ORIGINAL PAPER

Molecular Deformation Mechanisms in UHMWPE DuringTribological Loading in Artificial Joints

Mathias Christian Galetz • Uwe Glatzel

Received: 28 August 2009 / Accepted: 8 December 2009 / Published online: 8 January 2010

� Springer Science+Business Media, LLC 2010

Abstract No clear picture of the deformation and wear

mechanisms of Ultra High Molecular Weight Polyethylene

(UHMWPE) in artificial knee joints exists up to today. Tri-

bological tests were conducted under relevant loads and the

worn samples were extensively studied by XRD, DSC,

Raman, and SEM. It was shown that stresses close to the

surface in areas where high relative velocity in combination

with high normal loads are applied are most effective in

changing the microstructure and therefore most detrimental

to produce wear particles, while in the depth the same

deformed structure as under unidirectional cyclic loading is

found. A model was proposed which reflects the deformation

at different zones in the depth of a tribologically loaded

UHMWPE sample. This model shows amazing analogy to

the orientation of collagen fibrils in natural cartilage.

Keywords Biotribology � Wear mechanisms �Fatigue analysis

1 Introduction

Ultra High Molecular Weight Polyethylene (UHMWPE)

combined with Cobalt–Chromium–Molybdenum (CoCrMo)

alloys is used as a bearing couple in artificial knee joints since

40 years. Although knee arthroplasty has a remarkable

clinical track record, problems with wear and fatigue of

UHMWPE continue to limit the longevity of knee replace-

ments [1].

Microstructural changes due to frictional loading are

precursors to wear. These changes occur due to one of the

main wear mechanisms: abrasion, adhesion, or fatigue.

Abrasion is a cutting process of the polymer chains and

occurs when hard particles or asperities plough through the

polymer [2, 3]. Adhesion is also effective right at the

surface and can even lead to the development of a transfer

layer on the metallic counterpart due to high affinity of

metals to carbon and hydrogen. This is typical of polymers

like polyethylene with a smooth molecular profile and the

lack of side groups and kinks in the chain [4, 5]. The

formation of the transfer layer depends on the environment

and is more severe when tested in water, compared with

serum, but it has even been found on retrieved implants [6].

In the depth the mechanisms fatigue mechanisms occur.

UHMWPE tibial inserts are subjected to periodical physi-

ological loading during walking at peak stresses far greater

than the yield strength [7]. The relevant fatigue failure

mechanisms are associated with yielding and accumulated

plastic flow processes of UHMWPE [8, 9]. Implant

retrieval analysis suggests that patient weight, activity

level, and time of implantation are associated with the

severity of component damage [10], manifested as creep,

delaminations, pitting, and even cracks at and underneath

the surface [11, 12]. Wear as a result of creep leads to

changes in the form and dimensions of the tibia inlays as

well as molecular orientations without any loss of mass.

This in turn influences the kinematics of the knee joint and

therefore the tribological behavior [13]. It is known that

close to the surface a highly oriented layer often occurs in

UHMWPE during sliding [14–17]. Nevertheless, this can-

not explain the macroscopic deformations and is only

one aspect of molecular rearrangement. There still exists

no comprehensive picture of the deformation and wear

mechanisms in artificial knee joints, in spite of the

M. C. Galetz � U. Glatzel (&)

Metals & Alloys, University of Bayreuth, Bayreuth, Germany

e-mail: uwe.glatzel@uni-bayreuth.de

M. C. Galetz

e-mail: mathias.galetz@uni-bayreuth.de

123

Tribol Lett (2010) 38:1–13

DOI 10.1007/s11249-009-9563-y

numerous publications dealing with it. One reason might be

that under stress, several deformation mechanisms occur in

polyethylene. These mechanisms have mainly been inves-

tigated on high-density polyethylene (HDPE) [18, 19].

They can be transferred to UHMWPE with the limitation of

the influence of the spherulitic structure of HDPE that does

not exist in UHMWPE.

Basically, the interlamellar mechanisms in the amor-

phous phase and the intralamellar mechanisms in the crys-

talline fraction have to be distinguished.

