Note: Descriptions are shown in the official language in which they were submitted.
CA 02276154 1999-06-22
Nickel-based monocrysta.lline superalloy with
a high ti' solvus
The present invention concerns nickel-based superalloys
suitable in particular for manufacturing fixed and moving
monocrystalline gas turbine blades, and exhibiting a high
creep resistance at very high temperatures while retaining
good resistance to the combustion gas environment. These
alloys are particularly suited to applications in
aeronautical engines used to propel airplanes and
helicopters.
Nickel-based superalloys are materials with the highest
performance used today for the manufacture of fixed and
moving blades of the turbines of aeronautical gas turbine
engines. The work of ONERA in this field started at the
end of the 1970s and resulted, among other things, in the
filing of various patents of invention relating to
monocrystalline superalloys intended for different fields
of application : FR 2 503 188, FR 2 555 204, FR 2 557 598,
FR 2 599 757, FR 2 643 085 and FR 2 686 902.
The development of the performances of aeronautical gas
turbines, translated in terms of specific power and yield
and length of life, calls for the availability of alloys
for turbine blades exhibiting high temperature mechanical
properties (650 to 1150'C) and improved continuous
resistance to corrosion and hot oxidation. Extreme
operating conditions may in point of fact raise the metal
to temperatures in excess of 1100'C. In order to optimize
the resistance to hot corrosion and hot oxidation,
monocrystalline blades made of superalloy are moreover
generally coated with a protective deposit of the nickel
aluminide or MCrAlY alloy type. In order to prevent any
possible cracking or fracturing of these protective layers
under the effect of thermal cycling, which would adversely
effect the life of the parts, superalloys must however
exhibit high intrinsic oxidation and corrosion resistance.
CA 02276154 1999-06-22
In polycrystalline blades cast by conventional foundry
processes, a large part of the hot deformation during
service is produced in the region of the grain boundaries
which limits the life of the parts. The development of the
monocrystalline solidification. process has, by eliminating
grain boundaries, enabled the performance of nickel-based
superalloys to be increased spectacularly. In addition,
the process makes it possible to select the preferred
growth orientation of the monocrystalline part and hence to
select an optimum <001> orientation as regards creep
resistance and thermal fatigue: resistance which are the two
stress modes causing the greatest damage to turbine blades.
Successive improvements to the: mechanical performances, in
particular creep, of these superalloys for monocrystalline
blades have been made possible by optimizing their chemical
compositions. Indeed, apart from nickel, which is the major
constituent of these alloys, various addition elements
provide specific contributions to the properties of the
alloys. The functions of these: elements will be detailed
hereinafter. In the monocryst.alline superalloys covered by
the previously mentioned pater.its, the main addition
elements (weight concentratior.Ls at a level of a few
percent) have generally been selected from the following
list: chromium (Cr), cobalt (Co), molybdenum (Mo), tungsten
(W), aluminum (Al), titanium (Ti), tantalum (Ta) and
niobium (Nb). The elements Cr, Co and Mo and part of the W
participate mainly in hardenir.Lg the austenitic (y phase)
matrix where they enter into solution. The elements Al, Ti,
Ta and Nb promote precipitation in the y matrix of
hardening particles of a secor.Ld phase of the Ni3(Al, Ti, Ta,
Nb) type (T' phase). Minor elements (weight concentrations
below 0.5 %), such as silicon (Si) and hafnium (Hf), may
also be added in order to opti.mize the environmental
resistance as is demonstrated in FR 2 686 902.
CA 02276154 1999-06-22
Since the start of the 1980s, a large number of patents
dedicated to novel superalloy compositions for
monocrystalline blades have been filed worldwide. More
recent developments have consisted in particular of
incorporating the refractory elements rhenium (Re) and
ruthenium (Ru) in these alloys. These additions are aimed
especially at improving the high temperature creep
resistance of these monocrystalline superalloys while
preserving a stable microstructure at high temperatures as
regards the formation of particles of intermetallic phases
which are liable to bring about losses in the properties of
these alloys.
