polyimide membranes for gas separation

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POLYMER SCIENCE AND TECHNOLOGY

POLYIMIDES SYNTHESIS, APPLICATIONS AND RESEARCH

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POLYMER SCIENCE AND TECHNOLOGY

POLYIMIDES SYNTHESIS, APPLICATIONS AND RESEARCH

CLYDE MURPHY EDITOR

New York


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NOTICE TO THE READER The Publisher has taken reasonable care in the preparation of this book, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained in this book. The Publisher shall not be liable for any special, consequential, or exemplary damages resulting, in whole or in part, from the readers’ use of, or reliance upon, this material. Any parts of this book based on government reports are so indicated and copyright is claimed for those parts to the extent applicable to compilations of such works. Independent verification should be sought for any data, advice or recommendations contained in this book. In addition, no responsibility is assumed by the publisher for any injury and/or damage to persons or property arising from any methods, products, instructions, ideas or otherwise contained in this publication. This publication is designed to provide accurate and authoritative information with regard to the subject matter covered herein. It is sold with the clear understanding that the Publisher is not engaged in rendering legal or any other professional services. If legal or any other expert assistance is required, the services of a competent person should be sought. FROM A DECLARATION OF PARTICIPANTS JOINTLY ADOPTED BY A COMMITTEE OF THE AMERICAN BAR ASSOCIATION AND A COMMITTEE OF PUBLISHERS. Additional color graphics may be available in the e-book version of this book.

Library of Congress Cataloging-in-Publication Data ISBN: 978-1-53610-623-7

Published by Nova Science Publishers, Inc. † New York


CONTENTS Preface Chapter 1

Chapter 2

Chapter 3

Chapter 4

vii Polyimide Membranes for Gas Separation: Synthesis, Processing and Properties Xiao Yuan Chen, Nguyen Tien-Binh, Serge Kaliaguine and Denis Rodrigue Charge Behaviors on Direct-Fluorinated Polyimide Films and Polyimide Composites Boxue Du A Review on Recent Progress of Research and Applications for Polyimide Aerogels Jin-gang Liu, Xiu-min Zhang, Fei-xu Chen, Wang-shu Tong and Yi-He Zhang Advanced Fabrication and Multi-Properties of Polyimide Aerogels and Their Composites Hai M. Duong, Peng Liu and Daniel Jewell

Index

1

73

105

131 163



PREFACE Polyimide (PI) is an engineering polymer material with advantages of high temperature resistance, low temperature tolerance, radiation resistance, flexibility and excellent dielectric properties. This book presents research on the properties, composites and uses of polyimides. Chapter One focuses on the synthesis and structure-property relationships of neat polyimides, and provides a list of the different methods available to improve the properties of neat PI for gas separation. Chapter Two presents a study aimed at clarifying the effect of fluorination time on surface charge and space charge behaviors of fluorinated PI film and PI/Al2O3 composites. Chapter Three focuses on the latest research and developments on PI aerogels, including crosslinkers synthesis chemistry, PI aerogels synthesis chemistry, and their engineering applications. Chapter Four summarizes the current fabrication methods of the polyimide aerogels and their composites. Chapter 1 – Gas separation techniques using polymer membranes are emerging and gradually replacing traditional processes due to their high energy efficiency and small footprint which is highly suitable for applications like off-shore and small-scale operations. Today, the main interest in membrane separation is for carbon dioxide removal from natural gas or onsite nitrogen generation from air. On the other hand, large-scale applications require polymer materials having high gas permeability to minimize the need for high membrane surface area, as well as high selectivity to meet the requirements related to high purity (low contamination levels). Due to their excellent physicochemical properties such as good thermal stability, chemical resistance and processability, polyimides are one of the best candidates for the production of these membranes. The first part of this chapter focuses on the synthesis and structure-property relationships of neat polyimides. A wide


viii

Clyde Murphy

range of chemical structure can be produced since their synthesis is possible by starting with different combinations of dianhydrides and diamines commercially available or being developed in several research laboratories around the world. The relationships between gas permeation properties and polymer structure is presented and discussed. The second part presents a list of the different methods available to improve the properties of neat PI for gas separation. In particular, the preparation and applications of flat and hollow fiber membranes based on commercial PI, co-PI, crosslinked or PI blended with polymers of intrinsic microporosity (PIM) is presented. Other possibilities are physical or thermal modifications of PI such as thermally rearranged (TR) polymers, or the addition of organic and inorganic particles to produce mixed matrix membranes (MMM). Finally, an overview of the current properties of commercially available membranes is presented with openings for future development. Chapter 2 – Polyimide(PI) film, as a special type of engineering plastic film, is a kind of basic insulating material and is widely applied in the aerial, nuclear, microelectronic industry, turn to turn insulation and turn to ground insulation of inverter-fed motors. However, PI insulation encounters some serious problems in practice. The existence of surface charge and space charge has a great effect on breakdown characteristic and is the main reason leading to dielectric breakdown. Fluorination as change the chemical component in surface layer of polymers should give rise to the corresponding change in electrical properties of the surface layer thus influence the charge injection from electrodes when they are used as an insulator. Besides, the addition of nanoparticles into PI can improve the insulating properties compared with pure material. This chapter presents a study aimed at clarifying the effect of fluorination time on surface charge and space charge behaviors of fluorinated PI film and PI/Al2O3 composites. Obtained results show the dependence of the charge density as well as the charge decay rate upon the fluorination time of samples, varying as a function of the charge polarity and charging time. It is suggested that the fluorination can significantly improve the decay rate of the surface charge in PI films and the injection of space charge is also suppressed. In the study of PI/ Al2O3 composites, it is concluded that the fluorination can significantly improve the decay rate of the surface charge in PI films. Chapter 3 – Recent research and development of polyimide (PI) aerogels have been reviewed. PI aerogels possess both of the merits of conventional aromatic PIs and common polymer aerogels; thus have been widely investigated as components for high temperature, low dielectric constants, and low density applications. Up to now, they have found various applications in


Preface

ix

aerospace, aeronautical, microelectronic and optoelectronic fields. The current review coveres the latest research and development for PI aerogels, including crosslinkers syntheisis chemistry, PI aerogels synthesis chemistry, and their engineering applications. Especially, this review focuses on the applications of PI aerogels as high temperature resistant components for aerospace exploration; thermally and electrically insulating materials for electrical engineering; and low dielectric constant (low-k) interlayer dielectrics (ILDs) for microelctronic fabrications. Chapter 4 – Aerogels are extremely low-density materials with excellent thermal and acoustic insulating properties based on their high porosity and small pore diameter. This makes them attractive candidates for many aerospace applications, such as insulation for cryotanks and spacesuits, as well as more down-to-earth uses in construction, refrigeration, and pipe insulation. The main drawback that has prevented aerogels from having a broad commercial impact is their fragility. Conventional silica aerogels are fantastic insulators but crush easily and are difficult to work with. NASA’s Glenn Research Center has developed exceptionally strong polyimide aerogels that are up to 500 times stronger and have equivalent insulation ability to silica aerogels. As thin films, these polyimide aerogels are highly flexible, lightweight, and porous. Notably, the ability to fabricate the polyimide aerogels into thin films is a revolutionary advancement over silica aerogels. The innovation is technologically significant and unparalleled in the aerogel marketplace, as no other aerogel possesses the compressive and tensile strength of the Glenn polyimide aerogel while it simultaneously can be flexibly folded to contour to whatever shape is needed. In this book chapter, the authors summarize the current fabrication methods of the polyimide aerogels and their composites. The formation mechanisms, morphologies and multi-properties of polyimide aerogels and their composites are carefully discussed. This review chapter would be meaningful for exploiting the structures, properties and potential applications of the polyimide aerogels and their composites.



In: Polyimides Editor: Clyde Murphy

ISBN: 978-1-53610-596-4 © 2017 Nova Science Publishers, Inc.

Chapter 1

POLYIMIDE MEMBRANES FOR GAS SEPARATION: SYNTHESIS, PROCESSING AND PROPERTIES Xiao Yuan Chen1,2, Nguyen Tien-Binh1, Serge Kaliaguine1 and Denis Rodrigue1,* Department of Chemical Engineering, Université Laval, Quebec, Canada Centre National en Électrochimie et en Technologies Environnementales, Collège de Shawinigan, Shawinigan, Canada

1 2

ABSTRACT Gas separation techniques using polymer membranes are emerging and gradually replacing traditional processes due to their high energy efficiency and small footprint which is highly suitable for applications like off-shore and small-scale operations. Today, the main interest in membrane separation is for carbon dioxide removal from natural gas or onsite nitrogen generation from air. On the other hand, large-scale applications require polymer materials having high gas permeability to minimize the need for high membrane surface area, as well as high selectivity to meet the requirements related to high purity (low contamination levels). Due to their excellent physicochemical properties such as good thermal stability, chemical resistance and process-ability, polyimides are one of the best candidates for the production of these *

Corresponding author E-mail: [email protected].


2

Xiao Yuan Chen, Nguyen Tien-Binh, Serge Kaliaguine et al. membranes. The first part of this chapter focuses on the synthesis and structure-property relationships of neat polyimides. A wide range of chemical structure can be produced since their synthesis is possible by starting with different combinations of dianhydrides and diamines commercially available or being developed in several research laboratories around the world. The relationships between gas permeation properties and polymer structure is presented and discussed. The second part presents a list of the different methods available to improve the properties of neat PI for gas separation. In particular, the preparation and applications of flat and hollow fiber membranes based on commercial PI, co-PI, crosslinked or PI blended with polymers of intrinsic microporosity (PIM) is presented. Other possibilities are physical or thermal modifications of PI such as thermally rearranged (TR) polymers, or the addition of organic and inorganic particles to produce mixed matrix membranes (MMM). Finally, an overview of the current properties of commercially available membranes is presented with openings for future development.

Keywords: gas separation, polyimides, synthesis, flat membrane, hollow fiber membrane, mixed matrix membrane

1. INTRODUCTION Nowadays, polyimides (PI) are becoming common membrane materials for gas separation mostly because of several advantages. Firstly, PI have excellent thermal, chemical and mechanical properties. The glass transition temperature of most PI is above 300oC and their stability (the temperature to reach 5% weight loss) is usually above 400oC. Most of PI only dissolve in polar solvents. Therefore, PI membranes exhibit very high resistance to almost all chemical agents. Their heat resistance allows performing separations for long period of time at elevated temperatures [1]. Secondly, PI have excellent film-forming properties and can be easily used to prepare asymmetric structures (hollow fibers) or flat membranes which can be put into different configurations such as modules, spiral wounds and envelopes. Thirdly, PI is made from a polycondensation reaction between aromatic acid dianhydrides and aromatic diamines [2]. Due to the variety of acid dianhydrides and diamines commercially available or developed in laboratories around the world, PI based materials can be used to prepare a series of membrane with


Polyimide Membranes for Gas Separation

3

systematically different gas transport properties via the control of their chemical structures. Another feature of PI is that they exhibit higher permselectivity and permeability for gas separation than most common glassy polymers such as polysulfones, polyetherimide, polyethersulfone and polycarbonates [3]. Finaly, PI have high durability and can be used for long time periods under uniform conditions. Gas separation technologies based on PI membranes can be applied in several fields such as: separation of nitrogen and oxygen from air (O2/N2), separation of hydrogen from synthesis gas (syngas, H2/CH4, H2/CO2), in ammonia plants (H2/N2, H2/CH4) and from petroleum refining processes (H2/CO), purification of methane from biogas (CO2/CH4) and from natural gas (CO2/CH4), water vapor removal from natural gas and other gases (H2O/CH4), CO2 capture (CO2/N2), as well as helium recovery from rejected gas streams during natural gas processing (He/N2). PI membranes have been applied industrially at large scale. For example, UBE membranes (Japan) used polyimide hollow fiber membranes for H2 separation in 1989. Today, they are also to separate N2 from air [4]. MEDAL (Air Liquide, France) used polyimide hollow fibers membranes for CO2/CH4 separation in natural gas and biogas since 1994. Today, this company also uses PI membranes for O2/N2 and hydrogen recovery [5]. Evonik, with SEPURAN® Green polyimide hollow fibers, built a biogas plant in Switzerland with an upgrading capacity of 210 standard cubic meters per year in 2013. Rite (Japan) uses PI hollow fibers for three main applications: CO2/N2 (separation and capture of CO2 from combustion exhaust gases), CO2/CH4 (separation and capture of CO2 from natural gas) and CO2/H2 (separation and capture of CO2 from the Integrated Coal Gasification Combined Cycle (IGCC) process under high pressure. Alpha (Singapore) uses PI hollow fibers for CO2/CH4 separation. Fuji film developped PI flat membranes into a spiral wound configuration to efficiently (low energy consumption) separate CO2 from mixed gases to produce purified natural gas. This company also reports that compared to other separation methods (chemical absorption, physical absorption, etc.), this method makes it possible to reduce production facility size and energy cost. Finally, PI have been used for several other applications such as films, fibers, foams, binders, varnishes, matrix for composites, glues, adhesives and injection molded products. In this chaptre, a focus on membrane applications for gas separation is presented.


4

Xiao Yuan Chen, Nguyen Tien-Binh, Serge Kaliaguine et al.

2. POLYIMIDE SYNTHESIS The most popular synthesis route to obtain PI is the reaction of an aromatic dianhydride with an aromatic diamine (1:1 molar ratio) at room temperature in aprotic polar solvents such as N-methyl-pyrrolidone (NMP), N,N-dimethylacetamide (DMAc), N,N-dimethylformamide (DMF) and dimethyl-sulfoxide (DMSO) to form a soluble intermediate, polyamide acid, which is chemically or thermally converted to PI. The two-step synthesis route is presented in Figure 1.

Figure 1. General scheme for the two-step synthesis of PI.

The reaction is generally conducted at low temperature to favor the propagation of high molecular weight polyimide in the first step which is exothermic, while the imizidation step is proceeding either chemically at room temperature with some basic catalysts (isoquinoline, trimethylamine, pyridine), or thermally at higher temperature with azeotropic solvents (toluene, xylene) [6]. Several types of PI have been prepared for gas separation, yielding a good understanding on the relationships between PI molecular structures and their gas separation properties. It has been reported that PI having high chain stiffness, weak interchain interaction and loose chain packing tend to have high gas permeability [7-10]. This means that both dianhydrides and diamines components of the PI should not contain flexible links. Due to the glassy nature of these polymers, PI have a tendency to thermodynamically rearrange their chains reducing the number and sizes of free volume contributing to gas permeability. This phenomenon could be limited by inhibition of the free rotation inside diamine moieties around the imide links (or C-N bonds) as shown in Figure 2.


5


Figure 2. A comparison between the rotation of amine moieties in PI with and without ortho-substituted groups.

3. FUNDAMENTALS OF GAS TRANSPORT IN POLYMER MEMBRANES In membrane-based separations, a feed gas mixture is pressurized on the feed side of the membranes. Due to differences in physical properties such as kinetic diameter and condensability, the other side of the membrane will release a specific gas-enriched mixture called the permeate. The gas transport behavior through polymer membranes is generally explained by the solutiondiffusion [11, 12]. Gas molecules are absorbed in the membranes on the high pressure side (feed side), the absorbed molecules then diffuse through the membranes with a driving force being the concentration gradient. They finally desorb on the lower pressure side (permeate side). The permeability coefficient (or permeability), P, of a gas is defined as the product of its diffusivity, D, and solubility, S, in the membrane as:

P  D S

(1)

The permeability of a penetrant gas A (PA) through a polymer film is defined by:


6


PA 

NA l p2  p1

(2)

where NA is the steady state gas flux through the film, p2 and p1 are upstream and downstream pressures, respectively, and l is the film thickness. Permeability is generally reported with units of mol m m-2 s-1 Pa-1 or more conveniently in Barrer where 1 Barrer =10-10 cm-3 (STP) cm cm-2 s-1 cmHg-1. The permeability of a polymer to gases is dependent of the physical properties of the permeant gases (molecular size, shape, and polarity), membrane properties (physical and chemical structures) and interaction between the permeant gases and the polymer. Gas solubility generally increases with increasing gas condensability (or critical temperature). For example in most polymers, CO2 is more soluble than CH4, O2 and N2. If a polymer membrane is used to separate a gas mixture containing two components (A and B), where component A is more permeable than B, its separation efficiency can be characterized by the ideal selectivity, αA/B given by:

 A/ B 

PA DA S A   PB DB S B

(3)

DA/DB and SA/SB are the diffusivity selectivity and solubility selectivity, respectively. For a given gas pair, the solubility selectivity is almost constant for a wide range of chemically different polymers, while the diffusivity selectivity is highly function of polymer chain rigidity [13].

3.1. Gas Solubilities in Polymers PI belong to the class of glassy polymer for which the unrelaxed polymer chains create an excess amount of free volume within the bulk density which is absent in rubbery polymers. Therefore, the gas sorption capacity of glassy polymers is generally much higher than for their equivalent rubbery polymers counterparts. Gas dissolution in polymers is considered to be a two-step process: (1) condensation of the gas molecules to a liquid-like density and (2) mixing of the condensed penetrants within the polymer segments. Gas sorption in glassy


7


polymers was first proposed by Koros et al. [14-17] using the dual-mode model to explain the adsorption isotherm of gases being typically non-linear with pressure. The dual-mode model is mathematically described by Eqn(4). In this model, gas penetration in glassy polymers is a combination of the “dissolved mode” (gas dissolution in the equilibrium regions of the polymer following the Henry Law’s isotherm) and the “Langmuir mode” (gas dissolution in the non-equilibrium excess volume of glassy polymers).

C  CD  CH

(4)

where C is the total concentration of gas adsorbed in the polymer, CD is the dissolved mode penetrant concentration, CH is the penetrant concentration in the Langmuir mode. CD can be represented by a linear function of pressure, while CH is given by a Langmuir isotherm. The total concentraion C can be obtained as [18]:

C



kD p Dissolved mode

CH bp 1  bp

(5)

Langmuir mode

where kD is the Henry law’s constant, CH’ is the hole saturation constant or Langmuir sorption capacity and b is the Langmuir affinity parameter. CH’ is related to the unrelaxed exceed volume or fractional free volume of the glassy polymer. CH’ is a temperature-dependent parameter, but under the sorption isothermal conditions it is a constant characterizing the fractional free volume existing inside the glassy polymer. The gas sorption according to the Langmuir mode therefore increases with pressure and becomes saturated at high pressure ranges. Low pressure: CH 

High pressure: (p

CH bp CH  1  bp 1  1 bp 1 1)   0, CH  CH  const bp

(6)

The dual mode is well fitted to experimental data for gas sorption in most glassy polymers, which is typically concave at low pressures and linear at high


8


pressures. Figure 3 shows the sorption isotherm curves of CO2 and CH4 gases for Matrimid®, a glassy polymer [19], and PDMS, a rubbery polymer [20]. The gas sorption isotherms in PDMS rubbery polymer are linear with pressures indicating that the Langmuir mode is absent and the gas sorption is mostly following Henry’s law, the dissolved mode [21].

a

b Figure 3. Sorption isotherm at 35°C of (a) CO2 and (b) CH4 for Matrimid (▲), a glassy polymer [19] compared to PDMS (●), a rubbery polymer [20].



