© 2016 kathryn leigh harris - university of...
TRANSCRIPT
TRIBOCHEMICAL INTERACTIONS OF A PTFE/ALPHA ALUMINA COMPOSITE AT
THE SLIDING INTERFACE: A MECHANISM FOR ULTRA LOW WEAR
By
KATHRYN LEIGH HARRIS
A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY
UNIVERSITY OF FLORIDA
2016
© 2016 Kathryn Leigh Harris
To Infinity and Beyond
4
ACKNOWLEDGMENTS
I would like to thank our collaborators at DuPont for their support of this research. In
particular, Gregory S. Blackman and Christopher P. Junk have been invaluable for their generous
contribution of time and effort over the last few years. I would also like to thank the PTFE
enthusiasts whose efforts in the area preceded, coincided with and contributed to the work
presented here, including Professors David Burris and Brandon Krick, and Dr. Angela Pitenis. I
am forever grateful to all past and present members of the UF Tribology Lab for their dedication
in constructing a truly one of a kind experience in teamwork and friendship.
I thank my advisor, Dr. Gregory Sawyer for not only his academic guidance, but also for
the friendship, strength and personal support I have experienced under his leadership that have
known no equal, and are unlikely to.
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TABLE OF CONTENTS
page
ACKNOWLEDGMENTS ...............................................................................................................4
LIST OF TABLES ...........................................................................................................................7
LIST OF FIGURES .........................................................................................................................8
ABSTRACT ...................................................................................................................................11
CHAPTER
1 INTRODUCTION ..................................................................................................................13
2 BACKGROUND ....................................................................................................................15
The History of Polytetrafluoroethylene ..................................................................................15 The Wear of Neat PTFE .........................................................................................................17
Friction and Wear of PTFE Composites .................................................................................19 Tribofilms and Ultra-Low Wear .............................................................................................21 The Effects of Surface Roughness ..........................................................................................23
The Effects of α-Alumina Particle Morphology .....................................................................27 Tribochemistry ........................................................................................................................31
3 METHODS AND EXPERIMENTATION .............................................................................45
Materials and Sample Preparation ..........................................................................................45
Tribometer and Wear Test Design ..........................................................................................46 Wear Rate Calculation, Friction Measurements and Uncertainty ..........................................46
Stylus Profilometry .................................................................................................................48 Small Molecules Experiments ................................................................................................49
X-Ray Photoelectron Spectroscopy ........................................................................................49 Infrared Spectroscopy .............................................................................................................50 Etched PTFE Tests .................................................................................................................51
4 RESULTS AND DISCUSSION .............................................................................................54
Friction and Wear ...................................................................................................................54
Stylus Profilometry .................................................................................................................54 X-Ray Photoelectron Spectroscopy ........................................................................................54
Infrared Spectroscopy .............................................................................................................56
5 CONCLUSIONS ....................................................................................................................72
APPENDIX: FUTURE CONSIDERATIONS: COUNTERFACE EFFECTS..............................75
6
Wear and Friction Experiments on Additional Countersurfaces ............................................75
Infrared Spectroscopy and X-Ray Analysis of Transfer and Running Films .........................78 Surface and Sub-Surface Evolution of the Aluminum Countersamples ................................80
LIST OF REFERENCES ...............................................................................................................99
BIOGRAPHICAL SKETCH .......................................................................................................107
7
LIST OF TABLES
Table page
2-1 Tensile and compressive yield strengths of unfilled PTFE as reported by *Rae109 and
**DuPont110. ......................................................................................................................34
2-2 A summary of the particle size results, which vary by method, and the wear rate of
the PTFE composite containing them ................................................................................40
8
LIST OF FIGURES
Figure page
2-1 The polymerization reaction of PTFE................................................................................33
2-2 The PTFE molecule is helical in structure .........................................................................33
2-3 PTFE crystallizes to form lamellae rather than spherulites ...............................................34
2-4 A suite of wear experiments with neat PTFE in various conditions against 304
stainless steel ......................................................................................................................34
2-5 Time lapse images of the wear of unfilled PTFE illustrate the transfer, agglomeration
and growth of PTFE wear debris islands ...........................................................................35
2-6 The wear rates of various PTFE composites with 5 wt. % filler added are plotted
versus their average friction coefficients ...........................................................................36
2-7 A representative plot of the wear rate of a 5 wt. % α-alumina PTFE composite
illustrates the run-in, transition, and steady state behavior of the polymer in sliding .......37
2-8 A summary of Urueña’s transfer film wear study50 ...........................................................37
2-9 Urueña’s transfer film wear data measured using a stainless steel pin50 is compared
to Ye’s similar study57 .......................................................................................................38
2-10 A plot of the total wear rates of a PTFE/α-alumina composite against an array of
surfaces with prescribed angular roughness ......................................................................38
2-11 In a reproduced plot from Harris et al.,76 the total wear rate of the PTFE/α-alumina
composite is plotted against the roughness angle of the countersurface ...........................39
2-12 The wear rates of various PTFE/alumina composites are plotted versus the supplier
designated particle size of the fillers as described by Krick92 and Blanchet63 ...................39
2-13 SEM micrographs and SLS data demonstrate how BET data provided by particle
suppliers may not be an accurate representation of true particle size ................................41
2-14 X-Ray Microtomography adapted from Krick et al.’s 2015 particle size study92 .............42
2-15 TEM micrographs of the running films confirm that nanoscale alumina particles are
present at the sliding surface of the composite ..................................................................43
2-16 XPS spectra and optical images adapted from Krick43,47 indicate chemical changes at
the sliding interface ............................................................................................................44
2-17 The wear rate of a PTFE/α-alumina composite is plotted vs. relative humidity, and
the wear rate of another PTFE/α-alumina composite is plotted vs. pressure .....................44
9
3-1 Schematic of linear reciprocating tribometer with a flat-on-flat pin configuration, and
a six-axis load cell and several LVDTs for data acquisition. ............................................52
3-2 A schematic of the stripe test run using a PTFE/5 wt. % α-alumina composite pin
against 304 stainless steel with a lapped finish. .................................................................52
3-3 Stripe tests were performed on the etched surfaces of unfilled PTFE and a PTFE/α-
alumina composite .............................................................................................................53
4-1 Wear rates and friction coefficients for each experiment run as a part of the stripe test ...64
4-2 Stylus profiles of transfer films after 1 cycle, 100k cycles, and 1M cycles of
development .......................................................................................................................65
4-3 High resolution XPS of the transfer film in various stages of development .....................65
4-4 Infrared reflectance results from the transfer film on a stainless steel surface ..................66
4-5 A Hamaker solution for the attractive energy between a flat surface and a cylinder ........67
4-6 The carbonyl region of IR spectra of the tribofilms is compared to small molecule
model reactions with perfluorinated carboxylic acids .......................................................68
4-7 The chemical mechanism responsible for PTFE/α-alumina tribofilm formation and
adhesion .............................................................................................................................69
4-8 The wear rate of the PTFE/α-alumina composite and neat PTFE compared to that of
the submerged composite ...................................................................................................70
4-9 The wear rates of the etched and unmodified surfaces of unfilled PTFE and of the
PTFE/α-alumina composite are plotted versus the total sliding distance. .........................70
4-10 Optical images of the transfer and running films of the etched and unmodified
surfaces of unfilled PTFE and of the PTFE/α-alumina composite. ...................................71
5-1 Radical chemistry at the sliding interface proceeds despite mild conditions (low
speed, low nominal contact pressure, and low frictional temperature change) .................74
A-1 A summary of the wear results from PTFE/α-alumina composite experiments against
various metal countersurfaces. ...........................................................................................85
A-2 A summary of the friction results from PTFE/α-alumina composite experiments
against various metal countersurfaces. ..............................................................................86
A-3 A summary of the results of PTFE/α-alumina composite experiments against various
copper containing countersurfaces. ....................................................................................87
A-4 A summary of the results of PTFE/α-alumina composite experiments against lead,
gold and platinum ..............................................................................................................88
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A-5 Friction traces from the PTFE/α-alumina composite experiments against three
aluminum alloys and against stainless steel, and optical images of the transfer films ......89
A-6 FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three types of steel ................................................................................89
A-7 FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three different copper alloys .................................................................90
A-8 XPS spectra taken from the running films formed by sliding the PTFE/α-alumina
composite ...........................................................................................................................90
A-9 FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three aluminum alloys ..........................................................................91
A-10 Stylus profilometry traces taken across the center of transfer films formed by the
PTFE/α-alumina composite against three aluminum alloys ..............................................91
A-11 FT-IR spectra illustrate the changes in surface chemistry before and after attempting
to remove the transfer films ...............................................................................................92
A-12 FIB cross sections taken within the wear track and in nascent areas of the three
aluminum alloys tested ......................................................................................................93
A-13 A summary of the wear results of the PTFE/α-alumina composite against a number
of Al 6061 T6 countersurfaces ...........................................................................................94
A-14 Micrographs taken within FIB trenches in the nascent surfaces of five differently
prepared Al 6061 T6 surfaces. ...........................................................................................95
A-15 Hardness and modulus results from nanoindentations performed on the surfaces of
five differently prepared Al 6061 T6 surfaces. ..................................................................95
A-16 A series of images depicting the lapping and polishing process at ALSPI. ......................96
A-17 An EDS spectrum collected from a sample of the solids separated from the lapping
compound taken directly from the lapping wheel used by ALSPI. ...................................97
A-18 Backscattered electron micrographs and EDS maps of the subsurface (~1 mm depth)
of two Al 6061 T6 samples ................................................................................................98
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Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy
TRIBOCHEMICAL INTERACTIONS OF A PTFE/ALPHA ALUMINA COMPOSITE AT
THE SLIDING INTERFACE: A MECHANISM FOR ULTRA LOW WEAR
By
Kathryn Leigh Harris
May 2016
Chair: W. Gregory Sawyer
Major: Materials Science and Engineering
The wear and friction behavior of ultralow wear polytetrafluoroethylene (PTFE)/α-
alumina composites first described by Burris and Sawyer in 2006 has been studied intensively in
the years hence. The mechanisms responsible for the remarkable improvement in wear over
unfilled PTFE are not yet fully understood. The formation of tribofilms on the countersurface
and the running face of the polymer is crucial to the ultra-low wear behavior of the composite on
a metal countersurface. The complete chemical mechanism of transfer film formation and
adhesion, and its role in the exceptional wear performance has yet to be elucidated. Some debate
exists regarding the role of chemical interactions between the PTFE, the filler, and the metal
countersurface. Some have concluded that chemical changes are not an important part of the
ultralow wear mechanism in these materials at all.
A “stripe” test allowed comprehensive spectroscopic studies of PTFE/α-alumina transfer
films in various stages throughout development and led to a proposed mechanism which details
the initiation and adhesion of the tribofilms formed on both surfaces of the wear pair. PTFE
chains (carbon-carbon bonds) are broken mechanically during sliding and undergo a cascade of
reactions to produce carboxylate chain ends that chelate to the metal surface and to the surface of
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the porous, friable alumina filler particles. This tribochemical process forms a robust polymer-
on-polymer system that protects the steel countersurface from abrasion, and the polymer surface
from wear. The system is able to withstand hundreds of thousands, and possibly millions of
cycles of sliding with almost no wear of the polymer composite after an initial period of high
wear during run-in.
A mathematical model in support of the hypothesis of mechanical scission of carbon-
carbon bonds in the backbone of PTFE in simple sliding contact is detailed, using the Hamaker
model for van der Waals interactions between polymer fibrils and the countersurface (a cylinder
and a flat surface). The proven necessity of ambient moisture and oxygen is explained in the
mechanism, and model experiments using small molecules further support the assignment of
reactions in the proposed mechanism to the processes at the sliding interface.
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CHAPTER 1
INTRODUCTION
The value of tribological advancement is abundantly evident. The estimated cost of wear
and friction processes in the U.S. is estimated to be on the order of 6% of gross domestic product
(GDP)1 or approximately $900 billion in 2011 dollars. A reduction in friction coefficient of a
few percent can have a significant impact on efficiency. In addition to the substantial financial
burden generally attributed to wear and friction processes in power generation and
manufacturing, the lives and comfort of recipients of joint replacement implants depends on the
ability to engineer long-lasting sliding interfaces. An overarching goal of tribological research is
to find materials that will operate reliably at a steady and low friction coefficient while
maintaining an ultra-low wear rate. A comprehensive understanding of tribological materials
will be required as engineers and scientists strive to expand our influence into the frontiers of
space, the human body, and a host of applications demanding performance and durability.
Surface science is an integral element of tribology. Most surfaces must come into contact
with others during their useful lifetime. Surface interfaces are therefore perhaps the most critical
design element in machines and mechanical assemblies, as bulk properties are largely understood
and easily tested. Tribological phenomena at these interfaces are fundamental to system
performance, longevity, and reliability. The two most common metrics for tribological
interactions are friction coefficient and wear rate. The friction coefficient of a system, µ, is
defined as the ratio of the applied normal load to the lateral resistance to sliding. The wear rate,
K, is typically described as a volume loss per unit force per sliding distance, and is commonly
reported in units of mm3/(N∙m). For changes to be made to these metrics, changes must be made
to the surfaces involved. Friction and wear are not material properties, and vary as wildly as do
the applications of the systems in question. Numerous parameters, including choice of materials,
14
the applied contact pressures, sliding velocity and geometry and environmental conditions must
be considered in tribological design. Tribological interactions are also time-dependent and
frequently may rely on seemingly rare interactions that are a function of mechanics across all
length scales, of chemistry, of physics, and of all manner of materials phenomena.
