PROGRESS INHIGH POLYMER PLASTICS’ Presented before the Division of Paint, Varnish, and Plastics Chemistry at the l o e n d Meeting of the American Chemical Sociew, Atlantic City, N.J. (Pages 449 to 479)
Phase Transition and Elastic Behavior of High Polymers H. MARK Polytechnic Institute, Brooklyn,
N.Y.
the extended state i s less probable than a state of intermediate curl. W h i l e the‘alignment of the chains takes place, molecular attraction between different chains causes crystallization. The crystallized phase has a certain amount of elasticity, characterized by a high elastic modulus. An attempt i s made to coordinate the above elementary processes and t o propose an explanation for the stress-strain curve of high polymers consisting of long-chain molecules.
The elastic behavior of high polymers is generally acknowledged to be rather composite in nature. In the case of rubber at least four different elementary processes probably cooperate to produce the actual behavior of the material. At very low stresses the van der Waals cohesion of the different segments in the curled-up chains i s overcome, and the chains are brought into a more extended state. This type of elasticity has been first emphasized by E. Mack, The kinetic elasticity of long-chain molecules i s due to the fact that
takes place a t lower temperatures than fusion, and presumably indicates that below this transition point the individual segments of the chain molecules do not carry out rotational and vibrational movements as a whole, while they A first-order transition (characterized by a sudden change do carry out such internal Brownian movements above this of quantities such as specific volume, energy, and entropy) point. Hence, according to the ideas of Eyring (9),Kuhn which may be briefly called a “crystallization” (or fusion); the @I), and Meyer (M), they act as independent kinetic units crystallized areas formed in the course of this process are of the investigated system, Figure 2 is an interpretation of mostly elongated cylindrical or flat bundles of parallelized this concept. chainlike molecules, or segments of them, with an average The crystallites of many high polymers, if irradiated with length between 500 and 5000 A. units and an average diameter of 50-200 A. They do not have plane and sharp bounparallel monochromatic x-rays or electrons] produce diffracdaries, but they represent only regions of higher geometrical tion patterns of a certain intensity and sharpness; these patorder and presumably go over gradually into the mesomorterns, together with the chemical evidence] allow deductions phous and amorphous parts, which approach the structure of’ to be made as to the arrangement of the chains inside those a liquid. It seems that one and the same chain can extend crystallites and the arrangement of the atoms inside the chains. This kind of investigation has contributed greatly through such a crystallite] enter an amorphous area, extend over into another crystallite] etc. Figure 1 is an idealized toward elucidating the structure of natural polymers, and has sketch of this situation. been recently applied with A second-order transigreat success to long1 Two of the papers in this symposium have already been published in chain compounds, partion (characterized by a INDUBTRIAL AND ENGINEERINQ CHEMISTRY. H. I. Cramer’s “Industrial titularly by Fuller (11 ) Progress in Synthetio Rubberlike Polymers” appeared in February, page 243: ~ ~ ~ ~ “Vinylidene ~ c Plastios”, ~by W. C.~ ~ appeared ~ ~ by Baker ~ ! and and Chloride Goggin andfR. D. Lowry, and thermal expansion), in Maroh, page 327. workers. RESENT experimental knowledge indicates that linear P or moderately branched and cross-linked high polymers are capable of undergoing two types of phase transitions:
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I N D U S T R I A L A N D E N G I N E E R I N G CHEMISTRY
Figure I. High Polymer Showing Crystallized (C) and Amorphous ( A ) Areas
Thus, the crystallites are a useful tool for studying the fine structure and the texture of high polymers, but they also seem to be of considerable importance for their mechanical behavior. Experience shows that many processes, such as coagulation, swelling, quenching, annealing, stretching, or rolling of high polymers, are accompanied by either the production or the disappearance of crystallites, and it was soon recognized that both types of phase transitions are intimately connected with the mechanism of the mechanical deformation of certain high polymers. This paper will discuss two of these connections-namely, the influence of crystallization on the elastic behavior of rubber, and the influence of crystallization on the reinforcement during plastic flow.
