A study of polyimide thin films physical aging and plasticization behaviors

A STUDY OF POLYIMIDE THIN FILMS -
PHYSICAL AGING AND PLASTICIZATION BEHAVIORS







ZHOU CHUN










NATIONAL UNVERISTY OF SINGAPORE
2003
A STUDY OF POLYIMIDE THIN FILMS -
PHYSICAL AGING AND PLASTICIZATION BEHAVIORS

ZHOU CHUN
(B.Eng., BUAA)





A THESIS SUBMITTED
FOR THE DEGREE OF MASTER OF SCIENCE
DEPARTMENT OF CHEMISTRY
NATIONAL UNVERISTY OF SINGAPORE
2003
ACKNOWLEGEMENT

First of all, I would like to express my deepest appreciation and thanks to my
supervisors, Professor Neal Chung Tai-Shung, Professor Goh Suat Hong, Dr. Wang
Rong, and Dr. Liu Ye for their intellectually-stimulating guidance and invaluable
encouragement throughout my candidature.

I am grateful for the Research Scholarship from the National University of Singapore
(NUS) that enables me to pursue my M.Sci. degree. I am also indebted to the Institute
of Materials Research and Engineering (IMRE) of Singapore for the equipment and the
Top-up Award.

Thanks are also due to my fellow students and the researchers in our group, Mr. C.
Cao, Dr. S.L. Liu, Ms. M.L. Chng, Mr. D.F. Li, Mr. Y.C. Xiao, Mr. Y. Li, Mr. J.Y.
Xiong, Mr. K.Y. Wang, Mr. L. Shao, Mr. Z. Huang, Ms. P.S. Tin, Ms. W.F. Guo, Ms.
L.Y. Jiang, Ms. M.M. Teoh, and Ms. H.M. Guan for all the handy helps, technical
supports, invaluable discussion and suggestions. Special thanks are due to Dr. K.P.

Pramoda in IMRE for her very kind help in characterization assistance.

Last but not least, I am most grateful to my parent, brother, and my finacee, Miss N.
Li, for their absolute love, encouragement and support. This thesis would not have
existed without them.

i
TABLE OF CONTENTS
Page
ACKNOWLEDGEMENT i
TABLE OF CONTENTS ii
SUMMARY vi
NOMENCLATURE viii
LIST OF FIGURES xi
LIST OF TABLES

xix

CHAPTER 1 INTRODUCTION


1
1.1 Membrane and membrane-based gas separation 1
1.2 Transport mechanism of membranes 2
1.2.1 General principles 2
1.2.2 Solution-diffusion model 3
1.2.3 Sorption in glassy polymers – Dual mode sorption model

4
1.3 Membrane material selection and tailoring 5

1.4 CO
2
plasticization and physical aging of glassy polymer

7
1.5 Why thin films?

9
1.6 Goals and organization of this research

10


CHAPTER 2 LITERATURE REVIEW

12
2.1 The aging phenomenon of glassy polymers and the effect on gas

ii
separation membranes 12
2.1.1 Introduction 12
2.1.2 Non-equilibrium behavior of glassy polymers 13
2.1.2.1 Glass transition 13
2.1.2.2 Relaxation time distribution and cooperative relaxation 16
2.1.2.3 Secondary transition and the temperature range of physical aging 19
2.1.3 An overview of the effect of physical aging on gas separation
membranes
22
2.1.4 Experimental techniques in physical aging study of glassy
membranes

30
2.1.4.1 Mechanical properties, DSC, and PALS 31
2.1.4.2 Solid state NMR and ESR 36
2.2 CO
2
plasticization and anti-plasticization of gas separation
membranes
39
2.2.1 CO
2
plasticization 39
2.2.2 Suppression of CO
2
-induced plasticization 40

