Natural Gas Sweetening by Ultra-Microporous Polyimides Membranes

Dissertation by

Fahd Ibrahim Alghunaimi

In Partial Fulfillment of the Requirements

For the Degree of

Doctor of Philosophy

King Abdullah University of Science and Technology

Thuwal, Kingdom of Saudi Arabia

May, 2017 2

EXAMINATION COMMITTEE PAGE

The dissertation/thesis of Fahd Ibrahim Alghunaimi is approved by the examination committee.

Committee Chairperson: Ingo Pinnau, Advisor

Committee Members: Klaus-Viktor Peinemann, Yu Han, William J. Koros

3

© May, 2017

Fahd Ibrahim Alghunaimi

All Rights Reserved 4

ABSTRACT

Natural Gas Sweetening by Ultra-Microporous Polyimides Membranes

Fahd Ibrahim Alghunaimi

King Abdullah University of Science and Technology, 2017

Advisor: Ingo Pinnau

Most natural gas fields in Saudi Arabia contain around 10 mol.% . The present technology to remove carbon dioxide is performed by chemical absorption, which has many drawbacks. Alternatively, membrane-based gas separation technology has attracted great interest in recent years due to: (i) simple modular design, (ii) potential cost effectiveness, (iii) ease of scale-up, and (iv) environmental friendliness. The state-of-the- art membrane materials for natural gas sweetening are glassy cellulose acetate and polyimide, which were introduced in the 1980s. In the near future, the kingdom is planning to boost its production of natural gas for power generation and increase the feedstock for new petrochemical plants. Therefore, the kingdom and worldwide market has an urgent need for better membrane materials to remove carbon dioxide from raw natural gas.

The focus of this dissertation was to design new polyimide membrane materials for

CO2/CH4 separation exhibiting high permeability and high selectivity relative to the standard commercial materials tested under realistic mixed-gas feed conditions.

Furthermore, this study provided a fundamental understanding of structure/gas transport property relationships of triptycene-based PIM-polyimides. Optimally designed 5 intrinsically microporous polyimide (PIM-PIs) membranes in this work exhibited drastically increased CO2/CH4 selectivities of up to ~75. In addition, a novel triptycene- based hydroxyl-containing polyimide (TDA1-APAF) showed 5-fold higher permeabilities over benchmark commercial materials such as cellulose acetate.

Furthermore, this polyimide had a N2/CH4 selectivity of 2.3, thereby making it possible to simultaneously treat CO2- and N2-contaminated natural gas. Also, TDA1-APAF showed a CO2 permeability of 21 Barrer under binary 1:1 CO2/CH4 mixed-gas feed with a selectivity of 72 at a partial CO2 of 10 bar which are significantly better than cellulose triacetate. These results suggest that TDA1-APAF polyimide is an excellent candidate membrane material for removal of CO2 and N2 from natural gas. Moreover, based on the collected data for CO2/CH4 mixed-gas experiments from this work and previously published reports, a new mixed-gas 2017 CO2/CH4 permeability/selectivity upper bound curve was initiated to reflect the actual performance including plasticization phenomena at high feed pressure and 10 bar CO2 to simulate the real conditions of the wellhead pressure.

6

ACKNOWLEDGEMENTS

I would like to thank Allah for giving me this opportunity besides the strength and ability to overcome difficulties in successfully completing my journey as a Ph.D. student.

This Ph.D. was a true test of my capabilities as a student and researcher, and one of the more challenging tasks that I have undertaken. This work would not have been possible without the support and guidance of many people whom I would like to thank and acknowledge.

First and foremost, I would like to express my deepest appreciation to my advisor,

Prof. Ingo Pinnau, without his advice, guidance and support this dissertation would not have been possible. His wide knowledge and experience in the membrane field inspired my interest and love for membrane science. I admire his approach of setting high standards but trusting me to find a way to meet them – giving me a lot of space to think and work freely, but providing prompt feedback when it was critical. His guidance has advanced and sharpened my professional skills as an experimentalist, writer and presenter. I sincerely appreciate the time and effort he spent in analyzing and reviewing my data that helped me structure my thoughts and ideas into what turned out to be excellent published manuscripts.

I would like to thank Saudi Aramco to give me this opportunity to pursue my PhD study at KAUST. I would like also to dedicate my special thanks to my closest friend

Nasser Alaslai who has gone to great lengths in supporting me emotionally and morally during my difficult times. He always took an extreme interest in listening to my issues and suggested to me possibilities to overcome them. 7

I would like also to thank Dr. Bader Ghanem for his assistance throughout my research. He was an essential part of a feedback loop between the and synthesis labs, and synthesized all the polymers that are the subject of the studies in this dissertation. I have as well worked on a regular basis with Dr. Eric Litwiller who has been incredibly helpful in sharing his knowledge and expertise in how to use the mechanical tools that helped in building the state-of-the-art permeation apparatus on which I have conducted many of my research experiments. I am also grateful to have collaborated projects with great scientists including Dr. Yingge Wang, and Dr. Xiaohua

Ma. They were always willing to help and answer any of my questions related to chemistry. I would like as well to thank my group members Dr. Mohsin Mukaddam, Dr.

Octavio Salinas, Dr. Federico Pacheco, Zain and Mahmoud who were always available to discuss research problems, lab issues, and assist with certain concepts.

I am deeply honored to have my dissertation reviewed by top scientists Prof. Klaus-

Viktor Peinemann, Prof. Yu Han, and Prof. William J. Koros whose contributions to science have been remarkable.

Finally and most importantly, I would like to dedicate this Ph.D. to my family: brothers, sisters, my wife and my three kids Fajer, Faisal and Nehal. I sincerely thank everyone and hope that although the Ph.D. journey is complete, the friendships, sacrifices and accomplishments will continue in and out of the lab.

Fahd Ibrahim Alghunaimi

May 2017 8

TABLE OF CONTENTS

EXAMINATION COMMITTEE PAGE …………………………………………..……..2 COPYRIGHT PAGE ...... 3

ABSTRACT ...... 4

ACKNOWLEDGEMENTS ...... 6

TABLE OF CONTENTS ...... 8

LIST OF FIGURES ...... 15

LIST OF TABLES ...... 25

Chapter 1. Overview of Natural Gas Separation and Membrane Technology ...... 28

1.1. Natural Gas Overview ...... 28

1.2. Fields and Processes in Saudi Arabia ...... 32

1.3. Plant for CO2 and H2S Removal...... 38

1.4. Separation Technologies ...... 39

1.5. Membrane Technology ...... 41

1.6. Membranes for Gas Separation ...... 42

1.6.1. Membrane Developmental Milestones ...... 42

1.6.2. Potential Advantages ...... 44

1.6.3. Membrane Plant Design for CO2/CH4 Separation ...... 45

1.6.4. Commercial Materials for CO2/CH4 Separation ...... 47

1.6.5. Removal from Natural Gas ...... 51

1.7. Dissertation Goals and Hypotheses ...... 52

1.8. Dissertation Structure ...... 54

1.9. References ...... 59

Chapter 2. Gas Transport Fundamentals and Experimental Techniques ...... 66 9

2.1. Polymeric Membranes for Gas Separation...... 66

2.2. Gas Permeability ...... 69

2.3. Gas Selectivity...... 72

2.4. Gas Sorption in Glassy Polymers ...... 73

2.5. Principles of CO2/CH4 Membrane Separation ...... 76

2.6. Mixed-Gas Non-Idealities Transport ...... 76

2.6.1. Competitive Sorption ...... 77

2.6.2. Plasticization ...... 77

2.7. Experimental Methods and Techniques ...... 78

2.7.1. Evaluation of Microporosity in Polymers ...... 78

2.7.2. Pure-Gas Permeability Measurement System ...... 85

2.7.3. Mixed-Gas Measurement System ...... 88

2.7.4. Thermal Rearrangement Method ...... 91

2.7.5. Thermogravimetric Analysis (TGA)...... 92

2.7.6. Wide-Angle X-Ray Diffraction (WAXD) ...... 94

2.8. References ...... 97

Chapter 3. Materials and Literature Review of Microporous Polymeric Materials ...... 106

3.1. Microporous Materials ...... 106

3.2. Pore Size and Transport Mechanisms in PIMs ...... 109

3.3. Polymers of Intrinsic Microporosity for Gas Separation ...... 111

3.4. Improving Microporous Polyimides for Natural Gas Sweetening ...... 118

3.5. Triptycene-Based Polymers ...... 120

3.6. PIMs Physical Aging ...... 122 10

3.7. Plasticization in PIMs...... 123

3.8. Materials and Nomenclature ...... 125

3.8.1. Typical Synthetic Approach ...... 125

3.8.2. Research Focus ...... 126

3.8.3. Dianhydrides ...... 129

3.8.4. Diamines ...... 131

3.9. References ...... 133

Chapter 4. Gas Permeation and Physical Aging Properties of Iptycene Diamine-Based

Microporous Polyimides1...... 147

4.1. Abstract ...... 147

4.2. Introduction ...... 148

4.3. Experimental ...... 150

4.3.1. Polymer Characterization...... 150

4.3.2. Synthesis of Polyimides ...... 151

4.3.3. Polymer Film Preparation ...... 153

4.3.4. Pure- and Mixed-Gas Permeation Experiments ...... 153

4.4. Results and Discussion ...... 154

4.4.1. Physical Properties and Microstructures of 6FDA-DAT1 and 6FDA-

DAT2………………………………………………………………………………154

4.4.2. Pure-Gas Permeation Properties of 6FDA-DAT1 and 6FDA-DAT2 ...... 157

4.4.3. Pure-and Mixed-Gas CO2/CH4 Permeation Properties after Physical

Aging………………………………………………………………………………160

4.5. Conclusions ...... 164 11

4.6. Reference ...... 166

Chapter 5. Gas Transport Properties and Characterization of Polyimides of Intrinsic

Microporosity Based on Novel Triptycene Dianhydrides and 3,3ʹ-Dimethylnaphthidine

Diamine.1 ...... 171

5.1. Abstract ...... 171

5.2. Introduction ...... 172

5.3. Experimental ...... 175

5.3.1. Characterization Methods ...... 175

5.3.2. Membrane Fabrication ...... 176

5.3.3. Gas Permeation Experiments ...... 176

5.4. Results and Analysis ...... 177

5.4.1. Synthesis and Characterization of the TDA-DMN Polyimides ...... 177

5.4.2. Gas Transport Properties...... 182

5.5. Conclusions ...... 186

5.6. References ...... 187

Chapter 6. Gas Transport Properties and Characterization of Polyimides of Intrinsic

Microporosity Based on Non-Substituted- and Diethyl-Triptycene Dianhydrides and 3,3ʹ-

Dimethylnaphthidine Diamine...... 193

6.1. Abstract ...... 193

6.2. Introduction ...... 194

6.3. Experimental ...... 196

6.3.1. Characterization Methods ...... 196

6.3.2. Membrane Fabrication ...... 197 12

6.3.3. Gas Permeation Experiments ...... 197

6.4. Results and Analysis ...... 198

6.4.1. Synthesis and Characterization of the TDA-DMN Polyimides ...... 198

6.4.2. Gas Transport Properties...... 202

6.5. Conclusions ...... 209

6.6. References ...... 210

Chapter 7. Triptycene Dimethyl-Bridgehead Dianhydride-Based Intrinsically

Microporous Hydroxyl-Functionalized Polyimide for Dual CO2 and N2 Removal in

Natural Gas Upgrading Application1 ...... 213

7.1. Abstract ...... 213

7.2. Introduction ...... 214

7.3. Experimental ...... 218

7.3.1. Materials ...... 218

7.3.2. Polymer Characterization...... 219

7.3.3. Synthesis of TDA1-APAF Polymer ...... 219

7.3.4. Polymer Film Preparation ...... 220

7.3.5. Pure-Gas Permeation Experiments ...... 220

7.3.6. Mixed-Gas Permeation Experiments ...... 221

7.4. Results and Discussion ...... 221

7.4.1. Physical Properties and Microstructure of TDA1-APAF ...... 222

7.4.2. Fresh and Physically Aged Pure-Gas Permeation Properties ...... 225

7.4.3. Fresh and Physically Aged Mixed-Gas Permeation Properties ...... 228

7.5. Conclusions ...... 233 13

7.6. References ...... 235

Chapter 8. Synthesis and Gas Permeation Properties of a Novel Thermally-Rearranged

Polybenzoxazole Made from an Intrinsically Microporous Hydroxyl-Functionalized

Triptycene-Based Polyimide Precursor ...... 243

8.1. Abstract ...... 243

8.2. Introduction ...... 244

8.3. Experimental ...... 248

8.3.1. Materials ...... 248

8.3.2. Polymer Characterization of Pristine and Thermally Rearranged TDA1-

APAF Membranes ...... 248

8.3.3. Synthesis of TDA1-APAF, Polymer Film Preparation and Thermal

Conversion ...... 249

8.3.4. Pure-Gas Permeation Experiments ...... 250

8.3.5. Mixed-Gas Permeation Experiments ...... 251

8.4. Results and Discussion ...... 251

8.4.1. Thermal Rearrangement of TDA1-APAF Polyimide to Polybenzoxazole251

8.4.2. Characterization of Pristine TDA1-APAF and TR 460 Membranes ...... 255

8.4.3. Pure-Gas Permeation Properties of Fresh and Physically Aged TR 460

Membranes ...... 258

8.4.4. Mixed-Gas Permeation Properties of Thermally Rearranged TDA1-APAF

(TR 460) ...... 264

8.5. Conclusions ...... 267

8.6. References ...... 269 14

Chapter 9. Conclusions and Recommendations ...... 277

9.1. Introduction ...... 277

9.2. Dimethyl-Triptycene Building Block...... 277

9.3. Functionalized PIM-Polyimide ...... 278

9.4. Thermal Rearrangement of PIM-Polyimide ...... 280

9.5. Natural Gas Sweetening ...... 281

9.6. CO2/CH4 Mixed-Gas Performance and Proposed 2017 Upper Bound ...... 282

9.7. Recommendations for Future Work ...... 285

9.7.1. Copolymers ...... 285

9.7.2. Thin Films ...... 286

9.7.3. Multi-Component Mixtures ...... 287

9.7.4. Carbon Membranes...... 287

9.8. Publications ...... 288

9.9. References ...... 289

APPENDIX………………………………………………………………………….….292

15

LIST OF FIGURES

Fig. 1.1. Estimated world natural gas consumption in trillion standard cubic feet (TSCF) per year...... 29

Fig. 1.2. Estimated natural gas reserves in trillion cubic feet as of January 2016 [13]. ... 32

Fig. 1.3. Simplified diagram of the master gas system in Saudi Arabia...... 34

Fig. 1.4. Khursaniyah gas plant in Saudi Arabia (during the construction) [16]...... 35

Fig. 1.5. Simplified diagram of raw natural gas treatment and separation processes [17].

...... 35

Fig. 1.6. Simple schematic for an amine plant absorption process...... 39

Fig. 1.7. Statistics for total US energy consumption and energy consumed by separation processes (* a quad is a unit of energy equal to 1015 British thermal units. 1 BTU is about

0.0003 kilowatts-hours) [21]...... 40

Fig. 1.8. Simple schematic of a single-stage membrane separation process...... 41

Fig. 1.9. Scientific milestones of membrane-based gas separation [8]...... 43

Fig. 1.10. Potential membrane separation technology advantages...... 44

Fig. 1.11. One-stage plant design for CO2 removal from natural gas...... 45

Fig. 1.12. Two-stage plant design for CO2 removal from natural gas...... 46

Fig. 1.13. A system using perfluoropolymer membranes for CO2 removal and recovery from natural gas in Texas, USA...... 47

Fig. 1.14. 2008 Pure-gas upper bound correlation for CO2/CH4 separation (typical gas permeability/gas pair selectivity upper bound relationship) [41]...... 48

Fig. 1.15. High-performance membrane versus cellulose acetate for CO2/CH4 separation with two-stage plant design...... 50 16

Fig. 2.1. Typical gas permeation order in rubbery polymeric membranes [8]…………..67

Fig. 2.2. Relative molecular sizes (kinetic diameters) for natural gas components [8]….68

Fig. 2.3. Schematic illustrating the - mechanism……………………...69

Fig. 2.4. Schematic diagram depicting the dual-mode sorption model [25]……………..73

Fig. 2.5. The coefficient, S = c/p, decreases with hole saturation and reaches the asymptotic solubility limit, kD, as the pressure increases……………………………75

Fig. 2.6. Schematic illustrating the plasticization effect on glassy polymeric membranes……………………………………………………………………………….78

Fig. 2.7. Gas device for measuring BET surface area and pore size distribution in porous materials (Micromeritics, ASAP-2020)……………………………………….79

Fig. 2.8. Schematic representing BET gas adsorption technique (ASAP-2020) [48]…...79

Fig. 2.9. Schematic representing the adsorption of gas molecules onto the surface of a sample showing (a) the monolayer adsorption model assumed by the Langmuir theory and (b) multilayer adsorption model assumed by the BET theory [48]………………….81

Fig. 2.10. Schematic representing various adsorption isotherm types [54]……………...83

Fig. 2.11. Schematic design of a constant volume/variable pressure permeation system for pure-gas experiments…………………………………………………………………85

Fig. 2.12. Picture of the pure-gas permeation measurement system (constant volume/variable pressure permeation system)…………………………………………...86

Fig. 2.13. Graphical representation of the time lag technique in the constant- volume/variable pressure method………………………………………………………..88

Fig. 2.14. Schematic design of a constant-volume/variable-pressure permeation system for mixed-gas experiments……………………………………………………………….89 17

Fig. 2.15. Picture of the mixed-gas permeation measurement system (constant- volume/variable-pressure permeation system)…………………………………………..90

Fig. 2.16. Schematic of the Carbolite tubular furnace used for preparation of polybenzoxazole membranes from hydroxyl-functionalized polyimides………………..92

Fig. 2.17. Picture of the Carbolite tubular furnace, analyzer and control box…..92

Fig. 2.18. Picture of a thermogravimetric analyzer (TGA Q5000, TA Instruments)……93

Fig. 2.19. Schematic showing thermogravimetric analyzer (TGA) main components

[66]……………………………………………………………………………………….94

Fig. 2.20. Schematic representing main components of x-ray-diffraction technique……94

Fig. 3.1. Types of promising microporous materials. …………………………………..107

Fig. 3.2. Transport mechanism based on “pore” size...... 110

Fig. 3.3. The assumed microporous structure in non-porous membranes, membranes with isolated porosity, and membranes with interconnected porosity (i.e., PIMs)...... 111

Fig. 3.4. Chemical structure of PTMSP...... 111

Fig. 3.5. Chemical structure of PIM-1...... 112

Fig. 3.6. NLDFT analysis of N2 adsorption isotherms for PTMSP and PIM-1 [25]...... 113

Fig. 3.7. Chemical structure of PIM-PI-8...... 114

Fig. 3.8. Chemical structure of PIM-PI-10...... 114

Fig. 3.9. Chemical structure of 6FDA-TMPD...... 115

Fig. 3.10. Chemical structures of KAUST-PI-1 and KAUST-PI-7...... 116

Fig. 3.11. a) Nitrogen adsorption isotherms measured for PTMSP, PIM-1 and KAUST-

PI-1 at 77 K. b) NLDFT-analyzed pore-size distributions based on carbon slit-pore geometry, showing shifts to ultra-microporosity in KAUST-PI-1 [51]...... 116 18

Fig. 3.12. Typical polyimide synthesis procedure...... 120

Fig. 3.13. Triptycene molecule with high internal free volume [64]...... 121

Fig. 3.14. Correlation of polymer specific volume and before and after Tg

[27, 69]...... 122

Fig. 3.15. Triptycene building units used in this dissertation...... 127

Fig. 3.16. Chemical structures of triptycene and extended iptycene building units showing internal free volume (IFV) [100]...... 128

Fig. 3.17. Key structural features in the proposed novel polyimides...... 128

Fig. 4.1. Schematic of the chemical structure of triptycene and its internal free volume

(IFV)...... 149

Fig. 4.2. Schematic of the chemical structure of the extended iptycene demonstrating enhanced internal free volume (IFV)...... 149

Fig. 4.3. Thermogravimetric analysis (TGA) of 6FDA-DAT1 and 6FDA-DAT2...... 153

Fig. 4.4. Physisorption isotherms for 6FDA–DAT1 and 6FDA–DAT2 using N2 at -196 oC. Open symbols: adsorption; closed symbols: desorption...... 156

Fig. 4.5. Physisorption isotherms for 6FDA–DAT1 and 6FDA–DAT2 using CO2 at 0 °C.

Open symbols: adsorption; closed symbols: desorption...... 156

Fig. 4.6. NLDFT-based estimated pore size distribution obtained from N2 and CO2 isotherms for 6FDA–DAT1 and 6FDA–DAT2 assuming carbon slit-pore geometry. ... 157

Fig. 4.7. Pressure-dependence of pure- and mixed-gas CO2 permeabilities for 6FDA–

DAT1 and 6FDA–DAT2 (50:50 CO2:CH4 mixture; 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas...... 162 19

Fig. 4.8. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for 6FDA–

DAT1 and 6FDA–DAT2 (50:50 CO2:CH4 mixture, 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas...... 163

Fig. 4.9. Pressure-dependence of mixed-gas CO2/CH4 selectivities for 6FDA–DAT1 and

6FDA–DAT2 (50:50 CO2:CH4 mixture, 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas...... 164

Fig. 5.1. Structure of triptycene with indication of internal free volume (IFV); R = alkyl bridgehead substitutions [25-28]………………………………………………..……..173

Fig. 5.2. Structure of polyimides TDA1-DMN and TDAi3-DMN...... 174

Fig. 5.3. FTIR spectra of TDA1-DMN and TDAi3-DMN polyimides...... 177

Fig. 5.4. Thermogravimetric analysis (TGA) of TDA1-DMN and TDAi3-DMN...... 179

Fig. 5.5. Physisorption isotherms using N2 at -196 °C for TDA1-DMN (R=C1) and

TDAi3-DMN (R=i-C3)...... 180

Fig. 5.6. Physisorption isotherms using CO2 at 0 °C for TDA1-DMN (R=C1) and TDAi3-

DMN (R=i-C3)...... 180

Fig. 5.7. NLDFT-based estimated pore size distribution obtained from N2 and CO2 isotherms for TDA1-DMN and TDAi3-DMN assuming carbon slit-pore geometry. .... 181

Fig. 5.8. WAXD patterns of TDA1-DMN and TDAi3-DMN films...... 182

Fig. 5.9. Gas separation performance of TDA-DMN polyimides and PIM-PI-10 for (a)

O2/N2 and (b) H2/N2. Polymers are labeled as indicated in Fig. 3.6 and 5.2. The solid lines represent 2008 and 2015 permeability/selectivity tradeoffs [24, 36]...... 185 20

Fig. 5.10. Gas separation performance of TDA-DMN polyimides and PIM-PI-10 for

CO2/CH4. Polymers are labeled as indicated in Fig. 3.6 and 5.2. The solid lines represent

2008 permeability/selectivity tradeoffs [36]...... 185

Fig. 6.1. Structures for polyimides of intrinsic microporosity: a) non-extended EADA-

DMN (PIM-PI-12) and b) non-extended SBFDA-DMN……………………………….195

Fig. 6.2. Chemical structures of a) non-substituted-triptycene and b) diethyl-triptycene- building blocks...... 195

Fig. 6.3. Chemical structure of TDA0-DMN and TDA2-DMN polyimides...... 196

Fig. 6.4. Thermogravimetric analysis (TGA) of TDA0-DMN and TDA2-DMN...... 199

Fig. 6.5. Physisorption isotherms using N2 at -196 °C for TDA0-DMN, TDA1-DMN

(R=C1) and TDA2-DMN (R=C2)...... 200

Fig. 6.6. NLDFT-based estimated pore size distribution obtained from N2 at -196 °C isotherms (assuming carbon slit-pore geometry) for TDA0-DMN, TDA1-DMN (R=C1) and TDA2-DMN (R=C2)...... 200

Fig. 6.7. WAXD spectra of TDA0-DMN, TDA1-DMN and TDA2-DMN films...... 202

Fig. 6.8. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for CO2/CH4.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 permeability/selectivity tradeoffs [18]...... 205

Fig. 6.9. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for H2/N2.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 and

2015 permeability/selectivity tradeoffs [18, 19]...... 206 21

Fig. 6.10. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for O2/N2.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 and

2015 permeability/selectivity tradeoffs [18, 19]...... 207

Fig. 6.11. Triptycene building units of (a) non-substituted-, (b) extended iptycene, (c) dimethyl-substituted, (d) diethyl-substituted and (e) diisopropyl-substituted...... 208

Fig. 7.1. Structures of ladder PIM-1 and OH-functionalized PIM-PIs: 6FDA-APAF and

PIM-6FDA-OH………………………………………………………………………..217

Fig. 7.2. FTIR spectrum of TDA1-APAF polyimide...... 222

Fig. 7.3. Thermogravimetric analysis of TDA1-APAF polyimide film...... 223

Fig. 7.4. Physisorption isotherms for TDA1-APAF using N2 at -196 °C. Closed symbols: adsorption; open symbols: desorption...... 224

Fig. 7.5. WAXD curves of fresh and 250 days aged TDA1-APAF films...... 225

Fig. 7.6. Pressure-dependence of pure- and mixed-gas CO2 permeabilities for CTA [48], fresh and physically aged TDA1-APAF polyimide (1:1 CO2/CH4 mixture; 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas...... 229

Fig. 7.7. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for CTA [48], fresh and physically aged TDA1-APAF (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas...... 230

Fig. 7.8. Pressure-dependence of mixed-gas CO2/CH4 selectivities for CTA [48], fresh and physically aged TDA1-APAF samples (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye: open points, pure-gas; closed points, mixed-gas...... 231 22

Fig. 7.9. Gas separation performance of TDA1-APAF polyimide for (a) H2/CH4 and (b)

CO2/CH4. Open points: pure-gas feeds; closed points: 1:1 CO2/CH4 mixed-gas feeds.

Polymers are labeled as indicated in Fig. 7.1 and Scheme 7.1. ( ) refers to performance after physical aging in days. The solid lines represent 2008 and 2015 permeability/selectivity trade-offs [14, 25]...... 232

Fig. 7.10. CO2/CH4 mixed-gas permeability/selectivity trade-off curve for TDA1-APAF

(this study), KAUST-PI-1 [48], AO-PIM-1 [53], 6FDA-DAT1 [36], 6FDA-DAT2 [36],

TPDA-APAF [40], 6FDA-DAP [49], 6FDA-APAF [40], 6FDA-mPDA [49], TR6FDA-

HAB [50], TRPIM-6FDA-OH [51] and CTA [48]. All experiments were performed with a 50/50 (v/v) CO2/CH4 mixture at 20 bar feed pressure and 35°C using the constant volume/variable pressure technique...... 233

Fig. 8.1. Structures of PIM-1 and OH-functionalized PIM-PI (TDA1-APAF)………...246

Fig. 8.2. TGA/MS analysis of TDA1-APAF polyimide film...... 252

Fig. 8.3. Time-dependent TGA analysis of TDA1-APAF polyimide films held at isothermal of 400, 430 and 460 °C, respectively...... 253

Fig. 8.4. FTIR spectra for the pristine TDA1-APAF and thermally rearranged TR 460 film...... 256

Fig. 8.5. a) Physisorption isotherms for TDA1-APAF and TR 460 using N2 at -196 °C.

Closed symbols: adsorption; open symbols: desorption. b) NLDFT-based estimated pore size distribution obtained from N2 isotherms for TDA1-APAF and TR 460 assuming carbon slit-pore geometry...... 257

Fig. 8.6. Wide-angle x-ray diffraction spectra of fresh and aged TR 460 film samples. 261 23

Fig. 8.7. Pure-gas O2/N2 separation performance of previously reported APAF- polyimide-based PBOs [29, 30] and TDA1-APAF-derived TR 460 membrane (this study). The solid lines represent 2008 [10] and 2015 [45] permeability/selectivity trade- off curves...... 262

Fig. 8.8. Pressure-dependence of pure- and mixed-gas CO2 permeabilities of pristine

TDA1-APAF [21] and TR 460 (1:1 CO2/CH4 mixture; 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas...... 265

Fig. 8.9. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for pristine

[21] and TR 460 (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas...... 266

Fig. 8.10. Pressure-dependence of mixed-gas CO2/CH4 selectivities of pristine [21] and

TR 460 samples (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye: open points, pure-gas; closed points, mixed-gas...... 267

Fig. 9.1. Triptycene building units of (a) non-substituted-, (b) extended iptycene, (c) dimethyl-substituted, (d) diethyl-substituted and (e) diisopropyl-substituted………….278

Fig. 9.2. Gas separation performance of TDA1-DMN (this study), 6FDA-APAF [1] and

TDA1-APAF (this study) polyimides for CO2/ CH4. The solid lines represent 2008 permeability/selectivity trade-offs [2]...... 280

Fig. 9.3. Gas separation performance of CTA[4], pristine and TR TDA1-APAF (TR 460) membranes (this study) for CO2/CH4. The solid line represent 2008 permeability/selectivity trade-offs [2]...... 281

Fig. 9.4. CO2/CH4 permeability/selectivity trade-off curve for TDA1-APAF (this study),

TR TDA1-APAF (this study), KAUST-PI-1 [4], AO-PIM-1 [5], 6FDA-DAT1 (this 24 study), 6FDA-DAT2 (this study), TPDA-APAF [1], 6FDA-DAP [6], 6FDA-APAF [1],

6FDA-mPDA [6], TR6FDA-HAB [7], TRPIM-6FDA-OH [8] and CTA [4]. All experiments were performed under mixed-gas feed with a 50/50 (v/v) CO2/CH4 mixture at 20 bar feed pressure and 35°C using the constant volume/variable pressure technique.

...... 283

Fig. 9.5. Proposed CO2/CH4 mixed-gas upper bound (solid red line) based on CTA [4],

Matrimid [4], 6FDA-DAT1 (this study), 6FDA-DAT2 (this study), 6FDA-mPDA [6],

6FDA-DAP [6], 6FDA-DAR [6], 6FDA-APAF [1], TPDA-APAF [1], TPDA-ATAF [1],

PIM-1 [5], AO-PIM-1 [5], KAUST-PI-1 [4], KAUST-PI-5 [4], 6FDA-DAM:DABA [9],

TDA1-APAF (this study), TPDA-DAR [10] and TPDA-mPDA [10]. All data have the same conditions: 50/50 (v/v) CO2/CH4 feed mixture at 20 bar feed pressure and 35°C using the constant volume/variable pressure technique. The black dash lines represent

1991 and 2008 CO2/CH4 pure-gas upper bounds [2, 11]...... 284

25

LIST OF TABLES

Table 1.1. Typical composition of raw natural gas...... 30

Table 1.2. USA natural gas pipeline specifications...... 31

Table 1.3. Typical composition of raw natural gas in Saudi Arabia...... 33

Table 1.4. Pure-gas permeabilities and ideal pure-gas selectivities for cellulose acetate and Matrimid...... 49

Table 1.5. Comparison between CA and high-performance membrane for high-pressure mixed-gas CO2/CH4 separation...... 49

Table 2.1. Kinetic diameter for gases used in this research……………………………..71

Table 3.1. Glass transition temperatures for conventional low-free-volume and high-free volume PIMs………………………………………………………………………...….108

Table 3.2. Pure-gas permeabilities and ideal selectivities of various PIMs...... 117

Table 3.3. Nomenclature and chemical structures of 6FDA and PIM-type triptycene- based (TDA) dianhydrides that used in the polyimides synthesis...... 129

Table 3.4. Nomenclature and chemical structures of commercially available diamines that used in the polyimides synthesis...... 131

Table 4.1. Physical properties of 6FDA-DAT1 and 6FDA-DAT2……………………155

Table 4.2. Pure-gas permeabilities and ideal selectivities for 6FDA–DAT1 and 6FDA–

DAT2...... 158

Table 4.3. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for

6FDA-DAT1 and 6FDA-DAT2 based on time-lag method (35 °C; 2 bar)...... 160

Table 5.1. Molecular , thermal stability, and BET surface areas of TDA-DMN- based dimethyl- (R= C1) and diisopropyl- (R= i-C3) substituted PIM-PIs. …………..178 26

Table 5.2. Gas permeabilities and permselectivities for TDA-DMN-based polyimides at

2 bar and 35 °C...... 183

Table 6.1. Molecular weights, thermal stability, and BET surface areas of TDA-DMN-

Based PIM-PIs……………………………………………………………..…………...198

Table 6.2. Gas Permeabilities and permselectivities for TDA-DMN-based polyimides (2 bar and 35 °C) and other DMN-based PIM-PIs...... 203

Table 6.3. Summary for the effect of additional benzene ring and 9,10-dialkyl- substitutions to the triptycene moiety (Chapters 4, 5 and 6)...... 208

Table 7.1. Pure-gas permeabilities and ideal selectivities for fresh and physically aged of

TDA1-APAF membranes (at 2 bar; 35 °C)...... 226

Table 7.2. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for fresh and physically aged TDA1-APAF films...... 227

Table 7.3. Diffusivity selectivities and solubility selectivities of fresh and physically aged TDA1-APAF films (2 bar; 35 °C)...... 228

Table 8.1. Thermal rearrangement conditions and TR conversions for TDA1-APAF film samples in this study……………………………………………………………………254

Table 8.2. Pure-gas permeabilities and ideal selectivities for pristine TDA1-APAF and thermally rearranged TR 460 and TRC 460 PBOs compared with previously reported

APAF-based PBOs...... 259

Table 8.3. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for pristine and thermally rearranged (fresh and aged) TDA1-APAF membranes...... 263

Table 8.4. Diffusivity selectivities and solubility selectivities of pristine and thermally rearranged (fresh and aged) TDA1-APAF films (2 bar; 35 °C)...... 264 27

n Table 9.1. Overview of “Mixed-Gas Upper Bound” line parameters, where Pi = k αij

(i.e., Pi is permeability of i in Barrer, k is the front factor in Barrer, αij is the selectivity for i/j, and n is the slope), for key polymer membrane-based CO2/CH4 separations…..285

Table 9.2. Performance of PIMs near to or defining the new 2017 mixed-gas upper bound for key polymer membrane-based CO2/CH4 Separation...... 285

28

Chapter 1. Overview of Natural Gas Separation and Membrane Technology

This chapter provides an overview of worldwide natural gas consumption, production, future demand, compositions, separation technologies, required specifications, fields and main processes in Saudi Arabia. A summary of the commonly used absorption process (amine plant) technology for natural gas sweetening is included.

Moreover, this chapter reviews the membrane-based gas separation technology including the developmental milestones that headed this technology to commercialization at around

1980. Furthermore, it includes: (i) potential membrane advantages, (ii) membrane plant design, (iii) commercial materials for CO2/CH4 separation, (iv) nitrogen removal from natural gas, (v) dissertation goals and hypotheses and (vi) dissertation structure.

Currently, membrane technology has grown into an annual market of nearly $500 million due to major applications in air separations, separations and natural gas sweetening. Yet, removal of CO2 from natural gas using membranes has some limitations over competing technologies (such as absorption process) which require the development of new membrane materials with intrinsically better gas separation properties than those employed today such as cellulose acetate (CA).

1.1. Natural Gas Overview

Worldwide energy demand is projected to quickly expand caused by the continuing growth in global population. In the 2014 Annual Energy Outlook (AEO) report, total delivered energy consumption in the industrial sector was estimated to increase by 28% from 2012 to 2040 [1, 2]. Much of the growth will reflect natural gas use due to its relatively low carbon footprint, increased thermal efficiency, and cleaner burning benefits 29 compared to other fossil fuels. Natural gas combustion offers ~50% reductions in greenhouse gases over for the same amount of electricity, and it is therefore considered the cleanest-burning on the market today [3]. Natural gas displayed the fastest growing primary energy source with a doubling in worldwide production from

1980 to 2010 [4]. Currently, the consumption of natural gas worldwide has reached more than 100 trillion standard cubic feet per year [5]. It accounts for about 23% of the world’s energy consumption and the International Energy Agency predicted that the demand for natural gas will grow by approximately 44% through 2035 [6]. As shown in Fig. 1.1, which displays the future forecast of world natural gas consumption from 2008 to 2035

[7], the global demand is expected to grow to 169 trillion standard cubic feet per year.

200

150 Electric power 100 Industrial Buildings Transportation 50

0 2008 2015 2020 2025 2030 2035

Fig. 1.1. Estimated world natural gas consumption in trillion standard cubic feet (TSCF) per year.

Natural gas for household consumer usage consists essentially only of .

However, raw natural gas that is delivered from underground wells contains methane, 30 ethane, propane, butane, pentanes, carbon dioxide, hydrogen sulfide, nitrogen, helium and water together with traces of other compounds. Indeed, the raw natural gas changes significantly in composition from source to source; nevertheless, methane is the main component [5]. Table 1.1 shows a typical composition of the raw natural gas.

Table 1.1. Typical composition of raw natural gas.

Component %

CH4 70-90

C2H6 to C4H10 0-20

CO2 0-8

O2 0-0.2

N2 0-5

H2S 0-5

Generally, three types of wells are sources of raw natural gas which are fed to a gas processing plant:

wells - ‘associated gas’

 Condensate wells - 'associated gas'

 Gas wells - 'non-associated gas’ 31

In order to meet the pipeline specifications, all natural gas fields require some pre- treatment in the gas plant to remove the impurities. Table 1.2 shows the required natural gas U.S. pipeline specifications to transport and sell the methane [8].

Table 1.2. USA natural gas pipeline specifications.

Component Specifications

CH4 Balance

CO2 < 2 mol%

H2S < 4 ppm

H2O < 120 ppm

Hydrocarbons (C3+) 950-1050 Btu/scf

Inert gases (N2, He, etc.) < 4 mol%

The main objective of a gas processing plant is the removal of ethane, propane, butanes, pentanes, acid gases, nitrogen and water from methane. The traditional methods used for gas separations involve cryogenic distillation and absorption processes [9]. The associated higher of raw natural gas are typically referred to as 'natural gas liquids' (NGLs), which are very valuable by-products of natural gas processing. These

NGLs, which include ethane, propane, n-butane, iso-butane, and natural gasoline, are 32 sold separately for a variety of different uses, such as providing raw materials for production of petrochemicals and enhancing oil recovery in oil wells [10].

1.2. Fields and Processes in Saudi Arabia

In 1938, oil was discovered in the Kingdom of Saudi Arabia with around 255,000 million barrels of reserves, which accounts for approximately 25% of the world’s total reserves. In 1988, associated and non-associated natural gas reserves were estimated to be around 178 trillion standard cubic feet (SCF) [11]. In 2014, the natural gas reserves in the

Kingdom were estimated at 290.8 trillion SCF [12]. Currently, the Kingdom has total known gas reserves of around 297.6 trillion SCF, which are the fifth largest by county worldwide, as shown in Fig. 1.2.

Fig. 1.2. Estimated natural gas reserves in trillion cubic feet as of January 2016 [13].

In 2015, the gas production average rate per day was 11.6 billion SCF and the kingdom is planning to increase the production to more than 17 billion SCFD by 2020 and over the coming decade to beyond 20 billion SCFD. This will reduce the Kingdom’s 33 reliance on liquid fuel for electricity generation and to power seawater desalination plants. Moreover, greater natural gas volumes provide more feedstock for petrochemical and downstream value-added industries. Furthermore, the increased use of cleaner burning natural gas has environmental benefits in the form of lower emissions than the existing liquid fuels [14]. Lately, Saudi Aramco President and CEO Amin Nasser said:

“The increased gas production will mean more feedstock for industries to expand, and new ones to emerge that will help drive job creation, a key objective of Saudi Vision in

2030” [15].

Most of the produced gas in Saudi Arabia is associated with crude oil that contains significant amount of CO2 (around 10%) as shown in Table 1.3, which shows the typical raw natural gas composition in Saudi Arabia [11]. The present technology to remove

CO2 is by using absorption technology, which will be discussed in more detail in Section

1.3. Interestingly, membrane-based gas separation has many advantages over conventional technologies that could mitigate most of their drawbacks and replace the conventional separation processes.

Table 1.3. Typical composition of raw natural gas in Saudi Arabia.

Gas Associated (%) Non-associated (%)

CH4 62.77 69.01

C2H6 15.07 5.70

C3H8 6.64 2.30 34

C4H10 2.40 1.21

C5+ 1.12 0.90

H2S 2.80 5.02

CO2 9.20 3.46

N2 - 12.40

The Master Gas System (MGS) in Saudi Arabia consists of three main units: (i) gas-oil separating plants (GOSPs), (ii) gas plants and (iii) fractionation plants, as shown in Fig. 1.3 [11].

Fig. 1.3. Simplified diagram of the master gas system in Saudi Arabia. 35

Fig. 1.4 is a photo of a raw natural gas treatments and separations plant in Saudi

Arabia. It includes five main processes to remove various impurities from raw natural gas as shown in Fig. 1.5 and they are as follows:

Fig. 1.4. Khursaniyah gas plant in Saudi Arabia (during the construction) [16].

Fig. 1.5. Simplified diagram of raw natural gas treatment and separation processes [17]. 36

a) Oil and condensate removal

The separation of oil and condensate from natural gas is carried out by reducing the pressure in a three-phase separator. It is a closed vessel where the heavier liquids are separated from gases simply by gravity. For wells that have high-pressure gas along with light crude oil or condensate, specialized equipment is utilized, called the low- temperature separator (LTX). This separator uses pressure differentials to cool wet natural gas and separate oil and condensate. This pressure-temperature relationship can work in reverse to remove gas from liquid oil [18].

b) Sulfur and carbon dioxide removal (sweetening)

Sulfur is present in natural gas primarily as hydrogen sulfide (H2S) and mercaptanes.

Some natural gas wells contain significant amounts of sulfur and carbon dioxide (CO2), which is commonly called 'sour gas'. Natural gas is usually considered sour if the H2S content is higher than 5.7 milligrams per cubic meter. The main process for removing

H2S and CO2 from sour gas is by amine absorption, which is described in more detail in

Section 1.3. At very low , it is also possible to use a different technique which utilizes solid desiccants like iron sponges to remove both the H2S and CO2 [18].

c) Water removal

Removal of the associated water from natural gas involves a relatively simple separation and is usually performed at or near the wellhead. However, the removal of the that exists in natural gas involves a more complex treatment. This treatment consists of 'dehydrating' natural gas that is usually performed by either absorption or 37 adsorption. Absorption is used when a dehydrating agent, such as glycol, is applied to remove the water vapor by dehydration due its chemical affinity for water. On the other hand, adsorption occurs when the water vapor is condensed and collected in two or more adsorption towers on the surface of a solid desiccant, such as activated alumina or granular silica gel [17].

d) Nitrogen removal

When hydrogen sulfide and carbon dioxide are removed to acceptable levels, the treated gas is then sent to a nitrogen rejection unit (NRU) where the gas enters a series of passes through a column and brazed aluminum plates. The nitrogen is separated cryogenically or by using an absorbent solvent. The absorbed methane and hydrocarbons are removed from the solvent by reducing the pressure in multiple gas steps [18].

e) Separation of natural gas liquids

Typically, there are two methods to remove natural gas liquids (NGLs) from natural gas streams, which are complicated to operate, with moving parts, and with high capital and operating costs [10].