9

Figure 4. Sorption isotherm of CO2 in polymethyl methacrylate, a glassy polymer (Tg = 105°C) at different temperatures.

Since the unrelaxed volume in glassy polymers is a function of temperature, if a glassy polymer is heated close to its glass transition temperature (Tg), the Langmuir mode will become negligible. Figure 4 presents a typical example of the temperature evolution for CO2 sorption isotherms in a glassy polymer. The gas solubility decreases with temperature and becomes linear as the sorption isotherm is measured at temperatures above the glass transition temperature of the polymer. As described earlier, gas sorption in glassy polymers is more directly related to the packing density as well as the material’s thermal history. Because of the structure of glassy polymers, the unrelaxed volume decreases with time (aging phenomenon), particularly the polymer is stored at temperatures just below Tg, leading to a decrease in gas solubility and permeability. Gas solubility generally increases with increasing gas condensability (or critical temperature). For example, in most glassy polymers, CO2 is more soluble than CH4, O2 and N2 as shown in Figure 5. It has been expected that incorporation of functional groups into the polymer backbones could help to enhance the gas solubility [36, 37]. Although several functional groups have been incorporated in polymers for gas separation applications, no relationships between gas solubility and chemical functionality has been found yet. Figure 6 reports on the solubility data of several gases in a wide range of chemically different glassy polymers modified with amidoxime, hydroxyl and Trögerbase functional groups. It can be seen that the solubility is relatively unchanged for all these cases. But higher solubility for condensable gases was


10


observed for all the glassy polymers indicating that no specific polymer-gas interaction was related to the solubility.

Figure 5. Relationships between gas solubility and critical temperature measured at 35°C and 2 atm. PC = polycarbonate [16], PSF = polysulfone [14, 22], PI = HAB6FDA polyimide [23] and TR = thermally arranged polyimide [23].

Figure 6. Correlation between polymer functionality and gas solubility measured at 25°C and 1 atm. PIM-1 = polymer of intrinsic microporosity [24], AO-PIM-1 = amidoxime- PIM-1 [25], PI-TB = Tröger-base polyimide [26-29] and PI-OH = hydroxyl-PI [30-35].


11


3.2. Packing Density of PI An important property of polymer membranes for gas separation is the packing density of the polymer chains. This parameter is related to the free volume inside the membrane which is unoccupied by polymer molecules. Because of the rigid nature of glassy polymers, when the molecules are packed in the solid state, each of them often takes more space than its actual molecular volume. According to this concept, the packing density is quantified by the fractional free volume (FFV) and calculated as:

V f  (Vsp  1.3 Vw )

(7)

FFV  (V f / Vsp )

(8)

where Vf is the free volume, Vsp is the specific volume derived from the polymer density and Vw is the specific van de Waals volume calculated using the group contribution method of Bondi [38]. Typically, higher FFV values are producing higher gas permeabilities. Packing density and FFV are mainly controlled by chain stiffness, polymer chains local motion, side groups and chain bulkiness. Basically, each repeated unit of PI has four polar carbonyl groups organized on a planar structur as shown in Figure 7a. The van der Waals interaction between polymer chains makes the chain packing more efficient and decreases FFV, therefore, decreasing gas permeability. When bulky side groups are introduced into the diamine moieties (Figure 7b), they will twist these planar faces causing a structure contortion in PI yielding a weak chain packing and high FFV [39]. O

(b)

(a)

O

O

N O

N O

O N O

O R1 R2

R4

N

O

O

N O

O

R2

R1 R2

R3 O

O

R1

R4 N R3

R4

N

R4

R1

N R3

R2

O

R3 O

Figure 7. Schematic illustration of chain packing in PI: (a) without substituted side groups, polymer chains pack more efficiently, and (b) with substituted side groups, the chain packing is more difficult.


12


3.3. Effects of Diamine Structure The influence of diamines with bulky structured side-groups on FFV and gas separation properties has been intensively studied in the 1990s [7, 8, 40, 41]. Some examples are presented in Table 1 with typical values. The increasing number of methyl groups connected to the ortho positions of diamines mMPD < pDiMPD < DAM < TeMPD effectively inhibits intersegmental chain packing forming PI with higher FFV leading to higher gas permeability but lower selectivity than their lower FFV counterparts. Table 1. Gas permeability and selectivity for CH4, N2 and CO2 in a series of 6FDA-based PI, at 35°C and 10 atm [40] PI monomers Diamines

6FDA-Dianhydride

Permeability (Barrer) Ideal FFV selectivity CH4 N2 CO2 CO2/CH4 0.28 0.80 15.3 54.6 0.161

pPD 0.16

0.45

9.2

57.5

0.161

0.88

2.24

40.1

45.6

0.176

1.07

2.67

42.7

39.9

0.175

26.0

31.6

431

16.6

0.182

28.4

35.6

455.8 12.8

0.182

mPD

mMPD

pDiMPD

DAM

TeMPD


13


Previous studies have reported important structure-gas properties of PI guiding molecular design to upgrade the PI permeability without sacrifying selectivity. Polymer chain rigidity generally controls selectivity while interchain spacing and chain mobility govern permeability [42]. As seen in Table 1, the selectivities of PI made from meta-isomer diamines such as DAM and mMPD are higher than those based on para-isomer diamines namely TeMPD and pDiMPD, while permeabilities were quite similar. This is because the meta-links in diamines created PI with asymmetrical configurational backbones that were unfavorable for the free rotations around the imide bonds, hence the backbone becomes sterically stiffer than the corresponding symmetric PI produced from the para-isomer diamines. Another important element for polyimide gas properties is the chemistry and configuration of the internal-connecter groups. It has been expected that PI containing CO2-philic functional groups such as the amines could be good candidates for CO2 related separations, since these functional groups were believed to selectively interact with CO2, hence increasing both the CO2 permeability and selectivity. This is why Tanaka et al. [7] prepared a series of 6FDA-based PI with diamines made from a wide range of functional connector groups. The list of diamine structures studied is reported in Table 2. Table 2. The effects of the internal connector groups functionality on permeability, solubility and diffusivity of 6FDA-derivated PI, determined at 308 K and 10 atm H H2N

O C

H C

H2N

O

O S

H2N NH2

NH2

NH2

H3C CH3 DAF H

DAFO

N

N

NH2

H2N CDA

DDBT C2H5

NH2

H2N ECDA

Diamines PCO2 DCO2 SCO2 PCH4 PCO2/PCH4 DCO2/DCH4 SCO2/SCH4 DAFO 7.68 1.3 57 0.13 60 19 3.2 DAF 11.3 1.8 61 0.21 54 17 3.2 CDA 17.1 2.6 67 0.30 57 19 3.0 ECDA 33.8 5.3 63 1.03 33 12 2.8 DDBT 91.0 10.8 84 2.51 36 13 2.9 P in 10-10 cm3 (STP) cm-1 s-1 cmHg-1, D in 10-8 cm2 s-1, S in 10-3 cm3 (STP) cm-3 cmHg-1


14


Permeability for the 6FDA-based PI was in the following order: DAFO < DAF < CDA < ECDA < DDBT. The permeability variation is mainly coming from chnges in DCO2. CO2 diffusivity strongly varied from 1.3 up to 10.8, whereas the CO2 solubility fluctuated in a limited range from 57 to 84 for all the diamine moieties. Interestingly, the PI obtained from the diamines containing amine functional groups (CDA and ECDA) showed lower values for SCO2 and DCO2 than those of the PI prepared from the sulfonefunctionalized diamine (DDBT). This suggested that such a specific interaction between CO2 and the functional groups does not exist. This is in agreement with the quasi constant values of solubility selectivity of about 3.0 for CO2/CH4, as observed for the five PI based of functionalized or un functionalized polymers. There was no clear contribution of the functional groups on the PI separation properties. The highest permeability observed for DDBT-based PI was mainly attributed to the rigid, bulky and nonplanar structure of the diamine which turned perpendicular to the imide rings [7]. Bulky connector groups can strengthen the chain stiffness and inhibit chain packing density leading to increasing permeability. As seen in Table 3, polar connector groups such as –O‒ and –SO2‒ in diamines can increase the inter-chain dipole interactions or form hydrogen bonding significantly reducing the free volume. Therefore, the resulting PI have lower permeability but higher selectivity. On the other hand, nonpolar and bulky connectors, such as -C(CH3)2- and -C(CF3)2-, are believed to act as molecular spacers increasing the free volume and gas permeability of the polymers. PI with the polar connectors listed in Table 3 showed obviously lower permeability and higher selectivity than those prepared from the methyl-substituted diamines such as DAM or TeMPD in Table 1 suggesting that the bulky connectors induce a positive contribution to selectivity due to the improved backbone rigidity while the side-group substituents strongly enhance permeability because of increased inter-chain spacing. To improve permeability with limited selectivity loss, Guo et al. [44, 45] prepared a series of diamines having both bulky connectors and side-group substituents to obtain novel 6FDA-based PI showing advanced gas separation properties over commercial glassy polymers such as Matrimid®. As seen in Table 4, due to the increasing bulkiness of connectors, the pentiptycenestructured PI showed much higher permeability than the triptycene-structured PI, while their selectivity was similar. So the presence of polar connector groups such as –O‒ in PI structures is usually associated with unexpected inter-chain interactions and weaken the rigidity of the PI backbone due to freerotation around the connector groups, thus decreasing permeability and



15

selectivity of these PI. Therefore, significant efforts have been made to design non-functional group diamines with bulky and rigid bridging groups. A number of PI based having spiro-centrers [46-49], terphenylene [50], three dimensional triptycene [51-54] and Tröger base structures [26-29, 55-57] have been reported and exhibit high permeability and selectivity, even surpassing the well-known upper bound limits [58, 59]. Table 5 summarized the selected 6FDA-based PI for which the ones synthesized from diamine D showed the highest permeability due to the presence of four ortho-methyl groups as well as a rigid connector-Tröger base. The diamines E, F, and G with a similar structure, but having only two ortho-methyl groups, consequently led to PI having much lower permeabilities. The selectivity of PI made from diamines F and G was higher than those based on diamines A and B while their permeabilities were relatively comparable because the Tröger base link strictly prohibited any free rotation around the connectors, which is probably occurring in A and B. Table 3. The effect of internal connector groups on 6FDA-derivated PI [43]

Connector X

PCO2 (Barrer)

α (CO2/CH4)

-O-SO2-CH2-C(CH3)2-C(CF3)2-

17 16 19 30 64

49 47 45 43 40

In summary, in terms of diamine design, there are some important rules that should be kept in mind: (1) introducing ortho-substituted non-polar groups in diamines is among the most effective approach enhancing permeability because of the restricted rotational degree around the imide links, (2) polar substitutions or polar connectors tend to increase the inter-chain interactions or facilitate the intra-diamine flexibility respectively, therefore reducing permeability associated with a slightly increased selectivity, and (3) the bulky and rigid links are believed to act as molecular spacers and chain stiffeners that effectively limit intra-segmental mobility as well as the degree of polymer


16


chain packing. As a result, the PI based on Tröger base or triptycene links often have both higher permeability and selectivity. Table 4. 6FDA-derivated PI prepared from diamines with bulky connectors and side-group substituents

PCO2

Substituent R1 H CH3 CF3

α (CO2/CH4) 29 28 24

(Barrer) 73 55 132

PCO2

Substituent R2 H CH3 CF3

(Barrer) 14.4 9.0 19.7

α (CO2/CH4) 39 39 37

Table 5. Permeability and selectivity of selected 6FDA-PI flat membranes Permeability Ideal (Barrer) selectivity Ref. CH4 CO2 CO2/CH4

PI monomers Diamines

H2N

6FDA-Dianhydride

NH2

15.0

360

24.0

[50]

5.6

190

33.9

[50]

6.2

189

30.5

[51]

A

H2N

B

NH2

O O

F3C CF3

O O

NH2

H2N C

O

O


17

Polyimide Membranes for Gas Separation PI monomers Diamines

6FDA-Dianhydride NH2

N

Permeability Ideal (Barrer) selectivity Ref. CH4 CO2 CO2/CH4

160

1672

14.4

[55]

27

457

17

[26]

8

285

35.6

[29]

3.3

155

46.9

[29]

N

H2N

D

NH2

N N

H2N

E

H2N

N N

NH2

F NH2

N H2N

N G

3.4. Effect of Dianhydride Structure Strategies in modifying dianhydride structures are quite limited and more difficult than for diamines because of the highly reactive nature of the anhydride groups. Newly synthesized in laboratories as well as commercialized dianhydrides are therefore not as abundant as diamines. Table 6 compares the performance of PI prepared from 4,4′-oxydianiline diamine and the four most common dianhydrides. It can be seen that 6FDA-based PI have gas permeabilities of 5 to 25-folds higher with similar or half the selectivities than the others. Such an impressive performance is attributed to both -CF3- links in the 6FDA dianhydride structures, their bulkiness not only increases the spatial domain between polymer chains, referring to higher diffusivity, but also inhibits intra-segmental rotations within the dianhydride moieties, corresponding to better diffusivity selectivity. Therefore, 6FDA-


18


based PI often show a balanced gas separation behavior between high selectivity and moderate permeability. Due to the nature of the anhydride groups which are relatively polar and bulky, the incorporation of functional or bulky connectors into the dianhydride structures is less effective on enhancing the PI gas separation. But recent studies have been attempting to increase the rigidity of dianhydride structures. For example, Ghanem et al. [47] successfully obtained a series of extremely rigid PI, PI of intrinsic porosity called PIM-PI, whose dianhydride components included spiro-center links which strictly prohibited the self-rotation of dianhydrides. These PIM-PI exhibit high specific surface area with BET values, determined by nitrogen adsorption, above 500 m2/g. This interesting characteristic gives PIM-PI an improved permeability compared to more traditional PI usually having lower surface area. Table 6. Gas permeability and selectivity measured at 35°C and 10 atm, for CH4, N2, CO2 in several PI prepared from various dianhydrides [8] PI monomers Diamine

ODA

Permeability/Barrer Dianhydrides

Ideal FFV selectivity

CH4

N2

CO2 CO2/CH4

0.43

0.73 16.7 38.8

0.165

0.09

0.15 3.5

0.129

0.01

0.02 0.62 62

0.124

0.01

-

0.121

6FDA 38.9

PMDA

BTDA 0.64 64

BPDA


19

Polyimide Membranes for Gas Separation Table 7. The effect of diamines structures on the permeability and selectivity of PIM-PI determined at 30°C and 200-500 mbar of feed pressure using a barometric method [47] O O

O O

O

O

O O

O O Dianhydride

Diamines

Permeability (Barrer) CO2 CH4

α (CO2/CH4)

BET surface area (m2/g)

1100

77

14.3

680

210

9

23.3

500

520

27

19.3

471

420

20

21.0

486

510

27

18.9

485

3700

260

14.2

683

H2N NH2

F3C

H2N

CF3

NH2

F3C CF3

H2N

NH2 CF3

H2N

NH2 F3C

NH2 H2N

H2N

NH2


20


The spiro-center structure, in fact, is not completely rigid due to the presence of a single carbon joining the two cyclopentane rings that can show some flexibility. To further inhibit the segmental chain mobility particularly in the dianhydride moieties, which can improve selectivity, Ghanem et al. [52, 60] prepared ultra-microporous PI obtained from a novel triptycene structured dianhydride shown in Table 8. The contorted structure of the dianhydride is a non-planar and bridged tricyclic molecular site that additionally improves the PI rigidity while inhibiting the closed chain packing. Two isopropyl groups introduced in the dianhydride is expected to limit rotations around imide bonds. These molecular design of dianhydrides have substantially improved the gas separation properties that surpassed the upper bonds for most gas pairs such as O2/N2, H2/N2, H2/CH4 and even CO2/CH4. Table 8. The structure of 9,10-diisopropyl-triptycene dianhydride, permeability and selectivity of the obtained PI determined at 35°C and 2 bar [60]

O

O O

O

O

O

O

O

O

O Dianhydride

Permeability (Barrer)

Diamines

α (CO2/CH4)

BET surface area (m2/g)

CO2

CH4

2389

105

22.7

752

2071

101

19.7

-

H2N NH2

H2N

NH2

Although the results are very promising, there has been a limited number of new dianhydrides reported. Recently, Ma et al. [61] presented a novel spirobifluorene dianhydride (SBFDA) using a complex synthesis route using a polymerization step with a sterically hindered diamine (DMN) yielding a microporous PI. Because of the rigid 90° contorted center of the dianhydride,



21

the obtained polymer SBFDA-DMN exhibited a significant fraction of micropores mainly distributed at 7 and 15 Å. The structure of SBFDA-DMN, as well as its pure-gas separation property, are summarized in Table 9. The permeability was exceptionally high for all the gases with an acceptable selectivity, which were attributed to the self-locking rotation as well as the bulkiness of the spirobifluorene built-in dianhydride. Table 9. The structure of SBFDA-DMN polyimide and its permeability and selectivity determined at 35°C and 2 bar for a membrane soaked in methanol and then air-dried for 24 h [61] O O

N

N

O

O

Permeability (Barrer) H2 N2 O2 3342 369 1193

CO2 6674

CH4 581

α (CO2/CH4) 11.5

BET surface area (m2/g) 686

4. INHERITED LIMITATIONS OF POLYIMIDE MEMBRANES 4.1. Permeability-Selectivity Invert-Relationships A number of polymer categories have been examined for gas separation in searching of a suitable material having both high permeability and selectivity. In 1991, Robeson collected a large amount of data from hundred studies and proposed an upper bound limit showing a general trade-off between permeability and selectivity [59]. Recently, the upper bounds were revised using a larger and more recent database [58]. Since then, the Robeson’s upper bounds have been used as standards to determine the performance of any new materials designed. The upper bound for CO2/CH4 separation is shown Figure 8. It is clearly seen that only very few polymers have permeability-selectivity data close to or above the trade-off limit. Polymers having higher permeability usually exhibit lower selectivity and vice-versa.


22


Figure 8. Upper bound curves for the CO2/CH4 gas pair (TR = thermally rearranged polymers, PIMs = polymers of intrinsic porosity) [24, 25, 58, 59, 103, 106].

