AlllDl IIDTSS

r , NBS

PUBLICATIONS

NBSIR 82-2493 (FDA)

Relationship Between Morphology

and Mechanical Properties of

Ultra High Molecular Weight

Polyethylene

Second Annual Report for the Period

October 1, 1980 - September 30, 1981

Task 80-01, NBS-Bureau of Medical Devices Interagency Agreement

December 1981

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS

KATJONAL bukbau

OB STAKDAHUS

LTBBART MAY 2 5 1982

/

NBSIR 82-2493

RELATIONSHIP BETWEEN MORPHOLOGY

J. M. Crissman, Ph.D.

Principal Investigator

F. A. Khoury, Ph.D.

Project Leader

G. B. McKenna, Ph.D.

Project Leader

Second Annual Report for the Period

October 1, 1980 - September 30, 1981

Task 80-01, NBS-EMD Interagency Agreement

NBS-Bureau of Medical Devices Interagency Agreement

December 1981

AND MECHANICAL PROPERTIES OF

ULTRA HIGH MOLECULAR WEIGHT

POLYETHYLENE

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

NATIONAL BUREAU OF STANDARDS. Ernest Ambler. Director
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Abstract

This report describes work done during FY 1981 under task 80-01,

NBS-FDA/BMD (Bureau of Medical Devices) Interagency Agreement. The

report covers the second year of a four year project concerned with

the relationship between morphology and mechanical properties of ultra

high molecular weight polyethylene (UHMWPE) . Various aspects of the

UHMWPE being investigated include (1) the establishment of procedures

for compression molding (under vacuum) sheets and cylinders of the

polymer, (2) the influence of thermal history and molecular weight on

the time dependent mechanical behavior such as creep and stress relaxation

in both uniaxial extension and compression, (3) dynamic fatigue under

conditions of zero-tension sinusoidal loading, (4) environmental stresscrack resistance, and (5) the optical and electron microscopy of sheets

of the UHMWPE which have been deformed to varying degrees in uniaxial

extension.
ii

Table of Contents

Abstract ±

Table of Contents ii

1. Introduction 1

2. Polymer Morphology 2

2.1 Aspects of the Density (p) and Crystallinity

(x) of Compression Molded UHMWPE 2

2.2 Uniaxial stretching of UHMWPE 8

3. Mechanical Behavior Studies 13

3.1 Creep and Stress-Relaxation in Uniaxial

Extension 15

3.2 Creep and Stress-Relaxation in Simple

Compression 28

3.3 Fatigue under Creep and Cyclic Loading 31

4. Environmental Stress-Crack Resistance 39

5 . Summary 46

6 . Work Plan for FY 1982 52

7. References 54

8 . List of Figures 55
I. Introduction

This report describes work done during FY 1981 under task 80-01,

NBS-BMD (Bureau of Medical Devices) Interagency Agreement. The work plan

followed was basically that outlined on pages 34 and 35 of the First Annual

t Report [1]. In addition, we have examined various aspects of components of

a number of commercially manufactured orthopedic prostheses obtained for our

study by the Bureau of Medical Devices. These devices were acetabular cups

and tibial plateaux of various sizes and designs.

The work described in reference [1] was directed primarily toward two

areas of study. One was concerned with the establishment of procedures for

compression molding sheets of ultra high molecular weight polyethylene (UHMWPE)

suitable for mechanical behavior studies. As part of that phase of the work

three aspects of the raw polymer were examined, its morphological characteristics,

its melting behavior, and its crystallinity. The morphology of the molded

sheets was also examined. As a consequence of its very high melt viscosity, an

important feature of the UHMWPE is that both molded and machined prosthesis

components made from it retain a memory of the particulate nature of the raw

polymer.

The second principal area of study involved the characterization of

the mechanical behavior of compression molded sheets prepared under varied

thermal histories. Experiments in creep and stress relaxation in uniaxial

extension were begun, and failure studies under both static and sinusoidal

loading conditions were initiated. In addition, aspects of the deformation

and recovery of molded sheets subjected to uniaxial stretching were examined.

The work described in reference [1] was carried out on two different

lots of UHMWPE obtained from the same manufacturer. The two polymers were

t Numbers in brackets indicate references found at the end of this report.
2

Identified as UH-A and UH-B. Both had the same manufacturer specified

intrinsic viscosity of ~25 dl/g which corresponds to a molecular weight

6

* of approximately 4x10 . Since the two polymers were found to be so

similar in behavior no future work is planned for material UH-B. However

during FY 1981 we have extended our studies to a third UHMWPE obtained

from the same manufacturer. This material, which we shall identify as UH-C,

has a manufacturer specified intrinsic viscosity of =18 dl/g which

corresponds to a molecular weight of approximately 2x10^.

2. Polymer Morphology

2.1 Aspects of the Density (p) and Crystallinity (x)^ of Compression

Molded UHMl^^PE.

Among the features which are currently being investigated in

connection with the characterization of the dependence of the mechanical

behavior of the polymer on the structure and morphology of UHMWPE, is the

determination of the range of crystallinities which can be achieved in

compression molded samples of UHMWPE made from the two batches of raw

polymer whose limiting viscosity numbers are 25 dl/g and 18 dl/g respectively.

As indicated in the First Annual Report [1] the density of sheets

made from UH-A when molded at 200°C and cooled slowly to room temperature

(~1°C per minute) was in the range 0.934-0.936 g/cm^ (x=51%) . In order to

determine whether appreciably higher crystallinities could be attained in

sheets from this batch of raw polymer the effect of annealing in the

* See reference [1], page 5, for a brief discussion of the intrinsic

viscosity - molecular weight relationship in UHMWPE. Reference [2]

provides a detailed discussion of this relationship.

t Wt.% crystallinity.
the range 125®C on the density of slowly cooled samnles has since been

examined. The results, which are summarized in Table 1, show that the

maximum density which could be attained even after successive annealings

at 130°C (48 hrs.) and 135€ (96 hrs.) was only 0*941^ g/cm^ (x=63.4%).

Table 1

Effect of Annealing on the Density (p) and

Crystallinity (y) of Slowly Cooled Sheets of

Compression Molded UH-A UHMWPE .

Annealing Treatment Density (p) g/cm' Crystallinity (x) %

(i) Unannealed 0.934^ 58.7

(ii) 125'*C(48 hrs.) 0.9353 59.2

(iii) 125°C(48 hrs.)+130“c (48

hrs.)

0.936g 60.2

(iv) 125°C(48 hrs.)+130‘*C (48

+ 132°C(48 hrs.)

0.939^ 62.1

(v) 13o'*C(48 hrs.) + 135*c

(96 hrs.)

0.941g 63.4

* Annealing was carried out in N2 atmosphere

t Measured at 23‘^C in a water/ethanol density gradient column

Experiments have been started on samples made from the UHMWPE raw

polymer having a limiting viscosity number of 18 dl/g (UH-C) . Sheets molded

O O at 200 C and cooled slowly to room temperature (~1 C per minute) had a

O density of 0.9412 g/cm (x=63.1%). The density of these samples after

O o they were annealed at 130 C (72 hrs.) and then at 135 C (48 hrs.) was

0.952y (x=70.7%) which is appreciably higher than that of sheets of the

higher molecular weight polymer (UH-A) annealed under comparably conditions.

On the other hand, quenching the UH-C from the melt resulted in a density

at 0.925

q

g/cm^, a value which was only slightly higher than that for the

quenched UH-A. The densities of the UH-C samples, prepared under the three

different treatments described, are summarized in Table 2.
4 .

Table 2

Density (p) of Compression Molded Sheets

of UH-C UHMWPE Prepared Under Various

Thermal Histories.

Treatment 3 t Density (g/cm ) Crystallinity (v')%

Quenched

Slowly Cooled

Annealed at 130°C

(72 hours)

Followed by 135°C

(48 hours)*

0.925q 52.2

0-9412 63.1

0.952^ 70.7

*Annealing was carried out in N2 Atmosphere

tMeasured at 23°C in a water /ethanol density gradient column

In addition to indicating that higher densities can apparently be attained

with the lower molecular weight UH-C, which nevertheless still falls under

the appellation UHMWPE, the results summarized above provide some initial

guidelines for the preparation of sheets of UHMWPE of different molecular

weights but which have the same density and crystallinity. Comparison of

the mechanical properties of such sheets should provide insights into the

influence of molecular weight on the time dependent mechanical behavior of

UHMWPE as distinct from the effects of crystallinity.