Commonly, sliding surfaces are protected by the addition of a lubricious separating layer
to postpone damage to or modification of, the working surfaces. Fluid lubricants are useful in
systems operating at high enough speeds to promote hydrodynamic lubrication, but along with
greases are largely less desirable than solid lubricants in applications at higher operating
temperatures and contact pressures, lower speeds, vacuum environments, or that involve
reciprocations. Solid lubricants are in general less sensitive to contaminants, and safer for use in
biological applications. Polytetrafluoroethylene (PTFE) is used in a large number of tribological
applications because of its exceedingly low coefficient of friction.2–6 However, the neat polymer
has a very high wear rate (K~7x10-4 mm3/(N·m)).5–9 PTFE composites have been widely studied
because the inclusion of particulate and polymeric fillers improved the wear rate by several
orders of magnitude.6,8,10–13 Wear abatement of PTFE using fillers of various sizes and chemical
composition appears to be mechanistically unique: PTFE itself has been shown to wear in a
manner quite contrary to other engineering polymers, and the mechanism has been of interest for
decades.5,6,13–15
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CHAPTER 2
BACKGROUND
The History of Polytetrafluoroethylene
Polytetrafluoroethylene (PTFE) is a tough thermoplastic which tends to creep and wear in
part due to weak, largely dispersive intermolecular forces. Stronger intermolecular forces are
prevented by the low polarizability of the very strongly ionic C-F bonds present. The relative
weakness of these intermolecular forces contributes to the low surface tension, low wettability
and low friction of the polymer. The inertness and thermal stability of PTFE are also due in no
small part to the strength of the carbon-fluorine bond.16,17
The discovery of PTFE in 1938 by Plunkett was a serendipitous event that occurred while
it was still commonly believed that fully substituted ethylenes could not be polymerized.18
Plunkett noticed that an especially inert, waxy solid had formed under pressure in a valve on a
cylinder of tetrafluoroethylene (TFE), and patented his discovery and polymerization methods in
1941.19,20
Today, PTFE is produced commercially. Pyrolysis of chlorodifluoromethane yields
tetrafluoroethylene, which is polymerized into PTFE using various free radical initiators by
several methods (Figure 2-1). Granular PTFE is produced via aqueous suspension
polymerization and is cut or milled into particles tens of microns in size, while fine powder
resins are manufactured using an aqueous emulsion method and separation by agitation or the
addition of electrolytes. The fine powders (~200 nm) tend to agglomerate as particles (~500 µm),
which may be dispersed in lubricants and paste extruded. PTFE is also available in aqueous
dispersion with non-ionic surfactants that may be used as coatings after application and heat
treatment.21 PTFE radicals are most commonly terminated by combination or by acidic end
16
groups rather than by disproportionation, or removal of fluorine atoms to form a terminating
carbon-carbon double bond.22
The PTFE molecule itself adopts a helical structure, as well as increase in the C-C bond
angle past the usual 109° (Figure 2-2). This is assumed to be due to the overcrowding of fluorine
atoms that would occur in the zig-zag configuration of hydrocarbons.23 Because of this, the
molecules are cylindrical, or rodlike in shape. Below 19 °C, the helix has a repeat unit that sees a
180° twist per 13 carbon atoms (known sometimes as form II) and packs triclinically, and above
19 °C, the helix expands to rotate 180° per 15 carbon atoms (known sometimes as form IV).24,25
Above 30 °C, a further phase change occurs and the helix becomes more irregular, though it
maintains its lateral hexagonal packing up to its melt temperature, 327 °C.19,23
The high molecular weight of PTFE (ten to hundreds of thousands or even millions of
repeat units)26,27 and its high melt viscosity affect melt crystallization.16 Above the melt
temperature, PTFE forms a gel rather than a liquid, rendering it non-melt processible. As such, it
must be sintered above the melt temperature rather than extruded or molded. Uniquely, PTFE
crystallizes to form banded lamellae with interlamellar regions of non-crystalline polymer, as
proposed by Speerschneider in 1962, rather than the more typical spherulitic configuration of
other polymers28 as illustrated in Figure 2-3. The rate of cooling affects the lamellar thickness,
the percent crystallinity, and thus the density of the polymer. PTFE molecules within the
crystalline lamellae lie perpendicular to the length of a band, and parallel to the striations.26
Despite its exceptional thermal, chemical and frictional properties, neat PTFE is fairly
soft when unreinforced. Though orientation of the molecules and percent crystallinity do affect
the hardness, toughness, and tensile and compressive properties of PTFE somewhat, it is not
17
capable of supporting high loads without succumbing to creep or failure. Temperature dependent
tensile and compressive strengths of comparable PTFE resins are given in Table 2-1.
The Wear of Neat PTFE
The extremely low friction coefficient of PTFE as well as its chemical inertness, thermal
stability and low vapor pressure contribute to its singular utility as a solid lubricant. Because of
its low resistance to sliding, it has been widely studied by tribologists for decades. However, its
wear rate is markedly higher than that of other neat polymers, and the characteristic ‘island’
shape of the transferred material is fairly unique. The wear rate of neat PTFE is around 10-4
mm3/N∙m, independent of surface roughness or environment.
In 1952, Shooter and Tabor observed that the low friction of PTFE seemed independent
of the nature of the countersurface, and that it did not seem to be a result of lubrication by a self-
mated surface film.2 McLaren and Tabor attributed the speed and temperature sensitivity of
PTFE to viscoelastic origins.29 Two frictional regimes were observed by Makinson and Tabor in
1964.5 They concluded that at low speeds and/or high temperatures, friction was extremely low,
and that thin, aligned films were drawn from the surface asperities during sliding. At high speeds
and low temperatures, friction was increased and transfer of material was greater and thicker.
Furthermore, they were surprised to observe that the adhesion of the transferred material was
quite strong. From they concluded that the low friction coefficient arose from shear failure
between crystallites rather than from lack of adhesion to the surface. Steijn confirmed the
alignment of PTFE fibrils with the sliding direction and concluded that the chain slide past one
another with some ease.30 Pooley and Tabor then attributed the low friction to the lack of side
groups on the PTFE chain, and observed that crystallinity and morphology of the polymer
appeared to have almost no effect on friction and transfer.31
18
Makinson and Tabor’s conclusion regarding subsurface shear5 and its relation to the
friction and wear regimes of PTFE led to further studies regarding its seemingly contrary wear
mechanism. Tanaka proposed that because no banded structure was visible within even the
thickest wear platelets, gross wear occurred due to destruction of the banded structure proposed
by Speerschneider28 and that fibrillous wear debris was a result of slippage of crystalline slices in
adjacent bands, rather than of the removal of an entire section of a band.8 Uchiyama observed the
motion and growth of transferred PTFE fragments and noted their propensity to combine with
one another and also to transfer back to the sliding surface of the pin.32 In 1992, Blanchet and
Kennedy observed cracks that seemed to be propagating in the direction of sliding, and
concluded that in its high wear regime, PTFE must wear by delamination as subsurface defects
initiated cracks that could not support the shear stress applied during accelerated sliding.6
A suite of wear experiments with neat, compression molded PTFE pins (flat on flat) were
performed in order to illustrate its variable behavior and to observe the apparent mechanisms
first hand (Figure 2-4). All tests were run as 25.4 mm long reciprocations on ~120 nm Ra
stainless steel, and at 50.8 mm/s unless otherwise specified. Two tests against stainless steel in
lab air yielded the expected high speed wear rate of neat PTFE (10-4 mm3/N∙m) and formed the
characteristic patchy transfer film. Submerged in filtered water, the wear rate was nearly an order
of magnitude lower and the adhesion of the wear debris was greatly reduced, though the debris
that floated away from the track varied in size up to visible fractions of millimeters. A self-mated
test resulted in wear rates comparable to the submerged test. At 15 °C (below the 19 °C phase
transition – though no dependence of wear on phase is known) and under impinging dry nitrogen
to prevent condensation of water on the sliding surface, the wear rate of the PTFE pin was as
high as it was during the standard lab air tests, but the wear debris did not form the usual islands.
19
Instead, the debris was powdery and fine. A test was run at 35 °C (above the 30 °C phase
transition) and 1 mm/s in order to illustrate the speed transition in wear. Indeed, the sample ran
for 50,000 cycles and wore very little. The differences in the wear debris during the higher speed
tests suggest that the islands seen in the wear tracks of unfilled PTFE may require a certain
thermal energy input to form, and that they form from smaller, pre-transferred debris particles at
the sliding interface as sliding progresses rather than always delaminating in large pieces.
Time lapse photos of the transfer film of a neat PTFE pin run on 20 nm Ra nickel at the
same conditions as the previous tests (wear rate was again ~10-4 mm3/N∙m) illustrated the
evolution of the debris islands as the number of sliding passes increased. Images were taken after
every cycle. Every eighth image over the first 112 cycles of sliding are given in Figure 2-5.
Small, loosely adhered PTFE islands formed quickly, and could be seen to move around the wear
track before agglomerating into the large platelets characteristic of the high wear of PTFE. The
wear platelets remained mobile throughout the entire test (1000 cycles), and appeared to grow to
some critical size before being peeled off and removed as debris under the frictional shear stress.
Again it was evident that the large platelets had not delaminated at their final size. Furthermore,
small wear platelets were visible in the optical after just four cycles, so it is unlikely that all
PTFE wear debris is formed by the propagation of subsurface cracks. Indeed, surface plasmon
resonance experiments by Krick et al. confirmed that PTFE can transfer to a countersurface after
a single cycle of sliding.33
Friction and Wear of PTFE Composites
Filler particles and fibers of widely varying size and composition have long been used to
abate the high wear of PTFE, usually at the cost of a slight to moderate increase in the friction
coefficient.6–12 Micrometer-sized metal oxide particles (ZrO2 and TiO2) and glass fibers were
shown to each reduce the wear rate of PTFE by one or two orders of magnitude while
20
maintaining a relatively low friction coefficient (µ~0.2). It was then supposed that that particles
and fibers of moderate size were more effective at reducing wear than smaller powders and hard
particles.9 At the time, wear reduction via fillers was considered from a largely mechanical
standpoint, and it was presumed that fillers reduced wear by 1) supporting a large portion of the
normal load34,35, 2) arresting the propagation of cracks6,36,37, 3) disrupting or strengthening the
banded structure of the polymer8,9, 4) preventing initial transfer of the polymer9, or 5) modifying
the adhesion of the transfer film.38
Though hard, micrometer sized particles are successful in preventing some wear of
PTFE, they tend to abrade the countersurface, preventing the formation of a smooth transfer film
and creating a roughened sliding interface that exacerbates wear of the polymer. A shift to
smaller particles was therefore desirable to prevent countersurface damage. In recent decades,
various nanoparticles were indeed shown to be capable of reducing the wear rate of PTFE7,39. A
study by Li in 2001 experimented with reducing wear using ZnO nanoparticles40, finding that the
wear was reduced by close to two orders of magnitude, and Chen had similar results using
carbon nanotubes.41 Burris and Sawyer later discovered that the inclusion of 5 wt. % of a
particular α-phase alumina particle reduced wear by an additional factor of 100, a thousand fold
improvement over neat PTFE42, while also maintaining a lower friction coefficient than many
similar composites. It is logical that even low loadings of nanoparticles should have a marked
effect on the properties of the polymer matrix due to their high number density and surface area
compared to particles of greater size. The term “ultralow wear” PTFE composites describes these
exceptional materials with wear rates less than 10-6 mm3/ (N·m).40,42–50 Various other fillers
loaded into PTFE at 5 wt. % were demonstrated by Sawyer’s tribology group to have vastly
21
differing effects on the wear rate and friction coefficient of the polymer against lapped 304
stainless steel, though none matched the performance of the α-alumina composite (Figure 2-6).
Tribofilms and Ultra-Low Wear
Unfilled PTFE has long been shown to create transfer films of variable thickness and
covered, via mechanisms seemingly contrary to those of other, higher friction polymers.5,30,31 In
1993 Blanchet showed that PTFE transferred to surfaces in sliding, regardless of the chemical
cleanliness of the surfaces or the presence of oxides.10 The low wear of polymer nanocomposites
is in general associated with the formation of thinner, more robust and uniform transfer films on
the countersurface, and running films on the polymer surface.40,42–44,46,51–55 The three regimes of
film formation, run-in, transition, and steady state, are distinguished by morphological changes
to the transfer and running films, and changes in wear rate.45 An example of this behavior is
illustrated in Figure 2-7: the wear rate of a PTFE/alpha alumina composite decreases as sliding
continues and a tribofilms form.
Transfer films for composites of PTFE and α- and Г- phase alumina were observed by
Burris to fill in negative features of the countersurface during the run-in period of higher wear to
provide a self-mated polymer interface that proceeds at a steady state wear rate independent of
the original average roughness of the nascent surface, and the transfer films were seen to increase
in thickness and discontinuity with increasing wear rate of the polymer composites.51 In a set of
experiments that slid transfer films of neat PTFE, Burris also showed that the alignment with
sliding direction of fibrils in the transfer films was conducive to low wear, and that sliding
perpendicular to the aligned direction led to rapid failure of the films.
Because the ultra-low wear behavior of PTFE/alumina composites is always
accompanied by the formation and evolution of visible, reddish-brown tribofilms on both sliding
surfaces, and because the run-in wear of the composites is always higher in the period before
22
film formation, it follows that the films are crucial to the abatement of wear of the polymer. A
comprehensive study of some of the mechanical properties of the tribofilms was conducted by
Krick et al. in 2014, who documented the color change, wear transitions, the hardness and
modulus of the running films in various stages of formation, and the depth of the effects of the
running film on hardness and modulus.47 They confirmed that the transition from high to low
wear occurred in conjunction with the formation of the tribofilms, that the hardness and modulus
of the running film increased as wear progressed, and that these increases were greatest at the
near surface, but were detectable at indentation depths up to around 250 nm.
Following Krick et al.’s analysis of the running films, Urueña et al. studied the wear
resistance of the PTFE/α-alumina transfer films themselves by performing microtribometer
experiments on transfer films in various stages of development.50 A single PTFE/ α-alumina pin
was used to run a “stripe test”56 that decreased stroke length after every mass measurement in
order to expose areas of transfer film in seven intervals from one thousand to one million cycles.
