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192
ROLE
OF
GRAIN SIZE IN SCRATCH DAMAGE RESISTANCE IN
ZIRCONIAS AND SILICON NITRIDES
Seung Kun
Lee,
Robert
P.
Jensen, Michael
J.
Readey
Advanced Materials Technology
Caterpillar Inc.,
USA
ABSTRACT
Scratch damage
in
zirconias (Mg-PSZ, Y-TZP, and Ce-
TZP) and silicon nitrides (fine, medium, coarse grain)
are investigated. Scratch testing is carried out using a
conical diamond indenter. In all materials the damage
mode changes from smooth plastic deformation to
limited cracking with increasing scratch load: in Mg-
PSZ, plastic deformation is predominant at lower
loads, with microcracking at higher loads; in Y-TZP,
plastic deformation is predominant over the range of
the test loads-macrocracks initiate only at relatively
high loads, but penetrate to a relatively large depth;
again, Ce-TZP shows intermediate behavior, but with
cracking patterns closer to that of Mg-PSZ. Bending
tests on specimens subjected to scratch damage
indicate a relatively high damage tolerance in the Mg-
PSZ and Ce-TZP; Y-TZP shows the highest initial
strength, but suffers relatively large strength loss above
the critical load for cracking. In Y-TZP,
fine
Y-TZP
showed smooth scratch groove without cracking.
As
grain size increased, the scratch groove became
rougher with the development of microcracking. In
Si3N4, the critical load for a transition from plastic
deforamtion to microscracking decreased with
increasing grain size. Coarse Si3N4 has a higher
material removal rate than fine Si3N4. The material
removal mechanism is attributed
to
grain pullout in this
material. Implications concerning relative merits of
each material for wear properties, contact fatigue, and
machining damage are briefly discussed.
INTRODUCTION
Advanced ceramics such as zirconia and silicon nitride
have been identified as choices for sliding components
in
a variety of engineering applications, including
engine components (bearings, rollers, dies, tappets,
valves, fuel injectors), where contact, scratch, and wear
damage are critical factors for lifetime performance
[
1
-
41.
In zirconia ceramics, energy absorptive phase
transformation from metastable tetragonal phase to
monoclinic phase enhances fracture [5-91. Much
progress has been made during the past 25 years
in
this
kind of toughness enhancement
in
zirconias, by
microstructural control through modification of
selective additive phases (MgO, Y203, CeO) and heat
treatments
[
10-
131. In silicon nitride ceramics, on the
other hand, there has been a great success
in
improving
strength and toughness by controlling composition and
microstructure over the decade. Self-reinforced silicon
nitride with elongated grain structure provides an
excellent mechanical-reliability associated with high
toughness and strength
[
14,151. Properties like fatigue
[16] and machining and wear [17,18] have been
demonstrated to be strongly affected by such
microstructural factors.
Scratch testing using a translating sharp point provides
more direct information on potential processes
in
wear
and machining operations [17,18,19]. Data from this
test procedure may be expected to provide
complementary basic information on damage modes
in
engineering contact and sliding applications, as well as
on material removal
in
wear and machining and flaw
development in strength determination.
In this study we investigate scratch damage
in
selected
zirconia ceramics (Mg-PSZ, fine grained Y-TZP,
medium Y-TZP, coarse Y-TZP, and Ce-TZP) and
silicon nitride ceramics (fine grained Si3N4, medium
Si3N4, coarse Si3N4). The results demonstrate a need
for caution
in
materials selection for applications
in
severe contact conditions.
EXPERIMENTAL PROCEDURE
Materials
Commercially available zirconias and silicon nitrides
were used
in
this study.
In order to investigate effect
of grain size, the three types of Y-TZP ceramics with
different grain size were made by a routine ceramic
processing. The zirconia powders (3Y-TZP, Tosoh Co.
Tokyo, Japan) were compacted
in
a steel die of 50
mm
diameter at a pressure of 50 MPa, followed by sintering
at either 1450°C for
1
h, 150O0C for 2 h,
or
1600 OC for
3
h.
