U.S. patent application number 11/591561 was filed with the patent office on 2007-07-19 for ni-based superalloy having high oxidation resistance and gas turbine part.
This patent application is currently assigned to Hitachi, Ltd.. Invention is credited to Hiroyuki Doi, Akira Okayama, Tsuyoshi Takano, Hideki Tamaki, Akira Yoshinari.
Application Number | 20070163682 11/591561 |
Document ID | / |
Family ID | 33410564 |
Filed Date | 2007-07-19 |
United States Patent
Application |
20070163682 |
Kind Code |
A1 |
Tamaki; Hideki ; et
al. |
July 19, 2007 |
Ni-based superalloy having high oxidation resistance and gas
turbine part
Abstract
A Ni-based alloy hardened with the .gamma. phase, which is able
to exhibit not only superior strength at high temperatures, but
also excellent hot corrosion resistance and oxidation resistance at
high temperatures in spite of containing no Re or reducing the
amount of Re. The Ni-based superalloy contains, by weight, C: 0.01
to 0.5%, B: 0.01 to 0.04%, Hf: 0.1 to 2.5%, Co: 0.8 to 15%, Ta:
more than 0% but less than 8.5%, Cr: 1.5 to 16%, Mo: more than 0%
but less than 1.0%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al: 2.5 to 7%,
Nb: more than 0 % but less than 4%, V: 0 to less than 1.0%, Zr: 0
to less than 0.1%, Re: 0 to less than 9%, at least one of platinum
group elements: 0 to less than 0.5% in total, at least one of rare
earth elements: 0 to less than 0.1% in total, and the rest being Ni
except for unavoidable impurities.
Inventors: |
Tamaki; Hideki; (Hitachi,
JP) ; Yoshinari; Akira; (Hitachinaka, JP) ;
Okayama; Akira; (Hitachi, JP) ; Takano; Tsuyoshi;
(Hitachi, JP) ; Doi; Hiroyuki; (Tokai,
JP) |
Correspondence
Address: |
MATTINGLY, STANGER, MALUR & BRUNDIDGE, P.C.
1800 DIAGONAL ROAD
SUITE 370
ALEXANDRIA
VA
22314
US
|
Assignee: |
Hitachi, Ltd.
|
Family ID: |
33410564 |
Appl. No.: |
11/591561 |
Filed: |
November 2, 2006 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
10804065 |
Mar 19, 2004 |
7169241 |
|
|
11591561 |
Nov 2, 2006 |
|
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Current U.S.
Class: |
148/428 ;
420/444 |
Current CPC
Class: |
C22C 19/056 20130101;
C22C 19/057 20130101 |
Class at
Publication: |
148/428 ;
420/444 |
International
Class: |
C22C 19/05 20060101
C22C019/05 |
Foreign Application Data
Date |
Code |
Application Number |
May 9, 2003 |
JP |
2003-130966 |
Claims
1. A Ni-based superalloy having high oxidation resistance, the
superalloy being hardened by dispersing .gamma.' phases in a
.gamma.-phase matrix, wherein the superalloy contains, by weight,
C: 0.05 to 0.2%, B: 0.01 to 0.03%, Hf: 1.1 to 2.5%, Co: 9.7 to 15%,
Ta: 0.1 to 4.5%, Cr: 1.5 to 9%, Mo: 0.01 to 0.9%, W: 5 to 14%, Ti:
0.1 to 4.75%, Al: 4 to 7%, Nb: 0.1 to less than 4%, Re: 0.01 to
less than 9%, at least one of rare earth elements: 0 to less than
0.1% in total, and any of V, Zr and platinum group elements: not
more than 0.1.
2. A Ni-based superalloy according to claim 1, wherein Ti is in the
range of 0.1 to 0.45 percent by weight.
3. A Ni-based superalloy having high oxidation resistance, the
superalloy being hardened by dispersing .gamma.' phases in a
.gamma.-phase matrix, wherein the superalloy contains, by weight,
C: 0.01 to 0.5%, B: 0.01 to 0.03%, Hf: 1.1 to 2.5%, Co: 9.7 to 15%,
Ta: less than 8.5%, Cr: 1.5 to 16%, Mo: less than 1.0%, W: 5 to
14%, Ti: 0.1 to 4.75%, Al: 4 to 7%, Nb: less than 4%, Re: 0.01 to
less than 9%, at least one of platinum group elements: 0 to less
than 0.5% in total, at least one of rare earth elements: 0 to less
than 0.1% in total, and any of V and Zr: not more than 0.1, and
wherein a value obtained from a formula of (0.004.times.W content
(weight %)+0.004.times.2.times.Mo content (weight %)+0.004.times.Re
content (weight %))/(0.003.times.3.times.Ti content (weight
%)+0.006.times.Ta content (weight %)+0.006.times.2.times.Nb content
(weight %)) is in the range of 1.0 to 2.5.
4. A Ni-based superalloy according to claim 3, wherein the value
obtained from the formula of (0.004.times.W content (weight
%)+0.004.times.2.times.Mo content (weight %)+0.004.times.Re content
(weight %))/(0.003.times.3.times.Ti content (weight
%)+0.006.times.Ta content (weight %)+0.006.times.2.times.Nb content
(weight %)) is in the range of 1.5 to 2.0.
5. A Ni-based superalloy casting having high oxidation resistance,
wherein the superalloy according to claim 1 is cast by a
unidirectional solidifying process.
6. A Ni-based superalloy casting having high oxidation resistance,
wherein the superalloy according to claim 3 is cast by a
unidirectional solidifying process.
7. A gas turbine part made from the superalloy according to claim
1.
8. A gas turbine part made from the superalloy according to claim
3.
Description
CROSS-REFERENCES
[0001] This is a continuation application of U.S. Ser. No.
10/804,065, filed Mar. 19, 2004.
BACKGROUND OF THE INVENTION
[0002] 1. Field of the Invention
[0003] The present invention relates to a Ni-based superalloy
having excellent oxidation resistance at high temperatures, and a
gas turbine part made of the Ni-based superalloy. The Ni-based
superalloy of the present invention is suitable for use in rotor
blades and stator vanes of gas turbines.
[0004] 2. Description of the Related Art
[0005] The combustion gas temperature in gas turbines tends to
increase year by year for the purpose of increasing thermal
efficiency. Correspondingly, gas turbine parts have been required
to have more superior strength, hot corrosion resistance, and
oxidation resistance at high temperatures.
[0006] Hitherto, Ni-based superalloys hardened with
.gamma.'-precipitation have been used in rotor blades and stator
vanes of gas turbines. Also, improvements of material properties of
alloys have been made by employing various chemical compositions, a
variety of content ranges, and/or various methods for producing
castings (see, for example, JP,A 6-57359 (claims), JP,A 6-184685
(claims), and Japanese Patent No. 2905473 (claims)).
SUMMARY OF THE INVENTION
[0007] Ni-based superalloys developed for use in gas turbines for
airplane engines generally contain a large amount of expensive Re
with importance put on strength at high temperatures, and contain a
small amount of Cr that is effective in improving hot corrosion
resistance. On the other hand, Ni-based superalloys developed for
use in industrial gas turbines contain large amounts of Cr and Ti
with importance put on hot corrosion resistance, and contain a
small amount of expensive Re.
[0008] In the industrial gas turbines, however, it has also been
required to employ alloys having superior strength at high
temperatures, as well as excellent hot corrosion resistance and
oxidation resistance at high temperatures, from the viewpoint of
raising the combustion gas temperature and increasing the thermal
efficiency.
[0009] Accordingly, it is an object of the present invention to
provide a Ni-based superalloy which exhibits two different natures
of characteristics having been regarded as contradictory to each
other in the past, i.e., superior creep strength at high
temperatures and excellent hot corrosion resistance and oxidation
resistance at high temperatures, in spite of containing no
expensive Re or containing a small amount of Re.
