Fatigue Behavior of Acrylic Interpenetrating Polymer Networks

Chapter

13

Fatigue Behavior of Acrylic Interpenetrating Polymer Networks Tak Hur , John A. Manson , and Richard W. Hertzberg Downloaded by MICHIGAN STATE UNIV on February 18, 2015 | http://pubs.acs.org Publication Date: April 18, 1988 | doi: 10.1021/bk-1988-0367.ch013

1

1,2

3

1

Polymer Science and Engineering Program, Lehigh University, Bethlehem, PA 18015 Department of Chemistry, Lehigh University, Bethlehem, PA 18015 Department of Materials Science and Engineering, Lehigh University, Bethlehem, PA 18015 2

3

Interpenetrating polymer networks (IPNs) continue to excite both fundamental and technological interest. By combining elastomeric and brittle glassy phases it is often possible to obtain improved properties over a range of temperature and frequency. However, relatively little attention has been given to fatigue in IPNs, and to energy absorption in polyurethane rubber/poly(methyl methacrylate) (PU/PMMA) systems. In this paper it is shown that simultaneous interpenetrating networks based on an energy-absorbing PU are transparent, with a single broad glass-to-rubber transition indicating a microheterogeneous morphology. Fatigue resistance increases with the [PU] up to 50%, while energy absorption determined from dynamic properties and pendulum impact tests varies directly with the [PU]. The micromechanism of failure involves the generation of discontinuous growth bands associated with shear yielding rather than crazing. Interpenetrating polymer networks, both sequential (IPN) and simultaneous (SIN), continue to excite both fundamental and technological i n t e r e s t (1-21). For example, by combining elastomeric and b r i t t l e glassy phases, i t i s possible to obtain improved properties such as fatigue resistance, as w e l l as t e n s i l e , adhesive, and impact strengths. Although much research has emphasized crosslinked polystyrene as the glassy component, attention has also been given to crosslinked PMMA (4,5,8-13). Elastomeric phases have included a v a r i e t y of polymers such as rubbery a c r y l i c s , polyesters, and polyurethanes (4,6,10). Among these l a t t e r types, polyurethanes have become important as energy-absorbing materials i n orthopedic applications such as a t h l e t i c footwear, and sound and v i b r a t i o n damping (22-25). In such polyurethanes i t i s desirable to modify properties over a wide range of hardness and r e s i l i e n c e ; the polyurethanes themselves are of inherent i n t e r e s t as toughening phases for b r i t t l e 0097-6156/88/0367-0169$06.00/0 © 1988 American Chemical Society In Cross-Linked Polymers; Dickie, R., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 1988.

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170

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matrices. At the same time, although a fundamental understanding of fatigue behavior i s important to both s c i e n t i f i c inquiry and engineering a p p l i c a t i o n s (26), l i t t l e has been published on fatigue i n interpenetrating networks (21). For these reasons, new SINs based on a polyurethane (PU/PMMA) system were synthesized, and characterized with respect to properties including dynamic mechanical response, energy absorption, and fatigue crack propagation (FCP) behavior as a function of composition. In t h i s paper, u s e f u l combinations of these properties, and correlations between energy absorption and fatigue behavior, are demonstrated; the micromechanism of f a i l u r e i s also shown to involve discontinuous growth bands associated with shear y i e l d i n g rather than crazing. Experimental The polyurethane formulation involved a proprietary crosslinkable system based on poly(propylene g l y c o l ) and methylene diisocyanate (NC0/OH r a t i o = 1.0). For studies of v i s c o e l a s t i c , energy absorption, and fatigue behavior, the weight fractions of PMMA were 0, 0.25, 0.50, 0.75, and 1.0; f o r studies of t e n s i l e and tear strength, the r a t i o s were 0, 0.10, 0.20, 0.25, 0.30, and 0.40. Reactants were mixed at room temperature, degassed, poured into a mold, and cured at 60°C for 48 hr. Dynamic mechanical spectra were obtained at 110 Hz using an Autovibron u n i t , model DDV-IIIC; values of Τ were obtained from the temperatures corresponding to the maxima in^the loss modulus (E") peaks. The shear modulus G was also measured at room temperature using a Gehman t o r s i o n a l t e s t e r . Fatigue crack propagation tests were conducted as before (21) using a servohydraulic test machine under ambient atmosphere, compact-tension specimens, a sinusoidal frequency of 10 Hz, and an R value of 0.1 (R - min/max load). The crack length a was measured using a t r a v e l i n g microscope, and values of log da/dN, the crack growth rate per cycle, were plotted against log ΔΚ, where ΔΚ i s the range of the stress i n t e n s i t y factor K, a measure of the d r i v i n g force for crack extension. This p l o t r e f l e c t s the Paris equation (26,27): da/dN - Α ΔΚ exp(n), where A and η are constants. ΔΚ was calculated as ΔΚ = ΥΔσ/a, where Y i s a known geometrical factor and Δσ i s the range i n applied s t r e s s . The % energy absorption % e(abs) =

