ORT NO. DOT/MTB/OHMO-77/4

•ERFORMANCE OF PLASTIC PACKAGING FOR

HAZARDOUS MATERIALS TRANSPORTATION

PART III. STRESS CRACKING

JOHN M. CRISSMAN AND LOUIS J. ZAPAS

DECEMBER 1977

FINAL REPORT

Document is available to the U.S. public through

the National Technical Information Service,

Springfield, Virginia 22161,

Prepared for

U.S. DEPARTMENT OF TRANSPORTATION

MATERIALS TRANSPORTATION BUREAU

OFFICE OF HAZARDOUS MATERIALS OPERATIONS

WASHINGTON. D.C. 20590

NOTICE

This document is disseminated under the sponsorship of the

Department of Transportation in the interest of information

exchange. The United States Government assumes no liability

for its contents or use thereof.

PREFACE

This report is Part III of a series prepared for the Office of Hazardous Materials

Operations under Research Contract No. DOT AS-50074. Part I was concerned with mechanical

properties aspects of the performance of plastic packaging for the transport of hazardous

materials. Part II dealt with permeation aspects of performance in the transport of

hazardous materials. Part III addresses the problem of stress-cracking of large plastic

containers and discusses the mechanical properties aspects of this problem.

The authors wish to acknowledge the assistance and cooperation of the Office of

Hazardous Materials Operations Staff, in particular, Mr. Mario Glgllotti, for his continued

guidance and support throughout this project, and for the assistance and cooperation of

individual members of both the Plastic Drum Institute of the Society of the Plastics Industry

and the Manufacturing Chemists Association; as well as Mr. Paul Campbell of the Phillips

Petroleum Company. In addition, we are grateful to the editors of the Journal of the

Society of Plastics Engineers and the Materials Research and Standards Journal of the ASTM

for permission to reproduce several figures from journal articles.
' *!

-
TABLE OF CONTENTS

Page

1. Introduction 1

1.1 Review of Current Stress-Cracking Tests 1

2. Stress Cracking in Polyethylene 3

2.1 Molecular Weight and Molecular Weight Distribution 5

2.2 Branching 6

2.3 Density, Crystallinity, and Thermal History 7

2.4 Temperature 8

2.5 Prediction of Stress-Crack Performance 9

3. Experimental Studies 11

3.1 Time to Fail for Polyethylene in Uniaxial Extension in Air 12

3.2 Stress-Cracking of Polyethylene in Uniaxial Extension 23

3.3 Stress-Cracking of Polyethylene Subject to Biaxial Stress 32

3.4 Addivity of Damage in Polyethylene in the Presence of an Adverse 42

Environment

4. Summary and Recommendations 45

5. References 48
'
LIST OF TABLES

Page

Table 2.1 Typical environmental stress-cracking agents. 4

Table 2.2 Lifetime versus density in the presence of a stress-cracking 8 agent for polyethylene at 41 kgf/cm 2 (600 psi) and 353 K (60 °C).

Table 3.1.1 Polymer characteristics of two polyethylenes used in the 13

present study.

LIST OF FIGURES

Page

Figure 2.1 - Schematic representation of how molecular weight influences 5 stress-crack resistance for three different molecular weights of high

density polyethylene in nonyl phenoxypoly(ethyleneoxy)ethanol [MW(A) > MW(B) > MW(C)].

Figure 2.2 - Schematic representation of the behavior expected from branched 7 and unbranched polyethylene in the presence of a stress cracking agent such

as nonyl phenoxypoly (ethyl eneoxyl )ethanol

.

Figure 2.3 - Time to fail for commercial bottles of polyethylene as a function 10

of pressure and temperature. Data of Fulmer [17] courtesy of the SPE Journal.

Figure 2.4 - Data of figure 2.3 replotted in terms of reduced variables and 10

superposed to 60 °C. Data of Fulmer [17] courtesy of the SPE Journal. The

shift factors are indicated on the right vertical axis.

Figure 2.5 - Generalized master curve for environmental stress-cracking of 11 high density polyethylene. Data of Suezawa et al. [11]. Courtesy of the

Materials Research and Standards Journal of the ASTM.

Figure 3.1.1 - Effect of molecular weight on time to fail versus applied 12 stress in uniaxial extension for three samples of high density polyethylene

in air.

Figure 3.1.2 - Time to fail versus applied stress in uniaxial extension for 14

a high density linear polyethylene (sample D) in air.

Figure 3.1.3 - Comparison of time to fail versus applied stress in uniaxial 15

extension for two samples of high density linear polyethylene in air.

Figure 3.1.4 - Data of figure 3.1.2 replotted on log-log coordinates. Arrows 16

indicate specimens still under test.

Figure 3.1.5 - Master curve for sample D in air obtained by superposing the 18

data in figure 3.1.2 to 296 K (23 °C).

Figure 3.1.6 - Arrhenius plot of the shift factors determined from the master 19 curve shown in figure 3.1.5 and from creep data obtained for sample D in

air.

Figure 3.1.7 - Compliance J(t) versus time from creep to failure experiments 20

on sample D. Arrows indicate the onset of cold-drawing.

Figure 3.1.8 - Superposition of creep data shown for sample D in figure 21 3.1.7. Shift factors are indicated along the time axis.

Figure 3.1.9 - Time to fail versus applied stress in uniaxial extension for 22 an ethylene-hexene copolymer (sample E) in air.

Figure 3.1.10 - Data from figure 3.17 i replotted on log-log coordinates. 24 Open circles correspond to the time required to form a neck, filled circles

to the time of actual fracture. Arrows indicate specimens still under

test.

Figure 3.1.11 - Master curve for sample E in air obtained by superposing the

data from figure 3.1.9 to 296 K (23 °C). Symbols the same as in figure

3.1.10.

25
Page

Figure 3.1.12 - Superposition of creep compliance data for sample E in air. 26

Shift factors are indicated along the time axis.

Figure 3.1.13 - Arrhenius plot of the shift factors determined from the 27 master curve for sample E shown in figure 3.1.11.

Figure 3.2.1 - Schematic diagram of apparatus for determining stress- 28

crack resistance in uniaxial extension.

Figure 3.2.2 - Time to fail versus applied stress in uniaxial extension for 30

sample D in various stress-cracking agents. The solid lines represent the

data at 296 K and 330 K from figure 3.1.2.

Figure 3.2.3 - Master curve for the stress-crack behavior of sample D at 31 296 K. The solid line represents the data from figure 3.1.5 and the shift

factor for the 330 K data is indicated by the cross.

Figure 3.2.4 - Master curve for the stress-crack behavior of sample D at 33

330 K. The reference point for the room temperature shift factor indicated

by the cross is the upper left corner of the figure.

Figure 3.2.5 - Time to fail versus applied stress for sample E in various 34

stress-cracking agents. The solid lines represent the data at 296 K and

330 K from figure 3.1.9.

Figure 3.2.6 - Master curve for the stress-crack behavior of sample E at 35

296 K. The shift factor for the 330 K data is indicated by the cross.

Figure 3.2.7 - Master curve for the stress-crack behavior of sample E at 36

330 K. The reference point for the room temperature shift factor indicated

by the cross is the upper left corner of the figure.

Figure 3.3.1 - Schematic diagram of apparatus for determining stress-crack 37

behavior under biaxial stress conditions.

Figure 3.3.2 - Time to fail versus applied stress in uniaxial extension for a 39

high density linear polyethylene film. Open circles correspond to specimens

cut parallel to the extrusion direction, filled circles perpendicular, and

the triangles 45° to the extrusion direction.

Figure 3.3.3 - Time to fail versus internal pressure for polyethylene film 40

and bottles tested to failure under biaxial stress conditions in air and

in nonylphenoxypoly(ethyleneoxy)ethanol . Open symbols air, filled

symbols nonylphenoxypoly(ethyleneoxy)ethanol . The pressure scale

for the bottles is 20 times that shown.

Figure 3.3.4 - Creep failure data for polyethylene film under biaxial test 41 conditions. Internal pressures are as indicated and the arrows correspond

to the point of failure.

Figure 3.3.5 - Examples of polyethylene bottle failure in air and in 43

nonylphenoxypoly(ethyleneoxy)ethanol (a) in air, (b) in nonyl phenoxypoly-

(ethyleneoxy )ethanol . Numbers shown indicate internal bottle pressure

in psi

.
1. Introduction

In fiscal year 1976 under U. S. Department of Transportation (DOT) Contract AS-50074,

the National Bureau of Standards (NBS) Polymers Division initiated studies of mechanical

properties affecting plastic container reuse and of permeation properties of ladings through

plastics. This effort included a survey of the technical literature, an analysis of tests

given in 49 CFR (Code of Federal Regulations), a survey of other test methods, and laboratory studies on damage in polyethylene. The findings of all these investigations are presented in detail in references [l] 1 and [2]. The recommendations and conclusions made

in reference [1] are directed primarily toward the mechanical properties aspects of polyethylene containers used in the transportation of hazardous materials, while those in ref- erence [2] are concerned with permeation.

The present Final Report, also prepared for the U. S. Department of Transportation

Office of Hazardous Materials Operations under Contract AS-50074, describes additional

studies carried out in the NBS Polymers Division during fiscal year 1977. The emphasis of

this work has been directed toward providing background information and laboratory studies

on the subject of stress-cracking, the results of which will be pertinent to the establishment

of performance criteria.

For convenience, the report is divided into four sections:

° Section 1 will review stress-cracking tests currently in common use.

° Section 2 is intended to provide background material taken from the technical literature and germane to the subject of stress-cracking in polyethylene.

° Section 3 summarizes laboratory studies carried out in the NBS Polymers Division in

the areas of mechanical properties and stress-cracking.

° Section 4 contains recommendations for strengthening current DOT regulations in the areas of mechanical properties and stress-crack resistance, in particular paragraph

178.19 of 49 CFR (Specification 34).

1.1 Review of Current Stress-Cracking Tests

In reference [1], it was noted that one important area not addressed in the present

DOT regulations is stress cracking. Container performance under conditions in which the

lading may cause stress cracking reguires particular scrutiny in cases where paragraph

178.19 of 49 CFR [3] is applicable (Specification 34; reusable molded polyethylene container

for use with overpack). In order to maximize the expected service life of a "reusable"

container, it is necessary to determine what ladings are stress-cracking agents and what

level of stress-cracking agent is acceptable to be shipped and for how long or often.

Thus, there is a clear need to incorporate into DOT regulations a test, or tests, to determine

stress cracking severity.

Among the various stress-cracking tests presently in common usage are principly those

included in the American Society for Testing and Materials (ASTM) 1976 Annual Book of ASTM

Standards [4]. It will be useful to briefly review several of these tests, insofar as they

apply to plastics such as polyethylene.

(1) ASTM D1693-70: (Standard Method of Test for Environmental Stress-Cracking of Ethylene

Plastics).

This test is commonly referred to as the "bent strip test." Specimens in the form of

a strip are notched according to a specified procedure and then bent into a U shape in a holder. The specimen is then exposed to a surface-active agent either at 323 K (50 °C) or

373 K (100 °C), and the proportion of the total number of specimens that crack in a given

time is observed. From these data, the 50 percent failure point (F 50 ) is determined.

'Numbers in square brackets refer to references at the end of Part III.

1
However, criticism of this test has been made on at least two counts. In the first

place, it is known that the behavior of polymeric materials can be notch sensitive. That

is, the time to fail can be significantly different for notched specimens as opposed to unnotched specimens. Even though the notching procedure may be nominally the same for each

specimen, it is difficult to assure that differences will not occur from sample to sample

as a result of different operators.

The second point is that the strain on the specimen cannot be precisely specified.