1. Absorption Method

In this process, natural gas is passed through an absorption tower that is filled with absorbing oil that has high affinity for NGLs. The solution of absorption oil, propane, butanes, pentanes and heavier hydrocarbons is then heated above the boiling point of the

NGLs but below that of the absorbing oil for separation from methane. 38

2. Cryogenic Expander Process

The main concept of this process is to drop the natural gas temperature to approximately -84.4 °C by using a turbo expander process. The rapid temperature drop condenses ethane and heavier hydrocarbons; however, methane is maintained in gaseous form. After removal of the NGLs, their components are separated by fractionation. The fractionation process is based on the different boiling points of the hydrocarbons and consists mainly of distillation columns which are arranged in the following order:

1) Deethanizer column to separate the ethane from the NGL stream.

2) Depropanizer column to separate propane.

3) Debutanizer column to separate butane.

1.3. Amine Plant for CO2 and H2S Removal

The dominant technology used in natural gas sweetening is a reversible solvent absorption process (amine plant), as shown in Fig. 1.6, where the gas stream is bubbled through a tower containing amine solvent that absorbs acidic gases like CO2 and H2S with little methane loss [5, 19]. The main solution is regenerated by heat to remove the absorbed sulfur and carbon dioxide and then reused to treat more sour gas. Currently, monoethanolamine (MEA) and diethanolamine (DEA) are the two main amine absorbents used in this process. Sulfur can be sold and used if reduced to its elemental form. The elemental sulfur is a bright yellow powder and can often be seen in big storage facilities near gas treatment plants. The process used to recover sulfur is known as Claus process, which utilizes thermal and catalytic reactions to extract the elemental sulfur from the hydrogen sulfide solution [18]. 39

Fig. 1.6. Simple schematic for an amine plant absorption process.

The amine plant is commonly used in natural gas treatment plants, however, it has many drawbacks including: (i) high capital and operational costs, (ii) huge chemical consumption, (iii) frequent maintenance, (iv) large capital equipment area and (v) negative effects on environment [19].

1.4. Separation Technologies

Industrial separation processes account for ~ 45% of all energy used in petroleum refining and chemical plants [20]. Moreover, chemical separations in the USA account for ~50% of the industrial energy use and 10-15% of the nation’s total energy consumption. Most large-scale separation technologies, such as distillation, adsorption or absorption, require heat for separation. It has been estimated that alternative, more energy efficient separation processes, such as membrane-based gas separations, would use 90% less energy than required for distillation, as presented in Fig. 1.7 [21]. 40

Fig. 1.7. Statistics for total US energy consumption and energy consumed by separation processes (* a quad is a unit of energy equal to 1015 British thermal units. 1 BTU is about

0.0003 kilowatts-hours) [21].

Conventional technologies for natural gas separation such as amine absorption, cryogenic distillation and pressure-swing adsorption have been successfully applied to satisfy pipeline specification. However, they are energy intensive unit operations with huge chemical consumption and required large area for capital equipment [22].

Therefore, membrane-based gas separation technology is considered a viable and more energy efficient alternative to the conventional methods used for natural gas processing

[23-26]. Also, it offers many advantages in term of cost, simplicity and size [24, 27], as discussed in Section 1.6.2. 41

1.5. Membrane Technology

A can be defined as a layer that works as a selective barrier between two phases. It is impermeable or semi-permeable to specific solutes or molecules when exposed to the action of a driving . Some components are allowed passage by the membrane from the feed into a permeate stream, whereas others are rejected to the residue stream [28], as shown in Fig.1.8.

Fig. 1.8. Simple schematic of a single-stage membrane separation process.

The use of membranes for separation processes has grown quickly during the last 40 years. Reverse osmosis, nanofiltration, and are examples of large-scale membrane applications [29]. Fig. 1.8 shows a schematic representing a membrane separation process for natural gas processing. The high-pressure feed stream is separated to two outlet streams: (i) low pressure permeate and (ii) high pressure residue (retentate).

In natural gas sweetening, CH4 and higher hydrocarbons are present in the retentate, whereas CO2 is enriched in the permeate [9]. In a dense membrane, CO2 is always more permeable than CH4 because of the difference in the gas condensability and molecular sizes of the gas molecules, as presented in Chapter 2. 42

1.6. Membranes for Gas Separation

1.6.1. Membrane Developmental Milestones

The principles governing the separation of gases through membranes were recognized in the early 19th century with simple observations. In a short communication, Thomas

Graham reported that a bladder isolated in an atmosphere of carbon dioxide inflated by the passage of gas into the bladder, until it burst [30]. Then, he thoroughly studied the transport of gases across various diaphragms, which demonstrated the principles governing the separation of gases by membranes [31]. Later, the first systematic investigations on the ability of a material to separate different gases in a variety of polymers were performed by van Amerongen [32] and Barrer [33]. In particular Barrer addressed concepts of diffusion and solubility, gas flow in capillary systems and metals, and gas flow through polymers in his work [33]. However, researchers encountered serious obstacles in achieving adequately high fluxes across their thick isotropic films.

Reducing the thickness of simple isotropic films increased flux, as Graham also observed in 1866, but often led to pinhole defects that severely compromised selectivity. The first major milestone that made membrane separations attractive for industrial use occurred in the early 1960s when Loeb and Sourirajan discovered a simple procedure to form high- flux asymmetric membranes comprising a thin defect-free separating layer on top of a highly porous support by a phase inversion process [34]. Initially, the membrane was developed for reverse osmosis and the material was cellulose acetate (CA), a cheap and readily available polymer. The original formulations to produce high-performance RO membranes were later modified to their use for gas separation applications, specifically for CO2 removal from natural gas [19, 35, 36]. In 1980, Henis and Tripodi of Monsanto 43 introduced the first commercial Prism hydrogen-separating membrane using asymmetric hollow fibers that contained a thin coating of rubbery to repair any defects in the selective skin layer [37]. Since then, membrane-based gas separation systems have made incredible progress and have gained wider acceptance in a variety of applications. Currently, the use of membranes for gas separation processes is growing at a steady rate and the market in 2020 was estimated to be five times of that in

2000 [8]. Fig. 1.9 displays the important milestones in the history and scientific development of membrane gas separation technology [8].

Fig. 1.9. Scientific milestones of membrane-based gas separation [8]. 44

1.6.2. Potential Advantages

Membrane technology has been applied in recent years for separations such as natural gas upgrading, onsite nitrogen production from air and hydrogen recovery in competition with cryogenic separation, physical adsorption and chemical absorption processes [8, 10]. Membrane technology requires potentially less energy consumption since it does not require any phase change as compared to cryogenic separation. In addition, it is environmentally friendly because it does not consume toxic and corrosive chemicals for separation (for example, mono-ethanolamine (MEA) is used in the chemical absorption process). Hence, membrane gas separation allows for simpler system operation (simple modular design), cost effectiveness and can be accomplished with smaller footprints than chemical absorption. This is particularly important for use in remote applications such as offshore gas-processing platforms [10]. Fig. 1.10 summarizes the potential advantages of membrane technology.

Small Footprint

Low Installation Cost

Environment Friendly

No Phase Change

Simple, Continuous Operation

Fig. 1.10. Potential membrane separation technology advantages. 45

1.6.3. Membrane Plant Design for CO2/CH4 Separation

In general, optimum membrane plant design for CO2 removal from natural gas depends on many factors, such as (i) plant location, (ii) product location, (iii) product purity, (iv) product recovery and (v) total cost. The most commonly used membrane plant configuration for CO2 removal is based on a two-stage design to reduce the CO2 content of the product stream to less than 2%. Baker and Lokhandwala compared one- and two- stage process designs to treat 10 million standard cubic feet per day (MMSCFD) of gas containing 10% CO2 using cellulose acetate (CA) membranes. The simple one-stage design, shown in Fig. 1.11, could achieve the targeted goal of 2% CO2 in the residue product stream, but with unacceptably high CH4 loss of ~11.5%. The two-stage design

(Fig. 1.12) demonstrated that the goal is achievable with minimal CH4 loss around 1.5%; however, an additional recompression stage is required [5].

Fig. 1.11. One-stage plant design for CO2 removal from natural gas. 46

Fig. 1.12. Two-stage plant design for CO2 removal from natural gas.

Membrane Technology and Research, Inc. (MTR) has successfully deployed perfluoropolymer-based membrane systems to treat gas wells with feeds containing >

40% CO2 in Texas, U.S.A. (see Fig. 1.13) [38]. The system has successfully reduced the

CO2 content down to a pipeline-specific value of < 2% with more than 95% hydrocarbon recovery with a feed rate of a range between 1 to 300 MMSCFD. 47

Fig. 1.13. A system using perfluoropolymer membranes for CO2 removal and hydrocarbon recovery from natural gas in Texas, USA.

Hybrid systems have been suggested to reduce the cost of the conventional technologies [8, 26]. These systems consist of a membrane system and conventional amine scrubbing system to maintain the pipeline-specific purity. The gas plant in Mallet,

Texas, USA, which has been in operation since 1994, employed a hybrid system (Cynara

CA membranes + amine absorption unit) to remove CO2 from associated gas [39]. The membrane-based step of the separation process removes 70% of the CO2 while the final step provides the pipeline-specified purity using an amine absorption system. Overall, the process successfully reduced the equipment size and cost by 30%.

1.6.4. Commercial Materials for CO2/CH4 Separation

Current membrane systems are often not considered as feasible process for industrial

CO2/CH4 separation mainly due to a few drawbacks, which can include: (i) relatively low gas separation performance (low gas permeance and selectivity) and (ii) weakness to the harsh environment, e.g. sulfur compounds or suspended solids [40]. As a result, there is a 48 strong incentive to develop high-performance polymeric membrane. The ideal material for gas separation membranes should have the following properties: (i) high permeability,

(ii) high selectivity, (iii) good chemical resistance and (iv) stable long-term separation properties. Unfortunately, an inherent trade-off exists between permeability and selectivity, that is, highly permeable polymers show low selectivity and vice versa.

Robeson developed gas permeability/selectivity upper bound relationships (1991 and

2008) for a variety of gas pairs based on a large database of published pure-gas permeation data [41, 42] (Fig. 1.14).

Fig. 1.14. 2008 Pure-gas upper bound correlation for CO2/CH4 separation (typical gas permeability/gas pair selectivity upper bound relationship) [41].

As displayed previously in Fig. 1.9, the first membrane for CO2/CH4 separation was a dry cellulose acetate (CA) developed in the 1980s. To date, the dominant membrane materials are CA, polyimides (PI) and perfluoropolymers [19]. These polymers offer 49

good thermal and chemical stability with good mechanical properties; however, both CA

[19, 43, 44] and PI [45, 46] membranes provide only moderate selectivity and low

permeability as shown in Table 1.4.

Table 1.4. Pure-gas permeabilities and ideal pure-gas selectivities for cellulose acetate

and Matrimid.

Polymer Pure-gas permeability (Barrer) Ideal selectivity (휶)

H2 N2 O2 CH4 CO2 CO2/ CH4 O2/ N2

Cellulose acetate a 12 0.15 0.82 0.15 4.8 32 5.5

Matrimid [47] b 18 0.32 2.1 0.28 10 36 6.6

a Cellulose acetate (2.45 degree of acetylation), T=35 °C, 1 bar [44, 47].

b T=35 °C, O2, N2, CH4 at 2 bar [48], CO2 at 3.4 bar [49] and H2 at 4.1 bar [50].

Design of novel polymeric materials should target structures that provide an

enhancement in permeability and, more importantly, higher selectivities relative to

standard materials. Recently, Baker published a case study highlighting the potential

benefits that could be offered for the commercially attractive CO2/CH4 separation by

comparing the mixed-gas performance of a cellulose acetate membrane with an improved

high-performance membrane (Table 1.5) [8].

Table 1.5. Comparison between CA and high-performance membrane for high-pressure

mixed-gas CO2/CH4 separation. 50

Polymer (휶) CO2/ CH4 %CH4 loss of feed CO2 permeance (GPU)

Cellulose acetate 15 2.6 50

H. P. membrane 40 0.65 100

When the mixed-gas selectivity of the current industry standard membrane material

CA (CO2/CH4 = 15) was enhanced to 40 with a newly developed high-performance

membrane, the methane loss in the permeate was be reduced by 75% (from 2.6% to

0.65% of the feed). Furthermore, by increasing the CO2 permeance 2-fold, the required

membrane area can be reduced by 40% and the compressor load will be reduced by 35%

(Fig. 1.15) [8].

Fig. 1.15. High-performance membrane versus cellulose acetate for CO2/CH4 separation

with two-stage plant design. 51

The direction towards developing a high performance membrane with high selectivity for CO2/CH4 separation begins with an analysis of membrane material gas transport properties.

1.6.5. Nitrogen Removal from Natural Gas

It has been estimated that about 16% of all natural gas in the United States contains higher nitrogen content (up to 15%) than the allowable maximum value (< 3%) according to pipeline specifications [6, 51]. Also, the non-associated gas wells is Saudi Arabia contain around 12% nitrogen [11]. Removal of nitrogen by conventional separation technologies, such as cryogenic distillation, is extremely cost intensive [52]. One alternative technology is pressure-swing adsorption (PSA) using molecular sieves that preferentially adsorb nitrogen [53]. Excitingly, membrane processes can potentially offer more energy-efficient technology with low capital cost, small footprint, simple operation, and low maintenance, as well as minimal environmental impact [5, 19].

For this separation, glassy N2-selective or rubbery CH4-selective polymers can be used. However, efficient separation of N2 from CH4 is difficult to achieve because N2 is smaller than CH4 where diffusivity selectivity favors the permeation of N2; on the other hand, CH4 is more condensable than N2, which favors CH4/N2 solubility selectivity. As a result of these opposite effects, the overall N2/CH4 selectivities of polymeric membranes are generally low.

In principle, glassy membranes (N2-selective) should be more suitable for N2 removal because rubbery CH4-selective membranes require further recompression stages to recycle the low-pressure permeate which is making the process inefficient and costly 52

® [54]. At present, the best performing glassy polymers to date are Hyflon AD80,

® ® Hyflon AD60, Cytop , and a few 6FDA-based polyimides, which provide N2/CH4 selectivities of 2-3. Theoretically, polymeric gas separation membranes can be applied in many processing steps during upgrading of natural gas, including simultaneous removal of CO2 and N2 [55].

1.7. Dissertation Goals and Hypotheses

Today, natural gas is the fastest growing primary energy source in the world and the demand is expected to increase by almost 40% by 2035. Most raw natural gas sources require sweetening by removal of CO2 and H2S where amine absorption is the dominant separation process. Currently, large-scale sweetening of natural gas by membrane technology is somewhat limited due to its moderate performance compared to other competing technologies. However, membrane-based gas separation technology has attracted great interest in recent years due to: (i) simple modular design, (ii) potential cost effectiveness, (iii) ease of scale-up, and (iv) compact system size and environmental friendliness. The state-of-the-art membrane material for natural gas sweetening is glassy cellulose acetate, which was introduced in the 1980s and is still in commercial use despite very promising new materials developed on the laboratory scale.

A commercially attractive membrane material must be: (i) easily processable into a strong thin film that can be tightly packed into high surface area modules, (ii) chemically stable even with aggressive feed components and, most importantly (iii) characterized by high permeability and high selectivity. Pure-gas testing of novel membrane materials is the most common way to obtain their transport properties. However, in cases with 53

condensable feed gas components such as CO2, the pure-gas experiment neglects the competition between gases for available sorption sites as well as potential plasticization of the polymer induced by highly sorbing feed gas components. Thus, for realistic materials evaluation, mixed-gas experiments are required to simulate feed conditions of industrial streams.

The Kingdom of Saudi Arabia has proven natural gas reserves estimated to be around

297.6 trillion standard cubic feet and 60% of the fields contain about 10% of CO2. The present technology to remove CO2 is performed by chemical absorption, which has many drawbacks including high capital and operational costs, huge chemical consumption, frequent maintenance, large capital equipment area and negative effects on the environment. In the near future, Saudi Arabia is planning to boost its production of natural gas for power generation and increase the feedstock for new petrochemical plants.

This work was focused on natural gas sweetening by membrane technology, specifically CO2/CH4 separation. The kingdom and worldwide market has an urgent need for better membrane materials to remove carbon dioxide from raw natural gas. The objective of the work was to develop novel polyimides, which have high permeability and high selectivity relative to the standard materials tested under realistic mixed-gas feeds. Moreover, systematic structural variations of the polymers were used to identify key structural features to mitigate potential plasticization that may impact their separation performance. In this research different triptycene-based polyimides were investigated for

CO2/CH4 separation. The following hypotheses for this research were proposed: 54

1. Incorporating triptycene and extended iptycene moieties into a polymeric

backbone (as either dianhydride or diamine) will introduce internal free

volume (IFV) and create ultra-microporosity.

2. Replacing a triptycene building unit with an extended iptycene (in dianhydride

or diamine) will increase the internal Brunauer-Emmett-Teller (BET) surface

area of the polymer. This will lead to an increase in permeability and slight

decrease in selectivity.

3. Permeability and selectivity of triptycene-based dianhydrides or iptycene-

based dianhydrides can be controlled by substitution with alkyl groups of

various sizes and shapes while using the same diamine.

4. Permeability and selectivity of triptycene-based dianhydrides or iptycene-

based dianhydrides can be also controlled by varying the diamines.

5. A short interkink distance in the dianhydride will become highly contorted

and highly nonplanar and is likely to induce more inefficient chain packing

with ultra-microporosity.

1.8. Dissertation Structure

This dissertation is organized into nine chapters including this introductory Chapter

1, which has provided basic background information about natural gas separation and membrane technology. The main part of this research was to detail the key observations considered in the hypothesis of strategies for the rational design of triptycene-based PIM-

PIs, functionalized and thermally rearranged PIM-PIs with the goal of enhanced performance for membrane-based CO2/CH4 separation in natural gas sweetening 55 application. In general, all results are discussed with focus on structure/property relationships.

Chapter 2 provides an overview of polymeric materials, gas transport fundamentals and theoretical background to understand all experiments performed and conclusions deduced. A specific literature review is offered to provide related framework for discussions of the experimental results. Moreover, detailed explanations are provided for the experimental techniques, devices, and measurement principles in the investigations of relationships between the chemical structures, physical microstructures and gas transport properties of the polymers. Additionally, the characterization techniques used in this work, such as wide-angle x-ray diffraction (WAXD), thermogravimetric analysis (TGA) and Brunauer-Emmett-Teller (BET) sorption analysis are reviewed. These techniques were employed to obtain a fundamental understanding of the complex physical microstructure of the tested polymers and their unique gas transport properties.

Chapter 3 provides a literature review of PIMs materials and introduces the rationale for the design of novel triptycene-based PIM-polyimide (PIM-PI), functionalization and thermally rearrangement PIM-PIs subsequently discussed. A novel 9,10-dimethyl- triptycene moiety is presented as a structural motif in the polymeric backbone to enhance the gas separation performance.

Chapter 4 compares the gas transport properties of two 6FDA-dianhydride-based polyimides prepared from 2,6-diaminotriptycene (6FDA–DAT1) and its extended iptycene analog (6FDA–DAT2). The additional benzene ring on the extended triptycene moiety in 6FDA–DAT2 increases the free volume over 6FDA–DAT1 and reduces the 56

chain packing efficiency. Also, 6FDA–DAT2 exhibits a 75% increase in CO2 permeability compared to 6FDA–DAT1 coupled with a moderate decrease in CO2/CH4 selectivity. Interestingly, minimal physical aging was observed over 150 days for both polymers and attributed to the high internal free volume (IFV) of the shape-persistent iptycene geometries.

Chapter 5 compares the gas transport properties of two new intrinsically microporous polyimides that were obtained from novel triptycene-based dianhydrides containing dimethyl- or diisopropyl-bridgehead groups with a commercially available highly sterically hindered 3,3ʹ-dimethylnaphthidine (DMN) diamine monomer. The dimethyl bridgehead groups in the triptycene building block provided the DMN-based polyimide (TDA1-DMN) larger surface area than the diisopropyl-based polyimide

(TDAi3-DMN) with greater fraction of ultramicroporosity, as observed from N2 and CO2

NLDFT adsorption analysis and higher gas permeability and selectivity. Wide-angle x- ray diffraction (WAXD) measurements demonstrated that TDA1-DMN and TDAi3-

DMN exhibited a bimodal pore size distribution, where TDA1-DMN showed smaller d- spacing values and broader intensity peaks. Both TDA-DMN-based polyimides showed very high gas permeabilities with moderate selectivities.

Chapter 6 compares the gas transport properties two new intrinsically microporous polyimides that were obtained from novel of non-substituted and diethyl-triptycene-based dianhydrides with a commercially available highly sterically hindered 3,3ʹ- dimethylnaphthidine (DMN) diamine monomer. Chapter 4 showed that the polyimide with non-substituted triptycene building block had better CO2/CH4 separation performance than its extended iptycene analog and Chapter 5 explained that the 57 polyimide with dimethyl bridgehead groups in the triptycene building block provided better CO2 permeability and CO2/CH4 selectivity than the diisopropyl-based polyimide.

This is an extension of the two previous chapters with the conclusion that the polyimide containing 9,10-dimethyl (C1) bridgehead groups in the triptycene building block exhibited the best CO2/CH4 separation performance among others with the non- substituted, extended iptycene, diethyl- or diisopropyl-triptycene-based polyimides.

Chapter 7 summarizes the outstanding gas permeation properties of a high- performance hydroxyl-functionalized PIM-polyimide (TDA1-APAF) prepared from the

9,10-dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) and a commercially available 2,2-bis(3-amino-4-hydroxyphenyl)-hexafluoropropane (APAF) diamine monomer. A freshly prepared sample exhibited excellent gas permeation properties for

CO2/CH4 and H2/CH4. Furthermore, physical aging over 250 days resulted in significantly enhanced CO2/CH4 and H2/CH4 selectivities by 36% and 42%, respectively, with only ~ 25% loss in CO2 and H2 permeability. Aged TDA1-APAF exhibited 5-fold higher pure-gas CO2 permeability and two-fold higher CO2/CH4 permselectivity over conventional dense cellulose triacetate membranes at 2 bar. In addition, TDA1-APAF polyimide had good N2/CH4 selectivity, thereby making it potentially possible to bring natural gas with low, but unacceptable nitrogen content to pipeline specification. These results suggest that intrinsically microporous hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for removal of CO2 from natural gas and hydrogen purification in petrochemical refinery applications.

Chapter 8 shows that the novel triptycene-based PIM-PI (TDA1-APAF) membrane described in Chapter 7 can be thermally treated to enhance gas permeabilities. The 58 conversion of hydroxyl-containing polyimide into polybenzoxazole (PBO) was used by thermal rearrangement (TR) of the aromatic polymer chain with decarboxylation at elevated temperature. TR TDA1-APAF membrane displayed excellent O2 permeability coupled with good O2/N2 selectivity and high CO2 permeability coupled with moderate

CO2/CH4 selectivity. Interestingly, physical aging over 150 days resulted in enhanced

O2/N2 and CO2/CH4 selectivities. These results suggest that thermally-rearranged membranes from hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for CO2 removal from natural gas and air separation applications.

Finally, Chapter 9 addresses in a combining manner the principles assumed in the dissertation for the rational design of triptycene-based PIM-PIs, functionalized and thermally rearranged PIM-PIs exhibiting enhanced performance in natural gas sweetening application of membrane-based gas separation. Specific recommendations are given to: (i) further manipulate the described novel membrane materials in this work by either utilize copolymer or carbon molecular sieve materials and (ii) suggesting additional experiments to simulate the real conditions by using thinner films (e.g., 0.2 – 1

μm) or three components mixture feed tests.

59

1.9. References

[1] S. Yi, X. Ma, I. Pinnau, W.J. Koros, A high-performance hydroxyl-functionalized polymer of intrinsic microporosity for an environmentally attractive membrane-based approach to decontamination of sour natural gas, Journal of Materials Chemistry A, 3

(2015) 22794-22806.

[2] U.E.I. Administration, Annual Energy Outlook, in, US Department of Energy

Washington DC, 2014.

[3] https://www.aga.org/sites/default/files/natural_gas_facts_final.pdf.

[4] Global natural gas production doubled between 1980 and 2010, EIA, 2012.

[5] R.W. Baker, K. Lokhandwala, Natural gas processing with membranes: an overview,

Industrial & Engineering Chemistry Research, 47 (2008) 2109-2121.

[6] J. Kuo, K. Wang, C. Chen, Pros and cons of different nitrogen removal unit (NRU) technology, Journal of Natural Gas Science and Engineering, 7 (2012) 52-59.

[7] http://www.syntropolis.net/knowledgehub/wiki/natural-gas/, in.

[8] R.W. Baker, Future directions of membrane gas separation technology, Industrial &

Engineering Chemistry Research, 41 (2002) 1393-1411.

[9] A. Javaid, Membranes for solubility-based gas separation applications, Chemical

Engineering Journal, 112 (2005) 219-226.

[10] P. Bernardo, E. Drioli, G. Golemme, Membrane gas separation: a review/state of the art, Industrial & Engineering Chemistry Research, 48 (2009) 4638-4663. 60

[11] S.O. Duffuaa, J.A. Alzayer, M.A. Almarhoun, M.A. Alsaleh, A linear-programming model to evaluate gas availability for vital industries in Saudi Arabia, Journal of the

Operational Research Society, 43 (1992) 1035-1045.

[12] http://www.eia.gov/countries/country-data.cfm?fips=SA, in.

[13]https://emergingequity.org/2015/08/06/russias-energy-profile-analysis-of-the-worlds- largest-crude-oil-producer/.

[14]http://www.saudiaramco.com/en/home/our-business/sustaining-excellence/gas- development.html.

[15]http://www.saudiaramco.com/en/home/news-media/news/fadhili-gas-program- celebration.html.

[16]http://www.offshore-technology.com/contractors/design-engineering- construction/consolidated/consolidated2.html.

[17] S. Mokhatab, W.A. Poe, Handbook of natural gas transmission and processing, Gulf

Professional Publishing, 2012.

[18] http://www.naturalgas.org/naturalgas/processing_ng.asp.

[19] C.A. Scholes, G.W. Stevens, S.E. Kentish, Membrane gas separation applications in natural gas processing, Fuel, 96 (2012) 15-28.

[20]http://science.uwaterloo.ca/~mauriced/earth691duss/CO2_General%20CO2%20Sequ estration%20materilas/CO2_SeparationIssues_Issue55IETSArticle.pdf. 61

[21] D.S. Sholl, R.P. Lively, Seven chemical separations to change the world, Nature,

532 (2016) 435-437.

[22] A. Rojey, C. Jaffret, Natural gas: production, processing, transport, Editions

Technip, 1997.

[23] B.D. Bhide, S.A. Stern, Membrane processes for the removal of acid gases from natural-gas 2. effects of operating-conditions, economic-parameters, and membrane- properties, Journal of Membrane Science, 81 (1993) 239-252.

[24] A. Lee, H. Feldkirchner, S. Stern, A. Houde, J. Gamez, H. Meyer, Field tests of membrane modules for the separation of carbon dioxide from low-quality natural gas,

Gas Separation & Purification, 9 (1995) 35-43.

[25] F. Fournie, J. Agostini, Permeation membranes can efficiently replace conventional gas treatment processes, Journal of Petroleum Technology, 39 (1987) 707-712.

[26] R. McKee, M. Changela, G. Reading, CO2 removal-membrane plus amine,

Hydrocarbon Processing, 70 (1991) 63-65.

[27] Z.Y. Yeo, T.L. Chew, P.W. Zhu, A.R. Mohamed, S.-P. Chai, Conventional processes and membrane technology for carbon dioxide removal from natural gas: a review, Journal of Natural Gas Chemistry, 21 (2012) 282-298.

[28] R.W. Baker, Membrane Technology and Applications, John Wiley & Sons, Ltd.,

2004.

[29] M. Mulder, Basic Principles of Membrane Technology Second Edition, Kluwer

Academic Pub, 1996. 62

[30] T. Graham, Notice of the singular inflation of a bladder, Journal of Membrane

Science, 100 (1995) 9.

[31] T. Graham, On the absorption and dialytic separation of gases by colloid septa,

Philosophical Transactions of the Royal Society of London, 156 (1866) 399-439.

[32] G. Van Amerongen, Influence of structure of elastomers on their permeability to gases, Journal of Polymer Science, 5 (1950) 307-332.

[33] R.M. Barrer, Diffusion in and through Solids, CUP Archive, 1941.

[34] W.J. Koros, I. Pinnau, Membrane formation for gas separation processes, CRC

Press: Boca Raton, FL, 1994.

[35] S.E. Kentish, C.A. Scholes, G.W. Stevens, Carbon dioxide separation through polymeric membrane systems for flue gas applications, Recent Patents on Chemical

Engineering, 1 (2008) 52-66.

[36] L.S. White, Evolution of natural gas treatment with membrane systems, Membrane

Gas Separation, (2010) 313-332.

[37] J.M. Henis, M.K. Tripodi, A novel approach to gas separations using composite hollow fiber membranes, Separation Science and Technology, 15 (1980) 1059-1068.

[38] http://www.mtrinc.com/co2_removal.html.

[39] G. Blizzard, D. Parro, K. Hornback, Mallet gas processing facility uses membranes to efficiently separate CO2, Oil & Gas Journal, 103 (2005) 48-53. 63

[40] E. Drioli, G. Barbieri, Membrane engineering for the treatment of gases, Royal

Society Chemistry, 2011.

[41] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400.

[42] L.M. Robeson, Correlation of separation factor versus permeability for polymeric membranes, Journal of Membrane Science, 62 (1991) 165-185.

[43] M.D. Donohue, B.S. Minhas, S.Y. Lee, Permeation behavior of carbon-dioxide methane mixtures in cellulose-acetate membranes, Journal of Membrane Science, 42

(1989) 197-214.

[44] A.C. Puleo, D.R. Paul, S.S. Kelley, The effect of degree of acetylation on gas sorption and transport behavior in cellulose-acetate, Journal of Membrane Science, 47

(1989) 301-332.

[45] A. Bos, I. Pünt, M. Wessling, H. Strathmann, Plasticization-resistant glassy polyimide membranes for for CO2/CO4 separations, Separation and Purification

Technology, 14 (1998) 27-39.

[46] A. Bos, I. Pünt, M. Wessling, H. Strathmann, CO2-induced plasticization phenomena in glassy polymers, Journal of Membrane Science, 155 (1999) 67-78.

[47] D.F. Sanders, Z.P. Smith, R. Guo, L.M. Robeson, J.E. McGrath, D.R. Paul, B.D.

Freeman, Energy-efficient polymeric gas separation membranes for a sustainable future: a review, Polymer, 54 (2013) 4729-4761. 64

[48] Y. Huang, D.R. Paul, Effect of film thickness on the gas-permeation characteristics of glassy polymer membranes, Industrial & Engineering Chemistry Research, 46 (2007)

2342-2347.

[49] D.Q. Vu, W.J. Koros, S.J. Miller, Effect of condensable impurity in CO2/CH4 gas feeds on performance of mixed matrix membranes using carbon molecular sieves, Journal of Membrane Science, 221 (2003) 233-239.

[50] Y. Zhang, I.H. Musselman, J.P. Ferraris, K.J. Balkus Jr, Gas permeability properties of Matrimid® membranes containing the metal-organic framework Cu–BPY–HFS,

Journal of Membrane Science, 313 (2008) 170-181.

[51] T. Rufford, S. Smart, G. Watson, B. Graham, J. Boxall, J.D. Da Costa, E. May, The removal of CO2 and N2 from natural gas: a review of conventional and emerging process technologies, Journal of Petroleum Science and Engineering, 94 (2012) 123-154.

[52] B. Ohs, J. Lohaus, M. Wessling, Optimization of membrane based nitrogen removal from natural gas, Journal of Membrane Science, 498 (2016) 291-301.

[53] K.A. Lokhandwala, I. Pinnau, Z. He, K.D. Amo, A.R. DaCosta, J.G. Wijmans, R.W.

Baker, Membrane separation of nitrogen from natural gas: a case study from membrane synthesis to commercial deployment, Journal of Membrane Science, 346 (2010) 270-279.

[54] R.W. Baker, K.A. Lokhandwala, I. Pinnau, S. Segelke, Inventors; Membrane

Technology, Research, Inc., Assignee. Methane/nitrogen separation process. United

States patent US 5,669,958 (1997). 65

[55] T. Kim, W. Koros, G. Husk, K. O'Brien, “Reverse permselectivity” of N2 over CH4 in aromatic polyimides, Journal of Applied Polymer Science, 34 (1987) 1767-1771.

66

Chapter 2. Gas Transport Fundamentals and Experimental Techniques

This chapter explains the theoretical background for fundamental concepts related to gas transport across amorphous polymer membranes. Furthermore, it describes the experimental methods used to measure the gas sorption, pure- and mixed-gas permeation properties. Also, this chapter briefly introduces thermal rearrangement (TR) treatment, polymer characterization techniques using thermogravimetric analysis (TGA) and wide- angle x-ray diffraction (WAXD).

2.1. Polymeric Membranes for Gas Separation

Polymeric membranes have always been the main focus for industrial commercialization because of their cost-effectiveness, solution processability [1], adequate separation properties and ease of membrane manufacturing [2]. Similar to other membrane separation applications, such as water desalination, the use of polymeric membranes has also significantly increased in the gas separation field [3]. However, it is worth noting that despite the development of many novel high-performance polymers, only a few have achieved commercial success for gas separations. For example, in 2002, only about ten polymeric membranes were predominant in 90% of gas separation applications [4]. The low implementation of new membrane types is due to the highly demanding industrial criteria for membrane fabrication and properties [4, 5]. Ideal candidate polymeric materials for gas separation membranes should have the following properties:

1) The polymer must be solution processable. 67

2) Thin-film composite or asymmetric membranes must be fabricated with a defect-free

selective layer (~0.1-1 μm).

3) Both excellent gas permeability and high selectivity for the targeted separation.

4) Stable long-term performance over a period of 3 to 5 years.

Polymers in the rubbery state are above their glass transition temperature (Tg) and have high polymer chain mobility [6]. Accordingly, the polymer is elastic and the effect of molecular size of the permeating gases on relative mobility is relatively small.

Polydimethylsiloxane (PDMS) is an example of a rubbery polymer, which has high gas permeability due to high diffusion coefficients and its selectivity is controlled by differences in the gas solubility as indicated by gas condensability (Fig. 2.1) [7].

Fig. 2.1. Typical gas permeation order in rubbery polymeric membranes [8].

Because of their flexible nature, rubbery polymers exhibit low to moderate selectivity for CO2 removal from methane. Consequently, rubbery polymers are not utilized for acid gas treatment in natural gas separation applications [9].

When a polymer is kept below its Tg, the polymer will be in its glassy state. Glassy polymers are tough and rigid; therefore, they are often much less permeable to gases than 68 rubbery polymers, but their selectivity tends to be higher because of their better size- sieving capabilities [6, 8, 9]; therefore, gas separation is often dominated based on differences in molecular size, as displayed in Fig. 2.2. Typically, industrial membranes use glassy polymers for natural gas sweetening due to their high gas selectivity and good mechanical properties [7].

Fig. 2.2. Relative molecular sizes (kinetic diameters) for natural gas components [8].

As shown in Fig. 2.1 and 2.2, water is a much smaller and more condensable molecule than methane; therefore, it can easily be separated by both rubbery and glassy polymers. Nitrogen can be separated from methane by both rubbery and glassy polymers; however, in the case of glassy polymers, as N2 has smaller molecular size than methane it will be delivered as high-pressure product on the residue-side of the membrane system.

Oppositely, rubbery polymers are methane-selective, that is, they are more permeable for methane over nitrogen because methane is more condensable than nitrogen so the methane-rich permeate needs to be re-compressed to pipeline pressure which will add more cost to the separation process [8]. Hence, glassy polymers are more suitable for

N2/CH4 separation.

Generally, the selectivity is a function of the differences in diffusivity and solubility coefficients for two gases. Commonly, most gas separation membrane systems are based 69 on diffusivity selectivity; however, specific industrial and environmental applications for feed mixtures containing valuable organic vapors are driven by solubility selectivity [10].

2.2. Gas Permeability

Permeation of gases through a non-porous membrane (dense polymeric materials) is generally described by the solution-diffusion model [11-13]. According to this model, gas transport occurs in three successive steps: (i) sorption of the gas molecules on the upstream side of the membrane (high pressure side), then, (ii) diffusion across the membrane down a gradient (across thickness l), and finally (iii) desorption from the downstream side (low pressure side) [14], as presented in Fig. 2.3. Diffusion can be defined as the net transport of molecules from a higher concentration region to lower concentration region by random motion [15].

Fig. 2.3. Schematic illustrating the solution-diffusion mechanism.

The steady-state gas permeability through a membrane of thickness l is defined by

푁 푙 푃 = (2.1) 푝푢푝 − 푝푑표푤푛 70 where P is the gas permeability coefficient (cm3 (STP) cm/cm2 s cmHg), N is the steady-

3 2 state gas flux (cm (STP)/cm s) through the membrane, and pup and pdown are the upstream and downstream at the membrane interface, respectively.

The flux through the membrane is given by the expression [16]:

푑퐶 푁 = −퐷 (2.2) 푑푥 where D (cm2/s) is the effective diffusion coefficient in the polymer and C (cm3

푑퐶 (STP)/cm3 (polymer)) is the penetrant concentration in the membrane and is the 푑푥 concentration gradient of the gas across the membrane. Combining equations 2.1 and 2.2 and integrating across the membrane thickness gives:

퐶푢푝 − 퐶푑표푤푛 푃 = 퐷푒푓푓 (2.3) 푝푢푝 − 푝푑표푤푛

where 퐷푒푓푓 is the concentration-averaged effective diffusion coefficient and 퐶푢푝 and

퐶푑표푤푛 are the penetrant concentrations at the upstream and downstream side of the membrane, respectively. When the downstream pressure is much lower than the upstream pressure, the gas concentration in the membrane can be expressed as follows:

퐶 = 푆 . 푝 (2.4) where S is the solubility constant (cm3 (STP)/cm3 (polymer) cmHg) and p is gas pressure

(cmHg). Therefore, equation 2.3 can be simplified to:

퐶푢푝 푃 = 퐷푒푓푓 = 퐷푒푓푓 . 푆 (2.5) 푝푢푝 71

In polymer materials, gas permeation is mainly effected by a thermodynamic factor

(solubility factor) and kinetic factor (size-based factor) where gas diffusion coefficients

(D) increase when molecular size decreases and gas solubility coefficients (S) increase as condensability increases [8]. Therefore, the gas molecule size and condensability are dominating the permeation properties for a given membrane material. The permeability coefficient is usually expressed in Barrers, where 1 Barrer = 1×10-10 cm3(STP) cm/cm2 s cmHg. Table 2.1 shows the Breck kinetic diameters for six gases, which were used to study gas permeation properties in this research.

Table 2.1. Kinetic diameter for gases used in this research.

Gas Breck kinetic diameter, Å [17]

He 2.60

H2 2.89

CO2 3.30

O2 3.46

N2 3.64

CH4 3.80 72

2.3. Gas Selectivity

Selectivity is a measure of a membrane’s ability to separate two gases i and j, which is the ratio of their permeabilities, and is a measure of its effectiveness to achieve a high- purity product [13, 16, 18, 19]:

푖 푃푖 퐷푖 푆푖 훼푗 = = 푥 (2.6) 푃푗 퐷푗 푆푗

where 푃푖 and 푃푗 are the pure-gas permeability coefficients of gases i and j, respectively;

퐷 푖 is the diffusivity selectivity, defined as the ratio of the diffusion coefficients of gases i 퐷푗

푆 and j; and 푖 is the solubility selectivity, defined as the ratio of the solubility coefficients 푆푗 of the two gases. The mobility or diffusion selectivity is related to the ratio of the molecular sizes of the two gases and the sorption selectivity is related to their relative condensability. The diffusivity selectivity describes the size-sieving capability of the membrane and increases as the size difference between the penetrants increases. In contrast, the solubility selectivity describes the relative condensability of the penetrants and their interaction with the polymer and increases as the condensability difference between the penetrants increases. The sorption selectivity of non-interacting gases is similar for all amorphous polymers; however, the diffusion selectivity can vary significantly depending on the physical state of the polymer (rubbery or glassy state) [8].

It is worth noting that larger penetrants have lower diffusion coefficients and higher condensabilities. Thus, a trade-off exists between the two selectivity terms, which often oppose each other to reduce the overall selectivity. Generally, diffusivity selectivity dominates gas separation in glassy polymers because of their rigid structure. Meanwhile, 73 solubility selectivity dominates gas separation in rubbery polymers because of their flexible chain microstructure. The research reported in this dissertation was focused on glassy polyimides of microporosity (PIM-PIs).

2.4. Gas Sorption in Glassy Polymers

Gas sorption in glassy polymers differs significantly from that in rubbery polymers

[20]. The gas sorption in glassy polymers is expressed by the combination of Henry’s law and Langmuir expressions, which is called “dual-mode sorption theory” [21-24].

Fig. 2.4. Schematic diagram depicting the dual-mode sorption model [25].

As shown in Fig. 2.4, there are two different types of adsorption sites: (i) Henry type

(as observed in rubbery polymers) and (ii) Langmuir type adsorption sites (inner surfaces of microporous materials). Thus, the total quantity of sorbed gas at any given pressure equals the total sum of sorption in Henry’s and Langmuir modes. Henry’s law [25, 26] relates the gas concentration to the pressure linearly, as shown in equation 2.7:

퐶퐷 (푝) = 푘퐷 푝 (2.7)

where CD is assumed to be associated with the dense region of the polymer and 푘퐷 is the 74 solubility coefficient of the gas in the polymer.

The Langmuir isotherm, CH, describes the pore filling of micropores with a saturation

′ capacity 퐶 퐻 and the affinity constant b as follows:

퐶′ 푏푝 퐶 (푝) = 퐻 (2.8) 퐻 1 + 푏푝

Therefore, the pressure-dependent concentration for a glassy polymer is equal to the sum of Henry mode and Langmuir mode [23-25].

퐶′ 푏푝 퐶 (푝) = 퐶 + 퐶 = 푘 푝 + 퐻 (2.9) 퐷 퐻 퐷 1 + 푏푝

Obviously from Fig. 2.4, the dual-mode sorption at low pressure follows the

Langmuir mode (concave behavior), whereas at higher pressures the Henry mode becomes dominant (linear). Basically, the initial gas sorption in the glassy polymer occurs in the non-equilibrium excess free-volume portion of the polymer matrix, generally referred to as ‘Langmuir sites’ or ‘holes’ [27-32]. At a given temperature, a fixed number of holes is available in the glassy polymer (randomly distributed within the polymer) and upon hole saturation, the solubility coefficient approaches the asymptotic value of kD where gas molecules dissolve in the equilibrium dense portion of the polymer matrix – referred to as the ‘Henry’s mode’ or ‘dissolved mode’ of sorption as shown in

Fig. 2.5. 75

Fig. 2.5. The solubility coefficient, S = c/p, decreases with hole saturation and reaches the asymptotic solubility limit, kD, as the pressure increases.