CHAPTER 3 EXPERIMENTS
42
3. 1 Material synthesis and preparation of dense films 42
3.1.1 Materials 42
3.1.2 Preparation of dense membranes 43
3.1.3 Drying procedure and thermal history 44
3.2 Permeation measurements 46
3.3 Thickness acquisition by Scanning electron microscope (SEM) 50
3.4 Aging monitoring and CO
2
plasticization experiments of 6FDA- 52

iii
Durene dense membranes

3.4.1 Aging experiments 52
3.4.2 CO
2
plasticization experiments 52
3.5 Chemical cross-linking modification of 6FDA-Durene dense
membranes for the improvement of the resistance of CO
2
-induced
plasticization and suppressed aging process

52
3.5.1 Mechanism and procedure of the chemical cross-linking modification 53
3.5.2 FTIR Characterization of cross-linked 6FDA-Durene films 54

CHAPTER 4 GOVERNING EQUATION FOR PHYSICAL
AGING OF THICK AND THIN FULORO-
POLYIMIDE FILMS

56
4.1 Introduction 56
4.2 Derivation of the proposed equation 58
4.3 Results and Discussion 61


CHAPTER 5 Accelerated CO
2
Plasticization of Ultra-thin
Polyimide Films and the Effect of Surface
Chemical Cross-linking on Plasticization and
Physical Aging

68
5.1 Introduction 68

iv
5.2 Results and discussion 72
5.2.1 Effect of chemical cross-linking on physical aging 72
5.2.2 The accelerated CO
2
plasticization for thin films and the plasticization
resistance induced by cross-linking
74
5.3 Conclusions 81

CHAPTER 6 CONCLUSIONS

82
6.1 Experimental observation and theoretical aspects of the physical aging of
thick and thin polyimide films
82
6.2 Accelerated CO
2
plasticization of thin polyimide films and an effective
cross-linking modification to suppress plasticization and retard physical
aging
83
6.3 Comprehensive review of the effect of physical aging on glassy gas
separation membranes and remaining problems
83

REFERENCES
85



v
SUMMARY

A systematic research, which covers the characterization of the intrinsic gas
permeation properties, the physical aging process monitoring, the CO
2
plasticization
behavior evaluation of the dense 6FDA-Durene polyimide films of different thickness,
and finally the chemical cross-linking modification to withstand the plasticization of
CO
2
for CO
2
separation and retard the physical aging process, has been presented in
this thesis.

We attempted to study the effect of film thickness on the physical aging and the CO
2

plasticization behavior of the glassy polyimide membrane, because the asymmetric
membrane with a thin and dense separating layer has been widely applied in industrial
scale applications and is therefore of great interest, academically and industrially. In
addition, we proposed an easy and feasible chemical modification method to improve
the physical aging and CO
2

plasticization resistance of the membrane. The knowledge
of this has been proven to be critical for membrane based gas separation processes.

Specifically, this work investigated (i) the aging profile of 6FDA-Durene polyimide
dense films with different thickness, thus to correlate the aging of hollow fiber
containing a thin and dense selective layer with the aging of dense films of comparable
thickness; (ii) the CO
2
plasticization behaviors of 6FDA-Durene films with different
thickness; (iii) the effects of chemical cross-linking modification of 6FDA-Durene on
the aging and CO
2
plasticization behaviors.


vi
Finally, an accelerated physical aging process of the 6FDA fluoro-polyimide was
observed and employed to validate a proposed equation, derived from the molecular
mobility of polymer segments below the glass transition temperature of the polymer,
that serves to correlate the change of permeability as a function of time during the
physical aging process. Strongly thickness-dependent aging process was found by
employing pure O
2
and N
2
tests to monitor the change of gas permeation properties as
a function of aging time. Interestingly, an accelerated CO
2
plasticization indicates that
the conventionally defined “plasticization pressure” as an inherent material properties

measured from thick dense films is also strongly thickness dependent. Experimental
results suggest that chemically modified ultra-thin films show characteristics of
retarded aging process and significantly suppressed plasticization.















vii
NOMENCLATURE
A Effective area of the film (cm
2
)
b Langmuir affinity constant (atm
-1
)
C Local penetrant concentration in the film (cm
3
(SPT)/cm
3

(polymer))
C
1
Local penetrant concentration at the downstream side (cm
3
(STP) /
cm
3
(polymer))
C
2
Local penetrant concentration at the upstream side (cm
3
(STP) / cm
3

