Patent Grant
12152296
The invention relates generally to Materials science and engineering, metallurgy, and more particularly, to tungsten-free, low-density cobalt-based superalloys with stable γ′ (gamma-prime) precipitates, based on Co—Ta—V—Al and Co—Nb—V—Al systems, and method of producing the same.
The background description provided herein is for the purpose of generally presenting the context of the invention. The subject matter discussed in the background of the invention section should not be assumed to be prior art merely as a result of its mention in the background of the invention section. Similarly, a problem mentioned in the background of the invention section or associated with the subject matter of the background of the invention section should not be assumed to have been previously recognized in the prior art. The subject matter in the background of the invention section merely represents different approaches, which in and of themselves may also be inventions. Work of the presently named inventors, to the extent it is described in the background of the invention section, as well as aspects of the description that may not otherwise qualify as prior art at the time of filing, are neither expressly nor impliedly admitted as prior art against the invention.
To date, Ni-based superalloys based on the Ni—Al—Cr system are the preferred material for high-stress and high-temperature applications, such as turbine blades or disks, in both jet- and land-based turbines [44-46]. The high strength and creep resistance exhibited by these alloys stems from their γ+γ′ (gamma+gamma-prime) microstructure where the fcc-γ Ni matrix is strengthened with L12-ordered γ′ Ni3(Al,Ti) precipitates which are typically submicron to micron in size. State-of-the-art Ni-based superalloys contain upwards of ten additional alloying elements [44, 47, 48] which provide a variety of benefits: higher melting temperatures, solid-solution strengthening in both γ and γ′ phases, grain-boundary strengthening, γ′-coarsening resistance, as well as oxidation and corrosion resistance. In particular, much research has focused on the addition of refractory elements to provide solid-solution strengthening of the γ-matrix and γ′-precipitates (thus increasing resistance against dislocation slip in both phases) and to reduce diffusion-controlled process (thus slowing precipitate coarsening and dislocation climb in both phases). However, when added in excess, these elements form additional intermetallic phases which are detrimental to mechanical properties [1, 4, 6, 7]. An alternative to Ni-based superalloys, to achieve higher operating temperatures, are cobalt-based alloys which rely on the higher melting point of Co as compared to Ni (1495 vs. 1455° C.). Co-based alloy strengthened via carbide precipitation are widely used in low-stress, corrosive environments at elevated temperature. These alloys do not exhibit γ′ precipitates and are thus less strong and creep-resistant than γ+γ′ Ni-based superalloys [5-8]. However, Sato et al. [10] in 2006 discovered a new family of Co-base superalloys with a γ+γ′ microstructure analogous to that of Ni-based superalloy. These new superalloys, based on the Co—Al—W ternary system and containing coherent γ′-Co3(Al,W) precipitates, have high potential to surpass the performance of Ni-based superalloys. Some of the best Co—Al—W-based superalloys, with multiple alloying additions on par with commercial Ni-base superalloys, show high solidus and liquidus temperatures (50-150° C. greater than those of Ni-based superalloys) [7, 9], as well as high γ′ volume fractions (up to about 80%) and creep resistance comparable to Ni-based superalloys. However, some weaknesses of the Co—Al—W system that limit their potential are: (i) low solvus temperature of the γ′-phase (about 980° C.) [10, 13], and (ii) a high density (more than 9.0 g cm−3) due to their high W content [9]. Given the high density of Co—Al—W-based superalloys, there is strong interest in developing W-free Co-based superalloys. In 2015, Makineni et al. [16-19] demonstrated that Co-10Al-5Mo-2X (X=Nb and Ta) alloys show γ+γ′ structure, with a γ′ composition of Co3(Al,Mo,Nb) [16, 19] and Co3(Al,Mo,Nb) [18], markedly different from Co3(Al,W) in the previous alloys. These authors demonstrated that up to 30 at. % Ni can be added to replace Co. raising the solvus temperature from about 860 to about 950° C. in the Co(—Ni)—Al—Mo—Nb alloys. In modified alloys containing Ta, the solvus temperatures are about 920° C. and increase to about 1070° C. with Ni and Ti additions [18]. Initial studies suggested that the γ′ phase is stable at 800-900° C. for short aging times [19], however, after longer aging times, cellular and needle-like precipitation of the D019 phase are seen in these Co(—Ni)—Al—Mo—Nb alloys after 35 hours of aging at 800-950° C. The only alloy that did not seem to precipitate any Topologically Close Packed (TCP) phases after 100 hours of aging at 950° C. contains Ti, i.e., Co-30Ni-10Al-5Mo-2Nb-2Ti [17]. Since their discovery in 2015, these Co-based W-free superalloys have been the object of numerous investigation, since they could serve as a low-density alternative to W-bearing Co—Al—W-based superalloys [24, 51-53].
Other ternary Co-based systems have been recently identified to exhibit L12-ordered γ′ precipitates, e.g., Co—G—W [2], Co—Ti—Cr [20, 54], Co—Ti—Mo and Co—Ti—V [6, 21]. Co—Ge—W is a derivative from Co—Al—W (where Al is replaced by Ge) while the other three ternary systems are derivatives from the Co—Ti binary alloy that displays stable L12-ordered γ′ precipitates [54], where Ti is partially replaced by Cr, Mo or V, respectively. However, these alloys do not contain Al or (for the last two alloys) Cr, which are crucial for corrosion and oxidation resistance [44, 47, 49] and may also stabilize the γ′ phase [17, 20, 24, 51, 55]. Also, the last alloy system, Co—Ti—V, contains very high amount of V, which could be compromise hot corrosion resistance [56].
In the quest for new Co-based superalloy systems, computational studies implementing density functional theory (DFT) calculations and ThermoCalc/CALPHAD, coupled with experimental studies, are becoming increasingly capable to accelerate the discovery of new ternary systems and improve current alloys [12, 18, 19, 24, 22, 23, 57, 58]. Recently, Nyshadham et al. predicted the formation of stable γ′-L12 Co3(Nb0.5V0.5) and Co3(Ta0.5V0.5) precipitates in equilibrium with a γ-fcc Co—Nb—V/Ta matrix. These γ′-phases were not previously reported in the literature existing for the Co—Ta—V and Co—Nb—V systems. These two systems were studied by Ruan et al. and Wang et al. [60], respectively, who do not report the presence of a γ′-phase.
We experimentally explored alloys in the Co-rich corner of the ternary Co—Ta—V and Co—Nb—V systems [61]. We observed fine γ′ cuboidal precipitates with Co3(Ta0.76V0.24) and Co3(Nb0.81V0.19) composition in both systems which however were metastable, decomposing to: (i) C36 in the Co—Ta—V system and (ii) needle-like D019 in the Co—Nb—V system [61]. This metastability explains why Ruan et al. and Wang et al. did not observe these γ′-phases, since they used very long aging times to study equilibrium phases. This shows that these ternary γ′ phases are metastable and points to the need for their stabilization via additional alloying elements.
Therefore, a heretofore unaddressed need exists in the art to address the aforementioned deficiencies and inadequacies.
In one aspect, the invention relates to a cobalt based superalloy. In one embodiment, the superalloy includes a nominal composition comprising at least cobalt (Co), aluminum (Al), Z and vanadium (V), Z being at least one of tantalum (Ta) and niobium (Nb), which is processed such that the superalloy comprises γ and γ′ phases with stable γ+γ′ microstructures.
In certain embodiments, the nominal composition comprises one of
Co-a2Ni-b2Al-c2Ti-e2Nb-f2V-g2B-h2Cr, wherein a2 is in a range of about 0-40 at. %, b2 is in a range of about 2.5-10 at. %, c2 is in a range of about 0-4 at. %, e2 is in a range of about 2-4 at. %, f2 is in a range of about 1.5-6 at. %, g2 is in a range of about 0-1 at. %, h2 is in a range of about 0-20 at. %, and Co is in balance; and
In certain embodiments, the nominal composition comprises one of Co-5Al-1Ti-3Ta-3V; Co-6Al-3Ta-3V; Co-2.5Ni-5Al-1Ti-3Ta-3V; Co-2.5Ni-6Al-3Ta-3V; Co-10Ni-5Al-2Ti-3Ta-3V-0.04B; Co-10Ni-5Al-2Ti-3Ta-3V-0.04B-5Cr; Co-6Al-3Nb-3V; Co-5Al-1Ti-3Nb-3V; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-4Cr; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-8Cr; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-10Cr; Co-10Ni-7.5Al-3Ti-4.5Nb-4.5V-0.04B-4Cr; Co-10Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr; and Co-20Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr.
In certain embodiments, the nominal composition comprises Co-10Ni-5Al-2Ti-3Ta-3V-0.04B, and the γ′-nanoprecipitates have a composition of (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01).
In certain embodiments, the nominal composition further comprises less than 2 at. % of Mo and/or W.
In certain embodiments, the nominal composition further comprises one or more of C, O, Mn, Y, Fe, Si, B, Zr, Hf, Ru and Re.
In certain embodiments, the γ+γ′ microstructures are stable up to 1500 hours at a temperature of about 600-1100° C.
In certain embodiments, the γ′ phase is presented near grain boundaries with or without carbide and/or borides phases, beneficial for creep resistance.
In certain embodiments, less than 5 vol. % of other deleterious phases or no other deleterious phases other than the γ, γ′ carbide and/or borides phases are formed in the superalloy.
In certain embodiments, the superalloy is a tungsten-free and/or molybdenum-free cobalt based superalloy.
In another aspect, the invention relates to a method for producing a cobalt based superalloy. In one embodiment, the method includes providing a nominal composition comprising at least cobalt (Co), aluminum (Al), Z and vanadium (V), Z being at least one of tantalum (Ta) and niobium (Nb); and arc-melting the nominal composition under a partial Ar atmosphere to obtain an ingot; homogenizing the ingot at a first temperature for a first period of time, followed by water quenching; and performing aging heat treatment of the homogenized ingot at a second temperature for a second period of time, followed by water quenching, to form the cobalt-based superalloy that comprises γ and γ′ phases with stable γ+γ′ microstructures.
In certain embodiments, the ingots is melted a number of times under the partial Ar atmosphere and flipped between each melting cycle to improve homogeneity.