100 e(abs) [e(abs) + e(recovered)]

(1)

was determined i n several ways: (1) using a Zwick pendulum tester to obtain the % rebound r e s i l i e n c e R (% energy absorbed = 100-R); (2) using a computerized Instron tester (model 1332) to obtain hysteresis loops under c y c l i c compression at 10 Hz; and (3) using dynamic mechanical data to calculate the r a t i o of energy absorbed to energy input (per quarter c y c l e ) , given by π tan 6/(π tan 6 +2) (28, p. 606). Tensile and tear strengths were determined using ASTM standards D412 and D1004, respectively, at a crosshead speed of 0.42 mm/s (1 in/min) ; values reported are the average for 3 specimens. The e l a s t i c and i n e l a s t i c ( p l a s t i c ) components of the t o t a l elongation

In Cross-Linked Polymers; Dickie, R., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 1988.

171

Fatigue Behavior of Acrylic Networks

13. HUR ET AL.

at break were determined by measuring the r e t r a c t i o n of the specimens a f t e r rupture; an extensometer was used to ensure accuracy of the strains involved.


Results and Discussion Molecular and V i s c o e l a s t i c Characterization. The SIN samples were r e l a t i v e l y m i s c i b l e , as evidenced by transparency to the eye. Densities were 0.6-0.7% higher than predicted by a rule of mixtures; such d e n s i f i c a t i o n has been noted before (5). However, although the dynamic mechanical data shown i n Figure 1 show a single t r a n s i t i o n , the t r a n s i t i o n s of the SINs are broader than f o r the homopolymers. As shown i n Table 1 and Figure 1, the slopes of the complex modulus (E*) plots are lower f o r a l l the SINs than f o r the controls; the slope f o r the PU i s also lower than f o r the PMMA. The breadths of the tan δ and E curves (not shown) follow the same trend; as expected, the slopes and breadths of the tan δ peaks are inversely correlated. High transparency and a single t r a n s i t i o n have been noted f o r PU/polyacrylate SINs (19); i n some other PU/PMMA SINs, two loss peaks s h i f t e d s i g n i f i c a n t l y towards each other were noted (5). I n t e r e s t i n g l y , the widths of the tan δ peaks as a function of [PU] exhibit a maximum at a PU/PMMA r a t i o of 50/50; the maximum value i s very close to that reported by others (10), although at lower PU concentrations, values of tan δ for the IPNs of t h i s study are higher. tf

Table 1. Properties of PU/PMMA SINs PU/ PMMA

Τ g (°C)

tana (25°C)

a

E x 10" (MPa)

3

dlog

0.07 0.09 0.11 0.27 0.93

3.0 1.1 0.34 0.020 0.0028

e(abs)

a

1

121 74 35 3 -24

6

E* /dT

Cc" ) 0/100 25/75 50/50 75/25 100/0

ΔΤ,

0.055 0.029 0.026 0.026 0.048

°C 37 63 81 76 56

%

b

9.9 12.4 14.7 29.8 59.4

%

C

54 59 63 71 85

12 14 17 21 59

Ε and E* represent the t o r s i o n a l and complex modulus, respec t i v e l y , at 25°C. From dynamic spectra. ^From hysteresis loops. From pendulum impact t e s t . Breadth of tan δ peak at tan δ ( π ΐ 3 χ ) / 2 . Thus the SINs presumably have a f i n e - s c a l e , microheterogeneous character, while the PU i t s e l f has a r e l a t i v e l y broad d i s t r i b u t i o n of relaxation times r e l a t i v e to that of the PMMA. This o v e r a l l



172


Figure 1. (a) Complex modulus E* and (b) tanô f o r PU/PMMA SINs as a function of temperature.



