U.S. patent application number 10/994202 was filed with the patent office on 2005-10-13 for monocrystalline alloys with controlled partitioning.
Invention is credited to Feng, Qiang, Holmes, Peter, Konter, Maxim, Pollock, Tresa, Rowland, Laura.
Application Number | 20050224144 10/994202 |
Document ID | / |
Family ID | 35059341 |
Filed Date | 2005-10-13 |
United States Patent
Application |
20050224144 |
Kind Code |
A1 |
Pollock, Tresa ; et
al. |
October 13, 2005 |
Monocrystalline alloys with controlled partitioning
Abstract
Nickel-based superalloys, for fabrication of monocrystalline
turbine components to be used in industrial and aircraft turbine
engines, having the following composition (in wt %): 5.6-8.1% Al,
4.1-14.1% Ru, 6.1-9.9% Ta, 3.6-7.5% Re, and the remaining balance
Ni. The partitioning of alloying elements can be controlled to
achieve a wide range of precipitate shapes and exceptional
resistance to degradation under high temperature exposure
conditions.
Inventors: |
Pollock, Tresa; (Ann Arbor,
MI) ; Feng, Qiang; (Ann Arbor, MI) ; Rowland,
Laura; (Ann Arbor, MI) ; Konter, Maxim;
(Klingnau, CH) ; Holmes, Peter; (Winkel,
CH) |
Correspondence
Address: |
HARNESS, DICKEY & PIERCE, P.L.C.
P.O. BOX 828
BLOOMFIELD HILLS
MI
48303
US
|
Family ID: |
35059341 |
Appl. No.: |
10/994202 |
Filed: |
November 19, 2004 |
Related U.S. Patent Documents
|
|
|
|
|
|
Application
Number |
Filing Date |
Patent Number |
|
|
60537481 |
Jan 16, 2004 |
|
|
|
Current U.S.
Class: |
148/404 ;
148/428; 148/429 |
Current CPC
Class: |
F01D 25/005 20130101;
C22C 19/057 20130101; F05D 2300/607 20130101; F01D 5/28
20130101 |
Class at
Publication: |
148/404 ;
148/428; 148/429 |
International
Class: |
C22C 019/05 |
Claims
What is claimed is:
1. A superalloy comprising: about 5.6 to about 8.1 percent by
weight of aluminum (Al); about 4.1 to about 14.1 percent by weight
of ruthenium (Ru); about 6.1 to about 9.9 percent by weight of
tantalum (Ta); about 3.6 to about 7.5 percent by weight of rhenium
(Re); and nickel (Ni).
2. The superalloy according to claim 1, further comprising: about
0.1 to about 12.4 percent by weight of tungsten (W).
3. The superalloy according to claim 1, further comprising: about
0.1 to about 9.6 percent by weight of cobalt (Co).
4. The superalloy according to claim 1, further comprising: about
0.1 to about 6.9 percent by weight of chromium (Cr).
5. The superalloy according to claim 1, further comprising: about
0.1 to about 0.7 percent by weight of silicon (Si).
6. The superalloy according to claim 1, further comprising: about
0.1 to about 1.5 percent by weight of molybdenum (Mo).
7. The superalloy according to claim 1, further comprising: about
0.1 to about 0.8 percent by weight of titanium (Ti).
8. The superalloy according to claim 1, further comprising
spherical precipitates.
9. The superalloy according to claim 1 wherein said superalloy
forms a monocrystalline structure.
10. A turbine article comprising: a monocrystalline alloy having
about 5.6 to about 8.1 percent by weight of aluminum (Al), about
4.1 to about 14.1 percent by weight of ruthenium (Ru), about 6.1 to
about 9.9 percent by weight of tantalum (Ta), about 3.6 to about
7.5 percent by weight of rhenium (Re), and nickel (Ni).
11. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 12.4
percent by weight of tungsten (W).
12. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 9.6
percent by weight of cobalt (Co).
13. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 6.9
percent by weight of chromium (Cr).
14. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 0.7
percent by weight of silicon (Si).
15. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 1.5
percent by weight of molybdenum (Mo).
16. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises about 0.1 to about 0.8
percent by weight of titanium (Ti).
