Patent Application
20150368760
Improvements in internal combustion engine efficiency alone have the potential to increase passenger vehicle fuel economy by 25 to 40 percent and commercial vehicle fuel economy by 30 percent with a concomitant reduction in carbon dioxide emissions. Certain higher performance engines need higher temperature-capable valve materials due to increased exhaust gas temperatures, higher exhaust flow rates, higher cylinder pressures, and/or modified valve timings. Target temperatures for experimental engines are currently exceeding current 760° C. with the potential to reach 1000° C.
There is a critical need to develop materials that meet projected operational performance parameters but also are feasible with respect to cost constraints. In particular, new low-cost, valve alloys with improved properties at temperatures from 870 to 1000° C. are required for the next generation, high efficiency automotive and diesel engines.
Ni-based alloys are attractive candidates for improved valve materials. High temperature yield, tensile, and fatigue strengths have been identified as critical properties in determining the performance of these alloys in the valve application. In general, conventional Ni-based alloys are strengthened through a combination of solid solution strengthening and precipitation strengthening mechanisms with the latter needed to achieve higher strengths at higher temperatures. In one class of Ni-based superalloys, primary strengthening is obtained through the homogeneous precipitation of ordered, L12 structured, Ni3(X)-based intermetallic precipitates (where X can include Al, Ti, Nb, Ta or any combination of the foregoing) that are coherently embedded in a solid solution face centered cubic (FCC) matrix. In another class of Ni-based alloys, creep resistance is achieved through the precipitation of fine carbides (M23C6, M7C3, M6C where M is primarily Cr with substitution of Mo, W, for example) and carbonitrides (M(C, N) where M can include Nb, Ti, Hf, Ta or any combination of the foregoing for example) within the matrix, and larger carbides on grain boundaries to prevent grain boundary sliding. Moreover, high temperature oxidation resistance in these alloys is obtained through additions of Cr and Al. In other alloys, a combination of both types of precipitates may be used for optimum properties.
An evaluation of the microstructure of various Ni-based alloys and correlation with limited information on the fatigue properties that are available show that the amount (in terms of volume percent or weight percent) of the γ′ phase is likely to be a dominant factor in determining the performance of these alloys at high temperatures. Since the size of the strengthening precipitates is also critical, it is anticipated that the kinetics of coarsening this phase would also be influential in the long-term performance of the alloys in this application.
Several example commercial Ni-based alloy compositions are shown in Table 1. To obtain initial information on the microstructures of these alloys at equilibrium, thermodynamic calculations were carried out using JMatPro V6.1. Comparison of the results of the calculations showed that all alloys have a matrix of γ with the major strengthening phase as γ′. One or more carbide phases such as M23C6, MC, and M7C3 may also be present in different alloys. The primary difference between the microstructures of the various alloys is in the weight percent of the γ′ phase at a given temperature and the highest temperature at which the γ′ phase is stable in the different alloys.
Specific reference is made to U.S. Pat. No. 5,660,938, issued to Katsuaki Sato, et al. on Aug. 26, 1997 and entitled “Fe—Ni—Cr-Base Superalloy, Engine Valve and Knitted Mesh Supporter for Exhaust Gas Catalyzer.” An FE—Ni—Cr-base superalloy consists essentially of, by weight, up to 0.15% C, up to 1.0% Si, up to 3.0% Mn, 30 to 49% Ni, 10 to 18% Cr, 1.6 to 3.0% Al, one or more elements selected from Groups IVa and Va whose amount or total amount is 1.5 to 8.0%, the balance being Fe, optionally, minor amounts of other intentionally added elements, and unavoidable impurities. The optional other elements which can be intentionally added to or omitted from the alloy include Mo, W, Co, B, Mg, Ca, Re, Y and REM. The superalloy is suitable for forming engine valves, knitted mesh supporters for exhaust gas catalyzers and the like, and has excellent high-temperature strength and normal-temperature ductility after long-time heating, as well as sufficient oxidation resistance properties for these uses. The composition is required to satisfy the following Formulae (1) and (2) by atomic percent:
6.5≦Al+Ti+Zr+Hf+V+Nb+Ta≦10 (1)
0.45≦Al÷(Al+Ti+Zr+Hf+V+Nb+Ta)≦0.75 (2)
Specific reference is made to U.S. Pat. No. 6,372,181, issued to Michael G. Fahrmann, et al. on Apr. 16, 2002 and entitled “Low cost, Corrosion and Heat Resistant Alloy for Diesel Engine Valves.” A low cost, highly heat and corrosion resistant alloy useful for the manufacture of diesel engine components, particularly exhaust valves, comprises in % by weight about 0.15-0.65% C, 40-49% Ni, 18-22% Cr, 1.2-1.8% Al, 2-3% Ti, 0.9-7.8% Nb, not more than 1% Co and Mo each, the balance being essentially Fe and incidental impurities. The Ti:Al ratio is ≦2:1 and the Nb:C weight % ratio is within a range of 6:1 and 12:1. Ta may be substituted for Nb on an equiatomic basis.
