Toughened Plastics I - American Chemical Society

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20 Structure-Property Relations of High-Temperature Composite Polymer Matrices R. J. Morgan , R. J. Jurek , D. E. Larive , C. M. Tung , and T. Donnellan 1

1 2

3

1

1

2

3,4

Michigan Molecular Institute, Midland, MI 48640 Northrop Corporation, Hawthorne, CA 90250-3277 Naval Air Development Center, Warminster, PA 18974

The

structure-deformation-failure

mode-mechanical property

rela-

tions of high-temperature thermoplastic polyimide and thermoset bismaleimide (BMI) polymeric matrices and their composites will be discussed. In the case of polyimides, the effects of test temperature, thermal history, strain rate, type of filler, and filler volume fraction on structure-property

relations will be discussed. For BMIs we report

systematic Fourier transform infrared spectroscopy and differential scanning calorimetry studies of the cure reactions as a function of chemical composition and time-temperature

cure conditions and then

describe the resultant cross-linked network structure based on our understanding of the cure reactions. The optimization of the BMI matrix toughness will be considered in terms of network structure and process-induced matrix microcracking. We also describe optimization of composite prepreg,

lamination and postcure conditions based

on cure kinetics, and their relationship to the BMI

viscosity-time-

-temperature profiles. The critical processing-performance limitations of high-temperature polymer matrices will be critically discussed, and toughening approaches to address these limitations, such as toughness over a wide temperature range, will be presented.

I O L Y M E R MATRIX FIBROUS COMPOSITES that exhibit good composite " h o t wet" properties, processibility, and toughness (residual strength after impact) Current address: Grumman Corporation, Department 584, M.S. 802-26, Beth Page, Long Island, N Y 11714-3580.

4

0065-2393/93/0233-0493$06.00/0 © 1993 American Chemical Society

In Toughened Plastics I; Riew, C., et al.; Advances in Chemistry; American Chemical Society: Washington, DC, 1993.


494

RUBBER-TOUGHENED PLASTICS

will be required for future advanced military aircraft. High-temperature composite matrices such as thermoplastic polyimides exhibit both processing and toughness limitations, whereas thermoset bismaleimides are more pro cessible but exhibit a more brittle mechanical response. The two principal mechanisms for toughening carbon fiber-polymer matrix fibrous composites involve polymer matrix viscoelastic flow and crack deflection by the incorporation of microscopic second-phase additives into the polymer matrix. For high glass-transition temperature (T ) polymer matrices, high yield stresses at service environment conditions limit viscoelastic flow because of increasing yield stress with decreasing temperature below T . Optimization of viscoelastic flow in a cross-linked thermoset matrix additionally requires a fully reacted defect-free network (J). Networks that contain defects i n the form of unreacted groups initiate crack propagation early in the deformation process and these cracks truncate significant viscoelastic flow. Hence, there is a need to characterize the cure reactions and kinetics involved in the formation of high-temperature thermosets such as bismaleimides (BMIs) to optimize network structure for high toughness. A ductile, second-phase additive that deforms viscoelastically by crack bridging also can circumvent the yield stress-temperature dependence dilemma of the polymer matrix continuous phase (2). Small, submicrometer additives, with dimensions well below the critical matrix flaw size, can cause a tortuous crack propagation path i n the polymer matrix that results i n increased matrix fracture energy. Appropriately oriented organic microwhiskers that readily split longitudinally and are incorporated in the composite interfiber tow regions seem most promising from a high-fraeture-surface viewpoint. In this chapter we report on the following studies: g

g

1. The deformation, failure modes, and toughness of poly-4,4'oxydiphenylene pyromellitimide (POPPI: Structure I) and its

Γ

Ο Il

Ο II

Structure I. The chemical structure of poly-4,4-oxydiphenylene pyromellitimide (POPPI). ,


20.

MORGAN ET AL.

Structure-Property Relations of Polymer Matrices 495


composites as a function of test temperature (23-300 °C), thermal history (300-500 °C), strain rate, type of filler (short glass fibers or graphite flakes), and filler volume fraction (15-40 wt % graphite flakes). 2. The cure reaction characteristics: optimal prepregging, lamination, and postcure conditions of thermoset 4,4'-bismaleimidodiphenylmethane-0,0'-diallyl bisphenol A (BMPMD A B P A ; 1:1 molar) B M I (Matrimid 5292, Ciba-Geigy)-carbon fiber (AS4) composites. The chemical structure of the B M P M and D A B P A monomers are illustrated in Structure II. Toughening procedures by network structure optimization or the incorporation of second-phase crack deflectors or viscoelastic bridgers will be described.

