US3746581A - Zone annealing in dispersion strengthened materials - Google Patents

Zone annealing in dispersion strengthened materials Download PDF

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US3746581A
US3746581A US00221979A US3746581DA US3746581A US 3746581 A US3746581 A US 3746581A US 00221979 A US00221979 A US 00221979A US 3746581D A US3746581D A US 3746581DA US 3746581 A US3746581 A US 3746581A
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zone
alloy
temperature
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R Cairns
J Benjamin
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NAT NICKEL CO Inc
NATIONAL NICKEL CO INC US
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/052—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 40%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C32/00—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ
    • C22C32/001—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with only oxides
    • C22C32/0015—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with only oxides with only single oxides as main non-metallic constituents
    • C22C32/0026—Matrix based on Ni, Co, Cr or alloys thereof
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon

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  • a method for producing a hot worked high temperature heat resistant dispersion-strengthened superalloy characterized by a metallographic structure consisting essentially of large coarse grains having a preferred orientation relative to the axis of working, the method comprising subjecting a dispersion-strengthened alloy, e. g., superalloy, to zone annealing at an elevated germ 1nat1ve grain growth temperature to form coarse grams disposed in the working direction of the allloy and to achieve improved properties.
  • This invention relates to a powder metallurgy method for producing a preferred microstructure in a dispers onstrengthened heat resistant alloy, such as dispersionstrengthened age-hardenable superalloys and, in particular, to a method for producing hot worked dispersionstrengthened superalloy shapes from mechanically alloyed metal powder characterized by improved high temperature stress-rupture and creep properties.
  • the method comprises mixing a compressively deformable metallic powder with at least one other powdered material from the group consisting of a non-metallic material and another metallic material and dry milling the mixture under conditions of repeated mutual impact compression sufiiciently energetic to substantially reduce the thickness of at least the compressively deformable metallic constituents of the mixture and for a time sufiicient to produce non-pyrophoric Wrought composite particles which individually have substantially the composition of the mixture.
  • a dry charge of attritive elements e.g., nickel balls of plus 4 minus /2 inch average diameter
  • a powder mass of predetermined composition comprising a plurality of constituents, at least one of the constituents bemg a compressively deformable metal in an amount of at least by volume, with the remainder of the powder mass being at least one other constituent from the group consisting of a non-metal and another metal, the metals having a melting point of at least l000 K.
  • the volume ratio of the attritive elements to the powder mass is at least about 4:1 and, more advantageously, at least about 10:1.
  • the charge is then subjected to agitation milling, for example, in an attritor mill, under conditions in which a substantial portion of the attritive elements is maintained kinetically in a highly activated state of relative motion, whereby to cause the constituents to unite and form composite metal particles, the milling being continued until cold worked composite metal particles are produced characterized by markedly increased hardness (that is, the particles contain a substantial amount of stored energy) and further characterized by an internal structure in which the constituents are intimately inter- 3,746,581 Patented July 17, 1973 dispersed.
  • the particles which, in a preferred embodiment, are heavily cold worked to reach substantially the saturation hardness of the system involved, are subjected to a diffusion heat treatment, the intmiately interdispersed constituents diffuse one into the other rather quickly to produce a homogenized matrix.
  • the method is applicable to the production of dispersion-strengthened superalloys having a matrix composition normally very difiicult to produce by conventional powder metallurgy techniques, including alloys falling within the range of about 5% to 35% or even up to 60% chromium, about 0.5% to 6.5% aluminum, about 0.5% to 5.6% titanium, up to about 15% molybdenum, up to about 20% tungsten, up to about 10% columbium, up to about 10% tantalum, up to about 3% vanadium, up to about 2% manganese, up to about 2% silicon, up to about 0.75% carbon, up to about 0.1% boron, up to about 1% zirconium, up to about 0.2% magnesium, up to about 4% hafnium, up to about 35% iron, up to about 10% by volume of a refractory dispersoid, and the balance essentially nickel in an amount at least about 40% of the total composition.
  • cobalt can replace nickel. It
  • 131,761 by providing a confined batch of cold worked mechanically alloyed composite particles comprised of alloying constituents which when alloyed together provide a superalloy composition (which is preferably age-hardenable) and then hot working (e.g., hot extruding) the confined batch under correlated conditions of temperature, reduction ratio and strain rate such that when the resulting hot worked product is subsequently heated to an elevated germinative grain growth temperature, e.g., from about 2250 F. to below the incipient melting point of the alloy for a time up to about 4 hours or more, coarse grains are formed disposed in the working direction of the alloy.
  • an elevated germinative grain growth temperature e.g., from about 2250 F. to below the incipient melting point of the alloy for a time up to about 4 hours or more, coarse grains are formed disposed in the working direction of the alloy.
  • Another object is to provide a hot worked dispersionstrengthened superalloy shape having an improved metallurgical structure characterized by a substantially uniform distribution of coarse elongated grains disposed in the working direction of the alloy shape.
  • a further object is to provide a method of materially enhancing the mechanical properties of an extruded age-hardenable superalloy shape, such as age-hardenable dispersion-strengthened superalloys, by employing a relatively high temperature directional grain growth or annealing step.
