t/e o

c

i - /f3&

REPORTNO. DOT/MTB/OHMR-80/2

PERFORMANCE OF PLASTIC PACKAGING FOR

HAZARDOUS MATERIALS TRANSPORTATION

PART V. MECHANICAL PROPERTIES

JOHN M. CRISSMAN, LOUIS J. ZAPAS, AND

GORDON M. MARTIN

DECEMBER 1979

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 RESEARCH

WASHINGTON, O.C. 20590

Testing of An Oxygen Depletion

Detecting Device for Unvented

Gas Fueled Heaters

Esher R. Kweller

William Cuthrell

Product Performance Engineering Division

Center for Consumer Product Technology

National Engineering Laboratory

National Bureau of Standards

Washington, D.C. 20234

June 1 978

Issued January 1979

Prepared for Consumer Product Safety Commission
NOTICE

This document is disseminated under the sponsorship of the Department of Transportation

in the interest of information exchanger The United States Government assumes no liability

for its contents or use thereof.
PREFACE

This report is Part V in a series prepared for the Department of Transportation,

Office of Hazardous Materials Research under Research Contract Mo. DoT AS-50074. Parts

I and III were concerned with mechanical properties aspects of the performance of large

shipping containers used for the transport of hazardous materials. Parts II and IV dealt

with permeation aspects. Part V addresses the problems of stress-cracking in polyethylene

and mechanical properties aspects of crossl inkable polyethyl enes for rotational molding

used in the manufacture of large shipping containers.

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

Hazardous Materials Research Staff, in particular Mr. Mario Gigliotti. We are also

indebted to Mr. Paul Campbell of the Phillips Chemical Company for providing us with

sample materials.
.
TABLE OF CONTENTS

Page

1. Introduction 1

2. Environmental Stress-Cracking in Polyethylene 1

2.1 Uniaxial Creep - 1

2.2 Equal Biaxial Under Inflation 12

2.3 Bottle Inflation 20

2.4 Bent Strip Test 20

2.5 Summary 23

3. Crossl inkable Polyethylene Resins for Rotational Molding 26

3.1 Swell Ratio Tests 27

i

t 3.2 High Temperature Tensile Modulus Test 31

4.

Summary 35
.

w
LIST OF TABLES

Table 3.1 Comparison of Material Characteristics for Conventional

and Cross! inkable Polyethyl enes

Page

27

Table 3.1.1 Material Properties of Cross! inked Polyethylene Samples

Used for Swell Tests 29

Table 3.1.2 Results of Swell Ratio Tests for Samples of Crosslinked

Polyethylene 30

Table 3.2 Sample Characteristics of Crosslinked Rotationally Molded

Polyethylene Used for High Temperature Tensile Modulus Test 33

LIST OF FIGURES

Page

Figure 2.1.1 2

Figure 2.1.2

Figure 2.1.3

Figure 2.1.4

Figure 2.1.5

Figure 2.1.6

Figure 2.1.7

Figure 2.1.8

Figure 2.2.1

Figure 2.2.2

Figure 2.2.3

Log time to fail versus applied stress in uniaxial

creep for high density linear polyethylene (M, , - 99,000) in air. w

Uniaxial creep behavior for specimens of high density

linear polyethylene (Mw

= 99,000) in air.

Log time to fail versus applied stress in uniaxial

creep for an ethylene-hexene copolymer (Mw

= 211,000)

in air. w

Log time to fail versus log applied stress in uniaxial

creep for three high density linear polyethyl enes having

different weight average molecular weights.

Log time to fail versus applied stress in uniaxial creep

for high density linear polyethylene (M,, = 99,000) in

air (

) and in nonyl phenoxypoly(ethyleneoxy)ethanol

Log time to fail versus applied stress in uniaxial

creep for high density linear polyethylene in air

(

) and in dodecane (----).

Log time to fail versus applied stress in uniaxial creep

for an ethylene-hexene copolymer (M^ = 211,000) in air

(

) and in nonyl phenoxypoly(ethyleneoxy )ethanol

Log time to fail in adverse chemical environment versus

log time to fail in air for high density linear polyethylene (My

= 99,000) and an ethylene-hexene copolymer

(Mw

= 211,000).

Schematic diagram of apparatus for determining stresscracking behavior under equal biaxial (inflation)

deformations.

Radius of curvature versus percent elongation for

polyethylene sheets under inflation. The triangles

correspond to the ethylene-hexene copolymer. 0)0 -

2 9 6 K ; ©, a - 330K; (j) , [j] - 360K.

True stress versus time for high density linear polyethylene under inflation. The arrows indicate the

time at which failure occurred.

4

5

6

8

9

10

11

13

15

16
Figure 2.2.4

Figure 2.2.5

Figure 2.2.6

Figure 2.3.1

Figure 2.3.2

Figure 2.4.1

Figure 2.4.2

Figure 3.1

Figure 3.1.1

Figure 3.2.1

Figure 3.2.2

- Log time to fail versus constant applied pressure for

sheets of high density 1 inear polyethylene under

inflation in air (0,0) and in nonyl phenoxypoly-

(ethyleneoxy)ethanol (©,©).

- Log time to fail versus constant applied pressure for

sheets of an, ethylene-hexane copolymer under inflation

in air (0,0rO) and in nonyl phenoxypoly(ethyleneoxy)

ethanol (©,©,-^).

- Log time to fail in stress-cracking agent versus log

time to fail in air for polyethyl enes at two different

temperatures subjected to equal biaxial deformations

under inflation.

- Log time to fail versus internal bottle pressure for

bottles of high density linear polyethylene (M,, = 99,000) tested to failure under inflation in afr and

in stress-cracking agent.

- Log time to fail versus internal bottle pressure for

bottles of an ethylene-hexene copolymer (M^ = 211,000)

tested to failure under inflation in air and in stresscracking agent.

- Schematic diagram of apparatus for determining stresscrack resistance under conditions of a bent strip geometry

with added constant load.

- Log time to fail versus applied load for specimens of high

density linear polyethylene tested to failure in air

(0,A) and in nonylphenoxypoly(ethyleneoxy)ethanol

(©,A). The circles correspond to the bent strip geometry,

the triangles to uniaxial creep.

- Examples of recovery at elevated temperature after colddrawing of rotationally molded specimens of crossl inkable

polyethylene: A - undeformed specimen, B - cold-drawn

specimen, C - recovered specimen after 2 minutes at

423K (150°C)

.

- Correlation of swell ratio as determined by ASTM D2765

to that determined from dimensional changes for samples

of rotationally molded crossl inkabl e polyethylene having

different degrees of crosslinking.

- Log time to fail versus applied stress for different

samples of rotationally molded cross! inkabl e polyethylene

at 296 K.

- Stress relaxation behavior of rotationally molded crosslinked polyethylene.

Page

17

18

19

21

22

24

25

28

32

34

36
1. Introduction

This report is Part V in a series of reports prepared for the U. S. Department of

Transportation Office of Hazardous Materials Research (DoT OHMR) under DoT Contract

AS-50074. In fiscal 1976, the National Bureau of Standards (NBS) Polymer Science and

Standards Division initiated studies in the area of mechanical and permeation properties

of plastic materials used in the manufacture of large shipping containers for the transport

of hazardous materials. The findings of these investigations prior to the current reporting period can be found in references [1-4]*.

The present report sumnarizes experimental work done in the area of mechanical properties carried out since the completion of Part III [3]. The effort has been directed

primarily toward (1) a continuation of the studies initiated in Part III on the stresscrack resistance of polyethylene, and (2) an examination of different test methods for

determining the degree of crosslinking in cross! inkable polyethylenes for rotational molding.

2. Environmental Stress Cracking in Polyethylene

Part III in this series of reports [3] dealt in its entirety with the subject of

environmental stress cracking of ethylene plastics. The first section provided a discussion

of current ASTM test methods for determining environmental stress-crack resistance of

ethylene plastics [5]. The second section presented general background information on chemical and physical characteristics of polyethylene ,whi ch influence stress-crack resistance. In the third section, experimental work initiated in our laboratory in the area of

stress cracking in polyethylene was summarized. These studies were limited primarily to time to fail under conditions of static fatigue in uniaxial creep in both air and various

chemical environments.