In order to determine how the densities of the sheets of UH-A and UH-C

compare with those of manufactured prostheses, the densities of ten commercially

available unused and sterilized UHMWPE prostheses of different origins have

been determined. Most were labeled as having been Y”i^radlated for

sterlization and the dcsage was often specified as 2.5 Mrad minimum. In

some cases the dosage was not specified and in others the method of

sterilization was not given. Table 3 lists the devices and the sterilization

information provided.
Table 3

Summary of UHMWPE Commercial Prosthesis Components

Specimen Designation Device Type Method of Sterilization

1-A Acetabular Cup, 57mm O.D. Unknown

1-B Acetabular Cup, 45mm O.D. Unknown

2-C Acetabular Cup, 42mm O.D. Y-ray, 2 . 5 Mrad minimum

2-D Acetabular Cup, 58mm O.D. Y-ray, unknown dosage

2-E Tibial Plateau, Reg . , thick Y-ray, 2.5 Mrad.minimum

2-F Tibial Plateau, Small Y-ray, unknown dosage

3-G Tibial Plateau, Large Y-ray, 2.5 Mrad.

3-H Tibial Plateau, Small Y-ray, 2.5 Mrad.

4-1 Tibial Plateau, Large Unknown

4-J Acetabular Cup, 45mm O.D. Unknown

In each case density measurements were made on two pieces cut from

each prosthesis, one near a portion of the surface and the other from the

interior. The results are listed in Table 4. With the exception of the

pieces of specimen 2-E, in which case the difference in density between the

two pieces was ~.009 g/cm^, the differences in density between the two

pieces of the other prostheses were less than 0.003 g/cm . There were

however substantial differences in density among the prostheses whose

densities ranged from 0.923^ (Specimen A) to 0.952g g/cm (outer piece of

Specimen E). These values correspond respectively to that of the quenched

sheets of UH-A polymer and that of the slow cooled sheet of UH-C annealed

o ° successively at 130 C and 135 C (see above). In short, the densities of our

laboratory prepared sheets are certainly comparable to those of implants
used in the field. It should be borne in mind however that the implants

were subject to sterilization by irradiation, a treatment which may have

affected their densities [3,4].

Table 4

Density of Commercial Prostheses

Specimen Density

Surface

''(p) g/cm^

, Interior

Crystallinity (x)

Surface, Interior

1-A Acetabular Cup 0.923^, 0.923^ 51.2, 51.2

1~B Acetabular Cup 0.934^, 0.934g 58.5, 58.5

2-C Acetabular Cup 0.9352, 0.9343 59,1, 58.9

2-D Acetabular Cup 0.935q, 0.937^ 59.0, 60.7

2-E Tibial Plateau 0.952g, 0.944q 70.7, 65.0

2-F Tibial Plateau 0.950^, 0.9482 69.3, 67.8

3-G Tibial Plateau 0.937^, 0.936^ 60.8, 60.1

3-H Tibial Plateau 0.936q, 0.937^ 59.7, 60.4

4-1 Tibial Plateau 0.945g, 0.944g 66 . 2 , 65.5

4-J Acetabular Cup 0.942 , 0.942 64,3, 64.3

O * Measured at 23 C in a water/ethanol density gradient column.

Experiments have also been started on the preparation and characterization

of cylindrical specimens produced for use in compression tests. The following

procedure was established for vacuum molding of cylinders having a length o"f

2.54 cm and diameter of 1.27 cm.

1. The UHMWPE powder is Initially compacted in the mold

at room temperature using an arbor press.

The mold is placed in a platten press with light contact

pressure, the heat turned on, and vacuum applied.

2 .
7 .

3. The press is heated to 200°C (requires about 45 minutes)

and is held at that temperature for ten minutes.

4. The press is then turned off and the light contact pressure

released.

5. After the press has cooled for about 30 minutes to ~170°C a

plat ten pressure of approximately 20 MPa is applied.

6. The specimen is then slowly cooled in the mold under pressure

to below 90°C before removing the cylinder from the mold.

After removal from the mold the ends of the cylinders are

ground flat and parallel on a lathe using a high speed grinding

wheel.

In order to obtain some measure of material uniformity in the

molded cylinders one specimen of the UH-A polymer was sectioned and the

density of pieces taken from the interior and surface regions was

3 determined. The densities were 0.934^ and 0.934

q

g/cm respectively for

the two regions, which indicates a high degree of uniformity throughout

the specimen. The values were also very close to those determined for

the slowly cooled sheets of the UH-A polymer.

We have also Investigated the effect of molding cylinders

at a much lower temperature. For example, one cylinder was molded

using a maximum platten temperature of 165°C and following the same

procedures outlined above except for step 5. In this case the pressure

was applied immediately after the press was turned off. In Table 5

we compare the density of the complete sample, the density of an

annular ring of material, and that of a piece taken from the center

of the cylinder. As can be seen, there is a significantly higher

density in the core region than at the outer edge. Visual and optical
microscopic examination of cross sections of the cylinder revealed that

the interior region had a highly granular texture surrounded by a smooth

outer skin layer. This was most likely the result of a temperature

gradient in the mold, whereby incomplete melting occurred in the center,

and the center portion retained a more pronounced "grain memory" than

did the surface layer.

Table 5

Density Variation within a Cylinder of

UH-A Pol3rmer Slowly Cooled from 165^C

Sample Description

Bulk Specimen (Disc 2mm thick)

Outer Surface (2 mm thick annular ring)

Central Core (6.5 mm dia.plug, 2mm thick)

* Determined by hydrostatic weighing.

Density , g/cm~^

0.949

0.940

0.957

2.2 Uniaxial Stretching of UHMWPE

An aspect of the behavior of compression molded samples of

UHMWPE which is of interest in the context of implant performance is

their recovery from imposed strains below the breaking strain. Some

preliminary results carried out on slowly cooled samples of UH-A which

had been stretched at 37°C have been previously reported [1]. Data

have since been obtained for slowly cooled UH-A samples which were

stretched at 23 C and for slowly cooled samples of the lower molecular

weight UH-C samples which were stretched at 23°C and 37°C. The imposed

strains and corresponding residual strains after the samples were

released and stored at room temperature for twenty-four hours are

listed in Table 6. The specimens used in these experiments were

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

width of 0.32 cm and a thickness of ~0.1 cm. They were stretched
9 .

uniaxially at a constant rate of clamp separation (1 cm/min.)

corresponding to a specimen strain rate of ~0.02/min. The densities

2

of the UH-A and UH-C samples were 0.935 and 0.941 g/cm respectively.

As was the case for the UH-A samples stretched at 37°C [1] no

necking was observed during the stretching of the specimens in the present

experiments. This contrasts with lower molecular weight linear polyethylene (M rl90,000) which, under similar stretching conditions, undergo

w

necking at strains >0.2-0. 3. As can be seen from Table 6 both the UH-A

and UH-C samples exhibited considerable recovery from deformation from

the various strain levels imposed upon them. This latter behavior

distinguishes UHMWPE from lower molecular weight polyethylenes in which

recovery from straining is minimal xf neckxng occurs unless the specimen

is heated to a temperature near the melting point. It may be noted in

this connection that the UH-A and UH-C samples stretched at 23 C and 37 C

to strains of up to 2.00 contract in length, after being released and

stored for 24 hours at room temperature, by close to or more than half

the initial strain. The data in Table 6 also indicate that the fractional

3 Cirain recovery (relative to the imposed strain) which the samples undergo

is larger the smaller the imposed strain. The measurements reported in

Table 6 are the results of single tests at each strain, we shall therefore

not attempt any comparisons between the behavior of the UH-A and UH-C samples

or between the effects of temperature on both types of samples at this

stage, except to point out that the residual strains in samples subjected

to initial strains in the range 0.25—0.50 are about twice as large in the

UH-C samples as compared to the UH-A samples.
10 ,

Table 6

Residual Strain in Samples Stretched Uniaxially

at 23^C and 37*^0 after Unloading and Storage

at Room Temperature: for 24 hrs.

Residual Strain

Stretching Temperature 23°C 37°C

UH-A UH-C UH-A* UH-C

Imposed Strain

0.25 0.05 0.10 0.06 0.11

0.50 0.12 0.20 0.12 0.24

1.00 0.31 0.45 0.34 -

1.50 0.46 - 0.55 -

2.00 0.84 0.90 0.92 1.09

2.50 1.32 - - -

3.00 1.72 1.92 1.84** 1.86

3.50 - - 2.45 -

4.00 — - 2.80 2.98

* See Reference 1.

**We have previously reported a value of 1.15 [1] which, for reasons

which remain obscure, was much lower than expected when compared

with the residual strains obtained with the samples subjected to

lower and higher strains than 3.00. This new measurement is consistent

with the data corresponding to these other strains.