In certain embodiments, the method, prior to homogenizing the ingot, further comprises vacuum-encapsulating the ingot in a quartz ampoule.
In certain embodiments, the method, prior to performing aging heat treatment, further comprises vacuum-encapsulating the homogenized ingot in a quartz ampoule.
In certain embodiments, the first temperature is in a range of 900-1500° C., and the first period of time is in a range of 1-60 hours.
In certain embodiments, the second temperature is in a range of 600-1100° C., and the second period of time is in a range of 1-1500 hours.
In certain embodiments, the nominal composition comprises one of
In certain embodiments, the nominal composition comprises one of Co-5Al-1Ti-3Ta-3V; Co-6Al-3Ta-3V; Co-2.5Ni-5Al-1Ti-3Ta-3V; Co-2.5Ni-6Al-3Ta-3V; Co-10Ni-5Al-2Ti-3Ta-3V-0.04B; Co-10Ni-5Al-2Ti-3Ta-3V-0.04B-5Cr; Co-6Al-3Nb-3V; Co-5Al-1Ti-3Nb-3V; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-4Cr; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-8Cr; Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-10Cr; Co-10Ni-7.5Al-3Ti-4.5Nb-4.5V-0.04B-4Cr; Co-10Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr; and Co-20Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr.
In certain embodiments, the nominal composition comprises Co-10Ni-5Al-2Ti-3Ta-3V-0.04B, and the γ′-nanoprecipitates have a composition of (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01).
In certain embodiments, the nominal composition further comprises less than 2 at. % of Mo and/or W.
In certain embodiments, the nominal composition further comprises one or more of C, O, Mn, Y, Fe, Si, B, Zr, Hf, Ru and Re.
In certain embodiments, each element of the nominal composition has a purity of about 99.99%.
In certain embodiments, the γ+γ′ microstructures are stable up to 1500 hours at a temperature of about 600-1100° C.
In certain embodiments, the γ′ phase is presented near grain boundaries with or without carbide and/or borides phases, beneficial for creep resistance.
In certain embodiments, less than 5 vol. % of other deleterious phases or no other deleterious phases other than the γ, γ′ carbide and/or borides phases are formed in the superalloy.
In certain embodiments, the superalloy is a tungsten-free and/or molybdenum-free cobalt based superalloy.
These and other aspects of the invention will become apparent from the following description of the preferred embodiment taken in conjunction with the following drawings, although variations and modifications therein may be affected without departing from the spirit and scope of the novel concepts of the invention.
The following drawings form part of the present specification and are included to further demonstrate certain aspects of the invention. The invention may be better understood by reference to one or more of these drawings in combination with the detailed description of specific embodiments presented herein. The drawings described below are for illustration purposes only. The drawings are not intended to limit the scope of the present teachings in any way.
The present invention will now be described more fully hereinafter with reference to the accompanying drawings, in which exemplary embodiments of the present invention are shown. The present invention may, however, be embodied in many different forms and should not be construed as limited to the embodiments set forth herein. Rather, these embodiments are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the invention to those skilled in the art. Like reference numerals refer to like elements throughout.
The terms used in this specification generally have their ordinary meanings in the art, within the context of the invention, and in the specific context where each term is used. Certain terms that are used to describe the invention are discussed below, or elsewhere in the specification, to provide additional guidance to the practitioner regarding the description of the invention. For convenience, certain terms may be highlighted, for example using italics and/or quotation marks. The use of highlighting and/or capital letters has no influence on the scope and meaning of a term; the scope and meaning of a term are the same, in the same context, whether or not it is highlighted and/or in capital letters. It will be appreciated that the same thing can be said in more than one way. Consequently, alternative language and synonyms may be used for any one or more of the terms discussed herein, nor is any special significance to be placed upon whether or not a term is elaborated or discussed herein. Synonyms for certain terms are provided. A recital of one or more synonyms does not exclude the use of other synonyms. The use of examples anywhere in this specification, including examples of any terms discussed herein, is illustrative only and in no way limits the scope and meaning of the invention or of any exemplified term. Likewise, the invention is not limited to various embodiments given in this specification.
It will be understood that, although the terms first, second, third, etc. may be used herein to describe various elements, components, regions, layers and/or sections, these elements, components, regions, layers and/or sections should not be limited by these terms. These terms are only used to distinguish one element, component, region, layer or section from another element, component, region, layer or section. Thus, a first element, component, region, layer or section discussed below can be termed a second element, component, region, layer or section without departing from the teachings of the present invention.
It will be understood that, as used in the description herein and throughout the claims that follow, the meaning of “a”, “an”, and “the” includes plural reference unless the context clearly dictates otherwise. Also, it will be understood that when an element is referred to as being “on,” “attached” to, “connected” to, “coupled” with, “contacting.” etc., another element, it can be directly on, attached to, connected to, coupled with or contacting the other element or intervening elements may also be present. In contrast, when an element is referred to as being, for example, “directly on.” “directly attached” to, “directly connected” to, “directly coupled” with or “directly contacting” another element, there are no intervening elements present. It will also be appreciated by those of skill in the art that references to a structure or feature that is disposed “adjacent” to another feature may have portions that overlap or underlie the adjacent feature.
It will be further understood that the terms “comprises” and/or “comprising,” or “includes” and/or “including” or “has” and/or “having” when used in this specification specify the presence of stated features, regions, integers, steps, operations, elements, and/or components, but do not preclude the presence or addition of one or more other features, regions, integers, steps, operations, elements, components, and/or groups thereof.
Furthermore, relative terms, such as “lower” or “bottom” and “upper” or “top,” may be used herein to describe one element's relationship to another element as illustrated in the figures. It will be understood that relative terms are intended to encompass different orientations of the device in addition to the orientation shown in the figures. For example, if the device in one of the figures is turned over, elements described as being on the “lower” side of other elements would then be oriented on the “upper” sides of the other elements. The exemplary term “lower” can, therefore, encompass both an orientation of lower and upper, depending on the particular orientation of the figure. Similarly, if the device in one of the figures is turned over, elements described as “below” or “beneath” other elements would then be oriented “above” the other elements. The exemplary terms “below” or “beneath” can, therefore, encompass both an orientation of above and below.
Unless otherwise defined, all terms (including technical and scientific terms) used herein have the same meaning as commonly understood by one of ordinary skill in the art to which the present invention belongs. It will be further understood that terms, such as those defined in commonly used dictionaries, should be interpreted as having a meaning that is consistent with their meaning in the context of the relevant art and the present disclosure, and will not be interpreted in an idealized or overly formal sense unless expressly so defined herein.
As used in this disclosure, “around”, “about”, “approximately” or “substantially” shall generally mean within 20 percent, preferably within 10 percent, and more preferably within 5 percent of a given value or range. Numerical quantities given herein are approximate, meaning that the term “around”, “about”, “approximately” or “substantially” can be inferred if not expressly stated.
As used in this disclosure, the phrase “at least one of A, B, and C” should be construed to mean a logical (A or B or C), using a non-exclusive logical OR. As used herein, the term “and/or” includes any and all combinations of one or more of the associated listed items.
Embodiments of the invention are illustrated in detail hereinafter with reference to accompanying drawings. The description below is merely illustrative in nature and is in no way intended to limit the invention, its application, or uses. The broad teachings of the invention can be implemented in a variety of forms. Therefore, while this invention includes particular examples, the true scope of the invention should not be so limited since other modifications will become apparent upon a study of the drawings, the specification, and the following claims. For purposes of clarity, the same reference numbers will be used in the drawings to identify similar elements. It should be understood that one or more steps within a method may be executed in different order (or concurrently) without altering the principles of the invention.
We have recently found metastable γ-γ′ microstructures in the Co—Ta—V and Co—Nb—V systems, where γ′ precipitates are seen after short aging times of 2 hours, but they coarsen, lose their cuboidal shape, dissolve and are consumed by other deleterious phases at longer aging times, due to their metastable nature. Therefore, there is the need to achieve a stable γ′ phase in the Co—Ta—V and Co—Nb—V ternary systems.
In certain aspects, the invention discloses new families of tungsten-free, and/or molybdenum-free cobalt based superalloys with stable γ′ strengthening precipitates without formation of deleterious other phases, based on the Co—Ta—V and Co—Nb—V ternary systems by utilizing additions of aluminum (Al) and optionally titanium (Ti) to stabilize the γ′ phase in these ternary compositions while avoiding prior art where tungsten (W) and molybdenum (Mo) are used, or where very high V contents are used or where Al is missing. It is noted that, in future derivative alloys based on the ones disclosed in the disclosure, it might be desirable to add again small amounts of W and/or Mo, but not at the high levels found in prior art.
Additionally, in certain embodiments, nickel (Ni) is used to expand the γ′ phase field while boron (B) is added to provide solid solution and grain boundary strengthening. Chromium (Cr) is added to improve the corrosion resistance of the alloys and boron to strengthen grain boundaries, with both also providing solid solution strengthening. Alloys are arc-melted from high purity elements (Co, Ni, Al. Ta, Nb. V. Ti, Cr and B) under a partial Ar atmosphere, and then subjected to a solution heat treatment to obtain a homogeneous composition and to dissolve any phases formed during solidification/cooling. Millimeter size sections of the ingots are aged at elevated temperatures to precipitate the γ′ strengthening phase.
Furthermore, in certain embodiments, many other elements can be added to the new families of alloys, e.g., small quantities (<2 at. %) of Mo and/or W, as well as C, O, Mn, Y, Fe, Si, B. Zr. Hf, Ru, Re, etc., for oxidation-, corrosion-, coarsening-, fracture-, fatigue- and deformation resistance, as well as improved processability (e.g., casting and thermomechanical processing).
According to the invention, one family of the cobalt based superalloys is Co—Al—Ta—V-based alloys, where Ti, Cr and Ni can be further added. In one embodiment, a single-phases Co-based solid-solution (γ phase) is achieved after homogenization at about 1200-1300° C. After 1000 hours of aging at a temperature of about 850° C., the matrix only contain γ′ precipitates showing (a) fine cuboidal structure within γ grains and (b) coarsened structure at the γ grain boundaries (GB); no other deleterious phases are observed after homogenization and/or aging.