Section 2 of the present report describes additional experiments in which a variety

of methods have been examined for determining stress-crack resistance under biaxial

deformation histories. In the interest of continuity and ease of reading, we shall first

briefly review a portion of the earlier-work on uniaxial creep contained in Part III. In several instances, figures contained there have been reproduced here as well in order to facilitate the discussion. We review these results in order to illustrate a number of

important points regarding the failure behavior of polyethylenes in general. The same observations will carry over into the area of biaxial deformations to be discussed later.

Throughout this current study, we shall continue to concentrate on the same two high

density polyethylenes described in Part III. These two were chosen originally because of

their widely different stress-crack behavior as based on technical information provided by

the resin manufacturer. As determined by ASTM D1693 [5], Condition A, an Fjq value of

15-20 hours is quoted for the linear polymer, while that for the ethylene-hexene copolymer

is 450 hours.

2.1 Uniaxial Creep

Shown in Figure 2.1.1 are log time to fail versus applied stress (engineering stress)

data for -specimens of the linear polyethylene tested to failure in air at four different

temperatures. By failure is meant either the time at which necking of the specimen occurred

or the time at which the specimen fractured due to cracking. The tendency for the curves

to bend down at the highest levels of stress is a result of internal heating of the speci- men during the rapid creep. In principle, one would like to use data of this type in order

to predict the long time uniaxial creep behavior (failure time) for times well in excess of

those for which data were actually collected, for example times in excess of 108 seconds at

296K (23°C) . The ideal situation is one in which the data obtained at several elevated

temperatures can be superposed to some reference temperature by simply shifting the curves

along the time axis. The shift factors can then be plotted versus reciprocal temperature.

*Numbers in square brackets refer to references found at the back of this report.
TIME

TO

FAIL

(s)

5 10 15 20 25 30 35

STRESS (MPa)

Figure 2.1.1 - Log time to fail versus applied stress in uniaxial creep for high density

linear polyethylene (M^ = 99,000) in air.

2
and in the event that a straight line results, a temperature independent activation

energy can be determined. This information can then be used to predict the time to fail

for any other set of temperature and applied stress within the total range of temperatures

and stresses investigated. Unfortunately, the data in Figure 2.1.1 cannot be superposed

simply by a shift along the time axis (even when disregarding the bent down portions of

the curves). One reason is that more than one failure mechanism is present, and importantly the activation energies for the different mechanisms are not the same. For stresses

above about 15MPa the creep behavior becomes dominated by the viscoelastic properties of

the polymer. In this region the specimen creeps to a sufficiently large strain that an instability occurs. By instability we mean the point beyond which the material can no longer undergo a homogeneous deformation. Beyond this point morphology or structure

changes take place, for example void formation or localized yielding, and eventually necking of the specimen occurs. Based on a theoretical treatment by Zapas [6], Zapas and

Crissman [7] have shown experimentally that for linear polyethyl enes in the molecular

weight range (My) from 100,000 to 200,000 the instability point for uniaxial extension

occurs at a strain of about 10-12%. In fact, from creep data, it has been found possible

to predict, well in advance of any visible evidence, the point in time at which necking

will occur [7]. In terms of strain, necking generally occurs for linear polyethylene at

about 25 to 30%, depending upon the magnitude of the stress and to some extent on the

polymer type.

On the other hand, below about 12.5 MPa, the region of small applied stress, the

specimens fail by cracking. In this regime, the specimen develops a crack which propagates

through the specimen well before the global creep strain becomes sufficiently large to reach the point of instability. That this is true is 'illustrated by the creep behavior

shown in Figure 2.1.2 for several of the same specimens for which time to fail data are presented in Figure 2.1.1. At large loads the creep characteristically becomes increasingly rapid with log time up to the Doint where necking occurs. In each case where the load

was greater than about 14 MPa the creep strain at failure was greater than 12%. In the

region of small applied loads the creep curves tend to flatten so that over a large percentage of the total time to failure the global strain remains nearly constant. In this

event, the material more nearly approximates the behavior of an elastic material . It should be pointed out that (for linear polyethylenes) from the time that a crack first

becomes visible until fracture occurs, the elapsed time is only a very small fraction of

the total lifetime (time to failure). An interesting region is the one intermediate to the two extremes. For stresses between about 12.5 and 15 MPa (Figure 2.1.1) it cannot be

anticipated which failure mode will prevail, and in some instances specimens have been

observed to show both failure modes simultaneously.

Comparable time to fail versus applied stress data are shown for the ethylene-hexene

copolymer in Figure 2.1.3. One feature to be noted here is that the curves are steeper

than for the linear polymer. This behavior gives rise to the interesting result that at

high stresses (> 14MPa) the copolymer fails earlier at a given stress than does the linear

polymer, while at small stresses (< 14MPa) it has increasingly longer lifetime. Another

important point regarding the copolymer is that the mechanism which leads to instability

and eventual necking occurs even at very long times at room temperature (296K). We note

that the only specimen of this material for which cracking was observed was one requiring

2.8 years to fail. As will be seen presently, this behavior will have considerable bearing

on determining stress-crack resistance.

As a demonstration that the observed differences in behavior of the two polyethylenes

is not simply one of a difference in molecular weight, we- show in Figure 2.1.4 on log-log

coordinates the time to fail behavior for three different samples of linear polyethylene.

All three samples have very nearly the same number average molecular weight (15-16 x 104

) but different weight average molecular weights (9.9 - 19.2 x IQ 4

). The sample having an

My of 19.2 x 104 is nearly comparable to the ethylene-hexene copolymer for which My = 21.1 x 104 . Observe that at 296 K all three polymers show essentially the same behavior.

The small difference observable at the higher temperatures may reflect some difference in

crystallinity between sample types.

Up to this point we have considered only the behavior in air, which for convenience

will be regarded as the reference environment. It will be useful here to review briefly

some of the earlier work on the uniaxial creep behavior of these same two polymers in the

presence of stress-cracking agent or solvent. A description of the apparatus and procedures for carrying out the stress-cracking experiment in uniaxial creep have been given

in Part III [3] and will not be repeated here. In the discussion which follows, the

adverse environment was either a ten percent solution in distilled water of nonyl phenoxypoly

3
4

Figure

2.1.2

-

Uniaxial

creep

behavior

for

specimens

of

high

density

linear

polyethylene
TIME

TO

FAIL

(s)

0 5 10 15 20 25 30 35

STRESS (MPa)

Figure 2.1.3 - Log time to fail versus applied stress in uniaxial creep for an ethylenehexene copolymer (M^ = 211,000) in air.

5
2.5 5.0 7.5 10 15 20 30 40 STRESS (MPa)

Figure 2.1.4 - Log time to fail versus log applied stress in uniaxial creep for three high

density linear polyethyl enes having different weight average molecular weights.

6
(ethyl eneoxy)ethanol or the normal alkane, dodecane (Cn

2

H 26 ) - The first of these

(hereafter referred to simply as stress-cracking agent) is regarded to be a surface active

agent, whereas dodecane behaves as a swelling agent.

Results for the linear polymer in stress-cracking agent are shown in Figure 2.1.5

for two different temperatures . The solid lines represent the time to fail in air data

from Figure 2.1.1 and the solid lines with triangles the behavior in stress-cracking agent.

Above a certain stress, depending somewhat on the temperature, the curves for air and

stress-cracki ng agent merge. Although data are not shown in the figure, at stresses higher

than those for which data points are shown the behavior was the same in stress-cracking

agent as in air. The region where the two curves come together corresponds to the point at

which the failure mode changes from one of fracture by cracking to one of necking. Therefore, one conclusion to be drawn is that the failure behavior in the region associated with

the instability is unaffected by the presence of this type of stress-cracking agent.

