NiCrMoNb age hardenable alloy for creep-resistant high temperature applications, and methods of making

Abstract

Nickel alloys, methods of making nickel alloys, articles including the nickel alloys, uses of the alloys, and methods of treating nickel alloys are described. The inventive heat resistant structural materials are suitable for applications requiring high yield stress at room temperature and good creep strength at high temperatures, such as in gas turbines, steam turbines, fossil energy boilers, aero engines, power generation systems using fluids such as supercritical carbon dioxide (e.g., advanced ultra-supercritical power plants), concentrated solar power plants, nuclear power plants, molten salt reactors: turbine blades, casings, valves, heat exchangers and recuperators.

Claims

1. A nickel alloy, comprising: 19.0 to 22.5% Cr; 0.1 to 1.0% Al; 1.0 to 2.0% Ti; 3.0 to 5.0% Nb; 0.1 to 4.0% Ta; 4.0 to 7.0% Nb+Ta; 2.0 to 4.0% Fe; 0 to 0.06% C; 0 to 0.35% Mn; 7.0 to 9.5% Mo; 0 to 5.0% W; 0 to 0.20% Si; 0 to 0.05% V; 0 to 0.15% P; 0 to 0.010% S; and 0 to 0.010% B.

2. The alloy of claim 1 wherein the balance of the alloy comprises Ni and no more than 1% of all other elements.

3. The alloy of any of the preceding claims comprising at least 0.1% Ta, or at least 0.2% Ta, or at least 1.0% Ta, or 0.1 to 4.0% Ta; or none or consisting essentially of none of Co.

4. The alloy of any of the above claims comprising at least 3.5% Nb, or at least 4.0% Nb.

5. (canceled)

6. The alloy of any of the above claims comprising δ (and/or η) phase precipitates at the grain boundaries and, optionally, further comprising γ′ (or γ″ or γ′/γ″) precipitates distributed within a γ matrix grain interior.

7. The alloy of any of the above claims characterizable by a time to failure at 700° C. and 483 MPa of at least 230 h, or at least 270 h, or at least 300 h, or at least 400 h, or in the range of 230 to 750 h or in the range of 230 to 700 h.

8. The alloy of any of the above claims characterizable by a time to failure at 790° C. and 207 MPa of at least 320 h, or at least 400 h, or at least 500 h, or in the range of 230 to 650 h or in the range of 230 to 550 h.

9. The alloy of any of the above claims characterizable by a time to failure at 750° C. and 345 MPa of at least 180 h, or at least 200 h, or at least 250 h, or at least 300 h, or in the range of 180 to 500 h or in the range of 180 to 450 h.

10-12. (canceled)

13. The alloy of any of the above claims characterizable by ultimate tensile strength at 750° C. of at least 800 MPa, or at least 850 MPa, or at least 900 MPa, or in the range of 800 to 1050 MPa or 800 to 950 MPa.

14. The alloy of any of the above claims characterizable by ultimate tensile strength at 800° C. of at least 600 MPa, or at least 650 MPa, or at least 700 MPa, or in the range of 650 to 900 MPa or 650 to 800 MPa.

15. The alloy of any of the above claims characterizable by 0.2% yield stress at room temperature of at least 850 MPa, or at least 900 MPa, or at least 1000 MPa, or in the range of 800 to 1150 MPa or 800 to 1100 MPa.

16. The alloy of any of the above claims characterizable by 0.2% yield stress at 600° C. of at least 750 MPa, or at least 850 MPa, or at least 900 MPa, or in the range of 750 to 1050 MPa or 750 to 1000 MPa.

17. The alloy of any of the above claims characterizable by 0.2% yield stress at 700° C. of at least 700 MPa, or at least 800 MPa, or at least 850 MPa, or in the range of 700 to 1000 MPa or 700 to 950 MPa.

18. The alloy of any of the above claims characterizable by 0.2% yield stress at 750° C. of at least 700 MPa, or at least 750 MPa, or at least 800 MPa, or in the range of 700 to 850 MPa or 700 to 800 MPa.