Another family of the cobalt based superalloys is Co—Al—Nb—V—Ti based-alloys with Ni and Cr additions. In one embodiments, a single phase γ Co-based solid solution is achieved at 1200-1300° C., while the γ′ phase is seen near the GBs for Cr containing alloys. After one week of aging at about 850° C., a γ-γ′ microstructure is achieved with no other detrimental phases. Boron (0.01-1 at. %) is added for solid solution strengthening and GB strengthening. Chromium (Cr) and aluminum (Al) are added for oxidation and corrosion resistance, as well as precipitation and solid solution strengthening.
Yet another family of the cobalt based superalloys is a hybrid alloy containing both Ta and Nb, rather than one or the other, e.g., Co-20Ni-5Al-4Cr-3V-2Ti-1.5Nb-1.5Ta-0.08B. In one embodiment, the hybrid alloy shows single-phase γ Co-based solid-solution after homogenization, and γ′ precipitates after aging, with no other deleterious phases.
Generally, according to the invention, the cobalt based superalloy includes a nominal composition comprising at least Co, Al, Z and V, Z being at least one of Ta and Nb, which is processed such that the superalloy comprises γ and γ′ phases with stable γ+γ′ microstructures.
In certain embodiments, for the Co—Al—Ta—V-based alloys, the nominal composition comprises Co-a1Ni-b1Al-c1Ti-d1Ta-f1V-g1B-h1Cr, where a1 is in a range of about 0-40 at. %, b1 is in a range of about 2.5-10 at. %, c1 is in a range of about 0-4 at. %, d1 is in a range of about 2-4 at. %, f1 is in a range of about 1.5-6 at. %, g1 is in a range of about 0-0.1 at. %, h1 is in a range of about 0-20 at. %, and Co is in balance. In certain embodiments, the nominal composition comprises one of Co-5Al-1Ti-3Ta-3V, Co-6Al-3Ta-3V. Co-2.5Ni-5Al-1Ti-3Ta-3V, Co-2.5Ni-6Al-3Ta-3V, Co-10Ni-5Al-2Ti-3Ta-3V-0.04B, and Co-10Ni-5Al-2Ti-3Ta-3V-0.04B-5Cr, which are listed in Table 1 as 11, 12, 13, 14, NT1 and NT2, respectively.
In one embodiment, when the nominal composition comprises Co-10Ni-5Al-2Ti-3Ta-3V-0.04B, the γ′-nanoprecipitates in the cobalt based superalloy have a composition of (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01).
In certain embodiments, for the Co—Al—Nb—V—Ti based-alloys, the nominal composition comprises Co-a2Ni-b2Al-c2Ti-e2Nb-f2V-g2B-h2Cr, where a2 is in a range of about 0-40 at. %, b2 is in a range of about 2.5-10 at. %, c2 is in a range of about 0-4 at. %, e2 is in a range of about 2-4 at. %, f2 is in a range of about 1.5-6 at. %, g2 is in a range of about 0-0.1 at. %, h2 is in a range of about 0-20 at. %, and Co is in balance. In certain embodiments, the nominal composition comprises one of Co-6Al-3Nb-3V, Co-5Al-1Ti-3Nb-3V, Co-10Ni-5Al-2Ti-3Nb-3V-0.04B, Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-4Cr, Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-8Cr, Co-10Ni-5Al-2Ti-3Nb-3V-0.04B-10Cr, and Co-10Ni-7.5Al-3Ti-4.5Nb-4.5V-0.04B-4Cr, which are listed in Table 1 as NN1, NN2, NN3, NN4, NN5, NN6 and NN7, respectively.
In certain embodiments, for the hybrid alloy containing both Ta and Nb, the nominal composition comprises Co-a3Ni-b3Al-c3Ti-d3Ta-e3Nb-f3V-g3B-h3Cr, where a3 is in a range of about 0-40 at. %, b3 is in a range of about 2.5-10 at. %, c3 is in a range of about 0-4 at. %, d3 is in a range of about 2-4 at. %, e3 is in a range of about 2-4 at. %, f3 is in a range of about 1.5-6 at. %, g3 is in a range of about 0-0.1 at. %, h3 is in a range of about 0-20 at. %, and Co is in balance. In certain embodiments, the nominal composition comprises one of Co-10Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr, and Co-20Ni-5Al-2Ti-1.5Ta-1.5Nb-3V-0.04B-4Cr, which are listed in Table 1 as BNT10 and BNT20, respectively.
In certain embodiments, the nominal composition further comprises less than 2 at. % of Mo and/or W. In certain embodiments, the nominal composition may also include one or more of C, O, Mn, Y, Fe, Si, B, Zr, Hf, Ru and Re.
In certain embodiments, the γ+γ′ microstructures are stable up to 1500 hours at a temperature of about 600-1100° C.
In certain embodiments, the γ′ phase is presented near grain boundaries with or without carbide and/or borides phases, beneficial for creep resistance.
In certain embodiments, less than 5 vol. % of other deleterious phases or no other deleterious phases other than the γ, γ′ carbide and/or borides phases are formed in the superalloy.
In certain embodiments, the superalloy is a tungsten-free and/or molybdenum-free cobalt based superalloy.
In another aspect, the invention also discloses a method for producing a cobalt based superalloy. In one embodiment, the method includes providing a nominal composition comprising at least Co, Al, Z and V, Z being at least one of Ta and Nb; and arc-melting the nominal composition under a partial Ar atmosphere to obtain an ingot; homogenizing the ingot at a first temperature for a first period of time, followed water quenching; and performing aging heat treatment of the homogenized ingot at a second temperature for a second period of time, followed water quenching, to form the cobalt based superalloy that comprises γ and γ′ phases with stable γ+γ′ microstructures.
In certain embodiments, the ingots is melted a number of times under the partial Ar atmosphere and flipped between each melting cycle to improve homogeneity.
In certain embodiments, the method, prior to homogenizing the ingot, further comprises vacuum-encapsulating the ingot in a quartz ampoule.
In certain embodiments, the method, prior to performing aging heat treatment, further comprises vacuum-encapsulating the homogenized ingot in a quartz ampoule.
In certain embodiments, the first temperature is in a range of 900-1500° C., and the first period of time is in a range of 40-60 hours.
In certain embodiments, the second temperature is in a range of 600-1100° C., and the second period of time is in a range of 1-1500 hours.
In certain embodiments, each element of the nominal composition has a purity of about 99.99%.
The exemplary embodiments of the cobalt based superalloys are listed in Table 1, and their characterizations are listed in Table 2 and presented in
The cobalt based alloys according to the invention have, among other things, advantages over the existing alloys.
As compared to Ni-based superalloys (currently commercial), the cobalt based alloys of the invention have a higher melting point than nickel, which may increase turbine operating temperature, thus achieving higher efficiency. In addition, the cobalt based alloys have better corrosion resistance than nickel alloys.
As compared to W-containing Co-based superalloys based on Co—W—Al: the W-free cobalt based alloys of the invention achieve lower density enabling higher turbine rotations (and thus efficiency and power) for a given stress.
As compared to Mo-containing Co-based superalloys: the Mo-free cobalt based alloys of the invention do not have MoO3 formation and vaporization problems at temperatures as low as 500° C. Our Co—Al—Nb—V based alloys have very low density, given that Nb and V (8.6 and 6.0 g/cc) have densities lower than Mo (10.3 g/cc), and much lower than Ta (16.7 g/cc) or W (19.3 g/cc).
As compared to a Co-11Ti-15Cr superalloy (at. %), the cobalt based alloys of the invention contain Al for better oxidation/corrosion resistance, and V for low density. Also, the cobalt based alloys include Nb or Ta refractory elements which decrease γ′ coarsening rates and provide solid solution strengthening.
As compared to very recent a Co-5Ti-15V-2Al superalloy (at. %), the cobalt based alloys of the invention contain much less V (thus having better corrosion and oxidation resistance) and more Al (for better oxidation resistance and lower density). Also, the cobalt based alloys include Nb or Ta refractory elements which decrease γ′ coarsening rates and provide solid solution strengthening
The invention may find a wide spectrum of applications in, for example, parts for gas turbines, jet engines and diesel turbo engines, where one must combine high-temperature strength and stability as well as oxidation/corrosion resistance, and turbine blades and disks where one most combine low density, high-temperature creep resistance and good oxidation/corrosion/fatigue resistance.
These and other aspects of the present invention are further described below. Without intent to limit the scope of the invention, examples according to the embodiments of the present invention are given below. Note that titles or subtitles may be used in the examples for convenience of a reader, which in no way should limit the scope of the invention. Moreover, certain theories are proposed and disclosed herein; however, in no way they, whether they are right or wrong, should limit the scope of the invention so long as the invention is practiced according to the invention without regard for any particular theory or scheme of action.
γ+γ′ Microstructures in Co—Ta—V and Co—Nb—V Ternary Systems
In this exemplary example, the Co—Ta—V and Co—Nb—V ternary systems are investigated in a search for L12-ordered γ′ precipitation. Four alloys with nominal (at. %) composition Co-6Ta-6V, Co-5.4Ta-6.6V-xNi (x=0 and 10), and Co-6Nb-6V are arc-melted, homogenized at 1250° C., and aged at 900° C. for 2, 16 and 64 hours. Nanometric, cuboidal γ′ precipitates within a fcc-γ matrix are discovered in the Co—Ta—V system after aging for 2 hours, and in the Co—Nb—V system after cooling from homogenization. The compositions of these two new γ′-phases, as measured via atom probe tomography, are Co3(Ta0.76V0.24) and Co3(Nb0.65V0.35), respectively. Upon aging 900° C., the γ′ precipitates coarsen, dissolve and transform to lamellar C36-Co3(Ta,V) and needle-shape D019-Co3(Nb,V), measured as Co3(Nb0.81V0.19) by APT, respectively. This shows that these ternary γ′ phases are metastable and points to the need for their stabilization via additional alloying elements.