On the other hand, the failure mode involving fracture by cracking is influenced by

the presence of the stress-cracking agent, and in such a way as to suggest that the action

of the stress-cracking agent is one of a constant additive stress, in this case of about

2.2 MPa. It is an interesting feature that in air the time to fail curve is smooth overall,

showing no abrupt change in character as the failure mode changes. On log-log coordinates

(Figure 2.1.4) the behavior was very nearly linear over four or more decades in time.

However, in the presence of the stress-cracki ng agent, the two failure mechanisms become

quite distinct from one another.

As noted earlier, the creep strain at small stresses remains nearly constant over a high percentage of the total lifetime of the specimen, and the material behaves more in a manner similar to an elastic solid (strain independent of time). Such a behavior suggests

that the mechanism associated with cracking can be treated in a way reminiscent of an inorganic glass. In an earlier work [8] we have shown that the time to fail behavior for

the cracking mode can be represented rather well by the empirical relation developed some thirty years ago by Glathart and Preston [9] given by

log tg = A + B/a, (1)

where t„ is the time to fracture, a is the applied stress, and A and B are material

constants. This analysis leads to an activation energy for the fracture process of about

124 kJ/mole (29.6 kcal/mole).

By way of contrast, the behavior of the same linear polyethylene in the n-paraffin,

dodecane ( Cq

2

h 26^j s Presented in Figure 2.1.6. The dimensions used to calculate the stress

are based on the unswollen specimen. In order to avoid time effects, the specimens were

presoaked in the liquid for at least 16 hours prior to loading. Additional presoaking for

longer periods had no appreciable influence on the final results. The behavior in solvent

is markedly different than that in the stress-cracki ng agent. Whereas in stress-cracki ng

agent, only the region corresponding to fracture by cracking is affected, in solvent the

entire time to fail curve is shifted significantly. In solvent, triaxial stresses are introduced into the material due to swelling, and the effect appears to be comparable to that of an increase in temperature. In the case of amorphous polymers, it is known [10]

that a swelling agent both shifts the relaxation times and lowers the moduli. If thi

s^ same behavior holds true for semi crystal line polymers, one would then expect a shift along

both the time and stress axis for superposition, as would appear to be the case from the

results shown in Figure 2.1.6.

The behavior of the ethylene-hexene copolymer in stress-cracki ng agent is presented in

Figure 2.1.7. The principle difference in behavior between copolymer and the linear

polymer is that the point at which the time to fail in the stress-cracking agent deviates

from the oehavior in air occurs at longer times. This result is consistent with the

remark made earlier that the failure associated with instability and necking extends out

to much longer times for the copolymer than for the linear polymer. At 256K, quite long

times are required before the influence of the stress-cracking agent becomes apparent.

However, by raising the temperature to 330K (57°C) the deviation in behavior occurs at

relatively short times.

One interesting result of this work can be seen by plotting the log time to fail in

adverse environment versus the log time to fail in air, all data obtained at 330K, as is

shown in Figure 2.1.3. The dashed line denotes behavior that does not deviate from that

in air. The vertical distance from the dashed line can be interpreted as a measure of the

reduction of the static fatigue life brought about by the presence of an adverse chemical

environment. In stress-cracking agent the reduction in lifetime for the copolymer actually

7
TIME

TO

FAIL

(s)

Figure 2.1.5 - Log time to fail versus applied stress in uniaxial creep for high density

linear polyethylene (Mu

= 99,000) in air (

) and in nonyl phenoxypoly(ethl eneoxy) ethanol

( A )

8
TIME

TO

FAIL

(s)

5 10 15 20 25

STRESS (MN/m2

)

Figure 2.1.6 - Log time to fail versus applied stress in uniaxial creep for high density

linear polyethylene in air (

) and in dodecane ( ).

9
STRESS (fVUM/m2

)

Figure 2.1.7 - Log time to fail versus applied stress in uniaxial creep for an ethylenehexene copolymer (Mw = 211 ,000) in air (

) and in nonylphenoxypoly(ethyleneoxy)ethanol

(—4—). “

10
oo

CO

LO

CO

CM

1N3WN0HIAN3 1V3IW3H3 Nl HVd 01 3WI1 901

11

LOG

TIME

TO

FAIL

IN

AIR

Figure

2.1.8

-

Log

time

to

fail

in

adverse

chemical

environment

versus

log

time

to

fail

(M a

= r

2n r

ooo5 h

dGnS1ty

linear

P

o1

Vethylene

(M

W

=

99,000)

and

an

ethylene-hexene

copolymer
—

becomes greater at longer times than for the linear polyethylene. Yet due to the fact

that at small stresses the copolymer experiences a very much longer lifetime in air,

for a given stress level, more tim? is still required for the copolymer to fail in the

stress-cracking agent than for the linear polymer. The curve showing the behavior of the

linear polymer in dodecane demonstrates that the solvent strongly effects even the early

time behavior and that a significant reduction in lifetime occurs over the entire time

scale investigated.

2.2 Equal Biaxial Under Inflation

Up to this point we have considered the stress-crack behavior in uniaxial creep. In the next several sections we shall examine the stress-crack behavior of the same two high

density polyethyl enes under a variety of biaxial deformation histories. The first of

these to be considered is one involving equal biaxial deformations induced by inflation.

The experimental arrangement has been described in some detail in Section 3.3 of Part III

[3]. However, a number of experimental details need to be elaborated on here, so that an abbreviated description of the method will again be given here. The specimen, in the shape

of a flat sheet, is clamped at its outer edge as shown in the diagram in Figure 2.2.1. The

system is then pressureized at constant internal pressure until failure occurs. Failure

is detected either by sound or more often by the feel of escaping gas. For the geometry

shown in Figure 2.2.1, an equivalent, or radial, stress can be calculated at the pole of

the bubble from the expression

a

PrX

2 2h

where

a = radial stress component

P = internal pressure

r = radius of curvature at the pole

h = initial thickness of the sheet

X = stretch ratio given by 1 + e, where e is the strain

The equation is valid only in the vicinity of the pole, and for strains up to about 50%. At distances very far from the pole, or for strains larger than about 50%, the surface no longer approximates a sphere. In cases where a stress-cracking agent was present a small

amount of liquid was placed in the bottom of the pressure reservoir, then after pressurization the entire assembly was immediately inverted so 'that the stress-cracking agent was

in contact with the specimen only in its polar region. In the present experiment, the

stress-cracking agent was only in contact with the inner surface of the specimen. Had it

been in contact with both sides, the results to be described presently may have been different. One difficulty inherent in this method occurs when the applied internal pressure

is small (long times to fail). At low pressures the situation can arise whereby the radius

of curvature in the specimen becomes smaller in the region of the clamp than at the pole.

In this event, the failure may occur preferentially at the clamp, a region for which the

stress and strain are not determined. This problem was particularly troublesome when the

tests were carried out at high temperature in the presence of stress-cracking agent. At

high temperature, water vapor containing the stress-cracki ng agent was found to preferentially attack the region with convex curvature near the clamp, to the extent that in some cases the specimens cracked almost the entire way around the ring. For this reason,

only data for which the failure occurred in the region near the pole will be presented.

The specimens were prepared by compression molding flat sheets 15 cm in diameter and

nominally 0.060 cm in thickness. The temperature history during molding was controlled

according to the same procedures used to mold the uniaxial specimens. For a given speciment the thickness was determined at several locations within a region 1.5 cm in radius

about the center of the sheet. Typically, a variation in thickness of about 0.002 cm was observed, and for purposes of calculating the stress the minimum value recorded was used.

From specimen to specimen the minimum thickness varied by up to 0.005 cm. Therefore, in

12
<

LU

00

CD

<i CL. CL

o

o

CL

00

<

CD

13

Figure

2.2,1

-

Schematic

diagram

of

apparatus

for

determining

stress-cracking

behavior

under

equal

biaxial

(inflation)

deformations.
cases where the data are presented in terms of the applied internal pressure, the values

of pressure quoted have been normalized to a constant sheet thickness of 0.060 cm. The

strain was determined with the aid of two sets of india ink marks placed at right angles

to one another 1.5 cm from the center of the sheet.

