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Processing, thermal-oxidative stability and thermal cyclic fatigue of phenylethynyl-terminated polyimides
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Processing, thermal-oxidative stability and thermal cyclic fatigue of phenylethynyl-terminated polyimides
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Content
Processing, Thermal-Oxidative Stability and Thermal Cyclic Fatigue of
Phenylethynyl-Terminated Polyimides
by
Xiaochen Li
A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(Chemical Engineering)
May 2021
Copyright 2021 Xiaochen Li
ii
天 凉
好 个 秋
iii
DEDICATION
I dedicate this work to my parents and husband, who offered their unconditional love for me.
iv
ACKNOWLEDGEMENTS
I would like to thank my advisors Prof. Steven R. Nutt, his patient guidance and instructions. I’d
like to thank my colleges at USC M. C. Gill Composites, they have offered numerous help during
my Ph. D. I would like to send my appreciation to my parents, my husband, and my daughter, for
their love and care.
v
TABLE OF CONTENTS
Epigraph…………………………………………………..……………………………………...ii
Dedication………………………………………………………………………………………..iii
Acknowledgements………………………………………………………………….…...……...iv
List of Tables…………………………………………………………………………….……..viii
List of Figures………………………………………………………………………………..….ix
Abstract………………………………………………………………………….…………….…xi
Chapter 1. Imidization of phenylethynyl-terminated PMDA-type polyimide of the
Monomeric Reactant Process……..………… ……………………………..……...………..1
1.1 Introduction……..…………………………………………………...…..………………….1
1.2 Experiments……..…………………………………………………....……………………..3
1.2.1 Material Preparation ……..……………………………………………………....…………3
1.2.2 Thermogravimetric Analysis (TGA) …………………………………….……...….………3
1.2.3 Differential Scanning Calorimetry (DSC) ………………………………………….………3
1.2.4 Rheological Test…………………………………………...……………………….……….4
1.3 Results and Discussion……………...…………………………...………………...…………5
1.3.1 Volatile release during imidization and crosslinking………………………….…………….5
vi
1.3.2 Heat flow during imidization……………………………………….…...…………...……..10
1.3.3 Viscosity Profile during Imidization…………………………………………...…………..11
1.3.4 Different Chain Length…..………………………………..……….…...…………...……..14
1.3.5 Imidization Kinetics………………………………………..…..........................…………..15
1.4 Conclusion…………………………………………...……………………………………....17
Chapter 2. Thermal Oxidation of PEPA Terminated Polyimide…………………....………18
2.1 Introduction……..……………………………….…………………...……………………..18
2.2 Experiments……..…………………….……………………………...……………………..21
2.2.1 Material Preparation and Aging Condition …………..………………………...…………21
2.2.2 Weight Change …………………….……….…………………………….……...………...21
2.2.3 DMA …………………………………………….………………………………….……...22
2.2.4 Nanoindentation …………………………………………...……………………...……….22
2.2.5 FTIR ……………………………………………………….……………………….………23
2.2.6 XPS ………………………………………………………...……………...……….………23
2.3 Results and Discussion…………………………………………...…………………………25
2.3.1 Weight loss ………….……………………………………………….…………...………...25
2.3.2 Thermal Analysis ………………………………...……………………….……...………...28
2.3.3 Nanoindentation Measurements…………………………………………………….……...31
2.3.4 FTIR Analysis ……………………………………………...………………………………33
2.3.5 Surface XPS analysis ………………………………………………………………..……..42
2.4
Conclusion…………………………………………...……………………………………….....48
vii
Chapter 3. Thermal Oxidative Aging and Thermal Cycling of PETI-340M
Composite ..………………………………………………………………...……………………49
3.1 Introduction……..……………………………….…………………...……………………..50
3.2 Experiments……..…………………….……………………………...……………………..52
3.2.1 Material Preparation and Aging Condition …………..………………………...…………52
3.2.2 DMA …………………….……………………….……………………….……...………...52
3.2.3 Mechanical Tests ………………………….………….…………………………….……...53
3.2.4 Morphology Analysis …………………...………………………………………………….53
3.3 Results and Discussion …………………………………………...……………….………..53
3.3.1 Weight loss ………….……………………………………………..….…………...……… 53
3.3.2 Thermal Analysis ………………………………...………………….…….……...………..55
3.3.3 Morphology …………………………………………………….………………………….58
3.3.4 Mechanical Measurement …………………………...…..…………………………………64
3.4 Conclusion…………………………………………...……………………………...…….....69
B ib li ogr ap h y………… ………………… ……………. ……………………………………..…..71
viii
List of Tables
Table 1.1 Formulations of TriA-X with different diester ratio and isomerism……………..…….3
Table 1.2 Theoretical and accrual weight loss during imidization process of Tri-A resin………..8
Table 1.3 Melting temperature of different TriA-X formulations……………………………….12
Table 1.4 Imidization Kinetics of TriA-X formulations ……………………………………..…16
Table 2.1 Weight loss rate comparison for PMR-15 and TriA-X……………………………….26
Table 2.2 Area percentage of peaks for C 1s and O 1s of unaged and 316 ° C 1000 hr aged TriA
X………………………………………………………………………………………………….46
Table 3.1 Weight change per unit area for thermal oxidative aging of PETI-340M composite with
nonlinear curve fitting……………………………………………………………..……………..55
Table 3.2 Change of glass transition temperature for PETI-340M composites aged at 232 ° C and
288 ° C……………………………………………………………………………………………57
Table 3.3 Coefficients for the mechanical property degradation of PETI-340M composites aged at
232 ° C and 288 ° C………………………………………………………………………….……66
ix
List of Figures
Figure 1.1 Weight loss during imidization process of Tri-A resin……………………….………7
Figure 1.2 Volatile analysis of TriA-X imidization by TGA- mass-spectroscopy…………….....9
Figure 1.3. DSC of the imidization heat flow for different TriA-X formulations...........…..……11
Figure 1.4. Viscosity change during the imidization for different TriA-X formulations……..…12
Figure 1.5. Heat flux profile of powder mixture with different monomer ratio during imidization
using a dynamic heating rate of 10˚C/min ……………………………………………….…..….15
Figure 1.6. Tg of cured resin produced from imide oligomer with different n…………..……….15
Figure 2.1 Weight loss per unit area of TriA X during isothermal oxidative degradation at 288 ° C
and 316 ° C………………………………………………………………………….…………….25
Figure 2.2 Cross-section of TriA neat polymer a) Unaged; b-1)204˚C, 700 hr; b-2) 204˚C, 2000
hr; c-1) 288˚C, 700 hr; c-2) 288˚C, 1500 hr; d-1) 316˚C, 700 hr; d-2) 316 ˚C, 1500 hr …………28
Figure 2.3. a) tan(δ) of TriA X during aging at 288 °C, b) tan(δ) of bulk and interior TriA X after
aging ……………………………………………………………………………………………..30
Figure 2.4 a) Nanoindentation hardness of the unaged and aged resin surface and interior, b)
Nanoindentation modulus of the unaged and aged resin surface and interior……………………32
Figure 2.5 a) FTIR spectra of fresh TriA X vs. 316 ° C 500 hr aged resin surface and interior. b)
FTIR spectra of TriA X aged at different temperatures and times………………………………..35
Figure 2.6 FTIR spectra of fresh TriA X vs. resin surface after aging at 316 ° C for different
times………………………………………………………………………………………….…..36
Figure 2.7 Content of chemical groups of the 288 ° C 700 hr aged resin oxidized layer………….38
x
Figure 2.8. Instable groups of TriA X in thermal oxidative environment………………………..40
Figure 2.9. a) FTIR intensity of the released volatile and weight loss percentage during the TGA
dynamic ramp. b) FTIR spectra of the volatile released at 545 ° C and 625 ° C. …………………41
Figure 3.1 Weight change per unit area for thermal oxidative aging of PETI-340M composite with
nonlinear curve fitting……………………………………………………………………………54
Figure 3.2 Dynamic mechanical analysis of unaged and thermal oxidative aged PETI-340M. (a)
tan(δ); (b) Storage modulus………………………………………………………………………56
Figure 3.3 Microscopy image of cross-sections of unaged and aged PETI-340M Composites….58
Figure 3.4. A polished section of PETI-340M Composites aged at 288 ° C for 2500 hr…………59
Figure 3.5. Microscopy image of cross-sections of PETI-340M thermal-cycled for 400 cycles and
1600 cycles……………………………………………………………………………………….60
Figure 3.6. A crack of PETI-340M composites thermal-cycled for 1600 cycles…………………61
Figure 3.7. Micro-CT images of 2400 thermal cycled specimen. (a) 3-D geography; (b) an in-plan
cross-section………………………………………………………………………..…………….63
Figure 3.8. Mechanical change of thermal oxidized PETI-340M composites. (a) Combined loading
compressive strength; (b) Open-hole compressive strength; (c) Short-Beam Shear Strength…….65
xi
Abstract
Polyimides (PI’s) are widely used in high-temperature composite applications because of the high
glass transition temperatures and thermal stability. However, polyimides generally suffer from
poor processability, cracking induced by thermal shock, and thermal oxidation. A major limitation
of conventional polyimide resins is low toughness, which can lead to (premature) brittle failure.
However, conventional thermoset polyimides such as PMR-15, show low elongation-at-break due
to a rigid molecular structure and a low degree of polymerization of imide oligomer. The brittle
nature of the resin causes microcracks to propagate under repeated thermal cycling, which severely
limits service life. Moreover, the low molecular mobility at high temperature (>Tg) leads to
processing challenges and makes it difficult to obtain composite parts with low porosity. Another
concern with the polyimide chemistry is that methylene dianiline (MDA), which has been used in
traditional PMR polyimide formulations, is known to be carcinogenic.
In recent decades, 4-phenylethynylphthalic anhydride (PEPA) end-capped polyimides have been
developed to address such limitations, and these polymers generally offer wider process windows
and improved heat resistance, although they are not immune to oxidation and thermal shock
damage. A new class of PEPA-terminated PMDA-type polyimide (pyromellitic dianhydride),
denoted TriA X1, has been developed in an attempt to address the limitations of brittleness and
poor processability associated with most polyimides. The TriA X polyimide is derived from
PMDA and 2-phenyl-4, 4’-diaminodiphenyl ether (p-ODA). The pendent phenyl group of p-ODA
restricts molecular interaction and decreases the rotational flexibility of the chain. The distinctive
chemical structure of the polymer imparts a reduced elastic modulus during heating (from 109 Pa
to 107 Pa before and after glass transition), which facilitates the processing of carbon fiber
xii
composites.
1
Other PEPA-terminated polyimides have also been developed, for example, PETI-
340M, featured in their low viscosity during processing. However, for PEPA- terminated
composites, its imidization kinetics and the effect of different diester groups on the curing kinetics
were not fully understood. The thermal-oxidative stability and the stability under thermal cyclic
fatigue were not understood as well. Therefore, the motivation of this dissertation is to investigate
the imidization behavior of PEPA-terminated polyimide, and to evaluate the retention of
mechanical property in thermal-oxidative aging conditions and thermal cycling.
xiii
1
CHAPTER I
Imidization of Phenylethynyl-Terminated PMDA-type Polyimide of the Monomeric
Reactant Process (PMR)
In recent decades, a phenylethynyl-terminated PMDA-type polyimide, based on a KAPTON
backbone structure in PMR process, TriA-X, was developed.
1
This high-temperature polymer
exhibits greater elongation-to-break compared to existing polyimide systems. The objective of this
study was to compare the imidization behavior of TriA-X formulations with different diesters.
Dimethyl, diethyl, and diisopropyl esters of pyromellitic dianhydride (PMDA) have been applied
for synthesizing amorphous and asymmetric thermoset polyimide with phenyl-oxydianiline (P-
ODA) and monoethyl ester of phenylethynylphthalic anhydride (PEPA MEE) in PMR process.
Imidization was investigated by Thermogravic Analysis-Mass Spectrum (TGA-MS) and
Rheometer for monomer mixtures with different PMDA diesters. Diethyl esters of PMDA with
different para- and meta- isomer ratios were also compared for studying the isomer ratio’s effect
on cure behavior. Different imidization behavior was attributed to steric hindrance, stability of
intermediate, crystal lattice of formulations with different diesters.
