US2949355A - High temperature alloy - Google Patents

Description

Aug. 16, 1960 Filed July 27, 1955 Brinell Hardness Number w. w. DYRKACZ EI'AL 2,949,355

Solution Time (Hours) Fig. I

INVENTORS Wasil W. Dyrkucz Edward E. Reynolds Richard FyMocForlune and Richard K. Pitler Aug. 16, 1960 w. w. DYRKACZ ET AL 2, 4

' HIGH TEMPERATURE ALLOY Filed July 27, 1955 2 Sheets-Sheet 2 Hardness Number Brinell Treated (OF) Aging Timel6 Hours; Air Cooled Fig. 2

INVENTORS. Wusll W. Dyrkocz Edward E. Reynolds Richard R MucForlune and Richard K. Pitler I I I I Solution '00 I200 I300. I400 I500 I600 Unite States gheny Lndlnm Steel Corporation, Brackenridge, Pa, a corporation of Pennsylvania Filed July 27, 1955, Ser. No. 524,679-

5 Claims. (Cl. 75-426) This invention relates to improvements in austenitic iron base alloys, for use at elevated temperatures.

Recent trends in the aircraft and allied industries have shown the desirability of operating gas turbine engines and the like at higher temperatures to increase the efficiency thereof. This has resulted in the problem of finding an alloy which will have the necessary high stress level characteristics at the higher operating temperatures. In the past the industry has adverted to the use of iron base alloys containing high percentages of strategic alloying elements such as nickel and columbium or alternatively has employed the super alloys or non-iron base alloys. While some of these alloys have exhibited sufiicient strength at the desired operating temperatures, the strategic nature of the alloying elements as well as their high cost has in certain cases limited their commercial acceptance.

An object of this invention is to provide an austenitic iron-base nickel-free alloy which is capable of use at high stress levels at temperatures of up to about 1350 F.

Another object of this invention is to provide an austenitic iron base nickel-free alloy having manganese and chromium as essential alloying elements, with small amounts of molybdenum, vanadium and nitrogen contained therein, said alloy being suitable for use at high stress levels at temperatures up to about 1350 F.

A more specific object of this invention is to produce an austenitic iron base nickel-free alloy containing critical amounts of manganese, carbon, nitrogen, chromium, molybdenum and vanadium as essential alloying elements, said alloy being suitable for use at high stresses at tem peratures of up to about 1350 F.

These and other objects of this invention will become apparent from the following description when taken in conjunction with the accompanying drawings in which:

Figure 1 is a graph, the curves of which illustrate the eifect of the solution heat treatment temperature on the hardness of the alloy; and

Fig. 2 is a graph, the curves of which illustrate the effect of the aging temperature on the hardness of the alloy.

In its broader aspects the alloy of the invention comprises from about 0.20% to 0.35% carbon, from about 0.15% to about 075% silicon, from about 16.0% to about 20.0% manganese, from about 11.5% to about 13.5% chromium, from about 2.0% to about 4.0% molybdenum, from about 0.60% to about 0.95% vanadium, from about 0.10% to about 0.25% nitrogen and the balance substantially iron. It will be appreciated that where the balance is referred to as substantially all iron, it will of course include the incidental impurities usually found in the manufacture of said steels.

Each of the elements performs a specific function within the alloy. Carbon is used for imparting the requisite strength to this steel. It is an austenitizing element and contributes to maintaining a stable austenitic structure when at least 0.20% carbon is maintained therein. The

Patented Aug. 16, 1960 chief austenitizing element of this alloy is manganese and at least 16.0% is needed in order to insure a completely austenitic structure in the alloy. In order to impart sufiicient corrosion resistance especially in the presence of the combustion products of the fuels such as are used in gas turbines or other engines, at least 11.5% chromium is required. Amounts of greater than 13.5% chromium tend to produce delta ferrite which seriously affects the physical properties of this alloy. Molybdenum and vanadium impart additional strength to the alloy of this invention by strengthening the solid solution phase. Vanadium also contributes to precipitation hardening in the alloy. In addition to its strengthening effect on the alloy, nitrogen substantially contributes to the stabilization of the austenite phase of the alloy. It is desirable to maintain the silicon content at least 0.15% so that the alloy will be sufficiently deoxidized. Iron comprises the balance of the alloy with the exception of the incidental impurities such as copper, cobalt and nickel up to 2.0% each, phosphorus and sulfur up to 0.05% each, and other elements which are usually found in these alloys.

