Originally published July 1996
In the early days of its commercial development, polyethylene was widely considered to be inert to all liquids. The supposed inertness of this new material was immediately put to use, as one of the first polyethylene bottle applications was the packaging of concentrated hydrofluoric acid. At this point, the industry was suddenly flooded with reports of polyethylene failure. Polyethylene was reported to be unsatisfactory for cable use, and was found to crack violently on contact with methyl alcohol at room temperature. "Environmental stress cracking" was the term used to describe these failures. This term, however, was not officially defined until 1959, when J. B. Howard, who had pioneered research in the phenomenon, did so. Environmental stress cracking, according to Howard, is the "failure in surface-initiated brittle fraction of a polyethylene specimen or part under polyaxial stress in contact with a medium in the absence of which fracture does not occur under the same conditions of stress."1
Because the environmental stress cracking of polyethylene represented a serious industrial problem, a flurry of research followed in an attempt to make the material less susceptible to this type of failure. Studies conducted during this period included investigations of the effect of melt index,2 molecular-weight distribution,3 crystallinity,4 orientation,5 and rubber-particle toughening.6 Using information reported in these papers, designers in many cases have been able to elude the problem of environmental stress cracking by manipulating one or more of the above material parameters. Despite these empirical approaches, however, there is widespread ambiguity regarding the appropriate way to test for environmental stress-crack resistance (ESCR). There is also a need to clarify the fundamental molecular mechanism underlying this type of polymer failure.
The problem of environmental stress cracking can be an important one for many applications in the medical industry. These include medical labware, caps, closures, and implant components. This article attempts to provide a more fundamental understanding of the phenomenon, in the hope that solutions may become more evident.
A present limitation of ESCR testing is its inability to isolate the yield-stress property as a parameter independent of the polymer's mechanical properties. Thus, for example, constant-strain tests (such as the ASTM D 1693--the bent-strip test) have been criticized because of stiffness variations among specimens. Any such variations, of course, give rise to ambiguity when interpreting test results. Do differences among times- to-failure mirror a real difference in ESCR, or do they merely reflect the higher stress levels in the stiffer specimens?
A similar objection can be directed against constant-tensile-load testing. Although load is constant in this type of test, response to the load varies among materials. Therefore, specimen stiffness again figures as a complicating material property that can obscure ESCR as an independent parameter. Does a material fail quickly in this test as a result of its low ESCR, or because its low stiffness allows more deformation under the constant load?
A prime example of the confusion created by this situation can be seen when high- and low-density polyethylene are tested in the constant-strain, bent-strip ESCR test. What happens is that high-density polyethylene fails faster than low-density polyethylene. On the other hand, the same samples exhibit opposite effects in a constant-tensile-load ESCR test in which notched strips are subjected to a constant tensile load. Under the latter test conditions, low-density material cracks first (see Table I). (Tables and figures not yet available on-line.)
The reason for the difference in failure times between the two tests becomes clearer when one considers the influence of mechanical properties on a material's response to a load, as shown in Figure 1. Because of its relative stiffness, high-density polyethylene is stressed close to or beyond the yield point in a constant-strain test; cracking takes place in that portion of the bend in which the material is just below the yield strain. By contrast, low-density polyethylene does not come close to its yield point under the same test conditions. However, just the opposite occurs with the constant-tensile-load test, in which low-density polyethylene reaches its yield point more readily than high-density polyethylene, and is thus more susceptible to failure.
As the above discussion suggests, neither a constant-stress nor a constant-strain test provides a good criterion for discerning ESCR. Rather, the parameter that should be examined as the ordinate of an environmental stress cracking plot in a constant-load situation is percentage of yield stress or reduced stress. For a more realistic comparison of polyethylenes, this percentage should be kept constant, although the actual stress may differ widely among specimens.
The key parameter in constant-tensile-load testing is the so-called "ductile brittle transition," corresponding to that area of the stress-time plot in which a downward inflection, or "knee", becomes evident (see Figure 2). This inflection point represents the region of the curve in which ductile-creep-type deformation ends and brittle-stress-crack behavior begins. The later this transition occurs, the better the resistance of the material to environmental stress cracking.
When polyethylene is tested under constant load in the presence of a surfactant, it can be shown that the time to failure represents an acceleration of the intrinsic brittle-failure process that would take place in air at a later time. The better the ESCR, the better the resistance to slow crack growth in the absence of any obvious accelerating environment. The fracture surfaces of stress-crack failures in Igepal and in the lack of any accelerating environment are very similar (see Figure 3).
