enhance the barrier properties of

regenerated cellulose films [13, 16-18]. Extensive research has been conducted to

fabricate cellulose films from a variety of cellulose fibers by varieties of methods such as,

modification of regeneration process [9, 16], film surface modification [19, 20], layer-by-

layer coating [21-23], polymer blending, regenerated cellulose nanocomposites [24-27].

Sodium montmorillonite (MMT) is a nontoxic, inexpensive, chemically and thermally

stable, and commercially available nanoclay and is one of the most widely accepted

reinforcing materials in polymer nanocomposite industry. MMT, due to its specific

geometrical dimensions and high aspect ratios, possesses unique properties that make it

behave much differently from conventional clays or micro-clays. Regenerated

cellulose/nano-layer clay nanocomposites could be an innovative solution to improve

regenerated cellulose film’s weakness even at low nano-clay content.

Cerruti et al. (2008), investigated the morphologies, thermal and thermal

oxidative properties of cellulose/organo-modified montmorillonite (3% MMT)

nanocomposites prepared by dissolving cellulose in N-methylmorpholine-N-oxide water

solutions [28]. The homogeneous dispersion of clay in the cellulose matrix led to the

improvement of thermal, oxidative and mechanical properties; and the nanocomposites

182

also showed increased degradation temperatures compared to native cellulose.

Mahmoudian et al. (2012), studied the permeability of oxygen, water absorption,

morphology and mechanical properties of regenerated cellulose/montmorillonite

nanocomposites prepared in 1-butyl-3-methylimidazolium chloride using a casting

method. All the properties of nanocomposites with 6 wt% MMT loading were

dramatically enhanced compared to the unfilled regenerated cellulose. They claimed that

the decreases in barrier properties, thermal stability and tensile strength at a high MMT

content (8 wt%) were associated with the aggregation of MMT platelets. As well Yang et

al. (2014), recorded the mechanical strength and gas barrier water absorption curves of

cellulose films (with a viscosity average molecular weight of 8.6×104 gmol

−1) and

corresponding intercalated montmorillonite nanocomposites prepared from LiOH/urea

solutions. They found that with increasing clay content, the oxygen permeability of the

nanocomposite containing 15 wt% clay at 50 and 75 RH% was decreased by 42 and 33%,

respectively, compared to that of the native cellulose film. In addition, the tensile strength

of the nanocomposite films was 161 MPa and the Young’s modulus was 180% higher

than that of neat cellulose film. However, above 15 wt% clay content the increase in

water absorption and oxygen permeability was clearly obvious, probably because of

aggregation of silicate layers.

In this work, the mechanical properties, crystallinity, water vapor and grease

resistance of two grades of Cellulose fibers with different molecular weight and

crystallinity, as well as their nanocomposites with low MMT content were examined and

compared. However, to the best of our knowledge, despite an extensive search, no

detailed study on water vapor barrier and the crystallization of regenerated

183

cellulose/MMT nanocomposites, MMT clay/regenerated cellulose interactions, as well as

crystallinity, and molecular weight effect on water vapor barrier and mechanical

properties of regenerated cellulose were not found.

7.2. Experimental section

7.2.1. Materials

Two types of cellulose, microcrystalline cellulose (MCC) and white cotton linter

(CL), were obtained from Alfa aesar and Arnold Grummer USA, respectively. The

viscosity-average molecular weights of two samples are 3.8×104 and 17.3×10

4,

respectively, determined in cadoxen at 25 oC with a viscometric method. Commercially

available lithium hydroxide (LiOH 98%) and urea (98%) were of regent grade (from Alfa

aesar). Sodium montmorillonite (Cloisite Na+) with a cation exchange capacity (CEC) of

92.6 mequiv./100 g clay was kindly donated by Southern Clay Products (Gonzalez, TX).

All chemicals were used without further purification.

7.2.2. Preparation of cellulose nanocomposite films

The regenerated microcrystalline cellulose (RMCC) and regenerated cotton linter

(RCL) nanocomposites were prepared through a solution method with the addition of 0,

1, 3, 5 and, 10 wt% Cloisite Na+ (Na-MMT), which were labeled as RMMC, RMMC1,

RMMC3, RMMC5, and RMMC10 for microcrystalline cellulose nanocomposites, and

RCL, RCL1, RCL3, RCL5, RCL10 for regenerated cotton linter nanocomposites,

respectively.

184

For the fabrication of the nanocomposite film, the following process was

employed: First, appropriate amounts of nanoclays were dispersed in pre-prepared

LiOH/urea/H2O with a ratio of 5:15:80 (wt%) which was then stirred for 10 h at 10 °C

afterward mixture was stirred with sonication for an extra 2 hour at room temperature.

The solvent was filtered through a 0.4 µm filter before use and was pre-cooled to -15 °C

in a refrigerator. Then, a desired amount of cellulose was added immediately into pre-

cooled solvent, and stirred vigorously for 10 min at ambient temperature to obtain the

doped transparent cellulose nanocomposite. The resultant nanocomposite dope was

degassed by centrifugation at 7000 rpm for 15 min at 5 °C to completely remove bubbles

and cast on a glass plate with desirable thickness, and then immersed in acetone at 0 °C

[9] for 15 min to allow regeneration. The hydrogel regenerated cellulose sheet, was

mounted repeatedly on a polypropylene (PP) plate with adhesive tape to prevent

shrinkage and was thoroughly washed with running water and finally dried on PP plate at

room temperature. Figure 7-1 shows of the coins, cheese, berries and flower wrapped

with RCL10 and RMMC10. It seems the regenerated cellulose Na-MMT composites are

quite transparent for packaging and wrapping applications.

185

Figure ‎7-1. photographs of a) RCL10 and b) RMMC10 transparent film for wrapping and

packaging application


Fourier transform infrared spectrum

The regenerated cellulose and regenerated cellulose/clay nanocomposite films

were characterized using Fourier transform infrared spectroscope (FTIR); and the spectra

were scanned by the attenuated total reflection mode on Nicolet IS5 (Thermo Fisher

Scientific, USA) for 32 times with a resolution of 4 cm-1 from 4000 to 450cm-1.

X-ray diffraction

The crystallinity of the regenerated cellulose and regenerated nanocomposite

based films was investigated using an X-ray diffractometer (a Bruker D8 Advance) using

http://en.wikipedia.org/wiki/X-ray_Diffraction

186

a rotating anode generator and a wide-angle power goniometer. The patterns with Cu Kα

(the weighted average k is 0.15418 nm) at 40 kV and 30 mA and a SOL-XE energy-

dispersive detector were recorded with a scan rate of 1°/min over the 2θ range 2–40°.

The X-ray optics were the standard Bragg-Brentano para-focusing mode with the X-rays

diverging slit (1.00 mm) at the tube to strike the sample and then converging through an

anti-scatter receiving slit (1.00 mm) and a detector slit (0.20 mm).

The maximum intensities of cellulose II (regenerated cellulose based film)

crystalline region diffraction peaks are typically located at 2θ ≅ 12.2°, 19.9° and 22.2°

[with Miller indices of (1-10), (110) and (020), respectively]. Our cellulose II patterns,

suffer greatly from preferred orientations, and the intensities of the (110) and (020) peaks

are substantially diminished in comparison with samples with randomly oriented

crystallites. The crystalline Iβ form of native cellulose has typical diffraction peaks at 2

= 14.9, 16.5 and 22.7° which are assigned Miller indices of (1-10), (110), and (200),

respectively [29].

The crystallinities ( cx ) of native cellulose, the regenerated cellulose based films

and the nanocomposites (Table 1) were calculated by Origin 8.5 software according to

Park et al. [30] and Nam et al.[31].

The interlayer spacing (d-spacing) was calculated from the Bragg Equation (7.1).

sin2

wnd [7-1]

187

Where d is the interlayer distance of clay, λw is the wavelength of incident X-ray used, n

is the order of reflection, which is equal to 1 for the first order, and θ is the measured

angle.

Water vapor and gas transmission rates (WVTR)

WVTR measurements for all samples were performed on IGA-003 (Hiden-Isochema,

Warrington, UK) which consists of a high sensitivity microbalance (0.1 µg) accordance

to the methods described in TAPPI standard T464 om-12 (2012) and ASTM E96/E96M-

05 (2005). In this experiment, all samples (thickness 20, 50 and 90 µm) were pre-

conditioned on WVTR test measurement conditions, at 23 °C and 50 RH%, for 72 hours

to reach water adsorption equilibrium. The experimental set-up for measuring WVTR is

shown in Scheme 1. The measurements were carried out at 23 °C and 50±2 RH%, which

was achieved using saturated magnesium nitrate and flowing dry nitrogen gas at a flow

rate of 10 mL/min. After the permeation cell was placed in a chamber, the data was

collected after 1 h to allow transmission rate to reach a steady state. The weight

(including sample, permeation cell and salt solution) loss is proportional to the testing

time. At constant temperature and RH%, the WVTR can be calculated from the change in

weight of the container, test time, and the area of exposed cellulose film, as described by

Equation (7.2):

timeArea

changeWeightWVTR

. [7.2]

188

Scheme ‎7-1. IGA experimental set-up for WVTR

Oil barrier properties

The grease resistance of the cellulose films was evaluated according to a modified

standard method TAPPI T507 (Scheme 7-2) [3, 32]. The method involves the use of oleic

oil with an oil soluble red dye to measure the penetration of the oil through cellulose film.

Moisture free colored oil was prepared by mixing 100 ml of oil with 1.0 g of Oil Red O, a

red dye, and 5 g of anhydrous CaCl2 in a flask. 0.5 mL of the colored oil was applied to

pre-saturate the blotter paper sheet with colored oil. The assembly (Scheme 7-2) was

placed in an oven at 60 °C. Samples were removed every hour from 8h up to a week, to

measure the possible amount of oil that passed through the cellulose samples to the clean

blotters.

189

Scheme ‎7-2. Oil barrier properties set-up

Tensile strength measurement

The tensile strength (σb), modulus (E), and %elongation at break (εb) of the films

were determined using a universal tensile testing machine (Instron Model 4465,

Norwood, USA) at a stretching speed of 50 mm/min according to ASTM D882. Prior to

testing, all samples were conditioned at a temperature and humidity of 23 ± 1 °C and 50 ±

5 RH% for 72 h. The all films were cut into dumbbell-shaped samples. At least ten

dumbbell shaped samples (115.0mm×6.0mm) were tested for each film, and the average

thickness was recorded using a micrometer caliper prior to the test. Strain was assumed to

be the equal to the crosshead displacement divided by the original length of the sample.

190

Engineering stress was determined by dividing the load by the cross-sectional area of the

sample.


7.4.1. FTIR spectroscopy

In this work, FTIR spectroscopy was used to investigate and compare the

intermolecular interaction of the neat RMCC and RCL with their Na-MMT reinforced

nanocomposite films. Figures 7-2 and 7-3 display the FTIR spectra of RMCC, RCL and

their nanocomposite films in the region of 1500-1800 cm−1

and 800–1400 cm−1

,

respectively. The peak at 1651 cm-1

corresponds to the C–O stretching vibration of C–O–

H [33]. A gradual reduction in the intensity of this peak can be attributed to an increase in

the intermolecular interactions of the nanoclay and the cellulose C–O–H group (Figure 7-

2). In RMCC nanocomposite films, not only the C-O peak intensity decreased

dramatically, but the peak center was also shifted toward higher frequency value (from

1651 to 1661 cm-1

) in RMCC3 spectral. However, it was observed at higher clay loading,

the C-O peak was shifted backward gradually.

191


Na-MMT clay content in transmittance in the region of 1500-1800 cm−1

On the other hand, a peak at 840-850 cm-1

was observed only in RMCC film

spectra (Figure 7-3), which is related to the vibration of Al-O-C, and may indicate there

might be a mineral–organic interaction that occurred in this region through the hydroxyl

group of MCC and Al on the nanoclay [34]. Moreover, the intensity of bands at 840-850

cm−1

increased with increasing the Na-MMT content. It can be construed from the FTIR

results that intermolecular interaction via RMCC and Na-MMT are stronger that RCL

and Na-MMT.

192


Na-MMT clay content in transmittance in the region of 800-1400 cm−1

7.4.2. X-ray diffraction

The XRD analysis was carried out to distinguish the level of clay dispersion

within the regenerated cellulose matrix and also to specify the d-spacing between clay

platelets which is often referred to as gallery height [35, 36].

Figure 7-4 illustrate the XRD patterns the cellulose I (CL and MCC) and Cellulose II

(RMCC and RCL). Differences between the XRD patterns of cellulose I and cellulose II

193

indicate that the crystal structure transformation from cellulose I to cellulose II occurred

after the dissolution and regeneration process.

Figures 7-5 and 7-6 show the XRD patterns of a series of clay-containing

nanocomposite films in addition to those of native RCL and RMCC, respectively.

The Na-MMT showed a main peak [d(001)] at around 9.10° (2θ) with layer distance

equal to 0.96 nm (Table 7-1). As can be seen from Figures 7-5 and 7-6, the (001)

reflections shifted to lower angles (larger d-spacings) for all nanocomposites as compared

to native Na-MMT. Figures 7-5 and 7-6 also show the position of crystal peak of Na-

MMT was shifted more significantly in RMCC, compared to RCL nanocomposite films.