Various patents thus protect the fields of monocrystalline
superalloy compositions containing additions of one and/or
other of the elements Re and Ru, in particular US 4 719 080
(United Technologies Corporation), US 4 935 072 (Allied-
Signal Inc.), US 5 151 249 (General Electric), US 5 270 123
(General Electric) and US 5 482 789 (General Electric).
However, the information available as regards the
properties of these alloys is very limited and does not
permit a judgement to be made as to the industrial value of
these additions.
In France at the present time the monocrystalline
superalloys used are referred to as "first-generation", as
for example the grades AM1 and. MC2, both covered by patent
FR 2 557 598, and the alloy AM3 protected by the patent
FR 2 599 757. Among these, the alloy MC2 is considered as
the alloy with the highest performance as regards creep
resistance up to 1100 C. The future requirements of
engineers will call for the ability to have alloys for
blades available which have a higher performance than these
first-generation alloys. It is necessary in particular to
increase the maximum permissible temperatures for alloys
constituting turbine blades.
CA 02276154 1999-06-22
The object of the invention is thus to provide a novel
family of nickel-based monocrystalline superalloys
exhibiting improved creep resistance compared with that of
alloys exploited industrially, in particular at
temperatures above 1100 C, but: equally at lower
temperatures affecting various parts of the blades.
To this end, attempts have been made to introduce new
addition elements without penailizing other properties
essential for the good perforniance of these alloys, such as
density, hot corrosion and oxidation resistance and
microstructural stability.
An analysis of the state of the art as well as the results
of work carried out by the inventor quickly showed that
only alloys containing rhenium additions could make it
possible to exceed the creep r.esistance of the alloy MC2
above 1100'C. In order to cour.Lter-balance certain harmful
effects of rhenium (excessive density, microstructural
instability), it seems moreover advantageous to incorporate
ruthenium.
The invention concerns a nickel-based superalloy suitable
for the manufacture of turbo-engine parts by
monocrystalline solidification, wherein its mass
composition is as follows:
Cr . 3.5 to 7.5 %
Mo . 0 to 1.5 %
Re . 1.5 to 5.5 %
Ru . 0 to 5.5 %
W . 3.5 to 8.5 %
Al 5 to 6.5 %
Ti . 0 to 2.5 %
Ta . 4.5 to 9 %
Hf . 0.08 to 0.12 %
Si . 0.08 to 0.12 %
CA 02276154 1999-06-22
the complement to 100 % consisting of Ni and possible
impurities.
More particularly, the invention provides such a superalloy
having the following mass composition:
Cr . 3.5 to 5.5 %
Mo . 0 to 1.5 %
Re . 4.5 to 5.5 %
Ru . 2.5 to 5.5 %
W . 4.5 to 6.5 %
Al . 5 to 6.5 %
Ti . 0 to 1.5 %
Ta . 5 to 6.2 %
Hf . 0.08 to 0.12 %
Si . 0.08 to 0.12 %
the complement to 100 % consisting of Ni and possible
impurities.
Even more particularly, the mass composition of the
superalloy is as follows:
Cr . 3.5 to 5.5 %
Mo . 0 to 1.5 %
Re . 3.5 to 4.5 %
Ru . 3.5 to 5.5 %
W . 4.5 to 6.5 %
Al . 5.5 to 6.5 %
Ti . 0 to 1 %
Ta . 4.5 to 5.5 %
Hf . 0.08 to 0.12 %
Si . 0.08 to 0.12 %
the complement to 100 % consisting of Ni and possible
impurities.