Because of the sharp bend in the specimen, the outer surface may be highly strained in ten- sion, whereas the inner surface is in compression. If the material being tested is highly

crystalline or has a high flexural modulus, the specimen often fails while being bent into

the initial u-shape.

(2) ASTM D2561-70: (Standard Method of Test for Environmental Stress-Crack Resistance of

Blow-Molded Polyethylene Containers).

Three alternative procedures are included in this test. Procedure A consists of ex- posing any filled, sealed, blow-molded container to a potential stress-cracking agent

within the container at 333 K (60 °C).

Procedure B exposes a partially filled and sealed blow-molded standard container to the action of a specific known stress-cracki ng agent within the container at 333 K (60 °C), as well as to the action of this agent as an external environment.

Procedure C is similar to Procedure B except that a constant elevated internal pressure

is maintained [34.5 kPa (5 p s i ) ] in addition to elevated temperature.

In all three procedures the time to fail is observed and a minimum of 15 containers

tested. This test is most useful in isolating the weak point of the container and such

factors as the influence of the blow-molding conditions on the stress-crack resistance of

the container.

Perhaps the most serious criticism which has been directed at this test is that it is carried out at only one elevated pressure and temperature (Procedure C) and therefore

important information may be lost. For example, it has been reported that the failure

mechanism may be different depending upon the stress level and temperature [17]. In fact,

it has been reported [17] that one constant pressure bottle stress-cracking round robin

test was considered to be a total failure because a change in failure mechanism occurred

above the 34.5 kPa (5 psi) pressure employed in the test and none of the bottles failed in an acceptably short time.

(3) ASTM D2648-70: (Standard Recommended Practice for Measuring Time-to-Failure by Rupture

of Plastics Under Tension in Various Environments).

In this test, a range of fixed loads is applied in tension to specimens exposed to specific environments at constant temperature. The time to failure is then observed. Data obtained from tests of this type are useful in comparing the ability of materials to support

a constant applied load for long periods of time. In the presence of stress-cracking

agents, the lifetime may be significantly decreased under conditions of an applied stress. i

'

! One possible deficiency of this test, with regard to container design, is that the

stresses encountered in actual use nearly always contain an element of biaxial ity, or multiaxial ity, and, therefore, the data gathered under uniaxial conditions may not be appropriate. However, it has been suggested that a scaling may be possible between results

obtained from tensile tests in uniaxial extension and comparable data obtained from actual

container tests [17]. One major appeal of this latter test lies in its relative simplicity

compared to the other more elaborate or expensive tests mentioned previously.

f (4) ASTM D2951-71: (Standard Method of Test for Thermal Stress-Crack Resistance of Types

III and IV Polyethylene Plastics).

This test is similar to ASTM D1 693-70 to the extent that both tests are intended to be

used on materials designed for use as jacketing or insulation on wire and cable. Some

2
polyethylenes, in particular those of high density or low molecular weight, are sensitive

to embrittlement or cracking at elevated temperatures. As in D1693-70, the specimens are severely bent prior to testing, in this latter case wrapped spirally around a small diameter

metal rod, so that again the strain at the outer surface of the specimen may be large.

It is apparent that there remains a need to further develop a stress-cracking test

more suited to the testing of containers. Of the three environmental stress-cracking

tests just mentioned, 02561-70 is perhaps the most suitable test presently used since, in

principle. Procedure A can be applied to any size or type of container. However, this

test could be made to yield more quantitative information by requiring tests to be carried

out over a range of elevated pressures and temperatures. Some of the experimental work to

be described in Section 3 has been carried out with this point in mind.

2. Stress Cracking in Polyethylene

Section 2 of this report is intended to provide the DOT Office of Hazardous Materials

Operations with background information on stress cracking of plastics. The technical

literature concerned with this subject is extensive, and serves to underscore the importance

of the problem to commerce and industry. In the present report, we shall be concerned

solely with the environmental stress-cracking of polyethylene, since polyethylene has been

established as the most important commercial polymer used for the manufacture of plastic

containers. Included at the end of this report are a number of references [5-26] which in our view are representative of the technical literature as a whole. A cursory glance at

this work reveals that the major portion of research activity occurred during the 1 950 '

s

and early to mid 1 960 ' s . A fairly comprehensive review paper on the subject up to that

time appeared in 1959 [8] and a somewhat more recent article covering stress-cracking can

be found in reference [15]. It can also be observed that, more recently, the ever-increasing

pressure to use plastic materials in new and more demanding applications has led to renewed

interest in the subject.

The term, "environmental stress cracking," as it is commonly used, encompasses any one

of several forms of what shall be termed here premature failure of a material in times

shorter than experienced in an inert environment due to the presence of an adverse environment, all other factors being equal. Polyethylene gained rapid acceptance commercial ly in

many applications because initially it was believed to be relatively inert to a very broad

range of chemical compounds. One of its early uses, which remains so even yet, was in the

safe transport of hydrofluoric acid. However, it soon became apparent that, under certain

conditions, containers or parts failed after only very short service periods due to cracking.

One such observation was that polyethylene insulating material used in the manufacture of

wire and cable in many instances cracked very rapidly when treated with certain wetting

agents, detergents, or oil lubricants. Since that time the development of new more highly

stress-crack resistant polymer resins as well as improved processing technology has greatly

reduced the problem in many applications. Another response resulting from much of this

effort has been the adoption by ASTM of the several environmental stress-cracking tests

discussed in Section 1. However, where the reuse of large free-standing plastic containers

for hazardous materials transport is concerned, there remains a need to provide further

safeguards insofar as environmental stress-cracking problems may arise.

On a microscopic scale, the mechanism, or mechanisms, of stress-cracking are not well

understood. To cite one -vecy recent article [20], "Environmental stress-cracking in polyethylene has been and continues to be a major unknown in materials science." However, on

an empirical basis, a good deal of information is available, although there is as yet no general agreement as to what constitutes proper design criteria. There is agreement that,

in addition to the presence of an external agent, a large number of material parameters

exist which to a greater or lesser extent influence stress-crack behavior. Among these are molecular weight and molecular weight distribution, density, degree of crystall inity,

thermal history, branching, spherulitic structure (size), presence of crosslinks, and

ambient temperature. Common to most models which have been proposed for stress-cracking

are the following features: (1) the appearance of microvoids in the specimen; (2) a rate

process by which voids grow and eventually coalesce, and (3) the occurrence of plasticity

and "strain hardening" of material surrounding the voids.

3
As the term implies, "stress cracking" occurs in the presence of stress. Such a stress can come about either as a consequence of an applied deformation, in which case

stress relaxation occurs, or due to an applied load in which case the material creeps, or

a combination of both. In either case, the ultimate result is what we choose to call

premature failure, failure due not because of the imposition of the stress alone, but due

in large measure to the presence of an adverse environment which greatly accelerates the

process. It was mentioned previously that stress cracking in the broad sense encompasses

many forms of premature failure, however, a distinction is often made as to the particular

form. In the context of this report, two rather general classifications will be made:

(1) stress-cracking due to the presence of surface active agents, and (2) stress cracking

in the presence of a solvent or swelling agent. It should be noted that in both cases, chemical alteration of the polymer is not required in order to cause premature failure.

In the first instance, failure generally appears in the form of cracks, which eventually

enlarge to the point where rupture or fracture occurs. Therefore, fracture can occur well

below the yield point of the polymer. Substances which give rise to this form of failure

largely include polar compounds such as alcohols, and detergents or wetting agents. Such

materials do not necessarily solvate the polymer or attack it chemically, but initiate or enhance the growth of surface or internal defects in the presence of an applied stress.

It should be noted, however, that an externally applied stress is not necessary to cause

stress cracking. The phenomenon can also occur solely as a consequence of "built-in" or residual stress in the material which arises during processing.

Solvent-cracking, on the other hand involves actual solvation of the polymer. Many

liquids, such as organic solvents, are absorbed to a limited extent into the polymer,

causing swelling. For example, for one type of polyethylene it is reported [6] that at 298 K (25 °C) about ten weight-percent of n-dodecane (C 12 H 26 ) is absorbed, whereas at 333 K (60 °C) the amount absorbed is about thirty weight-percent. In solvent-cracking,

the cracking agent is absorbed into the polymer giving rise to internal triaxial stresses

which presumably weaken the material by interfering with the normal intermolecular bonding

forces. In Section 3 of this report, examples of both types of stress cracking will be

demonstrated for two different polyethyl enes, one a linear polymer, and the other an ethylene-hexene copolymer which is branched. It might also be noted that, as a rule,

those substances falling into the category of solvent-cracking agents also have high

permeation rates through polyethylene. In Table 2.1 are listed a number of typical

environmental stress-cracking agents. No attempt has been made in the list to distinguish

whether a particular chemical compound is a "surface active" stress-cracking agent or a "sol vent-crackfng" agent, since in some cases the distinction may not be clear cut. It can

be observed that several of the substances listed fall into the category of hazardous materia'

TABLE 2.1 Typical environmental stress-cracking agents*

Aliphatic and aromatic liquid hydrocarbons

Alkali hydroxides

A1 kanolamines

A1 cohol

s

Animal, vegetable, and mineral oils

Organic acids

Concentrated sulfuric acid

Ester-type plasticizers

Detergents, wetting agents, and soaps

Lubricants and mold release agents

Aldehydes and ketones

Plating bath additives

Sulfated and sulfonated alcohols

Depolymeri zed rubbers

Polybutenes and polyglycol esters

Sil icone oils

Sulfonated fluids

*List compiled from references [6] and [7].

4
The remainder of this section will be concerned with a brief description of how some

of the material parameters influence stress-cracking and of the direction taken by a number of researchers in attempting to predict time to fail due to stress cracking under

an arbitrary condition of stress, environment, and temperature. In preparing this section,

we have drawn heavily from the literature cited previously [5-25] and in some instances

from our own laboratory studies. Properties which influence stress-crack resistance

generaly fall into one of three categories; those involving the polymer chemistry, those

describing the physical state of the material, and external variables such as temperature.

All three can have a profound effect on lifetime behavior under stress in the presence of

an adverse environment.

2.1 Molecular Weight and Molecular Weight Distribution

Two characteristics of the polymer itself which are known to affect stress-crack

resistance are molecular weight and molecular weight distribution. As a general rule, it

is found that stress cracking is more severe the lower the molecular weight of the polyethylene. In terms of melt index, the higher the melt index, the poorer is the stresscrack resistance. This result is demonstrated schematically in figure 2.1 which shows a possible stress versus time to fail behavior for three different high density polyethylenes

in nonylphenoxypoly(ethyleneoxy)ethanol . Note that in the example of figure 2.1 the

functional form, or shape, of the failure curve is not drastically changed from one sample

to the other, but rather the major difference is an overall shift of the curves. However,

when comparing failure curves for different polymer types, this may not be the case.

Figure 2.1 - Schematic representation of how molecular weight influences stress-crack

resistance for three different molecular weights of high density polyethylene in nonyl

phenoxypoly(ethyleneoxy )ethanol [MW(A) > MW(B) > MW(C)].

5
A second parameter also related to molecular weight is molecular weight distribution,

or polydispersity. There is ample evidence that the presence of low molecular weight

material impairs stress-crack resistance. For example, it has been reported [9] that the

removal of most low molecular weight polymer from one sample, resulted in the average time

to crack increasing from 490 minutes to 648 minutes. In another instance, it was found

[15] that the removal of less than 10 percent of the lowest molecular-weight component

increased the cracking time from 6 hours to more than 2 months. Such behavior is one reason why melt index may not always be an effective measure of stress-crack resistance,

since it is most nearly a function of weight-average molecular weight and may not be sufficiently sensitive to a relatively large number of low molecular weight species present.

