US3925071A

US3925071A - Heat resistant alloys

Info

Publication number: US3925071A
Application number: US730226A
Authority: US
Inventors: Subrata Ghosh
Original assignee: Chrysler Corp
Current assignee: Old Carco LLC
Priority date: 1968-05-20
Filing date: 1968-05-20
Publication date: 1975-12-09
Anticipated expiration: 1992-12-09

Abstract

The development of a new family of alloys, having exceptional stability and high strength at temperatures beyond three-fourths of their absolute melting points, has been made possible by the discovery of an entirely new mechanism wherein supersaturation of the matrix occurs during heating as opposed to supersaturation during cooling, such as with conventional age hardening type alloys.

Description

United States Patent Ghosh 1 Dec. 9, 1975 [5 HEAT RESISTANT ALLOYS 3,174,851 3/1965 Buel'llfil' 75/170 [75] Inventor: Subrata Ghosh, Southfield, Mich. OTHER PUBLICATIONS [73 Assignee; Chryskr corporafion, Highland Metals Handbook, 1948 Edition, American Society park, i for Metals, pp. 1158, 1163, 1164, 1206, 1220, 1221, 1243, and 1245. [22] Filed: May 20, I968 [21] Appl. No.: 730,226 Primary Examir1erC. Lovell Attorney, Agent, or FirmTalburtt & Baldwin 52 U.S.Cl. 7 7 4 2; 148 32. l 1 5/1 0 1 8/3 72/1751 [57] ABSTRACT 51 int. Cl. C22c 19/00 The development of a new family of y having 53 Fi f Search 0 75/170, 17], 135, 147, ceptional stability and high strength at temperatures 75/122; 148/32.5, 32 beyond three-fourths of their absolute melting points, has been made possible by the discovery of an entirely [56] References cu new mechanism wherein supersaturation of the matrix UNITED STATES PATENTS occurs during heating as opposed to supersaturation during cooling, such as with conventional age harden- 2,570,193 10/1951 B|eber..... ing type alloys 2,910,356 10/1959 Grala 3,021,211 2/1962 Flinn 75/170 1 Claim, 8 Drawing Figures U.S. Patent Dec. 9, 1975 Sheet 1 of 3 K-FL 11/1302 l NIL/7!.

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IT TOW/145') US. Patent Dec.9, 1975 Sheet 2 of3 3,925,071

U.S. Patent Dec. 9, 1975 Sheet 3 of3 3,925,071

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HEAT RESISTANT ALLOYS BACKGROUND OF THE INVENTION The advent of the gas turbines and space exploration vehicles has imposed an ever-increasing demand for the development of superior high temperature resistant alloys for prolonged use as structural members. To meet this demand, metallurgists have constantly improved on their earlier achievements by new discoveries and innovations based on the fundamental principles of alloy strengthening in view of heat resistance.

The fundamental definition of heat resistance in the solid state is the ability to resist plastic deformation and failure under the action of stresses at high temperatures over a period of time. Plastic deformation, or dislocation glide, is inhibited by inherent structural means which interfere with the mobility of dislocations. To provide for these structural aspects, the metallurgists have taken recourse to the methods of: (1) solid solution, wherein the base metal is alloyed with other elements up to its saturation limit; (2) dispersion, wherein a second insoluble phase is uniformly dispersed within a pure metal or a solid solution matrix; (3) carbide for mation, wherein highly stable carbide structures are formed by the alloy components and distributed within the matrix in critical form and areas; and (4) precipitation, wherein a supersaturated solid solution is made, under controlled conditions, to reject the excess solutes in the form of second phase precipitates within the matrix in critical form and areas. The objective of providing interference to dislocation mobility may be achieved by either one or careful combination of the above methods, depending on the specific requirements.