194

Figure ‎7-4. XRD diffractograms of CL, MCC, RCL and RMCC

The main diffraction peak [d(001)] of RMCC with 1, 3, 5 and 10% nanoclay was

at around 4.28°, 4.6°, 5.28° and 5.62° (2θ), respectively. These peaks correspond to an

interlayer spacing of 2.06, 1.93, 1.68 and 1.58 nm, respectively (Table 1), which

indicated the entrance of RMCC into the interlayer galleries, an increase in d-spacing

between clay platelets, and the formation of nanocomposites with intercalated structure.

These results also put forward for consideration that the degree of intercalation decreased

with an increase in Na-MMT content from 1 to 10%.

195

These results also suggest that in RCL nanocomposites, the inter-platelet distance

of the Na-MMT does not change significantly compare to RMMC nanocomposite films.

This may be related to more inter molecular interaction of MCC with Na-MMT gallery;

and the finding is in agreement with the FITR results. In other words, the lower

molecular weight of MCC could have facilitated the penetration of dissolved MCC

molecules into the compact Na-MCC gallery and disrupted the crystalline structure of

nanoclay, leading to intercalated morphology with better and uniform dispersion of Na-

MMT in RMCC compare to RCL.

Figure ‎7-5. XRD patterns of a) RCL, b) RCL1, c) RCl3, d) RCL5, and e) RCL10 films

196

Figure ‎7-6. XRD patterns of a) RMCC, b) RMCC3, c) RMCC1, d) RMCC5, and e)

RMCC10 films

197

7.4.3. Crystallinity of regenerated cellulose and nanocomposite based films

The crystallinity of different native CL, MCC, regenerated cellulose based films

and the corresponding regenerated cellulose nanocomposite based films were calculated

according to Han et al. and Park et al. [30, 37] from XRD patterns and listed in Table 1.

The degrees of crystallization χc (%) are listed in Table 1. The crystallinity value

of the RCL and RMCC films are greatly lower than that of native CL and MCC (Figure

7-4), demonstrating that the true solvation destroys all intermolecular hydrogen bonds;

and the regeneration may fail to create new crystals as good as the original (native) ones

[38, 39]. The results showed that χc (%) values of RCL nanocomposite films decreased

gradually with the addition of Na-MMT in the RCL matrix.

The decrease of crystallinity may be attributed to the agglomeration of Na-MMT

occurring in the cellulose at higher clay loading, which retarded recrystallization during

regeneration process. For RMCC/Na-MMT nanocomposite films, as can be observed in

Table 1, the crystallinity of RMMC films is diminished slightly at Na-MMT content

higher than 3%; in another word, no significant difference was observed in the

crystallinity of RMCC with varying concentrations of nanoclay. While at 1% nanoclay

content (RMCC1) the crystallinity was slightly increased from 65.2% to 68.3%. This

result may suggest that a small amount of intercalated nanosized Na-MMT plates in

RMCC may act as nucleating centers and a slight preference for crystallization by

establishing a higher level of nucleation density[40].

198

Table ‎7.1. Experimental results of the regenerated cellulose from XRD paterns

Sample Main peak [d(001)]

related to Na-MMT

d-spacing (nm) Crystallinity χc

(%)

Na-MMT 8.70 1.05

MCC - - 82

RMCC - - 65.2

RMCC1 4.28 2.06 68.3

RMCC3 4.60 1.93 65.2

RMCC5 5.28 1.68 53.3

RMCC10 5.62 1.58 48.9

CL - - 71

RCL - - 49.2

RCL1 5.30 1.69 46.4

RCL3 5.90 1.47 32.1

RCL5 6.32 1.39 30.0

RCL10 6.70 1.29 27.1

7.4.4. Water vapor transmission and oil resistance

Table 7-2 presents water vapor permeability of neat RCL, RMCC and their

nanocomposite films with Na-MMT with different thicknesses (20, 50 and 90 µm) along

Cellephone (with 20 µm) and nano-fibrillated cellulose (NFC) film (with 40 µm), at 50±2

RH% and varies temperatures. RMCC and RCL demonstrate extremely higher resistance

against water vapor with lower WVTR than all the other films even with thickness of 20

µm (10 g/m2).

Figures 7-7a and 7-7b illustrate the WVTR of RMCC and RCL at various

thicknesses and temperatures, respectively. It seems that RMCC showed lower WVTR

than those from cotton linter (CL); and the water vapor resistance was proportional to the

thickness of films and inversely proportional to the temperature elevation.

It was reported that the water vapour transmission rates (WVTR) of polymers is

affiliated with the water solubility [41, 42], crystallinity [41, 43], and the effect of glass

199

transition temperature (Tg) [41]. Shogren, (1997), reported the water barrier property of

series bio-polyesters and found out that the water vapor resistance was negatively

correlated with lower crystallinities, higher free volumes, and higher polymer solubility

parameters. Duan and Thomas (2014), found out decreased linearly with increasing

crystallinity from 0 to 50% the WVTR of polylactic acid (PLA) decreased linearly and

the crystallites fundamentally acted as added fillers that inhibit gas and fluid to pass

through to polymer film by creating more tortuous path[43]. Tsuji et al also reported

almost identical footprint that the WVTR of PLA films decreased monotonically with

increasing crystallinity. Therefore, the high water vapor barrier of RMCC could be

explained by higher crystallinity of RMCC than RCL, which was concluded from XRD

results [44].

Figure ‎7-7. WVTR of RCL and RMCC at various thicknesses and temperature at 50

RH%

Furthermore, WVTR values of RCL and RMCC films with the different amounts

of Na-MMT loaded and at various thicknesses were determined, as shown in Figure 7-8.

The incorporation of Na-MMT, regardless of type regenerated cellulose film, has a

http://www-scopus-com.proxy.hil.unb.ca/authid/detail.url?authorId=7005440591&eid=2-s2.0-0031126418

200

significant effect on the WVTR at 50% RH and different temperatures, compared to

native regenerated cellulose.

As can be seen in Figure 7-8, the RMCC films showed lower water vapor

permeability than those from cotton linter regardless of the clay content, thickness and

temperature. The water permeability values for the nanocomposite films decreased with

the addition of clay into the RCL matrix, regardless of the film thickness. The water

vapor permeability of RCL nanocomposite films was reduced from 90, 158 and, 280

(g/m2/day) for neat RCL to 60, 110 and, 223 (g/m

2/day) for the films up to 5 wt% Na-

MMT loading with thicknesses of 20, 50 and 90 µm, respectively. By further increasing

the nanoclay content to 10 wt%, the water vapor permeability raised up to 77, 140 and,

245 (g/m2/day) with thickness of 20, 50 and 90 µm, respectively. The water vapor

permeability of the RMCC film was reduced from 65, 94, and, 234 (g/m2/day) to 43, 63,

and, 143 (g/m2/day) with the incorporation of up to 10 wt% Na-MMT with thickness of

20, 50 and 90 µm, respectively.

Overall, the water vapor resistance of a polymer film is enhanced with well-

dispersed additives with lower permeability. Furthermore, the addition of micro- and

nano-sized clay to the polymer matrix makes polymer films more tortuous, and thus

increases the pathways for water molecules to pass through the cellulose film. This could

explain the improved water vapour barrier properties of regenerated cellulose films after

the addition of Na-MMT[6, 18].

201

Figure ‎7-8. Water vapor permeability of RCL, RMCC and regenerated cellulose-Na-

MMT nanocomposite films with various Na-MMT contents measured at different

thickness at 50 RH% and 23 °C

As revealed by the XRD results, The Na-MMT clay is dispersed well in RMCC

matrix, compared to RCL matrix, which forced the water molecules to travel longer

distances across the films due to the formation of tortuous paths. On the other hand, the

crystallinity of the nanocomposite films was reduced by loading Na-MMT in RCL. The

reduction in crystallinity and the poor compatibility with Na-MMT may be the reason for

the decrease of water vapor barrier properties of the RCL nanocomposite films loaded

202

with more than 5% Na-MMT content while, the crystallinity in RMCC nanocomposites

are remained almost constant. This observation is also in line with the XRD data, i.e.,

besides other factors, the higher the degree of crystalline domains, the higher the barrier

performance, due to the physical impedance offered by crystal regions to the diffusion of

water molecules.

Furthermore, the regenerated cellulose films and nanocomposite films showed

excellent oil barrier properties, i.e. the red dyed oil did not penetrate films within the 7

day testing period at 60oC at all even 20 µm thick films which was probably due to tiny

porous structure of and highly hydrophilic structure of cellulose.

7.4.5. Mechanical Properties

The mechanical properties (tensile strength, Young modulus and elongation at

break) of both regenerated cellulose and theirs nanocomposite films were determined;

and are presented in Figure 7-9. The tensile strength of pristine RCL and RMCC is 160

and 107 MPa, respectively, which are in agreement with the results reported by others [9,

16, 25, 26].

Furthermore, it is shown that the tensile modulus and elongation at break of RCL

is higher than RMCC. It is generally observed, when the molecular weight of a certain

polymer goes up, the entanglement density will be increased. In another word, an

increase in molecular weight means that the average chain length per molecule increases,

which gives more entanglements, that is one of the main factors among others such as,

the crystallinity, chemical structure, glass transmission temperature, which affects the

mechanical properties especially impact strength of polymers [45, 46]. It means that the

203

increase of entanglement density could evidently improve the toughness, which may

explain the higher mechanical properties of neat RCL with more than four times higher

molecular weight than neat RMCC.

It is also observed that with increasing the loading of Na-MMT from 1 wt% to 5

wt% in RMCC and RCL matrix, as presented in Figure 7-9, the tensile strengths of RCL

and RMCC/Na-MMT nanocomposite films increased up to 219 MPa and 207 MPa,

respectively. This behaviour can be attributed not only to the high surface area, high

aspect ratio, and high elastic modulus of nanoclay but also to the strong interaction

between the hydrophilic Na-MMT and the hydrophilic polymer matrix (cellulose). Also

the intensive interphase interaction between the nanoclay and polymer matrix may

minimize stress concentration area and facilitate more uniform stress distribution, leading

to efficient load transfer to the nanoclay network[47].

204

Figure ‎7-9. Tensile strength, modulus and elongation at break of RCL and RMCC/Na-

MMT nanocomposite films at various clay content

As can be seen from Figure 7-9, RCL films show better tensile properties than

RMCC regardless of the nanoclay loading. Tensile strength and modulus of RCL

nanocomposites were improved gradually up to 3% nanoclay content; but no further

improvement was observed at higher nanoclay loading (> 3%). This may be due to the

possible agglomeration of Na-MMT layers in RCL matrix after a certain threshold

concentration is reached, which results in no further improvement of mechanical

properties. Furthermore, the tensile strength and modulus of RMCC nanocomposite films

were improved gradually in all clays contents except for 1% content, which can be

attributed to the better dispersion of Na-MMT into the RMCC matrix with no much effect

on crystallinity in comparison to RCL nanocomposite films. Surprisingly, the modulus

and tensile strength were enhanced dramatically by 13.5% (to 24.06MPa) with the

205

addition of Na-MMT at 1 wt%. This would be explainable not only by the best dispersion

of Na-MMT but also its role as a nucleating agent, which increased the crystallinity of

RMCC (see Table 1). Thus, these results clearly reveal the reinforcing property of Na-

MMT.

Table ‎7.2. Water vapor transmission rate of regenerated cellulose film at various

thickness, temperature and RH%

Samples Thickness (µm) g/m2 WVTR(g/m2/day)

50 RH% 10 °C

WVTR (g/m2/day)

50 RH% 23 °C

WVTR (g/m2/day)

50 RH% 38 °C

WVTR (g/m2/day)

90 RH% 38 °C

NFC 40 32 205±11*

371±18* 596±21

* 2700±33

*

Cellophane 20 10 189±4 321±3 530±7 1610±4

RCL

20 10 178±6 283±7 404±17 -

50 25 96±4 163±6 206±12 -

90 50 49±1 91±4 120±10 1340±21

RMCC

20 10 103±4 249±12 292±21 -

50 25 40±3 101±7 124±12 -

90 50 27±1 71±3 98±7 960±14

* Term “±” is the standard deviation

The change of elongation at break is different from the variation of tensile

strength, but the mechanism is more or less similar. However, a gradual decrease in

elongation of the RCL nanocomposites was observed with an increase in clay loading

more than 1 %. The embedment of the nanoclays into polymer matrix increases the stress

concentration, which may cause the brittleness in the polymer nanocomposites. However,

206

a sharp increase in elongation at break is observed when Na–MMT was incorporated into

RMCC at low concentration, about 1%. The elongation at break increased gradually up to

3% clay loading, and further addition of nanoclay led to the reduction of the strain. The

increase in elongation at break at low nanoclay content can be attributed to the better

dispersion of the nanoclay within the polymer matrix[48].

7.5. Conclusion

WVTR of transparent regenerated cellulose films from cotton linter and

microcrystalline cellulose and their nanocomposites in aqueous Urea/LiOH system has

been evaluated in detail in this article. WVTR values of RMCC and RCL exhibited

higher water vapor barrier properties at various thickness than NFC and cellphone film. It

was found the efficiency of the nano-clay addition depends on type of cellulose, cellulose

nona-clay interaction, the nano-clay content, and their level of dispersion. Incorporation

of nano-clay improved the mechanical and barrier properties of native RCL and RMCC

films. RMCC nanocomposite films exhibited better water vapor resistance (reduced

permeability) in comparison with RCL and its nanocomposite films. It was found all

regenerated films showed high penetration resistance to oil with no oil penetration at all.

