ClEA Gradient Polymers - American

23 Stress-Strain Behavior of P M M A / C l E A Gradient Polymers C. F. JASSO, S. D. HONG, and M. SHEN Downloaded by PEPPERDINE UNIV on August 23, 2017 | http://pubs.acs.org Publication Date: June 1, 1979 | doi: 10.1021/ba-1979-0176.ch023

1

2

Department of Chemical Engineering, University of California, Berkeley, C A 94720

Gradient polymers were prepared by diffusing the 2-chloro-ethyl acrylate (ClEA) monomer into crosslinked poly(methyl methacrylate). The resulting profile of the diffusion gradient is then fixed by polymerizing the monomer in situ by photochemical initiation. If the diffusion of the monomer is permitted to proceed to swelling equilibrium, then interpenetrating networks (IPN) are formed. The stress-strain curves of the gradient polymers are plasticlike in that they exhibit yield points and enhanced fracture strains. Those of interpenetrating networks at comparable chemical composition, on the other hand, show essentially rubbery behavior. Data may be interpreted in terms of the stress transfer mechanism or surface stabilization mechanism.

radient polymers are multicomponent polymers whose structure or composition is not homogeneous throughout the material, but varies as a function of position (I ). In other words, these polymers have gradients in their structures or compositions. In a previous paper (2), we have shown that it is possible to produce such materials by diffusing a guest monomer into a host polymer for a period of time sufficient for the establishment of a diffusion gradient profile. This profile is then "fixed" by polymerization of the monomer i n situ. The mechanical behavior of such gradient polymers was found to be quite different from the interCurrent address: Av. Mexico 2286-1, Guadalajara, Jalisco, Mexico. * Current address: Jet Propulsion Laboratory, Pasadena, CA 91103. 1

0-8412-0457-8/79/33-176-443$05.00/0 © 1979 American Chemical Society Cooper and Estes; Multiphase Polymers Advances in Chemistry; American Chemical Society: Washington, DC, 1979.

Downloaded by PEPPERDINE UNIV on August 23, 2017 | http://pubs.acs.org Publication Date: June 1, 1979 | doi: 10.1021/ba-1979-0176.ch023

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penetrating networks ( IPNs ) of comparable composition. O f course the latter material allows the monomer to reach a swelling equilibrium i n the host polymer and therefore has no gradient structure. In our preliminary report (2), w e have chosen poly (methyl metha crylate ) or P M M A as a host polymer and methyl acrylate as the guest monomer. They were both crosslinked by a divinyl acrylic monomer. However, because of the similarity i n the constitutions of these two components, it was not possible to establish the gradient profile through chemical analysis. I n this work, we have selected a halogenated acrylic monomer as the second component to be diffused into P M M A . B y analyz ing the halogen content, it was possible to determine the profiles of the gradient polymers. Stress-strain measurements of the samples were then carried out on these unique materials.

τ

ι

ι

ι

ι

t

(MlN)

r

Figure I. Uptake of 2-chloroethyl acrylate (percent by weight) by crosslinked poly(methyl methacrylate) at 60°C as a function of time of immersion

Cooper and Estes; Multiphase Polymers Advances in Chemistry; American Chemical Society: Washington, DC, 1979.


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Stress-Strain Behavior

Figure 2. Concentration profiles of poly(2-chloroethyl acrylate) in poly(methyl methacrylate) along the thickness (L ) dimension of the samples. (GRAD) Gradient polymer, (IPN) interpenetrating networks. 0

Experimental Monomers of methyl methacrylate and 2-chloroethyl acrylate (C1EA) were purchased from the Polysciences, Inc. They were distilled and subsequently mixed with 1.3% by weight of the crosslinking agent ( ethylene glycol dimethacrylate obtained from the J. T . Baker C o . ) and 1.9% by weight of the photosensitizer ( benzoin isobutyl ether supplied by the Stauffer Chemical Co. ). P M M A was first prepared by photopolymerization i n front of U V light for 48 hr. Samples were stored i n a vacuum oven at 60° C until a constant weight was achieved to remove remaining monomer. The cross-linked P M M A samples were then immersed i n the bath of C1EA monomer for various periods of time by monitoring the weight uptake (Figure 1). When the desired amount of monomer has been imbibed into the host polymer, the latter was removed, surface dried, and then immediately polymerized by U V radiation. IPNs were prepared in a similar manner except that after immersion in the monomer bath the sample was stored i n a sealed polymerization cell at 60°C for several days prior to polymerization. The profiles of these gradient polymers were determined by machining off the samples layer by layer. Shavings from each layer were then analyzed for chlorine content by combustion methods ( 3 ) . The results


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of these analyses are shown i n Figure 2, recalculated i n terms of mole percent of C1EA i n P M M A . A n Instron Universal Testing Machine M o d e l T M - S M was used to carry out the stress-strain measurements. A n environmental chamber, equipped with a Missimers P I T C temperature controller, was used i n providing constant temperatures to ± 0.5°C. Samples were of the approxi mate dimensions of 5.0 X 0.5 χ 0.1 c m . T o prevent slippage from the Instron clamps, special aluminum tabs were glued to the ends of the samples. 3


Results and

Discussion

Stress-strain curves for a series of gradient polymers of 2-chloroethyl acrylate i n poly (methyl methacrylate), which are designated as P M M A / Grad PC1EA, are shown i n Figure 3. Quantities i n the parentheses indi cate the mole percent of the acrylate. Also included are the stress-strain curves of pure P M M A and an interpenetrating network ( P M M A / I P N P C 1 E A ) . A l l of these measurements were carried out at 80° C and at strain rates of 2-3%/sec. First we note that the stress-strain curve of P M M A has a high initial slope (high elastic modulus) and fractures at Τ "

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1

r~

PMMA

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PC1EA

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(12.5)

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PCI Ε A

(29.7)

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JL 80

90

100

JL

J L

110

120

130

Figure 3. Stress-strain curves of three gradient polymers and one interpene trating network of poly(methyl methacrylate) with 2-chloroethyl acrylate at comparable strain rates of 2-3%/sec and same temperature of 80° C. The numerals in parentheses indicate concentrations (mole percent) of chloroethyl acrylate in poly(methyl methacrylate).