Structure II. The chemical structures of BMPM and DABPA monomers.


496

R U B B E R - T O U G H E N E D PLASTICS


Results and Discussion Thermoplastic Polyimides. The P O P P I polyimide fails in tension by massive crazing via cavitation and fibrillation rather than by the shearbanding mode, which would involve a larger volume of matrix viscoelastic flow. From tensile mechanical property tests of unfilled and filled P O P P I matrix we determined the following effects on polyimide toughness (toughness in these studies is defined as the area under a tensile stress-strain curve): 1. In Figure 1 we illustrate the toughness of the P O P P I matrix and its composites as a function of temperature. The toughness of all the polyimide materials decrease with increasing temperature. This decrease is a consequence of the lower strength of the craze fibrils with increasing temperature that results in the

0

SO

100

150 Temperature

200

250

300

(°C)

Figure 1. The toughness of POPPI matrix and its composites as a function of temperature.



20.

MORGAN ET AL.

Structure-Property Relations of Polymer Matrices

growth of larger craze cavities and produces associated higher stress concentrations. In the case of matrices that deform by shear banding rather than crazing, the polymer matrix ductility increases with increasing temperature (3). Toughness de creases with increasing filler volume fraction of the 1-5-μ m dimensional graphite flakes (Figure 1). Fracture topography observations indicate that the fillers inhibit the craze growth process. For larger dimensional fillers (glass fibers of ~ 10-μπι diameter), the filler decreases the composite toughness to a greater extent than the smaller graphite flakes (Figure 1) as a result of higher fabrication stress concentrations at the fiber-matrix interfacial region induced by the fiber-matrix thermal coefficient of expansion mismatches. 2. The toughness of P O P P I and its composites decreased with increasing strain rate because of the inhibition of the viscoelastic flow processes at faster testing speeds. 3. Thermal history also can modify the toughness of the poly imide composites. In the temperature regime near the P O P P I T (300-375 °C), the toughness increased after 24-h anneal as a result of relaxation of composite fabrication strains and any associated microvoids. However, exposure to temperatures above 400 °C in air results in oxidation-induced matrix crossfinking or chain scission and a corresponding decrease in ambient craze fibrillation capability and toughness. For exam ple, air exposure at 450 °C for 24 h results in a > 50% loss in ambient composite toughness. g

Thermoset Bismaleimides. BMI Cure Reaction Studies. System atic Fourier transform infrared spectroscopy (FTIR) and differential scanning calorimetry (DSC) studies were carried out on the B M P M - D A B P A B M I system as a function of chemical composition and temperature-time cure conditions. From the F T I R and D S C studies, together with published data by Ciba-Geigy researchers (4-6), we have arrived at the following conclusions: There are three principal temperature regimes associated with the cure reactions, namely: 1, 100-200 °C range; 2, 200-300 °C range; and 3, 300-350 °C range. In Figure 2 a D S C plot of a B M P M - D A B P A ( 1 : 1 molar)-AS4 carbon fiber composite prepreg exhibits exotherm peaks in each of the three regimes.


497


50

σ

w

w

a

G Ο

w W

W

G

Φ-

CD 00

ce Figure 2. DSC plot of BMPM~DABPA(l:l molar)-AM carbonfibercomposite prepreg after 30 days storage at Η Π — 13 °C, 10 °C/min in N atmosphere. CO 2


20.

MORGAN ET AL.

Structure-Property Relations of Polymer Matrices 499

In regime 1 (the 100-200 °C range), the "ene" reaction Ο

0

II C H

2

= C H - C H

C H

2

2

C

H

=

C

H

+

2

Ο

; H

2

= C H

C H

2

R

— C H — C H

1

(M —

R ~U

| J

2

II ο

C DABPA}


X

[|

II

C H

2

— C H '

II c

X

II ο C BMPM}

Ν

—R

2

Γ)