It was also the FTIR and XRD results indicated stronger filler–matrix interaction and

intercalated structure in RMCC nanocomposites compared to RCL nanocomposite films.

Acknowledgment

Authors are grateful to NSERC Strategic Network—Innovative Green Wood

Fibre Product (Canada), and NSF China (21466005) for funding.

207

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213

Chapter 8 Preparation and Characterization of surface modified

Exfoliated Na-Montmorillonite nanoclay with quaternized cellulose

Abstract

In this article, surface modification of Na-montmorillonite (Na-MMT) was carried

out by a series of ion exchange reaction with quaternized microcrystalline cellulose

(MCC) and Cotton linter (CL). The celluloses were rendered quaternized by grafting with

3-Chloro-2-hydroxypropyltrimethylammonium chloride (CHPTAC). Quaternary

celluloses with various charge densities (380 to 1080 µeq/g) were used to modify Na-

MMT. Fourier transform infrared spectroscopy (FTIR) patterns identified intercalation of

quaternized celluloses into the Na-MMT layers and/or adsorbed on the surface of Na-

MMT. The effects of quaternized celluloses on the surface characteristics, and moisture

adsorption properties of modified Na-MMT have also been studied. Furthermore, the zeta

potential, and surface topology and the bulk morphology of modified Na-MMT thin films

were investigated by AFM, TEM, and XRD respectively.

8.1. Introduction

Applications of nanolayer clay in industrial purposes depend upon their structural

variations, specific surface area and caution exchange capacity (CEC). Na-

Montmorillonite (Na-MMT) surface has extensive negative charges that allow the

adsorption of positively charged materials based on electrostatic attraction [1, 2].

In the last decade, there has been a growing interest towards to polycation-

modified clays (PMC), due to their potential as surface-modified substrates[3] for many

214

applications. Polycation polymers and PMCs have been extensively studied as retention

aids or flocculating agents in paper making[4], enhanced oil recovery in crude oil

extraction[5], Superabsorbent[6, 7], wound-healing accelerator in biomedicines[8-10],

catalysts[11] or adsorbents for organic pollutants[12], as chelating polymer for heavy

metals removal[13, 14], and in polymer nanocomposites preparation.[15] PMC with

small particle size has great flocculating activities and its flocculating activities make it

suitable in papermaking processes, as polymeric flocculants or retention aids.[4] The size

and morphology of these PMCs can be adjusted by adjusting various parameters such as,

carrier liquids, process and the molecular structure of polycations. Müller et al[16]

reported that the particle size and stability of PMCs are influenced by several parameters

such as, surface charge behaviour, particle size, polycation concentration, pH, molar

mass, preparation method, mixing order and ratio, etc.

Usually in Polymer industry, the platelet-shaped nanoclays with high aspect ratio

and high surface area, with a thickness of few nanometers and lengths of several hundred

nanometers are used to enhance series of polymer properties including mechanical, gas

barrier, thermal stability and processability[17-19]. Often different ratio, length and

structure of organic oligomer of quaternary ammonium surfactant compounds (with at

least one long aliphatic chain, C12–C18, which are cationic in nature,), are used for clay

modification with a view to increase the interlayer spacing and enhance polymer-clay

interaction[20, 21]. However, the key issue to improve the physical and chemical

properties of polymer lies in the proper dispersion of the nanoparticles in the polymer

matrix.

215

Despite having high potential, It seems that polycation-modified MMT has rarely been

used as a dispersant in the polymer nanocomposites[22]. Several researchers have

reported fabrication of clay nanocomposites with hydrophilic polymers (i.e starch, PVA,

proteins, cellulose, and etc.) without obtaining a high extent of intercalation or exfoliation

structure in spite of hydrophilic nature of unmodified clay[23-26]. It seems that the

affinity, between hydrophilic structures of polymer and natural clay even in present of

plasticizers, is not so high.[27-29] This low extent of exfoliation and homogeneous

dispersion as claimed is related to negative charges (−O− and −COO− groups) on

hydrophilic polymer that will repel the negatively charged surfaces of the silicate

layers.[30] In order to overcome this problem several authors have reported modification

of nanoclay with chitosan[6, 31, 32] or proteins[33] . Ho et al., (2012)[30] has reported

fabrication of trimethylammonium-modified nanofibrillated cellulose (NFC) with

different types of layered silicates by using a combination of high-shear homogenization,

pressure filtration and vacuum hot pressing.

The main objective of this research work was to develop a novel polymeric matrix

via surface modification of Na-montmorillonite (Na-MMT). At first, microcrystalline

cellulose (MCC) and α-Cellulose was quaternized using cationization agent. Then the

surface of Na-montmorillonite (Na-MMT) was modified according to the cation-

exchange capacity (CEC) of Na-MMT and quaternized cellulose matrices. So far, to our

best knowledge, there have not been any studies investigating the modification of Na-

MMT with MCC and α-Cellulose. The modified Na-MMT/cellulose matrices have been

characterized by FTIR spectroscopic analysis (FTIR), microscopic analyses (AFM,

216

TEM), and X-ray diffraction. The surface charge density, moisture adsorption, and zeta

potential were also measured.


8.2.1. Materials & Methods

Two types of cellulose, microcrystalline cellulose (MCC) and α-Cellulose, were

procured from Sigma-Aldrich with viscosity-average molecular weights (�̅�ƞ) of 3.8 ×

104 and 35.7 × 104, g/mol respectively in cadoxen at 25 °C. Commercially available

sodium hydroxide (NaOH 98%) and Urea (98%) were of regent grade (Alfa aesar,

Haverhill, MA, USA), 3-Chloro-2-hydroxypropyltrimethylammonium chloride

(CHPTAC) aqueous solution (60 wt %) was purchased from Sigma-Aldrich and Sodium

montmorillonite (Cloisite Na+) with a cation exchange capacity (CEC) of 92.6 mequiv.

/100 g clay was kindly donated by Southern Clay Products (Gonzalez, TX, USA). All

chemicals were used without further purification. Standard solutions of NaOH and HCl

were used for adjusting pH. The dialysis tube was obtained from Fisher Scientific

(Waltham, MA USA) and has an Mw cut-off of 2000 Dalton.

8.2.2. Synthesis of quaternized celluloses in NaOH/Urea Aqueous Solutions.

(Scheme 8-1)

A 2 wt % cellulose was added in pre cooled to -12 °C, NaOH/Urea aqueous (8:12:80 by

weight), solvent with vigorous stirring for 5 min at room temperature according to Cai et.

al.[34] In previous studies, 3-chloro-2-hydroxypropyl trimethyl ammonium chloride

(CHMAC) was used as cationization agents [35-37] for functionalization of xylan,

217

polyvinyl alcohol (PVA), etc. In the current study, CHMAC was used to functionalize the

cellulosic matrices. According to Song et al.[38] a certain amount of CHPTAC aqueous

solution was added drop wise into the 200 ml cellulose solution (obtained previously),

with the molar ratio of 3:1, 6:1, and 12:1 of CHPTAC to anhydroglucose units (AGU) in

MCC and α-cellulose. The quaternized microcrystalline cellulose samples were labeled

as, NM3, NM2, and NM1 (Table 8-1). Whereas, the quaternized α-cellulose samples

were labeled as, NC3, NC2, and NC1. The mixture was stirred with speed of 8000 rpm at

ambient temperature for 16 h. After 16 h, the reaction was stopped by neutraliztion with

aqueous HCl and dialyzed with regenerated cellulose tubes against distilled water for 7

days and finally freeze-dried (Christ Alpha 1-2, Osterode am Harz, Germany).

Table ‎8.1. Surface charge density of quaternized celluloses

Sample Cellulose CHPTAC to

AGU mole

ratio

Titration

standard

Surface charge

density µeq/g

NM1 MCC 12/1 PVSK 1052



NC1 α- Cellulose 12/1 PVSK 746



Na-MMT - - poly(DADMAC) 920

218

Scheme ‎8-1.Quaternization of Cellulose (QC)

8.2.3. Synthesis of quaternized celluloses Intercalated Na-MMT Adsorbent

(Scheme 8-2)

Modification of MMT was carried out by a cation exchange reaction.[3, 4, 39-41]

In a 500-ml three-neck flask equipped with continuous reflux, certain amount of Na-

MMT was mechanically stirred with 200 mL of deionized water at 80 °C for 3 h followed

by 1 h sonication to properly swell the layered montmorillonite. The pH of the dispersed

Na-MMT was adjusted to approximately 5 using with aqueous HCl. After that, the

required amount of Na-MMT (according to Table 8-2) was slowly poured into 100 ml of

pre-dissolved quaternized samples (with pre-adjust pH of around 5) samples and heated

at 80 °C for 12 h under mechanically stirred. The resultant mixture was then centrifuged

with a speed of 13000 rpm at 25 °C for 30 min, then the slurry was filtered and

repeatedly washed with 50 °C 1% acetic acid solution and afterward with deionized hot

water three times. Finally, the slurry was freeze-dried with lyophilizer (Christ Alpha 1-2,

Osterode am Harz, Germany) to obtain the khaki modified montmorillonite powder. The

219

recipes, remained ash of modified clays at 600 °C and coding of the poly-cation Na-

MMT samples are summarized in Table 8-2

Scheme ‎8-2. Poly-cation modified MMT (PMMT) with QC

Table ‎8.2. The recipes, remained ash of modified clays at 600 °C and coding of the poly-

cation Na-MMT samples

Sample C2NC1 CNC1 CNC2 CNC3

SCD of QC

/ Na-MMT

NC1/Na-MMT

(2/1)

NC1/Na-MMT

(1/1)

NC1/Na-MMT

(1/1)

NC3/Na-MMT

(1/1)

Ash remained

at 600 °C

38 wt% 50 wt% 40 wt% 36 wt%

Sample C2NM1 CNM1 CNM2 CNM3

SCD of QC

/ Na-MMT

NM1/Na-MMT

(2/1)

NM1/Na-MMT

(1/1)

NM2/Na-MMT

(1/1)

NM3/Na-MMT

(1/1)

Ash remained

at 600 °C

42 wt% 56 wt% 48 wt% 38 wt%

220


Determination of surface charge density

The surface charge density of the samples was determined using a Mutek Charge Titrator.

A quantity of 5.0 ml of the silica sample mixtures was added to the titration cell

containing 5.0 ml of deionized water and then titrated with the standard poly(DADMAC)

or PVSK solution (1.0 mN). The particle charge density was estimated from the volume

of the polymer solution required to reach the end point during titration (Table 8-1).

Fourier transform infrared spectrum

The regenerated cellulose and regenerated cellulose/clay nanocomposites films

were characterized using Fourier transform infrared spectroscopy (FT-IR) spectra. The

films were scanned by the attenuated total reflection mode on Nicolet IS5 (Thermo

Fisher, USA). Spectra were scanned 32 times with a resolution of 4 cm-1

from 4000 to

450cm-1

.

X-ray Diffraction

XRD patterns of tiny film of dissolved poly-cation-MMT clays in buffered

solution of 7 and 10 pH were recorded using on an XRD diffractometer (a Bruker D8

Advance spectrometer) with a rotating anode generator operating at current of 20 mA and

a voltage of 40 kV. The diffractograms were measured for 2θ, with a scan rate of 1°/min

over the 2𝜃 range 2–40. The interlayer spacing (d-spacing) was calculated from the

Bragg’s Equation (8.2).

2𝑑 𝑠𝑖𝑛𝜃 = 𝑛𝜆𝑤 [8.2]

http://en.wikipedia.org/wiki/X-ray_Diffraction

221

Where d is the interlayer distance of clay, λw is the wavelength used, n is the

order, which is equal to 1 for the first order, and θ is the measured angle.

Moisture adsorption

The moisture adsorption of modified clays was determined using Belsorp-Max

(BEL JAPAN, INC., Osaka, Japan). The Belsorp-max is an automatic moisture

adsorption /desorption volumetric measuring unit that measures absorption isotherm in

the relative humidity range RH% 10 to 100. The adsorption isotherm of the samples was

measured at 25 °C. The samples were dried for 48 h at 40 °C under vacuum before the

isotherm measurement.

Zeta Potential

Zeta potential of the samples was measured at 25 °C using ZetaPALS analyzer

(Brookhaven Instrument Corporation, USA). The potential was calculated by measuring

the electrophoretic mobility using the same instrument. The mobility value was converted

to potential using the Smoluchwski’s approximation and the reported value is an average

of 10 measurements.

AFM

Surface topology of quaternized cellulose and modified clay tiny film was

investigated using AFM. To determine the topology of the samples, atomic force

microscopy (AFM) measurements were performed using a Nanoscope IIIa from Veeco

Instruments Inc., Santa Barbara. Sample films were prepared by dissolving and

dispersing certain amount of quaternized cellulose and modified clays powder,

respectively, in buffer 7 solution and the mixtures were vigorously stirred for 1 h at room

temperature and those solutions and slurries were cast onto laminated petri dish with

222

polyethylene wrap. The slurries left in a fume hood for 3 day to evaporate the solvent and

kept in desiccator cabinet for a week before surface structure evaluation using AFM.

TEM

The samples for the Transmission electron microscopy (TEM) analysis was

prepared and stored in a desiccator for a week. Modified MMT films were embedded in

epoxy resin plates were first trimmed with iron knives subsequently; the ultrathin sections

were microtomed from these faces with a diamond knife and then examined using a

JEOL 2010 STEM, operated at an accelerating voltage of 200 keV.


