Intermediate strength alloys for high temperature service in liquid-salt cooled energy systems

Abstract

An alloy is composed essentially of, in terms of weight percent: 6 to 8.5 Cr, 5.5 to 13.5 Mo, 0.4 to 7.5 W, 1 to 2 Ti, 0.7 to 0.85 Mn, 0.05 to 0.3 Al, 0.08 to 0.5 C, 0 to 1 Nb, with the balance Ni, the alloy being characterized by, at 850° C., a yield strength of at least 25 Ksi, a tensile strength of at least 30 Ksi, a creep rupture life at 12 Ksi of at least 45 hours, and a corrosion rate, expressed in weight loss [g/(cm2 sec)]10−11 during a 1000 hour immersion in liquid FLiNaK at 850° C., in the range of 6 to 39.

Description

CROSS-REFERENCE TO RELATED APPLICATIONS

This patent application is related to U.S. patent application Ser. No. 13/833,357, now U.S. Pat. No. 9,377,245, entitled “High Strength Alloys for High Temperature Service in Liquid-Salt Cooled Energy Systems” filed on Mar. 15, 2013, the entire disclosure of which is incorporated herein by reference. Moreover, this patent application is related to U.S. patent application Ser. No. 13/958,672 entitled “Creep-Resistant, Cobalt-Containing Alloys for High Temperature, Liquid-Salt Heat Exchanger Systems” filed on Aug. 5, 2013, the entire disclosure of which is incorporated herein by reference. Moreover, this patent application is related to U.S. patent application Ser. No. 13/962,197, now U.S. Pat. No. 9,435,011, entitled “Creep-Resistant, Cobalt-Free Alloys for High Temperature, Liquid-Salt Heat Exchanger Systems” filed on Aug. 8, 2013, the entire disclosure of which is incorporated herein by reference. Moreover, this patent application is related to U.S. patent application Ser. No. 14/152,215 entitled “intermediate Strength Alloys for High Temperature Service in Liquid-Salt Cooled Energy Systems” filed on Jan. 10, 2014, the entire disclosure of which is incorporated herein by reference.

BACKGROUND OF THE INVENTION

An ever-increasing demand for higher system thermal efficiency has necessitated the operation of power generation cycles and heat conversion systems for chemical processes at progressively higher temperatures. As system operating temperatures are increased, fewer materials with acceptable mechanical properties and environmental compatibility are known. This dearth of materials is particularly acute in applications at temperatures above 700° C. and at significant stress levels where liquid fluoride salts are favored as heat transfer media because of their high thermal capacity and low vapor pressure. There is therefore a need for structural alloys for high-temperature heat transfer applications in order to enable increased thermal efficiency of energy conversion and transport systems thereby reducing system costs as well as reducing the waste heat rejected to the environment.

Fluoride salt cooled High temperature Reactors (FHRs) potentially have attractive performance and safety attributes. Defining features of FHRs include coated particle fuel, low-pressure fluoride salt cooling, and high-temperature heat production. The FHR heat transfer technology base is derived primarily from earlier molten salt reactors and their coated particle fuel is similar to that developed for high-temperature helium-cooled reactors. The excellent heat transfer characteristics of liquid fluoride salts enable full passive safety, at almost any power scale thereby enabling large power output reactors, with less massive piping and containment structures, and consequent economies of scale. FHRs potentially have improved economics, increased safety margins, and lower water usage characteristics than conventional water-cooled reactors.