Vol. 34, No. 4
If slightly vulcanized (0.5-1.0 per cent sulfur) rubber is extended under normal conditions ( Z O O C., final elongation around 700 per cent, rate of extension and relaxation around 100 per cent of the original length per minute), stress-strain curves of the general type shown in Figure 3 are obtained. The whole curve is inside the elastic limit; as soon as the stress is released, the sample returns t o its initial state. The first linear part of the curve (OA) corresponds to a modulus of elasticity between 106 and l o 7 dynes per sq. em. It is due to a gradual uncurling and disentangling of the randomly coiled chain molecules, and can be interpreted on the basis of the molecular theory of Mack (2s) and the kinetic theory of long-chain elasticity (16, 21, 32). X-ray patterns prove that there is little or no crystallization during this period of extension. The next part of the curve ( A B ) shows a distinct decrease of the elastic modulus down t o l o 4 dynes per sq. em.; this indicates that for some reason the material becomes more easily extensible than it mas a t the beginning of the experiment. It seems that some process is taking place inside the sample which supports the external force in producing the further extension of the specimen. This process is crystallization. At the end of part OA some of the chains, or segments of them, have undergone considerable straightening and alignment, so that they now approach the arrangement they would possess if they were part of a crystallite. In this state they may be considered as a crystallization nucleus or center, and the van der Waals forces between them (although not particularly strong owing to the hydrocarbon nature of rubber) will suffice to overcome the small entropy loss involved and will force them into the crystal lattice. In so doing, these intermolecular forces cause an additional straightening of the chains and produce an increase in length, dl, a t constant external stress; this is analogous t o the well-known action of the same intermolecular forces during the compression of a real gas, where they are responsible for the volume decrease, -dV, a t constant external pressure in the van der Waals pressure-volume diagram. In both cases a first-order phase transition takes place; the gas condenses to a liquid, and the rubber crystallizes, I n both cases latent heat is evolvedheat of condensation in the gas, heat of crystallization (about 4 calories per gram) in rubber.
Crystallization and Elastic Stress-Strain Curve of Rubber If one wants t o trace the influence of crystallization on the elastic behavior of high polymers, it seems appropriate to select natural rubber as the example. This substance is, perhaps, the most thoroughly investigated material among the long-chain polymers. Recently the lattice structure of the crystallized phase has been successfully investigated by Clark and co-workers ( 7 ) , Gehman and Field (10, I S ) , Meyer and Lotmar (26),and Morss (27), with the result that the internal structure and the mutual arrangement of the chains are known with fair approximation. The diffraction pattern of the amorphous phase was thoroughly studied by Warren (SI); it indicates the presence of irregularly folded and coiled chains with internal cis structure in respect to the existing double bonds. Recently, Bekkedahl and Wood (2, $3) made a complete thermodynamic analysis of rubber by measuring specific volume, energy, entropy, and heat capacity of the crystallized and amorphous phase as functions of temperature. This provides a valuable basi? for any kind of theoretical consideration.
Chains without rotation
Figure
I,
Chains with rotation
Concept of Internal Rotation of Chain Molecules
Alfrey ( 1 ) recently developed a differential equation for the stress-strain curve of rubber which reflects the above consideration. Let us start with a 1-gram rubber sample of length 20; let us extend it to length I , which is somewhere between A
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INDUSTRIAL AND ENGINEERING CHEMISTRY
and B on Figure 3. To do this, we have to apply a certain stress a; a t the same time a certain amount of the material (q grams) has crystallized. If we produce an additional elongation of dl, we can either increase the stress by do and keep q constant, or increase the amount of the crystallized phase by dq and keep u constant.
e
%
Figure
3 (Left).
ELONGATION
Typical Stress-Strain Curve of Rubber
Theoretical Stress-Strain Curves of High Polymers 1,crystallization favored; 2, crystallization prevented
Figure 4 (Right).
