Method for producing a nickel-base superalloy

A method of forming a Ni-base superalloy suitable for use as the material for gas turbine disks or the like has a composition containing, by weight, 0.01 to 0.15% of C, 15 to 22% of Cr, 3 to 6% of Mo, 3 to 6% of W, 5 to 15% of Co, 1.0 to 1.9% of Al, 1.5 to 3.0% of Ti, 3.0 to 6.0% of Ta, 0.001 to 0.020% of B and the balance substantially Ni except inevitable impurities. This alloy is produced using the conventional ingot making and a hot working process including working at a reducing ratio greater than or equal to 10%, first above the .gamma. solvus temperature, and then during cooling to the recrystallization temperature and then subjected to direct aging without solid-solution treatment. As a result, the alloy exhibits excellent strength properties well comparable to those of expensive alloys produced by powder metallurgy process.

BACKGROUND OF THE INVENTION 
1. Field of the Invention 
The present invention relates to an Ni-base superalloy (i.e., super heat 
resisting alloy) which is suitable for use as the material for disks or 
the like of a gas turbine, which can be hot worked and which has a high 
strength comparable to that of powder metallurgy alloy, and to a method 
for producing the same. 
2. Description of the Prior Art 
Current trends for greater output and higher efficiency of gas turbines 
naturally require that heat resisting parts of gas turbines operate under 
severer conditions. In case of disks of gas sturbines, an increasing 
demand exists for improvement in the mechanical strength of the material 
for disks rather than for a rise in the maximum withstandable temperature 
of the disks. Thus, the following two kinds of approaches have been made 
to increase the performance of turbine disk material. 
(1) Development of novel alloy having high .gamma.' phase content by powder 
metallurgy process. 
(2) Improvement in the strength of existing ingot alloy by thermomechanical 
treatment. 
As an example of Ni-base superalloy according to the approach (1), such a 
high strength alloy having .gamma.' phase content of about 50 vol.% as 
known under the name of RENE' 95 (RENE' being a trademark) or IN 100 (IN 
being a trademark) has been put to practical use in commercial base. 
The RENE' 95 is an alloy which is disclosed in Japanese Patent Examined 
Publication No. 46-22333. Initially, it was attempted to fabricate this 
alloy by the conventional ingot making and subsequent hot working process. 
This attempt, however, was unsuccessful because of difficulty in 
fabricating this alloy from the ingot material due to containing a large 
amount of .gamma.' phase, so that this alloy is fabricated only by powder 
metallurgy process at present. On the other hand, the IN 100 has been 
developed as a cast alloy from the beginning, so that no attempt has been 
made to commercially fabricate this alloy by the ingot making and hot 
working process. 
Further, Japanese Patent Unexamined Publication No. 63-114933 discloses an 
alloy which exhibits superior properties as a material for gas turbine 
disks. This alloy, however, also is a high .GAMMA.' alloy containing about 
45% of .gamma.' phase and, therefore, cannot be fabricated by the 
conventional ingot making and hot working process. 
Thus, an alloy having high .gamma.' phase content becomes impossible to be 
hot worked and, hence, is obliged to adopt powder metallurgy process. The 
powder metallurgy process, however, employs a number of steps so that the 
price is raised uneconomically. In addition, the powder metallurgy process 
tends to allow the product to contain oxides, impairing the reliability of 
the product. 
According to the approach (2) mentioned above, thermomechanical treatment, 
which is a combination of a hot working and a heat treatment, is effected 
on an Ni-base superalloy such as WASLOY (WASLOY being a trademark) 
or INCONEL 718 (INCONEL being a trademark), in order to achieve desired 
performance. Alloys obtained by such thermomechanical treatment exhibit 
mechanical properties which are excellent in comparison with conventional 
ingot alloys but are still inferior in comparison with those exhibited by 
supperalloys produced by the powder metallurgy process of the aforesaid 
approach (1). 
Further, Japanese Patent Unexamined Publication No. 63-145737 discloses an 
alloy which is said to be a high-strength ingot alloy having .gamma.' 
phase content of 45 vol.% and exhibiting superior hot workability. 
However, it is very difficult to hot work this alloy and an extremely high 
degree of forging technique is required due to the .gamma.' phase content 
which is much higher than that of existing ingot alloy. 
Considering merits and demerits of the aforementioned approaches (1) and 
(2) for increasing the performance of disk material, it is highly 
desirable to develop an alloy which can be produced by a process making 
use of existing production equipment, e.g., a process having the steps of 
conventional ingot making and hot working, and which has properties well 
comparable to those of alloys produced by powder metallurgy process, 
because such an alloy will enable inexpensive production of large-sized 
parts having high reliability. 
SUMMARY OF THE INVENTION 
Accordingly, an object of the present invention is to provide a high 
strength Ni-base superalloy which exhibits, despite a reduced .gamma.' 
phase content, a high strength well comparable to those of alloys produced 
by powder metallurgy process and which has excellent hot workability to 
enable easy production by conventional ingot making and hot working 
process. 