In the beginning, the deformation under stress is

restricted to the amorphous phase, because the glass-tran-

sition temperature of the amorphous phase of polyethylene

is far below room temperature (about -160 �C [1]),

resulting in a modulus and strength much lower than in the

crystalline domains. The activation energy for the defor-

mation of the amorphous phase is just 2–10% of the acti-

vation energy needed to induce chain-dislocation-slip

within the crystals [20]. It was shown that the deformation

of the amorphous phase is reversible up to a strain of 0.4%,

mainly due to the entropy-elastic behavior of the molecules

[21–25]. The main mechanisms in the amorphous regions

between the crystal lamellae are interlamellar shearing,

interlamellar stretching, and induced lamella stack rotation

[26].

At higher strains, additional mechanisms of deformation

occur in the amorphous phase. One mechanism that only

occurs under tension is the development of crazes. Espe-

cially at low temperatures or high deformation rates small

microvoids can be found. These microvoids are bridged

with molecules, but if the applied tensile load is sufficient,

these bridges elongate, break, and lead to cracks [27].

The further mechanisms at high strains all occur within

the crystalline part. Within the crystalline lamellae the

molecules are in a highly ordered state. In undeformed

polyethylene, these crystals have an orthorhombic structure

with the lattice parameters a = 740 pm, b = 493 pm, and

c = 254 pm [18].

The zigzag carbon chains of the molecules are parallel

to the crystal c-axis, which is perpendicular to the axis of

the lamellae. As a result, the long-chain molecules must be

folded back and forth many times on a plane perpendicular

to the axis of the lamellae [18]. At higher strain, the

deformation mechanisms in the crystals always act in

combination with the shear of the amorphous part [19]. In

the crystals, lattice shearing from orthorhombic to mono-

clinic configuration is found. Furthermore mechanical

twinning and dislocation motion takes place [22]. At high

strains, the critical resolved shear stress of dislocation

movement is exceeded and slip is induced. Dislocations in

polyethylene preferentially emerge around chain ends [28],

but can also occur in continuous chains as was shown by

[22].

The long-chain nature of polymer molecules requires

that the most preferred slip plane in polymer crystals

contains the molecular chain. Under this condition, the

polyethylene chains can remain unbroken through very

large deformations [29] and lead to a deformation-induced

orientation of the crystal structure in polyethylene. The

described mechanisms in the literature were mainly

investigated under the action of bulk stresses. It is inter-

esting, which and where these mechanisms occur in sam-

ples under relative motions and loads relevant for knee

joints.

This investigation of the orientation initiated by tribo-

logical loading in UHMWPE can also enhance the chances

to find beneficial material modifications.

2 Materials

The UHMWPE used in this study was surgical grade Chir-

ulen 1020 (POLY HI SOLIDUR�). The as-received, unaged,

non-sterile condition polyethylene bar was machined to the

required specimen geometry for the tribological test. Before

testing, the specimens were left for 24 h in air at ambient

temperature for curing. The counterpart material for the tri-

bological test was wrought CoCrMo28-6 (Zapp Medical

Alloys GmbH), machined, and polished to a surface rough-

ness Ra \ 0.1 lm, controlled by profilometric measurement

with a Perthometer C5D.

3 Methods

3.1 Self-Made Tribological Testing Device

A three-station wheel-on-flat testing device has been

developed (Fig. 1). In this tribological tester, the CoCrMo-

wheels have a diameter of 50 mm and the UHMWPE

samples a width of 15 mm, thickness of 10 mm and a

length of 80 mm, each weighing about 10 g. A single axis

twin peak Paul-type loading curve [30] was applied by

vertically pressing the testing chambers against the wheels.

The maximum load was 1.3 kN, which equals half of the

loading for a total knee joint during walking (ISO-14323-3)

and mimics one condyle per test chamber. Additionally the

flexion–extension motion of the knee was simulated by

the rotation angle of the wheel. A translation motion of the

flats reflected the anterior–posterior (A–P) motion of the

knee. In Fig. 2, a schematic drawing of the testing device is

shown, including the different loads applied. Figure 3 gives

an enlarged overview of the loading curves shown in

Fig. 2. The normal load is applied by pushing the test

chambers against the shaft with the wheels. To provide an

idea of the magnitude of the loading procedure, the relative

2 Tribol Lett (2010) 38:1–13

123

velocity between both bodies varies in the range of 118 and

-132 mm/s. The maximum pressure according the hertz-

ian equation is 14.5 MPa along with a contact width of

1.35 mm at peak load. A modulus of 200.000 MPa for

CoCrMo and 600 MPa for UHMWPE and a Poisson ratio

of 0.46 and 0.3 for the CoCrMo and the polyethylene,

respectively, were used for the calculation [31, 32].