Microtribometer tests using a stainless steel ball of radius 1.5 mm measured friction
coefficients in each of the exposed areas. The transfer film patches were said to be worn through
after a number of microtribometer cycles corresponding to a crossover in friction coefficient
from low and consistent to high and erratic. Stylus profilometry was used to measure the initial
thickness of each of the exposed patches, and the wear volume used to calculate the wear rate of
the transfer film in each test was determined by Scanning Electron Microscopy (SEM). The
results of these tests indicated that the transfer films did become more robust as development
progressed, but that the films themselves were not necessarily a low wear material. The wear
rates of the films were several orders of magnitude higher than those of the composite pins. The
results of the study are summarized in Figure 2-8.
23
Soon after these results were published, a study by Ye expounded upon them by
measuring the adhesive and cohesive strength of PTFE/α-alumina transfer films in tension, as
well as the sensitivity of the wear rate of the films to the surface energy of the probes used in
testing.57 Using the tensile strip methods of Agrawal and Raj and the Tresca failure criterion,58–60
Ye showed that the ratio of adhesion to shear strength of the films increased with increasing
transfer film development, and added that a film could not be considered to be adhesive until the
ratio was greater than one – a condition which was quickly reached by the films after several tens
of meters of sliding. Furthermore, it was discovered that against surface probes of low surface
energy (PTFE, HDPE, other polymers) that the wear rates of the films were in fact
extraordinarily low (10-8 to 10-9 mm3/N∙m), comparable to the steady state wear rates of the
composites once the self-mated films condition is established. Ye’s results are compared to
Urueña’s in Figure 2-9.
The omnipresence of tribofilms on the sliding surfaces of ultra-low wearing PTFE/α-
alumina nanocomposites has led to a number of forays into the mechanical changes undergone. It
is evident that the films play a crucial role in maintaining the low wear behavior of the system.
Nanoparticle fillers especially seem to facilitate the long life of the films as they are too small to
cause abrasive damage or third body wear within the contact.
The Effects of Surface Roughness
While counterface asperity shape may in some cases govern the friction coefficient,
asperity size and radius of curvature may govern wear.36,61,62 Blanchet et al. showed that surface
roughness and filler concentration strongly affect the wear rate of PTFE/alumina composites, and
that it transitioned from K~10-7 mm3/(N·m) to K~10-4 mm3/(N·m) past a critical roughness
threshold for a given composite composition.63 For the PTFE/alumina composite considered in
the following experiments, the roughness threshold was reported to be ~ 5 µm. Franklin et al.
24
showed that surface roughness aligned perpendicular to the sliding direction increased the wear
of a POM-20% PTFE composite with increasing surface roughness,64 and Friedrich showed that
unfilled PEEK sample wore more with increasing roughness aligned parallel to the sliding
direction.65
Recently, studies of the ultralow wearing PTFE/α-alumina composite system have
investigated the characteristics of the tribofilms formed during sliding rather than of the polymer
composites themselves. Third bodies (debris, transfer films and running films) affect the friction
and wear of solids, as has been documented for many material systems.53,66–75 The following
study by Harris et al. in 2015 investigated the link between the ability of third bodies to
accumulate on a countersample and the transition to ultra-low wear sliding.76
Stainless steel countersamples were prepared with aligned grooves defined by the angle
of the grooves relative to sliding direction, referred to as the “roughness angle”, θ. The roughness
angle varied from 0° (parallel to the sliding direction) to 90° (perpendicular to the sliding
direction). Directional roughness was prescribed using 600 grit silicon carbide sandpaper
adhered to a thick sheet of polydimethylsiloxane (PDMS), which was mounted to the stage of a
linear reciprocating tribometer.77 The stationary countersample was mounted to a jig that allowed
adjustment relative to the sliding direction in 5° increments. The aligned countersample was run
at 250 N for 200 reciprocating cycles, resulting in an average roughness (Ra) of 140 nm with a
standard deviation of 28 nm across all countersamples and an average groove width of ~60 µm.
No sample exceeded 220 nm Ra. A lapped finish (Ra ~150 nm) and a highly polished mirror
finish (Ra ~ 20 nm) were used as controls. All countersamples were washed with soap and water,
sonicated in methanol, and allowed to dry in laboratory air prior to experiments. Analysis of the
25
surface features of the metal countersamples was performed using a scanning white light
interferometer. Scan data were averaged over ten locations on each countersurface.
The wear rate of the PTFE/α-alumina composite was evaluated in sliding against six
different roughness angle surfaces, the lapped surface, and the polished surface. The total wear
rates at each mass measurement for each of the countersurface treatments is plotted in Figure 2-
10. All of the samples experienced transient wear processes, including a transition from an
initially higher wear rate to a lower final wear rate (run-in). The run-in transition in wear rate
was more dramatic for materials that reached ultralow wear rates (θ = 90°) than for those that
had high total wear rates (θ = 0°). Generally, the total wear rate and the steady state (final) wear
rate decreases as the roughness angle (θ) increases, producing the highest wear rate when sliding
parallel the roughness features and lowest wear when sliding perpendicular to the ridges. At
roughness angle 0°, the PTFE/α-alumina composite had a high steady state wear rate (9.6 × 10-6
mm3/(N·m)), and never reached ultra-low wear over the duration of the test. The steady state
wear rate for PTFE/α-alumina on the 90° countersample, however, was extremely low (K ~ 3.1 ×
10-8 mm3/(N·m)), lower than the lapped and the polished controls. The wear rate in the parallel
direction was around 300 times greater than the wear rate in the perpendicular direction.
Optical images and scanning electron micrographs of the wear tracks revealed transfer
film morphology and coverage. The 0° countersample developed a dark transfer film at the ends
of its stroke, but the center appeared largely bare. Electron micrographs of the zero degree
transfer film showed flaky patches of wear debris or transfer that appeared to be resting on top of
the roughness ridges. The 90° test formed a fairly uniform transfer film that covered most of the
sliding surface. The 45° film had characteristics of both the 90° and 0° countersamples and did
not develops darker orange regions or patches.
26
PTFE/alumina transfer films adhere strongly to steel countersurfaces. In the context of
debris mobility, the longer a third body (debris or tribofilm) remains in sliding contact, the
higher the probability of transfer film adhesion. By designing surfaces to retain or eject debris,
we can control the amount of mechanical work done on debris particles for a given tribological
system. On countersamples with roughness features aligned in the sliding direction, debris is
easily shed and wear rates remain high. When roughness features are perpendicular to the sliding
direction, debris may be trapped. This allows for increased sliding cycles over a given third
body, resulting in increased opportunities for adhesion and ultimately ultralow wear rates. On the
90° surface, retention of debris and formation of a stable transfer film appeared to produce steady
state wear rates significantly lower than on randomly rough surfaces of comparable Ra. Of
course, while the 90° surface had the largest capacity to retain debris, it required a greater run-in
wear volume to fill in the negative surface features before ultra-low wear was reached. The
results of this study suggested that fillers which abrade and damage to the metal surface create a
surface that promotes third body rejection and thus does not facilitate ultra-low wear.
Debris mobility may be modeled in several ways (Figure 2-11). Godet explored factors
that could affect the recirculation or ejection of third bodies from contact.74 Contact shape,
roughness and third body rheology effects were proposed based on this work78–80. Thereafter,
optical and chemical in situ studies by Wahl, Singer, and Chromik looked into the role of third
bodies.81–84 The study described in this chapter was specifically designed to control the mobility
of debris and its ejection from the wear track, and provided a simple model for debris mobility
based on the number of obstacles in the path of the debris. For the surfaces with aligned
roughness, these obstacles were the ridges. In a given sliding cycle, a mobile debris particle
crosses a number of ridges proportional to sin . A simple model considered the wear resistance
27
to increase with the number of ridges crossed during the stroke – the wear rate then decreased
with theta as1 sin .
A slightly more complex approach considered the local forces on debris and third bodies.
The local friction force is composed of a reaction force against the ridges and a debris migration
force along each ridge. With increasing θ, the migration force along the ridge decreases and the
reaction force increases. The ratio of the migration force to the reaction force resembles a local
friction coefficient. Wear rate appears to scale with the mobility factor, meaning that higher
mobility of debris leads to a higher wear rate. Real behavior of these wear systems depends on
many factors aside from debris mobility. Debris and transfer film retention likely depends on a
critical shear strength of the interface between countersurface and third bodies, dependent on the
size of surface features and roughness angles. At the critical shear stress, material could migrate
out of the grooves, increasing wear, resulting in sparse transfer films as seen on the higher
wearing samples.
The Effects of α-Alumina Particle Morphology
Granular PTFE 7C with 1-5 vol. % alumina particles added have been shown to reduce
the wear of the polymer to 1x10-7 mm3/(Nm) or lower [5-24].6,7,9,11,35,36,42,43,85–91 This was first
described by experiments using three different alumina fillers of manufacturer designated
particle sizes of “44 nm”, “80 nm” and “500 nm” [17, 19-24].35,42,43,51,63,89,90 The “80 nm”
alumina achieved steady state wear rates near 8x10-7 mm3/(N∙m) - lower than the other
composites by a factor of 100. No correlation between wear and particle size was observed.
Later, it was determined that the “44 nm” particles were spherical delta-gamma phase alumina,
and the others were alpha-phase.51 McElwain et al. and Blanchet et al. reported that the wear rate
of PTFE with 2.9 vol. % (~ 5 wt. %) alpha-phase alumina nanoparticles was two orders of
28
magnitude lower than with “microparticles”,35,63 and were also relying on the supplier designated
particle size.
One hypothesis regarding the effect of the addition of inorganic particles is that they
somehow interrupt the subsurface crack propagation of PTFE.35 Another hypothesis states that
wear reducing additives alter the crystalline structure of the PTFE into something tougher,
although so far there is no definitive evidence for such a mechanism.6,10 It was also proposed that
the banded structure of PTFE was disrupted, accounting changes in wear performance.8,9 In the
following summarized study,92 five alumina powders were used as fillers at 5 wt. % in Teflon®
PTFE 7C:
Alumina A-Alfa Aesar α-phase alumina powder with a supplier-specified approximate
particle size of 0.35 to 0.49 µm (Stock #42573, 99.95%)
Alumina B - Almatis calcined α -phase alumina powder (grade A 16 SG, 99.8%) with a
supplier-specified typical d50 particle size of 0.5 µm
Alumina C - Alfa Aesar α -phase alumina powder (Stock #44652, 99%) with a supplier-
specified approximate particle size of 60 nm, and confirmed to lead to ultra-low wear42
Alumina D - Alfa Aesar α -phase alumina powder (Stock #44653, 99%) with a supplier-
specified approximate particle size of 27-43 nm, also previously reported to produce
ultralow wear PTFE composites35,42,43,46–51,63,89,90,93–96
Alumina E - Nanostructured and Amorphous Materials, Inc. alpha-phase alumina powder
(Stock # 1015WW, 99.5%, according to the manufacturer, mostly alpha-phase with 5-
10% gamma phase) with a supplier-specified approximate particle size of 27-43 nm,
another filler confirmed to result in ultra-low wear of PTFE35,42,63,91
The wear rates achieved by PTFE samples filled with each of these types of alumina, as
well as previously observed values thereof are given in Figure 2-12. The results for PTFE filled
with Alumina A, C, and D agreed with Blanchet’s study. The composite filled with Alumina B
behaved unexpectedly given its supplier designated particle size was equal to that of Alumina A.
The publications mentioned before observed a large transition in wear rate appearing
above about 100 nm in particle size as reported by the vendor. Particle size distribution
29
measurements were conducted by Static Light Scattering (SLS), by the Brunauer, Emmett and
Teller method (BET), Scanning Electron Microscopy (SEM). The particle size measurements
supplied did not match the particle sizes measured independently. Results are detailed in Table
2-2 and are listed in increasing order by median particle size as determined by SLS. It was
assumed that the particle sizes provided by the vendors were measured using the BET method,
which assumes that the particles are spherical, dense, and monodisperse. Many of the alpha-
phase alumina particles used in these studies were discovered to be irregular in shape, and porous
by SLS and SEM, as demonstrated in Figure 2-13. For this reason, the assumptions for BET
particle size calculation would have been invalid.
In conjunction, SLS, BET surface area measurement, and microscopic examination of the
particles suggested that surface area and porosity of the particles increases from Alumina A to E.
It is commonly assumed that the large surface area of nanoparticles is one of the main sources of
wear reduction. However, in this case the internal porosity is not accessible by the PTFE
polymer chains in the bulk of the part. Furthermore, the transfer film topography results seemed
to contradict the SLS particle size measurements: The smallest particles by SLS, Alumina A,
resulted in abrasion of the steel countersample and wear scars significantly deeper than the
measured particle size. Alumina E had the largest particle size by static light scattering, but its
composite had the lowest wear rate and formed films that were significantly thinner than the
alumina particle size. Even so, a recent XPS study reported an alumina concentration between 1
and 5 at. % at the transfer film surface.48 It was thus hypothesized by the authors that the large
particles of Alumina E were friable and prone to breaking up due to mechanical stresses at the
sliding interface, preventing damage to the metal and the transfer film. Indeed, small, hard,
30
inorganic fillers in low wearing PTFE have been shown to accumulate at the polymer wear
surface and increase the abrasive wear of the steel surface.97
X-ray microtomography was used to investigate the near surface region of the
PTFE/Alumina E polymer composite after wear testing. Images acquired showed evidence that
alumina particles smaller than the particle size measured by SLS accumulated at the wear
interface, and also revealed that micrometer-sized particles remained in the bulk, as shown in
Figure 2-14. This result was strongly in support of the hypothesis that large friable particles in
the bulk broke up during sliding and that the resulting smaller particles built up at the interface.
The X-Ray Microtomography results were further confirmed by TEM micrographs,
shown in Figure 2-15. The TEM images are consistent with the X-ray microtomography results
in that both techniques give evidence of occasional well-distributed microscale filler particle, and
in that a distinctly different region within the top few micrometers of the worn polymer surface
contains fine scale features that Energy Dispersive Spectroscopy (EDS) in both the TEM and
SEM indicated are alumina.
