Interpenetrating Polymer Networks - American Chemical Society

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13 Phase Morphology of Simultaneous Interpenetrating Polymer Networks Effect of Differences of Component Solubility Parameters and Glass-Transition Temperatures H. L. Frisch and Peiguang Zhou Department of Chemistry, State University of New York at Albany, Albany, NY 12222

Recently we prepared two classes of simultaneous interpenetrating polymer networks (IPNs): the members of the first class always contained cross-linked poly(2,6-dimethyl-1,4-phenylene oxide) (PPO), a glassy material, as one network; the second class members always contained cross-linked aliphatic polycarbonate—urethane, a rubbery material, as one network. Except for polystyrene—PPO, the linear polymers that correspond to these IPNs were wholly immiscible. On the other hand, many but not all of the IPNs were fully miscible or showed large ranges of single-phase morphology (as seen from thermal, mechanical, or transmission electron microscopy studies). We approximately correlate the phase morphologies of these IPNs with differences in the solubility parameter and glass-transition temperatures of the linear homologs of the IPN components. All of the IPNs show an intermediate composition (or an intermediate composition in the single-phase region) in which the tensile stress to break exhibits a maximum. All the IPNs show superior physical properties to the corresponding physical blends of the linear polymers.

SIMULTANEOUS INTERPENETRATING POLYMER NETWORKS (1, 2) and related topological macromolecular isomers, such as the polymeric catenanes (3), involve at least two chemically different polymers predominantly held to0065-2393/94/0239-0269$06.00/0 © 1994 American Chemical Society


270

INTERPENETRATING POLYMER NETWORKS

gether by permanent entanglements (catenation) rather than direct grafting bonds between the different polymers. Simultaneous interpenetrating poly mer network (SIN) phase morphology has been reviewed in a number of references (4, 5). Binder and Frisch (6) pointed out the theoretical possibil ity that fully miscible binary interpenetrating polymer networks (IPNs) could be formed (at all compositions of the two different component cross-linked networks) from polymers that have completely immiscible linear chains. As yet no rigorous, complete statistical dynamical theory of I P N formation and stability has been established. Nonetheless, among the many factors on which such a theory may depend, there are two factors that certainly will play an important role. The first factor is the sign and magnitude of the excess Gibbs free energy of mixing of the linear polymer chains. Although this function is not directly available, a rough guide to its magnitude is provided by the following considerations: The order of magnitude of the entropy of mixing as given by the usual Flory expression for polymer mixtures is so small that the critical entity for van der Waals polymer mixtures with no strong H bonding, donor-acceptor, or acid-base interaction is played by the Van Laar-Hildenbrand-Scatchard-type enthalpy. A rough guide to the magnitude of the excess Gibbs free energy is provided by the square of the difference of the solubility parameters of the polymers in the mixture. The larger this quantity is, the larger will be the tendency to phase separate. We will utilize the overall solubility parameter value, δ, or, more significantly (if available), the dispersive, ô , the polar δ and the hydrogen bonding, δ , contributions (2). Even when δ has not been reported in the literature, we can crudely compute its value from group contributions as listed in reference 2. The better the match is (i.e., the smaller the difference) in δ or h , δ , and § of the component linear chains, the less positive is the excess Gibbs free energy of their mixing and, thus, we expect a smaller "thermodynamic" demixing tendency in the IPNs that produces either complete or microphase separa tion. The second factor is the magnitude of the chain mobilities of the I P N polymers and their differences. In the preparation of miscible simultaneous IPNs (SINs), one attempts to approximately match the kinetic rates of cross-linking of the individual networks (I). Phase separation would tend to be suppressed if the rate of " u p h i l l " diffusion were small compared to the effective rate of network cross-link formation. Suppressed phase separation topologically prevents the separation of the networks through permanent entanglements. Chains in "glassy" polymers move more sluggishly than chains of polymers with very low glass-transition temperatures, T . Therefore, poly mer chains with well matched T s ("rough match of chain mobilities"), particularly if both T s are glassy, should show a smaller "dynamical" demix ing tendency. We investigated two series of IPNs and our conclusions bear only on these investigations. References in the literature indicate that there are fewer d

Η

d

g

g

h


13.