  • Still another object is to provide a simple hot working and annealing or high temperature grain growth step for producing a dispersion strengthened age hardenable nickel-base superalloy shape having a substantially uniform distribution of coarse elongated grains across the cross section.
  • FIG. 1 is a plot on semi-logarithmic coordinates showing the relationship of extrusion ratio and extrusion temperature for hot working dispersion-strengthened nickelbase, age-hardenable superalloys, the area shown outlining the preferred combinations of extrusion ratio and extrusion temperature for subsequent application of the present invention
  • FIGS. 2A through 2D are reproductions of a series of photomacrographs taken at 2 times magnification comparing the effects of conventional annealing (FIG. 2A) and of zone annealling (FIGS. 2B, 2C and 2D) on the size and shape of coarse grains obtained in dispersionstrengthened superalloy bars; and
  • FIG. 3 is a photograph of a partially zone annealed bar in section to illustrate the effect of zone annealing upon an extruded superalloy bar.
  • the present invention is directed to a method for producing a hot worked dispersionstrengthened heat resistant alloy, e.g., superalloy, shape characterized by improved mechanical properties at elevated temperatures and by a metallographic structure consisting essentially of large coarse grains disposed in the direction of working of the hot worked shape.
  • a hot worked dispersionstrengthened heat resistant alloy e.g., superalloy
  • shape characterized by improved mechanical properties at elevated temperatures and by a metallographic structure consisting essentially of large coarse grains disposed in the direction of working of the hot worked shape.
  • One aspect of the invention resides in providing a batch of mechanically alloyed composite particles formed of constituents which, when alloyed together, provide a high temperature heat resistant alloy, e.g., an age-hardenable dispersion-strengthened superalloy, the composite particles being of substantially saturated hardness.
  • the foregoing composite particles are characterized metallographically by an internal structure comprising said constituents substantially intimately united and interdispersed.
  • a confined shape of the mechanically alloyed composite particles is hot worked in accordance with the invention at a temperature of over about 1600 F. and ranging up to about 2200 F. correlated to reduction ratios ranging broadly from over about 7 to about 35 (as more fully discussed hereinafter), and at a strain rate greater than a minimum value defined hereinafter such that when the resulting hot worked alloy is subsequently subjected to zone annealing at an elevated germinative grain growth or annealing temperature, elongated coarse grains are formed with the longitudinal axis thereof disposed in the working direction of the alloy shape.
  • zone annealing is meant the phenomenon by which coarse grains are formed by causing a hot zone at an elevated germinative grain growth temperature to traverse a hot worked shape from one end to the other in the direction of working so as to cause sequential growth of coarse elongated grains through the hot worked shape.
  • the alloy shape is one obtained by extruding a cylinder of 3.5 inches in diameter to a rod three-quarters of an inch in diameter
  • the coarse grains are elongated like fibers in the direction of working, that is, in the longitudinal direction of the rod.
  • the coarse grains may be plate-like in shape, the major axis of each grain being generally disposed in the direction of Working.
  • the foregoing method is unique in that the internal stored energy required for grain growth is introduced by the mechanical alloying process in producing the initial powder and by the thermomechanical processing conditions.
  • a single hot extrusion is sufiicient to effect the consolidation of a product capable of developing coarse grains at an elevated temperature to provide a metallographic structure characterized by large coarse grains elongated in the di- 4 rection of working.
  • zone annealing process is particularly advantageous over the grain growth technique described in copending application Ser. No. 131,761 in that the final product tends to exhibit additionally improved high temperature properties in combination with improved ductility as measured by percent elongation at failure of test specimens subjected to high temperature stress-rupture tests at various applied stresses.
  • grain growth is induced by a slowly moving hot zone traversing the hotworked consolidated shape, continuous growth of coarse grains in the direction of zone motion is generally favored. In this way, larger, more elongated grains are formed, and improved mechanical properties result.
  • the moving hot zone was produced by means of a resistance furnace travelling on a carriage surrounding a stationary ceramic tube which contained the hot worked specimens or bars to be treated.
  • the moving hot zone may be produced in other ways, such as by means of moving induction coils or other heating elements; or, alternatively the specimen itself may be moved through a hot zone of a stationary furnace.
  • the dispersoid added to the composition may be in the range of about 1% to about 5% or even up to about 10% by volume, e.g., nominally about 2.25 volume percent of particles preferably having average si-ze range of about to 500 or even about 1000 angstroms, e.g., about 300 angstroms.
  • the foregoing superallo'y is the dispersion-strengthened, hot extruded condition exhibits improved high temperature stress-rupture properties when it is preferably subjected to a grain coarsening heat treatment by zone annealing, at a temperature of at least about 2200 F. and preferably at a temperature of at least about 2300 F.
  • the alloy may be further heat treated and age-hardened.
  • the foregoing is achieved by advantageously controlling in combination the hot working reduction ratio (e.g., the extrusion ratio), the hot working or extrusion temperature and the hot working strain rate.