These heat treatments produced the “fine” (F-Y-
TZP), “medium” (M-Y-TZP), and “coarse” (C-Y-TZP)
microstructures. Microstructures of zirconias and
silicon nitrides are presented
in
Fig.1 and 2,
respectively.
Mechanical properties of zirconias and
silicon nitrides investigated
in
this study are
summarized
in
Table
1
and
2,
respectively. Mg-PSZ
had a large grain structure (grain size
=
58
pm)
and
higher toughness but lower strength (Table 1). Y-TZP
had a fine microstructure (grain size
=
0.4
pm)
and
higher strength and hardness but lower toughness. Ce-
showed intermediate toughness and lowest strength.
CFI3208 has a fine grain structure with
a
lower
toughness.
AS800 showed a coarse microstructure
with large elongated grains.
GS44
exhibited a bimodal
grain structure with high strength and toughness.
193
Table
1.
Characteristics
of
zirconia ceramics
Materials
Suppliers
Second Phase Grain size
Strength Toughness
CFI3208 Ceramic
for
Al,
Y
Fine (0.5)
730
f
34
5
.O
(P) (MPa) (MPa-m'n)
Industries
Hardness
14.8
(GPa)
GS44
AS800
Fig.2. Microstructures
of
CFI3208 (a), GS44 (c), and
AS800 silicon nitride (c).
Fig.
1.
Microstructures
of
Mg-PSZ
(a),
Ce-TZP
(b),
and Y-TZP (c).
Honeywell Al,
Y
Medium
(I.
1)
959
k
44
7.5 14.8
Honeywell
La
Coarse (3.2)
707
+_
39 7.2 14.2
194
Fig
1.
Top
and
side views
of
scratch damage
of
Mg-PSZ (a), Y-TZP
(b)
and
Ce-TZP
(c),
scratch load,
P
=
I00
N.
Scratch tests
Scratches were made using a sliding conical diamond
indenter with apex angle 120' and spherical tip radius
200
pm
(Automatic Scratch Tester, CSEM-
REVETEST, Neuchatel, Switzerland). In automatic
test mode, the load is increased continuously as
the
point translates across the surface. Both normal and
tangential forces were recorded. Normal loads ranged
from P
=
5
N to 130 N over about 20
mm
sliding
distance, at sliding speed 20 mdmin,
in
air.
Cross-
sectional area of the scratches were measured by a
surface profilometer, to obtain an estimate of the
volume removed per unit sliding distance. Some
additional scratch tests were made on selected
specimens, at prescribed constant loads.
Bonded-interface specimens
[
16,18,19,20] were used
to obtain section views through the indentation and
scratch damage zones
in
each material. Specimens
were cut into two half-blocks. The side surfaces of the
half-blocks were first polished and then clamped face-
to-face with an intervening
thin
layer of adhesive. The
top surfaces were then repolished. Indentations were
made along the surface traces of the bonded interfaces,
at load
P
=
3000 N. Scratch tests were made
perpendicular to these interfaces, at load
P
=
100
N.
The adhesive joining the interfaces was subsequently
dissolved
in
acetone. Separated half-blocks were gold
coated for top- and side-surface examination
in
Nomarski
illumination.
Four-point bending tests were conducted on bars 3
x
4
x
45
mm
(outer span
40
mm, inner span 20
mm)
that
had been subjected to scratch damage. The scratches
were made on the prospective tensile surfaces, with the
scratch axis parallel to the prospective bend axis, at
loads
P
=
20N
to
130
N.
The damage areas were
covered with a drop of dry silicon oil before flexure
and broken
in
fast fracture (<I0 ms), to minimize
environmental effects
in
the strength data.
All
broken
specimens were examined fractographically to locate
the source of failure, either scratch damage
or
extraneous flaws. Control strength tests were made on
unscratched specimens, to determine baseline
laboratory strengths.