(polymer))
C
D
Henry sorption concentration (cm
3
(STP) / cm
3
(polymer))
C
H
Langmuir sorption concentration (cm
3
(STP) / cm
3

(polymer))
c
H
’ Langmuir sorption capacity (cm
3
(STP) / cm
3
(polymer))
D Diffusion coefficient (cm
2
/s)
D
D
Average local measure of mobility of a penetrant in the
Henry site (cm
2
/s)
D
H
Average local measure of mobility of a penetrant in
Langmuir (H) environments (cm
2
/s)
dp/dt Rate of pressure in the low-pressure downstream
chamber (mmHg/sec)
K C
H
’b / k
D
D

k

Henry’ law constant ((cm
3
(STP)) / cm
3
(polymer) atm)
l
Membrane thickness (cm)
N Permeation flux (cm
3
/cm
2
-sec)
p Pressure (cm Hg)
P Permeability coefficient of a membrane to gas (1 barrer =1 x 10
-10


viii
cm
3
(STP)-cm / cm
2
-sec-cm Hg.)
p
0
Standard pressure of 1 atm or 76 cm Hg
p
1

Down stream pressure of the penetrants (cm Hg)
p
2
Upstream stream pressure of the penetrants (cm Hg)
p
∆

Pressure difference (cm Hg)
Q

Volumetric flow rate of gas at standard temperature
and pressure (cm
3
(STP) /sec)
R Universal gas constant
S Solubility coefficient (cm
3
(STP)/cm
3
(polymer)-cmHg)
T Absolute temperature of the measurement (K)
T
0
Standard temperature of 273.15K
T
g
Glass transition temperature of penetrant (K)
V Volume of the low-pressure chamber (cm
3
)

x Distance from the upstream side of the film to downstream (cm)
x
S
Local concentration of component 1 at the retentate side of
permeator
y Local concentration of component 1 at the permeate side of
permeator
BA /
α

Separation factor of a gas pair
*
/ BA
α

Ideal separation factor of a gas pair (permselectivity)
θ
Time lag (sec)
ρ
Density (g/cm
3
)

ix

Abbreviation


6FDA 2,2-Bis [3,4-dicarboxyphenyl] hexafluoropropane dianhydride
Durene 2,3,5,6-Teramethyl-1,4-phenylene diamine

DSC Differential scanning calorimetry
FTIR Fourier tansform infrared spectroscopy
NMP N-methyl pyrrolidone
SEM Scanning electron microscope
XRD Wide angle X-ray diffraction







x
LIST OF FIGURES
Figure 1.1 Literature data for CO
2
/CH
4
permselectivity versus CO
2

permeability (Robeson, 1991)

6
Figure 2.1 Schematic diagram of the glass transition of physical aging process

15
Figure 2.2 The relationship between the molecular mobility, M, the degree of
packing, and free volume (Struik, 1978)

16
Figure 2.3 Small-strain tensile creep curves of PVC quenched from 90 to 40
o
C
and annealed for various times (Struik, 1978)

17
Figure 2.4 Schematic Dynamic Mechanical Analysis (DMA) curve of
polymers

20
Figure 2.5 DMA curves of polycarbonate samples with different heat history
(Struik, 1978)

22
Figure 2.6 Effect of aging time, t, on the oxygen permeability coefficients for
BPA-BnzDCA films of the following thickness (McCaig and Paul,
2000)

28
Figure 2.7-a, 2.7-b The N
2
and He/ N
2
permeability coefficients as a function of
aging for PTMSP with different thickness (Dorkenoo and Pfromm,
2000)

29
Figure 2.8-a, 2.8-b Schematic diagram of thalpy and corresponding specific

heat changes for annealed (dashed line) and unannealed (solid line)
glasses on heating at rate r1 (Petrie, 1997)

34
Figure 2.9 Heating scans at 10 Kmin
-1
for polycarbonate samples cooled at the
rates indicated and immediately reheated in the DSC (Hutchinson et
al., 1999)