In accordance with one aspect of the present invention, the foregoing and other objects are achieved by an Fe—Ni—Cr alloy composed essentially of, in terms of weight percent: 1 to 3.5 Al, up to 2 Co, 15 to 19.5 Cr, up to 2 Cu, 23 to 40 Fe, up to 0.3 Hf, up to 4 Mn, 0.15 to 2 Mo, up to 0.15 Si, up to 1.05 Ta, 2.8 to 4.3 Ti, up to 0.5 W, up to 0.06 Zr, 0.02 to 0.15 C, 0.0001 to 0.007 N, balance Ni, wherein, in terms of atomic percent: 6.5 ≦A=Al+Ti+Zr+Hf+Ta≦10,0.33≦B=Al÷(Al+Ti+Zr+Hf+Ta)≦0.065, and 4≦C=(Fe+Cr)÷(Al+Ti+Zr+Hf+Ta)≦10, the alloy being essentially free of Nb and V.
For a better understanding of the present invention, together with other and further objects, advantages and capabilities thereof, reference is made to the following disclosure and appended claims in connection with the above-described drawings.
Computational thermodynamics was used to identify new, lower cost alloys with microstructure similar to the commercial alloys and having comparable properties. In contrast to the comparable, commercially available alloys with Ni+Co content greater 60 wt. %, Ni+Co content in the new alloys ranges from about 30 wt. % to 51 wt. % with the potential to achieve comparable properties. This implies that the alloys will be of lower cost with the potential to achieve targeted fatigue life. For example a well-known, commonly used valve alloy known as “Alloy 751” has about 71 wt. % Ni+Co as shown in Table 1.
Constraints in Alloy Development: The alloys used for valve materials should have high strength, good oxidation resistance, should have sufficient ductility at high temperatures to be shaped into valves. They should also have high volume fraction of γ′ to achieve strengths at high temperature along with the lowest possible coarsening rates to maintain strength for the longest period of time. The following elements are added to achieve the appropriate benefits:
Nickel: Primary addition, certain amount of nickel is required to achieve beneficial strength, and ductility properties. Higher the temperature of operation, greater is the amount of Ni required.
Iron: Addition of element minimizes cost of alloy. Provides solid solution strengthening. Too much addition can destabilize austenitic matrix.
Chromium: At least 15 wt. % is critically required in the compositions to ensure good oxidation resistance but limited to 20 wt. % to minimize formation of undesirable BCC phase or other brittle intermetallics.
Aluminum+Titanium: Provides primary strengthening through the formation of γ′ precipitates. Ratio of aluminum to other elements such as Ti, Nb, and Ta changes the high temperature stability of the γ′ precipitates, strengthening achievable for an average precipitate size, and the anti-phase boundary (APB) energy.
Niobium: Forms stable MC-type carbides, also can segregate to γ′ and affect high temperature stability and coarsening rate of γ′, affects APB energy, decreases creep rate due to precipitation of carbides.
Tantalum: Forms stable MC-type carbides, also can segregate to γ′ and affect high temperature stability and coarsening rate of γ′, lower average interdiffusion coefficient in the matrix, affects APB energy, decreases creep rate due to precipitation of carbides.
Molybdenum: Added for solid solution strengthening, also is the primary constituent in M6C carbides. Decreases average interdiffusion coefficient. Too much addition can result in the formation of undesirable, brittle intermetallic phases and can reduce oxidation resistance
Manganese: Stabilizes the austenitic matrix phase. Provides solid solution strengthening.
Silicon: Assists in high temperature oxidation resistance, a maximum of 1% Si may be added.
Carbon, Nitrogen: Required for the formation of carbide and carbo-nitride phases that can act as grain boundary pinning agents to minimize grain growth and to provide resistance to grain boundary sliding. Fine precipitation of carbides and carbonitrides can increase high temperature strength and creep resistance.
Copper: Stabilizes the austenitic matrix, provides solid solution strengthening.
Cobalt: Provides solid solution strengthening.
Tungsten: Provides solid solution strengthening and decreases average interdiffusion coefficient. Too much can result in the formation of brittle intermetallic phases.
Typically, Ni-based alloys are strengthened through a combination of solid solution strengthening, and precipitation strengthening. The primary advantage of solid solution strengthened alloys is microstructural stability. Since strengthening is primarily obtained through the presence of solute elements in solid solution that may be different in size, and chemical composition from the solvent and not through the presence of precipitates, microstructural changes such as coarsening of precipitates will not be relevant in determining the properties of these alloys. Furthermore, fabrication such as forming and welding operations are simpler due to solid-solution strengthening being the primary strengthening mechanism. However, solid solution strengthened alloys can be primarily used in applications that need relatively lower yield and tensile strengths and lower creep strength when compared to precipitation-strengthened alloys but require consistent properties for long periods of time. Thus the γ′-strengthened alloys provide the higher strength required for applications for which the solid solution strengthened alloys have insufficient strength. One disadvantage with γ′ alloys is that the strength decreases with time at temperature due to the coarsening of γ′ precipitates with time. The rate of loss of strength is directly related to the rate of growth of precipitates which increases with increase in temperature (which also results in an increase in interdiffusion coefficients).