Another object of the present invention is to provide a method for 
producing such a high strength Ni-base superalloy. 
The present inventors have conducted an intensive study on alloy 
compositions suitable for use as materials of gas turbine disks, as well 
as on production methods, and found that an Ni base superalloy having high 
strength well comparable to those of powder metallurgy alloys and 
excellent hot workability can be obtained with a specific alloy 
composition even though the .gamma.' phase content is reduced to less than 
40 vol.%. 
Namely, the present invention provides a hot workable Ni-base superalloy 
which can be produced by ingot making process and which is characterized 
by having excellent properties, in particular high strength, well 
comparable to those of alloys which, in an alloy system of this field, 
hitherto could not be obtained by ingot making and hot working process 
and, therefore, were produced by powder metallurgy process. 
According to a first aspect of the present invention, there is provided an 
Ni-base superalloy containing, by weight, 0.01 to 0.15% of C, 15 to 22% of 
Cr, 3 to 6% of Mo, 3 to 6% of W, 5 to 15% of Co, 1.0 to 1.9% of Al, 1.5 to 
3.0% of Ti, 3.0 to 6.0% of Ta, 0.001 to 0.020% of B and the balance 
substantially Ni except inevitable impurities. 
According to a second aspect of the present invention, there is provided an 
Ni-base superalloy containing, by weight, 0.01 to 0.05% of C, 17 to 19% of 
Cr, 4 to 5% of Mo, 4 to 5% of W, 8 to 12% of Co, 1.1 to 1.6% of Al, 2.1 to 
2.7% of Ti, 4.2 to 5.0% of Ta, 0.005 to 0.015% of B and the balance 
substantially Ni except inevitable impurities. 
According to a third aspect of the present invention, there is provided an 
Ni-base superalloy containing, by weight, 0.01 to 0.15% of C, 15 to 22% of 
Cr, 3 to 6% of Mo, 3 to 6% of W, 5 to 15% of Co, 1.0 to 1.9% of Al, 1.5 to 
3.0% of Ti, Ta and Nb in an amount which meets the conditions of 
3.0%.ltoreq.Ta+2Nb.ltoreq.6.0% and Ta.gtoreq.2Nb, 0.001 to 0.020% of B and 
the balance substantially Ni except inevitable impurities. 
According to a fourth aspect of the present invention, there is provided an 
Ni-base superalloy according to any one of the aforesaid first to third 
aspects, characterized in that the .gamma.' phase having a composition 
expressed by Ni.sub.3 (AlxTiyTaz) (x+y+z=1) or Ni.sub.3 (AlxTiyTazNbw) 
(x+y+z+w=1) is contained in an amount not greater than 40 vol.%, that the 
0.2% offset yield strength at 650.degree. C. is higher than 120 
kgf/mm.sup.2, and that the creep rupture time at 650.degree. C. and 100 
kgf/mm.sup.2 is longer than 80 hours. 
According to a fifth aspect of the present invention, there is provided a 
method for producing an Ni-base superalloy comprising the steps of: 
preparing an alloy according to any one of the aforesaid first to fourth 
aspects; subjecting said alloy to a final hot working in which said alloy 
is heated to and held at a temperature which is 20 to 100.degree. C. 
higher than the .gamma.' phase's solvus temperature and then hot worked at 
reduction ratio of 10% or greater during cooling to the recrystallization 
temperature and subsequently at reduction ratio of 10% or greater at 
temperatures lower than the recrystallization temperature; and directly 
aging said hot worked alloy at a temperature lower than 850.degree. C. 
without subjecting it to solid-solution heat treatment. 
According to the present invention, the contents of the respective alloying 
components are limited for the following reasons. 
In the present invention, C serves as a deoxidizer and, in addition, forms 
MC type carbides in combination with Ti, Ta and Nb. Further, when an aging 
without solid-solution heat treatment (hereafter, this aging is referred 
to "direct aging") is conducted, C discontinuously precipitates in grain 
boundaries M.sub.23 C.sub.6 type carbides composed mainly of Cr, thereby 
strengthening the grain boundaries and thus improving creep rupture 
properties. In order to obtain these effects, the C content should be 
0.01% at the smallest. On the other hand, any excessive C content 
exceeding 0.15% increases formation of primary carbides, thereby 
deteriorating the toughness. For these reasons, the C content should be 
limited to a range between 0.01 and 0.15%, preferably between 0.01 and 
0.05%. 
Cr is an element indispensable for obtaining oxidation resistance and 
corrosion resistance at high temperatures, and in order to meet oxidation 
resistance and corrosion resistance necessary for gas turbine disks, etc. 
the Cr content should be 15% at the smallest. On the other hand, if the Cr 
content exceeds 22% the structure becomes unstable and it becomes liable 
to form .sigma. phase, which is an brittle phase, in combination with Mo 
and W. For these reasons, the Cr content is limited to a range between 15 
and 22%, preferably between 17 and 19%. 