RESULTS
Zirconia ceramics
Figure
1
shows surface and side views of scratch
damage of the three materials from bonded-interface
specimens at
P
=
100
N. In Mg-PSZ (Fig.
1
a) a diffuse
microcrack zone extends well below the surface groove
into the subsurface. In the near-surface regions, these
microcracks have coalesced to produce extensive
chipping. In Y-TZP (Fig. Ib) no microcracking is
observed at all: instead, a deeply penetrating median
crack is evident immediately below the scratch. In Ce-
TZP
(Fig. lc) limited microcracking is observed, on a
smaller scale than
in
Mg-PSZ.
Figure
2
is a plot
of
volume removed per
unit
sliding
distance as a function of normal load for the three
materials, evaluated from surface profilometer traces
across the scratches. Mg-PSZ ultimately shows the
most severe wear rate, above load
75
N where surface
microcracking and chipping are manifest. Y-TZP
shows the lowest removal rate, over the entire range of
loads.
Figure
3
compares strength degradation data for
scratch damage on the three materials. In the Mg-PSZ
and Ce-TZP, strength falls off beyond the load for
extensive microcracking, but only slowly, indicating
high damage tolerance. Y-TZP, on the other hand,
show a typical brittle response; i.e. no perceptible
degradation
in
initially high strength up to about
80
N,
but with subsequent abrupt drop-off at higher loads,
195
highlighting the effectiveness of the median crack in
Fig. 3c.
I
I
I
I
Samh
load
(N)
Fig 2. Volume removed for Mg-PSZ, Ce-TZP, and Y-TZP,
as a function of load. Solid lines are empirical fits.
1200,
I
I
I
I
I
I
Fig.
3.
Strength
in
scratching
of
Mg-PSZ, Y-TZP and Ce-
TZP. Closed symbols represent failures from scratch origins,
open symbols failures from other flaw origins; shaded boxes
at left are strengths of as-polished specimens.
Fig.
4.
Surface views of scratch damage
in
in
F-Y-TZP (a),
M-Y-TZP (b), and C-Y-TZP (c), produced at a scratch load
P
Figure
4
presents surface views of scratch damage
in
F-Y-TZP, M-Y-TZP, and C-Y-TZP, produced at a
=
l00N.
scratch load
P
=
100
N.
F-Y-TZP shows a smother
surface track with
no
microcracking, indicative of
plastic deformation.
In
C-Y-TPZ,
on
the other hand,
extensive microcracking and larger debris (see arrows
in Fig. 4c) are evident around the scratch groove,
indicating brittle response in this material. M-Y-TZP
shows an intermediate feature. Scratch damage mode
depends strongly
on
grain size. Smaller grain size is
responsible for the smoother scratch track associated
mainly with plastic deformation in F-Y-TZP.
Fig
5.
SEM micrographs of surface views of scratch damage
in
CF13208 (a),
GS44
(b), and
AS800
silicon nitride (c).
196
Scratch
load
(N)
Fig.
6.
Material removal rate as function of scratch
load
in
silicon nitrides.
Silicon nitride ceramics
Figure
5
shows SEM micrographs of the surface views
of scratch damage
in
CF13208, GS44, and AS800
silicon nitrides. CFI3208 with a fine grain structure
exhibited a smooth scratch groove with less cracking.
AS800 with a coarse grain structure exhibited a rough
scratch track with grain pullout and dislodgement.
Microcracking and chipping are manifest
in
AS800,
indicating intensive interaction between material and
an indenter. GS44 showed intermediate cracking
behavior.
Figure
6
is a plot of volume removed per
unit
sliding
distance as a function of normal load for CH3208,
GS44, and
AS800,
evaluated from surface profilometer
traces across the scratches. AS800 ultimately shows
the most severe wear rate. CFI3208 and GS44 show
lower removal rates, consistant with scratch groove
in
Fig.
5.
DISCUSSION
We investigated the effect of grain size on scratch
damage mode
in
zirconias and silicon nitrides. Scratch
test techniques provide useful information to determine
scratch damage and to simulate machining damage
[
17,18,2
I].
It is also relevant for evaluating abrasive
wear resistance. Mg-PSZ
with
large grain size
(>50
pm)
shows the least smooth scratch track, with
extensive microcracking and associated chipping above
a critical load, resulting
in
a high wear rate at higher
loads. Y-TZP shows the smoothest track, and lowest
wear rate. Ce-TZP is again intermediate,
with
modest
microcracking.
