Note: Descriptions are shown in the official language in which they were submitted.
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A low friction coating for a cutting tool
The present invention relates to a composition
for use as a.coating on a cutting tool and, in
particular, to a low friction, wear-resistant coating
formed from a hard phase of, for example, TiN and a
soft, lubricating phase of, for example, MoS2.
Sputtered MoS2 films were developed some three
decades ago and have since been used principally for
space lubrication applications. Over the last five
years these films have been increasingly used as tool
coatings for cutting and shaping applications,
resulting in significant improvements to tool-life,
finish and productivity. Interest is now growing in
MoS2 tool coatings for dry machining applications,
where little or no lubricants and/or cooling fluids
are used. Although dry machining is now industrially
feasible using the latest generation of TiAIN
coatings, the additional use of an MoS2 overlayer has
been proposed to provide the lubricating effect
previously obtained by the use of environmentally-
harmful lubricant fluids.
The ever increasing demands placed on these
relatively soft films has resulted in considerable
effort to improve endurance properties. Deposition
parameters were first optimised for pure MoS2 films.
Later, the beneficial effect of metallic additives,
such as Ni, Pb and Ti, was applied to sputtered
coatings by co-deposition and in the form of
multilayers, bringing significant improvements to film
density, lubrication and endurance. In addition, the
multilayers display better resistance to oxidation,
which can degrade lubrication and endurance by
depleting superficial sulphur and undermining coating
adhesion. However, because the overlayers are still
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relatively soft, they are often consumed in the
critical break-in phase of a cutting tool and lose
their effectiveness soon afterwards.
TiB2 has been investigated as a hard matrix
material with a lubricating phase of MoS2. Whilst
hardness and low friction properties can generally be
obtained, such coatings can sometimes fail owing to
the brittleness of the TiH2 phase.
The present invention aims to address the
problems associated with the prior art and to provide
dense, relatively hard and wear-resistant coatings
with a sufficiently high enough concentration of MoS2
I5 or WSZ to obtain self lubrication. This is preferably
achieved by co-depositing a lubricating material with
a hard coating material.
Accordingly, in a first aspect the present
invention provides a composition for use as a coating
on a cutting tool, which composition comprises:
(i) a first material selected from TiNx, TiAINX,
TiAlYCrNx and CrNx, and
(ii) a second material selected from MoS= and WSZ,
wherein x is from 0.5 to 1.5, z is from 0.8 to 2.2 and
the atomic ratio (y) of Mo or W to Ti or Cr is from
0.1 to 0.8.
In a first preferred embodiment, x is from 0.5 to
1.5, z is from 1 to 2.2 and the atomic ratio (y) of Mo
or W to Ti or Cr is from 0.2 to 0.8. Advantageously,
x is from 0.7 to 1.3, more preferably 0.8 to 1.2,
still more preferably equal to or approximately equal
to 1 for even higher hardnesses and lower coefficients
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of friction. For the same reasons, z is preferably
from 1.2 to 2, more preferably from 1.4 to 1.8, still
mote preferably equal to or approximately equal to
1.6. Similarly, y is preferably from 0.2 to 0.5, more
preferably from 0.2 to 0.4, still more preferably
equal to or approximately equal to 0.3. The
composition according to this embodiment has a
hardness typically of at least 14 GPa, more typically
from 15 to 23 GPa, still more typically from 16 to 22
GPa.
In a second preferred embodiment, x is from 0.5
to 1.5, z is from 0.8 to 2.2 and the atomic ratio (y)
of Mo or W to Ti or Cr is between 0.1 to 0.2 when z is
from 0.8 to 2.2, and is from 0.1 to 0.8 when z is
between 0.8 and 1. x is advantageously from 0.7 to
1.3, more preferably 0.8 to 1.2, still more preferably
equal to or approximately equal to 1 for even higher
hardnesses and lower coefficients of friction. For
the same reasons z is preferably from 0.8 to 1.4.
