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NAT'L INST. OF STAND & TECH R.I.C.

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NBSIR 81-2209 (FDA)

Relationship Between Morphology

and Mechanical Properties of

Ultra High Molecular Weight

Polyethylene

G. B. McKenna, Ph.D.

Principal Investigator

F. A. Khoury, Ph.D.

Project Leader

J. M. Crissman, Ph.D.

Project Leader

First Annual Report for the Period

December 9, 1 980 - September 30, 1 980

Task 80-01, NBS-BMD Interagency Agreement

April 1981

^ °'

a \t*
AkA :
NATIONAL Hl/RCAU

Or STANDARDS

LIBRARY

DEC 7 198 1

NBSIR 81-2209 (FDA)

RELATIONSHIP BETWEEN MORPHOLOGY

AND MECHANICAL PROPERTIES OF

ULTRA HIGH MOLECULAR WEIGHT

POLYETHYLENE

G. B. McKenna, Ph.D.

Principal Investigator

F. A. Khoury, Ph.D.

Project Leader

J. M. Crissman, Ph.D.

Project Leader

First Annual Report for the Period

December 9, 1 980 - September 30, 1 980

Task 80-01, NBS-BMD Interagency Agreement

April 1981

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director
\ j* 'u'f <in -V. " •
Abstract

Aspects of the morphology of the constituent particles of the raw

ultra high molecular weight polyethylene (UHMWPE) used in this study have

been examined using scanning electron microscopy. In addition differential

scanning calorimetry has been used to determine the melting point (~413K)

and the crystallinity (78%) of the raw polymer. Protocols have been

established for the preparation of compression molded sheets of UHMWPE having

crystallinities of 50% and 60%. A memory of the particulate nature of the

raw polymer is retained in the sheets. There was a considerable variation in

the sizes of the birefringent structures in the sheets, the majority were extremely

small and barely visible under the polarizing microscope, the larger among these

structures appeared to be axialitic or incipiently spherulitic. The following

aspects of the mechanical behavior of the molded sheets have been examined:

Stress-strain behavior at constant rate of clamp separation, creep and single step

stress relaxation in uniaxial extension, and failure under both static and

sinusoidal loading conditions. The relaxation modulus of the lower density

material (quenched) was found to be smaller by a factor of two than that of the

higher density material (slow cooled) . The short time creep of the quenched

UHMWPE was significantly greater than that of the slow cooled polymer; however,

the limiting creep strain prior to failure was nearly the same for both materials.

Low cycle fatigue data obtained on the slow cooled UHMWPE suggests that it may

be possible to characterize the failure behavior of this polymer using a cumulative

damage rule. The lifetime of the UHMWPE under zero-tension sinusoidal loading

is 6-7 times longer than the lifetime in static (creep) tests in agreement with

the cumulative damage rule. Over the range of frequencies tested (0.0002-0.01 Hz)

limited data show no trend of lifetime dependence on test frequency. Results of

the failure testing indicate that the static lifetime of UHMWPE may be a lower
bound on the zero-tension fatigue lifetime. A study of the fine structure

changes resulting from uniaxial deformation under constant rate of clamp

separation has also been started. In addition an apparatus has been built for

the study of the equal biaxial deformation of UHMWPE sheets in air as well as

in the presence of stress-cracking agents. Experiments, using this equipment

are currently under way.
Table of Contents

Abstract i

Table of Contents iii

1. Introduction 1

2 . Background 2

3. Experimental Methods and Results 5

3.1 Raw Polymer Characterization 5

(a) Morphology 5

(b) Melting Behavior 7

(c) Crystallinity 8

3.2 Specimen Preparation and Characterization 8

(a) Molding Conditions and Specimen Densities

and Morphology 8

3.3 Mechanical Properties 14

(a) Stress-Strain Behavior at Constant Rate of

Clamp Separation 14

(b) Uniaxial Creep 19

(c) Stress Relaxation 20

(d) Failure Studies 23

(e) Equal-Biaxial Deformations Under Inflation and

Environmental Stress Cracking 27

(f) Miscellaneous 29

4. Summary 30

5. Work Plan for FY 1981 34

Appendix A 36

References 37

F igures 40
.
1. Introduction

Ultra high molecular weight polyethylene, hereafter referred to as

UHMWPE, is used in a variety of orthopedic implant devices because of its

chemical inertness and mechanical durability. In some applications, for

example the acetabular component of the total hip replacement, UHMWPE has

proven successful, and under proper clinical conditions may perform effectively for

15 to 20 years. However in other applications, such as knee and ankle

protheses, UHMWPE components have had only limited success. The increased

use of such devices in more youthful patients will require material

performance beyond 20 years. In addition to the materials aspect of overall

performance there is clearly a need for better design particularly for knee

and ankle protheses. At the present time there is little evidence of new and

better materials appearing in the marketplace. It remains to be determined

whether UHMWPE has been exploited to the limits of its capabilities. Insights

into what these limits may be requires detailed knowledge and understanding

of the time dependent mechanical behavior of the polymer and the influence of

factors such as molecular weight, crystallinity and morphology on its mechanical

behavior

.

The work being carried out under task 80-01, NBS-BMD Interagency

Agreement addresses principally the relationship between morphology and the

time dependent mechanical properties of UHMWPE. It is our view that the

accomplishment of the primary objectives outlined below will be of significance

both to the FDA Bureau of Medical Devices and to the medical device industry.

These objectives are:

(1) To investigate and determine the influence of polymer morphology on

the time dependent mechanical properties of UHMWPE, and thereby identify

the fine structure parameters which significantly affect long term

performance

.
(2) To obtain data on creep and stress relaxation (both in compression and

tension) behavior, as well as time dependent failure. Such

data will be useful as a basis for the design and evaluation of current

and future devices made from UHMWPE.

(3) To devise measurement techniques relevent to the development of

standard test methods for characterizing the time dependent mechanical

properties of implants based on UHMWPE.

2 . Background

One obvious factor which limits the expected lifetime of hip, knee, or

other joint protheses is the durability of the UHMWPE component. In clinical

use, failure of the polymer component is not a result of wear alone, but may

occur through creep [l]*, fatigue associated with wear [l,2], or even fracture

[3,4]. The length of time between device implantation and clinical failure

will depend upon the time dependent mechanical response of the material in

the clinical loading situation.

It is known that the time dependent mechanical response of semi-crv.stalline

polymers is highly affected by such structure or morphology related

features as degree of crystallinity, orientation, and spherulite size and

texture [5,6]. For example, in work performed in our laboratory it was found

that the stress relaxation behavior of UHMWPE, not unlike that of lower molecular

weight polyethylenes , is highly sensitive to annealing [7], Through annealing

of an as received sample the percent crystallinity (as determined from density

measurements) increased from 56% to 67%; at the same time the one second torsion

modulus increased from 280 MPa to 350 MPa. More importantly, the rate of

stress decay was significantly lower in the annealed material.

* Numbers in brackets refer to references located at the back of this report.
3

In the case of UHMWPE there are additional factors which must be

considered when studying the dependence of mechanical properties on

morphology. One such factor is the presence of a "grain memory" [7,8].

By grain memory we mean that the finished part has a fine texture which

reflects, in part, that of the nascent polymer. It is possible that the grain

memory affects performance. Another factor affecting performance is the

processing, that is, the temperature and pressure history used in manufacture.

A thermal history which brings about an improvement in one property may actually

result in the diminution of another. For example, increasing the degree of

crystallinity of low molecular weight polyethylene increases the modulus, yet

at the same time it decreases the resistance to brittle fracture and stress

cracking. Also, because of the very high molecular weight of UHMWPE the

melt viscosity is very high. The proper fusion of the polymer grains therefore

becomes an important aspect of processing. In a study of the impact fatigue

of various commercial grades of UHMWPE, Bhatega, Rieke, and Andrews [9]

reported that the poor performance of one material was associated with

incomplete fusion of the polymer powder. Finally, in a study of wear surfaces

[10] , cracking has been observed to occur at intergrain boundaries.

It is apparent that the time dependent mechanical behavior of UHMWPE

depends upon morphology and/or processing in a variety of ways. In order to

better understand this dependence we are studying the creep, stress relaxation,

and dynamic and static fatigue behavior of specimens of compression molded

material having significantly different degrees of crystallinity. At the

same time the morphology of these specimens is being examined as a function

of both the preparation conditions and deformation history. One objective of

this work is to determine the range of material property variability which
can be introduced by changes in crystallinity.

A word is in order concerning the types of tests commonly used to

evaluate the mechanical properties of polymers, for example those recommended

by ASTM. Such tests tend to measure short term behavior, and the results

obtained may not predict long term performance. Based on recent

work performed in our laboratory the following examples serve to illustrate

this point. In low molecular weight polyethylenes the failure mode in

uniaxial creep depends upon the magnitude of the applied stress. At high

stresses (short times to failure) the material undergoes large deformations

involving necking, cold-drawing, and eventual fracture. On the other hand

at low stresses (long times to fail) failure occurs by cracking and fracture

at small elongations [ 11 ] . In another study of the fatigue behavior of low

molecular weight polyethylenes it was found that the failure mode depended

upon the test frequency, changing from a ductile mode at low frequency to a

brittle-like mode at higher frequency [l2] . The behavior of UHMWPE under

conditions comparable to those just described remains to be determined. In the

case of the low molecular weight polyethylenes it was also found that the

lifetime was shorter under static loading than it was under cyclic loading

(for an equivalent peak load) . If this behavior is also true of the UHMWPE,

then the determination of fatigue life under a given load history is

simplified since the static measurements would represent a lower bound on the

fatigue life. Static testing is significantly easier to perform and is less

time consuming than is dynamic testing.