4.2. Plasticization The plasticization phenomenon is generally observed for glassy polymer membranes used for gas separations involving highly condensable gases such as CO2, H2S, H2O and hydrocarbons (C3+). It occurs when the partial pressure of condensable gases is high enough to disrupt the polymer chain packing and to promote segmental mobility that significantly decrease the membrane selectivity. For CO2/CH4 separation, which mostly uses glassy polymers, CO2induced plasticization accelerates the permeability of CH4, thus decreasing the CO2/CH4 selectivity. Bos et al. [62] investigated the relationships between plasticization and chemical strucures, as well as physical properties of 11 different glassy polymers. Although no corelation was found, the study reported that the polar goups in polymers tend to increase the CO2-induced plasticization because they have dipolar interaction with the gas. Duthie et al. [63] studied the CO2 permeability variations as a function of CO2 feed pressure at different operating temperatures and found that plasticization is more pronounced at lower temperatures.



23

4.3. Physical Aging It is well-known that glassy polymers uasually exist in a non-equilibrium state because their glass transition temperature are well above the room temperature. The relaxation of polymer chains toward an equilibrium is commonly refered to as physical aging. The process occurs at very low speed as it may take several months or years for micrometer thick membranes. Upon the relaxation process, the excess free volume in the polymer membranes will gradually decrease causing a steady decrease in their gas permeability [13]. The aging rate is associated to the polymer materials [63-70]. Highly permeable glassy polymers with loose chain packing density and high free volume experience more importantly the physical aging phenomenon, especially in the first-few days after their processing thus having significant permeabilty loss. Kim et al. [64] studied the physical aging of homo- and coPI containing carboxyl groups, but no clear corelation between the functional goups and the aging was found. However, the aging rates seemed to be correlated with the free volumes: the higher the permeability, the faster the aging rate is. Recently, Swaidan et al. [71] mornitored, over a period of 1-4 years, the aging phenomena in six “rigid” polymers of instrinsic microposity whose permeabilities were initially about several thousand Barrers. Physical aging was also found to occur in all the polymers despite their extremely ridid structures. Aging in glassy polymers is accepted as an unavoidable property, but techniques to reduce its effect on membrane gas separation are necessary.

5. MODIFICATIONS OF POLYIMIDES FOR GAS SEPARATION UPGRADING Although the relationships between gas separation properties and PI structures have been well-documented, but the synthesis of high performance polymers is generally costly and time-consuming. So a modification approach of the existing materials might be an alternative solution to overcome the limitations of separation performance.


24


5.1. Polymer Bending Polymer blending is an interesting route due to its simplicity, reproducibility, processability and low development cost. Blending enables to combine the advantages of different materials which are either miscible or immiscible. Miscible blends are generally more favorable for membrane fabrication since homogeneity is essential to produce membranes with uniform performance. Due to a good solubility in most common solvents, PI have been blended on a molecular level with a wide range of glassy polymers such as polycarbonate [72], polysulfone [73], polyethersulfone [74-76], polyaniline [77], and polybenzimidazole [78-80]. These polymers were expected to solve the problems related to PI swelling and plasticization caused by condensable gases such as CO2, hydrocarbons, and H2S in mixed gas separations. In a mixed gas conditions, glassy polymers, including PI, will be swelled or plasticized by the sorption of condensable gases (e.g., CO2) thereby increasing the permeation of the co-permeating gases (e.g., CH4, N2) in the mixture. The increased permeation of the slow gases in mixed-gas tests resulting in selectivity losses is an unexpected behavior and known as the plasticization phenomenon. Kapantaidakis et al. [73] observed anti-plasticization effects induced by polysulfone when it blended with a commercial PI (Matrimid®). Similarly, Bos et al. [81] found that a copolyimide (P84) can work as an antiplasticizer depressing the plasticization of Matrimid® as plasticization occurs due to the flexibility of glassy polymer chains at temperatures even much lower than their glass transition temperatures (Tg). To effectively inhibit the plasticization of the hosted polymers, the introduced polymer should have a rigid backbone and strong interactions with the hosted polymers. Polybenzimidazole (PBI), a high Tg glassy polymer having both donor and accepter hydrogen-bonding sites being capable of good interactions with PI [82-85], has been found as a good blending material for improving the performance of Matrimid® [79, 80, 83]. These studies generally observed enhanced CO2/CH4 selectivity for PBI-Matrimid blends, but at the expense of permeability compared with neat Matrimid because of the low permeation of PBI.

5.2. Mixed Matrix Membrane Mixed matrix membranes (MMM) or hybrid membranes are made of two components: an organic polymer used as the continuous phase and porous or



25

non-porous fillers used as the dispersed phase. The polymer is usually the main component (accounting for more than 80 wt.%), and controls the overall properties of MMM such as mechanical and chemical properties, as well as gas separation efficiency. The function of the fillers is to systematically enhance the separation properties of the neat polymers. MMM are generally prepared by a solution based on a suitable solvent used to dissolve the polymer forming a homogeneous polymer solution to which the fillers are added. The mixing of both phases is performed by means of mechanical stirring or ultrasonication to prevent particle agglomeration. The solvent is then removed by drying in ambient conditions or vacuum ovens. This is a straight-forward method applicable for a wide range of polymers and fillers, both being independently prepared before being mixed together. In the second approach, also known as the in situ method, the filler precursors are added directly to the polymer solution and the nano-particles are formed during the membrane formation. In this approach, the nano-particles precursors generally have good interactions with the polymer, thereby enabling good dispersion of the precursor and the resultant fillers. Some nano-particles have been obtained in the polymer medium by a sol-gel method including nano TiO2 [86, 87] or nano SiO2 [88-92] which were used as non-porous fillers. Due to the large surface area of nano-particles, they are expected to inhibit the polymer chains packing increasing the permeability. Inside the non-porous fillers, activated carbons [93] and ion exchanged zeolites [94] have a selective surface adsorption for CO2. Therefore, their incorporation increases the CO2/CH4 sepation of the MMM. Another category of porous filler includes zeolites, carbon molecular sieves and metal organic frameworks (MOF) of which the highly selective solubility and (or) diffusivity in the porous systems can be used to tailor the gas separation properties of MMM. Depending on the nature of the filler surface (hydrophilic or hydrophobic) and polymer materials used, the interaction between them may vary significantly leading to important changes in the MMM performance. To solve the polymer-filler incompatibility issue, recent studies focussed on three solutions: (i) using copolymers having hydrophilic functional groups that can form hydrogen bonding with the filler particles [37, 95-97], (ii) decreasing the filler particle sizes to the nano-scale so their interaction with the polymer phase will be improved [98-102], and (iii) thermo- or photo-treating the MMMs to induce covalent cross-linking between the polymer and filler [103-106]. A detail discussion could be found in our recent reviews [107, 108].


26


5.3. Inter-Chain Cross-Linking The plasticization resistance of a membrane can be improved by crosslinking of the carboxyl-containing PI using alkane-diols as cross-linking agents [109-111]. The cross-linking process includes two steps as described in Figure 9: (1) an esterification with an excess of alkane-diols mixed with a carboxylic-included polymer dissolved in a solvent where the esterification reactions can be accelerated at 140°C using p-toluenesulfonic acid as a catalyst to obtain an intermediate, called a monoester cross-linkable polymer (PDMC) [112], and (2) a transesterification in which the PDMC is used to cast membranes that are then annealed at high temperatures under vacuum to induce the inter-chain cross-linking reactions. Cross-linking offers a potential to improve not only the thermal and physical properties, but also gas separation properties. Hilock et al. [110] observed a 3-fold increase in permeability without any selectivity loss and good plasticization resistance up to 450 psi for the copolyimide 6FDA-DAM-DABA (3:2) when crosslinked with 1,3-propanediol. It should be noted that the transesterification reaction occurred at elevated temperatures which helped to increase the free volume due to the removal of pendant diol groups, therefore improving the gas permeability.

Figure 9. A two-step process to induce inter-chain cross-linking between a carboxylincluded polymer using 1,3-propanediol as the cross-linking agent.

The diol cross-linking agents used to establish ester linkages between coPI containing carboxylic acid pendant groups may cause unexpected side effects such as the unreacted cross-linking agents remained in the polymer acting as a



27

plasticizing agent decreasing permeability and selectivity. Additionally, ester links can potentially hydrolyze the acid gas feed streams which could reverse the cross-linking reactions and greatly reduce the membrane efficiency. Kraochvil and Koros [113], proposed a cross-linking mechanism without using cross-linking agents for COOH-contained coPI. The cross-linking occurred due to decarboxylation reactions of the COOH groups when the solidstate membranes were annealed at temperatures higher than Tg, the detail mechanism is given in Figure 10. Nevertheless, high temperature treatments are difficult to apply in membrane fabrication where the membranes are usually asymmetric: a compact skin layer coated on a porous support layer. The thermal treatment may result in densification and flexibility reduction of the support layer, as well as creating defects in the skin layer leading to lower mechanical strength and separation efficiency.

Figure 10. The mechanism of decarboxylation-induced cross-linking of copolyimide 6FDA-DABA-DAM (1:2) described by Kraochvil and Koros [113].


28


Another chemical cross-linking method being able to run at room temperatures was proposed by Chung et al. [114-117] using a diamine solution (1÷10% w/v in methanol) as cross-linking agent and where PI membranes were immersed for a period of time. Figure 11 presents a mechanism for the cross-linking agent that was believed to open the imide rings and establish inter-chain bridges by reacting with the carboxyl groups in two PI chains. The separation performance of the cross-linked membranes was sensitive to immersion time as longer immersion time produced higher cross-linking degree resulting in membranes having better plasticization-resistance, enhanced selectivity but significantly lower permeability than their precursors. The diamine cross-linking method, therefore, was generally applied for highly permeable PI with low selectivity. Several cross-linking agents successfully used for PI-based membranes were p-xylene diamine [117], m-xylene diamine [114], and ethylene diamine [116]. Compared to the diol cross-linking approach, the diamine has more advantages, in particular for hollow fibers because of its easy application: the immersion procedure in the diamine solution at room temperatures can be scaled-up easily. Although the diamine modification could enhance the selectivity, the lowchemical stability of the diamine agents (reaction with CO2 in the air or in feeding mixtures, and oxidation under sunlight), poses a concern for the longterm stability of the membranes treated with diamines. Additionally, the reaction of the diamine and the imide groups was found to be reversible [118]. When carried out at room temperature, the reaction may be incomplete resulting in a time-dependent performance of the treated membranes.

Figure 11. A proposed mechanism for the diamine-based cross-linking of PI.



29

5.4. Thermally-Rearranged Polymers Recently, a rigid polymer family derived from functionalized PI by the thermal post-treatment of membranes has been reported and designated as thermally rearranged (TR) polymers. The thermally induced rearrangement creates a new kind of free volume in TR polymers which increases the membranes permeability. TR membranes usually show outstanding separation properties as well as plasticization resistance for a wide range of gases. The TR concept was first proposed by Park and Lee groups [119] who used a post-synthesis treatment to obtain dense polybenzoxazole (PBO) and polybenzothiazole (PBT) membranes by annealing the precusor PI membranes at high temperature. To get TR polymers, the PI precusors must contain ortholinks positioned functional groups (–OH or –SH). The proposed mechanism of TR formation is described in Figure 12. Basically, thermal rearrangement is a conversion reaction in the solid membranes when the precursor membranes are annealed at temperatures higher than 350°C under inert atmospheres.

Figure 12. A proposed mechanism for the formation of TR membranes from functionalized PI.

TR polymers were shown to have very impressive separation performance, in particular for CO2/CH4 separations, with high permeability and selectivity due to their unique micropore structure obtained from the thermal rearrangement in solid-state membranes. The micropore distribution


30


of TR polymers could be further tuned by heat treatment protocols [120, 121], molecular structures of the precursor PI [31, 32, 120-122], imization routes [123], or even the membrane thickness [124]. Thermal rearrangement is believed to be a promising method to obtain polymer membranes suitable for gas separations, as the permeability-selectivity data for TR polymers were shown to surpass the Robeson’s upper bounds. Additonally, TR-based membranes show a strong plasticization resistance in mixed-gas permeation tests for CO2 partial pressure reaching 20 atm, at 35°C [125]. Data for selected TR polymers with good separation properties are presented in Table 10 indicating the polymer structure-dependant performance of TR polymers. Generally, TR polymers produced from monomers having rigid structures and bulky pendant groups tend to have high permeability. An optimization of the TR separation properties could be easily obtained by copolymerization using two or more diamines and dianhydrides to create several copolymer precursors. The TR-precursor copolymers were reported by Lee et al. [126] where an ordinary diamine and a hyroxyl-ortho diamine were used. The resulting copolymers made form two components, a polyimide (PI) and a hydroxyl-containing polyimide (HPI), the latter being transformed into TR polybenzoxazole (PBO). Variations of both diamine mole ratios will result in changes of the PBO domains in the TR copolymers. By preparing a series of copolymers, the study found a linear dependence of the TR copolymers on the PBO to PI ratio. Using the same concept, poly(benzoxazole-co-pyrrolone) TR copolymers (PBO-co-PPL) were also prepared [127]. These TR copolymers showed a good balance between permeability and selectivy which cannot be find in neat PBO and PPL. The PBO formation during the thermal rearrangement may create irregular cross-links with the PPL units yielding homogeneous membranes. As the PPL dormains are increasing, the selectivity of TR copolymers increases without significant permeability loss possibly because the PBO-PPL cross-links becomes more numerous. Although having very good separation efficiency, TR polymers have not yet been applied for large scale gas separation applications mainly due to their weak mechanical stability. Zhuang et al. [128] applied several treatments to improve the mechanical strength of TR polymers. They usued benzoxazole and benzimidazole incorporated diamines to prepare the TR precursors, while the thermal treatment was reduced to 1-2 hour at lower temperature (400°C). The obtained TR membranes exhibited tensile strengths of 71.4-113.9 MPa with elongations at break of 5.1-16.1%, but their gas separation performance was quite low.


31

Polyimide Membranes for Gas Separation Table 10. Permeability and selectivity data for selected TR-polymer flat membranes PI monomers

P (Barrer)

Diamines

Dianhydrides

H2N

αCO2/CH4 Ref.

CH4

CO2

NH2

35

1624 46

[129]

NH2

47

1591 34

[119]

46

1715 37

[119]

56

1160 20.7

[32]

34

675

20

[130]

85

1280 15

[119]

NH2

H2N DAB

HS H2N DABT

SH

CF3

F3C HO

OH

O H2N

Bis-APAF

F3C

O

CF3

NH2

O

O O

OH

O

6FDA

NH2

H2N HO HSBF

OH NH2

H2N HO

SP

OH NH2

H2N HO

O

O O

O

O

O

O

O

O HSBF

O Dianhydride


32


6. POLYIMIDE MEMBRANES (FLAT VS. HOLLOW FIBERS) As discussed above, PI is synthesized from acid anhydrides and diamines. The most common acid anhydrides are 6FDA, PMDA, BTDA, P3FDA, and BPDA, while several diamines can be used. PI, as a gas separation membrane material, can be divided in two categories: First, the commercial membranes including three common PI namely Kapton®(PMDA– ODA), P84 Polyimide (BTDA–TDI/MDI) and Matrimid® 5218 (BTDA– DAPI). The second category is composed of all the experimentally developed materials, for which a majority are based on 6FDA as the dianhydride.

6.1. Non Commercial Polyimides 6.1.1. Mixed Matrix Membranes Mixed matrix membrane represents a general approach used to improve the gas separation performance of polymer membranes. Neat polymer membranes have a permeability/selectivity trade-off and plasticization problems, whereas neat inorganic membranes easily crack because they are fragile and therefore difficult to put into a high surface area module. The idea to combine these two materials together by introducing fillers such as inorganic particles, mesoporous oxides, zeolites and MOFs into polymer matrices is expected to overcome the limitations represented by the Robeson upper bound limit and suppress the plasticization in the presence of CO2 or hydrocarbons. MMM are expected to have higher selectivity, permeability or both when compared to neat polymer membranes owing to the separation capacity of the particles. But the performances of MMM are not a simple sum of the intrinsic properties of the individual phases. Several parameters with complex interactions are controlling the performance of MMM. The transport properties of MMM are highly dependent on the nanoscale morphology of the filler/polymer interface.


Table 11. Gas separation performance for non-commercial PI flat mixed matrix membranes Fillers (loading, wt.%)

PI

Application

TiO2 Silica (SiO2) (up to 20) Silica (SiO2) (up to 20) MFI Silica (SiO2) (up to 35) TiO2 20 nm (up to 5) Cloisite1-10 Å (1, 3) Exfoliated zeolite Nu-6 (2) (up to 9.7) Exfoliated zeolite Nu-6(2) (up to 5.9) MCM-41and JDF-L1 (up to 12 and 10) Zeolite FAU/EMT (25) ZSM-2 (20)

HQDPA-DMMDA 6FDA-ODA/DAM (1/1)

H2/N2 CO2/CH4

ZSM-5 (up to 33) Zeolite L (20)

6FDA-TeMPD 6FDA-6FpDA-DABA

Test conditions (oC/bar) 35/1

6FDA-ODA/DAM (1/4)

Permeability (Barrer) (neat P to best value) PH2 = 3.81-14.1 PCO2 = 51.8-265

Selectivity (Neat S to max.)

Ref.

αH2/N2 = 167-187 αCO2/CH4 = 23.2-24.8

[86] [131]

PCO2 = 130-300

αCO2/CH4 = 24.2-34.2

6FDA-DAM

nC4/iC4

-

PnC4 = 3.7-7.8

αnC4/iC4 = 21-23

[132]

6FDA-ODA

CO2/CH4

21/6.8-13.7

PCO2 = 20.1-23

αCO2/CH4 = 44-62.6

[133]

6FDA-4MPD/DABA (4/1)

35/3.4

6FDA-4MPD/DABA (49/1)

H2/CH4 O2/N2 H2/CH4

PCO2 = 20.1-15.1 PH2 = 350-500 PO2 = 81.2-147 PH2 = 800-839

αCO2/CH4 = 44-63.2 αH2/CH4 = 30-37.9 αO2/N2 = 4.9-4.0 αH2/CH4 = 18-26.2

6FDA-4MPD/6FDA-DABA

H2/CH4

35/3.4

PH2 = 311-575

αH2/CH4 = 18.9-32

[135]

6FDA-ODA 6FDA-6FpDA-DABA

CO2/CH4 CO2/CH4 O2/N2 CO2/CH4 CO2/CH4

35/10 -

PCO2 = 16.5-40.0 PCO2 = 22.0-16.0 PO2 = 4.55-5.73 PCO2 = 1156-99.1 PCO2 = 21.1-20.1 PCO2 = 20.4-19.6 PCO2 = 19.0-18.3

αCO2/CH4 = 53.2-82 αCO2/CH4 = 30.2-24.2 αO2/N2 = 4.67-4.78 αCO2/CH4 = 18.6-20.2 αCO2/CH4 = 35.2-28.7 (4 bar) αCO2/CH4 = 40.8-39.2 (8 bar) αCO2/CH4 = 38.0-61.0 (12 bar)

[136] [137]

35/1 35/4,8,12


[134]

[138] [139]

Table 11. (Continued) Fillers (loading, wt.%)

PI

Application

ZIF-71(20)

6FDA-TeMPD

ZIF-8 (33.3)

6FDA-TeMPD

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

Test conditions (oC/bar) 35/3.5 35/3.5

6FDA-Durene (cross-linked) ZIF-8 (up to 30) ZIF-8 (up to 48) ZIF-8 (up to 40)

6FDA-Durene 6FDA-DAM 6FDA-TeMPD/DABA (9/1)

ZIF-8 (20)

6FDA-TeMPD/DABA (7/3)

ZIF-90A (15) ZIF-90B (15) Zr-BDC (25) NH2-Zr-BDC (25) Zn-BPDC (25) Cu-BTC (25) NH2-Cu-BTC (25) NH2-MIL-53(Al) (up to 15)

6FDA-DAM

Selectivity (Neat S to max.)