34
Figure 3.1 Chemical structure of 6FDA-Durene polyimide

43
Figure 3.2. Schematic diagram of dense membrane gas permeation test
apparatus

49
Figure 3.3 A typical curve of relationship between downstream pressure and
time

47
Figure 3.4 A typical series of SEM pictures of a thin film
(Scale bar: Upper: 10 µm, Middle: 1 µm, Bottom: 1 µm)

51
Figure 3.5 Mechanism of chemical cross-linking modification

53
Figure 3.6 FTIR spectra of 6FDA-Durene Films

(a) original samples; (b)-(d) samples immersed in 2%(w/w) p-
54

xi
xylene diamine solution for 0.5, 1 and 3 minutes at ambient
temperature, respectively

Figure 4.1-a O
2
permeability vs. aging time for films with different thickness

61
Figure 4.1-b Percentage change of O
2
permeability vs. aging time

63
Figure 4.2 The double logarithmic curve of Permeability and Aging time
for films with different thickness

63
Figure 4.3 The double logarithmic curve of Permeability and Aging time
for films with different thickness

66
Figure 4.4 O
2
/N
2
permselectivity vs. aging time for films with different

thickness

67
Figure 5.1 Percentage change of O
2
permeability vs. aging time
for different films

72
Figure 5.2 The change of selectivity coefficient of O
2
/N
2
as a function of
aging time for different films

74
Figure 5.3 CO
2
permeability as a function of exposure time

75
Figure 5.4 CO
2
permeability as a function of exposure time at feed pressure
of 8 atm for a cross-linked film at feed pressure of 8 atm for
different films

77
Figure 5.5 CO

2
permeability as a function of feed pressure for uncross-
linked films

77
Figure 5.6 CO
2
permeability as a function of feed pressure for cross-linked
and uncross-linked films

78






xviii
LIST OF TABLES
Table 2.1 Effect of vacuum oil vapor on the O
2
permeability coefficient of
PTMSP membranes (Nagai and Nakagawa, 1995)

24
Table 2.2 Lifetime spectrum parameters of PALS for PTMSP films
(Yampol’skii et al., 1993a)

36
Table 2.3 Solid state NMR results of aged and original PTMSP films

synthesized from different catalysts (Nagai et al., 1999)

38
Table 4.1 Values of B(T) and A for films with different thickness 65


xix
CHAPTER ONE
INTRODUCTION


1.1 Membrane and membrane-based gas separation

Membrane-based separation has appeared to be one of the promising and rapidly
growing areas in separation technology (Rousseau, 1987) because it is more
economical and energy-saving thus outweighs the traditional approaches like
cryogenic distillation that requires a phase change of the feed mixture. Most available
membrane-based separation processes are in the forms of gas separation, reverse
osmosis, microfiltration, ultrafiltration (Fane, 1984), liquid membranes, pervaporation
(Okada and Matsuura, 1991), dialysis and electrodialysis. The work presented here is
engaged in the membrane-based gas separation.

A membrane, principally a selective barrier, achieves a separation by allowing certain
components in a fluid mixture to pass through while rejecting others, thus resulting in
a preferential passage of certain components (Mulder, 1996). For an effective gas
separation process, the membrane materials shall be non-porous and have no defects
on a molecular level. The driving force for the permeation of gas penetrants through
the non-porous membranes is the chemical potential difference resulted from the
concentration difference at the upstream and downstream membrane sides (Koros and
Fleming, 1993). Separation is achieved as a consequence of the difference in the

relative transport rates of different penetrating gas molecules (i.e. components that
diffuse faster will be enriched in the permeate stream, while the other components will

1
become concentrated in the retentive stream). The transport properties of the non-
porous membranes have been widely characterized by investigating the permeability to
gases, the permselectivity to certain species over others, and the sorption of gases in
the materials.

1.2 Transport mechanism of membranes
1.2.1 General principles

For a homogeneous polymer dense membrane, the local flux N of a penetrant for one-
dimensional diffusion can be described by Fick’s law:
)(
x
C
DN
∂
∂
−=
(1.1)
where D is the diffusion coefficient, which might be a function of local concentration,
C, in the film, x refers to the distance from the upstream side of the film to
downstream side(x=0 upstream, x=L downstream). The driving force,
X
C
∂
∂
, represents

the gradient of penetrant concentration through the membrane, which can be replaced
by pressure gradient when Henry’s law holds.

