Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
1~3690S
BACK~ROUND OF THE INVENTION
The invention described herein was made in the
course of a contract with the Defense Advanced Research
Projects Agency of the U.S. Government.
Field of the Invention - This invention relates to
the field of the nickel base superalloys and processes
to enhance the properties of nickel base superalloys.
Descr~iption of the Prior Art - U.S. Patent,2,542,962
to Kinsey discloses a Ni-Al-Mo alloy with broad ranges
which encompass the alloy of the present invention. Kinsey
does not discuss the possibility of heat treatments~
U.S. Patent 3,617,397 to Maxwell, assigned to the
present assignee, discloses a cast nickel base superalloy
containing 8% Al and 18% Mo, however, the patent does not
disclose the specific processing steps which form a part
of the present invention.
U.S. Patent 4,012,241 and U.S. Patent 4,111,723
which are both assigned to the present assignee, relate to
directionally solidified eutectic compositions in the Ni-Al-Mo
system.
U.S. Patent 3,975,219 to Allen et al, also assigned
to the present assignee, describes a directional recrystalli-
zation process in which a hot worked nickel base superalloy
article is recrystallized in a thermal gradient to produce an
elongated grain microstructure.
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SUMMARY OF THE I NVE~TI ON
The present invention concerns an alloy composition
and process which can be employed together to produce an
article having exceptional high temperature mechanical
properties. The broad composition range is 5-13% Al,
5-30/O Mo, balance nickel. Preferably, additions of Ta are
made. The alloy composition is prepared in the form of a
homogeneous powder which is compacted to form a fine grain
article whose grain size is stabilized by gamma prime parti-
cles at the grain boundaries. This fine grained article is
passed through a thermal gradient so as to dissolve the gamma
prime particles and to permit grain growth to occur through
a directional secondary recrystallization effect. The grains
which result in the article are elongated in the direction of
relative motion between the article and the thermal gradient.
- The article is then solution treated and aged to achieve
mechanical properties at elevated temperatures which are
greatly in excess of t~e properties of currently used nickel
base superalloys.
More specifically, according to the invention, there
is provided a method for producing a high strength aligned
grain superalloy article from an alloy which contains from
5-10% Al and from 8-21% Mo, balance nickel, including the
steps of: a~ forming the alloy into a powder, b. compacting
the powder alloy into an article; c. heating the article at a
temperature between 2200F for a period of time in excess of
one hour, and the gamma prime solvus so as to form gamma prime
particles at the grain boundaries; d. progressively and
selectively heating the article by causing relative motion
between the article and a thermal gradient, said thermal
gradient having a hot end temperature which lies between the
gamma prime solvus temperature and the incipient melting
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temperature, so that the gamma prime particles at the grain
boundaries dissolve and directional grain growth occurs in
the direction of relative motion between the article and the
thermal gradient; whereby said superalloy article has an
aligned grain structure and at a temperature of 1800F a
stress in excess of 33 ksi is required to produce more than
1% of elongation in a time period of 100 hours.
The wrought recryst,allized powder metallurgy
article according to the invention comprises elongated
aligned grains of the gamma phase which contain a fine dis-
persion of the gamma prime phase, said article having a
composition of 5-10% Al, 8-21% Mo, up to 12% Ta, up to 1% Y,
balance essentially Ni; said article being capable of with-
standing stresses of up to 33 ksi at 1800F for 100 hours
without undergoing more than 1% elongation.
The foregoing and other objects, features and ad-
vantages of the present invention will become more apparent
in the light of the following detailed description of prefer~
red embodiments thereof as illustrated in the accompanying
drawings.
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1136905
BRIEF DESCRIPTION OF THE DRAWINGS
Fig. 1 is a photomicrograph, after extrusion, of
the invention alloy;
Fig. 2 is a photomicrograph of the invention alloy
after heat treatment in preparation for directional
secondary recrystallization;
Fig. 3 is a photomicrograph of the invention alloy
after directional secondary recrystallization;
Fig. 4 is a photomicrograph after solution heat
treatment and air cooling;
Fig. 5 is a creep curve showing time to 1% elonga-
tion for various combinations of stress and temperature;
Fig. 6 is a creep curve showing elongation as a
function of time at a specific temperature and stress;
Fig. 7 is a photomicrograph of the invention alloy
containing a low Ta level after solution treatment;
Fig. 8 is a photomicrograph of the invention alloy
(lcw Ta) after solution treatment and aging; and
Fig. 9 is a photomicrograph of the invention alloy
(high Ta) after solution treatment and aging.