I n general, both parameters u and q will change simultaneously, and we get the equation:
Quantities CY and p contain the modulus of elasticity of the amorphous phase, the heat of crystallization, and the entropy difference between the amorphous and crystallized phase. Putting these magnitudes into Equation 1, Alfrey showed that the flat part of the stress-strain curve can be successfully described in range AB, and that reasonable agreement with the experiment can be obtained. Beyond B the elastic modulus increases distinctly and reaches rather high values (up to 1O1Odynes per sq. cm.) in the neighborhood of end point C of the elastic extension curve. The reason is that the rubber crystallites have a much higher modulus (between 10’0 and IOf1dynes per sq. om.) than the amorphous phase, and that as soon as these crystallites predominate (q approaches unity), practically all the stress is taken up by the crystallized phase and an additional elongation, dl, is now possible only by stretching the rather stiff and rigid crystals. From this picture some idea may be obtained of how to change the typical rubber extension curve of Figure 3 by designing a synthetic polymer appropriately. If we favor crystallization, either by increasing the forces between the chains (OH, CONH, CC1 groups, etc.), or by making them fit more easily into the lattice (straight, not too flexible chains), OA will be shortened and may disappear completely (curve 1, Figure 4), and the increase of the modulus will be comparatively rapid. If we prevent crystallizations (weak intermolecular forces, highly branched or substituted chains), there will be no flat part A B (Figure 3 ) , and the curve will ultimately approach the shape of curve 2 in Figure 4. As a matter of fact, extension of rubber a t low temperatures (improved crystallization) approaches type 1, while high-temperature curves (inhibited crystallization) show a trend toward type 2. It must be pointed out that this way of taking into account the influence of crystallization on the elastic behavior of long.chain polymers is still crude and preliminary; it will have to be improved before it can be considered a satisfactory description of the interaction of the external stress and the intermolecular forces.
Influence of Phase Transitions on Plastic Behavior If extended beyond their elastic limit, high polymers exhibit complicated behavior. If the temperature is not too high or too low and the rate of loading is moderate, three types of deformation are generally observed as soon as a stress or shear is applied: (a) An instantaneous deformation, which is partly or as a whole reversible after the external force is removed; (b) a subsequent slow deformation, which slowly disappears partly or as a whole after the removal of the external influence; and (c) a permanent deformation, which does not disappear even after a long period. To have a clear picture of prevailing conditions, it may be well to review briefly a few fundamental facts and dehitions (4, 6, 19, 20,24). If a stress or shear is applied to a material such as platinum, quartz, glass, polystyrene, benzene, or water, two extreme cases of behavior are observed. I n one of them the external force produces a displacement, which disappears at once and completely, as soon as the material is unloaded. Suoh a system is called a n “elastic solid”. Examples are platinum (and many other metals), quartz (and many other ionic crystals), diamond, etc. I n the other extreme case the external force produces a velocity gradient, which leads to a deformation through flow. This deformation stays on indefinitely after the force ceases to act. Such a system is called a “liquid”. Examples are benzene, water, or mercury.
w
UNLOAD ED
IOy
LOADED
TIME
Figure 5.
Purely Elastic Behavior According to Hooke’s Law D = OA = CB = F / E where E = modulus of elasticity
To represent these two cases simply, the deformation, D , of the sample can be measured (elongation of a filament, deformation of a compressed cylinder, etc.) and plotted against time of observation. Figure 5 characterizes the elastic body. As soon as the stress or shear, F , is applied, deformation D takes place and remains constant as long as F acts. If F is removed, D disappears a t once. Such a system may be conveniently idealized by a steel spring (Figure 5). Hooke found that in many cases D is proportional to F ; the proportionality factor is the modulus of elasticity, E (or shear G) or Young’s modulus, and we speak of an ideally elastic body or of a material obeying Hooke’s law. Figure 6 describes the behavior of a liquid. As soon as the external force (shear), T , is applied, flow sets in and deformation D increases proportional to time. If the flow is stopped by the removal of T , D keeps its final value AC indefinitely. Such a system may be represented by a piston moving in a cylinder in which a liquid causes a certain friction (Figure 6). Newton found that frequently the velocity gradient of the flow is proportional to T ; we call the proportionality factor the “fluidity”, its reciprocal the “viscosity” of the liquid, and speak of a Newtonian flow. High polymers behave in almost all cases much less simply than elastic solids or liquids. To describe their properties it
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452
seems convenient to build up, gradually, more and more coniplicated cases combining elastic elements as shown in Figure 5 , and viscous elements as represented by Figure 6. This combination can be made by arranging the elements either in series or parallel. Maxwell (24) was the first to investigate such combinations. [If we arrange only elastic elements (having different moduli) in series, each of them is under the same stress and each extends according t o Di = F/E,; E; =
W
ior
Figure 6.