1. 1 Introduction
Recently, a new class of phenylethynyl-terminated PMDA type polyimides, denoted TriA X (TriA:
amorphous, asymmetric and aromatic, X: crosslinking), derived from PMDA and 2-phenyl-4, 4’-
diaminodiphenyl ether (p-ODA), was developed to effectively resolve the issues of brittleness and
poor processability
1
. In TriA X, the pendent phenyl group of p-ODA restricts molecular interaction
2
and decreases the rotational flexibility of the chain. For this resin system, a reduction in storage
modulus E’ during heating is achieved (from 109 Pa to 107 Pa), which greatly increases
processability, specifically for application in carbon fiber reinforced composites
2
. Using 4-
phenylethynyl phthalic anhydride (PEPA) as of end caps, the TriA X polyimide features a broader
processing window, as the crosslinking temperature for PEPA is greater than 320° C. Fundamental
thermal and mechanical properties of this PMDA/p-OPA polyimide resin, for example, glass
transition temperature, storage/loss modulus, and crystallization behavior, as reported by Miyauchi
and co-workers, have demonstrated that this resin is a potential candidate to replace conventional
polyimides in high-temperature composites
2
.
For the PMR process of TriA X, the first reaction is between the diester of pyromellitic dianhydride
(PMDA) and 2-phenyl-4,4’-diaminodiphenyl ether(p-ODA). Poly (amic acid) is formed with the
elimination of the corresponding alcohol. Imidization is followed by formation of the imide ring
while releasing the water molecule. In the step of formation of poly (amic acid), different type of
PMDA diester group, for example, dimethyl ester (DME), diethyl ester (DEE), and di-iso-propyl
ester (DiPE) has the different steric effect and different alkalinity, which leads to different reaction
kinetics and different process condition. Research regarding comparison and selection of ester
group hasn’t been done in previous research.
The initial step to synthesis the monomer mixture involves the esterification of PMDA. Both meta-
PMDA diester and para- PMDA diester can be generated by the esterification. Huang et al
3
reported that meta- PMDA diester has a denser and tighter packing in the unit cell than the para-
PMDA diester, which may lead to different reactivity. Volksen et al
4
showed that for the poly
(amic ethyl ester) concentrated solution, meta- isomer has lower viscosity than para- isomer.
Milevshaya et al showed that, in solution condition-, meta- poly (amic ester) geometry was more
3
difficult in attaining correct conformation for cyclo-imidization. But no study has been done on
the isomerism’s effect on polyimide resin system with PMR process.
5
In this study, we are going to investigate the imidization behavior of TriA X with different PMDA
diester groups and different isomer ratios by the PMR method. We will look into the volatile
release, heat flux, and viscosity of the resin during cure. Effect of electrophilicity, basicity and
steric hindrance of the diester group were found has influence on the curing process.
3
1.2 Experiments
1.2.1 Material Preparation
Three constituent chemicals - PMDA (Tokyo Kasei Kogyo, Japan), 2-phenyl-4,4’-
diaminodiphenyl ether (p-ODA) (Wakayama Seika Kogyo, Japan), and PEPA (Manac, Japan) -
were provided (Kaneka Corporation). The monomer mixture was prepared by mixing the above
chemicals into a hot ethanol solution, followed by vacuum drying. 4 types of PMDA with different
ester groups or different ratio of cis- trans- isomers was used for preparing the polyimide monomer
mixture: (1) PMDA DME, cis-:trans-=1:1, (2) PMDA DiPE, cis-:trans-=1:1, (3) PMDA DEE,
cis-:trans-=1:1, (4) PMDA DEE, cis-:trans-=3:1, shown in Table 1.1. The monomer mixture was
prepared by dissolving PMDA with p-ODA and PEPA in hot ethanol, followed by the vacuum
dry-up. The ratio of PMDA, p-ODA, and PEPA is 7:8:2, which was confirmed by the NMR spectra.
Formulations with different backbone length (n) was also prepared and shown in the following
table.
4
Formulation Ratio of
para:meta (T:C)
isomers
Name
N=7, DEE T:C=1:3 DEE T1C3
N=7, DEE T:C=1:1 DEE T1C1
N=7, DiPE T:C =1:1 DIPE
N=7, DME T:C=1:1 DME
N=4, DEE T:C=1:1 N=4
N=1, DEE T:C=1:1 N=1
N=0, DEE T:C=1:1 N=0
Table 1.1 Formulations of TriA-X with different diester ratio and isomerism
1.2.2 Thermogravimetric Analysis (TGA)
The thermogravimetric analysis (Q800, TA Instruments) was employed to study the weight loss
during dynamic and isothermal heating cycles. Dynamic TGA tests were conducted with ramp
rate 5 ˚C/min from 25˚C to 300 ˚C and holding at 300 ˚C for 60 min. The isothermal TGA tests
were conducted with a rapid temperature jump to the targeted temperature and hold for 90 min.
1.2.3 Differential Scanning Calorimetry (DSC)
The heat flux profile for imidization was obtained with DSC measurements (Q2000, TA
instrument). The DSC test was conducted dynamically from 25˚C to 300˚C with ramp rate 5˚C/min
in modulated mode. Nitrogen is used as purge gas under ambient pressure. Before the dynamic
run, a dynamic ramp from 25˚C to 100˚C with ramp rate 5˚C/min is applied to the sample as a pre-
process procedure to remove most of the solvent residue. A hole was drilled on the top of the pan
cover to prevent explosion in the DSC chamber due to the foaming effect of imidization.
1.2.4 Rheological Test
5
Rheological tests were conducted by an AR 2000 ex (TA Instrument) with parallel plates. A
hydraulic pump was used to press monomer mixture powders into pellets. A normal force control
at 1N was applied for all measurements to ensure the contacting of upper plate the lower plate
during the test period especially for the bubbling of imidization. Dynamic rheology tests were
conducted from 30˚C to 300˚C with ramp rate 5˚C/min followed by the isothermal hold at 300˚C.
1.3 Results and Discussion
1.3.1 Volatile release during imidization and crosslinking
The volatile release during molding of polyimide composites is crucial because imidization may
bring significant voids thus reducing the quality of the parts. Dynamic TGA tests with 5 ˚C /min
were conducted to determine the starting of imidization, the most rapid release of volatile, and the
required temperature to reach full imidization. The reaction scheme for the TriA-X formulation
with DEE diester is shown in Scheme 1.1. The scheme shows that water and ethanol molecules
were released during imidization.
6
Scheme 1.1. Scheme of imidization of TriA X (DEE).
The weight loss profiles are shown in Figure 1.1. The initial minor weight loss is due to remove
of the solvent residue in the mixture powder. The major weight loss was observed during 100˚C
to 250˚C. DME shows lowest imidization temperature around 120˚C. DIPE shows the highest
imidization temperature around 143˚C. Imidization temperature of DEE T1C3 (127.8˚C) was
observed to be almost the same compared with the imidization temperature of DEE T1C1
(127.9˚C).
The theoretical volatile generation percentage and the actual volatile generation percentage are
shown in table 1.2. The error of the theoretical value and actual value is assumed to be due to the
7
overlap of solvent release and imidization volatile at the early stage. All formulations reached fully
imidization during the dynamic ramp. No weight loss is observed at temperature higher than 350
˚C, which is the crosslinking stage of the PEPA endcap. All formulations reached fully imidization
during the dynamic ramp. No weight loss is observed at temperature higher than 350 ˚C, which is
the crosslinking stage of the PEPA endcap.
Figure 1.1. Weight loss during imidization process of Tri-A resin.
50 100 150 200 250 300 350 400
75
80
85
90
95
100
Weight (%)
Temperature ( C)
DME
DEE
DIPE
T1C3
8
Formulation
Initiation temp.
˚C
Theoretical Volatile
generation
%
Actual Volatile
generation
%
DME 98.9 17.36 16.97
DEE
T1C1
120.9 20.62 20.00
DIPE 133.5 23.65 24.67
DEE T1C3 120.3 20.62 20.92
Table 1.2. Theoretical and accrual weight loss during imidization process of Tri-A resin.
9
100 200 300 400 500
0
50
100
0
50
100
0
50
100
0
50
100
DME
Temperature ( C)
Methanol
Weight Loss Rate
Relative Volatile Release
DEE
Ethanol
Weight Loss Rate
DIPE
Iso-propanol
Weight Loss Rate
T1C3
Ethanol
Weight Loss Rate
Figure 1.2. Volatile analysis of TriA-X imidization by TGA- mass-spectroscopy
During imidization, the bond of the carbonyl group was broken by attacking the unshared pair
of electrons on the amine nitrogen, which leads to the formation of the intermediate with the π
electron moving to the carbonyl oxygen. During rearrangement, oxygen from the alkoxy group
attracted hydrogen from the amine and leave the polymer chain as R-OH. The formation of this
intermediate and releasing of R-OH can be affected by several factors. And the first dominating
factor is the electrophilicity of the carbonyl carbon of the ester group. The formulation with the R-
with the strongest electrophilicity, DME, will react at the lowest temperature. The other effect is
the hindrance of the R-group. Due to the larger hindrance of R-, The DIPE formulation needs the
highest temperature to initiate the reaction.
10
According to the reaction scheme, alcohol and water are released from the reaction. The weight
loss profiles are observed as one stage smooth curves which indicates alcohol and water are
generated from the reaction simultaneously, this was proved by the mass-spectra results, shown in
Figure 1.2 .The result shows that, water and the corresponding alcohol was released almost
simultaneously. Therefore, once the poly (amic acid) is formed, the chain will release the H 2O
molecule to form the imidized ring rapidly. Therefore, the formation of poly (amic acid) is the
rate-limiting step. This result provides an important information that vacuum has to be applied to
remove the volatile during the composites molding.
1.3.2 Heat flow during imidization
The heat flux profile of monomer mixture powder with different diester during imidization is
shown in Figure 1.3. The heat flux curves show that both the reaction between PEPA end cap with
p-ODA and PMDA diester with p-ODA are endothermic reactions starting from ~125 ˚C, which
corresponds to the TGA data. The initial transition is assumed to be related to the removal of
residue solvent and the realignment of p-ODA and PMDA diester molecules caused by
dehydration.
11
Figure 1.3. DSC of the imidization heat flow for different TriA-X formulations.
1.3.3 Viscosity Profile during Imidization
The viscosity of the resin during imidization is a crucial factor to produce void-free composites
parts since it can help design an optimized process of pressure, vacuum and temperature. In the
respect of chemistry, viscosity is also a key factor to track because the increase of mobility brings
difficulty for the formation of poly (amic acid) chain. Dynamic rheological tests with 5˚C/min
ramp rate were conducted to investigate the viscosity change.
During imidization in PMR process, with the increase of temperature, the monomer mixture first
melt, which is observed as a major drop in viscosity, and then undergoes imidization with
increasing of viscosity by chain formation. Fig. 1.4.The formation of poly (amic acid) becomes a
diffusion-controlled reaction with the increase of the chain length, whereas imidization is not
diffusion-controlled since it’s an intramolecular reaction. Initially, the melting temperature of the
monomer mixture is measured as Table 1.3. The Cis-rich monomer mixture shows a lower melting
50 100 150 200 250 300 350 400
-4
-2
0
2
4
6
8
Heat Flow (mW)
Temperature ( C)
DME
DEE T1C1
DIPE
DEE T1C3
12
temperature due to the lower crystallinity caused by the lower symmetry. The PMDA dimethyl
ester showing the highest melting temperature is also attributed to the higher crystallinity due to
the smaller steric hindrance of methyl group.
50 100 150 200 250 300
0.01
1
100
10000
1000000
|n*| (Pa.s)
Temperature ( C)
DME
DEE T1C1
DIPE
DEE T1C3
Figure 1.4. Viscosity change during the imidization for different TriA-X formulations.
Melting
Temperature
˚C
Starting of
Imidization
˚C
DME 89.1 121.7
DEE 76.1 125.5
DIPE 69.8 134.5
T1C3 71.0 124.5
Table 1.3 Melting temperature of different TriA-X formulations
13
According to the rheological data, the imidization temperature of DME is lower than other
formulations with a sharp peak. In contrast with DME and DEE, DIPE has the most spreading
peak. To understand the effect of ester group, the reaction mechanism of the formation of poly
(amic acid) is shown in Scheme 1.2. An active intermediate is formed by attack of a nitrogen
nucleophile to the carbon on the ester group. By this process, a weaker base, alcohol, is produced
by consuming of amine, which is a stronger base.