[0010] The present invention has been accomplished by conducting
studies on optimum balance among respective contents of elements
which are classified into three groups, i.e., elements Cr, Mo, W
and Re for strengthening primarily the .gamma. phase as a matrix of
the Ni-based superalloy, elements Ta, Ti and Nb for strengthening
primarily the .gamma.' phase as a precipitation hardening phase,
and elements C, B, Hf and Zr for strengthening primarily the grain
boundary, and by conducting closer studies on balance between the
.gamma.-phase strengthening elements and the .gamma.'-phase
strengthening elements in respective total contents.
[0011] The present invention resides in a Ni-based superalloy
having high oxidation resistance, the superalloy being hardened by
dispersing .gamma.' phases in a .gamma.-phase matrix, wherein the
superalloy contains, by weight, C: 0.01 to 0.5%, B: 0.01 to 0.04%,
Hf: 0.1 to 2.5%, Co: 0.8 to 15%, Ta: less than 8.5%, Cr: 1.5 to
16%, Mo: less than 1.0%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al: 2.5 to
7%, Nb: less than 4%, V: 0 to less than 1.0%, Zr: 0 to less than
0.1%, Re: 0 to less than 9%, at least one of platinum group
elements: 0 to less than 0.5% in total, and at least one of rare
earth elements: 0 to less than 0.1% in total. The other ingredient
is Ni except for unavoidable impurities, such as P and S, which are
mixed in the superalloy during the production stage.
[0012] In the present invention, the platinum group elements mean
Ru, Rh, Pd, Ir, and Pt. Among these elements, Ru is most
preferable. Also, the rare earth elements mean Sc, Y and
lanthanoid, i.e., La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er,
Tm, Yb and Lu. Among these elements, Y is most preferable.
[0013] In the Ni-based superalloy of the present invention, when
most importance is put on strength at high temperatures, the
superalloy preferably contains, by weight, C: 0.05 to 0.2%, B: 0.01
to 0.03%, Hf: 1.1 to 2.5%, Co: 9.7 to 15%, Ta: 0.1 to 4.5%, Cr: 1.5
to 9%, Mo: 0.01 to 0.9%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al: 4 to
7%, Nb: 0.1 to less than 4%, and Re: 0.01 to less than 9%. V and Zr
are intentionally not added and their contents are each held not
more than 0.005%. The rest is Ni along with the unavoidable
impurities.
[0014] When importance is put on oxidation resistance at high
temperatures of not lower than 1000.degree. C. in addition to
strength at high temperatures, the superalloy preferably contains,
by weight, C: 0.05 to 0.2%, B: 0.01 to 0.03%, Hf: 1.1 to 2.5%, Co:
9.7 to 15%, Ta: 0.1 to 4.5%, Cr: 1.5 to 9%, Mo: 0.01 to 0.9%, W: 5
to 14%, Ti: 0.1 to 0.45%, Al: 4 to 7%, Nb: 0.1 to less than 4%, Re:
0.01 to less than 9%, and at least one of rare earth elements: 0 to
less than 0.1% in total. V and Zr are intentionally not added and
their contents are each held not more than 0.005%. The rest is Ni
along with the unavoidable impurities.
[0015] When importance is put on hot corrosion resistance as well
while greater importance is put on strength at high temperatures,
the superalloy preferably contains, by weight, C: 0.05 to 0.2%, B:
0.01 to 0.03%, Hf: 1.1 to 2.5%, Co: 0.8 to 4.75%, Ta: 0.1 to 4.5%,
Cr: 1.5 to 9%, Mo: 0.01 to 0.9%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al:
4 to 7%, Nb: 0.1 to less than 4%, Re: 0.01 to less than 9%, at
least one of rare earth elements: 0 to less than 0.1% in total, and
any of V and Zr: not more than 0.005%. The rest is Ni along with
the unavoidable impurities.
[0016] When importance is put on oxidation resistance at high
temperatures of not lower than 1000.degree. C. in addition to
strength at high temperatures and hot corrosion resistance, the
superalloy preferably contains, by weight, C: 0.05 to 0.2%, B: 0.01
to 0.03%, Hf: 1.1 to 2.5%, Co: 0.8 to 4.75%, Ta: 0.1 to 4.5%, Cr:
1.5 to 9%, Mo: 0.01 to 0.9%, W: 5 to 14%, Ti: 0.1 to 0.45%, Al: 4
to 7%, Nb: 0.1 to less than 4%, Re: 0.01 to less than 9%, and at
least one of rare earth elements: 0 to less than 0.1% in total. V
and Zr are intentionally not added and their contents are each held
not more than 0.005%. The rest is Ni along with the unavoidable
impurities.
[0017] In practice, the Ni-based superalloy of the present
invention is used after performing only aging heat treatment
subsequent to casting without solution heat treatment, or after
performing solution heat treatment subsequent to casting and then
aging heat treatment.
[0018] The solution heat treatment is heat treatment for making the
.gamma.' phases dispersed in the solid-solution state in the
.gamma.-phase as the matrix. In the present invention, the solution
heat treatment may be replaced with partial solution heat treatment
in which only a part of the .gamma.' phases is brought into the
solid-solution state in the matrix.
[0019] Also, the aging heat treatment is heat treatment for
precipitating the .gamma.' phases. In the present invention, the
aging heat treatment may be performed plural times.
[0020] The solution heat treatment performed at high temperatures
is effective in increasing the strength at high temperatures, while
it imposes negative factors upon large-sized castings used in
industry-oriented gas turbines because of causing
re-crystallization, reducing the strength of grain boundary with
moving of the grain boundary, and hence pushing up the cost.
Accordingly, when superior strength at high temperatures is
required without performing the solution heat treatment, the
superalloy preferably contains, by weight, C: 0.01 to 0.5%, B: 0.01
to 0.03%, Hf: 1.1 to 2.5%, Co: 9.7 to 15%, Ta: less than 8.5%, Cr:
1.5 to 16%, Mo: less than 1.0%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al:
4 to 7%, Nb: less than 4%, Re: 0.01 to less than 9%, at least one
of platinum group elements: 0 to less than 0.5% in total, and at
least one of rare earth elements: 0 to less than 0.1% in total. V
and Zr are intentionally not added, and the rest is Ni along with
the unavoidable impurities. In addition, a value obtained from a
formula of (0.004.times.W content (weight %)+0.004.times.2.times.Mo
content (weight %)+0.004.times.Re content (weight
%))/(0.003.times.3.times.Ti content (weight %)+0.006.times.Ta
content (weight %)+0.006.times.2.times.Nb content (weight %)) is
preferably in the range of 1.0 to 2.5 and more preferably in the
range of 1.5 to 2.0.
[0021] When importance is put on hot corrosion resistance rather
than strength at high temperatures, the superalloy preferably
contains, by weight, C: 0.05 to 0.2%, B: 0.01 to 0.03%, Hf: 0.1 to
2.5%, Co: 0.8 to 15%, Ta: 0.1 to 4.5%, Cr: 9 to 16%, Mo: 0.01 to
0.3%, W: 5 to 14%, Ti: 0.1 to 4.75%, Al: 2.5 to 7%, Nb: 0.1 to less
than 4%, Re: 0 to less than 9%, and at least one of rare earth
elements: 0 to less than 0.1% in total. V, Zr and platinum group
elements are not contained. The rest is Ni along with the
unavoidable impurities.
[0022] When importance is put on hot corrosion resistance rather
than strength at high temperatures and importance is also put on
ductility, the superalloy preferably contains, by weight, C: 0.05
to 0.2%, B: 0.01 to 0.03%, Hf: 1.1 to 2.5%, Co: 0.8 to 15%, Ta: 0.1
to 4.5%, Cr: 9 to 16%, Mo: 0.01 to 0.3%, W: 5 to 14%, Ti: 0.1 to
4.75%, Al: 2.5 to 4.5%, Nb: 0.1 to less than 4%, Re: 0 to less than
9%, and at least one of rare earth elements: 0 to less than 0.1% in
total. V, Zr and platinum group elements are not contained. The
rest is Ni along with the unavoidable impurities.