Similarly, y is preferably from 0.12 to 0.18, more
preferably from 0.14 to 0.16, still more preferably
equal to or approximately equal to 0.15. The
composition according to this embodiment has a
hardness typically of at least 14 GPa, more preferably
from 22 to 32 GPa, still more preferably from 25 to 30
GPa.
For all embodiments of the present invention it
will be appreciated that the first material may
comprise more than one of TiNx, TiAlNx, TiAlYCrNxand
CrNx and the second material may comprise one or both
of MoSz and WSZ .
The composition according to all of the
embodiments of the present invention will generally
have a friction coefficient of from 0.08 to 0.3, more
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generally from 0.08 to 0.2, and a hardness typically
of at least 14 GPa.
In a third preferred embodiment of the present
invention, the composition comprises:
(i) a first material selected from TiNx, TiAINX,
TiAlYCrNx and CrNX, and
(ii) a second material selected from MoSZ and WSZ,
wherein x is equal to or approximately equal to 1, z
is equal to or approximately equal to 1.6, and the
atomic ratio (y) of Mo or W to Ti or Cr is from 0.2 to
0.4, and wherein the composition has a friction
coefficient of from 0.08 to 0.2, preferably from 0.08
to 0.12, and a hardness of from 16 to 22 GPa,
preferably from 18 to 22 GPa, more preferably from 20
to 22 GPa.
In a fourth preferred embodiment of the present
invention, the composition comprises:
(i) a first material selected from TiNx, TiAINX,
TiAlYCrNx and CrNx, and
(ii) a second material selected from MoSZ and WSZ,
wherein x is equal to or approximately equal to 1, z
is equal to or approximately equal to 1.1, and the
atomic ratio (y} of Mo or W to Ti or Cr is between 0.1
and 0.2, and wherein the composition has a friction
coefficient of from 0.08 to 0.2, preferably from 0.08
to 0.12, and a hardness of from 22 to 32 GPa,
preferably from 25 to 30 GPa, more preferably
approximately equal to 27 GPa.
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The microstructure of a coating according to the
present invention typically comprises a soft,
lubricating phase of MoSZ or WSZ in a hard matrix phase
of TiNx, TiAlNx, TiAlYCrNx or CrNx, with nanometre
grain size, typically of from 1 to 10 nm. The
presence of a 'lubricant reservoir' substantially
throughout the coating thickness allows lubrication to
be maintained even as the coating wears away. An
additional advantage arises because the MoSZ or WSZ,
phase is incorporated within a hard matrix and this
provides a degree of oxidation protection.
In the present invention, TiN, TiAlN, TiAlYCrN or
CrN are used as a hard phase because, in addition to
their outstanding mechanical properties (superhardness
and high toughness), these nitrides display relatively
good tribological properties compared with other hard
coatings. For example, TiN has a friction coefficient
typically quoted between 0.4 and 0.8 against steel, or
even less than 0.2 when rubbing against itself or
certain hard counter-faces, such as A1203.
The present invention also provides a wear-
resistant, self-lubricating coating comprising a
composition as herein described and, additionally, an
article having such a coating.
The present invention also provides a cutting
tool comprising a substrate having one or more
coatings thereon, wherein at least one of the said
coatings comprises a composition as herein described.
It will be appreciated that the substrate may have
more than one such coating, for example a first
coating of TiNx and MoSZ and a second coating of, for
example, TiAlNx and WSZ. It will alsa be appreciated
that compositional variations may exist within each
coating. For example, a gradient coating may be
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applied comprising TiNX and MoS=, wherein the MoSZ
content is zero (or at least approaching zero) at the
substrate surface and is continuously increased with
distance into the coating (and away from the
substrate) until the required level is reached.
The substrate material may be formed from any of
the conventional cutting tool materials, such as, for
example, a hard metal, a high speed steel or a cermet.