17. The turbine article according to claim 10 wherein said
monocrystalline alloy further comprises spherical precipitates.
18. A superalloy comprising: about 5.6 to about 6.3 percent by
weight of aluminum (Al); about 4.1 to about 9.7 percent by weight
of ruthenium (Ru); about 6.1 to about 9.9 percent by weight of
tantalum (Ta); about 3.6 to about 7.5 percent by weight of rhenium
(Re); and nickel (Ni).
19. The superalloy according to claim 18, further comprising: about
0.1 to about 12.4 percent by weight of tungsten (W).
20. The superalloy according to claim 18, further comprising: about
0.1 to about 9.6 percent by weight of cobalt (Co).
21. The superalloy according to claim 18, further comprising: about
0.1 to about 6.9 percent by weight of chromium (Cr).
22. The superalloy according to claim 18, further comprising: about
0.1 to about 0.7 percent by weight of silicon (Si).
23. The superalloy according to claim 18, further comprising: about
0.1 to about 1.5 percent by weight of molybdenum (Mo).
24. The superalloy according to claim 18, further comprising: about
0.1 to about 0.8 percent by weight of titanium (Ti).
25. The superalloy according to claim 18, further comprising
spherical precipitates.
26. The superalloy according to claim 18 wherein said superalloy
forms a monocrystalline structure.
Description
CROSS-REFERENCE TO RELATED APPLICATIONS
[0001] This application claims the benefit of U.S. Provisional
Application No. 60/537,481, filed on Jan. 16, 2004. The disclosure
of the above application is incorporated herein by reference.
FIELD OF THE INVENTION
[0002] The present invention relates to alloys and, more
particularly, relates to nickel-based superalloys for the
manufacture of monocrystalline structures.
BACKGROUND AND SUMMARY OF THE INVENTION
[0003] The present invention generally relates to advanced
materials for high temperature components in industrial power and
aircraft turbines and, specifically, monocrystalline superalloy
blades and vanes. To maximize the efficiency of these turbine
systems, the operating temperatures of blades and vanes must be
maximized to prevent damage and premature failure. By way of
background, it should be recognized that premature damage
accumulation may occur along grain boundaries when such components
are operated near their melting point. Accordingly, Bridgman-type
processes may be utilized to eliminate boundaries and, thus, permit
use of superalloys in monocrystalline form. At high temperatures,
monocrystalline blades and vanes undergo degradation due to creep,
phase instabilities, or oxidation and, consequently, must be
periodically replaced. It is desirable in many cases to minimize
these characteristics in order to maximize the useful life and
operating properties of turbines.
[0004] The addition of refractory alloying elements, such as
rhenium (Re) and tungsten (W), are desirable for improving the
maximum temperature capability of these monocrystalline alloys. As
a result, the addition of refractory alloying elements serves to
strengthen the monocrystal and, thus, delay the onset of creep
damage. However, conversely, high levels of refractory alloying
elements may lead to phase instabilities. One form of phase
instability is the formation of brittle topologically close packed
phases (TCPs). These phases form during long-term,
elevated-temperature exposures and tend to degrade mechanical
properties. To avoid precipitation of detrimental TCP phases during
service, low levels of chromium (Cr) are recommended. Low levels of
Cr, however, may result in poor oxidation and corrosion resistance.
That being said, it has recently been shown that the addition of
small amounts of ruthenium (Ru) decreases the propensity for the
precipitation of detrimental TCP phases. Another consequence of
refractory alloying additions is their tendency to cause a
breakdown of single crystal solidification. It is essential to
design alloys within composition ranges where it is possible to
produce them as monocrystals to avoid the disadvantages of the
prior art.
[0005] Phase instability may further occur in monocrystalline
alloys when the directional coarsening or "rafting" of the
Ni.sub.3Al-.gamma.' precipitates under the action of an externally
applied stress. Rafting is enhanced by high levels of Re, which
increase the lattice parameter of the y matrix to higher values
than the .gamma.' precipitate phase. In commercial monocrystalline
alloys stressed in tension along an applied stress A-A (see FIG.