The fuel and coolants for FHRs are suitable for operation at temperatures well in excess of the upper temperature limits of available structural alloys. A limiting factor in achieving the highest possible FHR core outlet temperatures and thus thermal efficiency is the availability of structural alloys having sufficient creep strength at the required temperatures combined with suitable fluoride salt chemical compatibility as well as ease of fabrication. Hastelloy® N (trademark owned by Haynes International, Inc.) (also known as Alloy N and INOR-8), developed at Oak Ridge National laboratory (ORNL) in the 1950s and 1960s, is currently a leading candidate FHR structural alloy for operations below 700° C. Hastelloy® N is limited to use in low stress applications to a maximum temperature of about 704° C. due to insufficient creep strength at higher temperatures, is limited to use in high stress applications such as steam generator tubes to about 600° C. due to insufficient creep strength at higher temperatures, is not fully qualified to current code requirements for high temperature reactors, and is challenging to fabricate due to its work hardening characteristics. There is therefore a need for corrosion-resistant nickel-based structural alloys designed to possess good creep resistance in liquid fluorides at higher temperatures in order to provide substantial improvements in FHR economics and performance. Calculations reveal that a net thermal efficiency of greater than 50% (as compared to about 33% net thermal efficiency of existing reactors) would be likely for FHRs using a high temperature structural alloy with concurrent reductions in capital costs, waste generation, fissile material requirements, and cooling water usage.

Other applications for these alloys include concentrated solar power (CSP), and processing equipment for fluoride environments. Molten-salt power towers are envisioned as operating in excess of 650° C. to achieve efficiency and cost targets. Temperatures of up to 700° C. are anticipated with the use of commercial supercritical steam turbines, and up to 800° C. with the use of supercritical CO₂ Brayton cycle system. Molten salts allow for the storage of solar energy and thus, the decoupling of solar energy collection from electricity generation. At the higher temperatures, molten fluoride salts offer the advantages of high thermal capacity, high heat transfer, and low vapor pressure. The development of materials with acceptable mechanical and molten salt corrosion resistance will allow for achieving the desired efficiency and cost targets.

Development of a high temperature structural alloy tailored to the specific high temperature strength and liquid salt corrosion resistance needs of liquid fluoride salt cooled-energy systems (especially FHRs) is contemplated to be of critical importance to ensuring feasibility and performance thereof. Simultaneously achieving creep resistance and liquid fluoride salt resistance at higher temperatures is challenging because conventional additions of certain alloying elements for achieving improved creep resistance and resistance to oxidation in air are detrimental to liquid fluoride salt resistance.

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(Al,Ti,Nb)-based intermetallic precipitates that are coherently embedded in a solid solution 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 is primarily Nb, or Ti, 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. Existing data (shown in FIG. 1) on liquid fluoride salt resistance of Ni-based alloys show that alloys containing aluminum, and substantial amounts of chromium have lower resistance to liquid fluoride salt. Commercial Nickel-based alloys with high strengths typically contain significant amounts of Cr (greater than 15 wt. % Cr) making them unsuitable for use in contact with liquid fluoride salts. Compositions (in weight %) of several commercially produced Ni-based alloys are shown in Table 1.

Hastelloy® N is an alloy that was designed to balance resistance to liquid fluoride salt corrosion with good creep properties at temperatures up to 704° C. This alloy is a Ni—Mo alloy containing additional alloying elements with solid solution strengthening being the primary strengthening mechanism; Hastelloy® N does not have γ′ precipitation strengthening. Its nominal composition is given as71Ni-7Cr-16Mo-5Fe*-1Si*-0.8Mn*-0.2Co*-0.35Cu*-0.5W*-0.35Al+Ti*-0.08C*where * indicates maximum allowed content of the indicated elements. Hastelloy® N generally consists of the following elements to provide the corresponding benefits:

Chromium: Added to ensure good oxidation resistance but minimized to keep liquid fluoride salt corrosion within acceptable limits. Also provides solid solution strengthening. Too much addition results in excessive attack by liquid fluoride salts.

Molybdenum: Principal strengthening addition for solid solution strengthening, provides good resistance to liquid fluoride salt, and results in lower interdiffusion coefficients. Also is the primary constituent in M₆C carbides. Too much addition can result in the formation of undesirable, brittle intermetallic phases.

Iron: Minimizes cost of alloy. Provides solid solution strengthening. Too much addition can destabilize austenitic matrix and decrease resistance to liquid fluoride salt.

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/or carbonitride phases that can act as grain boundary pinning agents to minimize grain growth and to provide resistance to grain boundary sliding. Fine precipitation of carbide and/or carbonitride phases 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 that can be deleterious to processability.

Aluminum+Titanium are not desirable in Hastelloy® N, in order to minimize corrosion by liquid salt. Combined wt. % of Al+Ti is typically kept to less than 0.35.