A layer consisting essentially of Ti, TiAl,
TiAlYCr, Cr, TiN, TiAlN, TiAlyCrN or CrN is
advantageously disposed between a surface of the
substrate and the one or more coatings, since this has
been found to improve adhesion between the coatings)
and the substrate. The choice of the material for
this intermediate layer is advantageously based on the
nitride constituent of the coating. For example, if
the coating comprises TiNX, then the intermediate
layer would preferably be formed from Ti or TiN. In a
preferred aspect, a first layer is provided consisting
essentially of Ti, TiAl, TiAlYCr or Cr and this layer
is disposed adjacent a surface of the substrate, and a
second layer consisting essentially of a nitride of
the material of the first layer, i.e. TiN, TiAlN,
TiAlYCrN or CrN, is disposed between the first layer
and the one or more coatings as herein described.
Again, the choice of the material for the first layer
is advantageously based on the nitride constituent of
the coating. For example, if the coating comprises
TiAINX, then the first layer would preferably be
formed from an alloy of Ti and A1. In this manner a
three-component adhesion-promoting underlayer may
advantageously be used. For example, a first layer of
Ti may be deposited on to the surface of the
substrate. Next, a layer of TiN may be deposited on
to the first layer. Next, a gradient layer of TiN and
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MoSz may be deposited, starting off with 0~ MoSZ (or at
least approaching zero) and then continuously
increasing the MoSz content to that of the functional
TiNx(MoSz)y self-lubricating layer as herein described.
Whilst the cutting tool according to the present
invention may be used with a lubricant and/or a
cooling fluid, it is particularly suited to dry-
machining applications, i.e. no lubricant or cooling
fluid, or micro-lubrication applications, where very
little lubricant and/or cooling fluid is/are required.
The present invention also provides for the use
of a composition as herein described as a low-
friction, wear-resistant coating. Whilst the coating
composition will generally be used in the manufacture
of cutting tools, it will be appreciated that it can
be used in any area requiring low-friction and good
wear-resistance.
In a another aspect of the present invention,
there is provided a process for forming a wear
resistant, low friction coating on a substrate, which
process comprises the steps of:
(i) providing a substrate
(ii) placing the substrate in a deposition chamber
having means to deposit first and second deposition
materials simultaneously on to a surface of the
substrate, wherein the first deposition material is
selected from TiN, TiAlN, TiAlYCrN and CrN and the
second deposition material is selected from MoS2 and
WS2;
(iii) co-depositing the first and second deposition
materials on to the surface of the substrate under
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conditions to form a coating having a composition as
herein described.
The deposition chamber may comprise a sputtering
system, for example an unbalanced do magnetron
sputtering system, in which the means to deposit the
first and second deposition materials simultaneously
preferably includes a target comprising the first
deposition material and a target comprising the second
deposition material. Alternatively, a composite
target may be used. In this way, Tin and MoS2, for
example, can be co-sputtered on to the surface of the
substrate. The composite target may take the form of
a mosaic target having, for example, one or more TiN
portions adjacent one or more MaS2 portions.
Magnetron sputter ion plating is described in detail
in GB-2 258 343.
The substrate is typically maintained at a
temperature of from 80 to 500°C, preferably from 100
to 300°C, more preferably from 150 to 250°C and still
more preferably approximately 200°C during deposition.
Deposition is preferably carried out in a vacuum or an
inert gas atmosphere. For example, an argon
atmosphere at a pressure of from 0.1 to 1 Pa Ar,
preferably from 0.4 to 1 Pa, has been found to be
suitable.
Deposition is generally carried out using a power
density typically in the range of from 5 to 20 Wcm2,
preferably from 5 to 15 Wcm2. Experiments have been
conducted using a target current of from 0.5 to 1.5 A
(equivalent to a target voltage of from 250 to 800 V)
and preferably approximately 1 A (equivalent to a
target voltage of approximately 520 V).
A coating thickness of typically from 0.5 to 5 ~m
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can be produced in less than an hour.
A bias voltage of from 0 to -200 V, preferably
from -50 to -150 V, more preferably from -75 to -125
V, is typically applied to the substrate during
deposition. A bias voltage of approximately -100 V
has been found to provide the best results.