The changes in crystalline orientation corresponding to increasing

residual strains in UH-A samples stretched at 37°C were described in a

previous report [1]. The wide angle X-ray diffraction patterns

exhibited by four samples of UH-C which had also been Initially stretched

at 37°C are shown in Fig.l(a-d). The imposed strains were 0.50, 2.00,

3.00 and 4.00. The corresponding residual strains (Table 6) after

storage at room temperature for 24 hrs. were 0.24, 1.09, 1.86 and 2.98

respectively. It should be noted that the residual strains in the

samples after six weeks storage, toward the end of which period the

diffraction patterns were recorded, had decreased to 0.20, 0.84, 1.75

and 2.92. Accordingly the latter figures will be referred to in identifying

these samples in what ensues.
11 .

Fig. 1. Wide angle x-ray diffraction patterns showing the changes in

orientation of the orthorhombic 110 and 200 reflections and

the triclinic 010^ reflection exhibited by specimens of slow

cooled UH-C sheets which had been stretched at 310K (37°C) to

strains of (a) 0.5, (b) 2.0, (c) 3.0, and (d) 4.0, and then

released and stored at room temperature. The residual strains

after storage for 24 hours were 0.24, 1.09, 1.86 and 2.98

respectively. The corresponding residual strains after six

weeks (at the end of which period these diffraction patterns

V7ere recorded) had decreased to 0.20, 0.84, 1.75, and 2.92.
The wide angle X-ray diffraction pattern of the unstrained UH-C

was similar to that of the UH-A[1]. The progressive changes in the

diffraction patterns with increasing residual strain illustrated in

Fig.l were also similar in character to those found previously in the

case of the UH-A samples [1]. As the residual strain increases there

is a progressive increase in the intensity of the orthorhombic 200

reflections in the equatorial region and a corresponding diminution

of the azimuthal spread of the arcs which are centered about the equator.

Simultaneously the orthorhombic 110 reflections become broadly split

with the most intense portions of the arcs tending to be closer to the

equatorial direction with increasing residual strain. The split 110

reflections eventually merge on the equator in the sample having a

residual strain 2.92 (Fig. Id). In addition to the features described

above, and paralleling the case of the UH-A samples [1] the diffraction

patterns of the stretched UH-C samples also exhibited reflections

associated with the metastable triclinic [5] form of polyethylene. The

most Intense of these reflections can be seen in each of the patterns

shown in Fig. 1 in which they are identified as 010^ (indexing according

to Turner-Jones [5]). The changes in the azimuthal dispositions of these

arcs with increasing residual strain paralleled very closely those of the

orthorhombic 110 arcs. The 010^ arcs are situated at the equator in the

sample having a residual strain of 2.92 (Fig. Id).

We are in the process of recording the small angle X-ray diffraction

characteristics of the same samples in order to compare them with those

previously obtained from the UH-A samples [1]. It can be reasonably

anticipated that the results will be similar to those found in the case

of the UH-A samples.

In addition to the experiments described above we have started to

examine in more detail the fine structure of sheets of UH-A and UH-C
13 .

which had different residual strains. Examination of the surfaces of

samples subjected to initial strains in the range 1. 0-3.0 under the

optical microscope and in a scanning electron microscope indicated that

stretching did not occur uniformly throughout the samples and,

correspondingly, that the residual strain was not uniform throughout the

released samples.

The manifestation of non-uniform deformation was evidenced in

particular in scanning electron micrographs of the surfaces of the

samples by the occurrence of regions (valleys) which exhibit an

apparent "fibrillar" appearance with the ’fibrils' oriented parallel

to the stretching direction. These regions of higher deformation

[see eg. marker X in Fig. 2] are interspersed between regions (hills) of

lower deformation which exhibit no evident fibrillar character (see e.g.

Y in Fig. 2). This apparent non-uniformity in the distribution of the

deformation in the sheets is on a scale (~tens of microns) which suggests

it might be associated with the memory of the coarse grain-like

character of the raw polymer particles which is retained in the

compression molded sheets [1]. It is tempting to speculate that

deformation of the polymer is initiated preferentially at the boundaries

between the fused raw polymer particles. This possibility will be

investigated further.

3. Mechanical Behavior Studies

The mechanical behavior studies described in the First Annual Report [1]

dealt primarily with the characterization of the UH-A material in uniaxial

extension. During FY 1981 this work was continued and now includes preliminary

results for simple compression. Also, as we noted earlier our study has been

expanded to include the UHMWPE (UH-C) material which has a lower molecular

weight.
14 .

Fig. 2. Scanning electron micrograph of a portion of the

surface of an UHMWPE sheet which was stretched at

37°C to a strain of 2 and then released and stored

at room temperature. (Residual strain after 24 hrs.,

0.92, stretching direction horizontal). Note apparent

’fibrillar' appearance of regions (x) , and the absence

of this feature in regions (y) . See text.

I

i
As will be seen in what follows distinct differences were observed

between the mechanical behavior of the UH-A and UH-C samples.

3.1 Creep and Stress Relaxation in Uniaxial Extension

Uniaxial creep data are shown in Figure 3 on log-log coordinates

for the slowly cooled UH-A, the polymer having the higher molecular

weight. The data cover times up to about one year. At the present time

only the specimen having an applied stress of 27.5 MPa has failed (by

fracture) as is indicated by the arrow in the diagram. Unlike lower

molecular weight polyethylenes, these samples of UHMWPE did not exhibit

the phenomenon of necking and cold-drawing, but rather extended

homogeneously (on a gross scale) to a maximum stretch of approximately

e=5. In the case of the annealed UH-A (data not shown), two specimens

tested at loads of 27.5 and 9 MPa showed essentially the same creep

behavior as the slowly cooled material under the same loads. The one

difference that did occur was that the annealed specimen loaded at

27.5 MPa failed at about one decade in time sooner than did the

slowly cooled material.

By way of further comparison, creep curves for the quenched

UH-A are presented in Figure 4. This material is more compliant than

the slowly cooled UH-A, as is evidenced by the fact that the time

taken to reach a strain of unity for the quenched sample is about one

tenth that for the slowly cooled sample. However the maximum elongation

obtained for the quenched UH-A (£•^5) is still essentially the same as that

for the slowly cooled or annealed materials. A most Interesting feature

of the creep behavior was exhibited by the quenched polymer under a load

of 9MPa. At the early times the creep rate was nearly constant. After

about one day the creep rate decreased and the creep curve appeared to be
16

Nivais

Figure

3.

Creep

strain

versus

time

for

the

slowly

cooled

UH-A.

The

numbers

indicate

the

magnitude

of

the

applied

engineering

stress

and

the

arrow

the

time

at

fracture

for

the

specimen

loaded

at

27.5

MPa.
17

W 13

Ui C

0) to

Nivais

Figure

4.

Creep

strain

versus

time

for

the

quenched

UH-A.

The

numb

indicate

the

magnitude

of

the

applied

engineering

stress

the

arrow

the

time

at

fracture

for

the

specimen

loaded

at

27.5

MPa.
18 .

•tending toward a plateau. However, after about one month, a sharp

increase occurred in the creep rate before again reaching a new plateau.

If the same phenomenon were to occur in the case of the slowly cooled

sample under the same load (9 MPa) , then because of its lower compliance

one would expect that the abrupt change in creep rate would manifest

itself after a much longer time. Based on a comparison of the creep data

for the quenched and slowly cooled specimens subjected to higher loads,

we can speculate that in about one to two more years there may occur a

sharp increase in the creep rate of this specimen followed by a new

plateau region at even longer times.

Creep curves have also been obtained for the slowly cooled UH-C,

and these results are shown in Figure 5. While they are similar in

overall appearance to the curves for the slowly cooled UH-A (Figure 3)

,

two distinct differences are worthy of mention. First, at the higher

loads (_> 11 MPa) and in the region of rapid creep (strains in the range

0.2 to 3.0) the maximum creep rate attained is larger for the UH-C

than for the UH-A. The second difference in behavior is that the

creep curves for the UH-C show less of a tendency to reach a plateau

value at the highest levels of strain. The UH-C continues to elongate

to an e as large as 7 or more, in contrast with e-5 for the UH-A.

In Figures 6 and 7 are shown the results of single step stressrelaxation experiments done on specimens of the slowly cooled UH-A

and UH-C. The same data are replotted as isochrones of true stress

versus strain in Figure 8 on log-log coordinates. While there is

very little difference in behavior between the two materials over most of the strain range examined, both materials exhibit in their
19

Figure

5.

Creep

strain

versus

time

for

the

slowly

cooled

UH-C.

The

numbers

indicate

the

magnitude

of

the

applied

engineering
20

O

(ed) ridi

in

O

SSddlS

Figure

6.

True

stress

versus

time

for

the

slowly

cooled

UH-A.