FRISCH AND ZHOU

Phase Morphology of Simultaneous IPNs

271

miscible IPNs than the corresponding blends of linear polymers (12). In the first G (for glass) series, one network was always composed of poly(2,6-dimethyl-l,4-phenylene oxide) (PPO), a glassy material at room temperature. The P P O was cross-linked by methyl bromination of P P O and condensation of the resultant bromide of P P O with a diamine such as ethylene diamine ( E D ) or hexamethylene diamine ( H M D ) . This condensation cross-linking of PPO did not interfere with (or was interfered with) the simultaneous free radical cross-linking employed to make the other network component of this G series of IPNs. The other free radical cross-linked networks were made from polystyrene (7) (PS), poly(methyl methacrylate) (8) ( P M M A ) , polybuta diene (9) (PB), polyurethane-polyacrylate (10) (PUA), and poly(dimethylsiloxane) (11) ( P D M S ) chains. P P O was chosen because (1) the linear and cross-linked P P O is glassy, (2) linear P P O and PS chains are miscible, and (3) polymers of other architectures of P P O and, say, P D M S (e.g., graft, block, and fully anionically mutually grafted material) can be easily made (11). In the second series, the R (for rubber) series, one network was always a polycarbonate-polyurethane (PCU). The P C U was prepared by cross-linking poly(l,6-hexanediol carbonate) with the biuret triisocyanate derived from hexamethylene diisocyanate via an addition reaction. This addition cross-lin king reaction does not interfere with (or is interfered with) the simultaneous free radical cross-linking reaction employed to make the other network component of the S I N . The other free radical cross-linked components of this R series of IPNs we made from P M M A (13), PB (14), poly(4-vinylpyridine) (15) (PVP), and PS (16). P C U was chosen for its low T and because it can produce polymers with other architectures (e.g., graft copolymers with PMMA). g

Experimental Details Experimental details of the synthesis and characterization of these IPNs can be found in references 7-11 and 13-16. Both the G and R SINs are made by simultaneously mixing and cross-linking the prepolymers and cross-hnking agent (possibly with catalysts) of the first network in homogeneous toluene solution with the difunctional monomer, polyfunctional cross-linker, and initiator of the second network. The molar masses between cross-links in these networks were typically of the order of 2 X 10 g/mol or a few multiples thereof. The materials were considered to be single-phased if they exhibited a single Τ from differential scanning calorimetry (DSC) or dynamic mechanical analysis (DMA) studies, exhibited no resolvable domains in high-magnification transmission electron mi croscopy (TEM) with appropriate staining (mercury stains for PPO, osmium tetroxide for PB, etc.), and had solid films ( l mm thick or thereabouts) that were transparent to visible light. Materials that exhibited two distinct T s (generally shifted inward from the T values of the pure cross-linked networks) were generally optically opaque and exhibited resolvable T E M domains and were thus classified as phase-separated (usually into two phases). Tensile strength and 3

g

g

272


elongation to break tests were carried out on a tensile tester (Instron) as a function of network composition. Essentially all IPNs exhibited an intermediate composition, y (in weight percent of one network component), at which a maximum tensile strength to break was observed (7-11, 13-17).


c

Ultimate Properties. The tensile strength and elongation to break were measured at room temperature on a tensile tester (Instron) at a crosshead speed of 5 in./min (ASTM D638). Differential Scanning Calorimetry and Dynamic Mechanical Analysis. The glass-transition temperatures (Ts) were determined by both DSC (DuPont 2910) and D M A (DuPont 983 D M A ; . DSC was calibrated using a 10-mg indium standard. DSC measurements were carried out on 10-mg samples from —100 to 200 °C at various scanning rates (e.g., 20 °C/min) under a nitrogen atmosphere. D M A testing was carried out from —100 to 200 °C (using a heating rate of 5 °C/min). Transmission Electron Microscopy. T E M micrographs were taken on a transmission electron microscope (Philips 300). The preparation of specimens for electron microscopy was as follows: All samples were slivered and the shavings were stained for 3 days in 4% aqueous osmium tetroxide. The samples were then rinsed briefly and dried, after which they were embedded in Spurr resin. Sections were cut with glass knives on an L.K.B. Ultratome III, were stained with uranyl acetate and lead citrate, and were then viewed. For the G series, we also stained the PPO with mercuric trifluoroacetate (17).