Mo is an element which dissolves into austenite phase so as to strengthen 
the matrix, thereby improving the strength at high temperatures. In order 
to obtain this effect, the Mo content should be 3% at the smallest. On the 
other hand, an excessive Mo content impairs the hot workability and, in 
addition, makes the structure unstable as Cr does, so that the upper limit 
of the Mo content is limited to 6%. Preferably, the Mo content is limited 
to a range between 4 and 5%. 
W is an element which dissolves into the matrix to thereby improve the 
tensile strength as Mo does. However, W exhibits a smaller diffusion rate 
than Mo because W has an atomic weight which is about two times that of 
Mo, so that W makes a greater contribution to the reduction in the creep 
rate than Mo, thereby improving the creep rupture life. In order to obtain 
the above effect, the W content should be 3% at the smallest. On the other 
hand, addition of W in excess of 6% adversely affects hot workability and 
stability of the structure as Mo does and undesirably increases the 
specific weight of the alloy. For these reasons, the W content is limited 
to a range between 3 and 6%, preferably between 4 and 5%. 
Co increases the amount of .gamma.' phase putting into solid solution at 
high temperature range so as to improve the hot workability. In order to 
obtain this effect, the Co content should be 5% at the smallest. However, 
an excessive Co content tends to cause precipitation of detrimental phases 
such as Laves phase or the like, so that the upper limit is limited to 
15%. Preferably, the Co content is limit to a range between 8 and 12%. 
Al is an indispensable element which allows precipitation of stable 
.gamma.' phase in combination with Ni, thereby obtaining the desired high 
temperature strength. In order to obtain this effect, the Al content 
should be 1.0% at the smallest. In the alloy of the present invention, in 
order to improve the high temperature strength, it is necessary that the 
lattice strain owing to precipitation of .gamma.' phase be increased by 
increasing the ratio {Ti+Ta(+Nb)}Al in the .gamma.' phase to thereby 
increase the lattice constant of the .gamma.' phase. To this end, the 
upper limit of the Al content is limited to 1.9%. Preferably, the Al 
content is limited to a range between 1.1 and 1.6%. 
Ti is an element which, like Al, allows precipitation of .gamma.' phase in 
combination with Ni, thereby increasing the high temperature strength. In 
order to obtain this effect, the Ti content should be 1.5% at the 
smallest. On the other hand, addition of Ti in excess of 3.0% 
inconveniently reduces the solid soluvility of Ta, which is an important 
element in the alloy of the present invention, into the .gamma.' phase, 
and undesirably allows precipitation of .eta. phase (Ni.sub.3 Ti) content 
is limited to a range between 1.5 and 3.0%, preferably between 2.1 and 
2.7%. 
One of the novel features of the alloy of the present invention over 
conventional alloys is based upon discovery of superior effect of Ta on 
creep rupture properties. In general, maximum operation temperature of 
disks of current gas turbines is around 650.degree. C., and Ta acts very 
effectively in such temperature range. Like Ti mentioned above, Ta 
dissolves into Al side of Ni.sub.3 Al, thereby increasing the lattice 
constant of .gamma.' phase and thus improving the tensile strength. 
Further, with respect to agglomerating rate of the .gamma.' phase, Ta has 
an effect of retarding grain growth of the .gamma.' phase at a temperature 
range of about 650.degree. C. because it has a larger atomic weight than 
another elements constituting the .gamma.' phase, so that it is effective 
for remarkably prolonging the creep rupture life. Ta belongs to the same 
group of the periodic table as Nb and has been considered to provide 
almost an equivalent effect on improvement of mechanical properites of 
Ni-base superalloy. The present inventors have found, however, that Ta 
produces, due to the fact that the atomic weight of Ta is two times that 
of Nb, a more advantageous effect on the agglomerating rate of .gamma.' 
phase than Nb and, hence, a greater effect in improving creep rupture 
strength. The present invention makes an effective use of this newly found 
advantage of Ta. 
In order to obtain the above effect, the Ta content should be 3.0% at the 
smallest. On the other hand, addition of Ta in excess of 6.0% adversely 
affects the hot workability and undesirably degrades ductility due to 
precipitation of the .delta. phase (Ni.sub.3 Ta). For these reasons, the 
Ta content is limited to a range between 3.0 and 6.0%, preferably between 
4.2 and 5.0%. 
Nb is an element belonging to the same group as Ta and produces a similar 
effect on improvement in the high temperature strength. The effect of Nb 
on improvement in the creep rupture life is not so remarkable as Ta. 
However, since Nb can be substituted with Ta at an atomic ratio up to 1:1 
without causing substantial degradation in the properties, the Nb content 
is limited to a range which meets the conditions of 
3.0.ltoreq.Ta+2Nb.ltoreq.6.0 and Ta.gtoreq.2Nb. 
B is effective, owing to its effect for strengthening the grain boundaries, 
in improving both high temperature strength and ductility. In order to 
obtain this effect, the B content should be 0.001% at the smallest. 