Finally, we recognize that wear is an important aspect of performance of

devices made from UHMWPE. Wear is a phenomenon which is influenced by a

number of factors including the time dependent mechanical properties of the material,

as well as the chemical and load environments [13] . Wear in prosthetic

components made from UHMWPE appears to be accompanied by both large scale

deformation and cracking of the surface region [14,15]. In our view, the wear

phenomenon cannot be understood until morphology-mechanical property relationships
5

are better established.

3 . Experimental Methods and Results

3 . 1 Raw Polymer Characterization

Unless otherwise stated the compression molded samples prepared

for the studies described in this report were all made from the same

batch of commercial implant grade ultrahigh molecular weight

polyethylene. For purposes of identification this batch will

henceforth be identified as UH-A. The raw polymer as received was in the

form of a powder (see below) , and to our knowledge contains no additives

such as antioxidant or stabilizers. The intrinsic viscosity (limiting

•k viscosity number) of UH-A is approximately 25 , which, based on the

manufacturer's method of estimating molecular weight from dilute solution

£ viscosity measurements, corresponds to a molecular weight of ~ 4x10 .

Three aspects of the raw polymer were examined namely, its morphological

characteristics, its melting behavior, and its crystallinity. These

features are described below:

(a) Morphology.

The polymer, as received, is in its as polymerized (i.e. nascent) state

and consists of fine particles. As indicated in Section 2, earlier studies

[7,8] of the morphology of both molded and machined prosthesis components

made from UHMWPE have shown that a memory of the particulate (or grain-like)

nature of the raw polymer is retained even though the polymer was heated

well above its melting point in the manufacturing process. This feature,

The determination of intrinsic viscosity from dilute solution viscosity

measurements is subject to a number of problems and uncertainties which

have been discussed in detail by Wagner and Dillon [l6j. The definition

°f "ultra high" molecular weight is under consideration by two ASTM

committees. Committee F-4 on Implant Devices and Committee D-20 on Plastics.

In the case of F-4 the criterion for use of the term 'ultra high' is that

the intrinsic viscocity should be « 15 or higher, whereas in the case of

it should be ~ 21 or greater. The UH-A meets both criteria.
6

which may be attributed to the extremely high viscosity of the

molten polymer, was readily observed in thin (~50 ym) cross-sections

examined under a light microscope at low magnifications using phase

contrast optics. Under those conditions, the boundaries between the

compacted nascent particles were seen throughout the cross-sections.

This observation emphasizes the importance of characterizing the

structure of UHMWPE products, not only to determine their crystallinity,

axialitic or spherulitic texture, and lamellar fine structure, but also

to monitor those aspects of their morphology which are directly related

to, and a consequence of, the particulate nature of the raw polymer and

its high melt viscosity. Knowledge of the morphology of the raw polymer

is therefore a necessary aspect of the present study since the nature of

the shapes, sizes, and fine textures, of nascent polyethylene particles

vary depending on the polymerization conditions [17,18].

Scanning electron micrographs showing nascent particles of UH-A are shown

in Figs, 1, 2a, and 3a, Portions of the particles in Figs. 2a, 3a, are

shown at higher magnifications in Figs. 2b, 2c, and Figs 3b, 3c respectively.

Prior to examination the particles had been sputtered in an Argon

atmosphere with an Au:Pd (.40%/ 60%) alloy in order to prevent charging

upon exposure to the electron beam.

Numerous particles can be seen in Figs, la, lb which illustrate their

irregular shapes and varying sizes. Their sizes generally range from ~ 80 ym

to ~ 300 ym. It is evident from these figures as well as Figs. 2a, 3a

that the surfaces of the particles are highly convoluted. At low magnification

(e.g. Fig. 1) the external appearance of the individual particles resembles

that of compact bunches of grapes. Examination at higher magnification
7

reveals that whereas portions of the surface of the particles are

relatively smooth, others exhibit a highly porous appearance, as can

be seen in Figs, 2a, 2b, and Figs. 3a, 3b. At even higher magnifications

(e.g. Figs 2c, 3c) the gaps or voids in the porous regions can be seen

to be bridged by thin fibrils which are 50 nm-100 nm in diameter. The

yarious origins of fibrils in nascent polyolefins have been reviewed bv

Marchessault et. al. [17]. It appears that in the case of the UH-A

particles, the bridging fibrils result from the deformation of the polymer

during polymerization as a consequence of the progressive expansion of the

particles caused by the continued generation of fresh polymer at the surface

of underlying encapsuled catalyst.

(b) Melting Behavior,

Differential scanning calorimetry (DSC) was used to examine the

melting behavior of the UH-A raw polymer. Unless otherwise specified the

heating rate used during the melting scan was 10°C/min. At this heating

rate the temperature correction for machine lag was determined to be Q.47K.

For this particular UHMWPE (i.e. UH-A) the temperature corresponding to

the maximum in the melting peak was 418. 3K (145°C) . Upon subsequent

cooling and remelting the temperature corresponding to the peak maximum in

the melting peak dropped to the 408K-411K (135°C-138°C) range depending

upon how rapidly the specimen was cooled following the first melting. It

is typical of crystalline polymers such as polyethylene that their melting

points vary according to their previous thermal and mechanical histories,

and are lower than the equilibrium melting point, i.e. the melting point

of an ideal extended chain crystal of the polymer (~419K, 146°C) [19] . The

observed melting point of the raw polymer was , however , unusually high. This

suggested the manifestation of ’superheating’ at the heating rate of 10°C/min.
8

In order to determine whether this was the case a series of experiments was

carried out in which specimens of fresh raw polymer were heated to various

temperatures at which partial melting occurred. The specimens were then held

for two hours at the chosen "dwell" temperature in each case after which

heating was continued. The dwell temperature corresponding to which no

further melting was observed upon subsequent heating was found to be

between 413K and 414K (.140°C-141°C) , thus indicating that 'superheating'

occurred at a heating rate of lOK/min and that the true melting point of

the raw polymer is at least 5K lower than 418. 3K.

(c) Crystallinity.

The crystallinity of the raw UH-A polymer was determined by DSC to

be 78% using the procedure outlined in Appendix A. As will be seen in the

following section this crystallinity is much higher than that in

compression molded sheets made from UH-A which were slowly cooled from

the molten state.

3 . 2 Specimen Preparation and Characterization

(a) Molding conditions and specimen densities and morphology.

Because of its very high molecular weight and consequent very

high melt viscosity the type of UHMWPE polymer used in this study cannot

be processed by conventional techniques such as injection molding or

standard extrusion. The sheets of polymer used in this study were

successfully prepared by compression molding. Initially, sheets were

molded according to the procedure recommended by the polymer manufacturer.

This method consists of a ten step process in which the polymer is heated

to temperatures as high as 478 to 493K (205 to 220°C) . However, the

sheets produced in this manner showed considerable yellowing around the

edges, indicative of degradation. As a result, a good deal of effort was

devoted to the design of a vacuum mold and the development of molding
| - 9 -

procedures by which specimens suitable for the mechanical

testing could be prepared in a reproducible manner. After

experimenting with different molding temperatures and pressures the

following standard molding procedure was chosen in order to

produce circular flat sheets approximately 15 cm in diameter and 0.1 cm

thick.

* *

(]_) The vacuum mold is filled with a predetermined amount

l| of UHMWPE powder and placed in the press. The heat

is then turned on, with only light contact pressure applied

I on the mold.

(2) When the platens reach 408K (135°C) a pressure of 2.4 MPa

(.350 PSI) is applied on the mold for a short period of

time and then released. The mold chamber containing the

polymer is then subjected to vacuum.

(3) The platens and mold are heated slowly to 473K (200°C)

and held at that temperature for ten minutes.

(.4) A pressure of 8.25 MPa (1240 PSI) is applied to the mold

and the heat turned off.

(5) The press is then allowed to cool to room temperature before

removing the mold. (Between 200°and 100°C the rate of cooling

in the press was approximately l°C/min.)

Specimens prepared in this manner had densities in the range 0.934-

3 0.936 g/cm . These densities, which were determined with a density

gradient column at 293. 3K, correspond to weight % crystallinities ,y , of

58%-60% (see Apendix A)

.

The bulk properties of conventional low molecular weight polyethylenes

are determined primarily by the crystallization conditions which govern the
10

crystallinity and the morphology of molded specimens. Although higher

dwell temperatures above the melting point, and longer dwell times at

these temperatures, tend to lead to a relative diminution of nucleation

density (nuclei/unit volume) during subsequent crystallization at low or

moderate undercoolings, the effect of melting conditions on the

crystallinity of samples subjected to the same crystallization histories

is negligible or secondary in the absence of any imposed extensional flow.

