Matrix Ductility and Toughening of Epoxy Resins - American Chemical

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of Epoxy Resins G. Levita Department of Chemical Engineering, Industrial Chemistry, and Materials Science, University of Pisa, Via Diotisalvi 2, 56100 Pisa, Italy

Epoxy resins can be modified by a number of reagents to produce solids that possess a variety of mechanical properties. Some of these reagents lead to the formation of two-phase structures that are much tougher than single-phase epoxies. Toughness depends largely on the physical and geometric properties of the separated phase. However, the properties of the rigid epoxy matrix are equally important, because the plastic processes responsible for toughening take place within the matrix. The ductility of epoxies can be varied in several ways. Some of them and the consequences for fracture resistance are considered in this chapter.

THE THEORETICAL STRENGTH, a , OF ELASTIC SOLIDS is given (I) by the t

simple model a ~ ( £ 7 / d ) , where £ , 7 , and d are the elastic modulus, the surface energy, and the interatomic distance, respectively. For rigid polymers a exceeds the experimental values by about one order of mag nitude. In the case of epoxy resins the difference between theoretical and experimental values can be significantly lower. The failure stress of resorcinol resins, for instance, is —160 M P a (2), and the ratio of theoretical to exper imental strength is about 2. Such a remarkable performance has been at tributed (2) to the high packing efficiency of epoxy networks because of strong segmental interactions that originate from hydrogen bonds. The number of mechanically active molecular segments consequently increases. The strength of polymers is controlled by the number of macromolecular chains that carry the external load. A n increase in this number leads to a lowered concentration of overstressed bonds. Overstressing is reduced when t

0 5

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0065-2393/89/0222-0093$07.50/0 © 1989 American Chemical Society

In Rubber-Toughened Plastics; Riew, C.; Advances in Chemistry; American Chemical Society: Washington, DC, 1989.

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stress-relief mechanisms, such as yielding, come into play. In the case of yielding, stress is transferred from overstressed regions (typically the crack tip) to adjacent zones, where it levels off to the yield value. If postyield deformability exists, an energy sink is also provided. Indeed, the only way to increase fracture resistance is to promote plastic processes. In the case of cross-linked polymers, these processes are severely restricted. As a con sequence, epoxy resins are rather brittle. The fracture energy, G , of glassy resins (~ 100-500 J/m ) is 10-40 times lower than that of polycarbonate, which ranks midway between brittle and tough polymers. I c

Downloaded by UCSF LIB CKM RSCS MGMT on November 20, 2014 | http://pubs.acs.org Publication Date: May 5, 1989 | doi: 10.1021/ba-1989-0222.ch004

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To activate plastic processes, molecular mobility must be increased. Unfortunately, this increase is accompanied by a decrease of elastic modulus, heat-distortion temperature, and hardness. Such a contradictory situation can be illustrated by the results of Mostovoy and Ripling (3, 4) who found that the fracture energy in amine- and anhydride-cured epoxies decreased as the modulus increased. The molecular mobility of interest here gives rise to conformational transitions such as those proposed by Robertson for yield ing (5). Short-range mobility, which is responsible for secondary relaxations (particularly β), can also play a role. Yielding behavior is of great importance in fracture. Specimen sizing is based on the value of the yield stress, which is controlled by the degree of plastic constraint. Plane stress or plane strain are limit conditions that can induce dramatic transitions from ductile to brittle behavior, as in the case of polycarbonate. Rubber toughening is the major source of crack-resistant polymers. This technology has been applied to almost every thermoplastic and also to ther mosets. A l l polymer systems consist of a rigid matrix containing a soft smallparticle phase whose behavior is supposed to be rubbery. Because the mor phologies and elastic constants of all systems are similar, one might expect the basic mechanisms of toughening to be, as a first approximation, the same for all polymers. In rubber-toughened polymers, in order to avoid loss of rigidity, the volume fraction of the dispersed phase is kept to the minimum value necessary for the attainment of a satisfactory level of toughness. Reinforcing mechanisms must exist within the matrix if they are to be active on a large scale. Thus the properties of the matrix polymer are at least as important as those of soft filler. A great deal of the fracture resistance of rubber-modified polymers comes from the deformation behavior of the ma trix. Figure 1 illustrates how the fracture energy of toughened thermoset plastics indeed depends upon the fracture energy of the matrix. A l l data are relative to blends containing about 10% rubber (mostly carboxyl-terminated butadiene-acrylonitrile copolymer (CTBN)) tested at nonimpact rates. The line in Figure 1 has no particular meaning except to represent an amplifi cation factor of 10. Although G can occasionally be as high as 1 k j / m , a very high value for cross-linked polymers, most G data for neat resins are grouped around 2

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Matrix Ductility and Toughening of Epoxy Resins


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Figure 1. Relationship between the fracture energy of some rubber-toughened thermosets, decomposite), and the fracture energy of the corresponding ma trix, Gic(matrix). Rubber content —10%. Data from refs. 6-13. 100-300 J / m . Different epoxy matrix properties can be obtained by a num 2

ber of approaches, including modification of cure conditions, use of chain extenders, and use of high-molecular-weight components. The effects of such techniques on fracture behavior are dealt with in this chapter.