The test was run 3 times to 3 million cycles with a fre-

quency of 1 Hz in distilled water at 37 �C. The rotational

movement about a vertical axis that is present in the natural

knee was omitted. As shown before, this reduces the wear

significantly [33–35]. The focus of the tests lies on the creep

behavior and deformation mechanisms. Low wear is bene-

ficial, because extensive wear complicates the determination

of the deformation processes close to the surface.

3.2 Dimensional Change

To evaluate the dimensional changes the wear scar depth

profile was determined with an Universal Surface Tester

(UST�, INNOWEP). The wear scar was scanned with a

steel needle, this needle moved with a constant velocity of

1 mm/s at a normal force of 1 mN.

3.3 Scanning Electron Microscopy

For scanning electron microscopy (SEM), samples were

sputter coated with gold and investigated using a Zeis-

s1540EsB Cross Beam. In order to resolve the crystalline

lamellae structure, the samples were etched by a technique

described and investigated in detail in [36]. Samples were

cleaned with acetone and ultrasonicated for 2 h in a

1.35 g/100-ml solution of potassium permanganate in 95–

98% pure sulfuric acid reagent. After removal from the

etchant, the samples were successively rinsed in ice-bath-

cooled sulfuric acid, 30% hydrogen peroxide, distilled

water, and acetone for 90 s each. One retrieved polyeth-

ylene inlay of a Miller-Galante-II implant was also

investigated to compare it with the in vitro tested samples.

Unfortunately, no medical history was known for this

implant.

3.4 X-ray Diffraction (XRD)

XRD was carried out using a Bruker 5005 diffractometer

with cobalt radiation (k = 179 pm). With a penetration

depth of about 1 mm the orientation is measured in the

polyethylene [37]. The diffractograms were scanned in 2hranges from 20� to 45� at a rate of 0.5�/min to determine

the orthorhombic (200), (110), and (020) peaks. From the

(200) and (020) peaks the crystal orientation of the trib-

ologically loaded samples was determined by scanning the

top view of the wear scar in loading direction as well as by

azimuthal scan and by irradiating the side of the sample, as

indicated in Fig. 13. Unfortunately, a detailed determina-

tion of a pole figure was not possible due to the problem of

the uneven surface profile after testing. The crystal orien-

tation was compared with samples tested in unidirectional

Fig. 1 Self-made tribological testing device

Fig. 2 Schematic drawing of the wheel on flat with the different

motions and loads

Fig. 3 Load, flexion/extension, and translation of the tribological

testing device

Tribol Lett (2010) 38:1–13 3

123

cyclic compression creep at 1 Hz in water at 23 �C and

23 MPa for 300.000 cycles, as described in detail else-

where [32]. One curve was derived in loading direction

(from the top) and the other perpendicular to the loading

direction (from the side).

3.5 Differential Scanning Calorimetry

Differential scanning calorimetry (DSC) was carried out

using a Mettler Toledo DSC/SDTA 821e on pristine

UHMWPE and wear debris. The debris was collected from

all tribological tests. Not enough material for the DSC

analysis could be collected in a single run. The samples

were heated under nitrogen atmosphere from 20 �C up to

180 �C at a rate of 10 �C/min, equilibrated for 3 min, and

were subsequently cooled down to 20 �C. After keeping

the temperature constant at 20 �C for 3 min, the samples

were again heated to 180 �C at a rate of 10 �C/min. From

the recorded thermograms, the degree of crystallinity was

calculated according to:

XC�DSC ¼DHC

DH0

ð1Þ

where DH0 = 291 J/g is taken as the enthalpy of fusion of

100% crystalline UHMWPE [1].