Reference may be had to Table I illustrating the range TABLE I [Composition (weight percent) Element General Preferred Range Range 0. 0.204) as 02541.30. Si 0 The alloy of this invention may be produced in any of the well known manners, for example, electric arc furnace melting. Predetermined quantities of scrap and/ or ferroalloys are placed in an electric arc furnace and are melted. Since the details of such melting operations are well known in the art, they will not be described here in detail. After the molten metal has attained the desired analysis, it is tapped into a ladle and cast therefrom into ingots. These ingots may then be processed in any of the Well known manners, for example, forging, pressing, rolling, extruding and the like into the form of a semi-finished mill product, for example, bar, billet, sheet, strip and the like. The semi-finished mill product may be readily fabricated into specific articles such as parts for gas turbines or other engines or apparatus and parts for use at high stress levels and at temperatures of up to about 1350 F.

In order to demonstrate more clearly the characteristics of the alloy of this invention reference may be had to Table II illustrating a number of alloys which were made and tested to illustrate the effect of the chemical composition on the physical properties of the alloy of this invention. It will be apparent by inspection that the alloying components are both within and outside the general range given in Table I. This was done in order to illustrate the effect of tthese components on the physical properties.

TABLE H [Chemical analysis (weight percent) .1

O Si Mn Cr Mo V N Fe 0. 28 0.42 16. 81 11. 85 0. 86 0.160 Bal. 0. 25 0. 19 17. 18 12. 21 2.64 0.80 0. 195 Bal. 0.35 0.21 17.18 12. 25 3.12 0. 86 0.189 Bal. 0. 4 0. 28 19. 98 12. 17 2. S4 0. 89 0. 164 Bill. 0. l 0. 48 16. 80 12. 37 3. 45 1. 04 0. 181 Bal. 0. 37 0. 16 18. 20 12.18 1. 74 0.95 0.156 Bal. 0. 30 0. 26 18. 40 12. 21 3.18 0.78 0. 020 Hal. 0. 30 0.24 17.36 12. 19 3. 04 0. 74 0. 138 132.1. 0. 31 0.25 17. 66 12. 04 2. 85 0.70 0. 165 1321. 0.30 0.26 17.66 11. 97 2. 88 0.68 0.153 Bel. 0. 27 0.25 17. 60 12. 07' 2. 81 0. 73 0.174 Ba l. 0. 31 0. 37 17.30 12. 2. 98 0. 74 0.224 Bal.

, Reference may be had to Table III showing the results of tests designed to illustrate the efiect of carbon, nitrogen and molybdenum on the tensile properties of the alloy of this invention when the given elements are varied within and outside of the ranges given in Table I.

TABLE III Efiect of carbon, nitrogen and molybdenum on tensile properties at room temperature 0027 Element Tensile oasei n1. Bed. of Alloy No. (percent) Strength Yield (percent) Area (p.s.i.) Strength (percent) (p.s.i.)

It can be seen by inspection of the test results listed in Table III that increasing the carbon content produces a substantial increase in the tensile and yield strengths without adversely afiecting the ductility of this alloy. This is apparent by comparing alloys D-251, D-2l8, R-48, R-47 and D-183 which illustrate that increasing the carbon content from 0.10% to 0.35% produces a corresponding increase in the tensile strength from 111,200 p.s.i. to 163,000 p.s.i. while at the same time increasing the ductility from 9% to 24% as measured by the percentage elongation. A substantial increase is also noted in the yield strength, namely an increase from 42,000 p.s.i. to 81,200 p.s.i. It is thus apparent that the carbon must be maintained within the given range in order to have the optimum combination of strength and elongation.

Referring now to alloys D-333, R-47, D-218, R-48, D-183 and D482 it can be seen that when the nitrogen is maintained Within the range given, it produces a definite strengthening effect on the alloy of this invention without adversely affecting the ductility. Thus by comparing the test data on alloy D-333 and alloy D-182 it can be seen that there is a definite increase in the tensile and yield strengths of from 121,000 p.s.i. to 152,500 p.s.i. and from 45,400 p.s.i. to 80,900 p.s.i., re-

spectively. While the ductility is lowered somewhat by the addition of nitrogen to the base composition, that is from 41% to 33%, it is to be noted that increasing the nitrogen content within the given range, for example from 0.153% to 0.195%, has increased the ductility as measured by the elongation from 30% to 33% so that the slight decrease in ductility occasioned by the addition of nitrogen is not too significant. Thus the nitrogen, in addition to providing for an austenitic structure within this alloy, produces a definite increase in the strength while at the same time maintaining an excellent ductility.