From the point of view of qualification tests, parameters such as molecular weight

and molecular weight distribution underscore the importance of running tests on pieces

taken from the finished product, as was pointed out in the previous report [1]. In order

to achieve a better finish on the final product or improved melt flow properties, it is common practice to alter the molecular weight distribution of the base polymer by the

addition of selected low molecular weight polymer. Even the addition of only a few percent

of such material can dramatically affect physical properties, including stress-crack

resistance. Moreover, the addition of color or stabilizing compounds is generally accom- plished by first blending the substance (pigment, carbon-black, etc.) with a lower molecular

weight polymer which is then added to the base resin at the time of processing. Also,

recycling of material left over during processing can lead to overall deterioration of the

polymer, since the recycled polymer may have suffered some degradation prior to its addition

to a later batch.

2.2 Branching

The third parameter associated with the polymer chemistry is branching. It is well

known, for example, that the introduction of branches along the polymer main chain can vastly improve stress-crack resistance. Many of the commercial polymer resins currently

in use, particularly those used for containers, are branched in one way or another. The

question as to just what distribution of branch points, the type of branching, or'branch

length will optimize stress-crack performance has been difficult to assess in view of the

experimental problems encountered in characterizing the polymer. Recently, new techniques

such as C

13 NMR (Nuclear Magnetic Resonance) offer hope in this direction.

There is agreement that it is primarily the shorter branches that control the cracking

behavior of olefinic polymer (in particular polyethylene) by virtue of their much greater

abundance. On the assumption that any given branch point is about equally disruptive,

irrespective of the length of the branch, the more abundent shorter branches are expected

to dominate in influencing the amount and structure of the crystalline phase.

Density has generally been regarded as one of the most useful measures for determining

branching. However, as we shall see presently, the density can be highly misleading. It

is pointed out in reference [15], "other facets of crystalline texture ascribable to the

effects of branching, but not reflected in density also exercise potent influence on cracking behavior. The size of the crystallites and spherul i tes‘ present appear to be

among the more important of these."

A comparison of behavior which might be expected for typical commercial branched and

unbranched polyethylene is shown schematically in figure 2.2. In Section 3 of this report

such behavior will be demonstrated by data obtained for two different polyethylenes. Note

that in the high stress region the linear polymer is superior in terms of time to fail,

whereas with decreasing stress the two curves cross so that at small stresses (the region

of stresses under normal operating conditions in transport), the branched material eventually becomes far superior. In the previous Final Report [1], such behavior was shown for

branched and unbranched polyethylene, but in the absence of stress-cracking agent.

6
*

TIME TO FAIL

Figure 2.2 - Schematic representation of the behavior expected from branched and unbranched

polyethylene in the presence of a stress-cracking agent such as nonyl phenoxypoly (ethyl eneoxyl )ethanol

.

2.3 Density, Crystallinity, and Thermal History

A number of parameters which describe physical state of the material are also influencial in determining stress-crack behavior. Two of the most important of these are density

and degree of crystallinity, but, because of their strong interdependency on one another

through thermal history, it is appropriate to consider them together. Noted already was

the extent to which branching influences both density and crystallinity and, in turn,

stress-crack resistance. However, even in the absence of branching, both the density and

crystal 1 inity can be altered significantly through differences in thermal history during

processing.

Also, it was alluded to earlier that density alone may be highly misleading in predicting stress-crack resistance. That this statement is true has been shown by Lander [26].

Shown in Table 2.2 are data for several specimens either similar or increasing in density.

In the absence of additional information, such as the melt index, there is no apparent

correlation to lifetime per se. Another complication is the fact that a number of these

specimens were branched. However, if the lifetimes in the presence of a stress-cracking

agent are compared under conditions of a constant melt index, then there is better correlation, the trend being that with increased density the lifetime increases.

Clearly, a problem exists when attempting to compare stress-crack behavior of different

polyethyl enes on a basis of density or crystallinity alone. There are many ways in which to arrive at the same density, while allowing other key parameters to vary. For example, a branched polyethylene can be slow cooled in such a manner as to arrive at a density of, say,

0.950 g/cm 3

. At the same time, a linear polyethylene (unbranched) can be rapidly cooled to the same density, yet their mechanical or stress-crack behavior may be significantly

different. Moreover, a given polytheylene resin can be processed under different thermal

conditions such that they have the same density but quite different spherulitic structure

(texture) and, consequently, stress-crack behavior [9],
Table 2.2 Lifetime versus density in the presence of a stress -cracking

for polyethylene at 41 kgf/cm 2 (600 psi) and 353K (60 °C). +

Density Melt Index Lifetime

(q/cm 3 (q/10 min) (Hours)

0.914 1.8 0.2

0.919 0.3 350.0

0.918 0.8 60.0

0.918 1.6 0.7

0.917 3.0 0.2

0.917 20.0 0.1

0.919 3.0 9.0

0.927 1.4 120.0

0.928 3.0 20.0

0.930 16.0 0.7

0.948 0.4 170.0

0.950 8.0 20.0

0.950 20.0 5.0

0.960 1.0 250.0

0.960 4.2 250.0

+ Data from Reference 26.

In addition to the parameters already cited, there are other important considerations

in the design of a new test, or tests, for stress-crack resistance. Howard [15] has

pointed out that data obtained via the bent-strip test (ASTM 1693) indicate that as the

density increases, the likelihood of premature failure also increases, all other factors

being equal. Yet the results of Lander [26], which were obtained from constant load

experiments done in extension, suggest the opposite behavior. Apparently, it is to be

expected that the results obtained depend in large measure on the type of test employed,

i.e., the type of deformation applied.

2.4 Temperature

Perhaps one of the most critical factors of all, insofar as stress cracking is concerne,

is the ambient temperature. Such a statement will be particularly applicable to the

transportation of hazardous materials where temperatures as high as 60 °C (140 °F) are not

unreasonable'. In reference [15], it is shown that for one type of polyethylene an increase

in temperature of about 24 °C (from about 38 °C to 62 °C) decreased the 50 percent failure

time (F 50 , ASTM D1693 test) by nearly a factor of 1000 in time. Numerous similar observations can be found throughout the technical literature [9,12,17,26].

However, some care must be exercised when attributing a decrease in lifetime at an elevated temperature to the presence of stress-cracking agent. A seeming deficiency of some results reported in the literature is that the only data presented are those for

which the experiment was done in the presence of a stress-cracking agent, and few, if any,

data are provided indicating behavior in an inert environment. In some instances, the

acceleration of failure due to increased temperature has been attributed primarily to presence of the stress-cracking agent. In the previous Final Report [1], it was shown

that for one linear high density polyethylene (Sample A) in air, an increase in temperature of 34 °C (from 23 °C to 57 °C) resulted in a decrease in time to fail by a factor of

600 at a stress of 150 kgf/cm 2

. For the same polymer at 97 °C and 75 kgf/cm 2

, a drop in failure time of nearly 5.5 decades from the room temperature value was recorded. Therefore, although the presence of stress-cracking agent may actually accelerate failure at higher temperatures, the primary effect may simply be one of an overall degradation of the

mechanical properties, in particular, a decrease in modulus. Experimental data showing

just this behavior will be presented in Section 3 of this report.

8
2.5 Prediction of Stress Crack Performance

Although the mechanism, or mechanisms, of stress-cracking may not be well understood

at the molecular level, some success can be achieved in predicting stress-crack behavior

of a given polymer when the conditions of stress, temperature and stress-cracking agent

are specified. Ideally, one wishes to obtain experimental data on the very long time

behavior under a wide variety of conditions (failure times of the order of years) in order

to predict lifetime use. However, from a practical point of view, such tests are not

feasible, especially where quality control in production is concerned; and in the absence

of such data one must then rely on accelerated test methods. Underlying the success of

accelerated testing are two assumptions. The first of these is that within the range of

temperature extremes expected in use, the mechanism, or mechanisms, of failure do not

change. By this, we mean that the functional form, or shape, of the time to failure

versus stress curve does not change significantly with temperature, but simply is shifted either vertically or horizontally, or both. The second assumption is that, in

addition to temperature, the accelerating effect (to failure) due to the presence of the

adverse environment also does not measureably alter the shape of the failure curve. This means that data collected on a variety of different stress-cracking agents can, for a given polymer under arbitrary conditions of stress and temperature, be superposed onto one master curve. In principle, the relative severity of a given stress-cracking agent can be

judged by specifying a set of shift factors which have as their reference point the master

curve. Then, the extent to which accelerated testing will be useful as a predictive tool

will depend in large measure on the amount of scatter in the data comprising the master

curve, i.e., the extent to which the predicted long time values can be bounded in an acceptably narrow band. Realistically, the best one can probably hope to do is to establish

a lower bound to the failure curve.

The one major unknown confronting this approach lies in predicting long time behavior

under conditions of a non-simple thermal or stress history. Such a situation will be of

particular importance in transportation where not only the temperature-stress history may

be complex, but also with regard to specification 34 where reuse is intended, it may be

difficult to assure that a container has not been exposed to more than one possible type

of stress-cracking agent. In such a case, as is the case with mechanical properties in

general, the concept of additivity of damage enters. Such an additivity of damage rule

was mentioned briefly in the previous report [1] and will be considered further in Section

3 of this report. Suppose that a container is exposed to a lading, a potential stresscracking agent, under conditions of temperature and stress severe enough that some damage

occurs. Such damage may not necessarily be detectable by ordinary means. Upon its return, even a very careful and thorough cleaning cannot insure that all of the previous lading

has been removed. In fact, the body of evidence to date suggests that if the conditions of

exposure have been sufficiently severe to cause damage to the material, then it will

probably not be possible to remove by ordinary cleaning a small amount of lading trapped

in or near microvoids, or other defects, which 'on a subsequent trip may again cause

further damage if the stress-temperature conditions are severe enough. That this is true

will be demonstrated for several simple histories in Section 3, the implication being that

a simple additivity of damage concept may not be applicable.

In Section 7 of reference [1], the influence of temperature on the time to failure

versus stress behavior was shown for several different polyethylenes. Those results were

obtained for specimens under constant load in uniaxial extension and in the absence of an adverse environment (see figure 7.1). In Section 3 of this report more extensive work

will be presented for two polyethylenes in both air and a variety of stress-cracking

liquids. However, before presentation of the experimental work, we shall cite two studies

chosen from the literature which will serve to point out the utility of accelerated test

methods insofar as stress-crack behavior is concerned.

In figure 2.3 are shown data obtained by Fulmer [17] for commercial bottles of

polyethylene tested to failure at different temperatures in nonylphenoxypoly(ethyleneoxy)-

ethanol . The bottles were tested according to a procedure similar to the ASTM D2561 test,

except that all the bottles were tested to actual failure at different levels of internal

pressure.

9
Figure 2.3 - Time to fail for commerical

bottles of polyethylene as a function of

pressure and temperature. Data of

Fulmer [17] courtesy of the SPE Journal.

Now, if the data of figure 2.3 are treated such that the pressure is divided by the

absolute temperature (reduced variable) and the resulting values at different temperatures

superposed, they appear as in figure 2.4. From figure 2.4, it can be seen that a reasonably

good superposition is obtained with an acceleration of a factor of nearly 100 in testing

time resulting from increasing the temperature from 60 to 97 °C.

Figure 2.4 - Data of Figure 2.3 replotted in terms of reduced variables and superposed

to 60 °C. Data of Fulmer [17] courtesy of

the SPE Journal. The shift factors are indicated by the right vertical axis.

c

E

I0 9

I0 4

IO»

10 *

10

TIME (min) ot 60*C| - \ £> 60*C

\

Q 73*C

O 85*C

• 97*C

- *\Q

- Y

SI*

"

c

E->Q ' ^

p ppp p

1

0 • 0

1 0 12 3

Pressure/ temperature , mm Hg/°K

10

Tempera

ture

,
Another example of an effort to establish a master curve for a variety of stresscracking agents is the work of Suezawa et al. [11]. Figure 2.5 shows their master curve

for one high density polyethylene in the presence of four different stress-cracking agents.