In the exemplary study, the formation of γ′ precipitates is experimentally demonstrated in the Co—Ta—V and the Co—Nb—V system. The ternary alloys have nominal compositions (at. %) of Co-6Nb-6V, Co-6Ta-6V, with equiatomic solute additions. Nyshadham et al. computationally predicted Co3(Nb0.5V0.5) and Co3(Ta0.5V0.5) are L12-ordered γ′ phases, which are not previously reported in phase diagrams in standard databases. Nyshadham et al. calculate that Co3(Nb0.5V0.5) and Co3(Ta0.5V0.5) have lower decomposition energy (19 and 18 meV/atom, respectively) and lower formation enthalpy (−156 and −189 meV, respectively) than the γ′-Co3(Al0.5W0.5) phase found in cobalt-base superalloy (66 meV/atom and −130 meV at T=0 K, respectively). They also predict that the two new L12-γ′ phases are in stable two-phase equilibrium with a fcc-γ Co-rich matrix, with a low lattice mismatch of 2%. In the exemplary example, two experimental alloys, as well as two modified ternary and quaternary alloys (Co-5.4Ta-6.6V-xNi with x=0 and 10 at. %), are aged at various times at 900° C. and their microstructures are studied via scanning electron microscopy (SEM) and atom probe tomography (APT) to search for the predicted γ′ phases. Additionally, the formation of non-γ′ phases is identified by microstructure analysis.
Button ingots of about 20 g with nominal composition Co-6Ta-6V, Co-5.4Ta-6.6V, Co-10Ni-5.38Ta-6.55V and Co-6Nb-6V (all compositions being given hereafter in at. %) were produced by arc melting of high purity Co (99.99%), Ta (99.99%), V (99.99%), Nb (99.99%) and Ni (99.99%), under a partial Ar atmosphere. The button ingots where remelted six times and flipped between each melting step to ensure homogenous distribution of the constituent elements in the alloys. Button ingots for Co-6Ta-6V and Co-6Nb-6V alloys were then homogenized in a high-vacuum furnace at 1250° C. for 48 hours and furnace-cooled to room temperature. Ingots for alloys Co-5.4Ta-6.6V and Co-10Ni-5.38Ta-6.55V where vacuum-encapsulated in quartz ampoules, homogenized at 1250° C. for 48 hours followed by water quenching. Sections of these ingots were vacuum-encapsulated in quartz ampoules and aged at 900° C. for 2, 16 and 64 hours, followed by water quenching.
Microstructure characterization and composition analysis were performed using SEM and energy dispersive spectroscopy (EDS) with a Hitachi SU8030 SEM equipped with an Oxford Aztec SDD EDS detector. As-homogenized and aged specimens used for SEM were grinded using 320, 400, 600, 800, 1200 grit SiC paper, polished using 6, 3, and 1 μm diamond suspension, and chemically etched at room temperature using a solution of 33% hydrochloric acid (12.1M), 33% acetic acid and 1% hydrofluoric acid volume in de-ionized water. SEM micrographs were taken from grains orientated close to a {100}-type plane.
To determine the partition behavior and the composition of the γ- and γ′ phases, 2 hour-aged samples of the Co-5.4Ta-6.6V and Co-6Nb-6V alloys were prepared into nanotip specimens for APT studies via a lift-out technique with a Ga+ dual-beam focused-ion beam (FIB) microscope using a FEI Helios Nanolab SEM/FIB. Areas displaying a γ+γ′ and a γ+D019 microstructure were extracted from the samples by creating rectangle-topped wedges using the FEI dual-beam FIB and attached to a Si micropost on a coupon. Sectioned wedges on the silicon micropost were then sharpened using Ga+ to a about 25 nm minimum radius. APT was performed using a Cameca local-electrode atom-probe (LEAP) 4000X-Si system (a 5000X-Si system was used for the 64-hour aged Co-6Nb-6V condition) with picosecond ultraviolet (wavelength=355 nm) laser with a specimen temperature of about 25 K, a pulse energy of 30 pJ, a 500-kHz pulse repetition rate and a 4% detection rate. Differential Scanning calorimetry (DSC) was performed on 2 hours-aged and on homogenized Co-6Nb-6V alloys from 300° C. to 500° C. (at a rate of 10°/min) using a Mettler-Toledo DSC822e. All phases presented here are studied based on morphology, microstructure and composition (using EDS and APT) in comparison with other published studies.
Microstructure Evolution for Co—Ta—V:
SEM micrographs for the Co-6Ta-6V alloy aged at 900° C. for 2 and 64 hours. A representative high-magnification micrograph is shown in portion a) of
Portion c) of
With the goal of stabilizing the γ′-L12 precipitates present at early aging times in the Co-6Ta-6V alloy, two new alloys were arc-melted with nominal compositions Co-5.4Ta-6.6V-xNi (x=0 and 10 at. %). The modified ternary alloy composition of Co-5.4Ta-6.6V corresponds to the SEM-EDS measurements obtained from the γ+γ′ regions of the Co-6Ta-6V aged for 2 hours at 900° C.
Apart from the alloys disclosed above, a button ingot of about 20 g with nominal composition Co-5.4Ta-6.6V-2.5Ni was produced using high purity Co (99.99%), Ta (99.99%), V (99.99%), and Ni (99.99%). The button ingot was arc melted under a partial Ar atmosphere and homogenized following the same procedure described above. Microstructure characterization was performed using SEM.
The Co-5.4Ta-6.6V-10Ni quaternary alloy shows a microstructure very similar to the Co-5.4Ta-6.6V alloy at 2 hours of aging, with the γ+γ′ microstructure present together with the globular C36 primary precipitates, as illustrated in
After short aging times, in both Ni-free and Ni-containing Co-5.4Ta-6.6V alloys, the L12-ordered γ′ phase cuboidal precipitates transform discontinuously into lamellar domains of γ-Co+C36, as illustrated in portions b) and c) of
Microstructure Evolution for Co-6Nb-6V:
Also present in the alloy are two-phase regions of fcc-γ+Co2(Nb,V) based on EDS and shown at higher magnification in portions a)-c) of
APT Study for Co-5.4Ta-6.6V:
APT Study for Co-6Nb-6V: Three-dimensional APT reconstructions for the Co-6Nb-6V alloy aged for 2 and 64 hours at 900° C. are shown in
The APT reconstructed tip for the 64 hours aged condition contains one needle-like precipitate, taken to be D019-Co3(Nb,V) based on composition as shown in
In the exemplary study, Co-rich ternary Co—Ta—V and Co—Nb—V systems are experimentally explored. Fine γ′ cuboidal precipitates in both Ta- and Nb-containing alloys were observed indeed, albeit with the presence of other phases: (i) C36 for the Co—Ta—V system and (ii) C15 and D019 for Co—Nb—V system. The Co—Ta—V alloy composition was refined in an attempt to prevent the formation of the undesirable C36 phases, and nickel was added with the same goal (as Ni is known to widen the γ-γ′ phase field in Co—Al—W alloys [29,30]); neither strategies succeeded and the γ′ phase in both alloy systems remained metastable. Additional computational efforts are now warranted to identify alloying elements that stabilize the γ′ phase in Co—Ta—V- and Co—Nb—V-based compositions, together with Al and Cr for oxidation and corrosion resistance.
Microstructure and phases in Co—Ta—V(—Ni) alloys: A previous study by Ruan et al. focused on the experimental microstructural investigation and phase equilibria in the Co—Ta—V system. Ruan et al. created experimental isothermal phase diagrams sections at 900, 1100, 1200 and 1300° C. based on their experimental study and the binary phase diagrams (Co—V, Co—Ta and Ta—V). They did not report the presence of an L12-ordered phase, consistent with our findings that the γ′ phase is metastable and their very long aging times to achieve equilibrium (up to 90 days). Recently, it was demonstrated via transmission electron microscopy (TEM) that a metastable L12-ordered γ′-Co3Ta phase exists in the Co—Ta binary system, for a Co-8.5Ta alloy aged at 1000° C. for 1 hour [27]. Furthermore, Nagel and Fultz [31] have reported the observation of a metastable L12-ordered γ′-Co3V phase, using in-situ neutron diffraction, produced by quenching in the Co—V binary system [31]. In the exemplary embodiments, we find ternary γ′ precipitates in the Co—Ta—V ternary system via metallographic/SEM, as shown in
The formation of lamellar domains by the discontinuous precipitation of two phase has been well studied over the years for different alloy systems [33-36], including Co-[15, 17, 37, 38] and Ni-based superalloys, and can be describe by the Turnbull theory of cellular precipitation [40]. The discontinuous transformation of the metastable microstructure occurs by both boundary precipitation and interfacial migration [39], with cell boundary diffusion most likely acting as the limiting growth rate factor [33-36]. The rapid transformation of the γ+γ′ microstructure indicates that the lamellar γ-Co+C36 is more energetically stable. Discontinuous precipitation into a lamellar microstructure (like those seen in our Co—Ta—V alloys) was reported after 35 hours at 800° C. by Makineni et al. [17] for W-free Co—Mo—Al—Nb superalloys with γ-γ′ microstructure.
Microstructure and phases in the Co-6Nb-6V alloy: Wang et al. [26] focused on the experimental microstructural investigation and phase equilibria in the Co—Nb—V system. The authors produced multiple isothermal phase diagram sections based on their experimental study at 1000 and 1200° C. and the binary phase diagrams (Co—V, Co—Nb and Nb—V). Like the previously mentioned study on the Co—Ta—V system [25], the authors did not report the presence of an L12-ordered phase for any of their compositions, consistent with their very long aging times (25 to 50 days). Interestingly, there is a phase with the Co3Nb stoichiometry in the Co—Nb binary phase diagram with the Laves MgNi2 type structure (C36) [42], which is however stable only between 1005 and 1249° C. Given that it has been demonstrated that the addition of Nb into another system (Co-10Al-5Mo-2Nb) has helped stabilized the γ′ phase [16, 19], it is plausible that this approach is also applicable to the Co—V system. Indeed, for our Co—Nb—V alloy, we find cuboidal γ′ precipitates in homogenized and furnace-cooled conditions. These samples are then aged at 900° C. and quenched. After 2 hours of aging, the γ′ precipitates transform into needle-like precipitates (taken to be D019) with a Widmanstätten morphology, comparable to what is seen during the transformation from fcc to D019 in Co—W alloys and, L12 to D019 in Co—Al—W and Co-10Al-5Mo-2Nb [17] alloys.