8.4.1. Chemical compositions and morphology

Figure 8-1 shows the FTIR spectra of quarterinzed celluloses modified Na-MMT clays

transmittance in the region of 450-1900 cm−1

. The peaks at 1479 and 1418 cm−1

represent

the scissoring band of -CH3 groups of the substituted CHMAC, and C-N stretching

vibration respectively. Moreover cellulose carboxyl group peak shifted to higher

wavelength and a new peak at 847 cm-1

, which it is related to the vibration of Al-O-R,

was observed in all modified clays spectra. This peak may indicate occurrence of a

possible mineral–organic reactions in this region through the R-N+ of quaternized

cellulose and Al-O on the nano-clay[42]. These changes in FTIR pattern confirm the

Poly-cation reaction.

223

Figure 8-1. FT-IR spectra of modified nanoclays from 450 to 1950 cm-1

XRD patterns (Figure 8-2) demonstrate the interlayer structure of polycation-

modified clay with quaternized cellulose at pH 6 and 10. The diffraction peak of all the

samples shifted significantly to the smaller 2θ angle compared to that of unmodified Na-

MMT and increased d-spacing is characteristics of exfoliated/intercalated layer

structures. Thus, samples at pH 6 showed a greater extent of the left shift corresponding

to a bigger d-spacing value than that of pH 10 treated samples (Table 8-3). These results

might be related to the lower solubility of quaternized cellulose at the alkaline media as

reported by Song et. al [38].

For samples at pH 10, the existence of two peaks implies the occurrence of

partially intercalated MMT. While the interlayer distances of C2NM1, C2NC1, and

224

CNM2 cannot be measured because no obvious diffraction peaks were observed in their

XRD patterns, indicating fully exfoliation of the silicate layers at pH 6.

These results also indicate that in quaternized cellulose with higher molecular weight

(NC1, NC2 and NC3), the inter-platelet distance of the modified Na-MMT changes less

compared to those modified with quaternized cellulose with lower molecular weight

(NM1, NM2, and NM3). In other words, the lower molecular weight of quaternized -

MCCs could have facilitated the penetration of water soluble molecules into the compact

Na-MCC gallery and disrupted the crystalline structure of nano-clay, leading to

exfoliated morphology with higher d-space

According to literature [43, 44], the amount of clay in the polymer is an

important parameter for the clay layers to be close to one another so this might be the

main reason why the d-space of CNC1 and CNM1 are lower than that of CNM2 and

CNC2, although the CEC of NM1 and NC1 are higher than that of NC2 and NM2

225

Figure ‎8-2. XRD of modified nanoclays at various pH, a) pH 6 and b) pH 10,

226

8.4.2. Water vapor adsorption of quartenized cellulose and modified clays powder

Water vapor adsorption of quartenized cellulose and modified clay powder in

Figure 8-3a and Figure 8-3b shows the moisture absorption study of quaternized

celluloses and the poly-cation modified clays in the range of 10 to 100 % RH. All

quaternized celluloses showed higher amount of water adsorption than that of polycation

modified clays for all the samples and at all humidity range (Fig. 8-2c to Fig. 8-2h).

While the poly-cation modified clays exhibited not only superior water vapor barrier with

weak substrate interaction but also enhanced moisture resistance which is even less than

that of MCC, α-cell and Na-MMT. The water vapor adsorption in poly-cation-MMT

clays are reduced systematically compared to that of the native celluloses and even Na-

MMT because of the presence of impermeable layered silicate in the poly-cation-MMT.

These exfoliated/intercalated layers increase the path-length for water molecules to

diffusion through the modified clay. Also, more hydrogen bonds may have formed

between the Na-MMT and the quartenized cellulose thus restricting the transmission of

water vapour molecules through the nanocomposite films. This intramolecular interaction

does not allow the establishment of tremendous hydrogen bonding with water. However,

it seems, at the higher relative humidity, the hydrogen bonds formed between the poly-

cation-MMT has caused a sharp increase in nanocomposites water adsorption content

which is still less than that of the quartenized cellulose. This phenomenon, reduction in

moisture absorption of this type of system is in good agreement with the results reported

by Perrotta, A. et. al.[45]. CNCs and NCs samples have the lower moisture adsorption

than the CMNs, and MNs respectively. CNC3 sample showed the lowest water vapor

adsorption that might be related to lowest CEC of NC3 (quaternized cellulose).

227

Figure ‎8-3. The isothermal moisture adsorption of quaternized celluloses and modified

nanoclay

228

8.4.3. Bulk morphology of polycation-clay

The TEM analysis of the C2NC1, CNM1 and CNC3 samples were carried out to

verify the state of dispersion of the polycation-clay layers into buffer media. Figure 8-5

shows the microstructural images of the C2NC1, CNM1 samples, obtained from TEM

analysis at pH 6 and 8. As shown in Figure 8-5 most of the polycation-clay layers were

randomly dispersed and intercalated into buffer regardless of pH. It seems lower pH

promoted the exfoliation morphology of surface modified clay. Sample C2NC1

illustrated high level of exfoliation regardless of pH than that of clay modified with NM1.

However, crystalline structure C2NC1 destroyed completely and clay layers arrange in

non-parallel order (fully exfoliated) at pH 6 (Fig. 8-6). The highly delaminated structure

of the C2NC1 and CNM1 layers can also be observed at the high magnification pattern

(Fig. 8-5b) revealed the absence of any aggregated spot. These results are in agreement

with XRD pattern observation.

.

229

Figure ‎8-4. TEM micrograph of diluted C2NC1 and C2NM1 at pH 6 and 8 a) C2NC1 at

pH8, b) C2NC1 at pH6 c) C2NC1 at pH6 (lower magnification), d) C2NM1 at pH8, e)

C2NM1 at pH 6, and f) C2NM1 at pH 6 (lower magnification)

230

Figure ‎8-5. Higher magnification TEM micrograph of diluted C2NC1 and C2NM1 at pH

6 and 8 a) C2NC1 at pH8, b) C2NC1 at pH6 c) C2NM1 at pH8, c) C2NM1 at pH 6

231

Figure ‎8-6. TEM of CNC3 tiny film at various pH; a) pH 6, b) pH 8, and c) pH 10

8.4.4. Surface topology of polycation-clay thin film

The AFM phase and 3D images of modified clay are presented in Figure 8-5 and

Figure 8-6. The average height and the equivalent mean roughness of the quaternized

cellulose and modified clay compounds are summarized in Table 8-3. It can be seen the

average height and the equivalent mean roughness of the quaternized cellulose samples

increased with decrease in the charge density.

As a result, the average height and the equivalent mean roughness values

obtained (Table 8-3) for the modified nanoclay samples are more than that of its

quaternized cellulose counterpart. The surface topology of clay modified with

quaternized MMC is different from that of clay modified with quaternized MMC in

general based on AFM images. It is observed that the order of the average height and

root mean squared roughness (RMS) were found to be CNC3>CNC2>CNC1>C2NC1

for clay modified quaternized α-cellulose and CNM3>CNM2>CNM1>C2NM1 for

clay modified quaternized MMC.

232

For polycation-modified clay, the surfaces became smooth with the increase in

charge density and surfactant loading of quaternized cellulose. The comb-like textures

on CNC3 and CNM3 samples were replaced with the smooth boundary on C2NM1 and

C2NC1. This reflects that the type of polycation cellulose molecules has a significant

influence on the surface topology of modified clay.

In addition, this reduction in surface roughness can be explained by the fact that

the surface modification reduced the filler-filler interaction and increased the d-spacing

of clay, which in turn reduced the formation of big agglomerates and improved the

filler dispersion As a result, the nanocomposite exhibited an intercalated/ exfoliated

structure in the buffer 7 media.

Figure ‎8-7. AFM 3D microroughness of a) C2NC1, b) CNC1 , c) CNC2,d) CNC3 ,e)

C2NM1 ,f) CNM1,g) CNM2, and h)CNM3 tiny film

233

Table ‎8.3. Summarized results from AFM test

Samples NM1 NM2 NM3 NC1 NC2 NC3

Rrms (root mean squared) (nm) 2.83 4.18 27.7 4.3 7.09 31.6

Ra (average surface roughness

)(nm)

2.27 2.6 22 3.33 6.0 25.9

Rmax (nm) 20.2 104 156 33.2 37.8 192

Samples C2NM1 CNM1 CNM2 CNM3 C2NC1 CNC1 CNC2 CNC3

Rrms (nm) 7.85 8.89 16.7 21.1 15.6 15.5 20.4 38.6

Ra (nm) 5.12 6.33 14.2 17 13.2 12.5 15.5 29.1

Rmax (nm) 237 62.5 100 118 67.1 78 148 227

234

Figure ‎8-8. AFM phase images of a) C2NM1, b) CNM2, c) CNM3,d) C2NC1, e) CNC2

and, f) CNC3

8.4.5. Zeta potential analyses at various pH values

The zeta potential of the quaternized cellulose, polycation-MMTs and Na-MMT

as a function of pH is presented in Figure 8-9. As shown in Figure 8-9a, the negative zeta

potential of Na-MMT is dependent on pH and it gradually decreased with the increase in

pH up to 10. This result is in agreement with those obtained by Durán et. al [46]. The ζ-

potential values for quaternized celluloses are plotted in Figure 8-9b at various pH at all

surface charge density value showed positive ζ-potential. In contrast to the Na-MMT, the

ζ-potential value of the quaternized cellulose slightly increased when the clay content

235

was increased and reached maximum at pH 6 before decreasing gradually with further

increase in pH. It seems that the ζ-potential value is proportional to surface charge

density as reported by Song et. al.[47]Through the poly-cation modification of Na-MMT

with quaternized cellulose, the surface charge of Na-MMT reversed from negative to

positive. The ζ-potential value of Polycation-modified clays showed the same trend as

was observed in quaternized celluloses. As can be seen in Figure 8-9c, the stability of

modified clay in buffers dropped more at alkaline media than acidic one. Thus the ζ-

potential value CNC3, CNC2 and CNM3 reduced from 18, 20.4 and 19.4 at pH 7 to 6.5,

8.6 and 8.5 at pH 14 (1 molar NaOH) , respectively, which is the sign of the incipient

instability in alkaline media. C2NM1 and C2CN1 showed higher ζ-potential value

compared to CNC1 and CNM1, respectively. It seems the double dosages of NC1 and

NC2 in clay modification increased positive charges on modified clay moderately, as

shown for C2NM1 and C2CN1 (Fig. 8-4(e)). It can be attributed to that the Na-MMT

external surface has been covered broadly in the bilayer formation when the quaternized

cellulose added over the CEC of Na-MMT

236

Figure ‎8-9. ξ potential analyses at various pH for a) Quaternized Cellulose samples b)

Modified-MMT, and c) Na-MMT

237

8.5. Conclusion

Poly-cation modified Na-MMT matrices were successfully synthesized from

ion-exchange reaction with quaternized cellulose samples. The XRD and TEM results

revealed that the surface modified-MMT exhibited an exfoliated or a highly intercalated

nanoclay structure around neutral pH. The results obtained from the moisture adsorption

and the dispersion analysis showed that the surface modification has made Na-MMT

more hydrophilic compared to that of the unmodified Na-MMT while, the moisture

adsorption of the modified was significantly lower compared to quaternized cellulose.

The ζ-potential values revealed that the CEC of polycation-MMT CEC reversed the

MMT surface charge from negative to positive regardless of quaternized cellulose type

and pH (from 1 to 14). Also according to the results obtained from AFM, the average

height and the equivalent mean roughness values obtained for modified clay compound

is more than that of its quaternized cellulose counterpart. The ζ-potential values

(negative value) also depend on pH and these values decreased above and below pH 6.

The ζ-potential of the CNC3, CNC2 and CNM3 samples in alkaline media decreased

and reached below 10 mV, suggesting possible instability of the suspensions in alkaline

media.

Acknowledgment

Authors are grateful to NSERC Strategic Network—Innovative Green Wood

Fibre Product (Canada),

238

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244

Chapter 9 WVTR, morphology and mechanical properties of

regenerated cotton linter nanocomposites reinforced with pre-exfoliated

Polycation-montmorillonite

Abstract:

Regenerated cellulose nanocomposite films were prepared blending a pre-

exfoliated and/or intercalated poly-cationic modified sodium montmorillonite (PMMT)

with native Cellulose (Cellulose I) using environmental friendly aqueous LiOH/urea

solvent. The XRD, TEM, analyses were carried out to characterize the structure,

morphology of the regenerated cellulose (Cellulose II)/PMMT nanocomposites. The

XRD and TEM patterns revealed an intercalated and/or exfoliated nanocomposite

structure. The presence of PMMT enhanced the water vapor transmission rate (WVTR)

of RCL. Due to the excellent interface interaction between Cellulose II and PMMT was

observed through TEM and XRD patterns with the significant reinforcing effects on

mechanical properties. Tensile strength and Young’s modulus of RCL films were

improved by 55.3% and 100%, respectively. The Cellulose II /PMMT based films

improved the water barrier properties of regenerated cotton linter based films by

decreasing the Water vapor transmission rates (WVTR) by 4 times. In general, the

transparent RCL nanocomposites based film with PMMT possessed fascinating water

vapor barriers and mechanical properties compare to neat regenerated cotton linter. The

results revealed that the PMMT modified with quaternized cellulose (QC) with medium

245

charged density induced better intercalation, water vapor barriers and mechanical

properties.