While these heretofore employed strengthening mechanisms have been used both singularly and in combination to provide the numerous high temperature alloys commercially available today, each is accompanied by limitations which resulted in the usetemperature of the alloy being restricted within onehalf to three-fourths of its absolute melting temperature (0.5-0.75 Tm). Thus, solid solution strengthening is limited to a narrow group of element combinations and their solid solubility ranges with respect to a particular base element. Dispersion strengthening, on the other hand, is difiicult to achieve due first to the lack of wide compatibility ranges between metallic matrices and basically non-metallic dispersion phases and secondly, due to the need for complicated powder metallurgy or expensive electro-deposition processes to introduce the dispersion phase. Carbide strengthening is limited to the increasing degree of brittleness introduced with higher percentages of carbides, the limited stability of the particular carbides formed in the alloy, as well as the compatibility ranges of the carbide structure with the matrix of the alloy.

One of the most important mechanisms employed by the metallurgist is precipitation hardening. The atoms of a pure, solid, crystalline element arrange themselves in a specific order in space which give rise to what is known as their unit cell. This basic cell structure may accept other atoms either in place of or in-between its own atoms resulting in a solid solution. Whether this acceptance of foreign" atoms takes place at random or to a certain limit, depends basically on the types of the host" and foreign atoms and the temperature. Usually, in combinations other than a stoichiometric chemical compound, immiscible mixtures or a continuous solid solution, this limit of acceptance or solid solubility decreases as the temperature is lowered below the solidus temperature. Thus, at a reduced temperature the system becomes supersaturated with the foreign solute atoms and assumes a state of lower free energy by rejecting some of these atoms from its structure and consequently becomes more stable. The rejected atoms, in turn, form a structural group of their own, with or without some of the host atoms, and gradually, through a step by step transitional atomic configura tion, appear in the matrix or host structure as second phase precipitates. The precipitates act as anchors to the dislocations and limit their mobility in the structure. However, due to the very fact that the precipitates are more stable at lower temperature and begin to dis solve in the matrix as their solid solubility increases with temperature, the anchoring of the dislocation gradually diminishes in strength and begin to yield to their advancing glide under stresses. Thus, the strengthening effects of the precipitates are eventually lost above the temperature where their diminishing size and number cannot cope with the increasing number of dis locations at stress leading to a yielding matrix. Consequently, all precipitation or age-hardening type superalloys are usable only up to a temperature range of 50-75% of their absolute melting points (0.5-0.75 Tm).

Heretofore, the approach which has been used to provide high temperature alloys is the production of alloy structures which are extremely hard and strong at low temperatures so that despite substantial reduction at elevated temperatures, the alloy is expected to retain sufficient strength for its intended use. Consequently, these alloys do not readily lend themselves to machin ing and other mechanical working at normal temperatures.

The present invention concentrates upon this basic weakness in the heretofore designed alloys for use at high temperatures and relates to alloys and method for their production which would provide superior high strength properties at temperatures up to of their absolute melting point (0.90 Tm).

SU MMARY An object of the present invention, therefore, is to provide a new mechanism for developing resistive strength in an alloy as the temperature of the alloy increases.

A further object of this invention is to provide new alloys having superior strength at temperatures in the range of from about one-half up to about ninety percent of their absolute melting point (0.9 Tm).

An additional object is to provide relatively inexpensive casting type alloys of low density which are remeltable without sacrificing desirable properties and which are composed of a minimum amount of costly or heavy elements.

Yet another object is to provide new high temperature alloys having a strengthening phase which is effective at temperatures close to the alloy's melting point, but which disappears on cooling so that the alloy can be readily worked at normal temperatures.

Other objects and advantages of the present invention will become apparent from a further reading of the description and the appended claims.

The above and other objects and advantages of this invention are obtained by my discovery that elements 3 of a complex alloy may be chosen such that, within specified composition ranges, supersaturation of the matrix takes place during heating as opposed to super saturation during cooling as in the case of conventional precipitation hardening.