The

data

were

obtained

from

single

step

stress-relaxation

experiments

in

extension

and

the

numbers

indicate

the

magnitude

of

the

applied

step

in

strain.
21

(Ed) ss3ais 3nyi

Figure

7.

True

stress

versus

time

for

the

slowly

cooled

UH-C.

The

data

were

obtained

from

single

step

stress-relaxation

experiments

in

extension

and

the

numbers

indicate

the

magnitude

of

the

applied

step

in

strain.
100.0

22

(edl/\l) SS3dlS BDUl

STRAIN

Figure

8.

Isochrones

of

true

stress

versus

strain

for

samples

UH-A

and

UH-C

slowly

cooled

from

the

melt

(from

single

step

stress-relaxation

data).

The

isochrones

are

for

times

of

from

10°

to

10^

seconds.
behavior at the larger strains (e>0.08) a departure from that obser^^ed

for lower molecular weight polyethylenes and many other semicrystalline

polymers. In a single step stress-relaxation experiment the specimen

is subjected to a sudden step in strain and the stress is monitored as

a function of time. For many semicrystalline polymers it is not possible

to obtain data in the region of large strain, since at strains larger

than about 0.08 to 0.10 the specimen either necks or fractures upon

application of the step in strain. Normally on a plot such as that shown

in Figure 8 the isochrones tend to flatten as the strain is increased.

When the point is reached where the isochrones become flat the material

undergoes an instability and necking or fracture results. In the case

of the two slowly cooled UHMWPE samples this condition is not reached.

Although the isochrones tend to flatten towards a plateau as the strain

is increased in the range from 0.01 to 0.10, they do not become flat, but

rather continue to rise at higher levels of strain, at least up to a

strain of unity.

Quenching of the UH-A polymer produced the stress-strain behavior

shown in Figure 9. Qualitatively these curves are similar to those

for the two slowly cooled materials. The main differences are that (1)

the isochrones for the quenched material show somewhat less of a

tendency to flatten in the region of strain from 0.01 to 0.30, and (2),

because of lower crystallinity, they are shifted to smaller stresses.

For example, at a strain of 0.01 and time of 10 seconds the stress for

the quenched UH-A (p=0.923 g/cm ) is smaller by a factor of 1.8 than

3 that for the slowly cooled UH-A (p=0.935 g/cm )

.

The effect of annealing on the stress-strain behavior of slow

cooled UH-A has been examined and is compared with the unannealed UH-A

in Fig. 10. Whereas for the slowly cooled material the isochrones

increase at all levels of strain, such is not the case for the annealed
24 /o <I g 

tj s 5 2 C "o « O S (0 O* *H

4) (U o • <

flC H

(0

fSw

4J

(0

(03

(0

(0

(0

(UM

4J

(0«3W

4J

U-lo

(0

0)oM

JCoo

(0 M

(d iJ

(d •o •0) e 13 o c

•H

cd <U ^ “

rH^

(U o M iH I(0 o

(0

u

0)u o

*J o 0) iH a 6 0) o

V) u-l

d) 4H

iH

O

00a

u »H QJ •H S *J OM U

U-l

O

<UU9

00

edl/\l ‘SSBdlS Bfldl
25

Oo o oo 0 T3

iH (U « fHo < o 1 o

C/)

a, (0 0

CO

_T3

(0 c

to • Mo

u-ico S 0 (j c o

•H ftO CO W)

V4 CSJ *j <3- <r

to CO <T>

(Ti • o O II

II to

So

4J

•rlcnc

(U 'O

V)

tn

(U

}-i

CO

4J

•rltoc

(U TO

ooc

0) 0)

:3 0

CMo 0

CO o

(Uu

C CM

T3 O (U M t3 § ^ CU O

CJ rH rH O O r^

to o o M O CM

to JQ

<UU3

bO

•H

^diAi ssBdis 3ndi
material. During application of the largest step the annealed

UH-A necked, and its behavior, shown in Figure 10, is more typical of

lower molecular weight polyethylenes . This same result was also found

to be true for a specimen of annealed UH-C. Thus, by increasing the

3 crystallinity of the UH-A from that of the slowly cooled state (-0.935 g/cm )

by only about 5% the nature of the stress-strain behavior, as determined

from stress-relaxation data, has been significantly altered. Note also

3

that although the slowly cooled UH-C (p=0.941 g/cm ) and annealed UH-A

3

(p=0.942 g/cm ) have approximately the same density, one material exhibits

necking whereas the other does not.

We have also carried out creep and stress relaxation experiments on

specimens machined from two of the commercial implant devices. The two

tibial plateaux labeled 3-G and 4-1 were the only ones from which large

enough specimens could be machined to fit our stress relaxation apparatus.

3 Material 3-G had a density of =0.937 g/cm which is about comparable to

that of our slowly cooled UH-A. Sample 4-1 had a density of =0.945

3 g/cm which is intermediate between our slowly cooled UH-C and

annealed UH-C. The behavior of these two commercial materials can

be summarized as follows. Material 3-G showed a creep and stress

relaxation behavior very similar to that of the slowly cooled UH-A (see

for example Figures 3 and 8) . The only exception was that in creep two

specimens tested at applied stresses of 25 and 16 MPa failed in times

2 7 of 6x10 and 1.1x10 seconds respectively, whereas the two specimens

of the slowly cooled UH-A tested at comparable stresses (Figure 3)

have not yet failed after more than one year (3.15xl0^s). In contrast

with the 3G material, material 4-1 exhibited a behavior more reminiscent

of the annealed UH-A or annealed UH-C (Figure 10). In creep, a specimen

under a stress of 16 MPa necked at a strain of about 30% and then

fractured in a time less than 1.5xl0^s. In stress relaxation necking

occurred during the application of a 20% step in strain.
27

In summary, uniaxial stretching of the UTM^PE, on a microscopic

scale, does not occur uniformly throughout the sample. Correspondingly,

upon release the residual strain in the sample is not uniform throughout.

Examination under the optical microscope and in a scanning electron

microscope reveals that there are regions of the surface (valleys) which

exhibit a "fibrillar" appearance with the ’fibrils' oriented parallel to

the stretching direction. These regions of higher deformation are

interspersed between regions (hills) of lower deformation which exhibit

no evident fibrillar character. The apparent non-uniformity in the

deformation of the sheets is on a scale (-tens of microns) which suggests

it may be associated with the memory of the coarse grain-like character

of the raw polymer particles.

The processing conditions can significantly influence the mechanical

behavior of UHMWPE in uniaxial extension. In creep and stress relaxation

tensile specimens of either quenched or slowly cooled UH-A and UH-C are

observed to deform uniformly, on a gross scale, up to the point of fracture.

This behavior departs from that found for lower molecular weight

polyethylenes which exhibit necking in creep at a strain greater than

20-30% and in stress relaxation necking occurs during the application of

a step in strain greater than about 8-10%. However, when annealed at a

temperature near the melting point, both the UH-A and UH-C do exhibit

necking in a manner reminiscent of lower molecular weight polyethylenes.

Therefore, under a specific set of deformation histories we see that by

increasing the crystallinity only a few percent through annealing

(by 5% in the case of the UH-A) the nature of the stress-strain behavior

has been completely changed. An interesting result is that the slowly

3 3

cooled UH-C (p=0.941 g/cm and the annealed UH-A (p=0.942 g/cm ) have

very nearly the same density, yet one material necks and the other one does

not. Therefore, density alone is not a reliable parameter for predicting

the mechanical behavior of UHMWPE.
28

3.2 Creep and Stress Relaxation in Simple Compression

An investigation of the creep and stress relaxation behavior of

UHMWPE in simple compression was started this year. Preliminary results

are presented below.

For the compression experiments molded cylinders of the slowly

cooled UH-A and UH-C were used. In the work to be described the maximum

strain was generally less than 15% which represents the limit which

could be achieved with our extensometer . A new apparatus is currently

under construction which will allow us to examine the creep behavior

both at longer times and at strains greater than 15%.

Creep data for slowly cooled UH-A and UH-C are presented in Figures 11

and 12 respectively. Overall, the behavior of both materials is rather

similar. For a given stress the creep rate is initially relatively low,

then increases somewhat with time reaching a maximum rate at around 100

seconds. Also, the curves are not quite parallel to one another and

tend to diverge at the longer times. In creep the principal difference

between the two polymers is that the UH-C, which is the more highly

crystalline of the two, is slightly less compliant. For example, at a

stress of 17.5 MPa and time of 3355 seconds the UH-A has crept to a

strain of 0.14 whereas the UH-C has reached a strain of only 0.10. In

comparing the results for compression with those for uniaxial extension

(Figures 3 and 5) it is found that within much of the region for which

the data overlap there is very little, if any, difference in creep behavior.

However, differences do appear at the highest levels of applied stress.