Y-TZP
is the most brittle
of
the three
materials, as reflected by the formation of a deeply
penetrating median crack
in
the subsurface region
above a critical load. Accordingly, Y-TZP is much
more susceptible to abrupt strength
loss
in
severe
scratching conditions; its strength drops from almost
twice to less than one half that of its counterparts over
the data range. Mg-PSZ and Ce-TZP, on the other
hand, while still subject to failure from scratch sites
above a threshold load, are much more damage
tolerant.
In Y-TZP, as grain size increased, the damage mode
in
Y-TZP changed from plastic deformation to
microcracking. A similar trend has been reported on
alumina
[
181.
This result indicates the important
implications on abrasive wear and machining. Grain
size plays a major role
in
the properties determined at
the microstructural level such as scratch, wear, and
machining. Large grains lead to a rougher scratch
track and a higher material removal rate. Again, high-
tougheness zirconia such as Mg-PSZ and C-Y-TZP
would not necessarily enhance wear properties, and
may even degrade them. F-Y-TZP with fine grains is
likely to exhibit better performance where high wear
resistance is an issue.
In silicon nitride, as grain size increased, scratch
damage mode changed from plastic deformation to
brittle microcracking. Grain pullout is evident
in
AS800, indicating severe scratch damage. AS800 with
coarse gains is expected to be less abrasive wear
resistance than CFI3208 and GS44 that have smaller
grain structures.
CONCLUSIONS
In all materials the damage mode changes from smooth
plastic deformation to limited cracking with increasing
scratch load. In Mg-PSZ, plastic deformation and
extensive distributed surface and subsurface
microcracking;
in
Y-TZP, a smooth scratch track with
limited plasticity and, above a critical load, deep
cracks;
in
Ce-TZP, similar to Mg-PSZ, but less
pronounced microcracking. Bending tests on
specimens subjected to scratch damage indicate a
relatively high damage tolerance
in
the Mg-PSZ and
Ce-TZP; Y-TZP shows the highest initial strength, but
suffers relatively large strength
loss
above the critical
load for cracking. In Y-TZP, fine Y-TZP showed
smooth scratch groove without cracking. As grain size
increased, the scratch groove became rougher with the
development of microcracking. In Si3N4, the critical
load for a transition from plastic deforamtion to
microscracking decreased
with
increasing grain size.
Coarse Si3N4 has a higher material removal rate than
fine Si3N4. The material removal mechanism is
attributed to grain pullout
in
coarse grain materials.
Acknowledgments
This research was sponsored by the Department of
Energy
in
USA, under Contract DE-FC05-970R22579.
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198
CERAMIC COATINGS WITH SOLID LUBRICANT ABILITY
FOR ENGINE APPLICATIONS
M. Buchmann,
R.
Gadow, A. Killinger
University of Stuttgart, Institute for Manufacturing Technologies of Ceramic
Components and Composites (IMTCCC)
Allmandring
5b,
D-70569
Stuttgart, Germany
ABSTRACT
Recent automotive engineering developments con-
cerning fuel consumption regulations and decreasing ma-
terial and manufacturing cost result in an increasing utili-
zation of light metal components for automotive applica-
tions. Significant weight savings
are
obtained by a modifi-
cation of the engine block material from cast iron to alu-
minum. Since all parts of a combustion engine interact
as
a system the single components must sustain the combus-
tion pressure and temperature as well as wear and friction
effects of moving surfaces in different environment and
lubrication. Approaches to increase efficiency and lifetime
of light metal engine blocks are protective thermally
sprayed coatings on cylinder bores. The used thermal
spray processes are high-energetic (Atmospheric Plasma
Spraying) and high-energetic hypersonic processes (High
Velocity Oxygen Fuel Spraying). The knowledge of the
mechanical and thermophysical properties of composite
materials is a key requirement for an optimized stable and
repeatable manufacturing process
as
well as for repro-
ducible high quality composites. This paper describes an
actual overview about the material screening, the manu-
facturing technology and the measured coating character-
istics. Residual stress measurements are performed and
the effects on the coating properties like hardness, friction
and wear is investigated.