Keywords: Poly-cation, exfoliation, bionanocoposites, Cellulose II, and WVTR

9.1. Introduction:

Montmorillonite (MMT) and its use as reinforcing, barriers improvement, fire

retardant and thermal stabilizer material in the field of nanotechnology has generated

significant interest now-a-days especially within the Biopolymer community such as

cellulose[1-4], starch[5-7] and chitosan[8-10]. However, there remain several challenges

such as, efficient intercalation of MMT, compatibilization of the nano layer clay with the

biopolymer matrix and development of suitable methods for processing these

nanocomposites, associated with the use of MMT with other polymer matrices [11, 12].

In general, the intercalated and exfoliated structures are two type of

nanocomposites morphology. The physical and chemical properties of polymer

nanocomposites with exfoliated structure are superior to ones with intercalated structure

[13, 14]. However, obtaining fully exfoliated clay nano-layer is extremely challenging in

most polymers matrix through the general blending process. Therefore, extension of

nanoclay dispersion in polymer matrix to reach the exfoliated morphology becomes

significantly intriguing from the scientific point of view [15-17].

In particular, the exfoliated nanocomposite demonstrates immense promise in

providing exceptional barrier properties, due to the presence of the clay layers, which

delays the diffusion of a gas molecule through the formation of complex tortuous paths

[16, 18, 19].

246

The achievement of dominant properties from polymer-clay nanocomposites are

not only through the molecular level dispersion of the clay but also through the strong

interactions between the polymer matrix and the clay layers. Therefore, the most feasible

and reasonable method to maximize the extent of exfoliation/intercalation of MMT is

surface modification of the nano-clays through ion-exchange, and poly-cation exchange

reaction with molecules that have the similar structure as the polymer matrix [18, 20].

This type of surface modification of clay not only requires the clay surface polarity to be

equivalent with the polarity of the polymer but also enlarges the distance between clay

layer platelets, which favors the penetration of polymer into the clay layer gallery.

Cellulose, a linear polysaccharide composed of (1, 4) linked b-D-glucopyranose

units, exhibits a number of desirable properties and has become one of the most

promising renewable polymeric materials. Cellulose fiber is widely used as packaging

materials in form of paper product, due to their favorable properties such as

biodegradability, sustainability, low cost and low environmental impact. However, the

hydrophilic and porous structure of the paper products made of cellulose fiber contributes

to the low water vapor resistance property of the paper-based packaging.

Over the years, several solvent systems for the preparation of regenerated

cellulose (Cellulose II) materials have been developed. Due to environmental concerns

lately, a low toxic solvent system known as aqueous alkali/Urea has been gaining interest

because of its potential to regenerate and chemically modify the cellulose[21, 22].

Aqueous alkali/Urea solvents pose excellent dissolubility of cellulose for a wide range of

molecular weight and crystallinity without significant degradation.

247

Cellulose II films exhibit excellent thermal, mechanical and barrier properties

specially water vapor barriers than the regular paper, cellophane commercialized

cellulose films, Nano-cellulose fibers (NCF) and nano-crystalline cellulose (NCC)

papers [23]. However, improvement of water-vapor barrier properties of the regenerated

cellulose films to be comparable with hydrophobic thermoplastic polymers requires extra

work.

Few authors have already published on regenerated cellulose based

nanocomposite films without obtaining a high extent of exfoliation [24-31]. They assume

that these compounds will have a high affinity, but according to the small-angle X-ray

diffraction results presented, it seems that mainly nano-clay is intercalated, resulting in a

low extent of exfoliation with less effect on water vapor resistance.

It seems despite of the hydrophilic nature of cellulose, not easy to disperse MMT

uniformly in regenerated cellulose matrix. This poor dispersion may be attributed to the

highly negative charges of hydroxyl and C-O- groups of cellulose which, ward off the

anionic surfaces of the silicate layers[32]. Consequently, much attention has been paid to

poly-cationic clay modification in order to improve the compatibility with a wider variety

of the hydrophilic matrices. Many methods have been accomplished for poly-cationic

clay surface modification, such as exchance reaction, poly-silylation, grafting,

polymerization, and use of coupling agent [20, 33-35].

In our previous work, series of novel intercalated poly-cationic modified clays

were prepared through exchange reaction of qurternized cellulose with sodium

montmorillonite.

248

To the best of our knowledge, no study has been done on the characterization of

highly exfoliated cellulose II with our new modified clay has been reported in the

literature. In order to develop cellulose film products with higher water vapor barrier,

mechanical properties and maintain the green features of packaging materials, we

intended to incorporate novel modified clay with quarternized cellulose into the cellulose

film products, fabricated via a solution process with aqueous alkali/Urea solution.


9.2.1. Materials

Three types of cellulose, α-Cellulose, and microcrystalline cellulose (MCC) were

purchased from Sigma-Aldrich and white cotton linter was purchased from Arnold

Grummer® USA. The viscosity-average molecular weights (�̅�ƞ) of celluloses was35.7 ×

104, 3.8 × 104 and, 17.3 × 104 g/mol, respectively, using cadoxen at 25 °C. 3-Chloro-2-

hydroxypropyltrimethylammonium chloride (CHPTAC) aqueous solution (60 wt %) was

purchased from Sigma-Aldrich. Commercially available sodium hydroxide (NaOH 98%)

and Urea (98%) were of regent grade were procured from Alfa aesar, Haverhill, MA,

USA), Sodium-montmorillonite (Cloisite Na+) with a cation exchange capacity (CEC) of

92.6 mequiv. /100 g clay was kindly donated by Southern Clay Products (Gonzalez, TX,

USA). The dialysis tube was purchased from Fisher Scientific (Waltham, MA USA) with

an MW cut-off of 1000 Dalton. All chemicals were used without further purification.

249

9.2.1.‎Quaternization‎of‎α-cellulose and microcrystalline cellulose

A 3 wt % cellulose was dissolved in pre cooled to -12 °C, NaOH/Urea aqueous

(8:12:80 by weight), according to Cai et. al.[36]. Then according to Song et al.[21] a

certain amount of 3-chloro-2-hydroxypropyl trimethyl ammonium chloride (CHMAC)

aqueous solution was used as cationization agents for quaternization of cellulose. Base on

the molar ratio of CHPTAC to anhydroglucose units (AGU) in cellulose 12:1, 6:1, and

3:1 quaternization cellulose samples were labelled as, NM1, NM2, and NM3 for

quaternizated MCC and NC1, NC2, and NC3 for quaternizated α-cellulose, respectively.

9.2.2.‎Synthesis‎of‎quaternized‎celluloses‎polycation‎modified‎montmorillonites‎

(PMMT)

Modification of MMT was carried out by a cation exchange reaction [20, 37-40].

Certain amount of Sodium-MMT was dispersed in 100 mL deionized water at 80 °C in a

250-ml three-neck flask equipped with reflux, and magnetically stirred for 3 hr at room

temperature, followed by sonication for 60 min. Certain amount of QC with the (1:1)

molar ratio of cations to pre-dispersed Na-MMT was dissolved in 50 mL deionized water

at 80 °C.

The QC solution was poured into the slurry dropwise. The pH of the dispersed

mixture was adjusted to approximately 5 using with aqueous HCl and mechanically

stirred at 80 °C for 12 h. The modified Na-MMT was collected by centrifuging a speed of

13000 rpm at 25 °C for 30 min. The slurry was filtered and washed three times with hot

water until the dispersion was free of chloride ions, as determined by an AgNO3 test.

Finally, the slurry was finally freeze-dried with lyophilizer (Christ Alpha 1-2, Osterode

250

am Harz, Germany) to obtain the khaki modified Montmorillonite powder. Modified clay

according to QC samples were labelled as, CNM1, CNM2, CNM3, CNC1, CNC2, and

CNC3, respectively.

9.2.3. Regenerated Cellulose based nanocomposites film preparation

The regenerated cotton linter (Cellulose II) nanocomposites were prepared

through a solution method with the addition of 6 wt% of each modified Na-MMT, which

were labeled as Cellulose II/CNC1, Cellulose II/CNC2, Cellulose II/CNC3, Cellulose

II/C2NC1, Cellulose II/CNM1, Cellulose II/CNM2, Cellulose II/CNM3. Appropriate

amounts of nanoclays were dispersed in pre-prepared LiOH/urea/H2O with a ratio of

5:15:80 (wt%) which was then stirred for 5 h at 10 °C followed by sonication at room

temperature for 1 h. The filtered suspension (through a 0.4 mm Millipore) was pre-cooled

to -15 °C in a refrigerator. A desired amount of cellulose was mixed immediately in pre-

cooled suspension with clay, and stirred vigorously for 10 min at room temperature. The

entrapped air bubbles in the obtained nanocomposite dope was degassed by

centrifugation at 7000 rpm for 15 min at 10 °C and coated on a glass plate, and then

immersed in acetone bath at 0 °C [41] for 15 min to allow regeneration. The Cellulose II

hydrogel sheet was mounted on a special Teflon plate with Magnetic Tape Roll to

prevent shrinkage and was thoroughly washed with running water and finally was dried

at room temperature for a day.

251


Fourier transform infrared (FTIR)

The spectroscopic analysis of regenerated cellulose and regenerated

cellulose/modified clays nanocomposite film was performed in a Spectrum 100 Series

with ATR attachment (PerkinElmer Ltd., Beaconsfield, Bucks; United Kingdom).

Spectra were analyzed using the spectrum software (version 3.02.01). The spectroscopic

analyses for the samples were performed within the spectral region of 650–4000 cm−1

with 32 scans recorded at a 4 cm−1

resolution. Before measurement, all samples were

dried under vacuum at 40 °C for 48 h to remove residual water.

Water‎vapor‎transmission‎rates‎(WVTR’s)

To direct measurements of water vapor transmission rates (WVTR’s) the IGA

permeation cell kit was used in accordance ASTM E 96/E96M-05 standard [20]. First all

samples were pre-conditioned at a present RH using salt solution, Mg (NO3)2 (52% RH),

and at 23 ˚C (Test condition) for three days to reach equilibrium. Then each film was

placed in in a chamber holder. The measurements were carried out at 23 °C and 50±2

RH%, which was achieved using saturated magnesium nitrate and flowing dry nitrogen

gas at a flow rate of 10 mL/min. to keep the outside environment at 0 RH%. The transfer

rates (WVTR’s) were calculated directly from the steady state rate of weight change of

the container, at a specified time interval, and the area of exposed film, as described by

Equation (9-1). Five replicates were made for each sample and the average values were

taken for analysis.

252

timeArea

changeWeightWVTR

. [9.1]

X-ray diffraction (XRD)

The crystallinity of the regenerated cellulose and regenerated nanocomposites films was

carried out in X-ray diffractometer (a Bruker D8 Advance spectrometer) using a rotating

anode generator and a wide-angle power goniometer. The patterns with Cu Ka not

filtered radiation (the weighted average k is 0.15418 nm) at 40 kV and 20 mA were

recorded with a scan rate of 1°/min over the 2θ range 2–40◦.

The crystallinities ( cx ) of native cellulose and the regenerated cellulose based

films and the nanocomposites were estimated from the areas responsible for (110), (1-10)

and (020) planes of diffraction peaks obtained using the Lorentz–Gaussian peak

separation method [42-45] according to the basic equation (9.2).

0

2

0

2

)(

)(

dssIS

dssIS

x

c

c [9.2]

Where s is the magnitude of the reciprocal-lattice vector and is given by s = (2sinh)/k, h

is one-half the scattering angle. k is the X-ray wavelength, I(s) is the intensity of coherent

X-ray scattering from a specimen, Ic(s) is the intensity of coherent X-ray scattering from

the crystalline region in the range 2 = 21–23.

253

Transmission electron microscope (TEM):

Transmission electron microscopy (TEM) analysis was examined using a JEOL

2011 STEM operated at an accelerating voltage of 200 keV to investigate the dispersion

of clay in cellulose matrix. Ultra-thin sections of the samples (approximately 60 nm

thick) were obtained at room temperature. Subsequently, the ultrathin sections were

microtomed from these faces with a diamond knife. The microtomed sections were

collected in a water-filled boat.

Mechanical properties

The tensile strength (σb), modulus (E), and %elongation at break (εb) of the

regenerated cellulose and regenerated cellulose composite films were determined using a

universal tensile testing machine (Instron Model 4465, Norwood, USA) with a cross-head

speed of 10 cm/min according to ASTM D882. The all nanocomposites films were cut

into dumbbell-shaped samples. Prior to testing, all samples were conditioned at a

temperature and humidity of 23±1 °C and 70±5 RH% for 72 h. For each sample a

minimum of 10 dumbbell shaped samples (115.0mm×6.0mm) were tested to determine

the average values, and the average thickness was recorded using a micrometer caliper

prior to the test.

The optical transmittance of the films

The optical transmittance of the films, a useful standard for the compatibility of

composite elements (Krause, 1972), was measured at wavelengths from 200 to 700 nm

using a 700 nm using an ultraviolet (UV)-visible spectrometer (Genesys 10-s, Thermo

Electron Corporation, USA)

254


9.4.1. Characterization of quarterinzed cellulose and quarterinzed celluloses

modified Na-MMT

FTIR spectroscopy is used to confirm the quarterinzation of celluloses and

exchange reaction of quarterinzed celluloses modified Na-MMT clays. Figures 9-2 and 9-

3 display the FTIR spectra of quarterinzed celluloses and quarterinzed celluloses

modified Na-MMT, respectively.