Ref

PCO2 = 805-2744 PO2 = 184-643 PH2 = 518-2136 PO2 = 108-450

αCO2/CH4 = 17.0-13.9 αO2/N2 = 3.35-3.18 αH2/CH4 = 17.9-15.3 αO2/N2 = 3.1-3.3 αH2/N2 = 14.9-16 αH2/CH4 = 14.4-20.3 αO2/N2 = ~8.6 αH2/N2 = ~14.1 αCO2/CH4 = 22.6-17.1 αC3H6/C3H8 = 12.4-31.0 αCO2/CH4 = 19.5-20.9 αC3H6/C3H8 = 11.7-27.4 αCO2/CH4 = 25.5-26.3 αC3H6/C3H8 = 11.7-27.5 αCO2/CH4 = 24-37 αCO2/CH4 = 24-34 αCO2/CH4 = 4.1-46.1 αCO2/CH4 = 44.1-51.6 αCO2/CH4 = 44.1-15.0 αCO2/CH4 = 44.1-51.2 αCO2/CH4 = 44.1-59.6 αH2/CH4 = 54.7-56.4 αCO2/CH4 = 35.1-36.9

[140]

PH2 = 52.1-283.5 PO2 = 2.0-5.0 CO2/CH4 C3H6/C3H8 CO2/CH4 C3H6/C3H8

25/2 35/2 35/10

CO2/CH4

35/2 35/10

6FDA-DSDA/4MPD

Permeability (Barrer) (neat P to best value)

H2/CH4 CO2/CH4

35/3

PCO2 = 1486-2185 PC3H6 = 15.7-56.2 PCO2 = 256-779 PC3H6 = 13.2-47.3 PCO2 = 158-698 PC3H6 = 10.2-42.7 PCO2 = 390-720 PCO2 = 390-590 PCO2 = 14.5-50.4 PCO2 = 14.5-13.7 PCO2 = 14.5-20.8 PCO2 = 14.5-21.8 PCO2 = 14.5-26.6 PCO2 = 57.9-66.5 PH2 = 90.1-100


[141]

[142] [98] [104]

[143] [144]

[95]


35

There are four different cases to characterize the interfacial structures between the polymer and particles. Case 1: interfacial void corresponds to the detachment of the polymer chains from the surface of the particles. This is due to the high stresses produced during the solvent evaporation step. Other possible reasons are the different thermal expansion coefficients between polymers and solid particles which may induce repulsive forces between the polymer and the fillers. This effect results in increased permeability. Case 2: rigidification corresponds to a lower mobility of the polymer chains in the vicinity or in contact with the rigid particles limiting their rearrangement. Therefore, an increase in selectivity with a large decrease in permeability is observed. A typical effect is a shift of the glass transition temperature (Tg) to a higher value. Case 3: pore blockage represents the partial or complete filling the particles’ pores by polymer chains. Pore blockage always decreases the gas permeability, while its effect on the selectivity depends on the nature of the particles used. Case 4: the ideal morphology corresponds to an homogeneous distribution of the particles with good interfacial contact. In Table 11, the data for MMM based on non-commercial PI are reported where most of the matrices are based on 6FDA. Form Table 11, 6FDA-ODA, 6FDA-Durene (TeMPD), 6FDA-DAM, 6FDA-4MPD 6FDA-TeMPD, copolymers of 6FDA-TeMPD-DAM, 6FDATeMPD/DABA, 6FDA-DSDA/4MPD were frequently used as polymer matrices, while zeolite, SiO2 and MOF as the fillers. Chen et al. [131, 136, 145-147] synthesized 6FDA-ODA, 6FDA-ODA-DAM, and combined their matrices with FAU-EMT zeolite, silica and MOF of MIL-53(AI), Zr-BDC, Cu-BTC, Zn-BPDC with amino-functionalized MOF as the fillers. The results indicated that Zr-BDC based MMM have CO2 permeability of 50 Barrer and CO2/CH4 selectivity of 46. Another interesting feature of NH2-MIL-53(Al) is a breathing phenomenon. MMM made from this “breathing” material are showing better CO2/CH4 separation factor with increasing pressure. Liu et al. [132] reported calcined-MFI and uncalcined silica MFI nanoparticles of 150 nm diameter with 6FDA-DAM MMM to separate nC4/iC4 (normal butane and isobutane). The results indicated nC4 permeability (with calcined silica) increased from 3.7 (neat 6FDA-DAM membrane) to 6.4 and 7.8 Barrer. The nC4/iC4 selectivity was almost constant (21 for 6FDA-DAM, 22 and 23 for MMM). Gorgojo et al. [134] prepared MMM from exfoliated zeolite Nu-6(2) with 6FDA-4MPD/DABA (4/1 and 49/1) copolyimide for O2/N2 and H2/CH4 separations. The authors reported that the formation of hydrogen bonds between the hydroxyl groups on the surface of the zeolite and the carboxyl


36


groups of the DABA group improved the separation performance for H2/CH4 due to the presence of zeolite [205]. The same group combined MCM-41and JDF-L1 as fillers in 6FDA-4MPD/DABA copolymer matrix for H2/CH4 separation. MCM-41 is an ordered mesoporous silica with 2-3 nm pores which is expected to increase permeability, while JDF-L1 is a layered microporous titanosilicate with a pore size of about 0.3 nm expected to improve gas selectivity. The best results indicated that the H2/CH4 selectivity was 32.0, while the H2 permeability was 440 Barrer. This has to be compared to the neat polymer having a H2 permeability of 130 Barrer and a H2/CH4 selectivity of 18.9 [135]. A cross-linkable copolymer 6FDA-DAM-DABA (3/2) using 1,3propanediol and a dispersion of zeolite SSZ-13 MMM was used for natural gas purification. The presence of SSZ-13 not only improved the CO2 permeability and CO2/CH4 selectivity, but also the resistance to CO2 plasticization up to 450 psia [148]. ZIF-8 is a commercially available MOF provided by Sigma-Aldrich, and one of the most studied ZIF. It has a structure similar to zeolites with larger pores of 11.6 Å connected through small apertures of 3.4 Å. It is thermally stable up to near 400oC. ZIF was incorporated in Matrimid, Ultem, polysulfone (PSF), polybenzimidazole (PBI), poly(1,4-phenylene ether-ether sulfone) (PEES) 6FDA-DAM, 6FDA-TeMPD [141, 142] and coplymer 6FDA-TeMPD/DABA [104]. Due to the density of ZIF-8 being below 1 g/cm3, the density of the MMM decreases with increasing ZIF-8 loading, while the Young’s modulus and Tg slightly increased. Zhang et al. [98] used ZIF-8/6FDA-DAM mixed matrix membranes for propylene/propane separations. The introduction of 200 nm diameter commercially available ZIF-8 substantially improved the C3H6 permeability, as well as C3H6/C3H8 selectivity (see Table 11). Askari and Chung [104] fabricated MMM using 6FDA-TrMPD/BABA (thermally cross-linkable at 200, 300 and 400oC) copolyimide (1/0, 7/3 and 9/1) with ZIF-8 for olefin/paraffin separation in the field of natural gas purification. The presence of ZIF-8 in 6FDA-TrMPD/BABA (9/10 polymer matrix improved the C3H6 permeability and C3H6/C3H8 selectivity. By increasing the temperature from 200 to 400oC, CO2 and C3H6 permeability increased from 158 to 698 Barrer with 20 wt.% ZIF-8, as well as C3H6/C3H8 increased from 11.7 to 27.5 Barrer. However, CO2/CH4 selectivity was unchanged for the 6FDA-TrMPD/BABA (7/30) copolymer. The authors studied the crosslinkable moiety (DABA) and annealing temperature effect on the resistance against plasticization. They observed that the permeabilities increased for any gases, while the selectivities



37

slightly decreased except for propylene/propane system with ZIF-8 in the polymer matrix compared to the neat polymer membrane. Incorporation of mesostructured oxides, zeolites and MOF in different PI as matrices was performed for several gas pairs separation, not only limited to CO2/CH4 and O2/N2, but also for hydrogen recovery and propylene/propane separations. From the data obtained, the zeolite surface needs to be modified using silane coupling agents to improve the adhesion between the polymerfiller interface. ZIF series are excellent materials for high filler loadings (up to 48 wt.%) to improve the gas permeability due to its very good compatiblity with the polymer phase without modification or compatibilizing agent. Another “breathing” MOF NH2-MIL-53(Al) was well incorporated in PI yielding the excellent advantage that it can be used for high pressure industrial application [95].

6.1.2. Hollow Fiber Membrane Hollow fiber membranes made from non-commercial PI were not often reported, due to the fact that lab-scale polyimide synthesis cannot yield large enough quantities. The asymmetric structure of hollow fiber membranes is produced by dry-wet spinning phase inversion using a manufacturing setup composed of two gear pumps, a spinning nozzle, a working tank, a coagulation bath, a steel rinsing bath (stainless isolated), a heating system, a rolling device and a filtration unit. The two metering pumps transfer precise quantities of the polymer solution dope and the bore fluid to a spinneret. Then, the polymer solution and non-solvent fluid go from the spinneret to an air-gap region for solvent evaporation. Finally, the take-up unit collects the fibers from the coagulation bath. Therefore, making one batch of hollow fiber requires at least 100 g of polymer, whereas less than 1 g is needed to make a flat membrane. Due to the limited availability of PI, most of the polymers used to make hollow fibers are commercial polymers. From the literature, there is about 10 types of 6FDA-based PI which were used to produce hollow fibers (Table 12). Koros’ group used 6FDA-DAM/DABA (4/1 and 3/2) copolymers to make hollow fiber membranes for O2/N2 and CO2/CH4 separations [113, 149]. The results indicated that 6FDA-DAM/DABA (3/2) membranes have excellent performance with CO2 permeances of 117 and 203 GPU, and CO2/CH4 selectivities of 37 and 30. When these membranes were crosslinked at 200°C for 2 h under vacuum, they reached a selectivity of 41, with a permeance of 58 GPU.


Table 12. Gas separation properties of non-commercial polyimide hollow fiber membranes Polymer 6FDA-DAM/ BABA(3:2)

Permeance (GPU) H2 CO2 36 117 58, 203

6FDA-DAM/ DABA (4:1) 6FDA-2,6-DAT 6FDA-HAB/ DAM (TR) 6FDA-APAF (TR)

105-1550

6FDA-ODA-NDA 6FDA-DAM 6FDA/BPDA-DAM 6FDA/BPDA-DDBT 120-220 6FDA-BAAF 6FDA-TeMPD-mPDA

O2

Selectivities H2/N2 O2/N2

63 22.5-28.0

4.6 4.7-5.1

100-128 19-55 162-300 97-560

16.0-24.6 2.4-9.5 15.1-28.8

4.9-5.8 3.8-4.9 4.3-5.7

31-2326 929-2471 63

251-498

498 44-75 370

7.3-22.7 9.3 158 82.8 9.5-15.6 28

10.6-30

2.2-4.2

71-81

4.7-5.2 4.82 3.9 3.93 5.4-6.4 4.6

H2/CH4

CO2/CH4 CO2/N2 34 37 41,30

40-58 55-63 59-74 11.3-131.5

14-37 36

113-140

41-45 40


28-32.5 44.6-75 17-28 8.8-20 15.9-17.8

23.7 25-27 23

Ref. [150] [151, 152] [153] [154] [155] [156] [114] [157] [158] [159] [160] [161] [162] [163] [164] [165] [166] [167]


39

Chung’s group prepared 6FDA-2,6-DAT copolymer hollow fiber membranes also to separate CO2/CH4 and O2/N2 separation. This material showed more selectivity than other PI for CO2/CH4 separation, while CO2 permeances can reach 100 GPU [23, 53, 72]. Lee’s group fabricated two PI of 6FDA-APAF [52, 87], and 6FDAHAB/DAM (TR) [88] hollow fiber membranes, with thermal rearrangement at 450oC. As expected, the gas separation performance was improved by thermal rearrangement at 450°C. Finally, three other copolymers of 6FDA/BPDA-DDBT, 6FDA-ODANDA 6FDA/BPDA-DAM, and 6FDA-DAM, 6FDA-BAAF were produced into hollow fiber membranes and showed good performance. Unfortunately, all these works were not sufficient to demonstrate the potential for industrial application.

6.2. Commercial Polyimides for Gas Separation Membranes The three commercial PI chemical structures are shown in Figure 13.

Figure 13. Chemical structure of the commercial PI: (A) Matrimid® 5218, (B) Kapton®, and (C) P84 polyimide.

6.2.1. Kapton® Kapton® is a commercial polyimide produced by DuPont. It is synthesized from PMDA (dianhydride) and ODA (diamine). It is a special material with very good chemical resistance; i.e. no organic solvents can dissolve this material. Kapton® does not melt or burn with the highest UL-94 flammability rating: V-0. The outstanding properties of Kapton® enables to work under both high (400oC) and low temperature (-269oC) where other


40


organic polymer materials would not. Kapton® HN films can be used for gas separation. From the “General Information” provided by DuPont, the permeabilitity (25 µm) of gases are as follows: PCO2 = 0.26, PO2 = 0.14, PH2 = 1.45, PN2 = 0.035 and PHe = 2.40 Barrer. One paper used a Kapton® film coated by a solution of 18 wt.% polyamic acid and 5 wt.% phenanthrene in DMAc at 70oC. These films were used for the separation of CO2 and ethanol. A CO2 permeability of 27x10-9 mol/m2.s.Pa and a CO2/ethanol selectivity of 8.7 at 100oC and 15 MPa were reported [168]. The same method was also used for CO2 and iso-octane separation with 12.8 selectivity at 150oC and 8-12 MPa [169]. Mensitieri et al. [170] studied dry and water saturated Kapton@ polyimide films of different thickness (13-50 µm) for oxygen sorption at 25oC. As mentioned above, no solvent can dissolve Kapton, so Kapton polyimide flat membrane was prepared only from polyamic acid (PAA). Sridhar et al. obtained the results of PCO2 = 8.5 and 1.5 Barrer, with a CO2/CH4 selectivity of 85 and 50.8 for single gas and mixed gas (5-2% CO2) cases at 40 bar for Kapton flat membranes, respectively. From the results obtained, the Kapton polyimide membrane was plasticized at higher feed pressure, resulting in a loss of selectivity [171]. The glass transition temperature (Tg) of Kapton® 100 HN film is between 360-410oC. The decomposition temperature (5% lost weight) is around 450oC in air, so this material is appropriate for the preparation of carbon membrane by carbonisation at high temperature. Petersen et al. used PAA solution casting on PTFE tube, imidization at 400oC, then the membrane was carbonised at 950oC. This carbon membrane showed very high selectivities (2000 for He/N2 and 1000 for H2/N2 at 0oC) [172]. Kim et al. [173] prepared carbon molecular sieve (CMS) membranes from the synthesis of BTDA-ODA, then pyrolyzed at 550, 700, and 800oC. From the single gas permeation experiments, the CMS membranes results are presented in Table 13 for 25oC and 1 atm. The results show that gas permeabilities decreased and selectivities increased with increasing soaking time and pyrolysis temperature. The CMS 800 membrane has He, CO2 and O2 permeabilities of 872, 176 and 61 Barrer, with He/N2, CO2/N2 and O2/N2 selectivities of 218, 44 and 15, respectively. Su et al. using commercially available Kapton® polyimide (100HN, Dupont), prepared a membrane with a thickness of 25 µm as the precursor for the carbonisation membranes. The best results were obtained from a carbonisation under vacuum at 1073 K yielding the lowest gas permeances but the highest ideal selectivity of 17.76 for O2/N2 [174]. The carbonisation



41

membranes treated under argon atmosphere at 1073 K have excellent selectivity of 477 for CO2/CH4. Obviously, hollow fibers and flat membranes cannot be made directly by casting or phase inverting using Kapton. The gas separation performance of the indirectly made films are not very high. However, as mentioned above, this special material it is a very good precursor for carbonisation membranes. Table 13. Gas permeation results for the pyrolytic CMS membranes as a function of pyrolysis temperature (°C) and soaking time Gas Permeability and selectivity PHe (Barrer) PCO2 (Barrer) PO2 (Barrer) PN2 (Barrer) He/N2 CO2/N2 O2/N2

CMS 500 (30 min) 2134 1535 509 56 38 27 9

CMS 700 (0 min) 2707 1989 632 72 38 28 9

CMS 700 (30 min) 1593 505 204 16 100 32 13

CMS 700 (60 min) 1108 350 136 9 123 39 15

CMS 800 (30 min) 872 176 61 4 218 44 15

6.2.2. P84 Polyimide P84 is a commercial (trade) name for a co-polymer produced by HP Polymer based on BTDA (3,3’,4,4’-benzophenone tetracarboxylic dianhydride) and PDA (80% methylphenylene-diamine), MDA (20% methylene diamine) reacted with isocyanates to produce fully imidized polyimide resins without the need for post-curing. The company Evonik also sells P84 polyimide powder or fibers. P84 glass transition temperature as determined from DSC is 315oC. Studies were prepared on the neat P84 flat and hollow fiber membrane [175], also carbon molecular sieve (CMS) membranes prepared based on P84 and P84 blended with other polymer (polyaniline) [77]. Barsema et al. [176] were the first to produce and characterize the permeation properties of dense flat and hollow fiber membranes of P84 for He, CO2, O2, and N2 at 25oC and 4 bar. The data are shown in Table 14. Table 14. Permeability (Barrer), permeance (GPU) and selectivity of P84 membranes Gases Flat membrane Hollow fiber Module

P/QH2 5.29 4.44

P/QCO2 P/QO2 0.99/1.05 0.24/0.22 2.2 0.41 0.65

P/QN2 H2/N2 0.024 106 0.047 0.056 78

CO2/N2 O2/N2 CO2/CH4 44.2 10.0 46.8 8.72 11.6 11.0


42


One patent from Air Liquide [177] reported some results on a dense film of P84 cast from a solution of 20% P84 and 80% NMP having CO2 permeability of 2.3 Barrer and CO2/CH4 selectivity of 47.1 at 20 bar and 35oC. Another patent from Praxair Technology inc. [178] described a hollow fibers module made from blending Matrimid and P84. It has an oxygen permeance of 18 GPU and a separation factor for oxygen/nitrogen of 6.5 under the tested conditions of 7 bar and 23°C for air separation. Tin et al. [175] used a P84 dense membrane carbonization process under three pyrolysis temperatures 550oC, 650oC and 800oC for CO2/CH4 separation. The permeability and ideal selectivity of CMP84-800 are 500 Barrer and 89, respectively. The developed carbon hollow fibers exhibit rather low H2 permeance values (8.2 GPU) with a highest H2/CH4 selectivity coefficient of 843 at 373 K [179]. A hollow fiber membrane from a blend of 12.5 wt.% P84 and 4.2 wt.% polyaniline (PANI) in solution was compared with another one using only 28.5 wt.% P84 in NMP. The results indicate that the blend exhibits higher gas permeability for all gases (For example PH2 from 1 to 76, and PCO2 from 0.2 to 19.6) and decreased selectivity (for example H2/N2 from 20.6 to 2.7, CO2/N2 from 3.6 to 0.7) at 25oC and 70 bar [179]. A membrane made from 15% SSZ-13 in P84 improved by 21% the CO2 permeability (from 1.8 to 1.93 Barrer) and 17% the CO2/N2 selectivity (from 25.9 to 29.7) of pure gases at 50 psi and 35oC compared to the neat membrane [180]. MMM were prepared with 20 wt.% MOF (Cu3BTC2, FeBTC and MIL53(Al)) as fillers in P84 for ethylene/ethane separation. The ethylene permeability of neat P84 and the three type of MMM were of 17, 17, 11, and 32 x10-18 mol m/(m2 s Pa), ethylene/ethane selectivity were of 4.1, 7.1, 5.0 and 3.9, under 5 bar, respectively. Therefore, the ethylene/ethane selectivity increased with the addition of Cu3BTC2 in P84 while keeping the same permeability. For the MOF of MIL-53, ethylene permeability was improved, but the selectivity was unchanged [181]. As a membrane material for gas separation, neat P84 has low permeability and high selectivity. But different modifications can be done for specific P84 application as a membrane for very high pressure conditions and also as a precursor for carbonisation at high temperature.