INTRODUCTION
Future engine design standards demand low environ-
mental impacts, enhance operation lifetime
as
well as
increased economy and energy efficiencies. Savings in
fuel consumption require an improved motor design to-
gether with a reduction of the total vehicle weight. Nowa-
days the mass proportion of the engine block on the total
automotive weight is in the range between
10
to
15
%.
Therefor light weight engine design offers
a
great poten-
tial for a successful mass reduction [I]. The state of the art
of light metal engines are cast aluminum crankcases. The
worse corrosion, wear and friction behaviour of aluminum
surfaces in a heavily loaded tribological system like cylin-
der liner
-
piston ring, make a wear and corrosion protec-
tion of the surface indispensable. Commonly used protec-
tion systems are integrated iron bushings, galvanic coat-
ings based on chromium
or
nickel, as well as hypereutec-
tic AlSi engine block alloys. Demands to reduce the pro-
duction process complexity, the process steps and manu-
facturing cost on the one side and to improve the opera-
tion behaviour, the environmental safety and recycling
process on the other side lead to the investigation of ther-
mally sprayed coatings
as
protective surfaces on light
metal alloys. Especially of tribological interest are coating
systems with low friction and wear coefficients. In future
tribological systems in engines will operate with a lubri-
cant film thickness lower than
1
pm
or
with biologically
degradable fluids or life-time lubrication or even under
dry condition. If traditional liquid lubricants cannot be
used, the tribological functions must be taken over by
material surfaces with solid lubricant capabilities [2].
THERMALLY SPRAYED COAT-
INGS FOR CYLINDER BORES
The thermal spray process offers the possibility
to apply a broad variety of metallurgical, cermet and ce-
ramic coatings on the surface of machine components,
even on systems with complex geometry. The coatings are
commonly used to improve the wear and corrosion resis-
tance, the operation temperature and thermal shock be-
haviour
or
to influence the electrical, magnetic and bio-
logical behaviour of the composite surface. The latest
technological features
are
high energetic, high velocity
coating systems. The HVOF process uses liquid fuels or
fuel
gases
for high energetic combustion with oxygen
(v-
-
300
-
600
ds;
emax
-
2.500
-
3.200 "C). It leads to
extremely dense coatings because of the high kinetic en-
ergy of the hot powder loaded gas jet. During the APS
process temperatures up to 20.000 "C
are
obtained, there-
for this system
is
mainly used for refractory materials.
Precedent to the coating process a grit blasting and de-
greasing of the substrate surface is performed. The rough-
ening of the surface with corundum of defined size im-
proves
the
mechanical adhesion of the coating and induces
compressive stresses into the substrate material. Following
to the coating deposition a mechanical (grinding, polish-
ing, honing) or thermal post-treatment of the coating sur-
face takes place.
In the case of cylinder bore coatings two principal
different methods can be taken into consideration. State of
the
art
for inside coatings
are
rotating plasma spray de-
vices. The engine block
is
totally fixed and the bore is
coated by means of a vertically moving rotating plasma
torch.
A
second very promising method is the deposition
of inside coatings with
a
fixed
APS
or
HVOF
spraying
199
gun and a rotating engine block. The main advantages of
the second technique are the improved coating micro-
structure by less porosity and coating adhesion, because
of
an enhanced jet propagation and thus particle velocity.
Coating studies are performed on aluminum tubes
(d
=
95 mm,
1
=
150 mm) as well as on 4 cylinder alumi-
num crankcases. During the coating process the tubes
or
engine blocks are totally fixed on a rotating table with a
rotation speed between 50
-
100 rpm. For the plasma
coating process an F1 inside torch is used with a spraying
distance of 45 mm perpendicular to the substrate surface.
With this system a maximum operation power of 25 kW
can be realized.
For
the plasma process a two axis feed
drive system is utilised. The coating is deposited in sev-
eral transitions with a constant vertical feed drive between
5
-
10
mm/s.