A complication does arise however in the case of the UHMWPE. As pointed

out in Sections 2 and 3 . la, previous studies have shown that a memory of the

particulate nature of the raw polymer is retained in compression molded

samples. Since one would expect that the mechanical properties of the

UHMWPE will be influenced by the effectiveness of the compaction, fusion,

and "adherence' of the raw polymer particles, the question which arises

is "what is a suitable maximum temperature at which comnression molding

should be carried out? ".This aspect was examined as part of the process which

led to the choice of the set of molding conditions described above. Sheets

of UHMWPE were molded using a maximum temperature which was varied from

463K to 503K with all other conditions held the same. The densities of

3 the sheets were in the range 0.934-0.936 g/cm . Furthermore creep and

single step stress relaxation measurements were made on samples from sheets

molded at up to 483K. Within experimental error the sheets molded at

different temperatures exhibited similar mechanical behavior. We have

chosen 473K as our molding temperature. It is sufficiently high to permit

proper flow of the material under pressure. At the same time there was no visual evidence of discoloration due to thermal degradation. This choice of

molding temperature must be considered as somewhat arbitrary. It is, however,

close to the low end of the molding temperature range recommended by the

polymer manufacturer. It should be noted that as part of the background work

leading to the choice of molding conditions, raw polymer samples were subjected.
11

in a D.S.C. apparatus, to thermal treatments simulating the molding and

crystallization procedures described above, i.e., the molding temperature

prior to cooling was varied in the range 463K-503K. The melting behavior

of the samples was found to be independent of the molding temperature used

in that range.

Cross-sections approximately 50]im thick were cut at right angles

to the plane of sheets made from the UH-A UHMWPE which had been molded

at various temperatures in the range 463K-503K and which had subsequently

been cooled slowly to room temperature as described above (as pointed out

earlier the crystallinities of the sheets were 58%-60%) . The cross-sections

were examined under a light microscope using phase contrast optics. They

were also examined between crossed polarizer and analyzer in a polarizing

microscope.

Phase contrast micrographs of cross-sections microtomed from sheets molded

at 463K, 473K and 503K are shown in Figs. 4(a,b,c). These figures clearly

illustrate that a memory of the original coarse irregularly shaped

particulate character of the raw polymer is retained in the molded

sheets even in the case when molding was carried out at ~ 90K above the

melting point of the raw UH-A polymer. The pointed markers in Fig. 4

identify portions of a few interparticle boundaries.

Micrographs of portions of cross-sections of the same sheets molded

at 463K, 473K, and 503K, are shown in Figs. 5, 6 and 7 respectively, all

of which figures are at a higher magnification than in Fig. 4. Figs.

5a, 6a and 7a are phase contrast micrographs whereas Figs. 5b, 6b, and 7b

represent, respectively, the same areas between crossed polarizer and

analyzer. Figures 5b, 6b and 7b show that there are substantial variations

in fine structure in each cross-section as evidenced by the differences in the

size of the birefringent structural units from one region to another in each

picture. In some regions the birefringent units are barely perceptible and
12

the corresponding areas exhibit a very fine grainy appearance under

the polarizing microscope.

Although large well developed spherulites were not observed

in these samples, rotation of the samples about the microscope axis

while keeping the crossed polarizer and analyzer fixed indicated that some

of the larger birefringent structural units were incipiently spherulitic or

axialitic in character. Detailed studies of the effects of molecular

weight and undercooling on the crystallization of polyethylene fractions

from the melt have been reported by Hoffman et . al . [20] and Maxfield

et. al [2l]. A discussion of the morphology of polyethylene fractions

is beyond the scope of this report. In the present context it suffices

to indicate that the observed fine structures exhibited by the crosssections in the polarizing microscope suggest that the absence of well

developed spherulites is a consequence of a high nucleation density

with the result that evolving spherulites impinged upon their nearest

k

neighbors before attaining 'mature' spherically symmetrical shapes.

The variations in the sizes of the birefringent structural units in the

cross-sections shown in Figs. 5b, 6b, and 7b, are indicative of

variations in nucleation density in each sample. Such variations may be

attributed to localized variations in the compaction and flow of the raw

polymer particles during the high temperature cycle of the compression

molding process.

In addition to the types of slow cooled sheets described above,

several sheets were prepared by rapidly quenching UHMWPE in the following

manner. A sheet which had already been molded from UH-A polymer at 473K

* Spherulites in polymers do not usually develop in a spherically

symmetrical manner from the moment of inception. See reference [22]

for a review and discussion of the nature of the early transitional

stages of evolution of spherulites.
13

and slow cooled was placed between chromed metal photographic plates

which were then placed in a press preheated to 473K. The plates were

maintained at 473K for ten minutes under slight pressure and were then

quickly removed from the press and immediately submerged into cold

3 water. Sheets produced in this manner had a density of about 0.923 g/cm

which corresponds to a weight per cent crystallinity of 51% as compared

to a value of 58%-60% for the slow cooled sheets. As was the case for

the slow cooled samples, examination of cross-sections of quenched sheets

under a light microscope revealed the retention of a memory of the

particulate nature the original raw polymer even though those sheets had

been subjected to an additional melting cycle. Furthermore, the quenched

samples also exhibited under the polarizing microscope variations in the

sizes of the birefringent structures in them.

«

In summary, protocols for preparing compression molded sheets of

UHMWPE having weight per cent crystallinities of close to either 50% or

60% have been established. A maximum molding temperature of 473K has been

selected. All the sheets were found to exhibit a memory of the particulate

nature of the raw polymer and, on a finer scale, substantial variations were

observed in the sizes of the birefringent structures in all the sheets.
14

3 . 3 Mechanical Properties

During the past year studies on mechanical properties were

started in the following areas: stress-strain behavior at constant

rate of clamp separation, creep and single step stress relaxation in uniaxial

extension, and failure under both static and sinusoidal loading conditions.

In addition preliminary work was done in equal biaxial deformation (under

inflation) and on stress crack resistance. The emphasis thus far has been

on the characterization of the mechanical behavior at room temperature of

specimens slowly cooled from the melt. Some measurements have however

been carried out on quenched sheets.

(a) Stress-Strain Behavior at Constant Rate of Clamp Separation

The specimens used for constant rate of clamp separation experiments were

dumbbell shaped having a length in the gage section of 1.3 cm, a width of

0.32 cm, and a thickness of approximately 0.10 cm. In one set of experiments

stress-strain curves were obtained for specimens slowly cooled from

one of a series of melt temperatures ranging from 463K to 503K. It was

found that the molding temperature had no significant influence on the

stress-strain curve, in agreement with the findings discussed in Section 2.

Curves representative of different rates of clamp separation at 296K are shown

in Figure 8; the arrows indicate the point at which fracture occurred.

These curves are typical of polyethylenes in general in that they are

rate dependent. For the range of rates used in this study the breaking

strength (stress at break) increased with increasing elongation rate whereas

the elongation at break decreased. One interesting feature of the behavior

of the UHMWPE, unlike that of lower molecular weight or some crosslinked

polyethylenes, is that it does not neck but rather elongates in a uniform

manner (on a macroscopic scale) right up to the point of fracture. Nusbaum
15

and Rose [23] claim to have observed necking in the fatigue testing of

radiation sterilized UHMWPE. Whether the necking occurred as a result of

the crosslinking or possibly was due to a significant reduction in

molecular weight because of irradiation is unknown.

An indication of the effect of temperature on the stress-strain

behavior of a slowly cooled specimen is illustrated in Fig. 8 in which a

curve determined at 310K (37°C) is also shown. Comparison with the curve

obtained at 296K at the same rate of clamp separation shows, not unexpectedly,

that the elongation at break increased, and the breaking strength decreased

with increasing temperature.

One aspect of the behavior of the slow cooled polymer sheets which is

of interest in the context of implant performance is the appreciable

recovery (contraction) exhibited by stretched samples when they are

released at various strains below the breaking strain, as well as after they

fail. For example the following observations were made in a series of

experiments carried out at 310K. Two samples were stretched at a clamp

separation rate of 0.04 in/min (corresponding to a strain rate of ~0.02/ min).

They both failed at a strain of 3.8. The broken samples were then kept at

room temperature for twenty-four hours. The lengths of the broken parts

were found to have contracted in that period to a residual deformation

(residual strain relative to original undeformed length) of 2.4 - 2.5.

In addition, a set of five samples from the same sheet were stretched at

the same rate to strains of 0,25, 0.50, 1.0, 2.0 and 3.0 respectivelv.

The samples were then released and kept at room temperature for twenty-four

hours after which time they had all contracted substantially. The

corresponding residual deformations were 0.06, 0.12, 0.34, 0.92, and 1.15.

The observations will eventually be compared with the behavior of annealed as well as quenched samples treated under similar conditions. Our overall objective

is to examine the effect of the crystallinity of UHMWPE on the recovery of

the polymer from uniaxial extension, and to determine the nature of the fine

structural changes associated with the stretching and recovery processes.
16

Wide angle and small angle x-ray diffraction methods and scanning

electron microscopy will be among the methods used to examine the nature

of the permanent and recoverable changes in the fine structure of the UHMWPE

caused by uniaxial extension. Information derived from such experiments

will also have a bearing on the structural changes which occur in creep,

single step stress relaxation, and fatigue experiments.

Initial observations derived from wide angle and small angle x-ray

diffraction patterns exhibited by the above mentioned slow cooled samples

which were uniaxially stretched at 310K to strains of 1.0, 2.0, and 3.0 and

then released are summarized below. As pointed out above the residual

strain in these samples after 24 hours at room temperature were 0.34, 0.92

and 1.15 respectively. The wide angle diffraction patterns in Fig. 9 (b,c,d

and the small angle patterns in Fig. 10(a,b,c) were obtained during a period

of two to three weeks later. The residual strains were not examined at the

times the x-ray patterns were recorded, they were however measured after

about eleven weeks and had by then decreased to 0.26, 0.72 and 0.91 respect!

