Note: Descriptions are shown in the official language in which they were submitted.
- 1 - 57,151
GA5 TURBINE BLADE ALLOY
BACKGROUND OF THE INVENTION
The current invention relates to an alloy suitable
for use in making gas turbine components, such as the rota~ing
blades in the turbine section of a gas turbine. More
specifically, the current invention concerns a nickel-based
alloy having a sufficiently high chromium content for good
corrosion resistance yet maintaining high strength when used
to make a directionally solidified turbine blade casting.
A gas turbine employs a plurality of rotating blades
in its turbine section. Such blades are exposed to gas at
temperatures in excess of 1100C (2000F) and subjected to
high stress. Consequently, the alloys from which such blades
are cast must, after suitable heat treatment, have very high
stress rupture skrength and sufficient metallurgical stability
; 15 to maintain this strength for many thousands of hours of
operation. Such alloys rnust also have sufficient ductility
to withstand the large thermal stresses imposed on turbine
blades. In addition, as a result of impurities in the fuel
and combustion air, the gases to which the blades are exposed
contain corrosive compounds, such as sulfides and chlorides.
Consequently, such blade alloys must also have good hish
temperature corrosion resistance, as well as oxidation
resistance.
Generally, high temperature corrosion resistance is
provided by the incorporation of substantial amounts of
chromium into turbine blade alloys. High chromium content
inhibits the basic fluxing of the alloy by forming a
continuous chromia scale khat is not susceptible to solution
_ ~ _ 57,151
and reprecipitation from a Na2SO4 melt, thereby providing an
effective barrier for the alloy. Moreover, chromium combines
with sulfur to form high melting point sulfides, thereby
inhibiting degradation due to sulfidation. As discussed in
R. Streiff and D.H. Boone, "Corrosion Resistant Modified
Aluminide Coating," Journal of Materials Engineering (1988),
a minimum level of 15% chromium is considered necessary for
good high temperature corrosion resistance.
One nickel-based alloy, used with some success for
a number of years in gas turbine blades, is manufactured by
the International Nickel Company and known commercially as IN-
738. A typical composition of IN-738 in weight percent, as
published in the American Society of Metal Handbook, 9th ed.,
volume 4, page 244, is as follows: Nickel 61, Chromium 16.0,
Cobalt 8.5, Molybdenum 1.7, Tungsten 2.5, Tantalum 1.7,
Niobium (aka Columbium) 0.9, Aluminum 3.4, Titanium 3.4,
Carbon .17, Boron 0.01, and Zirconium 0.10. In addition, the
manufacturer recommends that the electron vacancy number for
this alloy not exceed 2.36. This alloy is disclosed in U.S
Patent No. 3,459,545 (Bleber), hereby incorporated by
reference in its entirety.
It is known that directional solidification, whereby
a uniaxial grain structure is produced, increases both the
ability to withstand cyclical thermal stress and the stress
rupture strength of many nickel-based alloys -- for example,
U.S. Patent No. 4,519,979 (Shaw) discloses that directional
solidification of an alloy known commercially as IN-939 and
having a composition in weight percent of Carbon 0.15,
Chromium 22.5, Cobalt 19, Tungsten 2, Titanium 3.7, Aluminum
1.9, Tantalum 1.4, Niobium 1.0, Zirconium 0.1, Boron 0.01 and
the balance Nickel, increased the stress the rupture life of
the alloy fxom about 850 hours to 1370 hours at 870C and 200
N/mm2. The inventor has confirmed, however, that directional
solidification does little to increase the stress rupture
strength of the IN-738 alloy.