Alternatively, it may be envisaged to deposit the
hard matrix, TiN for example, reactively by
introducing a sufficient quantity of nitrogen into the
process gas whilst sputtering from a Ti target. It
may also be envisaged to reactively deposit the MoSZ,
for example, by reactive deposition using a Mo target
and the addition of a sulphur containing process gas
such as HzS .
Alternatively, an arc/sputter hybrid device may
be used to deposit, for example TiN and MoS2
simultaneously on the surface of the substrate. In
this manner, the hard phase of, for example, TiN, is
deposited from a titanium cathodic arc and MoS2, for
example, is simultaneously deposited by magnetron
sputtering. In this example, it is necessary to use a
mixed nitrogen/argon mixture. The nitrogen allows the
reactive deposition of TiN in conjunction with the Ti
arc source whilst the argon provides an efficient
sputtering gas for the MoS2. In effect, nitrogen is
not particularly effective for the sputtering of MoS2.
Furthermore, any other deposition process which
would allow the deposition of the first deposition
material in conjunction with magnetron sputtering of
the second deposition material, such as plasma-
assisted CVD or photon-induced deposition, may be
used.
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When a cathodic arc device is used to deposit the
hard phase, for example TiN, in conjunction with a
magnetron sputtering device to deposit the lubricating
material, for example MoS2, the arc current should be
adjusted so as to obtain a suitable TiN deposition
rate with respect to the MoS2 deposition rate in order
to produce a coating within the given composition.
Typically, the arc current is at least 50 A. In this
embodiment of the process, the pressure inside the
deposition chamber is provided by the nitrogen and
argon process gases. A pressure of from 0.1 to 1 Pa
is preferred with the volumic flow ratio of nitrogen
to argon preferably being approximately 10:1, so as to
provide a suitable compromise between the reactive
deposition of TiN with nitrogen and the sputtering of
MoS2 using argon. It will be appreciated by those
skilled in the art that whilst magnetron sputter
deposition is currently the most common method for
depositing MoSz, as arc technology improves, it may
also be possible to deposit the MoS2 using an MoS2 arc
source. Alternatively, it may also be possible to
deposit MoS2 reactively using a Mo arc source and a
sulphur containing process gas, such as HzS.
The present invention also provides a process for
machining, for example milling, a cast iron component,
which process comprises the step of machining a
component using a cutting tool as herein described.
The step of machining may be carried out in the
absence of a lubricant or a cooling fluid (dry
machining) or, alternatively, using a minimum amount
of a lubricant and/or a cooling fluid (micro-
lubrication).
It will be appreciated that machining may be
applied to a variety of materials as well as cast iron
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components, for example carbon steels, stainless
steels, aluminium, copper, and including alloys
thereof.
The presence of a lubricant reservoir
substantially throughout the coating thickness allows
lubrication to be maintained even as the coating wears
away.
The present invention will now be described
further, by way of example, with reference to the
following drawings, in which:
Figure 1 is a schematic representation of a co-
deposition process according to the present invention;
Figure 2 shows the variation of atomic ratios for
the overall composition TiNx(MoSZ)Y with substrate
position and substrate bias as measured by GDOES (x
and y), EDX (y and z) and XPS (x and z).
Figures 3(a) and (b) shows SEM cross-sections for
coatings deposited on Mo substrates in positions -30
at 0 (a) and -100 V (b) .
Figures 4(a) and (b) shows the variation in X-ray
diffraction spectra with substrate position for 0 V
bias (a) and -100 V bias (b). Each spectrum is
labelled with substrate position in mm.
Figure 5 shows the variation in hardness and
friction coefficient as a function of substrate
position for 0 V and -100 V bias. Friction values are
taken over the whole sliding distance, except for
biased coatings produced in positions -40 to -65, for
which the first 50 m of unstable friction are excluded
from the average calculation.
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Figure 6 are friction curves showing a typical
behaviour for biased coatings in positions -40 to -65,
characterised by an initial maximum in friction
(curves a, b and c), and friction behaviour typical of
all other coatings, characterised by practically
featureless low friction (curve d).