Figure 2.5 - Generalized master curve for environmental stress-cracking of high density

polyethylene. Data of Suezawa et al. [11]. Courtesy of the Materials Research and

Standards Journal of the ASTM.

In figure 2.5, at represents a generalized time to crack and Ba 0 a generalized

equivalent stress. In this work the biaxial ity of the imposed stress was achieved by

applying a torsional stress to the specimen (a cylinder) at the same time the specimen was loaded in tension. The ratio t/o represents the shear stress in torsion divided by the

tensile stress. As is evident, a rather good superposition can be obtained for a wide

range of conditions. In principle then, it should be possible from a master curve such as this to predict failure times for a nearly arbitrary state of stress, temperature, and

stress-cracking agent.

3. Experimental Studies

In fiscal year 1976, under DOT contract AS-50074, the NBS Polymers Division initiated

laboratory studies on damage in semicrystalline polymers. Much of this work was concerned

with the time to fail for polyethylene as a function of applied stress in uniaxial extension under conditions of varied temperature, molecular weight, and environment. In addition, the idea of additivity of damage was examined for several simple loading histories.

The work to be described in Section 3 of the present report represents a continuation and

further extension of experiments begun in fiscal year 1976. As a matter of convenience,

Section 3 has been subdivided into the following several headings:

(3.1) Time to failure for polyethylene in uniaxial extension in air.

(3.2) Stress cracking of polyethylene in uniaxial extension.

(3.3) Stress cracking of polyethylene subject to biaxial stress.

(3.4) Additivity of damage in polyethylene in the presence of an adverse environment.

The experimental procedures followed, unless otherwise specified, were the same as those described in reference [1]. New procedures or apparatus will be described in the

text where appropriate.

11
3.1 Time to Fail for Polyethylene in Uniaxial Extension in Air

In reference [1], log time to break versus applied stress data (uniaxial extension) were shown as a function of temperature for three different molecular weight high density

polyethylenes (figures 7.1, 7.2, and 7.3). Samples A and B were commercial linear polyethylenes, while sample C was the NBS Standard Reference Material (SRM) 1475 polymer. To better demonstrate the effect of molecular weight, we have replotted these data on loglog coordinates in figure 3.1.1.

Figure 3.1.1 - Effect of molecular weight on time to fail versus applied stress in uniaxial

extension for three samples of high density polyethylene in air.

12
A comment is appropriate here concerning the definition of breaking. In the previous

report, breaking, or failure, was defined simply as the occurrence either of fracture

(separation) or of neck formation followed by cold-drawing. By this definition, nearly

all the specimens from each of the three samples (see figure 3.1.1) tested at room temperature (296 K) failed; only specimens from sample A actually cold-drew to any extent

before breaking. At elevated temperatures, several specimens cold-drew, but did not fracture before the experiment was stopped. Therefore, in order to be consistent with the

previous report, we shall continue to apply the term "failure" to describe the data, with

the recognition that in many cases fracture did not occur.

To return to figure 3.1.1, it can be seen that for large stresses the behavior of all

three samples is not dramatically different as a function of molecular weight; and the

bounds determined by the experimental scatter in data points are reasonably narrow. However, as the stress decreases to small values, the divergence in lifetimes among the

three samples becomes much more pronounced, an indication that a significant improvement

in lifetime can be gained by increasing molecular weight (A > B > C). Note that at a given stress level (at small stresses), the scatter in the data point is less at the

elevated temperatures than is the case at room temperature. This behavior may result in

part from the difficulty encountered in maintaining a constant room temperature over long

periods of time, whereas, the elevated temperatures were controlled to within about

(0.05 °C). On the other hand, from the high degree of scatter in the data points and the

rather pronounced bending over of the curves for the highly crystalline low molecular

weight sample C, it is apparent that other factors, in addition to temperature, were important in determining lifetime.

More recently, experiments in uniaxial extension have been conducted on two additional

commercial ethylene polymers, and much of the remainder of work to be described in this

report will concern these two polymers. One sample, designated D, was selected because it was a linear high density polyethylene recommended for use in food packaging. This sample

has material characteristics very similar to sample A used earlier. The second sample,

designated E, was chosen because it is branched. This polymer was actually an ethylenehexene copolymer recommended by the manufacturer for use in blow molded items, such as large shipping containers, gasoline tanks, and chemical tanks. Some of the pertinent

material characteristics of both samples are given in Table 3.1.1.

Table 3.1.1 Polymer characteristics of two polyethylenes used in the present study.

Sample

Density Melt Flow 1 (gm/cm 3

) (gm/lOmin.) Structure

Environmental

Stress-Crack

Resistance 2 (Hours

)

D 0.964 0.75 3

E 0.950 10.

0

4 Linear 15-20

Branched 450

1 - ASTM Test D1 238-73

2 - ASTM Test D1 693-70

3 - Condition E of 1 4 - Condition F of 1

Experimental data for log time to fail versus applied stress in uniaxial extension

are shown for sample D at several temperatures in figure 3.1.2. Here, the term "time to fail" implies separation, for most of the specimens, with the exception of several at the

higher temperatures, broke almost immediately upon necking, or fractured prior to necking.

In overall behavior, sample D appears very similar to sample A [1]. For comparison, the

behavior of both is shown on log-log coordinates in figure 3.1.3. Overall, sample D is marginally superior in time to fail for a given stress and temperature, except at very

long times at room temperature (296 K) where they are nearly equal. Since sample D had a higher melt index than did A, one main difference in character may be that sample D was more crystalline, which could account for the observed overall lifetime superiority. The

higher molecular weight of sample A is reflected by the fact that for a given condition of

13
TIME

TO

FAIL

(s)

stress and temperature, sample D did not draw to the same extent as A before fracture

occurred.

Figure 3.1.2 - Time to fail versus applied stress in uniaxial extension for a high

density linear polyethylene (sample D) in air.

14
I0

1

1 L o SAMPLE A A - SAMPLE 0

10

7

Figure 3.1.3 - Comparison of time to fail versus applied stress in uniaxial extension

for two samples of high density linear polyethylene in air.

Another difference which should be noted is that both samples A and B used in the

earlier studies had experienced rather long shelf lives prior to molding (~13 years),

whereas sample D had just recently been manufactured. Therefore, it was not certain to what extent, if any, shelf life may be influential with respect to the properties under

investigation. Based upon the results indicated in figure 3.1.3, it would appear that age

probably had little, if any, effect, and that the primary difference was one of crystallinity and/or molecular weight. Nonetheless, in what follows, we have chosen sample D for further discussion since the experimental work carried out on this sample was more

extensive than for sample A.

For the case of creep to failure in uniaxial extension, the data in figure 3.1.2 have

been replotted on log-log coordinates in figure 3.1.4. The initial observation is that at

all four temperatures examined the data fall on a line that is very nearly linear. Two

15
exceptions can be noted, (1) at very high stresses, and (2) at room temperature for long

times to fail. With regard to the first exception, it was pointed out in the previous

report [1] that in this load region the creep occurs so rapidly that sufficient heat is generated (of the order of several degrees Celsius) to effectively lower the modulus of

the specimen resulting in earlier failure. At long times to fail at room temperature,

there appears to be a slight tendency for the curve to bend over. A similar trend was found for the highly crystalline samples B and C even when plotted on semilog paper.

Unfortunately, there are not yet enough data points at long times to tell whether a further erosion from linearity will occur as the loading becomes smaller and smaller. The

reader is reminded that 108 seconds represents about 3.2 years. The two points with

arrows correspond to specimens which have yet to fail.

Figure 3.1.4 - Data of figure 3.1.2 replotted on log-log coordinates. Arrows indicate

specimens still under test.

16
In addition to the high degree of linearity of the different curves in figure 3.1.4,

a second major feature is that they are parallel to one another as well. In Section 2,

the point was made, which is central to the issue of predicting lifetime behavior based on accelerated testing, is the assumption that test data obtained at different temperatures

can properly be superposed to provide one master curve. This means that the functional

form of the time to fail versus applied stress curves must not be substantively different

at different temperatures. The linear and parallel nature of the curves shown in figure

3.1.4 suggest that on semilog paper the data for sample D might obey an exponential

equation of the form

C Bo (3.1.1)

t^ = Ae

where t f

= time to fail;

o' = applied stress;

A and B are material constants;

C = an exponent to be determined.

Using equation (3.1.1), we have fit all four curves shown for sample D in figure 3.1.3,

and then, by shifting the curves obtained at each elevated temperature both vertically and

horizontally, all the data now appear on one master curve which is shown in figure 3.1.5.

The shift factors are shown by the crosses near the left ordinate axis, or time axis, in

the figure, and have as their reference point zero stress and 1 second. With the ex- ception of the high load range where the temperature rise during creep becomes a problem,

the superposition is remarkably good over nearly 8 decades in time. The solid line

represents a fit of the room temperature data to the equation

0.36

-3.6a

tg = (2.1 x 10 ll4 )e (sec.

)

where now the exponent of a represents a best fit to the room temperature data. In terms

of accelerated testing, say, for example, a test time of about 3 hours, an almost 6 decade

saving in time required to conduct the tests has been accomplished by raising the temperature from room temperature to 370 K (97 °C).

Once a set of time shift factors (a.[T]) has been determined, it is customary to plot

the natural logarithm of the shift factors versus reciprocal absolute temperature (1/T).

If the resulting line obeys a linear relationship, the behavior is said to be "Arrhenius"

and from the slope of the line an apparent activation energy can be calculated from the

expression

AH

a

t [T] = Be RT (3.1.2)

where a.[T] is the time shift factor, B is a constant, R is the gas constant (8.31 kJ/mole),

T is absolute temperature (degrees Kelvin), and aH the apparent activation energy. For sample D such a plot is given in figure 3.1.6 where the shift factors are those employed

in figure 3.1.5. It is clear that for this high density linear polyethylene the solid

circles can be represented quite well by a straight line, the slope of which yields an apparent activation energy of 150 kJ/mole (36 kcal/mole). Here, then, is one instance

where AH is constant over the entire range of temperature covered, which means that

between room temperature and about 100 °C one should be able to predict from figures 3.1.5

and 3.1.6, failure times in uniaxial extension for any arbitrary stress and temperature,

provided, of course, the conditions of specimen preparation are nominally the same.

Since the data in figure 3.1.2 were obtained from constant load experiments in creep,

the same information should be available by examination of the respective creep curves. Therefore, creep data were obtained for a number of specimens at different temperatures

at one level of stress in uniaxial extension (1.23 x 10 7 N/m 2 or 125 kgf/cm 2 ), and these

results are shown in figure 3.1.7. In a similar fashion to that employed in figure 3.1.5

17
TIME

TO

FAIL

(s)

Figure 3.1.5 - Master curve

figure 3.1.2 to 296 K (23

for sample D in air obtained

°C).

by superposing the data in

18
1000/T( K)

Figure 3.1.6 - Arrhenius plot of the shift factors determined from the master curve shown

in figure 3.1.5 and from creep data obtained for sample D in air.

to superpose the data, the creep curves in figure 3.1.7 were superposed onto the room tem- perature (296 K) creep curve to yield the master curve given in figure 3.1.8. Again, the

horizontal shift factors employed are those indicated along the time axis, and if plotted in figure 3.1.6 (crosses) yield the same straight line represented by the circles. Since both

the lines shown are essentially superposed on one another, the same apparent activation

energy is obtained by either method.