~

™ 2C %

II Ο

where

0

II ο

R>

(1) occurs slowly and there is evidence it may be reversible. Gel permeation chromatography indicates that only a small molecular weight increase occurs from 520 to 680 for the B M P M - D A B P A system after 16 h at 100 °C. Consistent with our viscosity-time-temperature observations, this reaction would only occur at significant rates above 180 °C. (In this lower temperature regime the presence of propenyl isomers rather than ally! isomers i n the diallyl bisphenol A will enhance chemical reactivity due to conjugation of the double bond with the phenyl ring.) In regime 2 (the 200-300 °C range), a number of chemical reactions occur in the following sequence: The "ene" reaction 1 occurs at a significant rate; then the B M P M homopolymerization, reaction 2 occurs: one

c=o + o=c Ν

c=o

Ν

—

CH—CH—CH—CH—

I . ,c o' N

N "

I I I cro^c c:o o' V

(2) or a longer chain "ene'" reaction product.


500


A D S C plot of the B M P M monomer alone exhibits a reaction peak centered in the 205-210 °C range. At higher temperatures both "ene" homopolymerization, reaction 3, and " e n e " - B M I double-bond cross-linking polymerization (Diels-Alder reaction), reaction 4, occur:


η

R

1

C H Z Z Z C H R

•

2

R ^ C H

C H R

^

2

(

3

)

where R and R are chain continuations of the "prepolymer" formed in the "ene" reaction; 1

2

0

II c C H =

C H

C H

2

—

C Η

Ν— R

2

0 Η

0 C H — C H

+

0

(4)

Η

where R and R are chain continuations of the "prepolymer" formed in the "ene" reaction and 1

2

0

C H

R

3

11

=

C H

0

or a longer chain "ene" reaction product.


20.

MORGAN E T AL.

501 Structure-Property Relations of Polymer Matrices

Reactions 3 and 4 are indicated by the D S C peak centered at 254 °C for the prepreg in Figure 2. A low-temperature shoulder exhibited by this peak is associated with B M P M homopolymerization. Dehydration of the hydroxy! groups of the D A B P A initiates above the 240 °C, enabling the formation of ether cross-links: HO

R

OH

+

HO

R

OH

•

HO

R

Ο

R

OH


(5) where R represents polymer chains formed from reactions 1-4. A D S C plot of the D A B P A monomer alone exhibits thermal activity above 240 °C that is associated with dehydration. Thermogravimetric analysis weight loss and F T I R studies confirm that ~ 75% of the hydroxyl groups undergo dehydration with an associated 2-wt % H 0 loss in this temperature region. In regime 3 (the 300-350 °C range), further cure occurs as the specimen test temperature exceeds the T of the resin. It is important from the network-toughness viewpoint that the reaction be allowed to proceed to completion, thus eliminating network defects in the form of unreacted groups. D S C studies as a function of postcure conditions indicate a postcure schedule of 2 h at 200 °C, 6 h at 250 °C, and 1 h at 300 °C would be optimal for full cure and minimization of composite fabrication stresses. A resultant T near 340 °C results from this postcure schedule. A l l postcure conditions in our studies at 250 °C and above do result in inherent microcracks in composite 0°, 90° laminate layers. 2

g

g

BMI Network Modeling. Based on our understanding of the B M P M - D A B P A cure reactions, network modeling studies have been carried out. The chemical structure of the B M P M - D A B P A "ene" adduct is illus trated in Structure III. The three double bonds capable of chain extension and cross-finking are designated A, B, and C in Structure III. Based on the cure reaction characteristics, geometric controlled sequential reactions for the trifunetional "ene" molecule, and symmetry conditions that allow the network to fully react, we show a proposed ring chain structure in Structure IV that would allow sequential polymerization of all three double bonds present in the "ene" molecule. The ring chains cross-link via A, B, or C "ene" double bonds (ring formation, therefore, would not occur at such cross-finks) or interchain etherification by hydroxyl dehydration. BMI Prepreg Development. Isothermal viscosity profiles in the 8 5 95 °C prepregging temperature range indicate that initial viscosities of the B M P M - D A B P A (1:1 molar) resin are in the 700-1200-cps range and only increase slowly with time. For example, the viscosity of the B M P M - D A B P A (1:1 molar) system only increases to 5000 cps after 1200 min at 91°, thus