Generally, it has been observed that the alloys that
derive the most improvement from directional solidification
$ ~ 3
_ 3 _ 57,151
have three characteristics -- (i) a relatively low chromium
content, (ii) a high gamma prime volume fraction and (iii) a
high solution temperature. Consequently, it has previously
been thought that the advantages of directional solidification
could be obtained for alloys such as IN-738 only by reducing
their chromium content -- see, for example, M. McLean,
"Directionally Solidified Materials for High Temperature
Service," published by Metals Society, London (1983), page
153. Thus, one popular alloy used in d.irectionally solidified
blades, known commercially as GTD-111, has a typical
composition in weight percent o~ Chromium 14.0, Cobalt 9.5,
Aluminum 3, Titanium 4.9, Tantalum 2.8, Tungsten 3.8,
Molybdenum 1.5, Boron 0.01, Carbon O.1 and the balance Nickel,
as disclosed by R. Viswanathan in "Damage Mechanisms and Life
Assessment of High-Temperature Components," published by the
American Society of Metals (1989). As can be seen, in GTD-lll
the benefits of directional solidification have been obtained
by reducing the 16% chromium level used in IN-738 to only 14%.
Unfortunately, as previously discussed, such relatively low
levels of chromium result in inadequate corrosion resistance.
Accordingly, it would be desirable to provide a
blade alloy that has high strength when directionally
solidified yet sufficient chromium to achieve adequate
corrosion resistance.
SUMMARY OF THE_INVENTION
It is an object of the current invention to provide
a nickel based alLoy suitable for use in casting the rotating
blades of a gas turbine and having high stress rupture
strength when directionally solidified, at least 15% hy weight
chromium for good corrosion resistance, and stability during
long term service.
This objectt as well as other objects, is
accomplished in a gas turbine having a rotating blade made
from an alloy comprising the following elements in weight
percent: Chromium 14.75 to 16.0, Cobalt 8.0 to 8.5, Aluminum
3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7
to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0, Carbon .05
3`~ ~
- 4 - 57,151
to .12, Columbium up to .5, Molybdenum up to 2.0, and the
balance Nickel. The alloy of the current invention may also
comprise impurities and incidental elements generally
associated with nickel-based alloys, such as Zirconium up to
.06 and Boron up to .015 percent by weight.
BRIEE DESCRIPTION OF THE DRAWINGS
Figure 1 is a graph of rupture time t, in hours,
versus stress a, in KSI (N/mm2), for both conventionally cast
(CC) and directionally solidified (DS) specimens of IN-738,
showing the effect of directional solidification on the stress
rupture strength of IN-738.
Figure 2 is an isometric view of a gas turbine
rotating blade.
Figure 3 is a bar chart of rupture time, t, in
thousands of hours, versus stress ~, in KSI (N/mm2), showing
the stress rupture life at 870C (1600F) and four stress
levels of (i) four heats of the alloy according to the current
invention, SAS1-SAS4, as directionally cast,
(ii) conventionally cast IN-738 and (iii) directionally
solidified IN-6203.
DESCRIPTION OF THE PREFERRED EMBODIMENT
Figure 1 shows a graph of time to rupture versus
stress at three temperature levels for both conventionally
cast -- that is, having an equia~ed grain structure -- and
directionally solidified IN-738 test specimens. As can be
seen, data at both the 870C (1600F) and 925C (1700~F)
temperature levels indicates that the rupture life of the
directionally solidified specimens is worse than that of the
conventionally cast specimen above a stress level of about 275
N/mm2 (40 KSI)o
There is shown in Figure 2 a rotating blade 1 used
in the turbine section of a gas turbine. The current
invention is directed to an alloy, referred to as OM 200, from
which such blades may be cast, especially using a directional
solidification casting process. Generally speaking, the
current invention is directed to an alloy comprising the
following elements in weight percent: Chromium 14.75 to 16.0,
- 5 - 57,151
Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3,
Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7,
Tungsten 2.0 to 4.0l Carbon .05 to .12, Columbium up to .5,
Molybdenum up to 2.0, and the balance Nickel. The alloy of
the current invention may also comprise impurities and
incidental elements generally associated with nickel-based
alloys, such as Zirconium up to .06 and Boron up to .015
percent by weight. Prefexably, the alloy of the current
invention consists essentially of the following elements in
weight percent: Chromium about 15.5, Cobalt about ~.0,
Aluminum about 4.0, Titanium abou$ 3.8, Aluminum plus Titanium
about 7O8~ Tantalum about 2.7, Tungsten about 2.6, Molybdenum
0.5, Carbon about 0.08 and the balance Nickel.
As previously discussed, both corrosion resistance
and strength are important properties for blade alloys.