19. The alloy of any of the above claims characterizable by 0.2% yield stress at 800° C. of at least 500 MPa, or at least 550 MPa, or at least 600 MPa, or in the range of 500 to 800 MPa or 500 to 700 MPa.

20-26. (canceled)

27. A method of heat treating a nickel alloy comprising Cr and Nb, comprising the sequential steps of treating the nickel alloy: 820±30° C. for 5 to 100 h followed by cooling to 750±15° C. at a rate of 1 to 5° C./min, and then holding at 750±15° C. for 2 to 15 h.

28. The method of claim 27 wherein the nickel alloy comprises the alloy of any of claims 1-21.

29-32. (canceled)

33. A solid nickel alloy comprising Cr; Al; Ti; Nb; Ta; Mo, comprising a chemical inhomogeneity of Cr; Al; Ti; Nb; Ta; and Mo that is less than 10%; and wherein chemical inhomogeneity is characterizable by a measurement using wavelength dispersive x-ray fluorescence on a polished cross-section the alloy and wherein any selected 10 μm diameter area within a 10 mm.sup.2 area of the polished surface has an elemental composition of each of Cr; Al; Ti; Nb; Ta; Mo that is within 10% of each element in the 10 mm.sup.2 area.

34-35. (canceled)

36. The alloy of any of claims 33-35 wherein the nickel alloy comprises the alloy of any of claims 1-21.

37. An article comprising the alloy of the any of the above claims.

38. (canceled)

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0020] FIG. 1 shows δ phase precipitates located at the grain boundaries of alloy F.

[0021] FIG. 2 shows elongation to failure as a function of creep life at 700° C. and 483 MPa for alloy STD, D, E, F with either SHT or HT aging heat treatment and alloys E1 and F2 with the HT aging heat treatment.

[0022] FIG. 3 shows elongation to failure as a function of creep life at 790° C. and 207 MPa for alloy STD with the SHT and HT aging heat treatment and alloys D, E, F, E1 and F2 with the HT aging heat treatment.

[0023] FIG. 4 shows elongation to failure as a function of creep life at 750° C. and 345 MPa for alloy STD with the SHT and HT aging heat treatment and alloys E1 and F2 with the HT aging heat treatment.

[0024] FIG. 5 shows tensile properties (UTS, 0.2% yield stress and elongation to failure) for Alloy STD, D, E, F, E1 and F2 with either standard (SHT) or HT aging heat treatment tested at room temperature (left) and 750° C. (right).

[0025] FIG. 6 shows tensile properties (UTS, 0.2% yield stress and elongation to failure) for Alloys STD and D with standard (SHT) aging heat treatment and Alloys E1 and F2 with HT aging tested over a range of temperatures from RT to 800° C.

[0026] FIG. 7 shows microstructure of (a) Alloy STD after the SHT, (b) alloy E after the first step of the HT aging treatment and (c) alloy F after the first step of the HT aging treatment.

[0027] FIG. 8 shows microstructure of (a) alloy F (b) E1 and (c) F2 after □ 100 h at 800° C. Note: the phases labelled were tentatively identified using energy dispersive spectroscopy and identification was not confirmed using more advanced techniques.

[0028] FIG. 9 shows results from a phase stability study showing alloys E and E1 exposed to 700° C. and 760° C. for 500 h, 1000 h and 2000 h.

[0029] FIG. 10 shows results from a phase stability study showing alloys F and F2 exposed to 700° C. and 760° C. for 500 h, 1000 h and 2000 h.

[0030] FIG. 11 shows oxidation testing results at 700° C. in air for Alloys STD, E1 and F2.

DETAILED DESCRIPTION OF THE INVENTION

[0031] In one aspect, the invention comprises a cast/wrought alloy of composition within the range of major elements listed in Table 2. (Note: In Table 1A and 1B, chemistry does not specify Ta additions nor does it specify a Nb+Ta guideline. Also, the balance of Fe in the NETL patent is set at between 2 and 4 wt. %, while in both patent examples Fe forms the balance, which may or may not fall within this range. Alternatively, Ni forms the balance in the NETL patent while it is fixed in each of the others. In addition, in the NETL patent a fixed level of B is specified and other minor differences in trace elements exist.)