Shown in Figure 2.2.2 on log-log coordinates is the radius of curvature (as determined

at the pole) versus percent elongation for two different linear polyethylenes and the

ethylene-hexene copolymer. The data for the branched material are shown by the triangles

in the upper portion of the figure. Measurements were carried out on both linear polymers

at three temperatures, 296K, 330K, and 360K (87°C). One interesting feature is that for

each polymer type, the behavior can be represented rather well by a unique linear relationship over the wide range of deformation and temperature Investigated. The primary difference between the two curves for the linear polyethylene samples is a slight difference

in slope. Since the true stress depends on both the radius of curvature and strain, under

conditions of a constant internal pressure one can use a master curve of this type to determine either quantity after having measured the other.

The complicated true stress behavior of the linear polyethylene at 296K is shown in

Figure 2.2.3 for different values of constant internal pressure. At the highest pressures

the stress goes through a minimum and then rises to a high value before failure occurs.

This region is the equivalent to that for uniaxial creep where the material experiences

an instability and eventually necks. In the equal biaxial case the material also becomes

unstable and may cold-draw to some extent before failing by the appearance of a crack.

On the other hand, at the low pressures the sheets fail by the appearance of a crack

before the material becomes globally unstable. In this region, the true stress decreases

monotonically, or tends to level off, up to the point of failure. At the smallest pressures the tendency is again one of the global strain, or true stress, being nearly constant

over a high percentage of the total lifetime.

In Figure 2.2.4 are shown results for the time to fail in air and stress-cracking

agent (on log coordinates) versus applied pressure for the sample of linear polyethylene

with = 99,000. Open symbols denote the behavior in air, and filled symbols that in

stress-cracking agent. It should be pointed out that similar measurements carried out on the linear polyethylene having on Mi^ of 192,000 yield almost identical results, to the

extent that both sets of data very nearly superpose. Note that in Figure 2.2.4 the independent variable is applied pressure and not stress. We have already seen that the true

stress behavior for this type of biaxial deformation is rather complicated (Figure 2.2.3)

and there appears at this point to be no straightforward way to represent all the time to

fail data in terms of the true stress. Cold-drawing occurs in the region above about 3 MPa at 296K while cracking takes pTace in the region below 3 MPa. Unlike the uniaxial

creep case, the overall curve representing the time to fail behavior, when represented in

terms of applied pressure, does show a distinct break where the change in failure mode

occurs. Again, the presence of the stress-cracking agent serves to make the abrupt change

more distinct. Note that the action of the stress-cra.cking agent is one of decreasing the

lifetime of the linear polyethylene (and the higher molecular weight linear polyethylene

as well) by a factor of about ten. The behavior of the ethylene-hexene copolymer under inflation is shown in Figure 2.2.5.

The curves representing the behavior in air are quite steep, to the extent that at 296K the

time required to collect data becomes excessive. A small letter "c" placed next to a data

point denoted failure by cracking. The remaining specimens failed by the mechanism of

cold-drawing. As can be seen, all of the specimens of this material tested at 296K failed

by cold-drawing. However, in the presence of the stress-cracking agent the failure mechanism associated with cracking is made much more emphatic to the extent that it is now

quite prominent even at room temperature. As was true for uniaxial extension, the use of

stress-cracking agent can be an effective means for separating out the mechanism of failure

due to cracking.

In the same manner employed earlier, we compare in Figure 2.2.6 the behavior of the

two polyethylenes under inflation by plotting the time to fail in stress-cracking agent

versus that in air on log-log coordinates. In the equal biaxial case, the rate at which

the lifetime is decreased by the presence of the stress-cracking agent is actually greater

for the copolymer over the entire time scale at 296K than for the linear polymer. At 330K,

the behavior is reminiscent of the uniaxial case where the two curves cross over. However,

in the equal biaxial case, the crossover point occurs nearly one decade earlier in time

than for uniaxial extension. We point out, however, that even though the rate a't which

the lifetime is decreased by the presence of the stress-cracking agent is greater for the

copolymer, it still experiences a longer lifetime for a given applied pressure. This is

14
RADIUS

OF

CURVATURE

(cm)

Fiqure 2.2.2 - Radius of curvature versus percent elongation for polyethylene sheets under

inflation. .The. triangl es correspond to the ethylene-hexene copolymer. O > * 296K ’ © ’ B - 330K; O , - 360K.

15
cn cm (ed

£

_oix) ssaais

16

Figure

2.2.3

-

True

stress

versus

time

for

high

density

linear

polyethylene

under

inflation.

The

arrows

indicate

the

time

at

which

failure

occurred.
(s) llVd 01 3IA1I1

17

APPLIED

PRESSURE

(MPa)

Figure

2.2.4

-

Log

time

to

fail

versus

constant

applied

pressure

for

sheets

of

high

density

linear

polyethylene

under

inflation

in

air

(0,0)

and

in

nonylphenoxypoly(ethyleneoxy)

ethanol

(©,©).
(S) 1SVJ 01 HiAlil

CO Q.

18

Figure

2.2.5

-

Log

time

to

fail

versus

constant

applied

pressure

for

sheets

of

an

ethylenehexene

copolymer

under

inflation

in

air

(0,y>rO}

and

in

nonylphenoxypoly(ethyleneoxy)

ethanol

(©,©r®-).
19

Figure

2.2.6

-

Log

time

to

fail

in

stress-cracking

agent

versus

log

time

to

fail

in

air

for

polyethylenes

at

two

different

temperatures

subjected

to

equal

biaxial

deformations

under

inflation.
because at small pressures the lifetime of the copolymer in air is very much greater than

is that of the linear polymer. In fact, if the data from Figures 2.2.4 and 2.2.5 are superposed on one another, it can be seen that for small pressures the lifetime of the

copolymer in stress-cracking agent is a factor of ten or more greater than that for the

linear polyethylene in stress-cracking agent. It would appear, therefore, that an important factor in determining stress-crack resistance is the ability of the material to perform in air.

2.3 Bottle Inflation

A second type of experiment involving biaxial deformations was bottle inflation at constant internal pressure. This method is reminiscent of ASTM 02561 Standard Method of

Test for Environmental Stress-Crack Resistance of Blow-Molded Polyethylene Containers,

Procedure C [5]. One difference is that the ASTM test is carried out at one pressure for

a period of up to 360 hours, whereas in the present work the bottles were tested to failure

at a series of different pressures. In cases where the stress-cracking agent was present,

the bottle was partially filled with solution (usually less than 1/4 full) and the bottle

then tumbled to insure that initially the entire inner surface was uniformly coated. Time

to failure versus applied constant internal pressure data were then collected at three

different temperatures.

Our results for 120 ml blow-molded bottles prepared from the same tv/o polyethyl enes already described are presented in Figures 2.3.1 and 2.3.2. As was true for uniaxial

extension and the equal biaxial under inflation experiment, the failure mode changes depending upon the magnitude of the applied pressure. For the highest pressures employed,

the bottles failed by cold-drawing, whereas at the lower pressures failure occurred as a result of cracking. As before, the lifetime of the copolymer bottle in stress-cracking

agent was significantly greater than that of the bottle prepared from the linear polymer.

Bottle inflation differs from the bubble type of test described in the last section

in that no portion of the bottle deforms equal biaxially. For example, in the region of

the side wall of the bottle the material creeps in such a way that the circumf erence

increases with time, while in the direction parallel to the length of the bottle very little deformation occurs. One interesting effect appeared with regard to the particular

bottles used in this work. In Figure 2.3.1, consider the data showing the behavior of the

linear polymer in stress-cracking agent" (filled symbols). Although these bottles failed by

cracking, the position in the bottle where the crack appeared depended on the magnitude of

the applied pressure. At the higher pressures, the crack occurred in the side wall of the

bottle, generally along or close to a mold parting line. This was the same region of the

bottle where failure due to cold-drawing occurred at still higher pressures. For the bottles used in this work, this location on the bottle was also the thinnest point. However,

at the lowest pressure investigated (at a given temperature) the crack occurred at the

chime of the bottle where the curvature and thickness were greatest. This observation

underscores the need to examine the stress-crack resustance of the entire container in

instances where internal pressure is an important consideration.