Figure 7 is a schematic top view of a deposition
chamber according to the present invention;
Figure 8 shows room temperature pin-on-disk test
results for TiN-MoSZ coatings as compared to TiN
standard coatings using either 6 mm diameter 100Cr6
steel or A1203 balls as counterfaces.
Figure 9 shows comparative lifetime test results
for uncoated HSS drills, TiN-coated drills and TiN-
MoS= coated drills.
Film deposition may be performed using an
unbalanced do magnetron sputtering system 1,
schematically represented in Figure 1 and described in
detail by P. Losbichler and C. Mitterer in Surf. Coat.
Technol., 97 (1997), 568-574. A composite TiN-MoS2
target 5 was used, made up of two halves cut from 150
mm diameter TiN and MoS2 targets, 6 and 7 respectively,
which were bonded to a water-cooled backing plate 8.
Substrates 10 were mechanically polished and degreased
stainless steel rectangles (10 x 20 mm). The
substrates 10 were fixed to a holder 11, located 6 cm
above the target 5 and placed at various positions
between -75 and + 75 mm (this number refers to the
horizontal distance between the centre of the
substrate 10 and the target's 5 central dividing line,
negative on the TiN side and positive on the MoS2
side). The chamber 2 was pumped down to a base
pressure of 10-' Pa with a turbo-molecular pump 3.
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Prior to film deposition, the target 5 was sputter
cleaned for 5 minutes behind a shutter 13 and
substrates 10 were sputter etched for 20 minutes with
an Ar pressure of 3.5 Pa and a do voltage of -1500 V.
Deposition was then carried out, with the substrate 10
temperature maintained at 200°C by a resistance heater
12, using 0.7 Pa Ar pressure, 1 A target 5 current
(around 520 V) and either 0 or -100 V substrate 10
bias. Deposition time was 45 minutes and resulted in
a coating thickness in the 0.5 to 3 ~cm range,
depending on substrate 10 position.
The overall coating chemistry and the phase
composition was investigated by XPS on a Cameca-
Nanoscan 50, incorporating a MAC2 semi-imaging
analyser, set at an energy resolution of 0.5 eV. The
unmonochromated MgKa source was operated at 12 kV and
30 mA. A rastered 3 keV, 0.2 /.c,A Ara ion beam was
employed to remove the surface oxide, the 0 is peak
being monitored until it no longer decreased. The
energy scale was calibrated using the Cu 2p3,2 and Au
4f"2 peaks at 932.67 and 83.98 eV respectively.
Quantification of the spectra was performed using
relative sensitivity factors determined from Ti Mo,
TiN and MoS2 standards. Due to the preferential
sputtering of S, the S content was estimated from the
Mo 3d5,2 (sulphide) peak position in the spectra from
the native surface.
Coating composition was also determined using
Glow Discharge Optical Emission Spectroscopy (GDOES),
calibrated using TiN and alloyed Mo-Ti standards, and
EDX analysis, X-ray microstructural analysis of the
samples was performed by Glancing Angle X-ray
Diffraction (GAXRD) using an unmonochromated copper
source at an incident angle of 1.0°, a high precision
beam collimation and sample positioning system, and a
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solid state detector to maximise the signal/noise
ratio.
Film morphology was studied using a Cambridge
Instruments Stereoscan 360 scanning electron
microscope (SEM).
Hardness and Young's modulus were determined by
an ultra-low load depth-sensing nanoindenter
(Nanoindenter II, from Nano Instruments Inc.)
described in detail by R. Gilmore, M.A. Baker; P.N.
Ginbson and W. Gissler in Surf. Coat, Technol. (in
press). A pin-on-disk tribometer (CSEM) was used to
determine the friction coefficient. Counterface
material was 6 mm diameter steel ball and track radius
was 3.5 mm. A normal load of 1 N and a sliding speed
of 0.05 m/s were used over a sliding distance of 250 m
in laboratory air and at an ambient temperature (24°C
tl) .