DESCRIPTION OF THE DRAWINGS FIG. I, shows a portion of the phase diagram of the binary nickel-aluminum system;

FIG. 2, is an enlarged view of a portion of the nickelaluminum phase diagram, the portion being indicated generally by the area designated as A" in FIG. I;

FIG. 3, illustrates a portion of the phase diagram of the binary copper aluminum system;

FIG. 4, illustrates a portion of the phase diagram of the binary aluminum-cobalt system;

FIG. 5, illustrates a portion of the phase diagram of the binary cobalt-tungsten system;

FIG. 6, illustrates a portion of the phase diagram of the binary copper-zinc system;

FIG. 7, illustrates the phase diagram at a temperature of [382 F. of the ternary nickel-aluminum-chromium system; and

FIG. 8, illustrates the phase diagram at a temperature of [832 F. of the ternary nickel-aluminum-chromium system.

DESCRIPTION OF THE INVENTION As mentioned above, heretofore known precipitation or age hardening type superalloys are only recommended for use at temperatures below about 75 percent of their absolute melting point. The reason for this being that the precipitates are stable at a temperature below the saturation limit, but gradually dissolve away as the temperature of the alloy increases. Also. the other strengthening sources begin to diminish gradually with the rise in temperature. Thus, it is seen that the major strength of the alloy at higher temperatures depends on the precipitates and those precipitates gradually become non-existent when higher temperatures are encountered.

In contrast with the above situation, I have now discovered that it is possible to develop alloys in which the dislocation anchoring precipitates will behave in an exactly opposite manner. Thus, at lower temperature, when there is no need for critical strengthening, the precipitates dissolve out; however, at higher temperatures. a reaction takes place within the aggregate of the host and alloying foreign atoms due to supersaturation of the matrix. This reaction results in the precipitation of an additional phase. Accordingly, the general softening effects of the high degree of heat energy is resisted by the appearance of precipitates at the increased temperature.

In utilizing the new strengthening mechanism of this invention to develop superior high temperature alloys having excellent strength characteristics up to about ninety percent of their melting point. resort is made first to the phase equilibrium diagrams from which an initial approximation of a base alloy composition may be made in which the alloy enters. upon being heated. into a solid state zone containing an additional phase. Thus. the method of the invention for producing alloys which exhibit superior strength properties at elevated temperatures comprises selecting first a base element and at least one alloying element which are capable of combining to form an alloy that undergoes a solid state to solid state phase transformation, and combining them to form an alloy which, upon heating, undergoes this transformation so as to form at least one additional phase at or below the temperature at which it will be used. It should be understood that this invention is not merely limited to alloy systems wherein the alloy composition is a single phase at room temperature and two phases at the desired use temperatures. It is only necessary in this invention that at least one additional solid state phase be formed at or near the desired use temperature. Accordingly, the strengthening mechanism of this invention is found, for example, in those alloys comprising two or more solid state phases at normal temperature and three or more solid state phases at use temperature, or alloys comprised of one phase when cooled and three or more solid state phases when heated.

Examples of binary systems wherein it is possible to combine elements according to the invention into a1- loys which form at least one additional phase upon heating include, but not limited to those shown in the following table. In this table, it will be seen that a binary alloy composed of about 72 to 73 atomic percent aluminum and 27 to 28 atomic percent cobalt will. when heated to a temperature in the range of about [832 to 2012 F. enter a phase region where an additional solid state phase will precipitate. Likewise, it is seen in the magnesium-zinc binary system that there are three distinct composition ranges in any one of which an alloy can be formed in which precipitation will occur upon the alloy being heated to a temperature in the range of about 572-662 F. It should be understood that the composition and temperature values shown in the following table are of a representative value and may de viate somewhat from actual limits.