The gas solubility in glassy polymers can be correlated with the fractional free volume (FFV), defined as the volume fraction of the polymer unoccupied by the polymer chain molecules [33]:

푣 − 푣 퐹퐹푉 = 0 (2.10) 푣

3 where 푣 is the polymer specific volume (cm /g) and 푣0 is the volume occupied by the polymer chain molecules [13, 15].

The solubility of gases depends on the relative affinity between the gas and polymer as well as the gas condensability, which is generally correlated with the Lennard-Jones force constant (ε/k), gas critical temperature (T ), or normal boiling point (T ) [13, 34]. c b

The relationship between the solubility and gas condensability is generally expressed as:

푙푛 푆 = 푎 + 푏푇푐 (2.11) 76

The constant a is a measure of the overall sorption capacity, while the slope b represents the increase in solubility with penetrant condensability. Gas solubility generally increases as gas condensability increases.

2.5. Principles of CO2/CH4 Membrane Separation

For CO /CH separation, CO (T = 217 K) is more condensable than CH (T = 113 2 4 2 b 4 b

K), hence, solubility selectivity always favors CO2 over CH4. However, the solubility selectivity is relatively constant with different structures of glassy polymers unless highly polar functionalities are introduced to the polymer backbone [19, 35]. Furthermore, CO2

(KD = 3.3 Å) is smaller than CH4 (KD = 3.8 Å) and, thus, diffusivity selectivity also favors

CO2. Accordingly, a promising route to produce better membrane materials for this separation is by fine-tuning the microporous glassy polymer structure to provide higher permeability compared to commercial polymers and higher selectivities that utilize the molecular size difference like molecular sieves.

2.6. Mixed-Gas Non-Idealities Transport

Membrane transport properties using pure-gas feeds are used to determine the ideal permeability, diffusion and solubility coefficients. However, for gas mixtures containing condensable components (such as CO2, C2+ hydrocarbons, water vapor etc.), the transport properties of one component are affected by co-permeation of the other feed gas components. In such cases, pure-gas properties can lead to severely inaccurate estimates of actual high-pressure mixed-gas separation performance [36]. For example, previous work showed that in CO2/CH4 mixed-gas separation experiments using cellulose acetate

(CA) membranes, the CO2 permeability decreased to about 50% relative to the pure-gas 77

value, whereas the permeability of CH4 tripled due to CO2-induced plasticization [37].

Consequently, the mixed-gas selectivity was reduced significantly due to two effects: (i) competitive sorption, which reduced CO2 permeability, and (ii) plasticization, which enhanced CH4 permeability.

2.6.1. Competitive Sorption

When two (or more) gases are simultaneously present in a mixture (like CO2 and

CH4), each gas competes for the available sorption sites in the glassy polymer membrane and reduces the other apparent gas solubility coefficient. As a result, the relative CO2 concentration is lower than its pure-gas concentration at an equivalent gas partial pressure

[16, 38]. Consequently, the mixed-gas CO2 permeability is lower than the pure-gas value due to the competitive sorption, which reduced the overall sorbed concentration of CO2 in the polymer across the entire pressure range. In the absence of plasticization, both gases are affected under mixed-gas conditions by competitive sorption effects [39, 40].

However, the preferential solubility of the more condensable CO2 often increases the polymer chain mobility, resulting in a relatively large increase in the diffusivity, and hence, permeability of the larger CH4 penetrant molecule [41, 42].

2.6.2. Plasticization

Sorption of condensable feed components such as CO2 into glassy polymeric membranes causes loosening of the stiff polymer chains [41]. Hence, dilation of the polymer morphology and the increased chain mobility will reduce its sieving ability. This effect called plasticization is detrimental to the performance of glassy polymers, which transport and discriminate between gas molecules based on molecular sizes. Large loss in 78

selectivity due to significant increase in CH4 permeability is often observed in CO2/CH4 mixed-gas experiments (Fig. 2.6). Researchers have introduced many techniques to mitigate plasticization such as polymer blending [43], chemical crosslinking [44], functionalization by (e.g. hydroxyl- and carboxyl groups) [45] and thermal treatment

[46].

Fig. 2.6. Schematic illustrating the plasticization effect on glassy polymeric membranes.

2.7. Experimental Methods and Techniques

2.7.1. Evaluation of Microporosity in Polymers

Characterization of the microporosity is significant to elucidate the gas transport properties of newly developed porous materials. Low-pressure physisorption, developed by Brunauer, Emmett and Teller (BET), (Fig. 2.7 and 2.8) is a technique to investigate the structure and surface areas of porous materials from their sorption isotherms [47]. 79

Fig. 2.7. Gas adsorption device for measuring BET surface area and pore size distribution in porous materials (Micromeritics, ASAP-2020).

Fig. 2.8. Schematic representing BET gas adsorption technique (ASAP-2020) [48].

Brunauer–Emmett–Teller (BET) theory is based on the physical adsorption of gas molecules on a solid surface and serves as the basis for an important analytical technique for the measurement of the specific surface area of materials [49]. The BET theory applies to systems of multi-layer adsorption, and typically utilizes probing gases that do not chemically react with material surfaces as adsorbates to quantify specific surface area. Nitrogen is most commonly used for surface probing by the BET method.

Therefore, standard BET analysis is usually conducted under cryogenic condition at the 80

normal boiling temperature of N2 (77 K). Alternative probing adsorbates are also utilized, allowing measurement of the surface area at different temperatures and with different probe scales, including argon and carbon dioxide [50]. BET theory is based on an equation for a sorption isotherm as follows:

푝 1 (퐶 − 1) 푝 = 푎 + 푎 (2.12) 푛푎(푝0 − 푝) 푛푚퐶 푛푚퐶 푝0

푝 where 푛 is the amount of gas adsorbed at the relative pressure ⁄ , 푝 is the saturation 푎 푝0 0

푎 vapor pressure of the gas, 푛푚 is the monolayer capacity and C is a constant which depends on the isotherm shape. A linear relation exists in a plot of the left hand side

푝 푝 against ⁄ . Typically this linearity occurs in the range of 0.01 < ⁄ < 0.35, but for 푝0 푝0 푝 microporous materials it is often 0.01 < ⁄ < 0.10 [51]. 푝0

BET theory is based on a simplified extension of the Langmuir mechanism to multilayer adsorption. Langmuir claimed that the rates of gas evaporation from and gas condensation onto a basic surface are proportional to the fractions of occupied and unoccupied area, respectively [52].

The Langmuir theory relates the monolayer adsorption of gas molecules (adsorbates) onto a solid surface (Fig. 2.9) to the gas pressure of a medium above the solid surface at fixed temperature as shown in equation 2.13, where θ is the fractional cover of the surface, p is the gas pressure and α is a constant.

훼 . 푝 휃 = (2.13) 1 + (훼 . 푝) 81

Fig. 2.9. Schematic representing the adsorption of gas molecules onto the surface of a sample showing (a) the monolayer adsorption model assumed by the Langmuir theory and (b) multilayer adsorption model assumed by the BET theory [48].

The assumptions of Langmuir theory are as follows:

o All surface sites have the same adsorption energy for the adsorbate.

o Adsorption of the solvent at one site occurs independently of adsorption at

adjacent sites.

o Activity of adsorbate is proportional to its concentration.

o Adsorbates form a monolayer.

o Each active site can be occupied by one adsorbate molecule only.

The Langmuir theory has a few faults that are addressed by the BET theory. The BET theory extends the Langmuir theory to multilayer adsorption (Fig. 2.9) with three extra assumptions:

o Gas molecules will physically adsorb on a solid in layers infinitely.

o The different adsorption layers do not interact. 82

o The theory can be applied to each layer.

Adsorption can be defined as the adhesion of molecules of gas to a surface. The amount of gas adsorbed depends on the exposed surface, temperature, gas pressure and strength of interaction between the gas and solid. In BET surface area analysis, nitrogen is usually used because of its availability in high purity and relatively small size. Because the interaction between gases and solids is usually weak, the surface is cooled using liquid N2 to obtain detectable amounts of adsorption. Known amounts of nitrogen gas are then released stepwise into the sample cell. Relative pressures less than are achieved by creating partial vacuum. Highly precise and accurate pressure transducers monitor the pressure changes due to the adsorption process. After the adsorption layers are formed, the sample is removed from the nitrogen atmosphere and heated to cause the adsorbed nitrogen to be released from the material and quantified.

The collected data is displayed in the form of a BET isotherm, which plots the amount of gas adsorbed as a function of the relative pressure [48]. The modern version of the

IUPAC classification scheme has defined six types of isotherms for gas/solid equilibria

[53], as shown in Fig. 2.10. It is based on a previous classification by Brunauer, which had five types of isotherms [54]. Type I isotherm is used to describe adsorption on microporous adsorbents. Types II and III describe adsorption on macroporous adsorbents.

Types IV and V are representing mono-and multilayer adsorption plus capillary condensation. Type VI, which was not included in the Brunauer classification, illustrates that the adsorption isotherms can have one or more steps [54]. 83

Fig. 2.10. Schematic representing various adsorption isotherm types [54].

The main components of the BET instrument include (1) degas port, (2) analysis port,

(3) a circulating bath for temperature regulation (e.g. liquid N2) and (4) vacuum pumps, including a turbo pump.

In this study, a typical experiment involved degassing a sample tube, then charging with nitrogen gas to 1 bar and loading with 50-100 mg of sample (polymer powder).

Then, an O-ring was installed on the tube to ensure full sealing. After that, the tube sample was connected to the degassing port and degassed under high vacuum with a heating temperature of 120 °C using a heating mantle for 16 hours. Then, the sample tube was backfilled with nitrogen and weighted to determine the of the degassed sample. Afterwards, the sample tube was connected at the sample port and a vacuum was applied for around 1 hour before starting the sorption experiment.

For nitrogen sorption, the sample tube was fitted with an isothermal Teflon jacket and the supplied Dewar vessel was filled with liquid nitrogen (77 K). The jacket creates a siphon of liquid nitrogen up the sample tube driven by evaporation at the top, maintaining 84 a layer of liquid nitrogen along the tube. After that, warm and cold free spaces were measured with helium. Then, adsorption and desorption isotherms were collected.

Equilibration times of 10 seconds between pressure readings were used and isotherms typically completed within 48 to 72 hours.

For carbon dioxide sorption, the Dewar was filled with a mixture of ethanol and water. The temperature was kept at 273 K by using a heat-exchange coil that was prepared and connected to a circulating bath. Free space was also measured with helium.

Then, adsorption and desorption isotherms were collected up to p/p0 = 0.03 (1 bar, 273 K) with equilibration intervals of 10 seconds. Experiments were typically completed within

14 hours.

Recently, various gas probes were used to study the pore size distribution of amorphous microporous polymers and network polymers [55-57]. This is because the gas adsorption measurements by using N2 as probe molecule at 77 K (-196 °C) are not fully representative of the physical state of the polymer as the typical gas transport testing is carried out at 35 °C [57]. Furthermore, the relatively large N2 molecule (KD = 3.64 Å) cannot completely access the ultramicropores (< 7 Å) due to diffusional restrains [58-60]; thus, the resulting data provide only partial information about the porous structure in the ultramicroporous range, which is of utmost importance to the sieving capability of a PIM for gas separations. Therefore, CO2 has been proposed as a smaller probe molecule (KD =

3.3 Å) at 273 K (0 °C) for ultramicroporosity to complement the data obtained from N2 sorption at 77 K (-196 °C) [61]. 85

2.7.2. Pure-Gas Permeability Measurement System

The pure-gas permeability through a dense polymeric membrane was determined using the constant-volume/variable- system [62, 63] which is a custom-made permeation system designed and built in the Advanced Membranes and

Porous Materials Center (AMPMC) at KAUST, as shown in Fig. 2.11 and Fig. 2.12. The system consists of: (1) fans for air circulation, (2) heaters, (3) permeation cell, (4) upstream volume, (5) downstream volume, (6) upstream pressure transducer, (7) downstream pressure transducer, and (8) sealed gas valves.

Fig. 2.11. Schematic design of a constant volume/variable pressure permeation system for pure-gas experiments. 86

Fig. 2.12. Picture of the pure-gas permeation measurement system (constant volume/variable pressure permeation system).

The permeation cell was a circular stainless steel holder purchased from Millipore

Corporation with an effective area of 13.8 cm2. The temperature inside the system was regulated by a temperature controller (Omega®) connected to the heater and two fans that circulated the air inside the permeation box. To start permeation experiment, first, the polymer film sample should be partially masked with an impermeable aluminum tape to expose a membrane area of approximately 1.0 cm2. The masked sample was then placed in the permeation cell, and another layer of aluminum tape was added to fix the masked sample in the cell. The interface between the aluminum tape and the membrane was carefully masked with epoxy glue to prevent any leakage through the openings. The glue was allowed to fully cure for approximately 16 to 24 hours. Thereafter, the cell was re- installed in the permeation system and exposed to vacuum from both the upstream and downstream sides for more than 24 hours at 35 °C. 87

After degassing the sample, the upstream and downstream valves were closed from the vacuum pump. After that, the change in permeate pressure with time (dp/dt) was measured to ensure that the leak rate of the system was below 10–7 torr/s. The upstream gas reservoir was then purged and pressurized with the desired feed gas, after which the feed gas valve was opened to introduce the test gas to the membrane. The permeate pressure was continuously monitored and the data were recorded until a steady-state dp/dt was achieved. The recorded data were utilized for further analysis to determine the diffusion (D) and solubility (S) coefficients values using the time lag method.

The pure-gas permeability of He, H2, N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar. The gas permeability was calculated by:

푉 푙 푑푝 푃 = 1010 푑 (2.14) 푃푢푝 퐴푅푇 푑푡

-10 3 2 where P is the permeability in Barrers (1 Barrer = 10 cm (STP cm/cm s cmHg), pup is the upstream pressure (cmHg), dp/dt is the steady-state permeate-side pressure over time

3 (cmHg/s), Vd is the calibrated permeate volume (cm ), l is the membrane thickness (cm),

A is the effective membrane area(cm2), T is the operating temperature (K), and R is the gas constant (0.278 cm3 cmHg /cm3(STP) K).

The apparent diffusion coefficient D (cm2/s) was calculated by D = l 2/6θ, where θ is the time lag of the permeability measurement and 푙 is the membrane thickness. The solubility coefficient S (cm3(STP)/cm3 cmHg) was then calculated from the solution- diffusion gas transport relationship:

S=P/D (2.15) 88

The time lag is referring to the time required for the gas to leave the transient state and establish equilibrium, as presented in Fig. 2.13. The time lag was calculated by extrapolating the pressure increase-time curve to the time axis after the equilibrium state was reached. To ensure that the diffusion coefficient calculations were reliable, steady- state readings over a minimum period of at least ten time lags were used.

Fig. 2.13. Graphical representation of the time lag technique in the constant- volume/variable pressure method.

2.7.3. Mixed-Gas Measurement System

Several techniques have been reported in the literature for testing the performance of polymeric membranes under mixed-gas feed conditions [63-65]. The constant- volume/variable-pressure technique was used in this research (Fig. 2.14 and Fig. 2.15).

The mixed-gas system consisted of: (1) fans for air circulation, (2) heaters, (3) permeation cell, (4) upstream volume, (5) downstream volume, (6) upstream pressure transducer, (7) downstream pressure transducer, (8) sealed gas valves, (9) gas 89 chromatograph and (10) exist port for residue in the permeation cell so that the feed gas could flow along the surface and exit. The stage-cut (permeate flow/feed flow) was maintained below 1% by purging the upstream of the permeation cell at an appropriate flow rate. This step ensured that the residue gas concentration was basically equal to the feed gas concentration throughout the permeation experiment and prevented concentration polarization at the upstream of the permeation cell.

Fig. 2.14. Schematic design of a constant-volume/variable-pressure permeation system for mixed-gas experiments. 90

Fig. 2.15. Picture of the mixed-gas permeation measurement system (constant- volume/variable-pressure permeation system).

CO2 and CH4 permeate concentrations were detected with a gas chromatograph (GC)

(Agilent 3000A Micro GC) equipped with a thermal conductivity detector (TCD) which was calibrated with several gases from Abdullah Hashim Group (AHG). A binary 50:50

CO2/CH4 mixture was used to determine the calibration constants for CO2 and CH4. The column temperature was set between 80 and 120 °C for all measurements and ultra-high purity helium was used as carrier gas.

The mixed-gas permeation measurements of all tested samples were performed at 35

°C using a method similar to that described by O’Brien et al. [63]. The feed gas mixture contained 50 vol.% CO2 and 50 vol.% CH4. The total feed pressure was varied between 4 and 30 bar. The mixed-gas permeability was calculated by:

10 푦퐶푂2푉푑퐿 푑푝 푃퐶푂2 = 10 (2.16) 푥퐶푂2푝푢푝퐴푅푇 푑푡

10 푦퐶퐻4푉푑퐿 푑푝 푃퐶퐻4 = 10 (2.17) 푥퐶퐻4푝푢푝퐴푅푇 푑푡 91 where y and x were the mole fractions in the permeate and feed, respectively. The GC was used to measure the permeate mole fraction (y) with a minimum of five runs that were conducted to confirm the accuracy of the analysis at each pressure. The mixed-gas

CO2/CH4 selectivity was obtained from:

푦퐶푂2/푦퐶퐻4 훼퐶푂2/퐶퐻4 = (2.18) 푥퐶푂2/푥퐶퐻4

2.7.4. Thermal Rearrangement Method

Hydroxyl-functionalized polyimide samples were converted by thermal rearrangement (TR) to polybenzoxazoles (PBOs) membranes by thermal treatment in a

Carbolite tubular furnace (Fig. 2.16 and Fig. 2.17). Heat treatment was conducted under nitrogen flow and polymer samples were sandwiched between two ceramic plates inside a ceramic boat. Metal spacers were inserted between the ceramic plates to allow nitrogen to reach the samples. The oxygen concentration in the furnace was continuously monitored with an O2 analyser (Cambridge Sensotech, Rapidox 3100). Ramping of the temperature did not begin until the O2 concentration was less than 2 ppm. The TR-PBO membranes were made at a rate of 10 °C/min to the targeted temperature where it was held isothermally for certain time to achieve essentially full conversion of the polyimide to

PBO. The required time for this process was determined by using thermogravimetric analysis (TGA). Finally, the furnace was shut off and the film was kept inside the oven to cool to room temperature under continuous nitrogen flow. 92

Fig. 2.16. Schematic of the Carbolite tubular furnace used for preparation of polybenzoxazole membranes from hydroxyl-functionalized polyimides.

Fig. 2.17. Picture of the Carbolite tubular furnace, oxygen analyzer and control box.

2.7.5. Thermogravimetric Analysis (TGA)

Thermogravimetric analysis (TGA) was conducted with a Q5000 instrument (TA

Instruments) to determine the thermal stability of polymers used in this study. TGA measures the amount and rate of weight change in a material, either as a function of temperature, or isothermally as a function of time, in an isolated atmosphere. Also, TGA was used in this research to ensure that cast and dried polymer films did not contain any traces of solvent prior to the permeability measurements. Moreover, TGA was used for 93 determination of the optimum thermal rearrangement treatment conditions (duration and temperature). The major components of TGA Q5000 (Fig. 2.18 and Fig. 2.19) are:

 The microbalance, which provides precise measurement of sample weight.

 The heating system (or infrared (IR) furnace), which controls the sample

temperature.

 The autosampler, which loads and unloads the sample to and from the balance.

 The mass flow controllers (MFCs), which control the purge gas to the furnace.

The thermal stability test of the membranes samples (film or powder) was measured in a nitrogen atmosphere using approximately 4-5 mg of sample, which were loaded onto a HT-platinum pan. Then, the sample was: (i) heated to 100 °C, (ii) kept isothermally for

30 minutes and finally, (iii) the temperature was ramped to 800 °C at 3 °C/min.

Fig. 2.18. Picture of a thermogravimetric analyzer (TGA Q5000, TA Instruments). 94

Fig. 2.19. Schematic showing thermogravimetric analyzer (TGA) main components [66].

2.7.6. Wide-Angle X-Ray Diffraction (WAXD)

X-ray diffraction is a non-destructive method for characterization of solid materials.

When x-rays are directed in solids they scatter in predictable patterns based upon the internal structure of the solid [67, 68]. The main components are source of x-ray, sample stage and detector (Fig. 2.20). This technique specifically refers to the analysis of Bragg peaks scattered to wide angles. The diffraction pattern generated allows determining the chemical composition or phase composition of the film, the texture of the film, the crystallite size and presence of film stress. According to this method, the sample is scanned and the scattering intensity is plotted as a function of the 2θ angle.

Fig. 2.20. Schematic representing main components of x-ray-diffraction technique. 95

Moreover, x-ray diffraction (XRD) is a powerful technique to study the atomic structure of a crystal to obtain the average d-spacing of a crystalline material using

Bragg’s law, n  = 2 d sin, where d is the spacing between the diffraction planes,  is the wavelength of the x-ray (=1.5418 Å for common lab x-ray source Cu kα radiation), and  is the incident angle. Due to the amorphous nature of most glassy polymeric materials used for gas separation membranes, a broad amorphous halo is often seen in the

XRD spectra; consequently only semi-quantitative information can be obtained compared to crystalline materials. Nevertheless, XRD spectra still yield important structure information for polymers [69].

For amorphous polymeric materials, the average d-spacing is often interpreted as the average intermolecular distances for polymer chains and can thus be used as a measure of the openness of the polymer matrix [69-72]. An increase in the intermolecular distance can result in an increase in free volume of the polymer and most likely an increase in gas permeability. For gas separation applications, most gas pairs of interest have molecular dimension generally less than 10 Å, therefore, only wide-angle x-ray diffraction

(WAXD) is needed to probe the d-spacing in this range. WAXD has been used extensively for polymeric gas separation membrane materials and carbon molecular sieves. The average d-spacing, shape and characteristic of the WAXD spectra has shown good correlation with the gas permeability and selectivity, although in some cases the interpretation is not straightforward [70-73].

WAXD spectra of the polymers studied in this work were collected on a Bruker D8

Advance diffractometer using a Bruker zero background holder. The scanning rate ranged 96 from 0.5 to 2 second per step, 0.02 degree per step, and 2 varying from 7 to 50o. Before measuring the sample spectra, a background spectra was collected using only the sample holder and it was then carefully subtracted from the sample spectra under the same experimental conditions. The average d-spacing values were then calculated using

Bragg’s law. The produced peaks were utilized to study the d-spacing distribution of the membrane materials and compare between the fresh and aged samples.

97

2.8. References

[1] Z.Y. Yeo, T.L. Chew, P.W. Zhu, A.R. Mohamed, S.-P. Chai, Conventional processes and membrane technology for carbon dioxide removal from natural gas: a review, Journal of Natural Gas Chemistry, 21 (2012) 282-298.

[2] A. Car, C. Stropnik, W. Yave, K.V. Peinemann, Tailor-made polymeric membranes based on segmented block copolymers for CO2 separation, Advanced Functional

Materials, 18 (2008) 2815-2823.

[3] S.A. Stern, Polymers for gas separations - the next decade, Journal of Membrane

Science, 94 (1994) 1-65.

[4] R.W. Baker, Future directions of membrane gas separation technology, Industrial &

Engineering Chemistry Research, 41 (2002) 1393-1411.

[5] R.W. Baker, B.T. Low, Gas separation membrane materials: a perspective,

Macromolecules, 47 (2014) 6999-7013.

[6] L.H. Sperling, Introduction to physical polymer science, John Wiley & Sons, 2005.

[7] P. Bernardo, E. Drioli, G. Golemme, Membrane gas separation: a review/state of the art, Industrial & Engineering Chemistry Research, 48 (2009) 4638-4663.

[8] R.W. Baker, K. Lokhandwala, Natural gas processing with membranes: an overview,

Industrial & Engineering Chemistry Research, 47 (2008) 2109-2121.

[9] J.M.S. Henis, M.K. Tripodi, The developing technology of gas separating membranes, Science, 220 (1983) 11-17. 98

[10] A. Javaid, Membranes for solubility-based gas separation applications, Chemical

Engineering Journal, 112 (2005) 219-226.

[11] T. Graham, On the absorption and dialytic separation of gases by colloid septa,

Philosophical Transactions of the Royal Society of London, 156 (1866) 399-439.

[12] J.G. Wijmans, R.W. Baker, The solution-diffusion model - a review, Journal of

Membrane Science, 107 (1995) 1-21.

[13] S. Matteucci, Y. Yampolskii, B.D. Freeman, I. Pinnau, Transport of gases and vapors in glassy and rubbery polymers, in Materials Science of Membranes for Gas and

Vapor Separation, Wiley, Chichester, England, (2006) 1-2.

[14] W.J. Koros, G.K. Fleming, Membrane-based gas separation, Journal of Membrane

Science, 83 (1993) 1-80.

[15] R.W. Baker, Membrane technology and applications, John Wiley & Sons, Ltd.,

2004.

[16] W.J. Koros, Simplified analysis of gas/polymer selective solubility behavior, Journal of Polymer Science: Polymer Physics Edition, 23 (1985) 1611-1628.

[17] D.W. Breck, molecular sieves: structure, chemistry and Use, Wiley, New

York, 636 (1974).

[18] J.M. Henis, M.K. Tripodi, Composite hollow fiber membranes for gas separation: the resistance model approach, Journal of Membrane Science, 8 (1981) 233-246. 99

[19] K. Ghosal, B.D. Freeman, Gas separation using polymer membranes: an overview,

Polymers for Advanced Technologies, 5 (1994) 673-697.

[20] J. Chiou, Y. Maeda, D.R. Paul, Gas and vapor sorption in polymers just below Tg,

Journal of Applied Polymer Science, 30 (1985) 4019-4029.

[21] R.E. Kesting, A. Fritzsche, Polymeric gas separation membranes, Wiley-

Interscience, 1993.

[22] C.L. Staiger, S.J. Pas, A.J. Hill, C.J. Cornelius, Gas separation, free volume distribution, and physical aging of a highly microporous spirobisindane polymer,

Chemistry of Materials, 20 (2008) 2606-2608.

[23] E. Sanders, W. Koros, H. Hopfenberg, V. Stannett, Mixed gas sorption in glassy polymers: equipment design considerations and preliminary results, Journal of Membrane

Science, 13 (1983) 161-174.

[24] R. Barrer, Diffusivities in glassy polymers for the dual mode sorption model, Journal of Membrane Science, 18 (1984) 25-35.

[25] P.M. Budd, N.B. McKeown, Highly permeable polymers for gas separation membranes, Polymer Chemistry, 1 (2010) 63-68.

[26] J. Guo, T.A. Barbari, A dual mode interpretation of the kinetics of penetrant-induced swelling and deswelling in a glassy polymer, Polymer, 51 (2010) 5145-5150.

[27] T. Barbari, R. Conforti, Recent theories of gas sorption in polymers, Polymers for

Advanced Technologies, 5 (1994) 698-707. 100

[28] R. Barrer, J. Barrie, J. Slater, Sorption and diffusion in ethyl cellulose. Part III. comparison between ethyl cellulose and rubber, Journal of Polymer Science, 27 (1958)

177-197.

[29] A.S. Michaels, W.R. Vieth, J.A. Barrie, Solution of gases in terephthalate, Journal of Applied Physics, 34 (1963) 1-12.

[30] W. Vieth, J. Howell, J. Hsieh, Dual sorption theory, Journal of Membrane Science, 1

(1976) 177-220.

[31] W. Vieth, H. Alcalay, A. Frabetti, Solution of gases in oriented poly(ethylene terephthalate), Journal of Applied Polymer Science, 8 (1964): 2125-2138.

[32] S. Stern, V. Saxena, Concentration-dependent transport of gases and vapors in glassy polymers, Journal of Membrane Science, 7 (1980) 47-59.

[33] P. Bernardo, E. Drioli, Membrane gas separation progresses for process intensification strategy in the petrochemical industry, Petroleum Chemistry, 50 (2010)

271-282.

[34] G. Van Amerongen, Influence of structure of elastomers on their permeability to gases, Journal of Polymer Science, 5 (1950) 307-332.

[35] T.C. Merkel, I. Pinnau, R. Prabhakar, B.D. Freeman, Gas and vapor transport properties of perfluoropolymers, in Materials Science of Membranes for Gas and Vapor

Separation, Wiley, Chichester, England, (2006) 251-270.

[36] E. Sanders, Penetrant-induced plasticization and gas permeation in glassy polymers,

Journal of Membrane Science, 37 (1988) 63-80. 101

[37] M.D. Donohue, B.S. Minhas, S.Y. Lee, Permeation behavior of carbon-dioxide methane mixtures in cellulose-acetate membranes, Journal of Membrane Science, 42

(1989) 197-214.

[38] W.J. Koros, R. Chern, V. Stannett, H. Hopfenberg, A model for permeation of mixed gases and vapors in glassy polymers, Journal of Polymer Science: Polymer

Physics Edition, 19 (1981) 1513-1530.

[39] R. Ash, R.M. Barrer, R.T. Lowson, Transport of single gases and of binary gas mixtures in a microporous carbon membrane, Journal of the Chemical Society, Faraday

Transactions 1: Physical Chemistry in Condensed Phases, 69 (1973) 2166-2178.

[40] R. Ash, R. Barrer, C. Pope, Flow of adsorbable gases and vapours in a microporous medium. II. Binary mixtures, in: Proceedings of the Royal Society of London A:

Mathematical, Physical and Engineering Sciences, The Royal Society, 1963, pp. 19-33.

[41] A. Bos, I. Pünt, M. Wessling, H. Strathmann, CO2-induced plasticization phenomena in glassy polymers, Journal of Membrane Science, 155 (1999) 67-78.

[42] E. Sanders, S. Jordan, R. Subramanian, Penetrant-plasticized permeation in polymethylmethacrylate, Journal of Membrane Science, 74 (1992) 29-36.

[43] A. Bos, I. Pünt, H. Strathmann, M. Wessling, Suppression of gas separation membrane plasticization by homogeneous polymer blending, AIChE Journal, 47 (2001)

1088-1093.

[44] C. Staudt-Bickel, W.J. Koros, Improvement of CO2/CH4 separation characteristics of polyimides by chemical crosslinking, Journal of Membrane Science, 155 (1999) 145-154. 102

[45] X.H. Ma, R. Swaidan, Y. Belmabkhout, Y.H. Zhu, E. Litwiller, M. Jouiad, I. Pinnau,

Y. Han, Synthesis and gas transport properties of hydroxyl-functionalized polyimides with intrinsic microporosity, Macromolecules, 45 (2012) 3841-3849.

[46] H.B. Park, C.H. Jung, Y.M. Lee, A.J. Hill, S.J. Pas, S.T. Mudie, E. Van Wagner,

B.D. Freeman, D.J. Cookson, Polymers with cavities tuned for fast selective transport of small molecules and ions, Science, 318 (2007) 254-258.

[47] K.S. Sing, Reporting physisorption data for gas/solid systems with special reference to the determination of surface area and porosity (Recommendations 1984), Pure and

Applied Chemistry, 57 (1985) 603-619.

[48] http://cnx.org/contents/9cBY4EHy@1/BET-Surface-Area-Analysis-of-N.

[49] S. Brunauer, P.H. Emmett, E. Teller, Adsorption of gases in multimolecular layers,

J. Am. Chem. Soc, 60 (1938) 309-319.

[50] D.A. Hanaor, M. Ghadiri, W. Chrzanowski, Y. Gan, Scalable surface area characterization by electrokinetic analysis of complex anion adsorption, Langmuir, 30

(2014) 15143-15152.

[51] R. Swaidan, Intrinsically microporous polymer membranes for high performance gas separation, Ph.D. dissertation, King Abdullah University of Science and Technology,

Saudi Arabia (2014).

[52] A.P. Gast, A.W. Adamson, Physical chemistry of surfaces, in, John Wiley and Sons,

Inc: New York, 1997.

[53] K.S. Sing, Physisorption of gases by carbon blacks, Carbon, 32 (1994) 1311-1317. 103

[54] M. Donohue, G. Aranovich, Classification of Gibbs adsorption isotherms, Advances in Colloid and Interface Science, 76 (1998) 137-152.

[55] J. Weber, O. Su, M. Antonietti, A. Thomas, Exploring polymers of intrinsic microporosity-microporous, soluble and polyimide, Macromolecular Rapid

Communications, 28 (2007) 1871-1876.

[56] J. Weber, J. Schmidt, A. Thomas, W. Böhlmann, Micropore analysis of polymer networks by gas sorption and 129Xe NMR spectroscopy: toward a better understanding of intrinsic microporosity, Langmuir, 26 (2010) 15650-15656.

[57] N. Ritter, I. Senkovska, S. Kaskel, J. Weber, Intrinsically microporous poly (imide) s: structure− porosity relationship studied by gas sorption and X-ray scattering,

Macromolecules, 44 (2011) 2025-2033.

[58] D. Lozano-Castello, D. Cazorla-Amoros, A. Linares-Solano, Usefulness of CO2 adsorption at 273 K for the characterization of porous carbons, Carbon, 42 (2004) 1233-

1242.

[59] D. Cazorla-Amorós, J. Alcaniz-Monge, M. De la Casa-Lillo, A. Linares-Solano,

CO2 as an adsorptive to characterize carbon molecular sieves and activated carbons,

Langmuir, 14 (1998) 4589-4596.

[60] K.M. Steel, W.J. Koros, Investigation of porosity of carbon materials and related effects on gas separation properties, Carbon, 41 (2003) 253-266.

[61] G. Amarasekera, M. Scarlett, D. Mainwaring, Micropore size distributions and specific interactions in , Fuel, 74 (1995) 115-118. 104

[62] D. Pye, H. Hoehn, M. Panar, Measurement of gas permeability of polymers. II. apparatus for determination of permeabilities of mixed gases and vapors, Journal of

Applied Polymer Science, 20 (1976) 287-301.

[63] K. O'Brien, W. Koros, T. Barbari, E. Sanders, A new technique for the measurement of multicomponent gas transport through polymeric films, Journal of Membrane Science,

29 (1986) 229-238.

[64] L.R. Smith, H. Czichos, T. Saito, L. Smith, Springer handbook of materials measurement methods, Springer, 2006.

[65] S. Damle, W.J. Koros, Permeation equipment for high-pressure gas separation membranes, Industrial & Engineering Chemistry Research, 42 (2003) 6389-6395.

[66] H. Song, K. Shah, E. Doroodchi, T. Wall, B. Moghtaderi, Analysis on chemical reaction kinetics of CuO/SiO2 oxygen carriers for chemical looping air separation,

Energy & Fuels, 28 (2013) 173-182.

[67] S.G. Podorov, N. Faleev, K. Pavlov, D. Paganin, S. Stepanov, E. Förster, A new approach to wide-angle dynamical X-ray diffraction by deformed crystals, Journal of

Applied Crystallography, 39 (2006) 652-655.

[68] S. Podorov, A. Nazarkin, Wide-angle x-ray diffraction theory versus classical dynamical theory, Recent Res. Devel. Optics, 7(2009) 11-17.

[69] R.L. Miller, R.F. Boyer, Regularities in x‐ray scattering patterns from amorphous polymers, Journal of Polymer Science: Polymer Physics Edition, 22 (1984) 2043-2050. 105

[70] E. Hensema, M. Mulder, C. Smolders, On the mechanism of gas transport in rigid polymer membranes, Journal of Applied Polymer Science, 49 (1993) 2081-2090.

[71] A. Shimazu, T. Miyazaki, K. Ikeda, Interpretation of d-spacing determined by wide angle X-ray scattering in 6FDA-based polyimide by molecular modeling, Journal of

Membrane Science, 166 (2000) 113-118.

[72] S. Stern, Y. Mi, H. Yamamoto, A.K.S. Clair, Structure/permeability relationships of polyimide membranes. Applications to the separation of gas mixtures, Journal of Polymer

Science Part B: Polymer Physics, 27 (1989) 1887-1909.

[73] X. Ma, R. Swaidan, B. Teng, H. Tan, O. Salinas, E. Litwiller, Y. Han, I. Pinnau,

Carbon molecular sieve gas separation membranes based on an intrinsically microporous polyimide precursor, Carbon, 62 (2013) 88-96.

106

Chapter 3. Materials and Literature Review of Microporous Polymeric Materials

This chapter provides a brief overview of (i) microporous materials, (ii) polymers of intrinsic microporosity (PIMs), (iii) pore size and transport mechanisms in PIMs membranes, (iv)_improving microporous polyimides for natural gas sweetening

(CO2/CH4 separation), (v) triptycene-based polymers, (vi) PIMs physical aging and (vii) plasticization in PIMs. In general, this work was based on collected data for nine polymeric membranes - seven polyimides (PIM-Polyimides) and two thermally rearranged polyimides (PBOs) - and its comparison to data for materials previously reported in the literature. The chemical structures are presented in this chapter for the dianhydrides and diamines that were used in the investigated polyimides. The employment of triptycene as structural building block motif was demonstrated in different PIM-PIs for the generation of advanced membrane materials for CO2/CH4 separation (natural gas sweetening).

3.1. Microporous Materials

Microporous materials, which contain pores less than 2.0 nm in size, are commonly used for gas separation, storage [1], gas adsorption [2, 3] and heterogeneous catalysis [4].

They conventionally consist of crystalline inorganic frameworks such as , metal organic frameworks (MOFs) and covalent organic frameworks (COFs) or an amorphous structure such as (Fig. 3.1) [5-11]. In addition, microporosity can be obtained from amorphous organic materials from recently developed intrinsically microporous polymers derived from a number of different monomers and polymerization reactions [12-17]. Polymers are the best materials to fabricate membranes, as explained 107 previously in Chapter 2, because of their low cost and the simplicity of processing them into different structures [18], whereas, zeolites, MOFs and COFs have a crystalline nature, which results in inflexibility and complicates their processability, obstructing their incorporation into the membrane gas separation industry. Amorphous, solution- processable polymers continued to be the best materials choice for membrane-based gas separation. They can be designed with intrinsic microporosity purely by advantage of chains that inefficiently pack in the solid state [19, 20].

Fig. 3.1. Types of promising microporous materials.

Polymers of intrinsic microporosity (PIMs) are particularly interesting for membranes because they have the potential to unite solution processability, mechanical flexibility and organic adjustability of commercially relevant polymers with the microporosity of inorganic materials. Intrinsic microporosity in polymers is defined as “a continuous network of interconnected intermolecular voids, which forms as a direct consequence of the shape and rigidity of the component macromolecules” [21, 22]. Most polymers have sufficient conformational flexibility to allow them to rearrange their shape so as to maximize intermolecular cohesive interactions and pack space efficiently. The target for

PIMs is to maximizing intrinsic microporosity and design polymers with highly rigid and 108 contorted molecular structures to provide “awkward” macromolecular shapes that cannot pack space efficiently [22].

Examples of PIMs that have emerged over the years in various classes, including polyacetylenes (e.g., poly(1-trimethylsilyl-1-propyne), PTMSP), ladder polymers (e.g.,

PIM-1) and polyimides (e.g., polyimide of intrinsic microporosity, PIM-PI-8, PIM-PI-10,

KAUST-PI-1 and KAUST-PI-7). PIMs are characterized by high glass transition temperatures (often not apparent below decomposition), high free volumes (up to ~30%) believed to be interconnected, high accessible surface areas (300-1000 m2 g-1), low densities and very high gas permeabilities relative to conventional materials, including those employed commercially like , polycarbonates and CA, due to largely increased solubility and diffusion contributions to transport. PIMs are sometimes categorized as “super-glassy” materials because of containing significantly larger excess free volumes than conventional low-free-volume polymers, as indicated in Table 3.1 [23-

26]. This is due to the design of torsion-resistant backbones equipped with bulky pendants (e.g., trimethylsilyl groups in PTMSP) and/or sites of contortion (e.g., spirobisindane moiety in PIM-1 and PIM-PI-8) that intentionally create free volume by obstructing chain packing.

Table 3.1. Glass transition temperatures for conventional low-free-volume and high-free volume PIMs.

o 3 3 Polymer Tg ( C) FFV (cm /cm polymer)

Polycarbonate 150 0.16

Polysulfone 186 0.16 109

PTMSP >250 0.25-0.30

PIM-1 >400 ~0.25

PIM-PI-8 >400 0.23

3.2. Pore Size and Transport Mechanisms in PIMs

The physical microstructure is the most significant parameter affecting transport through a polymer film [27-31]. The different transport mechanisms are mainly dominated by “pore” size [32, 33]. “Pore” in an amorphous polymer is defined as a region of free volume that either may vary in position or volume on the timescale of penetrant motion [34]. If the pore size is larger than the mean free path of a gas molecule

(i.e., the distance traveled by gas molecules between successive collisions), then convective flow occurs without any separation. As the pore size approaches the mean free path, then collisions with the wall result in Knudsen diffusion and the selectivity for gas

A over B is proportional to the inverse square root of the molecular weights of the gases.

As the size of the pore further decreases into the ultra-microporous region (< 7 Å), it approaches the molecular dimensions of gas molecules and vapors. More complex mechanisms can occur, depending on whether the penetrant involved is highly condensable or not, as shown in Fig. 3.2. PIMs have the potential to provide molecular sieving for gases because of their ultramicroporosity. The region of PIMs porosity is being powerfully researched for simultaneously boosting permeability and selectivity.

For example PTMSP, the earliest solution-processable amorphous PIM reported, has extremely high gas permeabilities coupled with very low gas pair selectivities relative to 110 conventional low-free-volume glassy polymers [35, 36] due to its high BET surface area and significant fraction of large microporous.

Fig. 3.2. Gas transport mechanisms for different membrane types based on “pore” size.

PIMs have complicated distributions in porosity as indicated by the broad amorphous peak often observed in WAXD spectra. Therefore, the concept of interconnected porosity is proposed and illustrated below in Fig. 3.3. Moreover, as explained earlier in Chapter

2, the chain motion in PIMs (glassy polymers) is restricted; hence, the diffusion coefficient is strongly dependent on penetrant size and shape [27, 28]. 111

Fig. 3.3. The assumed microporous structure in non-porous membranes, membranes with isolated porosity, and membranes with interconnected porosity (i.e., PIMs).

3.3. Polymers of Intrinsic Microporosity for Gas Separation

Masuda’s group from Kyoto University (Japan) discovered the first glassy, microporous polymer in 1983, poly(1-trimethylsilyl-1-propyne) (PTMSP) and its chemical structure is shown in Fig. 3.4 [19]. This polymer contains interconnected microporosity with large chain spacing that exhibits diffusion and permeability coefficients 103-106 times larger than those of conventional low-free-volume glassy polymers such as polysulfone, cellulose acetate and Matrimid polyimide [35]. Basically,

2 the high free volume (N2 BET surface area = 700-900 m /g) and large pore size in

PTMSP result in very low selectivity for gas separation applications. Therefore, PTMSP has not found any commercial use in gas separation membrane applications.

Fig. 3.4. Chemical structure of PTMSP. 112

In 2004, other types of ladder polymers of intrinsic microporosity (PIMs) were discovered by Budd’s and McKeown’s groups at the University of Manchester, UK.