It is important to note that ESCR is extremely sensitive to temperature. Crack-growth data generated on polyethylene suggests that for every 7°C increase in temperature, the crack-growth rate is doubled. Increased temperature can therefore be regarded as another type of crack-growth accelerator.
A Graphic Model for Failure
In order to understand environmental stress cracking on a molecular level, one must be able to visualize the structural variables that most directly influence the cracking. In this regard, a review of the fiber-deformation process in semicrystalline polymers presented in graphic form is helpful. The resulting model can then be used to contrast ductile behavior with environmental stress cracking generally observed at lower loading levels in polyethylene.
Although the exact conformation of polymer chains in lamellae is a point of some controversy,7,8 one can for purposes of this model make the simplifying assumption that these chains fold on one another. In conceptualizing the failure mechanism, it is important to consider the intercrystalline or amorphous polymer chains. Figure 4a is a simplified schematic showing three types of intercrystalline material: cilia, or chains suspended from the end of a crystalline chain; loose loops, or chains that begin and end in the same lamella; and tie molecules, or chains that begin and end in adjacent lamellae.
Ductile Deformation. If a tensile load is applied normal to the face of the lamellae, the tie molecules stretch, as shown in Figure 4b. (Note that to facilitate visualization of the model, the tie molecules are illustrated as continuous chains traversing adjacent lamellae. Although in much of the polymer literature tie molecules are indeed conceptualized in this way,9 other reports describe the molecules as entanglements of a number of separate chains.) At a certain point, however, the tie molecules can be pulled out no farther (see Figure 4c), and the lamellae break up into smaller units (see Figure 5a). (According to one explanation, these "mosaic blocks" are then directly incorporated into a new fiber morphology [see Figure 5b]).10 Because tie molecules essentially are the "cement" holding the lamellae "bricks" together, their integrity is critical in order for ductile-type behavior to occur.
Brittle Failure. Brittle-type slow-crack behavior takes place over longer periods of time at lower stress levels than does the ductile deformation discussed above. The first three steps in this process are similar to those illustrated in Figure 4. However, because the material is under a lower stress level, the force necessary to achieve large-scale fiber pullout is not attained. Therefore, the loading situation can be expected to remain as shown in Figure 4c for a relatively long time. However, under long-term, low-level stress, tie molecules can begin to untangle and relax, and after a certain period of time, most of the tie molecules untangle. Ultimately, the load cannot be supported by the few remaining tie molecules, and as a result the material fails in a brittle manner, passing from the state shown in Figure 4 to that of Figure 6.
From the constant-tensile-load data, it appears that immersion in a surfactant (Igepal solution, Rhone Poulenc) accelerates the brittle failure process: the molecular displacement associated with brittle failure seems to be enhanced in the presence of Igepal. Brittle failure in polyethylene is an intrinsic phenomenon, an effect that is merely accelerated by the presence of the surfactant. Interlamellar failure--the proposed rationale for environmental stress cracking in general--can be seen as a rate-dependent process. That is, given enough time at stresses below those inducing ductile failure, tie-molecule entanglements will relax, resulting in brittle failure regardless of the presence of any environmental "lubrication" or plasticization.
It would follow from this discussion that polyethylene materials containing relatively few tie molecules are more susceptible to the various brittle modes of failure. Conversely, materials with relatively high concentrations of tie molecules would be more resistant to these types of failures. However, it should be added that if the proportion of tie molecules to crystalline molecules is too high, the materials will display high ductility, but also very low stiffness.
Visualizing the mechanism of brittle failure in light of the model described in this article can help identify several important molecular parameters for optimizing ESCR in polyethylene. These parameters include:
* Molecular weight--The higher the molecular weight, the longer the polymer chains, resulting in more tie molecules as well as more effective tie-molecule entanglements. Because polymers are polydisperse, the entire molecular weight distribution of the material is a critical factor.
* Comonomer content--Because they contain a small amount of comono-mer, such as 1-butene or 1-hexene, medium- and low-density polyethylenes exhibit short branches that tend to inhibit crystallinity. A higher como-nomer concentration results in better brittle-fracture resistance, most likely because the portions of polymer chains with the longer branches (that is, 1-hexene or longer) do not enter the tightly packed lamellar lattice. Thus, the chains with these branches add to the intercrystalline tie-molecule material (see Figure 7).10
* Density/degree of crystallinity--It would be expected that the more crystalline the material, the fewer amorphous intercrystalline tie molecules that hold it together.
* Lamellar orientation--If the lamellae are predominantly oriented perpendicular to the tensile-stress direction, they would be more susceptible to interlamellar failure than if they were parallel to the stress. This effect would be minimized in the case of a spherulitic polyethylene, since the lamellae in spherulites are oriented radially.
