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.

Description

FIELD OF THE INVENTION

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

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.

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.

Phase instability may further occur in monocrystalline alloys when the directional coarsening or “rafting” of the Ni₃Al-γ′ 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 γ′ 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.

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.

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.

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

The present invention will become more fully understood from the detailed description and the accompanying drawings, wherein:

FIG. 1 a illustrates a high volume fraction of precipitates having spherical morphology in alloy UM-F11 according to the present invention;

FIG. 1 b 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° C. and 290 MPa;

FIG. 1 c schematically illustrates an applied stress A-A and faces normal and parallel to applied stress A-A;

FIG. 1 d 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° C. and 290 MPa;

FIG. 1 e illustrates “negative” rafting of a prior art alloy after 200 hours at 950° C. and 290 MPa;

FIG. 1 f schematically illustrates an applied stress A-A and faces normal and parallel to applied stress A-A;

FIG. 1 g illustrates “negative” rafting perpendicular to applied stress A-A of a prior art alloy after 200 hours at 950° C. and 290 MPa;

FIG. 2 a illustrates “positive rafting” in alloy UM-F18 at 950° C. and 290 MPa;

FIG. 2 b schematically illustrates an applied stress A-A and faces normal and parallel to applied stress A-A;

FIG. 2 c illustrates “positive rafting” in alloy UM-F18 at 950° C. and 290 MPa;

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;

FIG. 4 is a graph comparing creep properties of alloys of the present invention at 950° C. and 290 MPa with varying ranges of precipitate morphologies achieved by controlled partitioning;

FIG. 5 a illustrates “negative” rafting of alloy UM-F16 along a face normal to applied stress A-A;

FIG. 5 b schematically illustrates an applied stress A-A and faces normal and parallel to applied stress A-A;

FIG. 5 c illustrates “negative” rafting perpendicular to applied stress A-A of a prior art alloy after 200 hours at 950° C. and 290 MPa;

FIG. 6 is a graph comparing creep rupture properties of alloys of the present invention at 950° C. and 290 MPa;

FIG. 7 is a graph comparing yield strength retention in alloys of the present invention at 950° C. and 290 MPa with prior art alloy (MK4);

FIG. 8 is a graph illustrating cyclic oxidation of alloys of the present invention at 900° C.;

FIG. 9 is a graph illustrating cyclic oxidation of alloys of the present invention at 1100° C.;

FIG. 10 a illustrates the microstructure of alloy UM-F19 after 1500 hours at 950° C.; and

FIG. 10 b illustrates the microstructure of alloy UM-F20 after 3000 hours at 950° C.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

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.

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.

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 γ′ 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.

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.

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° 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.

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.

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.

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° 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× to 5× 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.

Turning now to FIG. 7, it can be seen plurality of alloys were first subjected to 1% creep straining at 950° 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.

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° C. and 1100° 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° 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.

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.

Claims

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.