We turn now to a discussion of sample E, the branched ethylene-hexene copolymer. From

the point of view of hazardous materials transportation, this polymer is of interest since

it is recommended by the manufacturer as the base resin for large free-standing containers.

Results of log time to fail versus stress in uniaxial extension are given for sample E in

figure 3.1.9. Now, however, in the case of the branched polymer, by failure is meant the

time to reach the onset of necking rather than actual fracture, since most specimens colddrew, but did not fracture prior to removing the load. This difference in behavior will be

shown presently for specimens tested to actual fracture at room temperature.

From a comparison of figures 3.1.2 and 3.1.9, it can be seen that for stresses above

about 150 kg/cm 2

, the branched polymer fails (necks) sooner at a given stress than does

the linear polymer, but below 150 kg/cm 2 it rapidly becomes superior. This same behavior

was observed previously [1] for sample A and the two commercial container materials which

were branched. Since sample E was somewhat lower in density than D and contained branching,

it almost certainly had lower crystallinity which could account for its poorer performance

at very large stresses in uniaxial extension.

19
I

i oLio 1

SAMPLE D a = 123 » 10

7 N/m2 O - 296K A - 313K

- 318K O - 325K

- 330K

O

+

o

10 2

TIME (*)

t o

0

°

<p

I

10 5

10 6

Figure 3.1.7 - Compliance J(t) versus time from creep to failure experiments on sample D. Arrows indicate the onset of cold-drawing.

|

20

inn

i

i

— i

run
100 cr

05O 10

1.0

SAMPLE D

cr = 1.23 x 10

7 N/m2

o - 296 K a - 313 K - 318 K O- 325 K + 330 K

„ %

+o

*

-

A OO

CPAO CPiAO O' A i CP P + A +

313 K 318 K J I L

325 K 330 K _Ll

10

' 10

' w

TIME (s)

10 : 10

l

Figure 3.1.8 - Superposition of creep data shown for sample D in figure 3.1.7. Shift

factors are indicated along the time axis.
TIME

TO

FAIL

(s)

Figure 3.1.9 - Time to fail versus applied stress in uniaxial extension for an ethylenehexene copolymer (sample E) in air.

>:

22
The behavior of the copolymer perhaps can better be seen by replotting the data of

figure 3.19 on log-log coordinates as in figure 3.1.10. Unlike sample D for which the data

can be represented quite well by straight lines, such is not the case here for sample E. Even on log-log coordinates there is a sharp curvature upwards at all three temperatures,

and it is clear that the data do not obey so simple a relationship as that for the linear

polymer. In figure 3.1.10, the open symbols now correspond to the observation of necking or cracking, whereas the filled symbols represent actual fracture. At room temperature, ob- serve that for a stress above about 200 kg/cm 2

, the specimens necked and then fractured

almost simultaneously. However, below 200 kg/cm 2

, specimens first necked and then cold-drew

to an extension ratio (x) of about 7 or 8. It is interesting that the region of stress

separating neck formation followed by fracture from that of neck formation followed by colddrawing is exceedingly narrow. This transition zone also happens to coincide with the point

at which the downward curvature appears in the failure curve. The several specimens which

cold-drew without breaking were then maintained under the same load until fracture occurred.

Note that for these specimens, fracture occurred approximately two decades in time later

than the point of global instability (necking). Although the ultimate extension ratio

reached before fracture depended upon the applied stress, it was generally between 12 and

14 .

It is again useful, from the point of view of accelerated testing, to superpose the

data in figure 3.1.9, in order to obtain a master curve from which the shift factors can be

used to determine an apparent activation energy. However, after several attempts, it was discovered that, although superposition was possible which was reasonable in appearance (see

figure 3.1.11), no suitable set of shift factors could be found which yielded a straight

line on an Arrhenius plot. Therefore, as in the case of sample D, a series of creep curves

were obtained for the branched polymer at different temperatures and same level of applied

stress. Several results are shown in figure 3.1.12, where the data have been superposed

onto the room temperature creep curve, and the horizontal shift factors are given on the

Arrhenius plot in figure 3.1.13. It is immediately apparent that even the creep results do

not yield a straight line and that this polymer, unlike sample D, cannot be characterized by

a unique AH. Nevertheless, in figure 3.1.11, we have superposed the 330 K data shown in

figure 3.1.9 onto the room temperature curve using the shift factor determined from the

creep experiment. Since no creep data were available at 373 K (97 °C), those points were

shifted in order to obtain the best fit in appearance, and for that reason, the shift factor

shown in figure 3.1.13 is indicated by the solid circle. Unfortunately, it was also discovered that the very same data could be shifted to fit essentially the same curve equally

well in appearance but with different shift factors. Therefore, an important point to be

made here with respect to accelerated testing is that some care must be exercised in determining a master curve, especially for polymers such as sample D for which the time to fail versus applied stress curve becomes very steep, in particular at high temperature and/or

long failure times. As we have seen before, this ambiguity emphasizes the necessity to have

at least some long-time failure results in order to have a greater degree of confidence in

predicting long time behavior.

3.2 Stress Cracking of Polyethylene in Unaxial Extension

One of the main objectives of this report is to examine the stress-cracking behavior of

polyethylene under a variety of imposed conditions, a primary goal being to provide information which will be useful in the establishment of an improved stress-cracking test pertinent to the transportation of hazardous materials in large reusable plastic containers.

From the discussion in Section 2, it is apparent that a large body of literature already

exists on the subject, and in some instances attempts have been made to establish master

design curves from which lifetime in the presence of a specified stress-cracking agent can

be predicted for fairly arbitrary conditions of stress and temperature. Sections 3.2 and

3.3 will summarize the results of a number of experiments carried out in our laboratory,

where both uniaxial and biaxial stresses have been employed in the presence of several known

stress-cracking agents. Section 3.2 will be concerned with stress-cracking behavior under

conditions of a uniaxially applied stress; experiments involving biaxial stresses will be

discussed in Section 3.3.

23
Figure 3.1.10 - Data from figure 3.1.9 replotted on lon-log coordinates. Open circles

correspond to the time required to form a neck, filled circles to the time of actual

fracture. Arrows indicate specimens still under test.

24
Figure 3.1.11 - Master curve for

figure 3.1.9 to 296 K (23 °C).

sample E in air obtained by superposing the data from

Symbols the same as in figure 3.1.10.

25
TIME (s)

Figure 3.1.12 - Superposition of creep compliance data for sample E in air. Shift factors

are indicated along the time axis.

26
15

SAMPLE E

1000/T(K)

Figure 3.1.13 - Arrhenius plot of the shift factors determined from the master curve for

sample E shown in figure 3.1.11.

Experiments similar to those described in Section 3.1 were carried out on the same two polyethylene samples (D and E), but in the presence of a stress-cracking agent. The experimental procedures employed here were essentially the same in principle as those described in

reference [9], and a diagram of the apparatus used is shown in figure 3.2.1. The desired

temperature was maintained by circulating a mixture of water and ethylene glycol from a controlled temperature bath through a double walled pyrex glass condenser surrounding the specimen.

The bottom was sealed with a rubber stopper to prevent leakage and the top covered

with aluminum foil so that only a small opening was left to allow the support filament to move unimpeded. All of the experimental data were obtained with the specimen submerged in

the stress-cracking agent rather than in the vapor, and the temperature was measured by

placing a thermometer in the liquid next to the center of the specimens. It was found typically that a temperature gradient of from 0.1 to 0.2 K occurred between the bottom and

top clamp. Fiducial marks were placed on the straight portion of the specimen in order to make creep measurements. It is worthy of note that most types of inks, one exception being

India ink, and other types of marking substances tested turned out to be stress-cracking

agents themselves, and even in air the specimens nearly always failed at one of the fiducial

marks.

27
Figure 3.2.1 - Schematic diagram of apparatus for determining stress-crack resistance in

uniaxial extension.

28
Shown in figure 3.2.2 are results for sample D in the presence of three known stresscracking agents. Nonylphenoxypoly(ethyleneoxy)ethanol was chosen since it is generally

regarded to be one of the most active of the so-called surface active stress-cracking

agents and is in widespread use in present stress-cracking tests, including ASTM tests. Both n-hexane (C 6 H 14 ) and n-dodecane (C 12 H 2 6) are normal alkane solvents known to swell

polyethylene. In all cases, the specimens were presoaked in the stress-cracking agent for

at least 16 hours prior to loading. It has been reported in the literature that in the case

of nonylphenoxypoly(ethyleneoxy)ethanol , presoaking the specimen has little, or no, effect

on the lifetime behavior, as was found to be the case here as well. However, for the two solvents there was a pronounced difference in lifetime at the relatively large stresses

depending upon the time of presoaking (up to times of about 16 hours). The same result was also found from weight gain measurements where after about 16 hours (overnight) the specimens

apparently absorbed very little additional solvent.

The solid lines in figure 3.3.2 correspond to the same solid lines drawn through the

data points shown in figure 3.1.2 for sample D at 296 K and 330 K in air. Observe first

that at room temperature, the presence of nonyl phenoxypoly(ethyleneoxy )ethanol has little or

no effect on lifetime at stresses above about 150 kgf/cm 2

; but below that value the lifetime

becomes increasingly shorter until reaching a point at low stresses where the data points

are displaced a nearly constant amount from the solid line. At the same time, the effect at 330 K (57 °C) is an approximately constant shift over the entire range of stresses tested,

the amount of shift being about the same as that for the long times at room temperature.

By comparison, the two solvents which are readily absorbed into polyethylene accelerate

failure much more rapidly. At a stress of 75 kgf/cm 2

, the presence of hexane or dodecane

has the same influence on time to fail at room temperature as it does running the same test

in air at 330 K (57 °C). Also, at 50 kgf/cm 2

, the presence of solvent at 330 K has decreased the lifetime by about four decades in time from that at room temperature in air,

which demonstrates the highly disastrous effect of having both solvent and elevated temperature present at the same time the specimen is under stress.

Another point to be made here is the following. Although the presence of nonyl phenoxypoly(ethyleneoxy)ethanol at 330 K decreased the lifetime at a stress of 100 kgf/cm 2 by a factor of nearly 1000 from that at 296 K in air, the overall effect is primarily one of

temperature and not one of stress-cracking agent. Only a modest decrease in time to fail

occurred due to the presence of nonylphenoxpoly(ethyleneoxy)ethanol . However, in the

following section, we shall see that for biaxial stress conditions the effect can be much

more dramatic. This is one reason why an effort is generally made in developing stresscrack tests to incorporate a biaxial stress component in order to enhance failure under

conditions more realistic of in-use conditions.

Now according to the same procedures used in Section 3.1, the data in figure 3.2.2 have

been superposed to yield the result given in figure 3.2.3. In figure 3.2.3 room temperature

(296 K) has been chosen as the reference temperature and the shift factor for the 330 K (57

°C) solid line is the same as the employed in figure 3.1.5. The data for the three stresscracking agents at 330 K have, in turn, been shifted correspondingly the same amount. As is evident, the stress-crack results obtained at 330 K correlate very well with the stresscrack behavior observed at room temperature. We point out that the shifting procedure

followed in figure 3.2.1 is somewhat different from that employed in a number of previous

studies, for example, references [11] or [17]. In those investigations, the superposition

was carried out with respect to a reference state specified not by temperature alone, but

also by adverse environment. One advantage in presenting the data as shown in figure 3.2.3

is that the relative severity of a given stress-cracking agent can readily be judged not only against other stress-cracking agents, but against the maximum lifetime possible in the

absence of an adverse environment. It is interesting to observe from figure 3.2.3 that, to the extent lifetime behavior at room temperature can be predicted realistically based on a superposition of accelerated test data obtained at higher temperatures (330 K), even at very

small stresses one must conclude that the presence of solvent will cause failure in a finite time.

