Note: Descriptions are shown in the official language in which they were submitted.
CA 02319695 2008-05-21
Alloy for Electrical Contacts and Electrodes and Method of
Making
FIELD OF THE INVENTION
This invention relates to materials for electrical contacts
and electrodes. In particular, this invention relates to
tungsten-copper composites.
BACKGROUND OF THE INVENTION
Alloy for electrical contacts and electrodes are
metallurgical composites based on heterogeneous systems
(pseudoalloys) of two or more components with vastly
different thermophysical properties. The properties of these
alloys represent optimal combinations of the component
characteristics and are required for operation in
applications such as high-current interrupters with gas or
oil as arc quenching media, electrical discharge machining,
spot welding, and other applications employing an electrical
discharge. Common to these applications is the electrical
arcing which occurs between the contacts or between an
electrode and a workpiece. For example, electrical contacts
serve as points of
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arc attachment at current switching. Despite the momentary
duration of switching, electric arcs in oil- or gas-filled
high-current interrupters fully develop into high temperature
plasma discharges. At the points of arc attachment, plasma
discharges generate electrodynamic forces and thermal fluxes
causing erosive wear of electrical contacts. Material erosion
reaches its peak in electrical contacts operating in an
oxidizing environment (e.g., air-blast high-current
interrupters) . To withstand this wear, electrical contact
materials must possess specific thermophysical properties.
Tungsten-copper composites are preferred materials for these
applications. However, electrical contacts made from unalloyed
W-Cu composites are prone to cracking in electrical discharge
environments. The problem appears to stem from the poor
thermal shock resistance of the composite. When the structural
continuity is partially or totally disrupted by the loss of
copper due to arc heating, the composite loses the ability to
undergo plastic deformation as a single structure. If
excessive thermal fluxes produced by the arc attachment to the
contact are not dissipated quickly, the thermal shock
generates high thermal stresses and cracks in the contact. For
composites with an average tungsten grain size of about 20 m,
cracking occurs after a period of more or less even ablation.
Tungsten-copper composites with an average tungsten grain size
of about 5 m show greater cracking, apparently resulting from
further sintering of the tungsten which causes significant
shrinkage and pore formation. The erosion rate and cracking
can increase considerably if the pore volume in the composite
material exceeds approximately 4%.
In addition to thermal shock resistance, the tungsten-copper
materials used in these applications should possess resistance
to loss of copper, resistance to erosion, and resistance to
corrosion. Conventional solutions to these requirements
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include alloying with 4-5 weight percent (wt.%) Ni and
maintaining a low pore volume.
Two basic powder metallurgical (P/M) techniques are used to
make W-Cu-Ni alloys for electrical contacts and electrodes:
sintering/infiltration and direct sintering.
Sintering/infiltration is a two-step manufacturing process
which consists of (1) pressing elemental W powder and
sintering the compact using one thermal cycle to form a
refractory component skeleton (or framework) with controlled
porosity, and (2) infiltrating the skeleton with the
eiectrically/thermally conductive Cu component using another
thermal cycle. The sintering/infiltration technique does not
allow fabrication of net-shape components and the use of fine
W powders (FSSS <5 m). In particular, fine W powders promote
iocalized densification in the W skeleton resulting in
partially closed porosity which cannot be infiltrated with Cu.
High-temperature sintering (above 1950-1500 C) of W in the
presence of liquid Ni promotes the growth of W grains in the
skeleton. Brittle W-Ni intermetallics form along the W grain
boundaries during high-temperature sintering of W in the
absence of Cu. This degrades the mechanical properties of the
W skeleton. In addition, it is difficult to assure uniformly
distributed contacts between the W and Ni particles when Ni is
used as a sintering aid for W powder.
Conventional direct sintering consists of blending and
compacting of the W, Cu, and Ni powders with an average
particle size of about 5 m. Then, depending on the Cu
content, the compacts are sintered at temperatures above or
below the melting point of Cu. Conventional direct sintering
approaches suffer from an inability to separately sinter a
strong W skeleton to act as the alloy backbone. In addition,
there are problems with (1) excessive coalescence and solid-
state sintering of Cu prior to melting, (2) excessive
coagulation of Cu at melting, (3) Cu bleedout from an
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improperly sintered W skeleton, (4) development of excessive
porosity (> 4%), (5) disintegration of the W skeleton, and (6)
loss of shape (slumping).