Besides the plasma process also the
HVOF
process
with the TOP
GUN@
system is investigated for the depo-
sition of inside coatings, compare figure
1.
The TOP
GUN@
system offers additionally the possibility to use
acetylene as fuel gas, performing sufficient high tem-
peratures to melt refractory materials. By means of a 7
axis robot system, including the regulation for the rotating
table, variable spraying angles (30'
-
90') as well as dif-
ferent spraying distances between 150 and 250 mm, de-
pendent of the used spray powders, can be flexibly pro-
grammed. The
HVOF
coating process is adjusted in a way
that the top of the cylinder bores, where the maximum
wear mechanisms are assumed, are coated perpendicular,
whereas the less loaded bottom regions are coated with an
decreasing spray angle of
30
degrees. By adapting the
feed rate of
HVOF
gun on the spraying angle a homoge-
neously coating thickness over the cylinder bore depth can
be reached.
Fig.
1
HVOF equipment for the inside coating of light
metal cylinder bores.
To avoid an overheating of the substrate during
HVOF
spraying and to improve the coating microstructure a
special liquid
COz
inside fan-shaped cooling system was
developed. With a mass flow of 1.7 kg
C02
per minute the
maximal process temperature during
HVOF
spraying are
lower than 250'C, measured pyrometrical on the alumi-
num tube outside. The use of a liquid
COz
system com-
bines the advantages
of
a very efficient and low cost
cooling system, compare figure 2.
Fig. 2 Left side
-
inside fan shaped
CO2
cooling system,
right side
-
pyrometrical measured temperature
distribution on the outside
of
an
AlSi
tube after the
coating process
MATERIAL SCREENING AND
CHARACTERISTICS
The material screening is focussed on the devel-
opment of high efficient coating systems for tribological
applications with regard to low friction coefficients and
wear rates as well as to high corrosion resistance. Beside
sliding speed, surface roughness and ambient conditions
like temperature and humidity the tribological behaviour
is mainly influenced by the used material systems and
their characteristics. Of special interest are materials with
solid lubricant properties
or
the capability to form lubri-
cious oxides under tribological conditions, like oxides of
titanium and vanadium as well as molybdenum and tung-
sten. Both types of materials can be characterized with
low shear strengths caused by planar defects which are
arranged in parallel crystallographic shear planes. For this
kind of materials low friction coefficients and wear rates
can be expected also under mixed and dry friction condi-
tions
[3].
To get a general material overview thirteen differ-
ent ceramic, cermet and metallurgical coating systems
were selected as coating material and investigated. In
combustion engines with dynamic thermal load applica-
tions the operation behaviour of layer composites is
mainly influenced by the thermophysical properties of the
different composite materials. The thermal expansion
coefficient
a
of the coatings represents a very important
material property influencing the lifetime of coated com-
ponents. Thermal expansion mismatches between sub-
strate and coating material result in the development of
residual stresses during thermomechanical load opera-
tions. Critical residual stresses caused by the manufactur-
ing process
in
addition with operation loads
can
lead to
component failure. Table
1
summarises the investigated
coating materials and the measured average values of
200
a
and the molar heat cp in the temperature range from
100
to
600
"C. Dilatometer measurements are performed on
powder samples
(ap)
and specially prepared spraycones
(@)
as reference 'quasi-bulk' material and compared with
the AlSi9Cu alloy data. The heat capacity cp is measured
by differential thermoanalysis (DTA, DSC).