13. HURETAL.


173

p a r t i a l m i s c i b i l i t y of a system that may not be inherently f u l l y miscible on a segmental scale may w e l l r e f l e c t the favorable effects of Η-bonding between the two constituents, s i m i l a r s o l u b i l i t y parameters (4), and of crosslinking-induced network constraints on gross phase separation. (The detailed morphology and the nature of the interactions between the PU and PMMA phases are under i n v e s t i g a t i o n by scanning and transmission electron microscopy and Fourier-transform infrared spectroscopy, respectively. So f a r , the size of the PU component appears to be i n the range of ^ 10-20 nm, a range consistent with the nearly-single-phase behaviors discussed above.) As expected, the Τ varies regularly with composition (Table 1, Figure 2); values f e l l §etween those predicted by the Fox and Pochan equations (1/T = w(A)/T (A) + w(B)/T (Β), and In Τ = w(A) In Τ (A) + w(B) In Τ (B?, respectively, where % i s the wt t i o n and A % n d Β refer to two homopolymers involved). The modulus at 25°C also varies with composition, and i s compared i n Figure 3 with the predictions of several models. These include models of Paul (29) ( s e r i e s - p a r a l l e l type), Davies (30) (especially suited to dual-phase c o n t i n u i t y ) , Budiansky (31) (phase interactions considered), and Hourston and Zia (14) (a modified Davies equation). The Davies and modified Davies equations are given by: G(exp a) = φ(A) G (A) (exp a) + φ(Β) G(B) (exp a ) , where G i s shear modulus, φ i s the volume f r a c t i o n of components A and B, and a has the values of 1/5 (30) or 1/10 (14), respectively. As shown i n Figure 3, the Davies and modified Davies equations give the best f i t to the data a v a i l a b l e , the former being best at high, and the l a t t e r at low, PMMA concentrations. Thus the behavior i s consistent with the existence of a s i g n i f i c a n t degree of fine-scale dual phase continuity i n t h i s system. Tensile and Tearing Behavior. As shown i n Figure 4, the t e n s i l e and tear strengths increased with increasing PMMA content; measurements of tear strength were not feasible at PMMA contents > 30%. In any case, the incorporation of PMMA at even r e l a t i v e l y low l e v e l s greatly improves the rather low strengths of the unmodified PU. As shown i n Figure 5, the o v e r a l l elongation at break exhibited a maximum at a concentration of 20% PMMA. At higher concentrations of PMMA, the r e l a t i v e contribution of p l a s t i c ( i r r e v e r s i b l e ) deformation increased with increasing PMMA content, while that of e l a s t i c (reversible) deformation decreased. Energy Absorption. Energy absorption [e(abs)] data are presented i n Table 1. C l e a r l y the energy absorption values increase d i r e c t l y with the PU content, and the values for the PU are very high indeed at room temperature. One would not expect absolute values of e(abs) obtained from such widely d i f f e r e n t tests involving quite variable states of stress and both l i n e a r and nonlinear v i s c o e l a s t i c behavior to agree. Even so, the values r e l a t i v e to that of the PU agree f a i r l y w e l l (within ± 7%) up to a PU/PMMA r a t i o of 50/50, and good correlations between tests are obtainable between r e s u l t s obtained over the whole range of composition. I t i s e s p e c i a l l y i n t e r e s t i n g that the value of e(abs) obtained from measurements of tan δ ( i n the range of l i n e a r v i s c o e l a s t i c i t y ) agree quite w e l l with values




174

WT. FRACTION, PU

Figure 2. Glass transition temperature of PU/PMMA SINs as a function of composition: —, Pochan equation; , Fox equation.

O

0.5

VOL. FRACTION,

1.0 P M M A

Figure 3. Shear modulus as a function of composition for PU/PMMA SINs. Models of (a) Paul (29); (b) Budiansky (31); (c) Davies (30); and (d) Hourston and Zia (14).



HUR ET AL.


0

10 20 30

175

40?%}

WT.FRACTION.PMMA

Figure 4. Tensile and tear strength function of composition. 200

of PU/PMMA SINs as a

r

WT. FRACTION, PMMA

Figure 5. Elongation behavior of PUR/PMMA SINs. Total elongation at break (·) i s broken down into both i r r e v e r s i b l e (•) and r e v e r s i b l e (•) contributions.