502



(Ο*

Structure III. The chemical structure of BMPM-DABPA "ene" adduct. indicating adequate viscosity-time profiles for prepregging in the 85-95 °C temperature range. The B M P M - D A B P A system was prepregged with adjustable wedge slit dies developed for 3000 and 12,000 filament-tow AS4 carbon fibers. The carbon fibers contain epoxide sizings that chemically degrade during postcure conditions of the Β M I composite and cause poor fiber-matrix interface integrity (7). Hence, during the prepreg process the sizing is burned off the fiber at 400 °C in air. B M P M - D A B P A ( 1 : 1 molar)-AS4 carbon fiber prepregs were consis tently fabricated to meet government specifications. BMI Composite Lamination. The optimum conditions to produce B M I composites from the prepreg using a programmable press (Tetrahedron) and miniclave were investigated. Dynamic mechanical analyses ( D M A ) plots exhibited broad minima in damping in the 160-190 °C temperature range that suggest a relatively constant viscosity for adequate time periods to achieve laminate consolidation and associated acceptable composite resin and void content. A variety of lamination conditions for B M P M - D A B P A ( 1 : 1 molar)-AS4 carbon fiber composite prepregs were studied and optimized. The resultant resin weights in the composite ranged from 15 to 31 wt %. The B M I resin



20.


MORGAN ET AL.

(C)

BMI Double Bond Homopolymerization

(A) and (Β) Έ Ν Ε

J

503

Η

Double Bond Polymerization

An "ENE" Molecule

•

Hydroxyl Group

Structure TV. Ring chain structure from "ene" double-bond polymerization. exhibited good flow characteristics at 177 °C and the resultant composites exhibited good surface appearance. BMI Composite Characterization. After postcure, the B M I compos ites were characterized to form a standard database to evaluate against toughening procedures. Scanning electron micrographs of the postcured fractured composite revealed a thoroughly and uniformly impregnated com posite. Short beam shear, tensile, and toughness ( G ) mechanical properties, together with ultrasonic nondestructive evaluation of the composites, were monitored. I c

BMI Toughening Procedures. The optimization of the postcure con ditions based on conflicting requirements of the desirability to attain a fully reacted B M I network and to minimize composite fabrication stresses associ ated with the mismatch of the fiber and matrix thermal coefficients of expansion were explored. Techniques to incorporate crack deflectors, such as organic microwhiskers of polyamide fibers (Kevlar 49; K-49) or finely divided fused silica, into the composite via the prepreg process were carried out. Our experience with the fluidized-bed approach led to variable, inconsistent third-phase



504

Table I. Average Mode-I G Toughness Values for BMPM-DABPA(l:l molar)-AS4 Carbon Fiber Composites I c

Standard Average Postcure Fiber Volume Average Thickness Condition Deviation in G Fraction (J/m ) (mm) (NorA) (J/M ) (%)


Additive in Resin (%)

Ic

2

a

2

0

2.50

Ν

65.7

259.9

43.2

0

2.71

Ν

63.5

301.6

32.1

0

2.75

Ν

—

324.8

34.3

0

2.65

Ν

63.8

326.9

16.2

0

2.65

A

63.8

331.1

6.0

P E I (3.7)

2.48

Ν

64.0

462.2

225.7 47.8

P E I (8-10)

3.34

A

66.7

397.3

P E I (7-8)

2.63

A

71.0

338.2

15.6

P E I (7-8)

2.95

A

65.0

437.3

63.7

K - 4 9 (8.6) K - 4 9 (8.6)

2.72

Ν

32.1

A

66.5 70.4

381.3

3.09

313.4

52.8

Ν indicates normal postcure 2 h at 200 °C and 6 h at 250 °C. A indicates additional postcure 2 h at 200 °C, 6 h at 250 °C, and 1 h at 300 °C.

a

additive concentrations as a result of time-dominated instabilities i n the fluidized-bed apparatus. Hence, we concentrated our efforts on the incorpo ration of crack deflectors into the prepreg via formation of initial slurries of the additive in the B M P M - D A B P A resin. Hence, the crack deflectors were incorporated into the B M P M - D A B P A mixture prior to prepregging to produce a slurry. The additives increased the resin viscosity by an order of magnitude, but prepregs were prepared successfully and the crack deflectors were uniformly distributed in the interfiber tow regions. For the introduction of viscoelastic crack bridgers into the composite we used polyetherimide (PEI) additive and the following procedure: The P E I is dissolved i n methylene chloride with the appropriate amount of D A B P A ( B M P M is not soluble in methylene chloride) and the methylene chloride is evaporated at 40 °C. Then the D A B P A - P E I mixture is mixed with the B M P M at 130 °C and prepregged. In Table I we summarize the composite toughness ( G ) values as a function of polyetherimide (PEI) additive, polyamide (Kevlar 49) powder, composite thickness, postcure conditions, and fiber volume fraction. Our principal conclusions from these data are as follows: I c