CA 02276154 1999-06-22
6
Three specific compositions of superalloys according to the
invention are given below:
Cr . 3.5 to 4.5 % 4.5 to 5.5 % 3.5 to 4.5 %
Mo . 0.5 to 1.5 % 0.5 to 1.5 %
Re . 3.5 to 4.5 % 3.5 to 4.5 % 4.5 to 5.5 %
Ru . 3.5 to 4.5 % 4.5 to 5.5 % 2.5 to 3.5 %
W . 4.5 to 5.5 % 5.5 to 6.5 % 5.5 to 6.5 %
Al 5.5 to 6.5 % 5.5 to 6.5 % 4.8 to 5.8 %
Ti . 0 to 1 % 0 to 1 % 0.5 to 1.5 %
Ta . 4.5 to 5.5 % 4.5 to 5.5 % 5.7 to 6.7 %
Hf . 0.08 to 0.12 % 0.08 to 0.12 % 0.08 to 0.12 %
Si . 0.08 to 0.12 % 0.08 to 0.12 % 0.08 to 0.12 %
the complement to 100 % consisting of Ni and possible
impurities.
The alloys according to the invention, prepared in the form
of monocrystals with a <001> orientation, have the
following properties:
- a density in all cases below 9 g.cm-3 and at best
8.8 g.cm-3, making it possible to minimize the mass of
monocrystalline blades and consequently to limit the
centrifugal stress acting on these blades and on the
turbine disc to which they are attached;
- an ability for homogenization by returning all the ry'
particles into solution, including the eutectic T/T'
phases;
- a high temperature for putting the hardening ry' phase
into solution, in all cases higher than that of previous
alloys containing neither rhenium nor ruthenium;
- the absence of brittle intermetallic phases which can
precipitate when high temperatures are maintained and which
CA 02276154 1999-06-22
7
are liable to bring about a reduction in the creep
resistance of the alloys and an. embrittlement of the
alloys;
- resistance to hot cyclic corrosion and hot cyclic
oxidation greater than that of previous alloys containing
neither rhenium nor ruthenium.
Obtaining all these properties simultaneously enables the
creep resistance of monocrystalline blades at very high
temperatures and their environniental resistance to be
optimized and therefore enables their life to be improved
as well as the performance of gas turbines.
The invention thus provides a unique combination of
properties for the alloys, which the state of the art was
not able to provide.
The alloys of the invention are! intended for the
manufacture of monocrystalline parts, i.e consisting of a
single metallurgical grain. Thi.s special structure is
obtained with the aid of a process for directed
solidification in a thermal gradient using a grain
selection device or a monocryst.alline seed at the start of
solidification.
After solidification, the superalloys consist essentially
of two phases: the austenitic y matrix is a nickel-based
solid solution in which the particles of the y' phase, an
intermetallic compound of which. the composition is based on
Ni3Al, precipitates during cooling in the solid state. The
addition elements are distributed in the two y and y'
phases but generally exhibit a particular affinity for one
or other of these two phases. Thus chromium, molybdenum,
rhenium and ruthenium are preferably distributed in the y
matrix, while aluminum, titanium and tantalum go
preferentially into the y' phase.
CA 02276154 1999-06-22
F!
In the crude monocrystalline solidification alloys, the
distribution of the particles of the hardening y' phase is
very heterogeneous in the volume of the monocrystal on
account of chemical segregations resulting from
solidification conditions specific to the process. The
microstructure is said to be d.endritic. The precipitates
are very fine in the core of the dendrites which solidify
first of all during cooling of the alloy and become larger
in the regions which then solidify from the center of the
dendrite. Moreover, at the end. of solidification, eutectic
phases consisting of massive particles of the y' phase
containing laminae of the y phase solidify in the regions
separating the dendrites.
Experiments however showed that the creep resistance of
nickel-based superalloys was optimized when the
distribution of the particles of the T' phase were
homogeneous in all the volume of the alloy with precipitate
sizes less than 1 micrometre, the optimum size of the
precipitates depending on the composition of the alloy.
The T' phase contained in the eutectic phases did not
contribute, in particular, to the hardening of alloys, and
the potential of the alloys for creep resistance is
therefore not totally exploited in the crude solidification
state. These massive blocks of the y/y' eutectic phase are
moreover preferred sites for the start of cracks during
cyclic stresses resulting from. thermal fatigue phenomena
due to the starting and stopping cycles of gas turbines.