FIG. 1 shows effects of alloying element additions on the depth of corrosion of Ni-alloys in 54.3LiF-41.0KF-11.2NaF-2.5UF₄ (mole percent) in a thermal convention loop operated between 815 and 650° C. (smaller depth of corrosion is better).

FIG. 2 shows the equilibrium phase fractions in Hastelloy® N as a function of temperature (nitrogen and boron are not included in the calculations). Note that solid solution strengthening and some carbide strengthening (through M₆C) are the primary strengthening mechanisms active in Hastelloy® N. This limits the strength and creep resistance of Hastelloy® N at high temperatures and restricts its useful temperatures to less than about 704° C. Components such as secondary heat exchangers need to withstand large pressure differences between salt on one side of the heat exchanger wall and a gaseous fluid at higher pressures on the other side. Such components hence need materials with high temperature strength greater than that of Hastelloy® N along with good resistance to salt, good oxidation resistance.

BRIEF SUMMARY OF THE INVENTION

In accordance with one aspect of the present invention, the foregoing and other objects are achieved by a new alloy that is composed essentially of, in terms of weight percent: 6 to 8.5 Cr, 5.5 to 13.5 Mo, 0.4 to 7.5 W, 1 to 2 Ti, 0.7 to 0.85 Mn, 0.05 to 0.3 Al, 0.08 to 0.5 C, 0 to 1 Nb, 0 to 0.01 Ta, with the balance Ni, the alloy being characterized by, at 850° C., a yield strength of at least 25 Ksi, a tensile strength of at least 30 Ksi, a creep rupture life at 12 Ksi of at least 45 hours, and a corrosion rate, expressed in weight loss [g/(cm² sec)]10⁻¹¹ during a 1000 hour immersion in liquid FLiNaK at 850° C., in the range of 6 to 39. The alloys can be essentially free of Fe, Co, and Hf, small amounts of which can be tolerated as impurities.

In the new alloys described herein, the range of Cr can be 6.6 to 7.22 weight percent, the range of Mo can be 5.76 to 12.79 weight percent, the range of W can be 0.44 to 6.32 weight percent, the range of Ti can be 1.15 to 1.22 weight percent, the range of Mn can be 0.75 to 0.77 weight percent, the range of Al can be 0.09 to 0.2 weight percent, and the range of C can be 0.2 to 0.47 weight percent.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a combination table and bar graph showing effects of alloying element additions on the depth of corrosion of Ni-alloys in 54.3LiF-41.0KF-11.2NaF-2.5UF₄ (mole percent) in a thermal convention loop operated between 815 and 650° C.

FIG. 2 is a graph showing phase equilibria for a typical composition of Hastelloy® N as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 3 is an expanded view of a portion of the graph shown in FIG. 2 to show details.

FIG. 4 is a graph showing phase equilibria for Alloy 1 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 5 is an expanded view of a portion of the graph shown in FIG. 4 to show details.

FIG. 6 is a graph showing phase equilibria for Alloy 2 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 7 is an expanded view of a portion of the graph shown in FIG. 6 to show details.

FIG. 8 is a graph showing phase equilibria for Alloy 6 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 9 is an expanded view of a portion of the graph shown in FIG. 8 to show details.

FIG. 10 is a graph showing phase equilibria for Alloy 6 mod 1 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 11 is an expanded view of a portion of the graph shown in FIG. 10 to show details.

FIG. 12 is a graph showing phase equilibria for Alloy 6 mod 5 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 13 is an expanded view of a portion of the graph shown in FIG. 12 to show details.

FIG. 14 is a graph showing phase equilibria for Alloy 25 as a function of temperature (nitrogen and boron are not included in the calculations).

FIG. 15 is an expanded view of a portion of the graph shown in FIG. 14 to show details.

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.

DETAILED DESCRIPTION OF THE INVENTION

New, essentially Fe-free, alloys primarily strengthened through carbides having improved high temperature strength and creep resistance; general composition limits are shown in Table 2. The primary strengthening in the new alloys is achieved through the precipitation of carbides along with solid solution strengthening. Moreover, the new alloys exhibit an advantageously lower average interdiffusion coefficient in the matrix. The skilled artisan will recognize that a lower interdiffusion rate results in, at high temperatures, lower coarsening rate of carbides, improved creep properties, lower oxidation rate, and lower corrosion rate.