29
Figure 3.2.2 - Time to fail versus applied stress in uniaxial extension for sample D in various stress-cracking agents. The solid lines represent the data at 296 K and 330 K from figure 3.1.2.

30
. STRESS (kgf/cm 2

)

Figure 3.2.3 - Master curve for the stress-crack behavior of sample D at 296 K. The solid

line represents the data from figure 3.1.5 and the shift factor for the 330 K data is

indicated by the cross.

31
From the standpoint of a qualification test a more useful procedure may be to superpose

the data not to room temperature as a reference temperature, but to a temperature corresponding approximately to the highest temperature expected in use. Such a procedure has been

followed by Fulmer [17] in bottle tests where the reference temperature was chosen to be 333

K (60 °C). In figure 3.2.4, the same data presented in figure 3.2.2 are also shown

superposed where the reference temperature is now 330 K (57 °C). The overall behavior is essentially the same as that seen in figure 3.2.3 except that the times to fail are now substantially reduced.

Analogous results to those shown in the previous three figures are presented in figures

3.2.5, 3.2.6, and 3.2.7 for Sample E, the branched copolymer. Whereas, in the case of the

linear polymer (sample D), the presence of nonylphenoxypoly(ethyleneoxy)ethanol had a significant effect on lifetime in uniaxial extension at room temperature, based on only limited

data there is no evidence that the nonylphenoxpoly(ethyleneoxy)ethanol had any influence on the branched copolymer at 296 K. It was pointed out in Section 2 that branching greatly

improves stress-crack resistance. That this is true, can readily be seen by comparing the

superposed data shown in figures 3.2.6 and 3.2.3. In fact, below about 150 kgf/cm 2

, it is apparent that the branched copolymer in nonylphenoxypoly(ethyleneoxy)ethanol is superior

even to the linear polymer in air.

As was true for sample D, the decrease in time to fail is more severe for the branched

copolymer in the presence of solvent than in nonylphenoxypoly(ethyleneoxy)ethanol and, as was also true in the earlier example, the net effect is primarily one simply of shifting the

entire curve from that expected in air. Note however that in figure 3.2.6, the curve for

n-dodecane (C 12 H 2 6) no longer coincides with that for n-hexane (C 6 H 14 ), as was the case in figure 3.2.3. Due to its volatility, no experiments were carried out at 330 K (57 °C) using

n-hexane, and no data are available for n-dodecane at 296 K. Nonetheless, it would appear

that the molecular weight of the solvent may be an important parameter in determining the

stress-crack behavior of the branched polymer. Similar behavior occurs in the presence of

silicon oils of different molecular weights [20], where the lower the molecular weight IS (higher viscosity), the more destructive the effect is.

3.3 Stress-Cracking of Polyethylene Subject to Biaxial Stress

In the two preceding subsections, experiments have been described in which specimens of

polyethylene were stressed uniaxial ly at constant load in the presence of various stresscracking agents. In the several paragraphs which follow, we shall further examine the

stress-crack behavior of one type of polyethylene when the applied stress is essentially

biaxial, as is more likely to be the case under actual in-use conditions. The method of

investigation chosen here is very similar in principle to that employed by Rivlin and

Saunders [27] to study the deformation of rubber. A diagram of the apparatus used is given

in figure 3.3.1. The specimen, in the shape of a flat sheet or film, is clamped between a top plate and base piece. In the present case, a 15.25 centimeter diameter hole (6 inches) was cut in the top plate and the top and bottom edges of the hole perimeter rounded and

tapered as shown in order to minimize curvature and eliminate sharp edges at the point of

contact of the sheet with the top plate. An o-ring located very close to the clamping point

was inserted to prevent loss of gas around the edge when the vessel was pressurized. The

gas inlet was connected via a pressure reducing valve either to a cylinder of nitrogen gas

or to an air line. Once the specimen was clamped in place, the system was pressurized and

maintained at a constant pressure until failure occurred. In the present experiment, failure was detected generally either by sound or more often by the feel of escaping gas. A more refined version of this test method would require a highly sensitive pressure indicator

installed in the gas line which, upon a slight decrease in pressure, would trigger a switch

to shut-off the gas supply and record the time to failure. In experiments where a stresscracking agent was present, failure was easily detected by the appearance of liquid or bubbles on the outer surface of the specimen. Initially a small amount of stress-cracking

agent was placed in the bottom of the pressure reservior and the specimen then clamped in place. Once pressurized, the entire assembly was immediately inverted to insure that the

stress-cracking agent remained near the center of the test specimen (at the lowest point of

the bubble) and did not contact the specimen near the o-ring seal. In cases where liquid was allowed to contact the specimen at or near the clamping point, failure often occurred

there rather than at the center of the bubble, an indication that the maximum stress may

32
SAMPLE D

0 50 100 150 200 250 300 350

STRESS (kgf/cm 2

)

Figure 3.2.4 - Master curve for the stress-crack behavior of sample D at 330 K. The

reference point for the room temperature shift factor indicated by the cross is the

upper left corner of the figure.

33
TIME

TO

FAIL

(s)

SAMPLE E

10 8

106 -

10 4 _

10

2

10 0

0 50 100 150 200 250 300 350

STRESS (kgf/cm 2

)

Figure 3.2.5 - agents. The

Time to fail versus applied stress for sample E in various stress-cracking

solid lines represent the data at 296 K and 330 K from figure 3.1.9.

34
Figure 3.2.6 - Master curve for the stress-crack behavior of sample E at 296 K. The

shift factor for the 330 K data is indicated by the cross.

35
SAMPLE E

Figure 3.2.7 - Master curve for the stress-crack behavior of sample E at 330 K. The

reference point for the room temperature shift factor indicated by the cross is the

upper left corner of the figure.

36
POLE

Figure 3.3.1 - Schematic diagram of apparatus for determining stress-crack behavior under

biaxial stress conditions.

occur in the region near the point of clamping. Even in air, some difficulty was encountered in obtaining times to fail greater than about 10 6 seconds, since at the lower pressures

nearly all of the specimens tested failed at the o-ring seal in times shorter than anticipated based on data obtained at the higher pressures. Since the specimens used in the

present experiments were thin films, it is also possible that a small amount of unknown

chemical leached out of the rubber o-ring which itself acted as a stress-cracking agent.

For the geometry shown in figure 3.3.1, an equivalent, or radial, stress can be

calculated at the pole of the bubble from the expression

Prx 2 T = “2F" (3.3.1)

Where

t = radial stress component;

P = internal pressure:

r = radius of curvature of the bubble at the pole;

h = initial thickness of the sheet; and

x = extension ratio given by 1+e where e is the strain.

Equation (3.3.1) is valid only in the vicinity of the pole; therefore, when a stress-cracking

agent is present, inverting the apparatus during testing is important in order to maximize

the expectancy that failure will occur in the polar region. It can also be seen from equation (3.3.1) that in order to convert from internal pressure to an equivalent radial stress

requires a knowledge of the radius of curvature as a function of internal pressure for each

specimen thickness used.

37
In the present study the specimens were all cut from one sheet of high density polyethylene film obtained from a commercial source. This polymer, designated hereafter as sample F, had a melt index of 0.75 as determined by the manufacturer (ASTM D1238), and

therefore, had the same value as that for sample D discussed earlier. The thickness of the

film was approximately 0.025 mm (.001 inch) and as supplied was in the form of a biaxial ly

oriented blown film with a blow-up ratio of 4:1.

Because the film was biaxial ly oriented, several specimens were cut parallel, perpendicular, and at 45° to the extrusion direction and tests made first in uniaxial extension. The results are presented in figure 3.3.2 and are drawn to the same log-log scale

employed in figure 3.1.3 which compares samples A and D. Little or no difference in time to fail occurred for the specimens cut parallel and perpendicular to the extrusion direction,

on indication that this particular film may be nearly equibiaxial ly oriented. The specimens

cut at 45° to the extrusion direction did fail 2 or 3 times sooner than did the specimens

cut along the orientation directions. When compared to the behavior of samples A and D in figure 3.1.3, the film actually is superior in terms of time to fail at stresses below about

150 kgf/cm 2. Whether or not this behavior reflects the orientation present or possibly

differences in polymer composition between sample D and F is uncertain, but orientation may

well be the predominent factor.

For the film specimens subjected to biaxial loading according to the procedure diagramed in figure 3.3.1, the experimental data are plotted in figure 3.3.3 (circles and

triangles). In air, there appears two well defined regions for the specimens tested at room temperature. At the higher end of the pressure scale, the slope of the failure curve is quite steep and essentially linear on log-log coordinates. In this region, the film failed

ductilly, that is, cold-drawing perpendicular to the extrusion direction occurred prior to the actual failure which was manifested in the form of cracks parallel to the extrusion direction. In the range of pressure below the break in the curve at times of about 104 seconds, there was no evidence of cold-drawing, but rather the appearance of crazes which

eventually developed into a crack failure. The pressure at which the break in the curve

occurs, and perhaps the fact that it occurs at all, will no doubt depend both on the amount

of orientation initially present in the film and on the initial specimen thickness.

In the presence of a stress-cracking agent such as nonylphenoxypoly(ethyleneoxy)ethanol

(solid circles in figure 3.3.3), the break in the curve becomes even more pronounced and occurs somewhat earlier in time (higher pressures). Interestingly enough, above a pressure

of about 2 x 104 pascals, the nonylphenoxypoly(ethyleneoxy)ethanol no longer influences the

time to fail, possibly because at times this short, the process becomes diffusion controlled

At lower pressures, the effect is much more dramatic. For example at a pressure of 0.86 x 104 pascals (1.25 psi), the time to fail is already shortened from that value in air by a factor of 40, and at 0.7 x 104 pascals (1.0 psi), the reduction is several decades. Again,

however, it should be pointed out that these differences may well depend on both the extent

of initial orientation and specimen thickness.

At 330 K, a similar behavior can be noted except that the time scale is significantly

reduced. The shape of the curve in the case of the stress-cracking agent (solid triangles

in figure 3.3.3) is the same overall as at 296 K. However, in air, the well defined break

in the curve which occurs at 296 K does not appear at 330 K, at least within the range of

pressures examined. Since in the presence of stress-cracking agent the shape of the failure

curve is the same at both temperatures, the data can be superposed in a similar manner to that done earlier in Section 3.1 or that done by Fulmer [17] in his bottle experiment. The

result is that if the natural logarithm of the shift factor is plotted versus the reciprocal

absolute temperature, an activation energy of only about 43 kJ/mole (= 10 kcal/mole) is obtained, a value which is very low compared to that reported by Fulmer for blow molded

bottles of branched polyethylene (= 22 kcal/mole). However, in the present case, there is insufficient data at other temperatures in order to place a very high degree of confidence

in the calculated activation energy.

Now, for several of the specimens tested in air at room temperature, creep measurements

were made as a function of time and these experiments are summarized in figure 3.3.4. The arrows indicate the point at which a crack opened allowing gas to escape. The strain, or elongation, was measured by placing fiducial marks of India ink about 1.5 centimeters on either side of the pole (center of the film), and marks were placed both parallel and

38
STRESS |kgf/cm 2

)

Figure 3.3.2 - Time to fail versus applied stress in uniaxial extension for a high density

linear polyethylene film. Open circles correspond to specimens cut parallel to the

extrusion direction, filled circles perpendicular, and the triangles 45° to the extrusion

direction.

39
PSI

1 2 3 4 5 6

INTERNAL PRESSURE « 10“ 4(pascals)

Figure 3.3.3 - Time to fail versus internal pressure for polyethylene film and bottles

tested to failure under biaxial stress conditions in air and in igepal. Open symbols

air, filled symbols igepal. The pressure scale for the bottles is 20 times that

shown.