BEFORE FLOW T
AFTER FLOW T
Purely Liquid Flow According to Newton's Law
D = z.t 11 where q = viscosit of liquid 7 = externarforce (shear) D = AC = BM
modulus of element number i. Total elongation D is the sum of all the Divalues. If the same elements are switched parallel, each of them attains the same elongation D, and the stresses Fj in the single elements are proportional to the E, values. Such combinations, as well as the combination of only viscous elements, do not lead to systems approaching the behavior of high polymers.] Figure 7 shows the arrangement in series. Both elements are subjected to the same stress; total deformation D is given by the addition of the individual deformations. If the load is applied, a sudden deformation DO takes place, the magnitude of which is given by the modulus of the elastic element. Then flow sets in and leads to an increase of D proportional t o time according to the fluidity (or viscosity) of the plastic element. If the load is removed, the spring jumps back immediately from B to C, which is equal to OA, and the constant elongation CC' which was produced by the flow, stays indefinitely. The material shows plasto-elasticity. Figure 8 shows the parallel arrangement of one plastic and one elastic element. Both elements have the same elongation a t any time, In the first moment if the system is loaded, the two elements (which we assume t o have the same cross section) will have to stand the same stress. Under its influence the spring starts to extend and the piston to flow. According to Hooke's law, the stress, which is taken up by the spring, increases proportionally t o its elongation, and therefore less and less force is left to move the piston. Finally all the stress is concentrated in the spring, and the piston stands still. The plastic element has gradually unloaded itself through flow and has shifted the stress to the spring, which a t the end of the experiment bears the whole load. Under the simple conditions of our example, D increases exponentially with time (Figure 8), just as in the case of the monomolecular formation of a chemical compound or a radioactive element. If the load is removed, the elastic energy which is now stored in the extended spring begins to work and gradually moves the piston back to its original position, which, however, is reached only asymptotically. We have a retarded elastic recovery, and get a straight line for log D as a function of time, whose inclination is a measure for the retarding action due t o the viscosity of the damping piston.
VOl. 34, No. 4
However, even the two composite systems of Figures T and 8 are not yet sufficient to represent the behavior of high polymers; but they are another step toward this goal. The flow curve of a typical high polymer (Figure 9) shows all the different elements of the two cases (Figure 7 and Figure 8) together-namely, a sudden jump, a retarded elastic extension, and a viscous flow. The next approximation therefore leads t o a system built up by two elastic and two plastic elements, having two different moduli, El and E2,and two different viscosities, v1 and v2, and being arranged parallel and in series, respectively. This case is shown in Figure 10; it has been frequently discussed in recent years, together with still more general systems by many authors, such as Bennewitz (S), Burgers (6),Holzmuller and Jenckel (18),Houwink (19),Leaderman (23) and others (6,lC,17,29, SO). The application of an external force immediately produces an extension of the upper (isolated) spring, as shown by OA = F / E I in the curve of Figure 10, and the piston just below (isolated plastic element) starts into motion with a constant velocity 7/v1. The other piston (coupled plastic element) also starts, a t first with velocity 7/72; but this velocity decreases asymptotically to zero because the stress is gradually transferred from the flowing piston to the extending spring. Reinforcement occurs along line A B . If a t instant t the load is taken away, then the isolated-elastic element will contract a t once by the amount OA. This is represented in Figure 10 by the distance BC. The piston below it does not move but stays where it was at instant t . The lower spring, however, starts contracting with a velocity that decreases t o zero exponentially as a result of the damping action of the plastic element switched parallel. Thus the system approaches point M' asymptotically and exhibits ideal and retarded elasticity back and forth as well as true flow. The curve in Figure 10, which shows its behavior graphically, resembles fairly well the type of experimental curves represented in Figure 9; therefore, we may use it in making the next step of approximation in describing the deformation of high polymeric materials.