In this case the material creeps more slowly in compression than it does

in extension. At an engineering stress of 16 MPa and time of 3300 seconds

we can estimate that the UH-A would creep by about 12% in compression

(Figure 11), whereas in extension it has crept about 20% (Figure 3). This
29

ujej^S dAissejdiuoo

Figure

11.

Creep

strain

versus

time

for

sample

UH-A

in

compression.

The

numbers

indicate

the

magnitude

of

the

applied

engineering

stress.
30

(0

e\ E

I

Figure

12.

Creep

strain

versus

time

for

sample

UH-C

in

compression.

The

numbers

indicate

the

magnitude

of

the

applied

engineering

stress.
31

difference may be a consequence of the fact that in compression the

true stress decreases with time, while in extension the opposite is true.

Stress-relaxation experiments have also been done on the UH-A and

UH-C in compression. The results are shown in Figures 13 and 14 in

which the values shown along the ordinate represent engineering stresses.

Although we do not show the plots here, when the data in Figures 13

and 14 are replotted as Isochrones of true stress versus strain, the

resulting stress-strain curves are essentially the same as those shown

in Figure 8 for uniaxial extension.

It is evident from the creep and stress-relaxation data in compression

just described, and those in extension, that UHMWPE does not exhibit a

linear viscoelastic behavior even at the smallest strains. If the behavior

were linear, then the isochrones plotted in Figure 8 would form a set of

parallel straight lines having a 45 degree slope. Another feature of

linear behavior is that it is possible from creep data to predict the

amount of recovery the material will experience upon being unloaded and

allowed to recover back toward its initial state. As an example of the

extent to which the UHMWPE polymer deviates from linear behavior we have

carried out a series of creep and recovery experiments on the UH-A material

using the same applied loads for which data are shown in Figure 11. The

results are presented in Figure 15 where the actual recovery data are

compared to the values predicted from linear viscoelasticity. As can be

seen, the actual recovery is significantly less than is the predicted recovery.

3.3 Fatigue under Creep and Cyclic Loading

The fatigue behavior under conditions of constant load (static

fatigue) and sinusoidal loading has been investigated for slowly cooled

O sheets of both the UH-A and UH-C at 296 C in air. In fatigue testing it

is common practice to subject specimens to a sinusoidal loading history

in tension of the form
100

32

o o o

^ d

edlAl ‘ssajis aAjssajduJOQ

'id’o

(/)

•N 0) E

Figure

13.

Engineering

stress

versus

time

for

the

slowly

cooled

UH-A

in

compression.

The

numbers

indicate

the

magnitude

of

the

applied

step

in

strain.
100

33

ediAl ‘SS3JIS dAjssejduioo

Figure

14.

Engineering

stress

versus

time

for

the

slowly

cooled

UH-C

in

compression.

The

numbers

indicate

the

magnitude

of

the

applied

step

in

strain.
34

4-1o

6 (0 O 01 M *H

4-1 U

0) 0) (0

(U

4-i n)

COVi

d> . .g CO S CO

>

0) *o

I <u ^ c

•o C

(U 0) rH *J o cu O 'C3 u

CO

CO iH& u

Q O

iH *H

CO >

CO <U rC

.c cu VI ^

O

O *H

•H VI JS CO > to iH O CU VI O

U

VI CO C *H

0) >

VIX Vi

0> CO

CU ^ H

CU CO S 60C

C CU O <U •H C

CO *H

CO 60

(UVis* 0 T3 O CU O *H c a.

•H (X

CO

CO VI 0) C JS

lU VI 0

•r( 4-1 Vl O

0)CV (U X TU o s

VI >» *H

Vl C

CU 60 > 2 O 0 o

cu (U Vl JC

VI 0 cu

CO VI

CO cv o

cu »H

o; *0 H C3 0 *H

ir>

cuVl3

60 •H

ujejis eAjssajdiuoo
35

a(t) = P+Q sin at ^ (1)

where P is the mean level and Q sin wt is the oscillatory component of

the applied engineering stress, ^ being 2-n times the test frequency. In

our experiments we have chosen one type of loading history often used in

fatigue tests, namely zero-tension sinusoidal loading, meaning that in

equation (1) P and Q are equal. For this loading history the maximum

strain in the specimen does not remain fixed, rather the specimen elongates

during the course of the experiment and may reach several times its

original length before fracture occurs.

As an aid in analizing fatigue data, various damage rules, or

failure criteria, have been proposed. We shall consider two here.

The first one, proposed initially by Bailey [6], is based on the

assumption that failure is the result of damage accumulation due to

stress alone. In integral form the time to fail, t„, is related to

the stress history, o(5), by the expression

TB[a(0]

= 1 , ( 2 )

o where Tg(a) is the time to fail function determined from experiments

done at constant stress, a, and C is a dummy variable in time. In

principle, if the function x„(a) is known, then the lifetime in any

D

other stress history can be predicted. In our own work and the work

of others [7-10] it is found that for experiments involving a

constant load history of the time to fail data can be described by an exponential function of stress of the form

T- (o) = Ae

B

-Bo

( 3 )
36

where A and B are constants determined experimentally. We can predict the

time to fail for the loading history described by equation (1) by

substituting equation (1) into equation (3) and then carrying out the

integation of equation (2). Such a procedure yields the following relation

. _ S (p)

H " I (BQ)

*

where tg is the fatigue lifetime, tg(P) is the lifetime in a constant

load experiment carried out at a mean stress P, and (BQ) is the zeroeth

order modified Bessel function. B, P and Q have the same meaning as in

equations (1) and (3). As can be seen, in this formulation the

fatigue lifetime is predicted to be Independent of the test frequency.

In Figure 16 the logarithm of the time to fail is plotted as a

function of the peak stress (P+Q) for specimens of the slowly cooled

UH-A tested at three different frequencies. Each data point represents

the average value of the logarithms of the failure times for a minimum

of six specimens. Also shown (solid line) are the results for specimens

tested to failure under constant load P. Two features may be noted.

First, it appears that the lifetime under constant load defines a lower

bound to the lifetime under cyclic loading, at least within the range of

frequencies covered in our experiments. Second, the behavior shown in

Figure 16 is clearly not that predicted by equation (4) since there is a

decided frequency dependence associated with the time to failure under

zero-tension sinusoidal loading. In short, the Bailey criterion does not

apply.

The second damage rule which we shall consider is one based on the

assumption that the failure process is "cycle dependent". By cycle

dependent is meant that the failure occurs after a constant number of cycles

independent of the frequency. This infers that the lifetime of a given
37

Peak Stress, MPa

Figure 16. Logarithm of the time to fail versus peak engineering stress

for specimens of the slowly cooled UH-A fatigued under zero- tension sinusoidal loading. The solid line corresponds to the

behavior observed for specimens tested to failure at a constant

load equal to one half the peak stress under sinusoidal loading.
Time

To

Fail,

seconds

38

Test Frequency, Hz

Figure 17. Frequency dependence of the fatigue lifetime of the slowly

cooled UH-A tested at four levels of peak stress. The dashed

line is the behavior which would be observed if the lifetime

were cycle dependent where it has been assumed that failure

occurred after 60 cycles.
sample is inversely proportional to test frequency (i.e., a Miner’s

Rule [7] type of damage accumulation). In instances where the behavior

can be described by such a rule routine fatigue testing can be

accelerated simply by increasing the test frequency. The frequency

dependence of the fatigue lifetime of the slowly cooled UH-A is shown

in Figure 17 for specimens tested at four levels of peak stress (P+Q)

.

Also, the behavior which would be observed if the lifetime were cycle

dependent is indicated by the dashed line where it has been assumed

that failure occurs after 60 cycles. This line is not unique, but will

shift up or down depending upon the assumed number of cycles to failure.

However, the slope of the line will be independent of the number of

cycles to failure. It can be seen from Figure 17 that the observed

behavior is also not "cycle dependent" since there is a clear difference

between the slopes of the experimental curves (full lines) and that

based on the "cycle dependent" rule (dashed line) , the slopes of the

former being smaller than the latter. In the case of the slowly cooled

UH-A we conclude that the fatigue behavior under zero-tension sinusoidal

loading lies between those predicted by the two damage rules considered.

It should also be pointed out that the behavior of the UHMWPE is not

necessarily characteristic of other polyethylenes. In similar tests

carried out in our laboratory on a lower molecular weight polyethylene

[7] the lifetime was found to increase with increasing test frequency, a

behavior which is opposite to that of the UHMWPE.

A. Environmental Stress-Crack Resistance

In addition to the mechanical studies reported above, we have been

investigating the environmental stress-crack resistance (e.s.c.r.) of

UHMWPE. A schematic of the apparatus being used is shown in Figure 18.

