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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: [email protected]
M. C. Galetz
e-mail: [email protected]
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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