UNLOADED
LOADED AFTZFWW
Figure 7.
Elastic and Plastic Elements in Series D = Do Dit OA = Do = F / E = BC = ideal elastic deformation
+
CC'
=
D1 = I .t = D1.t = MM' = deformation due 7
to flow
Presumably its molecular significance is that in such systems elastic elements, such as curled chain molecules or clusters of them (which exhibit kinetic elasticity or Mack elasticity), can be present in two different extreme situations: Either their extension does not involve the plastic displacement of any adjacent volume element (e. g., orientation of a crystallite or flow inside of an amorphous area containing chain ends), or their extension is possible only if such a proc-
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INDUSTRIAL AND ENGINEERING CHEMISTRY
ess of rearrangement takes place, since the elastic element is somehow connected with its surroundings (through-going chains, cross links, intermolecular forces, etc.). Equally, two extreme types of viscous elements are possible: I n one type the flow does not stress any elastic element in the immediate neighborhood (isolated plastic elements, which lead to flow), and in the other type they can flow only if any adjacent volume element is stressed and stores up elastic energy (combined plastic elements, which show retarded recovery). In reality all sorts of intermediate cases are t o be expected; and instead of two moduli and two viscosities, a distribution curve of both will occur. Fuoss and Kirkwood (12)and Busse (8) showed that such distribution curves are needed to describe satisfactorily the electrical behavior of high polymers. L O E D
&
F
TIME
Figure 8.
Elastic and Plastic Elements in Parallel,
- e-3)
D = F/E(I = D%(I E = elastic modulus of spring
- .-he)
7 = viscosity of cylinder liquid
This brief paper, however, will discuss only the simplified case of Figure 10. The deformation as a function of time is then given by the simple equation, D = Do
+ Dit + D2(1 - e-ht)
(2)
t o be true in some cases but not always. Systems which undergo phase transitions upon deformation will behave differently. I n treating them we have to remember that during extension elastic elements may change into plastic ones, and therefore that retardation and reinforcement will exceed the amount predicted by the elementary unloading mechanism in Equation 2. Decrease of crystallinity by fusion or swelling has the opposite effect; it decreases the viscosities of the plastic elements, and facilitates the restoring and recovering action of the stressed elastic portions. A simple way to find out the amount of such a reinforcement due t o a phase transition is t o measure the deviation of 2 A , Equation 3 from a O0 lo 2o 30 40 5 0 straight line in thelogaTIME rithmic diagram. In fact, Fi ure 9. Flow Curve of this equation sometimes leads to significant deviaTypical High Polymer tions which may be interpreted as being caused by additional crystallization. I n doing so, however, one must keep in mind that this is not the only possible explanation for such a behavior of the material, because Equation 2 can be considered only as a rough approximation for the actual conditions. It contains only two elastic and two plastic elements, while obviously a continuous distribution curve would be expected for both, and a more elaborate equation should be derived for the deformation of the sample as a function of time. Attempts in this direction have been undertaken (I), but the relations obtained have not yet been compared extensively with experimental data. On the other hand are facts which may favor the explanation of the excess of reinforcement by additional crystallization. First is the fact that x-ray patterns indicate a decreased diffuse intensity and an increased intensity of the interference spots after plastic deformation of such systems as vis-
E5
'y
;1 T
Here, Do represents the ideal elastic elongation, and Di accounts for the flow. These two characteristic constants (for the isolated elasticity and viscosity) can be directly and independently determined by measurements a t very short times (such as sound velocity, rebound, etc.) and at very long times (linear flow after reinforcement). The rest of Equation 2 represents the exponential reinforcement. The logarithm of y = Do Dlt Dz- D plotted against time gives a straight line, '
log
y =
+
+
log 4
- At
453
UNLZED
LOADED
(3)
whose intersect with the ordinate is In Dz while its slope is -X. Thus these two constants (retarded modulus and retarding viscosity) can be determined by the evaluation of the logarithmic graph, which corresponds to Equation 3,. provided the experimental values really lead to a straight line. Recent experiments of M. Harris and J. Press have shown that in certain cases (proteins, cellulose esters, or polyesters a t medium stresses) the above analysis leads to exponential reinforcement terms. Other systems, however, (viscose, rubber, or cellulose esters a t high elongations) show a reinforcement which considerably exceeds the one predicted by relation 3. This is where crystallization seems to enter the picture. The derivation of Equations 2 and 3 is based upon the assumption that a certain number of isolated and combined elastic and plastic elements are present a t the beginning of the experiment, and that their quantity and quality do not change during the deformation of the sample. This seems
Figure I O .