These materials have an inherent, interconnected porosity resulting from inefficient packing of polymer chains in the solid state [37, 38]. PIMs represent a new class of materials possessing microporosity, which has pores of less than 20 Å in diameter as defined by IUPAC. Ladder PIM-1 (Fig. 3.5) is one of these materials where a new design strategy was introduced to reduce chain mobility and increased chain spacing to produce glassy materials with improved performance exceeding the permeability/selectivity upper bounds for size-sieving-based applications like O2/N2 and H2/N2 [39]. PIM-1 is a spirobisindane-based benzodioxane ladder polymer which has excellent solution processability and thus good thin-film formation properties. It has large surface area (N2

BET surface area = 700-820 m2 g-1) that offers high permeability due to both high diffusivities and high . Besides that, the small micropores in addition to the chain rigidity result in moderate selectivities which makes PIM-1 a promising polymer for size-sieving membrane applications [24].

Fig. 3.5. Chemical structure of PIM-1.

Remarkably, both PTMSP and PIM-1 have almost similar N2 BET surface area; however, PTMSP has almost 20 times higher N2 permeability than PIM-1 [40]. In 113

addition, PTMSP has a CO2 permeability of ~ 18000 Barrer with CO2/CH4 selectivity of

4.3 [41], whereas, PIM-1 has lower CO2 permeability of ~ 3500 Barrer with higher

CO2/CH4 selectivity of about 10 [24]. This is because PTMSP has a larger fraction of bigger micropores than PIM-1 (> 10 Å), which results in higher N2 and CO2 permeabilities, as shown in Fig. 3.6. Furthermore, the large difference in CO2/CH4 selectivity results from larger fractions of smaller ultramicropores (pores less than 7 Å) in

PIM-1 than PTMSP.

Fig. 3.6. NLDFT analysis of N2 adsorption isotherms for PTMSP and PIM-1 [25].

Similar to PIM-1 design strategy, attempts have been made to introduce high free volume and microporosity into polyimides due to their thermal and chemical stability, solution processability, and mechanical strength [29, 42]. McKeown’s group successfully made PIM-PIs which has produced the most permeable polyimides with performance close to the upper bound for several important gas pairs [43-46]. PIM-PI-8 (Fig. 3.7) is 114 among the best performing in a series of such polymers synthesized by Ghanem et al.

[44]. Thereafter, Rogan et al. synthesized a new series of three novel PIM-PIs and PIM-

PI-10 (Fig. 3.8) exhibited the best gas performance among this series [47]. Both polymers

(PIM-PI-8 and PIM-PI-10) have excellent CO2 permeability and fair gas separation performance for CO2/CH4 compared to the conventional polyimides like Matrimid as presented in Table 3.2.

Fig. 3.7. Chemical structure of PIM-PI-8.

Fig. 3.8. Chemical structure of PIM-PI-10.

Generally, the novel PIM-type polyimides have higher gas permeability than conventional polyimides combined with moderate selectivities. N2 adsorption analysis at

2 -1 77 K showed that PIM-PI-10 had a N2 BET surface area of around 700 m g which was lower than PIM-1 as reported in the literature (700-820 m2 g-1) [47-50]. Also, permeability cofficients for all reported gases for PIM-PI-10 were almost 2-3 times lower in permeability than PIM-1. On the other hand, PIM-PI-10 permeabilities were 115 significantly higher than other high-free-volume polyimides such as 6FDA-TMPD (Fig.

3.9).

Fig. 3.9. Chemical structure of 6FDA-TMPD.

Smilar to PIM-PI-8 design strategy, free volume and microporosity were introduced into polyimides by replacing the spirobisindane with triptycene building unit which has a rigid three-dimensional geometry with a high energy barrier to molecular deformation makes void spaces between the aromatic bridges and internal to the triptycene moiety.

KAUST-PI-1 and KAUST-PI-7 (Fig. 3.10) were among the best in a series of such polymers synthesized and characterized by Ghanem et al. [51, 52]. The triptycene-based polyimide (KAUST-PI-1) showed significantly higher ultramicroporosity than PTMSP and PIM-1, as shown in Fig 3.11 [51], which improved CO2/CH4, H2/CH4 and O2/N2 selectivities (Table 3.2). 116

Fig. 3.10. Chemical structures of KAUST-PI-1 and KAUST-PI-7.

Fig. 3.11. a) Nitrogen adsorption isotherms measured for PTMSP, PIM-1 and KAUST-

PI-1 at 77 K. b) NLDFT-analyzed pore size distributions based on carbon slit-pore geometry, showing shifts to ultra-microporosity in KAUST-PI-1 [51].

The pure-gas permeabilities and ideal selectivities of candidate polymers for CO2/CH4 and O2/N2 separation discussed earlier in this section are summarized in Table 3.2. 117

Table 3.2. Pure-gas permeabilities and ideal selectivities of various PIMs.

Pure-gas Permeability (Barrer) Ideal Selectivity Polymer (훼)

He H2 N2 O2 CH4 CO2 CO2/ CH4 O2/ N2

PTMSPa [53] 4100 6900 2000 4000 4200 18000 4.3 2.0

PIM-1b [24] 1061 2332 238 786 360 3496 10 3.3

PIM-PI-8 [43] 660 1600 160 545 260 3700 14.2 3.4

PIM-PI-10 [47] 300 670 84 270 168 2154 12.8 3.2

6FDA-TMPD [47] - 549 36 122 28 440 15.7 3.4

KAUST-PI-1 [51] 1771 3983 107 627 105 2389 23 5.9

KAUST-PI-7 [52] 1371 3198 225 842 354 4391 12 3.7

a at T= 25 °C b at 35 °C and 4 atm

An analysis of the data in Table 3.2 reveals several basic strategies that can be drawn

from the designs of these PIMs, including the incorporation of bulky side groups on a

rigid backbone (PTMSP) and the integration of sites of contortion (PIM-1, PIM-PI-8 and

KAUST-PI-1), all which disturb efficient chain packing. PIMs provide a promising,

tunable platform for the design of materials with properties targeted in the above sections,

namely higher permeabilities with similar or higher selectivities compared to

conventional low free-volume, commercial polymers. Moreover, thermal treatment of

amorphous polymers is a traditional approach to make microporous materials. High

temperature pyrolysis produces highly selective and permeable amorphous carbons.

Another method to produce microporous polymers involves thermal treatment (> 400 oC)

of ortho-hydroxyl-functionalized polyimides, which results in a thermal rearrangement to 118 polybenzoxazoles, of which some have shown promising gas separation properties [54,

55]. More discussion about thermal treatment will be presented in Chapter 8.

3.4. Improving Microporous Polyimides for Natural Gas Sweetening

To date, there is no predictive quantitative formula available between polymer structure (such as PIM-polyimide) and gas transport properties to design a membrane material for a desired separation by molecular modeling. However, in 1988, Koros’ group introduced comprehensive qualitative criteria based on data for a series of aromatic polyimides [56]. They concluded that any change to the polymer repeat unit structure should balance gains/losses in intra-chain mobility and inter-chain spacing. Several examples emphasized that polyimides with much better performance can be obtained if both intra-chain mobility and inter-chain packing were controlled. Actually, most literature on the effects of the structures of the diamines and dianhydrides that were reacted to form polyimides were best rationalized by their influences on chain mobility and chain packing [28].

Strategies towards more permeable and selective polymers were developed and researchers primarily used measurable parameters like Tg, FFV, BET and d-spacing to establish structure/property relationships between chain-packing/chain-mobility and transport properties like P, D and S [27-29]. The Tg was taken as a measure of polymer

“rigidity” which was assumed to be crucial to disturbing packing and improving selectivity. However, it became apparent that the large-scale motions associated with a glass-rubber transition were not necessarily those that controlled diffusion in the glassy state. FFV was calculated indirectly from density data and applying group contribution. 119

However, it is only a rough representation of the amount of free volume. Similarly, d- spacing is only an estimation of the average inter-chain spacing approximated from the peak of a broad halo in the WAXD (wide angle x-ray diffraction) spectrum of an amorphous polymer.

The majority of the work aimed at development of novel organic polymeric materials with higher permeabilities and selectivities has focused on systematically varying the chemical structure to induce increases in diffusivity (penetrant mobility) and diffusivity selectivity (sieving) [28, 57]. Freeman and Robeson have established a theoretical foundation for the empirical upper bound relationship that demonstrated a vast contribution of diffusivity relative to solubility in the slope of the upper bound [31, 58].

PIMs, given their high free volumes and rigid backbones, have therefore gained huge interest as a platform for boosting diffusivity while maintaining or advancing diffusivity selectivity.

The targeted polymeric molecular sieve membrane materials for CO2/CH4 separation in this research were PIM-PIs. Unfortunately, the CO2/CH4 separation properties of most previously reported PIM-PIs were located below the Robeson CO2/CH4 upper bound with lower selectivity than the existing commercial materials [44]. However, by analyzing the structures of the current PIM-PIs, KAUST-PIs and 6FDA-polyimides, a general scheme for novel structure design can be formulated to further enhance permeability and selectivity compared to the existing CO2/CH4 commercial membrane materials. The main structural elements of previously designed polyimides were based on the spiro or triptycene center serving as the bulky kink in the structure and the single bond to the imide ring. It is proposed that the changes in the angle at the rigid contortion site in the 120 dianhydride or diamine can be utilized to improve gas separation performance by properly tuning the pore structure. The nature of the polycondensation reaction between a dianhydride and a diamine (Fig. 3.12) to form polyimides makes them adjustable to systematic structural changes and thus organized studies of structure/property relationships.

Fig. 3.12. Typical polyimide synthesis procedure.

3.5. Triptycene-Based Polymers

A new route towards introducing internal molecular free volume and rigidity into a polymer matrix can be achieved by incorporation of iptycene into the polymer chain.

Iptycenes are a unique class of a [2,2,2]-ring system in which the bridges are aromatic rings [59, 60]. The simplest member of this class is triptycene which was first synthesized in 1942 by Bartlett and later by Wittig et al. (Fig. 3.13) [61]. 121

Free volume and rigidity are intrinsic properties of the triptycene moiety due to the rigid three-dimensional geometry that produce a 120° angle [62]. A high energy barrier to molecular deformation creates void spaces between the aromatic bridges and the internal building block of the triptycene moiety which was defined as “internal free volume”

(IFV) by Swager’s group [60, 62]. Moreover, the non-planar bulky nature of the moiety prevents efficient packing of the polymer chains [63].

Fig. 3.13. Triptycene molecule with high internal free volume [64].

The Swager group was successful in synthesizing a series of amorphous, highly soluble 2,6-triptycene-diamine-based polyimides showing no apparent Tg with surface area ranging from 39 to 430 m2 g-1 [65]. One of these polyimides 6FDA-DATRI was formed between 2,6-diaminotriptycene and 4,4-(hexafluoroisopropylidene) diphtalic 122 dianhydride (6FDA). 6FDA-DATRI was first evaluated by Park’s group and showed promising gas separation properties [66]. In summary, incorporating of the iptycene moieties into a polymeric backbone is projected to introduce well-defined free volume by the bulky and rigid three-dimensional structures.

3.6. PIMs Physical Aging

Naturally, glassy polymers have non-equilibrium excess free volume (Fig. 3.14), which can slowly relax over time [67]. The relaxation of the excess free volume is referred to as physical aging phenomena and the mechanism is mainly a volumetric densification of the matrix caused by small-scale molecular motions (below Tg) and originates from the fact that glassy materials are generally out-of-equilibrium [68, 69].

Fig. 3.14. Correlation of polymer specific volume and temperature before and after Tg

[27, 69]. 123

Physical aging is commonly defined by the rate of change in specific volume 푣 of the polymer with time:

푑푣 푣 − 푣 = − ∞ (3.1) 푑푡 τ

The parameter τ is a characteristic timescale for relaxation of the specific volume in the polymer. Typically, more relative densification by aging occurs with increasing polymer free volume [70]. Glassy PIM-PIs are usually affected by time-dependent physical aging that impact their gas separation performance. Specifically, aging leads to reduction in the gas permeability, whereas selectivity typically increases.

3.7. Plasticization in PIMs

As explained in Chapter 2, plasticization is a macroscopic change in the physical properties of glassy polymers (e.g., polyphenylene oxide, polysulfone, polyetherimide, polycarbonate, cellulose acetate, poly(methyl methacrylate), polyimides) upon exposure to high CO2 pressures, which affect the membrane performance. Plasticization occurs when the CO2 concentration in a polymer is high enough to increase free volume and chain mobility [71]. The minimum pressure required to introduce a continuous increase in CO2 permeability as function of pressure is called the plasticization pressure. The evidence of plasticization exists in the gas transport properties [72-76] of the glassy polymers beyond a certain “plasticization pressure” including: (i) pressure-dependent increases in pure CO2 permeability, increases in mixed-gas CH4 permeability and reduction in mixed-gas CO2/CH4 selectivity; (ii) time-dependent increases in CO2 permeability at a given pressure and (iii) time dependent sorption [77]. To overcome the 124

methane loss in membrane-based CO2/CH4 separation, plasticization should be minimized. Successful strategies [78, 79] that solve CO2-induced plasticization are mainly to restrict chain mobility, including strong covalent connections between the chains resulting from chemical crosslinking [80-82] or high-temperature-based cross- linking by polybenzoxazole formation [54, 83] and pyrolysis [84, 85]. They also include increasing inter-chain interactions via the formation of charge-transfer complexes (CTCs) in polyimides and polyamide-imides [86-91] and the integration of polar moieties that engage in strong hydrogen bonding that increases the cohesive energy density (CED) of the microstructure [46, 90-94].

Essentially, the separation of CO2/CH4 mixtures by glassy polymers relies on a size- sieving mechanism, such that dilation of the polymer induced by plasticization compromises its size- and shape-discriminating ability and, hence, selectivity. However, plasticization is not an obstacle in commercial membrane use; in fact, state-of-the-art commercial cellulose acetate membranes plasticize during industrial use in natural gas treatment, resulting in mixed-gas CO2/CH4 selectivities that are effectively only 10-15

[95].

On the other hand, researchers [96] have considered plasticization-resistant membranes made from an emerging class of solution-processable materials known as polymers of intrinsic microporosity (PIMs) given their highly restricted intra-chain mobility [20, 22, 40, 97]. As discussed earlier in this chapter, PIMs comprise fused-ring and contorted backbones that pack inefficiently in the solid state to trap microporosity.

Consequently, they often show no Tg below their decomposition temperatures and tend to 125 exhibit lower mechanical flexibility. Moreover, the intermolecular interactions could limit CO2-induced dilation of the membrane.

3.8. Materials and Nomenclature

The previously reported PIM-PI platform has key structural features that can be varied by the modular nature of the polyimide synthesis involving a polycondensation reaction between a diamine and a dianhydride. Structural parameters that can be varied include the contortion site serving as the kink in the structure (i.e., spirobisindane or triptycene), the distance between kinks, the planarity of the ladder arm, and the flexibility around the single bonds. All these factors can significantly affect the extent of contortion and intra-chain rigidity, and consequently impact the chain packing and chain mobility which are crucial to gas transport across polymeric membranes.

3.8.1. Typical Synthetic Approach

In this work, PIM-polyimides were synthesized by a one-pot polycondensation reaction of a dianhydride and diamine with high molecular weight without the need to first cast the polyamic acid solution and then thermally imidize the membrane in a separate step [43, 51, 98, 99]. Briefly, equimolar amounts of the dianhydride and diamine monomers were stirred in freshly distilled m-cresol (7 mL) under a nitrogen atmosphere for 5 minutes to obtain a clear solution. The reaction temperature was raised gradually to

200 °C and the temperature was held for 4 hours under steady nitrogen flow. Water formed during the reaction was removed by addition of toluene. After cooling to 80 °C, the polyimide solution was precipitated in methanol (200 mL). The crude polymer was 126 filtered and dried in an oven at 120 °C. Purification of the polymer was done by two additional reprecipitations from a suitable solvent (commonly chloroform) into methanol.

3.8.2. Research Focus

This research was focused on studying the gas transport properties of novel microporous polyimides containing the five different triptycene moieties including: the non-substituted triptycene, its extended iptycene analog, and dimethyl-, diethyl- and diisopropyl-bridgehead substituted triptycene (Fig. 3.15). Ghanem et al. demonstrated that the gas adsorption properties of ladder-type polymeric networks containing triptycene can be tuned by the 9,10-bridgehead substituent groups [12]. Likely, gas transport properties could be adjusted by changing the substituent groups to the triptycene unit. 127

Fig. 3.15. Triptycene building units used in this work.

The five building blocks (Fig. 3.15) were utilized to evaluate the followings:

1. The effect of an additional benzene ring (stacked) on the extended iptycene moiety

(Fig. 3.16).

2. The effect of side groups to the triptycene moiety, specifically, several major

structural investigations were made on triptycene’s 9,10-bridgehead position in a non-

extended triptycene-based dianhydride (Fig. 3.17) by the substitution with alkyl

groups of various sizes and shapes while using the same diamine to correlate the

general effect of the R-group on gas transport and identify the most promising

structure for further investigations. The substituent should offer the ability to tune the

microstructure as evidenced by gas adsorption and transport properties in the

triptycene-derived PIM-PI [12, 52]. 128

Fig. 3.16. Chemical structures of triptycene and extended iptycene building units showing internal free volume (IFV) [100].

Fig. 3.17. Key structural features in the proposed novel polyimides.

Furthermore, with a fixed bridgehead substituent, the role of intrachain rigidity in mixed-gas CO2/CH4 plasticization behavior was investigated by using a variety of 129 diamines that systematically change the torsional flexibility of the single bond at the imide ring.

As mentioned earlier, the goal of this study was to develop PIM-PIs with an optimal combination of rigidity and adequate microporosity that produces molecular sieving behavior, where diffusivity selectivity is optimized and the detrimental behavior in the presence of plasticizing feed gases like CO2 is reduced. Besides that, physical aging assessments (long-term stability) were performed for all tested membrane materials to evaluate their stability performance over time. In addition, hydroxyl-functionalization and thermal rearrangement of one selected triptycene-based polyimide were utilized to further optimize the selectivity and permeability for CO2/CH4 separation application.

3.8.3. Dianhydrides

The materials used in this study include 6FDA- and triptycene-dianhydride-based polyimides (Table 3.3) and four commercially available diamines (Table 3.4). The structures and naming schemes for the dianhydrides used to synthesize the polyimides are summarized in Table 3.3. Full details of their synthesis and characterizations are provided in the literature [99-101].

Table 3.3. Nomenclature and chemical structures of 6FDA- and PIM-type triptycene- based (TDA) dianhydrides used in the polyimides synthesis.

130

Nomenclature Structure

6FDA

TDA0

TDA1

TDA2

131

TDAi3

3.8.4. Diamines

A series of commercial diamines, which were used in this research, are summarized along with their naming schemes in Table 3.4.

Table 3.4. Nomenclature and chemical structures of commercially available diamines used in the polyimides synthesis.

Nomenclatu Name Structure

re

DAT1 2,6-Diaminotriptycene

132

DAT2 2,6-Diaminobenzotriptycene

DMN 3,3′-Dimethyl-naphthidine

2,2-Bis(3-amino-4- APAF hydroxyphenyl)- hexafluoropropane

133

3.9. References

[1] J. Weitkamp, K.S.W. Sing, F. Schüth, Handbook of porous solids, Wiley-Vch, 2002.

[2] A.W. van den Berg, C.O. Areán, Materials for hydrogen storage: current research trends and perspectives, Chemical Communications, (2008) 668-681.

[3] M. Eddaoudi, J. Kim, N. Rosi, D. Vodak, J. Wachter, M. O'Keeffe, O.M. Yaghi,

Systematic design of pore size and functionality in isoreticular MOFs and their application in methane storage, Science, 295 (2002) 469-472.

[4] H.J. Mackintosh, P.M. Budd, N.B. McKeown, Catalysis by microporous phthalocyanine and porphyrin network polymers, Journal of Materials Chemistry, 18

(2008) 573-578.

[5] R. Xu, W. Pang, J. Yu, Q. Huo, J. Chen, Chemistry of zeolites and related porous materials: synthesis and structure, John Wiley & Sons, 2009.

[6] P.A. Wright, Microporous framework solids, Royal Society of Chemistry, 2008.

[7] N.L. Rosi, J. Eckert, M. Eddaoudi, D.T. Vodak, J. Kim, M. O'Keeffe, O.M. Yaghi,

Hydrogen storage in microporous metal-organic frameworks, Science, 300 (2003) 1127-

1129.

[8] R. Robson, Design and its limitations in the construction of bi-and poly-nuclear coordination complexes and coordination polymers (aka MOFs): a personal view, Dalton

Transactions, (2008) 5113-5131. 134

[9] A.P. Cote, A.I. Benin, N.W. Ockwig, M. O'Keeffe, A.J. Matzger, O.M. Yaghi,

Porous, crystalline, covalent organic frameworks, Science, 310 (2005) 1166-1170.

[10] A.P. Cote, H.M. El-Kaderi, H. Furukawa, J.R. Hunt, O.M. Yaghi, Reticular synthesis of microporous and mesoporous 2D covalent organic frameworks, Journal of the American Chemical Society, 129 (2007) 12914-12915.

[11] H.M. El-Kaderi, J.R. Hunt, J.L. Mendoza-Cortés, A.P. Côté, R.E. Taylor, M.

O'Keeffe, O.M. Yaghi, Designed synthesis of 3D covalent organic frameworks, Science,

316 (2007) 268-272.

[12] B.S. Ghanem, M. Hashem, K.D.M. Harris, K.J. Msayib, M.C. Xu, P.M. Budd, N.

Chaukura, D. Book, S. Tedds, A. Walton, N.B. McKeown, Triptycene-based polymers of intrinsic microporosity: organic materials that can be tailored for gas adsorption,

Macromolecules, 43 (2010) 5287-5294.

[13] J. Germain, J.M.J. Frechet, F. Svec, Nanoporous polymers for hydrogen storage,

Small 5, 10 (2009): 1098-1111.

[14] J. Germain, F. Svec, J.M.J. Frechet, Preparation of size-selective nanoporous polymer networks of aromatic rings: potential adsorbents for hydrogen storage,

Chemistry of Materials, 20 (2008) 7069-7076.

[15] S.W. Yuan, S. Kirklin, B. Dorney, D.J. Liu, L.P. Yu, Nanoporous polymers containing stereocontorted cores for hydrogen storage, Macromolecules, 42 (2009) 1554-

1559. 135

[16] N.B. McKeown, P.M. Budd, Exploitation of intrinsic microporosity in polymer- based materials, Macromolecules, 43 (2010) 5163-5176.

[17] R. Dawson, A.I. Cooper, D.J. Adams, Nanoporous organic polymer networks,

Progress in Polymer Science, 37 (2012) 530-563.

[18] F.Y. Li, Y.C. Xiao, T.S. Chung, S. Kawi, High-performance thermally self-cross- linked polymer of intrinsic microporosity (PIM-1) membranes for energy development,

Macromolecules, 45 (2012) 1427-1437.

[19] T. Masuda, E. Isobe, T. Higashimura, K. Takada, Poly[1-(trimethylsilyl)-1-propyne]

- a new high polymer synthesized with transition-metal catalysts and characterized by extremely high gas-permeability, Journal of the American Chemical Society, 105 (1983)

7473-7474.

[20] N.B. McKeown, P.M. Budd, Polymers of intrinsic microporosity (PIMs): organic materials for membrane separations, heterogeneous catalysis and hydrogen storage,

Chemical Society Reviews, 35 (2006) 675-683.

[21] O. Ilinitch, V. Fenelonov, A. Lapkin, L. Okkel, V. Terskikh, K. Zamaraev, Intrinsic microporosity and gas transport in polyphenylene oxide polymers, Microporous and

Mesoporous materials, 31 (1999) 97-110.

[22] N.B. McKeown, Polymers of intrinsic microporosity, ISRN Materials Science, 2012

(2012).

[23] L. Toy, K. Nagai, B. Freeman, I. Pinnau, Z. He, T. Masuda, M. Teraguchi, Y.P.

Yampolskii, Pure-gas and vapor permeation and sorption properties of poly [1-phenyl-2- 136

[p-(trimethylsilyl) phenyl] acetylene](PTMSDPA), Macromolecules, 33 (2000) 2516-

2524.

[24] S. Thomas, I. Pinnau, N.Y. Du, M.D. Guiver, Pure- and mixed-gas permeation properties of a microporous spirobisindane-based ladder polymer (PIM-1), Journal of

Membrane Science, 333 (2009) 125-131.

[25] R. Swaidan, Intrinsically microporous polymer membranes for high performance gas separation, Ph.D. dissertation, King Abdullah University of Science and Technology,

Saudi Arabia (2014).

[26] M. Aguilar‐Vega, D. Paul, Gas transport properties of polycarbonates and polysulfones with aromatic substitutions on the bisphenol connector group, Journal of

Polymer Science Part B: Polymer Physics, 31 (1993) 1599-1610.

[27] S. Matteucci, Y. Yampolskii, B.D. Freeman, I. Pinnau, Transport of gases and vapors in glassy and rubbery polymers, in Materials Science of Membranes for Gas and

Vapor Separation, Wiley, Chichester, UK (2006) 1-2.

[28] K. Ghosal, B.D. Freeman, Gas separation using polymer membranes: an overview,

Polymers for Advanced Technologies, 5 (1994) 673-697.

[29] K. Tanaka, K.-I. Okamoto, Structure and transport properties of polyimides as materials for gas and vapor membrane separation, John Wiley & Sons: West Sussex,

2006.

[30] W.J. Koros, G.K. Fleming, Membrane-based gas separation, Journal of Membrane

Science, 83 (1993) 1-80. 137

[31] B.D. Freeman, Basis of permeability/selectivity tradeoff relations in polymeric gas separation membranes, Macromolecules, 32 (1999) 375-380.

[32] T.C. Merkel, I. Pinnau, R. Prabhakar, B.D. Freeman, Gas and vapor transport properties of perfluoropolymers, Wiley: Chichester, England, 2006.

[33] R.W. Baker, Membrane technology and applications, John Wiley & Sons, Ltd.,

2004.

[34] J. Wijmans, R. Baker, The solution-diffusion model: a review, Journal of Membrane

Science, 107 (1995) 1-21.

[35] I. Pinnau, L. Toy, Transport of organic vapors through poly (1-trimethylsilyl-1- propyne), Journal of Membrane Science, 116 (1996) 199-209.

[36] C.L. Staiger, S.J. Pas, A.J. Hill, C.J. Cornelius, Gas separation, free volume distribution, and physical aging of a highly microporous spirobisindane polymer,

Chemistry of Materials, 20 (2008) 2606-2608.

[37] P.M. Budd, K.J. Msayib, C.E. Tattershall, B.S. Ghanem, K.J. Reynolds, N.B.

McKeown, D. Fritsch, Gas separation membranes from polymers of intrinsic microporosity, Journal of Membrane Science, 251 (2005) 263-269.

[38] P.M. Budd, N.B. McKeown, D. Fritsch, Free volume and intrinsic microporosity in polymers, Journal of Materials Chemistry, 15 (2005) 1977-1986.

[39] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400. 138

[40] P.M. Budd, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib, C.E.

Tattershall, Polymers of intrinsic microporosity (PIMs): robust, solution-processable, organic nanoporous materials, Chemical Communications, (2004) 230-231.

[41] L.M. Robeson, Correlation of separation factor versus permeability for polymeric membranes, Journal of Membrane Science, 62 (1991) 165-185.

[42] N.Y. Du, H.B. Park, M.M. Dal-Cin, M.D. Guiver, Advances in high permeability polymeric membrane materials for CO2 separations, Energy & Environmental Science, 5

(2012) 7306-7322.

[43] B.S. Ghanem, N.B. McKeown, P.M. Budd, J.D. Selbie, D. Fritsch, High- performance membranes from polyimides with intrinsic microporosity, Advanced

Materials, 20 (2008) 2766-2771.

[44] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[45] J. Weber, O. Su, M. Antonietti, A. Thomas, Exploring polymers of intrinsic microporosity- microporous, soluble polyamide and polyimide, Macromolecular Rapid

Communications, 28 (2007) 1871-1876.

[46] X.H. Ma, R. Swaidan, Y. Belmabkhout, Y.H. Zhu, E. Litwiller, M. Jouiad, I. Pinnau,

Y. Han, Synthesis and gas transport properties of hydroxyl-functionalized polyimides with intrinsic microporosity, Macromolecules, 45 (2012) 3841-3849. 139

[47] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820.

[48] P.M. Budd, E.S. Elabas, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib,

C.E. Tattershall, D. Wang, Solution-processed, organophilic membrane derived from a polymer of intrinsic microporosity, Advanced Materials, 16 (2004) 456-459.

[49] P.M. Budd, N.B. McKeown, B.S. Ghanem, K.J. Msayib, D. Fritsch, L. Starannikova,

N. Belov, O. Sanfirova, Y. Yampolskii, V. Shantarovich, Gas permeation parameters and other physicochemical properties of a polymer of intrinsic microporosity: polybenzodioxane PIM-1, Journal of Membrane Science, 325 (2008) 851-860.

[50] N.B. McKeown, B.S. Ghanem, K.J. Msayib, P.M. Budd, C.E. Tattershall, K.

Mahmood, S. Tan, D. Book, H.W. Langmi, A. Walton, Towards polymer-based hydrogen storage materials: engineering ultramicroporous cavities within polymers of intrinsic microporosity, Angewandte Chemie-International Edition, 45 (2006) 1804-1807.

[51] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra‐microporous triptycene‐ based polyimide membranes for high‐performance gas separation, Advanced Materials,

26 (2014) 3688-3692.

[52] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational design of intrinsically ultramicroporous polyimides containing bridgehead-substituted triptycene for highly selective and permeable gas separation membranes, Macromolecules, 47

(2014) 5104-5114. 140

[53] K. Takada, H. Matsuya, T. Masuda, T. Higashimura, Gas-permeability of polyacetylenes carrying substituents, Journal of Applied Polymer Science, 30 (1985)

1605-1616.

[54] H.B. Park, C.H. Jung, Y.M. Lee, A.J. Hill, S.J. Pas, S.T. Mudie, E. Van Wagner,

B.D. Freeman, D.J. Cookson, Polymers with cavities tuned for fast selective transport of small molecules and ions, Science, 318 (2007) 254-258.

[55] X. Ma, O. Salinas, E. Litwiller, I. Pinnau, Pristine and thermally-rearranged gas separation membranes from novel o-hydroxyl-functionalized spirobifluorene-based polyimides, Polymer Chemistry, 5 (2014) 6914-6922.

[56] T. Kim, W. Koros, G. Husk, K. O'Brien, Relationship between gas separation properties and chemical structure in a series of aromatic polyimides, Journal of

Membrane Science, 37 (1988) 45-62.

[57] J. McHattie, W.J. Koros, D.R. Paul, Gas transport properties of polysulphones: 1. role of symmetry of methyl group placement on bisphenol rings, Polymer, 32 (1991) 840-

850.

[58] L.M. Robeson, Z.P. Smith, B.D. Freeman, D.R. Paul, Contributions of diffusion and solubility selectivity to the upper bound analysis for glassy gas separation membranes,

Journal of Membrane Science, 453 (2014) 71-83.

[59] T.M. Swager, Iptycenes in the design of high performance polymers, Accounts of

Chemical Research, 41 (2008) 1181-1189. 141

[60] J.H. Chong, M.J. MacLachlan, Iptycenes in supramolecular and materials chemistry,

Chemical Society Reviews, 38 (2009) 3301-3315.

[61] N.T. Tsui, A.J. Paraskos, L. Torun, T.M. Swager, E.L. Thomas, Minimization of internal molecular free volume: a mechanism for the simultaneous enhancement of polymer stiffness, strength, and ductility, Macromolecules, 39 (2006) 3350-3358.

[62] T.M. Long, T.M. Swager, Minimization of free volume: alignment of triptycenes in liquid crystals and stretched polymers, Advanced Materials, 13 (2001) 601-604.

[63] Z. Chen, T.M. Swager, Synthesis and characterization of poly (2, 6-triptycene),

Macromolecules, 41 (2008) 6880-6885.

[64] C.-F. Chen, Novel triptycene-derived hosts: synthesis and their applications in supramolecular chemistry, Chemical Communications, 47 (2011) 1674-1688.

[65] S.A. Sydlik, Z. Chen, T.M. Swager, Triptycene polyimides: soluble polymers with high thermal stability and low refractive indices, Macromolecules, 44 (2011) 976-980.

[66] Y.J. Cho, H.B. Park, High performance polyimide with high internal free volume elements, Macromolecular Rapid Communications, 32 (2011) 579-586.

[67] A. Eisenberg, M. Shen, Recent advances in glass transitions in polymers, Rubber

Chemistry and Technology, 43 (1970) 156-170.

[68] P.H. Pfromm, The impact of physical aging of amorphous glassy polymers on gas separation membranes, Materials Science of Membranes for Gas and Vapour Separation,

(2006) 293-306. 142

[69] D. Cangialosi, V.M. Boucher, A. Alegría, J. Colmenero, Physical aging in polymers and polymer nanocomposites: recent results and open questions, Soft Matter, 9 (2013)

8619-8630.

[70] D.F. Sanders, Z.P. Smith, R. Guo, L.M. Robeson, J.E. McGrath, D.R. Paul, B.D.

Freeman, Energy-efficient polymeric gas separation membranes for a sustainable future: a review, Polymer, 54 (2013) 4729-4761.

[71] A. Bos, I. Pünt, M. Wessling, H. Strathmann, CO2-induced plasticization phenomena in glassy polymers, Journal of Membrane Science, 155 (1999) 67-78.

[72] J. Chiou, D.R. Paul, Effects of CO2 exposure on gas transport properties of glassy polymers, Journal of Membrane Science, 32 (1987) 195-205.

[73] M. Wessling, S. Schoeman, T. Van der Boomgaard, C. Smolders, Plasticization of gas separation membranes, Gas Separation & Purification, 5 (1991) 222-228.

[74] M. Wessling, I. Huisman, T. Boomgaard, C. Smolders, Time‐dependent permeation of carbon dioxide through a polyimide membrane above the plasticization pressure,

Journal of Applied Polymer Science, 58 (1995) 1959-1966.

[75] E. Sanders, Penetrant-induced plasticization and gas permeation in glassy polymers,

Journal of Membrane Science, 37 (1988) 63-80.

[76] E. Sanders, S. Jordan, R. Subramanian, Penetrant-plasticized permeation in polymethylmethacrylate, Journal of Membrane Science, 74 (1992) 29-36.

[77] K.-i. Okamoto, K. Tanaka, T. Shigematsu, H. Kita, A. Nakamura, Y. Kusuki,

Sorption and transport of carbon dioxide in a polyimide from 3, 3′, 4, 4′- 143 biphenyltetracarboxylic dianhydride and dimethyl-3, 7-diaminodibenzothiophene-5, 5′- dioxide, Polymer, 31 (1990) 673-678.

[78] N. Du, H.B. Park, M.M. Dal-Cin, M.D. Guiver, Advances in high permeability polymeric membrane materials for CO2 separations, Energy & Environmental Science, 5

(2012) 7306-7322.

[79] Y. Xiao, B.T. Low, S.S. Hosseini, T.S. Chung, D.R. Paul, The strategies of molecular architecture and modification of polyimide-based membranes for CO2 removal from natural gas-a review, Progress in Polymer Science, 34 (2009) 561-580.

[80] C. Staudt-Bickel, W.J. Koros, Improvement of CO2/CH4 separation characteristics of polyimides by chemical crosslinking, Journal of Membrane Science, 155 (1999) 145-154.

[81] J. Kim, W.J. Koros, D.R. Paul, Effects of CO2 exposure and physical aging on the gas permeability of thin 6FDA-based polyimide membranes: part 2. with crosslinking,

Journal of Membrane Science, 282 (2006) 32-43.

[82] W. Qiu, C.-C. Chen, L. Xu, L. Cui, D.R. Paul, W.J. Koros, Sub-Tg cross-linking of a polyimide membrane for enhanced CO2 plasticization resistance for natural gas separation, Macromolecules, 44 (2011) 6046-6056.

[83] R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, High pressure pure-and mixed-gas separation of CO2/CH4 by thermally-rearranged and carbon molecular sieve membranes derived from a polyimide of intrinsic microporosity, Journal of Membrane Science, 447

(2013) 387-394. 144

[84] D.Q. Vu, W.J. Koros, S.J. Miller, High pressure CO2/CH4 separation using carbon molecular sieve hollow fiber membranes, Industrial & Engineering Chemistry Research,

41 (2002) 367-380.

[85] D.Q. Vu, W.J. Koros, S.J. Miller, Effect of condensable impurity in CO2/CH4 gas feeds on performance of mixed matrix membranes using carbon molecular sieves, Journal of Membrane Science, 221 (2003) 233-239.

[86] A. Bos, I. Pünt, M. Wessling, H. Strathmann, Plasticization-resistant glassy polyimide membranes for for CO2/CO4 separations, Separation and Purification

Technology, 14 (1998) 27-39.

[87] X. Duthie, S. Kentish, S. Pas, A. Hill, C. Powell, K. Nagai, G. Stevens, G. Qiao,

Thermal treatment of dense polyimide membranes, Journal of Polymer Science Part B:

Polymer Physics, 46 (2008) 1879-1890.

[88] M. Das, W.J. Koros, Performance of 6FDA–6FpDA polyimide for propylene/propane separations, Journal of Membrane Science, 365 (2010) 399-408.

[89] J.T. Vaughn, W.J. Koros, J. Johnson, O. Karvan, Effect of thermal annealing on a novel polyamide–imide polymer membrane for aggressive acid gas separations, Journal of Membrane Science, 401 (2012) 163-174.

[90] H. Kawakami, M. Mikawa, S. Nagaoka, Gas transport properties in thermally cured aromatic polyimide membranes, Journal of Membrane Science, 118 (1996) 223-230. 145

[91] J. Vaughn, W. Koros, Effect of the amide bond diamine structure on the CO2, H2S, and CH4 transport properties of a series of novel 6FDA-based polyamide–imides for natural gas purification, Macromolecules, 45 (2012) 7036-7049.

[92] Y. Hirayama, T. Yoshinaga, Y. Kusuki, K. Ninomiya, T. Sakakibara, T. Tamari,

Relation of gas permeability with structure of aromatic polyimides II, Journal of

Membrane Science, 111 (1996) 183-192.

[93] R. Swaidan, B.S. Ghanem, E. Litwiller, I. Pinnau, Pure-and mixed-gas CO2/CH4 separation properties of PIM-1 and an amidoxime-functionalized PIM-1, Journal of

Membrane Science, 457 (2014) 95-102.

[94] W. Qiu, L. Xu, C.-C. Chen, D.R. Paul, W.J. Koros, Gas separation performance of

6FDA-based polyimides with different chemical structures, Polymer, 54 (2013) 6226-

6235.

[95] R.W. Baker, Future directions of membrane gas separation technology, Industrial &

Engineering Chemistry Research, 41 (2002) 1393-1411.

[96] R. Chern, W. Koros, E. Sanders, R. Yui, “Second component” effects in sorption and permeation of gases in glassy polymers, Journal of Membrane Science, 15 (1983)

157-169.

[97] P.M. Budd, N.B. McKeown, Highly permeable polymers for gas separation membranes, Polymer Chemistry, 1 (2010) 63-68.

[98] A.J. Bard, L.R. Faulkner, Fundamentals and applications, Electrochemical Methods,

2 (2001). 146

[99] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232.

[100] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[101] F. Alghunaimi, B. Ghanem, N. Alaslai, M. Mukaddam, I. Pinnau, Triptycene dimethyl-bridgehead dianhydride-based intrinsically microporous hydroxyl- functionalized polyimide for natural gas upgrading, Journal of Membrane Science, 520

(2016) 240-246.

147

Chapter 4. Gas Permeation and Physical Aging Properties of Iptycene Diamine-

Based Microporous Polyimides1

4.1. Abstract

The gas permeation properties of two 6FDA-dianhydride-based polyimides prepared from 2,6-diaminotriptycene (6FDA-DAT1) and its extended iptycene analog (6FDA-

DAT2) are reported. The additional benzene ring on the extended triptycene moiety in

6FDA–DAT2 increases the free volume over 6FDA-DAT1 and reduces the chain packing efficiency. The BET surface area based on nitrogen adsorption in 6FDA–DAT2 (450

2 2 m /g) is ~40% greater than that of 6FDA-DAT1 (320 m /g). 6FDA-DAT1 shows a CO2 permeability of 120 Barrer and CO2/CH4 selectivity of 38, whereas 6FDA-DAT2 exhibits a 75% increase in CO2 permeability to 210 Barrer coupled with a moderate decrease in selectivity (CO2/CH4 = 30). Interestingly, minimal physical aging was observed over 150 days for both polymers and attributed to the high internal free volume of the shape- persistent iptycene geometries. The aged polyimides maintained CO2/CH4 selectivities of

25-35 along with high CO2 permeabilities of 90-120 Barrer up to partial CO2 pressures of

10 bar of an aggressive 50:50 CO2:CH4 mixed-gas feed, suggesting potential application in membranes for natural gas sweetening.

1Portions of this chapter were adopted from Alghunaimi, F.; Ghanem, B.; Alaslai, N.; Swaidan, R.; Litwiller, E.; Pinnau, I. Journal of Membrane Science 2015 (490) 321-327. Alghunaimi performed gas permeation, sorption and characterization and wrote the manuscript. Ghanem synthesized the polymers. Litwiller, Alaslai and Swaidan directed construction of permeation apparatus and edited the manuscript. Pinnau supervised the work and edited the manuscript. 148

4.2. Introduction

Recently developed polymers of intrinsic microporosity (PIM) have demonstrated excellent gas separation performance often exceeding the latest 2008 upper bounds for important gas-pairs [1-7]. These amorphous glassy polymers are characterized by: (i) very high thermal stability; (ii) solubility in organic solvents; (iii) high BET surface area

(up to ~ 1000 m2/g); (iv) microporosity (pore size < 20 Å) as well as ultra-microporosity

(< 7 Å). One class of PIMs is based on high-free-volume aromatic polyimides (PIM-PIs) that contain rigid contortion sites either in their dianhydride and/or diamine moiety.

Examples of such molecular building blocks are: (i) spirobisindane dianhydrides [3, 8, 9] and diamines [10, 11], ethanoanthracene dianhydride [2, 12], spirobifluororene dianhydride [13] and diamines [14, 15], Tröger’s base diamines [16-18], and triptycene dianhydrides [6, 19] and diamines [11, 20, 21].

Triptycene is the simplest iptycene and frequently used in polymer building units due to its unique rigid and contorted structure with phenyl rings. Triptycene, as shown in Fig.

3.13 and 4.1, has a rigid 3-dimensional and fully aromatic structure which leads to poor polymer chain packing, and, in turn, results in high internal free volume (IFV) [22, 23]. 149

Fig. 4.1. Schematic of the chemical structure of triptycene and its internal free volume

(IFV).

An extended iptycene building block, shown in Fig. 3.13 and 4.2, demonstrates even larger IFV than triptycene due to the addition of a benzene ring to the transverse aromatic arm in the 9,10 position of the triptycene unit.