In short, the residual strains in the samples at the time the x-ray patterns

were taken were somewhat smaller than 0.34, 0.92 and 1.15, but certainly not

less than 0.26, 0.72 and 0.91 respectively. For the sake of brevity in the

ensuing discussion of the changes observed in the diffraction patterns with

increasing residual strain, we shall identify each of the samples by a twonumber code denoting the maximum strain imposed at 310K followed by a bracket

number corresponding to the residual strain after release for 24 hours at

room temperature. The manifestation of slow, long term decreases in the

residual strain in stretched samples after they are released is a feature

which will have to be closely monitored in future studies on the relationship

between structural changes and strain.

A wide angle x-ray diffraction pattern from an unstretched slow cooled

sheet is shown in Fig, 9a. Figs. 9, b,c,d are diffraction patterns obtained

from the 1.0(0.34), 2.0(0.92) and 3.0(1.15) samples respectively. The patter::

were recorded with a flat plate camera. Nickel filtered Cu K radiation

a
17

was used. The x-ray beam was perpendicular to the plane of the gage portion

of the deformed dumbbell shaped specimens. The intense 110 and 200

diffraction rings of the common orthorhombic crystalline form of polyethylene

are identified in Fig. 9a in which a diffuse halo, H, due to the amorphous

regions, can also be seen.

Confining ourselves to the azimuthal changes in the 110 and 200

orthorhombic reflections resulting from stretching, a comparison of

the four diffraction patterns shows that as the permanent strain in the

samples increases, there is a progressive increase in the intensification

of the 200 reflection in the equatorial region. Correspondingly the

110 reflections become broadly split with the most intense portions of

the azimuthally broad arcs tending to be closer to the equatorial

direction with increasing strain. Somewhat similar observations have

been previously reported for lower molecular weight polyethylenes [24,25].

In addition to the changes described above, reflections other than

those associated with the usual orthorhombic unit cell of polyethylene

were observed in the three stretched samples. The interplanar

spacing corresponding to these reflections agreed closely with spacings of

T Q.456nm, Q.38Qnm, and 0,355nm which are characteristic of a metastable

polymorphic form of polyethylene, which form has been variously described

as being triclinic [26] or monoclinic [27]. The indices of the planes

having these spacings in the triclinic unit cell proposed by Turner-Jones

are (010), (100), and (110) respectively. The corresponding reflections

will be identified in what follows as 010

t

, 100^, and 110 .

+ These spacings differ from the calculated (110) and (200) interplanar

spacings of the orthorhombic unit cell [28] which are 0.41nm and 0.37nm

respectively. Among the cases in which these three extra reflections

have been previously observed are cold worked samples of polyethylene

[29], and samples which were uniaxially extended to a 9.0 strain then

annealed at 100°C and subsequently compressed at right angles to the

initial stretching direction [27].
18

The most Intense among the extra reflections in the diffraction

patterns were the 010 . They can be seen very clearly in the patterns

of the 2,0(_Q.92) and 3.0(1.15) samples shown in Fig. 9c and Fig. 9d

respectively where they appear as arcs whose azimuthal orientations

correspond closely to those of the split 110 orthorhombic arcs. In

comparison, the 010 reflections are azimuthally much broader in the

case of the 1.0(Q,-34)_ sample (Fig. 9b) as is the case for the 110

orthorhombic reflections. It should be noted that in the diffraction

patterns of all three stretched samples the 010 arcs lie over a

diffuse halo associated with the amorphous regions in the specimens.

The 100 and 110 reflections are considerably weaker than the 010

t t t

and are barely if at all visible in Figs. 9b,c,d. They could however

be very faintly seen in the original photographic film the 100 being

situated radially between the intense 11Q and 200 orthorhombic reflections but

very close to the 200, and the (110)^ occurring at a larger radius than the

orthorhombic 200 arcs. Both the 100^_ and 110 appeared preferentially

arced about the equatorial direction. Whether they were split about that

direction could not be determined. Diffraction patterns recorded at larger

specimen to film distances may shed some light on this matter.

Small angle x-ray diffraction patterns were recorded from the same

samples as those used to obtain the wide angle patterns described above.

No diffraction maximum was observed in the unstretched slow cooled

sample. This may be attributed to either one of two factors, bearing in

mind that the lamellar thickness in melt crystallized polyethylene

is in the range 15nm to 90nm, being larger the higher the crystallization

temperature. First the equipment used to obtain the diffraction patterns

was a small angle camera using pinhole collimation. The upDer limit of
19

thickness measurement with this camera using nickel filtered Cu K

a radiation from a rotating anode source is ~30nm. It is conceivable that

the lamellae in samples slow cooled from 473K to room temperature are

thicker than 30nm. Alternatively it may be argued that because the samples

are continuously cooled there is a broad distribution in the thickness of

the lamellae with the result that no diffraction maximum will be observed

even if the lamellar thicknesses are <30nm due to the absence of regularly

periodic fluctuations in electron density distribution.

Small angle x-ray diffraction patterns obtained from the 1.0(0.34),

2.0(0.92) and 3.0(1.15) samples are shown in Fig. 10. The main feature

exhibited by these patterns is that as the residual strain increases

the patterns become four pointed in character (this feature

is barely evident in the 1.0(0.34) sample). In addition, the 2.0(0.92) and

3.0(1.15) samples exhibit a central horizontal streak indicating the presence

of long and comparatively narrow voids whose long dimension is oriented

parallel to the stretching direction. We shall not attempt to discuss these

patterns further in this report. They will be examined in the light of the

numerous previous small angle x-ray studies carried out on lower molecular

weight polyethylenes [see for example references 30-34] subjected to

relatively small strains. At present we will simply note that the patterns

in Fig. 10b and Fig. 10c indicate that the interfaces between the crystalline

and disordered regions in the 2.0(0.92) and 3.0(1.15) samples are appreciably

inclined relative to the stretching direction.

(b) Uniaxial Creep.

Uniaxial creep data have been obtained from a set of specimens which

were prepared by slow cooling in the press from 200°C. The creep strain

versus time is shown in Figure 11 for applied stresses (engineering stress)
20

in the range from 9 to 27.5MPa. The creep strain was determined with

the aid of a cathetometer by monitoring the separation, as a function of

time, of marks placed on the gage section of the specimen. The arrow

indicates the point at which the specimen stressed at 27.5MPa fractured,

and to date represents the only specimen in this set which has failed.

At the highest loads the creep strain approaches a maximum and nearly

constant value of about 5, whereas at the intermediate and small values of

applied load there does appear to be a decrease in the maximum value of

strain with decreasing load.

Quenching of the polymer has a significant effect on the creep behavior,

as can be seen in Figure 12 which shows a comparison of the creep response

for both the quenched and slowly cooled materials. Data are shown for

specimens which were subjected to two relatively large applied loads. The

principal difference in behavior is that the quenched material creeps

much more rapidly initially and then approaches its limiting value of strain

much sooner than does the slowly cooled material. Interestingly, in both

cases the limiting strain is about the same (~5) . A difference in

crystallinity of about 10% has brought about a quite significant change in

the short time creep behavior, while at the same time there is as yet no

indication that the failure behavior will be different in the two cases.

(c) Stress Relaxation

Another method used to characterize the time dependent mechanical

behavior of materials is stress relaxation. In this method the specimen is

subjected to a constant deformation and the stress is then monitored as a

function of time. Such data are presented in Figure 13 for the UHMWPE

polymer which was slowly cooled from 200°C. In the present case the

specimen was a long strip with straight sides and all the data shown in

Figure 13 were obtained from the same specimen. The procedure was as follows.
21

The specimen was initially subjected to the smallest value of strain

shown and the stress monitored for a period of three hours. The specimen

was then released and allowed to relax for a period of about 24 hours.

The new undeformed length was measured and the next larger step then applied.

This procedure was repeated until a strain of about 1.0 was achieved. In

the calculation of the true stress, shown as the ordinate in Figure 13, two

corrections have been made with respect to the original cross sectional area,

one due to the permanent set in the specimen as a result of the prior strain

histories, and the other a result of the current step being applied. By

permanent set we mean here that fraction of the earlier deformation remaining

in the specimen just prior to the application of a new step. For a specimen

deformed to a strain of 1.0 the permanent set may be as much as 0.20 or more.

If the specimen were allowed to relax for a longer period of time between

the application of steps, the recovery would be somewhat greater, and hence

the permanent set less. However as a matter of practicality we have

chosen as the recovery time a period of about one day.

In Figure 13 it can be seen that up to strains of about 0.35 the

rate of relaxation is nearly independent of the level of strain, i.e. the

curves are very nearly parallel to one another. This behavior suggests

that, to a fair approximation, at small strains the stress, cr, can be

represented as the product of a function of time, t, and a function of

strain, £, [i.e. a(e , t)=F (e) *H( t)] .