Results and Discussion The phase morphology of the G and R series of IPNs is summarized and compared to the corresponding blends of the linear polymers in Table I. A l l linear blends in Table I were immiscible except for the P P O - P S blend. The y is given for the P P O cross-linked with E D . Five of the IPNs made from completely immiscible blends ( P P O - P M M A , P P O - P U A , P P O - P B , P C U P M M A , and P C U - P B ) are completely miscible over the complete composi tion range, whereas two ( P C U - P S and P C U - P V P ) are fully miscible over half the P C U composition range. Only the P P O - P D M S system is microphase separated over essentially the whole composition range (with domain sizes ranging (18) from 0.01 to 0.2 μπι). c

We suggest that this behavior correlates well with the difference in solubility parameters (19) and glass-transition temperatures (20) of the linear chains that compose the networks. The values of these parameters are shown in Table II. As Table II suggests, the complete match in solubility parameters of P P O and PS is sufficient to produce fully miscible linear blends. The

13.

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273

Table I. Phase Morphology of the G and R Series of IPNs and that of the Corresponding Blend of Linear Polymers Composition Number of Range Distinct (wt%) T_ s


IPN PPO-PS PPO-PMMA PPO-PUA PPO-PB PPO-PDMS PCU-PMMA PCU-PB PCU-PS 50 PCU or less PCU-PVP 50 PCU or less a

Whole Whole Whole Whole Whole Whole Whole

1 1 1 1 2 1 1

100-60 PCU

1 2 1 2

100-60 PCU

Yc (wt%) 75 80 80 90 90 30 50 30

(PPO) (PPO) (PPO) (PPO) (PPO) (PCU) (PCU) (PCU)

70

(PCU)

Number of Phases In IPN

In Blends

1 1 1 1 2 1 1 1 2 1 2

1 2 2 2 2 2 2 2 2 2 2

a

The composition of the linear blends is thesame as the IPN.

Table II. T„ and Solubility Parameters of the Linear Polymer Chains Found in the IPNs Solubility Parameters [(cal/cm ) / ] 3

1 2

Polymer

h

PPO PVP PMMA PS PUA PCU PB PDMS 1

9.1-12.8 8.5-9.3 9.1° 8.8 8.1-8.6 7.3-7.6

8.6 9.6 9.2 8.6

3.0 5.1 5.0 3.0

2.0

8.8

2.5

1.2

4.2 2.0

Α

(°c) 211 142 105 100 -40 -95 -127

δ computed from group contributions (2).

perturbation produced by the cross-linking system is sufficiently small and both chain T s are significantly above room temperature so that the IPNs are completely miscible. The small mismatch in solubility parameters in the P P O - P M M A system prevents miscibility of the linear blends but still allows full miscibility in the IPNs (which, among other factors, are held together by long-range permanent topological entanglements). Clearly there is also an approximate "match" (i.e., small enough difference) i n the solubility parame ters and T s in this system. The large mismatch i n solubility parameters and T s of the P P O - P D M S system is where phase separation is found in both the blends and the IPNs over the whole composition range. The approximate g

g

g

274


match in the solubility parameters and the relatively small differences in T s are in accord with the full miscibility of the P C U - P B IPNs, even though the difference in solubility parameters prevents any miscibility of the linear blends of the somewhat polar P C U and nonpolar P B . P C U and P M M A are both somewhat polar and the solubility parameter mismatch is sufficiently small that the IPNs are fully miscible even though there is a significant mismatch in the T s; see Figure 1. It is not unreasonable to suppose the larger difference in T s of the P C U - P S and P C U - P V P results in miscibility over only a part of the full composition range. The phase morphology of the pseudo (or semi) IPNs is summarized in Table III. As expected the pseudo IPNs are generally phase separated as are the polymeric catenanes of P P O and P D M S . Figure 1 illustrates the difference in the D S C behavior of the (linear polymer) blend (two T s; phase separated) and the I P N (one T ; single phase) composed of 50-wt% P C U and 50-wt% P M M A as well as illustrates the fully transparent samples typically remain single phase with one T two and one-half years after preparation. The T of the I P N has decreased by about 15 °C. The original T s and tensile strength, τ, for I P N samples of P C U and P M M A are shown in Figure 2. Most samples exhibit such a decrease in the single T they possess. Almost all the pseudo (or semi) IPNs of P C U and P M M A as well as their linear blends are phase separated and exhibit two T s, as do the recently synthesized graft copolymer PMMA-gra/£-PCU and the fully gelled mutually grafted (cross-linked) P C U - P M M A network (21).