6.2.3. Matrimid®5218 Matrimid®5218 is currently one of the most studied and most commercially used polyimide membrane material. The Matrimid series is not as permeable as 6FDA-PI for gas permeation, but they have high gas selectivity [2]. Several studies have been published using this material to


43


prepare membranes for gas separation applications. The most recent studies focused on CO2/CH4 and O2/N2 separations, as well as for hydrogen recovery but one paper studied propane transport [182]. As discussed below, the majority of these works tried to improve the separation properties of these membranes. The literature values for single gas permeability and ideal selectivity of neat Matrimid under different conditions are summarized in Table 15. However, O2, N2 and CH4 permeability are very close even under different operation conditions. The selectivity of CO2/CH4 and O2/N2 are in the range of 34-40 and 6.6-7.6, respectively. Table 15. Permeabilities and ideal selectivities of various single gases of neat flat Matrimid®5218 membranes Test conditions (°C/bar) 35/10 22/4 25/1.4 35/2.3 35/4 50/6.5 35/2

Permeability (Barrer) H2 CO2 O2 N2

32.7 17.5 28.9 33.1 24.4

6.5 8.07 7.29 9.52 5.59 10.0 9.0

1.7 2.62 1.46 2.18 1.9

0.25 0.36 0.22 0.31 -

Selectivity CH4 C3H8 H2/N2 CO2/CH4 O2/N2 Ref.

0.19 0.23 0.21 0.24 0.34 0.15 0.36 0.25 0.22

90.8 79.5 93.2

97.6

34.0 35.2 34.8 39.5 37.3 27.8 40.9

6.6 7.2 6.8 7.1

7.6

[117] [183] [184] [182] [185] [186] [187]

Several articles deal with the preparation of flat MMM for gas separation. The added fillers are inorganic nanoparticles such as COK-12 silica [188], POSS®-Zn+2 [189], C60 [190], mesoporous structure materials as silica spheres [191], SO3H-surface functionalized MCM-41 [192] and pyrolyzed MCM-41 [193], as well as zeolite such as NaY [194], Ag+NAY [195], Sigma1 [196], Cu-BPY-HFS[197], MIL-53 [198], MIL-68 [199], MIL-125 [200], MIL-101 [201]. Vankelecom’s group [192] used ordered mesoporous silica spheres (MCM-41) and SO3H-surface functionalized MCM-41, as well as a material designated as CMS to produce MCM-41 pores with carbonaceous material obtained by the pyrolysis of furfuraldehyde in the silica pores at a temperature of 1073 K. Another type of mesostructured silica (COK-12) [188], MIL-125 [200], Cu3(BTC)2, ZIF-8 and MIL-53 [193] were also used as fillers in Matrimid®9725 MMM. The results showed that the –SO3H functionalized fillers had a better polymer-filler contact and a more homogenous dispersion throughout the polymer matrix compared to MCM-41. The CMS improved the


44


performance of MMM compared to MCM-41 with CO2 permeability increasing from 7 to 53 Barrer, while CO2/CH4 and CO2/N2 selectivities increasing from 35 to 42 and 34 to 38 respectively, at 35°C and 9 bar (see Table 16). Adding mesoporous COK-12 silica particles in the Matrimid® improved the CO2 permeability (from 7 to 15 Barrer) while keeping the CO2/CH4 (35) and CO2/N2 (34.2) selectivity constant when the filler loading varied from 10 to 30 wt.% at 10 bar and 25°C [188]. MIL-125(Ti) and NH2MIL-125 were used to make MMM based on Matrimid®9725 to enhance the performance of CO2/CH4 and CO2/N2 separation [200]. MMM with 30 wt.% loading of MIL-125(Ti) and NH2-MIL-125 have CO2 permeability of 27 and 50 Barrer respectively, but similar CO2/CH4 selectivity (37). Li et al. [189] developed nanocomposite MMM from nano-sized polyhedral oligomeric silsesquioxane POSS® Octa Amic Acid particles and Matrimid®5218. The gas separation performance of these MMM showed that all the permeability decreased while selectivity remained similar to the neat Matrimid value. Then, the membranes were treated with ZnCl2/MeOH solutions and the CO2/CH4 selectivity increased from 37 to 63 while for O2/N2 it increased from 7 to 9 after an ion binding treatment (20 wt.% POSS®– Matrimid®–0.3 M ZnCl2). Chung’s group [190] prepared MMM by physically blending a series of benzylamine-modified fullerene, C60 with neat Matrimid 5218. The results showed all gas permeability dropped with the presence of C60 molecules and the selectivity for He/N2 is increased from 87 to 106, while the selectivity for O2/N2 and CO2/CH4 was unchanged. Amooghin et al. [194, 195] prepared mixed matrix membranes with aminosilane grafted zeolite NaY and silane-modified zeolites (SM-NaY) [194] as well as NaY with Ag+ cation in Matrimid. SM-NaY enhanced the filler/polymer interfacial adhesion compared to NaY due to the silane groups of APDEMS. Therefore, the introduction NaY and SM-NaY improved the gas permeation for CO2/CH4 separation. At 15 wt.%, SM-NaY based membrane have higher CO2/CH4 selectivity (57.1) than NaY (43.3). The MMM based on NaY zeolite, having ion-exchanged Ag+ in Matrimid, had higher CO2 permeability (21.2 Barrer) and higher CO2/CH4 selectivity (60.1) compared to the other MMM. The effect of silver cation in NaY zeolite substantially improved the gas separation performance compared to neat Matrimid membranes (see details in Table 16).



45

Table 16. Gas separation performances of Matrimid based flat MMM Fillers Applications Test (loading, wt.%) conditions (oC/bar) COK-12 silica CO2/CH4 25/10 (up to 30) CO2/N2 POSS®-Zn+2 35/4 (up to 20) C60 (up to 10) CO2/CH4 23/0-25 O2/N2 He/N2 SO3H-MCMCO2/CH4 25/10 41 (up to 30) CO2/N2 Pyrolysis CO2/CH4 35/9 MCM-41 (up to CO2/N2 30) Silica spheres H2/CH4 35/1.7 CO2/N2 NaY zeolite CO2/CH4 35/2 (up to 20) Sm-NaY CO2/CH4 35/2 (up to 20) ZSM-5 H2/N2 25/1.4 (up to 30) H2/CH4 O2/N2 CO2/CH4 ZIF-8 CO2/CH4 22/4 (up to 30) CO2/N2 O2/N2 H2/N2 ZIF-8 H2/CH4 35/2.3 (up to 60) CO2/CH4 H2/C3H8 CO2/C3H8 Cu–BPY–HFS CO2/CH4 35/4 (up to 40) O2/N2 H2/N2 H2/CH4 Zn-IRMOF-1 CO2/CH4 50/6.5 (20) H2/CH4 Cu3(BTC)2 (30)

CO2/CH4 H2/CH4

50/6.5

Typical performance Permeability Selectivity (Barrer) PCO2=7-15 αCO2/CH4=35 αCO2/N4=34 PCO2=8.4-3.4 αCO2/CH4=36.4-62.8 PO2=2.1-0.99 αO2/N2=6.9-9.0 PCO2=7.1-3.4 αCO2/CH4=36-35 PO2=1.9-1.1 αO2/N2=6.8-6.8 PHe=25-17 αHe/N2=87-106 PCO2=8.4-12 αCO2/CH4=35-40 αCO2/N2=34-38 PCO2=8.4-52.6 αCO2/CH4=35-42 αCO2/N2=34-38

Ref.

PCO2=8.4-28 PH2=30-65 PCO2=8.4-17.5

αCO2/N2=26-40 αH2/CH4=132-164 αCO2/CH4=36.3-43.3

[191]

PCO2=8.4-9.7

αCO2/CH4=36.3-57.1

[194]

PH2=17.5-35.4 PCO2=8.0-15.0 PO2=1.46-2.82

αH2/N2=79.6-143 αH2/CH4=83.3-170 αO2/N2=6.6-10.4 αCO2/CH4=35-66 αCO2/CH4=35-25 αH2/CH4=143-148 αCO2/N2=22-17 αH2/N2=90-68 αCO2/CH4=35-124 αH2/CH4=143-427 αCO2/C3H8=28-50 αH2/C3H8=80-180 αO2/N2=6.6-6.3 αCO2/CH4=35-25 αH2/N2=79.6-54.8 αH2/CH4=83.3-45.4 αCO2/CH4=28.2-29.2 αH2/CH4=93.2-86.4

[197]

PCO2=8.8-29 PH2=33-112 PO2=2.6-10 PCO2=9.5-24.6 PH2=28.9-71.2 PO2=2.2-5.9 PH2=17.5-26.7 PCO2=7.3-15 PO2=1.5-3.1 PCO2=10-22.1 PH2=33.1-114 PCO2=10.022.1 PH2=33.1-66.9

αCO2/CH4=28.2-29.8 αH2/CH4=93.2-90.3


[188] [189] [201]

[192] [193]

[194]

[183]

[182]

[197]

[186]

46

Xiao Yuan Chen, Nguyen Tien-Binh, Serge Kaliaguine et al. Table 16. (Continued)

Fillers Applications Test (loading, wt.%) conditions (oC/bar) MOF-5 CO2/CH4 35/2 (up to 30) O2/N2 H2/N2 H2/CH4 Ag+ NaY CO2/CH4 35/2 zeolite (up to 20) MIL-101 CO2/CH4 35/10 (up to 30) CO2/N2 MIL-53(Al) CO2/CH4 25/3 (up to 25) MIL-125(Ti) CO2/CH4 35/9 (up to 30) CO2/N2 NH2-MIL125(Ti) (up to 30) MIL-68(Al) CO2/CH4 35/4 (10) Sigma-1 CO2/CH4 50/10 zeolite (up to 20) Sod-ZMOF CO2/CH4 35/4 (up to 20)

Typical performance Ref. Permeability Selectivity (Barrer) PH2=24.4-53.8 αO2/N2=7.6-7.9 [187] PCO2=9.0-20.2 αCO2/CH4=41.7-44.7 PO2=1.9-4.1 αH2/N2=97.6-103.5 αH2/CH4=113-120 PCO2=8.34-21.2 αCO2/CH4=36.2-60.1 [190]

PCO2=4.4-7.99

αCO2/CH4=35-56 αCO2/CH4=34-52 αCO2/CH4=35-38

[200]

αCO2/CH4=30-37 αCO2/N2=23-35 αCO2/CH4=30-50 αCO2/N2=23-35

[200]

PCO2=~176.6

αCO2/CH4=~-55.2

[199]

PCO2=4.5-9.6

αCO2/CH4=34-50

[196]

PCO2=5.6-12.2

αCO2/CH4=37.3-32.9 [185]

PCO2=9.0-15.0 PCO2=6-27 PCO2=6-50

[198]

Gheimasi et al. [196] prepared Sigma-1 zeolite loaded Matrimid MMM for CO2/CH4 separation. The result showed that CO2/CH4 selectivity were greatly improved by the filler particles compared to neat Matrimid membranes (from 34 to 64.1). However, the CO2 permeabilities of the MMM did not increase with increasing Sigma-1 loading. The authors explained this phenomenon by partial blockage of the zeolite pores by the polymer. ZSM-5 (Zeolite Socony Mobil–5) is an aluminosilicate zeolite belonging to the pentasil family of zeolites with uniform cubic shaped particles of 200 nm. It may be prepared with mesoporous structure composed of micropores (0.65 nm) and mesopores (2.7 nm). When ZSM-5 was mixed with Matrimid, the polymer chains were able to penetrate into the mesopores of ZSM-5, decreasing the possibility of non-selective macro-void between the zeolite and



47

Matrimid. The micropores of ZSM-5 can provide size and shape selectivity leading to improved ideal selectivities of these Matrimid based MMM [184]. ZIF-8 is a zeolitic imidazolate framework having a sodalite topology with a pore cavity of 11.6 Å and a theoretical pore aperture of 3.4 Å. Two papers showed that for MMM based on Matrimid 5218 and ZIF-8, all the gases permeabilities increased with increasing ZIF-8 loading up to 30 or 60 wt.%, but for the selectivities different results were observed. Song’s work showed that the selectivities decreased compared to the neat Matrimid with the exception of H2/CH4 [182, 183]. On the other hand, Ordoñez et al. [182] found that all the gas selectivities were improved with the introduction of ZIF-8. MMM were also prepared from microporous MOF of Cu–BPY–HFS crystals incorporated into Matrimid® [184]. The results showed that the introduction of this MOF allowed to improve all the gas permeabilities, but decreased their selectivities. The author considered that Cu–BPY–HFS had strong affinity towards CH4 favoring the permeation of CH4. As a result, the CH4/N2 selectivity increased from 0.95 to 1.21, while the CO2/CH4 and H2/CH4 selectivities decreased from 35 and 83 to 25 and 45, respectively. The data collected for Matrimid based MMM on CO2/CH4 separation are plotted in a standard representation in Figure.14: ideal selectivity as a function of CO2 permeability. The CO2 permeabilities are mostly in the 6-10 Barrer range, while the CO2/CH4 selectivities are between 34 and 40 for neat Matrimid membranes. With the introduction of fillers in MMM, the CO2 permeability can reach 40 (Zn-IRMOF-1) and 50 (MCM-41 and MIL) Barrer, which represents increases of 300-400% compared to the neat polymer membranes, with similar selectivity values. Therefore, the production of MMM is effective to improve the separation performance. On the other hand, CO2/CH4 selectivities can improve up to 60 (Poss-Zn, Sigma and ZSM-5) and 80-124 (ZIF-8) without loss in CO2 permeability (5-6 Barrer). One paper introduced a crosslinking agent in Matrimid flat membrane to suppress the plasticization phenomenon. Tin et al. [117] modified a flat Matrimid®5218 membrane by immersion in a 10% (w/v) of pxylenediamine methanol solution. The results showed that chemical crosslinking modification appeared to be an effective approach in suppressing the plasticization phenomenon, but did not improve the CO2/CH4 and O2/N2 separation performance. David et al. also studied the plasticization effect in flat [202] and hollow fiber [203] Matrimid®5218 membrane for hydrogen recovery from a postcombustion gas mixture (H2, CO, N2 and CO2). A small content of CO2 in the mixture had different effect for flat and hollow fiber membranes. For flat


48


membranes, adding 10% CO2 decreased the hydrogen permeability (42%) and H2/CO2 selectivity from 4.2 to 2.7. On the other hand, the hollow fiber membranes showed a better suppression of plasticization, but hydrogen permeance only decreased by 21% and H2/CO2 selectivity reached 3.1 for a H2/CO2 composition of 20/80 at 10 bar. For hollow fiber membranes, new preparation and post-treatments (subambient, high temperature or carbonisation) techniques were developed. Blends with other polymers such as PES [204-206], PSF [207], PBI [78], PIM1 [208], PEG and PEO-PDMS [209] were also studied. The asymmetric configuration for hollow fiber modules was preferred at an industrial level due to a more effective surface area per unit volume compared to other configurations. The separation performances for Matrimid hollow fiber membranes with six gases are summarized in Table 17. For example, Carruthers et al. [212] analyzed how to form defect-free, slightly defective, moderately defective and severely defective skin layers inside hollow fiber membranes. For defect-free membranes, the dense skin thickness was estimated from pure permeation of He, O2 and N2 and the results were in agreement with SEM images. 140 Neat Matrimid COK-12

CO2/CH4 selectivity

120

PossR -Zn+2 C60 SO3H-MCM-41, CMS

100

NaY zeolite Sigma-1 ZSM-5 ZIF-8 Cu-BPY-HFS Zn-IRMOF-1 MOF-5 MIL-53, 110, 125 (Ti) Cu3(BTC)2

80

60

40

20 0

10

20

30

40

50

CO2 permeability (Barrer)

Figure 14. Matrimid® based MMM performance for CO2/CH4 separation.