Fig. 4.2. Schematic of the chemical structure of the extended iptycene demonstrating enhanced internal free volume (IFV). 150

Swager’s group communicated the synthesis and characterization of a polyimide series based on 2,6-triptycene diamine (DAT1) and its extended iptycene diamine derivative (DAT2) for optical applications [11]. Cho et al. reported the synthesis, characterization and gas permeation properties of a polyimide based on the reaction of

4,4'-hexafluoroisopropylidene diphthalic anhydride (6FDA) and 2,6-diaminotriptycene

(DAT1) that exhibited good performance for gas separation [20].

In this chapter, the pure-gas permeability, diffusivity, and sorption properties of

6FDA-DAT1 and extended 6FDA-DAT2 are reported and analyzed with respect to their chemical structures. In addition, CO2/CH4 mixed-gas permeation experiments were performed and the effect of long-term physical aging on the gas permeation properties of the polymers was investigated. The polyimides were characterized by GPC and TGA.

Nitrogen and carbon dioxide adsorption experiments were performed to elucidate the micropore structure of the two PIM-PIs.

4.3. Experimental

4.3.1. Polymer Characterization

Fourier transform infrared (FT-IR) measurements were performed using a Varian

670-IR FT-IR spectrometer. Gel permeation chromatography (GPC, Viscotek) was carried out using chloroform as an eluent. Thermogravimetric analysis (TGA, TA Q-

5000) measurements were performed under nitrogen atmosphere. All TGA runs entailed a drying step at 100 °C for 30 minutes followed by a temperature ramp of 3 °C/min up to

800 °C. The BET surface area of the polymers was determined by N2 sorption at -196 °C using a Micromeritics ASAP-2020. Powder polymer samples were degassed under high 151 vacuum at 120 °C for 15 hours prior to analysis. Analysis of the pore size distributions was performed using the NLDFT (Non-Local Density Functional Theory) model using

N2 (at -196 °C) and CO2 (at 0 °C) sorption isotherms for carbon slit pore geometry provided by ASAP 2020 version 4.02 software

4.3.2. Synthesis of Polyimides

6FDA-DAT1 and 6FDA-DAT2 polyimides were synthesized via the cycloimidization reaction between equimolar amounts of 6FDA and the corresponding diamine (Scheme

4.1). To a dry 10 mL Schlenk tube equipped with nitrogen inlet and outlet were added the diamine monomer (1 mmol) and freshly distilled m-cresol (4 ml). After stirring for 5 min an equimolar amount of 6FDA (1.0 mmol) and isoquinoline (0.1 ml) was added. The mixture was stirred at ambient temperature for 15 minutes under a flow of nitrogen and then the temperature was raised gradually to 200 °C and kept at that temperature for 4 hours.

Scheme 4.1. Synthesis of 6FDA-DAT1 and 6FDA-DAT2. Polymer structures and their energy-minimized conformations are shown (Materials Studio 6.0, Accelrys). 152

After cooling, the polymer was precipitated by pouring the reaction mixture into an excess of methanol. The resulting fibrous polymer was collected by filtration and purified by reprecipitation from chloroform into methanol and then dried under vacuum at 120 °C for 20 h to give an off-white powder in quantitative yield.

6FDA-DAT1. Following the above general procedure, 6FDA-DAT1 was prepared from

-1 6FDA and DAT1 diamine. Analysis by GPC (CHCl3): Mn = 16200 g mol , Mw = 33300

-1 2 -1 g mol relative to polystyrene, Mw/Mn = 2.1. BET surface area (N2) = 320 m g , total

3 -1 pore volume = 0.24 cm g at p/po = 0.98, adsorption. TGA analysis: (N2), initial weight loss due to thermal degradation commences at Td = 520 °C (Fig. 4.3).

6FDA-DAT2. Following the above typical procedure, 6FDA-DAT2 was prepared from

-1 6FDA and DAT2 diamine. Analysis by GPC (CHCl3): Mn = 38000 g mol , Mw = 67000

-1 2 -1 g mol relative to polystyrene, Mw/Mn = 1.8. BET surface area (N2) = 450 m g , total

3 -1 pore volume = 0.57 cm g at p/po = 0.98, adsorption. TGA analysis: (N2), initial weight loss due to thermal degradation commences at Td = 510 °C (Fig. 4.3). 153

Fig. 4.3. Thermogravimetric analysis (TGA) of 6FDA-DAT1 and 6FDA-DAT2.

4.3.3. Polymer Film Preparation

Solutions of the polymers in chloroform (5 wt/vol%) were filtered through 0.45 μm polypropylene filters and isotropic films were obtained by slow evaporation of the solvent at room temperature from a levelled glass Petri dish. To remove any traces of residual solvent, the dry membranes were soaked in methanol for 10 h, air-dried, and then post-dried at 120 °C in a vacuum oven for 12 h. The resulting yellowish tough 6FDA–

DAT1 and 6FDA–DAT2 films with thickness of 65 ± 5 μm were used for gas permeability measurements.

4.3.4. Pure- and Mixed-Gas Permeation Experiments

The pure-gas permeability of 6FDA-DAT1 and 6FDA-DAT2 was determined by using the variable pressure/constant volume method. Each polymer film sample was degassed in the permeation cell under vacuum for at least 24 hours. The pure-gas permeability of He, H2, N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar. Pressure- 154

dependence of CO2 and CH4 permeability was determined from 2 to 15 bar. The gas permeability (P) was calculated by equation 2.14 where P is reported in Barrers (1 Barrer

= 10-10 cm3(STP·cm/cm2·s·cmHg). The ideal selectivity for a gas pair is given by equation 2.6.

The apparent diffusion coefficient D (cm2/s) was calculated by D = l2 /6 θ, where θ is the time lag of the permeability measurement and l is the membrane thickness. The solubility coefficient S (cm3(STP)/cm3·cmHg) was then calculated from the solution- diffusion gas transport relationship: S=P/D.

The mixed-gas permeation properties of the membranes were measured using a feed gas mixture of 1:1 CO2/CH4 at 35 °C with a setup similar to that described by O'Brien et al. [24]. Both polyimides were tested at total pressures of 4, 10, 20 and 30 bar, respectively. The mixed-gas permeabilities were calculated by equations 2.16 and 2.17.

Because the downstream pressure was negligible, the mixed-gas CO2/CH4 selectivity was obtained from equation 2.18.

4.4. Results and Discussion

4.4.1. Physical Properties and Microstructures of 6FDA-DAT1 and 6FDA-DAT2

The physical properties of the polyimides from our study and previous reports are listed in Table 4.1. The molecular weights and molecular weight distributions of 6FDA-

DAT-1 vary significantly depending on the reaction conditions used for synthesizing the polyimide.

155

Table 4.1. Physical properties of 6FDA-DAT1 and 6FDA-DAT2.

Polymer Mn (g/mol) PDI Td (°C) BET Surface Area

2 (N2) (m /g)

6FDA-DAT1a 21100 5.2 540 -

6FDA-DAT-1b 29000 1.8 531 68

6FDA-DAT1c 16200 2.1 520 320

6FDA-DAT2b 17000 1.6 540 430

6FDA-DAT2c 38000 1.8 510 450 a [20]; b [11]; c this study.

The influence of the additional benzene ring in the three-dimensional extended iptycene moiety can be assessed from N2 (-196 °C) and CO2 (0 °C) physisorption isotherms (Fig. 4.4 and 4.5). The corresponding NLDFT-derived pore size distributions

(PSDs) are shown in Fig. 4.4. The addition of a benzene ring increased the BET surface area from 320 m2/g for 6FDA–DAT1 to 450 m2/g for 6FDA–DAT2, which is expected based on the geometric structures of triptycene and its extended iptycene derivative (Fig.

4.1 and 4.2). Furthermore, N2 sorption isotherms indicate the development of larger microporosity in the 6FDA–DAT2 polymer. Previously, Swager’s group reported a BET surface area for 6FDA-DAT1 of only 68 m2/g; one possible reason for this lower surface area could be due to traces of residual solvent trapped in the micropores of the polymer. 156

Fig. 4.4. Physisorption isotherms for 6FDA–DAT1 and 6FDA–DAT2 using N2 at -196 oC. Open symbols: adsorption; closed symbols: desorption.

Fig. 4.5. Physisorption isotherms for 6FDA–DAT1 and 6FDA–DAT2 using CO2 at 0 °C.

Open symbols: adsorption; closed symbols: desorption. 157

The PSDs based on CO2 isotherms (Fig. 4.6) showed that the quantity of ultra- micropores (< 7 Å) is similar in both polymers. However, PSDs deduced from N2 isotherms qualitatively indicate that 6FDA–DAT2 contained an increased fraction of micropores larger than 7 Å compared to 6FDA–DAT1. Furthermore, pore sizes within the range of 8 to 12 Å in 6FDA–DAT1 shifted to sizes greater than 12 Å in 6FDA–

DAT2.

Fig. 4.6. NLDFT-based estimated pore size distribution obtained from N2 and CO2 isotherms for 6FDA–DAT1 and 6FDA–DAT2 assuming carbon slit-pore geometry.

4.4.2. Pure-Gas Permeation Properties of 6FDA-DAT1 and 6FDA-DAT2

The pure-gas permeation properties of 6FDA–DAT1 and 6FDA–DAT2 are shown in

Table 4.2. Gas permeabilities of 6FDA-DAT1 obtained in our study were lower and selectivities higher for all gas pairs compared to previously reported data by Cho et al.

[20]. Possible reasons for the different gas permeation properties may include: (i) 158

different polymerization protocols; (ii) different film formation, post-treatment and

drying conditions; (iii) different permeation test conditions. To obtain meaningful

structure/gas permeation property relationships between 6FDA-DAT1 and the extended

6FDA-DAT-2, the same film preparation and permeation test protocols were applied for

each polymer (Table 4.2). 6FDA–DAT2 exhibited higher permeability for all gases than

6FDA–DAT1 coupled with lower selectivity. For example, 6FDA-DAT1 showed a CO2

permeability of 120 Barrer and CO2/CH4 selectivity of 38, whereas 6FDA-DAT2

exhibited a 75% increase in CO2 permeability to 210 Barrer coupled with a moderate

decrease in selectivity (CO2/CH4 = 30). This result is qualitatively consistent with the

BET and PSD results discussed above. Adding a benzene ring to the iptycene moiety in

6FDA-DAT2 reduces efficient chain packing, which, in turn, leads to an increase in the

internal free volume (IFV), and, hence, increases gas permeability.

Table 4.2. Pure-gas permeabilities and ideal selectivities for 6FDA-DAT1 and 6FDA-

DAT2.

Pure-Gas Permeability (Barrer) Ideal Selectivity (훼)

Polymer He H2 N2 O2 CH4 CO2 CO2/CH4 N2/CH4 O2/N2

6FDA-DAT1a 198 257 8.1 39 6.2 189 30 1.3 4.8

Fresh 6FDA–DAT1b 161 198 4.7 25.4 3.2 120 38 1.5 5.4

Aged 6FDA-DAT1c (148) (170) (4.0) (21.4) (2.7) (102) (38) (1.5) (5.4)

Fresh 6FDA–DAT2b 204 281 9.0 43.3 7.1 210 30 1.3 4.8 159

Aged 6FDA–DAT2c (175) (229) (7.7) (40.7) (5.3) (160) (30) (1.4) (5.3)

Test and preparation conditions:

a T=35 °C, 1 bar; 2 hrs methanol soaked; dried under vacuum at 100 °C for 12 h [20].

b T=35 °C, 2 bar, 10 hrs methanol soaked; dried under vacuum at 120 °C for 12 h (this

study). Fresh = one day after drying.

c Same as b after 150 days of aging (this study).

The pure-gas diffusion (D) and solubility (S) coefficients of 6FDA-DAT1 and 6FDA-

DAT2 for N2, O2, CH4 and CO2 were calculated by the time-lag method and are shown in

Table 4.3. The higher gas permeabilities in 6FDA-DAT2 are generally caused by

increased diffusivity coefficients, whereas the solubility coefficients remain essentially

constant. Faster diffusion again indicates the formation of a more open and

interconnected microporosity by the bulkier DAT2 diamine (Fig. 4.6). Relative to ~60-

70% increases in the diffusivity coefficients of O2 (kD=3.46 Å), N2 (kD=3.64 Å) and CO2

(kD=3.3 Å), more than 100% increases were observed in that of the larger CH4 gas

(kD=3.8 Å) which is most impacted by the newly accessable diffusion pathways. Indeed,

the overall reduction in CO2/CH4 permselectivity from 38 in 6FDA-DAT1 to 30 in

6FDA-DAT2 (Table 4.2) is governed by a reduction in CO2/CH4 diffusivity selectivity

(i.e., D(CO2)/D(CH4)). 160

Table 4.3. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for

6FDA-DAT1 and 6FDA-DAT2 based on time-lag method (35 °C; 2 bar).

Diffusion coefficient Solubility coefficient

(10-8 cm2/s) (10-2 cm3(STP)/(cm3 cmHg))

Polymer N2 O2 CH4 CO2 N2 O2 CH4 CO2

6FDA-DAT1 [20] 5.5 23 1.5 13 1.5 1.7 4.1 14.5

6FDA–DAT1a 4.0 15 0.77 8.15 1.2 1.7 4.2 14.8

6FDA–DAT2a 6.4 24.1 1.61 14.4 1.4 1.8 4.4 14.6 a this study.

4.4.3. Pure-and Mixed-Gas CO2/CH4 Permeation Properties after Physical Aging

Microporous glassy polymers typically undergo physical aging or a physical densification of the polymer matrix in which the chains gradually assume a tighter packing arrangement [25, 26]. Non-equilibrium excess free volume resulting from casting procedures or methanol treatment is dissipated and large reductions in permeability along with increases in selectivity are often reported. Our group has previously demonstrated that thick films (~70-100 µm) undergo a rapid aging phase until an “aging knee” around

10-15 days at which quasi-steady-state transport properties may be observed even until

150 days [19]. Accordingly, 6FDA-DAT1 and 6FDA-DAT2 were aged for 150 days and their resulting pure-gas permeabilities and ideal selectivities are reported in Table 4.3.

Relative to the data for the freshly cast and dried films, roughly 15-20% reductions in gas 161 permeability coefficients were generally observed with negligible changes in selectivities.

It is suggested that the relatively flexible imide linkages (i.e., no ortho substituents in the diamine to sterically interfere with the nearby carbonyl groups and restrict rotation) to the

DAT1 and DAT2 diamines permit the chains to assume a well-packed arrangement soon after casting and drying. Thereafter, the main contributor to the free volume and microporosity characteristics observed in Figs. 4.4, 4.5 and 4.6 are the relatively “frozen” internal free volume pockets induced by the shape-persistent geometry of the triptycene moiety. Accordingly, minimal densification occurs with time.

Although an auspicious resilience against physical aging is evidenced in the time- dependence of the pure-gas transport data, sorption of condensable gases including CO2 can induce chain mobility, plasticize the matrix, alter the pore structure and significantly affect the separation properties observed in ideal pure-gas experiments. The pure- and mixed-gas permeabilities (Fig. 4.7 and 4.8) and selectivities (Fig. 4.9) were measured with a 1:1 CO2/CH4 mixture at increasing CO2 partial pressures (2, 5, 10 and 15 bar) for the aged polymers. The pure-gas CO2 permeabilities in 6FDA-DAT1 and 6FDA-DAT2 initially decreased with increasing feed pressure (Fig. 4.7). This can be rationalized by a decrease in CO2 solubility coefficients with pressure as is typically observed in glassy polymers. Aged 6FDA-DAT2 showed ~60% greater CO2 permeabilities attributed to its larger free volume. However, both polymers exhibit an up-turn in the CO2 permeability isotherms around ~10 bar attributed to plasticization of the matrices by the highly sorbing

CO2 gas. Moreover, competitive sorption occurs between CO2 and CH4 for available sorption sites and may reduce the overall sorbed concentration of CO2 in the polymer. 162

Across the pressure range, the mixed-gas CO2 permeabilities were thus lower than pure- gas CO2 permeabilities and did not exhibit an upturn at 10 bar.

Fig. 4.7. Pressure-dependence of pure- and mixed-gas CO2 permeabilities for 6FDA–

DAT1 and 6FDA–DAT2 (50:50 CO2:CH4 mixture; 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas.

Although the mixed-gas CO2 permeability isotherms do not upturn with pressure, Fig.

4.8 indicates a similar increase in mixed-gas CH4 permeabilities relative to the pure-gas

CH4 permeabilities from ~5 bar for both polymers. This is a direct indication of matrix dilation by sorbed CO2 molecules, which facilities diffusion of the larger CH4 gas and thereby increases its permeability. This is expected even in relatively rigid backbone polymers lacking sufficient physical or covalent interactions between chains [27-29]. 163

Fig. 4.8. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for 6FDA–

DAT1 and 6FDA–DAT2 (50:50 CO2:CH4 mixture, 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas.

Given the decrease in CO2 permeabilities and increase in CH4 permeabilities with pressure in the mixed-gas relative to the pure-gas experiments, the mixed-gas permselectivites (i.e., P(CO2)/P(CH4)) are lower than the pure-gas permselectivties (Fig.

4.9). In the context of typical natural gas sweetening conditions, where the partial pressure of CO2 is often near 5-10% in a ~60-70 bar feed, 6FDA-DAT1 exhibits a high selectivity of ~35 near the target of 40 set for membranes that would have the potential to replace current amine-absorption technology in the industry [30]. That is, high selectivity mitigates the extent of CH4 loss in the permeate, reducing the need for recompression of the product CH4 in a second stage and improving the economics of the separation [31].

While the higher permeabilities of the 6FDA-DAT2 polymer reduce membrane area requirements for a given feed flow, the use of the bulkier DAT2 diamine also reduces the 164 apparent selectivity to ~25. Overall, selectivities were similar to thick-film cellulose acetate membranes [32], where cellulose acetate is the commercial standard, but joined with 10-fold greater CO2 permeability.

Fig. 4.9. Pressure-dependence of mixed-gas CO2/CH4 selectivities for 6FDA–DAT1 and

6FDA–DAT2 (50:50 CO2:CH4 mixture, 35 °C). Lines are drawn to guide the eye: open symbols, pure-gas; closed symbols, mixed-gas.

4.5.Conclusions

In summary, two aromatic polyimides were synthesized by polycondensation reaction of 6FDA with 2,6-diaminotriptycene (6FDA–DAT1) and its bulkier, extended analog

(6FDA-DAT2). The effects of the additional benzene ring in the DAT2 diamine on the polymer microstructure and its gas transport properties were studied using low-pressure

N2 and CO2 physisorption and pure- and mixed-gas permeation experiments, respectively. An increase in free volume (~40% increase in BET surface area to 450 165 m2/g) was observed in 6FDA-DAT2. Also, NLDFT-based pore-size distributions indicated a bigger presence of large pores created by the bulkier DAT2 diamine.

Accordingly, the gas permeability coefficients were about 75% higher than in 6FDA-

DAT1 owing primarily to increases in gas diffusion coefficients. However, small gas- sieving pores were also maintained in 6FDA-DAT2 such that the ideal permselectivities were similar in both polymers. Moreover, the iptycene-based polymers demonstrated minimal physical aging over 150 days as evidenced by slight 15-20% drops in permeability coefficients with negligible changes in selectivities. In 1:1 CO2:CH4 mixed- gas experiments, both polymers exhibited upturns in mixed-gas CH4 permeabilities with pressure, suggesting CO2-induced plasticization. However, CO2:CH4 selectivities of 25-

35 were maintained near typical wellhead CO2 partial pressures and joined with high CO2 permeabilities exceeding 90 Barrer, deeming these polymers potential candidate materials for membrane-based gas separation when compared to the existing commercial polymers such as cellulose acetate.

166

4.6. Reference

[1] M. Carta, M. Croad, R. Malpass‐Evans, J.C. Jansen, P. Bernardo, G. Clarizia, K.

Friess, M. Lanč, N.B. McKeown, Triptycene induced enhancement of membrane gas selectivity for microporous Tröger's base polymers, Advanced Materials, 26 (2014) 3526-

3531.

[2] M. Carta, R. Malpass-Evans, M. Croad, Y. Rogan, J.C. Jansen, P. Bernardo, F.

Bazzarelli, N.B. McKeown, An efficient polymer molecular sieve for membrane gas separations, Science, 339 (2013) 303-307.

[3] B.S. Ghanem, N.B. McKeown, P.M. Budd, J.D. Selbie, D. Fritsch, High-performance membranes from polyimides with intrinsic microporosity, Advanced Materials, 20 (2008)

2766-2771.

[4] P.M. Budd, N.B. McKeown, Highly permeable polymers for gas separation membranes, Polymer Chemistry, 1 (2010) 63-68.

[5] P.M. Budd, K.J. Msayib, C.E. Tattershall, B.S. Ghanem, K.J. Reynolds, N.B.

McKeown, D. Fritsch, Gas separation membranes from polymers of intrinsic microporosity, Journal of Membrane Science, 251 (2005) 263-269.

[6] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra‐microporous triptycene‐based polyimide membranes for high‐performance gas separation, Advanced Materials, (2014).

[7] B.S. Ghanem, R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, Energy‐efficient hydrogen separation by AB‐type ladder‐polymer molecular sieves, Advanced Materials, 26 (2014)

6696-6700. 167

[8] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[9] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820.

[10] X.H. Ma, R. Swaidan, Y. Belmabkhout, Y.H. Zhu, E. Litwiller, M. Jouiad, I. Pinnau,

Y. Han, Synthesis and gas transport properties of hydroxyl-functionalized polyimides with intrinsic microporosity, Macromolecules, 45 (2012) 3841-3849.

[11] S.A. Sydlik, Z.H. Chen, T.M. Swager, Triptycene polyimides: soluble polymers with high thermal stability and low refractive indices, Macromolecules, 44 (2011) 976-980.

[12] Y. Rogan, R. Malpass-Evans, M. Carta, M. Lee, J.C. Jansen, P. Bernardo, G.

Clarizia, E. Tocci, K. Friess, M. Lanč, A highly permeable polyimide with enhanced selectivity for membrane gas separations, Journal of Materials Chemistry A, 2 (2014)

4874-4877.

[13] X. Ma, B. Ghanem, O. Salines, E. Litwiller, I. Pinnau, Synthesis and effect of physical aging on gas transport properties of a microporous polyimide derived from a novel spirobifluorene-based dianhydride, ACS Macro Letters, 4 (2015) 231-235. 168

[14] X. Ma, O. Salinas, E. Litwiller, I. Pinnau, Novel spirobifluorene-and dibromospirobifluorene-based polyimides of intrinsic microporosity for gas separation applications, Macromolecules, 46 (2013) 9618-9624.

[15] J. Weber, Q. Su, M. Antonietti, A. Thomas, Exploring polymers of intrinsic microporosity–microporous, soluble polyamide and polyimide, Macromolecular Rapid

Communications, 28 (2007) 1871-1876.

[16] Z. Wang, D. Wang, F. Zhang, J. Jin, Tröger’s base-based microporous polyimide membranes for high-performance gas separation, ACS Macro Letters, 3 (2014) 597-601.

[17] Z. Wang, D. Wang, J. Jin, Microporous polyimides with rationally designed chain structure achieving high performance for gas separation, Macromolecules, 47 (2014)

7477-7483.

[18] Y. Zhuang, J.G. Seong, Y.S. Do, H.J. Jo, Z. Cui, J. Lee, Y.M. Lee, M.D. Guiver,

Intrinsically microporous soluble polyimides incorporating Tröger’s base for membrane gas separation, Macromolecules, 47 (2014) 3254-3262.

[19] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational design of intrinsically ultramicroporous polyimides containing bridgehead-substituted triptycene for highly selective and permeable gas separation membranes, Macromolecules, 47

(2014) 5104-5114.

[20] Y.J. Cho, H.B. Park, High performance polyimide with high internal free volume elements, Macromolecular Rapid Communications, 32 (2011) 579-586. 169

[21] J.R. Wiegand, Z.P. Smith, Q. Liu, C.T. Patterson, B.D. Freeman, R. Guo, Synthesis and characterization of triptycene-based polyimides with tunable high fractional free volume for gas separation membranes, Journal of Materials Chemistry A, 2 (2014)

13309-13320.

[22] N.T. Tsui, A.J. Paraskos, L. Torun, T.M. Swager, E.L. Thomas, Minimization of internal molecular free volume: a mechanism for the simultaneous enhancement of polymer stiffness, strength, and ductility, Macromolecules, 39 (2006) 3350-3358.

[23] T.M. Swager, Iptycenes in the design of high performance polymers, Accounts of

Chemical Research, 41 (2008) 1181-1189.

[24] K.C. Obrien, W.J. Koros, T.A. Barbari, E.S. Sanders, A new technique for the measurement of multicomponent gas-transport through polymeric films, Journal of

Membrane Science, 29 (1986) 229-238.

[25] K.D. Dorkenoo, P.H. Pfromm, Accelerated physical aging of thin poly [1-

(trimethylsilyl)-1-propyne] films, Macromolecules, 33 (2000) 3747-3751.

[26] P.H. Pfromm, The impact of physical aging of amorphous glassy polymers on gas separation membranes, Materials Science of Membranes for Gas and Vapour Separation,

(2006) 293-306.

[27] R. Swaidan, B. Ghanem, E. Litwiller, I. Pinnau, Effects of hydroxyl- functionalization and sub-Tg thermal annealing on high pressure pure-and mixed-gas

CO2/CH4 separation by polyimide membranes based on 6FDA and triptycene-containing dianhydrides, Journal of Membrane Science, 475 (2015) 571-581. 170

[28] J.T. Vaughn, W.J. Koros, J. Johnson, O. Karvan, Effect of thermal annealing on a novel polyamide–imide polymer membrane for aggressive acid gas separations, Journal of Membrane Science, 401 (2012) 163-174.

[29] J.D. Wind, S.M. Sirard, D.R. Paul, P.F. Green, K.P. Johnston, W.J. Koros, Carbon dioxide-induced plasticization of polyimide membranes: pseudo-equilibrium relationships of diffusion, sorption, and swelling, Macromolecules, 36 (2003) 6433-6441.

[30] R.W. Baker, Future directions of membrane gas separation technology, Industrial &

Engineering Chemistry Research, 41 (2002) 1393-1411.

[31] B. Bhide, S. Stern, Membrane processes for the removal of acid gases from natural gas. II. effects of operating conditions, economic parameters, and membrane properties,

Journal of Membrane Science, 81 (1993) 239-252.

[32] A. Houde, B. Krishnakumar, S. Charati, S. Stern, Permeability of dense

(homogeneous) cellulose acetate membranes to methane, carbon dioxide, and their mixtures at elevated pressures, Journal of Applied Polymer Science, 62 (1996) 2181-

2192.

171

Chapter 5. Gas Transport Properties and Characterization of Polyimides of

Intrinsic Microporosity Based on Novel Triptycene Dianhydrides and 3,3ʹ-

Dimethylnaphthidine Diamine.1

5.1. Abstract

Two intrinsically microporous polyimides were obtained by polycondensation reaction of novel triptycene-based dianhydrides containing dimethyl- or diisopropyl bridgehead sidegroups with a commercially available highly sterically hindered 3,3ʹ- dimethylnaphthidine (DMN) diamine monomer. The dimethyl bridgehead groups in the triptycene building block provided the DMN-based polyimide (TDA1-DMN) with larger surface area (760 m2/g) than the diisopropyl-based polyimide (TDAi3-DMN) (680 m2/g), greater fraction of ultramicroporosity, as observed from N2 and CO2 NLDFT adsorption analysis, and higher gas permeability and selectivity. Wide-angle X-ray diffraction

(WAXD) measurements demonstrated that TDA1-DMN and TDAi3-DMN exhibited a bimodal pore size distribution, where TDA1-DMN showed smaller d-spacing values and broader intensity peaks. Both TDA-DMN-based polyimides showed very high gas permeabilities with moderate selectivities. For example, fresh TDA1-DMN exhibited an

O2 permeability of 783 Barrer coupled with an O2/N2 selectivity of 4.3 and CO2 permeability of 3700 Barrer with a CO2/CH4 selectivity of 17.1, values that surpassed the

2008 Robeson upper bounds.

1Portions of this chapter were adopted from Ghanem, B.; Alghunaimi, F.; Ma, X.; Alaslai, N.; Pinnau, I. Polymer 2016 (101) 225-232. Alghunaimi performed gas permeation, sorption, characterization and wrote the manuscript. Ghanem and Ma synthesized the polymers. Alaslai directed construction of permeation apparatus and edited the manuscript. Pinnau supervised the work and edited the manuscript. 172

5.2. Introduction

Triptycene-derived microporous network polymers and amorphous, soluble polymers of intrinsic microporosity (PIMs) are emerging classes of high-performance materials with potential applications in optoelectronics [1, 2], sensors [3], gas adsorption and storage [4-14] and membrane separations [15-23]. Specifically, triptycene-derived ladder

PIMs and polyimides of intrinsic microporosity (PIM-PIs) have shown exceptional performance for membrane-based gas separation applications, defining the latest 2015 permeability/selectivity upper bound relationships for O2/N2, H2/N2 and H2/CH4 [24].

Triptycene is a unique building block for high-performance microporous polymers as its

D3h symmetry with rigid Y-shaped structure containing arene units fused to a [2,2,2] octatriene bridgehead system provides intrinsic internal free volume (IFV), as shown in

Fig. 5.1 [25-28]. Further enhancement in free volume results from poor interchain packing due the bulky shape-persistent 3-D paddlewheel structure. The IFV and interchain free volume can be tailored by selection of the bridgehead substituents, which strongly influence the gas sorption, diffusion and gas permeation properties of triptycene polymers. For example, Ghanem et al. demonstrated in a series of triptycene network ladder polymers that alkyl substitutions on the 9,10-bridgehead can be used to systematically tune the BET surface area for enhanced gas sorption uptake [6].

Interestingly, short dimethyl- and bulky diisopropyl bridgehead groups gave the greatest microporosity with high BET surface areas of 1760 and 1601 m2/g, respectively. Longer alkyl side-chains gave rise to lower microporosity presumably due to their flexibility and partial blocking of the rigid triptycene framework [6]. 173

Fig. 5.1. Structure of triptycene with indication of internal free volume (IFV); R = alkyl bridgehead substitutions [25-28].

Similar results were reported for a series of solution-processable, amorphous 9,10- dialkyl-substituted triptycene dianhydride-based polyimides of intrinsic microporosity

(PIM-PIs) for membrane-based gas separation applications (KAUST-PI-series) [21]. The study demonstrated that tetramethylenediamine (TMPD)-based triptycene PIM-PI bearing branched diisopropyl (R=i-C3) bridgehead substituents (KAUST-PI-1) showed the best gas separation performance with properties significantly outperforming the 2008 permeability/selectivity upper bounds for several gas pairs due to its strong size-sieving properties [21]. Unfortunately, dimethyl-substituted triptycene TMPD-derived KAUST-

PI was insoluble and, therefore, a comparative study with the diisopropyl-based KAUST-

PI-1 could not be performed.

In this chapter, the synthesis of novel dimethyl- and diisopropyl-triptycene-based dianhydride monomers (TDAs) is reported, where reactive anhydride groups are fused directly to the triptycene unit, and commercially available DMN (Fig. 5.2). The bulky 174 diamine was chosen because previous studies showed that PIM-PIs made from DMN could be produced in high molecular weight, had good solubility in a variety of solvents and exhibited very high gas permeability [29-31]. The detailed synthetic route of dimethyl- (TDA1-DMN) and diisopropyl- (TDAi3-DMN)-based PIM-PIs is shown in

Scheme 5.1 [32]. It is important to note that this synthetic procedure differs from previously reported methods for producing triptycene dianhydrides [29, 30, 33].

Fig. 5.2. Structure of polyimides TDA1-DMN and TDAi3-DMN.

Full details of the characterization of the polymers (FTIR, BET, GPC, TGA and

WAXD) as well as the pure-gas transport properties for He, H2, N2, O2, CH4, and CO2 at

2 bar and 35 °C are reported. 175

Scheme 5.1. Synthetic route of triptycene-dianhydride-based polyimides TDA1DMN and

TDAi3-DMN [32].

5.3. Experimental

5.3.1. Characterization Methods

Gel permeation chromatography (GPC, Viscotek) analyses were carried out using chloroform as an eluent. Thermogravimetric analysis (TGA) was carried out under a nitrogen atmosphere (TA Q-5000). Fourier transform infrared (FT-IR) measurements were performed using a Varian 670-IR FT-IR spectrometer. X-ray diffraction (XRD) scattering was conducted at room temperature using a Bruker D8 Advance diffractometer operating in a 2θ range of 7º to 40º with a rate of 0.5 second per step and the d-spacings were calculated with Bragg’s law. The BET surface area of the polymers was determined by N2 sorption (-196 °C) using a Micromeritics ASAP-2020. Powder polymer samples were degassed under high vacuum at 120 °C for 16 h prior to analysis. Analysis of the 176

pore size distributions was performed by N2 and CO2 sorption using the NLDFT (non- local density functional theory) model for carbon slit pore geometry provided by ASAP

2020 version 4.02 software.

5.3.2. Membrane Fabrication

A 3-4% (w/v) polymer solution in chloroform was filtered and cast on a leveled glass

Petri dish followed by slow evaporation of the solvent over 2-3 days. To remove any traces of residual casting solvent, the resulting tough and flexible films were first dried at

120 °C for 12 h in a vacuum oven, then soaked in methanol for 24 h and finally dried at

120 °C for 24 h under high vacuum. Complete solvent removal was confirmed by TGA prior to the gas permeation tests. The films had thicknesses in the range of 85-100 µm as measured with a high-precision micrometer.

5.3.3. Gas Permeation Experiments

The pure-gas permeabilities of TDA-DMN polyimides were determined by using the variable pressure/constant volume method. Each polymer film sample was degassed in the permeation cell under vacuum for at least 24 hours. The pure-gas permeability of He,

H2, N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar. The gas permeability (P) was calculated by equation 2.14 where P is in Barrers (1 Barrer = 10-10 cm3(STP·cm/cm2·s·cmHg). The apparent diffusion coefficient D (cm2/s) was calculated by D = l2 /6 θ, where θ is the time lag of the permeability measurement and l is the membrane thickness. The ideal selectivity for a gas pair is given by equation 2.6. 177

5.4. Results and Analysis

5.4.1. Synthesis and Characterization of the TDA-DMN Polyimides

The two novel TDA1-DMN and TDAi3-DMN polyimides were synthesized via the one-step high-temperature cycloimidization reaction using an equimolar ratio of the commercially available DMN diamine and the newly designed triptycene dianhydrides in m-cresol [32]. The molecular structures of the polyimides were characterized by FTIR spectroscopy. The polyimides showed the characteristic absorption bands of the imide groups at ~ 1778, 1710 and 1387 cm-1 (Fig. 5.3). No signals associated with polyamic acid around 3200-3400 cm-1 were observed indicating complete imidization of the polymers.

Fig. 5.3. FTIR spectra of TDA1-DMN and TDAi3-DMN polyimides. 178

The chemical structures of the polyimides were further characterized by GPC, TGA, and

N2 and CO2 adsorption measurements. All polymers were readily soluble in organic solvents including chloroform from which transparent, tough and flexible isotropic membranes suitable for gas permeation tests were formed. The polymers exhibited high average molecular weights and narrow polydispersities, as evaluated by GPC using chloroform as eluent and polystyrene as standard (Table 5.1).

Table 5.1. Molecular weights, thermal stability, and BET surface areas of TDA-DMN- based dimethyl- (R= C1) and diisopropyl- (R= i-C3) substituted PIM-PIs.

Mn Td N2 BET S.A. Polymer R PDI (103 g mol-1) (°C) (m2 g-1)

TDA1-DMN C1 73 1.56 500 760

TDAi3-DMN i-C3 131 1.50 440 680

As expected, the TDA-DMN polyimides demonstrated high thermal stability with decomposition temperatures exceeding 400 °C as shown by the thermogravimetric analysis (TGA) in Fig. 5.4 and Table 5.1. 179

Fig. 5.4. Thermogravimetric analysis (TGA) of TDA1-DMN and TDAi3-DMN.

The effect of the 9,10-dimethyl- and diisopropyl substitutions on the three- dimensional triptycene moiety of the polyimides was assessed from N2 (-196 °C) and

CO2 (0 °C) physisorption isotherms (Fig. 5.5 and 5.6). The BET surface areas of TDA1-

DMN and TDAi3-DMN were 760 and 680 m2/g, respectively, indicating very high microporosity in the polyimides. The higher BET surface area obtained for dimethyl- bridgehead-containing TDA1-DMN in comparison to that of diisopropyl-bridgehead- substituted TDAi3-DMN was in qualitative agreement with earlier studies of dialkyl- substitutions on the microporosity of triptycene ladder network polymers [6]. 180

Fig. 5.5. Physisorption isotherms using N2 at -196 °C for TDA1-DMN (R=C1) and

TDAi3-DMN (R=i-C3).

Fig. 5.6. Physisorption isotherms using CO2 at 0 °C for TDA1-DMN (R=C1) and TDAi3-

DMN (R=i-C3). 181

The pore size distributions (PSDs) calculated using the NLDFT (Non-Local Density

Functional Theory) model (Fig. 5.7) clearly show the qualitative differences in the microporous texture of the TDA-DMN polyimides. TDA1-DMN had slightly greater fraction of ultramicropores (< 7 Å) compared to the diisopropyl-bridgehead-derived polyimide (TDAi3-DMN). Furthermore, TDA1-DMN polyimide exhibited a larger fraction of pores between ~7 to 12 Å. The qualitative assessment of the microporous structure of the PIM-PIs by BET and NLDFT analysis has a direct correlation with their gas permeation properties, as discussed below.

Fig. 5.7. NLDFT-based estimated pore size distribution obtained from N2 and CO2 isotherms for TDA1-DMN and TDAi3-DMN assuming carbon slit-pore geometry.

Wide-angle X-ray diffraction (WAXD) measurements were conducted on TDA1-

DMN and TDAi3-DMN polyimide films to investigate the effect of bridgehead substituents on the average polymer chain spacing. X-ray diffraction confirmed that both 182 materials were amorphous and had bimodal pore size distributions. TDA1-DMN exhibited two peaks centered at average d-spacings of 6.5 and 4.4 Å, whereas TDAi3-

DMN had two peaks of larger d-spacings of 7.2 and 4.7 Å. The WAXD results indicate clearly that the dimethyl-bridgehead-substituted TDA1-DMN contained a larger fraction of small ultramicroporous (< 5 Å) than its diisopropyl-based derivative TDAi3-DMN

(Fig. 5.8).

Fig. 5.8. WAXD patterns of TDA1-DMN and TDAi3-DMN films.

5.4.2. Gas Transport Properties

The pure-gas permeation properties of the two TDA-DMN based polyimide samples determined at 2 bar and 35 °C are reported in Table 5.2. To obtain meaningful structure/gas permeation property relationships between the dimethyl-and diisopropyl- 183

triptycene bridgehead substitutions, the same polymer film preparation and permeation

test protocols were applied for each polymer.

Table 5.2. Gas permeabilities and permselectivities for TDA-DMN-based polyimides at

2 bar and 35 °C.

Pure-gas permeability (Barrer*) Permselectivity (휶)

Polymer He H2 N2 O2 CH4 CO2 O2/N2 H2/N2 CO2/CH4

TDA1-DMNa 1182 3050 182 783 216 3700 4.3 16.7 17.1

Aged TDA1-DMNa (903) (2430) (134) (609) (158) (3000) (4.5) (18) (19)

TDAi3-DMNa 913 2233 160 594 211 3154 3.7 14.0 14.9

Aged TDAi3-DMNa (900) (2114) (130) (505) (170) (2670) (3.9) (16) (16)

PIM-PI-10b 300 670 84 270 168 2154 3.2 8.0 12.8

* 1 Barrer = 1x10-10 cm3(STP) cm/cm2 s cmHg.

a This study.

b Spirobisindane-based DMN polyimide reported by Rogan et al.[30] 184

The two new TDA-DMN polyimides exhibited very high gas permeabilities combined with moderate selectivities, similar to previously published data of structurally related DMN-based PIM-PIs with a spirobisindane- (PIM-PI-10) building block, as shown in Table 5.2 [30]. The new polyimides with 9,10-diakyltriptycene contortion sites exhibited both higher gas permeability and selectivity than PIM-PI-10 with spirobisindane contortion center as previously reported [4, 21]. Interestingly, the TDA polyimide with methyl bridgehead substituents showed better gas permeabilities and selectivities than that containing isopropyl side-chains, similar to the qualitative trends observed in BET surface area and WAXD experiments. The same general trend has previously been documented in studies on the effect of side-chains in high-free-volume di-substituted polyacetylenes containing linear- and branched-alkyl side-chains [6, 34,

35].

The permeability/selectivity trade-off behaviour of the two TDA-DMN polyimides for O2/N2, H2/N2 and CO2/CH4 separation is shown in Fig. 5.9a, 5.9b and 5.10. Data for the spirobisindane-DMN-based polyimide PIM-PI-10 are included for comparison. The

TDA-DMN polyimides exhibited gas separation performance close to or above the 2008

Robeson upper bounds [36]. Specifically, the TDA1-DMN polyimide demonstrated better performance for O2/N2 and H2/N2 performance surpassing the 2008 permeability/selectivity trade-off curves. However, their performance fell short for the more recently published 2015 upper bounds for O2/N2 and H2/N2 [24]. Furthermore,

TDA1-DMN polyimide exhibited better performance than PIM-PI-10 for CO2/CH4 exceeding the 2008 permeability/selectivity trade-off curves as shown in Fig. 5.10. 185

Fig. 5.9. Gas separation performance of TDA-DMN polyimides and PIM-PI-10 for (a)

O2/N2 and (b) H2/N2. Polymers are labeled as indicated in Fig. 3.6 and 5.2. The solid lines represent 2008 and 2015 permeability/selectivity tradeoffs [24, 36].

Fig. 5.10. Gas separation performance of TDA-DMN polyimides and PIM-PI-10 for

CO2/CH4. Polymers are labeled as indicated in Fig. 3.6 and 5.2. The solid lines represent

2008 permeability/selectivity tradeoffs [36]. 186

5.5. Conclusions

Two intrinsically microporous polyimides were obtained by reaction of 9,10-dialkyl

(C1 and i-C3) triptycene-based dianhydride monomers and DMN via a one-step high- temperature solution polycondensation reaction. All polyimides exhibited good solubility in common organic solvents, excellent thermal stability, and high free volume as confirmed by high BET surface areas. These two polyimides were synthesized to compare and evaluate the effect of the two side groups (methyl and diisopropyl) to the triptycene building block on gas transport properties as shown in Chapter-3 (Fig. 3.12).

Short methyl groups in the bridgehead of the triptycene building block provided the

TDA1-DMN polyimide with the highest BET surface area and greatest ultramicroporosity. This study further demonstrates that selection of the bridgehead substituents can be used to fine-tune the microporous texture, pore size distribution and, consequently, gas permeation properties of triptycene-based polyimides. The combination of high gas permeability and moderate selectivity places TDA1-DMN above the 2008 upper bounds for CO2/CH4, O2/N2 and H2/N2.