However, the B content exceeding 0.020% causes the initial melting 
temperature of the alloy of the present invention to be lowered, thereby 
deteriorating the hot workability. For these reasons, the B content is 
limited to a range between 0.001 and 0.020%, preferably between 0.005 and 
0.015%. 
Further, in many of Ni-base supperalloys, Zr is considered to be an element 
which, like B, strengthens the grain boundaries but Zr is fundamentally 
different from B in that it is a primary carbide former. The important 
feature in the alloy of the present invention resides in the fact that the 
grain boundaries are strengthened by precipitation of suitable amount of 
M.sub.23 C.sub.6 type carbides, so that in the alloy of the present 
invention no Zr is added, because if Zr were added the precipitation of 
the M.sub.23 C.sub.6 type carbides at the grain boundaries would be 
decreased. 
Ni is a basic element which constitutes an austenite matrix and a .gamma.' 
precipitation strengthening phase which is Ni.sub.3 (Al, Ti, Ta) or 
Ni.sub.3 (Al, Ti, Ta, Nb). 
Although inclusion of impurities such as Fe, Si, Mn, P, S, Mg, Ca, Zr and 
so forth is inevitable in the alloy of the present invention, such 
impurity elements may be contained if the contents of these elements meet 
the following conditions, because inclusion of such small amounts of 
impurity elements does not adversely affect the properties of the alloy. 
Fe.ltoreq.3.0% 
Si.ltoreq.0.5% 
Mn.ltoreq.1.0% 
P.ltoreq.0.03% 
S.ltoreq.0.03% 
Mg.ltoreq.0.02% 
Ca.ltoreq.0.02% 
Zr.ltoreq.0.01% 
In addition to the limitations on the content ranges of the respective 
elemens described above, in the alloy of the present invention the upper 
limit for the content of the .gamma.' phase composed of Ni in combination 
with Al, Ti and Ta or Ni in combination with Al, Ti, Ta and Nb is limited 
to 40 vol.%, in order to provide the alloys with an excellent hot 
workability when it is produced by the conventional ingot making and hot 
working process. It is possible to limit the .gamma.' phase content to 
less than 40 vol.% by controlling the amounts of the .gamma.' phase 
formers. 
Further, the alloy of the present invention can exhibit excellent 
properties applicable to the material for gas turbine disk, etc. by the 
production method mentioned below. Namely, the alloy of the present 
invention has recrystallization temperature in a range of 
1020.degree.-1050.degree. C. and thus exhibits excellent hot workability 
at temperatures higher than this recrystallization temperature. However, 
since the .gamma.' phase s solvus temperature (i.e., the temperature at 
which the .gamma.' phase completely dissolves into the matrix) of the 
alloy of the present invention is in a temperature range of 
1075.degree.-1120.degree. C., when the alloy is hot worked at a 
temperature higher than the recrystallization temperature but lower than 
the .gamma.' phase's solvus temperature it exhibits an excellent hot 
workability, but in this case nonuniform precipitation of the .gamma.' 
phase remains, so that it is undesirable from the viewpoints of structure 
and mechanical properties. Further, when the alloy is heated at a 
temperature higher than the .gamma.' phase's solvus temperature the 
nonuniformly precipitated .gamma.' phase is completely dissolved into the 
matrix and, as a result, the crystal grains become easy to grow, but in 
this case it exhibits a more excellent hot workability than when it is hot 
worked at a temperature higher than the recrystallization temperature but 
lower than the .gamma.' phase's solvus temperature and its microstructure 
after the hot working becomes uniform. For these reasons, at several heats 
in the initial stage of the hot working the alloy is plastically worked at 
a heating temperature higher than the .gamma.' phase's solvus temperature, 
at which it exhibits an extremely excellent hot workability, into a form 
approximating the desired shape in some extent and then, at an 
intermediate stage of the hot working, it is hot worked after having been 
heated for the purpose of grain refinement at a temperature range higher 
than the recrystallization temperature but lower than the .gamma.' phase's 
solvus temperature. Subsequently, it is heated for a short period of time 
in advance of the final hot working at a temperature which is 20.degree. 
to 100.degree. C. higher than the .gamma.' phase's solvus temperature so 
as to dissolve the nonuniformly precipitated .gamma.' phase into the 
matrix to thereby suppress as much as possible the growth of the crystal 
grains, and then it is finally hot worked. 