60

49

Polyimide Membranes for Gas Separation Table 17. Gas separation performance of Matrimid®5218 hollow fiber membranes Test conditions (oC/bar) 25/6.5 20/15 25/3.4-6.5 35/3.48 25/1 35/10 23/6

Permeance (GPU) H2

27.6

He CO2

O2

N2

122

8.4 8.3 6.08 79.5 16.9

1.1 1.6 0.90 13.8 2.6 0.28 0.9

11.8 28.5 421 86.3 7.0 20.9

4.0

CH4

He/N2

CO2/ CH4

110 0.35

34

2.5 0.21 97.0

34.5 33.3

O2/ N2

CO2/ Ref. N2

7.6 5.2 6.77 5.76 6.5

18 31.7 30.5 33.2

4.6

24.5

[15] [210] [211] [143] [208] [78] [209]

Visser et al. [213] prepared five type of hollow fiber membranes (pure Matrimid®5218, Matrimid/polyethersulfone 20/80, Matrimid/P84 50/50, cellulose acetate and PPO) to study the effect of plasticization and competitive sorption in CO2 and CH4 mixtures. The results showed that the blend of Matrimid and P84 can suppress plasticization while having higher CO2/CH4 selectivity (45) for the gas mixture (CO2/CH4 80/20) at 35oC and 16 bar without any modification. The selectivity of neat Matrimid blends with PES, CA and PPO were 32, 28, 10 and 10, respectively under the same conditions. Lee et al. [214] studied the plasticization phenomenon in presence of toluene and n-heptane as contaminants in CO2 and CH4 mixture for Matrimid hollow fiber. The results showed that the CO2 permeance and CO2/CH4 selectivity were reduced during exposure to toluene, n-heptane or their mixtures at 35oC and 50 to 400 psia with a 10/90 CO2/CH4 feed gas. For the annealed (220oC for 12-16 h under vacuum) hollow fiber membranes, the effect on gas separation performances was not clear. Liu et al. [132, 211] proposed a nodular selective skin layer for Matrimid hollow fiber, which improved the CO2 permeance while keeping the CO2/N2 selectivity almost the same as the dense-skin hollow fibers. In particular, these membranes displayed very high CO2/N2 selectivity (90.5) and CO2 permeance (63.3 GPU) at -20oC (the dense-skin values are 52.2 for selectivity and 16.6 GPU for permeance under the same conditions). They also worked on these nodular selective layer membranes by slightly increasing the quenching temperature (from 44.2 to 46.7oC) and decreasing the air gap (from 10 to 5 cm). The results (Table 18) showed that these nodular structures improved the gas permeance (CO2 and O2) up to 421 and 79.5 GPU, while the selectivity of CO2/N2 and O2/N2 were 30.5 and 5.8, respectively. The separation


50


performance was excellent with a CO2 permeance of 227 GPU and a CO2/N2 selectivity of 209 at -50oC. Kapantaidakis et al. [74] blended PES/PI (80/20, 50/50, 20/80) to improve the CO2/N2 separation performance. The results showed that PES/Matrimid (80/20) has a CO2 permeance of 60 GPU and a CO2/N2 selectivity of 40. Polybenzimidazole (PBI) is a rigid polymer with high thermal stability having a glass transition temperature of 420oC. PBI has lower permeability and higher selectivity compared to Matrimid. It is thus expected that a blend of Matrimid and PBI could improve the selectivity of H2/N2, CO2/CH4 and H2/CO2 and could also be used at higher temperature which would be useful for hydrogen recovery [78]. PIM-1 is a polymer of intrinsic microporosity (PIM-1) with very high permeability. Yong et al. [208] prepared defect-free hollow fiber based on PIM-1 and Matrimid and the results showed highly improved gas permeance (QCO2 from 86.3 to 153.4 GPU and 212.4 GPU) while slightly decreasing the selectivities. Li et al. [206] and Jiang et al. [205] both prepared two-layer hollow fiber membranes with Matrimid as the outer layer material and PES as the inner support material for gas separation. Ding et al. [207] also used Matrimid for the outer layer and PSF as the inner layer. This configuration was found to reduce the material costs because a very thin layer of the most expensive polymer was used. Since ethylene oxide units in poly(ethylene oxide) (PEO) and poly(ethylene glycol) (PEG) have a strong affinity towards CO2, it is expected that PEO and PEG can improve CO2 permeability and CO2/N2 selectivity. PDMS is also a polymer with high permeability. Hu et al. [209] added PEG and PEO-PDMS co-polymer in Matrimid hollow fiber membranes to improve on the gas separation performance. The results showed that CO2 permeance increased slightly while CO2/N2 selectivity improved from 24.5 to 50 and 56 with PEG 4% and PEO-PDMS 8%. Vu et al. [215, 216] prepared a carbon molecular sieves (CMS) flat membrane by pyrolysis of a Matrimid®5218 precursor at 800oC. This CMS flat membrane showed excellent performance with a CO2/CH4 selectivity of 200 and a CO2 permeability of 44 Barrer, but the O2/N2 selectivity (13.3) and O2 permeability (24 Barrer) at 35oC and 50 psia were much lower. They also prepared this type of CMS from Matrimid 5218 hollow fiber (pyrolysis temperature of 600oC) for the separation of CO2/CH4 (10/90) at high pressure (1000 psi). The results indicated that the CO2 permeance of CMS decreased to 11-13 Barrer for the Matrimid precursor fiber of 25-35, with an increase of CO2/CH4 selectivity from 35-40 to 69-83 at 24°C and a pressure range of 50200 psia (shell-side feed).


51

Polyimide Membranes for Gas Separation Table 18. A comparison of Matrimid blends with other polymers to produce hollow fiber membranes for gas separation Test Permeance (GPU) conditions H2 CO2 O2 N2 (oC/bar)

Polymer composition PES/Matrimid (20:80) PES/Matrimid (80:20) PBI/Matrimid (50:50) PIM-1/M(5/95) PIM-1/M(10/90) PIM-1/M(15/85) PEOPDMS(8%)+M PEG(4%)+M

CO2/ CH4 CH4

O2/ N2

CO2/ Ref. N2

25/3.9

40

1.0

40

25/3.9

60

1.53

39

0.07 0.05 48.0

30.9 [78]

35/10, 4

13.1 2.16

25/1 25/1 25/1

172.9 39.5 7.3 210.2 52.9 9.1 234.6 57.1 9.2

23/6

22.3

23/6

20

2.5

0.4

7.7 8.5 7.3

22.4 23.1 32.1

[204]

5.4 5.8 6.2

23.8 24.8 [208] 25.5

6.0

52.6

0.4

[215]

50

Among these three commercial PI, Kapton® and P84 membrane have small permeability, but they have high glass transition temperature of 400 and 315oC, respectively. Both materials can be used for carbonisation producing membranes with excellent performance for gas separation. Matrimid have some advantages such as being easily processed and blended with other polymers, even good interaction with different types of fillers. The development of hollow fiber preparation and post-treatment technologies are also very effective to improve the separation capacity of neat Matrimid hollow fiber membranes as they kept their selectivity which is important for industrial applications. Finally, Matrimid® based mixed matrix membranes can show improved gas separation properties compared to neat flat membranes.

CONCLUSION In this chapter, a review of several aspects related to polyimide applications in gas phase separations was presented. In particular, the focus was made on the synthesis and the production of hollow fibers and flat membranes. Several modifications were also presented to improve on the processability and the gas transport properties of the neat polymer. Nevertheless, several options are available to further improve the separation


52


performances of the matrix. To this end, the results were presented and discussed for polymer blending and the production of mixed matrix membranes (MMM). Other possibilities included post-treatment (chemical, physical and thermal treatment) as well as cross-linking, especially to get over some limitations such as the plasticization effect and the maximum operation pressure/temperature. Although a great deal of literature can be found on the subject for single gas and gas mixtures, more work is still being conducted to further improve on these properties and to develop materials up to the industrial scale.

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BIOGRAPHICAL SKETCHES Dr. Xiao Yuan Chen is a postdoctoral fellow at Université Laval. She is currently working on the processing and properties of polymer membranes for gas separation. Nguyen Tien-Binh is a Ph.D. student in chemical engineering at Université Laval. His main research areas are related to the synthesis and the modification of polymers for gas separation applications. Prof. Serge Kaliaguine is a professor of chemical engineering at Université Laval. For over 40 years, he has developed systems for reaction engineering and catalysis of chemical reaction. Prof. Denis Rodrigue is a professor of chemical engineering at Université Laval. For over 20 years, he has developed and optimized polymer processes for complex materials.





Chapter 2

CHARGE BEHAVIORS ON DIRECT-FLUORINATED POLYIMIDE FILMS AND POLYIMIDE COMPOSITES Boxue Du*, PhD School of Electrical Engineering and Automation Tianjin University, Tianjin, China

ABSTRACT Polyimide(PI) film, as a special type of engineering plastic film, is a kind of basic insulating material and is widely applied in the aerial, nuclear, microelectronic industry, turn to turn insulation and turn to ground insulation of inverter-fed motors. However, PI insulation encounters some serious problems in practice. The existence of surface charge and space charge has a great effect on breakdown characteristic and is the main reason leading to dielectric breakdown. Fluorination as change the chemical component in surface layer of polymers should give rise to the corresponding change in electrical properties of the surface layer thus influence the charge injection from electrodes when they are used as an insulator. Besides, the addition of nanoparticles into PI can improve the insulating properties compared with pure material. This chapter presents a study aimed at clarifying the effect of fluorination time on surface charge and space charge behaviors of fluorinated PI film and PI/Al2O3 composites. Obtained results show the dependence of the charge *

Corresponding Author: [email protected].


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Boxue Du density as well as the charge decay rate upon the fluorination time of samples, varying as a function of the charge polarity and charging time. It is suggested that the fluorination can significantly improve the decay rate of the surface charge in PI films and the injection of space charge is also suppressed. In the study of PI/ Al2O3 composites, it is concluded that the fluorination can significantly improve the decay rate of the surface charge in PI films.

1. INTRODUCTION In recent years, new electronic devices, such as cell-phone, digital camera and computer, have been quickly developed and produced. The flexible PrintCircuit Board was widely used in the electronic devices in order to reduce the size of devices. Polyimide (PI) that serves as a typical kind of engineering polymer material with advantages of high temperature resistance (400oC), low temperature tolerance (-269oC), radiation resistance, flexibility and excellent dielectric properties, has been used for the Print-Circuit Board [1-2]. For the flexible using, the size and weight of the electronic devices have been reduced every year, so it is necessary to reduce the thickness of the Print-Circuit Board used for electronic devices. Multilayer polyimide films have been used in order to reduce the thickness of the combination of the Print-Circuit Board. It is important to study the insulating property of the multilayer polyimide films. Many papers have investigated the insulation aging and failure mechanism of polyimide films. Kaufhold et al. has reported that the partial discharge was the main cause for insulation failure [3-4]. Bellomo et al. has found that the life of insulation materials used in variable-frequency motors was mainly affected by the partial discharge. Based on the partial discharge theory, some researchers thought that space charge played an important role in the insulation breakdown of polymer materials [5-6]. Fabiani et al. has investigated the degradation mechanisms occurring in magnet wires fed by power electronic waveforms through partial discharge and space charge measurements, as well as life tests [7]. Zhou et al. have studied the charge transport mechanism and space charge characteristics of polyimide films and reported that the introduction of nano-particles raised the number of traps in the dielectric, thereby forming stable space charge electric field, which effectively increases dielectric properties of corona-resistant film [8]. Kimura investigated the surface discharge of polyimide file and reported that surface charge of polyimide had an effect on the particle discharge [9]. However, little work has been reported on the chemical modification of


Charge Behaviors on Direct-Fluorinated Polyimide Films …

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polyimide nanocomposite film surface layer to modulate its electrical properties and so as to suppress the surface and space charge accumulation. Direct fluorination using fluorine gas, as one of the most effective approaches to the chemical modification of polymers, is widely developed from fundamental researches to industrial applications [10-11]. It has many advantages, of which it is worthwhile to be emphasized that the outstanding surface properties similar to the fluoropolymers can be obtained without changing the bulk characteristics of the starting polymer. Actually, the change in chemical component in the polymers surface layer by fluorination should give rise to the corresponding change in electrical properties of the surface layer, thus influence the charge injection from electrodes when they are used as an insulator. Therefore, this chapter will focus on understanding the effect of direct fluorination on dynamic behaviors of surface charge of polyimide film and dynamic behaviors of space charge in multilayered polyimide films. Furthermore, surface charge accumulation and charge decay of polyimide nanocomposite film will be discussed in the last part.

2. DYNAMIC BEHAVIOR OF SURFACE CHARGE ON DIRECT-FLUORINATED POLYIMIDE FILMS 2.1. Experiments The samples were prepared using commercially available pyromellitic dianhydride (PMDA) and 4, 4' -oxy dianiline (ODA). The polyimide film with the dimension 40 mm × 40 mm × 25 μm was applied in this experiment. Surface fluorination of the sheets was performed at about 328K (55℃) using a F2/N2 mixture with 20% F2 by volume. The duration of the fluorination was respectively 15, 30, 45 and 60min. After the treatment, the reactive gas mixture was purged from the vessel with nitrogen. Surface charge accumulation test was performed at room temperature with relative humidity of ~ 40%. The voltage of ±6 kV was applied between the needle and the grounded electrodes. The charging time was respectively 10, 20 and 30 minutes.


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2.2. Effect on Chemical Composition and Permittivity Surface chemical compositions of the fluorinated films together with the original (unfluorinated) film as a reference, were characterized by attenuated total reflection infrared (ATR-IR) analysis. It is well known that direct fluorination of polymer results in the disruption of C-H bond followed by a formation of C-F, C-F2 and C-F3 groups, as C-F bond is the strongest single bond in organic chemistry and has bond dissociation energy of 544 kJ/mol [12], much higher than the C-H bond (414 kJ/mol) [13]. Fluorination degree depends on the treatment conditions (pressure of the reactive mixture, fluorine partial pressure, treatment time and temperature) and polymer nature. Figure 1 shows the ATR-IR spectra of the original and surface fluorinated polyimide films. In the spectrum (a) of the original film, it can be found that the aromatic imide groups are evidenced by the absorption peak of C = O bonds at 1780, 1720, 720 cm-1 and the absorption peak of C-N bonds at 1373 cm-1 [14-15]. The absorptions at 1500, 1100, and 810 cm-1 are related to the stretching vibration of the conjugated C = C bonds in aromatic rings, the in-plane deformation vibration of the aromatic rings, and the in-plane deformation of phenyl−H, respectively [16]. It also can be observed that the characteristic absorption peaks of C-O-C bonds in the ODA at 1273 cm-1, which indicates that they were not completely consumed by the reaction. The spectrum (b) of fluorinated film shows that C = C, C = O, C−O−C and phenyl−H bonds are significantly reduced in strength or disappear by the fluorination, and simultaneously a very broad strong absorption band appears over the 950−1340 cm-1 region, which is ascribed to C−F, CF2 and CF3 groups. The results show clear evidence that fluorine is incorporated into the polyimide surface layer and that the C = C double bonds are saturated by the direct fluorination. Figure 2 shows the relationship between the relative permittivity and the fluorination time of polyimide film varying with the frequency at 20 oC. Almost all of the samples exhibit the same behavior except for different values under different frequencies. It is found that the dielectric permittivity decreases with increasing the fluorination time from 0 min to 45 min and increases from 45 min to 60 min. Simultaneously, it is noticed that the dielectric permittivity of the original sample is larger than other samples independent of the frequency. It can be assumed that the 2s and 2p orbit of fluorine atom is very close to the nucleus, which makes the fluorine element the smallest atom in all chemical elements except hydrogen atom [17]. So the electron polarizability of F element is very low, and the fluoroorganic has low



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dielectric permittivity. It also can be found that the relative permittivity of sample with 60 minutes fluorination time is higher than that with 45 minutes, which may be due to some chain scission of aromatic imide groups in polyimide occurs during the fluorination.

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Figure 2. Relationship between the relative permittivity and the frequency.


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2.3. Effect of Charging Method Surface charge accumulation is initiated from the corona discharge. Charges with the identical sign of applied voltage are encouraged by the electric field and move towards the upper surface of the sample in the air. Due to the presence of grid electrode, surface potential can be controlled so that all samples retain approximately the identical potential at the beginning of the decay test. The charge decay curves are shown in Figure 3. The decay of surface charge is fast at the initial period, whilst becomes slower with the lapse of time. The decay behavior of negative charge is similar to that of positive charge. The possible routes for charge decay are lateral charge spreading, transport through the bulk and recombination with ions of opposite sign in air. Considering the electrode arrangement and sample dimension, electric field generated between implanted charge and grounded electrode is higher than the field along the surface, and hence bulk transportation should play a dominant role in the decay. A few of charge is firstly injected in the electron hole of sample surface and then transferred through the conduction band of the samples as the high electric field. Most charge was accumulation in the surface of samples. The decay curves for samples charged under different time show distinguished features. Figures 3a and 3b show the variations in surface charge density at different charging time with original sample, it can be found that with original sample, the decay curves have same decrease tendency for both positive and negative charged. With increasing the charging time, the surface charge density increases. Figures 3c and 3d show that as the samples positive and negative charged for 20 minutes, the surface charge density are higher than that for 10 minutes, and have the same decay tendency. While as the sample charged for 30 minutes, the surface charge density is higher than that for 10 and 20 minutes at the beginning, but it decays faster than that for 10 and 20 minutes with the sample positive and negative charged. It is mainly because that the fluorinated surface layer existing in the fluorinated samples can significantly suppress the charge injecting in the samples. With increasing the charging time, in the force of electric field more charge accumulate in the sample surface, which make the initial surface charge increase with the increasing charging time. Once the source is turned off, the more surface charge accumulation in the sample surface decay fast with the electrostatic force of charge in the surface of samples. Although the surface charge density of the 30 minutes charged sample is larger than that for 10 and 20 minutes, more surface charge accumulated in the same fluorinated surface layer with that charged for 10 and



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20 minutes, the surface charge decay faster than that charged for 10 and 20 minutes. It can be found in Figure 3 that the decay curves flatten as the charging time is increased. This is primarily because the trapping for injecting charges is enhanced by increased charging time, and the implanted charge in turn inhibits charge transfer from surface to bulk, thereby prolonging the decay [18]. It is also observed that the charge density is general higher for negative polarity than for positive polarity in the same condition. This phenomenon reveals that more negative charges are accumulated on the surface than positive charges. It can be interpreted that the de-trapping process of a positive charge is easier than that of a negative charge.

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Figure 3. Relationship between the surface density and the decay time, (a) and (b) were original samples; (c) and (d) were fluorinate for 45 min.

2.4. Effect of Fluorination Time Figure 4 shows the relationship between initial surface charge density after charging process and the fluorination time. It is found that with the increasing of fluorination time from 0 minute to 45 minute, the initial surface charge density decreases and increases from 45 minute to 60 minute. It also can be observed that the initial surface charge density with sample negative charged is larger than that positive charged, mainly because it is easier to induce negative corona than positive corona at an identical absolute value of DC voltage [19-20]. The relationship between the decay rate and the



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Initial surface charge density (pC/mm 2)

fluorination time is shown in Figure 5. It can be found that the decay rate of fluorinated samples is much larger than that original sample in the same conditions, and the largest decay rate is about 90% appearing at the fluorination time 30 min with both positive and negative charged for thirty minutes. With the original sample positive charged for 10, 20 and 30 minutes, the decay rate is about 25% in decay time of 5 minute. With the original sample negative charged for 10 and 20 minutes, the decay rate is also about 25%, while is 40% when the sample was negative charged for 30 minutes. 3500

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Figure 4. Relationship between the initial surface charge density and the fluorination time after the charging process.