[0032] Additionally, and more importantly, the alloy is subjected to (1) NETL calculated homogenization cycle and (2) high temperature (HT) post-thermomechanical processing (TMP) aging heat treatment that enables the intentional precipitation of secondary phases (i.e., δ and/or η) at the grain boundaries to increase their resistance to deformation and damage and/or γ′ and/or γ″ precipitates in the grain interior to facilitate high room and high temperature yield stress and tensile strength.

TABLE-US-00003 TABLE 2 Composition range for the alloy, Ni and impurities balanced (wt. %). Element Cr Al Ti Nb Ta Nb + Ta Fe C Mn Mo W Si V P S B Ti/Al Min. 19.0 0.1 1.0 3.0 0.1 4.0 2.0 — — 7.0 — — — — — — 1.0 Max. 22.5 1.0 2.0 5.0 4.0 7.0 4.0 0.06 0.35 9.5 5.0 0.20 0.05 0.015 0.010 0.010 20.0
Several alloys were made following the details of the invention, as well as alloys outside the range of Table 2 (and closer to Tables 1A and 1B) for comparison and using melt processing techniques of an industrial nature. The compositions were designed from computational thermodynamic modeling performed using the software Thermo-Calc for phase diagrams and with the knowledge that Nb and Ta are elements that promote the formation of δ (Ni.sub.3Nb), η (Ni.sub.6AlNb) and γ″ (Ni.sub.3(Nb,Ta)).

[0033] Nine ingots of approximately 6.8 kg were made at NETL, namely Alloy STD, A, B, C, D, E, F, E1 and F2, with compositions listed in Table 3. The compositions of E, F, E1 and F2 are within the range of the invention (Table 2), while Alloy STD is the closest in chemistry to standard alloy 725. Alloys A, B, C, D were designed to investigate the effect of the Ti/Al ratio on tension and creep properties. The alloy compositions were composed of high-purity industry-grade melt materials with, or without, master alloys to incorporate Cr to the melt.

TABLE-US-00004 TABLE 3 Composition of the various heats from XRF analysis for the main elements, combustion analysis for C and calculated from the melt addition for B (wt. %) and calculated Ti/Al ratio. Alloy Cr Al Ti Nb Ta Fe C Mn Mo Si B Ni Ti/Al STD 21.19 0.24 1.82 3.44 0.16 3.87 0.039 0.04 7.08 0.018 0.003 Bal. 7.6 A 21.20 0.19 1.84 3.41 0.37 3.88 0.040 0.05 7.06 <0.010 0.003 Bal. 9.7 B 21.22 0.60 1.41 3.43 0.36 3.90 0.039 0.04 7.09 0.020 0.003 Bal. 2.4 C 21.23 0.68 1.29 3.44 0.39 3.88 0.040 0.05 7.09 0.014 0.003 Bal. 1.9 D 21.24 0.85 1.10 3.42 0.36 3.92 0.044 0.06 7.09 0.033 0.003 Bal. 1.3 E 21.15 0.89 1.15 3.38 3.46 3.88 0.038 0.04 7.02 0.021 0.003 Bal. 1.3 F 21.29 0.84 1.13 4.51 0.37 3.92 0.040 0.04 7.09 <0.010 0.003 Bal. 1.3 E1 20.51 0.20 1.87 3.35 3.55 3.83 0.049 0.05 7.03 <0.010 0.003 Bal. 9.4 F2 20.59 0.19 1.88 4.55 0.38 3.86 0.049 0.05 7.09 <0.010 0.003 Bal. 9.9