DESCRIPTION OF PREFERRED EMBODIMENTS
Unless otherwise noted~ all percentages listed
herein are weight percentages.
The present invention relates to a superalloy
article of specific composîtion and microstructure
which is the product of a specific process. The article
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113690S
has exceptional mechanical properties at elevat~d
temperature.
The basic alloy is a simple one based on the nickel-
aluminum-molybdenum system. The broad composition ranges
are from 5-13% Al, from 5-30% Mo, balance nickel. The
preferred ranges are 5-10% Al and 8-21% Mo. These ranges
overlap certain other alloys, in particular the alloys
disclosed in U. S. Patent 3,655,462. The present inven-
tion, however, achieves une~pected mechanical properties
through the combination of the alloy and the processing
sequence.
The basic ternary alloy has good short-term mechan-
ical properties at elevated temperatures but suffers
from microstructural instabilities which cause a marked
deterioration in properties after long-term exposures
at elevated temperatures. These instabilities include
the formation of massive gamma prime particles at the
grain boundaries and the precipitation of Mo from the
gamma phase. However, for certain applications the
basic alloy may be adequate. The addition of Ta to the
alloy in amounts of up to about 12% tends to stabilize
the microstructure and improve the long-term mechanical
properties at elevated temperatures. At least about
4% Ta appears necessary to significantly stabilize the
microstructure. Ta is a relatively dense element and
its incorporation into the alloy raises the density,
consequently a preferred Ta range is from about 4 to
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about 10%. The Ta is observed to largely substitute for
Al and when it is added the Al content may be reduced in
proportion to the atomic percent oE Ta added.
The major phases present in alloys of these compo-
sitions are the gamma and gamma prime phases. The gamma
phase (nickel solid solution~ is the matrix phase in
which are found discrete particles of the gamma prime
phase (Ni3Al) and other phases such as Ni2Mo which are
present as extremely fine dispersions. The gamma prime
phase is present in amounts of 40 to 70 volume percent.
The alloys have a gamma prime solvus temperature,
that temperature above which the gamma prime phase
dissolves into the gamma phase, which ranges from about
2320-2440F. The solidus and liquidus temperatures of
the alloys are very close together and range from about
2350-2470F. The solidus temperature preferably exceed
the gamma prime solvus temperature by at least 30F so
th~t the alloy may be solution treated without incipient
melting.
The microstructure achieved by the process to be
described below is polycrystalline microstructure and it
is well known that grain boundaries often adversely
affect the high temperature properties of materials. Such
adverse eEfects can be minimized through the addition o~
certain interstitial elements such as C and B. These
interstitial elements are believed to segregate to the
grain boundaries. Other noninterstitial elements such
li3690S
as Y, La and Ce produce similar effects. Particularly
outstandin~ results have been obtained with the addition
of about 150 ppm of Y and to a lesser extent with the
addition of abou-t .05% C to the present alloys. The
other elements have not proven to be as beneficial,
although they have not been extensively investigated.
Y, in amounts of up to 1%, and C in amounts of up to
about .1% are the preferred additions. B may be added
in amounts of up to about .05% and Zr in amounts of
about up to .1%. Other elements such as Ce and La may ^
be substituted in part for the Y addition.
Part or all of the Ta may be replaced by an
equiatomic amount of Cb, Ti or W. Likewise, part of the
; Mo may be replaced by an equiatomic amount of W or Re.
For certain applications, up to 5% Cr and up to 5% Co
might be added.
The preferred alloy consists of from 5-10% Al, from
8-21% Mo, from 4-12% Ta, up to 1% Y and up to .1% C.