The strengthening potential of γ′ is determined by various factors with the major factors being the volume fraction, size and particle size distribution, lattice parameter misfit between the γ and γ′ phases, and the antiphase boundary energy. The compositions of the alloys determine the wt. % of γ′ and compositions of the γ and γ phases as a function of temperature which affect the lattice parameter misfit, and antiphase boundary energy. The heat-treatment conditions determine the size and size distribution of the strengthening phase. Diffusion coefficients and lattice parameter misfit have a strong influence on the coarsening of the γ′ phase.
The alloys described herein were designed to: (1) maximize γ′ content at a temperature higher than prior alloys of this type and particularly at a temperature of 870° C., (2) maximize the strengthening potential of γ′ which is related to the compositions of the phases present at higher temperatures, (3) include elements that minimize the coarsening rate of γ′, and (4) precipitate small amounts of carbides for grain size control and creep minimization. Broadest constituent ranges for alloys of the present invention are set forth in Table 2. Some examples thereof are set forth in Table 3, with Alloy 751 for comparison.
Quantities A, B, and C are atomic percent values defined as follows (all in at. %):
A=Al+Ti+Zr+Hf+Ta (3)
B=Al÷(Al+Ti+Zr+Hf+Ta) (4)
C=Cr÷(Ni+Fe+Cr+Mn) (5)
The formulae are calculated in atomic %, and then converted to weight % for facilitation of manufacture. Quantity A generally represents an indication of the amount of γ precipitates that can form in the alloy compositions and must be in the range of 5.9 to 10.5, preferably in the range of 6 to 9, more preferably in the range of 7.5 to 8.5.
Quantity B generally represents an indication of a ratio of Al to other elements in γ′ precipitates that can form in the alloy compositions and must be in the range of 0.3 to 0.65, preferably in the range of 0.35 to 0.6, more preferably in the range of 0.4 to 0.55. In some compositions, a most preferred range is 0.44 to 0.46.
Quantity C represents a critical relationship between Cr and certain other elements in the alloy compositions. Quantity C generally represents an indication of the composition of the matrix (γ), and the lattice misfit between the matrix (γ) and the precipitate(γ′), and must be in the range of 0.17 to 0.23, preferably in the range of 0.18 to 0.022, more preferably in the range of 0.185 to 0.215, and most preferably in the range of 0.200 to 0.213.
Another characteristic that may be considered is the lattice misfit between γ and γ′, generally defined as
2(aγ′−aγ)/(aγ′+aγ) (6)
where aγ′ represents the lattice parameter of γ′ and aγ represents the lattice parameter of γ. The calculated value represents an indication of the contribution to hardening (e.g., yield and tensile strengths) from coherency strains between the precipitate and the matrix of the alloy composition. The lattice misfit for alloys of the present invention at 870° C. can be expected to fall within the range of −0.135% to +0.064%, and preferably in the range −0.02% and +0.02%, as shown in Table 6.
Alloys 4, 9, 16, 20, 34, 35, 161, 162, 163, 164, 200, 490-1, 490-4, 490-5, 490-6, shown in Table 3, were made using well known, conventional methods. Arc cast ingots were annealed at 1200° C. in an inert gas environment (vacuum can also be used). The ingots were then hot-rolled into plates for mechanical testing.
The alloys were heat-treated to achieve optimum combination of high strength and ductility. A solution annealing treatment was performed at 1121° C. for 4 hours followed by an aging treatment at 760° C. for 16 hours. Thus, all the alloys can be cast, heat-treated, and mechanically processed into plates and sheets. The skilled artisan will recognize that other, conventional heat-treatment schedules can be used.
Table 2 shows the compositions of the new alloys while specific examples are shown in Table 3.
Table 4 shows a summary of the volume fraction of the various alloys at 870° C. The wt. % of the primary strengthening phase γ′ varies from 13.0 to 24.0 wt. %.
Table 5 shows the yield strengths at room temperature and at 870° C. for the new alloys and the baseline alloy 751. Note that the new alloys have strengths about 26.22% to 71.04% better than that of the baseline alloy 751.
Table 6 shows the variation of quantities A, B, and C, and calculated lattice misfit between γ and γ′ at 870° C.
Tables 7 and 8 show the respective compositions of γ and γ′ in each invention alloy at 870° C., all in at. %. The data show that these compositions affect strength and oxidation properties of alloys at 870° C.
Although the primary target of current alloys is 870° C., the new alloys are also shown to have better properties at 800° C. than the alloys described in the Sato et al. patents referenced hereinabove.
Improved fatigue properties of selected newly developed alloys are shown in
While there has been shown and described what are at present considered to be examples of the invention, it will be obvious to those skilled in the art that various changes and modifications can be prepared therein without departing from the scope of the inventions defined by the appended claims.
The United States Government has rights in this invention pursuant to contract no. DE-AC05-00OR22725 between the United States Department of Energy and UT-Battelle, LLC.