[0034] The master alloys were used in Alloy STD to F and consisted of Ni-25Cr, Fe-24Cr and/or Ni-50Cr that were melted using vacuum induction melting (VIM) followed by electroslag remelting (ESR) to lower the concentration of tramp elements. The alloys were vacuum induction melted under Ar partial pressure to about 50° C. above the liquidus temperature predicted using Thermo-Calc. The liquid was then poured in a graphite mold with ceramic wash coat to prevent C pickup and solidified as cylindrical ingots. The ingots were then homogenized using the NETL approach, which assesses the chemistry and secondary dendrite arm spacing (SDAS) in a computational manner so as to minimize heterogeneity on a nanometer scale (see Jablonski and Hawk, J Mat Eng and Perf, V26, N1 pp 4-13 (2017)). The homogenization treatment schedule used is based on the chemistry and scale of the microstructure (SDAS) which is established by the section size of the cast ingot (Table 4). This homogenization heat treatment schedule incorporates thermodynamic factors as well as kinetic ones (i.e., time for specific critical elements to diffuse within the γ matrix) to complement microstructural considerations (i.e., spacing of dendrite arms in the solidified microstructure as well as the interdendritic spacing between those arms), and the sequence described encompasses homogeneity within the alloy from at least 1% to 10% of nominal for each element calculated using computational thermodynamic tools. The stepped approach to raising temperature during homogenization is done to avoid localized melting. After a time at lower temperature the chemistry has sufficiently homogenized to allow the raise in temperature incrementally which improves the time efficiency of homogenization while avoiding localized melting.

TABLE-US-00005 TABLE 4 Heat treatments required to obtain a certain residual inhomogeneity. Section Residual Inhomogeneity Size <10% <5% <1% Up to 5 in 1030° C./1 h + 1065° C./3 h + 1030° C./1 h + 1065° C./3 h + 1030° C./1 h + 1065° C./3 h + 1090° C./3 h + 1115° C./3 h + 1090° C./3 h + 1115° C./3 h + 1090° C./3 h + 1115° C./3 h + 1135° C./6 h + 1150° C./30 h 1135° C./6 h + 1150° C./48 h 1135° C./6 h + 1150° C. /72 h 5-8 in 1030° C./2 h +1065° C./6 h + 1030° C./2 h + 1065° C./6 h + 1030° C./2 h + 1065° C./6 h + 1090° C./6 h + 1115° C./8 h + 1090° C./6 h + 1115° C./8 h + 1090° C./6 h + 1115° C./8 h + 1135° C./12 h + 1150° C./50 h 1135° C./12 h + 1150° C./100 h 1135° C./12 h + 1150° C./200 h  >8 in 1030° C./4 h + 1065° C./12 h + 1030° C./4 h + 1065° C./12 h + 1030° C./4 h + 1065° C./12 h + 1090° C./12 h + 1115° C./12 h + 1090° C./12 h + 1115° C./12 h + 1090° C./12 h + 1115° C./12 h + 1135° C./24 h + 1150° C./100 h 1135° C./24 h + 1150° C./100 h 1135° C./24 h + 1150° C./100 h

[0035] The ingots were then hot worked at 1105±25° C. using steps of forging followed by hot rolling with reheating between each step to form 10 mm thick plates. The last reheat thermal cycle (following the last hot rolling step) was used as a solution heat treatment which lasted 400±100 seconds.

[0036] The TMP plates were then aged following the designed high temperature (HT) post-TMP aging heat treatment, which consisted of the following:

High Temperature (HT) Post-TMP Aging Heat Treatment:

[0037] (820±30° C. for 5 to 100 h) followed by cooling to (750±15° C. at 1 to 5° C./min) and then holding at (750±15° C. for 2 to 15 h).
The first thermal cycle in the heat treatment is used to allow sufficient time for the precipitation of specific grain boundary phases. The second step allows for precipitation (or additional precipitation) and growth of the γ′ and/or γ″ precipitates without disturbing the grain boundary phases. The cooling rate between each step was selected from a furnace programming standpoint. This cooling rate can be adjusted or simple furnace cooling can be employed. Argon gas fan cooling is used at the end of the heat treatment cycle to rapidly cool the plates to room temperature.

[0038] Mechanical testing was performed in tension between room temperature and 750° C. using the ASTM standard E-8 and in creep at 700° C./483 MPa and 790° C./207 MPa with additional conditions for some of the alloys using the ASTM standard E-139. The heat treatment conditions for the specimens tested are as follows: 1105° C. for 300 to 360 seconds for solution; 1030° C. for 1 h+1065° C. for 3 h+1090° C. for 3 h+1115° C. for 3 h+1135° C. for 6 h+1150° C. for 72 h for homogenization; and 800° C. for 20 h followed by cooling to 750° C. at 1 to 5° C./min and holding at 750° C. for 8 h.