The most important characteristic, permeability coefficient P, is defined as the flux N
normalized by pressure drop and membrane thickness, as shown below:
lpp
N
P
)(
12
−
=
(1.2)
where p
2
and p
1
are the pressures of penetrants at upstream and downstream sides,
separately, l denotes membrane thickness. (Staudt-Bickel and Koros, 1999)


2
1.2.2 Solution-diffusion model

The gas transport in most rubbery and glassy membranes can be explained by the
solution-diffusion model. There is no continuous transport passage for the gas
penetrants, but the thermally agitated motion of polymer chain segments generates
penetrant-scale transient gaps, thus the diffusion of penetrants from the feed stream to
permeate stream is achieved. Therefore, this prime “solution-diffusion” mechanism
consists of three steps, that is, the gas molecules in the upstream gas side (high-

pressure side) first sorb into the membrane surface, then diffuse across the membrane,
and finally desorb from the membrane surface on the downstream gas side (low-
pressure side) (Wijmans and Baker, 1995). As a result, the permeation coefficient can
be expressed as a product of a diffusion coefficient D and a solubility coefficient S:
DSP ⋅=
(1.3)
The solubility coefficient S is determined by the condensability of penetrants, the
interactions of polymer and penetrant, and the amount of free volume and its
distribution of the polymer; while the diffusion coefficient, D, is determined by the
packing and mobility of polymer segments as well as the size and shape of the
penetrating molecules.

The permselectivity is another critical parameter in membrane-based gas separation
processes, which is characterized by a separation factor α
A/B
in terms of the
downstream (y) and upstream (x
s
) mole fractions of two components A and B as
shown by the following equation:
BsAs
BA
BA
xx
yy
,,
/
/
/
=

α
(1.4)

3
In the case of negligible downstream pressure, α
A/B
is equal to the ideal separation
factor (permselectivity), α*
A/B
,

defined as the ratio of permeabilities of the two gases A
and B under mixed gas feeding conditions.
B
A
BA
P
P
=
*
/
α
(1.5)

1.2.3 Sorption in glassy polymers – Dual mode sorption model

The sorption of a certain kind of gas in glassy polymer membranes is a thermodynamic
process and the solubility coefficient is determined by (i) the inherent condensability of
the penetrant, (ii) the polymer–penetrant interactions, and (iii) the amount and
distribution of the excess free volume in the glassy polymer (Paul and Koros, 1976;

Koros et al, 1976; Chung and Kafchinski, 1997). As put forward by Koros (Koros,
1977), the equilibrium concentration of the sorbed gas in glassy polymers can be
described as a function of pressure, p, if the dual mode sorption model is to be applied:
bp
bpc
pkCCC
H
DHD
+
+=+=
1
'
(1.6)
where C is the total penetrant concentration, C
D
and C
H
represent the local
concentration of penetrant molecules sorbed in the Henry mode and Langmuir
environments, respectively. The parameters k
D
, c’
H
and b are the Henry law constant,
the Langmuir capacity constant, and the Langmuir affinity constant, separately, which
can be obtained by a non-linear least squares fitting of the sorption data. The Henry
sorption refers to the dissolution of a gas in the densified polymeric regions, while the
Langmuir sorption refers to the trapping of gas molecules in the unrelaxed volume of a
polymer matrix below glass transition temperature, T
g

. The Langmuir capacity, c’
H
,
represents the maximum amount of penetrant that can be sorbed in the Langmuir

4
environments of a glassy polymer. The Langmuir environments of a glassy polymer
refer to the excess free volume frozen in as a result of the non-equilibrium quenching
from the rubbery state to glassy state. Therefore, the Langmuir capacity of amorphous
polymers will disappear in rubbery state, that is, above Tg, or the transition time from
rubbery state to glassy state lasts long enough. The Langmuir affinity constant, b,
measures the rate of the sorption over desorp tion for penetrant in the Langmuir mode
(Koros et al, 1979). Thus by definition, the solubility coefficient is given as:
bp
bc
k
p
C
S
H
D
+
+==
1
'
(1.7)