This study posited that the best alumina particles for ultralow wear PTFE composites are
porous and micrometers in size, rather than dense nanoscale spheres, and that supplier-
designated particle sizes, likely based on BET surface area measurements, were unreliable. The
micrometer sized mesoporous-like filler materials are comparable to nanoparticles in surface
area, and have nanometer-scaled features. In general, wear rate of the composite decreased with
increasing BET surface area to average particle size ratios. It was agreed that multi-scale fillers
are ideal for such large wear reductions. The porosity of the alumina particles rendered them
mechanically friable, the micrometer-scale of the particles likely reduced subsurface
31
delamination of PTFE composites in the bulk to near surface region,47–49 and the nanometer-scale
alumina particle fragments were still assumed to stabilize and reinforce tribofilms.
Tribochemistry
Several studies have correlated chemical degradation of PTFE with changes in
mechanical properties98–101 and color.102 Chain scission and defluorination due to degradation
were suspected to lead to the presence of CF3, carbon-carbon double bonds, branched carbon
structures, COF and carboxylic acids groups, and crosslinks.103–107 XPS analysis by Scott Perry’s
group at the University of Florida found α-alumina concentrations in transfer films comparable
to the composites, and higher oxygen content than predicted which suggested oxidation of PTFE
occurred during sliding. Furthermore, a correlation of friction coefficient with oxygen content of
the transfer film was observed, though it could not be said which caused the other, or that
causation was even present. Based on this correlation, it was predicted that oxygen content was
lower at the ends of the wear tracks, where friction is lowest.108 Burris then suggested that
enough energy is absorbed in the thin layers during sliding to initiate low probability chemical
events. Further XPS analysis by Professor Perry’s group at the University of Florida indicated a
chemical change within the tribofilms as evidenced by the formation of a new peak at (288 eV)
and was expounded upon by Krick (Figure 2-16).43
Over several studies, Blanchet showed that e-beam irradiated PTFE composites
demonstrated increased wear resistance.109,110 Burris also showed that a chemically etched
unfilled PTFE sample was 100 times more wear resistant and had 10% higher friction than the
untreated polymer before the etched surface wore through. Furthermore, the etched polymer
displayed a peak at 288 eV corresponding to the peak seen by earlier XPS studies of transfer
films.108 This compelling evidence for tribochemical contribution to the ultra-low wear of a 5 wt.
% PTFE/α-alumina led Krick et al. to conduct a series of wear experiments in varied
32
environment to elucidate the effect of oxygen and water on the system.43 The composites wore
more with decreasing relative humidity, and performed slightly better in environments
containing oxygen than in those containing pure nitrogen. However, some oxygen was present
during all test below their detectable limit of 100 ppm. The wear of the composite was highest
submerged in water. They concluded that water is necessary for the ultra-low wear behavior of
PTFE/α-alumina composites, and that oxygen is likely necessary. The reduction in wear rate in
the presence of water suggested that a secondary wear mechanism was reduced by the formation
of thin, protective tribofilms. Tribofilms associated with ultra-low wear rates did not form in the
absence of water. Soon thereafter, Pitenis et al. conducted a similar suite of tests on the same
composite, varying the level of vacuum of the test environment.46 The wear rate of the composite
was highest in the lowest pressure system (~10-6 Torr), where a visible transfer film did not form,
and highest at 5 Torr, where it was suggested that the water to oxygen ratio was most favorable
for film formation. At the lowest pressure tested, oxygen was likely mostly absent, leaving a
mostly water environment, which was insufficient to promote film formation or ultra-low wear.
Figure 2-17 compiles Krick’s environmental data and Pitenis’ vacuum data, demonstrating the
dependence of wear behavior of the PTFE/α-alumina composite.
The importance of ambient species, in particular oxygen and water, to the longevity of
the composite suggested that the chemical mechanism for tribofilm formation must likewise
depend on their presence. The identification of this mechanism was of interest in order to fully
comprehend the role of the tribofilms in the ultra-low wear system created by the PTFE/α-
alumina composite, and would introduce the possibility of advancements in composite
production that could lead to increased part life by further facilitating the reactions involved.
33
Figure 2-1. The polymerization reaction of PTFE. Chlorodifluoromethane yields
tetrafluoroethylene and HCl after pyrolysis. Free radical polymerization of
tetrafluoroethylene creates the PTFE molecule.
Figure 2-2. The PTFE molecule is helical in structure. The twist of the helix is dependent on
temperature (two of its phase transitions occur at 19 and 30 °C).
34
Figure 2-3. PTFE crystallizes to form lamellae rather than spherulites. An illustration of the
banded structure of PTFE shows large lamellae running left to right with smaller
lamellae lying perpendicular between them. PTFE molecules within the crystalline
lamellae lie parallel to the striations.26
Figure 2-4. A suite of wear experiments with neat PTFE in various conditions against 304
stainless steel illustrates the speed dependent wear transition and the wear behavior of
the self-mated and submerged polymer.
Table 2-1. Tensile and compressive yield strengths of unfilled PTFE as reported by *Rae111 and
**DuPont112.
35
Figure 2-5. Time lapse images of the wear of unfilled PTFE illustrate the transfer, agglomeration
and growth of PTFE wear debris islands. Nascent steel is shown at the top left, cycle
1 in the second image, and cycle 8 in the third. Images proceed in 8 cycle intervals
thereafter, concluding with cycle 112 (bottom right). A scratch on the steel surface is
visible in the upper left corner of each image, and proves that each image was taken
in the same location as the wear debris evolved. Photos taken by author in Spring
2014.
36
Figure 2-6. The wear rates of various PTFE composites with 5 wt. % filler added are plotted
versus their average friction coefficients. The engineering ideal lies at the lower limit
of wear and friction coefficient.
37
Figure 2-7. A representative plot of the wear rate of a 5 wt. % α-alumina PTFE composite
illustrates the run-in, transition, and steady state behavior of the polymer in sliding.
Once very low wear rates are reached, they are maintained for hundreds of thousands
of sliding cycles.
Figure 2-8. A summary of Urueña’s transfer film wear study results that suggested the transfer
films formed by PTFE/α-alumina composites are not intrinsically low wear materials,
but are instead an integral part of a self-mated low wear system. This plot is a
reproduction from Urueña’s 2015 study.50
38
Figure 2-9. Urueña’s transfer film wear data measured using a stainless steel pin50 is compared to
Ye’s similar study which used a lower surface energy High Density Polyethylene
(HDPE) probe to wear away the transfer films.57 Wear rates of the composite pins
used to create the transfer films and the wear rates of the transfer films themselves are
superimposed. Ye’s study found that the PTFE/α-alumina transfer films were
exceptionally resistant to wear against low surface energy polymer probes.
Figure 2-10. A plot of the total wear rates of a PTFE/α-alumina composite against an array of
surfaces with prescribed angular roughness is reproduced from Harris’ et al.76 The
total wear rates decreased as the surface roughness angle approached 90°
(perpendicular to the sliding direction).
39
Figure 2-11. In a reproduced plot from Harris et al.,76 the total wear rate of the PTFE/α-alumina
composite is plotted against the roughness angle of the countersurface used in each
test. The two models proposed are superimposed for comparison.
Figure 2-12. The wear rates of various PTFE/alumina composites are plotted versus the supplier
designated particle size of the fillers as described by Krick92 and Blanchet.63 Krick’s
study goes on to show that the particle sizes assigned by the suppliers were largely
inaccurate.
40
Table 2-2. A summary of the particle size results, which vary by method, and the wear rate of the
PTFE composite containing them. The supplier designated particle sizes, BET
particle sizes (assuming dense spherical particles), and SLS particle sizes of each of
the types of alumina studied, and the wear rates and friction coefficients of the
composites molded at 5 wt. % of each of the types of alumina. A large discrepancy
was observed between the particle sizes estimated by the BET method and by SLS.
The lowest wear rate was achieved by the composite filled with alumina E, which had
the smallest particle size according to the supplier and according to BET, but by far
the largest particle size as measured using SLS. This table is comprised of data from
Krick’s 2015 particle size study.92
41
Figure 2-13. SEM micrographs and SLS data demonstrate how BET data provided by particle
suppliers may not be an accurate representation of true particle size. (a) SEM
Micrographs of alumina A showed that the particles were close in size to the
supplier’s designation. (b) SEM micrographs of alumina E illustrate the comparative
dimension of the particles to the much smaller size assigned by the supplier. (c) SLS
data for each of the alumina types confirms that the particle size distributions are in
some cases very different than the designations. Images and data are adapted from
Krick’s 2015 particle size study.92
42
Figure 2-14. X-Ray Microtomography adapted from Krick et al.’s 2015 particle size study92
shows the very small alumina particles that built up in the running film during sliding,
as compared to the more dispersed particles in the bulk.
43
Figure 2-15. TEM micrographs of the running films confirm that nanoscale alumina particles are
present at the sliding surface of the composite. The large, friable alumina E particles
break up at the interface and support the running surface of the polymer. Micrographs
courtesy of DuPont.
44
Figure 2-16. XPS spectra and optical images adapted from Krick43,47 indicate chemical changes
at the sliding interface. (a) XPS analysis of the bulk polymer and the tribofilms
formed during sliding indicate the creation of new chemical species during sliding.
(b) Evidence of chemical alteration of the surface is seen in the color change
undergone by the running surface of the composite.
Figure 2-17. The wear rate of a PTFE/α-alumina composite is plotted vs. relative humidity, and
the wear rate of another PTFE/α-alumina composite is plotted vs. pressure. Humidity
studies used alumina C as described above and are adapted from Krick.43 Vacuum
studies used alumina E as described above, and data is adapted from Pitenis.46
45
CHAPTER 3
METHODS AND EXPERIMENTATION
Materials and Sample Preparation
A PTFE/α-alumina composite was made using DuPont Teflon® PTFE 7C resin as the
matrix. The polymer was filled with 5 wt. % α-phase alumina (Nanostructured & Amorphous
Materials Inc., Stock#: 1015WW,45,46,48,96 (alumina E from the particle morphology study
described previously). The dry PTFE powder was combined with the alumina filler and
submerged in extra dry isopropanol to a total volume of 200 mL. The mixture was sonicated
using an ultrasonic horn (Branson Digital Sonicator 450 with a titanium tip at 40% amplitude for
three, one minute long sessions with 45 seconds rest between each). The dispersion was allowed
to dry completely in a fume hood for seven days. The dried powder mixture was then
compressed in a hydraulic press to approximately 100 MPa in a 440C stainless steel cylindrical
mold before being sintered in an oven. The heating process ramped upwards at 2 °C/min to 380
°C, where it was held for four hours and then cooled to room temperature. The sample was then
machined into a pin (6.3 x 6.3 x 12.7 mm). The square running faces of the pin were polished
with 800 grit silicon carbide sandpaper to an approximate average roughness (Ra) of 100 nm,
measured using a scanning white light interferometer. Finally, the sample was sonicated for 30
minutes in methanol and allowed to dry completely before testing.
The countersample used was a flat, rectangular (115 x 25 x 3.7 mm) plate of 304 stainless
steel finished by lapping (Ra ~ 150 nm) by Alabama Specialty Products Inc., the standard
running surface used in previous experiments with the PTFE/alumina composites.43,46,48,51,89,113
The countersample was cleaned with soap and water, rinsed with methanol, and allowed to dry
prior to experiments.
46
Tribometer and Wear Test Design
Wear testing was performed using a linear reciprocating tribometer (Figure 3-1) with flat-
on-flat sample geometry as described by Schmitz.77 A six-axis force transducer measured the
normal and frictional forces continuously throughout each test. The running face of the polymer
sample was loaded against the stainless steel countersample at 250 N, or a nominal contact
pressure of around 6.3 MPa (around 50% of the compressive yield strength of neat PTFE). The
contact pressure can be approximated using this nominal calculation because the polymer sample
wears into conformity with the flat countersurface within the first thousand cycles. Pre-existing
asperities are under much higher local contact pressures and wear away quickly. The steel
countersample was mounted to a linear ball-screw stage which reciprocated at 50.8 mm/s.
A “stripe test” 48,50,56 was performed in which the reciprocating stroke length decreased as
the number of sliding cycles increased at predetermined cycle intervals (Figure 3-2). The initial
reciprocating stroke length was 88.9 mm and was reduced by 10.2 mm after each experiment.
The final test was 27.9 mm long. This pattern of transfer film formation isolated and preserved
areas of the polymeric film over the various stages of development. These exposed areas allowed
physical and chemical analyses of the steps involved in transfer film evolution.
Wear Rate Calculation, Friction Measurements and Uncertainty
The wear rate, K, of each polymer sample is described as the volume lost per unit force
and distance and is recorded in units of mm3/N∙m (Equation 3-1), as described by Archard114.
The density, ρ, of each pin is calculated prior to testing using its measured dimensions, L, W1 and
W2, measured using clean calipers, and its initial mass, mi, measured using a Mettler Toledo scale
(Equation 3-2). Volume loss, ΔV, is then the mass loss, Δm, divided by the density (Equation 3-
3). Wear rates were calculated both as total volume lost per total distance slid (total wear rate),
and as volume lost per test per distance slid per test (test wear rate). The total wear rate includes
47
the run-in period of relatively high wear before the onset of ultra-low wear after stabilization of
the tribofilms and is calculated by considering the volume lost and distance traveled over all
previous tests. The test wear rate differentiates between early, higher wear rates, and the ultra-
low minimums reached after many sliding cycles by considering only the volume lost and
distance traveled during a single test. The steady state wear rate is defined as the wear rate
reached and maintained after the run-in period.
n
V
F D
(3-1)
1
i
i i
mLWW
(3-2)
mV
(3-3)
Volume calculations using vertical displacement measurements are generally less
accurate than those based on mass loss due to the possibility of creep of the polymer, or the
formation of transfer films of non-negligible thickness. The normal load is monitored throughout
the duration of each experiment. The mean value of the force, Fn, is then used to calculate the
wear rate of the polymer over the period for which it is averaged. The lateral displacement, d, of
the reciprocating stage was similarly recorded throughout each test and used to calculate the total
distance traveled, D.