The alloy previously described must be prepared
in the form of powder. This powder preparation process
is preferably one which involves the solidification of
molten metal and most preferably involves a high cooling
rate solidification process. Such a high cooling rate
helps to insure compositional homogeneity of the resulting
powder. All the work in connection with this invention
has employed a centrifugal atomization apparatus such as
that shown in U. S. Patents 4,025,249; 4,053,264 and
1~36905
4,078,873. The centrifugal atomization process involves
pouring the molten metal to be atomized on a rapidly rotating
disk. The molten metal is thrown off the edge of the disk in
a finely divided form. The particular process employed used
an extremely high rate of disk rotation, up to 35,000 rpm
and also employed auxiliary gas cooling which was provided by
a flowing curtain o~ helium gas surrounding the rotating disk
through which the atomized metal passed. The powder produced
by this process had an average particle size of 70 microns and
the rate of cooling was about 105-10 C per second. All of the
experimental work described herein employed powder of the type
described. Coarser powders or more slowly solidified powders
may require longer times at elevated temperatures to achieve
compositional homogeneity. The powder is produced and maintained
in a low oxygen environment so that the resultant powder contains
less than about 50 ppm of oxygen. The nitrogen level is also
maintained at or below 50 ppm.
The powder was compacted under conditions which caused
interparticle bounding and produced a void-free article of sub-
stantially theoretical density. Experimental work employed hotextrusion as the compaction technique. Stainless steel con-
tainers were evacuated, filled with the powder to be compacted
and sealed. The filled containers were preheated and extruded.
Extrusion was conducted at temperatures between 2200 and 2300F,
and extrusion ratios of from 6:1 to 43:1 have been successfully
employed. Fig. 1 shows the typical microstructure of the
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~3690:~
extruded powder. The exact nature of the compacting process
does not appear to be critical to the invention. other
compaction techniques including those performed at low
temperatures such as explosive compaction and those performed
at elevated temperatures where there is no significant metal
flow such as hot isostatic compaction may be employed. It
is important, however, that during the compaction process
the metal temperature is maintained below the gamma prime
solvus, so that grain growth is avoided and a fine grain
microstructure is maintained. The compacted grain size is
less than the particle size of the starting material. The
gamma prime second phase is effective to eliminate grain
boundary migration and thus maintain the initial grain size.
The material at this point has a fine uniform grain
size and it is observed that the microstructure is duplex.
The gamma phase grains have a grain size less than the start-
ing particle size. Finer particles of the gamma prime phase
are located at the grain boundaries and extremely fine
particles of gamma prime are located within the gamma phase
matrix. It is desirable that the amount of gamma prime at
the grain boundaries be maximized and this can be achieved
by a heat treatment step at a
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113~i90S
temperature bet~een about 2200F and the gamma prime
solvus temperature for a period of time in excess of
about one hour. The treatment used in the ~ork described
herein was 4 hours at 2300F. The heat treatment will
greatly increase the fraction of the gamma prime phase
found at the grain boundaries which is important in
achieving controllable and reproducible grain growth
in subse~uent process steps. This heat treatment step
forms a part of the preferred embodiment.
Of course, this heat treatment step may be combined
with other process steps, for example, if hot extrusion
is used to compact the powder, the extruded material
may be slowly cooled so that the period of time spent
between the gamma prime solvus temperature and about
2200F is sufficient to permit gamma prime phase growth
at grain boundaries.
The material at this point in the process sequence
has a uniform fine grain size and consists of grains of
the gamma phase which contain very fine gamma prime
particles on the order of .5-3 microns and larger gamma
prime phase particles on the order of 5-20 microns are
located at the grain boundaries. This microstructure
is shown in Fig. 2. The light colored phase is the
gamma prime phase. The fine grain size of ~his material
is stable at temperatures up to about the gamma prime
solvus temperature. More importantly, this fine grain
material displays superplastic behavior and may be
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~3~i905
deformed at temperatures from about 1800F to the gamma
prime solvus temperature with significant ductility and
a low flow stress.