Corrosion resistance in nickel-based alloys is provided
primarily by chromium. Nickel-based alloys used for gas
turbine components are strengthened by three mechanisms -- (i~
solid solution strengthening, (ii) strengthening resulting
from the presence of carbides and (iii) gamma prime
strengthening. Solid solution strengthening is provided by
molybdenuml chromium and tungsten and, to a lesser extent, by
cobalt, iron and vanadium. Gamma prime strengthening is
provided primarily by aluminum and titanium, which strengthen
the austenitic matrix through the precipitation of Ni3(Al
and/or Ti), an fcc intermetallic compound. In addition, the
aluminum in gamma prime can be replaced by tantalum and
columbium.
Since the strength of nickel-based super alloys
increases with increasing gamma prime volume fraction up to
about 60%, in the alloy according to the current invention the
content of aluminum, titanium, tantalum and columbium, which
tend to increase the gamma prime volume fraction, have been
controlled so as to achieve a hi~h gamma prime volume
fraction. The amount of gamma prime in an alloy can be
determined as discussed below with respect to the
determination of the electron vacancy number Nv. The amount
2Pi
6 57,151
of gamma prime, in weight percent, in four heats of the alloy
according to the current invention shown in Table I are
approximately 52, 54, 56 and 54, respectively. The
aforementioned preferred composition of the alloy according
to the curren-t invention has approxiamtely 56% gamma prime.
By comparison, the amount of gamma prime in IN-738 is
appro~imately 50% by weisht. (It should be noted that the
aforementioned variation in the gamma prime content among the
four heats of the alloy according to the current invention did
not adversely affect corrosion resistance or stability.)
As a result of operation at elevated tempexature,
nicXel-based alloys undergo microstructural chang~s. Such
changes include gamma prime coarsening, which ad~ersely
affects the strength of the alloy, and the transformation of
gamma prime into unwanted topologically close-packed secondary
phases, such as plate or needle-like sigma, eta, etc. The
formation of these plate-like phases adversely affects both
strength and toughness. Consequently, in order to ensure that
high strength and toughness are maintained for many thousands
of hours of operation at elevated temperature~ the composition
of the strengthening elements must be carefully balanced, as
explained below, so that the alloy has a certain degree of
microstructural stability.
It is well known that the electron vacancy number
Nv is a measure of the microstructural stability of the alloy.
The higher the ~alue of Nv the greater the tendency for the
formation of the aforementioned topologically close-packed
secondary phases -- see for example, the American Society of
Metal Handbook, 9th ed., volume 4, page 278. As disclosed in
the brochure entitled ~Alloy IN-738, Technical Data,"
published by The International Nickel Company and hereby
incorporated by reference in its entirety, the electron
vacancy number may be calculated from the equation Nv = 0 66
Ni + 1.71 Co + 2.66 Fe + 4.66 (Cr + Mo + W) + 6.~6 Zr. The
atomic percent of each element in the matri~ to be substituted
into this equation is determined by convexting the composition
from weight percen~ to atomic percent and assuming that (i)
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one-half the carbon forms MC in the order of TaC, CbC, TiC,
~ii) the remaining carhon forms M23C6 with the M comprising
twenty three atoms of Cr, (iii) boron is combined as Mo3B2,
(iii) gamma prime is Ni3(Al, Ti, Ta, Cb), and (iv) the
residual matrix consists of the atomic percent minus those
atoms contained in the carbides, the boride and the gamma
prime reaction so that the total of the remaining atomic
percentages gives the atomic concentration in the matrix.
Conversion of these atomic percentages to a 100 percent basis
gives the atomic percentage of each element to be substituted
into the equation above. Since it is desirable to maintain
a low value of Nv, in the alloy according to the current
invention the composition is adjusted so that, in addition to
obtaining high strength, the electron vacancy number of the
alloy does not exceed about 2.4. In the preferred composition
of the alloy the electron vacancy numbex is equal to about
2.4.