TABLE A Atomic Binary System Percent Temperature Range Elements of F. of

A B Element B Precipitation Aluminum Cobalt 27-28 1832-2012 Aluminum Iron 25-26 1832-2120 Aluminum Magnesium 44-45 572-734 Aluminum Titanium 50-51 2372-2660 Aluminum Tungsten 22-26 1832-2372 Antimony Tin 59.5- 212-464 Bismuth Lead 67-77 86-360 Cadmium Lidtium 16-25 302-626 Cadmium Nickel 17-21 752-932 Cerium Thorium 0.1-15 1238-1526 Chromium Tantalum 34-36 2912-3542 Cobalt Antimony 64-67 1 1 12-1652 Cobalt Osmium [-35 752-2822 Cobalt Rhenium [-25 752-2732 Cobalt Ruthenium [-33 842-2552 Cobalt Tungsten 7-13 1292-1922 Cobalt Vanadium 24-31 1922- l 958 Copper Aluminum l6-19.6 1040-2066 Copper Antimony 15-20 752-1022 Copper Cadmium 41.5-44.5 572-1022 Copper Gallium 16.3-18.6 1 148-1670 Copper Germanium 12-13 1 112-1472 Copper Magnesium 32-34 932-1292 Copper Silicon 17.5-18.5 752-1292 Copper Zinc 31.9-38.3 842- 1652 Gold Cadmium 27-34 932-1148 Gold Indium 15-21 932-1202 Gold Manganese 46-52 1022-1130 Indium Lead 27-31 266-338 Indium Thallium 43-58 68-302 Indium Tin 13-27 167-248 Iron Columbium 34-36 1832-2372 Iron Nitrogen 5.7-6.1 842-1202 Iron Tungsten 62-63 1832-2000 lron Tungsten 69-70 2600-2800 Iron Vanadium 43-52 2012-2246 Iron Zinc 69-709 1238-1436 Magnesium Cadmium 20-35 212-356 Magnesium Cobalt [-15 392-2012 Magnesium Indium 23-25 482-644 TABLE A-continued Atomic Binary System Percent Temperature Range Elements of F. of

A B Element B Precipitation Magnesium Lithium 17-175 572-1076 Magnesium Zinc 49.5-50.5 572-662 Magnesium Zinc 595-605 572-662 Magnesium Zinc 84-85 572-662 Manganese Nickel 47-55 1238-1400 Mercury Lead 66-69 122-302 Nickel Gallium 40-42 1112-1382 Nickel Germanium 23-25 1832-2102 Nickel Molybdenum 28-29 1472-1652 Nickel Molybdenum 34-35 1472-1652 Nickel Platinum 55-70 752-1 1 l2 Nickel Silicon 22-23 1652-2012 Nickel Vanadium 22-27 1 1 12-1832 Nickel Zinc 45.5-51.5 1256-1472 Palladium Lead 38-41 1346-1526 Palladium Zinc 32-34 1202-1292 Palladium Zinc 37-56 1 1 12-2102 Rhodium Tin 38-405 1652-4352 Silver Aluminum 21-26 392-842 Silver Aluminum 24.5-32.5 1 130-1328 Silver Cadmium 384-40 1 1 12-1360 Silver Magnesium 74-78 392-878 Silver Mercury 44-45 32-572 Silver Platinum -18 752-1742 Silver Tin 12-13 392-1292 Silver Tin 23.7-25.5 572-896 Silver Zinc 32-40 572-1292 Silver Zinc 58.5-61 752-1220 Silver Zinc 67-69 932-1 166 Thallium Cerium 55-56 1112-1652 Titanium Tungsten 0.2-9.0 1328-1616 Titanium Uranium 59-67 1292-1634 Cerium Yttrium l Yttrium 1) Lanthanum Yttrium 1 Magnesium Yttrium l) Thorium Zinc Cerium 10-1 3.5 932-1472 Zinc Cobalt 7-8 842-1040 Zinc Lithium 13-23 149-464 Zinc Magnesium 13-195 680-752 Zinc Magnesium 23-27 500-610 1 Existence of Inverse Precipitation Reaction in these Systems is Established, but Composition and Temperature are not well defined.

From the preceding table, it is seen that the mechanism of this invention is applicable to relatively high melting alloys such as iron-tungsten (2600-2800F.), medium temperature melting alloys as copper-zinc (840-l650F.), and low temperature systems such as bismuth-lead (86-360F.). Thus, it will be understood that this invention is not limited to alloys of high melting temperature, but instead is concerned with strengthening of the alloy at temperatures close to the melting point of the alloy, whatever that may be. It will likewise be understood that alloys based on this invention may very well be a binary alloy, but more often than not further alloying will be necessary to achieve the desired temperature strength.