1c), also known as a crystallographic orientation, directional
coarsening occurs in a manner that produces plate-shaped
precipitates oriented normal to the stressing applied stress A-A.
Turbine components are typically fabricated so that the major
stresses will be applied along a plane parallel to applied stress
A-A.
[0006] The present invention goes well beyond the prior art in the
use of higher levels of ruthenium (Ru) (up to about 14.1 wt %) to
control precipitate morphology and rafting behavior, suppress
precipitation of TCP phases, and improve creep properties. This is
possible through controlled partitioning, where differing amounts
of Ru affect the partitioning of elements in the alloy,
particularly the Re and W, to the gamma and gamma prime phases. The
exceptional aspect of the present invention is that alloys with
positive, zero, or negative misfit, no TCP phases and high levels
of Re can be designed. This is significant because rafting can be
completely suppressed or rafts parallel to or normal to the applied
tension applied stress A-A can form with zero, positive, or
negative misfit, respectively.
[0007] Furthermore, it has been demonstrated that with higher
levels of Ru, higher ratios of Cr/Re can be achieved,
simultaneously improving oxidation and creep behavior. Cr is
important in controlling partitioning. Since the three major
mechanisms of high temperature degradation (TCP phase formation,
creep damage, and oxidation) are improved, the alloys of the
present invention are capable of increasing the useful life and
temperature capability of critical turbine components.
[0008] Further areas of applicability of the present invention will
become apparent from the detailed description provided hereinafter.
It should be understood that the detailed description and specific
examples, while indicating the preferred embodiment of the
invention, are intended for purposes of illustration only and are
not intended to limit the scope of the invention.
BRIEF DESCRIPTION OF THE DRAWINGS
[0009] The present invention will become more fully understood from
the detailed description and the accompanying drawings,
wherein:
[0010] FIG. 1a illustrates a high volume fraction of precipitates
having spherical morphology in alloy UM-F11 according to the
present invention;
[0011] FIG. 1b illustrates a high volume fraction of precipitates
having spherical morphology in alloy UM-F11 along a face normal to
applied stress A-A after 125 hours at 950.degree. C. and 290
MPa;
[0012] FIG. 1c schematically illustrates an applied stress A-A and
faces normal and parallel to applied stress A-A;
[0013] FIG. 1d illustrates a high volume fraction of precipitates
having spherical morphology in alloy UM-F11 along a face parallel
to applied stress A-A after 125 hours at 950.degree. C. and 290
MPa;
[0014] FIG. 1e illustrates "negative" rafting of a prior art alloy
after 200 hours at 950.degree. C. and 290 MPa;
[0015] FIG. 1f schematically illustrates an applied stress A-A and
faces normal and parallel to applied stress A-A;
[0016] FIG. 1g illustrates "negative" rafting perpendicular to
applied stress A-A of a prior art alloy after 200 hours at
950.degree. C. and 290 MPa;
[0017] FIG. 2a illustrates "positive rafting" in alloy UM-F18 at
950.degree. C. and 290 MPa;
[0018] FIG. 2b schematically illustrates an applied stress A-A and
faces normal and parallel to applied stress A-A;
[0019] FIG. 2c illustrates "positive rafting" in alloy UM-F18 at
950.degree. C. and 290 MPa;
[0020] FIG. 3 is a graph comparing alloys of the present invention
with prior art alloy (MK4), which illustrates the acceleration of
creep rate in the prior art alloy following formation of raft
(after about 100 hours) and the deceleration of creep rate of the
alloys of the present invention;
[0021] FIG. 4 is a graph comparing creep properties of alloys of
the present invention at 950.degree. C. and 290 MPa with varying
ranges of precipitate morphologies achieved by controlled
partitioning;
[0022] FIG. 5a illustrates "negative" rafting of alloy UM-F16 along
a face normal to applied stress A-A;
[0023] FIG. 5b schematically illustrates an applied stress A-A and
faces normal and parallel to applied stress A-A;
[0024] FIG. 5c illustrates "negative" rafting perpendicular to
applied stress A-A of a prior art alloy after 200 hours at
950.degree. C. and 290 MPa;
[0025] FIG. 6 is a graph comparing creep rupture properties of
alloys of the present invention at 950.degree. C. and 290 MPa;
[0026] FIG. 7 is a graph comparing yield strength retention in
alloys of the present invention at 950.degree. C. and 290 MPa with
prior art alloy (MK4);
[0027] FIG. 8 is a graph illustrating cyclic oxidation of alloys of
the present invention at 900.degree. C.;
[0028] FIG. 9 is a graph illustrating cyclic oxidation of alloys of
the present invention at 1100.degree. C.;
[0029] FIG. 10a illustrates the microstructure of alloy UM-F19
after 1500 hours at 950.degree. C.; and
[0030] FIG. 10b illustrates the microstructure of alloy UM-F20
after 3000 hours at 950.degree. C.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0031] The following description of the preferred embodiments is
merely exemplary in nature and is in no way intended to limit the
invention, its application, or uses.