The results indicate that direct fluorination of polyimide film can significantly suppress surface charge accumulation and improve the decay rate, as shown in Figures 4 and 5. It is mainly because that the fluorine element has the largest electronegativity and makes the C-F bond highly polarized, which can form a fluorinated layer by surface fluorination in the samples. And it can be assumed that the fluorinate layer in the sample would capture electrons and form a shielding layer, which also shallows the deep traps. The trap distributions are shown in Figure 6, which was calculated according to the mathematical treatment in section 3, it is noted that the profiles of the distributions are similar, and the minimum value in trap depth appears at ~0.76 eV with fluorination time 45 min. A low energy level represents shallow traps in the sample that cannot obstruct the mobility of the trapped charge [21]. As


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charges are injected into the bulk only after the traps are filled [22], the surface potential is increased after the traps are filled. As the shallow traps, the charges easily move out from the traps, as a result of which the surface charge of fluorinated samples decay faster than that original sample [23]. 100

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Figure 5. Relationship between the decay rate and the fluorination time after the decay for 5 minute.



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Energy below conduction band (eV) Figure 6. Trap distribution at various fluorination times after applying a negative charge for 20 min.

2.5. Dissipation Time The dissipation time is defined as the period during which the total number of charges returns to 10% of the original charged value [24]. A long dissipation time indicates slow corresponding decay process. Figure 7a presents the changes of dissipation time. It can be found that the dissipation time increases with increasing the charging time. The dissipation time of a sample positive charged for ten minutes is 9.6 min, which is greater than that of 7.7 min for a negative charged sample; and the dissipation time of a sample negative charged for thirty minutes is 11.4 min, which is less than that of 9 min for a positive charged sample.


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Figure 7. The dissipation time in different conditions.

As shown in Figure 7b, the dissipation time decreases as the fluorination time is increased from 0 min to 45 min, while increases when the fluorination time is increased from 45 min to 60 min. In addition, it is found that the dissipation time for a negative charge is longer than that of a positive charge with the original sample, while the dissipation time of fluorination samples for a negative charge is shorter than that of a positive charge. It also can be observed that the dissipation time of original sample is much longer than that of fluorinated samples; even the longest dissipation time of the fluorination samples is less than 20 min, while the dissipation time of original sample is longer than 80 min. The results reveal that the fluorination can significantly improve the decay rate of the surface charge in polyimide films.



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3. EFFECT OF SURFACE FLUORINATION ON SPACE CHARGE BEHAVIOR IN MULTILAYERED POLYIMIDE FILMS 3.1. Experiments In this research, the duration of the fluorination was respectively 15, 30, 45 and 60 min. The multilayered polyimide films were prepared by stacking five films in layers with lubricant at the interfaces. We apply the pulsed electro-acoustic (PEA) method to get the space charge distributions in multilayered polyimide films and the applied voltage was 5000 V and the stressing time for each voltage is 60 min.

3.2. Effect on Space Charge Behavior We name the untreated sample and the samples fluorinated for 15, 30, 45, 60 minutes as PI0, PI15, PI30, PI45 and PI60 respectively. Space charge distributions of multilayered polyimide films are shown in Figure 8. Figure 8.a shows the space charge behavior of multilayered samples by stacking five PI0 films. It is observed that with the charging time 10 seconds, significant amount of homocharges accumulated at or near the cathode and anode. With the charging time for 1800 and 3600 seconds, a significant amount of charges accumulated at the interface near to the cathode followed by successive layers of negative, positive, negative and positive charges respectively, and then ending with a small positive homocharge near the anode. The maximum charge density in this case is 0.3 C/m3 at the interface near the cathode. With the charging time increasing, the maximum charge density increases. Figure 8.b shows the space charge behavior of multilayered sample by stacking five PI30 films. It can be found that the space charge distribution of multilayered film stacking with five fluorinated samples is quite different from that of multilayered film stacking with five untreated samples. The space charge behaviors of multilayered fluorinated films have the same tendency with the charging time increasing from 10 to 3600 s. A large amount of positive charges accumulate in the film near the anode and then followed by a significant amount of negative charges near the interface which is close to the anode. There is no positive charge accumulated in the multilayered sample except the film near anode.


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Figure 8. (Continued)


Charge Behaviors on Direct-Fluorinated Polyimide Films … 0.2

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180 200 220 240 260 275 Thickness (μm) (c) multilayered polyimide films by stacking five samples with fluorination time 60 minutes

Figure 8. Space charge distribution of multilayered polyimide films under dc 40 kV/mm.

The maximum charge density in this case is 0.21 C/m3 and with the increasing of charging time, the maximum increases. Figure 8.c shows the space charge behavior of multilayered sample by stacking five PI60 films. It is found that the space charge behavior of multilayered PI60 films has the similar tendency with the space charge behavior of multilayered PI30 films as shown in Figure 8.b. The negative charge accumulated in the sample by stacking five PI60 films was more close to the interface near the anode than that by stacking five PI30 films. Also the space charge behaviors of multilayered samples by stacking five PI15 and stacking five PI45 films, which are not shown in this research, are similar to the space charge behavior shown in Figures 8.b and 8.c. The space charge behaviors of multilayered samples by stacking fluorinated samples are obviously different from the behavior of multilayered sample by stacking untreated sample. According to MAXWELL Theory, the discontinuity of permittivity/conductivity ratio can induce interfacial charges at the interface, named MAXWELL-Wagner charges. As in this case, five treated films with the same permittivity/conductivity were stacked to form multilayered films, so the effect of permittivity/conductivity ratio on interface


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charges is very negligible. In this experiment, the charges generated from ionization of impurities in bulk can be neglected and the charges are mostly homocharges injected from electrodes. In common cases like the experimental conditions in this study, ion injection from electrode is impossible. The positive and negative charges injected to the samples are holes and electrons. It is also known that, charge transportation in material is influenced by traps depth, trapping characteristics and the conductivity. Our previous research has shown that the fluorination can not only decrease surface energy for polyimide film but also change the electrical properties, such as dielectric property, work function and the trap depth. We had test the conductivity of fluorinated film and the results tend to be higher than that of the untreated film. It is assumed that after the positive voltage was applied, the negative charges were injected by cathode and gradually accumulated in the untreated film as shown in Figure 8.a. However, for the fluorinated film, which has shallower trap and higher conductivity than untreated sample, the injected negative charges moved deeper into the multilayered samples and accumulated near the anode as the homocharges were accumulated in the film at or close to the anode. ×107

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Figure 9. (Continued)


Charge Behaviors on Direct-Fluorinated Polyimide Films … ×107


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Figure 9. Electrical field evolution in multilayered polyimide films under dc 40 kV/mm.


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The formation of space charge would affect the electrical field distribution. The electrical field distribution in the multilayered samples is shown in Figure 9. It is found that after voltage application, the electrical field built up quickly. As shown in Figure 9.a, for the multilayered samples by stacking five untreated samples, the electrical field is intense at the beginning of voltage application, because the charge injected from the electrodes and then accumulated near the electrodes at the very beginning. Then with the space charge moved and accumulated, the electrical field distorted gradually. However, with the multilayered samples by stacking five fluorinated films, the electrical field has the same distribution tendency from the beginning of the voltage application to the ending. It also can be found that the electric field of multilayered samples by stacking five fluorinated films is lower than the untreated films.

3.3. Effect on Space Charge Injection It can be found in Figure 8. that the maximum charge density of the multilayered sample by stacking five fluorinated samples is much smaller than that of the samples by stacking five untreated sample. The results indicate that the fluorination has effect on the space charge injection into the samples. The amount of space charge can be calculated by Q     x  Sdx d1

d2

(1)

where S is the effective sensor area, d1 and d2 are the beginning and ending positions of multilayered samples. The total charge amount of the multilayered samples under 5 kV were calculated and shown in Figure 10. It is shown that the total charge amount in multilayered samples by stacking with five untreated films is much more than the total charge amount in the samples by stacking with five fluorinated films. The results show that surface fluorination can significantly suppress the charge injection into samples, as the fluorination can change the surface chemical structure of films and change the trap depth.



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Charge amount (nC)

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15 30 45 Fluorination time of signal film in forming the multilayered samples (min)

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Figure 10. Relationship between the total charge amount of the multilayered samples and the fluorination time of signal film in forming the multilayered polyimide films with charging time 60 minutes.

4. SURFACE CHARGE ACCUMULATION AND DECAY ON DIRECT-FLUORINATED POLYIMIDE/AL2O3 NANOCOMPOSITES 4.1. Experiments In this research, the samples were prepared using commercially available pyromellitic dianhydride (PMDA) and 4, 4' -oxy dianiline (ODA), filled with nano-scale Al2O3 particles, which have an average diameter of 20 nm. The concentration of the nanoparticle was 0, 1, 3, 5, 7 wt%, respectively. The duration of the fluorination was 30 minutes. Besides, we performed the surface charge accumulation test at room temperature with relative humidity of ~ 40%. The voltage of ±6 kV was applied between the needle and the grounded electrodes.


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4.2. Effect of Charging Method Surface charge accumulation is initiated from the corona discharge. Charges with the identical sign of applied voltage are encouraged by the electric field and move towards the upper surface of the sample. Due to the presence of grid electrode, surface potential can be controlled so that all samples retain approximately the identical potential at the beginning of the decay test. The charge decay curves are shown in Figure 11. Surface charge density decreases with the decay time. The decay of surface charge is fast at the initial period, whilst becomes slower with the lapse of time. The decay behavior of negative charge is similar to that of positive charge. The possible routes for charge decay are lateral charge spreading, transport through the bulk and recombination with ions of opposite sign in air. Considering the electrode arrangement and sample dimension, electric field generated between implanted charge and grounded electrode is higher than the field along the surface, and hence bulk transportation should play a dominant role in the decay. The charge is firstly transferred from the surface to the bulk, and then undergoes trapping and de-trapping processes, leading to the decay of surface charge. The decay curves for samples charged under different time show distinguished features. The changes in surface charge decay of samples with nano-Al2O3 weight percent 3 wt%, caused by the variations in charging time, are shown in Figure 11. The decay curves flatten as the charging time is increased, showing that the charge becomes stabilized over time. This is primarily because the trapping for injecting charges is enhanced by increased charging time, and the implanted charge in turn inhibits charge transfer from surface to bulk, thereby prolonging the decay. It is also observed that the charge density is general higher for negative polarity than for positive polarity in the same condition. This phenomenon reveals that more negative charges are accumulated on the surface than positive charges. It can be interpreted that the de-trapping process of a positive charge is easier than that of a negative charge. Figure 11a shows the variations in surface charge density of samples without fluorination at different charging time; it can be found that the decay curves of the samples without fluorination have same decrease tendency for both positive and negative charged. With increasing the charging time, the surface charge density increases. Figure 11b shows that as the samples positive and negative charged for 5 and 10 minutes, the surface charge density decay fast in ten minutes of the decay time and then decay slowly after the ten



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minutes. While as the sample charged for 15 minutes, the surface charge density decays slower than that for 5 and 10 minutes.


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Figure11. Relationship between the surface charge density and the decay time of samples with nano-Al2O3 weight percent 3 wt%.

4.2. Effect of Fluorination As shown in Figure 11, the surface charge decay is fast at the beginning of the decay time. In order to observe the effect of fluorination and nanoparticles on samples decay rate directly, the average decay rate during the first five minutes of decay time is calculated. The relationships between the decay rate and the concentration of the samples with and without fluorination are shown in Figure 12. It is observed that the decay rate of a fluorinated sample is larger



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than that of the original sample and the largest decay rate is about 85% appearing at the concentration 0 wt% with both positive and negative charged for 10 mins. The results indicate that direct fluorination of polyimide film can significantly suppress surface charge accumulation and improve the decay rate. It is mainly because that the fluorine element has the largest electronegativity and makes the C-F bond highly polarized, which can form a fluorinated layer by surface fluorination in the samples. And it can be assumed that the fluorinate layer in the sample would capture electrons and form a shielding layer, which also shallows the deep traps. 90 (a) Positive charge: 10 min

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Figure 12. Relationship between the decay rate and the concentration after the decay for 5 minute.


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It is also observed that the decay rates decrease with increasing the concentration from 0 wt% to 3 wt% and increase with increasing the concentration from 3 wt% to 7 wt% for both negative and positive charged samples. The minimum value of the decay rate appears at 3 wt%. The results indicate that the nanoparticles restrain the charge transport through the bulk and influence the surface charge accumulation and decay rate. This shift may correlate with the way the various chemical structures that are influenced by nanoparticles. However, increasing of the concentration changes the injected charge distribution when interaction zones overlap. It was reported that the overlapping layers caused by the interaction zone’s surrounding nanoparticles will be more reactive than those of low concentration. When the interaction zones overlap, many conductive paths form through the overlap of the transition region in the bulk of nanocomposite [25], thereby reducing the charge trapping. Thus, the decay rates increase as the concentration increases from 3 wt% to 7 wt%. The trap distributions are shown in Figure 13. It is noted that the profiles of the distributions are similar, the minimum value in trap depth appears at ~0.76 eV with fluorinated pure sample. It also can be found that value of the sample with 3 wt% nanoparticles is largest for both original and fluorinated samples. The energy level of fluorinated samples is lower than that of samples without fluorination. A low energy level represents shallow traps in the sample that cannot obstruct the mobility of the trapped charge [23]. As charges are injected into the bulk only after the traps are filled, the surface potential increases after the traps are filled. Because of the shallow traps, the charges easily move out of the traps, as a result of which the surface charge of fluorinated samples decays faster than that of samples without fluorination. It can be speculated that the introduction of fluorine to the polyimide chemical structure can lower the traps.

4.3. Dissipation Time The dissipation time is defined as the period during which the total number of charges falls to 10% of the original charged value. Long dissipation time corresponds to slow surface charge decay. Figure 14 shows the relationship between the dissipation time and the concentration of Al2O3 in the untreated samples and the fluorinated ones respectively. It is found that the dissipation time increases with the concentration of Al2O3 increasing from 0


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wt% to 3 wt% and it decreases as the concentration increases from 3 wt% to 7 wt%. It is also observed that the dissipation time of the fluorinated samples is shorter than that of the untreated samples at the same concentration. 5

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Figure 14. Relationship between the dissipation time and the concentration without and with fluorination.

The minimum dissipation time is 0.15 hours, as seen in the fluorinated sample of 0 wt% concentration after being negatively charged (as shown in Figure 14b), and the maximum dissipation time is 4.5 hours, as seen in the untreated sample of 3 wt% concentration after negative charging (as shown in Figure 14a). The results reveal that the increasing trend of the dissipation time reverses at the Al2O3 concentration of 3 wt%. It also can be found that the fluorination can significantly improve the decay rate of the surface charge in polyimide nanocomposite films.



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CONCLUSION Effects of the direct fluorination on charge accumulation and charge decay of polyimide film and polyimide composites were investigated by using DC corona charging method and PEA method. The dynamic behaviors of charge were associated with the characteristics of surface states that are changed with the fluorinated layer of samples. The main conclusions can be summarized as follows, 1) The dielectric permittivity decreases with increasing the fluorination time from 0 min to 45 min and increases from 45 min to 60 min. The decay rate of fluorinated sample is much larger than that of original sample. The dissipation time of fluorinated sample is shorter than original sample, and the minimum dissipation time appears at fluorination time 45 minutes, which indicates that the fluorinate sample is harder to fail than original one. With the increase of charging time, the decay rate of original sample decreases, and the dissipation time increases. It suggests that the charge transfers from the surface to the bulk become slower with increasing the charging time. While the decay rate of fluorinated for 45 minute sample with the charging time 30 min is larger than that 10 and 20 min, which indicates that appropriate fluorination conditions can formed fluorinated layer without destroying the aromatic imide chemical structure in the polyimide film and significantly suppress the charge accumulation on the samples. The decay for both negative charge and positive charge is improved by the fluorination. Such improvement for negative charge is more remarkable than that for positive charge, and thereby the positive charge decays lower than the negative charge. 2) The space charge distribution of multilayered film by stacking with five fluorinated samples is quite different from that of multilayered film by stacking with five untreated samples. The space charge behaviors of multilayered fluorinated films have similar characteristics with the charging time from 10 seconds to 3600s. The conductivity of fluorinated film is higher than the untreated film. The amounts of space charge in multilayered polyimide films were calculated. The result shows that the total charge amount in multilayered samples by stacking with five untreated films is much larger than the total charge amount in the samples by stacking with


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Boxue Du five fluorinated films. It is indicated that surface fluorination can significantly suppress the charge injection into samples. 3) The dielectric permittivity and decay rate decrease with increasing the concentration from 0 wt% to 3 wt% and increase from 3 wt% to 7 wt%. The surface fluorination can lower the dielectric permittivity of samples. The decay rate of fluorinated sample is larger than the sample without fluorination. The dissipation time of fluorinated sample is shorter than the sample without fluorination, which indicates that fluorination can formed fluorinated layer without destroying the aromatic imide chemical structure in the polyimide nanocomposite film and significantly suppress the charge accumulation on the samples.

With the increase of charging time, the decay rate decreases, and the dissipation time increases. It suggests that the charges, which transfer from the surface to the bulk, become slower with increasing the charging time. The decay for both negative charge and positive charge is improved by the fluorination. Such improvement for negative charge is more remarkable than that for positive charge, and thereby the positive charge decays slower than the negative charge with the concentration of nano-Al2O3 0 wt%. While the decay for both negative charge and positive charge is restrained by the nanoparticles. Such restriction for negative charge is more remarkable than that for positive charge, and thereby the positive charge decays faster than the negative charge with the doped samples.

REFERENCES [1] [2]

[3]

S. Zelmat, M. L. Locatelli, T. Lebey and S. Diaham, “Investigations on high temperature polyimide potentialities for silicon carbide power device passivation,” Microelectr Eng., Vol. 83, pp. 51-54, 2006. S. Diaham, M.-L. Locatelli, T. Lebey, C. Raynaud, M. Lazar, H. Vang and D. Planson, “Polyimide Passivation Effect on High Voltage 4HSiCPiN Diode Breakdown Voltage,” Materials Sci. Forum, Vol. 615617, pp. 695-698 2009. M. Kaufhold, “Electrical stress and failure mechanism of the winding insulation in PWM-inverter-fed low-voltage induction motors,” IEEE Trans. Industrial Electronics, Vol. 47, pp. 396-402, 2000.