The method is a bent strip test, but it differs from the ASTM bent strip test

[11] in three ways: (1) All the specimens are constrained in such a way as
40
41

to maintain a constant geometry, (2) the specimens are subjected to a constant

applied load, and (3) the specimens are unnotched. The stress cracking liquid

being used is a ten percent solution of nonylphenoxypoly (ethyleneoxy ethanol)

in distilled water. Results for the UH-A and UH-C polymers prepared under

varied processing conditions are presented in Table 8. The tests were carried

O out at a temperature of 90 C and an applied load of 5 MPa. This rather high

temperature was chosen in order that data could be obtained in a reasonable

period of time.

Notwithstanding the scatter in the data shown in Table 7, it is

apparent that the e.s.c.r. of specimens made from UH-A and UH-C depend

upon the processing conditions, although there is no clear-cut

correlation of failure time with density. Based upon the available

data, annealing has reduced the e.s.c.r. of both materials substantially.

Yet in the case of the UH-A the e.s.c.r. of the quenched specimens falls

between that of the slowly cooled and that of the quenched materials.

Although there is a considerable spread in the failure times shown in

Table 8, even the shortest failure time for the UHMWPE (annealed UH-A)

is substantially greater than for lower molecular weight polyethylenes

slowly cooled from the melt. For example, a linear polyethylene (M^^-10^)

~ 5 .and an ethylene-hexene copolymer (M -2.2xlC ), under the same test w

conditions, had average failure times (average of ten specimens) of 1.0

and 4.4 hours respectively.
42

Table 7

Environmental Stress-Crack Resis tance* of

jWPE at 9.Q°C Prepared Under Va rious TbP-nnpl Histones

Material

Average

Failure

Time (Hrs.)

Shortest

Failure

Time (Hrs .

)

Longest

Failure

Time (Hrs.)

Density

(g/cm^)

UH-A (slowly

cooled)

tested in

vapor only

133 (a) 109 160 0.935

UH-A (slowly

cooled)

tested fully

submerged

207 (a) 156 256 0.935

UH-A (Quenched)t 151 (b) 82 308 0.923

UH-A (slowly cooled 69 (c)

and Annealed) -f- 69 69 0.942

UH-C (Slowly

Cooled) t

408 (d) 133 983 0.941

UH-C (Slowly 100 (c) 80 120 0.952

cooled and

Annealed) f * In a ten percent solution of nonylphenoxypoly (ethyleneoxy ethanol) in

distilled water.

+ Tested Fully Submerged.

(a) Average for five specimens

(b) Average for four specimens

(c) Average for two specimens

(d) Average for ten specimens

One consequence of operating at a temperature as high as 90 C is that

the stress-cracking agent is also present in the form of vapor if the

specimen is not fully submerged. It has long been suspected that in the

vapor phase the stress-cracking agent is even more effective in damaging

the polymer than it is in the liquid phase. In the case of the UHMWPE,

confirmation of this effect has been obtained (see Table 8). Of the ten

specimens of the slowly cooled UH-A tested five were tested in the vapor alone
43

and five were kept totally submerged. The average failure time for the five

in vapor was 133 hours, whereas for those kept submerged it was 207 hours.

0 ff0ct may become less important if the test is conducted at a lower

temperature where there may be less surfactant in the vapor.

o The use of an applied stress of 5 MPa and a temperature of 90 C produced

an interesting effect concerning the failure mode. An underlying feature

of a bent strip test is that the stress is concentrated in the bent region

of the specimen and that failure will consequently occur in that region.

Typically, a crack forms in the center of the strip at some point within the

bent portion and it then grows outward toward the edges. Just prior to failure

some cold-drawing may occur near the edges of the strip. For the two lower

molecular weight materials referred to above the point of failure always

occurred in the region of the bend irrespective of the temperature at which

the test was done. This was not the case, however, for the slowly cooled

UH-A and UH-C. For these two materials, which are softer than the more highly

crystalline lower molecular weight polymers, the applied stress of 5 MPa was

sufficient to cause significant elongation of the material during the tests

O at 90 C. Under these conditions stress-cracking , elongation, and failure

occurred preferentially along the straight unbent sections of the specimen.

The scanning electron micrographs in Fig. 19 serve to illustrate the

nature and extent of the damage which occurs in the unbent parts of samples

of UH-A during the course of the e.s.c.r. experiments. A view of portions of

the side surface (a) and the large surface (b) of a rectangular strip which

O had failed in the stress cracking agent at 90 C is shown in Fig. 19 (i) . A

mosaic of irregularly shaped blocks of different sizes which are interspersed

by deep ’canyons' can be seen at both surfaces. This appearance is typical

of all the surfaces of the strips. Another typical feature is that

examination of the base of the canyons [e.g. (c) in Fig. 19 (i) and Fig. 19 (ii)]
44

and

the

large

surface

’b'

of

a

rectangular

strip

which

had

failed

in

the

stress

cracking

agent

at

363

K.

Arrow

indicates

deformation

direction.

Note

the

mosaic

of

irregularly

shaped

blocks

of

different

sizes

in

both

surfaces.

The

blocks

are

interspersed

by

deep

canyons

such

as

’c'.

(ii)

A

view

at

higher

mangnif

ication

of

a

portion

of

(i)

.

See

text.
45

reveals an apparent ribbon-like fine texture oriented parallel to the stressing

direction. This latter feature suggests that the material underlying the

blocks is appreciably stretched, in contrast with the material in the blocks

themselves.

The nature of the successive stages (e.g. crazing, stretching, fracture) in

the evolution of the morphological features summarized above remains to be

characterized.

We have also carried out a number of tests on the slowly cooled UH-A

and UH-C at 75 C and an applied stress of 5 MPA. These results are

summarized in Table 8. In this case all eight specimens tested were kept

totally submerged in the stress-cracking agent and, unlike the situation at 90 C

the specimens showed very little tendency to elongate during the experiment

and the failure occurred in the region of the bend.

Finally, two specimens from the commercial implant 4-1 were tested at

363K. Of the ten devices examined this was the only one from which large

enough specimens could be machined to fit our test equipment. The density of

3 3 material 4-1 (0.945 g/cm ) lies between that of the annealed UH-A (0.942 g.cm )

and that of the annealed UH-C (0.952 g/cm ). The observed failure times for the

two specimens of 4—1 were 69.5 and 99.4 hours. The average failure time of

84.4 hours is also between those for the annealed UH-A (69 hours) and

annealed UH-C (100 hours). For these three material the e.s.c.r. is found to

be inversely proportional to the specimen density.

O

Table 8

O Stress-Crack Resistance of UHMWPE at 75 C

Material Average Failure Time (Hrs.)

UH-A (slowly cooled) 761 (a)

UH-C (slowly cooled) 906 (a)

* - in a ten percent solution of nonylphenoxy poly (ethyleneoxy ethanol)

in distilled water.

(a) average of 4 specimens.
46

5 . Summary

This report describes work done during FY 1981 under task 80-01,

NBS-FDA/BMD Interagency Agreement. The study is concerned with the

relationship between morphology and the mechanical behavior of ultra

high molecular weight polyethylene (UHMWPE) , which is the material

commonly used as one component of orthopedic implant devices. An

important aspect of this study is to determine the range of material

behavior variability which can be introduced by changes in the processing

conditions and choice of starting material. The project was begun in

FY 1980 and the initial years' activities are described in the First

Annual Report [1]. In the summary which follows we shall include a

restatement of several of the more important results which were reported

in the First Annual Report [1].

The materials being investigated are two commercial grade UHMW

polyethylenes which were obtained from the same manufacturer. Based

upon the manufacturer's method of estimating molecular weight from dilute

solution viscosity measurements they have molecular weights of approximately

6 6 4x10 and 2x10 . They have been designated UH-A and UH-C respectively.

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 prosthesis

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

O temperatures as high as 90c above the melting point. This observation

emphasizes the importance of characterizing the structure of UHMWPE products,

not only to determine their crystallinity, spherulite morphology, and

lamellar fine structure, but also to monitor those aspects of their morphology
47

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.

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 140 and 141 **C. The precent 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.

Procedures were established and molding equipment constructed for

compression molding (under vacuum) the powder into sheets ;:1 mm thick.