187

5.6. References

[1] S.A. Sydlik, Z. Chen, T.M. Swager, Triptycene polyimides: soluble polymers with high thermal stability and low refractive indices, Macromolecules, 44 (2011) 976-980.

[2] S.-H. Hsiao, H.-M. Wang, J.-S. Chou, W. Guo, T.-M. Lee, C.-M. Leu, C.-W. Su,

Triptycene poly (ether-imide) s with high solubility and optical transparency, Journal of

Polymer Research, 19 (2012) 1-12.

[3] C.-F. Chen, Y.-X. Ma, Iptycenes Chemistry: From Synthesis to Applications,

Springer Science & Business Media, 2012.

[4] B.S. Ghanem, K.J. Msayib, N.B. McKeown, K.D. Harris, Z. Pan, P.M. Budd, A.

Butler, J. Selbie, D. Book, A. Walton, A triptycene-based polymer of intrinsic microposity that displays enhanced surface area and hydrogen adsorption, Chemical

Communications, (2007) 67-69.

[5] C. Zhang, L.-H. Peng, B. Li, Y. Liu, P.-C. Zhu, Z. Wang, D.-H. Zhan, B. Tan, X.-L.

Yang, H.-B. Xu, Organic microporous polymer from a hexaphenylbenzene based triptycene monomer: synthesis and its gas storage properties, Polymer Chemistry, 4

(2013) 3663-3666.

[6] B.S. Ghanem, M. Hashem, K.D.M. Harris, K.J. Msayib, M.C. Xu, P.M. Budd, N.

Chaukura, D. Book, S. Tedds, A. Walton, N.B. McKeown, Triptycene-based polymers of intrinsic microporosity: organic materials that can be tailored for gas adsorption,

Macromolecules, 43 (2010) 5287-5294. 188

[7] C. Zhang, Y. Liu, B. Li, B. Tan, C.-F. Chen, H.-B. Xu, X.-L. Yang, Triptycene-based microporous polymers: synthesis and their gas storage properties, ACS Macro Letters, 1

(2011) 190-193.

[8] Y.-C. Zhao, Q.-Y. Cheng, D. Zhou, T. Wang, B.-H. Han, Preparation and characterization of triptycene-based microporous poly (benzimidazole) networks, Journal of Materials Chemistry, 22 (2012) 11509-11514.

[9] M.G. Rabbani, T.E. Reich, R.M. Kassab, K.T. Jackson, H.M. El-Kaderi, High CO2 uptake and selectivity by triptycene-derived benzimidazole-linked polymers, Chemical

Communications, 48 (2012) 1141-1143.

[10] C. Zhang, T.-L. Zhai, J.-J. Wang, Z. Wang, J.-M. Liu, B. Tan, X.-L. Yang, H.-B.

Xu, Triptycene-based microporous polyimides: synthesis and their high selectivity for

CO2 capture, Polymer, 55 (2014) 3642-3647.

[11] C. Zhang, P.-C. Zhu, L. Tan, J.-M. Liu, B. Tan, X.-L. Yang, H.-B. Xu, Triptycene- based hyper-cross-linked polymer sponge for gas storage and water treatment,

Macromolecules, 48 (2015) 8509-8514.

[12] C. Zhang, Z. Wang, J.-J. Wang, L. Tan, J.-M. Liu, B. Tan, X.-L. Yang, H.-B. Xu,

Synthesis and properties of triptycene-based microporous polymers, Polymer, 54 (2013)

6942-6946.

[13] Y. He, X. Zhu, Y. Li, C. Peng, J. Hu, H. Liu, Efficient CO2 capture by triptycene- based microporous organic polymer with functionalized modification, Microporous and

Mesoporous Materials, 214 (2015) 181-187. 189

[14] X.L. Lu, T.Y. Zhou, D. Wu, Q. Wen, X. Zhao, Q. Li, Q. Xiang, J.Q. Xu, Z.T. Li, A triptycene‐based porous organic polymer that exhibited high hydrogen and carbon dioxide storage capacities and excellent CO2/N2 selectivity, Chinese Journal of

Chemistry, 33 (2015) 539-544.

[15] Y.J. Cho, H.B. Park, High performance polyimide with high internal free volume elements, Macromolecular Rapid Communications, 32 (2011) 579-586.

[16] R. Swaidan, B. Ghanem, M. Al-Saeedi, E. Litwiller, I. Pinnau, Role of intrachain rigidity in the plasticization of intrinsically microporous triptycene-based polyimide membranes in mixed-gas CO2/CH4 separations, Macromolecules, 47 (2014) 7453-7462.

[17] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[18] J.R. Wiegand, Z.P. Smith, Q. Liu, C.T. Patterson, B.D. Freeman, R. Guo, Synthesis and characterization of triptycene-based polyimides with tunable high fractional free volume for gas separation membranes, Journal of Materials Chemistry A, 2 (2014)

13309-13320.

[19] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra‐microporous triptycene‐ based polyimide membranes for high‐performance gas separation, Advanced Materials,

26 (2014) 3688-3692. 190

[20] B.S. Ghanem, R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, Energy‐efficient hydrogen separation by AB‐type ladder‐polymer molecular sieves, Advanced Materials, 26 (2014)

6696-6700.

[21] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational design of intrinsically ultramicroporous polyimides containing bridgehead-substituted triptycene for highly selective and permeable gas separation membranes, Macromolecules, 47

(2014) 5104-5114.

[22] M. Carta, M. Croad, R. Malpass‐Evans, J.C. Jansen, P. Bernardo, G. Clarizia, K.

Friess, M. Lanč, N.B. McKeown, Triptycene induced enhancement of membrane gas selectivity for microporous Tröger's base polymers, Advanced Materials, 26 (2014) 3526-

3531.

[23] I. Rose, M. Carta, R. Malpass-Evans, M.-C. Ferrari, P. Bernardo, G. Clarizia, J.C.

Jansen, N.B. McKeown, Highly permeable benzotriptycene-based polymer of intrinsic microporosity, ACS Macro Letters, 4 (2015) 912-915.

[24] R. Swaidan, B. Ghanem, I. Pinnau, Fine-tuned intrinsically ultramicroporous polymers redefine the permeability/selectivity upper bounds of membrane-based air and hydrogen separations, ACS Macro Letters, 4 (2015) 947-951.

[25] T.M. Long, T.M. Swager, Minimization of free volume: alignment of triptycenes in liquid crystals and stretched polymers, Advanced Materials, 13 (2001) 601-604.

[26] N.G. White, M.J. MacLachlan, Soluble tetraaminotriptycene precursors, The Journal of Organic Chemistry, 80 (2015) 8390-8397. 191

[27] N.T. Tsui, A.J. Paraskos, L. Torun, T.M. Swager, E.L. Thomas, Minimization of internal molecular free volume: a mechanism for the simultaneous enhancement of polymer stiffness, strength, and ductility, Macromolecules, 39 (2006) 3350-3358.

[28] J.H. Chong, M.J. MacLachlan, Iptycenes in supramolecular and materials chemistry,

Chemical Society Reviews, 38 (2009) 3301-3315.

[29] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[30] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820.

[31] X. Ma, B. Ghanem, O. Salines, E. Litwiller, I. Pinnau, Synthesis and effect of physical aging on gas transport properties of a microporous polyimide derived from a novel spirobifluorene-based dianhydride, ACS Macro Letters, 4 (2015) 231-235.

[32] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232. 192

[33] Y.-H. Xiao, Y. Shao, X.-X. Ye, H. Cui, D.-L. Wang, X.-H. Zhou, S.-L. Sun, L.

Cheng, Microporous aromatic polyimides derived from triptycene-based dianhydride,

Chinese Chemical Letters, (2016).

[34] I. Pinnau, A. Morisato, Z. He, Influence of side-chain length on the gas permeation properties of poly (2-alkylacetylenes), Macromolecules, 37 (2004) 2823-2828.

[35] I. Pinnau, Z. He, A. Morisato, Synthesis and gas permeation properties of poly

(dialkylacetylenes) containing isopropyl-terminated side-chains, Journal of Membrane

Science, 241 (2004) 363-369.

[36] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400.

193

Chapter 6. Gas Transport Properties and Characterization of Polyimides of

Intrinsic Microporosity Based on Non-Substituted- and Diethyl-Triptycene

Dianhydrides and 3,3ʹ-Dimethylnaphthidine Diamine.

6.1. Abstract

Two new intrinsically microporous polyimides (TDA0-DMN and TDA2-DMN) were obtained by one-pot polycondensation reaction of novel non-substituted- and 9,10- diethyl-triptycene-based dianhydrides with a commercially available highly sterically hindered 3,3ʹ-dimethylnaphthidine (DMN) diamine monomer. The longer diethyl- substitution in TDA2-DMN reduced the surface area to 580 m2/g and ultramicroporosity compared to TDA0-DMN, which had a surface area of 700 m2/g. The dimethyl bridgehead groups in the triptycene building block in TDA1-DMN polyimide (Chapter

5) provide the largest surface area (760 m2/g), greatest fraction of ultramicroporosity, as observed from N2 and CO2 NLDFT adsorption analysis, and highest gas permeability compared to the non-substituted, diethyl- and diisopropyl-triptycene-based polyimides.

Wide-angle X-ray diffraction (WAXD) measurements demonstrated that TDA1-DMN and TDA2-DMN exhibited two peaks with a bimodal pore size distribution, whereas

TDA0-DMN showed one broad intensity peak. All TDA-DMN-based polyimides showed very high gas permeabilities with moderate selectivities. The dimethyl bridgehead groups in the triptycene building block polyimide exhibited the best CO2/CH4 separation performance among the non-substituted, extended iptycene, diethyl- and diisopropyl- triptycene-based polyimides. The dimethyl-triptycene building block will be used further in this research to design membrane materials for CO2/CH4 separation. 194

6.2. Introduction

Polymeric membranes are an attractive option for gas separation because of the adaptability of their chemistry, feasibility of synthesis, processability and cost effectiveness compared to the other membrane types. The main requirements for advanced polymeric membranes to efficiently separate gases are to enhance membrane performance and improve the chemical and thermal stability [1]. Glassy aromatic polyimides are able to solve these issues because of their high gas pair selectivity, excellent thermal stability, high chemical resistance and good mechanical properties [2].

Recently, considerable efforts have been devoted to the development of microporous aromatic polyimides by rational molecular design to create new membrane materials with improved gas separation properties [3-6]. For example, spirobisindane dianhydride

(SBIDA) derived polyimides of intrinsic microporosity (PIM-PIs), first reported by

Ghanem et al., showed remarkably enhanced permeability compared to conventional low- free-volume polyimides [7]. The most permeable PIM-PIs were obtained by polycondensation of extended (Fig. 3.7) and non-extended (Fig. 3.8) SBIDA with DMN, respectively (PIM-PI-8; PIM-PI-10) [4, 5]. However, the performance of these polyimides fell below the 2008 Robeson upper bounds because of their low to moderate selectivities. On the other hand, polyimides made from dianhydrides with non-extended enthanoanthracene- or spirobifluorene contortion sites with DMN (PIM-PI-12 and

SBFDA-DMN; Fig. 6.1 a and b) showed improved gas separation performance for several gas pairs, specifically after long-term physical aging [8, 9]. Recently, Pinnau’s group reported a series of 9,10-diisopropyltriptycene extended dianhydride-based 195 polyimides, which significantly outperformed the latest upper bounds for several gas pairs due to their strong size-sieving properties [3, 10].

Fig. 6.1. Structures for polyimides of intrinsic microporosity: a) non-extended EADA-

DMN (PIM-PI-12) and b) non-extended SBFDA-DMN.

Triptycene is the simplest iptycene and a unique building block for high-performance microporous polymers that provides intrinsic internal free volume (IFV), as shown in Fig.

3.16 [11-14]. In this chapter, two novel triptycene-based polyimides are reported using non-substituted- and 9,10-diethyl building blocks, as shown in Fig. 6.2.

Fig. 6.2. Chemical structures of a) non-substituted-triptycene and b) diethyl-triptycene- building blocks.

The gas transport properties for He, H2, N2, O2, CH4, and CO2 (2 bar, 35 °C) and full characterization details (BET, TGA and WAXD) for the polyimides (Fig. 6.3) are 196 reported. The bulky DMN diamine was chosen because previous studies showed that

PIM-PIs made from DMN could be produced in high molecular weight and had good solubility in a variety of solvents [4, 5, 9]. The gas permeation performance of TDA0-

DMN and TDA2-DMN was compared with TDA1-DMN that contained dimethyl bridgehead substitution (Chapter-5).

Fig. 6.3. Chemical structure of TDA0-DMN and TDA2-DMN polyimides.

6.3. Experimental

6.3.1. Characterization Methods

Thermogravimetric analysis (TGA) was carried out under a nitrogen atmosphere (TA

Q-5000). X-ray diffraction (XRD) scattering was conducted at room temperature using a 197

Bruker D8 Advance diffractometer operating in a 2θ range of 7º to 40º with a rate of 0.5 second per step and the d-spacings were calculated with Bragg’s law. The BET surface area of the polymers was determined by N2 sorption (-196 °C) using a Micromeritics

ASAP-2020. Powder polymer samples were degassed under high vacuum at 120 °C for

16 h prior to analysis. Analysis of the pore size distributions was performed by N2 sorption using the NLDFT (non-local density functional theory) model for carbon slit pore geometry provided by ASAP 2020 version 4.02 software.

6.3.2. Membrane Fabrication

A 3-4% (w/v) polymer solution in chloroform was filtered and cast on a levelled glass

Petri dish followed by slow evaporation of the solvent over 2-3 days. To remove any traces of residual casting solvent, the resulting tough and flexible films were first dried at

120 °C for 12 h in a vacuum oven, then soaked in methanol for 24 h and finally dried at

120 °C for 24 h under high vacuum. Complete solvent removal was confirmed by TGA prior to the gas permeation tests. The films had thicknesses in the range of 75-95 µm as measured with a high-precision micrometer.

6.3.3. Gas Permeation Experiments

The pure-gas permeabilities of TDA0-DMN and TDA2-DMN polyimides were determined by using the variable pressure/constant volume method. Each polymer film sample was degassed in the permeation cell under vacuum for at least 24 hours. The pure- gas permeability of He, H2, N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar. The gas permeability (P) was calculated by equation 2.14 where P is reported in Barrers (1

Barrer = 10-10 cm3(STP·cm/cm2·s·cmHg). The apparent diffusion coefficient D (cm2/s) 198 was calculated by D = l2 /6 θ, where θ is the time lag of the permeability measurement and l is the membrane thickness. The ideal selectivity for a gas pair is given by equation

2.6.

6.4. Results and Analysis

6.4.1. Synthesis and Characterization of the TDA-DMN Polyimides

The two novel TDA0-DMN and TDA2-DMN polyimides were synthesized via the one-step high-temperature cycloimidization reaction using an equimolar ratio of the commercially available DMN diamine and the newly designed triptycene dianhydrides in m-cresol [15]. As expected for polyimides, the TDA0-DMN and TDA2-DMN demonstrated high thermal stability with decomposition temperatures exceeding 400 °C as shown in Table 6.1 and Fig. 6.4.

Table 6.1. Molecular weights, thermal stability, and BET surface areas of TDA-DMN-

Based PIM-PIs.

a Td N2 BET S.A. Polymer Side group R (°C) (m2/g)

TDA0-DMN Non-substituted C0 510 700

TDA2-DMN Diethyl C2 470 580

a Onset of thermal decomposition. 199

Fig. 6.4. Thermogravimetric analysis (TGA) of TDA0-DMN and TDA2-DMN.

The effect of the non-substituted 9,10-dimethyl- (Chapter 5) and -diethyl- substitutions on the three-dimensional triptycene moiety of the polyimides was assessed from N2 (-196 °C) physisorption isotherms (Fig. 6.5). The BET surface areas of TDA0-

DMN, TDA1-DMN, and TDA2-DMN were 700, 760 and 580 m2/g, respectively, indicating very high microporosity in the polyimides. 200

Fig. 6.5. Physisorption isotherms using N2 at -196 °C for TDA0-DMN, TDA1-DMN

(R=C1) and TDA2-DMN (R=C2).

Fig. 6.6. NLDFT-based estimated pore size distribution obtained from N2 at -196 °C isotherms (assuming carbon slit-pore geometry) for TDA0-DMN, TDA1-DMN (R=C1) and TDA2-DMN (R=C2). 201

The pore size distributions (PSDs), calculated using the NLDFT (Non-Local Density

Functional Theory) model (Fig. 6.6), clearly show the gradient in the microporous texture of the TDA-DMN polyimides. TDA1-DMN had the largest fraction of ultramicropores (<

7 Å) compared to non-substituted polyimide or the polyimide containing the diethyl bridgehead moiety. The qualitative assessment of the microporous structure of the PIM-

PIs has a direct qualitative correlation with their gas permeation properties, as discussed below.

Wide-angle x-ray diffraction (WAXD) measurements were conducted on TDA0-

DMN and TDA2-DMN polyimide films. The spectra were then compared to TDA1-

DMN (Chapter 5) to investigate the effect of bridgehead substituents and their size on the average polymer chain spacing. X-ray diffraction confirmed that both materials were amorphous. TDA0-DMN has one wide peak centered at average d-spacing of 5.4 Å and

TDA2-DMN exhibited two peaks centered at average d-spacings of 7.2 and 5.0 Å, whereas TDA1-DMN had two peaks centered at average d-spacings of 6.5 and 4.4 Å.

Furthermore, the WAXD results indicate clearly that the dimethyl-bridgehead-substituted

TDA1-DMN contained a larger fraction of small ultramicroporous (< 4 Å) than its non- substituted analog (TDA0-DMN) and diethyl-based derivative (TDA2-DMN) (Fig. 6.7). 202

Fig. 6.7. WAXD spectra of TDA0-DMN, TDA1-DMN and TDA2-DMN films.

6.4.2. Gas Transport Properties

The pure-gas permeation properties of 1-day-old samples of the TDA-based polyimides determined at 2 bar and 35 °C are reported in Table 6.2. To obtain meaningful structure/gas permeation property relationships between the non-substituted-, dimethyl- and diethyl-triptycene bridgehead substitutions, the same polymer film preparation and permeation test protocols were applied for each polymer. 203

Table 6.2. Gas Permeabilities and permselectivities for TDA-DMN-based polyimides (2

bar and 35 °C) and other DMN-based PIM-PIs.

Pure-gas permeability (Barrer*) Permselectivity (훼)

Polymer He H2 N2 O2 CH4 CO2 O2/N2 H2/N2 CO2/CH4

TDA0-DMNa 833 2153 159 628 219 3291 3.9 14 15

TDA1-DMNb 1182 3047 182 783 216 3700 4.3 17 17

TDA2-DMNa 554 1230 56 229 63 1109 4.1 22 18

PIM-PI-10 [5] 300 670 84 270 168 2154 3.2 8.0 13

PIM-PI-12 [8] 1580 4230 369 1380 457 7340 3.7 11.5 16

SBFDA-DMN [9] - 2966 226 850 326 4700 3.8 13.1 14

* 1 Barrer = 1x10-10 cm3 (STP)·cm/cm2·s·cmHg.

a This study, 1-day old samples; b Chapter 5.

The TDA-DMN polyimide series exhibited very high gas permeabilities combined

with moderate selectivities, similar to previously published data of other DMN-based 204

PIM-PIs shown in Table 6.2 [5, 8, 9]. The permeabilities follow a rising trend with increasing polymer BET surface area: Pgas(TDA1-DMN) > Pgas(TDA0-DMN) >

Pgas(TDA2-DMN). The lowest permeability values were obtained for the diethyl-based

TDA2-DMN. The longer and more mobile diethyl bridgehead group probably occupied a larger fraction of the inter-chain polymer free volume elements relative to the non- substituted- or dimethyl-derivatives, leading to lower BET surface area and lower gas permeability. The same general trend has previously been documented in studies on the effect of side-chains in high-free-volume di-substituted polyacetylenes containing linear- and branched-alkyl side-chains [16, 17]. Interestingly, the O2/N2, H2/N2 and CO2/CH4 selectivities of the three TDA-DMN-based PIM-PIs were quite similar. Considering the relative trends of their NLDFT-derived pore size distributions (Fig. 6.6) and WAXD patterns (Fig. 6.7) it appears reasonable to suggest that: (i) differences in gas permeability as well as BET surface area were determined primarily by the amount of free volume and b) selectivities were similar due to only small variations in the fractions of ultramicropores (< 7 Å).

The permeability/selectivity trade-off behaviours of the three TDA-DMN polyimides for CO2/CH4, H2/N2 and O2/N2 separation are shown in Fig. 6.8, 6.9 and 6.10, respectively. Like previously reported DMN-based PIM-PIs, the TDA-DMN series exhibited gas separation performance close to or above the 2008 Robeson upper bounds

[18]. 205

Fig. 6.8. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for CO2/CH4.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 permeability/selectivity tradeoffs [18]. 206

Fig. 6.9. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for H2/N2.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 and

2015 permeability/selectivity tradeoffs [18, 19]. 207

Fig. 6.10. Gas separation performance of fresh (one day) TDA0-DMN, TDA1-DMN

(Chapter-5), TDA2-DMN polyimides and other DMN-based PIM-PIs for O2/N2.

Polymers are labeled as indicated in Fig. 5.2 and 6.3. The solid lines represent 2008 and

2015 permeability/selectivity tradeoffs [18, 19].

The collected data from Chapters 4, 5 and 6 for the gas transport measurements,

BET surface area and WAXD analysis of the five investigated triptycene building blocks

(non-substituted-, extended iptycene, dimethyl-, diethyl- and diisopropyl-substitutions), are summarized in Table 6.3. Based on this investigation it can be concluded that the 9,

10-dimethyl-substituted triptycene building block (Fig. 6.11) provides the optimum gas separation performance of the polyimide membrane for CO2/CH4 separation. This building block will be utilized further in this research to design novel membrane material for CO2/CH4 separation as presented in Chapter 7. 208

Table 6.3. Summary for the effect of additional benzene ring and 9,10-dialkyl- substitutions to the triptycene moiety (Chapters 4, 5 and 6).

Non-sub.-triptycene Extended Dimethyl-sub. Diethyl-sub. Diisopropyl-sub.

(a) compared to iptycene (b) (c) (d) (e)

Selectivity Decreased ~ no effect ~ no effect ~ no effect

Permeability Increased Increased Decreased ~ no effect

BET Surface Area Increased Increased Decreased ~ no effect

Fig. 6.11. Triptycene building units of (a) non-substituted-, (b) extended iptycene, (c) dimethyl-substituted, (d) diethyl-substituted and (e) diisopropyl-substituted. 209

6.5. Conclusions

Two intrinsically microporous polyimides were obtained by reaction of non- substituted- and 9,10-diethyl- (C2) triptycene-based dianhydride monomers and DMN via a one-step high-temperature solution polycondensation reaction. All polyimides exhibited good solubility in common organic solvents, excellent thermal stability, and high free volume as confirmed by high BET surface areas. This study is an extension for Chapter

5 to further demonstrate that selection of the bridgehead substituents can be used to fine- tune the microporous texture, pore size distribution and, consequently, gas permeation properties of triptycene-based polyimides. The dimethyl bridgehead groups in the triptycene building block provided the DMN-based polyimide (TDA1-DMN) with larger

2 surface area (760 m /g), greater fraction of ultramicroporosity, as observed from N2 and

CO2 NLDFT adsorption analysis, and higher gas permeability than TDA0-DMN and

TDA2-DMN. The non-substituted-triptycene-based polyimide (TDA0-DMN) showed similar selectivity but with lower permeability and BET surface area. The longer diethyl-

(TDA2-DMN) reduced the surface area and ultramicroporosity. The combination of high gas permeability and moderate selectivity places TDA1-DMN above the 2008 pure-gas permeability/selectivity upper bound for CO2/CH4. The 9,10-dimethyl (C1) bridgehead groups in the triptycene building block exhibited the best CO2/CH4 separation performance among others with the non-substituted, extended iptycene, diethyl- or diisopropyl-triptycene-based polyimides. 210

6.6. References

[1] J.L. Santiago-García, C. Álvarez, F. Sánchez, G. José, Gas transport properties of new aromatic polyimides based on 3, 8-diphenylpyrene-1, 2, 6, 7-tetracarboxylic dianhydride,

Journal of Membrane Science, 476 (2015) 442-448.

[2] D. Ayala, A. Lozano, J. De Abajo, C. Garcıa-Perez, J. De La Campa, K.-V.

Peinemann, B. Freeman, R. Prabhakar, Gas separation properties of aromatic polyimides,

Journal of Membrane Science, 215 (2003) 61-73.

[3] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational design of intrinsically ultramicroporous polyimides containing bridgehead-substituted triptycene for highly selective and permeable gas separation membranes, Macromolecules, 47

(2014) 5104-5114.

[4] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[5] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820. 211

[6] B.S. Ghanem, N.B. McKeown, P.M. Budd, D. Fritsch, Polymers of intrinsic microporosity derived from bis (phenazyl) monomers, Macromolecules, 41 (2008) 1640-

1646.

[7] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[8] Y. Rogan, R. Malpass-Evans, M. Carta, M. Lee, J.C. Jansen, P. Bernardo, G. Clarizia,

E. Tocci, K. Friess, M. Lanč, A highly permeable polyimide with enhanced selectivity for membrane gas separations, Journal of Materials Chemistry A, 2 (2014) 4874-4877.

[9] X. Ma, B. Ghanem, O. Salines, E. Litwiller, I. Pinnau, Synthesis and effect of physical aging on gas transport properties of a microporous polyimide derived from a novel spirobifluorene-based dianhydride, ACS Macro Letters, 4 (2015) 231-235.

[10] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra‐microporous triptycene‐ based polyimide membranes for high‐performance gas separation, Advanced Materials,

26 (2014) 3688-3692.

[11] T.M. Long, T.M. Swager, Minimization of free volume: alignment of triptycenes in liquid crystals and stretched polymers, Advanced Materials, 13 (2001) 601-604.

[12] N.G. White, M.J. MacLachlan, Soluble tetraaminotriptycene precursors, The Journal of Organic Chemistry, 80 (2015) 8390-8397. 212

[13] N.T. Tsui, A.J. Paraskos, L. Torun, T.M. Swager, E.L. Thomas, Minimization of internal molecular free volume: a mechanism for the simultaneous enhancement of polymer stiffness, strength, and ductility, Macromolecules, 39 (2006) 3350-3358.

[14] J.H. Chong, M.J. MacLachlan, Iptycenes in supramolecular and materials chemistry,

Chemical Society Reviews, 38 (2009) 3301-3315.

[15] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232.

[16] I. Pinnau, A. Morisato, Z. He, Influence of side-chain length on the gas permeation properties of poly (2-alkylacetylenes), Macromolecules, 37 (2004) 2823-2828.

[17] I. Pinnau, Z. He, A. Morisato, Synthesis and gas permeation properties of poly

(dialkylacetylenes) containing isopropyl-terminated side-chains, Journal of Membrane

Science, 241 (2004) 363-369.

[18] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400.

[19] R. Swaidan, B. Ghanem, I. Pinnau, Fine-tuned intrinsically ultramicroporous polymers redefine the permeability/selectivity upper bounds of membrane-based air and hydrogen separations, ACS Macro Letters, 4 (2015) 947-951.

213

Chapter 7. Triptycene Dimethyl-Bridgehead Dianhydride-Based Intrinsically

Microporous Hydroxyl-Functionalized Polyimide for Dual CO2 and N2 Removal in

Natural Gas Upgrading Application1

7.1. Abstract

The gas permeation properties of a high-performance hydroxyl-functionalized PIM- polyimide (TDA1-APAF) prepared from a novel 9,10-dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) and a commercially available 2,2-bis(3-amino-4- hydroxyphenyl)-hexafluoropropane (APAF) diamine monomer are reported. The microporous polymer had a BET surface area based on nitrogen adsorption of 260 m2/g.

A freshly prepared sample exhibited excellent gas permeation properties: (i) CO2 permeability of 40 Barrer coupled with a CO2/CH4 selectivity of 55 and (ii) H2 permeability of 94 Barrer with a H2/CH4 selectivity of 129. Physical aging over 250 days resulted in significantly enhanced CO2/CH4 and H2/CH4 selectivity of 75 and 183, respectively with only ~ 25% loss in CO2 and H2 permeability. Aged TDA1-APAF exhibited 5-fold higher pure-gas CO2 permeability (30 Barrer) and 2-fold higher

CO2/CH4 permselectivity over conventional dense cellulose triacetate membranes at 2 bar. In addition, TDA1-APAF polyimide had a N2/CH4 selectivity of 2.3, thereby making it possible to simultaneously treat CO2- and nitrogen-contaminated natural gas. Gas

1Portions of this chapter were adopted from Alghunaimi, F.; Ghanem, B.; Alaslai, N.; Mukaddam, M.; Pinnau, I. Journal of Membrane Science 2016 (520) 240-246. Alghunaimi performed gas permeation, sorption and characterization and wrote the manuscript. Ghanem synthesized the polymers. Litwiller, Alaslai and Swaidan directed construction of permeation apparatus and edited the manuscript. Pinnau supervised the work and edited the manuscript [1]. 214

mixture permeation experiments with a 1:1 CO2/CH4 feed mixture demonstrated higher mixed- than pure-gas selectivity and plasticization resistance up to 30 bar. These results suggest that intrinsically microporous hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for removal of CO2 and N2 from natural gas and hydrogen purification in petrochemical refinery applications.

7.2. Introduction

Worldwide demand for energy is projected to quickly expand caused by the continuing growth in global population. In the 2014 Annual Energy Outlook (AEO) report, total delivered energy consumption in the industrial sector was estimated to increase by 28% from 2012 to 2040 [2, 3]. Much of the growth will reflect natural gas use due to its relatively low carbon footprint, increased thermal efficiency, and cleaner burning benefits compared to other fossil fuels. Currently, natural gas accounts for about

23% of the world’s energy consumption and the International Energy Agency predicts that the demand for natural gas will grow by approximately 44% through 2035 [4]. In spite of this promising projection, raw natural gas typically contains significant amounts of impurities such as water vapor, carbon dioxide (CO2), hydrogen sulfide (H2S), nitrogen (N2) and other inert gases, which must be removed before pipeline delivery to end-users. Currently, amine absorption dominates acid gas separation technology, which reduces the CO2 content to meet the pipeline requirement (< 2%). However, absorption processes present environmental concerns besides high capital and maintenance costs of the large scrubber units [5]. It has been estimated that about 16% of all natural gas in the

United States contains higher nitrogen content (up to 15%) than the allowable maximum value (< 3%) according to pipeline specifications [4, 6]. Removal of nitrogen by 215 conventional separation technologies, such as cryogenic distillation, is extremely cost intensive [7]. One alternative technology is pressure-swing adsorption (PSA) using molecular sieves that preferentially adsorb nitrogen [8]. In principle, polymeric gas separation membranes can be applied in many processing steps during upgrading of natural gas, including simultaneous removal of CO2 and N2 [9]. Membrane processes can potentially offer more energy-efficient technology with low capital cost, small footprint, simple operation, and low maintenance, as well as minimal environmental impact [10,

11].

The main requirements for advanced polymeric membrane materials are: (i) high permeability, (ii) high selectivity, (iii) long-term durability and (iv) processability into thin-films [12]. In 1991, Robeson reported an inherent trade-off between permeability (P) and selectivity (α), that is, high permeability polymers typically have low selectivity and vice versa [13]. An updated version was reported in 2008 [14] where the trade-off curves moved upwards primarily due to the development of perfluorinated glassy polymers [15] and ladder polymers of intrinsic microporosity (PIMs) with pores < 20 Å [16-18].

The concept of PIMs was first reported by Mckeown and Budd in 2004 using spirobisindane-based ladder polymers [16-18]. For example, PIM-1 (Fig. 7.1) showed high permeability with moderate selectivity for the separation of O2/N2 and CO2/N2 [18-

21]. Thereafter, significant advances were achieved with the development of ladder polymers derived from AB-type triptycene- (TPIM-1 and TPIM-2) [22] and Tröger’s base-building blocks [23, 24] that significantly outperformed all polymers listed on the

2008 Robeson upper bound for O2/N2, H2/N2 and H2/CH4 separation [25]. Polyimides of intrinsic microporosity (PIM-PIs) designed with contorted and inflexible backbones 216 resulted also in membrane materials with promising gas separation properties [26-34].

Ghanem et al. [30, 35] and Rogan et al. [31] reported spirobisindane dianhydride-based

PIM-PIs which displayed significantly improved permeability among all known polyimides. The first PIM-PI derived from a triptycene diamine showed higher selectivity but at the cost of lower permeability; e.g. 6FDA-2,6-diaminotriptycene (6FDA-DAT1) had a CO2 permeability of 120 Barrer and a CO2/CH4 selectivity of 38 [36]. Recently,

Ghanem et al. and Swaidan et al. [26, 37] reported a series of PIM-PIs made from 9,10- bridgehead-substituted triptycene dianhydrides with significantly improved gas separation performance properties, which surpassed the 2008 upper bound for several gas pairs due to their strong size-sieving ultramicroporous structures. However, their applicability for efficient CO2/CH4 separation was limited by their relatively low selectivity.

Previous work demonstrated that 6FDA-derived PI membranes made from hydroxyl- containing diamines resulted in polyimides, such as 6FDA-APAF (Fig. 7.1), with the highest CO2/CH4 selectivities reported to date; however, the CO2 permeability of these hydroxyl-functionalized polyimides was relatively low (< 10 Barrer) [38, 39]. Recently, our group synthesized the first PIM-PI containing hydroxyl groups (PIM-6FDA-OH)

(Fig. 7.1), which exhibited notable performance for CO2/CH4 separation [32].

Furthermore, it was demonstrated that a PIM-PI-OH made from 9,10-triisopropyl- bridgehead triptycene dianhydride (TPDA) and 2,2-bis(3-amino-4-hydroxy-phenyl)- hexafluoropropane (APAF) exhibited enhanced CO2/CH4 selectivity relative to a TPDA- based polyimide derived from a related, purely hydrocarbon-based diamine (5,5’-

(hexafluoroisopropylidene)-di-o-toluidine (ATAF)). The effect of hydroxyl- 217 functionalization was elucidated with fluorescence spectroscopy, which clearly indicated the formation of a strong charge transfer complex (CTC). As a result, the OH-based polyimide exhibited more efficient intermolecular chain packing which led to lower CO2 sorption capacity but significantly enhanced gas sieving capability, as indicated by high diffusion selectivity [40].

Fig. 7.1. Structures of ladder PIM-1 and OH-functionalized PIM-PIs: 6FDA-APAF and

PIM-6FDA-OH.

In this chapter, the gas transport properties of a newly developed high-performance triptycene-based hydroxyl-functionalized PIM-PI membrane material are reported which was synthesized by a one-step polycondensation reaction between a commercial hydroxyl-containing diamine monomer (APAF) and a new 9,10-dimethyl-2,3,6,7- triptycene tetracarboxylic dianhydride (TDA1) monomer, as shown in Scheme 7.1. The 218 polyimide was fully characterized by FTIR, GPC, TGA, XRD, and BET surface area measurements.

Scheme 7.1. Synthesis of TDA1-APAF polyimide.

Additionally, pure-and mixed-gas permeation properties of fresh and physically aged membranes were analyzed to identify key structure/property relationships that could guide the rational design of highly selective and permeable PIM-PIs for natural gas sweetening and petrochemical refinery applications.

7.3. Experimental

7.3.1. Materials

The compounds 1,2-dimethoxybenzene, 2-aminobenzoic acid, trifluoromethanesulfonic anhydride, acetic anhydride, 1,1′- bis(diphenylphosphino)ferrocene, tris(dibenzylideneacetone)dipalladium(0), boron tribromide (BBr3), zinc cyanide, potassium hydroxide, and isoquinoline were obtained from Aldrich and used as received. 9,10-Dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) monomer was synthesized as described in the literature [41]. The monomer 2,2′-bis(3-amino-4-hydroxyphenyl)-hexafluoropropane (APAF) was purchased from Aldrich and purified by vacuum sublimation at 220 °C. m-Cresol was distilled 219 under reduced pressure and stored under nitrogen in the dark over 4 Å molecular sieves.

All other solvents were obtained from various commercial sources and used as received.

7.3.2. Polymer Characterization

Fourier transform infrared (FTIR) measurements were carried out using a Varian 670-

IR FTIR spectrometer. Gel permeation chromatography (GPC, Agilent 1200) was done using tetrahydrofuran (THF) as an eluent. Thermogravimetric analysis (TGA, TA Q-

5000) measurement was performed under nitrogen atmosphere with heating rate of 3

°C/min up to 800 °C. The BET surface area of the polymer was determined by N2 sorption at -196 °C using a Micromeritics ASAP-2020. The polymer powder was degassed under high vacuum at 150 °C for 16 hours prior to analysis. Maximum pore volume was identified at p/po< 0.96 of the N2 isotherm.

7.3.3. Synthesis of TDA1-APAF Polymer

A mixture of TDA1 (0.23 g, 0.54 mmol), APAF (0.2 g, 0.54 mmol) and 3 ml m-cresol was stirred in a Schlenk tube for 30 minutes under nitrogen at room temperature. The mixture was heated at 80 °C for 1 hour and catalytic amount of isoquinoline was added and the temperature was raised gradually to 200 °C and kept at that temperature until the mixture became very viscous. Water formed during imidization was removed with a stream of nitrogen. After cooling, the polyimide solution was added into methanol (250 ml) and the crude fibrous polymer was filtered, washed with methanol and dried. Re- precipitation by a THF–methanol mixture was carried out twice for further purification and the polymer was finally dried at 120 °C in a vacuum oven for 24 hours to give 86% yield as shown in the literature [1]. GPC (THF): Mn = 42,000 g/ mol, Mw = 106,000 220

g/mol relative to polystyrene, Mw/Mn = 2.5. TGA analysis: (nitrogen), thermal

2 degradation commences at Td ~ 480 °C. BET surface area = 260 m /g, total pore volume

3 = 0.301 cm /g at (p/p0 = 0.96, adsorption).

7.3.4. Polymer Film Preparation

A polymer solution (5 wt/vol%) in THF was filtered through 0.45 μm polypropylene filters, poured onto a flat glass Petri dish and slowly evaporated at room temperature for one day. The obtained membrane was dried at 120 °C. To remove any traces of residual solvent, the membrane was soaked in methanol for 10 h, air-dried, and then post-dried at

250 °C in a vacuum oven for 12 h. Tough films with thickness of 70 ± 5 μm were used for gas permeability measurements. Prior to the gas permeation tests, TGA experiments were performed to confirm that the polyimide film was solvent-free. Film thickness and effective area for gas permeation measurements were determined by a digital micrometer and scanner, respectively.

7.3.5. Pure-Gas Permeation Experiments

The pure-gas permeabilities of TDA1-APAF were determined by using the variable pressure/constant volume method. The polymer film sample was degassed in the permeation cell under vacuum for at least 24 hours. The pure-gas permeability of He, H2,

N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar. The gas permeability (P) was calculated by equation 2.14 where P is in Barrers (1 Barrer = 10-10 cm3(STP·cm/cm2·s·cmHg). The apparent diffusion coefficient D (cm2/s) was calculated by D = l2 /6 θ, where θ is the time lag of the permeability measurement and l is the membrane thickness. The ideal selectivity for a gas pair is given by equation 2.6. 221

7.3.6. Mixed-Gas Permeation Experiments

The mixed-gas permeation measurements of TDA1-APAF were performed at 35 °C using a custom-designed mixed-gas permeation system similar to that described by

O’Brien et al. [42]. The feed gas mixture contained 50 vol.% CH4/50 vol.% CO2 and the total feed pressure was varied between 4 and 30 bar; the permeate pressure was less than

0.01 bar. The permeate flow rate to feed flow rate, i.e. the stage-cut, was set at 0.01.

Applying these conditions, the residue composition was essentially equal to that of the feed gas. CO2 and CH4 permeate concentrations were detected with a gas chromatograph

(Agilent 3000A Micro GC) equipped with a thermal conductivity detector.

7.4. Results and Discussion

Our previous study on polyimides (Chapter-5) derived from a new 9,10-dimethyl-

2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) monomer and 3,3ʹ- dimethylnaphthidine (DMN) showed that the triptycene building block with bridgehead methyl side groups offers a polyimide with high BET surface area of 760 m2 g-1 and very high gas permeabilities (e.g. CO2 = 3700 Barrer) but with only moderate selectivities

(e.g. CO2/CH4 selectivity = 17). In this chapter, a new ortho-hydroxyl-functionalized polyimide (TDA1-APAF) was designed to enhance selectivities, while maintaining high gas permeability. The new polyimide was synthesized from TDA1 dianhydride [41] and a commercial hydroxyl-containing diamine monomer (APAF). The chemical structure of the polyimide was confirmed by FTIR and 1H NMR spectroscopy, as shown in the literature [1]. The FTIR spectrum of the polyimide film (Fig. 7.2) contained absorption 222 bands at approximately 1779 and 1710 cm-1 (C=O asymmetric and symmetric stretching),

1374 cm-1 (C-N stretching) and 846 cm-1 (imide ring deformation).

Fig. 7.2. FTIR spectrum of TDA1-APAF polyimide.

The polymer exhibited high molecular weight (Mw= 106,000 g mol-1) as determined by GPC and good solubility in THF from which dense films were cast for permeation studies.

7.4.1. Physical Properties and Microstructure of TDA1-APAF

The polymer showed high thermal stability in N2 atmosphere with an onset decomposition temperature of 480 °C determined by thermal gravimetric analysis, as shown in Fig. 7.3. 223

Fig. 7.3. Thermogravimetric analysis of TDA1-APAF polyimide film.

The BET surface area of TDA1-APAF derived from N2 sorption isotherm at -196 °C was 260 m2/g (Fig. 7.4). 224

Fig. 7.4. Physisorption isotherms for TDA1-APAF using N2 at -196 °C. Closed symbols: adsorption; open symbols: desorption.

Wide-angle x-ray diffraction (WAXD) measurements were conducted on a fresh

TDA1-APAF and a 250 days aged membrane to investigate the effect of changes in the average polymer chain packing on gas transport properties. Both membranes exhibited an amorphous halo centered at around d ≈ 5.3 Å (Fig. 7.5). In comparison with the fresh sample, the physically aged sample had one additional shoulder diffraction peak at d ≈

3.7 Å. This result implies that the aged polymer contained a new fraction of smaller pores in the low ultramicroporous range. This shift was expected to result in higher selectivities associated with lower gas permeabilities. 225

Fig. 7.5. WAXD curves of fresh and 250 days aged TDA1-APAF films.

7.4.2. Fresh and Physically Aged Pure-Gas Permeation Properties

Pure-gas permeation experiments were performed on a fresh, one-day old sample of

TDA1-APAF at 2 bar and 35 °C (Table 7.1). The gas permeability decreased in the order: H2 > He > CO2 > O2 > N2 > CH4 which confirmed the molecular sieving behavior of the polymer, by the fact that the H2 permeability was larger than that for CO2, opposite of the trend normally observed for high BET surface area PIMs materials [19].