By comparing the test results given for alloys D-116, D 267, D-182 and D-l83, it can be seen that the molybdenum also contributes substantially to the strength of the alloy. It is noted that the strengths of the alloys containing no molybdenum, for example D-116, are low compared to those of molybdenum containing alloys such as alloys D-267, D-182 and D-18'3. Thus it is seen by the test results recorded in Table III that an increase of from 142,200 p.s.i. to 163,000 p.s.i. in the tensile strength from 65,600 p.s.i. to 81,200 p.s.i. in the yield strength and from 14% to 24% in the elongation is accomplished by increasing the molybdenum content from 1.74% to 3.12%. It is thus apparent that the carbon, nitrogen and molybdenum contribute substantially to the strength of the alloy of this invention as measured by its tensile properties.

In all true engineering applications, the measurement of the room temperature tensile properties will not give a true criteria for predicting the performance of metals at elevated temperatures. It is therefore apparent that recourse must be had to a measurement of these alloys under superimposed conditions of the type and magnitude that the alloys may encounter in actual operation. The stress-rupture test fulfills these conditions to a large extent and is the generally accepted evaluation criteria of alloys for high temperature service. Reference may be had to Table IV illustrating the stress rupture proper ties of the alloy of this invention.

TABLE IV Efiect of carbon, nitrogen and molybdenum on stress ruptu're properties at 1200 F.

Element Stress Rupture El. Red. of Alloy No. (percent) (p.s.i.) Time (percent) Area (Hrs) (percent) By inspection of the test results given in Table IV it is seen that by increasing the carbon content a significant increase in the time to produce rupture has occurred when these alloys are tested at 1200 F. and under a stress of 50,000 p.s.i. Thus, by comparing alloys D-25-1 with alloy D-218 it is seen that an increase of from 0.10% carbon to 0.24% carbon has produced an increase in the time to produce rupture of from 3 hours to 116 hours.

A further increase is also noted when comparing alloy D-218 with alloy 13-183 which shows that by increasing the carbon content from 0.24% to 0.35% the time to produce rupture has been increased. Carbon contents above 0.35 result in low ductility.

The efiect of nitrogen on the rupture time is also illustrated in Table IV by comparing alloys D433, R-45, R-47, R48 and R-S 8. It is seen when comparing alloy D-333 with alloy R-45 that increasing the nitrogen content from 0.029% to 0.138% has produced an increase in the rupture time of from 41 hours to 97 hours. It is to be noted in this respect, however, that the magnitude of the rupture time is substantially less than that indicated for the series of the alloys illustrating the eifect of carbon on the time to produce rupture. This difierence is due to a different stress level (55,000 p.s.i. as compared to 50,000 p.s.i. which was used in the carbon series) applied to the series of alloys illustrating the effects of nitrogen on the stress rupture properties. A corresponding increase in the rupture time is also evidenced by comparing alloy R-45 with alloy R-58 which indicates that increasing the nitrogen content from 0.138% to 0.224% has produced an increase in the time to produce rupture from 97 hours to 128 hours. It is thus evident that the nitrogen has a definite beneficial effect upon the stress rupture properties of the alloy of this invention when maintained within the range given in Table I.

Comparing the test results given for alloys D-116, D267, D-182 and D-183, it is apparent that increasing the molybdenum content produces a definite increase in the time to produce rupture. This effect is especially noticeable when comparing alloy D116 with alloy D- 267 and comparing alloy D-267 with alloy D-183 which illustrate that by increasing the molybdenum content from 0% to 1.74% and from 1.74% to 3.12% the time to produce rupture has been increased from 20 hours to 138 hours and from 138 hours to 240 hours, respectively. In producing the alloy of this invention, it was found that the upper limit of the critical amount of molybdenum set forth in Table I must not be exceeded since molybdenum contents in excess of 4.0% have an adverse effect on the austenitic stability of the alloy. It is apparent that carbon, nitrogen and molybdenum produce a definite and substantial efiect on the rupture life of the alloys of this invention. Judging from the criterion of ductility, however, the carbon, nitrogen and molybdenum contents must be maintained Within the ranges given within Table I.