Computational design was used to ensure that formation of brittle intermetallic phases that form in the new alloys is very low or zero weight % in the operating temperature range of contemplated greatest interest (750 to 950° C.). In the alloys, small amounts of carbide formers such as Ti, Nb, and Ta have been added to form carbides and the carbon levels have been increased as compared to Hastelloy® N. Carbides such as, for example, MC, M₂₃C₆, or a combination thereof are present in these alloys in addition to the M₆C carbides that are formed in Hastelloy® N.

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 as compared to precipitation-strengthened alloys but require consistent properties for a very long period of time (25-80 years).

In contrast, the carbide strengthened alloys described herein provide the higher strength required for applications for which the solid solution strengthened alloys have insufficient strength and also provide improved creep strength. One disadvantage with conventional carbide strengthened alloys is that the strength decreases with time at temperature due to the coarsening of the carbide precipitates. 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 addition of sufficient amount of Mo, W, and/or Ta to the alloys of the present invention ensures that the interdiffusion coefficient is kept as low as reasonably possible and the coarsening rates are low, thus retaining properties for an extended period of time.

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 Hastelloy® N for comparison. It is contemplated that alloys of the present invention may contain up to 5% Fe with concomitant reduction in some beneficial properties, such as creep resistance and oxidation resistance.

EXAMPLES

Alloys 1, 2, 6, 6 Mod 1, 6 Mod 5, and 25, shown in Table 3, were made using well known, conventional methods. Vacuum 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. A solution annealing treatment was performed at 1150° C. for 1 hour. 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.

FIGS. 3-15 show the results from equilibrium calculations obtained from the computational thermodynamics software JMatPro v 6.2. Actual compositions were used for all the calculations.

Table 4 shows equilibrium wt. % of phases present in alloys at 850° C. The primary strengthening phase is predicted to be carbides. Two different types of carbides are predicted to be present with the primary carbides being M₆C and M₂₃C₆. Total carbides can be present in a range of 5.4 to 16.84 wt. %. M₆C carbides can be present in a range of 1.9 to 16.14 wt. %. M₂₃C₆ carbides can be present in a range of up to 3.5 wt. %. Ratio of M₂₃C₆ to total carbides can be up to 0.64.

Yield and tensile strengths have been measured at 850° C. and compared with the baseline properties of Hastelloy® N and are shown in Table 5. Typical yield strengths of alloys of the present invention are contemplated to be at least 25 Ksi, preferably at least 30 Ksi. Typical tensile strengths of alloys of the present invention are contemplated to be at least 30 Ksi, preferably at least 34 Ksi.

Creep rupture life has been measured in the solution annealed condition at 850° C. at a stress level of 12 Ksi with as shown in Table 6. Creep rupture lives of alloys of the present invention are contemplated to be at least 45 hours, preferably at least 47 hours. It was surprising and unexpected that creep rupture lives of Alloy 6 and Alloy 6 Mod 5 were determined to be in the thousands of hours.

Table 4 and Table 6 show that, in general, creep rupture resistance usually increases with increasing ratio of M₂₃C₆ to total carbide contents and is best when this ratio is about 0.64. Exceptions to this trend can be seen in the experimental results.

Resistances to liquid salt corrosion were measured by placing the alloy specimens of measured dimensions and weight in sealed molybdenum capsules in contact with a fixed amount of FLiNaK, a liquid salt heat exchange medium. The molybdenum capsules were enclosed in outer capsule to minimize high temperature air oxidation and heated in a furnace at 850° C. for 1,000 hours. After exposure, the capsules were opened and the specimens cleaned, weighed and their dimension measured. Corrosion resistance to liquid fluoride salt was evaluated based on normalized weight change and metallography and scanning electron microscopy. Results presented in Table 7 demonstrate that these alloys all have corrosion rates slightly higher than that of Hastelloy® N in these isothermal tests but with significantly improved mechanical properties. Thus a balance has been struck between improved mechanical properties and resistance to attack by liquid fluoride salt. Note that the aluminum and chromium has been kept at minimum required levels without adversely affecting the oxidation resistance and the salt resistance of these alloys. Typical corrosion rates of alloys of the present invention, expressed in weight loss [g/(cm² sec)]×10⁻¹¹ during a 1000 hour immersion in liquid FLiNaK at 850° C., are contemplated to be in the range of about 6 to about 39. Thus a balance has been struck between improved mechanical properties and resistance to attack by liquid fluoride salt.