The compositions of the alloys of the invention have been
selected so as to be able to obtain biphase -y/T'
structures, consisting of a hcmogeneous precipitation of y'
particles in aT matrix resulting from the monocrystalline
solidification steps and the heat treatments detailed
below. In order to achieve th.is optimized microstructure,
it is first of all necessary to apply a heat treatment
designed to dissolve the -y' phase precipitates contained in
CA 02276154 1999-06-22
9
the dendrites and to eliminate the solidified eutectic
phase between the dendrites. Solution of the y'
precipitates is achieved when the heat treatment
temperature reaches the temperature of the y' solvus (the
temperature at which the y' phase precipitates are put into
solution) which is a characteristic of the chemical
composition of the alloy. In practice, the value of the y'
solvus varies periodically in the crude monocrystalline
solidification alloy in relation to the local chemistry of
the alloy. Accordingly, the 1' solvus increases within the
core of the dendrite towards the inter-dendritic regions on
account of chemical segregations, until the initial fusion
temperature of the eutectic y' phase is reached, this phase
being the last solid formed during cooling of the alloy
from the liquid state. This initial fusion temperature is
in practice similar to the solidus temperature (initial
fusion temperature) of the alloy. The homogenization
treatment temperature must therefore remain below the
solidus temperature.
In practice it has been possible to obtain complete
solution of the y' precipitates and the eutectics y/y' in
the alloys of the invention by virtue of the application of
sequences of heat treatments including prior homogenization
of the dendritic structures. This sequence of heat
treatments comprises a first pre-homogenization treatment
of 3 hours at a temperature of' between 1300 and 1310 C,
then a progressive increase of' 3 0 C at a rate of 3 C. h-1,
before a new step of 3 hours at a temperature of between
1330 and 1340 C, it being necessary to carry out final
cooling at a rate such that tY:Le final size of the y' phase
precipitates is less than 300 nm. The entirety of the
eutectic y/y' phases is thus eliminated. It was possible
to obtain this result for all the alloys of the invention.
The sequence of heat treatments which has just been
described is an example enabling the anticipated result to
be obtained. This does not exclude the possibility of
CA 02276154 1999-06-22
obtaining a similar result by using another sequence of
heat treatments, the result of the treatment being more
important than the manner in which it is arrived at. The
important thing is to have demonstrated the possibility of
5 obtaining such a result in the case of the alloys of the
invention.
The alloys of the invention were tested after they had been
subjected to a sequence of homogenization treatments and
10 treatments for putting the T' phase into solution as
described above, and then to two annealing heat treatments
enabling the size and volume fraction of the y' phase
precipitates to be fixed. A first annealing consists of a
treatment of 4 to 16 hours at a temperature of between 1050
and 1150 C enabling the size of the y' phase precipitates
to be fixed at between 300 and. 500 nm. A second annealing
treatment consists of a treatment of 15 to 25 hours at a
temperature of between 850 and. 870'C enabling the fraction
of the precipitated T' phase to be optimized. These
annealing treatments are compatible with the diffusion
treatments of protective coatings and brazing treatments
generally applied to monocrystalline turbine blades during
their manufacture. Micrographic examination shows that the
ry' phase precipitates have a roughly cubic shape and
represent a volume fraction of at least 70 % in the alloy.
They are contained in the y matrix which appears in the
form of fine channels between these precipitates.
The high temperature resistance to creep is greater the
higher the volume fraction of the hardening ry' phase
precipitated in the alloy. At ambient temperature, the
alloys of the invention contain a volume fraction close to
70 %. When the temperature increases from ambient
temperature, the y' phase dissolves progressively in the y
matrix, slowly up to about 1000'C and then more rapidly
above 1000 C. When the T' solvus temperature is exceeded
the ry' precipitates are then totally dissolved. The
CA 02276154 1999-06-22
11
reduction in the volume fraction of the ti' phase when the
temperature increases is one of the causes of the reduction
of the creep resistance of superalloys.