Additionally, by measuring the melting point, informa-

tion about the thickness and stability of the crystals could

be obtained. The crystal thickness lc is linked to the melting

point according to the Thomson–Gibbs equation [38]:

T ¼ T0 � 1� 2re

lcDH0

� �ð2Þ

T0 = 418.7 K is the melting temperature of a hypothetical

crystal of unlimited size, without any influence of surface

energy. The folding energy re = 93 mJ/m2 is a measure-

ment of the surface energy of the crystal ends, where the

chains are folded. DH0 is again the enthalpy of fusion of

100% crystalline UHMWPE.

3.6 Raman Spectroscopy

Raman spectra were recorded in back scattering conditions

using a Jobin Yvon Horiba LABRAM spectrograph at

9100 magnification. The laser was a monochromatic HeNe

laser with a wavelength of 632 nm. At such a configuration

the Raman signal was found to reach its maximum inten-

sity in polyethylene at a depth of 3 lm [39]. Therefore, the

Raman spectrum derives from the first few microns of the

sample.

In polyethylene, a simple two-phase model cannot

account for the observations from Raman measurements

and other methods [40, 41]. A third intermediate phase has

been proposed that is preferentially located at the crystallite

fold surface boundary. It has been associated with the

presence of chain loops and entangled chain segments in

this region [40, 42].

This fact complicates the quantitative analysis of

UHMWPE. The method originally developed by [40] and

confirmed by recent measurements [42] is used. According

to this method, the crystallinity XC-Raman can be calculated,

using:

XC�Raman ¼I1416

0:45I 1295þ1303ð Þð3Þ

I1416 is the 1416 cm-1 Raman band area assigned to the

orthorhombic crystal; I(1295?1303) is the area of an internal

standard band. It is sum of the two CH2 twisting bands, the

1303 cm-1 bands belongs to the amorphous part and the

1295 cm-1 band is assigned to the crystalline as well as to the

intermediate phase. The signal at 1303 cm-1 has a

bandwidth of about 30 cm-1 and the signal at 1295 cm-1

has a bandwidth of 7 cm-1. To determine the amorphous

fraction in polyethylene, usually the band at 1080 cm-1 is

used and the interphase is calculated by subtracting the

crystalline and amorphous fraction from 100%. Unfor-

tunately, this band is hardly visible in the ultra high

molecular type [43]. Another possibility is to split the two

internal bands and calculate the amorphous part XA-Raman

[40]:

XA�Raman ¼I1303

Ið1295þ1303Þð4Þ

For the calculation of the crystallinity, all spectra were

baseline corrected, normalized to unit area, and then fitted

using a combined Gauss-Lorentz sum function analysis to

evaluate peak intensities. The classification and assignment

of all different bands in polyethylene is shown in Table 1.

4 Results

4.1 Wear

After 3-million cycles no gravimetric average weight

change of the samples within the accuracy of measurement

was found, after correction for the weight gain of unloaded

samples placed in water. Nevertheless, a small amount of

debris in the water was visible by the naked eye as well as a

transfer layer of polyethylene on the CoCrMo wheels. In

Fig. 4, the deformation of the UHMWPE sample as a

function of cycles is shown. After 100,000 cycles the

samples already show a significant permanent deformation,

which further increases with a decelerating deformation

rate in the course of the test. On all samples a prominent

blister was observed, always in the same area of the sample

(Fig. 4).

4 Tribol Lett (2010) 38:1–13

123

4.2 Scanning Electron Microscopy (SEM)

Without etching only a general smoothening of the surface

could be seen, when investigated by SEM. After removing

the amorphous part by etching as described earlier, the

molecular structure could be discovered. Figure 5 shows

the typical lamellae structure of the etched pristine

UHMWPE. In the wear scar two typical modified struc-

tures were found, see Figs. 6 and 7. It was revealed that in

most parts of the wear track the structure of the lamellae

was completely destroyed, Fig. 6. Additionally highly

strained areas with microfibrils aligned in the direction of

friction could be found (Fig. 7). The same fibrils have been

observed earlier [44]. These oriented fibrils could fan-out

and turn into debris, an observation that is confirmed by the

appearance of a typical wear particle, Fig. 8. In the area of

the blister this orientation was most pronounced on all

samples. In lower magnification it was observed that on

one side the blister is attached to the bulk, on the other it is

loose (Fig. 9). Before etching, the different orientations

could not be observed and the whole wear track seemed

covered with a polymer layer. Furthermore, after etching, a

spherulitic type of excrescence was observed (Fig. 10).