[0039] Observation of the microstructure of the alloys reveals the formation of various precipitate phases varying as a function of the alloy composition. For example, δ (and/or η) phase precipitates along the grain boundaries following the HT aging treatment are shown in alloy F and represented in FIG. 1 (identified using energy dispersive spectroscopy). This is a desirable microstructural feature, that is, the formation of designed small precipitates (i.e., δ and/or η) along the boundaries that remain stable under loading conditions particularly at high temperatures.

[0040] Alloy STD, closest in composition to standard alloy 725, was evaluated with the standard aging heat treatment as specified by the manufacturer to compare to the mechanical properties of the invention (alloys with compositions in the range of Table 2 with the HT aging treatment). The standard aging heat treatment specified in the Special Metals datasheet for alloy 725 consists of the following:

Standard heat treatment (SHT): (732° C. for 8 h) followed by cooling to (621° C. at 56° C./h) and holding at (621° C. for 8 h)
Specimens from Alloy STD, A, B, C, D, E and F were tested after SHT while other specimens from Alloy STD, D, E, F, E1 and F2 were tested after HT aging treatment. Specimens E, F, E1 and F2 contain higher levels of Nb and Ta (within Table 2) and were optimized for the HT aging treatment.

[0041] The creep properties of the alloys are listed in Table 5 and some of the results are discussed below using figures to facilitate the observation of various trends.

Note: Alloys A, B and C will not be discussed as they were part of a previous investigation where the Ti/Al ratio was investigated. See (https://link.springer.com/chapter/10.1007%2F978-3-319-89480-5 26).
From Table 3, the alloys of invention have either a low Ti/Al ratio (E, F, based on D) or high Ti/Al (E1, F2, similar to the standard 725-like Alloy STD) which was employed to favor the precipitation of either γ′ or γ″ within the grains and potentially either δ or η along the grain boundaries since increasing Al promotes formation of γ′ and Ti promotes formation of η. Phase stability is discussed at the end of this section.

TABLE-US-00006 TABLE 5 Creep properties of the various alloys Temperature Stress Time to LMP Elongation to Alloy Aging (° C.) (MPa) Failure (h) (C = 20)/1000 Failure (%) STD Standard (SHT) 700 483 194 21.686 9.0 750 345 173 22.750 7.7 790 207 316 23.917 7.8 HT 700 483 271 21.827 13.8 750 345 154 22.698 10.5 790 207 288 23.874 10.0 A Standard (SHT) 700 483 217 21.733 8.9 B Standard (SHT) 700 483 153 21.586 5.5 C Standard (SHT) 700 483 122 21.490 4.3 D Standard (SHT) 700 483 80 21.312 1.7 HT 700 483 231 21.760 5.7 790 207 355 23.971 14.2 E Standard (SHT) 700 483 104 21.423 0.6 HT 700 483 490 22.078 3.9 790 207 462 24.093 12.3 F Standard (SHT) 700 483 171 21.633 1.4 HT 700 483 432 22.024 5.1 790 207 393 24.018 9.5 El HT 700 483 690 22.222 7.1 725 414 457 22.615 7.6 750 345 414 23.137 8.0 790 207 548 24.171 8.6 F2 HT 700 483 409 22.001 6.7 725 414 424 22.582 7.3 750 345 310 23.009 9.4 790 207 384 24.007 9.8
The creep results at 700° C. and 483 MPa are compared for all alloys that were either aged using the SHT or the HT heat treatment in FIG. 2. Note: alloys E1 and F2 were only tested in the HT condition. From FIG. 2, the HT aging treatment resulted in dramatic increases in both the creep life and creep ductility of the alloys. Using the HT aging treatment resulted in the following improvements compared to the standard aging treatment: [0042] The life of Alloy STD increased by 40% and its ductility by 53%. [0043] The life of alloy D increased by 189% and its ductility by 235%. [0044] The life of alloy E increased by 371% and its ductility by 550%. [0045] The life of alloy F increased by 153% and its ductility by 264%.
Comparing to Alloy STD in its standard aging heat treatment (SHT, or condition most representative of standard alloy 725/625 PLUS): [0046] The HT aging heat treatment resulted in an improvement of creep life by 40% and ductility by 53%. [0047] The HT aging treatment combined with Nb addition resulted in an improvement of creep life of >110%. [0048] The HT aging treatment combined with Nb addition and low Ti/Al ratio resulted in an improvement of creep life of >120%. [0049] The HT aging treatment combined with Ta addition resulted in an improvement of creep life of >250% with slightly lower ductility. [0050] The HT aging treatment combined with Ta addition and low Ti/Al ratio resulted in an improvement of creep life of >130%.
These results underline improvements of the invention, that is, HT aging heat treatment combined with the elevated concentrations of Nb and/or Ta result is superior creep performance measured as creep life.
Note: The low Ti/Al alloys have lower ductility in general (D, E, F) compared to the high Ti/Al ratio alloys STD, E1 and F2.
A second creep condition was investigated for most alloy/heat treatment conditions and consisted of 790° C. and 207 MPa. The results are represented in FIG. 3. Comparison with Alloy STD (with SHT aging) showed that chemistry variations combined with HT aging extended each alloy variant creep life. The creep life increased by 73% in alloy E1 with a slight increase in ductility of 10% while the creep life increased by 46% in alloy E with a 58% increase in creep ductility. The creep life increased by 22% in alloy F2 with a 26% increase in creep ductility while the creep life increased by 24% in alloy F with a 22% increase in creep ductility.