Scheme 1.2. Formation of the poly (amic acid)
From this scheme, we identified two important factors affecting the reaction rate. The first factor
is the steric hindrance of the hydrocarbon group on the ester group, which in our case are methyl,
ethyl and isopropyl. Isopropyl, the hydrocarbon group with the largest steric hindrance, brings
most difficulty for the attacking compare to methyl and ethyl, thus it requiring higher energy to
stimulate the reaction. The second factor is the basicity of the released alcohol. Isopropanol, with
higher basicity than methanol and ethanol due to the stronger nucleophilicity of isopropyl ,is more
difficult to be released. The high steric hindrance of the isopropyl and high basicity of isopropanol
together make the intermediate requiring higher energy to be generated and more stable to exist,
14
which gives DIPE a higher imidization temperature and more spreading peak in the viscosity
profile.
1.3.4 Different Chain Length
The heat flux profile of monomer mixture powder with different monomer ratio during imidization
is shown in Figure 1.5. The heat flux curves show that both the reaction between PEPA end cap
with p-ODA and PMDA diester with p-ODA are endothermic reactions starting from ~125 ˚C,
which corresponds to the TGA data. However, obviously, the heat of the imidization reaction
between PEPA end cap and p-ODA is greater than the heat absorption of the imidization between
p-ODA and PMDA diester. The peak temperature of imidization is listed as n = 0: 149 ° C, n =
1:151 ° C, n = 4: 153 ° C, and n = 7: 153 ° C. The trend in the starting temperature of imidization
also matches the TGA curves, that is, PMDA diester requires higher temperature to react with p-
ODA than PEPA does.
The Tg values of cured resin generated from imide oligomer with different degrees of
polymerization, n = 1, n = 4, and n = 7, were measured by DMA and are shown in Figure 1.6. The
Tg values of the three formulations are n = 1 > n = 4 > n = 7, because the value of Tg increases
with increase in degree of crosslinking. With higher weight percentage of PEPA end cap, n = 1
shows the greatest degree of crosslinking which leads to the highest Tg. However, with a lower
degree of polymerization of the imide oligomer, n = 1 could show more brittleness compared to n
= 4 and n = 7, a subject of future investigation.
15
Figure 1.5. Heat flux profile of powder mixture with different monomer ratio during imidization using a dynamic
heating rate of 10˚C/min
Figure 1.6. Tg of cured resin produced from imide oligomer with different n.
1.3.5 Imidization Kinetics
The activation energy (Ea) and the pre-exponential factor (A) were calculated from the TGA result,
modeled by the Arrhenius equation, Equation 1.1. Both first order and second order reactions were
16
considered and fitted, and both case obtained R
2
higher than 0.9, Table 1.3. The kinetics result is
consistent with the DSC and TGA result, that larger ester group requires higher energy to activate
the reaction, and cis-rich formulation requires higher energy to activate the reaction. N1, and N4
showed different activation energy with N7, because the curing kinetics is also influenced by the
condition of crystalline of the monomer mixture.
𝑑𝑎
𝑑𝑡
= 𝐴 𝑒 𝑥 𝑝 ( −
𝐸 𝑎 𝑅𝑇
) ( 1 − 𝑎 )
𝑛 (1.1)
Formulation
1
st
Order
Ea
kJ/mol
1
st
Order
A
k s
-1
1
st
Order
R
2
2
nd
Order
Ea
kJ/mol
2
nd
Order
A
k s
-1
2
nd
Order
R
2
T1C1 DEE 53.8 14.758 0.988 59.3 85.74 0.938
T1C3 DEE 76.3 7025.1 0.98 92.26 986723.1 0.988
T1C1 DIPE 63.13 85.8 0.958 75.09 3085.949 0.976
T1C1 DME 55.69 53298.3 0.94 63.66 645.384 0.965
N4 58.98 74.369 0.951 67.201 1088.16 0.961
N1 51.22 7.943 0.951 60.88 185.16 0.971
Table 1.3 Imidization Kinetics of TriA-X formulations
17
1.4 Conclusion
The TriA X polyimide examined in this study features a wide processing window that will
undoubtedly be an appealing aspect of the new resin for high-temperature composite applications.
The enlarged process window stems from the characteristic asymmetric repeating units and PEPA
end caps of the structure. These features afford more time (and/or higher temperatures) for
imidization, a necessary step before crosslinking. The expanded process window in turn can
facilitate the production of composite parts with low void contents. Furthermore, the ability to
imidize at higher temperatures can be used to reduce viscosity and open up the possibility for
applying resin transfer molding (RTM) or other out-of-autoclave processing methods for
composites manufacturing.
Imidization studies also show that the reaction between p-ODA and PEPA has a slower reaction
rate than the reaction rate of p-ODA and PMDA diester. The formulations with imide oligomers
of higher degree of polymerization show higher cure temperature due to the lower mobility of the
chain. Polyimides with larger degrees of polymerization show lower values of glass transition
temperature in general. Formulations with different degrees of imidization can be chosen to suit
different applications. Furthermore, the TriA X polyimide is also a low-toxicity resin without
MDA or other toxic volatiles.
Process window and viscosity was invested for Tri-A X with different diester group. The result
shows that the changing of diester group results in difference in imidization temperature and
imidization kinetics. The isomerism, in our study, didn’t show a big influence on the cure behavior.
Imidization occurred by a two-stage process in which an initial rapid phase was followed by a
slower one. This work provide a guideline for selecting proper alcohol for the esterification step
for PMR polyimide resin
18
CHAPTER 2: Thermal Oxidation of PEPA Terminated Polyimide
The thermal oxidative stability of a 4-phenylethynylphthalic anhydride (PEPA) end-capped
polyimide was investigated. A surface reaction layer formed due to oxidization during thermal
aging and grew in thickness with increasing aging time. Analysis of the surface layer revealed a
partial loss of aromatic ring, ether linkage, and imide linkage in the aged polymer. The partial loss
of the imide linkage and ether linkage in the surface layer was corroborated by the observed release
of CO and CO2 reaction products. The oxidized layer exhibited discoloration and an increase in
glass transition temperature Tg. The surface discoloration was attributed to formation of the
conjugated unsaturated or aromatic carbonyl groups and/or the charge transfer complex. Interior
regions of the oxidized samples were largely unaffected, except for a more compact molecular
configuration. Compared with aging at high temperatures (288 ˚C, 316 ˚C), aging at 204 ˚C
produced similar chemical changes in the cured polymer, albeit with a low degradation rate. The
oxidative stability of the imide polymer surpassed that of conventional nadic-end-capped
polyimides due to the greater intrinsic thermal stability of the PEPA end-cap. The unusual
oxidative stability of the polyimide, combined with superior mechanical properties, warrant
consideration as a composite matrix for future applications in high temperature service conditions.
2.1 Introduction
Oxidative reactions in polyimides can potentially cause problems that include cracking, brittleness,
chemical changes, all of which can compromise retention of mechanical properties. In this work,
we explore and report the thermal oxidative stability of a new polyimide and associated composites.
19
In particular, we identify chemical changes associated with anticipated high-temperature service
conditions in oxidizing ambient.
Chemical changes occurring during polyimide oxidation are not fully understood, due in part to
the complexity of the oxidation reaction. Incomplete understanding of oxidation mechanisms
limits efforts to redesign and optimize the molecular structure of PI. Thus, for the TriA X PI,
understanding the thermal oxidative stability is necessary to evaluate performance as a high-
temperature material. The objective of this study is to determine the mechanisms of thermal
oxidative degradation of the polyimide TriA X. We investigate the degradation effects of thermal
oxidation, including weight loss, formation of an oxidized layer, and microcrack formation. The
reaction products generated during thermal oxidation are also investigated to provide insights into
the physical changes.
Reports on the oxidation behavior of PI formulations have shown that an oxidized surface layer
typically forms. For example, one PI long used in the aerospace industry is PMR-15, (PMR stands
for polymerized monomeric reactant). During thermal oxidation of PMR-15
6
, a reaction occurs
that results in evolution of three distinct regions - an oxidized shell, an active reaction zone, and a
non-oxidized core
7
. Analysis of the oxidized surface layer of PMR-15 with norbornenyl-end-cap
showed that the nadic end-cap skeleton was virtually undetectable
8
. The methylene moiety
oxidized to ketone, causing shrinkage and stiffening of the chain. Other possible degradation
pathways have been proposed that involve degradation of the nadic end-cap moiety and the
methylene dianiline (MDA) moiety
9
. However, the mechanism of weight loss has not been
identified, and the possible thermal degradation of the non-oxidized core was not investigated. In
related work, the oxidation behavior of PI composites was investigated. For example, Xie et al
studied the oxidative stability of CF-PI composites (AFR 700B/T650-35) using thermogravimetric
20
analysis-mass spectroscopy (TGA-MS)
10
. By analyzing gaseous species generated during heating,
they proposed multiple possible chain breakdown sites, mostly related to the end norbornene group,
the end aniline group, the –CF3-C-CF3- and the imide ring. The use of TG-MS provided real-time
detection of release of degradation products. However, aging mechanisms in long-term high-
temperature applications may differ from the reaction mechanisms detected during the more
aggressive conditions of accelerated pyrolysis in TGA.
The present PI (TriA X) features a distinctive molecular structure that is likely to impart oxidation
mechanisms and kinetics unlike nadic end-cap PI’s. In particular, the PEPA end-cap of TriA X
does not contain carbon with sp
3
hybridization, unlike the norbornenyl-end-cap often employed in
polyimides. After crosslinking, the bond dissociation energy (BDE) of the alkenyl phenyl bond is
~ 452 kJ/mol
11
, rendering the end-cap resistant to pyrolysis and release of free radicals. For the
backbone, the (TriA X) molecule also exhibits greater heat resistance than PMR-15 by virtue of
the replacement of the methylene bridge in the diamine with the biphenyl ester croup. The
oxidation-susceptible bond in the TriA X backbone is most likely the imide bond, which possesses
a low BDE (~270 kJ/mol)
12
. Thus, due to the distinctive backbone structure and end-cap species,
TriA X PI is expected to exhibit oxidation behavior unlike other PI’s. These differences may lead
to distinct oxidation kinetics, a different oxidization growth mechanism, and different effects on
mechanical properties
13
.
In this work, we identify reaction mechanisms that occur during the thermal oxidation of the
PEPA-terminated PI (TriA X). We also show that the PEPA-terminated PI molecule exhibits
superior thermal stability compared with nadic end-capped PI (PMR-15), exhibiting less rapid
weight loss and an absence of aging-induced cracks. The thermal stability stems from the intrinsic
stability of the PEPA end-cap, which exceeds that of the PI backbone and the nadic end-cap. Unlike
21
PMR-15, a single distinct oxidized layer forms during thermal oxidation, and the reaction
responsible is attributed to breakage and rearrangement of the imide ring, ether group and the
PEPA end-cap. Aging at moderate temperature was also conducted to compare with aging at higher
temperature typically employed in accelerated aging studies, and the oxidation mechanism were
consistent in both regimes.
2.2 Experiments
2.2.1Material Preparation and Aging Condition
The chemicals PMDA (Tokyo Kasei Kogyo, Japan), PEPA (Manac, Japan) and 2-phenyl-4, 4’-
diaminodiphenyl ether (p-ODA) (Wakayama Seika Kogyo, Japan), were obtained from Kaneka
Corporation. The monomer mixture was prepared by mixing the chemicals in a hot ethanol solution,
followed by vacuum drying. After preliminary screening, the ratio of PMDA, p-ODA, and PEPA
was chosen to be 7:8:2 (degree of polymerization, n=7), and confirmed by NMR measurements.
The PI panel used for the aging study was produced using a hot press with a metal mold. The
monomer mixture was first converted to oligomer by heating to 280˚C for 60 min. The resulting
oligomer was ground to fine powder and placed in a steel mold 3 mm deep. The oligomer was
heated to 370 ˚C for 150 min to achieve crosslinking.
Prior to oxidation, all samples were preheated to 200 ˚C for 1 hour to remove moisture. Oxidative
aging treatments up to 2000 hours were performed in an air-circulated oven set to 204, 288 and
316 ˚C to accelerate aging. The latter aging temperatures were chosen for comparison with
reported data on PMR-15, while 204 ˚C was chosen to approximate anticipated service
temperatures. Aging treatments were interrupted periodically to remove samples for measurements.