Experimental Methods and Materials A series of commercial diglycidyl ether of bisphenol A ( D G E B A ) resins of 400-4000 molecular weight (M W) were used. The epoxy prepolymer with the lowest molecular weight (Epon 828, M W 390, Shell) was a liquid at room temperature. The mediumto high-molecular-weight resins—Eposir 7161 ( M W 980), 7170-P ( M W 1640), and 7180 ( M W 3700) by Società Italiana Résine (SIR) were solid (melting point 50, 80, and 110 °C). The hardeners were piperidine (PIP, boiling point 106 °C) and diami-



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nodiphenyl sulfone (DDS, melting point 175 °C). PIP was used at a concentration of 5 phr. DDS was employed with a little excess (5%) over stoichiometry to com pensate for its sluggish reactivity. The reactive rubbers were either amine-terminated butadiene-acrylonitrile co polymer (ATBN) or CTBN (BF Goodrich). ATBN 1300 x 21 and ATBN 1300x16 were used for the D G E B A - P I P formulations, with acrylonitrile (AN) contents of 10 and 16%, respectively. In the case of D G E B A - D D S formulations, only the CTBN 1300 x 13 was used (AN 27%). In all blends, the rubber level was 10% by weight. Technical-grade bisphenol A (BPA) was used as the chain extender. Details about sample preparation and cure conditions of formulations based on the 828 resin can be found in refs. 6-8. High-molecular-weight resins were first melted at 180 °C before adding DDS. The mixtures were then gently stirred until they clarified. The viscous liquids were poured into aluminum trays and cured under vacuum (to ensure outgassing) at 180 °C for 12 h. Rubber-modified formulations were obtained by adding CTBN after the hardener dissolved into the resin. Fracture tests were carried out with a three-point bend assembly. Samples had span-to-width ratios of 4 or 8. The critical stress-intensity factor, K , was obtained by the slope of a versus l/YVdi plots (σ is stress at the onset of crack propagation, Y is a geometrical parameter, and a is crack length). G was obtained by the rela tionship ic

c

0

i c

cJ^fJ^

(1)

where ν is Poisson ratio and Ε is elastic modulus. Dumbell specimens were used for tensile tests. Scanning electron microscopy (SEM) was carried out (ISI DS180 and J E O L T-300 microscopes) on samples broken at room temperature. Surfaces were gold coated. Shear moduli were measured with a torsion tester according to ASTM (American Society for Testing and Materials) D 1043 61T.

Discussion Effect o f C u r e C o n d i t i o n s . The development of good mechanical properties requires the attainment of a sufficiently high degree of conversion, although the highest fracture resistance develops when the conversion is somewhat lower than the maximum achievable (14). The cross-link process is delayed by an increase in viscosity, and it can be brought close to the topological limit only by opposing the reduction of molecular mobility. In creased conversion is typically obtained by curing or postcuring at high temperatures, which does not necessarily result in better mechanical prop erties. Curing at different temperatures can in fact lead to diverse network topologies. The effect of cure temperature (i ) on the fracture properties of toughened epoxies has been studied by Levita et al. (6) and by Butta et al. (7). cure

Figure 2 shows that the fracture toughness of 828-PIP increases on increasing t . Networks obtained at high temperature are apparently looser and more flexible, as suggested by the progressive reduction of the glasstransition temperature (T ) and elastic modulus in the rubbery plateau (Fig ures 3 and 4). This condition could result from hardener volatilization (15, cure

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CURE TEMPERATURE (C) Figure 2. Fracture toughness, K , of the 828-PIP resin as a function of cure temperature. /c

16). Domenici et al. (17), however, did not observe any weight loss during the polymerization at 120 °C. Oxidation of piperidine, which can also occur, could diminish the cross-link efficiency of the hardener. It has frequently been observed (14, 18-20) that glasses obtained at low temperature have higher moduli, in spite of their lower cross-link density. The data in Figure 4 agree with such a finding. The effect has been attributed to the higher free-volume content in glasses cured at high temperature. The addition of a small amount of reactive rubbers such as A T B N s can result in a substantial increase in fracture resistance. Unfortunately, the dependence of K on cure temperature becomes more pronounced, as shown in Figure 5 for two ATBNs. The presence of the rubber does not bring about toughening if the cure temperature is too low. In the examined formulations, fracture resistance develops only when t is higher than 100 °C. The fact that K substantially increases when t positively influences the properties of the matrix (Figure 2) supports the simple idea that matrix and composite ductilities are strongly related. I c

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CURE TEMPERATURE (O Figure 3. Glass-transition temperature of the 828-PIP resin as a function of cure temperature.