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20

30

1

1

1

1

Γ

40

50

60

70

80

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100

c (%) Figure 4.

Same as Figure 3 except that the strain rates are between 18-25%/ sec and that the PMMA data are not included

about 6 % strain. Upon introduction of a 5 % PC1EA gradient, there is a slight decrease of the initial slope but a dramatic increase i n fracture strain ( ~ 65% ). I n addition, a pronounced yield region is observed around 5 % strain. Further increases in PC1EA contents to 12.5 and 29.7% show a continuing decrease i n the elastic moduli but even higher fracture strains were attained (85 and 130%, respectively ). Both of the latter gradient polymers, however, still exhibit the yielding behavior and are plasticlike in mechanical properties. In contrast, the I P N at 29.8% PC1EA content is rubbery and fractures at slightly greater than 100% strain. Apparently the presence of the gradient structure enables the samples to retain the plasticlike properties without sacrificing the enhanced ability to withstand deformation. The effects of increasing strain rates on P M M A / G r a d PC1EA and P M M A / I P N PC1EA are shown i n Figure 4. Increased strain rates are seen to increase the yield stresses as well as stress levels i n the plateau regions of the gradient polymers but to decrease the fracture strains. F o r

American Chemical Society Library 1155 16th St. N. W. Cooper and Estes; Multiphase Polymers Washington, D. C. 20036

Advances in Chemistry; American Chemical Society: Washington, DC, 1979.


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the interpenetrating networks of comparable composition, there is no observable yielding, but the stress levels are increased by higher strain rates. However, there also appear to be decreases in the fracture strains. As expected, the effects of decreasing strain rates are just the opposite (Figures 5 and 6 ) . Again IPNs behave essentially as rubbers, while the gradient polymers exhibit plasticlike properties. The effect of temperature on the stress-strain properties of P M M A , its gradient polymers with various compositions, and an I P N are shown in Figures 7, 8, and 9. These experiments were performed at various strain rates at 60°C. Comparison with the 80°C data shows that the main effects of temperature are to increase the stress levels i n the plateau regions at lower temperatures without significant differences i n other aspects. In our previous paper ( 2 ) , we proposed a possible mechanism to interpret the stress-strain behavior of gradient polymers. W e perceived the gradient polymer as consisting of infinite number of layers of varying compositions. Upon deformation, the macroscopic strain is the same for the entire sample. Because of the fact that the moduli of the various

Figure 5.

Same as Figure 3 except that the strain rates are between 3-6 X 10'Wo/sec



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Figure 6.

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Same as Figure 3 except that the strain rates are between 3-4 X 10~ %/sec

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130

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PMMA/GRAD PCIEA (29.7)

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1—»—η

Same as Figure 6 except that the temperature is 60°C

30

2

1

MMA/GRAD PCIEA (12.5)

1

P M M A / I P N PCIEA (29.8) € = 2.6 x I0" % S E C '

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Figure 7.

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JASSO E T AL.



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MULTIPHASE

"Τ 1 1 / PMMA ' c = 1.9% SEC"

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1

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1

POLYMERS

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1

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PMMA/GRAD PCIEA (12.5)


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1

PMMA/GRAD PCIEA (29.7) é = 2% SEC"

PMMA/GRAD PCIEA (29.8) c = 2.7% SEC"

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20

30

40

L 50

J 60

I 70

I 80

L 90

100

1

110

120

€ (%)

Figure 9.

Same as Figure 3 except that the temperature is 60°C

layers are different owing to their differences i n composition, those layers with higher moduli must sustain greater stresses (for the same strain). According to Eyring's stress-biased activated rate theory of yielding (4,5), the barrier height for a molecular segment to jump i n the forward direction is reduced by the applied stress. As a consequence, the higher modulus layers i n the gradient polymer should show greater tendency to yield because of the greater stress biases they have than those with lower moduli. This mechanism w i l l be inoperative for interpenetrating networks because of their uniform composition. A n alternative mechanism for the observed high fracture strain may be the reduction in surface imperfections for the gradient polymers. T h e surfaces of these materials must be more résistent to fracture because of the relatively high loading of rubbery phases. Thus, under stress they are likely to craze or crack, which would have initiated fracture for the sample as a whole. T h e validity of either mechanism, however, must await verification. In conclusion, we find that gradient polymers produced b y diffusion polymerization generally show enhanced fracture strain, while retaining their plasticlike properties. This behavior appears to be a consequence


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of the gradient structure, rather than being attributable to the compositions alone. Interpenetrating networks with comparable composition do not possess the unique mechanical behavior of the gradient polymers. Acknowledgment This work was supported b y the Office of Naval Research.


Literature Cited

RECEIVED

1. 2. 3. 4. 5.

Shen, M., Bever, M. B., J. Mat. Sci. (1972) 7, 741. Akovali, G., Bikyar, K., Shen, M., J. Appl. Polym. Sci. (1976) 20, 2419. Sundberg, O. E., Royer, G. L., Anal. Chem. (1946) 18, 719. Eyring, H., J. Chem. Phys. (1936) 4, 283. Matz, D. J., Guldemond, W. G., Cooper, S. L., J. Polym. Sci., Polym. Phys. Ed. (1972) 10, 1917.

April 14, 1978.


ClEA Gradient Polymers - American

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