In sum, in the exemplary embodiments, we report the discovery, in ternary Co—Nb—V and Co—Ta—V alloys, of metastable γ′ cuboidal nanosize precipitates present within a terminal γ-fcc Co-rich solid solution. These γ′ precipitates are formed after homogenization and aging at 900° C. in alloys with nominal compositions (at. %): Co-6Nb-6V and Co-6Ta-6V, as well as Co-5.4Ta-6.6V (composition refined based on EDS measurements of the matrix of the Co-6Ta-6V alloy) and Co-5.4Ta-6.6V-10Ni (as Ni is known to stabilize γ′ in Co—Al—W alloys). The phases present are: γ, γ′ and C36-Co3(Ta,V) for the Co—Ta—V alloys; Y. γ′, C15-Co2(Nb,V) and D019-Co3(Nb,V) for the Co—Nb—V alloy.
According the exemplary study, the following conclusions are drawn:
Micron-size C36 spheroidal particles, consistent with dendritic solidification, are present within a γ matrix in the Co-6Ta-6V alloy and, to a lesser extent, in the Co-5.4Ta-6.6V-0/10Ni alloys. In the Co-6Nb-6V alloy, regions of micron-size C15 phase within a γ matrix (expected to have formed on solidification) coexist with regions with fine γ+γ′ microstructure (expected to have formed during furnace-cooling after homogenization).
SEM-based microstructure analysis shows that cuboidal, nanometric γ′ precipitates are present after short aging times (2 hours at 900° C.) in all alloys. Degradation of the cuboidal shape at 900° C., accompanied by coarsening and dissolution, occurs after 2 hours for Co—Nb—V and 16 hours for Co—Ta—V alloys.
The composition of the cuboidal γ′ precipitates, as measured by APT after 2 hours of aging, is Co3(Ta0.76V0.24) and Co3(Nb0.65V0.35) for alloys Co-5.4Ta-6.6V and Co-6Nb-6V, respectively.
The γ′ phase in the Co—Ta—V alloys is metastable at 900° C., as it is consumed via the discontinuous precipitation of lamellar γ-Co+C36-Co3(Ta,V). Transformation starts after 16 hours and is mostly completed after 64 hours.
Nickel addition of 10 at. % does stabilize the γ′ phase in Co-5.4Ta-6.6V and in fact promotes the formation of lamellar γ-Co+C36-Co3Ta at the expense of the γ+γ′ microstructure.
The γ′ phase in the Co-6Nb-6V alloy is formed during furnace cooling after homogenization at 1250° C. and is also metastable at 900° C. After short aging times (2 hours) at that temperature, the γ′ phase is consumed by the growth of D019 needles with Widmanstätten morphology and Co3(Nb0.81V0.19) composition, as measured by APT after 64 hours of aging at 900° C.
Effects of Al, Ti and Cr Additions on γ-γ′ Microstructure of W-Free Co—Ta—V-Based Superalloys
This exemplary study shows that the recently-discovered metastable γ′-Co3(Ta0.76V0.24) phase formed on aging in a Co-6Ta-6V (at. %) ternary alloy can be stabilized by partial replacement of Ta and V with Al and Ti. In two alloys with composition Co-6Al-3Ta-3V and Co-5Al-3Ta-3V-1Ti with γ+γ′ microstructure, the γ′-precipitates remain stable for up to 168 hours at 850 and 900° C., with no precipitation of additional phases. Adding Ni and B and doubling the Ti concentration produces a γ/γ′ superalloy, Co-10Ni-5Al-3Ta-3V-2Ti-0.04B (at. %), with γ′ precipitates which are stable up to 6 weeks of aging at 850° C., while slowly coarsening and coalescing from cuboidal to elongated shapes. After 1 day of aging at 850° C. the γ′ nanoprecipitates have (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01) composition, with Al and Ti replacing at the same rate both Ta and V in the original metastable Co3(Ta0.76V0.24) phase. To improve oxidation resistance, 4% Cr is added to the new superalloy, resulting in a somewhat higher volume fraction of finer cuboidal γ′ precipitates after one week of aging at 850° C., but no other deleterious phases. These W- and Mo-free γ/γ′ superalloys show good creep resistance at 850° C., on a par with two other recent Co-base γ/γ′ superalloys: (i) Co-9W-9Al-8Cr (at. %) which has higher density due to its high W content, and (ii) Co-30Ni-10Al-5Mo-2Nb (at. %) which has lower density (as it is W-free) but contains triple the Ni concentration.
In the exemplary study, we identify alloying elements that stabilize the metastable γ′ phase in Co—Ta—V-based alloys and to find optimal concentrations of the constituent elements to achieve a γ/γ′ superalloy with good creep-, coarsening- and oxidation resistance. In certain embodiments, for new alloying elements, we focus here on (i) Al and Ti, to partition to and stabilize the γ′ phase, (ii) Ni, to expand the γ+γ′ the phase field, (iii) Cr, to provide additional oxidation and corrosion resistance, and (iv) B, to provide solid-solution and grain-boundary strengthening in Co-based superalloys. We demonstrate that the γ′ phase is stabilized for a range of modified alloys containing some or all these elements upon aging at 850 and 900° C. and we show that an optimized γ/γ′ superalloy, Co-10Ni-5Al-3Ta-3V-2Ti-0.04B, has excellent creep resistance at 850° C., and can be further modified by 4% Cr while maintaining its γ+γ′ microstructure.
Ingots about 10 g in mass with nominal compositions Co-(2−x)Ni-(6−y)Al-3Ta-3V−yTi (x=0 or 2 and y=0 or 1) and Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4), where all compositions hereafter are given in at. %, were produced by arc-melting of high purity Co (99.99%) Al (99.99%), Ta (99.95%), V (99.95%), B (99.99%), Ti (99.95%) and Ni (99.99%). The ingots were melted five times under a partial Ar atmosphere and flipped between each melt cycle to improve homogeneity. The alloys were then vacuum-encapsulated in quartz ampoules and homogenized at 1200-1250° C. for 48 hours, followed by water quenching. The ingots were cut into about 2×2×6 mm specimens which were vacuum-encapsulated in quartz ampoules for aging heat-treatments: (i) Co-(2−x)Ni-(6−y)Al-3Ta-3V−yTi (x=0 or 2 and y=0 or 1) alloys were aged at 850 and 900° C. for 7.5 and 168 hours (1 week); (ii) Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4) alloys were aged at 850° C. for 168, 500 and 1000 hours (1, 3 and 6 weeks). All alloys were water-quenched following the aging heat treatment.
Alloy cross-sections were imaged via scanning electron microscopy (SEM) using a Hitachi SU8030 SEM and a FEI Quanta 650 SEM, both equipped with an Oxford Aztec silicon drift detector (SDD) energy dispersive spectrometer (EDS). Solutionized specimens, with and without subsequent aging, were ground using 320, 400, 600, 800, 1200 grit SiC paper, polished using 6, 3, and 1 μm diamond suspension, and chemically etched at room temperature using Carapella's re-agent. ImageJ was used to quantify γ′ precipitates in Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4) alloys. A large number of precipitates for each aging condition, between 500 and 750 over 5 SEM micrographs, were hand-traced and their area A was used to calculate the mean circular equivalent radius (<R(t)>=√{square root over (A/π)}). Also, the γ′-precipitate area fraction (φ) was determined using ImageJ thresholding and averaged over 11 micrographs. In an isotropic microstructure, this area fraction would be equal to the volume fraction. However, given the regular arrangement of the precipitates, this assumption can lead to an overestimation of the precipitate volume fractions. Nevertheless, changes in area fractions would be equivalent to changes of similar magnitudes in the volume fractions. Moreover, the Local Thickness function of BoneJ (ImageJ plugin) was used to determine the thickness (taken to be the smallest dimension between width and height of a precipitate) of the γ′-precipitate, to consider directional coarsening and coalescence.
Vickers microhardness tests were conducted at ambient temperature on polished cross-sections of the Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4) alloys using a Struers Duramin-5 microhardness tester with an applied load of 3 kg and a 5 s dwell time. Each hardness value is the average of 20 measurements performed over multiple grains. High-temperature compression creep was performed on the Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4) alloy in air at 850° C. for different stresses, using a dead-low creep frame. Creep samples were machined in the form of cylinders (6 mm in height and 3 mm in diameter) by wire electron discharge machining (EDM) and aged for 168 hours at 850° C., followed by water quenching. A creep specimen, water-quenched immediately after unloading, was sectioned parallel and perpendicular to the applied load. Sample preparation for SEM imaging was done in the same manner as for the aged stress-free specimens.
APT was used to determine the elemental partition behavior and composition of the γ- and γ′ phases in Co-10Ni-5Al-3Ta-3V-2Ti-0.04B aged for 24 hours at 850° C. Nanotip specimens were prepared via an FIB lift-out technique using a Ga+ dual-beam in a FEI Helios Nanolab SEM/FIB instrument. Regions of the representative γ-γ′ microstructure were extracted by creating rectangle-topped wedges using the dual-beam FIB which were attached to a Si micropost on a coupon and sharpened using Ga+ to a about 40 nm minimum tip radius. APT was then performed using a Cameca LEAP 5000X-Si system with a picosecond ultraviolet (wavelength=355 nm) laser with a specimen temperature of about 25 K, a pulse energy of 25 pJ, a 500-kHz pulse repetition rate and a 4% detection rate.
DSC was performed on Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B (x=0 and 4) alloys using a NETZSCH STA 449 F3 Jupiter instrument to determine γ′ solvus temperature. Before DSC measurements, homogenized samples were aged for 24 hours at 850° C. followed by water quenching, and then segmented into 20-50 mg pieces. During DSC experiments, the samples were heated from 25 to 1300° C. (at 20° C./min), under an Ar atmosphere, and held for 1 hour at 1300° C. to allow dissolution of γ′. Cooling, was done from 1300 to 25° C. (at the same rate of 20° C./min) and held at 25° C. for 10 minutes. This cycle was repeated three times.