2.4 Bent Strip Test

In this section, we shall consider a third type of experiment, one which employs a bent strip geometry. A diagram of the apparatus is given in Figure 2.4.1. The specimen,

in the form of a strip, is bent around a metal cylinder and is clamped as shown. The

metal cylinder is anchored to a base and the clamp is connected via two pulleys to an applied constant load. The diameter of the cylinder was 1.11 cm (7/16") and the thickness

of the specimen was nominally 0.10 cm (0.04"). In cases where the tests were conducted

in stress-cracking agent, the specimen assembly was submerged in the liquid maintained at

330K (57°C). Since the failure always occurs in the region of the specimen in contact

with the cylinder, there is no need to submerge the specimen completely.

Preliminary results obtained for the linear polymer, both in air and in stress-cracking agent, are shown in Figure 2.4.2. Also included for comparison purposes are the data

obtained for the same material in uniaxial creep. The non-homogeneous deformation introduced into the specimen simply by bending it around the cylinder has a pronounced effect

on the behavior at loads below about 7.5 MPa, both in air and stress-cracking agent. Note,

however, that as the load is increased, the behavior of the bent strip becomes increasingly

like that of the uniaxial specimen. Also, it can be seen that even for a condition of

zero applied load the bent strip specimens fail within a relatively short time. Under

20
Figure 2.3.1 - Log time to fail versus internal bottle pressure for bottles of high

density linear polyethylene (Mw

= 99,000) tested to failure under inflation in air and

in stress-cracking agent.

21
TIME

TO

FAIL

(s)

Figure 2.3.2 - Log time to fail versus internal bottle pressure for bottles of

hexene copolymer (M^ = 211,000) tested to failure under inflation in air and ii cracking agent.

an ethylenestress22
zero applied load, the tines to fail in stress-cracking agent are, on average, not greatly

different than those quoted by the resin manufacturer for specimens of the same polymer

tested according to ASTM D1693 Standard Method of Test for Environmental Stress-Cracking

of Ethylene Plastics, Condition A (F5 Q = 15-20 hours) [5]. In the limit of zero applied

load, the primary difference between the two methods reduces to one of degree of severity

in the bent portion of the specimen, and a notched specimen is used. Also, it should be

pointed out that the ASTM test is conducted at 323K (50°C) rather than at 330K (57°C) as was the case here.

The extent to which failure is accelerated by bending the specimen can be controlled

to a degree by changing three parameters: the diameter of the cylinder, specimen thickness,

and magnitude of the applied load. For example, the thinner the specimen is made or the

larger the diameter of the cylinder, the closer the behavior will be to that for uniaxial

creep. For specimens which were only one half the nominal thickness of 0.10 cm a significant bending up of the failure curve was observed from that shown in Figure 2.4.2. Also,

the scatter in the data becomes decidedly worse as the zero applied load condition is

approached. Therefore, it would appear that one advantage to the use of an applied constant

load is improved reproducibi 1 ity. The second obvious advantage is one of time savings in

that the time required to produce failure can be reduced.

2.5 Summary

In the manufacture of large polyethylene shipping containers, a test for environmental

stress-crack resistance can be useful at one or more stages of production. Prior to processing, such a test provides both a means for screening raw materials and as a quality

control measure, for monitoring batch to batch differences in the polymer resin. Following

fabrication, the test method may be used to establish performance characteristics of the

finished container. In the latter case, the test involves either the cutting of specimens

directly from the container and subjecting of them to a stress-crack resistance test such

as ASTM D1693 [5], or the container as a whole is subjected to some type of accelerated

holding test. Tests carried out on specimens cut from the side wall or ends of the con- tainer yield valuable information concerning the effects of processing on the material

behavior of those protions of the container. However, because of complicated geometry

and/or conditions of processing, the behavior of such specimens may not reflect completely

the behavior of the container as a whole. It was noted in Section 2.4 that bottles tested

to failure under inflation in a stress-cracking agent failed at different locations on the

bottle depending upon the magnitude of the internal pressure. Therefore, testing of the

entire container under conditions of varied pressure and temperature may be important. For

this purpose, a test of the type described in ASTM 02561 [5] can be applied, since in

principle it can be adapted to any size or shape container.

In the work just described, we have investigated the stress-crack behavior of two high density polyethyl enes under four different deformation histories. These included

uniaxial creep, equal biaxial under inflation, bottle inflation, and bent strip geometries.

An important observation was that all four methods yield essentially the same information

regarding stress-crack resistance. Another point made is that a primary factor in determining stress-crack resistance appears to be the ability of the material to perform in air,

and not necessarily the rate at which failure is accelerated by the stress-cracking agent

relative to the behavior in air. Finally, we note that unnotched specimens were used in

all four methods examined.

As a test method, uniaxial creep represents a well controlled, easy to perform, experiment, but long testing times are required to obtain the desired information. A test

employing equal biaxial deformations under inflation greatly emphasizes the failure mechanism associated with crack growth, and is, therefore, a highly sensitive measure of

stress-crack resistance. However, disadvantages are (.1) flat sheets of material are required, ( 2 ) the state of stress in the specimen as a function of position and time is

complex, and ( 3 ) the specimens do not always fail in the polar region, especially at elevated temperatures, and the results may tend to be biased. Bottle tests are useful in cases where the geometry of the finished product is suitable for inflation. This type of

test is valuable for determining the weakest point. However, the test should probably be

carried out at more than one pressure, since it has been demonstrated that in the presence

of a stress-cracking agent, the location of the failure can vary with the pressure.

As a general test method for determining environmental stress-crack resistance, the

method involving a bent strip geometry described in Section 2.4 appears to combine several

of the more desirable features of the various other tests investigated. The experiment is

23
Figure 2.4.1 - Schematic diagram of apparatus for determining stress-crack resistance under

conditions of a bent strip geometry with added constant load.

24
versus applied load for specimens of high density linear

polyethylene tested to failure in air (O.A) and in nonylphenoxypoly(ethyleneoxy)ethanol

r«> yi i The rirrles corresDond to the bent strip geometry, the triangles to umaxia,

Figure 2,4.2 - Log time to fail

(©,A).

creep.

25
easier to perform and involves a less complicated state of stress than does the equal

biaxial deformation under inflation experiment. At the same time, the bent strip geometry

can significantly reduce testing times compared to that required for uniaxial creep tests.

Based on preliminary results, it is our view that this particular method has several

advantages over the bent strip test currently in widespread use. Under conditions of an uncontrolled bend geometry, the deformation in the region of the bend can be quite severe.

As a function of time, the severity of the bend will be influenced by the ability of the

specimen to relax, which, in turn, will depend on such factors as crystallinity. In the

present method, the bend geometry is controlled for all specimens and constant in time. The

imposition of a small constant applied stress appears to have two effects, one of reducing

the experimental scatter in the data, and the other of accelerating the failure and thus

reducing testing time.

The type of test under consideration here is most suited to situations where the speciments can either be specially molded or cut directly from the processed part. With reference to large shipping containers, this means the machining of specimens either from the

side wall or the ends. However, as noted earlier, because of complex container geometry

and/or conditions of processing, testing of the entire container under inflation may also

be important. The most complete information can be gained by conducting both types of test.