Microporous glassy polymers undergo a physical aging process in which the chains slowly turn into a more tightly packed arrangement towards their equilibrium glassy state

[43, 44]. Thus, pure-gas permeation experiments were performed on the aged sample after 250 days at 2 bar and 35 °C (Table 7.1). The permeabilities for all gases decreased

by only ~ 25% loss for H2 and CO2 (푃퐻2= 73 Barrer and 푃퐶푂2= 30 Barrer) but selectivities increased significantly for H2/CH4 and CO2/CH4 with values of 183 and 75, respectively. 226

Moreover, the aged sample exhibited (푃푁2= 0.9 Barrer) and N2/CH4 of 2.3 which is

double the cellulose triacetate (N2/CH4 = 1.1) selectivity with 4-fold higher permeability

[45]. Table 7.1 shows the data from this work in comparison to previously reported

literature data for three commercial glassy membrane materials, that is, cellulose

triacetate, Matrimid and polysulfone.

Table 7.1. Pure-gas permeabilities and ideal selectivities for fresh and physically aged of

TDA1-APAF membranes (at 2 bar; 35 °C).

Pure-Gas Permeability (Barrer) Ideal Selectivity (훼)

Polymer He H2 N2 O2 CH4 CO2 CO2/CH4 H2/CH4 N2/CH4 O2/N2

TDA1-APAFa 92 94 1.5 8.5 0.73 40 55 129 2.1 5.7

Aged TDA1-APAFb (70) (73) (0.90) (6.4) (0.40) (30) (75) (183) (2.3) (7.1)

Cellulose acetate (DS 19.6 15.5 0.23 1.46 0.20 6.6 33 78 1.1 6.3 2.85) [45]

Matrimid [46] - 17.5 0.22 1.46 0.21 7.3 35 83 1.1 6.6

Polysulfone [47] 13 14 0.25 1.40 0.25 5.6 22 56 1.0 5.6

a membrane was soaked in methanol for 10 hrs; dried under vacuum at 250 °C for 12 hrs.

Tested after 1 day. 227 b tested after 250 days.

The pure-gas diffusion (D) and solubility (S) coefficients of the fresh and physically aged TDA1-APAF for N2, O2, CH4 and CO2 were calculated by the time-lag method and are shown in Table 7.2. As expected, the tighter aged membrane displayed lower D and

S values compared to the fresh sample. Concurrently, the enhanced ultramicroporosity in the aged TDA1-APAF sample resulted in improved size-sieving properties and, therefore, higher O2/N2, N2/CH4 and CO2/CH4 diffusivity selectivity values than those of the fresh

PIM-PI sample (Table 7.3).

Table 7.2. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for fresh and physically aged TDA1-APAF films.

Diffusion coefficient Solubility coefficient

(10-8 cm2/s) (10-2 cm3(STP)/(cm3 cmHg))

Polymer N2 O2 CH4 CO2 N2 O2 CH4 CO2

TDA1-APAFa 0.82 4.5 0.12 1.9 1.8 1.9 6.1 21.1

Aged TDA1-APAFb (0.60) (4.0) (0.08) (1.6) (1.5) (1.6) (5.0) (18.5)

a tested after 1 day. b tested after 250 days. 228

Table 7.3. Diffusivity selectivities and solubility selectivities of fresh and physically

aged TDA1-APAF films (2 bar; 35 °C).

Diffusivity selectivity (α)D Solubility selectivity (α)S

Polymer O2/N2 N2/CH4 CO2/CH4 O2/N2 N2/CH4 CO2/CH4

TDA1-APAFa 5.5 6.83 15.8 1.06 0.30 3.5

Aged TDA1-APAFb (6.7) (7.5) (20) (1.07) (0.30) (3.7)

a tested after 1 day.

b tested after 250 days.

7.4.3. Fresh and Physically Aged Mixed-Gas Permeation Properties

Commonly, high sorption uptake of condensable gases like CO2 can dilate the

polymer matrix, induce increased chain mobility, modify the pore structure and,

therefore, significantly affect the gas transport properties of glassy polymers. The pure-

and mixed-gas permeabilities (Figs. 7.6 and 7.7) and selectivities (Fig. 7.8) were

measured using a 1:1 CO2/CH4 feed mixture with increasing CO2 partial pressure (2, 5, 7,

10, 12 and 15 bar) for fresh and aged TDA1-APAF samples. The pure-gas CO2

permeabilities in both fresh and aged membranes decreased with increasing feed pressure

up to 15 bar (Fig. 7.6), similar to the behavior of related PIM-PIs previously reported [36,

40, 48, 49]. This can be rationalized by a decrease in CO2 solubility coefficient with 229 pressure as is typically observed in glassy polymers. In addition, competitive sorption occurs between CO2 and CH4 for available sorption sites which may reduce the overall sorbed concentration of CO2 in the polymer. As a result, the mixed-gas CO2 permeabilities were lower than the pure-gas CO2 permeability values determined at the same partial pressure.

Fig. 7.6. Pressure-dependence of pure- and mixed-gas CO2 permeabilities for CTA [48], fresh and physically aged TDA1-APAF polyimide (1:1 CO2/CH4 mixture; 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas.

Interestingly, the mixed-gas CH4 permeability was also lower than the pure-gas value over the entire pressure range, potentially due to partial blocking of the micropores by preferential sorption of CO2, as shown in Fig. 7.7. Similar behavior has previously been observed for related PIM-PIs [40] and polybenzoxazoles derived from thermally- rearranged polyimides [50, 51]. 230

Fig. 7.7. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for CTA [48], fresh and physically aged TDA1-APAF (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas.

As a result, the mixed-gas permselectivites of CO2/CH4 were slightly higher than those determined under pure-gas conditions (Fig. 7.8) due to a reduction of the mixed-gas

CH4 permeability by co-permeation of CO2, as previously observed in TPDA-APAF, and some thermally-rearranged polyimides [50, 51]. Under typical natural gas sweetening conditions, where the partial pressure of CO2 is often near 5-10% in a 60-70 bar feed, the aged membrane exhibited a very high selectivity of ~75. This high selectivity significantly benefits the membrane process economics as it mitigates the amount of CH4 loss into the permeate, reducing the need for recompression of the product CH4 in a second stage and improving the economics of the separation [52]. The commercial material (cellulose triacetate) mixed-gas properties were included in Figs. 7.6, 7.7 and 7.8 for comparison purposes. 231

Fig. 7.8. Pressure-dependence of mixed-gas CO2/CH4 selectivities for CTA [48], fresh and physically aged TDA1-APAF samples (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye: open points, pure-gas; closed points, mixed-gas.

The permeability/selectivity trade-off behavior of TDA1-APAF polyimide for

H2/CH4 and CO2/CH4 separation (Fig. 7.9a-b) is similar to the previously reported

APAF-based PIM-PI, where the gas separation performance was close or on the 2008

Robeson upper bounds [14]. Furthermore, the mixed-gas performance of the fresh and aged TDA1-APAF polyimide membrane for CO2/CH4 separation is compared to related

PIMs and TR polymers evaluated under the same test conditions (35 °C; 10 bar CO2 partial pressure) as presented in Fig. 7.10. The aged intrinsically microporous TDA1-

2 APAF polyimide (surface area of 260 m /g) exhibited almost similar mixed-gas CO2/CH4 selectivity with two- to four-fold higher CO2 permeability compared to low free volume

6FDA-based hydroxyl-functionalized polyimides (6FDA-APAF and 6FDA-DAP). 232

Fig. 7.9. Gas separation performance of TDA1-APAF polyimide for (a) H2/CH4 and (b)

CO2/CH4. Open points: pure-gas feeds; closed points: 1:1 CO2/CH4 mixed-gas feeds.

Polymers are labeled as indicated in Fig. 7.1 and Scheme 7.1. ( ) refers to performance after physical aging in days. The solid lines represent 2008 and 2015 permeability/selectivity trade-offs [14, 25]. 233

Fig. 7.10. CO2/CH4 mixed-gas permeability/selectivity trade-off curve for TDA1-APAF

(this study), KAUST-PI-1 [48], AO-PIM-1 [53], 6FDA-DAT1 [36], 6FDA-DAT2 [36],

TPDA-APAF [40], 6FDA-DAP [49], 6FDA-APAF [40], 6FDA-mPDA [49], TR6FDA-

HAB [50], TRPIM-6FDA-OH [51] and CTA [48]. All experiments were performed with a 50/50 (v/v) CO2/CH4 mixture at 20 bar feed pressure and 35°C using the constant volume/variable pressure technique.

7.5. Conclusions

In this chapter, a new TDA1-APAF polyimide, synthesized from 9,10-dimethyl-

2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) and a commercial hydroxyl- diamine (APAF) via one-step high-temperature solution imidization reaction, is reported.

This polymer exhibited high molecular weight, good solubility and high thermal stability.

The pure-gas permeation data showed that introducing hydroxyl groups to the polyimide leads to a significant increase in permselectivity due to a large increase in diffusivity 234

selectivity for a variety of gas pairs, most notably H2/CH4 and CO2/CH4. The triptycene- based hydroxyl-containing polyimide showed a CO2 permeability of 21 under binary 1:1

CO2/CH4 mixed-gas feed with a selectivity of 72 at a partial CO2 pressure of 10 bar.

These properties are significantly better than those of cellulose triacetate, which exhibits

푃퐶푂2= 8 Barrer and CO2/CH4 selectivity of 25 when tested under the same conditions.

Moreover, aged TDA1-APAF polyimide had a N2/CH4 selectivity of 2.3, thereby making it possible to simultaneously treat CO2- and nitrogen-contaminated natural gas. These dramatically enhanced properties make TDA1-APAF polyimide an excellent candidate material for development of asymmetric or thin-film composite membranes for industrial natural gas sweetening.

235

7.6. References

[1] F. Alghunaimi, B. Ghanem, N. Alaslai, M. Mukaddam, I. Pinnau, Triptycene dimethyl-bridgehead dianhydride-based intrinsically microporous hydroxyl- functionalized polyimide for natural gas upgrading, Journal of Membrane Science, 520

(2016) 240-246.

[2] S. Yi, X. Ma, I. Pinnau, W.J. Koros, A high-performance hydroxyl-functionalized polymer of intrinsic microporosity for an environmentally attractive membrane-based approach to decontamination of sour natural gas, Journal of Materials Chemistry A, 3

(2015) 22794-22806.

[3] T.E. Rufford, S. Smart, G.C. Watson, B. Graham, J. Boxall, J.D. Da Costa, E. May,

The removal of CO2 and N2 from natural gas: a review of conventional and emerging process technologies, Journal of Petroleum Science and Engineering, 94 (2012) 123-154.

[4] J. Kuo, K. Wang, C. Chen, Pros and cons of different nitrogen removal unit (NRU) technology, Journal of Natural Gas Science and Engineering, 7 (2012) 52-59.

[5] B. Bhide, A. Voskericyan, S. Stern, Hybrid processes for the removal of acid gases from natural gas, Journal of Membrane Science, 140 (1998) 27-49.

[6] T. Rufford, S. Smart, G. Watson, B. Graham, J. Boxall, J.D. da Costa, E. May, The removal of CO2 and N2 from natural gas: a review of conventional and emerging process technologies, Journal of Petroleum Science and Engineering, 94 (2012) 123-154.

[7] B. Ohs, J. Lohaus, M. Wessling, Optimization of membrane based nitrogen removal from natural gas, Journal of Membrane Science, 498 (2016) 291-301. 236

[8] K.A. Lokhandwala, I. Pinnau, Z. He, K.D. Amo, A.R. DaCosta, J.G. Wijmans, R.W.

Baker, Membrane separation of nitrogen from natural gas: a case study from membrane synthesis to commercial deployment, Journal of Membrane Science, 346 (2010) 270-279.

[9] T. Kim, W. Koros, G. Husk, K. O'brien, “Reverse permselectivity” of N2 over CH4 in aromatic polyimides, Journal of Applied Polymer Science, 34 (1987) 1767-1771.

[10] R.W. Baker, K. Lokhandwala, Natural gas processing with membranes: an overview, Industrial & Engineering Chemistry Research, 47 (2008) 2109-2121.

[11] C.A. Scholes, G.W. Stevens, S.E. Kentish, Membrane gas separation applications in natural gas processing, Fuel, 96 (2012) 15-28.

[12] J.L. Santiago-García, C. Álvarez, F. Sánchez, G. José, Gas transport properties of new aromatic polyimides based on 3, 8-diphenylpyrene-1, 2, 6, 7-tetracarboxylic dianhydride, Journal of Membrane Science, 476 (2015) 442-448.

[13] L.M. Robeson, Correlation of separation factor versus permeability for polymeric membranes, Journal of Membrane Science, 62 (1991) 165-185.

[14] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400.

[15] D.F. Sanders, Z.P. Smith, R. Guo, L.M. Robeson, J.E. McGrath, D.R. Paul, B.D.

Freeman, Energy-efficient polymeric gas separation membranes for a sustainable future: a review, Polymer, 54 (2013) 4729-4761. 237

[16] P.M. Budd, E.S. Elabas, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib,

C.E. Tattershall, D. Wang, Solution-processed, organophilic membrane derived from a polymer of intrinsic microporosity, Advanced Materials, 16 (2004) 456-459.

[17] P.M. Budd, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib, C.E.

Tattershall, Polymers of intrinsic microporosity (PIMs): robust, solution-processable, organic nanoporous materials, Chemical Communications, (2004) 230-231.

[18] N.B. McKeown, P.M. Budd, Polymers of intrinsic microporosity (PIMs): organic materials for membrane separations, heterogeneous catalysis and hydrogen storage,

Chemical Society Reviews, 35 (2006) 675-683.

[19] P.M. Budd, K.J. Msayib, C.E. Tattershall, B.S. Ghanem, K.J. Reynolds, N.B.

McKeown, D. Fritsch, Gas separation membranes from polymers of intrinsic microporosity, Journal of Membrane Science, 251 (2005) 263-269.

[20] P.M. Budd, N.B. McKeown, B.S. Ghanem, K.J. Msayib, D. Fritsch, L. Starannikova,

N. Belov, O. Sanfirova, Y. Yampolskii, V. Shantarovich, Gas permeation parameters and other physicochemical properties of a polymer of intrinsic microporosity: polybenzodioxane PIM-1, Journal of Membrane Science, 325 (2008) 851-860.

[21] P.M. Budd, N.B. McKeown, D. Fritsch, Free volume and intrinsic microporosity in polymers, Journal of Materials Chemistry, 15 (2005) 1977-1986.

[22] B.S. Ghanem, R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, Energy‐efficient hydrogen separation by AB‐type ladder‐polymer molecular sieves, Advanced Materials, 26 (2014)

6696-6700. 238

[23] M. Carta, M. Croad, R. Malpass‐Evans, J.C. Jansen, P. Bernardo, G. Clarizia, K.

Friess, M. Lanč, N.B. McKeown, Triptycene induced enhancement of membrane gas selectivity for microporous Tröger's base polymers, Advanced Materials, 26 (2014) 3526-

3531.

[24] I. Rose, M. Carta, R. Malpass-Evans, M.-C. Ferrari, P. Bernardo, G. Clarizia, J.C.

Jansen, N.B. McKeown, Highly permeable benzotriptycene-based polymer of intrinsic microporosity, ACS Macro Letters, 4 (2015) 912-915.

[25] R. Swaidan, B. Ghanem, I. Pinnau, Fine-tuned intrinsically ultramicroporous polymers redefine the permeability/selectivity upper bounds of membrane-based air and hydrogen separations, ACS Macro Letters, 4 (2015) 947-951.

[26] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra‐microporous triptycene‐ based polyimide membranes for high‐performance gas separation, Advanced Materials,

26 (2014) 3688-3692.

[27] N. Ritter, I. Senkovska, S. Kaskel, J. Weber, Intrinsically microporous poly(imide)s: structure−porosity relationship studied by gas sorption and X-ray scattering,

Macromolecules, 44 (2011) 2025-2033.

[28] J. Weber, O. Su, M. Antonietti, A. Thomas, Exploring polymers of intrinsic microporosity-microporous, soluble polyamide and polyimide, Macromolecular Rapid

Communications, 28 (2007) 1871-1876.

[29] Y. Rogan, R. Malpass-Evans, M. Carta, M. Lee, J.C. Jansen, P. Bernardo, G.

Clarizia, E. Tocci, K. Friess, M. Lanč, A highly permeable polyimide with enhanced 239 selectivity for membrane gas separations, Journal of Materials Chemistry A, 2 (2014)

4874-4877.

[30] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[31] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820.

[32] X.H. Ma, R. Swaidan, Y. Belmabkhout, Y.H. Zhu, E. Litwiller, M. Jouiad, I. Pinnau,

Y. Han, Synthesis and gas transport properties of hydroxyl-functionalized polyimides with intrinsic microporosity, Macromolecules, 45 (2012) 3841-3849.

[33] X. Ma, O. Salinas, E. Litwiller, I. Pinnau, Novel spirobifluorene-and dibromospirobifluorene-based polyimides of intrinsic microporosity for gas separation applications, Macromolecules, 46 (2013) 9618-9624.

[34] Y. Zhuang, J.G. Seong, Y.S. Do, H.J. Jo, Z. Cui, J. Lee, Y.M. Lee, M.D. Guiver,

Intrinsically microporous soluble polyimides incorporating Tröger’s base for membrane gas separation, Macromolecules, 47 (2014) 3254-3262. 240

[35] B.S. Ghanem, N.B. McKeown, P.M. Budd, J.D. Selbie, D. Fritsch, High- performance membranes from polyimides with intrinsic microporosity, Advanced

Materials, 20 (2008) 2766-2771.

[36] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[37] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational design of intrinsically ultramicroporous polyimides containing bridgehead-substituted triptycene for highly selective and permeable gas separation membranes, Macromolecules, 47

(2014) 5104-5114.

[38] C.H. Jung, Y.M. Lee, Gas permeation properties of hydroxyl-group containing polyimide membranes, Macromolecular Research, 16 (2008) 555-560.

[39] S.A. Stern, H. Kawakami, A.Y. Houde, G. Zhou, Material and process for separating carbon dioxide from methane, Google Patents, 1997.

[40] R. Swaidan, B. Ghanem, E. Litwiller, I. Pinnau, Effects of hydroxyl- functionalization and sub-Tg thermal annealing on high pressure pure-and mixed-gas

CO2/CH4 separation by polyimide membranes based on 6FDA and triptycene-containing dianhydrides, Journal of Membrane Science, 475 (2015) 571-581.

[41] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232. 241

[42] K. O'Brien, W. Koros, T. Barbari, E. Sanders, A new technique for the measurement of multicomponent gas transport through polymeric films, Journal of Membrane Science,

29 (1986) 229-238.

[43] K.D. Dorkenoo, P.H. Pfromm, Accelerated physical aging of thin poly[1-

(trimethylsilyl)-1-propyne] films, Macromolecules, 33 (2000) 3747-3751.

[44] P.H. Pfromm, The impact of physical aging of amorphous glassy polymers on gas separation membranes, Materials Science of Membranes for Gas and Vapour Separation,

(2006) 293-306.

[45] A.C. Puleo, D.R. Paul, S.S. Kelley, The effect of degree of acetylation on gas sorption and transport behavior in cellulose-acetate, Journal of Membrane Science, 47

(1989) 301-332.

[46] Y. Zhang, I.H. Musselman, J.P. Ferraris, K.J. Balkus, Gas permeability properties of

Matrimid® membranes containing the metal-organic framework Cu–BPY–HFS, Journal of Membrane Science, 313 (2008) 170-181.

[47] C. Aitken, W. Koros, D. Paul, Effect of structural symmetry on gas transport properties of polysulfones, Macromolecules, 25 (1992) 3424-3434.

[48] R. Swaidan, B. Ghanem, M. Al-Saeedi, E. Litwiller, I. Pinnau, Role of intrachain rigidity in the plasticization of intrinsically microporous triptycene-based polyimide membranes in mixed-gas CO2/CH4 separations, Macromolecules, 47 (2014) 7453-7462.

[49] N. Alaslai, B. Ghanem, F. Alghunaimi, E. Litwiller, I. Pinnau, Pure-and mixed-gas permeation properties of highly selective and plasticization resistant hydroxyl-diamine- 242

based 6FDA polyimides for CO2/CH4 separation, Journal of Membrane Science, 505

(2016) 100–107.

[50] K.L. Gleason, Z.P. Smith, Q. Liu, D.R. Paul, B.D. Freeman, Pure-and mixed-gas permeation of CO2 and CH4 in thermally rearranged polymers based on 3, 3′-dihydroxy-

4, 4′-diamino-biphenyl (HAB) and 2, 2′-bis-(3, 4-dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane Science, 475 (2015) 204-214.

[51] R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, High pressure pure-and mixed-gas separation of CO2/CH4 by thermally-rearranged and carbon molecular sieve membranes derived from a polyimide of intrinsic microporosity, Journal of Membrane Science, 447

(2013) 387-394.

[52] J.D. Wind, S.M. Sirard, D.R. Paul, P.F. Green, K.P. Johnston, W.J. Koros, Carbon dioxide-induced plasticization of polyimide membranes: pseudo-equilibrium relationships of diffusion, sorption, and swelling, Macromolecules, 36 (2003) 6433-6441.

[53] R. Swaidan, B.S. Ghanem, E. Litwiller, I. Pinnau, Pure-and mixed-gas CO2/CH4 separation properties of PIM-1 and an amidoxime-functionalized PIM-1, Journal of

Membrane Science, 457 (2014) 95-102.

243

Chapter 8. Synthesis and Gas Permeation Properties of a Novel Thermally-

Rearranged Polybenzoxazole Made from an Intrinsically Microporous Hydroxyl-

Functionalized Triptycene-Based Polyimide Precursor

8.1. Abstract

A hydroxyl-functionalized triptycene-based polyimide of intrinsic microporosity

(TDA1-APAF) (Chapter 7) was converted to a polybenzoxazole (PBO) by heat treatment at 460 °C under nitrogen atmosphere. TDA1-APAF treated for 15 minutes (TR

460) resulted in a PBO conversion of 95% based on a theoretical weight loss of 11.7 wt.% of the polyimide precursor. The BET surface area of the TR 460 (680 m2/g) was significantly higher than that of the TDA1-APAF polyimide (260 m2/g) as determined by nitrogen adsorption at -196 °C. Heating TDA1-APAF for 30 minutes (TRC 460) resulted in a weight loss of 13.5 wt.%, indicating full conversion to PBO and partial main-chain degradation. The TR 460 membrane displayed excellent O2 permeability of 311 Barrer coupled with an O2/N2 selectivity of 5.4 and CO2 permeability of 1328 Barrer with a

CO2/CH4 selectivity of 27. Interestingly, physical aging over 150 days resulted in enhanced O2/N2 selectivity of 6.3 with an O2 permeability of 185 Barrer. The novel triptycene-based TR 460 PBO outperformed all previously reported APAF-polyimide- based PBOs with gas permeation performance close to recently reported polymers located on the 2015 O2/N2 upper bound. Based on this study, thermally-rearranged membranes from hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for air separation, specifically in applications where space and weight of membrane systems are of utmost importance such as nitrogen production for inert atmospheres in fuel lines and tanks on aircrafts and off-shore oil- or natural gas 244 platforms. Mixed-gas permeation experiments also demonstrated good performance of the TR 460 membrane for natural gas sweetening with a CO2 permeability of ~1,000

Barrer and CO2/CH4 selectivity of 22 at a typical CO2 wellhead partial pressure of 10 bar.

8.2. Introduction

Membrane-based gas separation is a well-established industrial technology with great potential for a wide diversity of large-scale industrial applications, such as natural gas sweetening, hydrogen recovery, CO2 separation from flue gas, and air separation [1-3]. It was estimated in 2002 that ~50% of the membrane-based gas separation market was based on onsite nitrogen production from air [3]. The success of membranes in air separation can be attributed to the ubiquitous clean feed gas, a simple single-stage process design and the delivery of the nitrogen product gas at feed pressure. Commercial polymeric membrane materials, including tetrabromo-polycarbonate, Matrimid® polyimide and polysulfone, are offering selectivities of ~6 to 7 that can produce 99% nitrogen but with low oxygen permeability of ~1-2 Barrer [3]. Currently, poly(phenylene oxide) is the most permeable polymer membrane used for commercial air separation with an oxygen permeability of 16 Barrer and O2/N2 selectivity of 4.9 [4]. It is interesting to note that these commercial air separation membrane materials fall far below the O2/N2 permeability/selectivity trade-off curve reported by Robeson in 1991 [5]. Hence, there is a strong incentive for development of advanced polymers that exhibit better performance for air separation. Specifically, in applications where space and membrane system weight limitations play a major role, such as nitrogen blanketing of fuel tanks on aircrafts and oil and natural gas offshore platforms, require polymer membranes with higher permeability 245

without significant loss of O2/N2 selectivity compared to commercial air separation membranes.

Polymers of intrinsic microporosity (PIMs), first introduced by Budd and McKeown in 2004, are a rapidly expanding class of amorphous glassy ladder polymers [6, 7]. PIMs have an inherent, interconnected porosity with micropores of less than 20 Å resulting from inefficient packing of polymer chains in the solid state [8, 9]. PIM-1 (Fig. 8.1), the prototype of this materials class is a spirobisindane-based ladder polymer, which has a

Brunauer-Teller-Emmett (BET) surface of ~860 m2/g, excellent solution processability and thus good thin-film formation properties. Furthermore, PIM-1 and a related bis(phenazyl)-derived spirobisindane ladder polymer (PIM-7) demonstrated very high gas permeability and moderate selectivity for a variety of gas pairs, defining the 2008 upper bound for O2/N2 separation [10].

More recently, intrinsically microporous polyimides (PIM-PIs) were developed with performance close to the 2008 upper bounds for several important gas pairs [11-14].

Ghanem et al. [11, 12] and Rogan et al. [15] reported the first series of PIM-PIs based on spirobisindane dianhydrides and various diamines, which displayed excellent gas permeability with moderate selectivity. Thereafter, PIM-PIs with enhanced selectivity were prepared with newly designed sterically hindered monomers based on spirobifluorene-, ethanoanthracene-, Tröger’s base-, and triptycene-building blocks [12-

20]. Recently, our group synthesized TDA1-APAF, a hydroxyl-functionalized triptycene- based PIM-PI (Fig. 8.1), which exhibited notable performance for O2/N2 and CO2/CH4 separation [21]. 246

Fig. 8.1. Structures of PIM-1 and OH-functionalized PIM-PI (TDA1-APAF).

An alternative method to introduce microporosity into glassy polymers is based on the formation of polybenzoxazoles (PBO) by a high-temperature (> 400 °C) decarboxylation and rearrangement reaction of polyimide precursors bearing hydroxyl- functional groups in ortho-position to the imide linkages [22]. In 2007, Park et al. reported gas permeation properties of a family of thermally-rearranged (TR) PBO polymers, which exhibited: (i) exceptional CO2 permeability and good CO2/CH4 selectivity, (ii) resistance to plasticization, and (iii) excellent chemical stability [22-24].

Using soluble hydroxyl-functionalized polyimide precursors allows fabrication of hollow fibers using conventional technology and subsequent conversion of the polyimides to

PBOs by thermal treatment with desirable transport properties [25, 26]. Most TR-derived

PBO polymers are based on low-free-volume polyimide precursors made from 4,4'-

(hexafluoroisopropylidene)diphthalic anhydride (6-FDA) and 2,2′-bis(3-amino-4- hydroxyphenyl)-hexafluoropropane (APAF) or 3,3'-dihydroxy-4,4'-diamino-biphenyl

(HAB) [23, 24, 27-32]. So far, only a few reports described the use of hydroxyl- functionalized intrinsically microporous polyimides (PIM-PIs) as precursors for the formation of PBOs, including spirobisindane- and spirobifluorene-based building blocks

[33-36]. Recently, a PBO derived from a low-free-volume triptycene dianhydride/APAF- 247 based polyimide (TPHI) showed commendable gas separation performance. The PBO made by heat treatment of TPHI at 400 °C for 2 hours (PBO conversion of 87% based on theoretical weight loss of 9.7 wt.% for full conversion) showed CO2, N2, and CH4 permeabilities of 320, 16 and 8.3 Barrer, respectively, combined with CO2/N2 and

CO2/CH4 selectivities of 20 and 39 [32]. Unfortunately, O2 permeability data were not reported in this work.

In this chapter, the synthesis (Scheme 8.1) and gas transport properties of a TR membrane, prepared from a recently reported intrinsically microporous triptycene-based hydroxyl-functionalized PIM-PI [21], were reported. The pristine polyimide was prepared by a one-step polycondensation reaction between a commercial hydroxyl- containing diamine monomer (APAF) and 9,10-dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) monomer, as previously reported by our group [21,

37]. The TR membranes were characterized by FTIR, BET, XRD and TGA measurements. Additionally, pure-and mixed-gas permeation properties of pristine

TDA1-APAF and TDA1-APAF-derived TR membranes (fresh and physically aged) were performed to demonstrate their potential as highly permeable and selective membranes for air and natural gas separation applications.

Scheme 8.1. Synthesis of TDA1-APAF-derived TR membrane. 248

8.3. Experimental

8.3.1. Materials

The compounds 1,2-dimethoxybenzene, 2-aminobenzoic acid, trifluoromethanesulfonic anhydride, acetic anhydride, 1,1′- bis(diphenylphosphino)ferrocene, tris(dibenzylideneacetone) dipalladium(0), boron tribromide (BBr3), zinc cyanide, potassium hydroxide, and isoquinoline were obtained from Aldrich and used as received. 9,10-Dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) monomer was synthesized as previously described [37]. The monomer 2,2′-bis(3-amino-4-hydroxyphenyl)-hexafluoropropane (APAF) was purchased from Aldrich and purified by vacuum sublimation at 220 °C. m-Cresol was distilled under reduced pressure and stored under nitrogen in the dark over 4 Å molecular sieves.

All other solvents were obtained from various commercial sources and used as received

[21].

8.3.2. Polymer Characterization of Pristine and Thermally Rearranged TDA1-

APAF Membranes

Fourier transform infrared (FTIR) measurements were carried out using a Varian 670-

IR spectrometer. Thermogravimetric analysis (TGA) was performed with a TG 209 F1

(Netzsch) coupled with a NETZSCH QMS 403 C Aëolos® mass spectrometer (MS) under nitrogen atmosphere with a heating rate of 5 °C/min up to 800 °C. The BET surface areas of the pristine TDA1-APAF and TR 460 were determined by N2 sorption at

-196 °C using a Micromeritics ASAP-2020. Powder samples were degassed under high vacuum at 120 °C for 15h prior to analysis. Analysis of the pore size distributions was 249

performed using the NLDFT (Non-Local Density Functional Theory) model using N2 (at

-196 °C) sorption isotherms for carbon-slit pore geometry provided by ASAP-2020 version 4.02 software. Wide-angle x-ray diffraction (WAXD) patterns of the fresh and aged TR 460 films were recorded in the reflection mode at room temperature on a Bruker

D8 Advance diffractometer (wavelength λ = 1.54 Å). The average d-spacing values of the polymers were calculated using Bragg’s law in the 2θ range of 7 to 40°.

8.3.3. Synthesis of TDA1-APAF, Polymer Film Preparation and Thermal

Conversion

A mixture of TDA1 (0.23 g, 0.54 mmol), APAF (0.2 g, 0.54 mmol) and 3 ml m-cresol was stirred in a Schlenk tube for 30 minutes under nitrogen at room temperature. The mixture was heated at 80 °C for 1 hour and 3 drops of isoquinoline were added as catalyst. The temperature was then raised gradually to 200 °C and kept at that temperature until the mixture became very viscous. Water formed during imidization was removed with a stream of nitrogen. After cooling, the polyimide solution was added into methanol (250 ml) and the crude fibrous polymer was filtered, washed with methanol and dried. Re-precipitation by a THF–methanol mixture was carried out twice for further purification and the polymer was finally dried at 120 °C in a vacuum oven for 24 hours to give 86% yield as previously reported [21]. A polymer solution (5 wt/vol%) in THF was filtered through 0.45 μm polypropylene filters, poured onto a flat glass Petri dish and slowly evaporated at room temperature for one day. The obtained membrane was dried at

120 °C. To remove any traces of residual solvent, the membrane was soaked in methanol for 10 h, air-dried, and then post-dried at 250 °C in a vacuum oven for 12 h. 250

Polyimide (TDA1-APAF) film samples were converted to PBOs membranes by thermal treatment in a Carbolite tubular furnace. Pyrolysis was conducted under nitrogen flow and polymer samples were sandwiched between two ceramic plates. Metal spacers were inserted between the ceramic plates to allow nitrogen to continuously flush the samples. The oxygen concentration in the furnace was monitored with an O2 analyser

(Cambridge Sensotech, Rapidox 3100). Ramping of the treatment temperature did not begin until the O2 concentration was less than 2 ppm. The thermally rearranged membranes were made at a rate of 10 °C/min from room temperature to 460 °C where it was held isothermally for 15 and 30 minutes to achieve conversion of the polyimide to

PBO. The furnace was then shut off and the film was left to cool to room temperature under continuous nitrogen flow. Film thickness and effective area for gas permeation measurements were determined by a digital micrometer and scanner, respectively. PBO films with thickness of 70 ± 5 μm were used for gas permeation experiments. Samples for long-term physical aging studies were stored in a desiccator under vacuum.

8.3.4. Pure-Gas Permeation Experiments

The pure-gas permeabilities of thermally rearranged (TR) TDA1-APAF membranes were determined by using the variable pressure/constant volume method. The polymer film sample was degassed in the permeation cell under vacuum for at least 24 hours. The pure-gas permeability of He, H2, N2, O2, CH4, and CO2 was measured at 35 °C and 2 bar.

The gas permeability (P) was calculated by equation 2.14 where P is in Barrers (1 Barrer

= 10-10 cm3(STP·cm/cm2·s·cmHg). The apparent diffusion coefficient D (cm2/s) was calculated by D = l2 /6 θ, where θ is the time lag of the permeability measurement and l is the membrane thickness. The solubility coefficient S (cm3(STP)/cm3 cmHg) was then 251 calculated from the solution-diffusion gas transport relationship: S = P/D. The ideal selectivity for a gas pair is given by equation 2.6.

8.3.5. Mixed-Gas Permeation Experiments

The mixed-gas permeation measurements of TR 460 were performed at 35 °C using a custom-designed mixed-gas permeation system similar to that described by O’Brien et al.

[39]. The feed gas mixture contained 50 vol.% CH4/50 vol.% CO2 and the total feed pressure was varied between 4 and 30 bar; the permeate pressure was less than 0.01 bar.

The permeate flow rate to feed flow rate, i.e. the stage-cut, was set at 0.01. Applying these conditions, the residue composition was essentially equal to that of the feed gas and concentration polarization effects could be neglected. CO2 and CH4 permeate concentrations were detected with a gas chromatograph (Agilent 3000A Micro GC) equipped with a thermal conductivity detector. The mixed-gas permeabilities were calculated by equations 2.16 and 2.17. Because the downstream pressure was negligible, the mixed-gas CO2/CH4 selectivity was obtained from equation 2.18.

8.4. Results and Discussion

8.4.1. Thermal Rearrangement of TDA1-APAF Polyimide to Polybenzoxazole

The thermal behavior of TDA1-APAF was investigated by TGA-MS to set up optimum thermal conditions for conversion of the polyimide to the corresponding TR- derived PBO. As shown in Fig. 8.2, a suitable TR range was identified approximately between 400 and 500 °C. It is interesting to note the TGA weight loss curve of the

TDA1-APAF did not show the typical two-step decomposition profile commonly observed for hydroxyl-functionalized polyimides, where the first step occurs due to 252

evolution of CO2 during the formation of PBO and the second step results from main- chain degradation. The evolution of CO2 during the heating cycle was clearly identified by mass spectroscopy and followed the general trend of the weight-loss derivative curve

(Fig. 8.2).

Fig. 8.2. TGA/MS analysis of TDA1-APAF polyimide film.

Additional TGA experiments were performed on three samples with different holding temperatures of 400, 430 and 460 °C, respectively. Initially, the PIM-PI film samples were heated at a rate of 5 °C/min and then kept isothermally for one hour at each temperature while continuously tracing the weight reduction, as shown in Fig. 8.3. The weight loss of samples made at 400, 430 and 460 °C for 1 hour were 3.0, 5.0 and 13.3 wt.% respectively. Based on these results it appeared that the heat treatments at 400 and 253

430 °C over a duration of one hour were insufficient to obtain high conversion of the

TDA1-APAF to the PBO. Alternatively, a one-hour treatment at 460 °C resulted in formation of a partially degraded PBO membrane as indicated by the higher degree of conversion (13.3 wt.%) compared to the theoretical maximum PBO conversion value of

11.7 wt.%. Therefore, heat treatment at 460 °C with shorter holding times of 15 (~ 69% conversion) and 30 minutes (~ 91% conversion), respectively, was chosen for formation of TDA1-APAF-derived TR membranes.

Fig. 8.3. Time-dependent TGA analysis of TDA1-APAF polyimide films held at isothermal temperatures of 400, 430 and 460 °C, respectively.

TR conversion, as presented in Table 8.1, is defined as the ratio of the actual mass loss after cooling from the treatment temperature to ambient temperature upon thermal rearrangement to the theoretical mass loss [25]:

254

퐴푐푡푢푎푙 푀푎푠푠 퐿표푠푠 % 푇푅 퐶표푛푣푒푟푠𝑖표푛 = 푥 100 푇ℎ푒표푟푒푡𝑖푐푎푙 푀푎푠푠 퐿표푠푠

This calculation was based on the assumption that any polyimide segment that underwent thermal rearrangement was converted to its corresponding PBO segment. The implications and limitations of this methodology have been discussed by Sanders et al.

[38]. The total weight loss of the two TDA1-APAF membrane samples after thermal treatment in the Carbolite tubular furnace at 460 °C for 15 and 30 minutes, and subsequent cooling to room temperature, was 11.1 and 13.5 wt.%, respectively. Clearly, the 15-minute treatment with cooling cycle to room temperature gave the optimum TR conversion of 95% (referred to as TR 460), whereas the 30-minute treatment resulted in some main-chain degradation as indicated by an additional weight loss of 1.8% relative to the theoretical value of 11.7 wt.% for 100% PBO conversion (referred to as TRC 460).

Full characterization of the TRC 460 membrane was outside the main scope of this work; therefore, only pure-gas permeation data are provided below to show the effect of partial carbonization on the gas transport properties of the PBO membranes.

Table 8.1. Thermal rearrangement conditions and TR conversions for TDA1-APAF film samples in this study.

Time at Rearrangement TR conversion Sample rearrangement temperature (°C) (%) temperature (min)

TR 460 460 15 95 255

TRC 460 460 30 >100*

* Additional weight loss of 1.8% due to main-chain degradation.

8.4.2. Characterization of Pristine TDA1-APAF and TR 460 Membranes

Our previous study on the pristine TDA1-APAF polyimide synthesized from 9,10- dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) [37] and a commercially available 2,2-bis(3-amino-4-hydroxyphenyl)-hexafluoropropane (APAF) diamine monomer offered a PIM-PI with moderately high BET surface area of 260 m2/g. The

TDA1-APAF membrane had O2 and CO2 permeabilities of 8.5 and 40 Barrer coupled with O2/N2 and CO2/CH4 selectivites of 5.7 and 55, respectively. In this study, TDA1-

APAF was used as precursor polymer for the formation of a PIM-triptycene-based PBO.

The FTIR spectra for the pristine TDA1-APAF and TR 460, shown in Fig. 8.4, confirmed the conversion of the polyimide to the PBO as the TR 460 film showed essentially no peaks around 3400 cm-1, which indicated the absence of hydroxyl groups after thermal treatment. Moreover, the TDA1-APAF precursor polymer had characteristic imide peaks at approximately 1779 and 1710 cm-1 which were significantly reduced for the thermally rearranged sample while two characteristic PBO peaks at ~1450 and 1103 cm-1 emerged, as noted in previous studies [23, 40-42]. 256

Fig. 8.4. FTIR spectra for the pristine TDA1-APAF and thermally rearranged TR 460 film.

The changes in the transformation of the pristine TDA1-APAF polyimide to its TR

460 derivative can be assessed from N2 (-196 °C) physisorption isotherms (Fig. 8.5a).

The BET surface area of the TDA1-APAF increased significantly from 260 to 680 m2 g-1 by thermal rearrangement to the TR 460. The corresponding NLDFT-derived pore size distributions (PSDs) are shown in Fig. 8.5b. The PSDs based on N2 isotherms showed that the quantity of micropores (>7 Å) increased significantly in TR 460, which correlates qualitatively well with significantly higher gas permeability as compared to the pristine

TDA1-APAF, as discussed below. Interestingly, TR 460 also contained a larger fraction of ultramicropores < 7 Å compared to TDA1-APAF, which is expected to have a positive effect on gas pair selectivity.

257

Fig. 8.5. a) Physisorption isotherms for TDA1-APAF and TR 460 using N2 at -196 °C.

Closed symbols: adsorption; open symbols: desorption. b) NLDFT-based estimated pore size distribution obtained from N2 isotherms for TDA1-APAF and TR 460 assuming carbon slit-pore geometry. 258

8.4.3. Pure-Gas Permeation Properties of Fresh and Physically Aged TR 460

Membranes

Pure-gas permeation experiments at 2 bar and 35 °C were performed on two freshly prepared thermally treated film samples of TR 460 (15 min) and TRC 460 (30 min), as shown in Table 8.2. As expected, the fresh thermally treated samples showed a significant increase in gas permeabilities and a decrease in pure-gas selectivities for all gas pairs relative to the pristine TDA1-APAF polyimide sample. For example, the CO2 permeability of TDA1-APAF of 40 Barrer increased 33-fold for the 15-minute heated TR

460 sample to 1328 Barrer coupled with a large decrease in CO2/CH4 selectivity from 55 for the pristine polyimide to 27 for the TR 460 sample. This result was caused by an increase in BET surface area (Fig. 8.5a) and shift to a significantly larger fraction of pores > 7 Å in the TR 460 sample, as indicated by the PSD shown in Fig. 8.5b. It is noteworthy that the TR 460 membrane showed only a small decrease in O2/N2 selectivity to 5.4 from 5.7 for the pristine polyimide, while the O2 permeability increased 37-fold from 8.5 to 311 Barrer. The partially degraded TRC 460 film exhibited significantly higher gas permeabilities than the pure PBO TR 460 film with only slightly lower gas pair selectivities.

Typically, glassy polymers undergo a physical aging process in which the chains slowly turn into a more tightly packed arrangement towards their equilibrium glassy state

[43, 44]. Thus, pure-gas permeation experiments were also performed on an aged TR 460 sample after 150 days at 2 bar and 35 °C (Table 8.2). The permeabilities for all gases decreased for the aged TR 460 membrane; for example, the O2 permeability decreased from 311 (fresh TR sample) to 185 Barrer (aged TR sample) but O2/N2 selectivity 259

increased significantly to 6.3 from 5.4 for the fresh TR sample. Furthermore, the aged TR

membrane exhibited an unexpected increase in O2/N2 selectivity relative to that of the

pristine polyimide precursor.