The error in each measured quantity is as follows: u(m) = 0.02 mg, u(L) = 0.02 mm,
u(Fn) = 0.2 N, and u(d) = 0.04 mm/mm. The error in the total distance traveled is
( ) ( )u D D u d , and the error in the mass loss is ( ) 2( ( ))u m u m , according to the law of
propagation of uncertainty. Further propagation of these errors leads to expressions for the
uncertainty in the measured volume loss, V, as given by (Equation 3-4), as well as for the product
48
of the average normal load and total distance traveled (the denominator of the expression for the
wear rate, K), given by Equation 3-5.
2
2 2
2 2i i
2
i i
V ΔmVu(ΔV) = u(Δm) + u(m)
m m
2
2 2 2 2
i 1 i 2 1 2
i
Δm+ u(L) (LW ) +(LW ) +(W W )
m
(3-4)
2 2 22 2
n n nu(F D) = u(F ) (D) + u(D) (F ) (3-5)
The uncertainty in ΔV and in Fn∙D is used to perform 1000 Monte Carlo simulations of
each wear test. In each simulation, possible values for each parameter are generated randomly
within plus or minus three times the value of the calculated uncertainty. The wear rate is then
calculated for each of the simulations, and an average and a standard deviation thereof is
obtained for each experiment.
Friction coefficients were calculated as the quotient of the average frictional force (measured
continuously throughout each test, and averaged at the end of each test, and the average normal
force77. The associated uncertainties in the friction coefficient and wear rate calculations are
described in greater detail in Schmitz et al.77,115
Stylus Profilometry
Stylus profilometry was used to map transfer film topography (KLA Tencor P-16 with a
2 µm radius probe and 5 mg load). Fifty 12 mm line scans were acquired to map 2 mm x 12 mm
sections in each exposed area of transfer film development. The nearby bare stainless steel
surface was used as a zero reference to shift the profiles according to a common baseline. The
average thickness of each exposed section of the film was calculated as the difference between
49
the nascent metal zero reference and the center of each film section scanned in the 2 x 12 mm
sections described above.
Small Molecules Experiments
Pressed polymer disks were created at DuPont for IR comparison to the tribofilms formed
during the wear test. The same α-alumina used to create the PTFE composite pin was pre-dried
and heated in a flask in an oil bath at 150 °C and 20 milliTorr for five hours. The flask was
backfilled with dry nitrogen and transferred into a nitrogen-atmosphere drybox. In a separate
flask in the drybox, 10 grams of the alumina were combined with tridecafluoroheptanoic acid
(TCI America, 98.0%, 0.23 g), and dry isopropyl alcohol (Fisher Scientific, 30 mL). The mixture
was stirred in the drybox, transferred to a fume hood and reduced in vacuo (150 Torr) in a water
bath at 28 °C in order to minimize vaporization of the acid. The acid/alumina mixture was added
at 5 wt. % to 10.0 grams of pre-dried PTFE and mixed for 18 hours on a roller mill. The polymer
mixture was cold-pressed under 2.5-3 tons pressure at ambient temperature into 13 mm circular
disks using approximately 100 mg of material in a hydraulic press. The resulting PTFE films
were around 355 µm thick.
X-Ray Photoelectron Spectroscopy
X-ray photoelectron spectroscopy (XPS) was used to analyze each area of exposed
transfer film. XPS was performed at DuPont using a Physical Electronics Quantera Scanning
ESCA Microprobe with a focused (100 um) monochromatic Al K-alpha X-ray (1486.6 eV) beam
at 18 kV and 100 W. The electron energy analyzer was operated in constant energy mode with a
pass energy of 55 and a 0.2 eV step size between points for high resolution spectra (energy
resolution of the system was approximately 0.84 eV using the Ag3d5 peak). Atomic percent
concentrations were normalized to 100%. A dual electron and argon ion beam system was used
for charge compensation.
50
Infrared Spectroscopy
The infrared spectra used to analyze the metal surfaces and tribofilms were obtained by
collaborators at DuPont using a Nicolet 6700 FT-IR spectrometer with a Continuum Microscope
(Thermo Fisher Scientific) in reflectance mode. The area of analysis was a 100 µm square.
Background spectra were collected away from the fluoropolymer wear track in a clean area of
the steel countersample. Transfer film spectra were obtained by reflectance at three spots along
the centerline at the midpoint of each of the seven exposed transfer film areas. This triplicate
analysis yielded consistent results for each of the areas inspected.
The cumulative wear debris at the end of the one million cycle region was analyzed using
attenuated total reflectance infrared (ATR-IR) on a Golden Gate (Specac) horizontal diamond
ATR unit. Background spectra were collected with a clean diamond surface. Transfer film
residue spectra were collected after the film areas were analyzed and removed from the diamond
surface, but before the crystal was cleaned with ethanol. No pressure was applied from the
clamping device for residue spectra collection. Spectra were corrected for the wavelength
dependence of penetration depth to closely resemble transmission spectra.
Transmission spectra of the pressed polymer films (PTFE 7C/α-alumina/C6F13COOH and
PTFE 7C /α-alumina) were obtained at DuPont using a Nicolet Magna 560 FT-IR spectrometer
(Thermo Fisher Scientific). A background spectrum was collected using an empty film card of
the same type used to mount the films. Acquired spectra were converted to absorbance spectra
for comparison. The C-F overtone peak near 2365 cm-1 was used as a guide to detect changes in
this region, after spectral subtraction of a PTFE control film. Due to the intense C-F stretch
region of these thick disk samples, the region below ~1320 cm-1 was distorted.
51
Etched PTFE Tests
Unfilled DuPont Teflon® PTFE 7C and a PTFE 7C/α-alumina composite filled with 5
wt. % α-phase alumina (Nanostructured & Amorphous Materials Inc., Stock#: 1015WW,45,46,48,96
(alumina E from the particle morphology study described previously) samples were prepared as
described above. These samples were machined into pins approximately 6.3 x 6.3 x 12 mm in
dimension and polished using 800 grit SiC sandpaper. One each of the filled and unfilled
polymer samples were then etched using Fluoroetch® (Acton Co.), which modifies and oxidizes
the polymer surface so that it contains carboxylic acids116, and caused the exposed surfaces to
turn a dark, chocolate brown. The samples were not polished, in order to preserve the surface
treatment on the running faces of the pins. Before testing, the pins were sonicated in methanol
for 30 minutes and allowed to dry overnight. A stripe test48,50,56 was performed using the etched
surfaces of the PTFE and the α-alumina composite (Figure 3-3). The running films of each of
these samples were removed, and a stripe test was then performed using each of the inner
surfaces of the unfilled PTFE and the α-alumina composite. The PTFE/α-alumina composite ran
for 1k, 10k, and 100k cycles, and the unfilled PTFE samples ran for 10, 100 and 1k cycles. All
samples slid at 50.8 mm/s under a 250 N normal load. Wear track lengths for all samples were 41
mm, 31 mm and 19 mm. Wear rates were calculated as described above.
52
Figure 3-1. Schematic of linear reciprocating tribometer with a flat-on-flat pin configuration, and
a six-axis load cell and several LVDTs for data acquisition.
Figure 3-2. A schematic of the stripe test run using a PTFE/5 wt. % α-alumina composite pin
against 304 stainless steel with a lapped finish. After each test, the pin was massed,
and the track length of the following test was shortened in order to expose a series of
transfer films in various stages of development. Schematic is adapted from Harris
2015.49
53
Figure 3-3. Stripe tests were performed on the etched surfaces of unfilled PTFE and a PTFE/α-
alumina composite. The running films were removed, and two further stripe tests
were performed on the inner, un-modified polymer surfaces.
54
CHAPTER 4
RESULTS AND DISCUSSION
Friction and Wear
The wear and friction performance of the PTFE/α-alumina composite was consistent with
previous studies of similar composite materials.43,44,46–48,63,93 Over the first 10k cycles, a run-in
period of moderately high wear was observed followed by a decrease in wear rate over the next
100k cycles to less than 10-6 mm3/(N·m) (Figure 4-1a). The coefficient of friction also decreased
slightly over the course of the experiment (Figure 4-1b), remaining near µ~0.19.
Stylus Profilometry
Figure 4-2 illustrates the dramatic changes undergone by the transfer film as it develops.
The measured root mean squared roughness (Rrms) of the stainless steel substrate was between
140 and 190 nm. Under 10k cycles, the average transfer film height was on the order of the
roughness of the surface. The average height of the transfer film after 100k cycles was
approximately 200 nm, with a local Rrms around 300 nm. The 1 million cycle film rose
approximately 1,000 nm above the baseline, with a local Rrms around 610 nm. Film profiles
shown in Figure 4-2 display the changing character of the roughness throughout the development
of the film. The most significant topographical changes in the transfer film occurred between
100k and 1 million cycles as the PTFE composite filled in the countersample grooves caused by
lapping and polishing.
X-Ray Photoelectron Spectroscopy
XPS (Figure 4-3) spectra provided a chemical time lapse of the transfer film formation.
Fluoropolymer transfer was detected after one cycle of sliding as evidenced by a peak around
292 eV. First cycle transfer of unfilled PTFE has previously been measured by surface plasmon
resonance33 and Auger117 techniques. In lab air, stainless steel countersurfaces are always coated
55
with adventitious hydrocarbon contaminants and mixed oxides. A sputter depth profile was
performed to determine the approximate thickness of the contaminant layer and to measure the
composition of the bulk metal. After less than five minutes of sputtering, carbon and oxygen
compositions dropped below 10%. The contamination thickness was therefore likely less than 10
nm. At the end of the profile, the composition was 71% iron, 16% chromium, and 5% nickel -
consistent with bulk stainless steel given measurement uncertainties. In the XPS spectra of the
carbon region (shown in Figure 4-3a), the CF2 peak near 292 eV grew and the C-C/ C-H peak
near 285 eV shrank with as sliding cycles increased. This suggested that fluoropolymer coverage
increased within the transfer film region as sliding progressed. It also suggests that the surface is
cleaned of contaminants by the sliding composite, possibly due to the hard alumina filler
particles. This is further supported by observations of metals and oxides on the wear surface of
the polymer pin.
Figure 4-3b illustrates the changes in atomic concentration data from XPS experiments
on the transfer film as a function of sliding cycles. Transfer in the first cycle was succeeded by a
fluctuating fluoropolymer signal as debris was generated and ejected from the wear track as
wear. The concentrations of metal species from the steel (Fe, Cr) gradually decrease as the
transfer film thickens and the fluorine concentration rises. Alumina is first detected at 100 cycles
and reaches concentrations comparable to those in the bulk composite in the next few thousand
cycles. The wear rate drops to a steady state between 10k and 100k cycles as the protective
tribofilms grow. The small CF2 peak observed by XPS for the “0 cycle” experiment is probably
due to contamination from fluorocarbon wear debris on the surface of the steel counter sample
outside of the worn region.
56
Infrared Spectroscopy
Fourier transform infrared (FTIR) spectroscopy performed within the exposed areas of
transfer film formed during the 1, 100k and 1M cycle tests, and on the cumulative wear debris
allowed for detailed chemical analysis of the evolution of the tribofilms and therefore of the wear
system. The spectral analysis is summarized in Figure 4-4. Spectra collected within the one cycle
transfer film area verified that the fluoropolymer had been immediately transferred to the metal.
However, the spectra also revealed an unusual set of peaks in the C-F region consisting of the
typical PTFE peaks at 1203 and 1149 cm-1 along with a new peak at 1253 cm-1 (Figure 4-4a).
This additional peak was originally derived by Moynihan in 1959 from first principles
calculations44 but has rarely been observed experimentally. It has previously been observed in
analysis of a PTFE powder film,118 for a PTFE sample run against a polyethylene film,119 and for
a PTFE sample run against 304 stainless steel.120 Lauer attributed some variations in intensity of
the peak near 1250 cm-1 to a stretching mode of the polymer molecule. He noted that the
intensity varied with the alignment of helical PTFE chains relative to the detector when using a
polarizer.120 This suggested that aligned PTFE chains were transferred to steel after a single
sliding pass, which agrees with surface plasmon resonance results from Krick et al.33, XPS
results from Uçar121, and the current study.
The extremely high molecular weight (~ tens of millions g/mol27) of PTFE means that the
granular polymer is very unlikely to transfer an entire chain to the metal surface due to the high
entanglement of the chains. Thus, for transfer to occur on a small scale prior to large scale
delamination, scission of the C-C backbone must occur. A mathematical argument can be made
to support the bond breaking and transfer of PTFE chains in the first cycle of sliding by
considering the intramolecular forces between the polymer chains and the metal countersurface.
PTFE fibrils have been shown to quickly and preferentially align in the direction of
57
sliding.8,13,122–125 Thus, the model proposed is based on a balance between the sum of all the van
der Waals attractions between an aligned PTFE fibril and the metal surface, and the force
required to break an aligned PTFE fibril (Figure 4-5). Values for the model were chosen
conservatively so as to make it as unlikely as possible that fibrils would break rather than slide
across the surface.