[0051] Another creep condition was considered for alloys E1 and F2 following the HT aging heat treatment with comparison to alloy STD and consisted of 750° C. and 345 MPa. The results are represented in FIG. 4. Comparison with Alloy STD (with SHT aging) showed that chemistry variations combined with HT aging extended each alloy variant creep life. The creep life increased by 139% in alloy E1 with a 4% increase in creep ductility while the creep life increased by 79% in alloy F2 with a 22% increase in creep ductility.

[0052] Results from tension testing are compiled in Table 6. Similar to the description of the creep results, figures are provided below to better illustrate the trends.

TABLE-US-00007 TABLE 6 Tensile properties of the various alloys 0.2% Yield Elongation Temperature UTS Stress to Alloy Aging (° C.) (MPa) (MPa) failure (%) STD Standard (SHT) 24 1,224 809 43.1 200 1,135 752 40.7 400 1,053 730 43.9 500 1,022 723 40.5 600 978 734 37.5 650 1,074 748 24.6 700 882 650 24.7 750 794 622 16.5 HT 24 1,246 812 32.3 750 790 664 38.3 A Standard (SHT) 24 1,282 861 39.6 750 806 646 8.0 B Standard (SHT) 24 1,150 722 48.0 750 710 538 14.0 C Standard (SHT) 24 1,149 724 47.4 750 705 535 14.4 D Standard (SHT) 24 1,186 749 43.0 200 1,092 685 46.5 400 1,007 671 44.4 500 988 670 44.5 600 966 659 39.9 650 930 644 18.3 700 800 606 12.5 750 698 566 8.3 HT 24 1,222 722 35.4 750 822 629 20.2 E Standard (SHT) 24 1,224 838 42.9 750 741 646 7.1 HT 24 1,169 898 4.5 750 943 774 9.5 F Standard (SHT) 24 1,190 795 47.8 750 706 595 7.8 HT 24 1,253 865 15.5 750 899 735 10.5 E1 HT 24 1,412 1,073 23.7 600 1,305 976 28.2 700 1,074 916 18.3 750 883 788 15.2 800 728 626 19.1 F2 HT 24 1,384 1,027 32.1 600 1,278 947 32.9 700 1,025 862 24.9 750 845 740 17.1 800 684 587 22.0
Comparison of the results from tension testing of Alloy STD, D, E, F, E1 and F2 are shown in FIG. 5 for specimens that received either the SHT or the HT aging heat treatment. Note: alloys E1 and F2 were only tested in the HT condition. Results from tests at room temperature show general improvements in ultimate tensile strength (UTS) and 0.2% yield stress (YS) for the specimens aged using the HT aging treatment minus a couple exceptions. The ductility was generally lower but satisfactory as the test condition is in excess of the targeted applications. Particularly, alloys E1 and F2 generally outperformed all others. Compared to Alloy STD with the standard aging heat treatment (SHT), the UTS was increased by 15% and the 0.2% YS by 33% in E1 with HT aging while keeping an elongation to failure above 20%.
At 750° C., results are consistently in favor of the HT aging heat treatment. The results reveal increases in UTS, YS and elongation to failure due to the HT aging heat treatment, particularly, for alloys E, F, E1 and F2 that were optimized by adding Nb and/or Ta. Compared to Alloy STD with the SHT aging treatment (i.e., the condition closest to standard INCONEL 725), the HT aging treatment and modified chemistry resulted in average increases of 12% in UTS and 22% in YS.