3. Crossl inkable Polyethylene Resins for Rotational Molding

Several years ago, a new class of crossl inkable polyethylene resins for rotational

molding entered the marketplace. Such materials are currently coming under the scrutiny

of the DoT OHMO as a result of requests for special exemptions for their use in the area of

fabrication of large shipping containers for hazardous materials. The rotational molding of

crossl inkabl e polyethylene represents a different approach to container fabrication than

that which we shall refer to as "conventional" polyethylene. The physical properties of

conventional polyethylene are determined primarily during the manufacture of the polymer

resin. Although the processing represents an important step which can influence the performance of the finished product, these properties are determined principally by the

polymer chemistry, including factors such as molecular weight distribution, and extent of

branching. For crossl inkabl e polyethylenes , on the otherhand, both the polymer chemistry

and physical properties are highly dependent upon the processing, more specifically on the

crosslinking step. Properties such as stress-crack resistance, impact strength, and fracture toughness are critically dependent on the degree of crosslinking.

Since the degree of crosslinking clearly represents a key parameter to determining

performance characteristics of crossl inkabl e polyethylene, test methods should be able to directly correlate the two. Several types of test for which it is believed a correlation

can be established include the following:

(1) low temperature impact resistance

(2) percentage of extractables

(3) swelling ratio

(4) high temperature tensile modulus

(5) stress-crack resistance

Test number (1) is recommended by a resin manufacturer as a quality control measure to establish a proper degree of crosslinking. The apparent advantage of a test of this type

is one of time savings. It also requires very little in the way of costly equipment and

can be carried out in the field on specimens of nearly any size or shape. In the absence

of quantitative data establishing a correlation between degree of crosslinking and impact

resistance, we cannot comment on this test. Tests (2) and (3) are widely used to determine

degree of crosslinking and are incorporated into ASTM D2765 Standard Method of 'Test for

Degree of Crosslinking in Ethylene Plastics as Determined by Solvent Extraction [5J. The

main criticism of this test is that it requires in excess of 48 hours to perform. Tests

(4) and (5) have been suggested as methods which should be highly sensitive to degree of

crosslinking. However, based on work described earlier in this report, it must be anticipated that test (5) for stress-crack resistance will be highly time consuming.

In the next several sections we shall summarize experimental work carried out in our laboratory concerned with tests (2) through (4). Experiments concerned with test (5) for

stress-crack resistance have been initiated but the work has not yet progressed sufficiently

far to be included in this report. For convenient reference, we have included in Table 3.1

a comparison of several material characteristics between crosslinkable polyethylene and

26
conventional high density polyethylene.

Table 3.1

Comparisoh of Material Characteristics

HOPE Crosslinkable Polyethylene

Chemical nature

Linear or lightly branched chains Cross! inked network

Melt flow properties

Resin can be characterized by Melt flow is meaningless

its melt flow index

Potential for chemical

changes during fabrication

Substantial. Long cycles at high

temperatures are required to distribute the resin and to effect

the proper degree of crosslinking

Limited. Heating, injection, and

blow molding cycles are generally

relatively short

Fabrication process

Blow molding (in general) Rotational molding

Applicable ASTM specifications

D1248 (type III) - Specification for D2647 - Specification for CrossPlastics Molding and Extrusion Materials linkable Ethylene Plastics

Nature of 'Attack by Solvents

Dissolves Swells

Recovery at high temperature after plastic -deformation

Partial recovery or melting Essentially complete recovery

depending on temperature

One interesting feature of crosslinkable polyethylene is listed in the last item in

Table 3.1, i.e. the ability to recover after large scale deformation. Conventional polyethylenes cold-drawn in uniaxial extension will, when heated to just below the melting

point, show partial recovery toward the original configuration. Crosslinkable polyethylenes, on the other hand, can show essentially complete recovery. This is demonstrated in

Figure 3.1 for specimens from three different samples of crosslinkable polyethyl enes

.

Characteristics of the three samples are given in the following section. In each case,

the specimens labeled A represent the original specimens as cut. Specimens B were subjected to a constant load of 1.23 x 10? Pa and were then allowed to cold-draw in creep to their natural draw ratio (A. = 5 at 296K). The drawn specimens were then placed in an oven

at 423K (150°) for about 2 minutes. The recovered specimens are designated as C. It will

be seen in the following section that, based on swell ratio, the three samples have widely

different degrees of crosslinking, yet their drawing and recovery behavior is essentially

the same.

3.1 Swell Ratio Tests

One of the most widely used tests for establishing degree of crosslinking in polyethylene is a test of the type described in ASTM D2765 Standard Test Method for Degree of

27
Figure 3.1 - Examples of recovery at elevated temperature after cold-drawing of rotation

ally molded specimens of cross! inkabl e polyethylene: A - undeformed specimen, B - colddrawn specimen, C - recovered specimen after 2 minutes at 423K (150°C).

28
Crosslinking in Ethylene Plastics as Determined by Solvent Extraction [5]. As noted

earlier, this test yields two quantities, the swell ratio and percentage extractabl es

.

Both quantities are regarded as representing a measure of degree of crosslinking achieved

by the processor. In this section, we shall examine by two different procedures the

behavior of eight samples of crosslinked polyethyl ene. All eight samples were reportedly

prepared by the rotational molding technique. Listed in Table 3.1 are the density and

percent gel values for each. With the exception of Sample 3, all the samples were obtained

from a resin manufacturer who provided the values for percent gel. Gel is the insoluble,

non-extractabl e portion of the resin that is produced by crosslinking. Sample 3 was provided by a commercial fabricator, and the resin was stated to be nominally the same as Sample 1. All density measurements were determined in our laboratory by hydrostatic weighing in distilled water. The error values shown for the density represent an estimate based

upon analysis of our measurement technique. It should be observed that, while the density

is not a highly sensitive measure of the degree of crosslinking, there is a trend toward

higher density with decreased percent gel. This behavior is a likely consequence of greater crystal 1 inity in those samples which are more lightly crosslinked.

Table 3.1.1

Material Properties of Samples Used for Swell Tests

Sample

Number

12

34

56

78

Percent Gel Density at 296 . 2

K

(23.2°C)

(d/em3

)

>60

>90

50

68

8182

88

0.941 + .001 0.940 + .001

0.940 + .001

0.945 + .001

0.945 + .001

0.942 + .001

0.942 + .001

0.942 + .001

Two different types of swell tests were carried out on each sample. In one method, a minimum of three specimens, in most cases four, were cut from different parts of each

sample and were tested according to the procedures outlined in ASTM D2765 (Method C).

Using this method, both the swell ratio and percentage extractables were determined. In a second set of experiments, a minimum of two specimens were tested as follows. Small rec- tangular shaped specimens were machined so as to be uniform in dimensions. The length,

width, and thickness of each specimen were then measured in the dry state with a micrometer.

Each specimen was placed on an aluminum foil tray in a bottle and submerged in xylene to

a volume ratio of solvent to specimen of about 500. The bottle, in turn, was placed in a hot oil bath maintained at 383K (110°C) for a period of 24 hours. Consistent with the

ASTM procedure, it was found that after 24 hours no further swelling of the specimen occurred. Periodically, the specimens were turned over to insure that no sticking to the foil

occurred and to allow each face equal access to solvent. This procedure was found necessary in order to provide uniform test conditions for all materials, since specimens that

are lightly crosslinked will creep under their own weight if suspended vertically in the

solvent. After 24 hours, the specimen was removed from the xylene and its dimensions

measured immediately to the nearest 0.5 millimeter with a ruler. In the event the specimen

did not s we 1 1 uniformly, the dimensions used for the purpose of calculating the swollen

volume represent average values. Although no rate of swelling measurements have as yet

been performed, one possible advantage of this second type of test is that specimens improperly crosslinked may be sufficiently swollen after only a few hours so as to be

considered unacceptable for the intended use. Results for the two types of swell tests are tabulated in Table 3.1.2. For those

specimens tested according to the ASTM procedure both the swell ratio and percent extractables are included. In the case of Sample 4 (50% gel), we were unsuccessful in several

attempts to determine a swell ratio based on volume change since the swollen material did

not maintain dimensional stability. Therefore, tests based on dimensional changes cannot

be expected to be reliable for samples with a percent gel of less than 6074-70% . It is also

true that for percent gel levels below 60% the ASTM procedure for determining percent

29
extractables becomes more difficult since the test procedure relies upon recovering all

the non-extractabl e material from the solvent prior to the drying step. Note that even the

most highly cross! inked sample swelled by a factor of ten or more and that about 15% of the

material was extracted.