Table 8 shows the relationship between the susceptibility to corrosion of the alloys shown by liquid fluoride salts, specifically FLiNaK with the Mo Equivalent, defined as

$\mathrm{Mo}\mathrm{Equivalent}=\%\mathrm{Mo}+1.15\times \frac{183.84\left(\mathrm{Atomic}\mathrm{Weight}\mathrm{of}W\right)}{95.95\left(\mathrm{Atomic}\mathrm{Weight}\mathrm{of}\mathrm{Mo}\right)}\times \%W$ where % refers to atomic percent of the element present in the alloy. 183.84 is the atomic wt. of W 95.95 is the atomic wt. of Mo. It has been observed that for these alloys the Mo Equivalent should be in the range of 4 to 9 for good resistance to liquid fluoride salts, specifically FLiNaK.

Table 8 shows the corrosion susceptibility index which quantifies the susceptibility to corrosion of the alloys shown in Table 3 by liquid fluoride salts, specifically FLiNaK. Corrosion susceptibility index is generally defined as follows:

$\mathrm{CSI}=\frac{\%\mathrm{Al}+\%\mathrm{Cr}+\%\mathrm{Ti}+\%\mathrm{Nb}+\%\mathrm{Hf}+\%\mathrm{Ta}}{\%\mathrm{Ni}+\%\mathrm{Fe}+\%\mathrm{Co}+\%\mathrm{Mn}+\%\mathrm{Mo}+\%W+\%\mathrm{Re}+\%\mathrm{Ru}}$ where % refers to atomic percent of the element present in the alloy. It has been observed that for these alloys, CSI should be no less than about 0.1 and no greater than about 0.14 in addition to maintaining the elements in the preferred ranges. This results in the optimum combination of mechanical properties (high temperature strength and creep resistance) and resistance to fluoride salts.

Tables 1-8 follow.

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.

TABLE 1
Compositions of several commercial Ni-based alloys (in weight %).
Alloy	C	Si	Mn	Al	Co	Cr	Cu	Fe	Mo	Nb	Ni	Ta	Ti	W	Zr
X750	0.03	0.09	0.08	0.68	0.04	15.7	0.08	8.03	—	0.86	Bal	0.01	2.56	—	—
Nimonic 80A	0.08	0.1	0.06	1.44	0.05	19.6	0.03	0.53	—	—	Bal	—	2.53	—	—
IN 751	0.03	0.09	0.08	1.2	0.04	15.7	0.08	8.03	—	0.86	Bal	0.01	2.56	—	—
Nimonic 90	0.07	0.18	0.07	1.4	16.1	19.4	0.04	0.51	0.09	0.02	Bal	—	2.4	—	0.07
Waspaloy	0.03	0.03	0.03	1.28	12.5	19.3	0.02	1.56	4.2	—	Bal	—	2.97	—	0.05
Rene 41	0.06	0.01	0.01	1.6	10.6	18.4	0.01	0.2	9.9	—	Bal	—	3.2	—	—
Udimet 520	0.04	0.05	0.01	2.0	11.7	18.6	0.01	0.59	6.35	—	Bal	—	3.0
Udimet 720	0.01	0.01	0.01	2.5	14.8	15.9	0.01	0.12	3.0	0.01	Bal	—	5.14	1.23	0.03
Alloy 617	0.07	0	0	1.2	12.5	22	0	1	9	0	54	0	0.3	0	0