176

o b t a i n e d from the pendulum impact t e s t s . T h u s , as shown i n F i g u r e 6, the v a l u e s o f e ( a b s ) o b t a i n e d from h y s t e r e s i s loops p a r a l l e l t h o s e o b t a i n e d by a v e r a g i n g the r e s u l t s f o r the pendulum and dynamic m e c h a n i c a l t e s t s . In each case the dependence o f e ( a b s ) on c o m p o s i t i o n f o l l o w s a l o w e r - b o u n d r u l e o f m i x t u r e s ( E q u a t i o n 2) rather w e l l . ςτΝΠ


Kaos,

biJN;

=

c(

a

b

s

,

PU) χ e ( b s , a

[φ( ΜΑ) e ( a b s , PU) + ΡΜ

PMMA)

Φ (PU)

m

e( bs, a

PMMA)]

Κ

Δ

)

Fatigue Crack P r o p a g a t i o n . FCP d a t a were o b t a i n a b l e o n l y w i t h PU/PMMA r a t i o s o f 0 / 1 0 0 , 2 5 / 7 5 , and 5 0 / 5 0 , a t l e a s t w i t h the specimen and t e s t c o n d i t i o n s u s e d ; a t h i g h e r c o n c e n t r a t i o n s i t was not p o s s i b l e to o b t a i n s t a b l e c r a c k g r o w t h . T h i s i s not u n e x p e c t e d , f o r the FCP b e h a v i o r o f r u b b e r - t o u g h e n e d p l a s t i c s i s b e l i e v e d t o be dominated by a dynamic b a l a n c e between the a b i l i t y o f the r u b b e r to g e n e r a t e e n e r g y - d i s s i p a t i n g p r o c e s s e s a t the c r a c k t i p and s o f t e n i n g of the b u l k polymer ( 2 6 ) . In t h i s c a s e , the s i g n i f i c a n t d e c r e a s e i n modulus as the PU c o n t e n t increases e v i d e n t l y overcame the b e n e f i c i a l e f f e c t s o f energy d i s s i p a t i o n . Specimens specifically d e s i g n e d f o r e l a s t o m e r s w i l l have to be used to e v a l u a t e the h i g h l y e l a s t o m e r i c systems. W i t h r e s p e c t to the t h r e e specimens t h a t c o u l d be t e s t e d , the FCP b e h a v i o r f o l l o w e d the P a r i s law ( F i g u r e 7) o v e r the range o f Δ Κ s t u d i e d , and the FCP r a t e s a t c o n s t a n t Δ κ v a r i e d i n v e r s e l y w i t h the PU c o n t e n t , w h i l e the maximum v a l u e o f Δ κ a t t a i n a b l e (K* ) v a r i e d d i r e c t l y (Figure 8). Thus a t a c o n s t a n t v a l u e o f Δ κ , the FCP r a t e of PMMA was r e d u c e d by more than an o r d e r o f magnitude f o r the 50/50 SIN (at Δ κ - 0.6 MPa Λη) and o v e r h a l f an o r d e r o f magnitude as f i n a l f a i l u r e was a p p r o a c h e d . A t the same t i m e , the d r i v i n g f o r c e for c r a c k e x t e n s i o n Δ κ * (the v a l u e o^ Δ κ c o r r e s p o n d i n g to an a r b i t r a r y c r a c k s p e e d , i n t h i s case 10 mm/cycle) was i n c r e a s e d from 0.7 to 0.9 as the PU c o n t e n t was i n c r e a s e d to 50%. A l t h o u g h the v a l u e s o f ε ( a b s ) do n o t c o r r e s p o n d to v a l u e s o f the s t r a i n energy r e l e a s e r a t e G (where K = E G ) , the f a t i g u e p a r a m e t e r s Δ κ * and d a / d N at constant Δ κ c o r r e l a t e w e l l w i t h B ( a b s ) (Figure 9). Also, d i v i s i o n o f ( Δ κ * ) by Ε g i v e s a f i g u r e o f m e r i t a n a l o g o u s to the s t r a i n energy r e l e a s e r a t e t h a t i n c r e a s e s from 0.17 to 2.36 as the PU c o n t e n t i n c r e a s e s , i n q u a l i t a t i v e agreement w i t h the i n c r e a s e i n e(abs) . With r e s p e c t t o micromechanisms o f f a i l u r e , a t low v a l u e s o f ΔΚ, d i s c o n t i n u o u s growth bands whose s p a c i n g s c o r r e s p o n d to many c y c l e s o f l o a d i n g were o b s e r v e d ( 2 6 , 3 2 ) . F i g u r e 10 shows the e f f e c t of c o m p o s i t i o n on r , the s p a c i n g o f the b a n d s , the y i e l d s t r e s s ο ( e s t i m a t e d from the Dugdale r e l a t i o n s h i p , r = ^ K / 8 σ ) , and tlie number o f c y c l e s p e r b a n d . F i g u r e 11 shows t h a t ^rSie tre^nd o f band s i z e as a f u n c t i o n o f Δ κ resembles t h a t o f a t y p i c a l ABS ( a c r y l o n i t r i l e - b u t a d i e n e - s t y r e n e t e r p o l y m e r ) , and t h a t the s l o p e o f the c u r v e conforms to the Dugdale p r e d i c t i o n . W h i l e such bands a r e often a s s o c i a t e d w i t h c r a z i n g , shear y i e l d i n g i s expected i n a c r o s s l i n k e d system. Indeed the e s t i m a t e d v a l u e o f σ f o r the PMMA i s 77 MPa, a v a l u e more t y p i c a l o f y i e l d , r a t h e r tnan c r a z i n g , stresses. At h i g h v a l u e s o f Δ κ , diamond and f e a t h e r m a r k i n g s a r e s e e n ; such m a r k i n g s have been r e l a t e d t o the o c c u r r e n c e o f m u l t i p h a s e f r a c t u r e and c r a c k f o r k i n g i n a h e t e r o g e n e o u s polymer 2