g

g

g

g

g

g

g

g

g

g

Temperature (°C)

Figure 1. DSC spectra of the full IPN ofPMMA-PCU [50-wt% PCU; run date, 1 October 1992 (DSC V .OB Dupont 2000)] (a); the same full IPN of PMMA-PCU [50-wt% PCU; run date, 17 April 1990 (DSC V .OB Dupont 2100)] (b); and a linear blend of PMMA and PCU (50-wt% PCU) (c). 4

4

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275


Table III. Phase Morphology of the Pseudo (or Semi) IPNs of R and G Series and the PPO-PDMS Polymeric Catenanes Pseudo IPNs


a

Number ofT s

Phase Domains

Number of Phases

2 2 2 l 2 2 2

Present Present Present Present Present Present Present

2 2 2 2 2 2 2

g

C-PPO-L-PMMA and L-PPO-C-PMMA C-PPO-L-PB and L-PPO-CPB C-PPO-L-PDMS and L-PPO-C-PDMS Polymeric catenanes of PPO-PDMS L-PCU-C-PMMA C-PCU-L-PB and L-PCU-C-PB C-PCU-L-PVP and L-PCU-C-PVP a

c

b

C

C is cross-linked and L is linear. The melting point of pure cyclic PDMS. Also two TL s.

I PMMA

ι 20

ι 40 PCU

ι

ι

60

80

I PCU

(wt%)

Figure 2. The ^ass-transition temperatures (solid line) and tensile strength (dashed line) of the full IPNs ofPCU-PMMA. In contrast, the P C U - P V P IPNs are only single phase if the weight percentage of P C U is larger than 50%; Figure 3 shows the T s and the τ of this system. Figure 4 shows the transmission electron micrographs of the pure cross-linked P C U and PVP, the single-phase I P N (of 70-wt% P C U ) , the phase-separated I P N (of 50-wt% P C U ) , and the phase-separated blend of the linear polymers. Figure 2 illustrates the existence of a (absolute) maximum in τ with composition, y , expressed as weight percent (in this case) of P C U . W e g

c

276



Tg(°C)

PCU (wt)% Figure 3. The glass-transition temperatures (dotted lines) and tensile strength (solid line) of the full IPNs ofPCU-PVP. previously proposed a simple plausibility argument to explain why such maxima in τ can be observed (18); this argument will be discussed next.

Theoretical Details A mean field plausibility argument explains why a maximum in tensile strength can sometimes be observed. The cohesive component of the tensile stress to break, T , of a one-phase system of two components that are not permanently entangled with each other is generally a monotone function of composition, say, the weight fraction of the first component, w. For initial mathematical simplicity we take this monotone relation as linear and write COH

T

coh

=

a

(1)

a, a + b > 0

+

with a the T of the second network and a + b the T of the first network. In an I P N there is another source of ultimate "strength" due to the permanent mutual entanglement of the component networks, T . We take the total tensile strength to break τ to be additive; that is, COH

COH

ENT

T

=

T

coh + ent T

(2)

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277

Little is gained by taking a more general linear combination for τ because it can be assumed that the proportionality factors of this combination are already factors in the constants that appear in T and T . To obtain a crude, mean field estimate of T , we consider first a special I P N composed of chains of two isotopic species, all of identical molar mass between cross-links, that meet at the vertices of a cubic lattice (six chains of each separate network are cross-linked there). We further assume that end effects and effects due to the thickness of the chains are negligible. Clearly the only effective mechanism by which such chains can interpenetrate extensively to form an I P N of two coherent, connected networks is if the species consist of two interpenetrating cubic lattices, w is the weight fraction of the first isotopic species and 1 — w is the weight fraction of the second species. The center of the unit cell of the second network can be displaced anywhere within the unit cell of the first network in this (cesium chloride) structure network. The probability of effective entanglements in this I P N is P(w) = w(l — w), by symmetry and neglecting higher order terms in w. P(w) has a maximum value at w = 1/2 when dP(w)/dw vanishes. We expect that in this case we can take coh

ent


ent

c

T

= AP(w)

ent

= Aw(l

A>0

- w);