The Hamaker solution for the attractive energy between a flat surface and a cylinder is a
function of the radius of the fibril, R, the length of the cylinder (or here, the fibril), L, the
separation distance between the fibril and the surface, d, and the Hamaker constant, A12 (the
theoretical value of the Hamaker constant for a PTFE-silica interaction (7.6 × 10−20 J) is used
here126). Expanded PTFE, in which PTFE filaments are highly aligned, exhibits a fibril tensile
strength of approximately 400-700 MPa. The energy equation is differentiated to yield the
attractive force between the fibril and the countersurface, Fadh, per unit length (Eq 4-1). This
attractive force, Fadh, multiplied by the friction coefficient of PTFE (~0.1 – a lower bound
considering the friction coefficient of the composite is nearly twice as high) equals the tensile
force applied to an aligned fibril in contact over the length L during sliding. This force is then set
equal to the ultimate tensile strength of the fibril, σf. Eq 4-2 allows us to solve for a critical fibril
length Lc (Eq 4-3). The interpretation is that any fibril in contact with the countersurface over a
length Lc or greater may be broken in sliding. This hypothesis is supported in the literature by
Makinson in 19645 and Brainard in 1973117. Additionally, evidence has previously been
published for the transfer of oriented films of PTFE onto a glass (silica) substrate during sliding
contact.127 It is evident that forces on aligned fibrils due to adhesion and friction are more than
sufficient (at entirely reasonable values of Lc) to break fibrils, and therefore C-C bonds at the
sliding surface.
58
𝐹𝑎𝑑ℎ
𝐿=
𝐴12√𝑅
8√2∙𝑑52
(4-1)
µ𝐹𝑎𝑑ℎ = 𝜎𝑓 ∙ 𝜋𝑅2 (4-2)
𝐿𝑐~36𝜎𝑓
µ𝐴12∙ 𝑑
5
2 ∙ 𝑅3
2 (4-3)
In further support of instant transfer of PTFE, a control spectrum was obtained from the
an inner surface of the PTFE composite – accessed by slicing the sample with a razor blade -
with no sliding history, but the same thermal and environmental history. When the diamond ATR
crystal was held in place against the nascent surface, the spectrum contained the usual IR peaks
for bulk PTFE (Figure 4-5c i) at 1203 and 1149 cm-1. Once the diamond ATR crystal was
removed from contact with the sample, an IR spectrum was re-acquired in air before the crystal
was cleaned. Peaks in the C-F region were still visible, and in addition the spectrum included the
additional absorbance peak at 1253 cm-1 (Figure 4-5c ii) - identical to the one from the single
cycle sliding experiment (Figure 4-5c iii). PTFE chains must therefore have transferred from the
surface of the composite to the ATR crystal even after arguably static contact. The same three
peaks were also obtained from similar residual analysis of the ATR crystal after static contact
with an as-molded PTFE surface, which eliminated the razor cut as the source of the transferred
polymer chains. The transfer of PTFE to metal surfaces after static contact in high vacuum has
thus far only been observed by Auger photoelectron spectroscopy (AES).117
The IR spectra collected within the 100k cycle region or transfer film development
revealed new, broad peaks at 3388, 1650 and 1432 cm-1 in addition to the previously observed
PTFE backbone peaks at 1203 and 1149 cm-1. Within the 1M cycle region, these new peaks
dominate the spectrum (Figure 4-4b). Carboxylic acid end groups are often observed in
perfluoropolymers128. Their associated peaks are relatively sharp and are assigned to a mixture of
59
both monomer (1775 cm-1) and dimer (1813 cm-1) forms. However, similar, much broader and
lower frequency carbonyl peaks at 1655 and 1441 cm-1 have been reported for instances where
fluorinated carboxylic acids chelated to metals as shown by Kajdas et al.129,130 In Kajdas’ study,
perfluorooctanoic acid (C7F15COOH) was coated onto a steel surface and then heated. A
reproduced spectrum is provided in Figure 4-6 a. Figure 4-6 c illustrates the similarity between
the IR spectrum from Kajdas’ chelated perfluoro- acid and the IR spectrum taken within the 1M
cycle film region, which suggests that the species present are identical, and that the broken PTFE
chains within the transfer film must be adhered to the metal surface in the same manner.
In a separate experiment at DuPont, a small molecule model compound (C6F13COOH)
and the same α-alumina were pre-mixed and dispersed in PTFE 7C for analysis by transmission
IR (Figure 4-6b). The resulting spectrum is very similar to that of the ATR-IR spectrum of the
running film (Figure 4-6d), which is in turn quite similar to that of the transfer film, suggesting
similar chemical changes occur in all. A control experiment of similar design without dispersed
α-alumina produced a spectrum that instead displayed the sharper monomer and dimer Rf-COOH
acid peaks at 1813 and 1775 cm-1, respectively. It is therefore suggested that the carboxylate
ends of PTFE chain ends chelate not only to the steel surface under the transfer film, but also to
the surfaces of the alumina filler particles. In fact, hydrocarbon carboxylic acids are already
known to react with the amphoteric surface of alumina particles. 131 It is not unreasonable, then,
for perfluorinated carboxylic acids (much stronger Brønsted acids than their hydrocarbon
analogs132) to react with and chelate to the alumina surface, even in the absence of excess heat.
ATR-IR analysis (Figure 4-4c) of the cumulative wear debris formed the end of the 1M
cycle track (shown schematically in Figure 3-2a) revealed large absorbance peaks for the PTFE
backbone, smaller monomer and dimer carboxylic acids peaks, and small peaks corresponding to
60
the chelated salts. The cumulative wear debris is composed of polymer fragments shed during
the wear process. Thus, the chemical intermediates (carboxylic acids) were observed here in
addition to chelated salts (to the alumina particles) because further chemical modification or
chelation was not possible at all of the reaction sites after ejection from the wear track.
The results of the comparative IR analysis of the formation and evolution of the
tribofilms formed in this ultra-low wear polymer composite system led to the chemical
mechanism proposed in Figure 4-5. Figure 4-7 a & b illustrate the first step of the process: the
mechanochemical scission of the carbon-carbon bond in PTFE5,117 , which forms reactive
perfluoroalkyl radicals. The following steps of the mechanism (Figure 4-7 b-e) are identical to
the process of e-beam irradiation of high molecular weight PTFE in ambient air.14,103,133 The
perfluoroalkyl radicals then react with atmospheric oxygen to form a peroxy radical (Figure 4-
7c), which quickly decompose into more stable acyl fluoride end groups (Figure 4-7d). Acyl
fluoride end groups are unstable toward water and therefore hydrolyze in ambient humidity to
form carboxylic acids (Figure 4-7e). The dependence of the wear rate of this system on humidity
and vacuum environments has been described in previous chapters.43,46 HF is produced128 during
these steps, and likely goes on to form metal fluorides at the surface of the countersample. The
carboxylic acids chelate to the steel countersurface (Figure 4-7f) as evidenced by the IR studies
described. This strongly adheres them to the metal surface to form the robust transfer film that is
so closely associated with the ultra-low wear behavior of the composite.44,45
Jintang and Hongxin134 presented a similar mechanism for the mechanochemistry of
PTFE at the sliding interface, but do not mention the critical carboxylate groups which are the
source of the adhesion between the transfer film and the metal countersurface. The chelation of
carboxylate ended PTFE chains to the alumina fillers at the interface also reinforces the running
61
surface of the polymer47,92, which increases the wear resistance of the polymer, and likely
reduces creep in the system. The tribochemical interaction between the PTFE/α-alumina
composite that breaks the C-C bonds is the first step in a complex cascade of events. Low wear
in this instance appears to be a property of the system created during sliding, and is heavily
dependent on the environment. Although conditions at the sliding interface are much milder than
what is required for thermal bond cleavage in PTFE,135 similar reaction products are detected in
both the running and transfer films at the interface (Figure 4-6). At low sliding speed, low
nominal contact pressure, and frictional heating not more than ~1° C removed from ambient
temperature, it is the coupling of mechanical and chemical effects that facilitates the formation of
the interfacial tribofilms necessary to maintain low wear over hundreds of thousands and even
millions of sliding cycles.
A study published by Khare around the same time as these results further investigated the
effects of environmental composition and counterface temperature on the chemical and
morphological response of the composite and its associated tribofilms.136 Khare et al. found that
across a range of humidity and oxygen contents, the composite always ran in to a low-wear
value, but in the case of the dry environments the wear of the composite increased again after
run-in. At high temperatures, wear was lower, the transfer film was thinner, and the carboxylate
signal was reduced. It is likely that increased temperature drove the reactions kinetically while
being insufficient to remove all of the necessary intermediate species at the interface (water,
oxygen). A wear experiment of the same type as many previously mentioned (flat on flat,
reciprocating over 25.4 mm at 50.8 mm/s at around 6 MPa) submerged the sliding interface of
the PTFE/α-alumina composite in question in distilled water. The sample ran for ten thousand
cycles at a wear rate nearly as high as unfilled PTFE, and three orders of magnitude higher than
62
the dry composite after the same number of cycles (Figure 4-8). A transfer film never formed,
visible adhesion of debris was minimal to nonexistent and no color change to reddish brown was
observed. The reactions necessary for stable film formation could not proceed in the submerged
environment, and thus the wear of the sample remained very high
Khare et al. observed the formation of carboxylate salts in both high wear and low wear
systems, and found no correlation between the oxygen content of the test chamber and the steady
state wear rate of the composite. However, they could not rule out the presence of trace amounts
of oxygen that could still be sufficient for the mechanism to proceed as suggested here.
Furthermore, their IR studies assigned peaks at 1315 and 1360 cm-1 to shortened chain lengths
within the polymer, supporting the hypothesis of mechanical chain scission described above, and
interface pairing experiments showed (in agreement with Bahadur and Tabor36) that an existing
transfer film did not reduce the wear of a fresh composite pin, and that an existing running film
significantly reduced the run-in volume against a fresh countersurface. The reinforcement of the
composite due to the chelation of PTFE chains to the alumina filler particles must then be the
dominant source of wear abatement, and the transfer films are an effect thereof.
It was then supposed that if a precursor to the running film could be formed at on the
sliding surface of the composite that the run in period of high wear might be abated or avoided
entirely. Samples of unfilled PTFE and of the PTFE/α-alumina composite were etched and tested
as described previously. Wear results from the stripe tests are given in Figure 4-9. The etched
surface of the unfilled PTFE sample lowered the wear rate by almost two orders of magnitude
over the first ten cycles (~0.5 m) of sliding, but quickly wore through over the next hundred
meters. The wear rate then returned to a higher value, closer to that of the unmodified PTFE
sample. The wear rate of the etched composite over the first 1000 cycles (~50 m) was reduced by
63
nearly an order of magnitude compared to the inner, unmodified surface, and over the next 10k
cycles, the etched sample was observed to have lost no mass within the error of the balance. The
final, steady state wear rate of both the etched and unmodified composite surfaces was very
similar, but the etching treatment did indeed appear to greatly reduce the run-in volume during
the beginning of the test.
The results of these wear tests strongly correlate with the appearance of the transfer and
running films formed (Figure 4-10). The highest wearing sample, the unfilled, unmodified PTFE,
left large wear platelets in the patchy transfer pattern characteristic of PTFE sliding at more than
~10mm/s.5 The etched surface of the unfilled PTFE performed slightly better, and the debris
within the wear track resembles the brown transfer film characteristic of the composite at the
metal surface, but is covered with PTFE wear platelets that formed once the etched film was
worn through and the polymer resumed increased wear. The edges of the etched surface not on
the leading or trailing faces were not worn through, but this is possibly due to a leveling issue
between the polymer pin and the metal countersurface. The wear results should not be
compromised however, as the wear rate is normalized by the supported load and such the wear
area is reduced but the nominal contact pressure is increased. The difference in contact pressure
should not be enough to invoke any significant dependence of wear rate on load that may exist.
The unmodified inner surface of the PTFE/α-alumina composite performed as expected
and formed the characteristic brown transfer film. The appearance of the running film changed
very little over the duration of the stripe test. Most notably, the etched surface of the composite
formed a much thinner transfer film, as supported by the markedly lower run-in wear volume.
Again, the running surface of the polymer hardly changed in appearance throughout the duration
of the stripe test, but based on the presence of any visible transfer film it is possible that the
64
etched surface layer may have been mostly or completely replaced by a running film formed in-
situ. It is likely that the carboxylic acid end groups in the etched layer will have immediately
chelated to the alumina particles at the near surface. During the run-in period, the polymer is
reinforced as the large, friable alumina particles break up, further increasing the concentration of
virtual cross-links in the running film. Transfer to the countersurface is minimal during this
period, but a thin layer of chelated PTFE chains still appears to fill in the negative features of the
steel and provide a softer, reinforced polymer surface for the polymer pin to slide against. The
results of these etching experiments support Khare’s observation that a pre-existing running film
greatly reduces run-in of the composite, and confirm that the introduction of carboxylic acid
chain ends at the sliding interface speed up the transition to ultra-low wear.
Figure 4-1. Wear rates and friction coefficients for each experiment run as a part of the stripe
test. The wear rate of the composite behaved as expected and ran in to ultra-low wear
despite the alteration of the wear track length (a), and decreased through a run-in and
transition period to a steady state on the order of 10-7 mm3/N∙m. The friction
coefficient (b) increased over the first hundred cycles and then ran in to around 0.17
over the million cycles of sliding.
65
Figure 4-2. Stylus profiles of transfer films after 1 cycle, 100k cycles, and 1M cycles of
development. Most transfer film growth took place between 100k and 1M cycles.
Each line scan provided is an average of 50 line scans at 0.04 mm intervals.
Figure 4-3. High resolution XPS of the transfer film in various stages of development
demonstrate some of the chemical changes present in the films, and the evolution of
atomic concentration as film formation progresses. a) C1s spectra from exposed areas
of transfer film after 0, 1, 100k and 1M sliding cycles and b) atomic concentrations in
regions of transfer film development after 0, 1, 100k and 1M sliding cycles compared
to bare stainless steel.
66
Figure 4-4. Infrared reflectance results from the transfer film on a stainless steel surface after (a)
one cycle of sliding, (b) 100k and 1M cycles, and (c) ATR-IR spectrum of cumulative
wear debris. (Reproduced from Harris 201549)
67
Figure 4-5. A Hamaker solution for the attractive energy between a flat surface and a cylinder is
used to support the hypothesis that PTFE fibrils break and transfer during the first
cycle of sliding. (a) A PTFE fibril of average radius R and separation distance d from
a countersample contacts the sliding countersurface over a length L. The fibril
experiences an adhesive force due to van der Waals interactions at the surface (b) A
plot of the critical fibril length Lc as it changes with average fibril radius R for
separations d=1, 2 and 10 Å demonstrates the increase in Lc as R and d increase.