More specifically, the alloy material to be worked, which has been heated 
to a temperature which is 20 to 100.degree. C. higher than the .gamma.' 
phase's solvus temperature prior to the final hot working, is worked at a 
reduction ratio of 10% or greater in the course of cooling to the 
recrystallization temperature, and subsequently worked at a reduction 
ratio of 10% or greater at a temperature lower than the recrystallization 
temperature so as to refine the crystal grains and impart a sufficient 
work strain. Incidentally, the term "reduction ratio" is used in this 
specification to mean the degree of the working effected on the alloy 
material. When the working is effected to reduce the cross-sectional area 
while increasing the length of the alloy material, the reduction ratio is 
expressed as follows: 
EQU {(A-a)/A}.times.100% 
where A and a respectively represent the cross-sectional area before and 
after the working. On the other hand, when the working is effected to 
reduce the length of the alloy material while increasing the 
cross-sectional area, i.e., an upset forging, the reduction ratio is 
expressed as follows: 
EQU {(L-l/L}.times.100% 
where L represents the original length of the material while l represents 
the length after the working. 
When the heating temperature exceeds the temperature range which is 
20.degree. to 100.degree. C. higher than the .gamma.' phase's solvus 
temperature the coarsening of the crystal grains is promoted and, on the 
other hand, when it is too low the .gamma.' phase is not completely 
dissolved into the matrix. In contrast to this, when the reduction ratio 
of the working effected during cooling to the recrystallization 
temperature is less than 10% it is impossible to satisfactorily refine the 
crystal grains and, on the other hand, when the reduction ratio of the 
working effected at temperatures lower than the recrystallization 
temperature is less than 10% the work strain becomes insufficient, so that 
it becomes impossible to obtain the desired strength. For these reasons, 
the reduction ratio is limited to 10% or greater. 
Further, with respect to the heat treatment, a direct aging is effected 
without solid-solution heat treatment, in order to make use of the 
strengthening effect obtained in the crystal grains and grain boundaries 
owing to the work strain derived from the hot working. Since the aging has 
to be conducted at a temperature range in which the effect of the work 
strain is not extinguished, the upper limit temperature for the aging is 
limited to 850.degree. C. One of the purposes of the aging is to cause a 
sufficient precipitation of fine .gamma.' phase in the grains, while 
another purpose is to precipitate M.sub.23 C.sub.6 type carbides at the 
grain boundaries. In the case of direct aging the M.sub.23 C.sub.6 type 
carbides are more easily precipitated at the grain boundaries in 
comparison with aging conducted after a solid-solution heat treatment and, 
in addition, they are precipitated in discontinuous and granular form, 
thereby strengthening the grain boundaries and greatly contributing to the 
improvement in the creep rupture life.

DESCRIPTION OF THE PREFERRED EMBODIMENTS 
EXAMPLE 1 
Each of the alloys of compositions shown in Table 1 was melted in a vacuum 
induction melting furnace and casted into a ingot of 10 kg. The ingot was 
soaked at 1200.degree. C. for 20 hours and forged into a 30 mm square rod. 
The forging was conducted in four heats, wherein the first and fourth 
heats were executed by heating at 1150.degree. C., while the second and 
third heats were executed by heating in the temperature range between 
1050.degree. C. and 1070 .degree. C. In the fourth heat, the working was 
executed at a reduction ratio of 25% in the temperature range between 
1150.degree. C. and 1030.degree. C. and, further, at a reduction ratio of 
15% in the temperature range between 1030.degree. C. and 980.degree. C. 
The alloys according to the present invention and the comparison alloys 
Nos. 21, 22 and 24 exhibited excellent hot workability, but the comparison 
alloy No. 23 whose .gamma.' phase content is 41.8 vol.% was cracked during 
the forging and the forging was stopped. 
In this Example, although forging was adopted as the hot working, it is 
needless to say that hot rolling may be adopted. 
TABLE 1 
__________________________________________________________________________ 
.gamma.' phase 
Hot* 
Alloy 
Chemical composition (wt. %) content 
worka- 
No. C Cr Mo W Co Al Ti Nb Ta B Ni (vol. %) 
bility 
Remarks 
__________________________________________________________________________ 
1 0.033 
18.1 
4.59 
4.70 
10.5 
1.35 
2.38 
-- 4.80 
0.010 
Bal. 
30.2 .smallcircle. 
Alloy of 
invention 
2 0.033 
18.5 
4.65 
4.50 
10.4 
1.73 
1.92 
-- 4.00 
0.010 
" 29.9 .smallcircle. 
Alloy of 
invention 
3 0.030 
18.4 
4.59 
4.61 
10.5 
1.34 
2.46 
0.95 
2.79 
0.010 
" 30.0 .smallcircle. 
Alloy of 
invention 
4 0.029 
18.2 
4.65 
4.50 
10.5 
1.34 
2.36 
-- 4.46 
0.009 
" 29.5 .smallcircle. 
Alloy of 
invention 
5 0.033 
18.1 
4.61 
4.67 
10.4 
1.31 
2.34 
-- 4.77 
0.010 
" 29.6 .smallcircle. 
Alloy of 
invention 
6 0.032 
18.0 
4.54 
4.86 
10.5 
1.25 
2.34 
-- 4.70 
0.010 
" 29.6 .smallcircle. 
Alloy of 
invention 
7 0.033 
18.1 
4.62 
4.69 
6.3 
1.33 
2.41 
-- 4.75 
0.010 
" 30.2 .smallcircle. 