2.2.2 Weight Change
22
Ten PI panels were cut to 100× 10× 3mm coupons for weight measurements. After preheating, the
dimensions of the coupons were measured and repeated 10 times to determine the total surface
area, and the initial weight prior to aging was measured using a balance with 0.01mg accuracy.
Coupons were placed in an air-circulated oven for thermal oxidative aging, and removed
periodically for weight measurements. Samples were placed in a desiccator to prevent moisture
absorption during cooling. The weight change per unit area was calculated for each of the 10
coupons. Thin polished sections (0.3-0.4 mm) were prepared for microscopy observation using
transmitted light (Keyence VHX-5000).
2.2.3 DMA
The glass transition temperature, Tg, was determined by conducting dynamic mechanical analysis
tests (DMA 2980, TA Instruments). A dynamic ramp of 5˚C/min was performed from 30˚C to
500˚C. Three groups of samples were tested: (1) unaged neat PI (TriA X) samples, (2) neat PI
samples (3× 10× 30mm) aged in an air-circulated oven for 20, 50, 100, 250, 500, 1000, 1500, 2000
hours, and (3) neat PI samples aged in an air-circulated oven for 20, 50, 100, 250, 500, 1000, 1500,
2000 hours with each exposed surface polished to at least 0.5mm. To measure the thickness of the
oxidized layer, polished sections of unaged and aged samples were prepared and inspected using
light microscopy.
2.2.4 Nanoindentation
Indentation tests on the aged specimens were carried out (Hysitron TI950 TriboIndenter,
Minneapolis, MN) equipped with a diamond Berkovich tip. Tests were performed at room
temperature using the continuous strain rate method, with dh/dt = (dh/dt)/h = 0.075 s
-1
where h is
the penetration depth of the indentor in the polymer substrate. The maximum penetration depth
23
was h=2 μm and the load was Pmax=5000 μN. Special care was taken to avoid viscous effects on
the load-displacement curve by performing a hold-step while maintaining constant load, and
measuring the required properties during a subsequent unload. For each material sample, at least
10 indentations were carried out in different locations of the specimen. The material apparent
hardness H and the elastic modulus E were determined from the load-displacement curve in the
unloading cycle after the hold period using geometrical area of the Berkovich indenter and the
elastic compliance. Thus, artifacts due to polymer pile-up surrounding the tip masking the true
contact area were avoided in the analysis.
2.2.5 FTIR
Attenuated total reflection (ATR)-FTIR (Nicolet 4700) was employed to obtain IR spectra from
aged samples. Sections were prepared from (a) the thermally oxidized surface, and (b) the interior
of the PI sample, and both were polished to films < 100 μm thick. Wavenumbers from 4000 cm-1
to 400 cm-1 were collected, and 64 scans were performed for each sample. Coupled TGA-FTIR
experiments were performed in N2 using a heating rate of 10 ˚C/min up to 850 ˚C to detect and
identify reaction products as they evolved. Scanning FTIR (Hyperion, Bruker) was conducted on
the cross-section of the aged TriA X with reflective light. IR spectra were acquired from the surface
to the interior every 30 μm. The typical light spot diameter was 20-50 μm.
2.2.6 XPS
X-ray photoelectron spectroscopy (XPS, Kratos Axis Ultra DLD) was used to determine the
composition of PI films before and after thermo-oxidative aging. Thin-film polished sections were
prepared from the exposed surface and the sample interior. The oxidized surface and the interior
were analyzed using an Al anode under 10
-9
- 10
-8
torr vacuum with sampling depth 2-10 nm. For
each sample, a survey spectrum was acquired to obtain overall composition, followed by three
24
high-resolution scans for carbon, oxygen and nitrogen. For O 1s, the range for binding energy was
chosen as 550 - 525eV, and for C 1s, 310 - 215eV. 15 sweeps were conducted for each range.
Scheme 2.1. Molecular structure of the TriA X oligomer and reaction scheme of crosslinking of
PEPA end cap
25
2.3 Results and Discussion
2.3.1 Weight loss
The weight loss of the PI with aging time at 316˚C and 288 ˚C, was shown in Figure 2.1. The
rate of weight loss during aging, Table 2.1 shows that, at 316˚C (0.739 mg*mm
2
*hr
-1
) was ~
4.2× greater than the weight loss rate for aging at 288˚C (0.176 mg*mm
2
*hr
-1
). Both rates were
much less than the rate of weight loss values reported previously for PMR-15 (~0.35
mg*mm
2
*hr
-1
for 288˚C
14
, ~8.5 mg*mm
2
*hr
-1
for 316˚C
15
), indicating greater thermal stability.
Figure 2.1. Weight loss per unit area of TriA X during isothermal oxidative degradation at
288 ° C and 316 ° C
26
Weight Loss Rate
Aging Temperature
PMR-15 TriA-X
mg*mm
2
*hr
-1
mg*mm
2
*hr
-1
288˚C ~0.35
14
0.176
316˚C ~8.5
15
0.739
Table 2.1 Weight loss rate comparison for PMR-15 and TriA-X
The surprisingly low rate of weight loss observed can be attributed primarily to the absence of
surface and bulk cracks in the PI here (TriA X). Cracking and crazing are commonly reported
during thermal oxidation and thermal cycling of PI’s and PI composites
14-17
. In addition, voids
reportedly appear in oxidized layers of PMR-15 aged at 288 and 316˚C
18
. Microcracks and voids
exacerbate surface reactions in PI’s, providing pathways for accelerated ingress of oxygen,
increasing the rate of weight loss. Thus, during thermal oxidation of conventional PI (e.g., PMR-
15, 316˚C in air), the weight loss shows a power-law dependence on aging time (exponent = 1.63),
which is attributed to an increase in exposed surface area caused by crazing, cracking, and pitting
15
. However, during oxidation at 288 and 316˚C, the PI here (TriA X) shows a quasi-linear weight
loss relationship with time, and no cracking was observed at the surface or in the bulk. Note that,
although weight loss for 316˚C data appear to be linear, close inspection shows an apparent
acceleration of mass loss rate near 400h, which would indicate a breakdown of a passive protection.
The end-cap of the PI molecule affects the rate of weight loss during oxidation. Polyimides with
norbornenyl end-caps (such as PMR-15) undergo more rapid weight loss compared with PI’s
sharing the same backbone
9
. Meador et al provided evidence that the most rapid large initial loss
is due to degradation of the nadic-end-cap
9
. In contrast, for nadic-end-capped polyimides, the
initial rate of weight loss decreases as the relative amount of nadic decreases
19
. Note that PMR-
15, where n=2, is expected to have greater cross-linking density than TriA-X (n=7). For TriA-X,
27
n=7 was studied due to its balance of mechanical and thermal performance. Many factors could
account for the lower weight loss rate. Such as absence of microcrack and lower cross-linking
density. It also can be attributed to the greater thermal stability of the PEPA end-cap than the nadic
end cap. Degradation of the TriA X backbone and PEPA end-cap is discussed in subsequent
sections.
Cross-sections of the unaged and thermally oxidized TriA X PI samples after aging at different
conditions are shown in Fig. 2. Aged at 288˚C and 316˚C, the surface of the aged polymer shows
discoloration when illuminated with transmitted light, and the discoloration is attributed to the
generation and accumulation of highly conjugated groups or other chromophores
20
. Aging at
204˚C produced no distinct discolored layer was observed even after 2000 hr. The oxidized layer
thickness in the sample aged at 316˚C is less than the layer thickness produced by aging at 288˚C,
which is a typical diffusion/reaction result
21
. The oxidized layer that grows rapidly during high
aging temperature retards further diffusion of O2. Aging at 204˚C did not produce a detectable
oxidized layer, possibly because the critical oxidized layer thickness (Lc)
21
at this temperature
exceeded half the sample thickness (0.5*L), and/or because the process at this temperature is
limited more by the reaction rate than by diffusion.
28
Figure 2.2. Cross-section of TriA neat polymer a) Unaged; b-1)204˚C, 700 hr; b-2) 204˚C, 2000
hr; c-1) 288˚C, 700 hr; c-2) 288˚C, 1500 hr; d-1) 316˚C, 700 hr; d-2) 316 ˚C, 1500 hr
The absence of cracks in the aged polymer is largely responsible for the lower rate of weight loss
observed. In the following sections, the discolored region will be referred to as the oxidized layer,
while the translucent interior region will be referred to simply as the interior. The entire sample
(interior plus oxidized layer) will be referred as the bulk sample.
2.3.2 Thermal Analysis
DMA experiments were performed to measure the Tg (via tan(δ)) for unaged PI and PI thermally
oxidized at 288 ° C for different times, and results are shown in Fig. 2.3. In general, Tg increased
with aging time (Fig. 2.3 a), and similar increases in Tg have been reported after thermal oxidation
of PMR-15
16
. In Fig.2.3 a, at 500 h, the tan(δ) peak became broader and less symmetric with
aging time and temperature, while the peak centroid shifted to higher temperature. After further
aging for 1000 and 2000 h, the tan(δ) peak exhibited a shoulder at the lower-temperature flank,
29
and longer aging times caused apparent peak separation, with a second peak developing near
325 ° C.
30
Figure 2.3. a) tan(δ) of TriA X during aging at 288 ° C, b) tan(δ) of bulk and interior TriA X after
aging
Two possible explanations were considered for the peak separation observed in Fig. 3a – (1) β
relaxation of the oxidized PI, and (2) formation of a second phase during aging. To distinguish
between these two possibilities, DMA samples were prepared from the aged bulk polymer and
from the aged interior (with oxidized layer removed), and the results are shown in Fig. 3b.
As shown in Fig. 3b, removing the oxidized surface layer also removed the broad high-temperature
tan(δ) peak (>350 ° C). Thus, the low-temperature peak (at ~330° C) must originate from the
interior of the sample, while the broad peak at ~350 ° C corresponds to the discolored oxidized
31
layer. Furthermore, the interior and the oxidized layer exhibit different thermal mechanical
behavior. No obvious change of the value of Tg for the interior was observed after aging, indicating
that the chemical structure here was largely unaffected by aging (see Fig. 3b). The author also
observed an increase in G’, which indicates the interior undergoes physical aging such as reduction
in free-volume and decrease of mobility. In contrast, a notable change is apparent the oxidized
layer. In Fig. 3b, the broadening of the α relaxation peak at high temperature indicates the growth
of a thicker oxidized layer with different chemical structure, wider molecular weight distribution,
or more compact packing. In section 3.4, we attribute the stronger intra/inter molecular interactions
between the polymer chains to the loss of pendant phenyl groups of p-ODA
22
, and/or the
generation of oxidative products.
2.3.3 Nanoindentation Measurements
The apparent hardness H (indentation load divided by imprint projected area) and elastic modulus
E of the unaged polymer, aged polymer on the surface exposed to aging, and aged polymer interior
are shown in Fig. 4. The larger standard deviation obtained for the aged surface tests is attributed
to the sample preparation procedure, which left initial roughness on the surface. Nevertheless, both
the elastic modulus and the hardness increased in tests carried out on the specimen surface, while
the specimen interior showed almost no change with aging time, as shown in Figures 2.4(a) and
(b). The trends for surface and interior hardness converge to a single value as aging time decreases,
corresponding to the unaged polymer. Note that the elastic modulus of the unaged TriA X (≈4.70
GPa) as measured by indentation is remarkably similar to the nanoindentation elastic modulus of
PMR-15 (≈4.72 GPa)
23
.
32
Figure 2.4. a) Nanoindentation hardness of the unaged and aged resin surface and interior, b)
Nanoindentation modulus of the unaged and aged resin surface and interior
After aging, both hardness and modulus on the surface increased. This increase indicates the
generation of a hard oxidized surface layer on the polymer. The surface layer features strong
33
intermolecular interactions that increase with aging time. In contrast, both hardness and modulus
remain nearly constant for interior regions of the sample, with only a slight decrease. With thermal
oxidative aging of PMR-15 and other thermosets, an increase in modulus of the oxidized layer is
commonly observed
18, 23
. In PMR-15, the average elastic modulus in the oxidized region is
insensitive to aging time
23
. However, for TriA X, a continuous increase in hardness with aging
time was observed, (Fig. 4), indicating a continuous change in the chemical composition of the
surface. To clarify the reaction that occurs during thermal oxidation, FTIR and XPS were used to
analyze the chemical composition.