These excrescences are a typical artifact of the etching

process. These bumps cover larger areas even on pristine

polyethylene the longer the etching time and the higher the

etching temperature [36]. Nevertheless, it is interesting that

these excrescences clearly aligned along the loading

direction (Fig. 11). It is remarkable that these aligned

excrescences were not only found on the samples tested in

the self-made tribological testing device, but were also

found on the retrieved component. On the samples tested in

the tribological testing device, the entire wear scar was

covered with strongly aligned excrescences parallel to the

sliding direction. On the retrieved knee inlay only some

areas showed this orientation, while the rest is covered with

circular bumps (Fig. 12).

4.3 X-ray Diffraction

Figure 13 shows the two different directions of irradiation

for the XRD measurements. The curves were recorded by

directly irradiating the wear scar from the top and from the

side of the sample. In Fig. 14, the respective 2-h scans of

the tribological tested sample after 3-million cycles in

Table 1 Classification and assignment of different bands observed in polyethylene [43]

Raman

band

1064 1080 1131 1170 1295

Mode tas(C–C) antisymmetric

stretching

t(C–C)

stretching

ts(C–C) symmetric

stretching

q(C–C) rocking s(C–H) twisting

Phase Crystalline, amorphous Amorphous Crystalline, amorphous Crystalline,

amorphous

Crystalline,

interphase

Raman band 1303 1370 1416 1440 1460

Mode s(C–H) twisting x(CH2) wagging d(CH2) bending d(CH2) bending d(CH2) bending

Phase Amorphous Crystalline, amorphous Crystalline Amorphous (interphase) Amorphous

Fig. 4 Wear scar profile after 0.1-, 1-, and 3-million cycles

Fig. 5 Lamellae of pristine UHMWPE visible after etching (SEM)

Tribol Lett (2010) 38:1–13 5

123

comparison with the reference spectrum of the undeformed

pristine UHMWPE are shown. For clarification the curves

are parallel shifted in vertical direction. The peaks belong

to the indicated crystalline planes and the intensity gave a

hint of the orientation. The measuring setup was the same

for all samples. That left fluctuations in the X-ray intensity

as the only source for errors, when the curves were com-

pared. In Fig. 15, the signal from the sample is shown that

Fig. 6 Destroyed lamellae at spot, where a wear particle detached

(etched, SEM)

Fig. 7 Microfibrils after etching of the amorphous phase that blurred

the non-etched surface in the wear scar after testing

Fig. 8 Microfibrillar appearance of a typical wear particle

Fig. 9 Typical appearance of the blister (etched, SEM)

Fig. 10 Spherulitic type of excrescence, which was observed on

pristine UHMWPE after etching

Fig. 11 Aligned excrescences on the wear scar of the in vitro tested

sample after etching

6 Tribol Lett (2010) 38:1–13

123

has been cyclically compressed for 0.3-million cycles in

water at 23 �C and 23 MPa.

The intensity of the monocline (010)m-peak was

observable in all samples, but was not significantly affected

by loading. However, the intensity of the orientation of the

different orthorhombic crystal peaks differed significantly.

A clear difference in the intensity of the (200)o and (020)o

could be noticed between the different samples and direc-

tions of irradiation, providing information about the ori-

entation of the orthorhombic lattice. In both, the cyclically

and tribologically loaded samples, the orientation of the

a- as well as the b-axes indicates that the (020)o plane

normal is oriented orthogonally to the direction of the

applied normal load. The (200)o plane normal is oriented

in a parallel way. The orientation of the crystals in the

depth of the tribological tested sample is shown in the inset

of Fig. 13. The configuration of the molecules within these

crystals was indicated by the zigzag lines. Azimuthal scans

of the (200) peak from the top and side of the samples did

not show any further orientation at the surface.