Another representation of the single step stress relaxation data is the

one shown in Figure 14 where the data are shown as 120, 960 and 10,000

second isochrones of true stress vs. strain. Hereafter such a representation

will be referred to as an (isochronal) stress-strain curve. The points

shown as open symbols refer to the data taken from Figure 13. Note that the

stress-strain behavior is nonlinear over the entire range of strains examined.
22

A linear region would be observed in the stress-strain curve only by

operating at strains below 0.001. By linear behavior we mean that the

modulus is independent of the level of strain and on a log-log plot an

isochronal stress-strain curve will be a straight line having a slope of

unity. A feature which sets the UHMWPE apart from conventional low

molecular weight polyethylenes is that we were able to apply a step in strain

as large as 1.0. For linear polyethylenes having molecular weights below

200,000 we have observed that the stress-strain curve flattens at strains above

about 0.10. The point at which the curve becomes flat corresponds to a

yield point or point at which the material becomes unstable and the

phenomenon of necking may occur. For these materials, if one attempts to

apply a rapid step in strain to a value larger than about 0.10 the specimen

fails. The observation that the UHMWPE does not neck is reflected in

Figure 14 by the fact that the stress-strain curve does not flatten but

continues to rise at all levels of strain. [Similar data for a specimen

o o which was slowly cooled from 165 C rather than 200 C are also shown in

Figure 14. It is evident that there is no significant difference in the

behavior of the two .]

A comparison between the slowly cooled and quenched materials is

shown in Figure 15. A significant difference can be seen in stress

relaxation behavior between the slow cooled and quenched materials. The

isochrones for the quenched material fall well below those for the slowly

cooled polymer, and for a given strain the stress (and correspondingly

the modulus) is smaller by as much as a factor of two. The isochrones for

the quenched material can probably be considered to represent a lower bound

for the relaxation behavior of compression molded samples made from the UH-A

polymer

.
23

(d) Failure Studies

The failure behavior of UHMWPE is currently being studied under

both static (constant loading) and sinusoidal loading conditions. There

is a difficulty in conducting time to fail experiments on the UH-A material,

because for comparable loading conditions it exhibits a much longer

lifetime than do lower molecular weight polyethylenes . Consequently very

long testing times are required to collect the desired data, and so far

all the data collected have been obtained on specimens subjected to large

loads which results in relatively short times to fail (<100 days)

As an aid in analyzing the time to fail data let us first consider an

additivity of damage criterion of the type first proposed by Bailey [35].

This criterion is based on the assumption that material failure is a result

of damage accumulation due to stress. That is to say, damage which a

material experiences as a result of a particular stress history is not

recovered, but accumulates in time to some point at which failure occurs.

The simple example shown schematically in Figure 16 will serve to demonstrate

the idea. Suppose that a material up to time t has a history free of stress.

At time t , it is subjected to a stress a for a period of time At . At time

o o o

t^, it experiences a new stress for time At^, and so on until at time t^

a stress, a^, is imposed for a period of time At^. The cumulative time of

damage or fractional time to failure, is then

where the r (a ) are the times to fail for a constant stress a . Note

B n n that in the time intervals where the stress is zero there is no contribution

Atn

f(t +At ) n n ( 1 )

to the damage. This means that in the interval from t=0 to t=t the

total damage f(t,) is simply ^ t o/T (a n ) . However, it should be pointed out

1 D U
24

that the same damage could occur if a smaller stress were applied over

the entire time interval from t=0 to t=t^. This leads to the concept of

equivalency of stress histories for which the same total damage occurs.

According to the damage criterion, the condition corresponding to failure

is f(t +At ) = 1. n n

In integral form the criterion can be expressed as.

d K = 1

( 2 )

where the integration is carried out over the time variable, £, and

T„[<J(5)] now represents the time to fail function determined from experiments

D carried out at constant stress. Each increment in time, d£ , is weighted

inversely as the lifetime, T (a) , which the material would have had under

D

a constant stress CT. The upper limit in the integral is then t^, the time

D to fail for the actual stress history being applied. An advantage of a

damage rule of this type is that, in principle, if one can determine the

functional form of T (a) from constant load experiments, then the lifetime

D under other loading histories can be predicted from equation (2) . For

example, in earlier work on lower molecular weight polyethylenes it was

found [12,36] that over several decades in time the time to fail behavior

could be described by the relationships

X

B

(0) = A 6

(3)

where A and B are constants that depend upon material properties, and a

is the applied stress (engineering stress) . For the case where the loading

history is sinusoidal

0(0 = P + Q sin (go O (A)

P represents the mean stress and Q the amplitude of the oscillatory component
25

of stress, and the time to fail is predicted by equation (2) to be

Vp)

I 0 (BQ)

(5)

where t (P) is the time to fail for a constant stress, P, and Io(BQ)

the zero order modified Bessel function.

In the case of the lower molecular weight polyethylenes , it was

found that for simple loading histories [ 36 J such as those described by

equation (1) the damage rule was valid within the experimental error.

However, for sinusoidal loading conditions [l2j the lifetimes predicted

from equation (5) were shorter than the experimentally determined values.

Thus for the polyethylenes described in reference [12] , the lifetimes

predicted using the cumulative damage rule can be considered to represent

a lower bound to the actual lifetimes observed experimentally.

Our experience with polymers in general suggests that an important

aspect of time dependent failure is the frequency dependence of the fatigue

lifetime under sinusoidal loading conditions. One feature of the damage rule

described above is that for sinusoidal loading conditions the predicted

lifetime is independent of the test frequency, as can be seen from equation

(.5)

Coleman [37,38] has developed the mathematics of additivity of damage

rules for polymer fibers rather extensively. One of the damage rules

which he has developed corresponds to the "cycle dependent" fatigue behavior

which is often assumed in engineering practice. In cycle dependent

fatigue, the damage accumulates due to cycling of the material rather than

due to duration under load. In such a case failure would occur after a

given number of cycles, N-, and the lifetime, t , would vary inversely with

r B test frequency, f:

N,

t„ =

f ( 6 )
26

In the testing of a lower molecular weight polyethylene [12

]

we have found that neither equation (5) nor equation (6) predicts the

lifetime of the material. The prediction from equation (5) is

conservative (i.e. observed lifetime is longer than that predicted) while

equation (6) greatly overestimates lifetime. This result is of significance

in laboratory testing and indicates the importance of determining the

frequency dependence of fatigue lifetime. The following illustrates why this

is so

.

A common assumption in engineering practice is that fatigue failure

is a cycle dependent phenomenon, not a time dependent one. As a result,

common laboratory practice calls for "accelerating" fatigue tests by

increasing the test frequency, counting the number of cycles to failure and

using this number as a measure of the fatigue resistance of the material.

Thus, for example, if an "accelerated" test at 35Hz were to determine that

g the lifetime of a UHMWPE component is 10 cycles, then the expected

g in vivo lifetime (assuming a frequency of 1 Hz) would also be 10 cycles, or

3 years. If, however, the time to fail of the material were independent

of test frequency, as predicted by equation 5, then the lifetime at 35 Hz of

8 6

10 cycles would correspond to a lifetime of 2.9x10 cycles at 1 Hz or

a time to fail of only 33 days.

The failure criteria described above then provide a useful framework

within which to consider the time dependent failure behavior of UHMWPE. We

have chosen to study two important aspects of material failure, namely:

(1) The time to fail behavior under both static and cyclic loading conditions

(2) The dependence of the time to fail on test frequency under cyclic loading

Our results to date for specimens failed at constant load (creep data)

and in zero-tension sinusoidal loading at 0.002 Hz (fatigue data) are shown in
27

Figure 17. The solid and dashed lines represent linear regression curves

of the form given by equation (.3) . It should be pointed out that these

curves are based on relatively few data points which exhibit considerable

scatter (unlike the lower molecular weight polyethylenes [12,36] which show

very little scatter in the time to fail data) . The data indicate that the

fatigue lifetimes under sinusoidal loading are approximately 6-7 times

greater than the static lifetimes. This result is in close agreement with

the fatigue lifetime predicted from equation (5) using the regression

equation fitted to the creep data.

We are also in the process of doing fatigue experiments at 0.0002

and 0.01 Hz to investigate the effect, if any, of test frequency on the

material lifetime. Preliminary results do not show any clear trend toward

a dependence of lifetime on frequency.

A comment is in order concerning the choice of frequencies used in

this study to characterize the fatigue behavior of UHMWPE. We have chosen

test frequencies in the 0.0002 to 0.01 Hz range, whereas in the clinical

environment frequencies in the 1 Hz range are more typical. In order to

conduct fatigue tests within an acceptable length of time it is necessary

to carry out the testing at relatively high levels of stress. At such

stress levels the amount of hysteretic heating can be substantial, thus

greatly biasing the test results. We have therefore chosen somewhat lower

frequencies in order to assure that the tests are conducted under nominally

isothermal conditions.

(e) Equal-Biaxial Deformations Under Inflation and Environmental

Stress Cracking

Work in these areas has just begun. The disc inflation apparatus used

to achieve equal-biaxial deformation histories is shown schematically in
28

Figure 17. A sheet of material is clamped as shown and a constant

pressure is then applied to one side of the specimen. The deformation

is equal-biaxial only in the immediate vicinity of the pole. The true

stress at the pole can be calculated for a given stretch using the

following relationship [39]:

- f- -T

(7)

where a is the true stress, p is the applied pressure, r is the radius

of curvature measured at the pole, h is the initial thickness of the

sheet, and X is the stretch determined at the pole.

The environmental stress-crack resistance of the UHMWPE is being

investigated by two methods. In one method a small amount of the

stress-cracking agent (in the present case a ten percent solution of

nonylphenoxy-poly (ethyleneoxy) ethanol in distilled water) is introduced

into the inflation apparatus prior to pressurization. As soon as the

apparatus is pressurized it is inverted so that the liquid is in contact

with the specimen only in the region near the pole.