Enhanced sintering of W-Cu is strongly influenced by the
formation of a Cu-based liquid phase above 1083 C. Ni and Cu
have an unlimited mutual solubility which in combination with
partial solubility of W in Ni (38 wt.% of W in Ni at 1100 C)
greatly improves the wetting of W by Cu and eliminates Cu
bleedout. The sintered density, strength, and microhardness
increase linearly with Ni additions. The affinity of Ni for
both Cu and W introduces a solution-reprecipitation mechanism
for sintering W. Operation of this mechanism reaches a
significant level at Ni concentrations in the alloy of at
least 2 wt.%.
The Cu-Ni liquid phase is chemically active with respect to W.
It begins dissolving W and forming a Cu-Ni-W matrix. Due to
the limited solubility of W in Ni, the concentration of
dissolved W in the matrix eventually reaches an equilibrium
level. Formation of the Cu-Ni-W matrix turns on the solution-
reprecipitation mechanism which affects the sintering of W.
The matrix acts as a W carrier by dissolving tiny W particles
and necks, and transporting and redepositing W onto surfaces
of larger particles causing their further growth. This
thermodynamically warranted process is governed by kinetic
parameters such as the concentration of Ni in the matrix, the
size of W particles, and the temperature. Microstructure and
mechanical properties of W-Cu-Ni alloys produced by a
solution-reprecipitation mechanism are strongly influenced by
two metallurgical phenomena, the Kirkendall effect and Ostwald
ripening. Higher diffusion rates of Cu and W in Ni, compared
to those for Ni in Cu and W, result in formation of pores and
voids (the Kirkendall effect) which may not be totally
eliminated by sintering. Coarsening and spheroidization of W
particles in the presence of an active Cu-Ni liquid phase
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(Ostwald ripening) may lead to porosity, disintegration of
the W skeleton, and loss of shape (slumping) of the sintered
material. Due to the above effects, alloys made by a
solution-reprecipitation sintering technique are very
sensitive to processing conditions. Even slight changes in
sintering temperature will cause a drastic reduction in
strength and ductility of these alloys. Depending on the
shape and size of pores, significant fluctuations of strength
and ductility are observed even though the sintered density
of the alloys may be within 99% of theoretical density (TD).
SUMMARY OF THE INVENTION
It is desirable to obviate the disadvantages of the prior
art.
It is also desirable to provide a W-Cu-Ni alloy having
thermophysical properties suitable for use in electrical
contacts and electrodes.
It is further desirable to provide a powder blend for forming
a W-Cu-Ni alloy having thermophysical properties suitable for
use in electrical contacts and electrodes.
It is also desirable to provide a method for direct sintering
a W-Cu-Ni alloy which substantially eliminates the formation
of brittle intermetallics and slumping during sintering.
In accordance with one aspect of the present invention, there
is provided a powder blend for making a W-Cu-Ni alloy
comprising a W-Cu composite powder and a nickel powder, the
W-Cu composite powder comprising individual particles having
a tungsten phase and a copper phase wherein the tungsten
phase substantially encapsulates the copper phase, wherein
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the copper to nickel weight ratio is from about 4.0:1 to
about 4.2:1. The W-Cu-Ni alloy may comprise about 3 to about
6 weight percent of nickel.
In accordance with another aspect of the present invention,
there is provided a powder blend for making a W-Cu-Ni alloy
comprising a W-Cu composite powder and a nickel powder, the
W-Cu composite powder comprising individual particles having
a tungsten phase and a copper phase wherein the tungsten
1o phase substantially encapsulates the copper phase, wherein
the nickel powder comprises from about 3 to about 6 weight
percent of the blend.
In accordance with another aspect of the present invention,
there is provided a W-Cu-Ni alloy comprising a sintered
tungsten skeleton containing a Cu-Ni matrix, the alloy having
no brittle intermetallics, wherein the weight ratio of copper
to nickel is from about 4.0:1 to about 4.2:1. The W-Cu-Ni
alloy may comprise about 3 to about 6 weight percent of
nickel.