Sheets were prepared by various thermal treatments which ranged from

quenching of the polymer from 200°C into ice water to slow cooling it from

200°C follow by annealing at a temperature as high as 135°C. The densities

and percent crystallinities of the various samples were as follows.
48

Treatment Density (g/cm~^) Wt.% Crystallinity (y)

UH-A (Quenched)

UH-A (Slowly Cooled)

UH-A (Annealed at 130°C

for 72 hrs. followed by

135°C for 48 hours)

UH-C (Quenched)

UH-C (Slowly Cooled)

UH-C (Annealed at 130°C

for 72 hrs. followed by

135°C for 48 hours)

0.923 51.0

0.935 59.0

0.942 63.4

0.925 52.2

0.941 63.1

0.953 70.7

We have examined pieces cut from ten commercially available

unused and sterilized UHMWPE prosthesis components of different

origins. Their densities varied from a low value of 0.924 g/cm"^

to a high value of 0.953 g/cm^. This large variation in density

among the commercial products is comparable to that found among

our sample preparations which range from the quenched UH-A to

the annealed UH-C. However, no direct comparisons can be made

since all ten of the commercial products had been sterilized by

y-irradlation which is known to affect the density [10].

Examination of cross-section of our molded 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 slowly 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

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

O in sheets compression molded at 230 C.

An aspect of the behavior of UHMWPE which is of interest in the context of

implant performance is its recovery from imposed strains below the breaking

strain. Unlike lower molecular weight polyethylenes, which exhibit necking

when stretched in uniaxial extension, the slowly cooled UH-A and UH-C elongate

uniformly on a gross scale up until the point of fracture. A series of

experiments was carried out in which specimens of the slowly cooled UH-A and

UH-C were stretched to strains of from 0.25 to 4.0. After stretching, the
49

specimens were released and allowed to relax at 23°C for a period of

24 hours. The specimens which were stretched initially to strains of

up to 2.0 all recovered after 24 hours by close to or more than half

the initial strain. Examination of the surfaces of samples subjected

to Initial strains in the range 1. 0-3.0 under the optical microscope

and in a scanning electron microscope indicated that stretching did

not occur uniformly throughout the samples and, correspondingly, that

the residual strain was not uniform throughout the released samples.

The manifestation of non-uniform deformation was evidenced in

particular in scanning electron micrographs of the surfaces of the

samples by the occurrence of regions which exhibit an apparent

"fibrillar" appearance with the ’fibrils' oriented parallel to the

stretching direction. These regions of higher deformation are

interspersed between regions of lower deformation which exhibit no

evident fibrillar character. This apparent non-uniformity in the

distribution of the deformation in the sheets is on a scale (-tens of

microns) which suggests it might be associated with the memory of the

coarse grain-like character of the raw polymer particles which is

retained in the compression molded sheets. It is tempting to speculate

that deformation of the polymer is concentrated preferentially at the

boundaries between the fused raw polymer particles.

The Influence of the different thermal treatments on the mechanical

behavior of UHM-JPE has been investigated with experiments in uniaxial

creep and single step stress relaxation in extension. As noted earlier,

in uniaxial extension both the slowly cooled UH-A and UH-C do not exhibit

necking, rather, they extend uniformly on a macroscopic scale up to the point

of fracture. The overall creep behavior of the two slowly cooled

materials was found to be similar, except that the UH-C creeps somewhat
50

faster at large strains and creeps to a maximum elongation at fracture

of about e=7 as compared to e=5 for the UH-A. The stress strain

behavior (true stress), as determined from stress-relaxation data, was

also found to be similar for strains of up to e=l. However, after

annealing both materials did exhibit necking in stress relaxation

experiments. In the case of the UH-A increasing the crystallinity by

only 5% through annealing had the effect of significantly altering the

nature of the stress-strain behavior. It was also observed that,

O although the slowly cooled UH-C (p=0.941 g/cm ) and annealed UH-A

3

(p=0.942 g/cm ) have very nearly the same density, one material exhibited

necking while the other did not. Therefore, density alone is not a

reliable parameter for predicting the mechanical behavior of UHMWPE

in the uniaxial extension.

Experiments were also started on the preparation of cylindrical

specimens for use in compression tests. Procedures were established

for compression molding (under vacuum) cylinders having a length of 2.54 cm

and diameter of 1.27 cm. In order to obtain some measure of material

uniformity in the molded cylinders one specimen of the UH-A polymer was

sectioned and the density of pieces taken from the interior and surface

3 regions was determined. The densities were 0.9342 and 0.934q g/cm

respectively for the two regions, which indicates a high degree of

uniformity throughout the specimen. The values are also very close to

those determined for the slowly cooled sheets of the UH-A polymer. At

this time only preliminary creep and stress relaxation data have been

obtained for cylinders of slowly cooled UH-A and UH-C. An apparatus is

currently under construction which will allow us to examine the creep

behavior at long times and large strains.
51

The fatigue lifetime behavior of the UHMWPE has been studied

under conditions of constant load (static fatigue) and zero-tension

sinusoidal loading. For this loading history the maximum strain

in the specimen does not remain fixed, rather the specimen elongates

during the course of the experiment and may reach several times its

original length before fracture occurs. As an aid in analyzing the

fatigue data we have considered two damage rules. One is based on the

assumption that failure is a result of damage accumulation due to stress

alone. In this formulation the fatigue lifetime is predicted to be

independent of the test frequency. The second damage rule is based

on the assumption that the failure process is "cycle dependent". Cycle

dependent failure implies that the failure occurs after a constant

number of cycles independent of the test frequency. Fatigue data have

been obtained at different levels of peak stress and test frequency for

samples of the slowly cooled UH-A. The observed fatigue lifetime is

found not to obey either damage rule. The fatigue lifetime exhibits a

clear frequency dependence and the slopes of the curves representing the

fatigue lifetime versus test frequency are smaller than those which are

based on the cycle dependent rule. Also, the fatigue behavior of the

UHMWPE is not necessarily characteristic of other polyethylenes. In

similar tests carried out in our laboratory on lower molecular weight

polyethylene [7] the lifetime was found to increase with increasing test

frequency, whereas for the UHMWPE the opposite is found to be true.

In addition to the studies of morphology and mechanical behavior we

have investigated the environmental stress crack resistance (e.s.c.r.) of

UHMWPE. The results for specimens of the UH-A and UH-C cut from sheets

prepared under the varied processing conditions described earlier can be

summarized as follows. The e.s.c.r. of both polymers depends upon the
52

processing condltons, although there is no clear-cut correlation of

the failure time with density. Based upon the available data,

annealing reduces the e.s.c.r. of both materials substantially. Yet

in the case of the UH-A the e.s.c.r. of the quenched specimens falls

between that of the slowly cooled and that of the annealed

materials. Although there is a considerable spread in the failure

times depending upon the processing conditions even the shortest failure

times for the UHMWPE are substantially greater than for lower molecular

weight polyethylenes slow cooled from the melt. Scanning electron

micrographs of portions of the specimen after failure indicate a

mosaic of irregularly shaped blocks of different sizes which are

interspersed by deep canyons. Examination of the base of the canyons

reveals an apparent ribbon-like fine texture oriented parallel to the

stressing direction. This latter feature suggests that the material

underlying the blocks is appreciably stretched, in contrast with the

material in the blocks themselves.

6. Work Plan for FY 1982

1. Experimental work begun in the area of uniaxial creep in

O extension and compression at 23 C will be continued. Attention will be

given to the construction of an apparatus for the purpose of investigating

the long term creep behavior in compression. Of particular Interest here

will be the influence of thermal history on the long time behavior in

compression. The creep experiments in uniaxial extension will be

O extended to a higher temperature (37 C) and to an environment of saline

solution in addition to air.

2. In the area of morphology studies we shall continue to investigate

changes in morphology associated with deformation. This work will include

the examination with wide and low angle x-rays of the polymer at various
53

stages of deformation under load. It is of interest to ascertain

whether differences in morphology can be identified for samples

prepared under varied thermal histories, but subjected to the same

deformation history. Another important aspect of this study is to

determine the extent to which the memory of the particulate nature of

the raw polymer may influence the deformation process.

3. Work already begun on the environmental stress-crack resistance

of UHMWPE will be continued. At the present time both the UH-A and

O O UH-C polymers have been examined at 75 C and 90 C. The present plan is

to continue these tests at a lower temperature and the tests may be

expanded to Include liquids other than the standard stress-cracking

liquid presently being used.

4. Experiments concerned with the effect of yii^i^^diation on the

morphology and mechanical behavior will be initiated. For this phase

of the work specimens having different densities will be Y“irradiated

at varied doses bracketing the range commonly used for sterilization

purposes. The mechanical behavior of the irradiated polymer will then

be compared to that of the unirradiated material. If deviations occur,

examination of the morphology of the irradiated material will be carried

out for specimens in the deformed state.
54

References

[1] G. E. McKenna, F. Khoury, and J. M. Crissman, "Relationship Between

Morphology and Mechanical Properties of Ultra High Molecular Weight

Polyethylene", First Annual Report to the Bureau of Medical Devices,

Task 80-01, NBS-BMD Interagency Agreement, October, 1980.