Alloy of 
invention 
8 0.029 
18.0 
4.55 
4.62 
13.5 
1.37 
2.30 
-- 4.55 
0.011 
" 30.0 .smallcircle. 
Alloy of 
invention 
9 0.029 
18.9 
3.55 
3.40 
11.0 
1.35 
2.40 
-- 4.59 
0.009 
" 30.1 .smallcircle. 
Alloy of 
invention 
10 0.033 
17.8 
5.45 
5.59 
10.3 
1.30 
2.35 
-- 4.52 
0.011 
" 30.0 .smallcircle. 
Alloy of 
invention 
11 0.030 
21.2 
4.61 
4.66 
10.4 
1.35 
2.39 
-- 4.68 
0.010 
" 30.5 .smallcircle. 
Alloy of 
invention 
12 0.029 
19.1 
4.81 
4.46 
11.0 
1.40 
2.31 
0.30 
4.11 
0.011 
" 30.5 .smallcircle. 
Alloy of 
invention 
13 0.032 
18.0 
4.60 
4.38 
10.2 
1.55 
2.65 
-- 5.13 
0.009 
" 34.3 .smallcircle. 
Alloy of 
invention 
21 0.033 
18.9 
4.72 
4.48 
10.9 
1.37 
3.80 
-- -- 0.010 
" 30.0 .smallcircle. 
Comparison 
alloy 
22 0.031 
18.6 
4.72 
4.51 
10.8 
1.36 
2.73 
2.62 
-- 0.011 
" 31.7 .smallcircle. 
Comparison 
alloy 
23 0.035 
18.0 
4.30 
4.35 
11.0 
1.75 
3.20 
-- 7.00 
0.009 
" 41.8 x Comparison 
alloy 
24 0.032 
17.9 
4.58 
4.94 
10.6 
1.22 
2.31 
-- 4.89 
0.010 
" 29.5 .smallcircle. 
Comparison 
alloy 
(Zr: 0.05) 
__________________________________________________________________________ 
*Note: Marks .smallcircle. and x represent, respectively, nonoccurrence o 
cracking and occurrence of cracking during forging. 
EXAMPLE 2 
Tables 2 and 3 show influence of a heat treatment on tensile properties and 
creep rupture properties of the alloy No. 2 of the present invention. In 
the solid-solution heat treatment, the alloy was heated to and held at 
1000.degree. C. for 2 hours followed by oil quenching. The aging treatment 
was conducted in two steps: namely, heating at 650.degree. C. for 24 hours 
followed by air cooling and heating at 760.degree. C. for 8 hours followed 
by air cooling. From Table 2, it will be seen that the alloy material 
subjected to direct aging exhibits, both at room temperature and 
650.degree. C., 0.2% offset yield strength and tensile strength which are 
improved by only about 10% over those of the alloy material subjected to 
aging after a solid-solution heat treatment, but from Table 3 it will be 
seen that the alloy material subjected to direct aging exhibits much 
excellent property in its creep rupture life over that of the alloy 
material subjected to aging after a solid-solution heat treatment. 
Further, it will be seen that the alloy material subjected to direct aging 
exhibits excellent values also in its elongation and reduction of area. 
TABLE 2 
______________________________________ 
Heat Tensile properties 
treat- 0.2% off- 
Al- ment Test set yield 
Tensile Elonga- 
loy condi- temp. strength 
strength 
tion (2) 
No. tion (.degree.C.) 
(kgf/mm.sup.2) 
(kgf/mm.sup.2) 
(%) (%) 
______________________________________ 
2 Direct Room 146.8 168.1 11.2 20.5 
aging temp. 
650 128.0 154.3 14.6 18.8 
(1) Room 137.0 159.9 15.7 31.6 
temp. 
650 118.9 146.6 17.3 17.7 
______________________________________ 
(1): Solidsolution heat treatment + aging 
(2): Reduction of area 
TABLE 3 
______________________________________ 
Creep rupture 
properties 
Heat Test condition Elon- 
Alloy treatment 
Temp. Stress Life gation 
(2) 
No. condition 
(.degree.C.) 
(kgf/mm.sup.2) 
(hours) 
(%) (%) 
______________________________________ 
2 Direct 650 100 93.2 20.8 20.9 
aging 
(1) " " 42.9 6.1 9.5 
______________________________________ 
(1): Solidsolution heat treatment + aging 
(2): Reduction of area 
EXAMPLE 3 
Alloy Nos. 1 to 13, 21, 22 and 24 produced in Example 1 were subjected to 
direct aging and their tensile properties were tested at room temperature, 
650.degree. C., 705.degree. C. and 760.degree. C., and the results thereof 
are shown in Table 4. Both the alloys of the invention and the comparison 
alloys exhibit very excellent values in their offset yield strength, 
tensile strength and elongation at room temperature, 650.degree. C. and 
705.degree. C. 