• The average G values of 330 J / m for the unmodified com posites were similar for both a 250 and 300 °C final postcure temperature. However, the 300 °C postcured composite had a higher T (near 350 °C) with no decrease i n the 23 °C G value. The higher 300 °C postcure produced two counter acting effects on the G composite toughness values: Higher composite fabrication stresses, and probably associated microc2

I c

g

I c

I c


20.

MORGAN E T AL.


rack formation, as a result of higher processing temperatures caused a decrease in G composite toughness, whereas the higher postcure temperature increased the B M I network per fection and increased the G composite toughness. I c

I c

• For fiber volume fractions of approximately 65%, the incorpo ration of P E I increased the G values 32% to 440 J / m for a 300 °C postcured composite. Individual data points from the G tests were recorded as high as 900 J / m , which suggests excellent potential for this toughening procedure, but also suggests the probability of inhomogeneous distribution of the P E I within the composite test specimens. 2

I c

2


I c

• For a fiber volume fraction of 66.5%, a 15% increase in G was achieved by the incorporation of polyamide (Kevlar 49) powder.

I c

• There is a general downward trend in G values with increas ing fiber volume fraction. From the available data illustrated in Table I for similar additive concentrations and postcure condi tions, a 1 % increase in fiber volume fraction results in a 5 % decrease in G composite fracture toughness. This observation identifies the need to optimize composite toughness at the expense of stiffness as a function of fiber volume fraction. I c

I c

Summary Thermoplastic polyimide composite matrices are relatively tough as a result of crazing. However, their toughness is decreased with increasing temperature, strain rate, and fillers. Β M i s , however, are more brittle and require network structure optimization and the incorporation of additive toughening agents. The cure reactions of B M P M - D A B P A bismaleimide are complex. In the 100-200 °C range, the B M P M and D A B P A react to form an "ene" prepolymer. In the 200-300 °C range, the "ene" molecule polymerizes and cross-links by (1) cooperative reactions between the three double bonds in the "ene" molecule and (2) dehydration through hydroxyl groups, with associated moisture evolution. In the 200-250 °C range, full cure is not attained because of diffusion restrictions associated with the glassy state. Full cure is nearly attained after a 1-h 300 °C postcure that results in a final T of near 350 °C. The cross-linked network structure is complex as a result of the pentafunctional nature of the "ene" prepolymer. Optimum prepreg, lamination, and postcure conditions were ascertained for processing B M P M - D A B P A Q : molar)-AS4 carbon fiber composites and for the introduction of toughening agents, namely, polyetherimide (PEI) particulates and polyamide (Kevlar 49) whiskers. g


505

506


G composite toughness values are enhanced by incorporation of P E I particulates and polyamide (Kevlar 49) powder into the B M I composite matrix. G composite values decrease with increasing fiber volume fraction, which requires composite toughness to be optimized at the expense of stiffness. I c

I c


Acknowledgment This work was partially supported by N o r t h r o p / N A D C Contract No. N62269-88-C-0254.

References 1. Morgan, R. J.; Kong, F. M.; Walkup, C. M . Polymer 1984, 25,

375.

2. Morgan, R. J. Proc. Annual ESD/ASM Advanced Composites Conf.; ASM Interna tional: Metals Park, O H , 1986; p 179.

3. Morgan, R. J.; Walkup, C. M . J. Appl. Polym. Sci. 1987, 34, 37. 4. Chaudhari, M.; Galvin, T.; King, J. SAMPE J. July / August 1985, 17.

5. Lee, B. H.; Chaudhari, Μ. Α.; Blyakhman, V. Polymer News 1988, 13, 297. 6. Zahir, S.; Chaudhari, Μ. Α.; King, J. Makromol. Chem., Makromol. Symp. 1989, 25, 141. 7. DePruneda, J. H . S.; Morgan, R. J. J. Mater. Sci. 1990, 25, 4776. RECEIVED

for review March 6, 1991.

ACCEPTED

revised manuscript June 8,

1992.


Toughened Plastics I - American Chemical Society

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