[0053] Comparison over the range of tension testing temperatures is provided in FIG. 6 for alloys STD and D with the SHT aging and alloys E1 and F2 with the HT aging. Specifically, the UTS and 0.2% YS increased by 22% and 41% in alloy E1 and 16% and 33% in alloy F2 compared to the standard alloy.

[0054] As previously mentioned, the invention may use a designed HT aging treatment, preferably combined with chemistry adjustments and a homogenization heat treatment schedule to a 1% of nominal in order to promote the formation of precipitate phases (i.e., δ and/or η) uniformly along the grain boundaries and/or complex matrix precipitate (γ″ or γ″ or γ′/γ″) formation uniformly within the grains. In Alloy STD after the SHT aging treatment (FIG. 7a), the microstructure consists of grain boundary carbides (gray and white) and very fine γ′ precipitates. These precipitates are also distributed within the γ matrix grain interior. However, using the high temperature first step (here 800° C. for 20 h) in the HT aging treatment promotes the formation of precipitates such as δ and/or η along the grain boundaries when high levels of Nb or Ta are used in the alloy, as observed in alloy E and F in FIGS. 7b and 7c. The δ and/or η phases reinforce the grain boundaries by adding resistance to operative deformation mechanisms. Their presence has also been reported to diminish, or arrest, crack propagation.

[0055] Furthermore, the first step of the HT aging treatment can be adjusted, either by extending the time at the aging temperature or modifying the aging treatment temperature and holding time. For example, the time at 800° C. was changed to 100 h for alloys F, E1 and F2 to observe its microstructure. In doing so, the precipitation of δ phase (tentatively identified using energy dispersive spectroscopy) was further promoted, see for example, FIG. 8a.

Finally, adjustments to the Ti/Al ratio can be used to control the type of matrix precipitate (γ′ or γ″ or γ′/γ″) strengthening desired. For instance, F (low Ti/Al ratio) contained γ′ precipitates within the matrix grains that greatly contribute to its strength while E1 and F2 showed γ″ matrix precipitates as observed in FIGS. 8b and c.

[0056] An important consideration associated with the use of various precipitate phases along the grain boundaries to improve mechanical properties is related to phase stability. It is preferable for the precipitates to remain stable, i.e., to show limited coarsening or morphology or composition changes with increasing exposure time or temperature. Phase stability can be controlled from a combination of chemistry changes within the range of Table 2 and/or different temperature or time for the first step of the HT aging heat treatment (within the specified range). Results from a phase stability study on alloys E and E1 are shown in FIG. 9 and alloys F and F2 in FIG. 10 for exposure at 700° C. and 760° C. for 500 h, 1000 h and 2000 h. Alloys E1, F and F2 appeared more stable than E as less growth of the grain boundary phases towards the matrix was shown with increasing exposure time at 700° C. The results also revealed that the alloys are unstable at 760° C., thus this temperature is unsuitable for use of the alloys for extended times.