1.3 Membrane material selection and tailoring

Compared with rubbers and crystalline / semi-crystalline polymers, glassy polymers

have emerged as the preferred materials for gas separation for the advantageous
combination of permselectivity and permeation properties (because the chain mobility
in crystalline / semi-crystalline polymers is relatively small in the highly ordered and
restricted crystalline structure, and the chains are too mobile in rubbers to achieve a
good selectivity, while amorphous polymer stands in between). For most available
polymers, the characteristics of permeability and selectivity are generally contradictive
in nature: a trade-off between permeability and permselectivity, i.e. high
permselectivity is coupled with low permeability and vice versa, has been commonly
observed. This can be illustrated by a typical famous Robeson’s “upper bound” curve
of permeability and permselectivity for CO
2
/ CH
4
separation of various polymers as
shown in Figure1.1 (Robeson, 1991). It is clearly shown that glassy polymers possess
much lower permeabilities but relatively high permselectivities than rubbers.

5

Figure 1.1 Literature data for CO
2
/CH
4
permselectivity versus CO
2
permeability
(Robeson, 1991)

Certainly, this “upper bound” curve is not fixed. Numerous efforts have been put in
pushing the limits of current available polymers such as molecular design and tailoring,

polymer blending, inorganic-organic mixing etc With the development of polymer
material, the “upper bound” curves gradually move up. The objective of membrane
material selection is to look for the polymer candidate beyond the “upper bound”.


Since the transport properties of polymeric membrane are attributed to the combination
of the contribution from several factors: 1.total free volume; 2. distribution of free
volume; 3. intersegmental resistance to chain motions; 4. intrasegmental resistance to
chain motions (Coleman and Koros, 1994), two principles are generally considered for
membrane material selection (Coleman et al, 1993; Coleman and Koros, 1994): 1: A
family of polymer materials will tend to increase permeability while maintaining
permselectivity through the structural alterations, which inhibit chain packing with an

6
inhibition to rotational mobility of flexible linkages on the polymer backbone; 2: A
family of polymer materials will tend to decrease permeability with desirable increases
in permselectivity through the structural alterations, which significant suppress the
segmental mobility while causing only small changes in chain packing.

Among most available polymers such as polysulfone, polyimides, polyamides,
polycarbonates, polyetherimide and sulfonated polysulfone, 6FDA
(hexafluorodianhydride)-based polyimides have attracted much attention for gas
separation due to both impressive gas performance with many other desirable
properties such as spinnability, thermal and chemical stability and mechanical strength
as compared with non-fluoropolyimides (Coleman et al, 1993; Coleman and Koros,
1994; Costello and Koros, 1995; Kawakami et al, 1997; Zimmerman et al, 1998;
Staudt-Bickel and Koros, 1999, 2000; Zimmerman and Koros, 1999a, 1999b). The
6FDA-based polyimides possess better gas performance with high permeability and
permselectivity because their rigid structure contain bulk groups of (CF
3

), by which
the efficient packing is inhibited and local segment mobility is reduced. For the
advantages and prospects in large-scale application in industry, the 6FDA-Durene has
been chosen to study in this work.

1.4 CO
2
plasticization and physical aging of glassy polymer

An important application of gas separation membranes is to remove acid gas from
natural gas. Natural gas is a complex mixture that contains the desirable components
such as hydrocarbons and some unpleasant components such as CO
2
, H
2
S and water

7
vapor. Not only are the acid gases corrosive to pipelines but also they reduce the
energy content of the natural gas.