[0023] When importance is put on hot corrosion resistance and a
cost reduction is intended, the superalloy preferably contains, by
weight, C: 0.05 to 0.2%, B: 0.01 to 0.03%, Hf: 0.1 to 2.5%, Co: 0.8
to 15%, Ta: less than 0.5%, Cr: 9 to 16%, Mo: 0.01 to 0.3%, W: 5 to
14%, Ti: 2 to 4.75%, Al: 2.5 to less than 4%, Nb: 0.75 to less than
4%, and at least one of rare earth elements: 0 to less than 0.1% in
total. V and Zr are intentionally not added. The rest is Ni along
with the unavoidable impurities.
[0024] When much importance is put on hot corrosion resistance, the
superalloy preferably contains, by weight, C: 0.05 to 0.2%, B: 0.01
to 0.03%, Hf: 0.1 to 2.5%, Co: 0.8 to 15%, Ta: less than 0.5%, Cr:
more than 13% but not more than 16%, Mo: 0.01 to 0.3%, W: 5 to 14%,
Ti: 2 to 4.75%, Al: 2.5 to less than 4%, and Nb: 2 to less than 4%.
V and Zr are intentionally not added. The rest is Ni along with the
unavoidable impurities.
[0025] When importance is put on hot corrosion resistance and an
alloy balanced between micro-structural stability and oxidation
resistance at high temperatures is intended, the superalloy
preferably contains, by weight, C: 0.05 to 0.2%, B: 0.01 to 0.03%,
Hf: 0.1 to 2.5%, Co: 0.8 to 15%, Ta: 0.1 to 4.5%, Cr: 9 to 16%, Mo:
0.01 to 0.3%, W: 5 to 14%, Ti: 2 to 4.75%, Al: 2.5 to less than
4.5%, Nb: 0.1 to less than 4%, Re: 0 to less than 9%, and at least
one of rare earth elements: 0 to less than 0.1% in total. V and Zr
are intentionally not added. The rest is Ni along with the
unavoidable impurities. In addition, a value obtained from a
formula of (3.8.times.Ti content (weight %)+2.times.Nb content
(weight %)+Ta content (weight %))/(2.times.Mo content (weight %)+W
content (weight %)+Re content (weight %)) is in the range of 1.6 to
2.8, and a value obtained from a formula of (3.8.times.Ti content
(weight %)+3.5.times.Cr content (weight %))/(6.8.times.Al content
(weight %)) is in the range of 1.8 to 3.1.
[0026] Further, according to the present invention, a casting made
of the Ni-based superalloy set forth above is provided. In
particular, a unidirectionally solidified casting is provided. The
Ni-based superalloy casting according to the present invention is
suitable as a high-temperature part for gas turbines, and it is
suitably used as a rotor blade or a stator vane for industrial gas
turbines.
[0027] The effects and proper content ranges of the individual
elements will be described below.
[0028] C forms MC-type carbides with Hf, Ta, Nb, Ti, etc. and forms
M.sub.23C.sub.6-- and M.sub.6C-type carbides with Cr, W, Mo, etc.,
thereby impeding movement of the grain boundary at high
temperatures, thereby impeding movement of the grain boundary at
high temperatures and strengthening the grain boundary. In order to
obtain that effect, the superalloy is required to contain C of not
less than at least 0.01% at minimum by weight and preferably not
less than 0.05%. If the C content is increased, the elements
effective in strengthening of the .gamma. phase and the .gamma.'
phase in the solid-solution state are captured by carbides and the
alloy strength at high temperatures reduces. Accordingly, the upper
limit of the C content must be restricted to 0.5 percent by weight.
When importance is put on the strength at high temperatures, it is
desired that the upper limit of the C content be set to 0.2 percent
by weight.
[0029] B has the effect of filling a non-aligned portion of the
grain boundary and increasing the bonding force of the grain
boundary. The superalloy is required to contain at least 0.01
percent by weight. However, because B noticeably lowers the melting
point of the Ni-based superalloy, the B content must be restricted
to 0.04 percent by weight at maximum. From the viewpoint of
stabilizing the strength at high temperatures, it is desired that
the upper limit of the B content be set to 0.03 percent by
weight.
[0030] Hf segregates at the grain boundary and develops the effect
of increasing ductility of the grain boundary. However, if the
alloy strength is increased, the strength of the grain boundary is
relatively reduced and the alloy ductility is noticeably reduced.
The presence of Hf is effective in preventing such a phenomenon.
Hence, the superalloy is required to contain Hf of at least 0.1
percent by weight and, in particularly, preferably not less than
1.1 percent by weight. Excessive addition of Hf, however, lowers
the alloy melting point as with B. For that reason, the upper limit
of the Hf content must be set to 2.5 percent by weight.
[0031] Co has the effect of lowering the solution temperature of
the .gamma.' phase and enabling the solution heat treatment to be
more easily performed. Particularly, when used in partial solution,
Co has the effect of increasing the volume fraction of solutioned
area even at low heat-treatment temperatures. Also, even when used
in the production process not including the solution heat
treatment, addition of Co has the effect of lowering the
precipitation temperature of the .gamma.' phase and increasing an
area in which the .gamma.' phase having an excellent shape is
precipitated. Those effects also contribute to increasing the
strength at high temperatures. In order to develop those effects,
the superalloy is required to contain Co of not less than at least
0.8 percent by weight. In the case of producing the superalloy with
much importance put on the strength at high temperatures, the Co
content is preferably not less than 9.7 percent by weight.
Excessive addition of Co, however, makes the .gamma.' phase
instable and rather leads to a reduction of the strength. For that
reason, the Co content must be held not more than 15 percent by
weight at maximum. Additionally, because of Co reducing hot
corrosion resistance, the Co content is in the range of not more
than 4.75 percent by weight when hot corrosion resistance is
required and the Cr content is less than 9 percent by weight.
[0032] Ta is an element very effective in strengthening of the
.gamma.' phase in the solid-solution state. In order to obtain
superior strength at high temperatures without performing the
solution heat treatment, an absolute value of mismatch in the
lattice constants between the .gamma.' phase and the .gamma. phase
must be kept small, and the Ta content must be more than 0% but
less than 8.5 percent by weight. From the standpoint of minimizing
the mismatch in the lattice constants, the Ta content is preferably
not more than 4.5 percent by weight. Because of Ta being an
expensive element, it is desired when importance is put on the cost
that the Ta content be less than 0.5 percent by weight while the Nb
content be increased. Replacement of a part of Ta with Nb rather
improves hot corrosion resistance.
[0033] In contrast with Ta, W is effective in strengthening of
primarily the .gamma. phase in the solid-solution state. To keep
small an absolute value of mismatch in the lattice constants
between the .gamma.' phase and the .gamma. phase, the superalloy is
required to contain W of not less than at least 5 percent by
weight. Excessive addition of W, however, deteriorates the phase
stability of the alloy and leads to precipitation of a deleterious
phase, e.g., a TCP phase, thereby noticeably reducing the hot
corrosion resistance. For that reason, the W content must be
restricted to 14 percent by weight at maximum.
[0034] Mo is an element belonging to the same group as W and has
substantially the same effect as W. In order to obtain superior
strength at high temperatures, it is desired that the superalloy
contain Mo of not less than 0.01 percent by weight. However, the
inventors have confirmed that addition of Mo deteriorates hot
corrosion resistance in combustion environment to far larger extent
than addition of W. Therefore, the Mo content in the superalloy of
the present invention is selected to be less than 1.0 percent by
weight at maximum and preferably not more than 0.9 percent by
weight. When much importance is put on hot corrosion resistance,
the Mo content is preferably not more than 0.3 percent by
weight.
[0035] As with W and Mo, Re is effective in strengthening of
primarily the .gamma. phase in the solid-solution state. Also, Re
deteriorates hot corrosion resistance in combustion environment,
but its influence is smaller than Mo and W. Re is therefore an
element very effective in realizing hot corrosion resistance and
strength at high temperatures in balance. However, Re is
distributed into the .gamma.' phase at a much lower rate and hence
tends to affect the phase stability. For that reason, the Re
content must be less than 9 percent by weight at maximum. Because
of Re being a very expensive element, it is desired that Re be
added at the necessary least content when the superalloy is used in
large-sized industrial gas turbines. When importance is put on the
cost, Re can be excluded from the added elements.