40
LO

CD

CO

o

CMo

N0I1V9N013 %

41

TIME

(s)

Figure

3.3.4.

Creep

to

Failure

Data

for

Polyethylene

Film

under

Biaxial

Test

Conditions,

Internal

Pressures

Are

as

Indicated

and

the

Arrows

Correspond

to

the

Point

of

Failure.
perpendicular to the extrusion direction. The position of the marks was then followed in time with a traveling microscope. In most all cases, a difference in percent elongation was observed between the two directions, the percent elongation generally being greater in the

direction parallel to the extrusion direction. It was also true that the amount of difference in percent elongation diminished with decreased pressure, as might be expected. In figure 3.3.4, the points shown represent the average percent elongation for the two directions and, except for the two specimens for which bars are shown, the extremes were at most

only a small amount greater than the size of the data points themselves.

One interesting aspect to the creep data is that crazes first became visible at a strain of about 4-5 percent. This means that on a microscopic scale they must occur at even smaller strain. Also, it can be noted that as a general rule the crazes first appeared in

time by a factor of from 10 to 100 earlier than the actual failure. Unfortunately, no creep

data were obtained in the presence of nonyl phenoxypoly(ethyl eneoxy )ethanol due to the impracticality of making creep measurements with the apparatus inverted. The question then

remains as to whether the effect of the stress-cracking agent is principly one of inducing

the formation of crazes at an earlier stage (smaller strain), or is one of greatly reducing

the time to failure after the crazes form. From the point of view of stress-cracking, one important consideration in container design, where reuse is intended, may be to reduce the

maximum allowable strain expected under conditions where the container remains filled over a long period of time, to a level below that where crazes form easily.

Finallywith respect to this film, it should be noted that in the presence of the

solvents, n-hexane or n-dodecane, the permeation rate through the film was so great under

even the smallest pressures applied that no failure occurred before all the liquid had

passed through the film.

In another set of experiments, blow molded bottles (Boston Round) having a capacity of

227 cm 3 (8 oz.) and produced from the same polymer as sample D were tested to failure at 296 K containing either air or nonyl phenoxypoly(ethyl eneoxy )ethanol solution. These results

are also presented in figure 3.3.3 (open and solid squares), where now a factor of 20 scale

change has been applied along the pressure axis. Unlike the film, no break appears in the

failure curve for the bottles in air at 296 K, whereas in the stress-cracking agent a similar break does occur which is also reminiscent of the behavior reported by Fulmer [17] for

blow molded detergent bottles tested to failure in Igepal .

Some interesting comparisons can be made between the behavior exhibited by the bottles

and that found from the bubble test on film. In the film, the location of the failure was more or less random with respect to the pole position; however, in the bottle test all the

bottle failed at the same point regardless of whether or not the failure was ductile, or whether the experiment was done in air or in stress-cracking agent. An example of this is shown in figure 3.3.5 for several bottles of sample D tested to failure in air at different

pressures. In each case the final crack opening occurred at the same point, near the midpoint of the bottle and along a seam parting line. This was true whether the bottle first

cold-drew preceding failure (in a circumferential direction), or cracking occurred with no prior cold-drawing visible. Thus, a pressurized container test of this type may be

advantageous from the point of view of container design, but from a material properties

aspect is less so unless the container can be made highly uniform overall.

3.4 Additivity of Damage in Polyethylene in the Presence of an Adverse Environment

In Section 7.5 of reference [1], the concept of additivity of damage was introduced.

For cases in which the environment is relatively inert, such as air, it was shown that the

cumulative damage could be expressed simply as a sum of fractional times to failure, providing the stress and temperature histories were known. Further, it was demonstrated that

information gained from experiments involving a constant rate of loading in uniaxial ex- tension could be used to predict the time to fail for conditions of static loading. In transportation, where container reuse may be an issue, a cumulative damage criterion is of some interest; for, as it was suggested previously, if the average environment which a container experiences during normal in-use conditions can be reasonably estimated, then

a qualification test for reuse based on a cumulative damage principle might be appropriate.

42
i

Figure 3.3.5 - Examples of polyethylene bottle failure in air and in igepal. (a) in air,

(b) in nonyl phenoxy poly (ethyl eneoxy) ethanol . Numbers shown indicate internal bottle

pressure in psi.

M»lfl

43
However, in the previous work both the temperature and environment (air) were held fixed,

and whether or not additivity of damage remains a valid concept in the presence of an adverse environment, which possibly may change with time as well, is yet to be resolved.

We shall report here the results of experiments in which a series of very simple

stress, temperature, and environmental histories have been applied to specimens of polyethylene. Sample D, the linear high density polyethylene, was chosen since the effect of

stress-cracking agents on this polymer was found to be more dramatic than was the case for

the branched polymer. As will be seen, the behavior of this polymer becomes more difficult

to predict based on additivity of damage when the additional parameter of an adverse environment is added.

In one set of experiments two specimens were heated to 330 K (57 °C) and then loaded in

uniaxial extension to a stress of 75 kgf/cm 2 for a length of time equal to about one half

the expected total time to failure as determined from the 330 K data given in figure 3.1.2.

At that point, the load was removed and a solution of nonyl phenoxypoly(ethyleneoxy )ethanol

in water was added to one specimen and n-dodecane to the other. Both specimens were held

overnight at 330 K in the liquid before the load was reapplied and the time from that point

on to the actual failure was recorded. In this case, the fractional time to failure is

defined simply as time to failure in stress-cracking agent after reloading divided by the

time to failure determined for the same stress from figure 3.2.2. The result was that the

total fractional time for the specimen loaded in air and then in nonyl phenoxypoly-

(ethyl eneoxy)ethanol was 0.93, and correspondingly for the specimen first loaded in air and

then in n-dodecane, it was 1.2. According to the additivity of damage criterion employed

earlier, the fractional times to failure should sum to unity at the time of actual failure.

For a second set of specimens the reverse procedure was followed, that is the initial

portion of the experiment was done in stress-cracking agent and the final portion in air.

However, one important difference in the later set of experiments was that after the initial

loading period in stress-cracking agent, each specimen was cleaned as thoroughly as possible

and then allowed to dry overnight at 330 K before being reloaded. In the case of the nonylphenoxypoly(ethyl eneoxy)ethanol , the specimen was scrubbed in distilled water with a medium

bristle brush prior to drying, whereas the specimen exposed to n-docecane was first washed

in a mild soap and water solution and then rinsed in distilled water. For these specimens,

the total fractional time to failure with the nonyl phenoxypoly(ethyl eneoxy )ethanol was 0.71

and 0.69 for the n-dodecane.

Since only one specimen was tested for each set of conditions imposed, no meaningful

statistical analysis is possible. However, for the first two specimens with total fractional times of 0.93 and 1.2, their departure from unity must be considered within expectation when account is taken of the scatter typical of the data points shown in figures 3.1.3

(330 K) and 3.2.2 (330 K). On the other hand, for the latter two specimens, the fractional

times to failure observed during the second portion of the experiment are quite short (0.21

for the specimen initially in nonyl phenoxypoly(ethyl eneoxy )ethanol and 0.19 for that initially in n-dodecane compared to the expected value of 0.5). In fact, based on the results « I in figure 3.2.2 (330 K), one would conclude that removing the specimen from the nonyl phenoxypoly(ethyl eneoxy jethanol and then washing and drying it had little, if any, effect since the

total time to failure was about what it would have been if left in the nonyl phenoxypoly-

(ethyl eneoxy)ethanol from the outset. For the specimen tested initially in n-docecane, a similar statement applies except that the effect is less severe. Therefore, we can infer,

albeit based on rather limited data, that it may be difficult, if at all possible, to remove completely material adsorbed or absorbed while under stress simply by washing and drying the

specimen. Essentially this same conclusion was reached in reference [9], where particular

reference was made with respect to the presence of nonyl phenoxypoly(ethyl eneoxy )ethanol . As

a result, it may be true that an additivity of damage criteria will at best provide a bound,

either upper or lower, to lifetime prediction whenever environmental factors in addition to temperature and stress are important. Of course, as a predictive tool, or in terms of

qualification procedures, the knowledge of an upper bound is of little, or no, practical

value.

44
Another area of concern, where the reuse of shipping containers may be involved, is

the possibility that the adverse environment may change as a function of time. One can

envisage situations in which a container designated for reuse is exposed to more than one

type of lading, either accidentally or otherwise, and the effects due to stress-cracking

nay be compounded, thus further eroding its useful service life. As an example of this

possibility, two specimens of sample D were presoaked, one in nonyl phenoxypoly(ethyl eneoxy)-

ethanol and the other in n-dodecane, for a period of one day at 330 K. Each specimen was

then loaded to 75 kgf/cm 2 for a length of time corresponding to one half the expected time

to failure. The load was then removed and both specimens cleaned as before. They were

then presoaked overnight at 330 K, but in the reverse sense from before (n-dodecane and

nonyl phenoxypoly(ethyl eneoxy)ethanol , respectively), and again reloaded at the same stress

jntil failure occurred. The resulting total fractional times to failure were respectively

1.3 and 1.1. Therefore, it would appear that for the two combinations tested no further

degradation in mechanical integrity occurred over and above that expected for either

stress-cracking agent separately. This is not to say that for other combinations of

1,1 surface active stress-cracking agents or solvents an effect will not appear.

4. Summary and Recommendations

In Part I of the previous Final Report [1], the specifications and qualification

tests contained in paragraph 178.19 of 49 CFR (Specification 34) were reviewed, the result

Deing that several areas were identified where, insofar as mechanical properties are

important, the present test requirements were judged to be inadequate. Specific recommendations were made to upgrade several of the qualification tests, primarily by redesigning

them to provide more quantitative information. Among the several recommendations made

were: (1) the necessity for container manufacturers and/or shippers to carry out performance

tests on specimens taken from the finished container; (2) incorporating into qualification

tests the idea of a test to failure; and (3) the establishment of retest procedures. Also

noted was one important area not addressed in Specification 34, namely that of stresscracki ng

.

The principle focus of the present report has been directed toward the problem of

stress-cracking in polyethylene, and an effort has been made to delineate the more important

parameters influencing stress-cracking and to provide an indication of the directions

previously taken to establish test methodology for determining stress-crack resistance.

Section 1 dealt with a brief review of stress-cracking tests presently in common use, and

these were found to be primarily the tests included by the ASTM as standard test methods

or recommended practices. Section 2 was intended to provide the DOT Office of Hazardous

Materials Operations with general background information on the subject of stress cracking,

and to point out several of the more important parameters which influence stress-crack

resistance. Section 3 presented a summary of experimental work carried out in the NBS

Polymers Division during the contract period. This work included time to fail experiments

carried out under both uniaxial and biaxial stress conditions in both air and known

stress-cracking agents.

In the final section several recommendations are made covering two areas not presently

addressed in 49 CFR 178.19, namely stress-cracking and reuse.

As already noted, DOT regulations do not currently require a test for stress-crack

resistance. From discussions with industry personnel and a survey of the recent technical

literature (see Section 2), it is clear that stress-cracking remains a problem of concern

to industry. Furthermore, based on existing data available in the technical literature

and on experiments conducted in our own laboratory (see section 3), it appears that the

greatest utility to conducting tests for stress-crack resistance lies in: (1) the selection

of a polymer resin with good stress-crack resistance, and (2) establishing the relative

severity of different products as stress-cracking agents by comparing them to a known

stress-cracking agent. From a practical viewpoint, one of the most obvious and fruitful

approaches to minimizing failures in the field due to stress cracking is to optimize

stress-crack resistance at the outset by proper testing at the time of manufacture.