2.3.4 FTIR Analysis
Chemical changes in PI’s during thermal oxidation can be identified using FTIR spectroscopy of
aged samples. In previous research, chemical changes during oxidative degradation and thermal
cycling of polyimides with norbornenyl end-caps were monitored by tracking changes in FTIR
spectra. Based on the results, investigators proposed conversion of methylene from the
methylenedianiline moiety to a carbonyl group
9, 24
. Degradation of the O=C-N moiety in the imide
ring, and degradation of C=C in the unreacted nadic end-cap, were also proposed as possible
degradation reactions
9, 24
. However, the present PI (TriA X) lacks a methylene group in the
monomer, and the end-cap has no C-C and C=C bonds, and thus must follow a different pathway
to oxidative degradation. FTIR spectra acquired from samples of unaged and aged PI, both surface
and interior, showed the progressive change in chemistry during thermal oxidation (Figs. 5 and 6).
Spectra from the polymer surface (Fig. 5 a) show a decrease in a set of peaks (arrows) relative to
the spectra of the aged polymer interior, which are nearly unchanged. Fig. 5 b shows that the
spectrum after aging at 288˚C shows less change compared with aging at 316˚C for the same time,
34
while the spectrum after aging at 204˚C shows only a slight decrease at 1218 and 695 cm-
1
.
(Compare Figs. 5a and 5b)
35
Figure 2.5. a) FTIR spectra of fresh TriA X vs. 316 ° C 500 hr aged resin surface and interior. b)
FTIR spectra of TriA X aged at different temperatures and times.
FTIR spectra from the unaged surface show bands generated from the imide ring, appearing at
1776, 1720, 1360, and 721 cm
-1
(Fig. 6). Bands at 1776 cm
-1
and 1720 cm
-1
correspond to the in-
phase and out-of-phase stretch of the imide carbonyl. After oxidation, FTIR spectra show that the
peak at 1720 cm
-1
overlapped the carbonyl peak of the oxidation product. Hay et al and Meador et
al
9, 25
assigned the band at 721 cm
-1
(719 cm
-1
), (which decreased with aging), to the out-of-plane
bending of the C=C on the unreacted nadic end-cap. However, we assigned this peak to deviational
vibration of the imide ring because of its absence in the spectrum from monomer powder, and the
proportional relationship with peaks at 1776 and 1360 cm
-1
. The observed slight decrease in
36
amplitude of the imide ring-related peaks reflects the instability of the imide group during aging
at 288 and 316˚C.
Figure 2.6. FTIR spectra of fresh TriA X vs. resin surface after aging at 316 ° C for different
times.
Figure 2.6 also shows the effects of aging time on FTIR spectra acquired from the oxidized surface.
With increasing aging time, there is a slight but steady decrease in band amplitude relative to the
aromatic ring (Ar). Skeletal vibration of the benzene ring resulted in strong bands at 1500 and
1478 cm
-1
(labeled with red *) and these bands decreased with increasing oxidation time. C-H
bending peaks of the mono-substituted aromatic ring were observed at 767 and 695 cm
-1
(circle
mark). The mono-Ar peak comes from both the PEPA end-cap and p-ODA. The band at 817 cm
-1
(arrow) was assigned to the C-H bending of the para-disubstituted aromatic ring. Some of the weak
37
overtone bands of the aromatic ring skeletal vibration were obscured by the strong C=O bands in
the region of 1600-1800 cm
-1
. The peak at 1218 cm
-1
(asterisk) was assigned to the phase stretch
of Ar-O-Ar
26, 27
. Compared with the imide band at 721 cm
-1
, bands at 1218, 817, 767 and 695 cm
-
1
show clear decreases during oxidation, indicating lower thermal oxidative stability of the Ar-Ar,
Ar-O-Ar and Ar-N bonds compared to the imide ring at 316˚C. After protracted aging at lower
temperature (204˚C for 3500hr), slight of decrease in relative intensity of Ar-O-Ar bond and mono-
Ar peak were observed, which may come from the loss of PEPA end-cap or p-ODA. Thus, at
higher aging temperatures, (288˚C and 316˚C), the p-ODA, and PEPA end-cap showed lower
stability compared with imide related groups.
Loss of the mono-substituted aromatic ring on the p-ODA moiety can account for the observed
discoloration of the surface layer. Before aging, this benzene ring caused steric hindrance to reduce
formation of inter/intra chain charge transfer complex (CTC) between nitrogen (electron donor)
and the carbonyl group (electron acceptor)
28
. By removing this benzene pendant from the
backbone, the inter/intra chain packing becomes stronger by formation of a more rigid CTC,
accounting for the darker color of the aged surface layer. The more rigid packing also accounts for
the increase in hardness and modulus of the aged surface layer.
Oxidation products were identified from specific peaks which developed after aging at 316˚C (Fig.
6). For example, the peak at 1607 cm
-1
was assigned to the stretching of the conjugated C=O group
generated by oxidation
9, 24
. Bands at 1218 and 1089 cm
-1
were attributed to stretching of the C-O
in the Ar-O-Ar moiety. However, instead of decreasing proportionally with the band at 1218 cm
-
1
, the band at 1089 cm
-1
increased and broadened slightly, then finally became greater in amplitude
than the band at 1218 cm
-1
. The increase and broadening of the strong peak at 1089 cm
-1
was
attributed to generation of alkoxy groups from the ester or acid of oxidation products, which
38
overlapped with the ether C-O peaks in the spectrum. The bands at 1196 and 1607 cm
-1
were
assigned to acyl stretching and C=O bending, leading to the conclusion that the oxidation products
were primarily unsaturated esters. Ketone is proposed as another possible oxidation product, due
to the appearance of the strong band near the 1710 cm
-1
region.
FTIR spectra generally do not provide clear indications of post-curing during thermal oxidation.
In the TriA X PI, there was no sp
1
C-H in the alkynes group of the end-cap before aging, nor was
there a sp
2
C-H in the cross-linked end-cap. The bands of C≡C and C=C overlapped with CO2
peaks and the existing strong peaks of benzene. Consequently, changes in chemical composition
of the end-cap, such as post-crosslinking and degradation, were not apparent. The interior of the
aged samples yielded spectra much like those of PMR-15
9
, with no sign of crosslinking (because
of the absence of C-H bands). Bands for nitriles and alkynes, which were detected as intermediates
during the UV laser-induced decomposition and pyrolysis of Kapton
29
, were not observed after
thermal oxidation of TriA X.
39
Figure 2.7. Content of chemical groups of the 288 ° C 700 hr aged resin oxidized layer
Scanning FTIR revealed progressive degradation from the polymer surface to the interior, as
shown in Fig. 2.7, which plots the concentration of chemical groups as a function of sub-surface
distance. To compare degradation of different chemical groups, the intensities of 5 peaks - 1720
cm
-1
(-C=O), 1360 (C-N of the imide ring), 1218 cm
-1
(ether linkage), 695 cm
-1
(mono-Ar) 817
cm
-1
(para-Ar) were measured and compared with the intensities of the peaks of the interior. Fig.
7 shows that the mono-substituted aromatic ring which comes from the PEPA end-cap and ether
linkage has the lowest stability, followed by the ether linkage with the next lowest stability. The
intensity for the C=O group decreased less than the C-N group because of the generation of
oxidative product, which also contained C=O groups. The similarity of the curves for the para-
aromatic ring and the C-N peak indicates that these two groups degraded almost simultaneously.
The oxidized layer for this sample (aged at 288°C for 700 hr) was ~ 800 μm thick, and this value
is consistent with that of the discolored region. The correlation indicates that the discoloration can
be attributed to progressive chemical degradation (which is accompanied by increases in hardness
and Tg).
40
Figure 2.8. Instable groups of TriA X in thermal oxidative environment
FTIR analysis shows the instable chemical groups in the polymer chain and end-cap, shown in
Figure 2.8. Fang et al
30
proposed multiple phenylethynyl reactions and further reactions for the
curing of polyene and aromatic ring structures for the PEPA end-cap To further assess the thermal
stability of the chemical bonds in TriA X, dynamic TGA was performed, and the volatiles were
analyzed by FTIR (Fig. 9a). The intensity curve shows that the PI started to decompose in N 2 at
~500 ° C, and reached the highest decomposition rate at ~576 ° C. The FTIR spectra from volatiles
released at 545 ° C and 625 ° C during dynamic ramp represent the early and later stages of
decomposition, discussed next.
41
Figure 2.9. a) FTIR intensity of the released volatile and weight loss percentage during the TGA
dynamic ramp. b) FTIR spectra of the volatile released at 545 ° C and 625 ° C.
The primary volatiles generated at 545 ° C were identified as CO (2182 cm
-1
, 2111 cm
-1
) and CO2
(2380-2280 cm
-1
, 669 cm
-1
, red curve in Fig. 2.9 b). During pyrolysis in the absence of O2, CO and
CO2 are generated from the decomposition and rearrangement of the imide ring and ether linkage
of the p-ODA. In contrast, the spectrum acquired at 625 ° C exhibited bands associated with
unsaturated alkyl groups, including bands at 3016 cm
-1
from sp
2
C-H stretching, and at 965 cm
-1
930 cm
-1
and 714 cm
-1
from sp
2
C-H bending (black curve in Fig. 2.9 b). Alkene is the primary
released species, although the broad band at 3110 - 2990 cm
-1
indicates the possible existence of
an aromatic ring. The fluctuating peaks in the high wavenumber region (>3500 cm
-1
) reflect
42
constant generation of H2O molecules from the beginning of the ramp, initially as absorbed
moisture, then later as a decomposition product in the high-temperature region. The small peak at
3333 cm
-1
, can be assigned to small amounts of amines, particularly C6H5-NH2 and CH2=CH-NH2.
Analysis of the TGA-FTIR data reveals that in N2, the ether linkage and imide ring on the polymer
backbone have lower thermal stability, which is consistent with the scanning FTIR spectra and
with previous reports on pyrolysis and carbonization of Kapton
29, 31
and AFR 700B
10
. The
generation of C6H5-NH2 and CH2=CH-NH2 stems from the decomposition of -Ar-CO-N(CO)-Ar-
and -Ar-O-Ar- bonds, which have relatively low bond dislocation energies (BDE’s) and thus lower
thermal stability. The BDE for Ar-O-Ar is 329.7 kJ/mol
4
, while the BDE’s for Ar-Ar and Ar-vinyl
are ~ 450 - 460 kJ/mol
11
. Thus, in the absence of O2, bonds with lower BDE have higher tendency
to break.
The small amounts of alkene and aromatic groups detected by IR are possible degradation products
of the end-cap, but may also arise from decomposition of benzene rings within p-ODA of the
backbone. Even if the alkene and aromatic groups arise from decomposition of the PEPA end-cap,
the IR intensity of these species is much weaker than the IR intensity of CO 2 and CO, especially
in the early stage of weight loss. These observations indicate that in the absence of O2, the thermal
stability of the PEPA end-cap is greater than that of the backbone. However, for nadic-end-capped
polyimides, the end-cap-related species (cyclopentadiene and CH3OH) exhibit greater relative
intensity compared to backbone-generated species
10
. Thus, the TGA-FTIR results indicate that in
O2-free environments, PEPA end-cap has greater thermal oxidative stability than both the
polyimide backbone and the conventional norbornenyl end-cap.
2.3.5 Surface XPS analysis
43
To clarify the change in chemical composition during oxidation and to identify possible
nonvolatile oxidation products, XPS spectra of C 1s and O 1s were acquired from the unaged
polymer, the aged surface, and the aged interior of the TriA X polymer. In evaluating the XPS
spectra, it is useful to consider the ratio of the total area of C 1s over the total area of O 1s (C/O
ratio). Using this metric, and considering the effects of oxidation at 316° C for 1000 hours, the C/O
ratio decreased from 1.43 to 1.10 in spectra from the sample interior. In contrast, in spectra from
the oxidized layer, the C/O ratio was 0.99 after oxidation. The decrease in C/O ratio reflects the
relative increase in the amount of O during oxidation, as well as the loss of C-related groups.