Shaded circles are in place to represent fibrils of various aspect ratios. (c) (i) IR
spectra of bulk PTFE, (ii) the residue spectrum after pulling away from the bulk
PTFE, and (iii) the 1 cycle transfer film. (Adapted from Harris 201549).
68
Figure 4-6. The carbonyl region of IR spectra of the tribofilms is compared to small molecule
model reactions with perfluorinated carboxylic acids: (a) the IR spectrum
published by Kajdas and Przedlacki130 ascribed to the chelated salt of
perfluorooctanoic acid on a steel countersurface. (b) the IR spectrum collected
from a pressed PTFE film filled with a dispersion of alumina particles and
perfluoroheptanoic acid (c) the IR reflectance spectrum acquired within the 1M
cycle transfer film, and (d) ATR-IR spectrum of the running surface of the
polymer composite pin after 1M cycles. (Adapted from Harris 201549)
69
Figure 4-7. The chemical mechanism responsible for PTFE/α-alumina tribofilm formation and
adhesion. Chain scission of PTFE (a) is caused during sliding due to frictional forces
and adhesive forces between the PTFE fibrils and the metal countersurface (b).
Environmental oxygen reacts with the radicals created (c). The unstable end groups
decompose to form acyl fluorides (d). Moisture in the air hydrolyzes these groups,
forming carboxylic acids (e). Carboxylic acid end groups react with and chelate to the
surface of the metal and alumina particles (f). Figure adapted from Harris 2015.49
70
Figure 4-8. The wear rate of the PTFE/α-alumina composite and neat PTFE compared to that of
the submerged composite. The wear rate of the composite was three orders of
magnitude higher against 304 stainless steel when submerged in distilled water than
against the same type of countersurface in laboratory air.
Figure 4-9. The wear rates of the etched and unmodified surfaces of unfilled PTFE and of the
PTFE/α-alumina composite are plotted versus the total sliding distance. Sliding
distance is used rather than the number of sliding cycles because of the variable
length of the wear track during a stripe test. The mass loss of the PTFE/α-alumina
composite after the second (10k cycle) test was measured to be zero, which cannot be
represented on a log-log plot and is therefore omitted from the figure.
71
Figure 4-10. Optical images of the transfer and running films of the etched and unmodified
surfaces of unfilled PTFE and of the PTFE/α-alumina composite. Transfer film
images are displayed alongside images of the running films of each of the samples
after the end of each stripe test (10k cycles total for the unfilled PTFE samples, and
100k cycles total for the composite samples). Photos taken by author in Spring 2016.
72
CHAPTER 5
CONCLUSIONS
Transfer of PTFE to the metal countersurface occurs after a single cycle of sliding, and
even after static contact, as evidenced by FTIR (Figure 4-4). The polymer/alumina composite
then gently abrades/cleanses the adventitious carbon and surface oxides from the surface of the
steel as it slides. Without removal of this contaminant layer and native oxides, the chelation of
the carboxyl terminated fluoropolymer would be much less likely, which would limit the
adhesion of the tribofilm. The hypothesis of mechanochemical chain scission is supported by a
force balance model based on van der Waals interactions between the polymer and the metal
countersurface, described in Figure 4-5. The wear rate of the composite was high during the run-
in period and did not fall below 10-6 mm3/(N·m) until the 100k cycle test. The change in wear
rate from high to low coincided with the appearance of IR peaks at 1650 and 1432 cm-1 in the
transfer and running film spectra (Figure 4-4b, Figure 4-6c and d) which are indicative of the
presence of chelated carboxylate polymer chain ends (Figure 4-7f). The concentration of
chelated salts increased significantly during the 1M cycle test (Figure 4-4b) as more and more
chains broke and reacted with the metal and Al2O3 surfaces. Carboxylic acid chain ends that did
not chelate to a surface were ejected as wear debris, and were the source of the carboxylic acid
monomers and dimers identified by IR of the cumulative wear debris (Figure 4-4c).
The formation of an ultralow wear PTFE transfer film on 304 stainless steel involves the
chemical interaction between the polymer composite, the embedded alumina particles, the
ambient atmosphere, and the metal countersurface. This is a cycle-dependent process (under
reciprocating conditions) that relies on the mechanical input of energy to cause chain scission,
which in turn initiates the radical-driven mechanism of transfer film formation. The reactions
proposed in this mechanism have been observed previously - in non-tribological settings – but in
73
this context they neatly describe the chemistry behind the observed wear behavior of this
PTFE/α-alumina composite. Figure 5-1 summarizes this ultralow wear system, which arises from
a complex set of variables that together allow chemical modification of both sliding surfaces
under relatively mild conditions. The ultra-low wear behavior of the PTFE/α-alumina composite
described is driven by reactions within the running film, and is accompanied by similar reactions
with the countersurface, though the latter are an effect rather than a cause of the low wear
behavior. High wear composites remove debris quickly from contact, but in reinforced, ultra-low
wear systems, debris is retained and may react with species present.
It is shown in this work that PTFE, though widely regarded as an inert, environmentally
insensitive polymer is in fact quite reactive within the sliding contact. An explanation for chain
scission of PTFE under mild conditions has introduced a new perspective regarding the behavior
of ultra-low wear PTFE/α-alumina composites. The following steps of the mechanism have been
understood in the context of fluoropolymers, but were never applied to the wear process because
the presence of the initial carbon radical due to the scission of the carbon backbone had not been
considered. The performance of these ultra-low wear composites cannot be attributed to any
single component of the system, but the chemical bonds formed during sliding between the
polymer chains, the alumina particles and the countersurface are crucial to the formation and
longevity of the tribofilms. This discovery has not only changed the industry’s understanding of
these composites, but introduces the possibility of simple surface modifications to the
composites themselves rather than to their countersurfaces that further increase their usable
lifetime.
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Figure 5-1. Radical chemistry at the sliding interface proceeds despite mild conditions (low
speed, low nominal contact pressure, and low frictional temperature change). The
circulation, rather than ejection, of debris between the transfer and running films is
likely key to the high cycle maintenance of ultralow wear. Adapted from Harris
2015.49
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APPENDIX
FUTURE CONSIDERATIONS: COUNTERFACE EFFECTS
Wear and Friction Experiments on Additional Countersurfaces
Throughout the studies on the relationship between PTFE, the α-alumina filler and the
stainless steel countersurface, which has become the standard for polymer wear testing, many
wear and friction studies were performed on alternate metal countersurfaces in order to elucidate
more characteristics of the polymer/metal system. Wear and friction results for a suite of
experiments spanning multiple years is presented in Figure A-1 and Figure A-2. Experiments
were all performed using the same polymer pin dimensions, sliding speeds and track lengths as
usual (6.3x6.3 mm polymer running surface, 250 N normal load, 50.8 mm/s on a 25.4 mm long
wear track). The error in measured wear rate is left out of the plots due to the logarithmic scale
and density of data, but in the case of tests run out to 100,000 cycles or more, the error in
measured wear rates during steady state is frequently, if not always very small, and usually an
order of magnitude less than the measured wear rate.
Against most steels, which were lapped and polished to average roughnesses around 150
nm, the measured wear of the composite behaved as expected, with a period of run-in followed
by steady state low wear between 10-8 and 10-5 mm3/N∙m, as illustrated in Figure A-1. The
friction coefficients of the polymer/steel systems were measured to be mostly around 0.2.
Strangely, in the case of the high strength 4340 countersample, the friction coefficient increased
dramatically after 10,000 cycles, but the wear rate of the composite remained at a lot, steady
state value. The cause of this behavior is unknown, but it seems to highlight the robust nature of
the tribofilms formed, as the wear performance was not compromised.
Three copper countersurfaces, pure copper (CDA110), beryllium copper (CDA 172) and
naval brass (C464) were selected for testing due to their wide range of mechanical properties and
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because copper was perceived as a fairly reactive metal that could possibly facilitate reactions at
the interface, which could increase the chance of stable film formation. Metal surfaces were
lapped and polished to ~150 nm Ra. The wear results are summarized in Figure A-1. Both the
pure copper and the naval brass performed relatively poorly, with composite wear rates in the
range of 10-5, while the beryllium copper rivalled the performance of 304 stainless steel. The
friction was highest during the beryllium copper test, and especially low against pure copper. A
wear summary of the copper tests is provided in Figure A-3 in conjunction with friction traces
for each, and photos of the polymer and metal wear surfaces. On the poor performing pure
copper surface, the polymer film is patchy and a greenish tint at the surface suggests the presence
of copper carbonates. Furthermore, the wear debris at the track ends also appears to be filled with
green copper carbonate and black copper oxides, and the countersurface is abraded. The running
film appears poorly developed and likely never completely formed under gross wear conditions.
The beryllium copper films were well formed, and the naval brass countersurface was abraded
and accompanied by a good deal of black wear debris, likely degraded PTFE and copper oxide.
The friction trace of the beryllium copper test also includes drops from around 0.3 to around 0.2
at cycle numbers corresponding to mass measurements, for which the polymer was removed
from contact with the transfer film. It is speculated that this may be due to water adsorption
occurring at the film surfaces during these periods.
Friction and wear tests were also performed against 24k gold, pure platinum, and pure
lead. The platinum and gold surfaces were smooth, mirrored surfaces when they arrived and
were not polished further. The lead countersample was polished by hand to around 400 nm Ra.
Wear and friction results are given in Figures A-1 and A-2 respectively. The gold countersurface
itself failed before reaching 10,000 cycles. Because of counterface failure, the later wear rate
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includes mass uptake from gold transfer, and the later friction coefficient includes the ploughing
force to yield the gold. A more detailed summary of the experiment is provided by Figure A-4.
No transfer film could be observed on the countersurface after its failure, and a great deal of gold
transferred to the polymer running surface, visible in the optical images in Figure A-4b. Most
interestingly, the polymer running surface not obscured by significant gold transfer appeared
bluish in tint. It is possible to create visibly blue solution of large (~100nm) gold nanoparticles,
or by adding excess salt to a solution of smaller gold nanoparticles, and a blue solution of gold
‘nano-urchins’ is also possible137. Chemical interaction between the composite components and
the gold countersurface was not expected. The platinum countersurface held up during the wear
test, although the ultimate performance of the composite was just shy of very low wear, never
dipping into the 10-7 mm3/N∙m range. Increases in friction to 0.2 at mass measurement locations
are visible in the platinum friction and are again attributed to water adsorption at the films. The
lead countersurface held up surprisingly well for around 150,000 cycles. The wear of the
polymer during this period was also fairly low, but not as low as is usually observed against
stainless steel. The transfer film that formed before the failure of the lead countersurface
appeared to be extremely thin – thin enough to appear iridescent. The wear track before and after
failure is given in Figure A-4b. The polymer surface from the lead test was also discolored as
during the gold test, but the color observed was yellow. The yellow tint persisted after the failure
of the countersurface during which the polymer pin was also misshaped and subject to lead
transfer. Both platinum and lead were thought to be more reactive than the gold, and therefore
were considered more likely candidates for stable film formation.
Three aluminum alloys (6061 T6, 7075 T6, 1100) were chosen as countersurfaces for the
PTFE composite as a negative control, based on the classically poor tribological performance of
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aluminum surfaces which is normally abated using a surface coating, particulate reinforcement
of the metal, or anodizing138–140. Surprisingly, all three of the alloys outperformed the stainless
steel countersurface used for reference with respect to the wear of the composite (Figure A-1) at
a slightly higher coefficient of friction (Figure A-2). Traces of the friction coefficient of the
composite against the three aluminum alloys and against stainless steel are provided in Figure A-
5, along with optical images of the transfer films. The films in each case appeared robust and
uniform, and were all gray or white in color.
Infrared Spectroscopy and X-Ray Analysis of Transfer and Running Films
The transfer films formed on C1018 and M2, the worst and best performing steels,
respectively, were analyzed using FT-IR in reflectance mode in the same manner as the 304
stainless countersurfaces described in previous chapters. A comparison of the resulting spectra is
given in Figure A-6 with arbitrary intensity. All spectra taken from the wear tracks against steels
show C-F stretch near 1200 cm-1 and carboxylate salt C=O peaks near 1434 and 1655 cm-1 on
C1018, and 1430 and 1670 cm-1 on M2. Water is indicated by a broad O-H peak near 3300 cm-1.
These are likely to indicate hydrated perfluoro carboxylate salt material(s). A broad peak below
900 cm-1 is likely due to metal-oxide M-O stretch, probably from aluminum oxide. Minor peaks
near 949, 993, and 1311 cm-1 in the M2 spectrum are unassigned at this time. It is important to
note that the presence of chelated polymer chains at the metal surface does not in itself predict
low wear of the polymer composite. The polymer chains may chelate not only to the metal
surface, but also to the filler particles, and chelation is observed in the wear debris as well as in
the transfer and running films.
Representative IR spectra of the transfer films on CDA 110 and CDA 172 are provided
(with arbitrary intensity) in Figure A-7, again in comparison to the stainless steel transfer film.
Both copper spectra show C-F stretch near 1200 cm-1 and carboxylate salt C=O peaks. The
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higher frequency C=O peak on pure copper show a sharp location near 1615 cm-1 in most spectra
and a broader location in the 1650-1660 cm-1 range. The sharper peak could be from the salt and
the broader peak from water of hydration. The second carboxylate peak is near 1435 cm-1.
Carboxylate peaks on the beryllium copper countersurface are seen near 1435 and 1655 cm-1.