Table 3.1 .2

Sample

Swell Ratio

ASTM 2765

Results of Swell

from

Ratio Tests

Percentage Extractable

from ASTM 2765

Swell Ratio from

Volume Chanqe

26.1 23.6

1 26.9 24.5 19.5

26.9 23.3 19.4

26.9 23.1

12.5 15.7

2* 12.9 17.4 10.2

18.0 14.8 11.9

18.6 15.2

51.6 21.3

50.0 •34.0

3 49.2 22.3 36.3

46.3 19.9 41 .2 43.8 25.0

40.7 21.4

48.4 53.9

4 44.6 34.9

44.7 41.6

46.8 56.1

44.4 39.8

5 43.3 36.3 33.9

44.4 36.7 32.6

47.0 40.5

33.9 26.4 25.9

6 33.4 25.3 24.3

30.8 24.0 25.2

31 .8 22.9 25.2

35.0 22.4 30.1

7 34.5 19.8 31 .1 37.7 24.3 34.8

30
37.8 23.9 33.6

32.6 21.4 30.4

8 33.2 21 .0 27.3

34.5 22.5 33.6

33.8

*Pieces from two different samples

Another pertinent observation regarding the results shown in Table 3.1.2 is that the

swell ratio as determined by volume change is, in most cases, significantly less than the

value found using A STM D2765. This result is due, at least in part, to the fact that in

ASTM D2765 the swell ratio is defined to be the volume of polymer in the gel plus the

volume of absorbed xylene divided by the volume of polymer in the gel, whereas in the test

based on dimensional changes the denominator is the sum of both the volume of polymer in

the gel plus the volume of polymer that is extractable, that is

swell ratio (ASTM) = v

g + v

s ,

swell ratio (dimensional change) = v

g + v

s , V_ +

9 ®

where represent the total initial volume.

In order to examine more closely the correlation between data obtained from the two

tests we show in Figure 3.1.1 the swell ratio as determined by ASTM D2765 plotted versus

the swell ratio found from dimensional changes. The vertical and horizontal lines with

bars represent the spread in values for each sample and the crossover point corresponds to the arithmetic mean for the two directions. The data suggest that there is nearly a linear correlation between the values for swell ratio as determined by the two types of

test. Although the number of specimens tested from each sample was too small for any

meaningful statistical analysis, we point out that, for the data given in Table 3.1.2, the

variances, on average, are smallest for the method employing dimensional changes and

largest for the percentage extractable measurements. We conclude, therefore, that the

determination of the swell ratio by careful measurement of dimensional changes can yield

comparable information concerning degree of crosslinking to that obtained with ASTM D2765,

and with equal reproducibility.

3.2 High Temperature Tensile Modulus

It has been suggested that the high temperature tensile modulus may be useful as a measure of degree of crosslinking in crosslinked polyethyl enes . Such a test would have

appeal from the viewpoint of time required to perform the test compared to swell tests.

In both conventional and crosslinked polyethylene the tensile properties are determined

in large measure by the crystal! inity , the more crystalline the material the higher the

tensile modulus. At temperatures below the melting point, the tensile modulus will be

rather insensitive to changes in degree of crosslinking. However, at temperatures above

the crystalline melting point, the mechanical integrity of the crosslinked material will

be maintained solely by the crosslinked network, which does not melt, and it is anticipated that the tensile modulus of the network alone will be sensitive to the extent of crosslinking. The remainder of this section will be concerned with experiments carried out in our laboratory in an effort to determine the feasibility of such a test as a means of

establishing degree of crosslinking.

Shown in Table 3.2 are several sample characteristics for sample 1-3 described earlier.

31
SWELL

RATIO

(A

STM

D2765)

50

Figure 3.1.1 - Correlation of swell ratio as determined by ASTM D2765 to that determined

from dimensional changes for samples of rotationally molded crossl inkable polyethylene

having different degree of crosslinking.

32
Table 3.2

Sample Characteristics of Crosslinked Rotationally. Molded Polyethylene

l Sample Density

3 (gm/cm )

esr

(3)

(hrs.

)

(4) Flexural Modulus '

(kPa)

1 0. 937-0. 940^

(0.941 ) (2)

200 760,000

2 0.930-0.933^

(0 . 942) (2)

>1000 . 690,000

3 0.940^ undetermined undetermi ned

(1) ASTM D1505 - As determined by resin manufacturer

(2) As determined in our laboratory by hydrostatic weighing - averaging of six specimens

(3) ASTM D1693 Standard Method of Test for Environmental Stress-Cracking of Ethylene

Plastics, Condition A - As determined by resin manufacturer

(4) ASTM D790 - As determined by resin manufacturer

In one set of experiments, specimens were prepared by first machining flat sheets

from each of the three samples. Dumbbell shaped specimens were then cut from each sheet

with a die so that the final dimensions of the straight portion of the specimen were approximately 2.75 cm in length, 0.32 cm in width, and 0.20 cm in thickness. These specimens were then tested to failure in uniaxial creep under static load conditions in a manner identical to that described in Parts I and III of this series of reports [1,3]. To be

consistent with terminology used in the previous reports we shall again define failure to be the time at which the specimen either necked or fractured due to cracking, whichever

occurred first. Results for these tests are presented in Figure 3.2.1 which shows the time

to fail in logarithmic coordinates plotted versus the apolied load, all data taken at 2 9 6 K

.

As is evident, there is little, if any_ difference in behavior for all three samples. It was also true that no significant difference occurred for specimens cut at either 45° or

90° to the direction for which the specimens shown in Figure 3.2.1 were cut. Compared to conventional polyethyl enes , the curve shown in Figure 3.2.1 for the crosslinked samples is even steeper than that for the ethylene-hexene copolymer given in Figure 2.1.3. We remarked

earlier in this section that the tensile modulus is not likely to depend critically upon the

degree of crosslinking at temperatures below the crystalline melting point. It is apparent

that the time to fail versus applied load behavior is also insensitive to the degree of

crosslinking even though the swelling ratio differs by as much as a factor of four for these

three samples.

Next, the tensile modulus (Young's modulus) was determined at 433K (160°C) for specimens taken from samples 2 and 3 described in Table 3.2. Of the three samples, these two

exhibited the greatest difference in swelling ratio and presumably, therefore, have the

greatest difference in degree of cresslinking. Measurements of the tensile modulus were accomplished with the aid of an Instron Tensile Tester Model 1130* equipped with a 500

gram load cell. In order to carry out the measurements at high temperatures an oven was specially constructed so that both the grips and specimen were well within the controlled

temperature zone. Temperature was measured with a temperature transducer device placed as close as possible to, but not touching, the specimen. The temperature gradient over the

length of the specimen was determined to be no more than about 2K (2°C). For this set of

measurements, the specimens were in the form of straight strips rather than dumbbell shaped

as in the time to fail study.

After mounting the specimen in the grips and setting the oven in place, the specimen

was slowly heated to 423K (160°C), a temperature well above the highest melting temperature

for the crystalline material. Due to thermal expansion of the polymer during heating,

particularly through the melting region, some difficulty was encountered in maintaining the

specimen in a zero load condition prior to running the test. On our machine, this could be

*Certain commercial materials and equipment are identified in this report in order to specify adequately the experimental procedure. In no case does such identification imply rec- ommendation or endorsement by the National Bureau of Standards, nor does it imply necessarily the best available for the purpose.

33
Figure

mol ded

3.2.1 - Log time to fail versus applied stress for different samples of rotationally

cross! inkabl e polyethylene at 296K. A - Sample 1, - Sample 2, Q - Sample 3.