While not wishing to be constrained by theory,
the following comments are made in arr attempt to throw
some light on the results.
Coating composition as a function of substrate
position and bias voltage is shown in Figure 2.
Because coatings from position +20 to +75 tended to
rapidly spall off once removed from the vacuum
chamber, characterisation results are available for
substrate positions -75 to +15 only. The MoSZ content
decreases progressively from a maximum value of
approximately 66~ at position +15 to a lower Limit of
around 18~ at position -75. The MoSZ content is
calculated as 100 x 1/(1 + y) for the overall
composition TiNx(MoSZ)Y. The TiN is generally
stoichiometric with only very slight variation with
substrate position and no measurable effect due to
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substrate bias. The MoS2 phase is slightly sub-
stoichiometric at about MoSl,s, showing relatively
little variation with substrate position or bias.
SEM cross sections of selected coatings
(corresponding to max hardness) are shown in Figures
3(a) and (b) for 0 V and -100 V bias. The films
appear to be, dense with good adhesion and a fine
grained structure.
Figure 4 shows GAXRD spectra for the various
coatings with superimposed MoSZ and TiNx phases to
separate. Where easily observed, the TiN phase is
obviously nanocrystalline, and it shaws a more ordered
structure under the TiN half of the target (from
positions -15 to -75) with an average crystallite size
of the order of several nm. There is a significant
shift of the TiN peaks to lower angles for all
samples, representing an increase in lattice
parameter. This might be attributable to the
incorporation of Mo and/or S in the cubic structure,
though the observed shift is more probably caused by N
being substituted by S, since it is known that Mo
substitution does not produce a large change in the
lattice parameter. There appears to be a rapid
widening of the peak width from position -45 to
position -25. This is assumed to be due to a rapid
decrease of the TiN grain size owing to competition
with the formation of extended MoSZ zones which
becomes more likely the nearer the substrate is to the
MoS2 half of the target. A diffraction spectrum
typical of sputter-deposited MoS2 is evident at
position -15. The structure of the MoS2 which gives
rise to such a spectrum has been described as
consisting of randomly stacked 001 planes with
fluctuations of the distance in the c direction
between these planes. There is very little extension
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of the structure in the a-b direction. As such, this
represents neither a proper nanocrystalline nor a real
amorphous structure. It is likely that in such a
structure a considerable concentration of Ti and N
atoms could be accommodated, and this may further
account for the sudden reduction in the TiN
crystallite size which is observed in the GAXRD
spectrum when the MoS2 structure starts to form, since
Ti or N atoms arriving at MoS2 growth zones will
easily become trapped there. For coatings of low MoSx
content, there is a lack of any obvious MoSz 002 peak.
The shape and XPS binding energies of the Ti
2p3,2, N ls, S 2p3,2 and Mo 3d5,2 peaks for the TiN and
MoS2 standards (respectively 455.3, 397.4, 162.1 and
229.2 eV) are in good agreement with the corresponding
average peak positions for all of the co-sputtered
samples (respectively 455.3 ~ 0.1, 397.3 ~ 0.1, 162.1
~ 0.1 and 228.9 ~ 0.15 eV), indicating in general the
presence of separate TiNx and MoSZ phases, also
indicated by the GAXRD results. The lower XPS Mo peak
position reflects the sub-stoichiometry of the MoS2
phase. The XPS spectra do not allow confirmation of
the suspected substitution of Mo and/or S into the TiN
lattice suggested by the GAXRD results possibly
because of peak overlapping for S in the form of MoS2
and substituted in the TiN lattice. As a reference
for the S 2p3,2 peak position, TiS has a quoted binding
energy of 163.5 eV, but has an h.c.p. structure.