To illustrate this behavior, a one-inch bar of
extruded RSR 104 material (composition shown in Table I)
was heat treated for 4 hours at 2300F and then hot
rolled. Several samples were hot rolled at temperatures
oE 2200 and 2300F. The samples were red~ced from a
starting thickness of 1.0 inch to a final thickness of
.030 inch. Between 12 to 15 passes were employed to
achieve the total reduction with a reduction per pass
of from .030 to .080 inch. After each pass, the material
was returned to a furnace maintained at 2200-2300F for
a period of time sufficient to bring the material tempera-
ture back up to the starting temperature. No difficulty
was encountered in hot rolling this material and this
success is somewhat surprising in view of the notorious
difficulties encountered in hot working conventional
superalloys. This behavior has important commercial
~0 implications since it will make possible fabrication of
intricate shapes at low cost. The deformation step is
of course optional and not required.
The next step is a directional grain growth step
which makes use of the phenomena known as secondary
recrystallization. An apparatus is required which can
produce a large temperature difference over a small
distance (a steep gradient). The hot end of the gradient
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exceeds t~e ga~a pri~.e solvus t~,perature DUt iS less
th2n the incipient melting te~perature. Th~ article is
sLo~71y ~oved relative to the thermal gr2dient so that it
is progressively and selectively heated to ~ te~.perature
bet~,een the ga~.a primQ sol~Jus temperatùre and the
incipient r,elting temperature. As the article is heated
within this temperature range, the ga~.a pri~e particl~s
which have stabilized the grain boundarie~ dissolve
permitting grain gro~th to occur. Grain grot~th occurs,
the driving force being a reduction in grain bou~dary
area. The grains will grow in the direction of relat.ive
motion between the article and the ther~al gradient and
the resultant grains will be elongated having a major
dimJension ~7hich is at least ten times the minor dimension.
The thermal gradient should be as steep as possiDle
in order to ~aximiz2 the directiona,lit~- o~ the grains
and mini~i~e the possibi'ity of laterzl grain gro~.7th and
resultant transverse grain boundaries. The ther~al
gradient should be at least 20F per inc'n ~.ea~ured at
the ga~lm2 prime solvus, but prererably gr~ater. The
experimenta~ work reported herein was done with a thermal
0~
/; ,,,.-j,fgradient of 150 to 300F per inch1. The Oradient T~as
obtained by using induction heatinD with a graphit2
susceptor through ~hich the article passes The raLe at
which the article c2n be ~oved rel2tive -o the ther~al
gradient 2ppears to be on the order o~ 2bcut one-hal
inch per hour. For obvious co~.~.erci~l re~sons, it is
3 ~
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~36905
desirable to move the workpiece relative to the gradient
at as rapid a rate as possible.
Following the directional secondary recrystallization
step, the material will have a rather coarse gamma prime
microstructure, within the elongated grains, as shown
in Fig. 3. This is a consequence of the slow cooling
rate through the gamma prime solvus temperature which
permits a large amount of gamma prime phase nucleation
and growth. It is preferred that as a final step the
entire article be solution heat treated at a temperature
between the gamma prime solvus temperature and the
incipient melting temperature for a period of time suffi-
cient to produce a solid solution, then rapidly cooled
and aged to produce a refined gamma prime structure.
The conditions experimentally used were 4 hours above
the gamma prime solvus followed by an air cool to room
temperature and the article was then reheated to 1950F
for 4 hours and 1600F for 12 hours. Those skilled in
the art will appreciate that these conditions may be
~0 widely varied and that they may be combined with other
operations such as coating operations which involve
heating the article. The effect of the heat treatment
is to produce a very refined dispersion of gamma prime
part;cles. Fig. 4 shows the microstructure after the
aging step. The refined microstructure contributes to
the improved mechanical properties.
Table I gives the compositions of three experimental
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alloys, processed according to the invention, denoted
as RSR alloys and two commercial prior art superalloys,
PWA 1422 and PWA 1419. RSR 104 alloy is the basic
ternary Ni-Al-Mo. RSR 166 is the same alloy composition
as RSR 104 with the addition of 150 ppm Y. RSR 143 is
a Ni-Al-Mo alloy with 6% Ta added. PWA 1422 is one of
the strongest commercially used nickel base superalloys.