(3)

with A the maximum contribution to the tensile strength that can come from entanglements. This entanglement contribution to the stress must be positive (hence A > O). In the general case of an I P N formed from nonequivalent chains with different molar masses between cross-links of possibly different functionality, we introduce in roughly the same approximation an effective entanglement weight fraction of network 1, φ, by defining two entanglement efficiencies, e and e , that take into account the different abilities of the two chains to encounter one another in such a way that loops of the two networks are permanently entangled. We set x

2

φ = e^/\e^Jo 4- e ( l — w)] ; 2

1 - φ = e (l 2

- w)/\e w Y

+ e (l 2

- w)]

(4)

and take now, in sufficient approximation, P(w)

= φ(1 - φ)

(5)

and x

ent

= AP(w)

= Αφ(1 - φ )

(6)

Note that this choice correctly makes P(w) vanish for w = 0 or w = 1 and reduces eq 5 correctly to P(w) = w(l — w) for the isotopic I P N for which

Figure 4. Transmission electron micrographs of the full IPNs, linear blend, and pure cross-linked components of PCU and PVP: Pure cross-linked poly(vinylpyridine) (x 221,000) (a); pure cross-linked polycarbonate-polyurethane ( X38,000) (b); the full IPN of PCU-PVP (70-wt% PCU; X31,200) (c).



FRISCH AND ZHOU Phase Morphology of Simultaneous IPNs

ft*

I U

OH

^3

1

280


e = e . Substituting eqs 5 and 1 into eq 2 yields l

2

τ = a + bw + Αφ(1 - φ)

(7)

To find a possible maximum value of τ for 0 < w < 1, we equate the first derivative of τ with respect to w (given by eq 7) to zero. Setting Δ = e — e , σ = e + e , and κ = A/b, we find that the composition w — y for which τ reaches maximum, y , must satisfy the cubic relation in y:


x

1

2

2

c

Ay 3

3

+ 3e

Δ */ + (3β|Δ - Ke e a)y 2

2

2

l

2

+ e {e\ 2

+ κ^β ) = 0

(8)

2

For sufficiently small Δ and large κ, so that terms of order Δ and Δ / κ be neglected, 2

2

3

can

t/ = |[(1 + 1/K) - |(1 - 3 / Κ ) ( Δ Α ) ] 2

c

2

(9)

where w is the w value that maximizes P(w). Finally, a somewhat better engineering approximation to the monotone T than the linear relation given by eq 1 is c

COH

T

COH

= a + bw ;

a, a + b > 0,

a

(10)

with α an emperical parameter, where α > 0. Introducing this relation into eq 2, we have τ = a + bw

a

+ Αφ(1 - φ )

(11)

where φ still is given by eq 4. The only parameters involved are the tensile stress to break of the pure network components, τ(ιυ = 0) = a and T(W = 1) = a + b, and three previously defined dimensionless parameters α, κ, and r = e /e Without loss of generality, we assume b > 0 and rearrange eq 11: 2

v

τ/fe - a/b

= w

a

+ κπι>(1 - w)/[w

+ r ( l - w)]

2

(12)

where r = e /e 2

l

(13)

Figure 5 shows plot of (τ/fo — a/b) versus w for various values of α, κ, and r that exhibit maxima or no maxima. Figure 6 shows the maxima in τ with P P O composition for the G series of IPNs (7-11) that resemble roughly the theoretical dependence. Figure 6 also represents the behavior of the R series


Figure 5. Plot of r /h — a/b versus w given by eq 1, b > 0, for the dimensionless parameter values a — 1 /2,1,2, r = 0.2,1,10, and κ = 0.1, 1, 10. h

282



T.S. to

weight % PPO (w) Figure 6. The tensile strength to break, τ, versus weight percent PPO for the G series of IPNs. (cf. Figure 2). The P P O - P U A curve shown in Figure 6 is for a sample that has a P P O network cross-linked with H M D .