[0057] Oxidation testing was performed on alloys STD, E1 and F2 to assess if the changes in chemistry of the invention to improve high temperature mechanical properties had an effect on the oxidation resistance of the alloy. Results for testing in air at 700° C. are shown in FIG. 11 in the form of specific mass change as a function of exposure time at temperature. Alloy F2 behaved in a similar manner than the standard alloy (STD) while E1 showed lower mass gains overall. Therefore, the compositional changes did not have a negative effect on the oxidation resistance of the alloy under the testing conditions investigated.

[0058] Key Takeaways:

A) The invention provides (1) high-temperature (HT) aging heat treatment, (2) homogenization heat treatment and (3) preferred material's chemistry range (Table 2), the entirety of which produces (4) designed specific microstructures consisting of desirable precipitates along the grain boundaries between matrix grains with either γ′ or γ″ or γ′/γ″ strengthening precipitates in the matrix grains, to enable corrosion resistant Ni-based superalloy in applications that require superior tension properties and creep resistance at temperatures above 650° C.
B) The homogenization heat treatment provides an initial uniform chemistry throughout.
C) The HT aging heat treatment employs a high temperature, long term, first step thermal exposure that enables the precipitation of secondary phases, such as (but not limited to) δ and/or η, along the grain boundaries, which serves to restrict grain boundary sliding and enhance creep performance.
D) The HT heat treatment also allows for a complex precipitation of strengthening phases (γ′ or γ″ or γ′/γ″) within the grains interior that can be combined with the grain boundary precipitates.
E) The chemistry of the alloys can be adjusted within the range of Table 2 to favor the precipitation of γ′ or γ″ or γ′/γ″, particularly using variations in the Al, Ti, Nb and Ta concentrations.
F) In particular, elevated levels of Nb and/or Ta favor the precipitation of δ precipitates along the grain boundaries.
Variations within the scope of the invention include but are not limited to:
Changes in composition: The composition of the alloy can be varied within the range specified in Table 2. A minimum Nb+Ta concentration was set for the formation of necessary grain boundary phases to obtain the improved properties. Ti and Al concentrations encompass Ti/Al ratios that can be tailored to form desired precipitate phases within the grains.
Homogenization heat treatment: The homogenization heat treatment used in the making the invention alloy was optimized using computational design. However, other combinations of time/temperature and number of segments can be utilized, resulting in different levels of homogenization (i.e., ±10, 5 or 1% or better). The homogenization heat treatment was specified in Table 4.
Solution heat treatment: The solution heat treatment used was 1105° C. for 5-7 min. However, this temperature can be adjusted (lower or higher) as well as the holding time to dissolve all secondary phases that may have formed during processing. Thus, the solution heat treatment was expressed as follow: 1105±25° C. for 400±100 sec.
Aging heat treatment: Aging trials revealed significant variations in the size or fraction of precipitate phases such as γ′ or γ″ or γ′/γ″ and carbides and more importantly, the grain boundary phases. The two-step aging can be designed using different temperatures with various holding times, heating and cooling rates. The first step is needed to promote the formation of grain boundary phases, such as δ and/or η, while the second step allows for additional precipitation and growth of the γ′ or γ″ or γ′/γ″ precipitates. The HT aging treatment is summarized below: (820±30° C. for 5 to 100 h) followed by cooling to (750±15° C. at 1 to 5° C./min) and then holding at (750±15° C. for 2 to 15 h).

[0059] The invention comprises heat resistant structural materials for applications requiring high yield stress at room temperature and good creep strength at high temperatures such as articles found in gas turbines, steam turbines, fossil energy boilers, etc., including compressor blades, turbine blades, turbine discs and spacers, turbine vanes, combustion cans, etc. The invention also comprises structural materials for components in aero engines (gas fan turbines) or power generation systems using fluids such as supercritical carbon dioxide (e.g., advanced ultra-supercritical power plants), concentrated solar power plants, nuclear power plants, molten salt reactors: turbine blades, casings, valves, heat exchangers and recuperators, etc. The invention further includes any of the articles mentioned above comprising any of the materials described herein.

[0060] The invention may comprise any of the materials, conditions, properties, or other features mentioned above. Throughout the descriptions, “%” means weight %.