[0032] With initial reference to Table 1, a plurality of
embodiments are illustrated that are within the scope of the
present invention. However, it should be appreciated that these
examples are non-limiting and, thus, additional compositions may be
used or the values enumerated modified.
[0033] A first preferred embodiment defined by the principles of
the present invention include a class of high refractory content
single crystals with spherical precipitates that exhibit no rafting
when subjected to external stresses. All current commercial single
crystal alloys possess microstructures with .gamma.' cuboidal
precipitates that arise due to lattice misfit between the matrix
and precipitates. This misfit occurs due to strong partitioning of
the Re and W to the gamma matrix phase. When subjected to tensile
stresses along the applied stress A-A (see FIG. 1c) at high
temperatures, the initially cube-shaped precipitates coarsen
considerably and evolve to long plate-shaped precipitates that are
oriented with their broad faces normal to the applied tensile
stress A-A. This process is known as "rafting". This change in the
structure of the material causes a change in material properties
during service and may result in a weakening of the material.
Rafting can be suppressed if partitioning of elements to the
precipitates is changed in a manner to achieve spherical
precipitates, which have negligible lattice misfit.
[0034] With brief reference to FIG. 1c, a portion of an alloy is
illustrated having applied tensile stress A-A, a face 10 normal or
transverse to stress A-A, and a face 20 being generally parallel to
stress A-A.
[0035] As seen in FIG. 1a, spherical precipitates are present in
the solution treated and aged condition for Alloy UM-F11 having a
composition as set forth in Table 1. FIGS. 1b and 1d demonstrate a
lack of rafting after 125 hours at 950.degree. C. and 290 MPa along
transverse face 10 and parallel face 20. For comparison, rafting
under the same imposed temperature and stress in a prior art alloy
defined in U.S. Pat. No. 5,888,451, is illustrated in FIGS. 1e and
1g. As can be seen in FIGS. 1e and 1g, the prior art alloy fails to
define spherical precipitates and, thus, may suffer from the
disadvantages enumerated above in connection with additional prior
art.
[0036] Likewise, alloy UM-F9 of the present invention results in
spherical precipitates with no rafting following application of
temperature and stress. It should be emphasized that stable,
spherical precipitates have never before been reported in strong,
Re-containing alloys. This stabilization of precipitate morphology
under stress occurs in response to a low ratio of Cr/Ru and high
ratio of Ru/(Re+W), from about 0-0.4 and about 0.7-1.2, (in wt %),
respectively. Within this composition range, the alloys can be
solidified as monocrystals using conventional Bridgman growth
techniques.
[0037] In another embodiment of the present invention, rafts in a
Re-containing alloy align parallel to the direction of the applied
tensile stress A-A. An example of this is illustrated in FIGS. 2a
and 2b for Alloy UM-F18 stressed in tension along the applied
stress A-A. Rafting in this orientation in a Re or W containing
alloy has not been reported before, due to unrealized regimes for
control of element partitioning. Controlled partitioning to achieve
this "positive" rafting requires intermediate ratios of Cr/Ru and
Ru/(Re+W). Again, within this composition range, the alloys can be
solidified as monocrystals using conventional Bridgman growth
techniques.