Figure 9-2 shows the FTIR spectra of quarterinzed celluloses in the region of 450-

1900 cm−1

. The new peaks at 1479 and 1418 cm−1

ascribe the scissoring band of -CH3

groups, and C-N stretching vibration of the substituted CHMAC, respectively. Moreover

cellulose C-O-R group peak shifted slightly toward higher values. These peaks suggest

the presence of quarterinzation reaction in the cellulose.

Figure ‎9-1. FT-IR spectra of QC from 1000 to 1750 cm-1 wavenumber

255

On the other hand, extra bands were observed in all modified clays spectra at 469

and 847 cm-1

compared to quarterinzed celluloses, MMC, and Na-MMT spectrums,

which are ascribed to the Si-O-R stretching and vibration of Al-O-R ,respectively. These

peaks may indicate occurrence of a possible mineral–organic reaction or interaction of

quaternized cellulose’s R-N+ group and Al-O or Si-O on the nano-clay.

Figure 9-2. FT-IR spectra of modified clays from

450 950 cm-1 wavenumber

256

9.4.2. XRD analysis of polycation modified Na-MMT (Filler Structure in

Nanocomposites.)

As illustrated in Figure 9-3 the diffraction peak in the XRD pattern of

nanocomposites was narrower and with the first order Bragg's diffraction.

According to our previous work it is obvious the diffractions did not significantly

shift to lower angles as compared to most the modified clay. However, diffraction peak of

clay in Cellulose II/CNM1, Cellulose II/CNM2, and Cellulose II/CNM3 almost remain in

same positions around 2θ=4.2 as shown in Figure 9-4, confirming that intercalation

between modified clay and Cellulose II matrix is not depend upon charge density of

quarenized cellulose were used for modification Na-MMT. This phenomenon as

described by Song ea. al.[21] can be related to lower solubility quarenized cellulose in

alkaline media either lower compatibility of Cellulose II with quaterized-MMT. The

diffractions of clay in Cellulose II/CNCs systems shifted to lower degree in comparison

with Cellulose II/CNMs. On the other hand no diffraction peak in the XRD pattern

suggests that the Cellulose II/CNC3 nanocomposite have an exfoliated morphology.

sin2

wnd [9-3]

257

Figure 9-3. XRD of Cellulose II, Na-MMT, and Cellulose II nanocomposites

9.4.3. TEM morphology analysis of nanocomposites

To obtain the visual observation and dedication of the morphology of

nanocomposite the TEM were used. TEM micrographs affirmed that the modified clays

were distributed uniformly in the cellulose II matrix. As can be seen in Figure 9-3 the

intercalated either partially exfoliated structure of modified nano-clay in cellulose II

matrix were further confirmed by TEM analysis of the nanocomposite films.

Nevertheless, this result from TEM is supported by XRD measurements.

Moreover, at 6 wt% of CNC3 (sample with no diffraction peak in XRD pattern)

nanoclay content, a high level of exfoliation with a random dispersion reflecting excellent

homogeneity in the CNC3 nanoclay dispersion is observed in Figure 9-2 which is in

258

agreement with XRD result. Consequently, Cellulose II/CNC2 sample (Figure 9-4b) can

be labelled in terms of TEM micrograph as the sample with homogenizes and partial

exfoliated morphology. Figures 9-4b and 9-4c illustrate the morphology pattern of

Cellulose II/CNM2 and Cellulose II/CNM1 nanocomposites, respectively. Both

nanocomposites TEM photograph show, the intercalated and, moreover, some

agglomerated structure can also be found in these cases.

Figure 9-4. TEM micrograph of a) Cellulose II/CNC3, b)Cellulose II/CNC2, c) Cellulose

II/CNM3, and d) Cellulose II/CNM1 nanocomposites

259

9.4.4. Crystallinity of Cellulose II based films and its nanocpmposites

The crystallinity values of the Cellulose II and Cellulose/CNCs and Cellulose

II/CNM nanocomposites obtained from XRD experiments and computed according to

equation 9.1 are listed in Table 1. It is clear that the addition of quaternized cellulose

modified clay in Cellulose II matrix accelerated the nucleation of biopolymer crystals

regardless of modified clay types. As listed in Table 1 the crystallinity of nanocomposites

almost double after adding CNC2, CNC3, and CNM3 and reached to 65, 63, 63,

respectively in comparisons with cellulose II based film.

The increment of crystallinity with the addition of modified clay suggests that

quarternized cellulose modified clay act as nucleating centers and a significant preference

for rapid crystallization by establishing a higher level of nucleation density due to either

intercalated or partially exfoliated morphology of modified clay as confirmed by TEM

and XRD. Furthermore, it seems modified clay stimulate the crystallization rate, leading

to the narrower crystallization peak at 2θ≅20 degree from XRD pattern. It is well known

that increment in crystallites may enhance mechanical and barrier properties [46, 47].

Crystalline sites act as clays that restrain the diffusion of gas and fluid especially water

vapor molecule through the polymeric film by creating more tortuous path[48].

260

Table ‎9.1. Experimental results of the cellulose II, cellulose II nanocomposites from XRD patterns

Sample d(001)

d-spacing

(nm)

Crystal plane

(110)

Crystal plane

(020)

Crystallinity χc (%)

Na-MMT 8.70 1.05

Cellulose I - - 14.82 22.64 64

Cellulose II - - 12.48 20.91 37

Cellulose II/CNC1 4.08 2.1 12.34 20.72 46

Cellulose II/CNC2 3.34 2.6 12.29 20.94 65

Cellulose II/CNC3 - - 12.34 20.79 63

Cellulose II/CNM1 4.42 1.99 12.34 20.68 60

Cellulose II/CNM2 4.28 2.07 12.32 20.81 56

Cellulose II/CNM3 4.26 2.08 12.24 20.68 63

9.4.5. Mechanical properties

Tensile strength and Young’s modulus were estimated from the initial slope of the

stress strain curve (Figure 9-5) (between 0.05% and 1% strain) using the linear regression

method. It was shown that (Figure 9-5), as expected, regardless of the nanoclay type, and

surface treatment, the addition of modified MMTs into nanocomposites improved tensile

strength, modulus parameters, and increased the elongation at break compared to native

cellulose II based film. Enhancement in mechanical properties is attributed not only to the

high aspect ratio and high surface area of modified clay but also to the intense interaction

between the modified clay and the structure of cellulose in molecular scale. Also the

drastic interphase interaction between the modified clay and cellulose II matrix can

261

diminish stress concentration area and simplify more uniform stress distribution, leading

to synergistic load transfer to the nanolayer clay gallery

Figure 9-5 reveals, that in general the mechanical properties of the

nanocomposites were impressed by the modified clay nature and the molecular weight of

quaternary ammonium cellulose used for surface treatment of clay.

The tensile strength, Young’s modulus of Cellulose II/CNC nanocomposites are

superior to those of nanocomposites made from CNM modified clay.

It is also observed that with decreasing surface charge density of quaternary

ammonium α-cellulose modified Na-MMT nanocomposites, as presented in Figure 9-5,

the tensile strengths of Cellulose II/CNCs nanocomposites were increased from 86 MPa

for pure cellulose II film up to 197 Mpa, in Cellulose II/CNC3 nanocomposite. The same

phenomenon was observed in Cellulose II/CNMs system and tensile strength increased

47% as compared to pure regenerated cellulose based film.

Furthermore, the addition of modified clays to Cellulose II matrix also led to a

significant increase in Young’s modulus, from 2.08±0.21 GPa for pure Cellulose II to -

5.76±0.45 GPa for Cellulose II/CNC3 nanocomposite at 6 wt% modified clay content. It

seems CNC3 and CNC2 were modified with quaternized cellulose with lowest surface

charge density and higher molecular weight show more improvement in tensile strength

and Young’s modulus of nanocomposites as compared to the rest of quaternized cellulose

with higher surface charge density and lower molecular weight. However, cellulose

II/CNC3 and cellulose II/CNC2 illustrated the highest tensile strength and Young’s

modulus among all the nanocomposites.

262

It seems that the more increment in tensile strengths and Young’s modulus of

Cellulose II/CNC system can be attributed to intensive interfacial bonding between the

CNCs clay and the cellulose II matrix that facilitates the entrance of biopolymer

molecules into the nanolayer gallery to increase the surface area with homogeneous

dispersion of nanoclay in biopolymer matrix. This is a serious property in

nanocomposites as stress concentrations can develop at the nanoclay-matrix interphase

because of the discrepancy in material properties between the clay and matrix. However

the high level dispersion of nanoclay with exfoliated structure in polymer matrix can

reduce stress concentration area and simplify more monotonic stress repartition [49].

The elongation of Cellulose/CNCs nanocomposites diminished sharply regardless

of type of modification, as compared to pristine Cellulose II. While elongation at break of

Cellulose II treated with quaternized MCC modified clay increased slightly. Especially it

should be emphasized; these results are in contrast to the results reported by others for

Cellulose II/Na-MMT nanocomposites [24, 26]. Increment in elongation at break of

nanocomposites comparison with pure Cellulose II may be attributed to plasticizing

effect of quaternized MMC [50].

263

Figure 9-5. Tensile strength, modulus and elongation at break of Cellulose II, A)

Cellulose II/CNC1, B) Cellulose II/CNC2, C) Cellulose II/CNC3, D) Cellulose II/CNM1,

Cellulose II.CNM2, and Cellulose II/CNM3 at various thickness

9.4.5. Water vapor mass transfer properties

The WVTR measurement was aimed to investigate the influence of the type of

clay, order of exfoliation and crystallinity on the WVTR of nanocomposites. The steady-

state water vapor transmission rates of Cellulose II film and its nanocomposites with 20,

40, and 90 µm thickness were measured at a temperature of 23 oC and constant 50%

relative humidity, according to Equation (9-2), as can be seen in Figure 9-6. A general

trend was observed in WVTR value regardless nanoclay type. It was found from WVTR

data that there is a significant reduction in WVTR with adding 6 wt% modified nano-clay

content as compared to pristine Cellulose II films. The effect of nano-clay on the barrier

264

properties of polymers can be quantified by the tortuous path theory due to Nielsen [6].

Hence the sharp reduction in water vapor permeability is due to the effect of the nano-

clay with exfoliated structure. Moreover as discussed in XRD section, addition of

modified clay has the significant effect in increment in crystallization of nanocomposites

in comparison with pure Cellulose II. Polymer crystallinity significantly affects the

permeability. The WVTR value is inversely proportional to degree of crystallinity[51]. In

general, the lower WVTR’s in crystalline regions are attributed to the lower diffusivity of

water molecules. The higher crystalline spot creates a more tortuous path for the transport

of water vapor molecules. Thereby, in this part two factors have been considered to

evaluate the permeability reduction, crystallinity and level of clay platelet exfoliation.

As illustrated in Figure 9-6 lowest WVTR values are found for films obtained

from nanocomposites contain CNC3 and CNC2 which is comparable with Poly lactic

acid (PLA)[14]. This great reduction in WVTR associated with Cellulose II/CNC3 and

Cellulose II/CNC2 nanocomposites samples might be attributed to exfoliated

morphology. These results demonstrate that well dispersed nanoclay with well layered

and high aspect ratio forces water molecules to travel a longer distance across the

cellulose II nanocomposites as it follows a tortuous path around the exfoliated clay.

However, Cellulose/CNM1, Cellulose/CNM2, and Cellulose/CNM3,

Cellulose/CNC1 nanocomposites in spite of similar morphology as concluded from XRD

patterns, illustrate the distinctive water vapor resistance characteristics. The order of

water vapor resistance was found to be Cellulose/CNM3~ Cellulose/CNM1>

Cellulose/CNM2>Cellulose/CNC1 which is directly proportional to their degree of

crystallinity.

265

However, according to WVTR results, crystallinity degree, and level of

exfoliation, in this study it can be concluded that the level of exfoliation plays the most

important role in WVTR value.

Moreover WVTR values were inversely proportional to the thickness of samples

with increasing film thickness from 20 to 90 µm the WVTR of pristine Cellulose II

reduced from 279 to 82 g/m2/day and for the best sample (Cellulose II/CNC3) reduced

from 94 to 19 g/m2/day at RH 50% and 23 °C.

Figure 9-6. WVTR (at RH 50% and 23 °C) of Cellulose II, A) Cellulose II/CNC1, B)

Cellulose II/CNC2, C) Cellulose II/CNC3, D) Cellulose II/CNM1, Cellulose II.CNM2,

and Cellulose II/CNM3 at various thickness

Optical transparency in the wavenumber visible light can provide great potentials

for optical applications. Figure 9-7 demonstrates the optical transmittance of the

266

Cellulose II and Cellulose II nanocomposite films. Figure 9-7 illustrates the well optical

transparency of the nanocomposite film irrespective of modified clay type and remained

similar to Cellulose II film. It seems the modified clay has almost no effect on

transparency of cellulose nanacomposites and the optical transmittance of the films at 600

nm remained between 80 and 90%. This phenomenon can be related to the homogenous

dispersion of nanoclay in cellulose II matrix which causes no optical refraction and

scattering when the nanoclay was added to cellulose II matrix.

Figure 9-7.Transparency of Cellulose II and its nanocomposites film at visible

wavelength light with a thickness of 20 µm

267

9.5. Conclusion

In this work the effects of various surface treatments on morphology of Na-MMT

with different cationic-cellulose have investigated. Cellulose II-based nanocomposites

film with series modified MMT were fabricated through solution blending using

alkaline/urea/water mixture at low temperature.