34
done only by manual adjustment of the crosshead. This problem would be much less troublesome.if the tests were carried out on a machine having a load control capability. Once the

specimen had come to temperature equil ibrium, the gage length was determined by measurement of the grip separation and the tensile test was then run. The tensile modulus was determined from the. initial slope of the stress-strain curve. Ideally, at small deformations, there is a linear portion of the stress-strain curve for which the initial slope

represents one measure of the tensile modulus. However, in the present case, the stressstrain curves were nonl inear, even at small deformations, so that it was difficult to determine the initial slope. Therefore, for the type of materials under consideration

here, this particular method for determining the tensile modulus must be considered as rather imprecise. For samples 2 and 3, the values obtained for the modulus at 433K by

this method were 7.5 x 10^ N/m^ and 5.5 x 10^ N/m^ respectively. Considering the large

difference in swelling ratio for these two materials it must be concluded that even the

modulus as determined at high temperatures is rather insensitive to the amount of cross1 inking.

As a result of the small difference in modulus observed for the two samples, and the

high degree of nonlinearity in the stress-strain curve, it was decided to examine these

same materials in stress relaxation. The results of this effort are presented in Figure

3 . 2.2 which shows the stress plotted versus log time for times up to 1Q3 seconds (% 15

minutes). Also included for comparison are two additional cross! inked materials, one an irradiation crosslinked polyethylene, and a vulcanized natural rubber. The data should not

be compared quantitatively since the levels of strain at which the experiments were con- ducted were not the same for each sample. The point to be made is that all four samples of

crosslinked polyethylene show considerable relaxation over the entire time scale covered.

In order to properly determine the modulus of the crosslinked network, the stress value

should plateau, as it did for the vulcanized rubber after only a few seconds. The plateau

region should be sensitive to the degree of crosslinking present. For the four samples of

crosslinked polyethylene examined, the stress either will plateau at much longer times than

those investigated, or will continue to relax indefinitely. In either case, this type of

measurement does not appear practical as a test method, both with regard to sensitivity or time savings.

4. Summary

We have investigated the failure behavior of several types of commercial polyethylene

over a wide range of stress and deformation histories, both in air and in stress-cracking

agent. It has been shown that the failure mechanism at long times is different from that

that short failure times, and, in order to predict long time behavior based on accelerated

tests, testing should be done in the region of stress and for deformation appropriate to the failure mode. It has further been demonstrated that a primary factor in determining

stress-crack resistance is the ability of the polymer to perform in air, and not neces- sarily the rate at which failure is accelerated due to the presence of the stress-cracking

agent. From the variety of experiments performed (i.e. uniaxial creep, equal biaxial under

inflation, bottle inflation, and bent strip geometry) it is apparent that essentially the same information with regard to stress-crack resistance can be obtained from all the methods

employed. On this basis we have proposed one possible general test method which combines

a number of the more desirable features of the various methods examined, and is a method

which we believe has several advantages over ASTM D1693.

In the case of crossl inkable polyethylenes for rotational molding, two possible tests

for determining the degree of crosslinking have been examined, tests which might realize

a significant time savings over that required to carry out current swell tests. It has

been found that a test based on the determination of the high temperature tensile modulus

is impractical on several counts. On the other hand, the swell ratio as determined from

dimensional changes appears to give an equally reliable measure cf the degree of cross- linking as current standard test procedures for determining swell ratio or percent extractables and requires less time to perform. Finally, any activity toward the adoption of test

methods such as those developed through this work should be stimulated from within the consensus standards community.

35
Figure 3.2.2 - Stress relaxation behavior of rotationally molded crosslinked polyethylene.

36
References

[1] J. M. Crissman, C. M. Guttman, and L. J. Zapas, "Performance of Plastic Packaging for

Hazardous Materials Transportation Part I. Mechanical Properties", Report No. D0T/MTB/0HM0-76/4 , available from National Technical Information Service, Springfield,

VA 22151.

[2] J. D. Barnes and G. M. Martin, "Performance of Plastic Packaging for Hazardous Materials

Transportation Part II. Permeation", Report No. D0T/MTB/0HM0-76/5, available from

National Technical Information Service, Springfield, VA 22151.

[3] J. M. Crissman and L. J. Zapas, "Performance of Plastic Packaging for Hazardous

Materials Transportation Part III. Stress Cracking", Report No. D0T/MTB/0HM0-77/4,

available from National Technical Information Service, Springfield, VA 22151.

[4] J. D. Barnes, G. M. Martin, and F. L. McCrackin, "Performance of Plastic Packaging for

Hazardous Materials Transportation Part IV. Standardizing Permeation Measurements",

Report No. DOT/MTB/OHMO-77/5, available from National Technical Information Service,

Springfield, VA 22151.

[5] ASTM 1978 Annual Book of ASTM Standards, Parts 35 and 36, American Society for Testing

and Materials, 1916 Race Street, Philadelphia, PA 19103.

[6] Unpublished work.

[7] L. J. Zapas and J. M. Crissman, Polymer Engineering and Science 1_9, No. 2, 104 (1979).

[8] J. M. Crissman and L. J. Zapas, ACS Symposium Series, No. 95, Durability of Macromolecular Materials, R. K. Eby, Editor, 289 (1979).

[9] J. L. Glathart and F. W. Preston, J. Appl . Phys. V7, 189 (1946).

[10]

J. D. Ferry, Viscoelastic Properties of Polymers, Chapter*17, Second Edition, John

Wiley and Sons, New York, 1970.

37

Technical Report Documentation Pag

1 . Report No. D0T/MTB/0HM0-80/2

2. Government Accession No. 3. Recipient’s Catalog No.

4. Title and Subtitle PERFORMANCE OF PLASTIC PACKAGING FOR HAZARDOUS

MATERIALS TRANSPORTATION: PART V MECHANICAL

PROPERTIES

5. Report Date December, 1979

6. Performing Organization Code

NBSIR 79-1938

3. Performing Organi zation^Report No.

7. Author's) John M. Crissman, Louis J. Zapas and Gordon M. Martin

9. Performing Organization Name and Address National Bureau of Standards

U. S. Department of Commerce

Washington, D. C. 20234

10. Work Unit No. (TRAIS)

11. Contract or Grant No. DOT AS-50074

13. Type of Report and Period Covered Final

Oct. 1978 - Oct. 1979

12. Sponsoring Agency Name and Address U. S. Department of Transportation

Materials Transportation Bureau

Office of Hazardous Materials Research

Washington, D. C. 20590

14. Sponsoring Agency Code

15. Supplementary Motes

16. AbstractThis report represents Part V in a series of reports prepared for the U. S.

Department of Transportation Office of Hazardous Materials Research under DoT

Contract AS-50074. The report summarizes experimental work done in the area of

mechanical properties _ of polymeric materials used in the manufacture of large

shipping containers for the transport of hazardous materials. The effort has been

directly toward (1) a continuation of studies on the stress-crack resistance of

different types of polyethylene, and (2) an examination of various test methods

for determining the degree of crosslinking in crossl inkabl e polyethyl enes for

rotational molding. The stress-crack resistance of polyethylene has been investig;

ed for four widely different deformation histories, including uniaxial creep, equa'

biaxial under inflation, bottle inflation, and bent strip geometry. On a basis of

information gained from these studies, a test for stress-crack resistance is proposed which combines desirable features from the various methods examined. In the

area of crossl inkable polyethylenes , two possible tests for determining degree of

crosslinking have been examined. One method involving the determination of the

high temperature tensile modulus was found to be impractical, while a swell test

based on dimensional changes was found to yield equally reliable results to those

obtained by current standard test methods.

17. Key Words Crossl inkabl e polyethylene; failure

mechanisms; high temperature tensile

modulus; instability under deformation;

mechanical properties; polyethylene; stress- Springfield, VA 22151

crack resistance.

13. Distribution Statement Documentation is available to the

U. S. public through the National

Technical Information Service,

19. Security Clossif. (of this report) Unci assified

20. Security Clossif. (of thi s page)

Unclassified

21* No. of P ages 38

22. Price

Form DOT F 1700.7 (8-72) Reproduction of completed page authorized