[0038] An additional embodiment of the present invention
illustrates that if partitioning can be controlled, creep
acceleration and strength degradation as a result of rafting can be
avoided. Ruthenium additions permit these objectives to be achieved
in Re and W-containing alloys. FIG. 3 illustrates creep curves (not
all have reached rupture) for alloys with cuboidal precipitates
(UM-F16, UM-F19, UM-F20, UM-F27) that raft in the conventional
"negative" sense in comparison to the non-rafting alloys (UM-F9,
UM-F11). Creep in the prior art alloy MK-4 accelerates with the
formation of rafts, while in the Ru-containing alloys the creep
rate is still decreasing as the rafts form. FIG. 4 compares the
creep properties of a range of Ru-containing alloys. It is
important to note the improved creep properties of the rafted
Ru-containing alloys compared to the non-rafted Ru-containing
alloys. The rafted structure present after 200 hours of creep at
950.degree. C. and 290 MPa in UM-F16 is illustrated in FIGS. 5a and
5b. Furthermore, FIG. 6 illustrates improved creep properties of
the conventionally rafted Ru-containing alloys compared to the
prior art alloy MK-4, which possesses similar levels of Re and W
but no Ru. The creep rupture life of the Ru-containing alloys is a
factor of 2.times. to 5.times. longer than prior art. These high
strength, creep resistant rafting alloys can be achieved with high
ratios of Cr/Ru and lower ratios Ru/(Re+W). As seen in FIG. 4,
alloys with intermediate precipitate shapes also have intermediate
creep properties, due to intermediate partitioning, which
demonstrates that a range of behavior can be designed into the
alloys.
[0039] Turning now to FIG. 7, it can be seen plurality of alloys
were first subjected to 1% creep straining at 950.degree. C. and
290 MPa. Room temperature tensile tests were then conducted on the
crept specimens and compared to the tensile properties of the
material in the virgin state. The non-Ru prior art alloy MK-4
suffers approximately 30% degradation in strength due to the high
temperature creep exposure, while the Ru-containing alloys UM-F9,
UM-F16, UM-F19, UM-F20 and UM-F22 are either strengthened by the
high temperature creep exposure or are negligibly affected. It is
important to note that this absence of strength degradation is
present for positive, negative, and non-rafting alloys. This
feature of these alloys is very important to the performance of
turbine blades and vanes since they experience creep deformation in
service. Again, within this composition range, the alloys can be
solidified as monocrystals using conventional Bridgman growth
techniques.
[0040] In another embodiment of the present invention, high
oxidation resistance is combined with high creep strength and a
high resistance to TCP phase precipitation in Ru-containing alloys.
FIGS. 8 and 9 show cyclic oxidation properties of selected alloys
compared to the prior art alloy MK-4. Achieving improved creep
properties and higher temperature capability in monocrystalline
alloys is nearly always associated with a degradation in cyclic
oxidation behavior. Ideally, the monocrystal will neither lose or
gain weight during elevated temperature cycling. In FIG. 6, alloys
UM-F16, UM-F19 and UM-F20 display this desirable behavior at both
900.degree. C. and 1100.degree. C. and are comparable to the prior
art alloy MK-4. Combining high oxidation resistance with high creep
resistance requires intermediate to high levels of Ru (3.5-6 at %)
and high levels of Cr (8 at %/6.7 wt %). These high levels of Cr in
monocrystal alloys typically result in microstructural
instabilities and precipitation of a significant volume fraction of
detrimental TCP phases in non-Ru alloys. FIGS. 1a-1d, 2a-2c, and 4
demonstrate the absence of TCPs in positive, negative and
non-rafting alloys after 100-200 h. of creep. FIGS. 10a and 10b
illustrate the microstructures of UM-F19 and UM-F20 after 1500 and
3000 hours of exposure at 950.degree. C., with a complete absence
of any TCP phase instabilities. The alloys examined in the study
were exceptionally resistant to this form of degradation. The new
discovery in this embodiment is that Ru enables high levels of Cr
to be added to improve oxidation resistance without the onset of
TCP-type phase instabilities.
[0041] The description of the invention is merely exemplary in
nature and, thus, variations that do not depart from the gist of
the invention are intended to be within the scope of the invention.
Such variations are not to be regarded as a departure from the
spirit and scope of the invention.
* * * * *