In accordance with another aspect of the present invention,
there is provided a method for forming a W-Cu-Ni alloy
comprising: (a) forming a powder blend of a W-Cu composite
powder and a nickel powder, the W-Cu composite powder
comprising individual particles having a tungsten phase and a
copper phase wherein the tungsten phase substantially
encapsulates the copper phase; (b) pressing the powder to
form a compact; and (c) sintering the compact to form a W-Cu-
Ni alloy. The blend may have a copper to nickel weight ratio
of from about 4.0:1 to about 4.2:1. The blend may comprise
about 3 to about 6 weight percent of nickel.
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There is disclosed a powder blend for making a W-Cu-Ni alloy
comprising a W-Cu composite powder and a nickel powder, the
W-Cu composite powder comprising individual particles having
a tungsten phase and a copper phase wherein the tungsten
phase substantially encapsulates the copper phase.
There is further disclosed a W-Cu-Ni alloy comprising a
sintered tungsten skeleton containing a Cu-Ni matrix, the
alloy having no brittle intermetallics.
BRIEF DESCRIPTION OF THE DRAWINGS
Figs. 1-10 are optical micrographs (500X) of the
microstructure of a W-Cu-Ni alloy in various stages of the
sintering cycle.
Fig. 11 is an optical micrograph (500X) of the microstructure
of a presintered compact of the W-Cu-Ni alloy.
Figs. 12-14 are SEM micrographs (1000X) of the
microstructures of three W-Cu-Ni alloys having different
compositions.
DESCRIPTION OF THE PREFERRED EMBODIMENTS
The tungsten-copper (W-Cu) composite powders used in this
invention are described in U.S. Patent No. 5,956,560 issued
9/21/99 to Dorfman et al. These powders contain individual
particles each having a tungsten phase and a copper phase
wehrein the tungsten pahse substantially encapsulates the
copper phase. The W-Cu composite powder is mechanically
blended with a nickel powder to form a powder blend which is
then pressed and sintered to form the W-Cu-Ni alloy. The
unique distribution of the tungsten and copper phases in the
individual particles of the W-Cu composite powder combined
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with the nickel phase presents two important technical
advantages in achieving the direct sintering of the compacted
powder blend to form net-shape articles of the W-Cu-Ni alloy.
The first technical advantage is the ability to selectively
activate the sintering of a W framework in-situ in the
presence of, and without interference by, the Cu phase. In
the W-Cu composite powders used herein, the W and Cu phase
are not a mechanical mixture of elemental powders. Instead,
the individual particles of the W-Cu composite powder are
dual phase particles comprised of a tungsten phase which
substantially encapsulates a copper phase. Because the
tungsten phase exists on the exterior of the particles, the
Cu phase does not interfere with the formation of
predominantly W-W and W-Ni-W contacts upon compaction of the
powder blend. Due to the submicron size and high sintering
activity of the W phase, these contacts facilitate in-situ
sintering of a W framework in the solid-state sintering
region (950-1050 C) prior to the melting of the Cu. Also, by
removing diffusion barriers and improving the W-W mass
transport, the Ni phase selectively activates the solid-state
sintering of the W phase. Compared to sintering without
activation, the
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shrinkage of the compact increases several times which results
in a much more rigid W framework.
The second technical advantage is the ability to provide, as
~ two separate steps of the same sintering cycle, conditions for
(i) formation of a Cu-Ni melt, and (ii) the controlled
modification of the W framework by the Cu-Ni melt to impart
ductility (yield and elongation) to the W-Cu-Ni alloy. At
temperatures equal to or above the melting point of Cu
(1083 C), liquid Cu promotes complete densification of the W
framework by enhancing the W particle rearrangement and shape
accommodation. Formation of a Cu-Ni melt also begins at the
above temperatures with all of the Ni gradually going into the
melt due to unlimited mutual solubility of Cu and Ni. The Cu-
Ni melt is kept inside the W framework by capillary forces.