[0036] Cr forms a protective layer of Cr.sub.2O.sub.3, and it is an
essential element for maintaining hot corrosion resistance of the
Ni-based superalloy. Accordingly, the superalloy is required to
contain Cr of at least 1.5 percent by weight. When importance is
put on hot corrosion resistance, the Cr content is preferably not
less than 9 percent by weight, and when much importance is put on
hot corrosion resistance, the Cr content is preferably not less
than 13 percent by weight.
[0037] Excessive addition of Cr, however, deteriorates the phase
stability of the alloy and leads to precipitation of a deleterious
phase, e.g., a TCP phase, as with W. For that reason, the upper
limit of the Cr content must be restricted to 16 percent by weight.
When it is required to increase the W and Re contents for the
purpose of enhancing the strength at high temperatures, the Cr
content is preferably selected to be not more than 9 percent by
weight.
[0038] Al is an essential element for forming Ni.sub.3Al as the
.gamma.' phase, and the superalloy is required to contain Al of not
less than at least 2.5 percent by weight. When increasing a volume
fraction of the .gamma.' phase with importance put on strength at
high temperatures, the Al content is preferably not less than 4
percent by weight. Further, Al forms a protective layer of
Al.sub.2O.sub.3, thereby improving oxidation resistance and hot
corrosion resistance. Excessive addition of Al, however,
deteriorates the strengthening of the .gamma.' phase in the
solid-solution state and rather reduces the strength at high
temperatures. For that reason, the Al content must be held not more
than 7 percent by weight at maximum. When the Cr content is
increased with importance put on hot corrosion resistance, the Al
content is selected to be preferably in the range of 2.5 to 4.5
percent by weight and more preferably in the range of 2.5 to less
than 4 percent by weight.
[0039] Ti has the effect of preventing formation of composite
oxides of Cr and Al, thereby improving corrosion resistance of the
superalloy. Accordingly, the superalloy is required to contain Ti
of at least 0.1 percent by weight. When more importance is put on
hot corrosion resistance, the Ti content is preferably not less
than 2 percent by weight. Excessive addition of Ti, however,
impedes stability of the .gamma.' phase and deteriorates oxidation
resistance at high temperatures. For that reason, the Ti content
must be held at 4.75 percent by weight at maximum. If the Ti
content is increased, this requires the content of the other
.gamma.'-phase strengthening element, i.e., Ta, to be reduced
correspondingly, thus resulting in a reduction of the alloy
strength. Therefore, when importance is put on both strength at
high temperatures and oxidation resistance at high temperature of
not lower than 1000.degree. C., the Ti content is preferably
selected to be not more than 0.45 percent by weight.
[0040] Nb has, though not so effective as Ti, the effect of
preventing formation of composite oxides of Cr and Al, thereby
improving hot corrosion resistance of the superalloy. Also, Nb has
the effect of strengthening the .gamma.' phase in the
solid-solution state. This effect of Nb is smaller than that of Ta,
but greater than that of Ti. Accordingly, Nb is an element
effective in improving the hot corrosion resistance without
reducing the strength at high temperatures. The minimum content of
Nb may be an appreciable value. In order to effectively develop the
above-mentioned effects, however, the Nb content is preferably not
more than at least 0.1 percent by weight. When importance is put on
both the hot corrosion resistance and the cost and the Ta content
is not more than 0.5 percent by weight, the Nb content is selected
to be preferably not less than 0.75 percent by weight and more
preferably not less than 2 percent by weight. On the other hand, in
order to maintain the phase stability of the .gamma.' phase, the
upper limit of the Nb content must be restricted to less than 4
percent by weight.
[0041] Zr has the similar effect to that of Hf. However, because
addition of Zr significantly lowers the melting point of the
Ni-based superalloy, the Zr content must be held less than 0.1
percent by weight even when added. On the other hand, it has been
confirmed that Zr added in such a content range rather deteriorates
ductility of the grain boundary. Accordingly, it is most desired in
the superalloy of the present invention that Zr be intentionally
not added and the Zr content be held as low as possible at 0.005
percent by weight or below.
[0042] Addition of V lowers a limitation in solid solution of Ta
and Nb and hence leads to a reduction of the strength at high
temperatures. Also, addition of V significantly deteriorates hot
corrosion resistance. For those reasons, when added, the V content
is held less than 1.0 percent by weight and, whenever possible, not
more than 0.005 percent by weight. If possible, it is desired that
V be not added.
[0043] The rare earth elements increase adhesion of a protective
layer of Al.sub.2O.sub.3 and greatly improves oxidation resistance.
Addition of the rare earth elements, however, significantly lowers
the melting point of the Ni-based superalloy. For that reason, the
content of the rare earth elements is preferably selected to be in
the range of 0 to less than 0.1 percent by weight. The rare earth
elements mean group-3A elements of the Periodic Table, which
include Y, Sc, lanthanoid such as La and Ce, and actinoid such as
Ac.
[0044] The platinum group elements have the action of widening a
limitation in solid solution of the elements which are contained in
the superalloy and are effective in enhancing the strength at high
temperatures, such as W or Re. Because of being very expensive, the
content of the platinum group elements is selected to be less than
0.5 percent by weight. It is desired that, whenever possible, the
content of the platinum group elements be held not more than 0.005
percent by weight. The platinum group elements can be excluded from
the added elements.
[0045] The formula of (0.004.times.W content (weight
%)+0.004.times.2.times.Mo content (weight %)+0.004.times.Re content
(weight %))/(0.003.times.3.times.Ti content (weight
%)+0.006.times.Ta content (weight %)+0.006.times.2.times.Nb content
(weight %)) (numerical value obtained from this formula will be
referred to as a parameter 1 hereinafter) gives an index ratio
representing how much the lattice constants of the .gamma. phase
and the .gamma.' phase are increased respectively by the elements
(such as W, Mo and Re) for strengthening primarily the .gamma.
phase and the elements (such as Ti, Ta and Nb) for strengthening
primarily the .gamma.' phase. A coefficient prefixed to each
element represents how much the element increases the lattice
constant of the .gamma.- or .gamma.' phase per 1 atom % (unit:
10.sup.-1 nm/at %). Further, because the coefficients are set on
the premise that the mass numbers of Ta, W and Re are regarded as
substantially equal to each other, the Nb, Mo and Ti contents are
each multiplied by an additional coefficient corresponding to the
mass number ratio of W to Nb, Mo or Ti. The mismatch in the lattice
constants between the .gamma. phase and the .gamma.' phase can be
estimated based on the parameter 1, and the range where the
mismatch in the lattice constants can be kept satisfactory at high
temperatures corresponds to the range of the parameter 1 from 1.0
to 2.5. If the mismatch in the lattice constants is smaller than
1.0, the lattice constant of the .gamma.' phase would be too large,
and if it is larger than 2.5, the lattice constant of the .gamma.
phase would be too large. In either case, the mismatch in the
lattice constants cannot be kept satisfactory. In the range where
the mismatch in the lattice constants is satisfactory, the .gamma.'
phase is stable and, even in the as-cast condition, the
.gamma.'-phase holds the cubic shape. Accordingly, the superalloy
exhibits superior strength at high temperatures without the
solution heat treatment. Even in the case of using the superalloy
in the partial solid-solution state, it is also important to
control the parameter 1 because the shape of the .gamma.' phase in
the as-cast condition causes influences. Industrial gas turbines
have larger sizes than gas turbines for airplane engines. In the
industrial gas turbines, therefore, excessive residual stresses are
generated during the casting, and recrystallization tends to occur
in the subsequent solution heat treatment. Also, the strength of
the grain boundary in a unidirectionally solidified material
reduces with movement of the grain boundary to a larger extent as
the temperature of the solution heat treatment rises and the
treatment time is prolonged. For that reason, an alloy used in
high-temperature parts of the industrial gas turbines is preferably
one that is able to exhibit superior strength at high temperatures
without the solution heat treatment or with the partial
solid-solution heat treatment in which the temperature of the
solution heat treatment is set as low as possible and the treatment
time is set as short as possible.