In the second method a test being developed for ASTM Committee D-20.10

on plastics is used. This test is based on a bent strip geometry. These

tests are performed at different temperatures in the 333 to 393K range.

Unlike low molecular weight polyethyl enes the lifetime of the UHMWPE

in stress cracking agent is extremely long. For example, one sheet of the

UHMWPE being tested in the inflation apparatus at a pressure of 0.48 MPa

(70 PSI) has not failed in nearly seven months whereas a similar sheet of

a low molecular weight polyethylene failed almost immediately. Thus far,

none of the specimens being tested either in air or in stress-cracking

agent, by either method, has failed. It appears that the pressurizing
29

system for the inflation apparatus will have to be redesigned in order

to operate at higher pressures, or alternatively the tests will have

to be carried out at higher temperatures.

(f) Miscellaneous

In the Second Quarterly Report [40] we compared the creep behavior

of the UHMWPE polymer used in the present study to that of specimens

molded from another batch of UHMWPE polymer (UH-B) obtained from the

same manufacturer several years earlier. It was noted that the creep

compliance and time to failure behavior was quite different for the two

materials even though the intrinsic viscosities of the raw polymers

and densities of the molded sheets were essentially the same. The specimens

made from the two batches of polymer were, however, molded under different

conditions. The molding temperature was 473K (200°C) for the UH-A

polymer and it was 428K (155°C) for the UH-B. Since then we have prepared

0 molded sheets from the UH-B polymer using the same molding (47 3K) and

slow cooling procedure outlined in Section 2. The creep and stress

relaxation data were within 5% of the values for the UH-A material

presently being used. The greater creep compliance and shorter time to

fail exhibited by the UH-B specimens molded at 428K can be reasonably

attributed to poor interparticle adhesion at the lower molding temperature.
30

4 . Summary

Work carried out in FY 1980 consisted of two phases, the first of which

was concerned with the preparation of compression molded sheets of UHMWPE

suitable for studying the mechanical properties of the polymer. Features of

the fine structure of the raw polymer as well as of the compression molded

sheets were also examined. In the second phase investigations were started

on various aspects of the mechanical properties of UHMWPE, namely stress-strain

behavior at constant rate of clamp separation, creep and single step stress

relaxation in uniaxial extension, and failure under both static and sinusoidal

loading conditions. A study of fine structure changes resulting from uniaxial

deformation under constant rate of clamp separation has also been started. In

addition, preliminary work was done on the equal biaxial deformation (under

inflation) and the stress crack resistance of UHMWPE.

Three aspects of the raw polymer have been examined, its morphological

characteristics, its melting behavior, and its crystallinity. The polymer, as

received, is in its as-polymerized state and consists of fine particles. Because

of its very high melt viscosity both molded and machined prothesis components

made from it retain a memory of the particulate nature of the raw polymer. This

was also true for specimens molded in our laboratory at temperatures as high as

90K above the melting point. This observation emphasizes the importance of

characterizing the structure of UHMWPE products, not only to determine their

crystallinity, but also to monitor those aspects of their morphology which are

directly related to, and a consequence of, the particulate nature of the raw

polymer. Since the nature of the shapes, sizes, and fine textures of nascent

polyethylene particles vary depending upon the polymerization conditions,

knowledge of the morphology of the raw polymer is a necessary aspect of the

present study.
31

The raw polymer particles were examined in a scanning electron

microscope. The particles, which are irregularly shaped, and which differ

in size in the range 80ym - 300ym, exhibit at low magnifications an external

appearance resembling that of compact clumps of grapes. Examination at high

magnifications revealed regions which were highly porous. The voids in

these regions were bridged by microfibrils, 50nm - lOOnm in diameter, which probably

result from the deformation of the polymer during polymerization as a consequence

of the progressive expansion of the particles caused by the continued generation

of fresh polymer at the surface of underlying encapsuled catalyst.

Differential scanning calorimetry (DSC) was used to examine the melting

behavior and crystallinity of the raw polymer. The melting point was found to

be between 413 and 414K (140-141°C) . The percent crystallinity, as determined from

the area under the melting curve, was 78%, which represents a value significantly

higher than can be expected for the same material compression molded at

temperatures above the melting point.

Considerable time was devoted to the design and construction of vacuum

molding equipment and to the establishment of procedures for molding laboratory

specimens of the UHMWPE. Protocols were established for preparing compression

molded sheets having weight % crystallinities of close to either 50% or 60%

3 3 (densities close to 0.923 g/cm or 0.936 g . cm ) by respectively rapidly cooling

or slowly cooling polymer melted at 473K.

Examination of cross-sections of the sheets under a light microscope

using phase contrast optics revealed the retention of a memory of the original

particulate nature of the raw polymer in the sheets. Examination of crosssections of the slow cooled sheets under a polarizing light microscope showed

that there was a wide variation in the sizes of the birefringent structures in

them. The larger among these structures appeared to be axialitic or spherulitic
- 32

in character. Retention of a memory of the raw polymer particles and

the wide variation in the sizes of the birefringent structures were also observed

in sheets compression molded at 503K.

3 The deformation and failure of slow cooled (p=0.936 g/cm ) samples of UHMWPE

subjected to uniaxial extension have been examined in constant rate of clamp

separation, creep, stress relaxation, and sinusoidal loading experiments. Only

the creep and stress relaxation characteristics of rapidly cooled samples

3

(p=0.923 g/cm ) have been examined so far.

Constant rate of clamp separation experiments were carried out on slow

cooled samples at 296K and 310K at clamp separation rates in the range 0.02 in. /min

to 2.0 in. /min. Both the modulus and the strength at break decreased with

decreasing rate of deformation and with increasing temperature as is the case for

for lower molecular weight linear polyethylenes . In contrast with the behavior of

the lower molecular weight polyethylenes the slow cooled sheets of UHMWPE did

not exhibit the phenomenon of necking. The specimens underwent uniform drawing

along their length up to the point of failure. Constant rate of clamp separation

experiments on slow cooled samples were also carried out in which specimens were

stretched to various predetermined strains below the breaking strain and were

then released. These specimens, which exhibited substantial recovery from the

imposed strain, are being used as models for the study of the morphological

changes associated with the deformation of UHMWPE. A preliminary examination of

such specimens has been carried out using wide angle and small angle x-ray

diffraction.

Stress relaxation and creep experiments have been carried out at 296K on

both slow cooled and quenched samples. Although the crystallinities of these

samples (~60%, ~51% respectively) differed by only ~10%, there were significant

differences between their stress relaxation behaviors and between their creep
33

behaviors. At strains ranging from 0.001 to 1.0 and over times from

60s to 10,000s the relaxation modulus of the slow cooled samples was

approximately twice as large as that of the rapidly quenched ones. In the

creep experiments the slow cooled samples exhibited a much greater resistance

to creep than the quenched ones during the early stages of deformation. The

limiting creep elongation (strain at failure), however, was approximately the

same (~5) for both types of samples.
- 34

5 . Work Plan for FY 1981

The work described in this report represents the initial stages of a

project concerned with the characterization of the morphology and mechanical

behavior of UHMWPE. In particular we are interested in, (1) determining the

influence of processing conditions (thermal history) on the morphology and

mechanical behavior of UHMWPE and (2) studying the changes in morphology which

result from mechanical deformation. During FY 1980 we concentrated our efforts

on developing molding apparatus and procedures for producing sheets of UHMWPE

which are suitable for preparation of specimens for uniaxial (dumbbell, strip)

and equal biaxial (circular sheets) test specimens. In addition considerable

effort was spent in characterizing the morphology and mechanical behavior of

UHMWPE, which was slowly cooled from the melt. The experiments, to characterize

the mechanical behvaior of UHMWPE have been largely carried out at room

temperature and in uniaxial extension.

During FY 1981 the work presented in this report will be extended to obtain

additional data at room temperature. In addition, work will be initiated to

study the mechanical behavior of UHMWPE at higher temperatures and in other modes

of deformation. The following outlines the work which is proposed for FY 1981.

1. Experimental work begun in the areas of uniaxial creep, stress

relaxation and static and dynamic fatigue will be continued and

extended to higher temperature [310K (37°C)]. The creep behavior

will be characterized in 0.15N saline solution as well as laboratory

air. In addition, the effect of morphology as changed by varying

specimen fabrication procedures (quenching, annealing, etc.) on

mechanical behavior will continue to be examined.
35

2. Work already initiated to study the equal biaxial creep behavior

and stress crack resistance of UHMWPE will be continued. Because

of the long lifetime of UHMWPE in stress-cracking fluid the stress

crack tests will be carried out at elevated temperatures.

3. Experiments will be initiated to study the compressive creep and

compressive stress relaxation behavior of UHMWPE. Initial work will

involve the development of apparatus and procedures for compression

molding cylinders of UHMWPE.

4. The use of specimens deformed under uniaxial testing conditions as a

model for studying morphological changes associated with deformation

will continue. This work will be extended to examine features associated

with creep and cyclic loading of the polymer in addition to the constant

rate of elongation experiments examined in FY 1980.
36

Appendix A

The percent crystallinity of polyethylene can be determined either

from density measurements or from Differential Scanning Calorimetry (DSC)

.