It is to be noted that the foregoing rupture times have been based on the socalled smooth bar test and thus do not give any indication of the notch rupture sensitivity of the alloys of this invention. The test used to measure the notch bar rupture life consists of machining a V- notch into a standard rupture specimen at the reduced section of the specimen which in this instance had an outside diameter of 0.275. The V-notch is a 60 notch positioned at the center of the reduced section, and has a 0.005 radius at the base of the notch and a 0.195" diameter across the base of the notch. The notch is produced by wet grinding and has a geometry which is such as to produce a theoretical stress concentration factor (K of 4.2 for the notch bar in tension. This type of test is used as a measure of notch sensitivity in rupture life, the criterion used to evaluate the characteristic of freedom from notch sensitivity being that the notched bar rupture life be at least as long as the standard or smooth bar rupture life under a given stress and temperature.

While the higher carbon and nitrogen contents seem to increase the sensitivity to notch rupture, that is, decrease the time to produce rupture in a notched bar, the higher molybdenum content increases the time to produce rupture. The beneficial eifect of molybdenum in its higher contents is somewhat ofiset by the higher carbon 6 and nitrogen contents. However, the overall effect on the notch rupture sensitivity of these alloys is greatly affected by heat treatment to which the alloy of this invention is subjected as will be described more fully hereinafter.

Since the alloy of this invention gains its strength in part from precipitation hardening, it must be subjected to a heat treatment in order to obtain the optimum prop erties therein. The heat treatment found to be successful with the alloy of this invention has been to first subject the alloy to a solution heat treatment at a temperature in the range between 1800 F. and 2100 F. for a time period of about one hour. The effect of the solution heat treatment temperature is illustrated by reference to Fig. 1, curves 10, 12 and 14 of which illustrate the variations in the hardness with respect to time at different solution heat treatment temperatures of 1750 F. and 1900 F. and 2050 B, respectively. It is evident from curve 10 of Fig. 1 that even after four hours at the solution heat treatment temperature of 1750 F. the alloy of this invention has not responded to heat treatment as evidenced by the high hardness of almost 300 BHN. Note also that there still is a slope to curve 10 which indicates that the alloy is still not in its solution heat treated softest form. On the other hand, solution heat treatment temperatures of 1900 F. and 2050 F. for a time period of as little as one hour are sufl icient as illustrated by curves 12 and 14, respectively, to render the alloy in its soft ductile form. It is to be noted from curves 12 and 14 that there is no significant decrease in the hardness after holding the alloys at the respective solution temperatures for more than one hour. Therefore it is apparent that a temperature of at least 1800 F. for about one hour is needed in order to render the alloy in its soft ductile condition. Temperatures in excess of 2100 F. are not suited for the heat treatment of the alloy of this invention because of the possibility of producing delta ferrite and/or burning or incipient melting of the alloy at the grain boundaries.

In order to maintain the alloy in the soft ductile condition after solution heat treatment and thus take full advantage of the aging characteristics of this alloy, it has been found desirable to rapidly cool the alloy of this invention from the solution heat treatment temperature. It has been found in all cases that when these alloys are solution heat treated within the temperature range given hereinbefore for a time period of about one hour and water quenched therefrom, these alloys have always possessed a hardness of less than 275 BHN.

In order to obtain the desired hardness, strength and ductility, the solution treated alloy is aged at a temperature in the range between 1200 F. and 1500 F. for a time period of between 10 and 20 hours. The effect of the aging treatment is to provide the alloy with sufficient hardness and strength without adversely atfecting the ductility and especially decreasing the sensitivity to the notch rupture. It was found that when alloy R-46 which had been previously solution treated at temperature of 1900 F. for one hour and then Water quenched and had a hardness of 248 BHN was aged for 16 hours at 1300 F., the hardness of the alloy increased to 293 BHN. Similar results were obtained after alloy R-46 was subjected to a solution treatment of 2050 F. for one hour, followed by a water quench and then aged 12 hours at 1300 F. in that the hardness increased from 223 BHN to 311 BHN.