Possibly a better reference is f.c.c. MnS (Mn having
an electronegativity very similar to Ti) for which the
binding energy is quoted as 162.0 eV, which would lead
to a peak overlapping that of MoS2. Even for the
lowest MoSZ contents (18~), it seems reasonable to
conclude that an MoSz phase exists since only a small
part could substitute into the TiN lattice. At first
sight, this observation appears to be in contradiction
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with the corresponding GAXRD spectra which shows a
lack of any obvious 002 peak at position -45 and
positions further from the MoS2 half, despite the high
scattering power of Mo. One possible explanation of
this would be a complete loss of correlation in the c-
direction of the MoSZ structure and this may mean that
the MoSZ, which is not incorporated in the TiN
structure, is reduced to single layers, presumably
located between the TiN grains.
ZO
Figure 5 shows the hardness and friction results.
Hardness passes through a maximum around position -40,
for both the biased and non-biased coatings with
maximum average values of about 20 and 17 GPa
respectively. The increase in hardness between
position +15 and -90 can be attributed to the
improving crystal~structure of the TiN phase as the
MoSz content diminishes. The systematically higher
hardness values in the case of -100 V bias for
positions left of 20 may be attributable to the
optimum deposition conditions for the TiN phase. It
is, however, not clear why the hardness then
diminishes for positions to the left of -40. It would
rather be expected that hardness continues to increase
or stabilise as MoSZ content decreases. It does not
appear to be an artefact of the nanoindentation
technique which may start to measure substrate
hardness as coating thickness falls below 10~ of
penetration depth: the results obtained with 200 nm
penetration depth were cross-checked with 50 nm
results and were in good agreement. It may be that
the fall in hardness is principally related to some
microstructural change due to changing deposition
parameters close to the target's edge related to some
shadowing at the target's extremity.
Friction remained low for all studied
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compositions. Low friction was obtained despite MoSZ
under-stoichiometry and an apparently random basal
plane orientation, though this is not in contradiction
with the literature, where it is reported that
stoichiometries above approximately 1.2 are usually
lubricating, and MoS2 is able to preferentially re-
orient under the action of friction with basal planes
parallel to sliding direction. The average friction
coefficient was generally close to 0.1 and showed no
direct correlation with hardness. Friction curves
were typically flat (Figure 6(d)) with the notable
exception of the biased coatings deposited in
positions -40 to -65. For these coatings, the
friction coefficient passed through a large and
distinctive initial maximum as high as 1.1 (Figures
6(a-c)). The friction coefficient then fell to around
0.2 after about 50 m and remained relatively stable
for the remaining 200 m of the test. For samples
displaying this atypical behaviour, the friction
values shown in Figure 5 are averages taken over the
final 200 m to exclude the initial instabilities. The
distinctive shape of these curves were reproducible
and thought perhaps to be of some significance.
Initial friction maxima are often associated with
transfer film formation and the resemblance to
friction curves reported for TiN rubbing against steel
under fretting conditions is significant. It may be
that the observed friction maximum corresponds to the
formation of a third body from the TiN coating which
is transformed under the action of friction to a sub-
stoichiometric form of rutile, resulting in the
observed low final friction. It may be that for our
biased samples produced in positions -40 to -65, which
have the best formed TiN lattice and highest hardness,
friction is governed by the TiN phase in the initial
stages until the tribo-assisted formation of a
lubricious oxide transfer layer allows the friction
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coefficient to return to a relatively low value.
However, this behaviour apparently disappears at even
lower MoS2 contents, possibly due to the fact that the
hardness drops away.
The comments above are provided in an attempt to
explain the results and are not intended to constrain
the present invention by theory.
In another example, film deposition was performed
using the system schematically represented in Figure
7. The deposition facility 20 was composed of two
opposing arc evaporation Ti sources 25 and two
opposing MoS2 magnetron targets 30, which could be
operated independently. The samples (drills and flat
substances) were positioned in a carousel-type sample
holder 35 with three rotational axes. The drills were
of the type HSS DIN 338 of 6 mm diameter and 100 mm
length and the flat samples were hard metal discs of
diameter 23 mm. The flat samples were later used for
mechanical and structural characterisation (such as
glancing angle X-ray diffractometry, EDX,
nanoindentation, pin-on-disk testing, scratch testing)
whilst the drills were used for field tests.