As previously discussed, the aluminum in gamma prime
can be replaced by columbium and/or tantalum, as well as
titanium. However, since replacement of Al by Ti, Cb or ~a
in Ni3Al adversely affects the misfit between the austenitic
matrix and the gamma prime particles, the stability of these
compounds in order of decreasing stability is Ni3Al, Ni3Ti and
Ni3Cb(orTa). As a result, the titanium/aluminum ratio plays
a major role in gamma prime coarsening. The titanium/aluminum
ratio also plays a major role in the transformation of gamma
prime into the aforementioned unwanted plate-like phases.
Accordingly, in order to minimize gamma prime coarsening and
to improve the stabilit~ of the gamma prime phase, the
titanium/aluminum ratio is ~enerally maintained below 2:1 (by
weight). The transformation of gamma prime into unwanted
needle or plate-like phases can also be retarded by the
addition of tungsten.
To assess the stability of the alloy according to
the current invention, the shoulder and grip sections of creep
rupture specimens tested at 871C (1600F) and 205.9 N/mm2 (30
XSI) were examined metallographically for sigma, eta and mu
2'~
- 8 - 57J151
phases. It was found that the aging of these specimens at
871C, under stress/ for up to 4000 hours did not produce any
undesirable phases, thereby suggesting that the alloy of the
current invention is stable.
In the alloy according to the current invention, the
amounts of aluminum, titanium, tantalum, columbium and
molybdenum have been balanced so as to attain high strength
when the alloy is directionally solidified while maintaining
good microstructural stability. Significantly, this result
has been achieved without the need to reduce the chromium
content, and, -therefore, without impairing corrosion
xesistance, as had heretofore been thought necessary by those
skilled in the art. Specifically, in the alloy according to
the current invention, the content of aluminum and titanium
has been increased, when compared with IN-738, to a minimum
of 3.4% for each, with the minimum combined aluminum plus
titanium content being 7.7%. The maximum amounts of aluminum
and titanium are 4.0% and 4.3%, respectively, with the maximum
aluminum plus titanium content being 8.3%. In addition, the
amounts of columbium and molybdenum have been reduced so that
the optimum composition includes no columbium and only 0.5%
molybdenum. At most only 0.5% columbium and 2.0~ molybdenum
are permitted in the alloy of the current invention.
Chromium, however, has been maintained in the range of 14.75%
to 16%, so that adequate corrosion resistance is maintained.
Tungsten is maintained in the range of 2.0% to 4.0% and
tantalum in the range of 1.75% to 2.7%. Zirconium and boron
are limited to 0.06% and 0.015~, respectively, with none of
either of these elements being present in the most preferred
composition. Carbon is maintained in the 0.05% to 0.12%
range. Moreover, as previously discussed, in the alloy
according to the current invention, the elements are adjusted
wîthin the aforementioned ranges so that the eleckron vacancy
number is maintained at no more than about 2.4, thereby
ensuring that adequate microstructural stability is achieved.
As a result of the foregoing, the strength of the
alloy according to the current invention when cast by a
9 ~ 57,151
directional solidification process is high, despite its
relatively high chromium content. Significantly, good
microstructural stability of the alloy according to the
current invention has been achieved by considerably reducing
the levels of columbium and molybdenum, when compared to IN-
738, so that the amounts of aluminum and titanium can be
increased without driving the electron vacancy number too
high.
Four heats of the alloy according to the current
invention, designated SASl through SAS4, were prepared using
the compositions in weight percent shown in Table I.
Specimens were cast from these heats 80 as to be directionally
solidified and were given the standard heat treatment
suggested for IN-738 -- that is, a 1121C (2050F~ solution
heat treatment for two hours followed by a 843C (1550F)
aging treatment for twenty four hours. These specimens were
subjected to a variety of tests and the results compared to
conventionally cast IN-738 and directionally solidified IN-
6203. IN-6203 is a nickel-based alloy having a nominal
composition in weight percent of Chromium 22.0, Cobalt 19.0,
Aluminum 2.3, Titanium 3.5, Tantalum 1.10, Columbium 0.80:
Tungsten 2.00, Boron 0.01, Carbon 0.15, Zirconium 0.10
Hafnium 0.75 and the balance Nickel.