The alloy of this invention may also be strengthened by hot-cold working and/or by cold working. In hotcold working, the alloy of this invention is heated to a temperature below the recrystallization temperature, which in this instance is in the range of between 1000 F. and 1500 F., and is worked at this temperature. Since the working temperature is below the recrystallization temperature, cold Work is imparted to the alloy and the strength properties are increased, especially the yield strength. The effect of the hot-cold work on the alloy of this invention is not lost when the alloy is used in service because the temperature of the intended use of this alloy is below the recrystallization temperature and recrystallization is very sluggish at these temperatures. Since the alloy of this invention is an austenitic alloy its mechanical properties can also be increased through cold working as is well known in the art.

Reference may be had to the graph ofFig. 2, the curves of which illustrate the eifect of aging at temperatures of 1100 F. to 1600" F. for 16 hours on the hardness of 10 alloy R-46 after solution heat treatment at different temperatures. Curve 20 shows the effect of aging at the diiferent temperatures on the hardness of alloy R-46 after a solution heat treatment at 1750 F. for one hour followed by water quenching. It is apparent from curve that there is little change in the hardness with respect to the difierent aging temperatures when the low solution heat treatment temperature of 1750 F. is employed. On the other hand, curve 22 of Fig. 2 shows the efiect of the different aging temperatures on the hardness after 0 a solution heat treatment at 1900 F. for one hour followed by a water quenching therefrom and clearly illustrates the substantial increase in the hardness as the aging temperature increases from 1100 F. to 1500 F. A similar result is also shown by curve 24 which shows the 25 effect of the ditferent aging temperatures after a solution heat treatment at 2050 F. for one hour followed by water quenching. The high solution heat treatment temperature of 2050 F. coupled with an aging temperature above 1400 F. causes the alloy to tend to overage and lose some of its hardness and strength. Thus if the higher solution heat treatment temperatures are used, it is desirable to limit the aging temperatures to about 1400 F. in order to obtain the desired optimum between hardness, strength and ductility.

The most significant and beneficial eflect which may be accomplished by way of heat treatment is illustrated by reference to Table V showing the efiect of the heat treatment on the average tensile and stress-rupture properties of the alloy of this invention.

TABLE V Efiect of heat treatment 0.02% Tensile Properties Brineil Tensile Offset El. Red. of

Treatment Hard- Strength Yield (percent) Area ness (p.s.l.) Strength (percent) (p.s.i.)

STRESS RUP'IURE PROPERTIES (l,200

Notch Smooth El. Red. of

Bar Bar (Per- Area (Hrs) (Hrs) cent) (Fercent) A 121 74 14 25 B 36 77 8 l4 From the test results given in Table V it can be seen that the higher solution heat treatment temperature is eifective to produce slightly higher tensile and yield strengths than the lower solution heat treatment temperature of 1900 F. However the lower solution heat treatment temperature, that is, 1900 F., produces a remarkable increase in the time to produce rupture in the notched bar as compared to the smooth bar. It is apparent that where the considerations in the engineering application of alloy of this invention demand a high notched rupture life, the lower solution heat treatment temperature #of 1900 F. is preferred. On the other hand, if the considerations are only the tensile properties and hardness, then the higher solution heat treatment of 2050 F. is preferred since the resulting alloy will possess higher hardness and slightly higher tensile and yield strengths without a detrimental eifect on the ductility of the alloy of this invention. For certain applications requiring very high hardness, tensile and yield strengths, the alloy may be cold worked up to 50%. This is usually done in the solutionheat treated condition and may or may not be followed by aging. Hot-cold working in the temperature range of 1000 F. to 1500 F. may also be used to improve properties.

In order to more clearly demonstrate the advantages of this invention, reference may be had to the following. An alloy of this invention of the analysis given hereinbefore, for example that given for alloy D-l83, was made and fabricated into a semi-finished mill product, for example, bar stock. This bar stock was then hot formed to the general shape of the component and heat treated by solution heat treating at a temperature of 1900 F. for one hour followed by a quenching in water and aging at about 1300 F. for about 16 hours and thereafter air cooled. The component, such as a turbine bucket, was then machined to final tolerances. The resulting turbine bucket possessed the desired strength and hardness and at the same time had a high notch rupture ductility insuring long life in operation as a component of a gas turbine.