A. Microstructure Evolution
Effect of Al and Ti Additions on γ′ Phase Stability: Metastable γ′-precipitates were first reported in Co-6Ta-6V after furnace-cooling from homogenization at 1300° C. as well as subsequent aging for 2-16 hours at 900° C. [61]; however, they were consumed by discontinuous precipitation (DP) of a C36 phase with Co3X composition after 64 hours aging at 900° C. To stabilize the γ′-phase in this ternary alloy, half the Ta and V concentrations (3 at. % each) are replaced with Al, with a cumulative amount of 6 at. %, leading to the first alloy studied here, Co-6Al-3Ta-3V. After solution heat-treatment at 1200° C. for 48 hours, the microstructure is single-phase γ, with no γ′ precipitates in grains or at grain boundaries. A γ+γ′ microstructure is formed upon aging at 850 and 900° C., as illustrated in
To achieve a higher γ′-volume fraction, 1 at. % of Al in the first, quaternary alloy is replaced by 1 at. % Ti, leading to the second, quinary alloy, Co-5Al-3Ta-3V-1Ti. After solution heat-treatment at 1200° C. for 48 hours, the alloy shows a precipitate-free γ-phase. A two-phase γ+γ′ microstructure is present after aging at 850 and 900° C. for 7.5 and 168 hours, as shown in
Finally, in both the first and second alloys, 2 at. % Ni was added for potential improvements in the γ′-stability. These modified alloys, Co-2Ni-6Al-3Ta-3V and Co-2Ni-5Al-3Ta-3V-1Ti, were aged for the same temperatures and times (850 and 900° C. for 7.5 and 168 hours) and are shown in
Alloying Element Increases (Ni and Ti) and Additions (Cr and B): In a next series of alloys, the concentration of Ni was increased from 2 to 10 at. % and the Ti concentration was doubled from 1 to 2 at. %, with the aim to further raise γ′-stability and volume fraction, and to decrease density (for Ti additions). Also, B micro-additions (0.04 at. %) were made to achieve grain-boundary- and solid-solution strengthening. Finally, Cr was added to improve oxidation and corrosion resistance, and help decrease the alloy density. Because Cr is known to create additional undesirable phases (e.g., the D019 phase in Co-7Al-7W-(10, 13)Cr and μ (D85) phase in Co-7Al-7W-(21,17)Cr alloys [38]), two alloys were cast, one Cr-free and the other with 4 at. % Cr, with nominal compositions: Co-10Ni-5Al-xCr-3Ta-3V-2Ti-0.04B at. % (x=0 and 4, hereafter referred to as 0Cr- and 4Cr-alloys), corresponding to a wt. % composition of Co-9.8Ni-xCr-2.24Al-9.0Ta-2.5V-1.6Ti-0.007B (x=0 and 3.5).
Comparing the 0Cr and 4Cr alloys at short aging times, as shown in
Previous research has shown that Cr additions to Co-based superalloys contribute to a decrease in γ/γ′ lattice mismatch. As the lattice mismatch decreases, the shape of the coherent γ′ precipitates evolves from cuboidal with sharp corners to cuboidal with rounded corners to spherical. Recent research by Zenk et al. [20, 65], has shown that, in Co-12Ti, a high misfit contributes to the formation of elongated γ′ precipitates with an irregular, non-smooth interface with the γ-matrix. Cr additions decreases the lattice misfit resulting in cuboidal γ′ precipitates with sharp corners and a smooth interface with the γ-matrix [20]. A similar effect is seen in our 0Cr and 4Cr alloys, with γ′-precipitates becoming more cuboidal in the latter alloy after aging for 168 hours at 850° C., as shown in
B. Atom Probe Tomography (APT) Study
An APT volume reconstruction is shown in
As summarized in Table 7, the compositions of the γ- and γ′-phases are measured as Co-8.3Ni-4.9Al-1.2Ta-2.8V-1.4Ti-0B and Co-13Ni-10.5Ta-5.6Al-4.9Ti-3.8V-0.2B, respectively. The γ′-phase stoichiometry is then (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01), assuming that Ni and Co exclusively occupy the A-sublattice of γ′-A3B and that the other five elements exclusively occupy the B sublattice. For comparison, the γ′-phase composition in the ternary Co-6Ta-6V alloy is Co3(Ta0.76V0.24) [61]. A similar Ta/V≈3 ratio thus exists in the γ′-phases of both alloys, indicating that Al and Ti, added in the present 0Cr alloy, show no strong preference for replacing either Ta or V in the ternary Co3(Ta0.76V0.24) phase.
The γ′ volume fraction for this alloy can be found based on elemental concentrations for each phase and bulk tip composition as measured by LEAP (Table 7), using the lever rule:
Ciγ′φ+Ciγ(1−φ)=CiBulk (1)
where φ is the γ′ volume fraction and Ciγ, Ciγ′, and CiBulk are the concentrations of species i (i=Co, Ni, Al, Ta,V, Ti and B) in the γ-phase, γ′-phase and bulk tip, respectively. Rewriting Eq, (1), the γ′ volume fraction is:
φ=(CiBulk−Ciγ)/(Ciγ′−Ciγ) (2)
Thus, the γ′ volume fraction can be obtained by plotting (CiBulk−Ciγ) against (Ciγ′−Ciγ) for each element i and fitting a linear regression whose slope is equal to q. Using this approach, as shown in
C. γ′ Solvus Temperature and Mass Density
The solvus temperature of the 0Cr and 4Cr alloys was measured as 965° C. for both alloys, visible as a sharp exothermic peak in the DSC cooling curve,
D. Mechanical Properties
Hardness at Ambient Temperature:
Creep properties at 850° C.:
where A is a constant, Q is the activation energy, Rg is the ideal gas constant, and T is the absolute temperature. Also shown in
Interestingly, no significant difference in creep resistance is seen between our 0Cr and 4Cr alloys. This is contrary to what was reported by Povstugar et al. where adding 4 and 8 at. % Cr into the ternary Co-9Al-9W caused a dramatic increase in the steady-state strain rate, by about 1 and about 2 orders of magnitude, respectively. The results shown here are more similar to those reported by Chung et al. who studied Co-30Ni-11Al-2Ti-5.5W-2.5Ta-0.1B-xCr (x=0, 4, 8 and 12) alloys with very high (about 98%) γ′ fraction, where Cr additions had little effect on creep resistance. Similarly, a study by Ng et al. of Co-30Ni-7Al-4Ti-3Mo-2W-1Nb-1Ta-0.1B-xCr (x=0, 4, 8 and 12) alloys with 70-80% γ′ fraction shows that Cr does not affect the creep strength of the alloys in a significant matter. Thus, it is likely that Cr does not affect the shear resistance of the L12 precipitates of our 0Cr and 4Cr alloys.
The post-creep microstructure of our 0Cr and 4Cr alloys, after having accumulated about 10% strain during 154 and 44 hours of deformation, respectively, is shown in
In the cross-sections shown in portions b) and f) of
In sum, we report here the stabilization of the metastable γ′-Co3(Ta0.76V0.24) phase, recently found as precipitates in the Co-6Ta-6V (at. %) ternary alloy [61], and the development of a W- and Mo-free γ+γ′ superalloy with good coarsening- and creep resistance (as measured experimentally) and improved oxidation resistance (as expected from substantial Al and Cr contents). The following six alloys were developed, in order of complexity.
At first, the Co-6Ta-6V alloy is modified by replacing half its Ta and V with Al. The resulting Co-6Al-3Ta-3V alloy shows a γ+γ′ microstructure free of other phases, up to the longest aging time of 168 hours studied at 850 and 900° C. By contrast, the γ′ phase of the original ternary Co-6Ta-6V alloy was studied up to 64 hours at 900° C., when more than 50% of the γ′ phase had transformed into lamellar C36.
In a second alloy, 1% Al is replaced with 1% Ti. The resulting Co-5Al-3Ta-3V-1Ti alloy shows no difference in microstructure aging evolution, as compared to the first, Ti-free alloy, indicating that Ti can be added without destabilizing the γ′ phase.
A third and fourth alloy are created by adding 2% Ni to the above compositions. They also show no difference in microstructure after aging at 850 and 900° C. for 168 hours, as compared to the first and second Ni-free alloys. This opens the door to Ni additions.
Based on the previous Ni and Ti additions, a fifth alloy is made by adding 10% Ni and 0.04% B to the above second alloy and by doubling its Ti concentration. This Co-10Ni-5Al-3Ta-3V-2Ti-0.04B alloy shows a stable γ-γ′ microstructure up to 1000 hours (6 weeks) of aging at 850° C. No other phases are present in grains or at grain boundaries, which are decorated with coarsened γ′ precipitates. This is a very significant improvement as compared to the original Co-6Ta-6V in which the metastable γ′ phase transformed to C36 (Co3(Ta,V)) after only about 2-16 hours of aging at 900° C.
For Co-10Ni-5Al-3Ta-3V-2Ti-0.04B aged for short times at 850° C. (24 hours), the composition of the cuboidal γ′-nanoprecipitates is measured by atom probe tomography as (Co0.83Ni0.17)3(Ta0.42Al0.23 Ti0.19V0.15B0.01); Al and Ti thus replace at the same rate both Ta and V in the ternary Co3(Ta0.76V0.24).
The peak hardness of Co-10Ni-5Al-3Ta-3V-2Ti-0.04B (after 1-3 weeks at 850° C.) is greater than most W-free Co-based superalloys and some Co—Al—W-based superalloys. This may be due to strong B partitioning to the γ′-phase (and a concomitant lack of borides at the grain boundaries).
The creep resistance of Co-10Ni-5Al-3Ta-3V-2Ti-0.04B at 850° C., is higher than W-free Co-30Ni-10Al-5Mo-2Nb, and comparable to W-bearing Co-9Al-9W-8Cr alloys.
A 4% Cr addition (to improve oxidation resistance) leads to the sixth alloy. Co-10Ni-5Al-4Cr-3Ta-3V-2Ti-0.04B, which shows a slight increase in γ′ volume fraction and a slight decrease in γ′ size in the early stage (168 hours) of aging at 850° C. At longer aging time, the γ′ precipitate fraction, size and shape are undistinguishable from to those of the non-Cr containing alloy, and so are the hardness evolution and the creep resistance at 850° C.
The foregoing description of the exemplary embodiments of the present invention has been presented only for the purposes of illustration and description and is not intended to be exhaustive or to limit the invention to the precise forms disclosed. Many modifications and variations are possible in light of the above teaching.