TABLE 2
Compositions of new alloys (analyzed compositions in wt. %)
	Element	Minimum wt. %	Maximum wt. %
	Cr	6	8.5
	Mo	5.5	13.5
	W	0.4	7.5
	Ti	1	2
	Mn	0.7	0.85
	Al	0.05	0.3
	Nb	0	1
	Ta	0	0.01
	C	0.08	0.5
	Ni	Balance
	Co	Essentially 0
	Hf	Essentially 0
	Fe	Essentially 0

TABLE 3
Compositions of new alloys compared to Hastelloy ® N (analyzed compositions in wt. %)
Alloy	Ni	Fe	Al	Co	Cr	Mn	Mo	Ti	Nb	Hf	Ta	W	C	B**	N**	Total
Hastelloy ® N*	68.7	5		0.2	7	0.8	16	0	—	0	0	0.5	0.08	0.01	—	100
Alloy 1	77.6794	0	0.1	0	6.73	0.77	12.77	1.2	0	0	0	0.5	0.25	0.0004	0.0002	100
Alloy 2	77.2594	0	0.2	0	6.74	0.77	12.79	1.19	0	0	0	0.58	0.47	0.0004	0.0002	100
Alloy 6	84.6588	0	0.1	0	6.6	0.75	6.01	1.19	0	0	0	0.45	0.24	0.001	0.0002	100
Alloy 6 Mod 1	82.6683	0.01	0.11	0	7.22	0.76	7.06	1.47	0	0	0.01	0.44	0.25	0	0.0017	100
Alloy 6 Mod 5	80.4393	0.05	0.09	0	6.78	0.75	5.76	1.08	0	0	0	4.81	0.24	0	0.0007	100
Alloy 25	76.4397	0.01	0.1	0	6.88	0.77	6.43	1.91	0.94	0	0	6.32	0.2	0	0.0003	100
*Nominal composition shown, Hastelloy ® N also contains 1 Si, 0.35 Cu, 0.5 max of Al + Ti
**Boron and Nitrogen are not included in the equilibrium calculations

TABLE 4
Equilibrium wt. % of Phases Present in Alloys at 850° C.
	Wt. %	Wt. %	Wt. %	Total	M23C6/Total
Alloy	γ	M6C	M23C6	Carbides	Carbides
Hastelloy ® N	98.77	1.23	0	1.23	0
Alloy 1	90.25	9.75	0	9.75	0
Alloy 2	83.17	16.13	0.70	16.83	0.04
Alloy 6	94.60	1.92	3.48	5.40	0.64
Alloy 6 Mod 1	93.02	4.83	2.15	6.98	0.31
Alloy 6 Mod 5	93.95	3.19	2.86	6.05	0.47
Alloy 25	92.47	7.53	0	7.53	0.0

TABLE 5
Yield and Tensile Strengths of Alloys at 850°
C. and Improvement over the baseline alloys Alloy N.
			% Improvement
Alloy	Yield Strength	Tensile strength	in Tensile
Hastelloy ® N	35.29	45.70	0
Alloy 1	40.24	42.32	14
Alloy 2	38.46	40.27	8.9
Alloy 6	31.10	35.03	−11.9
Alloy 6 Mod 1	41.9	53.7	17.5
Alloy 6 Mod 5	42	56.8	24.3
Alloy 25	38.47	43.33	9

TABLE 6
Creep rupture lives of alloys at 850° C., at a stress
of 12 Ksi and improvement over the base alloy Alloy N.
		% Improvement in
Alloy	Creep Rupture Life	creep rupture life
Hastelloy ®	3.77 (average of 3 tests)	0
Alloy 1	102.7	2624
Alloy 2	48.05	1175
Alloy 6	3984	105576
Alloy 6 Mod 1	53.4	1316
Alloy 6 Mod 5	2285.2	60515
Alloy 25	110.6	2834

TABLE 7
Corrosion Rate (Weight Loss) Measured During a
1000 hour immersion in liquid FLiNaK at 850° C.
Alloy         Corrosion rate [g/(cm2sec)]10−11
Hastelloy ® N 1.21
Alloy 1       26.34
Alloy 2       26.04
Alloy 6       38.84
Alloy 6 Mod 1 34.03*
Alloy 6 Mod 5 27.26*
Alloy 25      7.09
*Calculated from Mo equivalent