PWA 1419 is a prior art nickel base superalloy containing
Ni, Mo, Al and C. This alloy is claimed in U. S. Patent
3,655,462. Both PWA 1422 and 1419 are used in the
directional solidified form as described in U. S. Patent
3,260,505. Directional solidification is a casting pro-
cess which produces a microstructure containing elongated
grains. The properties discussed below were obtained
from directionally solidified samples. Fig. 5 is a
graph which shows the creep properties of the various
superalloys in Table I. Fig. 5 shows the time required
to reach 1% elongation in creep as a function of stress
and temperature. Time and temperature are combined in
2~ the form of the Larson-Miller parameter, where T is
absolute temperature and t is time in hours. For example,
if one wished to determine the stress level required to
produce 1% creep deformation at 1800F in 100 hours, one
would determine the Larson-Miller parameter to be 49.7.
Then referring to Fig. 5 it can be seen that the stress
level required for RSR 143 material exceeds 40 ksi while
the stress level required for P~A 1422 material, a very
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~1136905
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1136~05
strong prior art superalloy, is only about 25 ksi. Even
more interesting is that the stress required to produce
1% creep in PWA 1419 is less than 20 ksi. Since PWA
1419 has a composition whicX is very similar to the RSR
alloys, this figure demonstrates the significance of the
novel processing steps employed. The difference between
the curves for RSR 104 (no Ta) and RSR 143 (6% Ta)
demonstrates the effect of Ta additions and the effect
is seen to be most pronounced at high values of the
Larson-Miller para~leter, corresponding to long exposure
times and/or high exposure temperatures.
Fig. 6 shows creep deformation behavior of the RSR
alloys and PW~ 1422 presented in a somewhat different
fashion. The curves show elongation as a function of
time at 1900F with an applied load of 14.2 ksi. (The
curve for RSR 143 is extrapolated from data obtained at
1900F and 30 ksi.) It can be seen that P~dA 1422 reaches
1% creep in about 100 hours and fails in about 250 hours.
The basic ternary RSR alloy, RSR 104, requires about 700
hours to reach 1% creep and fails in about 900 hours, a
signiicant improvement over PWA 1422. A dramatic
increase in creep li~e is seen to result when 6% Ta is
added to the basic alloy and this is shown as RSR 143
which requires more than 16,000 hours to reach 1% creep
and somewhat more than 17,000 hours to fail. These
creep results are truly outstanding when compared with
the prior art superalloys, however, both RSR 104 and
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1~L3690S
RSR 143 show somewhat low total creep ductilities. The
addition of 150 ppm of Y to the basic RSR 104 composition
results in alloy RSR 166 and the addition of this small
amount of Y raises the total creep ductility from about
3.3% to about 9.3%, a very significant increase. It is
believed that this improvement by the addition of Y
would be generally observed throughout the composition
ranges previously described.
Further evidence for the beneficial effect oE Ta on
the microstructural stability of the present alloys can
be seen in Figs. 7, 8 and 9. Fig. 7 shows alloy RSR 144
(composition similar to alloy RSR 143 but containing 3%
Ta and 7% Al) after a solution treatment at 2400F and
a rapid cooling step. The grain boundaries are seen to
contain small amounts of the light colored gamma prime
phase. The microstructure of RSR 143 after a similar
solution treatment is very similar. Fig. 8 shows the
microstructure of RSR 144 solution treated at 2400F
and aged for 50 hours at 2000F. Massive amounts of
the gamma prime phase can be seen at the grain boundaries
after this aging step and it is believed that this pre-
cipitation of the gamma prime phase at the grain bound-
aries adversely affects the high temperature mechanical
properties. Fig. 9 shows the microstructure of RSR 143
after a 2400F solution treatment and 50 hours at 2000F.
The difference between Fig. 8 and Fig. 9 is dramatic
since the massive gamma prime phase which is evident in
1136905
Fig. 8 is completely absent from Fig. 9. Thus, it is
evident that a Ta level somewhere between 3 and 6 weight
percent, e.g. 4%, is effective in suppressing precipita-
tion of the gamma prime phase at the grain boundaries in
the alloys of the invention.
Although this invention has been shown and described
with respect to a preferred embodiment thereof, it should
be understood by those skilled in the art that various
changes and omissions in the form and detail thereof may
be made therein ~ithout departing from the spirit and
scope of the invention.
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