In the present work the density was determined with a water-ethanol density

gradient column at 296. 3K. The weight percent crystallinity is calculated

with the following equation,

P„ Cp-pJ

X (percent crystallinity) = — Xioo ( 1 ) p (p

c -pa

)

where p is the measured density of the polymer, p a is the density of the

amorphous polyethylene, and p c is the density of a polyethylene crystal. In

the present instance values of p a =0.855 g.cmJ [4l] and p c =1.0 g.cm

J [28]

were used.

The percent crystallinity, as determined from DSC, is defined by the

following relationship.

X

= AH

AH

o

( 2 )

where AH is the heat of fusion determined from the area under the melting

peak and AHQ is the heat of fusion for a purely crystalline specimen. In the

present case a value of 71.43 cal/g was used for AH Q [40 ]

-

AH is determined with the following equation:

T> ATT" AH = 0.2614 cal/g.

Wb

where R = heating rate (m cal/°c)

A = area under the melting peak (mm”)

K = a calibration constant

W = mass of the specimen (mg)

S = recorder speed (mm/min)
37

References

1. J. M. Dowling, J, R, Atkinson, D. Dowson, and J. Charnley,

"The Characteristics of Acetabular Cups Worn in the Human Body",

J. Bone Joint Surg., 60-B(3), 375 (1978).

2. J. R. Atkinson, K. J. Brown and D. Dowson, "The Wear of High

Molecular Weight Polyethylene. Part I", J. Lubrication Technology,

100(2), 208 (1978).

3. E. A. Salvati, T. M. Wright, A. H. Burstein, and B. Jacobs,

"Fracture of Polyethylene Acetabular Cups", J. Bone Joint Surg.,

61-A (8) , 1239 (1979)

4. B. Weightman, D. P. Isherwood and S. A. V. Swanson, "The Fracture

of Ultrahigh Molecular Weight Polyethylene in the Human Body",

J. Biomed. Mater. Res., 1^3, 669 (1979).

5. P. D. Ritchie, ed. Physics of Plastics , Iliffe Books, Ltd.,

London, 1965.

6. F. W. Bilmeyer, Jr., Textbook of Polymer Science , Interscience

Publishers, New York, 1966.

7. F. Khoury, H. L. Wagner, L. J. Zapas, and J. P. Colson, "Aspects

of the Characterization of Ultra High Molecular Weight Polyethylene",

paper presented at the Ninth International Biomaterials Symposium,

New Orleans, 1977.

8. D. F. Gibbons, J. M. Anderson, R. L. Martin and T. Nelson, "Wear and

Degradation of Retrieved Ultra High Molecular Weight Polyethylene and

Other Polymeric Implants", Corrosion and Degradation of Implant

Materials, ASTM STP684, B. C. Syrett and A. Acharya Eds., ASTM 1979,

pp 20-40.

9. S. K. Bhateja, J. K. Rieke and E. H. Andrews, "Impact Fatigue Response

of Ultra-High Molecular Weight Linear Polyethylene", J. Matls. Sci . ,14

,

2103 (1979).

10. R. M. Rose, A. Crugnola, M. Ries, W. R. Cimino, I. Paul, and E. L.

Radin, "On the Origins of High In Vivo Wear Rates in Polyethylene

Components of Total Joint Prostheses", Clinical Orthopedics and

Related Research, 145 , 277 (1979)

11.

J. M. Crissman and L. J. Zapas, "Creep Failure and Fracture of

Polyethylene in Uniaxial Extension", Polym. Eng. and Sci., 1J?, 99 (1979).

12. G. B. McKenna and R. W. Penn, "Time Dependent Failure of Poly (methyl

methacrylate) and Polyethylene", Polymer, _21_, 213 (1980).

13. J. K. Lancaster, "Basic Mechanisms of Friction and Wear of Polymers",

Plastics and Polymers, 297 (Dec. 1973).
- 38

14. P. S. Walker, "The Measurement and Effects of Friction and Wear

in Artificial Hip Joints", J. Biomed. Mater. Res. Symposium, 4_, 327 (1973) .

15. B. 0. Weightman, I. L. Paul, R. M. Rose, S. R. Simon and E. L. Radin,

"A Comparative Study of Total Hip Replacement Prostheses",

J. Biomechanics, 6_, 299 (1973) .

16. H. L. Wagner and J. G. Dillon, "A Specification Method for Ultra

High Molecular Weight Polyethylene for Implant Use", J. Biomed.

Mater. Res., 13, 821 (1979).

17. R. H. Marchessault , B. Fisa, and H. D. Chanzy, "Nascent morphology of

polyolefins", Crit. Rev. Macromol. Sci. 315 (1972).

18. B. Wunderlich, "Macromolecular Physics, Vol. 2", Academic Press, N.Y.

1976, see pp 282-293.

19. For a discussion of the equilibrium melting point of polyethylene see

for example, J. D. Hoffman, G. T. Davis, and J. I. Lauritzen Jr.,

"The Rate of Crystallization of Linear Polymers with Chain Folding",

Chapter 7 (see page 567), in Treatise on Solid State Chemistry Vol. 3,

N. B. Hannag Ed., Plenum Press N.Y. (1976).

20. J. D. Hoffman, G. S. Ross, L. Frolen and J. I. Lauritzen, Jr.,

"On the growth Rate of Spherulites and Axialites from the Melt in

Polyethylene Fractions: Regime I and Regime II Crystallization",

J. Res. Nat. Bur. Stds. (U»S.), 79A, 671 (1975).

21. J. Maxfield and L. Mandelkern, "Crystallinity, Supermolecular Structure

and Thermodynamic Properties of Linear Polyethylene Fractions",

Macromolecules, 1£, 1141 (1977)

.

22. F. Khoury and E. Passaglia, "The Morphology of Crystalline Synthetic

Polymers", in Treatise on Solid State Chemistry, N. B. Hannay ed

.

(Plenum Press, New York), Vol. 3_, 335 (1976). (See pp 466-472).

23. H. J, Nusbaum and R. M. Rose, "The Effects of Radiation Sterilization

on the Properties of Ultra High Molecular Weight Polyethylene",

J. Biomed. Mater. Res. L3, 557 (1979).

24. I. L. Hay and A. Keller, "Polymer Deformation in Terms of Spherulites",

Kolloid Z. u Z. Polymere, 204 , 43 (1965).

25. A. Keller, "Unusual Orientation Phenomena in Polyethylene Interpreted in

Terms of the Morphology", J. Polymer Sci. 1_5, 31 (1955).

26. A. Turner Jones, "The Triclinic Crystal Form of Polymethylenes and

Polyethylene", J. Polymer Sci., 62_, 553 (1962).

27. K. Tanaka, T. Seto, and T. Hara, "Crystal Structure of a New Form of High

Density Polyethylene", J. Phys . Soc. Japan, 17_, 873 (1962).
- 39

28. C. W. Bunn, "The Crystal Structure of long-chain paraffin hydrocarbons.

The shape of the Cl^ group", Trans. Faraday Soc . J9, 482 (1939).

29. P. W. Teare and D. R. Holmes, "Extra Reflections in the X-ray Diffraction

Pattern of Polyethylenes and Polymethylenes", J. Polymer Sci. 24^ 496

(1957).

30. N. Kasai and M. Kakudo, "Fine Texture in Necking Portions of Cold-Drawn

Polyethylene", J. Polymer Sci. A 2, 1955 (1964).

31. R. Corneliussen and A. Peterlin, "The influence of Temperature on the

Plastic Deformation of Polyethylene", Makromol. Chem. 105 , 193 (1967).

32. V. A. Marikhin and L. P. Myasinkova, "The Role of Interf ibrillar Tie

Molecules in Drawing of Polymers", J. Polymer Sci: Polymer Symp. 58 , 97 (1977).

33. P. B. Bowden and R. J. Young, "Deformation Mechanisms in Crystalline

Polymers", J. Matls. Res. 9^, 2034 (1974).

34. L. E. Alexander, "X-ray Diffraction Methods in Polymer Science", WileyInterscience, New York (1969)

35. J. Bailey, "An attempt to Correlate Some Tensile Strength Measurements on

Glass: III", Glass Industry, 2(3, 94 (1939).

36. „J, M, Crissman, C, M, Guttman, and L, J, Zapas, "Performance of Plastic

Packaging for Hazardous Materials Transportation. Part I: Mechanical

Properties", Report No. DoT/MTB/0HM0-76/4 , October, 1976.

37. Coleman, B.D. "Application of the Theory of Absolute Reaction Rates to the

Creep Failure of Polymeric Filaments", J. Polymer Sci., 2(3, 447(1956).

38. Coleman, B. D., "Statistics and Time Dependence of Mechanical Breakdown in

Fibers", J. Appl . Phys. 29, 968 (.1958).

39. J. E. Adkins and R. S. Rivlin, "Large Elastic Deformations of Isotropic

Materials: IX. The Deformation of Thin Shells", Phil. Trans. Royal Soc.

London, A244 , 505 (1952).

40. G. B. McKenna, F. Khoury and J. M. Crissman, Second Quarterly Report

Jan. -April, 1980, Task 80-01, NBS-BMD Interagency Agreement.

41. G. T. Davis and R. K. Eby, "Glass Transition of Polyethylene: Volume

Relaxation", J. Appl. Phys, 44_, 4274 (1973).