A120fli02-60/40
Cr203
Cr203/Ti02-60/40
AISi9Cu
1
-
I
28.5
1
0.91
1
5.97 5.4 0.92
7.21 7.4 0.71
7.84
7.5
0.73
Ti02
1
8.4
I
9.4
1
0.73
1
MOO^
AI203/ZrO2-60/40
Cr2C9/NiCr-80/20
5.0
5.8 0.73
6.9 7.2 0.83
9.1
10.7
0.55
Cr20fli02-75/25
I
7.8
1
7.5
I
0.74
I
(Ti,Mo)(C,N)/NiCr
CrB/NiCr-75/25
6.0 9.4 0.42
10.7 10.5 0.44
MoICr-Steel-70130
Mo
I
4.7
I
5.0
I
0.24
I
9.75 0.30
Mo/NiCrBSi-30/70
I
-
I
12.2
I
0.42
1
Table
1
Material screening and thermophysical material
properties
COATING CHARACTERIZATION
The quality of thermally sprayed coatings with
regard to the residual stress situation, microstructure,
surface roughness, coating porosity, hardness and me-
chanical properties can widely be varied by tuning the
equivalent spraying parameters like energy supply, sub-
strate preheating and simultaneous process cooling. Dur-
ing thermal spraying the spray powder is partially
or
even
fully molten within milliseconds, accelerated to high ve-
locities and propelled onto the substrate surface.
The
more
or
less lamellar microstructure of many thermally sprayed
coatings show different phases, due to the rapid solidifi-
cation and quenching processes, a macro and micropo-
rosity as well as oxidation areas between the single splats
if metals are applied in contact with the atmosphere. The
correlation between microstructure and coating properties
has to be known to optimize the spraying parameters.
RESIDUAL STRESSES AND BONDING
STRENGTH
Residual stresses arise during the thermal spray
process and influence significantly the coating quality and
composite performance. Critical residual stresses can
cause failure of coatings in form of delamination and
buckling effects in the coating and the interface as well as
they
cause
perpendicular cracking if the stress level
reaches the ultimate strength of the coating
or
lead to
plastic strain in the interface. Tensile residual stresses
in
the coating reduce the component lifetime since this fa-
vours crack formation and propagation. Furthermore con-
densed corrosive products can penetrate the coating
through microcracks, destabilize the coating and attack the
substrate material
or
the interface layer. Tensile stresses in
the coating propagate stress corrosion cracks. Also the
adhesion between coating and substrate
is
mainly influ-
enced by the residual stresses in the interface. The final
residual stress situation of thermally coated components is
superimposed by several individual stress mechanisms.
The reasons for residual stresses during manufacturing are
temperature gradients in material combinations with
originally incompatible thermophysical properties as well
as mechanical loads which occur during substrate pre-
processing, thermal spraying and finally composite post-
processing.
With the hole drilling method, more exact the cir-
cular micromilling method, the residual stresses in com-
ponents are determined partially destructive. In several
drilling and milling processes a circular microhole is
brought step by step
(5
-
10
pm) into the component sur-
face. The residual stresses
are
relieved due to this material
removal, deform the surface around the hole and are
measured as relaxed strains
E
at the surface by means of
high precision strain gauges. Out of this strains the in
plane stresses
ox, oy
are incrementally determined by
Hooke's law
[4].
The residual stresses of APS and HVOF deposited
Ti02 cylinder bore coatings, sprayed with different cool-
ing rates are measured up to a drilling depth of
0.4
mm,
see
figure
3.
The coating thickness is constant
160
pm.
Due to the low particle velocities
(-
50
-
150
ds)
and the
large degree of totally melted particles the occurring re-
sidual stresses
of
APS deposited coatings are mainly in-
fluenced by the temporary temperature distribution in the
layer composite and the thermophysical material proper-
ties. Because of the high particle velocities and the low
degree of particle fusion during HVOF spraying, here the
thermal stresses are superimposed by additional compres-
sive stresses induced by the high particle impact energy. It
can be noted that the kinetic particle energy can be con-
trolled via the total gas flow rate as a result of the oxygen
/
fuel ratio.
For
the APS process tensile stresses in the coating
and the interface can be measured with increasing cooling
rates by C02 and air jets. For decreasing cooling rates,
only air cooling resp. without cooling, compressive
stresses can be detected. Decreasing cooling rates cause a
higher uniform temperature level
in
the composite. Be-
cause of the mismatch in the thermal expansion coeffi-
cients between coating
q
and substrate
as
higher com-
pressive thermal stresses arise in coating and interface.
For
the HVOF deposited coatings compressive stresses in
coating and interface can be measured. Due to the larger
20
1