Samples were cleaned in an ultrasonic bath in a
benzene/alcohol solution. Before deposition, the
chamber was evacuated to a residual pressure of 5x10-5
mbar and then heated to 250°C for 30 minutes using
infra red heaters 40, 41 and 42. Coating was then
performed as follows.
i) An ion etching process was performed in a
substantially pure Ar atmosphere at a pressure of
approximately 0.01 mbar.
ii) To enhance adhesion, a ~ 5 nm Ti layer was first
deposited, also in a pure 0.01 mbar Ar
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atmosphere, followed by a ~ 100 nm TiN coating
deposited reactively in a substantially pure
nitrogen atmosphere of approximately 0.001 mbar
using approximately 70 A on the Ti arc source.
iii) Then the MoS2 sputter source was operated in an
Ar/NZ atmosphere using flow rates of
approximately 50 and 500 standard cubic
centimetres per minute respectively to maintain
the pressure at approximately 0.01 mbar. The
power to the MoS2 sputter target was
progressively increased from approximately 300 to
1500 Watts whilst decreasing the Ti arc current
to approximately 55 A, to create a 300 nm
gradient layer of increasing MoS., content.
iv) Finally a coating of about three microns was
produced by maintaining the Ti arc current and
sputter power at approximately 5.5 A and about
1500 Watts respectively. The bias voltage on the
carousel was held at approximately -100 volts.
For comparison, pure TiN coatings were also
prepared using steps i) and ii) above, whereby TiN
deposition was continued so as to produce a 3 micron
coating.
The chemical composition of the co-deposited
coatings was determined using EDX analysis and found
to be (TiN) o.e, (MoSZ) 0.13 wherein z is in the range of
0.8 < z < 1.4.
The coating was found to display a f.c.c.
structure typical for TiN with a slightly enlarged
lattice parameter as determined by Glancing Angle X-
Ray Diffraction.
The hardness of the coatings was determined to be
about 27 GPa with a nanoindenter, approximately the
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same as that measured for the pure TiN coatings. The
nanoindenter was calibrated with a Si 111 wafer,
assuming for this material a hardness of 11 GPa.
The adhesion was determined with a scratch
tester. A critical load value of 105 N was found.
The friction coefficient was measured with a pin-
on-disk tester using a steel (100Cr6) or A1203 sphere
of diameter 6 mm as a counterface at a sliding speed
of 0.1 m/s (track radius 8 mm) and a load of 5 N at
room temperature and with a relative humidity of about
40~. Figure 8 displays the friction coefficient of a
TiN and a TiN-MoS2 coating as function of the
distance. After a short running-in phase the friction
coefficient of the TiN-MoSz coating assumes a constant
value of less than 0.2 for both types of counterface
in comparison to 0.8 for the TiN coating.
All drills have been subjected to a field test
under dry machining conditions on a CNC milling
machine of type Deckel P2A by drilling holes under the
following conditions:
feed rate = 0.32 mm/rev,
speed = 1600 rpm,
hole depth = 27 mm,
work piece = carbon steel C35.
The lifetime was determined by the number of
holes drilled before the drill broke. In order to
facilitate the running-in phase the first five holes
were drilled with a depth of 20 mm. The results are
shown in Figure 9, where the lifetimes of several TiN-
MoS2 coated drills are compared with uncoated drills
and TiN-coated drills. The expected increase in
lifetime is observed for the TiN coated drills with
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respect to the uncoated drills by a factor of
approximately four. A further increase in lifetime by
about 25~ was observed for the TiN-MoSz coated drills.
The present invention provides a process for
producing dense, well adhering coatings which can
combine a hardness exceeding about 20 GPa with a
friction coefficient of less than about 0.2.
Optimum hardness/friction properties were
obtained for films deposited using approximately -100
V substrate bias.
The TiN-MoS2 system has potential for producing
well-adhering films comparable with standard TiN tool
coatings by the use of a TiN underlayer then graded
interlayer.