Elastic and Plastic Elements in Series and in Parallel Combined D = Do Dlt DZ(1 e-h*) Do = OA = 8C = ideal elastic deformation CC' = Dz (1 - e - h t ) = retarded elastic deformation C'C" Dit = VISCOUS flow
+
+
-
cose or cellulose acetate. Furthermore, Hermans and Kratky (16), who studied the plastic deformation of cellulose xanthate, found that the amount and anisotropy of swelling are clearly affected by plastic deformation of the sample; and it is known generally that stretching decreases swelling and chemical reactivity. Nevertheless, not enough experimental material is available to prove that the excess of reinforcement due to deforma-
454
INDUSTRIAL AND ENGINEERING CHEMISTRY
tion in certain cases is due solely to phase transition. There is some indication that crystallization contributes to this effect, but further studies of flow curves and investigations of the extended samples with x-rays, by swelling experiments, or by the new method of Nickerson (28) will be necessary to clear the situation. As to the second-order phase transition of high polymers, it seems that the freezing in of the internal Brownian movement is connected with the brittle point of the material. Particularly, polyhydrocarbons and polyesters show a distinct tendency to pass from a rubberlike state at higher temperatures to a brittle (glassy) state as soon as a certain rather sharp temperature limit is reached. Bekkedahl and Xood (2, 53) determined this transition point for rubber to be -80" C.; others (3, 18, 19) found that the brittle point of polystyrene is around $80" C. The paper of Carswell, Hayes, and Nason (page 454) gives additional convincing experimental evidence of this point. Summarizing, our present experimental knowledge suggests a distinct influence of both types of phase transitions on the mechanical behavior of high polymers, although it is not yet possible to connect quantitatively the amount of converted material with the magnitude of the mechanical effect.
Literature Cited (1) rllfrey, T., Rubber Chem. Tech., 14, 525 (1941). (2) . . Bekkedahl, N., J . Research Natl. Bur. Standards, 13, 411 (1934); 23, 571 (1939). (3) Bennewitz, K., and Rotger, H., P h y s i k Z.,40, 416 (1939). (4) Bingham, E. C., "Fluidity and Plasticity", p. 217 (1922).