The embodiments were chosen and described in order to explain the principles of the invention and their practical application so as to activate others skilled in the art to utilize the invention and various embodiments and with various modifications as are suited to the particular use contemplated. Alternative embodiments will become apparent to those skilled in the art to which the present invention pertains without departing from its spirit and scope. Accordingly, the scope of the present invention is defined by the appended claims rather than the foregoing description and the exemplary embodiments described therein.
This application claims priority to and the benefit of, pursuant to 35 U.S.C. § 119(e), of U.S. Provisional Patent Application Ser. No. 62/674,780, filed May 22, 2018, entitled “TUNGSTEN-FREE, LOW-DENSITY COBALT-BASED SUPERALLOYS WITH GAMMA PRIME PRECIPITATES, BASED ON CO—TA—V—AL AND CO—NB—V—AL SYSTEMS”, by David C. Dunand and Fernando L. Reyes Tirado, which is incorporated herein by reference in its entirety. Some references, which may include patents, patent applications and various publications, are cited in a reference list and discussed in the description of this invention. The citation and/or discussion of such references is provided merely to clarify the description of the invention and is not an admission that any such reference is “prior art” to the invention described herein. All references cited and discussed in this specification are incorporated herein by reference in their entireties and to the same extent as if each reference was individually incorporated by reference. In terms of notation, hereinafter, “[n]” represents the nth reference cited in the reference list. For example, [61] represents the 61th reference cited in the reference list, namely, F. L. Reyes Tirado, J. Perrin Toinin, D. C. Dunand, γ+γ′ microstructures in the Co—Ta—V and Co—Nb—V ternary systems, Acta Mater. 151 (2018) 137-148.
This invention was made with government support under 70NANB14H012 awarded by the National Institute of Standards and Technology. The government has certain rights in the invention.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/US2019/033450 | 5/22/2019 | WO |
| Publishing Document | Publishing Date | Country | Kind |
|---|---|---|---|
| WO2019/226731 | 11/28/2019 | WO | A |
| Number | Name | Date | Kind |
|---|---|---|---|
| 20120312426 | Suzuki et al. | Dec 2012 | A1 |
| 20150354031 | Gehrmann et al. | Dec 2015 | A1 |
| 20170037498 | Makineni et al. | Feb 2017 | A1 |
| Number | Date | Country |
|---|---|---|
| 104630569 | May 2015 | CN |
| 5144270 | Feb 2013 | JP |
| Entry |
|---|
| J.R. Davis, Nickel, Cobalt, and Their Alloys, ASM International, 2000. |
| H. Chinen, J. Sato, T. Omori, K. Oikawa, I. Ohnuma, R. Kainuma, K. Ishida, New ternary compound Co3(Ge, W) with L12 structure, Scr. Mater. 56 (2007) 141-143. |
| K. Shinagawa, T. Omori, J. Sato, K. Oikawa, I. Ohnuma, R. Kainuma, K. Ishida, Phase Equilibria and Microstructure on γ′ Phase in Co—Ni—Al—W System, Mater. Trans. 49 (2008) 1474-1479. |
| L. Zhu, C. Wei, hours. Qi, L. Jiang, Z. Jin, J.C. Zhao, Experimental investigation of phase equilibria in the Co-rich part of the Co—Al—X (X=W, Mo, Nb, Ni, Ta) ternary systems using diffusion multiples, J. Alloys Compd. 691 (2017) 110-118. |
| A. Suzuki, hours. Inui, T.M. Pollock, L12-Strengthened Cobalt-Base Superalloys, Annu. Rev. Mater. Res. 45 (2015) 345-368. |
| J.J. Ruan, X.J. Liu, S.Y. Yang, W.W. Xu, T. Omori, T. Yang, B. Deng, hours.X. Jiang, C.P. Wang, R. Kainuma, K. Ishida, Novel Co—Ti—V-base superalloys reinforced by L12-ordered γ′ phase, Intermetallics. 92 (2018) 126-132. |
| P.J. Bocchini, C.K. Sudbrack, D.J. Sauza, R.D. Noebe, D.N. Seidman, D.C. Dunand, Effect of tungsten concentration on microstructures of Co—10Ni—6Al—(0,2,4,6)W—6Ti (at%) cobalt-based superalloys, Mater. Sci. Eng. A. 700 (2017) 481-486. |
| P.J. Bocchini, C.K. Sudbrack, R.D. Noebe, D.C. Dunand, D.N. Seidman, Microstructural and creep properties of boron- and zirconium-containing cobalt-based superalloys, Mater. Sci. Eng. A. 682 (2017) 260-269. |
| P.J. Bocchini, C.K. Sudbrack, R.D. Noebe, D.C. Dunand, D.N. Seidman, Effects of titanium substitutions for aluminum and tungsten in Co—10Ni—9Al—9W (at%) superalloys, Mater. Sci. Eng. A. 705 (2017) 122-132. |
| J. Sato, Cobalt-Base High-Temperature Alloys, Science (80-). 312 (2006) 90-91. |
| P.J. Bocchini, E.A. Lass, K.W. Moon, M.E. Williams, C.E. Campbell, U.R. Kattner, D.C. Dunand, D.N. Seidman, Atom-probe tomographic study of γ/γ′ interfaces and compositions in an aged Co—Al—W superalloy, Scr. Mater. 68 (2013) 563-566. |
| J.E. Saal, C. Wolverton, Thermodynamic stability of Co—AL—W L12 γ′, Acta Mater. 61 (2013) 2330-2338. |
| M. Ooshima, K. Tanaka, N.L. Okamoto, K. Kishida, hours. Inui, Effects of quaternary alloying elements on the γ′ solvus temperature of Co—Al—W based alloys with fcc/L12 two-phase microstructures, J. Alloys Compd. 508 (2010) 71-78. |
| T. Omori, K. Oikawa, J. Sato, I. Ohnuma, U.R. Kattner, R. Kainuma, K. Ishida, Partition behavior of alloying elements and phase transformation temperatures in Co—Al—W-base quaternary systems, Intermetallics. 32 (2013) 274-283. |
| D.J. Sauza, P.J. Bocchini, D.C. Dunand, D.N. Seidman, Influence of ruthenium on microstructural evolution in a model Co—Al—W superalloy, Acta Mater. 117 (2016) 135-145. |
| S.K. Makineni, B. Nithin, K. Chattopadhyay, A new tungsten-free γ-γ′ Co—Al—Mo—Nb-based superalloy, Scr. Mater. 98 (2015) 36-39. |
| S.K. Makineni, B. Nithin, D. Palanisamy, K. Chattopadhyay, Phase evolution and crystallography of precipitates during decomposition of new “tungsten-free” Co(Ni)—Mo—Al—Nb γ-γ′ superalloys at elevated temperatures, J. Mater. Sci. 51 (2016) 7843-7860. |
| S.K. Makineni, A. Samanta, T. Rojhirunsakool, T. Alam, B. Nithin, A.K. Singh, R. Banerjee, K. Chattopadhyay, A new class of high strength high temperature Cobalt based γ-γ′ Co—Mo—Al alloys stabilized with Ta addition, Acta Mater. 97 (2015) 29-40. |
| S.K. Makineni, B. Nithin, K. Chattopadhyay, Synthesis of a new tungsten-free γ-γ′ Cobalt-based superalloy by tuning alloying additions, Acta Mater. 85 (2015) 85-94. |
| C.H. Zenk, I. Povstugar, R. Li, F. Rinaldi, S. Neumeier, D. Raabe, M. Göken, A novel type of Co—Ti—Cr-base γ/γ′ superalloys with low mass density, Acta Mater. 135 (2017) 244-251. |
| J.J. Ruan, C.P. Wang, C.C. Zhao, S.Y. Yang, T. Yang, X.J. Liu, Experimental investigation of phase equilibria and microstructure in the Co—Ti—V ternary system, Intermetallics. 49 (2014) 121-131. |
| S. Kirklin, J.E. Saal, V.I. Hegde, C. Wolverton, High-throughput computational search for strengthening precipitates in alloys, Acta Mater. 102 (2016) 125-135. |
| C. Nyshadham, C. Oses, J.E. Hansen, I. Takeuchi, S. Curtarolo, G.L.W. Hart, A computational high-throughput search for new ternary superalloys, Acta Mater. 122 (2017) 438-447. |
| B. Nithin, A. Samanta, S.K. Makineni, T. Alam, P. Pandey, A.K. Singh, R. Banerjee, K. Chattopadhyay, Effect of Cr addition on γ-γ′ cobalt-based Co—Mo—Al—Ta class of superalloys: a combined experimental and computational study, J. Mater. Sci. 52 (2017) 11036-11047. |
| J.J. Ruan, C.P. Wang, S.Y. Yang, T. Omori, T. Yang, Y. Kimura, X.J. Liu, R. Kainuma, K. Ishida, Experimental Investigations of microstructures and phase equilibria in the Co—V—Ta ternary system, J. Alloys Compd. 664 (2016) 141-148. |
| C.P. Wang, S. Yang, S.Y. Yang, D. Wang, J.J. Ruan, J. Li, X.J. Liu, Experimental Investigation of the Phase Equilibria in the Co—Nb—V Ternary System, J. Phase Equilibria Diffus. 36 (2015) 592-598. |
| K. Shinagawa, hours. Chinen, T. Omori, K. Oikawa, I. Ohnuma, K. Ishida, R. Kainuma, Phase equilibria and thermodynamic calculation of the Co—Ta binary system, Intermetallics. 49 (2014) 87-97. |
| P.A. Carvalho, P.M. Bronsveld, B.J. Kooi, J.T.M. De Hosson, On the fcc → D0 19 transformation in Co—W alloys, Acta Mater. 50 (2002) 4511-4526. |