In Fig. 16, the peak ratio of the (200)o and (020)o peaks of

the tribologically loaded samples to the respective peaks of

the pristine UHMWPE as a function of the maximum wear

scar depth is shown. It shows that the degree of orientation in

the depth of the samples increased with deformation during

tribological loading. The results are in good accordance

with [45]. They found the same orientation for HDPE as a

result of a tribological rolling test, but without sliding.

4.4 Differential Scanning Calorimetry (DSC)

Figure 17 shows the first- and second-heating curves

determined by differential scanning calorimetry for pristine

UHMWPE and wear debris. Again, for better comparison

the curves are vertically shifted.

When the first- and the second-heating curve are com-

pared, it can be seen that the melting temperature of both

decreased. Because the samples were subjected to high

temperature during the first run, chain scission was induced.

In the second run, the samples had been melted and crys-

tallized again. Strain orientation effects can be revealed by

comparing the shape of the curve to that of the first run. Any

Fig. 12 Aligned excrescences in the wear scar of the retrieved MG-II

knee inlay

Fig. 13 Direction of X-ray irradiation, together with the determined

orientation of the orthorhombic crystal structure under the wear scar

after tribological loading

Fig. 14 X-ray diffraction pattern of the tribologically loaded sample

after 3-million cycles at 37 �C

Fig. 15 X-ray diffraction pattern of a sample reciprocally loaded

under unidirectional compression load

Tribol Lett (2010) 38:1–13 7

123

strain-induced crystallization had been erased in the second

run.

For the worn material compared to the pristine

UHMWPE, the melting point was found to be lower and

the crystallinity was largely increased, summarized in

Table 2. According to the Thompson–Gibbs equation,

Eq. 2, the lower melting point observed for the debris is

also an indication of thinner and less stable crystallites.

Figure 18 shows a typical Raman spectrum for the

minimum of the wear scar on the UHMWPE samples, run

in the tribological test to 3-million cycles at 37 �C in

comparison with pristine UHMWPE. Each curve is the

average of five measurements. From the complex CH2

bending vibrations between 1400 and 1500 cm-1, the

signal at 1416 cm-1 decreased after frictional loading

when compared to the unworn material. This band is

assigned to the crystalline fraction. The other two bands at

1440 and 1460 cm-1 increase.

The signals from different spots of the wear scar

appeared similar. The tendency toward an increase in

amorphous material at the expense of the crystalline frac-

tion is obvious along the wear scar, see Fig. 19. The typical

changes in the spectrum were already observed after 0.1-

million cycles of tribological testing. This means that the

molecular structure within the first microns on the surface

was changed at the beginning of the test. It did not alter

with further test duration.

5 Discussion

The Raman spectroscopy revealed a reduced crystallinity at

the surface after tribological loading. When investigated by

SEM, this layer was also found. It blurred the surface

appearance before etching, but could be removed com-

pletely during the etching process, what supports the

observation from the Raman spectroscopy that it is rather

amorphous. This film is ascribed to the intimate contact

of the UHMWPE with the counterpart. The adhesive

Fig. 16 Ratio of the (200)o and (020)o peaks of the tribologically

loaded samples to the pristine UHMWPE as a function of the

maximum wear scar depth

Fig. 17 DSC plot of different UHMWPE samples—first and second

(dashed) heating

Table 2 Crystallinity, according to DSC-measurements of the pris-

tine UHMWPE and the debris, according to Eq. 1

Heating Debris Pristine

1. Heating

Crystallinity (%) 61,8 47,0

Tpeak (�C) 134,5 137,9

2. Heating

Crystallinity (%) 63,0 43,0

Tpeak (�C) 132,1 135,2

Fig. 18 Typical Raman spectrum of the wear scars

8 Tribol Lett (2010) 38:1–13

123

interaction between the bearing-materials can be quite high

as proven by the occurrence of the transfer layer to the

CoCrMo wheels [46] and the ordered crystalline structure

within the polymer is destroyed.

Underneath the film, a strong microfibrillar orientation

was observed, shown in Fig. 7. This orientation of the

molecules in the zone close to the surface had been

observed earlier [47, 48]. To generate such a microfibrillar

structure, high stresses have to occur close to the surface.