To understand the changes in binding energy of C 1s, five constituent peaks were separated and
plotted, as shown in Fig. 2. 10 a. The first of these peaks (peak 1), the broad peak at 284.6 eV,
includes (a) the sp
2
carbon from the alkene groups of the crosslinked end-cap, (b) aromatic rings
that are not attached to an imide ring, and (c) aromatic carbon of the p-ODA moiety, not directly
connected with N or O.
31-34
. The second peak (peak 2), at 285.2 eV, arises from (a) carbon atoms
shared by the imide ring and the aromatic ring on the end-cap, (b) from the C-N bond, and (c) the
unsubstituted C in the PMDA aromatic ring
35
. The third peak (285.7 eV, peak 3) corresponds to
the aromatic carbon connected with the imide ring in the backbone. This carbon has a ~0.5 eV
shift compared to the aromatic carbon connecting with the imide ring in the end-cap, because of
the stronger conjugation effect of the additional imide ring
35
. This peak also corresponds to the
secondary carbonyl carbon from the oxidation product. The fourth peak (286.4 eV, peak 4)
corresponds to the C-O carbon of the ether linkage, or to alkoxy carbon from the oxidation product.
Finally, the peak at 288.6 eV (peak 5) corresponds to the carbonyl carbon from the imide ring
33,
35
.
44
45
Figure 2.10. a) Peak separation for C 1s of the aged interior, b) C 1s of fresh TriA X and 316 ° C
1000 hr aged resin surface and interior
As shown in Fig. 2.10 b, XPS spectra from the oxidized surface showed distinct differences from
the interior. In particular, the shift of the carbonyl peak from 288.6 eV to 288.2 eV reflects
consumption of imide rings and generation of oxidation products on the polymer surface. The
binding energy at 288.2 eV points to ketone, aldehyde, and amine as possible oxidation products,
and reduces the possibility of acid (usually this has a binding energy ~289.3 eV). The decrease in
amplitude of the peak at 285.2 eV also indicates degradation of the imide ring. The increase in the
amplitude of the 285.7 eV peak is attributed to overlap of the secondary C1s for ketone or aldehyde
(~285.4 eV), both possible oxidation products
35
. Combined with the IR analysis of nonvolatile
oxidation products, ketone emerges as the most likely oxidation product. In particular, the
carbonyl group conjugated with an aromatic ring, alkene group, and other unsaturated groups on
the polymer chain, constitute the chromophore for discoloration of the oxidized surface layer
36
.
The C 1s spectrum from the aged sample interior shows much less differences compared with
spectra from the aged surface, as indicated in Table .2. In the sample interior, the peak from the
carbonyl carbon of the imide ring (at 288.6 eV) remains constant in integrated intensity, an
indication of the stability of the imide group (Fig. 10 b). The relative decrease in amplitude of
285.2 eV and 285.7 eV compared with the sp
2
carbon (284.6 eV) may reflect a shift of binding
energy due to formation of a CTC.
46
Table 2.2. Area percentage of peaks for C 1s and O 1s of unaged and 316 ° C 1000 hr aged TriA
X
C 1s O 1s
BE (eV) 284.6 285.2 285.7 286.4 288.6 531.5 532.9
Area
percentage
(%)
Theoreti
cal
52.6 14.7 12.0 6.9 13.8 66.7 33.3
Unaged 55.9 20.5 7.9 9.2 6.5 72.4 27.6
Aged
interior
65.86 12.9 5.3 9.9 6.1 69.1 30.9
Aged
surface
65.4 7.9 9.0 11.9
5.8
(288.2e
V)
94.1 6.0
47
Figure 2.11. a) Peak separation for O 1s of the aged interior b) O 1s of fresh TriA X and 316 ° C
1000 hr aged resin surface and interior.
Spectra from the aged surface show that the O 1s peak separated into two peaks (Fig. 2.11 a). The
first peak (531.5 eV) is carbonyl, while the second (532.9 eV) is O in the ether group. In Fig. 2.11
b, the total O content increased after aging, and thus both peaks increased in intensity. The relative
amplitude of the 532.9 eV peak for the aged interior increased slightly. However, due to the large
peak width of O 1s, the binding energy of C-O with a large group of functional groups overlaps in
this range, preventing identification of the specific functionality
35
. On the aged surface, aromatic
ether bonds were broken, and C=O became the major species (~94%) resulting from generation of
aromatic carbonyl or alkene carbonyl. These species may have formed the conjugated
48
chromophore in the oxidized layer according to the second discoloration mechanism described
earlier.
2.4 Conclusions
The thermal oxidation study of TriA X polyimide provide understanding of molecular-level
reaction mechanisms, such as the degradation of ether moiety and the end-cap. During thermal
oxidative aging of TriA X at high temperatures, the PEPA end-cap and p-ODA moiety is less
stable than the imide ring. PEPA end-cap, which is relatively stable in N2, were largely affected
by existence of O2. This provides useful insight to guide further polyimide development and
modification of polyimides. Discoloration of the oxidized layer resulted from chemical
degradation was observed and investigated. The rigid oxidized surface layer, comprised of
unsaturated conjugated ketone chromophore, showed higher hardness and glass transition
temperature.
The reactions and kinetics at service temperatures may very well differ from those observed at the
higher temperatures typically used for accelerated aging studies. Thus, there is an inherent risk in
drawing conclusions from accelerated aging experiments. From the present investigation, aging
kinetics accelerated with aging temperature, although the chemical changes that occurred during
aging at 288 ° C and 316 ° C aged were unchanged. Aging at 204˚C, closer to anticipated service
temperatures, shows a much slower but similar oxidation reaction and mechanism.
49
TriA X exhibits greater resistance to thermal oxidation than the nadic-end-capped PI, PMR-15,
with less weight loss and no cracking. The thermal stability of TriA X derives from the intrinsic
stability of the PEPA end-cap relative to the nadic end-cap, both in inert and oxidative conditions.
These features, coupled with the unusually high ductility of TriA X, indicate potential to expand
the property space for high-temperature polymers, affording greater flexibility to designers of
composite parts for high-temperature service. The non-crystalline molecular structure
22
and PEPA
end-cap are likely to be preferred in future polyimide polymer development. Furthermore, the
principles used to design TriA X formulations also can be applied to guide development of other
high-performance polymers, such as bismaleimides and cyanate ester resin, to increase thermal
stability and ductility.
CHAPTER 3: Thermal Oxidative Aging and Thermal Cycling of PETI-340M
Composites
Polyimide composites (PETI-340M) were fabricated and subjected to high-temperature aging and thermal
cycling to evaluate resistance to degradation. Mechanical degradation mechanisms and kinetics depended
on aging temperature. Aging at 232 °C resulted in strength loss due to polymer degradation, while intra-
tow cracking was the dominant mechanism during aging at 288 °C. Composite panels subjected to thermal
cycling fatigue (-54 ° C to 232 ° C) retained mechanical properties without microcracking. However, in
regions containing pre-existing fabrication-induced defects (primarily voids), intra-tow micro-cracks were
observed after thermal cycling. Unlike some polyimide composites (PMR-15), oxidative aging effects
during thermal cycling were negligible. The thermal oxidative stability and the retention of mechanical
performance after thermal cycling indicates potential for long-term high-temperature structural applications.
50
3.1 Introduction
Oxidative reactions and thermal cycling fatigue in polyimide composites can cause cracking, embrittlement,
and chemical changes, all of which compromise retention of mechanical properties.
7, 8, 15, 17
In this work,
we explore and report the effects of thermal oxidative aging and thermal cycling on the mechanical
properties and microstructure of PETI-340M composites.
Interest in composite materials for aerospace applications involving harsh environments has driven the
development of polyimide composites. Polyimides (PI’s) are widely used in high-temperature composite
applications because of the high glass transition temperature (T g) and thermal stability. However, PMR-15,
the long-standing nadic-end-capped polyimide, generally suffers from poor processability, cracking
induced by thermal shock, and thermal oxidation. Long exposure to high temperatures can cause damage
and loss of strength in such composites. For example, thermal cycling during orbit can cause low amplitude
thermal fatigue and induce damage, including microcracks and delamination.
16, 37
PETI-340M, a phenylethynyl terminated polyimide (PETI) resin, features relatively low oligomer
molecular weight and compatibility with RTM processing (resin transform molding). However, the thermal
oxidative stability and thermal fatigue resistance have not been investigated or reported. The objective of
the present study was to determine the microstructural mechanisms involved in thermal oxidative
degradation and thermal cycling of PETI-340M composites. We investigated the degradation effects on
matrix-dominated mechanical properties, including compressive and interlaminar shear strength. Crack
formation during thermal oxidation and thermal cycling were also investigated using light microscopy and
micro-CT. To simulate anticipated service environments, aging was performed at 232 °C and 288 °C.
Much effort has been devoted to understanding the stability of polyimide composites subjected to thermal
cycling fatigue.
38-43
In conventional polyimide composites, such as PMR-15, as thermal fatigue progresses,
the crack density (crack/length of the specimen) increases asymptotically and approaches an equilibrium
value, while the coefficient of thermal expansion (CTE) decreases. Microcracks begin as intra-ply cracks
51
and grow to inter-ply cracks as the number of cycles increased, leading ultimately to delamination. Matrix-
dominated mechanical properties, particularly compressive and interlaminar shear strength, are adversely
affected by thermal fatigue and thermal aging. Extensional and flexural stiffness are fiber-dominated
properties, and typically show negligible effects of thermal cycling.
Much of the literature on high-temperature polyimide composites focuses on one formulation - PMR-15.
Owens et al. studied the thermal cycling of PMR-15 composites and showed that microcracks initially
occurred in the outer plies of laminates until a stress-relieved plateau was reached.
43
Continued cycling
resulted in cracking of the inner plies, and isothermal aging after thermal cycling produced additional
microcracking of inner plies. Tompkins et al. showed that in PMR-15 composites, thermal cycling induced
transverse microcracks and delamination.
40
The compression and interlaminar shear strengths of
unidirectional laminates decreased after cycling only when tests were performed at 316 ° C. However, the
interlaminar shear strength (ISS) of the quasi-isotropic laminate was affected significantly when tested at
both room temperature and 316 ° C. The results indicated the ISS of the quasi-isotropic laminate depended
on the density and extent of microstructural damage, while the ISS of the unidirectional laminate depended
on both the testing temperature and the damage induced by thermal cycling.
40
Zrida et al. compared two
thermal cycles with a common minimum temperature and reported that the highest temperature in the cycle
strongly influenced damage development.
38
They also showed that composites were only marginally
affected in the high-temperature test, except for panel edges and surface plies that were exposed to air. They
also showed that voids in the laminates acted as stress concentrators but also led to larger stress relaxation
after crack initiation, which delayed the appearance of new cracks and arrested the growth of existing cracks.
In this study, we investigated the effects of thermal oxidative aging and thermal cycling fatigue on matrix-
dominated mechanical properties of PETI-340M composites. To better understand the dominant factors
affecting mechanical performance, we conducted weight loss measurements, dynamic mechanical analysis,
light microscopy, and micro-CT measurements. The experiments provided a clear understanding of crack
initiation, propagation, and the effects on mechanical properties.
52
This research revealed different aging mechanisms for the two different thermal oxidative temperatures.
Even with a loss in strength, aging at 232 °C did not result in aging-induced micro-cracks, while aging at
288 °C produced micro-cracking. Micro-cracking induced by thermal cycling was observed only on
specimens that possessed fabrication-induced voids or defects. In general, the PETI polyimide showed high
thermal stability with retention of mechanical performance.
3.2 Experiments
3.2.1 Material Preparation and Aging Condition
A CF-polyimide prepreg was selected for this study (PETI-340M/HTS40 3K 8HS). The thermoset
polyimide matrix system, originally developed for RTM (resin transform molding), featured a low
processing viscosity. The laminate configuration was quasi-isotropic {0/+45/-45/90} 4s, and panels were
fabricated 2.6 mm thick by 400 mm in length and width. Panels were placed in a vacuum chamber for over
7 days to eliminate moisture before thermal cycling and thermal aging.
Panels prepared for thermal cycling were placed on a mechanically driven sliding tray, and thermal cycling
was conducted at -54 ° C to 232 ° C, The holding time at each temperature was 15 min. Specimens were
removed every 400 cycles to inspect for cracks and for mechanical testing. Thermal aging specimens were
heated in an air-circulated oven at 232 ° C and 288 ° C.