However, the presence of maxima in Figure 5 indicates that other pa rameters, such as the morphology and the internal structure of the soft phase, are to be considered. In order to ensure the formation of finely sized rubbery domains, the liquid rubber has to be soluble in the epoxy prepolymer. In addition, phase separation has to take place before gelation, when the vis cosity is sufficiently high to prevent particle coalescence, but still low enough to favor nucleation and growth. Most phase separation takes place before gelation, although some localized change in composition occurs after gela tion (21). Figures 6 and 7 illustrate the morphological modifications brought about by curing A T B N blends at different temperatures. When f is too low, phase separation is suppressed and the material is transparent, even after postcuring at elevated temperature. Increasing t brings about an enlarge ment of particle size and a decrease of the number of particles. A compar ison of data in Figures 5 and 8 suggests a strong dependence of K on particle dimensions. The effect of particle size on the deformation be havior of C T B N - D G E B A systems was studied by Sultan and McGarry (22), who showed that an enlargement of particles was paralleled by an increase in fracture energy. The dimensions reported in that paper were smaller than those observed in A T B N formulations, probably a consequence of the diverse A N content of the rubbers. The apparent volume fraction of the segregated rubber was found to be higher than the volume of the added rubber. cure

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The relationship between the volume of added rubber and that of the separated phase is complex in toughened polymers. In high-impact polysty rene, for instance, 6% of polybutadiene can give rise to more than 50% of soft phase (23). A n increase in the volume of the soft phase over that of the added rubber has also been observed in some toughened epoxies (8, 9, 24). There are two reasons for this observation. First, the particles can incorporate some matrix material as a rigid filler. The presence of extraneous material in the rubber particles has been reported repeatedly (8, 9, 25). Second,



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CURE TEMPERATURE (C) Figure 5. Effect of cure temperature on the fracture toughness of 828-PIP-1300X21 and (·) 828-PIP-1300 X16 formulations.

(•)

some resin can molecularly dissolve into the particles as a result of phase equilibria (21,26). This resin can react with the terminal groups of the rubber and yield copolymers that increase its strength. Thus the deformation be havior of the particles is influenced by both the amount of dissolved resin and the extent of rubber-resin reactions. Sayre et al. (27) have shown that the fracture toughness of A T B N - and CTBN-modifled epoxies decreases by exposing the materials to radiation that excessively increases the cross-link density of the rubber. Figures 9 and 10 reveal the different internal structures of particles formed at different temperatures. Changing the A N content of the rubber influences the solubility param eter and, consequently, the final morphology. Rowe et al. (28) showed that the average particle size decreases on increasing the A N content of C T B N rubbers. Data in Figure 8 indicate that the effect is similar with A T B N s .



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Figure 6. SEM of the 828-PIP-1300 X 21 formulation cured at 15 ° C .

Figure 7. SEM of the 828-PIP-1300 X 21 formulation cured at 120 ° C .


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CURE TEMPERATURE (C) Figure 8. Average particle diameter as a function of cure temperature for (•) 828-PIP-1300X21 and (·) 828-PIP-1300 X16 formulations.

A T B N 1300 X 21 gives rise to particles about twice the size of particles orig inated by A T B N 1300 Χ 16. This result contradicts the previous statement about the relationship between K

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and particle diameter. That conclusion

cannot be generalized because it is not expected to be valid when a new parameter, the A N content, comes into play. The A N content apparently overrides the importance of particle size. The morphology of G T B N - 8 2 8 is similarly modified by the cure tem perature and the A N level. Again, phase separation does not take place if the cure temperature is too low.

Chain

Extension.

The mechanical properties of linear polymers

markedly depend on the average chain length (i.e., on molecular weight)



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Figure 9. SEM of the 828-PIP-1300X21 formulation cured at 60 °C.

Figure 10. SEM of the 828-PIP-1300X21 formulation cured at 160 °C.