One of the major benefits of the invention is to increase
substantially the y' solvus temperature so as to preserve a
high volume fraction of the y' phase at temperatures above
1100 C and thus to obtain a very high creep resistance at
these temperatures. The invention therefore concerns so-
called "high -y' solvus" alloys exhibiting a very high creep
resistance above 1100'C. Experience acquired by the
inventor in the field has shown that increases in the
concentrations of Al, Ti, Ta, Mo and W bring about an
increase in the ry' solvus. On the other hand, additions of
the elements Cr and Co lead to a reduction in the
temperature of the ry' solvus. As regards rhenium and
ruthenium, previous work has not explicitly drawn any
conclusions as to their specific action on the ry' solvus
temperature.
Increases in the concentration of elements increasing the
T' solvus may however lead to effects which can harm the
properties of the alloys. Thus too high a concentration of
the elements Al, Ti and Ta leads to the formation of an
excessive quantity of y/y' eutectic phases during
solidification of these alloys. These phases can then no
longer be totally eliminated by subsequent heat treatments,
which harms the homogeneity of the alloy and consequently
its creep resistance. Moreover, the concentration of Ta
should be limited since this element has a high atomic mass
and penalizes the alloys from the point of view of density.
The elements Mo and W also have a beneficial effect on the
y' solvus but these elements are heavy, in particular W,
and their concentration must be controlled so as not to
increase the density of the alloys excessively.
CA 02276154 1999-06-22
1.2
In addition, the solubility oi: these elements in the y
matrix is limited, in the same way as that of rhenium and
to a lesser extent those of cobalt and chromium, which may
lead to the precipitation of brittle intermetallic phases
of the a, , P or Laves phase type. The presence of these
phases, referred to as topoloclically close-packed (T.C.P.)
may bring about a loss of inechanical properties in
superalloys where they precipitate. Obtaining alloys that
are not liable to form these brittle intermetallic phases
is one of the main arguments of prior patents on
monocrystalline superalloys.
A reduction in the concentrations of the elements Cr and Co
brings about a reduction in the temperature of the y'
solvus. Thus, one of the maiii ideas of the invention is to
refrain from any addition of Co whose effect on the creep
resistance of the superalloys is small compared with that
of the other addition elements. On the other hand,
chromium has been kept since its presence is indispensable
for maintaining good hot corrosion resistance.
Examples of the invention detailed below show that the
objective of obtaining high-solvus alloys has been achieved
by virtue of a judicious choice of chemical compositions
taking into account the consicierations which have just been
expounded.
Apart from the optimization of the volume fraction and the
solvus temperature of the y' phase, an improvement to the
creep resistance of monocrystalline superalloys can be
obtained by increasing the concentrations of the refractory
elements Mo, W, Re and Ta which play an important role in
the solid solution hardening of the -y and y' phases. These
heavy elements moreover retarci all the elementary
mechanisms that are controlled by the diffusion of atoms,
which has beneficial consequences on the creep resistance
CA 02276154 1999-06-22
1.3
of the alloys. The addition of rhenium in particular
limits the growth of y' phase particles during periods
where high temperatures are maintained, a phenomenon which
participates in the degradation of the mechanical
properties of superalloys with time. Moreover, an increase
in the concentrations of refractory elements retards the
heat-activated movement of dislocations which propagate
deformation in superalloys, which has the effect of
reducing the creep rate.
The concentrations of refractory elements should however be
carefully balanced so as not t:o increase the density of the
alloys excessively.
When the elements W and Mo are present in too high a
concentration, they have a harmful effect on the oxidation
and corrosion resistance of monocrystalline superalloys
while the presence of rhenium does not penalize the
environmental resistance of these alloys.