The effect of QC-MMT clays have been studied on morphology, mechanical, and

barrier properties of Cellulose II. It was shown that, depending on the nano-clay type, and

surface treatment, the addition of modified MMT into nanocomposites improved tensile

strength, modulus parameters, and increased the elongation at break compared to native

cellulose II based film regardless of type of modification.

The type quaternary ammonium cellulose as surface modifiers was the main

factor in level of clay dispersion in cellulose II matrix. QC-MMTs were modified with

quaternary ammonium cellulose with lower quaternary ammonium showed better and

homogeneous dispersion with exfoliated morphology. The exfoliated morphology was

confirmed with XRD and TEM observations. The modified clay with quaternary

ammonium α-cellulose regardless of content of quaternary ammonium enhanced the

mechanical properties of cellulose II nanocomposites more as compared with ones

modified with quaternary ammonium microcrystalline cellulose. The poly-cationic

surface-modified MMT allowed having significant improvements both in crystallinity

and in water vapor barrier properties of Cellulose II nanocomposites.

In particular, the Cellulose II/CNC3 and Cellulose II/CNC2 nanocomposite film

showed a tensile strength and Young’s modulus than rest of the Cellulose II

nanocomposite films. The water vapor permeability of Cellulose II film reduced with

268

addition of QC-MMT up to 4 times at 50% RH and 23 oC. These improvements in

mechanical and water vapor permeability of Cellulose II are attributed to the high aspect

ratio of the nanoclay platelets and exfoliated structure of nanoclay in the nanocomposites.

The general behavior of the showcased Cellulose II-based QC-MMT

nanocomposites illustrates assurance for high-performance materials in packaging

industries especially in dairy, meat and bakery packaging.

Acknowledgment

Authors are grateful to NSERC Strategic Network—Innovative Green Wood Fibre

Product (Canada),

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Chapter 10 Conclusions, and Future Perspectives

The key objective of this dissertation, pursued through comprehensive studies, to

enhance the water vapor resistance properties of cellulose-based green packaging

materials. To this end, various nano-biocomposites used to enhance water vapor barrier

of cellulose-based packaging were systematically developed and thoroughly

characterized. The findings from the current research offer encouraging prospects in

terms of how to substantially increase the water vapor resistance properties of cellulose-

based paper and film through a relatively simple process.

This chapter is comprised of two main sections, the first of which contains a

concise overview of the results attained from the characterization of PLA and PHBV and

their nanocomposites coated on paper substrate.

The second section provides a summary of the results obtained from the

characterization of the regenerated cellulose and regenerated cellulose nanocomposites

based film with Sodium-montmorillonite (Na-MMT) and a series of novel modified

montmorillonite that were prepared via solution blending in aqueous alkaline/urea solvent

at low temperatures.

10.1. Summary of the preparation and characterization of biopolymers

nanocomposites coated paper to enhance water vapor resistance cellulose-based

material in packaging application

276

10.1.1. Summary of the preparation and characterization of PLA nanocomposites

coated on paper

According to the literature review, the water vapor barrier of PLA-based

nanocomposites and clay dispersion are some of the major hindrances to the development

of PLA-clay nanocomposites with improved thermal and barrier properties in packaging

and paper coating applications. To remedy this problem, the current study initially sought

to control the barrier properties of nano-clay dispersion in PLA nanocomposites. In an

effort to improve the barrier properties of paper, polylactic acid (PLA) and their

nanocomposites with two different types of nano-clay and various clay content were

coated on paper via a solution blending process using different coating methods. The

coating processes carried out to control the hydrophilic surface and cover the porous

structure of paper in order to improve the water vapor resistance of paper products. To

achieve this objective, two different nano-clays, Cloisite®30B MMT clay and the

Cloisite® 30B clay modified with POSS (M-CS30B), were used. The central aim of this

experiment was to examine the ability of these nanocomposites to enhance the water

vapor barrier properties of paper. In line with the stated goals of this research, the water

vapor transmission rate (WVTR) was also studied using the microbalance gravimetric

technique (IGA-Hiden apparatus). This technique provided direct measurements of water

vapor transmission rates (WVTRs) on paper or coated paper with PLA nanocomposites

via paper coating using bar coating and hot press coating methods. XRD revealed the

dispersion or phase behavior of the M-CS30B in PLA matrix. The intercalation of clay

with PLA was further demonstrated using TEM. Overall, this experiment revealed that

the hot press coating method is an effective approach to improving the water vapor

277

resistance of coated papers. The paper coated with nanocomposites diminished the

WVTR value by 74% [26 g/(m2 day)], as compared to paper coated with neat PLA.

10.1.2. Summary of preparation and characterization of PHBV nanocomposites and

water vapor barrier properties of nanocomposite-coated paper

Consistent with the literature review of well known biopolymers, neat PHBV

film has demonstrated the lowest water vapor transmission rates (WVTR). Despite such

characteristics, the water vapor transmision rate of PHBV is still considerably higher

than most synthetic thermoplastic polymers like PE, PP, and PET. There are, nontheless,

drawbacks such as low deformability interface and melting temperature close to the

degradation temperature that prohibit its widespread application in industrial settings.

Accordingly, solution intercalation methods were developed for preparing novel

nanocomposites based on the PHBV biodegradable polyester and layered silicates

(specifically two organically modified clays, Cloisite®30B (CS30B), and PHB-g-CS30B

amended with β-butyrolactone were intercalated in the Cloisite®30B clay by a ring-

opening polymerization).

The coating of PHBV on base paper was performed using two different methods.

The more efficient method for lowering WVTR values was utilized for the coating of

PHBV/nanocomposites on the paper.

The findings reveal that the coating method and clay exfoliation were the most

important factors impacting water vapor permeability. The water vapor barrier of the

coated papers improved significantly once the surface of the base substrate was pre-spray

278

coated with PHBV suspension prior to the laminate coating of PHBV film with a hot

press.

The morphology of clay and nanocomposites was investigated using X-ray

diffraction (XRD). Additionally, the state of dispersion and the phase behavior of the clay

in PHBV matrix were observed using a transmission electron microscope (TEM). The

papers coated with exfoliated PHBV/nanocomposites exhibited even lower WVTR

values.

Compared to counterpart nanoclay (CS30B) and pristine PHBV, the PHB-g-

CS30B nanoclay had a greater effect on the water barrier properties of the PHBV, even at

all nanoclay content with exfoliated structure. In contrast nanocomposites based on

PHBV with various CS30B showed either partially exfoliated or intercalated structures of

CS30B in the matrix at low concentrations (less than 3wt%). However, at higher CS30B

loading, an aggregated structure is obtained. WVTR of PHBV nanocomposite based film

with 10 wt% PHB-g-CS30B clay reduced from 67 g/m2/day (at 90 RH % and 38 °C) and

16.7 g/m2/day (at 50 RH % and 23 °C) for neat PHBV to 26.4 g/m

2/day and 6.4 g/m

2/day

respectively, due to the increase in the effective path length for diffusion as revealed by

the XRD and TEM results.

Moreover, the strong hydrophobic nature of PHB-g-CS30B clay nanocomposites

at all clay content displayed a lower WVTR, water adsorption, and lower water contact-

angle. Surface topology of PHBV and PHBV nanocomposites based film AFM images

revealed that with the addition of PHB-g-CS30B, the strength of agglomerated clay was

reduced. In similar vein, results obtained from AFM indicate that the presence of CS30B

at the surface of PHBV increases surface roughness. This is primarily due to its larger

279

particle size as well as increased aggregation at the surface, which in turn leads to higher

water contact angle and a significantly higher WVTR.

Furthermore, from the thermal characterizations of PHBV/CS30B

nanocomposites using DSC, the fact that the crystalline phase is impervious to nano-

clays and does not function as a nucleating agent has been determined. It has also been

established that free, intercalated and aggregated nanoplatelets coexist within the

nanocomposites. This is while in PHBV/PHB-g-CS30B, modified nano-clay acted as a

nucleating agent at low clay content. Compared to their pristine counterparts,

nanocomposites exhibit lower melting temperature in the entire concentration range

irrespective of nanoclay types. Moreover, the glass transition temperature of the PHBV

of the nanocomposite decreases independent of nanoclay concentration and crystallinity

of the system reduced at high nanoclay load. Meanwhile, nanocomposite containing

PHB-g-CS30B demonstrated improved thermal stability and thermal properties in

general at all concentration compared with nanocomposite contained CS30B.

In the frequency sweep experiment, the storage modulus and complex viscosity

increased following the incorporation of the modified clay in PHBV both in the glassy

and the rubbery regions. This procedure is pertinent to most thermoplastic applications

such as coating, film blowing, and extrusion process. For example, the storage modulus

value increases from order of 10 Pa to 103 Pa at the log frequency value of 10 Hz with

the addition of the modified clay to PHBV. On the contrary, CS30B exhibited little or

no effect on the complex viscosity of PHBV. Despite an increase in the loading of

Cloisite®30B, no significant changes were observed in the storage modulus compared to

pristine PHBV.

280

In a nutshell, the current study elucidates the importance of interfacial

interaction and exfoliated structure. It can thus be concluded that the proper surface

modification can play a constructive role by providing an excellent dispersion of

nanoclay in the PHBV matrix by allowing for good interfacial interaction with the

desired properties. Pre-exfoliated nanoclay not only led to a lower melting point and

considerably higher thermal stability for PHBV, it also enhanced the narrow melt

processing window of PHBV with significant growth witnessed in the storage modulus

and complex viscosity. Additionally, pre-intercalated nanoclay and more compatible

clay provided high water vapor barrier for PHBV matrix, comparable to synthetic fossil-

based polymer film (according to Appendix I comparable with PP and LDPE).

10.2. Summary of the preparation and characterization of regenerated cellulose-

based film nanocomposites.

This section proposes the application of a simple yet environmentally friendly

method to fabricate Sodium montmorillonite (Na-MMT)/regenerated cellulose based

films (with two type of cellulose) using an aqueous solution of LiOH/urea as the

processing solvent. TEM and XRD revealed that the Sodium montmorillonite was

intercalated uniformly dispersed along the surface of the regenerated cellulose film. FTIR

measurements also indicated the existence of notably stronger hydrogen bonding

interactions between the Na-MMT and the microcrystalline cellulose matrix compared to

cotton linter. This experiment resulted in substantial improvement in the water vapor

barrier properties of the cellulose nanocomposite films. According to the results, all

regenerated films showed high penetration resistance to oil without any oil penetration

281

being reported. Moreover, following the addition of 5wt% Na-MMT to cotton linter, its

tensile strength and Young's modulus, were enhanced by about 57 and 73%, respectively.

This is considerable improvement compared to the pristine regenerated cotton linter film.

Pre-exfoliated and pH sensitive Polycation MMT (PMMT) were successfully

synthesized from the ion-exchange reaction within the six quaternized cellulose structures

that were previously modified by 3-Chloro-2-hydroxypropyltrimethylammonium

chloride. In line with this, various systems of Poly-cation montmorillonite and cellulose

regenerated based films were thoroughly examined. The interaction achieved between

cellulose II and different Poly-cation -MMT types can be categorized as exceptional. The

well exfoliated morphology achieved in this experiment was confirmed by XRD and

TEM observations. The poly-cationic surface-modified MMT generated significant

improvements with regards to crystallinity and water vapor barrier properties of Cellulose

II nanocomposites in comparison with regenerated cellulose/Na-MMT nanocomposites as

previously reported in Chapter 8. These materials also exhibited excellent mechanical

properties (tensile strength and E-modulus) compared to their counterpart, Cellulose

II/Na-MMT nanocomposite. Cellulose II/ PMMT nanocomposite also showed water

vapor barrier improvement, virtually 30 times greater than commercial copy paper. The

water vapor barrier improvement was shown to be four times higher than blank

regenerated cotton linter based film and twice as much as the regenerated cotton liner/Na-

MMT nanocomposite, similar to PLA film at 50% RH and 23 ˚C. Essentially, the general

behavior of the showcased Cellulose II-based poly-cationic MMT nanocomposites seals

the fate of these biopolymers as high-performance materials in the packaging industry,

particularly in relation to dairy, meat, and bakery packaging.

282

10.3. Recommendations for future research

In the current research, the water vapor barrier properties of various coated papers

with PLA and PHBV and their nanocomposites with montmorillonite clay and

regenerated cellulose nanocomposites with montmorillonite based films and different

degrees of thickness have been evaluated.

First and foremost, this study recommends that researchers embrace the growing

need to further examine transport parameters, namely water vapor diffusion coefficient

and water vapor transmission rates and permeability of cellulose-based nanocomposites

film and exfoliated nano-biocomposites materials from already evaluated experimental

measurements. This involves the characterization of the mechanism of water vapor

transport properties in functional-modified clay with PLA and PHBV nanocomposites

and regenerated cellulose nanocomposites. The molding work on mass transport

properties and kinetics should also be thoroughly investigated to develop a more vivid

understanding of these materials in terms of their contribution to packaging applications.

In keeping with the first part of this study, future research can contribute to

practice by focusing on preparing the melt blending nanocomposite polylactic acid (PLA)

and PHBV reinforced with modified and unmodified preprepared nano-clays in industrial

process. This is aimed at improving the developed method to a point where it can be

applied in commercial settings and result in the widespread industrial use of these

polymers. In this lieu, scrutinizing the bio-degradability of prepared samples can be an

intriguing consideration for future research.