In the area of the blister this orientation is most pro-

nounced. In Figs. 20 and 21, the area where the blister

occurred is compared to the loading parameters normal

load and normal load multiplied by the absolute value of

the relative velocity over the translation. The translation

motion reflects the wear scar length. The arrows in Figs. 20

and 21 indicate the direction of the wheel when it passes

over the UHMWPE flat and the small number give the

percentage of one loading cycle (compare to Fig. 2). It is

obvious that the blister and therefore the highest orientation

are not caused by high loads alone (Fig. 20). The blister is

generated in an area that is loaded twice during one cycle

with high loads, in combination with high relative veloci-

ties (high wheel slip) (Fig. 21). The high friction energy

can even increase the temperature locally [49], what

facilitates the orientation of the polymer.

The high deformation necessary to induce the microfi-

brils is only possible because of the tie-molecules that

connect the different crystallites. The superior ultimate

strain of UHMWPE, compared to lower molecular weight

polyethylene, is one explanation for the success of this

version of polyethylene in artificial joints. Nevertheless,

high stretching involves a degradation of the bonds

between adjacent microfibrils and the material below. The

highly strained material seems to induce a heterogeneous

nucleation for the redeposition of the etched, prior amor-

phous, polyethylene during the etching process and leads to

the observed aligned excrescences (Figs. 11, 12).

Fig. 19 Phase composition

derived from the surface by

Raman spectroscopy with 3-lm

penetration depth at different

spots of the wear scar after

3-million cycles

Fig. 20 Load distribution plotted against the translation of the

UHMWPE sample

Fig. 21 Absolute value of the normal load multiplied by the relative

velocity plotted against the translation of the UHMWPE sample

Tribol Lett (2010) 38:1–13 9

123

During deformation the stress is not equally applied

between the molecules due to their different chain length or

degree of stretching. Chain scission occurs. This leads to

further stress concentrations on the remaining molecules

and leaves highly reactive free radicals [50]. This finding is

confirmed by the DSC measurement of the debris that

shows higher crystallinity. The second DSC run resembles

the first. Therefore, the higher crystallinity is not only an

effect due to drawing, but also due to a decrease in

molecular weight. Additionally, the lower melting point

shows a degradation of the molecular weight (Fig. 17).

This explains why the highest degree of oxidation is always

found underneath the surface [51]. The crystallinity right at

very surface (1–3 lm) that was determined by Raman

spectroscopy decreases in contrast to that of the debris that

usually were 10 or more microns thick lamellae. This

shows that the large majority of wear particles does not

derive from the very top of the bearing surface, but is

generated when microfibrils get loose. This is affirmed by

the appearance of the typical wear particles in Fig. 8.

When the material deforms plastically into microfibrils,

the underlying material is highly sheared. Figure 6 shows

this area. The microfibrils were only sparsely connected to

the underlying material and can be easily removed by the

etching process. These observations also explain why wear

is much higher when multidirectional sliding is applied.

Then shear stresses additionally occur in-between adjacent

lamellae [34].

The comparable surface appearance after etching of the

samples from the tribological testing device (Fig. 11) and a

retrieved Miller-Galante-II implant (Fig. 12) is an indica-

tion that at least in some parts of the sample the wear

mechanism is similar. Unidirectional deformed material

seems to act as a nucleus for the alignment of the observed

excrescences.

Although some wear in the form of debris and transfer

was observed, gravimetric measurements were unable to

reveal a reliable weight loss in comparison to unloaded

samples put in water at the corresponding temperature.

With the wear rate that low, most of the contact defor-

mation reported in Fig. 4 is due to creep.

In the orientation measurement with XRD neither the

amorphous film in the first 1–3 lm nor the orientation of

the microfibrils that extent to a depth of 10–20 lm under

the surface could be determined. Instead in the depth of the

samples, the orientation of the a- and b-axes of the

orthorhombic crystal was obtained by XRD. The orienta-

tion of the c-axis is perpendicular to it and corresponds

directly to the direction of lamellar alignment. It is the

same after reciprocal compression fatigue and tribological

loading therefore the same mechanisms are believed to

occur under frictional as well as under reciprocal com-

pression loading.