All test specimens were water-jet cut from the thermal cycled or thermal oxidative aged panels at room
temperature and stored in a desiccator at 20°C for over 7 days to remove moisture before mechanical tests
were performed.
3.2.2 DMA
The glass transition temperature, T g. was measured by dynamic mechanical analysis (Q800, TA
Instruments). A dynamic ramp of 5 ° C/min was used to heat from 30° C to 400°C, and T g was determined
53
by both the tan(δ) peak and the 𝐺 ′
transition. Two groups of samples were tested: (1) unaged composites,
and (2) composites aged at 232 ° C and 288 ° C.
3.2.3 Mechanical Tests
Compression tests were conducted to determine combined loading compression (CLC) and open-hole
compression (OHC) strength of the unaged composites, aged composites, and thermal-cycled composites.
CLC tests were performed according to ASTM D6641, with a loading rate of 1.3 mm/min, and digital image
correlation was used to map displacements (Aramis, Trilion). OHC tests were performed using a Boeing-
modified testing fixture, a loading rate of 1.3 mm/min, and a test procedure derived from ASTM D6484.
Short beam shear (SBS) tests were conducted on unaged, thermal cycled, and thermal oxidized composites
according to ASTM D2344. Six samples were tested per group using a loading rate of 1 mm/min. Thermal
oxidative aged composites specimens were tested with a load frame (5567, INSTRON 5567, Norwood, MA,
USA) equipped with an environmental chamber.
3.2.4 Microstructural Analysis
Polished sections were prepared and analyzed by light microscopy and micro-computed tomography
(micro-CT). To minimize the possibility of additional cracks from cutting, more than 1 mm was removed
by grinding from each surface. Specimens for light microscopy were polished using sandpaper (P240, P400,
P600, P1200) followed by alumina suspension (5um and 0.3um). Micro-CT scans were performed (XT H
225ST, Nikon, Japan) using Mo-K α incident radiation with λ= 0.71 A
ͦ . Voltage/intensity was set at 60
kV/220 mA to achieve resolution below 4 μm/pixel.
3.3 Results and Discussion
3.3.1 Weight Loss
54
Figure 3.1 plots weight loss per unit area for composites as a function of aging time at 232 and 288 ° C. The
curves were produced by nonlinear regression of the experimental data to an allometric power curve,
Equation 3.1. The fitted coefficient and the comparable coefficient for PMR-15 are shown in Table 3.1.
The exponent, B, increases with aging temperature and is influenced significantly by crazing and cracking
of the specimen surfaces (B=1 corresponds to uncracked surfaces).
15
Figure 3.1: Weight change per unit area for thermal oxidative aging of PETI-340M composite
with nonlinear curve fitting.
Comparison of the coefficients of PETI-340M composites and PMR-15 composites shows that
PETI-340M composites have much higher resistance in the thermal oxidative aging environment.
For example, after aged at 288 ° C for 4000 hours, the weight loss per unit area for PMR-15 is
0.644 mg/cm
2
,
15
and the weight loss per unit area for PETI-340M composites is 0.0228 mg/cm
2
.
The results in Table 3.1 shows that, at 232 ° C, the surface of the material is craze-free because the
coefficient B is 1.035. However, for aging at 288 ° C, B was increased to 2.068, indicating the
generating of surface cracks. For 288 ° C aging, PETI-340M and PMR-15 composites have similar
B (2.068 vs. 1.93), however, for PETI-340M composites, the other two coefficients are much lower
55
than PMR-15 (A: 8.072*10
-10
vs. 7.166*10
-8,
C: 1.177*10
-4
vs. 0.00205), which causes the low
weight loss rate for PETI-340M composites.
Weight Loss per Unit Area=At
B
+C (3.1)
A=S1k1+S2k2 (3.2)
Where
S1=Area of composites surface
S2=Area of cutting with fibers exposed on the surface
Material Aging
Temperature
A B C
PETI-340M
Composites
232 ° C 4.427*10
-8
1.035 -4,257*10
-6
288 ° C 8.072*10
-10
2.068 1.177*10
-4
PMR-15
Neat Resin
15
288 ° C 4.914*10
-1
0.98 0.25
288 ° C 7.166*10
-8
1.93 0.00205
Table 3.1. Coefficients from empirical curve fitting; PETI-340M composites, PMR-15 neat resin,
and PMR-15 composites
Note that A is a surface-dependent coefficient.
8, 15
For the quasi-isotropic PETI-340M laminates, the
expression for A is shown in Equation 2 where k 1 and k 2 are weight-loss fluxes. S 1 is the area of the un-
machined composite surface, and corresponds to much higher k 1, especially at the lower aging temperature
(T<288 ° C)
15
. Considering the dimension difference of PETI-340M composites (80 × 80 mm) and PMR-
15 composites (25 × 76 mm), weight loss per unit area of PETI-340M composites could be even smaller
compared to PMR-15 composites.
3.3.2 Thermal Analysis
56
A prior report on thermal oxidative aging of 4-phenylethynyl phthalic anhydride (PEPA-) end-capped
polyimide revealed asymmetric broadening of the tan(δ) peak at high aging temperature (316 °C),
indicating the growth of a thick oxidized layer with more compact packing and wider molecular weight
distribution
44
. This peak broadening has not reported for aging of PMR-15. However, similar peak
broadening and shouldering were observed for PETI-340M composites aged at 288 ° C (indicated by red
arrow), although it was not observed after aging at 232 ° C, Figure 3.2a.
Figure 3.2. Dynamic mechanical analysis of unaged and thermal oxidative aged PETI-340M. (a)
tan(δ); (b) Storage modulus.
Measured changes in T g are shown in Table 3.2, and indicate different aging mechanisms at 232 ° C and
288 ° C. T g was measured using the peak of the tan(δ) curve, and by the tangents of 𝐺 ′
. Aging at the lower
temperature showed a decrease in T g, a result of chain scission. However, aging at 288 ° C led to formation
of a condensed oxidized phase, and T g increased, a change that occurred in the early stage of aging (< 25
hr). Aging at 288 ° C for 1000 hr caused changes detected by DMA. Because of a shouldering effect, a
single T g could not be identified by the peak of the tan(δ) curve after aging at 288 ° C. A low-temperature
peak in the tan(δ) curve (at ~360° C) originated from the interior of the sample, while a broad peak at >
400 ° C corresponded to the oxidized layer.
44
Compact packing of the surface layer was attributed to the
loss of pendant phenyl groups
22
, and the generation of oxidative products. These factors led to stronger
(a) (b)
57
intra/inter molecular interactions between the polymer chains. The slight increase in 𝐺 ′
, Figure 3.2b, was
attributed to the reduction in free-volume and decrease of molecule mobility.
44
Similar changes in T g after aging was also reported for PMR-15 composites.
14
PMR-15 composites showed
an increase of T g when the aging temperature was >260 ° C, but a decrease in T g when the aging temperature
was 204 ° C. The disparities indicated that when conducting accelerated aging studies for polyimide
composites (such as PETI-340M, PMR-14, and TriA X), the aging temperature must be selected
appropriately to match the temperature of the anticipated application. Although aging conducted at higher
aging temperatures (> 288 ° C) reduce the time for experiments, such tests may also activate a different
aging mechanism compared to the service temperature.
Material Aging
Condition
Aging Time Tg
(Measured by
the peak of
tan(δ))
Tg
(Measured by
G’ transition)
°C hour °C °C
PETI-340M
Composites
Unaged 0 349.89 327.66
232 25 348.75 334.05
232 500 352.37 325.86
232 1000 341.11 305.91
232 1500 339.38 314.43
288 25 356.01 332.56
288 750 \ 339.98
288 1000 \ 340.53
288 1500 \ 338.72
Table 3.2. Change of glass transition temperature for PETI-340M composites aged at 232 ° C
and 288 ° C.
58
3.3.3 Morphology
Polished sections of unaged and thermal aged PETI-340M composites are shown in Figure 3.3. Fresh
composites were void-free and crack-free. Similarly, no crack or voids were observed for the specimen
aged at 232 ° C for 3500 hr, which is consistent with the weight loss exponent (B=1.035).
Figure 3.3. Microscopy image of cross-sections of unaged and aged PETI-340M Composites.
Aging at 288 ° C for 2000 hr produced microcracks in the surface plies of molded composites. Matrix
microcracks extended parallel to the fibers within tows, although they were initially intra-tow within surface
plies, and later propagated to inner plies. After aging for 3000 hr at 288 ° C, cracks were observed evenly
59
in all plies, and remained intra-tow. Crack generation and propagation from the surface to the inner plies
was reflected in the weight loss exponent coefficient (B=2.068).
Figure 3.4. A polished section of PETI-340M Composites aged at 288 ° C for 2500 hr.
In addition to near-surface cracks, micro-voids started to appear in resin-rich regions after 2000 hr aging,
Figure 3.4. The micro-voids formed in surface plies, and were evenly distributed in resin-rich regions. The
generation of micro-voids was associated with matrix weight loss and shrinkage caused by oxidation
reactions. The cracks did not propagate beyond fiber tow boundaries during aging at 288 ° C for up to 3000
hr. The cross-ply architecture of the fabric presented orthogonal tows at tow borders which acted as crack
arrestors. However, delamination occurred when long cracks extended across fiber-rich regions. When
aging time exceeded 4000 hr, fibers were damaged and often pulled out during polishing, preventing clear
sections.
60
Figure 3.5. Microscopy image of cross-sections of PETI-340M thermal-cycled for 400 cycles
and 1600 cycles.
The thermal cycling that occurs in orbit can be considered as low amplitude thermal fatigue, in which
microcracks begin as intraply cracks, grow into interply cracks, and finally cause delamination. In such
cases, microcracks arise from thermal cycling, appearing initially in surface plies, until a stress-relieved
state is achieved.
43
For PMR-15 composites, the crack density approaches an equilibrium value, with the
coefficient of thermal expansion reduced by 40% after 500 cycles between -156 ° C and 316 ° C.
45
In the
present study, however, a smaller temperature range, -54 ° C to 232 ° C, was used for PETI-340M to simulate
anticipated service environments.
Two primary types of damage arise during thermal cycling of composites - transverse microcracks and
delamination.
40
Transverse microcracks (TVMs) develop when the internal stress exceeds the transverse
strength of a lamina due to mismatch between thermomechanical properties of fiber and matrix, and
between lamina of different fiber orientations. Delamination develops when the thermally cyclic stress
generated exceeds the interlaminar strength of the laminate. The stress caused by thermal cycling cannot
61
exceed 50% of the tensile strength of PEPA-terminated polyimide for this thermal cycling condition.
39
Thus,
the polymer matrix is expected to resist microcracking during the thermal cycles used (-54 ° C to 232 ° C).
TVMs were observed after 400, 1600, and 2400 thermal cycles, but not after 800, 1200, and 2000 thermal
cycles. PETI-340M composites after 400 and 1600 thermal cycles are shown in Figure 3.5. Some minor
cracks appeared in the surface ply after 400 cycles. However, this type of crack only appeared on one side
of the laminate. The appearance of TVM is stochastic and not directly dependent on number of cycles. Thus,
these cracks are attributed to local stress concentrations arising during fabrication or in the sample mounting
press.
Figure 3.6. A crack of PETI-340M composites thermal-cycled for 1600 cycles.
A typical microcrack after 1600 thermal cycles is shown in Figure 3.6. Such micro-cracks were usually
observed within a single fiber tow. However, unlike the intra-tow cracks in thermally aged samples, which
did not extend across the tow boundary, cracks induced by thermal cyclic fatigue penetrated the adjoining
resin-rich area, as shown in Figure 3.6. Micro-voids were not observed in resin-rich regions. However,
fabrication voids/defects were observed within fiber tows, and these appeared to serve as initiation sites for
micro-cracks during thermal cycling. The second damage mode, delamination, was not observed in PETI-
340M composites for this cycling condition.
Based on observations and analysis of micro-crack formation during thermal cycling, crack initiation stems
from voids produced during fabrication. However, once a crack initiates, it relieves local stresses, which
62
delays formation of new cracks and arrests the growth of existing cracks.
38
Void-free laminates endure
thermal cyclic fatigue without discernible damage in the temperature range chosen (-54 ° C to 232° C).
However, in regions with voids, micro-cracks will initiate at the defect and propagate during cycling.