Moreover, the refractory element Ru, within the context of
the invention, is of value in that it has a density half
that of rhenium. Work carried out by the inventor in this
field shows that Ru promotes t;he precipitation of brittle
intermetallic phases to a lesser extent than does rhenium.
The alloys according to the irivention also include
simultaneous additions of silicon and hafnium. Such
additions make it possible to optimize the hot oxidation
resistance of the alloys by inlproving the adhesion of the
protective alumina layer formed at high temperatures.
Alloys according to the invent:ion were prepared and
solidified in the form of monocrystals with a <001>
crystallographic orientation and were tested. This
crystallographic orientation was that usually selected for
the directed solidification of: monocrystalline turbine
CA 02276154 1999-06-22
14
blades. It confers on these parts an optimum combination
of creep resistance, heat fatigue resistance and mechanical
fatigue resistance.
By way of example, the nominal chemical compositions (% by
weight) of a few alloys of the: invention are collected in
table I, together with that of the reference alloy MC2
described in FR 2 557 598. This alloy serves as a
reference since it is, to the inventor's knowledge, the one
with the highest creep performance among the alloys
containing neither rhenium nor ruthenium.
Table I
Alloy Ni Co Cr Mo W Re Ru Al TTi Ta Si Hf
MC2 Base 5 8 2 8 - - 5 1.5 6 - -
MC820 Base - 5 1 F 2 - 5.5 1 6 0.1 0.1
- - -T-
MC533 Base - 7 - L5F3 3 6 - 6 0.1 0.1
MC440 Base - 5 1 4 4 - 5.5 - 9 0.1 0.1
MC722 Base - 4.5 1 7 2.5 2.5 5.8 - 6 0.1 0.1
MC623 Base - 6 1 F 2 3 5.7 0.5 5.5 0.1 0.1
MC622 Base - 5.5 1 761 2.5 2 5.9 0.5 5 0.1 0.1
MC544 Base - 4 1 5 4 4 6 0.5 5 0.1 0.1
MC645 Base - 5 - 6 4 5 6 0.5 5 0.1 0.1
MC653 Base - 4 1 6 F5-
3 5.3 1 6.2 0.1 0.1
The values of the densities of these alloys were measured
and are given in Table II. These values were in all cases
less than 8.95, and for the most part were less than 8.8.
They thus satisfied the set objective.
CA 02276154 1999-06-22
Table II
Alloy Density Ty.SOIVõS
(g. crri ) ( C)
MC2 8.62 1266
5 MC820 8.78 1300
MC533 8.64 1292
MC440 8.85 1304
MC722 8.82 1300
MC623 8.71 1294
10 MC622 8.68 1298
MC544 8.75 1292
MC645 8.75 1320
MC653 8.93 1308
15 In the raw monocrystallization state, these alloys exhibit
variable -y/-y' eutectic fractions but the application of
homogenization treatments such as previously described
makes it possible to put the T' phase precipitates back
into solution completely and to eliminate the T/T' eutectic
phases without bringing about local melting of the alloys.
The T' solvus temperatures were measured by dilatometric
thermal analysis on specimens of previously homogenized
alloys. The values of the y' solvus have been reported in
Table II. The value of the y' of solvus of the alloy MC2
measured under similar conditions is also given for
comparison in table II. The ry' solvus temperatures of the
alloys of the invention were always greater than that of
the reference alloy MC2, differences varying between 26 and
54 C according to the alloys.
Tensile creep tests were carried out on specimens machined
from various alloys of the invention into monocrystalline
bars with a <001> orientation. The bars had been
CA 02276154 1999-06-22
16
previously homogenized and then annealed according to the
procedures described previousl.y. The values of the times
to rupture for different creep conditions and for various
alloys of the invention are compared in table III with the
values obtained under the same: conditions on the
monocrystalline reference alloy MC2.