2

2

2



13. HURET AL.


0

Lu_0

•

•

» • 0.5

•

•

•

» 1.0

WT. FRACTION, PU Figure 6. Energy absorption i n PU/PMMA SINs as a function of composition: · , from hysteresis loops; Ο , average of r e s u l t s from pendulum and dynamic mechanical t e s t s ; -, lower-bound t h e o r e t i c a l curve.

ι

.4

ι

ι

.6

.8

ι

1.0 Δ K, MPa/m

Figure 7. Fatigue crack growth rates as a function of composition f o r PU/PMMA SINs. Ratios are f o r PU/PMMA contents.


177



178

O

0.2

0.4

0.6

WT. FRACTION, P U

Figure 8. E f f e c t s of composition of PU/PMMA SINs on (a) FCP rate (at ΔΚ - 0.6 MPa) and (b) K' (at da/dN - 3X 10~ mm/ cycle). J

ENERGY ABSORPTION. %

Figure 9. Effects of % e n e r g ^ a b s o r p t i o n e(abs) on FCP n/cycle) and (b) da/dN (ΔΚ 10 parameters (a) ΔΚ* (at da/dN 0.6 MPa).



13. HURETAL.


O

WT.

0.2 0.4 FRACTION,PU

179

0.6

Figure 10. E f f e c t s of composition of PUR/PMMA SINs on: (a) r , the size of DGB spacings; (b) σ , y i e l d stress; and (c) log N, the number of cycles per band. y

0.5

0.2

0.8

1.0

2.0

ΔΚ, MPa Figure 11. Comparison of r , the length of discontinuous growth bands as a function of ΔΚ, f o r PU/PMMA SINs with values f o r several other polymers (26).



180

subjected to more applied energy than i s required to propagate a stable crack (33). Acknowledgment The authors wish to acknowledge p a r t i a l support from the Ben Franklin Program, NET/ATC (State of Pennsylvania) and from the National Science Foundation, Grant No. 8106489. Assistance with polymer formulation was provided by Dr. F r i t z H o s t e t t l e r , Polymer Dynamics, Inc., Bethlehem, PA.


Literature Cited 1. 2. 3.