| D.J. Sauza, P.J. Bocchini, D.C. Dunand, D.N. Seidman, Influence of ruthenium on microstructural evolution in a model [Formula presented] superalloy, Acta Mater. 117 (2016) 135-145. |
| S. Meher, L.J. Carroll, T.M. Pollock, M.C. Carroll, Solute partitioning in multi-component γ/γ′ Co—Ni-base superalloys with near-zero lattice misfit, Scr. Mater. 113 (2016) 185-189. |
| L.. Nagel, B. Fultz, Phase Equilibria of Co3V, J. Phase Equilibria. 18 (1997) 21-23. |
| E.A. Lass, M.E. Williams, C.E. Campbell, K.-W. Moon, U.R. Kattner, γ′ Phase Stability and Phase Equilibrium in Ternary Co—Al—W at 900° C., J. Phase Equilibria Diffus. 35 (2014) 711-723. |
| S. Lee, K. Lee, T. Chuang, Discontinuous coarsening of discontinuous precipitates in a Co-6 at.% Mo alloy, Mater. Sci. Eng. A. 251 (1998) 135-141. |
| V. Ramaswamy, E.P. Butler, P.R. Swann, Direct observation of discontinuous precipitation in Al-28at%Zn, J. Microsc. 97 (1973) 259-268. |
| E.P. Butler, V. Ramaswamy, P.R. Swann, In Situ Observation of Cellular Precipitation in an Al-28 at % Zn Alloy by High Voltage Electron Microscopy, Acta Metall. 21 (1973) 517-524. |
| K.N. Braszczyńska-Malik, Discontinuous and continuous precipitation in magnesium-aluminium type alloys, J. Alloys Compd. 477 (2009) 870-876. |
| A. Bauer, S. Neumeier, F. Pyczak, M. Goken, Creep Strength and Microstructure of Polycrystalline Gamma Prime-Strengthened Cobalt-Base Superalloys, Superalloys 2012. (2012) 695-703. |
| H.Y. Yan, V.A. Vorontsov, D. Dye, Alloying effects in polycrystalline γ′ strengthened Co—Al—W base alloys, Intermetallics. 48 (2014) 44-53. |
| J.D. Nystrom, T.M. Pollock, W.H. Murphy, A. Garg, Discontinuous cellular precipitation in a high-refractory nickel-base superalloy, Metall. Mater. Trans. A. 28 (1997) 2443-2452. |
| D. Turnbull, Theory of cellular precipitation, Acta Metall. 3 (1955) 55-63. doi:10.1016/0001-6160(55)90012-2. |
| F. Stein, D. Jiang, M. Palm, G. Sauthoff, D. Grüner, G. Kreiner, Experimental reinvestigation of the Co—Nb phase diagram, Intermetallics. 16 (2008) 785-792. |
| S. Kobayashi, Y. Tsukamoto, T. Takasugi, hours. Chinen, T. Omori, K. Ishida, S. Zaefferer, Determination of phase equilibria in the Co-rich Co—Al—W ternary system with a diffusion-couple technique, Intermetallics. 17 (2009) 1085-1089. |
| W. Huang, Y.A. Chang, A thermodynamic analysis of the NiAl system, Intermetallics. 6 (1998) 487-498. |
| P.J. Bocchini, C.K. Sudbrack, R.D. Noebe, D.N. Seidman, Temporal Evolution of a Model Co—Al—W Superalloy Aged at 650° C. and 750° C., Acta Mater. 159 (2018) 197-208. |
| T.M. Pollock, S. Tin, Nickel-Based Superalloys for Advanced Turbine Engines: Chemistry, Microstructure and Properties, J. Propuls. Power. 22 (2006) 361-374. |
| K.E. Yoon, R.D. Noebe, D.N. Seidman, Effects of rhenium addition on the temporal evolution of the nanostructure and chemistry of a model Ni—Cr—Al superalloy . I: Experimental observations, Acta Mater. 55 (2007) 1145-1157. |
| P. Pandey, S.K. Makineni, A. Samanta, A. Sharma, S.M. Das, B. Nithin, C. Srivastava, A.K. Singh, D. Raabe, B. Gault, K. Chattopadhyay, Elemental site occupancy in the L12 A3B ordered intermetallic phase in Co-based superalloys and its influence on the microstructure, Acta Mater. (2018). |
| A. Tomaszewska, T. Mikuszewski, G. Moskal, D. Migas, Primary microstructure, microsegregation and precipitates characterization of an as-cast new type γ-γ′ Co—Al—Mo—Nb cobalt-based superalloy, J. Alloys Compd. 750 (2018) 741-749. |
| E.A. Lass, D.J. Sauza, D.C. Dunand, D.N. Seidman, Multicomponent γ′-strengthened Co-based superalloys with increased solvus temperatures and reduced mass densities, Acta Mater. 147 (2018) 284-295. |
| H.J. Im, S.K. Makineni, B. Gault, F. Stein, D. Raabe, P.P. Choi, Elemental partitioning and site-occupancy in γ/γ′ forming Co—Ti—Mo and Co—Ti—Cr alloys, Scr. Mater. 154 (2018) 159-162. |
| KIPO (ISR/KR), “International Search Report for PCT/US2019/033450”, Korea, Aug. 23, 2019. |
| X.P. Tan, J.L. Liu, T. Jin, X.F. Sun, Z.Q. Hu, Influence of Cr addition on microstructure of a 5% Re-containing single crystal nickel-based superalloy, Trans. Nonferrous Met. Soc. China (English Ed. 21 (2011) 1004-1008. |
| N.S. Patel, V. Pavlík, M. Boča, High-Temperature Corrosion Behavior of Superalloys in Molten Salts—A Review, Crit. Rev. Solid State Mater. Sci. 42 (2017) 83-97. |
| W.W. Xu, S.L. Shang, C.P. Wang, T.Q. Gang, Y.F. Huang, L.J. Chen, X.J. Liu, Z.K. Liu, Accelerating exploitation of Co—Al-based superalloys from theoretical study, Mater. Des. 142 (2018) 139-148. |
| W.Y. Wang, F. Xue, Y. Zhang, S.L. Shang, Y. Wang, K.A. Darling, L.J. Kecskes, J. Li, X. Hui, Q. Feng, Z.K. Liu, Atomic and electronic basis for solutes strengthened (010) anti-phase boundary of L12Co3(Al, Tm): A comprehensive first-principles study, Acta Mater. 145 (2018) 30-40. |
| F.L. Reyes Tirado, J. Perrin Toinin, D.C. Dunand, γ+γ′ microstructures in the Co—Ta—V and Co—Nb—V ternary systems, Acta Mater. 151 (2018) 137-148. |
| L. Wang, M. Oehring, Y. Liu, U. Lorenz, F. Pyczak, Site occupancy of alloying elements in the L12 structure determined by channeling enhanced microanalysis in γ/γ′ Co—9Al—9W—2X alloys, Acta Mater. 162 (2019) 176-188. |
| I. Povstugar, P.-P. Choi, S. Neumeier, A. Bauer, C.H. Zenk, M. Göken, D. Raabe, Elemental partitioning and mechanical properties of Ti- and Ta-containing Co—Al—W-base superalloys studied by atom probe tomography and nanoindentation, Acta Mater. 78 (2014) 78-85. |
| F. Xue, hours.J. Zhou, X.F. Ding, M.L. Wang, Q. Feng, Improved high temperature γ′ stability of Co—Al—W-base alloys containing Ti and Ta, Mater. Lett. 112 (2013) 215-218. |
| C.H. Zenk, S. Neumeier, hours.J. Stone, M. Göken, Mechanical properties and lattice misfit of γ/γ′ strengthened Co-base superalloys in the Co—W—Al—Ti quaternary system, Intermetallics. 55 (2014) 28-39. doi:10.1016/J.INTERMET.2014.07.006. |
| S. Meher, hours.-Y. Yan, S. Nag, D. Dye, R. Banerjee, Solute partitioning and site preference in γ/γ′ cobalt-base alloys, Scr. Mater. 67 (2012) 850-853. |
| M. Kolb, L.P. Freund, F. Fischer, I. Povstugar, S.K. Makineni, B. Gault, D. Raabe, J. Müller, E. Spiecker, S. Neumeier, M. Göken, On the grain boundary strengthening effect of boron in γ/γ′ Cobalt-base superalloys, Acta Mater. 145 (2018) 247-254. |
| M. Jin, N. Miao, W. Zhao, J. Zhou, Q. Du, Z. Sun, Structural stability and mechanical properties of Co3(Al, M) (M=Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, W) compounds, Comput. Mater. Sci. 148 (2018) 27-37. doi:10.1016/J.COMMATSCI.2018.02.015. |
| I. Povstugar, C.H. Zenk, R. Li, P.-P. Choi, S. Neumeier, O. Dolotko, M. Hoelzel, M. Göken, D. Raabe, Elemental partitioning, lattice misfit and creep behaviour of Cr containing γ ′ strengthened Co base superalloys, Mater. Sci. Technol. 32 (2016) 220-225. |
| M. Kolb, C.H. Zenk, A. Kirzinger, I. Povstugar, D. Raabe, S. Neumeier, M. Göken, Influence of rhenium on c9-strengthened cobalt-base superalloys, (2018). |
| Q. Liu, J. Coakley, D.N. Seidman, D.C. Dunand, Precipitate Evolution and Creep Behavior of a W-Free Co-based Superalloy, Metall. Mater. Trans. A. 47 (2016) 6090-6096. |
| Y. Li, F. Pyczak, J. Paul, M. Oehring, U. Lorenz, Z. Yao, Y. Ning, Rafting of γ′ precipitates in a Co—9Al—9W superalloy during compressive creep, Mater. Sci. Eng. A. 719 (2018) 43-48. |
| F. Pyczak, A. Bauer, M. Göken, S. Neumeier, U. Lorenz, M. Oehring, N. Schell, A. Schreyer, A. Stark, F. Symanzik, Plastic deformation mechanisms in a crept L12 hardened Co-base superalloy, Mater. Sci. Eng. A. 571 (2013) 13-18. |
| Y. Li, F. Pyczak, J. Paul, M. Oehring, U. Lorenz, Z. Yao, Microstructure evolution in L12 hardened Co-base superalloys during creep, J. Mater. Res. 32 (2017) 4522-4530. |
| T. Murakumo, T. Kobayashi, Y. Koizumi, hours. Harada, Creep behaviour of Ni-base single-crystal superalloys with various γ′ volume fraction, Acta Mater. 52 (2004) 3737-3744. |
| Number | Date | Country | |
|---|---|---|---|
| 20210207255 A1 | Jul 2021 | US |
| Number | Date | Country | |
|---|---|---|---|
| 62674780 | May 2018 | US |