Table III
Alloy Creep conditions/life in hours
T=760 C T=950 C T=1050 C T=1100 C T=1150 C
a=840 MPa Q=300 MPa Q=150 MPa Q=130 MPa v=100 MPa
MC2 369 198 485 156 5.6
MC820 386 205 439 168 105
MC533 561 298 401 151 52
MC440 154 162 198 102 52
MC722 118 274 248 87 109
MC623 455 222 289 126 62
MC622 175 232 257 129 117
MC544 162 458 486 199 151
MC645 2105 404 499 171 185
MC653 1153 456 726 216 194
All the alloys of the examples showed a life in creep at
1150'C very much greater than that of the reference alloy
MC2. The ratio between the lives varied between
approximately 9 and 33. This result was in agreement with
the main set objective. The clain in life at this
temperature was spectacular anLd is attributed, at least in
part, to the significant increase in the ry' solvus
temperature in the alloys of the invention compared with
the reference alloy MC2.
For the other test conditions, the alloys of the invention
showed variable lives which could be greater than those of
CA 02276154 1999-06-22
17
the reference alloy MC2 according to the alloy and the
temperature considered. Remarkable results were obtained
in particular at 950 C and at 760 C in the case of certain
alloys of the invention.
The highest performance alloys were the alloys MC544, MC645
and MC653. They exhibited lives in creep at least equal
to, and generally greater than, that of the alloy MC2
within all the temperature interval considered except the
alloy MC544 at 760 C. The larcfest gains in life were
obtained at 950 and 1150 C.
Cyclic oxidation tests at 1100 C were carried out in air on
specimens of superalloys of the invention homogenized and
annealed according to the proce:dures previously described.
Each test cycle comprised a constant temperature at 1100 C
followed bv cooling to ambient temperature. The behaviour
in cyclic oxidation of the various alloys are illustrated
in the graphs of figures la and. lb where variations in the
density are given (loss in mass per unit area) of the
samples as a function of the number of one hour oxidation
cycles. Tests were carried out under the same conditions
on the reference alloy MC2. The oxidation resistance of a
superalloy improved as its density variation became lower.
All the alloys of the invention thus showed a cyclic
oxidation resistance greater than that of reference alloy
MC2.
Cyclic corrosion tests were carried out at 850 C on
specimens of alloys of the invention and the reference
alloy MC2. The specimens had been previously homogenized
and annealed according to the procedures previously
described. Each cycle comprised a constant temperature for
one hour at 850 C followed by cooling to ambient
temperature. The samples were contaminated with Na2SOd
(0.5 mg.cm'Z) every 50 hours. 'Variations in the density of
the alloy specimens are given as a function of the number
CA 02276154 1999-06-22
18
of cycles in the graphs of figures 2a and 2b. The
behaviour in corrosion was considered as satisfactory when
the mass of the specimen hardly varied at all. This was
the incubation period. An accelerated corrosion stage took
place at the end of the incubation stage. This accelerated
corrosion resulted more often in a rapid increase in mass
corresponding to the formation of corrosion products. The
graphs show a mediocre behaviou.r for the reference alloy
MC2 for which the accelerated corrosion stage took place
rapidly. The alloys of the invention showed incubation
stages with variable durations, but in all cases were
longer than that characterizing the reference alloy MC2,
which demonstrated a better resistance to cyclic corrosion.
The microstructures of the alloys of the invention were
checked at the end of the isothermal ageing treatments of
200 hours at 1050 C and at the end of creep tests taken to
rupture at 760, 950, 1050, 1100 and 1150 C so as to check
their microstructural stability as regards precipitation of
undesirable intermetallic phases of the a, or Laves phase
type. Only the alloy MC820 showed needle-like particles of
a phase rich in rhenium at the end of the ageing treatment
of 200 hours at 1050 C as well as at the end of creep tests
to rupture at 1050 and 1100 C. These particles were
localized in the cores of the dendrites, where rhenium
separated preferentially during the process of directed
solidification. All the other alloys of the invention
cited in table I were free from particles of undesirable
phases rich in rhenium at the e:nd of the ageing treatments
and creep tests.