[0046] Accordingly, the Ni-based superalloy with the parameter 1
being in the range of 1.5 to 2.5 is suitable for use as the
high-temperature parts of the industrial gas turbines. When
importance is put on particularly the strength at high
temperatures, the parameter 1 is preferably set in the range of 1.0
to 2.0.
[0047] The formula of (3.8.times.Ti content (weight %)+2.times.Nb
content (weight %)+Ta content (weight %))/(2.times.Mo content
(weight %)+W content (weight %)+Re content (weight %)) (numerical
value obtained from this formula will be referred to as a parameter
2 hereinafter) represents an atom % ratio of the elements (such as
Ti, Nb and Ta) for strengthening primarily the .gamma.' phase and
the elements (such as Mo, W and Re) for strengthening primarily the
.gamma. phase. A small value of the parameter 2 indicates that the
proportions of Mo and W adversely affecting hot corrosion
resistance are relatively large, and hot corrosion resistance tends
to deteriorate. In the case of the parameter 2 having a large
value, i.e., in the case of Ti, Nb and Ta being added in larger
amount, because these elements serve as .eta.-phase forming
elements, the .eta. phase becomes more stable than the .gamma.'
phase and the alloy strength tends to lower. To obtain excellent
hot corrosion resistance, therefore, the parameter 2 is required to
be not less than 1.6. On the other hand, to obtain superior
strength at high temperatures while keeping the .gamma.' phase
stable, the parameter 2 is required to be not more than 2.8.
[0048] The formula of (3.8.times.Ti content (weight %)+3.5.times.Cr
content (weight %))/(6.8.times.Al content (weight %)) (numerical
value obtained from this formula will be referred to as a parameter
3 hereinafter) represents an influence upon formation of an oxide
layer effective for enhancement of hot corrosion resistance. The
oxide layer is preferably formed as a multilayer made of
Cr.sub.2O.sub.3, TiO.sub.2 and Al.sub.2O.sub.3 in this order from
an outermost layer, while a care should be paid so as to avoid
formation of a composite oxide layer of those three elements. If
the parameter 3 exceeds below 1.8, a composite oxide layer made of
primarily Al and having lower protective effect would tend to form
with a reduction of the ratio of Cr and Ti to Al, thus resulting in
deterioration of hot corrosion resistance. On the other hand, if
the parameter 3 exceeds above 3.1, a stable protective layer of
Al.sub.2O.sub.3 would become hard to form with a reduction of the
ratio of Al to Cr and Ti, thus similarly resulting in deterioration
of hot corrosion resistance. For those reasons, the parameter 3 is
preferably set in the range of 1.8 to 3.1.
BRIEF DESCRIPTION OF THE DRAWINGS
[0049] FIG. 1 is a graph showing the result of creep rupture tests
made on an alloy group not subjected to the solution heat
treatment;
[0050] FIG. 2 is a characteristic graph showing the relationship
between the parameter 1 and the creep rupture time;
[0051] FIG. 3 is a graph showing the result of creep rupture tests
made on an alloy group subjected to the solution heat
treatment;
[0052] FIG. 4 is a characteristic graph showing the relationship
between the parameter 2 and the creep rupture time;
[0053] FIG. 5 is a graph showing the result of hot corrosion
resistance evaluation, based on burner rig tests, of the alloy
group subjected to the solution heat treatment;
[0054] FIG. 6 is a graph showing the result of hot corrosion
resistance evaluation, based on burner rig tests, of the alloy
group subjected to the solution heat treatment, the results being
rearranged with respect to the parameter 2 and the parameter 3;
[0055] FIG. 7 is a characteristic graph showing an effect of the Zr
content upon ductility of the grain boundary;
[0056] FIG. 8 is a characteristic graph showing an effect of the Hf
content upon ductility of the grain boundary;
[0057] FIG. 9 is a characteristic graph showing an effect of the C
content upon ductility of the grain boundary;
[0058] FIG. 10 is a characteristic graph showing an effect of the B
content upon ductility of the grain boundary;
[0059] FIG. 11 is a graph showing the result of hot corrosion
resistance evaluation, based on burner rig tests, of the alloy
group not subjected to the solution heat treatment;
[0060] FIG. 12 is a characteristic graph showing the relationship
between the Mo content and the amount of weight change after the
burner rig test;
[0061] FIG. 13 is a characteristic graph showing the relationship
between the Co content and the amount of weight change after the
burner rig test;
[0062] FIG. 14 is a characteristic graph showing the relationship
between the Nb content and the amount of weight change after the
burner rig test;
[0063] FIG. 15 is a graph showing the result of oxidation
resistance tests; and
[0064] FIG. 16 is a characteristic graph showing the relationship
between the Ti content and the amount of weight change after the
oxidation resistance tests for the alloy group subjected to the
solution heat treatment.
DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0065] Table 1, given below, lists chemical compositions and heat
treatment conditions of alloys according to the present invention
and comparative alloys employed in experiments which were conducted
during the process for accomplishing the present invention. The
alloys were classified into two groups, i.e., one in which the
alloys were subjected to the solution heat treatment and then the
aging heat treatment, and the other in which the alloys were
subjected to the aging heat treatment only with omission of the
solution heat treatment. The alloys subjected to the solution heat
treatment are of the type that importance is put on hot corrosion
resistance rather than strength at high temperatures, and the
alloys not subjected to the solution heat treatment are of the type
that importance is put on strength at high temperatures. Designing
the alloy so as to have superior strength at high temperatures
without the solution heat treatment is advantageous in preventing
recrystallization during the solution heat treatment and cutting
the cost required for the solution heat treatment.
[0066] The alloys listed in Table 1 were cast with a mold-drawing
unidirectional solidifying process by using master ingots having
respective compositions adjusted in advance. After the casting, the
alloys were subjected to the heat treatments under the conditions
shown in Table 1. Then, specimens for evaluation tests were sampled
from the alloys by mechanical machining. The specimens for
evaluation tests were each in the form of a unidirectionally
solidified slab having sizes of 100 mm.times.15 mm.times.230 mm.
The creep rupture time shown in Table 2 was evaluated under
conditions of 850.degree. C.-40 kgf/mm.sup.2 or 982.degree. C.-14
kgf/mm.sup.2. The hot corrosion resistance was evaluated based on a
weight change resulting from carrying out a burner rig test at
900.degree. C. and repeating it 7 hours.times.5 cycles. Light oil
containing 0.04 mass % sulfur was employed as fuel for the burner
rig test, and for the purpose of accelerating hot corrosion, a
solution of 1 mass % NaCl was sprayed into combustion gas at a rate
of 30 cc/min.
[0067] Furthermore, the oxidation resistance was evaluated based on
a weight change resulting from heating the specimen in the
atmosphere at a rate of 1100.degree. C./20 h and repeating the
heating 15 cycles. TABLE-US-00001 TABLE 1 Chemical composition (wt
%) alloy Cr Al Ti W Mo Ta Nb Re Co Hf Zr C B Y Ru Ni alloy C1 6.25
5.65 0.40 8.61 0.55 3.30 0.21 2.98 10.50 1.40 0 0.07 0.015 0.01 --
Bal. alloy C2 6.13 5.88 0.40 8.47 0.48 3.10 0.22 2.90 4.50 1.40 0
0.07 0.015 0.01 -- Bal. alloy C3 6.25 5.65 0.70 8.61 0.55 2.60 0.21
2.95 10.60 1.40 0 0.07 0.015 0.01 -- Bal. alloy C4 6.15 5.89 0.70
8.49 0.50 3.15 0.23 2.91 4.49 1.40 0 0.07 0.015 0.01 -- Bal. alloy
187 6.03 5.73 0.70 8.55 0.50 2.30 0.58 2.97 10.30 1.50 0 0.07 0.015
0.01 -- Bal. alloy 188 6.07 5.76 0.71 8.60 0.50 1.14 1.17 2.98
10.20 1.45 0 0.07 0.015 0.01 -- Bal. alloy 189 6.10 5.73 0.71 8.64
0.51 0.20 1.77 3.00 9.85 1.43 0 0.07 0.015 0.01 -- Bal. alloy 190
6.01 5.71 0.23 8.52 0.50 2.26 1.15 2.96 9.98 1.41 0 0.07 0.015 0.01
-- Bal. alloy 191 6.08 5.75 0.23 8.57 0.50 1.14 1.75 2.97 9.87 1.49
0 0.07 0.015 0.01 -- Bal. alloy 192 6.08 5.78 0.23 8.62 0.51 0.20
2.34 2.88 10.05 1.37 0 0.08 0.015 0.01 -- Bal. alloy 193 6.00 5.70
0.23 8.98 0.25 2.26 1.15 2.95 10.01 1.38 0 0.07 0.015 0.01 -- Bal.