From the foregoing, it is seen that this invention is applicable to many and varied alloy systems and, accordingly, for purposes of clarity in the ensuing discussion of the invention, reference will be made only to the nickel-aluminum binary system. P10. 1, illustrates a portion of the nickel-aluminum binary system containing from about 50 to 100 weight percent nickel, the balance being aluminum. It will be noted that the alloys containing from approximately 83.0 to 85.5 weight percent nickel and 17 to 14.5 weight percent aluminum are composed of a single solid state phase which is the intermetallic compound Ni Al at room temperature. However, when heated, a solid state to solid state transformation reaction occurs and the alloy moves into a binary phase. This can be seen more clearly with reference to FIG. 2, wherein the single solid state phase region of Ni Al is hatched. Line 1 denotes the left hand boundry of this single phase region and it is seen that if an alloy composition such as A is heated in excess of about 2200F. the transformation line 1 is crossed and the alloy moves in a zone where an additional solid state phase of NiAl is found. This means that the alloy which is essentially percent Ni,-,Al at lower temperatures will begin to precipitate NiAl at a temperature, which will vary with the composition of the alloy. Likewise, it is seen that alloy compositions such as B and C when heated to approximately 2500F. will cross transformation line 2 in which event an additional solid state phase (nickel solid solution, 7 will commence to precipitate. Accordingly, it is seen that by careful selection of elements, it is possible to form an alloy which resists softening at elevated temperatures by virtue of bringing into the alloy structure new dislocation anchoring precipitates. Of course, this may not mean that the alloy becomes stronger than when it is at lower temperature; however, its heat resistance in comparison with heretofore known alloys is greatly increased since, in contrast to conventional precipitation hardened alloys wherein the precipitates dissolve on heating, here the precipitates are actually formed on heating.

As mentioned earlier, the Inverse Precipitation" mechanism of this invention can be used in many and varied alloy systems, By way of illustration reference is made to FIG. 3, showing a portion of the copperaluminum binary system in which it is seen that an alloy having a composition in the range of about 7.5 weight percent aluminum-92.5 weight percent copper (composition C) to about 9.8 weight percent aluminum-90.2 weight percent copper (composition D) will pass through line 3 upon heating. and undergo a transition from the single solid state a phase to the binary solid state a+B phase in which 5 will be a precipitate.

FIG. 4, illustrates a portion of the aluminum-cobalt binary phase system. In this system, alloy compositions of from about 46.6 to 48 weight percent cobalt-52 to 53.4 weight percent aluminum will, upon being heated to a temperature of about 2020F., pass through the transformation line 4 and change, as seen with respect to composition E, from the single solid 5 phase to the binary 8+ 6 in which 6 will form the precipitate.

FIG. 5, illustrates a portion of the cobalt-tungsten binary system and it is seen that alloys having a composition in the range of about 44 weight percent tungsten- 56 weight percent cobalt (composition F) to about 48 weight percent tungsten-52 weight percent cobalt (composition G) will, when heated to a temperature higher than the transformation lines 5 or 5' undergo a transition from the single solid 1 phase to a binary phase of 7 plus B or y plus 8 in which B or 8 respectively, will form the precipitate. FIG. 6 illustrates a portion of the copper-zinc phase system. Here, again alloys having compositions in the range of about 32.5 weight percent zinc-67.5 weight percent copper (composition H) to about 38 weight percent zinc-62 weight percent copper (composition J) will cross the transformation line 6 and enter the solid phase B B region wherein B will form the precipitate. These illustrations indicate the feasibility range of the lnverse Precipitation mechanism in the particular alloy system and development of a particular alloy can be achieved with routine experimentation.

The temperature at which the additional solid phase formation begins will, as seen from the drawings, depend on the composition of the alloy. For most applications, it is preferable to prepare an alloy composition in which the additional solid state phase appears when the alloy is heated to a temperature in the range of from about A to of its absolute melting point (Tm). Absolute melting point being the melting temperature plus 460 in the Fahrenheit scale (Rankine) or plus 273 in the Centigrade scale (Kelvin).