Thus a model is favored, which was proposed by Gal-

eski et al. [52] for unidirectional loading and is illustrated

in Fig. 22. It involves a reorientation and rearrangement of

crystals under compression and ends up with a new long-

range order. In this rearrangement, the chain-slip processes

in the crystals act in combination with the shear deforma-

tion of the amorphous phase. The observed X-ray pattern of

the tribologically tested samples proves that this deforma-

tion occurs under physiologically relevant loads.

In summary, the orientation of the crystals in UHMWPE

after tribological loading resembles very much the orien-

tation of collagen fibrils in human articular cartilage,

observed by Benninghoff [53] and later approved by many

others and shown in Fig. 23. The amorphous layer at the

very top has no equivalence to natural cartilage, but

amazingly the collagen fibrils and the polymer lamellae

underneath show the same orientation. For comparison in

Fig. 24 an overview of the different crystal orientation in

the wear scar and the corresponding estimated depths are

shown. The problem in our test is that the normal load and

relative velocity changed along the wear scar. The degree

and depth distribution of microstructural deformation

changes along the wear track, as do the loading parameters;

Fig. 22 Deformation model for polyethylene under compression

loading [52]

Fig. 23 Orientation of the collagen fibrils in human cartilage [59]

10 Tribol Lett (2010) 38:1–13

123

further investigation is needed under constant loading

conditions to clarify the interactions. Nevertheless, the

general mechanisms and orientation are the same along the

wear track.

When these different zones of deformation are com-

pared with the general stresses under a slider (Fig. 25), it is

obvious how the microstructural changes were generated.

At the surface adhesion and therefore high shear stresses

destroy the crystalline structure of the polymer.

A slider and especially a slipping wheel also evoke

tangential forces. These stresses lead to highly oriented

microfibrils. These tensile stresses seem to be more

effective in changing the microstructure and therefore more

detrimental than the compression stress ahead of the slider

and the shear stresses underneath the slider. The tensile

stresses may also induce crazes under non-unidirectional

loading and therefore wear particles.

Under the orientated layer and a barely oriented transi-

tion zone mainly reciprocal compression stresses occur,

that lead to an orientation perpendicular to that close to the

surface. In-between is a zone where high shear is applied

according to the Hertzian model. Because of the friction

force at the surface the maximum shear occurs not as deep

as proposed by the Herzian equations. Nevertheless, it

occurs at a certain depth were the polymer can hardly adapt

and is even weakened due to the transition between the

microfibrillar orientation and the compression-induced

orientation.

The good agreement with results, reported for other

semi-crystalline polymers such as high-density polyethyl-

ene [45], polytetraflourpolyethylene [54], or polyamide [55,

56] suggests that the observed mechanisms are in fact

generic of many linear semi-crystalline polymers. The

prevalence of the different particular mechanisms instead

depends on microstructural properties such as molecular

weight or chemical composition and the loading conditions.

For the processing of UHMWPE, two improvements

can be proposed from our findings. A better wear behavior

might be achieved by preconditioning the samples by

exposure to compression loads at a temperature close to

the melting point. This is meant to create the required

orientation to reduce creep and shear within the material

during tribological loading. That approach had already

been tested with promising results [57]. Additionally, a

certain degree of crosslinking is beneficial to restrict the

huge plastic deformation close to the surface and reduce

the molecular orientation. A high molecular mismatch

between the certain zones is detrimental as it weakens the

strength of the polymer. This approach was proven to be

beneficial by the decreased wear behavior of crosslinked

material [58].

6 Conclusions

– In linear sliding wear is very low due to microfibrillar

orientation at the surface.

– Stresses due to high wheel slip in combination with

high normal loads close to the surface seem most

effective in changing the microstructure and therefore

most detrimental and effective in producing wear

particles.

– In the depth the same deformation as under unidirec-

tional reciprocal loading was found.

– The crystalline lamellae play a crucial role in the

deformation and wear mechanism of UHMWPE.

– A model (Fig. 22) was proposed which reflects the

deformation in different depth of a tribologically

loaded UHMWPE.

– The crystalline lamellae under tribological loading

arrange similar to the orientation of collagen fibrils in

natural cartilage.

Fig. 24 Model of the orientation in the UHMWPE that derives under

frictional loading (thicknesses not drawn to scale)

Fig. 25 Distribution of stresses under a slider

Tribol Lett (2010) 38:1–13 11

123

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