Since W is partially soluble in Ni, the dissolved Ni increases
the diffusion rate of W into the Cu-Ni melt facilitating the
transport process of W into the melt and the modification of
the W framework by the melt. The Cu-Ni melt begins dissolving
W and forming a Cu-Ni-W matrix. Formation of a Cu-Ni-W matrix
turns on the solution-reprecipitation mechanism for the W
phase. The matrix acts as a W carrier by dissolving tiny W
particles and necks and, after becoming saturated with W,
redepositing the dissolved W on larger W particles causing
2E their further growth. The Kirkendall effect manifests itself
by porosity development and coagulation of pores; these pores
eventually disappear in the course of controlled exposure of
the composite material to the solution-reprecipitation
process. The process exerts a profound effect on the
contiguity (W grain interface), size, and morphology of W
particles forming the framework. By rounding off the
particles and increasing their size, it loosens up the bond
between W particles in the framework thereby weakening it and
imparting ductility (yield and elongation) to the sintered W-
Cu-Ni alloy. If allowed to proceed in an uncontrolled manner,
the solution-reprecipitation process, by the effect of Ostwald
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ripening; will eventually lead to complete disintegration of
the W framework (zero contiguity of the W phase) and loss of
shape (slumping). Therefore, by varying the time-temperature
parameters of the solution-reprecipitation process, the
properties of the W framework and of the sintered composite W-
Cu-Ni product can be modified from being hard and brittle
(after activated solid-state sintering), to having moderate
mechanical strength and ductility (after controlled exposure
to the solution-reprecipitation process), to being weak and
prone to slumping (after the effect of Ostwald ripening)
The following parameters and materials are preferred for the
making of the W-Cu-Ni alloy of this invention.
A. Tungsten-copper composite powder
Copper content: about 10 to about 25 wt.%
Median particle size: about 0.5 to about 20 m
Thickness of the tungsten phase on the particles: about 0.1 to
about 0.2 m
B. Nickel Powder
Median particle size: about 1 to about 15 m
C. Powder Blend
Weight ratio of copper to nickel: about 4.0:1 to about 4.2:1
The more preferred Cu:Ni. weight ratio is 4.1:1. This ratio is
based on an Cu3_8Ni solid solution which has a known XRD
pattern. (A Cu:Ni atomic ratio of 3.8:1 has a weight ratio of
4.1:1.) The Cu:Ni ratio is kept within the narrow range in
order to verify by XRD that no brittle intermetallics have
been formed in the alloy.
The preferred range of weight ratios for the alloy
compositions is from 82 wt.% W and 18 wt.% (Cu:Ni = 4.1:1) to
70 wt.% Wand 30 wt.% (Cu:Ni = 4.1:1).
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The amount of Ni preferably comprises about 3 to about 6 wt.$
of the blend and, more preferably, about 4 wt.%.
D. Pressing and Sintering
Compaction pressure: about 45 to about 70 ksi
Sintering temperature: about 1180 C to about 1200 C; more
preferably about 1190 C
Sintered density: 99 0.5% theoretical density; more
preferably at least about 99% TD
E. Alloy Properties
Elongation: about 2% to about 20%; more preferably about 3%
to about 5%.
Calculated average W particle size: about 2.5 to about
15.0 um.
Contiguity of tungsten skeleton: about 15% to about 30%.
The following non-limiting examples are presented.
EXAMPLES
Tungsten-copper composite powders and a nickel powder with a
median particle size of 8.8 pm were used in preparation of
the feedstocks for making three different composition of the
W-Cu-Ni alloy. A solid lubricant (0.5 wt.% of AcrawaxTM C
produced by Lonza Co. in Fair Lawn, New Jersey) was added to
powder feedstocks to improve pressibility. Feedstocks were
prepared in a V-blender with an intensifier bar by blending
an 8 kg powder batch for 60 minutes. Powder ratios and alloy
compositions are provided in Table 1. The Cu:Ni weight ration
was maintained at 4.1:1.