TABLE 8
Composition of alloys in at. % and the calculation of the Corrosion Susceptibility Index (CSI)
Alloy	Ni	Fe	Al	Co	Cr	Mn	Mo	Ti	Nb	Hf	Ta	W	C	Mo eq.	CSI
Hastelloy ® N	75.735	4.443	0	0.157	7.473	0.594	10.34	0	0	0	0	0.02	0.154	—	0.0819
Alloy 1	80.098	0	0.224	0	7.833	0.848	8.055	1.517	0	0	0	0.165	1.26	8.42	0.1074
Alloy 2	78.925	0	0.444	0	7.772	0.84	7.993	1.491	0	0	0	0.189	2.346	8.41	0.1104
Alloy 6	85.017	0	0.218	0	7.481	0.805	3.692	1.465	0	0	0	0.144	1.178	4.01	0.1022
Alloy 6 Mod 1	83.196	0.0106	0.241	0.0	8.202	0.817	4.347	1.814	0.0	0	0.0033	0.141	1.229	4.66	0.1159
Alloy 6 Mod 5	83.186	0.054	0.202	0	7.915	0.829	3.644	1.369	0	0	0	1.588	1.213	7.14	0.1062
Alloy 25	80.362	0.011	0.229	0	8.164	0.865	4.135	2.462	0.624	0	0	2.121	1.027	8.81	0.1312

Claims

1. An alloy consisting essentially of, in terms of weight percent: Cr 6 to 8.5Mo 5.5 to 13.5W 0.4 to 7.5Ti 1 to 2Mn 0.7 to 0.85Al 0.05 to 0.3C 0.08 to 0.5Nb 0 to 1Ta 0 to 0.01Ni balancesaid alloy being characterized by, at 850° C., a yield strength of at least 25 Ksi, a tensile strength of at least 30 Ksi, a creep rupture life at 12 Ksi of at least 45 hours, and a corrosion rate, expressed in weight loss [g/(cm2 sec)]10−11 during a 1000 hour immersion in liquid FLiNaK at 850° C., in the range of 6 to 39, wherein said alloy is devoid of γ′ precipitates, and is further characterized by a corrosion susceptibility index of from 0.10 to 0.14, and wherein said alloy is essentially Fe free, Co free, and Hf free.

2. An alloy in accordance with claim 1 wherein said alloy is essentially Nb free.

3. An alloy in accordance with claim 1 wherein said alloy is essentially Ta free.

4. An alloy in accordance with claim 1 wherein the range of Cr is 6.6 to 7.22 weight percent.

5. An alloy in accordance with claim 1 wherein the range of Mo is 5.76 to 12.79 weight percent.

6. An alloy in accordance with claim 1 wherein the range of W is 0.44 to 6.32 weight percent.

7. An alloy in accordance with claim 1 wherein the range of Ti is 1.15 to 1.22 weight percent.

8. An alloy in accordance with claim 1 wherein the range of Mn is 0.75 to 0.77 weight percent.

9. An alloy in accordance with claim 1 wherein the range of Al is 0.09 to 0.2 weight percent.

10. An alloy in accordance with claim 1 wherein the range of C is 0.2 to 0.47 weight percent.

11. An alloy in accordance with claim 1 wherein said alloy is further characterized by 5.4 to 16.84 wt. % total carbides.

12. An alloy in accordance with claim 11 wherein said alloy is further characterized by 1.92 to 16.14 wt. % M6C carbides.

13. An alloy in accordance with claim 11 wherein said alloy is further characterized by up to 3.48 wt. % M23C6 carbides.

14. An alloy in accordance with claim 1 wherein said yield strength is at least 30 Ksi.

15. An alloy in accordance with claim 1 wherein said tensile strength is at least 34 Ksi.

16. An alloy in accordance with claim 1 wherein said creep rupture life is at least 47 hours.

17. An alloy in accordance with claim 1 wherein said alloy is further characterized by a Mo equivalent of no less than 4, wherein:

18. An alloy in accordance with claim 1 wherein the range of C is from 0.24 to 0.5 weight percent.