alloy 194 6.03 5.73 0.23 9.03 0.25 1.13 1.75 2.97 9.97 1.37 0 0.07
0.015 0.01 -- Bal. alloy 195 6.07 5.76 0.23 9.08 0.25 0.20 2.34
2.98 9.96 1.41 0 0.08 0.015 0.01 -- Bal. alloy 196 6.07 5.57 0.80
9.08 2.59 2.89 0.21 1.50 9.88 1.28 0 0.07 0.015 0.01 -- Bal. alloy
197 6.25 5.54 0.80 10.10 1.73 2.88 0.22 1.49 9.87 1.33 0 0.08 0.015
0.01 -- Bal. alloy 200 6.20 5.75 0.70 7.50 0.85 4.48 0.21 2.95
10.00 1.40 0 0.09 0.016 0.01 -- Bal. alloy 201 6.20 5.60 1.20 7.70
0.85 4.30 0.24 2.95 10.00 1.40 0 0.05 0.020 0.01 -- Bal. alloy 298
6.05 5.71 0.70 8.56 0.50 2.25 0.50 2.95 10.00 0.11 0.005 0.10 0.010
0.01 -- Bal. alloy 299 6.05 5.77 0.68 8.54 0.50 2.27 0.56 2.99
10.05 0.12 0.012 0.10 0.010 0.01 -- Bal. alloy 300 6.08 5.75 0.68
8.55 0.50 2.28 0.56 2.98 10.05 0.12 0.016 0.10 0.010 0.01 -- Bal.
alloy 301 6.03 5.71 0.69 8.51 0.47 2.25 0.50 2.95 10.06 0.11 0.060
0.10 0.010 0.01 -- Bal. alloy 302 6.05 S.65 0.70 8.38 0.48 2.21
0.48 2.95 10.01 0 0 0.10 0.010 0.01 -- Bal. alloy 303 6.07 5.68
0.70 8.58 0.42 2.23 0.54 2.95 10.05 0.50 0 0.10 0.010 0.01 -- Bal.
alloy 304 6.10 5.45 0.70 8.33 0.49 2.21 0.49 2.95 10.03 1.02 0 0.10
0.010 0.01 -- Bal. alloy 305 6.08 5.41 0.68 8.34 0.47 2.26 0.58
2.96 10.04 1.41 0 0.10 0.010 0.01 -- Bal. alloy 306 6.03 5.33 0.75
8.31 0.52 2.24 0.62 2.99 10.00 0.11 0 0 0.010 0.01 -- Bal. alloy
307 6.07 5.40 0.71 8.55 0.50 2.29 0.60 2.99 10.05 0.12 0 0.04 0.010
0.01 -- Bal. alloy 308 6.10 5.77 0.73 8.58 0.51 2.28 0.60 2.97
10.05 0.12 0 0.10 0 0.01 -- Bal. alloy 309 6.05 5.73 0.73 8.56 0.50
2.26 0.58 2.97 10.06 0.11 0 0.10 0.007 0.01 -- Bal. alloy 310 6.03
5.75 0.70 8.58 0.50 2.27 0.59 2.96 10.05 0.11 0 0.10 0.030 0.01 --
Bal. alloy Y62 7.10 5.10 0.11 8.80 0.80 8.90 0.80 1.40 1.00 0.25 0
0.07 0.020 0 -- Bal. alloy Y62 Y 7.10 5.10 0.12 8.80 0.80 8.90 0.80
1.40 1.00 0.25 0 0.07 0.020 0 -- Bal. alloy Y63 7.10 5.10 0.12 8.80
0.80 8.90 0.80 1.40 1.00 0.25 0 0.07 0.020 0 -- Bal. alloy MM 8.25
5.56 0.75 9.65 0.50 3.05 0 0 9.50 1.40 0.007 0.07 0.015 0 -- Bal.
alloy 241 8.66 5.21 0.00 7.37 0.17 11.81 0.12 0.00 10.00 1.12 0
0.07 0.015 0.01 -- Bal. alloy 242 8.86 5.33 1.26 7.54 0.17 9.11
0.13 0.00 10.01 1.13 0 0.07 0.015 0.01 -- Bal. alloy 245 8.96 4.80
2.33 7.64 0.18 6.72 0.14 0.00 10.02 1.15 0 0.07 0.015 0.01 -- Bal.
alloy 246 9.05 4.22 3.41 7.67 0.18 4.23 0.15 0.00 10.05 1.17 0 0.07
0.015 0.01 -- Bal. alloy 247 9.19 3.69 4.57 7.83 0.18 1.76 0.16
0.00 9.97 1.18 0 0.07 0.015 0.01 -- Bal. alloy 248 9.27 3.72 4.60
7.89 0.18 0.05 1.01 0.00 9.75 1.12 0 0.07 0.015 0.01 -- Bal. alloy
251 9.75 5.25 0.00 6.37 0.17 11.91 0.14 0.00 9.76 1.11 0 0.07 0.015
0.01 -- Bal. alloy 252 9.97 5.38 1.27 6.51 0.17 9.19 0.12 0.00 9.77
1.12 0 0.07 0.015 0.01 -- Bal. alloy 255 10.10 4.84 2.35 6.59 0.18
6.78 0.11 0.00 9.78 1.11 0 0.07 0.015 0.01 -- Bal. alloy 256 10.22
4.29 3.47 6.68 0.18 4.31 0.11 0.00 9.79 1.13 0 0.07 0.015 0.01 --
Bal. alloy 257 10.35 3.73 4.61 6.76 0.18 1.77 0.10 0.00 9.80 1.14 0
0.07 0.015 0.01 -- Bal. alloy 258 10.27 4.92 2.39 6.70 0.18 3.45
1.87 0.00 9.85 1.12 0 0.07 0.015 0.01 -- Bal. alloy 259 10.33 4.34
3.50 6.75 0.18 2.18 1.21 0.00 9.88 1.11 0 0.07 0.015 0.01 -- Bal.
alloy 260 10.44 3.76 4.65 6.82 0.18 0.05 1.01 0.00 9.96 1.12 0 0.07
0.015 0.01 -- Bal. alloy 281 13.02 2.93 4.80 11.69 0.16 0.05 0.11
0.00 9.71 1.11 0 0.07 0.015 0.01 0.2 Bal. alloy 282 13.02 2.93 4.81
9.64 0.16 2.02 0.11 0.00 9.73 1.12 0 0.07 0.015 0.01 -- Bal. alloy
283 13.03 2.93 4.81 7.59 0.16 4.04 0.12 0.00 9.71 1.14 0 0.07 0.015
0.01 -- Bal. alloy 284 13.03 2.94 4.81 5.54 0.16 6.06 0.11 0.00
9.77 1.11 0 0.07 0.015 0.01 -- Bal. alloy 285 13.07 3.40 4.02 7.62
0.16 4.05 0.12 0.00 9.78 1.12 0 0.07 0.015 0.01 -- Bal. alloy 286
13.12 3.86 3.23 7.64 0.16 4.06 0.11 0.00 9.79 1.12 0 0.07 0.015
0.01 -- Bal. alloy 287 13.16 4.33 2.43 7.67 0.16 4.08 0.11 0.00
9.80 1.13 0 0.07 0.015 0.01 -- Bal. alloy 289 12.93 3.36 2.39 7.54
0.16 4.01 1.93 0.00 9.77 1.11 0 0.07 0.015 0.01 -- Bal. alloy 290
13.38 3.94 3.29 7.80 0.16 0.05 2.12 0.00 9.72 1.11 0 0.07 0.015
0.01 -- Bal. alloy 291 13.43 4.42 2.48 7.83 0.17 0.05 2.13 0.00
9.73 1.11 0 0.07 0.015 0.01 -- Bal. alloy R 14.00 3.00 5.00 4.00
4.00 0.00 0 0 9.50 0 0.030 0.17 0.015 0 -- Bal. alloy G 14.00 3.00
4.90 3.80 1.50 2.80 0 0 9.50 0 0 0.10 0.010 0 -- Bal. Heat
treatment Parameter 1 Parameter 2 Parameter 3 Solution heat
treatment Aging 1.89 0.41 0.61 -- 1080.degree. C./4 h + 871.degree.