There are no special techniques nor equipment needed to produce the alloy of this invention. It is significant that there is a noticeable absence of highly strategic alloying elements. In this respect the economies of the alloy in its heat treated form indicate a highly desirable product.

We claim:

1. A precipitation hardened austenitic iron base alloy suitable for use at high stress levels and temperatures of up to 1350 F. and consisting of from 0.20% to 0.35% carbon, from 0.15% to 0.75% silicon, from 16.0% to 20.0% manganese, from 11.5% to 13.5% chromium, from 2.0% to 4.0% molybdenum, from 0.60% to 0.95% vanadium, from 0.10% to 0.25% nitrogen and the balance iron with incidental impurities, said alloy having been heat treated by a solution heat treatment at a temperature in the range between 1800 F. and 2100 F. for about one hour and aged at a temperature in the range between 1200 F. and 1500 F. for a time period of from 10 to 20 hours.

2. A precipitation hardened austenitic iron base alloy suitable for use at high stress levels and temperatures of up to 1350 F. 'and consisting of from 0.25 to 0.30% carbon, from 0.15% to 0.50% silicon, from 17.0% to 19.0% manganese, from 12.0% to 13.0% chromium, from 2.5% to 3.25% molybdenum, from 0.7% to 0.9% vanadium, from 0.15% to 0.22% nitrogen and the balance iron with incidental impurities, said alloy having been heat treated by a solution heat treatment at a temperature in the range between 1900 F. and 2050 F. for about one hour and aged at a temperature in the range between 1300 F. and 1500 F. for a time period of about 16 hours.

3. A precipitation hardened iron base alloy suitable for use at high stress levels and temperatures of up to 1350 F. and consisting of about 0.35% carbon, about 0.21% silicon, about 17.2% manganese, about 12.25% chromium, about 3.12% molybdenum, about 0.86% vanadium, and about 0.19% nitrogen and the balance iron with incidental impurities, said alloy having been heat treated by a solution heat treatment at a temperature of about 1900 F. for about one hour, water quenched, aged at a temperature of about 1300 F. for a time period of about 16 hours and air cooled.

4. An article of manufacture suitable for use at temper-atures of up to 1350 F. at high stress levels and in the presence of combustion products of fuels, the article comprising an austenitic precipitation hardening iron base heat treated alloy containing 0.20% to 0.35% carbon,

0.15% to 0.75 silicon, 16.0% to 20.0% manganese, 11.5% to 13.5% chromium, 2.0% to 4.0% molybdenum, 0.60% to 0.95% vanadium, 0.10% to 0.25% nitrogen and the balance iron with incidental impurities, the article having been heat treated by a solution heat treatment at a temperature in the range between 1800 F. and 2100 F. and aged at a temperature in the range between 1200 F. and 1500 F. for a time period between 10 and 20 hours.

5. A precipitation hardenable austenitic iron base alloy consisting of 0.30% to 0.35% carbon, 0.20% to 0.75% silicon, 16.0% to 20.0% manganese, 11.5% to 13.5% chromium, 2.0% to 4.0% molybdenum, 0.60% to 0.95% vanadium, 0.10% to 0.25 nitrogen and the balance iron with incidental impurities.

10 References Cited in the file of this patent UNITED STATES PATENTS 2,562,854 Binder July 31, 1951 2,686,116 Schemph Aug. 10, 1954- 2,698,785 Jennings Jan. 4, 1955 2,711,959 De Long et a1. June 28, 1955 FOREIGN PATENTS 152,291 Austria Jan. 25, 1938 884,908 France Aug. 31, 1943 OTHER REFERENCES Stainless Iron and Steel, vol. 1, pages 138 and 139. Published in 1951 by Chapman and Hall, London, England.

Claims (1)

1. 5. A PRECIPITATION HARDENABLE AUSTENITIC IRON BASE ALLOY CONSISTING OF 0.03% TO 0.35% CARBON, 0.20% TO 0.75% SILICON, 16.0% TO 20.0% MANGANESE, 11.5% TO 13.5% CHRONIUM, 2.0% TO 4.0% MOLYBDENUM, 0.60% TO 0.95% VANADIUM, 3.10% TO 0.25% NITROGEN AND THE BALANCE IRON WITH INCIDENTAL IMPURITIES.

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