PMR-15 composites exhibit a different mechanism for crack-formation, in which micro-cracks initiate in
surface plies, then propagate to inner plies. Because of the low thermal oxidative stability of PMR-15,
thermal oxidation is conflated with thermal cycling fatigue, and thus crack growth follows a pattern similar
to thermal aging, which progresses from surface plies to inner plies. In contrast, PETI-340M is thermally
stable at the cycling temperature (-54 ° C to 232 ° C), and consequently oxidative aging plays a negligible
role during thermal cycling.
38
63
Figure 3.7. Micro-CT images of 2400 thermal cycled specimen. (a) 3-D geography; (b) an in-
plan cross-section.
Micro-CT images for unaged and thermally cycled PETI-340M composites were analyzed to identify
micro-cracks induced by thermal fatigue. The unaged composites were void-free structure, while after 2400
cycles, cracks appeared, as shown in Figure 3.7 (yellow frames).
Figure 3.7a shows a 3-D image of a transverse section after thermal cycling. In the image, white/gray
shading corresponds to low-density regions. The two large grey slabs represent air near the sample, while
light gray streaks within the dark core region are internal cracks, which extend primarily in directions
parallel to the plies. Figure 3.7b shows an in-plane section. Unlike Figure 3.7a, Figure 3.7b clearly reveals
the fiber tow architecture (darker shades represent matrix regions with lower density). The lighter gray
64
regions represent fiber tows, while the dark lines along the fiber tows in the yellow squares are cracks
extending parallel the fibers. Note that in this case, the cracks appear to extend within the same tow,
indicating that it may be the same crack.
Cracks appeared in cross-ply (non-crimp) regions throughout the sample thickness, which is consistent with
the crack morphology observed by light microscopy. The micro-cracking is attributed to stress
concentrations around fabrication-induced defects, such as incomplete impregnation, which is stochastic in
nature. Thus, the observed crack density is independent of the number of thermal cycles, and micro-
cracking appears largely in the middle plies, where incomplete impregnation usually occurs. Voids have
two primary effects on thermal cycling damage. First, they serve as stress concentrators and crack initiation
sites, but once initiated, the cracks relax local stresses, and thus delay initiation of new cracks.
44, 46
The
effect of this type of crack mechanism on mechanical behavior is described in a subsequent section.
3.3.4 Mechanical Properties Measurements
Thermal Oxidative Aging
There have been multiple investigations regarding the thermo-oxidative stability (TOS) of PMR-15
composites, and these studies report two types of thermo-oxidative transformations during thermal aging,
including crack initiation at cut edges and surface cracking at molded surfaces.
16
In production, polyimide
composites typically are molded into parts, and surface cracks are the dominating factor in mechanical
degradation. To better simulate production conditions, after aging, composite edges were trimmed 38 mm
and machined to coupons from the middle of plates for mechanical testing.
Compression properties are generally considered to be matrix-dominated, and they play a critical role in
service because of failures caused by out-of-plane forces. Both combined loading compression (CLC) and
open-hole compression (OHC) were performed. OHC provides a controlled simulation of a natural defect
65
or a fabrication artifact from fasteners. Compression failure can be susceptible to such features, which
increase the probability of buckling.
Figure 3.8. Mechanical change of thermal oxidized PETI-340M composites. (a) Combined loading
compressive strength; (b) Open-hole compressive strength; (c) Short-Beam Shear Strength.
66
Interlaminar shear strength (ISS) as a function of aging time was measured by conducting short beam shear
tests (Fig. 8). The decline in strength with aging time was attributed to degradation of the matrix and the
fiber-matrix interfacial bonding.
40
During SBS tests, specimens also are susceptible to local transverse
forces near the loading pins. Thus, detrimental interlaminar shear stresses that develop at local
discontinuities can cause local damage and premature failure. In this study, SBS tests showed transverse
tensile failure, fiber fracture by buckling, and interply delamination in valid tests.
47
In Figure 3.8, both CLC and OHC results showed a decrease in compressive strength after aging at 232 and
288 ° C. Aging at 288 ° C for 2500h caused a 36% decrease in OHC, which matched the decrease in CLC
strength (38 %). Aging at 288 ° C for 11000 h caused only a 13.7 % decrease in OHC strength, and the
decrease in CLC strength was slightly more (22.5 %). The rate of strength loss with aging time is described
by Equation 3 below, where S c can be compressive strength or SBS strength. The fitted m and n are shown
in Table 3.3.
Sc= -m*t+n (3)
Mechanical Test Aging
Temperature
m n
CLC
232 ° C 0.0115 509.6
288 ° C 0.0683 502.0
OHC
232 ° C 0.0045 337.3
288 ° C 0.0437 325.8
SBS
232 ° C 0.00226 92.3
288 ° C 0.0184 90.78
Table 3.3. Coefficients for the mechanical property degradation of PETI-340M composites aged
at 232 ° C and 288 ° C.
The rate of CLC strength loss during aging at 288° C was ~6× faster than for aging at 232° C. However, the
rate of OHC strength loss was ~10× faster at 288 ° C, while the rate of SBS strength loss was 8.1 faster. The
67
difference in the rates of CLC and OHC strength loss was attributed to the fact that the matrix became more
rigid and brittle during aging (particularly at 288 ° C), which increased the crack initiation and propagation
rates during hole drilling during OHC sample preparation. The different m value for CLC and SBS strength
arose because SBS strength was more sensitive to fiber-matrix interfacial strength, which also decreased
during aging.
40
Chemical changes in the matrix induced by thermal oxidative aging and crack formation at the surface were
the primary factors governing strength loss in polyimide composites. Indeed, the two aging temperatures
activated different mechanisms of mechanical degradation. The DMA results showed an increase in storage
modulus 𝐺 ′
after aging at 288 ° C, but a decrease in 𝐺 ′
after aging at 232 ° C. The decrease in 𝐺 ′
(and
associated …) resulted in the mechanical degradation observed from aging at 232° C. In contrast, the
formation and propagation of micro-cracks from surface plies to inner plies caused the mechanical
degradation observed from aging at 288 °C. At this temperature, post-curing and oxidation of the matrix,
which changed the 𝐺 ′
and T g, was secondary.
In PMR-15 composites, the degradation in compression strength resulting from thermal oxidation occurred
by thermal (time-dependent) and oxidative (weight-loss) mechanisms
16
. The same was true for PETI-340M
composites studied here, where the rate of weight loss rate correlated to the retention of compressive
strength.
16
However, aging of PMR-15 composites also indicated that the weight loss behavior was not a
reliable indicator of mechanical behavior for aging below 260 ° C. At higher aging temperatures (>260 ° C),
the concentration of microcracks in the materials was much greater than for aging at lower temperature
(204 ° C),
16
although the compression strength was similar.
Matrix-dominated properties such as CLC and SBS strength, are most likely to manifest effects of thermal
cycling (fiber-dominated properties, such as tensile strength and modulus, show little decrease with thermal
cycling, despite initiation of microcracks.
43
) SBS tests were performed at two temperatures (20°C (RT) and
232 ° C), to assess the effects of test temperature on ISS strength. When tested at RT, the ISS of
68
unidirectional {0} 20 PMR-15 composites was unaffected by thermal cycling, despite the fact that cracking
and delamination were observed. However, when tested at 316 ° C, ISS decreased by 15% after 500 cycles.
40
Figure 3.9. Mechanical change of thermal cycled PETI-340M composites. (a) Combined loading
compressive strength; (b) Short-Beam Shear Strength.
Strength values as a function of number of thermal cycles is shown in Figure 3.9. The CLC strength showed
a slow increase after 400 thermal cycles, Fig 4a, while SBS strength remained largely unaffected at both
test temperatures. Micro-cracking in surface plies caused an initial but slight decrease in CLC strength.
However, such cracks relaxed local stresses and also arrested crack growth, limiting crack propagation.
Subsequently, post-curing became the dominating factor, causing an increase in CLC strength. However,
the SBS strength showed no increase after 400 cycles, because ISS is more sensitive to cracking within
fiber-rich regions.
44, 46
69
Unlike PETI-340M composites, PMR-15 composites exhibit microcracking in inner plies during thermal
cycling. The cracking reduces the matrix-dominated residual strength at both 20 and 232 ° C, although the
effect is less pronounced at 232 ° C, and no change in modulus is observed.
43
In contrast, PETI-340M
composites show no change in matrix properties after thermal cycling, and minor cracks are arrested by the
cross-ply weave geometry. The retention of strength after thermal cycling can be attributed to the superior
thermal stability of the matrix, combined with the woven fiber architecture. In general, PETI-340M
composites exhibit greater retention of mechanical properties after aging and thermal cycling compared to
PMR-15 composites.
3.4 Conclusion
Aging mechanisms were identified for PETI-340M composites aged at two temperatures. Strength loss after
aging at 232° C was attributed to matrix weight loss caused by chemical reactions resulting in chain scission,
with a decrease in T g. For aging at 288° C, degradation was caused by both matrix oxidation, (and an
increasing T g), as well as progressive micro-cracking. ISS was more sensitive to micro-cracks and
delamination in the middle plies compared to unnotched compressive strength. Future investigations of
thermal oxidative stability of composites should consider judicious selection of aging temperatures that
reflect the intended service temperatures, and select mechanical test methods accordingly.
During thermal fatigue, micro-cracking of PETI-340M composites initiated at fabrication defects, which
concentrated stress locally during thermal cycling. Despite micro-cracking, PETI-340M composites
retained matrix-dominated properties to a greater degree compared to PMR-15 composites. The superior
strength retention was attributed to the greater oxidative stability and the greater tensile strength of the
PEPA-terminated polyimide matrix.
39
Void-free composites (PEPA-terminated polyimide PETI-340M and
TriA-X composites) show no cracks or voids after thermal cycling. The superior strength retention of
70
PEPA-terminated polyimide composites indicates potential for long-term structural use in aerospace
applications.
71
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Abstract (if available)
Abstract
Polyimides (PI’s) are widely used in high-temperature composite applications because of the high glass transition temperatures and thermal stability. However, polyimides generally suffer from poor processability, cracking induced by thermal shock, and thermal oxidation. A major limitation of conventional polyimide resins is low toughness, which can lead to (premature) brittle failure. However, conventional thermoset polyimides such as PMR-15, show low elongation-at-break due to a rigid molecular structure and a low degree of polymerization of imide oligomer. The brittle nature of the resin causes microcracks to propagate under repeated thermal cycling, which severely limits service life. Moreover, the low molecular mobility at high temperature (>Tg) leads to processing challenges and makes it difficult to obtain composite parts with low porosity. Another concern with the polyimide chemistry is that methylene dianiline (MDA), which has been used in traditional PMR polyimide formulations, is known to be carcinogenic. ❧ In recent decades, 4-phenylethynylphthalic anhydride (PEPA) end-capped polyimides have been developed to address such limitations, and these polymers generally offer wider process windows and improved heat resistance, although they are not immune to oxidation and thermal shock damage. A new class of PEPA-terminated PMDA-type polyimide (pyromellitic dianhydride), denoted TriA X1, has been developed in an attempt to address the limitations of brittleness and poor processability associated with most polyimides. The TriA X polyimide is derived from PMDA and 2-phenyl-4, 4’-diaminodiphenyl ether (p-ODA). The pendent phenyl group of p-ODA restricts molecular interaction and decreases the rotational flexibility of the chain. The distinctive chemical structure of the polymer imparts a reduced elastic modulus during heating (from 109 Pa to 107 Pa before and after glass transition), which facilitates the processing of carbon fiber composites. Other PEPA-terminated polyimides have also been developed, for example, PETI-340M, featured in their low viscosity during processing. However, for PEPA-terminated composites, its imidization kinetics and the effect of different diester groups on the curing kinetics were not fully understood. The thermal-oxidative stability and the stability under thermal cyclic fatigue were not understood as well. Therefore, the motivation of this dissertation is to investigate the imidization behavior of PEPA-terminated polyimide, and to evaluate the retention of mechanical property in thermal-oxidative aging conditions and thermal cycling.
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Creator
Li, Xiaochen
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Core Title
Processing, thermal-oxidative stability and thermal cyclic fatigue of phenylethynyl-terminated polyimides
School
Viterbi School of Engineering
Degree
Doctor of Philosophy
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Chemical Engineering
Publication Date
04/25/2021
Defense Date
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