C./20 h 1.92 0.41 0.57 1.95 0.45 0.64 1.68 0.51 0.60 1.75 0.49 0.61
1.74 0.49 0.61 1.66 0.51 0.62 1.67 0.44 0.56 1.65 0.44 0.57 1.57
0.46 0.56 1.66 0.44 0.56 1.65 0.44 0.56 1.58 0.46 0.56 2.18 0.40
0.64 -- 1080.degree. C./4 h + 871.degree. C./20 h 2.08 0.42 0.66
1.30 0.62 0.62 1.17 0.76 0.69 1.83 0.47 0.61 1250.degree. C./4 h
1.79 0.48 0.61 1.79 0.48 0.61 1.82 0.47 0.61 1.83 0.47 0.62 1.78
0.48 0.62 1.82 0.48 0.65 1.74 0.49 0.65 1.68 0.51 0.66 1.73 0.49
0.65 1.73 0.50 0.61 1.74 0.49 0.61 1.75 0.49 0.61 0.73 0.93 0.73
1280.degree. C./4 h 1080.degree. C./4 h + 871.degree. C./20 h 0.73
0.93 0.73 0.73 0.93 0.73 -- 1.59 0.55 0.84 1232.degree. C./2 h
982.degree. C./5 h + 871.degree. C./20 h 0.43 1.56 0.86
1200.degree. C./4 h 1080.degree. C./4 h + 871.degree. C./20 h 0.45
1.79 0.99 0.47 1.99 1.23 0.49 2.18 1.56 0.51 2.37 1.97 0.51 2.37
1.97 0.37 1.82 0.95 1200.degree. C./4 h 1080.degree. C./4 h +
871.degree. C./20 h 0.39 2.08 1.09 0.41 2.29 1.35 0.43 2.52 1.68
0.45 2.73 2.12 0.40 2.30 1.35 0.42 2.52 1.68 0.44 2.75 2.12 0.86
1.54 3.20 1200.degree. C./2 h 1080.degree. C./4 h + 871.degree.
C./20 h 0.59 2.06 3.20 0.40 2.85 3.20 0.26 4.19 3.20 0.45 2.46 2.64
0.51 2.08 2.21 0.60 1.69 1.88 0.42 2.16 2.38 0.52 2.07 2.21 0.61
1.68 1.88 0.85 1.58 3.33 1204.degree. C./2 h 1093.degree. C./4 h +
1051.degree. C./2 h + 843.degree. C./16 h 0.38 3.15 3.31
1121.degree. C./2 h 843.degree. C./2 h Parameter 1: (0.004 .times.
W(wt %) + 0.004 .times. 2 .times. Mo(wt %) + 0.004 .times. Re(wt
%))/(0.003 .times. 3.75 .times. Ti(wt %) + 0.006 .times. Ta(wt %) +
0.006 .times. 2 .times. Nb(wt %)) Parameter 2: (3.8 .times. Ti(wt
%) + 2 .times. Nb(wt %) + Ta(wt %))/(2 .times. Mo(wt %) + W(wt %) +
Re(wt %)) Parameter 3: (3.5 .times. Cr(wt %) + 3.8 .times. Ti(wt
%))/(6.8 .times. Al(wt %))
[0068] TABLE-US-00002 TABLE 2 Creep rupture time (h) Parameter
Parameter Parameter 982.degree. C.-14 850.degree. C.-40 Burner rig
test result Oxidation test result alloy 1 2 3 kgf/mm.sup.2
kgf/mm.sup.2 Weight change (mg/cm.sup.2) Weight change
(mg/cm.sup.3) alloy C1 1.89 0.41 0.61 1239 -- 20.15 -0.5 alloy C2
1.92 0.41 0.57 1176 -- 11.29 -0.5 alloy C3 1.95 0.45 0.64 1240 --
19.85 -0.9 alloy C4 1.88 0.46 0.60 1153 -- 10.50 -0.8 alloy 187
1.75 0.49 0.61 1224 -- 25.47 -0.7 alloy 188 1.74 0.49 0.61 1186 --
21.17 -0.8 alloy 189 1.66 0.51 0.62 1091 -- 13.06 -0.9 alloy 190
1.67 0.44 0.56 1154 -- 28.20 -0.3 alloy 191 1.65 0.44 0.57 1072 --
28.04 -0.4 alloy 192 1.57 0.46 0.56 1070 -- 15.61 -0.3 alloy 193
1.66 0.44 0.56 1236 -- 23.18 -0.4 alloy 194 1.65 0.44 0.56 1177 --
23.61 -0.4 alloy 195 1.58 0.46 0.56 1176 -- 14.71 -0.3 alloy 196
2.18 0.40 0.64 1004 -- 35.09 -1.2 alloy 197 2.08 0.42 0.66 875 --
32.30 -1.3 alloy 200 1.30 0.62 0.62 1024 -- -- -- alloy 201 1.17
0.76 0.69 858 -- -- -- alloy 298 1.83 0.47 0.61 -- -- -- -- alloy
299 1.79 0.48 0.61 -- -- -- -- alloy 300 1.79 0.48 0.61 -- -- -- --
alloy 301 1.82 0.47 0.61 -- -- -- -- alloy 302 1.83 0.47 0.62 -- --
-- -- alloy 303 1.78 0.48 0.62 -- -- -- -- alloy 304 1.82 0.48 0.65
-- -- -- -- alloy 305 1.74 0.49 0.65 -- -- -- -- alloy 306 1.68
0.51 0.66 -- -- -- -- alloy 307 1.73 0.49 0.65 -- -- -- -- alloy
308 1.73 0.50 0.61 -- -- -- -- alloy 309 1.74 0.49 0.61 -- -- -- --
alloy 310 1.75 0.49 0.61 -- -- -- -- alloy Y62 0.73 0.93 0.73 4220
-- 0.20 -4.7 alloy Y63 0.73 0.93 0.73 590 -- -- -- alloy MM 1.59
0.55 0.84 725 965 35.67 -18.0 alloy 241 0.43 1.56 0.86 -- 793 37.80
-- alloy 242 0.45 1.79 0.99 -- 731 35.20 -- alloy 245 0.47 1.99
1.23 -- 688 20.22 -- alloy 246 0.49 2.18 1.56 -- 654 10.30 -- alloy
247 0.51 2.37 1.97 -- 633 0.33 -- alloy 248 0.51 2.37 1.97 -- 615
0.21 -- alloy 251 0.37 1.82 0.95 -- 677 33.80 -- alloy 252 0.39
2.08 1.09 -- 675 28.50 -- alloy 255 0.41 2.29 1.35 -- 653 16.60 --
alloy 256 0.43 2.52 1.68 -- 648 10.50 -- alloy 257 0.45 2.73 2.12
-- 620 0.35 -- alloy 258 0.40 2.30 1.35 -- 637 10.30 -- alloy 259
0.42 2.52 1.68 -- 645 11.30 -- alloy 260 0.44 2.75 2.12 -- 625 0.33
-- alloy 281 0.86 1.54 3.20 -- 680 18.30 -200.5 alloy 282 0.59 2.06
3.20 -- 550 15.30 -120.8 alloy 283 0.40 2.85 3.20 -- 280 0.55 -78.8
alloy 284 0.26 4.19 3.20 -- 3 0.31 -38.7 alloy 285 0.45 2.46 2.64
-- 320 0.32 -30.8 alloy 286 0.51 2.08 2.21 -- 404 0.16 -25.7 alloy
287 0.60 1.69 1.88 -- 409 0.21 -23.8 alloy 289 0.42 2.16 2.38 --
538 0.25 -25.5 alloy 290 0.52 2.07 2.21 -- 541 0.18 -30.5 alloy 291
0.61 1.68 1.88 -- 546 0.19 -24.8 alloy R 0.85 1.58 3.33 -- 292
14.04 -137.8 alloy G 0.38 3.15 3.31 -- 92 0.15 -53.3 Parameter 1:
(0.004 .times. W(wt %) + 0.004 .times. 2 .times. Mo(wt %) + 0.004
.times. Re(wt %))/(0.003 .times. 3.75 .times. Ti(wt %) + 0.006
.times. Ta(wt %) + 0.006 .times. 2 .times. Nb(wt %)) Parameter 2:
(3.8 .times. Ti(wt %) + 2 .times. Nb(wt %) + Ta(wt %))/(2 .times.