The CO
2
induced plasticization refers to the increase of CO
2
permeability as a function
of feed pressure (Bos et al, 1999; Ismail and Lorna, 2002). According to the solution-
diffusion mechanism, the permeation of penetrants in glassy polymer membranes is
controlled by two aspects: solution and diffusion (Paul and Yampol’skii, 1994; Koros

and Fleming, 1993; Stern, 1994; Koros, 1977). In most glassy polymers, the diffusion
coefficient has more contribution to the permeability. Being a kinetic factor, the
diffusion coefficient is correlated to the packing and motion of polymer segments, and
the size and shape of penetrating molecules. As a plasticizer, CO
2
may either swell up
the interstitial place among polymer chains, which brings up a larger free volume or /
and enhance segmental and side groups mobility (Paul and Yampol’skii, 1994; Koros
and Fleming, 1993; Staudt-Bickel and Koros, 1999; Bos et al, 1999; Ismail and Lorna,
2002; Koros, 1977, Bos et al, 1998a; Wessling et al, 1991). Though the CO
2
-induced
plasticization accelerates the diffusion of penetrants, simultaneously, it severely
deteriorates the gas permselectivity of CO
2
over other gases. Therefore, many efforts
have been put into diminishing the effect of plasticization caused by CO
2
on the
membrane separation performance.

Physical aging phenomenon is not a negligible factor for most of glassy polymeric
membranes because it will lower their gas performances. Physical aging of glassy
polymers stems from the non-equilibrium nature of the glassy states towards
equilibrium, which is associated with the volume relaxation of polymers below Tg.
Consequently, the segmental mobility of the polymer chains is reduced (Kovacs, 1958;
Chan and Paul, 1980; Chow, 1984; Bartos et al, 1990; Hutchinson, 1995; McCaig and

8
Paul, 2000). Many materials properties such as viscoelastic, mechanical, electrical and

calorific properties, will change with the increase in storage time of polymer
membranes under the conditions of constant temperature, zero stress and without
external forces (Struik, 1978; Aref-Azar et al, 1983; Carfagna, et al, 1988; Vigier and
Tatibouet, 1993; Hill et al, 1990; Bradshaw and Brinson, 1997). In the meantime,
physical aging can also dramatically deteriorate the gas permeability of glassy
polymeric membranes.

The solutions to these two issues are vital to the wide application of membrane-based
gas separation. Additionally, the fact that the penetrants might act as “lubricant” to the
segmental adjustment of chains is also worthy of consideration.

1.5 Why thin films?

Besides the property-oriented molecular tailoring, the introduction of the asymmetric
membranes, which typically consist of a thin selective layer and a porous support layer,
has been a breakthrough towards high productivity and high membrane area ratio in
membrane development because the productivity of this kind of membranes is
inversely related to the thickness of the effective layer. Hollow fibers of glassy
polymers with an ultra-thin selective layer have been extensively employed for large-
scale industrial membrane-based gas separation processes because of the high gas
permeance and nearly intrinsic gas selectivity. However, the physical aging,
characterizing in the drastic gas permeance or flux drop, has been commonly observed
in the selective layer (Rezac et al., 1993; Chung and Teoh, 1999; Chung and Kafchinski,
1996; Pinnau, 1991). Our works were tailored to examine the aging profile of 6FDA-

9
Durene polyimide dense films with different thickness, thus to correlate the aging of
hollow fiber containing a thin and dense selective layer with the aging of dense films of
comparable thickness.

Though extensive works have been carried out for plasticization study, most of the
observations were obtained from the thick dense films (typically around 50 µm). For
example, Bos extensively studied the plasticization behavior of thick dense films of
commercial polyimide Matrimid 5218 (Bos, 1996). Up to date, there are few reports
on the plasticization behavior of thin dense films, which is similar to the case of the
plasticization of thin layer of asymmetric membranes, and is suitable for the study that
seeks to understand and suppress the plasticization behavior. In fact, that has
motivated us to conduct the direct thin film plasticization studies, and as a result
strikingly different plasticization behaviors of films of different thickness and possible
mechanisms are presented here.

1.6 Goals and organization of this research

The goals of this work were (i) to examine the aging profile of 6FDA-Durene
polyimide dense films with different thickness, thus to correlate the aging of hollow
fiber containing a thin and dense selective layer with the aging of dense films of
comparable thickness; (ii) to investigate the CO
2
plasticization behaviors of 6FDA-
Durene films with different thickness; (iii) to study the effects of chemical cross-
linking modification of 6FDA-Durene on the aging and CO
2
plasticization behaviors.


10