It is apparent from the drawings that many of the alloy composition ranges in which it was possible to obtain high temperature precipitation are rather narrow if only a binary system is considered. Although it is true that a binary alloy, prepared according to this mechanism may be suitable for certain high temperature application, the critical effects of adding further alloying elements, particularly for expansion of the high temperature precipitation zone as well as additional strengthening is a logical step to consider. Accordingly, as in conventional alloys, the use of additional elements for alloying is employed. It was found that through proper selection of the alloying elements it is possible to control the phase regions so as to increase the range of alloy compositions in which the high temperature precipitation mechanism of this invention can be em ployed, as well as, alter the transformation temperatures for a particular base composition. For example, FIG. 7 illustrates the ternary nickel-aluminumchromium phase diagram at a temperature of approximately l382 F. it will be noted that the single phase region of Ni Al is considerably larger than in the binary system of FIGS. 1 and 2 up to a critical amount of chromium addition. Accordingly, this means that the range of alloy compositions which will undergo transition, upon heating, to form an additional solid state phase is considerably increased. Thus, an alloy of composition X in FIG. 7 is in the lOO percent Ni Al phase region at l382 F. However, when the temperature of the ternary alloy X of FIG. 7 is raised to about 1832 F. the phase regions shift and the composition X is now in the binary phase region of Ni Al and NiAl as seen in FIG. 8. This means that at this temperature NiAl precipitates will appear in the alloy structure. Similarly, alloys of Y" and Z compositions will precipitate y and a B respectively as the higher temperature is reached.

Based on the concept as outlined above, a series of high temperature superalloys containing up to 20 elements and comprising about 7 to 17 weight percent aluminum and about 65 to 85.5 weight percent nickel have been developed starting with a binary nickelaluminum base composition which is then alloyed with other suitable elements. The preferred elements for use in alloying quantities with a base nickel-aluminum system are chromium, niobium, carbon, titanium, cobalt, molybdenum, tungsten, tantalum, boron, silicon, vanadium, beryllum, nitorgen, rare earths, yttrium, zirconium, copper, hafnium, rhenium, oxygen, manganese and iron. These alloying elements in addition to widening the composition zone where the high temperature precipitation mechanism of this invention is attainable, strengthen the alloy in conventional fashion such as by solid solution, conventional precipitation, carbide strengthening, etc. Accordingly, the alloys provided by this invention derive their high strength from all the conventional strengthening mechanisms and, in addition, offer exceptional strength at temperatures close to their melting points by means of high temperature precipitation. In other words, the strengthening at lower temperatures is derived from the conventional sources of alloy strengthening, and at temperatures when such conventional sources tend to become exhausted, the high temperature precipitation takes place to maintain their strength and further their use temperature beyond three-fourth (0.75 Tm) of their melting point. Thus, the new strengthening mechanism of this invention supplements and does not merely replace conventional strengthening mechanisms.

It should also be recognized that the high temperature precipitation mechanism of this invention makes it feasible to obtain a stable dispersion phase in a casting type alloy. For example, in the nickel-aluminum base alloys, oxygen can be reacted with the NiAl precipitate to form a stable dispersion of aluminum oxide according to the equation: 6NiAl (ppt) 30 2Ni Al 2Al 0 The oxygen can be supplied from dissolved oxygen in the alloy or by diffusion from the outside.

The following examples illustrate alloys prepared in accordance with this invention:

EXAMPLE 1 An alloy according to the present invention was prepared having the following nominal analysis:

Nickel 79.4 weight percent Aluminum 134 weight percent Chromium 7.2 weight percent This alloy was prepared by simultaneously charging all the elements into an induction furnace enclosed within a vacuum chamber, and then evacuating the chamber to produce pressure in the furnace of approximately 1-10 microns. The furnace was then heated and the charge melted, forming a molten mass having a temperature between 25002 600 F. The heating was continued until the melt reached a temperature of 3200 F. so as to ensure homogenization. After about 5 minutes at this temperature, the melt was cooled to about 3050 F. and poured into a pre-heated mold located inside the vacuum chamber so as to form test bars. The test bar castings were withdrawn from the vacuum chamber about eight minutes after casting and allowed to cool in air. The test bars were approximately 4 inches long and /2 inch in diameter at the end portions with a reduced A inch diameter center neck portion of l-V2 inches.