Charge Behaviors on Direct-Fluorinated Polyimide Films … [4]

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R. J. Beeckman. “Studies on Magnet Wire Degradation with Inverter Driven Motors,” IEEE Electr. Electronics Insulation Conference and Electrical Manufacturing and Coil Winding Conf., New York, USA, pp. 383-387, 1997. J. R. Bellomo, P. Castelan amd T. Lebey, “The Effect of Pulsed Voltages on Dielectric Material Properties,” IEEE Trans. Dielectr. Electr. Insul., Vol. 6, No. 1, pp. 20-26, 1999. J. R. Bellomo, T. Iebey, J. M. Oraision and F. Peltier. “Electrical Aging of Stator Insulation of Low Voltage Rotating Machines Supplied by Inverters,” IEEE Int’l. Sympos. Electr. Insul., New York, USA, pp. 210213, 1996. D. Fabiani, G. C. Montanari, A. Cavallini and G. Mazzanti, “Relation between space charge accumulation and partial discharge activity in enameled wires under PWM-like voltage waveforms”, IEEE Trans. Dielectr. Electr. Insul., Vol. 11, pp. 393-405, 2004. L. R. Zhou, G. N. Wu, B. Gao, K. Zhou, J. Liu, K. J. Cao and L. J. Zhou, “Study on Charge Transport Mechanism and Space Charge Characteristics of Polyimide Films,” IEEE Trans. Dielectr. Electr. Insul, Vol. 16, No. 4, pp.1143-1149, 2009. K. Kimura, S. Itaya and M. Hikita, “Partial Discharge Behavior on PI Film with Micro Gap under Steep-front Step Voltage,” IEEE Int’l. Sympos. Electr. Insul., Boston, MA, USA, pp. 1-4, 2002. P. Kharitonov, “Direct fluorination of polymers-From fundamental research to industrial applications,” Prog. Org. Coat., Vol. 61, pp. 192204, 2008. Tressaud, E. Durand, C. Labrugere, A. P. Kharitonov, and L.N. Kharitonova, “Modification of surface properties of carbonbased and polymeric materials through fluorination routes: From fundamental research to industrial applications,” J. Fluorine Chem., Vol., 128, pp. 378-391, 2007. D. M. Lemal, “Perspective on fluorocarbon chemistry,” J. Org. Chem., Vol. 69, pp. 1-11, 2004. S. J. Blanksby and G. B. Ellison, “Bond dissociation energies of organic molecules,” Acct. Chem. Res., Vol. 36, pp. 255-263, 2003. S. Mathews, I. Kim, and Chang-Sik Ha, “Synthesis, Characterization, and Properties of Fully Aliphatic Polyimides and Their Derivatives for Microelectronics and Optoelectronics Applications,” Macromolecular Research, Vol. 15, No. 2, pp. 114-128 2007.


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[15] R.G. Pethe, C.M. Carlin, H.H. Patterson, and W.N. Unertld, “Effect of Dose Stoichiometry on the Structure of Vapor-deposited Polyimide Thin Films”, Polym. Sci., Polym. Chem., Vol. 8, No. 12, pp. 3218-3228, 1993. [16] Cherdoud-Chihani, M. Mouzali, and M.J.M. Abadie, “Study of crosslinking acid copolymer/DGEBA systems by FTIR,” J. Appl. Polym. Sci., Vol. 87, pp. 2033-2051, 2003. [17] R.E. Banks, B. E. Smart and J.C. Tatlow, Characteristics of C-F Systems, in: Organofluorine Chemistry: Principles and Commercial Applications, Plenum Press, New York, pp. 57-88, 1994. [18] F. P. Wang, Z.F. Xia, X.l. Qiu, J. Shen, X.P. Zhang, and Z.L. An, “Piezoelectric Properties and Charge Dynamics in Poly(vinylidene fluoride-hexafluoropropylene) Copolymer Films with Different Content of HFP,” IEEE Trans. Dielectr. Electr. Insul., Vol. 13, No. 5, pp. 11321139, 2006. [19] Y. Gao, B.X. Du, Z.L. Ma and X.H. Zhu, “Decay Behavior of Surface Charge on Gamma-Ray Irradiated Epoxy Resin,” IEEE Int’l. Conf. Solid Dielectr.(ICSD), pp. 1-4, 2010. [20] X. Du, J. W. Zhang and Y. Gao, “Effects of TiO2 Particles on Surface Charge of Epoxy Nanocomposites,” IEEE Trans. Dielectr. Electr. Insul., Vol. 19, No. 3, pp. 755-762, 2012. [21] G. Chen, Z. Xu and L. X. Zhang, “Measurement of Surface Potential Decay of Corona Charged Polymer Films using the Pulsed Electroacoustic Method,” Measur. Sci. Tech., Vol. 18, pp. 1453-1458, 2007. [22] H. Berlepsch, “Interpretation of Surface Potential Kinetics in HDPE by a Trapping Model,” J. Phys. D. Appl. Phys., Vol. 18, No. 6, pp. 11551170, 1985. [23] X. Du and Y. Gao, “Gamma-ray Irradiation Inhibiting Surface Charge Accumulation on Polyethylene,” IEEE Trans. Dielectr. Electr. Insul., Vol. 16, No. 3, pp. 876-881, 2009. [24] G. Chen, Y. Tanaka, T. Takada and L. Zhong, “Effect of Polyethylene Interface on Space Charge Formation,” IEEE Trans. Dielectr. Electr. Insul., Vol. 11, No. 1, pp. 113-121, 2004. [25] P. Molinié, “Measuring and Modeling Transient Insulator Response to Charging: the Contribution of Surface Potential Studies,” IEEE Trans. Dielectr. Electr. Insul., Vol. 12, No. 5, pp. 939-950, 2005. [26] B. X. Du, J. Li, W. Du, “Surface Charge Accumulation and Decay on Direct-fluorinated Polyimide/Al2O3 Nanocomposites,” IEEE Trans. Dielectr. Electr. Insul., Vol. 20, No. 5, pp. 1764-1771, 2013.



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Chapter 3

A REVIEW ON RECENT PROGRESS OF RESEARCH AND APPLICATIONS FOR POLYIMIDE AEROGELS Jin-gang Liu1,, Xiu-min Zhang2,†, Fei-xu Chen1, Wang-shu Tong1 and Yi-He Zhang1 School of Materials Science and Technology, China University of Geosciences, Beijing, China 2 School of Electrical Engineering, Beijing Jiaotong University, Beijing, China 1

ABSTRACT Recent research and development of polyimide (PI) aerogels have been reviewed. PI aerogels possess both of the merits of conventional aromatic PIs and common polymer aerogels; thus have been widely investigated as components for high temperature, low dielectric constants, and low density applications. Up to now, they have found various applications in aerospace, aeronautical, microelectronic and optoelectronic fields. The current review coveres the latest research and development for PI aerogels, including crosslinkers syntheisis chemistry, PI aerogels synthesis chemistry, and their engineering applications. Especially, this review focuses on the applications of PI aerogels as high  †

Corresponding Author: [email protected] (Liu). [email protected] (Zhang).


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Jin-gang Liu, Xiu-min Zhang, Fei-xu Chen et al. temperature resistant components for aerospace exploration; thermally and electrically insulating materials for electrical engineering; and low dielectric constant (low-k) interlayer dielectrics (ILDs) for microelctronic fabrications.

Keywords: polyimide aerogels, synthesis; thermal properties, low dielectric constant, dielectrics

INTRODUCTION Aerogels represent a class of solid-state materials derived from gels in which the liquid component of the gels has been replaced with gas [1, 2]. Thus, aerogels usually possess porosities higher than 80%, even up to 99%. This structural characteristic endows aerogels various remarkable properties, including extremely low density (close to air), low thermal conductivity, low dielectric constant (close to 1.0), high specific surface area, high acoustic attenuation, and high thermal resistance. Therefore, aerogels show promise for a wide range of applications [3]. According to the structural compositions of the aerogels, they can be roughly divided into three types, that is, inorganic aerogels, organic aerogels, and inorganic/organic hybrid ones. This term aerogel was reported by Kistler in 1930s, who first successfully developed inorganic silica aerogels from silica gel via supercritical fluid drying procedure [4]. Inorganic silica aerogel is usually nicknamed “frozen smoke” or “solid smoke” or “blue smoke” due to its translucent nature and the way light scatters in the material. However, the fragile nature for the material makes it difficult to be handled in its monolithic form, which greatly limiting its practical applications. Recently, organic aerogels have attracted increasing attention due to their intrinsic low density, low thermal conductivity, high thermal insulation, low dielectric constants, and especially superior mechanical properties, low cost, and easy availability compared with their inorganic counterparts [5]. Figure 1 compares the typical characteristics of inorganic and organic aerogels. Purely organic aerogels were first reported by Pekala and coworkers in 1989, which were based on a framework consisting of a resorcinolformaldehyde (RF) resin [6]. The RF aerogels exhibited low thermal conductivity of 0.012 W/m K at 0.16 g/cm3. Since then, various organic


A Review on Recent Progress of Research and Applications …

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aerogels, such as polyurethane [7], polyurea [8], polystyrene [9], amide [10], and polydicyclopentadiene [11] aerogels have been widely investigated in the literature and have found various applications in high-tech fields. However, common organic aerogels usually suffer from their low thermal and dimensional stability at elevated temperatures; thus cannot meet the severe demands of high temperature applications, such as thermal protecting coatings or interlayer electrically insulation systems for microelectronic device fabrications [12]. In order to prevent the thermal deformation of the aerogels in the above applications, high temperature resistant organic aerogels are highly desired.

Figure 1. Classification of aerogels and their typical characteristics.

Polyimides (PIs) area class of high temperature resistant polymers and have been widely used in aerospace, electrical, microelectronic, and optoelectronic industry for nearly half a century due to their excellent combined thermal, mechanical, and dielectric properties [13]. Thus, it can be anticipated that the PIs in the form of aerogels might exhibit good comprehensive properties. Very recently, the research and development of a series of functional PI aerogels for electronic and space explorations have been reported. A variety of novel PI aerogels have been developed. In the current paper, the recent R&D for PI aerogels, including their synthesis chemistry and potential applications in microelectronic engineering, electrical engineering, and aerospace engineering were reviewed.


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POLYIMIDE AEROGEL SYNTHESIS CHEMISTRY Monomers Synthesis Chemistry The characteristics of the starting monomers, including dianhydrides, diamines, and multi-functional crosslinkers (or end-cappers) are the key components for developing PI aerogels because that most of the features and functions for the final PI aerogels, such as thermal stability, mechanical properties (strength and toughness), and dielectric properties (dielectric constant and dissipation factor) are usually achieved via the functionalization of the starting monomers. On the other hand, the cost of the PI aerogels is, to some extent, determined by the starting materials. Typical commercially available monomers for PI aerogels synthesis are shown in Figure 2. The dianhydride monomers usually include pyromelliticdianhydride (PMDA), 3,3',4,4'-biphenyltetracarboxylic dianhydride (BPDA), 3,3',4,4'-benzophenetetracarboxylic dianhydride (BTDA), and 1,2,3,4-cyclobutanetetracarboxylic dianhydride(CBDA). The aromatic diamines include para-phenylenediamine (PPD), 4,4'-oxydianline (ODA), bisaniline-p-xylidene (BAX), 2,2'-dimethylbenzidine (DMBZ), 2,2'-bis(trifluoromethyl)benzidine (TFDB), 2-(4-aminophenyl)-5-aminobenzimidazole (4-APBI), 2-(3-aminophenyl)-5-aminobenzimidazole(3-APBI), 5-amino-2-(4-aminophenyl)benzoxazole (APBO), and N-[(heptaisobutylPOSS)propyl]-3,5-diaminobenzamide (DA-POSS). The multifunctional crosslinkers include 1,3,5-tris(4amino-phenoxy)benzene (TAB), 4-(4'-aminophenyl)-2,6-bis(4''-aminophenyl) pyridine (TAPP), octa(aminophenyl) silsesquioxane (OAPS), and so on. The common characteristics for these monomers are the rigid-rod and conjugated units in their molecular structures, including biphenyl, cyclobutane, benzimidazole, benzoxazole, ortriphenylpyridine moieties. This molecular feature is beneficial for inducing the gel formation during the preparation of the PI wet gels inpolar aprotic solvents, such as N,N-dimethylforamide (DMF), N,N-dimethylacetamide (DMAc), or N-methyl-2-pyrroridinone (NMP). In order to adjust the mechanical properties of the PI aerogels, flexible linkages, such as ether and carbonyl are also utilized. For the crosslinkers, multifunctional amine compounds are usually adopted due to their relatively easy availability and low synthesis cost. The synthesis of dianhydrides and diamines monomers has been well reported in the literature [14]. Thus, the synthesis chemistry for the rarely reported crosslinkers will be introduced in the current paper.



Figure 2. Typical monomers for PI aerogels synthesis.


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Jin-gang Liu, Xiu-min Zhang, Fei-xu Chen et al. Table 1. Common crosslinkers for PI aerogel synthesis

Crosslinkers tris(4-aminophenyl)amine (TAPA)

Chemical structure NH2

H 2N

N

tris(3-aminophenyl) phosphine oxide(TAPO) P O

1,3,5-tris(4aminophenyl)benzene (TAPB)

260

[16]

263

[17]

277.3

[18]

88-89

[19, 20]

not detected

[21]

NH2

NH2

H 2N

NH2

4-(4'-aminophenyl)-2,6bis(4''-aminophenyl) pyridine (TAPP)

NH2

H 2N

NH2

N

1,3,5-tris(4-aminophenoxy) benzene (TAB)

NH2

O H2 N

octa(aminophenyl) silsesquioxane (OAPS)

ref [15]

NH2

NH2

H2 N

Melting point (oC) 244.92

O

R R

O Si O Si

R O Si Si R O O

O

NH2

O

R Si O Si RO O Si O R Si O R

NH2 R=

Table 1 shows the chemical structures for amino-type crosslinkers reported in the literature and Figure 3 summarizes their typical synthetic pathways. It can be seen that, mild reaction conditions are often used in the synthesis of the crosslinkers. In most cases, the total synthesis yields for these compounds are moderate to high. For instance, Xu and coworkers synthesized TAPA by reducing the corresponding trinitro compound using hydrazine hydrate, catalyzed by palladium/C [15]. The crude product was easily purified by recrystallization with ethanol with a yield higher than 80%. Wang and coworkers reported the phosphorus-containing triamine, TAPO by a two-step synthesis reaction [16]. First, the nitration reaction of triphenylphosphine oxide was carried out in the presence of nitric acid and concentrated sulfuric



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acid to afford the corresponding trinitro compound. Then, the trinitro compound was reduced by stannous chloride (SnCl2) to afford the target TAPO as light brown crystalline solid. He and coworkers prepared the rigid-rod TAPB by reduction of the corresponding trinitro compound with hydrazine hydrate and ferric chloride (FeCl3) [17]. The trinitro compound was synthesized by the dehydration reaction of 4-nitroacetophenone under the catalysis of trifluoromethanesulfonic acid. Shen and coworkers reported another rigid-rod triphenylpyridine-containing triamine compound, TAPP from the reaction of 4-nitroacetophenone and 4-nitrobenzaldhyde, followed by the catalytic hydrogenation of Pd/C catalyst and hydrazine hydrate [18]. Takeichi and coworkers reported the synthesis chemistry of TAB usingphloroglucinol and 4-fluoronitrobenzene as the staring materials, followed by reduction of the trinitro intermediate with Na2S [19]. Similarly, the same compound was synthesized from phloroglucinol and 4chloronitrobenzene, followed by reduction of the trinitro compound with SnCl2 [20]. Besides the trifunctionalcrosslinkers mentioned above, other multifunctional crosslinkers have also been reported in the literature. For instance, octa-functional OAPS crosslinker has been reported by Wu and coworkers by a simple procedure involving nitration of octaphenylsilsesquioxane, followed by mild reduction with Pd/C catalyst and hydrazine hydrate [21]. Abundant multi-functional crosslinkers developed in the literature make it possible to develop crosslinked PI aerogels with enhanced mechanical and thermal properties. A variety of PI aerogels have been synthesized by using the crosslinkers shown in Table 1.

PI Aerogels Synthesis Chemistry and Manufacturing Process PI aerogels can be divided into two types, which are linear and crosslinked according to their different microstructures. The synthesis pathway for these two types of PI aerogels is a bit different. Linear PI aerogel has been currently achieved commercialization by Aspen Aerogels, Inc., USA. The company obtained the linear PI aerogel synthesis patent in 2006 [22]. According to the synthesis procedure, PMDA was first polymerized with ODA in NMP at an equal molar ratio. Then, acetic anhydride and pyridine was added to the obtained PAA solution, resulting in the formation of PI wet gel. Then, the precipitated PI gel was further heating at 190oC in NMP to facilitate the imidization. After that, the NMP solvent was replaced by ethanol. Finally, the


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PI wet gel containing ethanol was dried in ScCO2 to afford the final PI aerogel. The density, pore size and specific surface area parameters for the PI aerogel can be adjusted through the solid content of the starting PAA solution. For example, when the solid content of PAA solution was 2 wt%, the PI aerogel showed a density of 0.02 g/cm3, the average pore size of 28.2 nm, and the specific surface area of 1063 m2/g. When the solid content is 5wt%, the obtained PI aerogel exhibited the density, average pore size, and the specific surface area of 0.03 g/cm3,40 nm and 1328 m2/g, respectively.

Figure 3. Synthesis chemistry of crosslinkers for PI aerogels synthesis.

Leventis, et al. synthesized monolithic linear PI aerogels by ScCO2 drying the PI wet gels synthesized from PMDA and 4,4'-methylenediphenyldiisocyanate (MDI) instead of diamine at room temperature, followed by the



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imidization reaction of the obtained seven-member intermediate by expelling carbon dioxide at 60-90oC [23]. The molar ratio of PMDA and MDI was 1:1 and no crosslinking agent was used. The derived PI aerogel showed similar Branuaer-Emmet-Teller (BET) surface area (300-400 m2/g) with their chemically identical PI aerogel derived from PMDA and 4,4'methylenedianiline (MDA) by conventional route. Although the cost of this synthesis pathway from dianhydride and diisocyanate at relatively low processing temperature is lower than that of the conventional route for linear PI aerogel synthesis at high temperature, the obtained aerogel usually showed inferior mechanical properties. In addition, the volume shrinkage of the wet PI gels was usually higher than 30%. In summary, for linear PI aerogels, volume shrinkage as large as 60% might occur during the supercritical drying process due to the linear nature of their molecular structures. However, the shrinkage value is usually less than 20% for crosslinked PI aerogels due to the three dimensional network of their structures. Thus, crosslinked PI aerogels have been widely investigated recently. Up to now, the synthesis chemistry for crosslinked PI aerogels has been well established in the literature, as shown in Figure 4. First, excessive dianhydrides monomer (m stands for the mole number of dianhydride) was reacted with diamines monomer (n stands for the molenumberof diamine, n