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Table 1
Alloy W-Cu Composite Ni Powder Alloy Composition
Powder (g per lOOg
Copper Median of W-Cu W (wt.%) (Cu:Ni = 4.1:1)
(wt.%) Size ( m) Powder) (wt.%)
A 15 6.9 3.66 82 18
B 20 13.9 4.9 76.3 23.7
c 25 16.1 6.1 70.7 29.3
Sintering tests were conducted with the use of powder
feedstock to make the alloy B. A quantity of about 7.5 kg of
the feedstock was isostatically pressed at 45 ksi into a slug
with a green density of about 56% of theoretical density (TD)
and approximate dimensions of 3.75 inches in diameter by 4.75
inches long. Dewaxing and sintering were carried out in a
tube furnace with flowing dry hydrogen. The rate of
temperature increase was 2 C/minute. The slug was dewaxed at
450 C for 4 hours and presintered at 1000 C for 4 hours. The
presence of nickel dramatically increased the solid-state
sintering of W-Cu composite powder. A linear shrinkage of
about 20% was observed compared to a linear shrinkage of about
5% without nickel.
The presintered sluq was cut longitudinally into eight
sections weighing about 930g each. These sections were used
in systematic liquid-phase sintering tests. Optical
microscopy (OM), Scanning Electron Microscopy (SEM), Energy
Dispersive X-ray Spectroscopy (EDS), X-Ray Diffraction (XRD),
and other standard methods of physical testing, such as yield
strength (YS), ultimate tensile strenoth (UTS), transverse
rupture strength (TRS), hardness, etc., were used in alloy
characterization. The test data are presented in Table 2 and
in Figures 1 to 10 which correspond to the samples in Table 2.
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Figure 11 shows the microstructure of the presintered slug.
The presintered slug had the lowest sintered density (92.9%
TD) and finest size W grains. This is characteristic of the
material produced by solid-state sintering.
The evolution of the microstructure and properties of the W-
Cu-Ni alloy made by liquid-phase sintering is controlled by
the temperature and time parameters of the sintering cycle
which, in turn, control the solution-reprecipitation mechanism
and growth of W particles. The concentration of W in the Cu-
Ni-W matrix increases with holding time from about 0.8 wt.% to
an equilibrium level of about 2 wt.% to about 2.2 wt.% for the
range of processing temperatures (Table 2). Three distinct
microstructure-property ranges are observed.
In samples 1-5 (Figs. 1-5), the tungsten particles were fine
and highly interconnected. The grain count was [500-1200] x103
grains/mm 2 yielding a calculated average particle diameter of
1-1.6 m. This range is associated with formation of a strong
tungsten framework and a gradual increase in the density (up
to 95-96%TD) and mechanical properties of the sintered alloy
(UTS, TRS, hardness). However, the alloy remains brittle under
these conditions. Even a fivefold increase in residence time
did not produce any significant change in the material's
2 properties.
In samples 6-9 (Fig. 6-9), the tungsten particles were medium
size, less interconnected, and partially rounded. The grain
count was [60-90] x103grains/mm ` yielding a calculated average
particle diameter of 2.6-4.6 m. This range is characterized
by a continued increase in the sintered density (up to
97.6%TD), the elimination of pores and voids including those
generated by the Kirkendall effect (Figures 7 and 8), and the
appearance of elongation (ductility) in the alloy (up to 10%
~5 elongation in sample 8) The presence of Ni lowered the
electrical conductivity to less than 18% IACS.
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In sample 10 (Fig. 10) , the particles were rounded and weakly
connected. The grain count was below 60 x 10' grains/mm 2
yielding a calculated average particle diameter of above 4.6
m. The particle rounding is a result of Ostwald ripening and
caused a substantial disintegration of the W framework, a
lowering of the mechanical properties and density, and
slumping of the sample during sintering.
-13-
CA 02319695 2000-09-15
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CA 02319695 2000-09-15
99-2-925
YH'1' Y~ N'1'
Referring to the test data in Table 2, it is obvious that
alloy elongation resulted from a controlled trade off between
the UTS and ductility of the material. UTS and hardness were
gradually built up in the samples 1-5. Ostwald ripening
increased the particle diameter, decreased the grain count,
and lowered the contiguity of the W framework in samples 6-8.