42. M. G. Broadhurst, "Thermodynamic Properties of Polyethylene Predicted

from Paraffin Data, J. Research Natl. Bur. Std. 67A, 233 (1963).
Fig. 1 Scanning electron micrographs illustrating the

particles of raw ultra high molecular weight

polyethylene UH-A at two different magnifications.
41 -

**V f.

Fig, 2 Scanning electron micrograph showing particles P,Q,R,S

of raw UH-A polymer. The region below that marked X

of particle P in (a) is shown a two higher magnifications

in (b) and (c) respectively. Note the fibrils in (b) and

(c) in which the arrows identify the same region in both

pictures
- 42 -

c Fig. 3 (a) Scanning electron micrograph of a particle of raw UH-A

polymer. The region to the right of the area Y is shown at

two higher magnifications in (b) and (c) respectively. Note

the fibrils in (b) and (c) in which the arrows identify the same region in both pictures.
43

do

•H

CO

CO •H§

P-l u

u

<u 00 <u d 4= •H CO

CO3 XI w <U cI cc

d>

CO

4Jd

•HO

CU

COu

CU

>ud

<UCUoo

COo

1-1 CJ

•rl6 1

4-1 4=00

•uu

ccj cu

V4

<Uo

cu

Std

r-1 -H ,Q 00 dd

*Hd

cudu1 CO -d

CO ^I CO

*0 vO

co <r o

V-l /-N U d

00d

CU S

aj o

1-1 rC

OO to o

1-1 o coo B *r1

•Udo

•H

14U •Hd

eodeSO4 d

4J o

CU *H O CO

CO

4-1 CU

COd

1-1 4Jdoo

cu

COd4=Cu

uo

cu 4C

4Jd

<u

<uS

4-1 0) 40

CO

<u •H

lud

x)ddo

-Q

CU 40

Cu 4-1 gO M—

O O

(Ueo

COo

4-4

00

Tl

pH
- 44 -

Fig. 5 Light micrographs of an area in a cross-section of a sheet of UH-A polymer molded at 463K,(a) Phase contrast

optics, (b) crossed polarizer and analyzer. Note the

variation in the sizes of the birefringent structures

in (b) . The fine variations in contrast in (a) are associated with the birefringent structures in the sample.
- 45 -

Fig, 6 Light micrographs of an area in a cross-section of a sheet of

UH<-A polymer molded at 473K. (a) Phase contrast optics, (b)

crossed polarizer and analyzer. Note the variation in the sizes

of the birefringent structures in (b) . The fine variations in

contrast in (a) are associated with the birefringent structures

in the sample.
- 46

Fig. 7 Light micrographs of an area in a cross-section of a sheet

of UH-A polymer molded at 503K. (a) Phase contrast optics,

(b) crossed polarizer and analyzer. Note the variation in

the sizes of the birefringent structures in (b) . The fine

variations in contrast in (a) are associated with the

birefringent structures in the sample.
- 47

Fig, 8 Engineering stress versus strain for specimens of UHMWPE

elongated at different rates of clamp separation. UH-A

slowly cooled from 473K (200°C)

,
~ 48 -

C d

Fig. 9 (a) Wide angle x-ray diffraction pattern from an unstretched sheet of

UH-A which was molded at 473K and then slow cooled. The other diffraction

patterns were exhibited by specimens (from the same sheet) which had been

stretched at 310K (37°C) to strains of (b) 1.0, (c) 2.0 and (d) 3.0, and

then released and stored at room temperature. The residual strains in

the specimens after storage at room temperature for twenty-four hours were

0.34, 0.92, and 1.15 respectively. The corresponding residual strains

after eleven weeks were 0.26, 0.72 and 0.92. The diffraction patterns were recorded about two to three weeks after stretching (see text).

Stretching direction vertical. Letter H in (a) denotes ’amorphous' halo.

Reflections 010 t are associated with the triclinic crystal form [26] of

polyethylene

.

200

2 00

110
- 49 -

\

a

b

c

Fig. 10 Small angle x-ray diffraction patterns from the same stretched specimens

as those used to obtain the wide angle patterns in Fig. 9b, 9c and 9d

The specimens were originally stretched at 310K to strains of (a) 1.0,

(b) 2.0, and (c) 3.0. The residual strains in the specimens after they

were released and stored at room temperature for 24 hours were 0.34, 0.92

and 1.15 respectively. The corresponding residual strains after eleven

weeks were 0.26, 0.72 and 0.92. The diffraction patterns were recorded

about two to three weeks after stretching (see text). Stretching direction

vertical. (Specimen to film distance 52cm; magnification of photographs

compared to original film XI. 5).
- 50 -

(/>

111 2

NIVH1S

Uniaxial

creep

for

specimens

of

UHMWPE

(UH-A)

slowly

cooled

from

473K

(200°C)

.

The

engineering

stress

is

indicated

by

the

numbers

shown.
51

CO

LU

2

I—

N I V dlS

Fig.

12

Comparison

of

uniaxial

creep

for

specimens

of

UHMWPE

(UH-A)

slowly

cooled

and

quenched

from

473K

(200°C)

.

Open

symbols

-

slowly

cooled,

filled

symbols

-

quenched;

circles

-

27.5

MPa,

squares

-

21

MPa

engineering

stress.
52

(Ed) SSBdlS 3HH1

Fig.

13

True

stress

versus

time

for

UHMV7PE

(UH-A)

slowly

cooled

from

473K

(200

C)

,

The

data

were

obtained

from

single

step

stress-relaxation

experiments

carried

out

on

one

specimen

and

in

the

order

of

increasing

strain.
53

( e d) SS 3dlS dfldl

Fig?

14

Isochronal

true

stress

versus

strain

for

specimens

of

UHMWPE

(UH-A)

slowly

cooled

from

473K

(200°C)

-

open

symbols,

and

438K

(165°C)

-

filled

symbols.

Crosses

are

for

samples

of

UHMWPE

(UH-B)

slowly

cooled

from

473K

(200°C)
TRUE

STRESS

(Pa)

10

8

10

7

10

10

'

0

A

0

o ° 1

°

O A

n

O

A

0 3

a

A

O

A

0 • A •>0-120 sec. —

A a;a-960

Bp- 10,000

J L_ L

1 .001 .01

STRAIN

0.1 1.0

Fig. 15 Comparison of Isochronal Stress-Strain Data for UHMWPE (UH-A)

specimens slowly cooled (open symbols) and quenched (filled

symbols) from 473K (200°C)

.
55

Fig.

16

Schematic

representation

of

a

possible

loading

history

applied

to

a

specimen

of

UHMWPE.

See

text

for

explanation

of

how

the

load

during

each

time

interval

A

t.

would

contribute

to

the

damage

accumulating

in

the

material.

1
Time

To

Fail,

56

Fig, 17 Time to fail versus peak engineering stress for specimens of

UHMWPE (UH-A) in creep (static fatigue) and zero-tension

sinusoidal fatigue loading histories. The solid and dashed

lines represent the linear regression curves for the data points

shown. Samples were slowly cooled from the molding temperature,

0-438K(i65°C); X-463K (190 6 C) ;Q , V - 473K (200°C)

.
57

U.S. DfcIPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO.

NBSIR 81-2209

2. Performing Organ. Report No. 3. Publication Date

December 1980

4. TITLE AND SUBTITLE

Relationship Between Morphology and Mechanical Properties of Ultra High

Molecular Weight Polyethylene

5. AUTHOR(S) J G.B. McKenna, F. A. Khoury and J.M. Crissman

6. PERFORMING ORGANIZATION (If joint or other than N BS. see instructions)

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

7. Contract/Grant No.

8.

Type of Report & Period-Covered

Annual Report

Dec. 9, 1979 - Sept . 30,15^6

9.

SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS t Street. City. State, ZIP)

Food and Drug Administration

Bureau of Medical Devices

3257 Georgia Avenue

Silver Spring, Md, 20910

10.

SUPPLEMENTARY NOTES

(321 Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11.

ABSTRACT (A 200-word or less factual summary of most significant in formation. If document includes a significant

bibliography or literature survey, mention it here)

Aspects of the morphology of the constituent particles of the raw ultra high molecular

weight polyethylene (UHMWPE) used in this study have been examined using scanning

electron microscopy. In addition differential scanning calorimetry has oeen used to

determine the melting point (~413K) and the crystallinity (78%) of the raw polymer.

Protocols have been established for the preparation of compression molded sheets of

UHMWPE having crystallinities of 50% and 60%. A memory of the particulate nature of

the raw polymer is retained in the sheets. The following aspects of the mechanical

behavior of the molded sheets have been examined! Stress—strain behavior at constant

rate of clamp separation, creep and single step stress relaxation in uniaxial extension

and failure under both static and sinusoidal loading conditions. The relaxation

modulus of the lower density material (quenched) was found to be smaller by a factor

or two than that of the higher density material (slow cooled) . The short time creep

of the quenched UHMWPE was significantly greater than that of the slow cooled polymer;

however, the limiting creep strain prior to failure was nearly the same for both

materials. Low cycle fatigue data obtained on the slow cooled UHMWPE suggests that it

may be possible to characterize the failure behavior of this polymer using a cumulative damage rule. The lifetime of the UHMWPE under zero-tension sinusoidal

loading is 6-7 times longer than the lifetime in static (creep) tests in agreement wit:

the cumulative damage rule. '

12.

KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) Ultra High Molecular Weight Polyethylene, UHMWPE, Morphology, Polymer,

Mechanical Properties, Creep, Stress Relaxation, Fatigue, Failure.

13.

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