TABLE I
Element SASl SAS2 SAS3 AS4
Cr 16.00 16.00 14.75 16.00
Co 8.50 8.00 8.00 8.00
Al 3.40 3.40 3.60 4.00
Ti 4.30 4.30 4.30 3.40
Al+Ti 7.70 7.70 7.90 7.40
W 2.60 2.60 2.60 2.60
Mo 1.75 0.90 1.00 0.60
Ta 1.75 2.70 2.70 2.70
Cb 0.00 0.00 0.50 0.00
C 0.09 0.09 0.09 0.09
Zr 0.04 0.04 0.04 0.04
B 0.01 0.01 0.01 0.01
Ni Bal Bal Bal Bal
The results of low fatigue testing at 871C (1600F)
of specimens from heats SAS1, SAS3 and SAS4 are shown in Table
II and indicate that the cyclic life of the alloy according
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to the current invention is superior to IN-738. The results
of impact testing of three specimens from heat SASl -- taken
from the tip, mid-section and bottom of th~ cast test bar,
respectively - are shown in Table III and reveal that the
impact strength of the alloy of the current invention is
comparable to that of IN-738. Note that IN-738 is provided
in both low and high carbon forms and the data is reported
accordingly where appropriate. In its low carbon form the
carbon content is reduced from 0.17% to 0.11% and zirconium
is reduced from 0.10~ to 0.05%t be weight, compared to
conventional high carbon IN-738.
The results of stress rupture tests are shown in
Tables IV and V and Figure 3, wherein the rupture life of
specimens from each of the four heats are compared with those
of conventionally cast IN-738 and directionally solidified IN-
6203. As can be seen, the data shows that the stress rupture
life of the alloy according to the curren-t invention, when
directionally solidified, is significantly superior to both
that of both IN-738 and IN-6203 -- for example, at 69 N/mm2
(10 KSI), the alloy of the current invention can provide the
same rupture life as IN-738 and IN-6203 at a temperature about
56C (100F) higher. When applied to a gas turbine blade,
such increased metal temperature capability allows the turbine
to operate at higher gas temperatures, ~hereby significantly
improving its performance.
~ABLE I I
Total Strain Cycles to Fail~lre
Range IN-738LC SAS1 SAS3 SAS4
0.012 - 1412 1442 1061
0.01 317 2990 3973 2852
0.0085 862 5342 6673 11200
0.0075 - 16220 50173+ 10504
0.007 1510
0.00~ 1819
0.005 2503
0.004 10153 - - -
2. ~ ~
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T~BLE III
Temp., C _ Impact Energy~ Joules
IN-738HC IN-738LC SASl
Tip Mid Bot
5 24 6.3 8.8 6.8 - 8.1
260 6.9 - 7.9 11.316.7
538 11.7 - 7.5 9.612.1
649 7.6 10.0 6.3 8.714.1
760 7.1 8.5 6.0 7.39.9
10871 8.1 9.8 9.6 9.114.4
Table IV
Temp., C Time to Rupture~ Hours @ 69 N/mm2 ~10 KSI)
_N-738 IN-6203 SASl
15 954 4043
968 2~96
982 1437 2752 8400+
996 820 1359 6450+
1010 - 485 3745
Table V
S$resæ Time to Rupture, Hours @ 871C ~
N/mm2 (KSI ) IN-738 IN-6203SASl SAS3 SAS4
275.8 (40)188 169 403 633 317
24~.3 (35)420 683 1165 1397 776
25206.9 (30) 1143 1~44 3272 4211 2676
172.4 (25)3177 - 8~00~8800~8247
Turbine blades cast from the alloy of the current
invention are advantageously made by vacuum-induction melting
and vacuum casting using a directional solidification process.
Directional solidification causes th~ grain boundaries to be
oriented substantially parallel to the principal stress axis
of the blade with almost no grain boundaries orient~d normal
to the principal stress axis. Techniques for directional
solidification are well known in the art -- see, for example,
U.S. Patent Nos. 3,260,505 (Ver Snyder)l 3,494,709 (Piearcey)
and 3,897,815 (Smashey), hereby incorporated by reference in
their en~irety.