Medical Design Considerations
For the medical design community, a number of practical considerations can be drawn from the above discussion. For example, the stiffness of polyethylene is directly related to degree of crystallinity, as measured by density. Designers have a tendency to prefer higher-density material as a means of downgauging parts. However, it is clear that lower density results in better ESCR. Those who specify these materials will need to understand this property trade-off as they define the optimum polyethylene for a specific application.
Another critical trade-off involves molecular weight. In injection molding, for example, a higher-molecular-weight material produces higher melt viscosity, which leads to slower processing and cycle time. As we have seen, however, the higher molecular weight is desirable for ESCR. As a result, there is a need to balance processability and crack resistance for the application at hand. In many cases, resin manufacturers have used bimodal distribution of molecular weight as a means of resolving this dilemma.
As mentioned previously, temperature has a profound effect on ESCR. Crack resistance of polyethylene parts in implants cannot be determined at room temperature if one hopes to get an accurate picture of the service life of the part in situ. As a first, rough approximation of service life, accelerated testing can take place using the rule of thumb that crack-growth rate doubles for every 7°C (11°F) increase in temperature. Rate-process or Arrhenius testing can be used to quantify this accelerating effect.
Finally, it is very difficult to determine a priori precisely which environments are true environmental-stress-crack agents for polyethylene. Alcohols, silicone oils, and surfactants are the most commonly known agents, but this list is by no means exhaustive. For specific cases in which the effect of an environment on polyethylene has not been documented, testing is required to identify any possible environmental-stress-cracking effect.
It is important to note that solvents can also affect polyethylene, but in an entirely different manner than environmental-stress-cracking agents. Solvents such as hexane have a tendency to soften polyethylene, thereby lowering the stiffness of the material. In contrast, environmental-stress-cracking agents do not appreciably diffuse into polyethylene, and stiffness is unaffected. In the case of solvent attack, the best solution is to actually increase the density of the polyethylene so that the penetration of solvent is slowed. Because the mechanisms of failure for solvent attack and environmental stress cracking are so different, the means of optimizing material parameters are different as well.
Environmental stress cracking in polyethylene takes place because of interlamellar failure, which in turn is caused by relaxation of tie molecules. Resistance to this mode of failure can be best gauged through constant-tensile-load testing, while taking into account the differences in yield point among various polyethylenes. The structural variables that most directly influence ESCR include crystallinity, molecular-weight distribution, branch length, and lamellar orientation.
Some of the material in this paper appeared in Chapter 16 of Failure of Plastics, Brostow W, and Corneliussen RD (eds), Munich/Vienna, Carl Hanser Verlag, 1986.
1. Howard JB, "A Review of Stress Cracking in Polyethylene," SPE J, 15:397, 1959.
2. Dukes WA, "The Endurance of Polyethylene under Constant Tension While Immersed in Igepal," British Plastics, 34:123, 1961.
3. Herman JN, and Biesenberger JA, "Molecular-Weight Distribution and Environmental Stress Cracking of Linear Polyethylene," Poly Eng Sci, 6:341, 1966.
4. Howard JB, and Martin WM, "Effects of Thermal History on Some Properties of Polyethylene," SPE J, 16:407, 1960.
5. Howard JB, "Fracture of Polymers: Long-Term Phenomena," in Encyclopedia of Polymer Science and Technology, vol 7, New York, Wiley Interscience p 261, 1967.
6. Spenadel L, "Effect of Rubber on Environmental Stress Crack Resistance of Polyethylene," J Appl Poly Sci, 16:2375, 1972.
7. Yoon DY, and Flory PJ, "Small-Angle Neutron Scattering by Semicrystalline Polyethylene," Polymer, 18:509, 1977.
8. DiMarzio EA, and Guttman CM, "Three Statistical Mechanical Arguments That Favor Chain Folding in Polymer Systems of Lamellar Morphology," Polymer, 21:118, 1980.
9. Peterlin A, "Morphology and Fracture of Drawn Semicrystalline Polymers," J Macromol Sci Phys, B8:83, 1973.
10. Hannon JM, "Microscopic Aspects of Polyethylene at Stress Below the Yield Point," J Appl Poly Sci, 18: 3761, 1974.
Arnold Lustiger, PhD, is a staff scientist at Exxon Research and Engineering (Annandale, NJ), where he conducts research on morphology-property relationships in polyethylene. He was previously employed at Battelle Memorial Institute (Columbus, OH) and at AT&T Bell Laboratories (Murray Hill, NJ). Lustiger has consulted for numerous companies in the area of polyethylene stress cracking and has published widely on the topic.