[2] H. L. Wagner and J. G. Dillon, "A Specification Method for Ultra High

Molecular Weight Polyethylene for Implant Use", J. Biomed. Matls. Res.,

13, 821 (1979).

[3] R. J. Roe, E. S. Grood, R. Shastri, C. A. Gosselin, and F. R. Noyes,

"Effect of radiation sterilization and aging on ultra high molecular

weight polyethylene", J. Biomed. Mat. Res., 209 (1981).

[4] H. J. Nusbaum and R. M. Rose, "The Effects of Radiation Sterlization on

the Properties of Ultrahigh Molecular Weight Polyethylene", J. Biomed.

Mat. Res., U, 557 (1979).

[5] A. Turner Jones, "The Triclinic Crystal Form of Polymethylenes and

Polyethylene", J. Polymer Sci., 533 (1962).

[6] J. Bailey, "An Attempt to Correlate Some Tensile Strength Measurements

on Glass: III", Glass Industry, 94 (1939).

[7] G. B. McKenna and R. W. Penn, "Time Dependent Fatigue in Poly (methylmethacrylate)

and Polyethylene", Polymer, 213 (1980).

[8] J. M. Crissman and L. J. Zapas, "Creep Failure and Fracture of Polyethylene in

Uniaxial Extension", Polym. Eng. and Sci., 19, 99 (1979).

[9] G. B. McKenna and R. W. Penn, "Time Dependent Failure of a Polyolefin Rubber

Candidate Material for Blood Pump Applications", J. Biomed. Matls. Res., 14 , 689 (1980).

[10] B. D. Coleman, "Application of the Theory of Absolute Reaction Rates to

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

[11] ASTM D 1693-70, "Standard Method of Test for Environmental Stress-Cracking of

Ethylene Plastics", 1976 Annual Book of ASTM Standards.
55

Figure List

Fig. 1. Wide Angle X-ray Diffraction patterns of slow cooled sheets of

UH-C polymer which were stretched unlaxlally at 37°C to various

strains, and then stored at room temperature. The specimens were

stretched to strains of (a) 0.50; (b) 2.00; (c) 3.00, (d) 4.00.

The residual strains after storage at room temperature for twnety-four

hours were 0.24, 1.09, 1.86 and 2.98 respectively. The residual strains

after six weeks storage, toward the end of which period the diffraction

patterns were recorded had decreased to 0.20, 0.84, 1.75, 2.92

respectively. [Stretching direction vertical; the patterns were

recorded with a flat plate camera using CuKoc radiation]

.

Fig. 2. (a) Low magnification scanning electron micrograph of a portion of the

surface of a slow cooled UHMWPE sheet (Polymer UH-A) which was

stretched to a strain of 2 and then released and stored at room

temperature. The residual strains after 24 hrs. and after eleven weeks

were 0.92 and 0.72 respectively. Stretching direction horizontal.

(b) A view at higher magnification of a portion of the area shown In

(a). Note the apparent 'fibrillar’ appearance of the regions (valleys)

denoted X and the smoother appearance of the regions (hills) denoted Y.

The significance of the faint strlatlons which are orthogonal to the

'fibrils' In the X-reglons and can be seen In (b) remains obscure at

this stage. (Stretching direction horizontal).

Fig. 3. Creep In uniaxial extension at 296K of sample UH-A slowly cooled

from the melt (density = 0.934 g.cm3). Numbers Indicate the magnitude

of the applied engineering stress expressed In MPa. The arrow Indicates

the time at which fracture occurred for the specimen loaded at 27.5 MPa.

Fig. 4. Creep In uniaxial extension at 296K of sample UH-A quenched from the melt

(density = 0.923 g/cm^). Numbers Indicate the magnitude of the applied

engineering stress expressed In MPa. The arrow Indicates the time at

which fracture occurred for the specimen loaded at 27.5 MPa.

Fig. 5. Creep In uniaxial extension at 296K of sample UH-C slowly cooled from

the melt (density = 0.941 g/cm^). Numbers Indicate the magnitude of

the applied engineering stress expressed In MPa.

Fig. 6. True stress versus time for sample UH-A slowly cooled from the melt.

The data were obtained from single step stress relaxation experiments

and the numbers Indicate the magnitude of the strain.

Fig. 7. True stress versus time for sample UH-C slowly cooled from the melt.

The data were obtained from single step stress relaxation experiments

and the numbers Indicate the magnitude of the strain.

Fig. 8. Isochones of true stress versus strain for samples UH-A and UH-C slowly

cooled from the melt. The data were taken from Figures 6 and 7.

Numbers Indicate the time In seconds.
56

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

Fig.

9. Isochrones of true stress versus strain for sample UH-A quenched

from the melt. Numbers indicate the time in seconds for a particular

isochrone.

10. Isochrones of true stress versus strain for sample UH-A slowly cooled

from the melt - 0, and annealed - 0. Numbers indicate the time in

seconds for a particular isochrone.

11. Creep in simple compression of sample UH-C slowly cooled from the

melt. Numbers indicate the applied engineering stress.

12. Creep in simple compression of sample UH-C slowly cooled from the melt.

Numbers indicate the applied engineering stress.

13. Stress-relaxation behavior in compression of slowly cooled UH-A. The

values of stress shown along the ordinate indicate engineering stress

and the numbers at the right various levels of strain.

14. Stress-relaxation behavior in compression of slowly cooled UH-C. The

values of stress shown along the ordinate indicate engineering stress

and the numbers at the right various levels of stress.

15. Recovery data for simple UH-A compared to the recovery predicted from

linear viscoelasticity. The numbers indicate the magnitude of the

applied engineering stress.

16. Logarithm of the time to failure versus peak engineering stress for

sample UH-A under sinusoidal cyclic loading.

17. Comparison of the frequency dependence of the fatigue lifetime of

sample UH-A to that predicted for cycle dependent behavior.

18. Schematic representation of apparatus used to determine environmental

stress-crack resistance.

19. (1) Scanning electron micrograph showing the side (a) and the large

surface (b) of a rectangular strip of compression molded UHMWPE

(Polymer UH-A, slowly cooled from the melt) which had failed under

tension (5 MPa) while immersed in a 10% aqueous solution of stress

cracking agent [nonylphenoxy poly (ethyleneoxy ethanol)] at 90°C.

The base (c) of a deep valley between blocks is sho^>m at a higher

magnification in (ii). Note the irregularly shaped blocks of different

sizes situated throug.h the surface of the strip in (i). Note also the

apparent ribbon-like appearance of the material at the base of the

*'valley" (e) in (ii). The arrows in (i) and (ii) point the direction

parallel to the stressing direction.
NBS-n4A IREV. 2 - 80 ) U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO.

NBSIR 82-2493

2. Performing Organ. Report No. 3. Publication Date

MAY 1982

4

.

TITLE AND SUBTITLE

NBS-BMD Interagency Agreement, Task 80-01, Second Annual Report

"Relationship Between Morphology and Mechanical Properties of UHMWP"

5. AUTHOR(S)

John M. Crissman, F. A. Khoury, and G. B. McKenna

6. PERFORMING ORGANIZATION (If joint or other than N BS, see instructions) 7. Contract/Grant No. NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

8. Type of Report & Period Covered

Annual Report

Oct. 1,1980-Sept. 30, 1981

9.

SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State, ZIP)

Food and Drug Administration

Bureau of Medical Devices

3257 Georgia Ave.

Silver Spring, Md. 20910

10.

SUPPLEMENTARY NOTES

Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11.

ABSTRACT (A 200~word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here)

This report describes work done during FY 1981 under task 80-01, NBS-FDA/BMD

(Bureau of Medical Devices) Interagency Agreement. The report covers the second

year of a four year project concerned with the relatship between morphology and

mechanical properties of ultra high molecular weight polyethylene (UHMWPE) . Various

aspects of the UHMWPE being investigated include (1) the establishment of procedures

for compression molding (under vacuum) sheets and cylinders of the polymer, (2) the

influence of thermal history and molecular weight on the time dependent mechanical

behavior such as creep and stress relaxation in both uniaxial extension and compressioii,

(3) dynamic fatigue under conditions of zero-tension sinusoidal loading, (4)

environmental stress-crack resistance, and (5) the optical and electron microscopy of

sheets of the UHMWPE which have been deformed to varying degrees in uniaxial extension.

12. KEY WORDS (S/x to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) Creep, Fatigue, Morphology, Polyethylene, Stress-Crack Resistance, Stress-Relaxation,

Ultra High Molecular Weight.

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PRINTED PAGES

pX~l Uni imited

1 1 For Official Distribution. Do Not Release to NTIS 60

r 1 Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C.

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QQ Order From National Technical Information Service (NTIS), Springfield, VA. 22161 $9.00

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