In Table 5 there are shown creep rupture properties of the alloy materials 
subjected to direct aging, under the creep test condition of 650.degree. 
C. and 100 kgf/mm.sup.2. 
However, with respect to the alloys Nos. 1 and 5 of the present invention, 
their creep rupture properties under the creep test condition of 
705.degree. C. and 75 kgf/mm.sup.2 are also shown in Table 5. It will be 
seen that the comparison alloys exhibit the tensile properties equivalent 
to those of the alloys of the present invention, but they are much 
inferior in their creep rupture life. The comparison alloy No. 21 exhibits 
a creep rupture life which is as short as 22.3 hours, because it does not 
contain Ta and Ni at all. Further, the comparison alloy No. 22 exhibits a 
creep rupture life of 61.8 hours owing to the effect of Nb, thus providing 
a remarkable improvement in comparison with the comparison alloy No. 21, 
nevertheless this improved creep rupture life is still inferior to those 
exhibited by the alloy of the present invention. A comparison alloy No. 
24, which has a composition very similar to that of the alloy No. 1 but 
contains 0.05% of Zr, caused a notch rupture in a short time of 13.7 
hours, and from this fact it will be seen that addition of very small 
amount of Zr exerts an unfavorable effect on the creep rupture properties 
in the alloy of the present invention. 
TABLE 4 
__________________________________________________________________________ 
0.2% offset yield strength 
Tensile strength 
(kgf/mm.sup.2) (kgf/mm.sup.2) Elongation (%) 
Alloy 
Room Room Room 
No. temp. 
650.degree. C. 
705.degree. C. 
760.degree. C. 
temp. 
650.degree. C. 
705.degree. C. 
760.degree. C. 
temp. 
650.degree. C. 
705.degree. C. 
760.degree. C. 
Remarks 
__________________________________________________________________________ 
1 142.8 
123.6 
117.7 
97.8 
161.5 
151.5 
137.8 
117.7 
16.1 
12.6 
20.9 23.8 
Alloy of 
invention 
2 146.8 
128.0 
-- 87.2 
168.1 
154.3 
-- 110.4 
11.2 
14.6 
-- 38.4 
Alloy of 
invention 
3 145.2 
125.8 
-- 93.1 
163.0 
153.1 
-- 112.7 
13.4 
15.0 
-- 25.1 
Alloy of 
invention 
4 152.1 
128.6 
-- 92.8 
163.6 
159.6 
-- 113.9 
13.6 
24.0 
-- 28.3 
Alloy of 
invention 
5 136.8 
122.5 
117.6 
103.4 
156.6 
143.7 
136.7 
119.0 
13.2 
13.8 
16.7 14.2 
Alloy of 
invention 
6 147.6 
126.0 
124.7 
-- 168.0 
157.3 
142.4 
-- 15.9 
16.5 
12.2 -- Alloy of 
invention 
7 145.2 
125.1 
-- -- 163.7 
152.3 
-- -- 14.0 
15.1 
-- -- Alloy of 
invention 
8 143.8 
127.2 
-- -- 163.3 
153.5 
-- -- 14.3 
14.6 
-- -- Alloy of 
invention 
9 142.0 
123.0 
-- -- 158.0 
150.7 
-- -- 14.6 
15.6 
-- -- Alloy of 
invention 
10 142.0 
123.0 
-- -- 166.3 
154.1 
-- -- 12.0 
13.1 
-- -- Alloy of 
invention 
11 145.0 
125.3 
-- -- 167.0 
153.5 
-- -- 13.4 
15.1 
-- -- Alloy of 
invention 
12 143.1 
127.5 
-- -- 165.4 
153.3 
-- -- 13.4 
14.2 
-- -- Alloy of 
invention 
13 154.0 
129.9 
-- -- 169.5 
162.4 
-- -- 9.7 
10.1 
-- -- Alloy of 
invention 
21 146.7 
125.4 
-- 83.1 
170.3 
153.2 
-- 104.7 
16.5 
40.7 
-- 49.0 
Comparison 
alloy 
22 -- 120.5 
-- -- -- 151.5 
-- -- -- 30.3 
-- -- Comparison 
alloy 
24 142.0 
127.1 
124.4 
-- 168.4 
160.6 
141.8 
-- 15.0 
14.6 
7.4 -- Comparison 
alloy 
__________________________________________________________________________ 
TABLE 5 
______________________________________ 
Creep rupture 
properties 
Test condition Elon- 
Alloy Temp. Stress Life gation 
* 
No. (.degree.C.) 