The gamma prime distribution depends on heat
treatment, as well as composition. The standard heat
treatment for nickel-based alloys such as IN-738 -~ i.e., a
solution treatment followed by an aging treatment -- produces
duplex gamma prime comprised of coarse, cuboidal primary gamma
prime and fine, spherical gamma prime in approximately equal
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amounts. The coarse gamma prime is undesolved gamma prime
that did not go into solution during the solution treatment.
Hence, the amount of coarse gamma prime present in the alloy
depends on the degree by which the solution temperature is
below the ga~a prime solvus temperature, at which all of the
gamma prime goes into solution -- that is, the lower the
solution temperature r the greater the amount of coarse gamma
prime. The fine gamma prime forms during the aging treatment,
the amount depending on the amount of gamma prime that did not
go into solution during solution treatment.
Directionally solidified components are cooled at
a slower rate than conventionally cast components. As a
result, directional solidification produces coarser primary
gamma prime so that the volume fraction of cuboidal gamma
prime is higher and that of spherical gamma prime is lower
than in directionally cast components. Unfortunately, a high
volume fraction of coarse gamma prime has a deleterious effect
on strength. This discrepancy in gamma prime distribution can
be minimized by optimizing the heat treatment, particularly
the solution treatment, based on the gamma prime solvus and
incipient melting temperatures, both of which depend on alloy
composition. Generally, it is desirable that the heat
treatment be performed at a temperature between the solvus
temperature and the incipient melting temperature. Hence, in
or~er to obtain the full benefits of heat treatment, the
solvus temperature should be considerably below the incipient
melting temperature. Unfortunately, although aluminum,
titanium, tantalum and columbium increase the gamma prime
volume fraction, and therefore, strength, as previously
discussed, they also have the effect of raising the solvus
temperature and decreasing the incipient melting temperature,
thereby narrowing the heat treatment window.
The gamma prime solvus and incipient melting
temperatures for three melts of the alloy according to the
current invention are shown in Table VI. These temperatures
were determined using the differential thermal analysis and
gradient bar -- i.e., metallography -- method, in which the
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bar was exposed to various temperatures in the 1066C (1950F)
~o 1427C t2300F) temperature range for four hours and then
fan cooled. As can be seen, the solvus temperature varies
from 1211C to 1229C. By comparison, the solvus temperature
for IN-738 is approximately 1204C (2200F). As previously
mentioned, IN-738 is typically given a 1121C (2050F)
solution heat treatment for two hours followed by a 843C
(1550F) aging treatment for twenty four hours. Various heat
treatment regimes for the alloy according to the current
invention, based on its incipient melting and solvus
temperatures, are discussed further below.
TABLE VI
SASl SAS3 SAS4
Incipient melting temp. C 1241 1229 1229
15Gamma prime solvus temp., C 1211 1211 1229
According to the current invention, the as~cast
blades may be heat treated in any of four ways -- (i) solution
treating for ~ hours at 1121C (2050F), followed by aging for
24 hours at 843C (1550F), (ii) solution treating for 4 hours
at 1149C (2100F), followed by agins for 24 hours at 843C
(1550F), (iii) solution treating Eor 4 hours at 1204C
(2200F), resolution treating for 2 hours at 1121C (2050F)
followed by aging for 24 hours at 843C (1550F), and
(iv) solution treating for 4 hours a~ 1204C (2200F),
resolution treating for 4 hours at 1149C (2100F), followed
by aging for 24 hours at 843C (1550F). These heat treatment
optimizes mechanical strength by creating a duplex gamma prime
structure in a gamma matrix and discrete chrome-carbides at
the grain boundaries. The use of higher solution temperatures
will decrease the amount of coarse gamma prime and increase
the amount of fine gamma prime, hence, further strengthening
the alloy.
Although the current invention has been described
with reference to an alloy for casting turbine blades, it
should be understood that the alloy according to the current
invention may also be used to form other components that
require high strength at elevated temperature and good
corrosion resistance. Accordingly, the present invention may
~ 182r~
- 14 - 57, 151
be embodied in other specific forms without departing from the
spirit or essential attributes thereof and, accordingly,
reference should be made to the appended claims, rather than
to the foregoing specification, as indicating the scope of the
invention.