Therefore, as a qualification procedure, the DOT should require under 49 CFR 178.19-7 a test for stress-crack resistance. Test data should be submitted to the DOT indicating

stress-crack resistance of the container material to a prospective lading and to a known

45
I

stress-cracking agent. Such tests should be carried out at the time of initial production

and whenever any change is made in the material or processing technique.

We interject that the simple knowledge that a substance is a stress-cracking agent

should not necessarily exclude it from shipment in a particular container (a stress-cracking

agent being defined here as any substance when in contact with the polymer causes failure in

a time shorter than that in air, all other factors being equal). If it can be demonstrated

that the failure time in the presence of a particular lading and under relatively severe environmental conditions is much longer than any expected service life or time in transportation, then there would appear to be no strong justification for rejection on that basis

alone. At some point, however, a judgment has to be made as to the suitability of shipping

the product in a given container. The criterion to be invoked in such a decision can only

be made based on a knowledge of industry wide experience and from evaluation of the expected

transportation environment. From the standpoint of the DOT Office of Hazardous Materials

Operations, two valuable aids in this regard should be: (1) the information available through

the submission of test data on stress-crack resistance, and (2) an up-to-date, complete as possible, record of field failures. We note parenthetically that the DOT Hazardous Materials

Incident Report (Item F) makes provision for the reporting of at least 16 different types

of packaging failures. None of these pertain specifically to failure due to stress-cracking,

yet nearly all of them can either be a result of, or be accelerated by, exposure to a lading

which at the same time Is a stress-cracking agent. Cumulative damage to a container caused

by a stress-cracking agent will, in general, not be visible even by careful visual inspection. However, the container may have been sufficiently weakened at some point that it will

fail under conditions which normally would not cause failure. In most Instances, the man in

the field who observes the failure will not be in a position to judge whether or not the

failure was a result of prior internal damage due to stress-cracking. However, in cases where the incidence of field failures with a particular lading, or type of lading, falls

well outside any statistically likely incidence rate for the Industry as a whole, a knowledge

of the severity of a lading, or type of lading, as a stress-cracking agent may be valuable

in re-evaluating the suitability of transporting that lading In the container in question.

Now as to just which stress-cracking test is the proper one for Inclusion In DOT regulations remains open to discussion. In Section 1 of this report, it was observed that

the most often used tests for stress-crack resistance are those currently included in the

ASTM Book of ASTM Standards [4], It was further noted that considerable criticism has been

directed toward several of these tests, and recent discussions with members of the plastic

container Industry further substantiates the general viewpoint that even the most widely

accepted tests are not considered to the wholly satisfactory. Clearly, there remains a need

to develop a performance oriented stress-cracking test which can be carried out In the

laboratory on specimens taken directly from a finished container. Much of the work described

in Section 3 of this report was undertaken to establish a data base from which such a test

can be developed. However, further work Is required in several areas in order to ascertain

which type of test is the most appropriate. For example, it has been suggested [17] that a scaling may be possible between results obtained from tensile tests in uniaxial extension

and comparable data obtained from actual container tests. If true, then the actual testing

can be greatly simplified from that required under conditions of biaxial stresses. However

the evidence that such a scaling exists between the two types of tests is not well documented

In spite of their apparent shortcomings it is recommended that until a new test for

stress-crack resistance becomes available, DOT regulations should include one of the presently accepted test methods for determining stress-crack resistance. As a polymer resin

test, one intended to aid in the selection of a polymer to optimize stress-crack resistance

and for determining the relative severity of various products as stress-cracking agents,

tests such as ASTM 2561-70 and ASTM 2648-70, or a test similar to that described In Section

3.3 of this report appear to be the most appropriate. In order to provide more quantitative

information, a test such as ASTM 2561-70 should be carried out at more than one internal

pressure

.

Where reuse is a primary concern, it is recognized that the requirement of a stresscracking test on a continuing basis to a given production lot of containers may be impractical since: (1) the retest procedures may become sufficiently costly that the entire operation becomes uneconomical, and (2) not all of the containers may be returned to the shipper

at the same time. From the same production lot one container may experience a much shorter

46
time in the field than another, possibly under unknown conditions, so that the difficult

question arises as to just how the containers should be chosen for retesting and at what

intervals they should be retested. Presumably, a deterioration in the performance characteristics of a container (cumulative damage) due to exposure to a lading which is a stresscracking agent, if sufficiently severe, will be reflected upon subsequent testing by an inability to meet the minimum retest specifications. It is, therefore, recommended that

some provision be made in DOT regulations for retesting of containers on a periodic basis.

A retest procedure can be established by choosing a subset of tests contained in 49 CFR

178.19-7 and requiring that containers: (1) be selected at random from a given production

lot that has been returned after some specified number of trips, and (2) be submitted to the

chosen tests. The three most relevent tests would appear to be a drop test, hydrostatic

pressure test, and a static compression test.

We know also from experiments carried out in our laboratory that the application of

only a small stress over a very long period of time can result in failure, and that in the

presence of stress-cracking agent, the amount of deformation (strain) in the container can

be very small at failure (5 percent or less). Therefore, it is important that the dimensions of the container remain constant, or nearly so, over the expected service life of the

container. By placing fiducial marks near the top, center, and bottom of the container, and

on opposite sides, the circumference of the container can easily be monitored periodically

with a tape measure. If a container has been subjected to either a large internal pressure

or compressive force for any length of time, damage done to the container may be reflected

by an irreversible increase in the circumference of the container. Further exposure to even moderate environmental conditions may then be sufficient to cause failure, especially if

the lading is known to be a stress-cracking agent. Periodic inspection of the container

dimensions should provide one relatively easy and non-time consuming method for detecting

potentially hazardous situations. For example, a container which is found to have increased

irreversibly in circumference by, say, 5 percent or more will probably have difficulty

passing the qualification tests required under specification 34. Such a test does not insure that a container will not fail as a result of stress-cracking, or otherwise, since

the failure may occur at the chime pr near a closure, but does provide a means of screening

lots of containers which may have experienced unusually severe conditions resulting in

incipient damage.
5. REFERENCES

1

[1 ] NBSIR 76-11 68 ( R ) , Performance of Plastic Packaging for Hazardous Materials Transportation Part I. Mechanical Properties , J. M. Crissman, C. M. Guttman, and

L. J. Zapas, Report prepared for the U.S. Department of Transportation, Office

of Hazardous Materials Operations, Washington, D.C. 20590, under Contract Number

DOT AS-50074 , October 1976.

[2] NBSIR 76-1 1 63(R) , Performance of Plastic Packaging for Hazardous Materials

Transporation Part II. Permeation , J. D. Barnes and G. M. Martin. Report prepared

for the U.S. Department of Transportation, Office of Hazardous Materials Operations,

Washington, D.C. 20590, under Contract Number DOT AS-50074, October 1976.

[3] Code of Federal Regulations 49 CFR, Transportation Parts 100 to 199 (1974), U.S. Government Printing Office, Washington, D.C. 20402.

[4] ASTM 1976 Annual Book of ASTM Standards Parts 35 and 36, American Society for Testing

and Materials, 1916 Race Street, Philadelphia, PA 19103.

[5] R. B. Richard, Trans. Faraday Soc. 42^, 10 (1946).

[6] J. B. DeCoste, F. S. Malm, and V. T. Waller, Ind. Eng. Chem., 4_3 , 117 (1951).

[7] F. J. Bockhoff and J. A. Nuemann, SPE Journal K), 17 (May 1954).

[8] J. B. Howard, SPE Journal 1_5, 397 (May 1 959).

[9] P. Hittmair and R. Ullman, J. Appl . Polymer Sci . 6, No. 19, 1 (1962).

[10] R. A. Isaksen, S. Newman, and R. J. Clark, J. Appl. Polymer Sci. 7_, 51 5 (1963).

[11] Y. Suezawa, H. Hojo, T. Ikeda, and Y. Okamura, Mater. Res. Std. 5 , 55 (1965).

[12] G. E. Fulmer, Trans. Soc. Rheol . 9:2, 121 (1965).

[13] G. E. Fulmer, Trans. Soc. Pheol . 10:2, 501 (1966).

[14] H. R. Larsen, Mater. Res. Stds. 6, No. 7, 334 (1966).

[15] J. B. Howard, Fracture, Longterm Phenomena, Encyclopedia of Poly. Sci. and Tech.,

Vol . 7, 261 (1967), N. M. Bikales Editor, W11 ey-Interscience Pub.

[16] G. E. Fulmer, Polymer Eng. & Sci. _7> 1 (1967).

[17] G. E. Fulmer, SPE Journal 24, No. 10, 108 (1968).

[18] G. E. Fulmer, Proc. Fifth Int. Congress Rheol. Vol. 1, 607 (1969), Univ. Press,

Baltimore, Md., Shigeharu Onogi, Editor.

[19] 32nd Annual Technical Conf. Soc. of Plastics Eng. Preprints, 217 (May 1974).

[20] M. E. R. Shanahan and J. Schultz, J. Poly. Sci. T4B, 1567 (1976).

[21] P. D. Calvert, Nature 26j>, 13 (1977).

[22] P. D. Frayer, P. Po-Luk Ton; and W. W. Dreher, Polymer Eng. Sci. 1_7, No. 1 , 27 (1 977).

[23] M. J. Hannon, J. Appl. Polymer Sci. 1_8, 3761 (1 974).

[24] L. F. Henry III, J. Testing Evaluation 4, No. 4, 263 (1976).

[25] D. B. Hunkar, Modern Packaging Mag., November 1976.

48
[26] L. L. Lander, SPE Journal 16, 1329 (1960).

'27] R. S. Rivlin and D. W. Saunders, Trans, of Royal Soc. London, No. 865, Vol . 243,

7^1 1 1 QM \
I

IL

technical I'eojrr Documentation

1 . Report No. 2. Government Accession No.

DOT/ MTB/OHMO- 77/4

4. Title ond Subtitle Performance of Plastic Packaging for Hazardous

Materials Transportation: Part III. Stress

Cracking

7. Author's) Johm M. Crissman & Louis J. Zapas

9. Performing Orgom zohon Name and Address National Bureau of Standards

U. S. Department of Commerce

Washington, D. C. 20234

) 2. Sponsoring Agency Nome ond Address

U. S. Department of Transportation

Materials Transportation Bureau

Office of Hazardous Materials Operations

Wa<;hi ngt.nn , n. f.. 20530

3. Recipient's Cotclog No.

5. Report Oo»e December 1977

6. Performing Organization Code

-NESI R 77-1 4 TQ8 Performing Organization Report No.

10. Work Unit No. ( T R AIS)

11. Contract or Grant No. DOT AS-50074

13. Type of Report ord Period Covered

Final 6/76-6/77

14. Sponsoring Agency Code MTH-1

1 ^ . Supplementary Notes

16. Abstroct

This report is Part III of a series prepared for the Office of Hazardous

Materials Operations under Research Contract No. AS-50074. Parts I and II dealt with mechanical properties and permeation aspects of the performance

of plastic packaging for the transport of hazardous materials. Part III describes studies carried out in the NBS Polymers Division during fiscal

1977. The emphasis of this work has been directed toward providing background

information and laboratory studies on the subject of stress-cracking, the

results of which will be pertinent to the establishment of performance criteria.

i/. Keyword, Hazardous materials; mechanical

properties; inethods of test, plastic

packaging; polyethylene; stress

cracking; transportation.

1? Security Cl'tssif. ( u I this repoil)

Unclassified

18. Distribution Statement

20. S ocunty Classif. (of this pogo)

Unci assi f ied

21. No. of Poges 22. Pries

49

Form DOT F 1700.7 ( 8 - 72 ) Reproducer i of completed poy/t authorized