Table 8.2. Pure-gas permeabilities and ideal selectivities for pristine TDA1-APAF and

thermally rearranged TR 460 and TRC 460 PBOs compared with previously reported

APAF-based PBOs.

Pure-gas permeability (Barrer) Ideal selectivity (휶)

Polymer He H2 N2 O2 CH4 CO2 O2/N2 H2/N2 CO2/CH4

TDA1-APAF [21] 92 94 1.5 8.5 0.73 40 5.7 63 55

TR 460a 703 1547 58 311 49.4 1328 5.4 27 27

Aged TR 460b (664) (1304) (29.4) (185) (20.1) (699) (6.3) (44) (35)

TRC 460a 1013 2264 109 511 99.5 2386 4.7 21 24

TR 6FDA-APAF [27] 356 408 19 81 12 398 4.3 22 34

TR 6FDA-APAF [28] 269 294 12.6 52.5 7.5 261 4.2 23 35

TR BPDA-bisAPAF [24] - 444 20 93 15 597 4.7 22 40 260

TR PMDA-bisAPAF [30] - 635 34 148 23 952 4.4 19 41

TR BPDA-bisAPAF [31] - 1228 48 220 41 1014 4.6 26 25

TPHI-TR-400 [32] - 520 16 - 8.3 320 - 32 39

a This study. Tested after 1 day (at 2 bar; 35 °C).

b This study. Tested after 150 days (at 2 bar; 35 °C).

To shed some light on the differences in chain packing of the fresh and aged TR 460

film samples, wide-angle x-ray diffraction (WAXD) measurements were conducted, as

shown in Fig. 8.6. Both samples showed broad amorphous peaks at low scattering angles

2 from ~10 to 20°; however, a clear shift of the spectra to higher scattering angles 2

(lower average d-spacing) was observed for the aged TR 460 sample. In addition, the

intensity of the shoulder peak around 22° (d-spacing of ~ 4 Å) was significantly enhanced

in the aged TR 460 sample. Both trends are qualitatively consistent with densification of

the aged sample and the corresponding loss in permeability and increase in selectivity

compared to the fresh TR 460 sample. 261

Fig. 8.6. Wide-angle x-ray diffraction spectra of fresh and aged TR 460 film samples.

The pure-gas permeation properties of the TR 460 and TRC 460 membranes are also compared to data of previously reported APAF-polyimide-derived PBOs in Table 8.2.

Clearly, the performance of the newly developed triptycene-based TR 460 membrane was superior to other APAF-derived TR membranes for O2/N2 separation, as shown in Fig.

8.7. In fact, its performance exceeded the 2008 upper bound [10] and was close to the more recently reported 2015 O2/N2 upper bound [45]. 262

Fig. 8.7. Pure-gas O2/N2 separation performance of previously reported APAF- polyimide-based PBOs [29, 30] and TDA1-APAF-derived TR 460 membrane (this study). The solid lines represent 2008 [10] and 2015 [45] permeability/selectivity trade- off curves.

The pure-gas diffusivity (D) values of the pristine TDA1-APAF and TR 460 films for

N2, O2, CH4 and CO2 were calculated by the time-lag method and are presented in Table

8.3 together with the solubility (S) coefficients. The increased gas permeabilities of the fresh TR 460 sample resulted from significantly enhanced diffusion coefficients compared to the values of the fresh TDA1-APAF polyimide, whereas much smaller increases were observed in their gas solubilities. The reduction in CO2/CH4 permselectivity from 55 for the pristine TDA1-APAF polyimide to 27 for the TR 460

(Table 8.2) occurred due to a decrease in the CO2/CH4 diffusion selectivity from 15.8 to

6.9, as shown in Table 8.4. As expected, the tighter aged TR 460 membrane displayed lower D and S values compared to the fresh sample (Table 8.3). Concurrently, the aged 263

TR 460 sample yielded enhanced size-sieving properties and, therefore, higher O2/N2,

N2/CH4 and CO2/CH4 diffusivity selectivity values than those of the fresh sample (Table

8.4).

Table 8.3. Pure-gas diffusion and solubility coefficients of N2, O2, CH4 and CO2 for pristine and thermally rearranged (fresh and aged) TDA1-APAF membranes.

Diffusion coefficient Solubility coefficient

(10-8 cm2/s) (10-2 cm3(STP)/(cm3 cmHg))

Polymer N2 O2 CH4 CO2 N2 O2 CH4 CO2

TDA1-APAFa 0.82 4.5 0.12 1.9 1.8 1.9 6.1 21.1

Fresh TR 460b 17.3 88.2 7.21 49.6 3.35 3.53 6.85 26.8

Aged TR 460c (10.15) (56.7) (3.94) (35.7) (2.9) (3.26) (5.1) (19.5)

a Tested after 1 day [21]. b This study. TR at 460 °C (15min). Tested after 1 day. c This study. TR at 460 °C (15min). Tested after 150 days.

264

Table 8.4. Diffusivity selectivities and solubility selectivities of pristine and thermally

rearranged (fresh and aged) TDA1-APAF films (2 bar; 35 °C).

Diffusivity selectivity (α)D Solubility selectivity (α)S

Polymer O2/N2 N2/CH4 CO2/CH4 O2/N2 N2/CH4 CO2/CH4

TDA1-APAFa 5.5 6.83 15.8 1.06 0.30 3.5

Fresh TR TDA1-APAFb 5.1 2.40 6.9 1.05 0.49 3.91

Aged TR TDA1-APAFc 5.59 2.58 9.1 1.12 0.57 3.82

a tested after 1 day [21].

b this study. TR at 460 °C (15min). Tested after 1 day.

c this study. TR at 460 °C (15min). Tested after 150 days.

8.4.4. Mixed-Gas Permeation Properties of Thermally Rearranged TDA1-APAF

(TR 460)

Commonly, high sorption uptake of condensable gases like CO2 can dilate glassy

polymers, induce increased chain mobility, modify the pore structure and, therefore,

significantly affect their gas transport properties. The pure- and mixed-gas permeabilities

(Figs. 8.8 and 8.9) and selectivities (Fig. 8.10) were measured using a 1:1 CO2/CH4 feed

mixture with increasing CO2 partial pressure (2, 5, 7, 10, 12 and 15 bar) for pristine and 265

thermally rearranged TDA1-APAF samples (TR 460). The pure-gas CO2 permeabilities in both pristine and the TR membrane decreased with increasing CO2 partial pressure up to 15 bar (Fig. 8.8), similar to the behavior of previously reported related PIM-PIs [46-

49]. This can be rationalized by a decrease in CO2 solubility coefficient with pressure as is typically observed in glassy polymers due to their dual-mode sorption behavior. In addition, under mixed-gas permeation conditions competitive sorption occurs between

CO2 and CH4 for available sorption sites, which may reduce the overall sorbed concentration of each gas in the polymer. As a result, the mixed-gas CO2 permeability was lower than the pure-gas CO2 permeability value determined at the same partial pressure.

Fig. 8.8. Pressure-dependence of pure- and mixed-gas CO2 permeabilities of pristine

TDA1-APAF [21] and TR 460 (1:1 CO2/CH4 mixture; 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas. 266

The mixed-gas CH4 permeability was also lower than the pure-gas value up to a partial pressure of ~10 bar, as shown in Fig. 8.9. However, the mixed-gas CH4 permeability was continuously increasing over the entire range pressure and surpassed the pure-gas value possibly due to some plasticization effects. The mixed-gas permselectivites of CO2/CH4 at partial CO2 pressures up to ~ 8 bar were slightly higher or equal to those determined under pure-gas conditions (Fig. 8.10) due to a reduction of the mixed-gas CH4 permeability by co-permeation of CO2. This mixed-gas permeation

“blocking” effect of the less permeable feed gas component by co-permeation of CO2 was previously observed for some ladder PIMs [50], PIM-PIs [21, 46] and PBOs obtained by thermal rearrangement of polyimides [34, 51].

Fig. 8.9. Pressure-dependence of pure- and mixed-gas CH4 permeabilities for pristine

[21] and TR 460 (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye. Open points: pure-gas; closed points: mixed-gas. 267

Fig. 8.10. Pressure-dependence of mixed-gas CO2/CH4 selectivities of pristine [21] and

TR 460 samples (1:1 CO2/CH4 mixture, 35 °C). Lines are drawn to guide the eye: open points, pure-gas; closed points, mixed-gas.

8.5. Conclusions

In this work, a new TR membrane was prepared from a triptycene-based TDA1-

APAF polyimide derived from 9,10-dimethyl-2,3,6,7-triptycene tetracarboxylic dianhydride (TDA1) and a commercial dihydroxyl-diamine (APAF) via one-step high- temperature solution imidization reaction. Pure-gas permeation data demonstrated that a

TR membrane heated at 460 °C for 15 minutes (TR 460) showed significant increase in permeability due to large increases in diffusion coefficients for a variety of gases. A freshly tested TR 460 membrane showed excellent O2 permeability of 311 Barrer coupled with O2/N2 selectivity of 5.4 and CO2 permeability of 1328 Barrer with CO2/CH4 selectivity of 27. Physically aged TR 460 membrane samples tested over a period of 150 268

days showed a decrease in gas permeability (푃푂2= 185 Barrer) but with an increased

O2/N2 selectivity of 6.3, which surpassed the 2008 upper bound and was close to the performance of other high-performance state-of-the-art PIM and PIM-PI membrane materials recently reported for the 2015 O2/N2 upper bound. Therefore, the PIM-PI- triptycene-based TR 460 membrane could find potential use in nitrogen production from air for a variety of demanding applications in the aircraft- and oil- and natural gas industries. The utilization of the TR membrane for CO2 removal from natural gas is also commendable due to its high mixed-gas CO2 permeability of ~1,000 Barrer combined with CO2/CH4 selectivity of 22 at a typical CO2 wellhead partial pressure of 10 bar.

269

8.6. References

[1] P. Bernardo, E. Drioli, G. Golemme, Membrane gas separation: a review/state of the art, Industrial & Engineering Chemistry Research, 48 (2009) 4638-4663.

[2] R.W. Baker, K. Lokhandwala, Natural gas processing with membranes: an overview,

Industrial & Engineering Chemistry Research, 47 (2008) 2109-2121.

[3] R.W. Baker, Future directions of membrane gas separation technology, Industrial &

Engineering Chemistry Research, 41 (2002) 1393-1411.

[4] Y. Huang, D. Paul, Physical aging of thin glassy polymer films monitored by gas permeability, Polymer, 45 (2004) 8377-8393.

[5] L.M. Robeson, Correlation of separation factor versus permeability for polymeric membranes, Journal of Membrane Science, 62 (1991) 165-185.

[6] P.M. Budd, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib, C.E.

Tattershall, Polymers of intrinsic microporosity (PIMs): robust, solution-processable, organic nanoporous materials, Chemical Communications, (2004) 230-231.

[7] P.M. Budd, E.S. Elabas, B.S. Ghanem, S. Makhseed, N.B. McKeown, K.J. Msayib,

C.E. Tattershall, D. Wang, Solution-processed, organophilic membrane derived from a polymer of intrinsic microporosity, Advanced Materials, 16 (2004) 456-459.

[8] P.M. Budd, K.J. Msayib, C.E. Tattershall, B.S. Ghanem, K.J. Reynolds, N.B.

McKeown, D. Fritsch, Gas separation membranes from polymers of intrinsic microporosity, Journal of Membrane Science, 251 (2005) 263-269. 270

[9] P.M. Budd, N.B. McKeown, D. Fritsch, Free volume and intrinsic microporosity in polymers, Journal of Materials Chemistry, 15 (2005) 1977-1986.

[10] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320

(2008) 390-400.

[11] B.S. Ghanem, N.B. McKeown, P.M. Budd, J.D. Selbie, D. Fritsch, High- performance membranes from polyimides with intrinsic microporosity, Advanced

Materials, 20 (2008) 2766-2771.

[12] B.S. Ghanem, N.B. McKeown, P.M. Budd, N.M. Al-Harbi, D. Fritsch, K. Heinrich,

L. Starannikova, A. Tokarev, Y. Yampolskii, Synthesis, characterization, and gas permeation properties of a novel group of polymers with intrinsic microporosity: PIM- polyimides, Macromolecules, 42 (2009) 7881-7888.

[13] J. Weber, O. Su, M. Antonietti, A. Thomas, Exploring polymers of intrinsic microporosity-microporous, soluble polyamide and polyimide, Macromolecular Rapid

Communications, 28 (2007) 1871-1876.

[14] X.H. Ma, R. Swaidan, Y. Belmabkhout, Y.H. Zhu, E. Litwiller, M. Jouiad, I. Pinnau,

Y. Han, Synthesis and gas transport properties of hydroxyl-functionalized polyimides with intrinsic microporosity, Macromolecules, 45 (2012) 3841-3849.

[15] Y. Rogan, L. Starannikova, V. Ryzhikh, Y. Yampolskii, P. Bernardo, F. Bazzarelli,

J.C. Jansen, N.B. McKeown, Synthesis and gas permeation properties of novel spirobisindane-based polyimides of intrinsic microporosity, Polymer Chemistry, 4 (2013)

3813-3820. 271

[16] B.S. Ghanem, R. Swaidan, E. Litwiller, I. Pinnau, Ultra-microporous triptycene- based polyimide membranes for high-performance gas separation, Advanced Materials,

26 (2014) 3688-3692.

[17] N. Ritter, I. Senkovska, S. Kaskel, J. Weber, Intrinsically microporous poly(imide)s: structure-porosity relationship studied by gas sorption and X-ray scattering,

Macromolecules, 44 (2011) 2025-2033.

[18] Y. Rogan, R. Malpass-Evans, M. Carta, M. Lee, J.C. Jansen, P. Bernardo, G.

Clarizia, E. Tocci, K. Friess, M. Lanč, A highly permeable polyimide with enhanced selectivity for membrane gas separations, Journal of Materials Chemistry A, 2 (2014)

4874-4877.

[19] X. Ma, O. Salinas, E. Litwiller, I. Pinnau, Novel spirobifluorene- and dibromospirobifluorene-based polyimides of intrinsic microporosity for gas separation applications, Macromolecules, 46 (2013) 9618-9624.

[20] Y. Zhuang, J.G. Seong, Y.S. Do, H.J. Jo, Z. Cui, J. Lee, Y.M. Lee, M.D. Guiver,

Intrinsically microporous soluble polyimides incorporating Tröger’s base for membrane gas separation, Macromolecules, 47 (2014) 3254-3262.

[21] F. Alghunaimi, B. Ghanem, N. Alaslai, M. Mukaddam, I. Pinnau, Triptycene dimethyl-bridgehead dianhydride-based intrinsically microporous hydroxyl- functionalized polyimide for natural gas upgrading, Journal of Membrane Science, 520

(2016) 240-246. 272

[22] D.F. Sanders, Z.P. Smith, C.P. Ribeiro, R. Guo, J.E. McGrath, D.R. Paul, B.D.

Freeman, Gas permeability, diffusivity, and free volume of thermally rearranged polymers based on 3, 3′-dihydroxy-4, 4′-diamino-biphenyl (HAB) and 2, 2′-bis-(3, 4- dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane

Science, 409 (2012) 232-241.

[23] H.B. Park, C.H. Jung, Y.M. Lee, A.J. Hill, S.J. Pas, S.T. Mudie, E. Van Wagner,

B.D. Freeman, D.J. Cookson, Polymers with cavities tuned for fast selective transport of small molecules and ions, Science, 318 (2007) 254-258.

[24] H.B. Park, S.H. Han, C.H. Jung, Y.M. Lee, A.J. Hill, Thermally rearranged (TR) polymer membranes for CO2 separation, Journal of Membrane Science, 359 (2010) 11-

24.

[25] Q. Liu, D.R. Paul, B.D. Freeman, Gas permeation and mechanical properties of thermally rearranged (TR) copolyimides, Polymer, 82 (2016) 378-391.

[26] G.L. Tullos, J.M. Powers, S.J. Jeskey, L.J. Mathias, Thermal conversion of hydroxy- containing imides to benzoxazoles: polymer and model compound study,

Macromolecules, 32 (1999) 3598-3612.

[27] S.H. Han, N. Misdan, S. Kim, C.M. Doherty, A.J. Hill, Y.M. Lee, Thermally rearranged (TR) polybenzoxazole: effects of diverse imidization routes on physical properties and gas transport behaviors, Macromolecules, 43 (2010) 7657-7667. 273

[28] M. Calle, C.M. Doherty, A.J. Hill, Y.M. Lee, Cross-linked thermally rearranged poly (benzoxazole-co-imide) membranes for gas separation, Macromolecules, 46 (2013)

8179-8189.

[29] W. Liu, W. Xie, Acetate-functional thermally rearranged polyimides based on 2, 2- bis (3-amino-4-hydroxyphenyl) hexafluoropropane and various dianhydrides for gas separations, Industrial & Engineering Chemistry Research, 53 (2013) 871-879.

[30] S. Kim, Y.M. Lee, Rigid and microporous polymers for gas separation membranes,

Progress in Polymer Science, 43 (2015) 1-32.

[31] C.H. Jung, J.E. Lee, S.H. Han, H.B. Park, Y.M. Lee, Highly permeable and selective poly (benzoxazole-co-imide) membranes for gas separation, Journal of Membrane

Science, 350 (2010) 301-309.

[32] S. Luo, J. Liu, H. Lin, B.A. Kazanowska, M.D. Hunckler, R.K. Roeder, R. Guo,

Preparation and gas transport properties of triptycene-containing polybenzoxazole

(PBO)-based polymers derived from thermal rearrangement (TR) and thermal cyclodehydration (TC) processes, Journal of Materials Chemistry A, 4 (2016) 17050-

17062.

[33] S. Li, H.J. Jo, S.H. Han, C.H. Park, S. Kim, P.M. Budd, Y.M. Lee, Mechanically robust thermally rearranged (TR) polymer membranes with spirobisindane for gas separation, Journal of Membrane Science, 434 (2013) 137-147.

[34] R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, High pressure pure-and mixed-gas separation of CO2/CH4 by thermally-rearranged and carbon molecular sieve membranes 274 derived from a polyimide of intrinsic microporosity, Journal of Membrane Science, 447

(2013) 387-394.

[35] H. Shamsipur, B.A. Dawood, P.M. Budd, P. Bernardo, G. Clarizia, J.C. Jansen,

Thermally rearrangeable PIM-polyimides for gas separation membranes,

Macromolecules, 47 (2014) 5595-5606.

[36] X. Ma, O. Salinas, E. Litwiller, I. Pinnau, Pristine and thermally-rearranged gas separation membranes from novel o-hydroxyl-functionalized spirobifluorene-based polyimides, Polymer Chemistry, 5 (2014) 6914-6922.

[37] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232.

[38] D.F. Sanders, R. Guo, Z.P. Smith, Q. Liu, K.A. Stevens, J.E. McGrath, D.R. Paul,

B.D. Freeman, Influence of polyimide precursor synthesis route and ortho-position functional group on thermally rearranged (TR) polymer properties: conversion and free volume, Polymer, 55 (2014) 1636-1647.

[39] K. O'Brien, W. Koros, T. Barbari, E. Sanders, A new technique for the measurement of multicomponent gas transport through polymeric films, Journal of Membrane Science,

29 (1986) 229-238.

[40] X. Ma, R. Swaidan, B. Teng, H. Tan, O. Salinas, E. Litwiller, Y. Han, I. Pinnau,

Carbon molecular sieve gas separation membranes based on an intrinsically microporous polyimide precursor, Carbon, 62 (2013) 88-96. 275

[41] R. Guo, D.F. Sanders, Z.P. Smith, B.D. Freeman, D.R. Paul, J.E. McGrath,

Synthesis and characterization of thermally rearranged (TR) polymers: effect of glass transition temperature of aromatic poly (hydroxyimide) precursors on TR process and gas permeation properties, Journal of Materials Chemistry A, 1 (2013) 6063-6072.

[42] M. Calle, A.E. Lozano, Y.M. Lee, Formation of thermally rearranged (TR) polybenzoxazoles: effect of synthesis routes and polymer form, European Polymer

Journal, 48 (2012) 1313-1322.

[43] K.D. Dorkenoo, P.H. Pfromm, Accelerated physical aging of thin poly[1-

(trimethylsilyl)-1-propyne] films, Macromolecules, 33 (2000) 3747-3751.

[44] P.H. Pfromm, The impact of physical aging of amorphous glassy polymers on gas separation membranes, in Materials Science of Membranes for Gas and Vapor

Separation, Y. Yampolskii, I. Pinnau, B.D. Freeman (Eds.), John Wiley & Sons, (2006)

293-306.

[45] R. Swaidan, B. Ghanem, I. Pinnau, Fine-tuned intrinsically ultramicroporous polymers redefine the permeability/selectivity upper bounds of membrane-based air and hydrogen separations, ACS Macro Letters, 4 (2015) 947-951.

[46] R. Swaidan, B. Ghanem, E. Litwiller, I. Pinnau, Effects of hydroxyl- functionalization and sub-Tg thermal annealing on high pressure pure-and mixed-gas

CO2/CH4 separation by polyimide membranes based on 6FDA and triptycene-containing dianhydrides, Journal of Membrane Science, 475 (2015) 571-581. 276

[47] R. Swaidan, B. Ghanem, M. Al-Saeedi, E. Litwiller, I. Pinnau, Role of intrachain rigidity in the plasticization of intrinsically microporous triptycene-based polyimide membranes in mixed-gas CO2/CH4 separations, Macromolecules, 47 (2014) 7453-7462.

[48] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[49] N. Alaslai, B. Ghanem, F. Alghunaimi, E. Litwiller, I. Pinnau, Pure-and mixed-gas permeation properties of highly selective and plasticization resistant hydroxyl-diamine- based 6FDA polyimides for CO2/CH4 separation, Journal of Membrane Science, (2016).

[50] N. Du, H.B. Park, G.P. Robertson, M.M. Dal-Cin, T. Visser, L. Scoles, M.D. Guiver,

Polymer nanosieve membranes for CO2-capture applications, Nature Materials, 10 (2011)

372-375.

[51] K.L. Gleason, Z.P. Smith, Q. Liu, D.R. Paul, B.D. Freeman, Pure-and mixed-gas permeation of CO2 and CH4 in thermally rearranged polymers based on 3, 3′-dihydroxy-

4, 4′-diamino-biphenyl (HAB) and 2, 2′-bis-(3, 4-dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane Science, 475 (2015) 204-214.

277

Chapter 9. Conclusions and Recommendations

9.1. Introduction

The objective of this research was to design high-performance polymeric membrane materials for CO2/CH4 separation. Chapters 4 through 7 highlighted design principles that established the potential of PIM-PIs materials to be tuned for highly permeable and highly selective gas transport relative to the standard commercial membrane materials such as cellulose acetate. Sorption isotherms (from N2 and CO2), XRD peaks, TGA analysis, pure-and mixed-gas experiments were utilized to rationalize the different structure/property relationships observed for separations involving streams free of condensable gases (He, H2, N2 and O2) and those involving condensable gases (CH4 and

CO2). This chapter provides a brief overview of the outcomes discovered in this research in a joining manner. In addition, some recommendations are presented for further materials development.

9.2. Dimethyl-Triptycene Building Block

Five triptycene-based building blocks (Fig. 9.1) were utilized in the research

(incorporated in polyimide membranes) to evaluate and determine the optimum moiety that could be used in designing membrane materials for CO2/CH4 separation. Based on the collected data (Chapters 4 to 6) from the gas transport measurements, BET surface area and WAXD analysis, the dimethyl-triptycene-based building block (Fig. 9.1 c) exhibited the best CO2/CH4 separation performance among others with the non- substituted, extended iptycene, diethyl- or diisopropyl-triptycene-based polyimides. This 278 building block enhanced the BET surface area, provided greater fraction of ultramicroporosity and simultaneously improved the gas permeability.

Fig. 9.1. Triptycene building units of (a) non-substituted-, (b) extended iptycene, (c) dimethyl-substituted, (d) diethyl-substituted and (e) diisopropyl-substituted.

9.3. Functionalized PIM-Polyimide

TDA1-DMN polyimide which was derived from new 9,10-dimethyl-2,3,6,7- triptycene tetracarboxylic dianhydride (TDA1) monomer and 3,3ʹ-dimethylnaphthidine

(DMN) showed that the triptycene building block with methyl bridgehead offered a polyimide with high BET surface area of 760 m2/g and very high gas permeabilities (e.g.

CO2 = 3700 Barrer) but with only moderate selectivities (e.g. CO2/CH4 selectivity = 17) as shown in Chapter 5. Previous work demonstrated that 6FDA-derived PI membranes made from hydroxyl-containing diamines resulted in polyimides, such as 6FDA-APAF 279

with the highest CO2/CH4 selectivities reported to date; however, the CO2 permeability of these hydroxyl-functionalized polyimides was relatively low (< 10 Barrer) [1]. Therefore, in this dissertation a new ortho-hydroxyl-functionalized polyimide (TDA1-APAF) was designed with high selectivity (CO2/CH4 = 55) and moderate gas permeability (CO2 = 40

Barrer) due to the presence of a triptycene building block with dimethyl substitution

(Chapter 7). This new polyimide was synthesized from TDA1 dianhydride and a commercial hydroxyl-containing diamine monomer (APAF).

The pure-gas permeation data showed that introducing hydroxyl groups to the polyimide led to a significant increase in permselectivity due to a large increase in diffusivity selectivity for a variety of gas pairs, most notably H2/CH4 and CO2/CH4 (Fig.

9.2). The microporous polymer TDA1-APAF had a BET surface area based on nitrogen adsorption of 260 m2/g. Physical aging over 250 days resulted in significantly enhanced

CO2/CH4 selectivity to 75, respectively with only ~ 25% loss in CO2 and H2 permeability.

Aged TDA1-APAF exhibited 5-fold higher pure-gas CO2 permeability (30 Barrer) and 2- fold higher CO2/CH4 permselectivity over conventional dense cellulose triacetate membranes at 2 bar. Furthermore, the triptycene-based hydroxyl-containing polyimide showed a CO2 permeability of 21 under binary 1:1 CO2/CH4 mixed-gas feed with a selectivity of 72 at a partial CO2 pressure of 10 bar. These properties are significantly

better those of cellulose triacetate, which exhibits 푃퐶푂2= 8 Barrer and CO2/CH4 selectivity of 25 when tested under the same conditions. 280

Fig. 9.2. Gas separation performance of TDA1-DMN (this study), 6FDA-APAF [1] and

TDA1-APAF (this study) polyimides for CO2/CH4. The solid line represents the 2008

CO2/CH4 permeability/selectivity trade-off curve [2].

9.4. Thermal Rearrangement of PIM-Polyimide

An alternative method to introduce microporosity into glassy polymers is based on the formation of polybenzoxazoles (PBO) by a high-temperature (> 400 °C) decarboxylation and rearrangement reaction of polyimide precursors bearing functional hydroxyl groups in ortho-position to the imide linkages [3]. The hydroxyl-functionalized triptycene-based polyimide of intrinsic microporosity (TDA1-APAF) was converted to a polybenzoxazole (PBO) (Chapter 8) by heating treatment at 460 °C under nitrogen atmosphere. TDA1-APAF polyimide treated for 15 minutes (TR 460) resulted in a PBO conversion of 95% based on a theoretical weight loss of 11.7 wt.% of the polyimide precursor. Heating TDA1-APAF for 30 minutes (TRC 460) resulted in a weight loss of

13.5 wt.%, indicating full conversion to PBO and partial main chain degradation. The TR 281

460 membrane displayed excellent CO2 permeability of 1328 Barrer with a CO2/CH4 selectivity of 27 (Fig. 9.3). Interestingly, physical aging over 150 days resulted in an enhanced CO2/CH4 selectivity of 35 coupled with CO2 permeability of 699 Barrer. These results suggest that thermally-rearranged membranes from hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for removal of

CO2 on natural gas sweetening and air separation applications where the physically aged sample resulted in an O2/N2 selectivity of 6.3 with an O2 permeability of 185 Barrer.

Fig. 9.3. Gas separation performance of CTA [4], pristine and TR TDA1-APAF (TR 460) membranes (this study) for CO2/CH4. The solid line represents the 2008 CO2/CH4 permeability/selectivity trade-off curve [2].

9.5. Natural Gas Sweetening

In Chapter 1, the targeted material for membrane-based natural gas sweetening was identified as a CO2/CH4 selectivity of 40 under mixed-gas feeds (where selectivity 282

minimizes highly costly CH4 losses). For the PIM-PIs in this work, selectivities varied drastically between ~15 for TDAi3-DMN (Chapter 5) and ~75 for the aged TDA1-

APAF (Chapter 7). In addition, the aged TDA1-APAF showed 5-fold higher permeabilities over conventional polyimides (e.g., 6FDA-APAF) [1] and benchmark commercial materials (e.g., CA and Matrimid®). Furthermore, this polyimide had a

N2/CH4 selectivity of 2.3, thereby making it possible to simultaneously treat CO2- and nitrogen-contaminated natural gas. These results suggest that intrinsically microporous hydroxyl-functionalized triptycene-based polyimides are promising candidate membrane materials for removal of CO2 and N2 from natural gas.

9.6. CO2/CH4 Mixed-Gas Performance and Proposed 2017 Upper Bound

Chapters 4, 7 and 8 include 1:1 CO2/CH4 mixed-gas experiments for many novel polyimides and thermally rearranged membranes. Those results were compared to the literature data tested under the same conditions as shown in Fig. 9.4. The permeability/selectivity trade-off curve in this figure was initiated by Robeson which is based on pure-gas data for CO2 and CH4. However, the mixture of condensable gases may be affected by plasticization phenomena, which can significantly reduce the separation performance. Moreover, dual-mode and competitive mixed-gas sorption effects can significantly reduce the CO2 permeability. Therefore, the actual, more desirable trade-off curve for CO2/CH4 separation should be based on mixed-gas experiments to reflect the actual performance of polymer membrane materials at high feed pressure with up to 10 bar of CO2 partial pressure to simulate the real conditions of typical natural gas wellhead pressure. 283

Fig. 9.4. Mixed-gas CO2/CH4 permeability/selectivity data for TDA1-APAF (this study),

TR TDA1-APAF (this study), KAUST-PI-1 [4], AO-PIM-1 [5], 6FDA-DAT1 (this study), 6FDA-DAT2 (this study), TPDA-APAF [1], 6FDA-DAP [6], 6FDA-APAF [1],

6FDA-mPDA [6], TR6FDA-HAB [7], TRPIM-6FDA-OH [8] and CTA [4]. All experiments were performed under mixed-gas feed with a 50/50 (v/v) CO2/CH4 mixture at 20 bar feed pressure and 35°C using the constant volume/variable pressure technique.

Based on the collected data for CO2/CH4 mixed-gas experiments from this work and literature data, a new mixed-gas upper bound curve was constructed, as shown in Fig.

9.5. These membranes are: CTA [4], Matrimid [4], 6FDA-DAT1 (this study), 6FDA-

DAT2 (this study), 6FDA-mPDA [6], 6FDA-DAP [6], 6FDA-DAR [6], 6FDA-APAF

[1], TPDA-APAF [1], TPDA-ATAF [1], PIM-1 [5], AO-PIM-1 [5], KAUST-PI-1 [4],

KAUST-PI-5 [4], 6FDA-DAM:DABA [9], TDA1-APAF (this study), TPDA-DAR [10] and TPDA-mPDA [10]. 284

Fig. 9.5. Proposed CO2/CH4 mixed-gas upper bound (solid red line) based on CTA [4],

Matrimid [4], 6FDA-DAT1 (this study), 6FDA-DAT2 (this study), 6FDA-mPDA [6],

6FDA-DAP [6], 6FDA-DAR [6], 6FDA-APAF [1], TPDA-APAF [1], TPDA-ATAF [1],

PIM-1 [5], AO-PIM-1 [5], KAUST-PI-1 [4], KAUST-PI-5 [4], 6FDA-DAM:DABA [9],

TDA1-APAF (this study), TPDA-DAR [10] and TPDA-mPDA [10]. All data were generated at the same conditions: 50/50 (v/v) CO2/CH4 feed mixture at 20 bar feed pressure and 35 °C using the constant volume/variable pressure technique. The black dash lines represent 1991 and 2008 CO2/CH4 pure-gas upper bounds [2, 11].

Parameters for the new 2017 CO2/CH4 mixed-gas upper bounds are provided in

Table 9.1. For validation, the permeabilities and selectivities of the PIMs near to or defining the 2017 upper bounds are listed in Table 9.2. 285

n Table 9.1. Overview of “Mixed-Gas Upper Bound” line parameters, where Pi = k αij

(i.e., Pi is permeability of i in Barrer, k is the front factor in Barrer, αij is the selectivity for i/j, and n is the slope), for key polymer membrane-based CO2/CH4 separations.

Mixed-gas 2017 (this

Pure-gas 1991 [11] Pure-gas 2008 [2] study)

Gas pair k (Barrer) n k (Barrer) n k (Barrer) n

CO2/CH4 1,073,700 -2.6246 5,369,140 -2.636 5,000,000 -2.9

Table 9.2. Performance of PIMs near to or defining the new 2017 mixed-gas upper bound for key polymer membrane-based CO2/CH4 separation.

Polymer CO2 permeability (Barrer) CO2/CH4 selectivity

Aged TDA1-APAF (this

study) 21 72

TPDA-DAR [10] 133 38

AO-PIM-1 [5] 780 21

KAUST-PI-1 [4] 2800 13

9.7. Recommendations for Future Work

9.7.1. Copolymers

Different structural features can be built into a polyimide by simply reacting various diamines and/or dianhydrides in a one-pot synthesis. By varying the structures of the 286 diamines and/or dianhydrides and their ratios, copolymers can be synthesized. For example, it was shown in Chapter 5 that TDA1-DMN showed very high gas permeability with moderate selectivity. The work summarized in Chapter 7 demonstrated that hydroxyl-functionalized diamines (e.g., APAF) provided significantly higher CO2/CH4 selectivity than DMN diamine. Accordingly, a copolymer between the

TDA1 dianhydride and DMN and APAF diamines could be formed to reach a balance in permeability and selectivity under mixed-gas environments. Furthermore, co-polymers typically offer significantly enhanced solubility as compared to homo-polymers, which could affect their processability into either integerally-skinned asymmetric membranes made by phase inversion or thin-film composite membranes by solution coating onto porous substrates.

9.7.2. Thin Films

All gas transport experiments in this dissertation were determined using thick films

60-90 μm in thickness. As discussed in Chapters 2 and 3, physical aging and plasticization in polymer membranes can be dependent on film thickness and are typically inflated for thinner films (e.g., 0.2 – 1 μm). To date however, few studies have addressed the thin-film permeation properties of PIM-PIs. Accordingly, it is recommended to prepare thin-film composites using the most promising polymers from this work (such as TDA1-APAF) in flat-sheet or hollow fiber geometries and examine how thickness may affect the design principles developed using thick films. This practical information would be necessary for a further assessment of the commercial potential of PIM-based membranes for natural gas sweetening application.. 287

9.7.3. Multi-Component Mixtures

Field conditions in natural gas processing are relatively different compared to those typically used for membrane materials evaluation in laboratories. This work has demonstrated the potential of TDA1-APAF (Chapter-7) as highly selective and plasticization-resistant membrane for CO2 removal under dry feed conditions. However, the presence of water in natural gas could affect the separation performance of polymer membranes. Therefore, multi-component mixtures could be used to evaluate the ability of

PIM-PIs to maintain high CO2/CH4 selectivities while extracting undesirable CO2 gas.

Also, TDA1-APAF (Chapter 7) should be evaluated under C2+ hydrocarbon-containing

CO2/CH4 mixtures to reassess the design principles under more practical industrial conditions.

9.7.4. Carbon Molecular Sieve Membranes

Carbon molecular sieve (CMS) membranes have been proposed as very promising candidate materials for gas separations [12]. CMS are highly porous materials and possess a distribution of pore sizes with ultramicroporous pore openings having similar dimensions as the molecular sizes of gas molecules. Carbon membranes are usually produced by pyrolysis of several different polymeric materials in hollow-fiber geometry, including cellulose acetate and polyimides. Generally, the pyrolysis temperatures can range from 500 to 1000 °C in inert or vacuum environments [13]. So far, the best performing carbon membranes for gas separation applications were produced by pyrolysis of aromatic polyimides [12]. The PIM-polyimides materials in this work are potentially good candidates to prepare carbon molecular sieve hollow fibers membranes. 288

9.8. Publications

The following is a list of journal publications that were accomplished during my

Ph.D. study:

1) Gas permeation and physical aging properties of iptycene diamine-based microporous

polyimides. Journal of Membrane Science 2015 (490) 321-327 [14].

2) Triptycene dimethyl-bridgehead dianhydride-based intrinsically microporous

hydroxyl-functionalized polyimide for natural gas upgrading, Journal of Membrane

Science 2016 (520) 240-246 [15].

3) Synthesis and characterization of novel triptycene dianhydrides and polyimides of

intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer 2016 (101) 225-

232 [16].

4) New phenazine-containing ladder polymer of intrinsic microporosity from a

spirobisindane-based AB-type monomer, RSC Advances 2016 (83) 79625-79630 [17].

5) Pure-and mixed-gas permeation properties of highly selective and plasticization

resistant hydroxyl-diamine-based 6FDA polyimides for CO2/CH4 separation, Journal

of Membrane Science 505 (2016): 100-107 [6].

6) High-performance intrinsically microporous dihydroxyl-functionalized triptycene-

based polyimide for natural gas separation, Polymer 91 (2016): 128-135 [10].

7) Synthesis and gas permeation properties of a novel thermally-rearranged

polybenzoxazole made from an intrinsically microporous hydroxyl-functionalized

triptycene-based polyimide precursor (Submitted to Polymer)

289

9.9. References

[1] R. Swaidan, B. Ghanem, E. Litwiller, I. Pinnau, Effects of hydroxyl-functionalization and sub-Tg thermal annealing on high pressure pure-and mixed-gas CO2/CH4 separation by polyimide membranes based on 6FDA and triptycene-containing dianhydrides,

Journal of Membrane Science, 475 (2015) 571-581.

[2] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320 (2008)

390-400.

[3] D.F. Sanders, Z.P. Smith, C.P. Ribeiro, R. Guo, J.E. McGrath, D.R. Paul, B.D.

Freeman, Gas permeability, diffusivity, and free volume of thermally rearranged polymers based on 3, 3′-dihydroxy-4, 4′-diamino-biphenyl (HAB) and 2, 2′-bis-(3, 4- dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane

Science, 409 (2012) 232-241.

[4] R. Swaidan, B. Ghanem, M. Al-Saeedi, E. Litwiller, I. Pinnau, Role of intrachain rigidity in the plasticization of intrinsically microporous triptycene-based polyimide membranes in mixed-gas CO2/CH4 separations, Macromolecules, 47 (2014) 7453-7462.

[5] R. Swaidan, B.S. Ghanem, E. Litwiller, I. Pinnau, Pure-and mixed-gas CO2/CH4 separation properties of PIM-1 and an amidoxime-functionalized PIM-1, Journal of

Membrane Science, 457 (2014) 95-102.

[6] N. Alaslai, B. Ghanem, F. Alghunaimi, E. Litwiller, I. Pinnau, Pure-and mixed-gas permeation properties of highly selective and plasticization resistant hydroxyl-diamine- 290

based 6FDA polyimides for CO2/CH4 separation, Journal of Membrane Science, 505

(2016) 100–107.

[7] K.L. Gleason, Z.P. Smith, Q. Liu, D.R. Paul, B.D. Freeman, Pure-and mixed-gas permeation of CO2 and CH4 in thermally rearranged polymers based on 3, 3′-dihydroxy-

4, 4′-diamino-biphenyl (HAB) and 2, 2′-bis-(3, 4-dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane Science, 475 (2015) 204-214.

[8] R. Swaidan, X. Ma, E. Litwiller, I. Pinnau, High pressure pure-and mixed-gas separation of CO2/CH4 by thermally-rearranged and carbon molecular sieve membranes derived from a polyimide of intrinsic microporosity, Journal of Membrane Science, 447

(2013) 387-394.

[9] J.D. Wind, D.R. Paul, W.J. Koros, Natural gas permeation in polyimide membranes,

Journal of Membrane Science, 228 (2004) 227-236.

[10] N. Alaslai, B. Ghanem, F. Alghunaimi, I. Pinnau, High-performance intrinsically microporous dihydroxyl-functionalized triptycene-based polyimide for natural gas separation, Polymer, 91 (2016) 128-135.

[11] L.M. Robeson, Correlation of separation factor versus permeability for polymeric membranes, Journal of Membrane Science, 62 (1991) 165-185.

[12] C.W. Jones, W.J. Koros, Carbon molecular sieve gas separation membranes-I. preparation and characterization based on polyimide precursors, Carbon, 32 (1994) 1419-

1425. 291

[13] D.Q. Vu, W.J. Koros, S.J. Miller, High pressure CO2/CH4 separation using carbon molecular sieve hollow fiber membranes, Industrial & Engineering Chemistry Research,

41 (2002) 367-380.

[14] F. Alghunaimi, B. Ghanem, N. Alaslai, R. Swaidan, E. Litwiller, I. Pinnau, Gas permeation and physical aging properties of iptycene diamine-based microporous polyimides, Journal of Membrane Science, 490 (2015) 321-327.

[15] F. Alghunaimi, B. Ghanem, N. Alaslai, M. Mukaddam, I. Pinnau, Triptycene dimethyl-bridgehead dianhydride-based intrinsically microporous hydroxyl- functionalized polyimide for natural gas upgrading, Journal of Membrane Science, 520

(2016) 240-246.

[16] B. Ghanem, F. Alghunaimi, X. Ma, N. Alaslai, I. Pinnau, Synthesis and characterization of novel triptycene dianhydrides and polyimides of intrinsic microporosity based on 3, 3ʹ-dimethylnaphthidine, Polymer, 101 (2016) 225-232.

[17] B. Ghanem, F. Alghunaimi, N. Alaslai, X. Ma, I. Pinnau, New phenazine-containing ladder polymer of intrinsic microporosity from a spirobisindane-based AB-type monomer, RSC Advances, 6 (2016) 79625-79630.

292

APPENDIX

Chapter 3

1) Co-permeation Phenomena in Mixed-Gas Experiment

The co-permeation in glassy polymer is the transport of two different gases

through membrane at the same time. The transport properties of one component are

affected by co-permeation of the other feed gas. For example, previous work showed

that in CO2/CH4 mixed-gas separation experiments using cellulose acetate (CA)

membranes, the CO2 permeability decreased to about 50% relative to the pure-gas

value. The mixed-gas CO2 permeability is lower than the pure-gas value due to the

competitive sorption, which reduced the overall sorbed concentration of CO2 in the

polymer across the entire pressure range.

Chapter 4

Table 1. Diffusivity selectivities and solubility selectivities for 6FDA-DAT1 and 6FD- DAT2 membranes (2 bar; 35 °C).

Diffusivity selectivity (α)D Solubility selectivity (α)S

Polymer O2/N2 N2/CH4 CO2/CH4 O2/N2 N2/CH4 CO2/CH4 6FDA-DAT1 3.8 5.2 11 1.4 0.3 3.5 6FDA-DAT2 3.8 4.0 8.9 1.27 0.3 3.3