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until the dependence almost vanishes above a certain level. In the case of cross-linked polymers the relevant parameter is the cross-link density, that is, the number of chemical junctions per unit volume or the average mo lecular weight of the segments among cross-links (M ). The formation of junctions has the important effect of limiting molecular mobility, particularly long-range mobility, to such an extent that many cross-linked glasses do not even yield. As a consequence, the fracture energy dramatically drops. The mirror-like fracture surfaces of highly cross-linked polymers indicate that plastic processes are almost totally inhibited. Downloaded by UCSF LIB CKM RSCS MGMT on November 20, 2014 | http://pubs.acs.org Publication Date: May 5, 1989 | doi: 10.1021/ba-1989-0222.ch004

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A n easy way to moderate brittleness is therefore to increase M , which can be done easily by using high-molecular-weight flexible hardeners such as polyaminesjor polysulfides. Unfortunately, the increase of fracture energy is often not very important and does not compensate for the deterioration of other properties. A wiser approach consists in spacing out the junctions with stiff molecules. In this way modulus and distortion temperature are not unduly reduced. The diminished cross-linked density favors postyield deformability and constitutes an effective energy sink. c

Bisphenol A (BPA) is a basic chemical for the production of D G E B A resins, whose molecular weight is determined by the ratio of BPA to epichlorohydrin (29). The general formula of D G E B A s can be written as shown in structure I, with η ranging from 0 to 20 in commercial grades. B P A is by far the most important modifier of D G E B A resins. High-molecular-weight resins can be obtained from low-molecular-weight products by reacting them with BPA in the presence of a basic catalyst (30). A n immediate molecular weight increase can be achieved by mixing low-molecular-weight epoxies with B P A and curing with polyamines, as in the case of some adhesives that are curable at room temperature. Two competitive processes have been shown to take place: chain extension and cross-linking (31, 32). The ratio between the rate constants of such processes, k /k , depends on the nature of the catalyst used. The variation with time of epoxy group concentration is schematically shown in Figure 11. Products of rather high molecular weight ( M W >4000) can form before cross-linking starts. 1

2

BPA is also used in rubber-modified formulations (8, 9, 11, 33-38) be cause it magnifies the effect of the rubber. Riew et al. (35) observed that the fracture energy of a 5% C T B N - D G E B A system increased by a factor of

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REACTION TIME Figure 11. Extent of epoxide reaction vs. time. (Reproduced with permission from ref 32. Copyright 1973 Wiley.) 4 when 27% B P A was added. Moreover, good results have been obtained with the use of nonreactive rubbers, either liquid or solid. For this reason BPA has been used frequently in toughened systems. Levita et al. (8) spe cifically investigated the role of BPA. Figure 12 shows how the fracture toughness of the low-molecular-weight 828 resin increases on adding BPA. The effect is similar to that brought about by the cure temperature in Figure 2. The important difference is that X is not adversely affected by the addition of BPA (Figure 13). The effect on tensile behavior is even more important. The unmodified D G E B A is sub stantially brittle because of a lack of postyield deformability (Figure 14). O n the contrary, at 30% BPA yielding is followed by a consistent softening and by the formation and propagation of a neck. The deformation at break, —0.08, by far underestimates the real de formability of the material, because it was obtained simply by specimen elongation to the original length ratio. A more realistic estimate was obtained by observing the behavior of tiny air bubbles dispersed within the test pieces. The unmodified D G E B A broke soon after the formation of a shear band. Bubbles located outside the band retained the original shape, whereas those lying within the band developed cracks at the equator but did not show any appreciable departure from the spherical shape (Figure 15). O n the contrary, bubbles located in the sheared region of the 30% BPA blend became ellipsoidal in shape, with axial ratios greater than 2 (Figure g


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16). The local deformation at the bubble equator was estimated to be higher than 40%. Such a high network deformability easily gives rise to crack blunt ing in rubber-modified formulations. A further indication that the increase of K in Figure 12 entirely relies on postyield processes is the fact that yield stress is independent of BPA content for both neat and rubber-modified resins (Figure 17). I c

Rubber-modified formulations greatly benefit from the greater ductility produced by BPA. Fracture toughness steadily increases when BPA is added (Figure 18). The higher toughness of ternary B P A - D G E B A - r u b b e r blends over that of binary D G E B A - r u b b e r blends was supposed to arise substan tially from modifications of soft-phase geometric parameters. Riew et al. (35) proposed that the effect of B P A was production of a


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Figure 13. Shear moduli (5 s) vs. temperature for neat and BPA-modified (25%) 828-PIP resin.

bimodal particle-size distribution. Small particles (

Matrix Ductility and Toughening of Epoxy Resins - American Chemical

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