(kgf/mm.sup.2) 
(hours) 
(%) (%) Remarks 
______________________________________ 
1 650 100 133.8 13.9 14.9 Alloy of 
invention 
2 " " 93.2 20.8 20.9 Alloy of 
invention 
3 " " 91.8 17.1 19.4 Alloy of 
invention 
4 " " 111.1 19.2 23.6 Alloy of 
invention 
5 " " 114.5 8.3 12.8 Alloy of 
invention 
6 " " 143.5 14.5 16.1 Alloy of 
invention 
7 " " 110.5 15.2 17.9 Alloy of 
invention 
8 " " 117.3 16.3 19.1 Alloy of 
invention 
9 " " 105.5 16.1 18.8 Alloy of 
invention 
10 " " 120.3 10.7 12.1 Alloy of 
invention 
11 " " 106.5 15.8 18.0 Alloy of 
invention 
12 " " 110.2 18.9 20.0 Alloy of 
invention 
13 " " 150.9 7.0 10.1 Alloy of 
invention 
21 " " 22.3 26.9 50.9 Com- 
parison 
alloy 
22 " " 61.8 8.3 12.8 Com- 
parison 
alloy 
24 " " 13.7 Notch rupture 
Com- 
parison 
alloy 
1 705 75 87.9 23.9 40.0 Alloy of 
invention 
5 " " 116.1 20.0 31.5 Alloy of 
invention 
______________________________________ 
*Reduction of area 
Next, in comparing mutually the alloys of the present invention, the alloys 
Nos. 1, 4 and 5 exhibit longer creep rupture life in comparison with the 
alloys Nos. 2 and 3 by virtue of containing greater amount of Ta. However, 
the alloy No. 2 containing 4.0% of Ta and the alloy No. 12 in which an 
amount of Ta corresponding to 13 atomic% of that in No. 1 is substituted 
with Nb as well as the alloy No. 3 in which an amount of Ta corresponding 
to 40 atomic% of that in No. 1 is substituted with Nb exhibit shorter 
creep rupture life than the alloy No. 1, but they exhibit the fully 
satisfactory properties. The alloys Nos. 7 and 8 exhibit stable properties 
regardless of the change in the Co content. The alloy No. 10, when 
compared with the alloy No. 9 having smaller contents of Mo and W, 
exhibits greater tensile strength and creep rupture life, but its 
ductility is somewhat smaller than the alloy No. 9. The alloy No. 11 
having a greater Cr content than the alloys Nos. 1, 4, 5 and 6 exhibits 
properties which are quite acceptable. The alloy No. 13 having a 
comparatively large .gamma.' phase content of 34.3 vol.% exhibits 
excellent hot workability, as well as improved tensile strength and creep 
rupture life, but is ductility is somewhat inferior to those of other 
alloys of the present invention. 
FIG. 1 shows tensile properties (0.2% offset yield strength and elongation) 
of the alloy No. 1 of the present invention in comparison with those of 
conventional alloys Nos. 31, 32 and 33, while FIG. 2 shows 100-hour creep 
rupture strength of the alloy No. 1 of the present invention in comparison 
with those of the conventional alloys Nos. 31, 32 and 33. The conventional 
alloy No. 31 is RENE' 95 
(0.06C-13Cr-8Co-3.5Mo-3.5W-2.5Ti-3.5Nb-0.05Zr-0.01B-Bal.Ni) which is 
considered to be the best one presently available by powder metallurgy 
process. The alloy No. 32 is INCONEL 718 (0.05C-19Cr-3Mo-0 
8Ti-0.5Al-5Nb-18Fe-Bal.Ni) subjected to a thermomechanical treatment. The 
alloy No. 33 is INCONEL 718 subjected to no thermomechanical treatment. 
The values concerning the alloys Nos. 31 and 33 were extracted from a 
catalog (3rd edition, July 1977) of International Nickel Company, Inc., 
while the values concerning the alloy No. 32 were extracted from a 
literature "F. Turner and H. S. von Harrach: Materials Sci. and Tech., 
1986, 2, 733-740". However, with respect to the alloys Nos. 1 and 32, the 
values shown in FIG. 2 are those obtained by extrapolating the rupture 
time to 100 hours with the aid of Larson-Miller parameter. 
From FIG. 1, it will be seen that the alloy of the present invention 
exhibits, at temperatures up to 705.degree. C., the 0.2% offset yield 
strength substantially equivalent to that of the alloy No. 31 and much 
superior to that of the alloy No. 33 and, in addition, it exhibits, at 
650.degree. C., the strength much higher than that of the alloy No. 32. 
Further, the alloy of the present invention exhibits excellent property 
with respect also to elongation. Referring now to FIG. 2, the 100-hour 
creep rupture strength exhibited by the alloy of the present invention at 
temperatures up to 705.degree. C. is substantially equal to that of the 
alloy No. 31 which is a powder metallurgy alloy. Thus, the alloy of the 
present invention is much superior to conventional alloys produced by the 
ingot making and hot working process also in the aspect of creep rupture 
strength. 
As has been described, according to the alloy of the present invention and 
the method for producing the same, it becomes possible to attain a 
strength level demanded by the material for turbine disks or the like, 
which has hitherto been obtained solely by powder metallurgy process, by 
using the conventional ingot making and hot working process, so that the 
present invention greatly contributes to improvement in the reliability of 
the parts such as gas turbine disks, as well as to reduction in the cost 
of production of such parts.