Summary and Conclusions Completely miscible, one-phase IPNs can be made from components whose linear polymers are wholly immiscible (e.g., P P O - P M M A , P P O - P U A , P P O PB, P C U - P M M A , and P C U - P B ) . In come cases, one-phase IPNs that are miscible only over a portion of the composition range result (e.g., PCU-natural rubber, P C U - P S , and P C U - P V P ) , whereas other IPNs are two-phased in the whole composition range (e.g., P P O - P D M S ) . As previously suggested (19, 20), the differences in solubility parameter and T s roughly correlate with the phase morphology of fully or partially miscible IPNs even in compositions of polymers that form immiscible linear blends. These IPNs can have superior properties (e.g., ultimate mechanical strength) compared to the g

13. FRISCH AND ZHOU


283

pure cross-linked networks or any blend of the corresponding linear poly mers.

Acknowledgment Interpenetrating Polymer Networks Downloaded from pubs.acs.org by UNIV OF CALIFORNIA SAN DIEGO on 08/28/15. For personal use only.

This work was supported by the National Science Foundation under Grant DMR-90-23541.

References 1. Sperling, L. H . Interpenetrating Polymer Networks and Related Materials; Plenum: New York, 1981. 2. Olabisi, O.; Robeson, L. M . ; Shaw, M . T. Polymer-Polymer Miscibility; Aca demic: New York, 1976. 3. Huang, W. Frisch, H . L.; Hua, Y. H.; Semlyen, J. Z. J. Polym. Sci., Polym. Chem. Ed. 1990, 28, 1807. 4. Recent Developments in Polymethanes and Interpenetrating Polymer Networks; Frisch, K. C., Jr., Ed.; Technomics: Lancaster, PA, 1988. 5. Polymer Alloys; Klempner, D.; Frisch, K. C., Eds.; Plenum: New York, 1977. 6. Binder, K.; Frisch, H . L. J. Chem. Phys. 1984, 81, 2126. 7. Frisch, H . L.; Klempner, D.; Yoon, H . K.; Frisch, K. C. Macromolecules 1980, 13, 1016. 8. Singh, S.; Ghiradella, H.; Frisch, H . L. Macromolecules 1990, 23, 375. 9. Frisch, H . L.; Hua, Y. H . Macromolecules 1989, 22, 91. 10. Mengnjoh, P. C.; Frisch, H . L. J. Polym. Sci., Polym. Chem. Ed. 1989,27, 3363; Frisch, H . L.; Mengnjoh, P. C. J. Polym. Sci., Polym. Lett. Ed. 1989, 27, 285. 11. Frisch, H . L.; Gebreyes, K.; Frisch, K. C. J. Polym. Sci., Polym. Chem. Ed. 1988, 26, 2589; Gebreyes, K.; Frisch, H . L. J. Polym. Sci., Polym. Chem. Ed. 1988, 26, 3391; Huang, W.; Frisch, H . L. Macromol. Chem. Phys. Suppl. 1989, 15, 137. 12. Briber, R. M.; Bauer, B. J. Macromolecules 1991, 24, 1899. 13. Frisch, H . L.; Zhou, P.; Frisch, K. C.; Xiao, X. H . ; Huang, W.; Ghiradella, H . J. Polym. Sci., Polym. Chem. Ed. 1991, 29, 1031-1038. 14. Frisch, H . L.; Zhou, P. J. Polym. Sci., Polym. Chem. Ed. 1992, 30, 2794. 15. Zhou, P.; Frisch, H . L. J. Polym. Sci., Polym. Chem. Ed. 1992, 30, 835. 16. Zhou, P.; Frisch, H . L. J. Polym. Sci., Polym. Chem. Ed. 1992, 30, 887. 17. Hobbs, S. Y.; Watkins, V. H.; Russell R. R. J. Polym. Sci., Polym. Phys. Ed. 1980, 18, 393. 18. Frisch, H . L.; Huang, M . W. In Siloxane Polymers; Semylen, J. Α.; Clarson, S., Eds.; Prentice-Hall, Englewood Cliffs, NJ, 1993. 19. Frisch, H . L. Mater. Res. Soc. Symp. Proc. 1990, 171, 231. 20. de Barros, G. G.; Huang, M . W.; Frisch, H . L. J. Appl. Polym. Sci. 1992, 44, 255. 21. Zhou, P.; Frisch, H . L. J. Polym. Sci., Polym. Chem. Ed. 1992, 30, 2577. ;

RECEIVED for review September 2 0 , 1 9 9 1 . ACCEPTED revised October 1 5 , 1 9 9 2 .

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Interpenetrating Polymer Networks - American Chemical Society

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