Vol. 34, No. 4
Brillouin, M., Am. chim. phys., [ 7 ] 14, 311 (1898). Burgers, J. M., "First Report on Viscosity and Plasticity", 2nd ed., pp. 5 and 73 (1939). Clark, G. L., IYD.ENG.CHEX.,31, 1379 (1939). Davies, J. M.,Miller, R. F., and Busse, 1%'. F., J . Am. Chem. Soc., 63, 361 (1941). Eyring, H., Powell, R. E., and Roseveare, W. E., IKD. ENG. CHEM.,33,430 (1941); J . Am. Chem. Soc., 62, 3113 (1940). Field, J. E., J . Applied P h y s . , 12, 23 (1941). Fuller, C. S., Chem. Reu., 26, 143 (1940). Fuoss, R. M., and Kirkwood, J. G., J . Am. C'hem. Soc., 63, 385 (1941). Gehman, S. D., and Field, J. E., J . Applied Phys., 10, 564 (1939). Gemant, A,, Naturwzssenschafteia, 23, 406 (1935). EXQ.CHEX.,33, 624 (1941). Guth, E., and James, H. M., IND. Hermans, P. J., and Kratky, O., Kolloid-Z., 86, 246; 89, 345, 349 (1939). Hohenemser, K., and Frazer, UT.,J . Rheol., 3, 16 (1932). Holzmiiller, W.,and Jenckel, E., 2. physik. Chem., A186, 359 (1940). Houwink, R., "Elasticity, Plasticity and Structure of Matter", 1-II.. Q27
Jeffreys, H., "The Earth", 2nd ed., p. 263 (1929). Kuhn, Vi'., Kolloid-Z., 68, 2 (1934). Leaderman, H., Testile Research, 11, 171 (1941). Mack. E . . J . Phus. Chem.. 41.221 11937). Maxwell, J. C., k h i l . Mag., [4],35, 134 (1868). Meyer, K . H., NaturzLissenschaften, 16, 781 (1928). Meyer, K. H., and Lotmar, W., Helv. Clcim. Acta, 19, 68 (1936). Morss, H. 9., J . Am Chem. Soc., 60, 237 (1938). Nickerson, R. P., IND.ENG.CHEM.,33, 1022 (1941). Poole, H . J., Trans. Faraday SOC.,21, 114 (1925). 2. physrk. Chem., 47, 671 (1892). Vogt, W., Warren, E. B., and Simard, G. L., J . Am. Chem. Soc., 58, 507 (1936). Wohlisch, E., Kolloid-Z., 89, 239 (1939). Wood, L. A., Proc. Rubber Tech. Conf., London, 1938, 933.
Physical Properties of Polystyrene as Influenced by Temperature Preparation of Test Specimens and Methods of Testing
T. S. CARSWELL, R. F. HAYES, AND H.K. NASON Monsanto Chemical Company, Springfield,
Mass.
HE physical properties of plastics are markedly inT fluenced by ambient temperature, but comparatively few quantitative data on this subject have been published for polystyrene. The variation of flexural strength and deflection a t break for compression-molded polystyrene and for a number of other compression-molded plastic materials at temperatures from -70" to f200" C, (-94' to +392" F.) was described by Nitsche and Salewski (18). Jenckel and Lagally (6) determined the tensile strength of extruded polystyrene filaments at 30" t o 60" C. (86" to 140" F.). The elongation a t 20" to 90" C. (68" to 194" F.) was reported for extruded and racked polystyrene foil by Muller (11). Since t h e mechanical properties of plastic materials are profoundly influenced by the methods used in preparing the test specimens, data on such properties are meaningless unless the details of preparation are also known. This paper describes variations in some of the mechanical properties of injection-molded polystyrene over the range froin -75" to +loo" C. (-103' to +212" F.). This method was chosen because injection molding is by far the most commonly used commercial process for the fabrication of polystyrene.
Test specimens were prepared from three grades of polystyrene whose average molecular weights were, respectively, 60,000, 95,000, and 115,000. Molecular weights were calculated from the viscosities of 0.2 per cent solutions of the polymers in toluene by means of the Staudinger relation (18):
where T~~ = specific viscosity of solution c = concentration of solution, unit moles of polymer 1
M ICm
liter
= =
molecular weight the constant 1.8 X lo-'
The specimens were injection-molded on a Reed-Prentice injection-molding machine of 2-ounce capacity. The test specimen for the tensile and flexural tests mas of the dumbbell type (Figure 1) with a 4-inch length of uniform l / k X '/4 inch cross section. These specimens were injection-molded at a heater temperature of 440' F., a ram pressure of 1100-1250 pounds per square inch, and a total cycle of 50 seconds. The specimen for the impact and hardness tests was a 5-inch X l/2 inch cross section. Injection-molding conlength of ditions were: heater temperature 450" F. ; ram pressure, 1200 pounds per square inch; total cycle, 60 seconds. A11