This resulted in weakening of the W framework (appearance of a
yield point and YS; attainment of the highest TRS value) and
the highest material ductility in sample 8. Further weakening
of the W framework in samples 9 and 10 led to slumping of the
alloy.
The microstructure of sample 8 (Figure 8) is typical of a W
heavy alloy produced by liquid phase sintering, only it is
much finer. The only peaks determined by XRD were those which
belonged to W and the Cu3,8Ni solid solution. To the extent of
sensitivity of the XRD method, it can be concluded that
formation of brittle intermetallics was eliminated in the
process of manufacturing the alloy.
Feedstock quantities of alloy compositions A, B and C (about
4.3 - 4.5 kg) were pressed as described for samples 1-10. To
improve the uniformity of heat transfer to green compacts,
they were dewaxed and sintered in pure alumina sand using the
same conditions. The sintering produced slugs with
approximate dimensions of 1.75 inches in diameter x 7.5 inches
long. Sintering cycles included three isothermal holds. The
first was at 1000 C for hydrogen cleaning of oxygen from the
powder compacts. The next was at 1100 C for removing oxygen
from the molten copper, formation of a Cu-Ni solid solution,
and presintering of the compacts. The final hold was at the
sintering temperature. Final sinterino conditions for the
alloy were optimized for obtaining the highest elongation
without slumping of the alloys. The data in Table 3 represent
averages of six determinations on separate samples of the
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alloy. Figs. 12, 13 and 14 are SEM photomicrographs of the
microstructure of alloys A, B, and C, respectively.
Table 3
Alloy Composition Sintering YS UTS Elongation %TD $IACS
(wt.%) Cycle, (Mpa) (Mpa) ($)1
C - Hours
A 82W - 1000 C, 4 hr
18(Cu:Ni = 1100 C, 4 hr 553 605 2.8 98.8 19.5
4.1:1) 1200 C, 4 hr
B 76.3W - 1000 C, 4 hr
23.7(Cu:Ni = 1100 C, 9.hr 571 700 15 99.1 17.4
4.1:1) 1190 C, 3 hr
C 70.7W - 1000 C, 4 hr
29.3(Cu:Ni = 1100 C, 4 hr 602 685 4.5 98.9 15.6
4.1:1) 1180 C, 1 hr
'Tension Testing of Metallic Materials, ASTM Test Method E-8
Each sintered alloy exhibited a very fine, homogeneous
microstructure. The ranges for the W grain count and the
calculated average grain size were 12.4 x 103 to 39.8 x 103
grains/mm 2 and 10.2 m to 5.7 m, respectively. Despite the
small size of particles, the contiguity of the W framework was
effectively lowered to a level of 18-27 percent in order to
obtain substantial elongation without slumping. The
elongation exhibited for alloy B was similar to that of W
heavy alloys which have a substantially coarser microstructure
(grain sizes from 30 m to 100 m). =
The arc erosion rate in SF6 of alloy A was compared with the
erosion rate of a W-15Cu pseudoalloy made from a W-Cu
composite powder having 15 wt.% Cu and no Ni and two other
conventional electrical contact materials consistina of
infiltrated W-Cu pseudoalloys. The purpose of the test was to
assess the applicability of the alloys as electrical contact
materials in high power interrupters. The change in mass of
both the anode and the cathode were recorded and volume
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burnoff for both contacts was determined based on the density
of the material. The arcing behavior of the Alloy A in SF6 was
similar to that of the reference materials but Alloy A
exhibited lower erosion rates. Furthermore, Alloy A
demonstrated an erosion rate that was practically linear as a
function of current density and, at the higher current
densities, was within the range of the best conventional
electrical contact materials previously tested in an SF6
environment. Alloy A also demonstrated a very consistent
performance in air. Arced contacts showed no structural
disintearation of the recrystallized surface material in spite
of moderate cracks which are characteristic of all electrical
contact materials operating in air. The erosion rates of Alloy
A were even lower than other conventional contact materials in
SF6 which is a less harsh environment than air.
While there has been shown and described what are at the
present considered the preferred embodiments of the invention,
it will be obvious to those skilled in the art that various
changes and modifications may be made therein without
departina from the scope of the invention as defined by the
appended claims.
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