It was determined that from room temperature to approximately 1560 F. the alloy was the single Ni Al phase and that above 1560 F. NiAl began to precipitate. At about 2l00 F. the amount of NiAl reached about 15 to 20 percent of the matrix. The alloy had a theoretical density of 0.24 pounds per cubic inch and a Rockwell C (R hardness of 32-35.

The following data were obtained through tests conducted at room temperatures and elevated temperatures on the above prepared alloy.

Stress Rupture Life At: l500F/30,000 psi l 1 Hours 3.0% Elongation EXAMPLE 2 An alloy of the following nominal analysis was prepared as set forth in Example I,

Nickel Aluminum Chromium Titanium 74.8 weight percent 1 L weight percent i2.5 weight percent 1.7 weight percent This alloy had a theoretical density of 0.260 pounds per This alloy had a theoretical density of 0.278 pounds per cubic inch and a Rockwell hardness R of 30-35.

The following data were obtained through tests conducted on the above prepared alloy.

cubic inch and a Rockwell hardness R of 32-35.

Stress Rupture Life At'.

l500F/30.000 psi 59 Hours 6% Elongation Stress Rupture Life At:

i100F/20.00o si meow/20.000 psi 1800F/l5,000 psi wow/10.000 psi TABLE 4 Temp. Ult. 02% 7( "F Tensile Stress Yield Elongation 70 l 18 I02 7 i800 67 64 3 470 Hours-7 0% Elongation 62 Hours-2.0% Elongation 2i] Hours-5.0% Elongation l36 Hours-2.0% Elongation EXAMPLE 5 An alloy of the following nominal analysis was pre 0 pared as set forth in Example 1:

Nickel Aluminum Chromium Titanium Cobalt Molybdenum Tantalum Niobium Tungsten Carbon Boron Rare Earths Zirconium 70.42 weight percent 7.82 weight percent 9.27 weight percent L82 weight percent 3.47 weight percent 2.93 weight percent 1.92 weight percent 0.25 weight percent 0.50 weight percent 0.50 weight percent 005 weight percent 0.05 weight percent L00 weight percent This alloy had a theoretical density of 0.275 pounds per cubic inch and a Rockwell hardness R of 28-36. Melting of this alloy began at 2300F. (2760 in Absolute Rankine Scale). No deterioration in properties was noted in this alloy on remelted and re-cast condition. The following data were obtained through tests conducted on the above prepared alloy.

EXAMPLE 3 An alloy of the following nominal analysis was prepared as set forth in EXAMPLE 1:

Nickel 7 l.l5 weight percent Aluminum 9.50 weight percent Chromium 9.65 weight percent Titanium 2.54 weight percent Cobalt 3.13 weight percent Molybdenum 3.05 weight percent Tantalum 0.95 weight percent Boron .03 weight percent This alloy had a theoretical density of 0.269 pounds per cubic inch and a Rockwell hardness R of 30-35.

TABLE 3 Temp. Ult. 0.2%

"F Tensile Stress Yield Elongation 70 I23 I l 1 3 I800 65 63 2 Stress Rupture Life At:

l500F/25,000 psi 1837 Hours 2% Elongation l500F/30,000 psi 1000 Hours 2% Elongation l500F/40,000 psi 410 Hours 1% Elongation l700F/20,000 psi 109 Hours 1% Elongation EXAMPLE 4 An alloy of the following nominal analysis was prepared as set forth in Example 1:

Nickel 71.30 weight percent Aluminum 7.93 weight percent Chromium 9.38 weight percent Titanium 1.85 weight percent Cobalt 3.5l weight percent Molybdenum 2.97 weight percent Tantalum 1.94 weight percent Niobium 0.25 weight percent Tungsten 0.50 weight percent Carbon 0.30 weight percent Boron 0.05 weight percent Rare Earths 0.05 weight percent Stress Rupture Life At:

1800F/20,000 si l800F/l5.000 psi 2000Fl 7,500 psi TABLE 5 Temp. Ult. 0.2%

F Tensile Stress Yield Elongation I37 Hours-7.0% Elongation 548 Hours-7.0% Elongation 326 Hours-4.0% Elongation The development of the alloys in the foregoing exam- The validity of the concept that an equilibrium phase transformation in the solid state at high temperatures may be utilized to strengthen an alloy for high temperature applications is, therefore, proven through these examples. As noted earlier, the applications of this concept may be wide and varied, depending on the requirement for a particular use, for alloys of different base elements. in any event, if the alloying elements are 1 l chosen as outlined in the foregoing. so that high temperature precipitation is involved, the alloy will outperform other high temperature alloys of similar compositions but based only on the conventional strengthening mechanisms.

According to the foregoing it will be apparent that the objectives of this invention, namely to develop a superior, remeltable, inexpensive casting type high temperature resistant alloy with temperature capability up to about 0.90 Tm, but containing only a small quantity of expensive heavy elements, have been achieved. in a direct comparison of properties and attributes of the alloy of Example 5, it was found that the cost of raw materials and manufacture of one pound of this alloy will be the lowest of all heretofore known high temperature superalloys; its density will be the lowest of all heretofore known high temperature superalloys, and pound per pound, its strength will be superior to the best of the heretofore known commercial superalloys. Its temperature capability will be about 0.90 Tm as against about 0.75 Tm for commercial alloys.

I Claim:

1. A temperature resistant metallic alloy consisting essentially of about 65.0 to 85.5 weight per cent nickel and about 7 to 17 weight per cent aluminum which combines to form an alloy composition, said alloy composition having a matrix which is a stable solid solution when at a temperature less than about half of the absolute melting point of said alloy, said matrix when at a temperature above about half and within two-thirds of the absolute melting point of said alloy changing from said solid solution to a supersaturated solid solution such that said matrix due to said supersaturation above said temperature of about half of said absolute melting point forms at least one additional solid state phase in precipitate form to provide strengthening of the alloy matrix above said temperature of about half of said absolute melting point, said alloy including a stable aluminum oxide dispersion phase which will remain upon cooling, said dispersion phase being formed by reacting gaseous oxygen with said additional solid state phase,

Claims

1. A TEMPERATURE RESISTANT METALLIC ALLOY CONSISTING ESSENTIALLY OF ABOUT 65.0 TO 85.5 WEIGHT PER CENT NICKEL AND ABOUT 7 TO 17 WEIGHT PER CENT ALUMINUM WHICH COMBINES TO FORM AN ALLOY COMPOSITION, SAID ALLOY COMPOSITION HAVING A MATRIX WHICH IS A STABLE SOLID SOLUTION WHEN AT A TEMPERATURE LESS THAN ABOUT HALF OF THE ABSOLUTE MELTING POINT OF SAID ALLOY, SAID MATRIX WHEN AT A TEMPERATURE ABOVE ABOUT HALF AND WITHIN TWO-THIRDS OF THE ABSOLUTE MELTING POINT OF SAID ALLOY CHANGING FROM SAID SOLID SOLUTION TO A SUPERSATURATED SOLID SOLUTION SUCH THAT SAID MATRIX DUE TO SAID SUPERSATURATUON ABOVE SAID TEMPERATURE OF ABOUT HALF OF SAID ABSOLUTE MELTING POINT FORMS AT LEAST ONE ADDITIONAL SOLID STATE PHASE IN PRECIPITATE FORM TO PROVIDE STRENGTHENING OF THE ALLOY MATRIX ABOVE SAID TEMPERATURE OF ABOUT HALF OF SAID ABSOLUTEE MELTING POINT, SAID ALLOY INCLUDING A STABLE ALUMINUM OXIDE DISPERSION PHASE WHICH WILL REMAIN UPON COOLING, SAID DISPERSION PHASE BEING FORMED BY REACTING GASEOUS OXYGEN WITH SAID ADDITIONAL SOLID STAGE PHASE.