Third: To enhance barrier properties regenerated cellulose based film the multi-

functional systems have to be developed in parallel with nanocomposites film such as

283

surface hydrophobization, layer-by-layer coating and plasma treatment with

fluorocarbon, aluminum oxide, and silica.

Fourth: For future research, it would be interesting to analyze the effect of carbon

nanotube, modified carbon nanotube directly, and lignin on regenerated cellulose

properties. Prospective research in this field can potentially enable scientists to

systematically explore the effect of crystallinity, pore size, and pore size distribution on

the WVTR of regenerated cellulose.

284

Appendices

Appendix-I: Effect of super-hydrophobicity on WVTR of hand-sheet

Figure AI-1. Surface modification of hand-sheet with POSS

285

Table AI.1. WVTR and Contact angle of hydrophobic hand shit

Sample

Contact angle

Coat g/m2

*WVTR

(g/m2/day)

Control sample Hand sheet

(Soft wood)

- - -610

Hydrophobic Hand-sheet

140o

- 420

PLA

85o

27

70

*Temp 23 °C and RH 50%

Figure AI-2 Contact angle of surface modified hand-sheet

286

Appendix-II: Summary of WVTR of various material have been measured using IGA

Table AII.1. Summary of WVTR of various material have been measured using IGA

Sample: Thickness

(µm)

WVTR2

(g/m2/day)

WVTR1

(g/m2/day)

calendered paper (60 g/m2) 90 4150±31 942±12



Writing Paper (80 g/m2) 90 3120±17 660±15

R-Cotton Linter film (50 g/m2) 90 1340±21 88±4

R-Microcrystalline Cellulose Film (50 g/m2) 90 960±14 65±3

R-Microcrystalline Cellulose Film (25 g/m2) 50 - 94±5

R-Microcrystalline Cellulose Film (10 g/m2) 20 - 234±16

NFC nano-paper 40 2700±23 370±12

Cellophane 20 1900±16 320±5

HDPE (20 g/m2) 20 19±1 -

LDPE (19 g/m2) 20 33±1 -

PP (19 g/m2) 20 22±2 -

PHVB/PHB-g-CS30B) (20 g/m2) 16 25±1 7.5±0.5

PET (28 g/m2) 20 103±4 -

PVC (28 g/m2) 20 49±2 -

PHVB (20 g/m2) 16 67±1 15±1

PLA (25 g/m2)1 20

1 528±9

1 57±2

1

R-Cotton liner/modified nanoclay (25 g/m2) 50 - 63±1

R-Cotton liner/modified nanoclay (50 g/m2) 90 564±12 31±0

Aluminum foil - ~0 -

Appendix-III: Primary study on effect of pore size on WVTR

In general, copier paper, and nano cellulose paper such as NFC and NCC have

larger pores (>5- 10 µm). In contrast, the regenerated cellulose film have pore diameter

less than 50 nm and denser structure.[1] It is well known that the large pores control the

WVTR in copier paper that may create direct pathways for water vapor to penetrate

1Test condition: RH 50% @ 23 °C

2Test condition: RH 90% @ 38 °C

287

through paper substrate, a consequence in decreasing water vapor barrier properties

compared to the regenerated cellulose film.

It has been found that the pore size and porosity of the regenerated cellulose film

decreased consistently with increasing cellulose concentration[1, 2] and reduced

regeneration temperature[1]. Therefore, pore size is a key factor that has influence on

WVTR. In addition, the pore size distribution could also impact the water vapor

transmission rate. For ex. narrow pore size distribution exhibits improved barrier

properties with increased resistance to flow of water vapor across the thickness of film at

STP. According to Liu, S. et al [1],RMCC films with varied pore size with slight

modifications were prepared. All samples were regenerated in acetone instead of acid

bath. The filtration velocity method were used to measure the average pore diameter

𝐷𝑚 (μm) and calculated using the following equations 10.1 and 10.2 [3]

𝐷𝑚 = √(2.9 − 1.75𝜀)32ƞ𝑙𝑄

𝜀𝐴∆𝑃 [AIII. 1]

𝜀 =(𝑊𝑤𝑒𝑡 − 𝑊𝑑𝑟𝑦)

𝜌𝑤 × 𝐴 × 𝑙 [AIII. 2]

where ƞ is the absolute viscosity of water (8.9 × 10−4

Pa·s), l is the thickness of the

membrane, Q is the volume of the permeate water per unit time (m3/s), ∆𝑃 is the pressure

difference, A is the effective area of the of the cellulose film (1.21 cm2), ε is the cellulose

membrane porosity, t is the time of measurement, 𝑊𝑤𝑒𝑡 and 𝑊𝑑𝑟𝑦 are the weight of wet

and dry membranes (g), respectively; 𝜌𝑤 is the density of cellulose (g/cm3).

The correlation between different pore size RMCC film and water vapor

transmission rate is shown in Figure 10-1 while pore size were increased, the WVTR of

288

the RMCC films gradually increased with the same behavior for different RMCC film

thickness. The result suggests that WVTR is inversely proportional to the pore size.

Therefore, pore size reduction makes a higher resistance to water vapor pass through the

regenerated cellulose film resulting in a decrease in water vapor transmission rate. Thus,

RMCC film (90 µm) with different pore size showed WVTR in between 92 g.m-2

.day-1

to

58 g.m-2

.day-1

i.e., from highest to lowest pore size at 23 oC and 50 % RH.

Figure AIII-1. WVTR of RMCC at various pore sizes at 23 oC and 50 RH%

References

289

[1] S. Liu, L. Zhang, Effects of polymer concentration and coagulation temperature on

the properties of regenerated cellulose films prepared from LiOH/urea solution,

Cellulose, 16 (2009) 189-198.

[2] V.P. Khare, A.R. Greenberg, S.S. Kelley, H. Pilath, I.J. Roh, J. Tyber, Synthesis and

characterization of dense and porous cellulose films, J. Appl. Polym. Sci., 105 (2007)

1228-1236.

[3] M. Hossein Razzaghi, A. Safekordi, M. Tavakolmoghadam, F. Rekabdar, M.

Hemmati, Morphological and separation performance study of PVDF/CA blend

membranes, J. Membr. Sci., 470 (2014) 547-557.

Curriculum Vitae

Candidate’s full name: Madjid Farmahini Farahani

Universities attended (with dates and degrees obtained):

University of New Brunswick, 2010-2015, PhD, Chemical Engineering

University of Tehran, 2004-2006, M.Sc., Polymer Engineering.

Isfahan University of Technology, 1999-2004, B.Sc., Chemical Engineering–Polymer

Industry.

Publications:

Madjid Farmahini-Farahani, Huning Xiao, and Yi Zhao, “PLA nanocomposites

containinng modified nanoclay with synergistic Barrier to Water Vapour for

coated paper” J Appl Polymer Sci, 2014,131, 40952

Madjid Farmahini-Farahani, Alemayehu H. Bedane, Yuanfeng Pang, Huining

Xiao, Mladen Eić and Felipe Chibante “ Cellulose/nanoclay composite films with

high water vapor resistance and mechanical strength”, Cellulose, 2015, Volume

22, Pages 3941-3953

Madjid Farmahini-Farahani, Huining Xiao, Avik Khan and Yi Zhang

“Preparation and characterization of exfoliated PHBV nanocomposites to enhance

water vapor barriers of calendared paper ”, Ind. Eng. Chem. Res., 2015, 54,

Pages 11277–11284

Madjid Farmahini-Farahani, Avik Khan, Peng Lu, and Huining. Xiao, “Surface

morphological analysis and water vapor barrier of Poly(3-hydroxybutyrate-co-3-

hydroxyvalerate) nano-composite:” Under revision, Applied Clay Science

Madjid Farmahini-Farahani, Avik Khan and, Huining. Xiao “Preparation and

Characterization of surface modified Exfoliated Montmorillonite nanoclay with

quaternized cellulose” Under revision Applied Materials & Interfaces

Alemayehu H Bedane, Mladen Eić, Madjid Farmahini-Farahani, Huining Xiao

“Water vapor transport properties of regenerated cellulose and nanofibrillated

cellulose films” Journal of Membrane Science Volume 493, Pages 46–57

Alemayehu H Bedane, Huining Xiao, Mladen Eić, Madjid Farmahini-Farahani

“Structural and thermodynamic characterization of modified cellulose fiber-based

http://www-scopus-com.proxy.hil.unb.ca/record/display.url?eid=2-s2.0-84920749647&origin=resultslist&sort=plf-f&src=s&st1=paper+coat&sid=F6E80327383609213F49024E87D991A3.mw4ft95QGjz1tIFG9A1uw%3a90&sot=b&sdt=b&sl=17&s=TITLE%28paper+coat%29&relpos=0&relpos=0&citeCnt=0&searchTerm=TITLE%28paper+coat%29

http://www.sciencedirect.com/science/journal/03767388

http://www.sciencedirect.com/science/journal/03767388/493/supp/C

materials and related interactions with water vapor” Applied Surface Science

Volume 351, Pages 725–737

Madjid Farmahini-Farahani, Avik. khan, Alemayehu H Bedane, M. Eic, F.

Chibante and H. Xiao “WVTR and morphology of regenerated cotton linter

nanocomposites reinforced with exfoliated Polycation-montmorillonite” In-

progress

Madjid Farmahini-Farahani, Seyed-Hassan Jafari, Hosein Ali Khonakdar, Ahmad

Yavari, and Amir Melati, “FTIR and thermal analyses of intermolecular

interactions in poly(trimethylene terephthalate)/phenoxy blends” Polym. Eng. &

Sci. 2011, 51, 518–52

Mohammad Tarameshlou, Seyed-Hassan Jafari, Hossein Ali Khonakdar, Afsane

Fakhravar, and Madjid Farmahini-Farahani, “PET-based nanocomposites made

by reactive and remodified clays” Iranian Polymer Journal 2010, 19, 521-529

Amir Mellati, H. Attar, and Madjid Farmahini Farahani, “Microencapsulation of

Saccharomyces cerevisiae using a Novel Sol-Gel Method and Investigate on its

Bioactivity” Asian Journal of Biotechnology 2010, 2, 127-132

Madjid Farmahini-Farahani, Seyed Hassan Jafari, Hossein Ali Khonakdar, Frank

Böhme, Hartmut Komber, Ahmad Yavari, and Mohammad Tarameshlou

“Investigation of exchange reactions and rheological response of reactive blends

of poly (trimethylene terephthalate) and phenoxy resin” Polym. Int. 2008, 57,

612-617

Madjid Farmahini-Farahani, Seyed Hassan Jafari, Hossein Ali Khonakdar, Frank

Böhme, Ahmad Yavari, and Mohammad Tarameshlou “Investigation of thermal

decomposition behavior and kinetic analysis of PTT/Phenoxy Blends” Journal of

Applied polymer science, 2008, 110, 2924–2931.

M. Tarameshlou, S.H. Jafari, H.A. Khonakdar, M. Farmahini-Farahani and S.

Ahmadian “Synthesis of exfoliated polyamide 6, 6/organically modified

montmorillonite nanocomposites by in situ interfacial polymerization” Polymer

Composites 2007, 28, 733- 738

Refereed Conference Publications

http://www.sciencedirect.com/science/journal/01694332

http://www.sciencedirect.com/science/journal/01694332/351/supp/C

https://www.researchgate.net/journal/1026-1265_Iranian_Polymer_Journal

http://scialert.net/jindex.php?issn=1996-0700

Effects of nano biopolymer coils on sand dune stabilization and dust controlling,

4th International Colloids Conference, June 15-18, 2014, Madrid, and Spain.

Characteristics and Water Vapor Barrier Properties of Exfoliated PHBV

Nanocomposites Coated on Paper, 63rd Canadian Chemical Engineering

Conference, October 20-23, 2013, Fredericton, NB, Canada

Effects of Biopolymer Coils on Sand Dune Stabilization and Dust Controlling,

63rd Canadian Chemical Engineering Conference, October 20-23, 2013,

Fredericton, NB, Canada

Microencapsulation of living saccharomyces serevisiae by a new Sol-Gel Method,

SBE's 4th International Conference on Bioengineering and Nanotechnology

(ICBN) July 22-24, 2008, Dublin, Ireland

Theoretical and experimental Investigations of Tensile Properties of COC/POE

and COC/EVA blends; World Polymer Congress – MACRO 2006, 41st

International Symposium on Macromolecules; 16th - 21st July, 2006, Rio de

Janeiro - RJ, Brazil

Seminars, talks and presentations:

WVTR of Fully Exfoliated PHBV Nanocomposite Coated on Paper Semi-Annual

Meeting held at FPInnovations (VANLab), November 13 – 15th, 2012,

Vancouver, Canada

Paper Coated with Clay-Based Nanocomposites of PLA and PHBV for

Synergistic Barrier to Water Vapor Transfer: Green Fiber semi-annual meeting

May 1-3, 2012, NAV Centre Cornwall, Ontario

POSS-modified bentonite/PLA composites for coated paper with low WVTR 2nd

IGWFBN Annual meeting – Aug 30 – Sept 1, 2011, Mont Gabriel, QC

Patents:

Preparation of Poly (ethylene terephthalate / modified Montmorillonite

Nanocomposite via melt mixing” Iranian Patent No. 38204, December 30, 2006

Preparation of Polyamide 6, 6 /Montmorillonite Nanocomposite via Interfacial

Polycondensation” Iranian Patent No. 38205, December 30, 2006

Documents

Bio-nanocomposites for enhancing water vapor barrier of