Mo(wt %) + W(wt %) + Re(wt %)) Parameter 3: (3.5 .times. Cr(wt %) +
3.8 .times. Ti(wt %))/(6.8 .times. Al(wt %))
[0069] FIG. 1 shows the result of creep rupture tests made on the
alloy group evaluated without being subjected to the solution heat
treatment. In the creep rupture test, the specimen was sampled in a
direction parallel to the solidifying direction, i.e., a direction
parallel to the grain boundary. FIG. 2 shows the relationship
between the parameter 1 and the creep rupture time. As seen from
the results shown in FIGS. 1 and 2, the alloys with the parameter 1
being in the range of 1.0 to 2.5 exhibit superior creep rupture
strength even when not subjected to the solution heat treatment,
while the alloys with the parameter 1 being outside the
above-mentioned range exhibit superior creep rupture strength when
subjected to the solution heat treatment as well, but their creep
rupture strength is significantly reduced when subjected to the
aging heat treatment only.
[0070] FIG. 3 shows the result of creep rupture tests made on the
alloy group subjected to the solution heat treatment. Also in this
test, the specimen was sampled in a direction parallel to the
solidifying direction, i.e., a direction parallel to the grain
boundary. FIG. 4 shows the relationship between the parameter 2 and
the creep rupture time. As seen from FIG. 4, the creep rupture time
is shortened as the parameter 2 increases. The reason is that, when
the parameter 2 exceeds above 2.8, stability of the .gamma.' phase
is lost and precipitation of the .eta. phase is started.
[0071] FIG. 5 shows the result of hot corrosion resistance
evaluation, based on burner rig tests at 900.degree. C. (7
hours.times.5 cycles), of the alloy group subjected to the solution
heat treatment. FIG. 6 shows the hot corrosion resistance
evaluation results after being rearranged with respect to the
parameter 2 and the parameter 3. From the standpoint of creep
rupture strength, the parameter 2 is preferably set as small as
possible. From the standpoint of hot corrosion resistance, however,
the parameter 2 is preferably set as large as possible as seen from
the results of FIG. 6. To obtain not only excellent hot corrosion
resistance, but also satisfactory creep rupture strength,
therefore, it is desired that the parameter 2 be in the range of
1.6 to 2.8 and the parameter 3 be in the range of 1.8 to 3.2.
[0072] FIG. 7 shows the result of studying an effect of Zr upon
ductility of the grain boundary. Each specimen was sampled from the
unidirectionally solidified slab mentioned above, and it was
subjected to the solution heat treatment of 1250.degree. C./4 h/AC
and then two steps of the aging heat treatment, i.e., 1080.degree.
C./4.h/AC+871.degree. C./20 h/AC. The specimen was sampled in a
direction perpendicular to the solidifying direction, i.e., a
direction perpendicular to the grain boundary. After the heat
treatments, the specimen was subjected to a tension test at
800.degree. C., and the effect of Zr upon ductility of the grain
boundary was studied based on the elongation resulting from the
tension test. As seen from FIG. 7, maximum ductility is obtained
when Zr is not added.
[0073] FIG. 8 shows the result of studying an effect of Hf upon
ductility of the grain boundary. As in the above case of studying
the effect of Zr, each specimen was sampled from the
unidirectionally solidified slab mentioned above, and it was
subjected to the solution heat treatment of 1250.degree. C./4 h/AC
and then two steps of the aging heat treatment, i.e., 1080.degree.
C./4 h/AC+871.degree. C./20 h/AC. The specimen was sampled in the
direction perpendicular to the solidifying direction. After the
heat treatments, the specimen was subjected to a tension test at
800.degree. C., and the effect of Hf upon ductility of the grain
boundary was studied based on the elongation resulting from the
tension test. As seen from FIG. 8, unlike Zr, Hf is remarkably
effective in improving the ductility of the crystal grain
boundary.
[0074] FIG. 9 shows the result of studying an effect of C upon
ductility of the grain boundary. Each specimen was sampled from the
unidirectionally solidified slab mentioned above, and it was
subjected to the solution heat treatment of 1250.degree. C./4 h/AC
and then two steps of the aging heat treatment, i.e., 1080.degree.
C./4 h/AC+871.degree. C./20 h/AC. The specimen was sampled in the
direction perpendicular to the solidifying direction. After the
heat treatments, the specimen was subjected to a tension test at
800.degree. C., and the effect of C upon ductility of the grain
boundary was studied based on the elongation resulting from the
tension test. As seen from the result of FIG. 9, C is remarkably
effective in improving the ductility of the grain boundary.
[0075] FIG. 10 shows the result of studying an effect of B upon
ductility of the grain boundary. Each specimen was sampled from the
unidirectionally solidified slab mentioned above, and it was
subjected to the solution heat treatment of 1250.degree. C./4 h/AC
and then two steps of the aging heat treatment, i.e., 1080.degree.
C./4 h/AC+871.degree. C./20 h/AC. The specimen was sampled in the
direction perpendicular to the solidifying direction. After the
heat treatments, the specimen was subjected to a tension test at
800.degree. C., and the effect of B upon ductility of the grain
boundary was studied based on the elongation resulting from the
tension test. As seen from the result of FIG. 10, B is remarkably
effective in improving the ductility of the crystal grain
boundary.
[0076] FIG. 11 shows the result of hot corrosion resistance
evaluation, based on burner rig tests, of the alloy group not
subjected to the solution heat treatment. FIG. 12 shows the
relationship between the Mo content and the amount of weight change
after the burner rig test. As seen from the results of FIGS. 11 and
12, hot corrosion resistance is improved with a reduction of the Mo
content.
[0077] Further, FIG. 13 shows the relationship between the Co
content and the amount of weight change after the burner rig test.
As seen from the result of FIG. 13, hot corrosion resistance is
improved with a reduction of the Co content.
[0078] FIG. 14 shows the relationship between the Nb content and
the amount of weight change after the burner rig test. As seen from
the result of FIG. 14, Nb is effective in improving hot corrosion
resistance.
[0079] FIG. 15 shows the result of oxidation resistance tests, and
FIG. 16 shows the relationship between the Ti content and the
amount of weight change after the oxidation resistance test for the
alloy group subjected to the solution heat treatment. As seen from
the results of FIGS. 15 and 16, oxidation resistance is improved
with a reduction of the Ti content.
[0080] Thus, the present invention can provide a Ni-based
superalloy which is able to exhibit not only superior strength at
high temperatures, but also excellent hot corrosion resistance and
oxidation resistance at high temperatures in spite of containing no
Re or reducing the amount of Re.
* * * * *