4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

Manson, J. Α.; Sperling, L. H. Polymer Blends and Composites; Plenum: New York, 1976; Chap. 3. Sperling, L. H. Interpenetrating Polymer Networks and Related Materials; Plenum: New York; 1981. Sperling, L. H., ACS In Multicomponent Polymer Materials; Paul, D. R.; Sperling, L. H., Eds.; ACS Advances in Chemistry Series No. 211, American Chemical Society: Washington, D.C., 1986; p. 21. Kim, S. C.; Klempner, D.; Frisch, K. C.; Radigan, W.; Frisch, H. L. Macromolecules 1976, 9, 258. Kim, S. C.; Klempner, D.; Frisch, K. C.; Frisch, H. L. Macromolecules 1976, 9, 263, 1187. Allen, G.; Bowden, M. J.; Todd, S. M.; Blundell, D. J.; Jeffs, G. M.; Davies, W. E.A. Polymer 1974, 15, 28. Kim, S. C.; Klempner, D.; Frisch, K. C.; Frisch, H. L. Macromolecules 1977, 10, 1187. Djomo, H.; Morin, Α.; Damyanidu, M.; Meyer, G. C. Polymer 1983, 24, 65. Djomo, H.; Widmaier, J. M.; Meyer, G. C. Polymer 1983 24, 1415. Morin, Α.; Djomo, H.; Meyer, G. C. Polym. Eng. Sci. 1983, 23, 394. Hermant, S.; Damyanidu, M.; Meyer, G. C. Polymer 1983, 24, 1419. Hermant, J.; Meyer, G. C. Eur. Polym. J., 1984, 20, 85. Lee, D. S.; Kim, S. C. Macromolecules, 1984, 17, 268, 2193. Hourston, D. J.; Zia, Y. J. Appl. Polym. Sci. 1983, 28, 3745. Hourston, D. J.; Zia, Y. J. Appl. Polym. Sci. 1984, 29, 2951. Kim, S. C.; Klempner, D.; Frisch, K. C.; Frisch, H. L.; Ghiradella, H. Polym. Eng. Sci. 1975, 15, 339. Frisch, H. L.; Frisch, K. C.; Klempner, D. Polym. Eng. Sci. 1974, 14, 646. Meyer, G. C.; Mehrenberger, P. Y. Eur. Polym. J. 1977, 13, 383. Frisch, K. C.; Klempner, D.; Migdal, S.; Frisch, H. L.; Ghiradella, H. Polym. Eng. Sci. 1974, 14, 76. Devia, N.; Sperling, L. H.; Manson, J. Α.; Conde, A. Polym. Eng. Sci. 1979, 19, 878.



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21. Qureshi, S.; Manson, J. Α.; Sperling, L. H.; Murphy, C. J. In Polymer Applications of Renewable Resource Materials; Carraher, C. E.; Sperling, L. H. Eds.; Plenum: New York, 1983. 22. Light, L. H.; McLellan, G. E.; Klenerman, L. J. Biomech. 1980, 13, 477. 23. Lin, J. S. Ph.D. dissertation, Lehigh University, 1986. 24. Lin, J. S.; Manson, J. A. In Proc. 45th Ann. Tech. Conf., Society of Plastics Engineers: Los Angeles, May, 1987, p. 478. 25. Newton, C. H.; Connelly, G. M.; Lewis, J. E.; Manson, J. A. In Proc. 39th ACEMB Conference, Baltimore, September, 1986, p. 38. 26. Hertzberg, R. W.; Manson, J. A. Fatigue in Engineering Plastics; Academic: New York, 1980. 27. Paris, P.C. In Proc. 10th Sagamore Army Res. Conf.; Syracuse University Press: Syracuse, 1964; p. 107. 28. Ferry, W. D. Viscoelastic Properties of Polymers, 2nd ed., Wiley: New York, 1970. 29. Paul, B. Trans. AIME 1980, 218, 36. 30. Davies, W. E. A. J. Phys. (D) 1971, 4, 318. 31. Budiansky, B. J. Mech. Phys. Solids 1965, 13, 223. 32. Skibo, M. D.; Manson, J. Α.; Webler, S. M.; Hertzberg, R. W.; Collins, E. A. In Durability of Macromolecular Materials; Eby, R. K.; Ed.; ACS Advances in Chemistry Series No. 95, American Chemical Society: Washington, D.C., 1979; p. 311. 33. Andrews, Ε. H. Fracture of Polymers; American Elsevier: New York; 1968. RECEIVED February 8, 1988


Fatigue Behavior of Acrylic Interpenetrating Polymer Networks

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