Water is indicated in both spectra by a broad O-H peak near 3400 cm-1, consistent with hydrated
perfluoro carboxylate salt materials. Again, a broad peak below 900 cm-1 is likely due to metal-
oxide M-O stretch, probably from aluminum oxide. A sharp peak near 3643 cm-1 in most spectra
may be from the water of hydration of an inorganic component. A broad peak below 900 cm-1 is
likely due to metal-oxide M-O stretch, probably from aluminum oxide. Minor peaks near 1320
and 1363 cm-1 are consistent and are likely due to the salt and may have to do with an oxalate-
like structure but are not assigned at this time. Again, the presence of carboxylate salts suggests
chain scission in PTFE and reactivity at the scission sites, but is not a predictor of low wear.
XPS spectra (C1s) were collected from the running films created during the
lead/platinum/gold tests, and are given with arbitrary intensity in Figure A-8 in comparison to
that of nascent PTFE and that of a running film formed against 304 stainless steel. In all cases,
chain scission and end group chemistry appear to be present. However, precise bonding
information regarding chelation is not available from this analysis, and IR spectra are not yet
available.
In the case of all three transfer films on the aluminum alloys, C-F stretch near 1200 cm-1,
water (the broad O-H peak near 3250 cm-1) and carboxylate salt peaks (1425 and 1672 cm-1 in
the case of 6061, 1425 and 1655 in the case of 7075, and 1417 and 1657 cm-1 in the case of
1100) are present in the FT-IR spectra, shown with arbitrary intensity in Figure A-9. As is the
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case with most of the other transfer films, these results are consistent with the presence of a
hydrated perfluoro carboxylate salt material.
Surface and Sub-Surface Evolution of the Aluminum Countersamples
Further analysis of the aluminum transfer films and countersamples was performed
because of the remarkable nature of the wear results, and results were considered in comparison
to the case of 304 stainless steel. Against each of the three aluminum countersurfaces, a thin,
stable transfer film had formed, the polymer composite reached ultra-low wear, and contrary to
expectation the aluminum countersamples were not subjected to delamination or other visible
mechanical failure. Stylus profilometry of the transfer films shows more variability across the
aluminum samples than against stainless steel (Figure A-10), and it appears that in the case of the
6061 T6 countersurface, the polymer composite wore into the metal surface as the low wear self-
mated film system formed.
The interaction between the polymer composite and the aluminum countersurfaces was
then assumed to extend into the subsurface of the metal. Removal of the polymer film allows
inspection of the changes undergone by the metal surfaces during testing. The transfer films
formed by polymer composites against stainless steel are frequently easily removed using boiling
water, and the same process was applied to two of the aluminum countersurfaces (6061 T6 and
Al 7075 T6). IR spectra taken in various areas within the wear regions after boiling water
treatment are provided for 304 stainless steel, Al 6061 T6, and Al 7075 T6 in Figure A-11.
Within the direction change region (characteristic of reciprocating tests) and the center of the
wear track of the stainless steel countersample, no significant peaks were detected, indicating no
residual fluorinated material on the metal surface.
A spectrum taken within the striated wear track area of the 6061 T6 countersurface
showed no fluorinated material, but spectra collected in the direction-change area and the end of
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the travel adjacent to the direction change region revealed the persisting presence of a hydrated
perfluoro carboxylate salt. A peak near 1049 cm-1 in the direction change area is visible, and may
have to do with aluminum chemistry. A spectrum from the striated wear track area of the 7075
T6 countersurface provides weak evidence for fluorinated material, probably the carboxylate
salt. The end-region spectrum taken at the end of the wear track area adjacent to the direction
change region shows no significant fluorine presence. A spectrum taken within the 7075 T6
wear track includes a new peak near 1065 cm-1 that may have to do with aluminum chemistry. A
peak at 1065 cm-1 is associated in vaccine science with a structural hydroxide environment of
poorly crystalline boehmite, an aluminum oxide hydroxide mineral that may have to do with the
boiling treatment.141 A spectrum from the direction-change region of the 7075 T6 countersurface
clearly indicates the persisting presence of a hydrated perfluoro carboxylate salt, and includes the
1065 cm-1 peak. The C-F peaks near 1200 cm-1 were more consistent in frequency compared to
the 6061 T6 countersurface. The O-H region above 3000 cm-1 is variable in intensity for both
aluminum countersurfaces analyzed, and may have to do with hydration of aluminum, hydroxide
formation and/or hydrated carboxylate salt when present.
In order to further investigate the interaction of the composite and the transfer film with
the aluminum countersurface, a Focused Ion Beam (FIB) microscope was used to cut trenches
within and outside of the wear track. Images acquired at the University of Florida are given in
Figure A-12a (within the wear track) and A-12b (in the nascent metal away from the wear track),
and images acquired later at the Center for Integrated Nanotechnology (CINT) in Albuquerque
are provided in Figure A-12c. Within nearly every FIB trench, distinct and unusual features are
visible in the region 1-2 µm beneath the surface. Grain size appears greatly reduced at the near
surface. In the trenches taken within the wear tracks, it is presumed that the black stripe at the
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surface below the protective platinum layer and the carbon coat is the polymer composite’s
transfer film. The large circular feature beneath the surface of the 1100 is of particular note, and
may suggest the presence of polymer penetration into the surface of the metal, which could
explain the persistence of fluorine signal after the removal of the visible transfer film with
boiling water. However, because the near surface appears affected within the nascent metals, the
features visible cannot be confidently attributed in full or in part to the sliding composite contact.
In the months following the FIB investigation, it was still thought that the polymer
composite was somehow uniquely suited for ultra-low wear sliding contact against aluminum
countersurfaces, and many further wear tests were conducted against aluminum 6061 T6
countersurfaces from various sources. Results of the wear tests are summarized in Figure A-13.
The original, high performance aluminum countersurfaces were delivered as ‘lapped’, at a
quoted average roughness of around 150 nm. Attempts in lab to replicate the success of the
original surface via wet polishing or lapping against a cast iron wheel with silicon carbide/oil
lapping compound produced only countersurfaces against which the wear of the composite was
orders of magnitude higher than against the original samples, seemingly regardless of measured
average surface roughness. It was apparent at this juncture that only the aluminum
countersurfaces prepared by the original supplier, Alabama Specialty Products Incorporated
(ALSPI), were consistently achieving ultra-low wear with the PTFE/α-alumina composite,
despite attempts to reproduce their polishing process exactly. Indeed, aluminum samples from
other suppliers that were sent to ALSPI to be lapped also held up in wear tests and saw the
polymer composite reach ultra-low wear conditions for hundreds of thousands of cycles.
It was then supposed that perhaps some component of their lapping process was unique to
their location. Samples of the ALSPI lapped 6061 T6 countersurfaces and of the poor-performing
83
coupons from other sources and subjected to various surface treatments were shipped to the
University of Illinois at Urbana-Champaign for further FIB analysis for comparison, and for
nanoindentation as an initial check for any changes to mechanical properties of the surface due to
the ALSPI lapping process. The FIB images taken are shown in Figure A-14, and confirm a large
difference between the surfaces produced at ALSPI and the others tested. The images from the as
received, unpolished surfaces (samples 1 and 2) show no distinct features at the near surface, and
those of polished or milled (samples 4 and 5) show some grain refinement. However, the ALSPI
lapped countersurface (sample 3) again was seen to possess an unusually disrupted structure. It is
important to note that the features in the ALSPI lapped surfaces illustrated here are seen in a
large fraction of the trenches milled. Nanoindentation (Berkovich tip, 5 mN) was performed on
each countersurface and the hardness values and elastic moduli reported in Figure A-15 are the
average results from nine indents per countersurface. No significant difference in hardness or
modulus was found between the surfaces tested despite the large disparity in tribological
performance, although a brief Energy Dispersive X-Ray Spectroscopy (EDS) scan indicated the
presence of iron in the surface of the ALSPI sample (3), and not in sample 1.
An in-person visit to ALSPI allowed for closer inspection of the lapping process that was
producing such remarkable aluminum countersamples. The lapping wheel used was cast iron,
and coated with thick, oil lapping compound that was not cleaned or changed between different
metal samples. The aluminum samples were lapped for around ten minutes under only the dead
weight load of the sample holder, wiped clean with a dry rag, and polished briefly and gently by
hand using 600 grit silicon carbide sandpaper that was well used previously to smooth mostly
iron samples. Figure A-16 includes several images of the process. Samples were collected of the
lapping compound, which was later suspended in ethanol and filtered for solids to be analyzed
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via EDS. One resulting spectrum is provided in Figure A-17, and indicates the presence of not
only the original components of the lapping compound (Si, C, O), but an extremely large amount
of contamination by metals processed in the facility (Co, Cr, Fe, Al). One resulting hypothesis
was that the lapping process embedded iron from the lapping wheel (which wears over time, with
use) and other metal particles left in the lapping compound into the relatively soft and malleable
surface of the aluminum, acting as reinforcement, though nano-indentation revealed no hardness
or modulus effects.
Collaborators at DuPont collected EDS maps of the near surface in FIB trenches milled in
a 6061 T6 sample lapped at ALSPI, and in another subjected only to hand polishing on wet
silicon carbide sandpaper. Backscatter micrographs and the corresponding EDS maps in Figure
A-18 show that both samples contain iron (consistent with the composition of the alloy), but the
distribution thereof is quite different between them. In the hand polished sample, against which
the wear of the PTFE/α-alumina composite was quite high, the distribution of iron is regular and
random. In the ALSPI lapped sample, against which the wear of the composite was very low, the
iron is distributed throughout the sample, but is also concentrated in horizontal, parallel tracks
around 10 µm apart. It is presumed that this is due to the lapping process, but the mechanism is
as yet unknown. However, the nano-crystallization of the surface of aluminum and other
crystalline in a rubbing contact has been described for decades as the Beilby layer, the depth of
which is dependent on the ‘vigor’ of the polishing action.142–147 Recently, improvements to the
wear resistance of aluminum via mechanical surface modification were described by Nimura in
in the case of fretting in oil.148 It is possible that correctly replicating the surface modifications
seen in aluminum samples lapped by ALSPI, a simple, low cost method of preparing aluminum
85
surfaces for dry tribological contact with polymer composites could be established, reducing the
need for coatings or more complex surface treatments that are commonly used today.
Figure A-1. A summary of the wear results from PTFE/α-alumina composite experiments against
various metal countersurfaces.
86
Figure A-2. A summary of the friction results from PTFE/α-alumina composite experiments
against various metal countersurfaces.
87
Figure A-3. A summary of the results of PTFE/α-alumina composite experiments against various
copper containing countersurfaces. a) The total wear rate as sliding progressed, b)
optical images of the wear tracks and running films and c) friction traces spanning the
length of each experiment. Photos taken by author in 2014.
88
Figure A-4. A summary of the results of PTFE/α-alumina composite experiments against lead,
gold and platinum. a) The total wear rate as sliding progressed, b) optical images of
the wear tracks and running films and c) friction traces spanning the length of each
experiment. Photos taken by author in 2014.
89
Figure A-5. Friction traces from the PTFE/α-alumina composite experiments against three
aluminum alloys and against stainless steel, and optical images of the transfer films
from the aluminum experiments. Photos taken by author in 2014.
Figure A-6. FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three types of steel are compared to the spectrum taken within a
transfer film formed on stainless steel.
90
Figure A-7. FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three different copper alloys are compared to the spectrum taken
within a transfer film formed on stainless steel.
Figure A-8. XPS spectra taken from the running films formed by sliding the PTFE/α-alumina
composite against platinum, gold, lead, and stainless steel are compared to a spectrum
taken from the nascent polymer surface.
91
Figure A-9. FT-IR spectra taken within the transfer films formed by the PTFE/α-alumina
composite against three aluminum alloys are compared to the spectrum taken within a
transfer film formed on stainless steel.
Figure A-10. Stylus profilometry traces taken across the center of transfer films formed by the
PTFE/α-alumina composite against three aluminum alloys are compared to the profile
taken across a transfer film formed on stainless steel.
92
Figure A-11. FT-IR spectra illustrate the changes in surface chemistry before and after
attempting to remove the transfer films from 304 stainless steel, Al 6061 T6 and Al
7075 T6 with boiling water
93
Figure A-12. FIB cross sections taken within the wear track and in nascent areas of the three
aluminum alloys tested. a) micrographs acquired within the wear tracks of all three
alloys, b) micrographs acquired in nascent areas of the three alloys, and c)
micrographs acquired within the wear track of an Al 6061 T6 countersurface. Images
in a & b were collected by Nick Rudowski at the University of Florida in Fall 2013.
Images in c were collected at CINT in Albuquerque, NM by the author in Spring
2014.
94
Figure A-13. A summary of the wear results of the PTFE/α-alumina composite against a number
of Al 6061 T6 countersurfaces purchased from various locations and prepared in
different ways.
95
Figure A-14. Micrographs taken within FIB trenches in the nascent surfaces of five differently
prepared Al 6061 T6 surfaces. Micrographs collected by Matthew Bresin in Spring
2015 at the University of Illinois at Urbana-Champaign.
Figure A-15. Hardness and modulus results from nanoindentations performed on the surfaces of
five differently prepared Al 6061 T6 surfaces. No significant difference is observed.
96
Figure A-16. A series of images depicting the lapping and polishing process at ALSPI. Photos
taken by the author in Summer 2015 at Alabama Specialty Products Inc. in Munford,
AL.
97
Figure A-17. An EDS spectrum collected from a sample of the solids separated from the lapping
compound taken directly from the lapping wheel used by ALSPI.
98
Figure A-18. Backscattered electron micrographs and EDS maps of the subsurface (~1 mm
depth) of two Al 6061 T6 samples, one untreated, and one lapped and polished by
ALSPI. The images on the left were taken from a FIB trench in an untreated sample,
and the images on the right were taken from a FIB trench in a sample lapped and
polished by ALSPI.
99
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BIOGRAPHICAL SKETCH
Kathryn Harris completed her undergraduate degree in materials science with a minor in
chemistry at the University of Florida in 2011. She continued with her graduate studies under Dr.
W. G. Sawyer in the UF Tribology Lab. She very seriously considered dedicating this document
to her cat.