  • the coarse grains are characterized by preferred orientation after germinative grain growth treatment. In the case of an extruded rod-like product, there is an increase in grain size of at least 100 fold in the longest direction.
  • the reduction ratio is determined by the original cross section area of the shape before working divided by the cross section of the final product produced therefrom after working. For example, a shape or billet of 3.5 inches in diameter hot worked (e.g., hot rolled, hot press forged or hot extruded) to a final diameter of about five-eighths of an inch undergoes a change in cross section corresponding to a reduction ratio of about 31.4: 1.
  • the extrusion temperature i.e., the temperature to which the material is heated for extrusion, for uniform results should advantageously not be less than about 17-60 F. and may range up to about 2200 F.
  • An important advantage of the invention resides in the use of mechanically alloyed metal particles of substantially saturation hardness.
  • metal powder By using such metal powder in the process of the invention, large coarse elongated grains can be produced uniformly across substantially the whole cross section of the final product. This is an unexpected improvement, considering that in normal extrusion processes, the grain size after recrystallization may be different across the cross section due to strain gradients varying from a maximum at the outside surface of the hot worked product to a minimum at the center thereof.
  • thermomechanical energy component ranging from 1.793 to 2.250 (preferably at least 2.028)
  • a nickel-titanium-aluminum master alloy is first prepared by vacuum induction melting. The resulting ingot is crushed and ground to minus 200 mesh powder.
  • the powder (Powder A) contains 72.93% nickel, 16.72% titanium, 7.75% aluminum, 1.55% iron, 0.62% copper, 0.033% carbon, 0.050% A1 0 and 0.036 TiO About 14.9 weight percent of this powder is blended with 63.7% carbonyl nickel powder having a Fisher subsieve size of about 5 to 7 microns, 19.8% chromium powder having a particle size passing 100 mesh, 00.25% of a Ni-28% zirconium master alloy passing 200 mesh, 0 .04% of a Ni-17% boron master alloy passing 200 mesh and about 1.33% by Weight of yttria of particle size of about 350 A.
  • the microstructure of the particles making up the powder is characterized by nearly complete homogeneity, when viewed optically at 250 diameters, comprising each of the constituents substantially intimately united and dispersed.
  • This material it was found that increasing the time of milling at 132 r.p.m. from 20 to 40 hours markedly improves homogeneity of the mechanically alloyed powder to the point that fragments of the starting ingredients become practically indistinguishable upon optical examination at 250 diameters.
  • the structural homogeneity obtained about 20 hours milling at 182 r.p.m. is about the same as that obtained upon 40 hours milling at 132 r.p.m.
  • the coordinates at the intersections of the lines bounding the closed curve KLMNOK on FIG. 1 are as follows with extrusion ratio and extrusion temperature being given respectively in each case: K(7; 1600 F.); L(9.5; 1600 F.); M(35; 2020 F.); N(35; 2210 F.) and 0(7; 1700 F.).
  • results obtained with the heat treatment of the invention are generally superior to those obtained with the conventional heat treatment in two respects: the stress-rupture lives are higher and the corresponding ductility as determined by percent elongation or percent reduction in area at failure is greater.
  • the results in relation to Alloy No. 2 indicate that the procedure of the invention enabled raising the 100 hour rupture life at 1900 F. of this material from less than 15 k.s.i. as conventionally processed to greater than 17 'k.s.i. after zone annealing. Bars extruded under conditions near the line KON of FIG. 1 will be improved mainly in ductility as compared to conventionally annealed material.
  • a rupture life of 106 hours with an elongation of 16% and a reduction in area of 35% was determined at 1900 F. under an applied stress of 15 k.s.i.; and, at 1400 F., under an applied stress of 30 k.s.i., a rupture life of 1.5 hours with an elongation of 45% and a reduction in area of 62% resulted.
  • the throttle setting for the extrusions in the case of each of Alloys 1 to 4 was on the 750-ton press.
  • the moving hot zone was applied at a rate of 1.75 inches per hour.
  • the actual heating rate and time at temperature of the bar being heated will depend on the size of the bar being heated, the size of the hot zone, the temperature profile of the hot zone, and the rate of relative movement between the hot zone and the bar being heated (one or the other may move).
  • the required heating rate is easily obtainable by experimentation.
  • the heating of the extruded bar to the germinative grain growth temperature should not be so slow in the range of about 1900" F. to 2200 F. (referred to as the sub grain-growth energy annealing range), that stress or energy annealing occurs to the extent that the stored energy in the work being treated is substantially removed. Stress or energy annealing is relatively slow at temperatures up to about 1900 F.
  • the residence time during the heating up period above 1900 F. and below the grain growth temperature is important and should not be so long, e.g., in excess of 30 minutes, that the metal bar is substantially stress or energy annealed. When the residence time in the temperature region between 1900 F. and the grain growth temperature is too long, generally fine grains are obtained when the specimen involved is finally heated to the germinative grain growth temperature.
  • the heating rate through the temperature range below the germinative grain growth temperature should be sufficiently rapid to avoid substantial stress and energy annealing before the bar reaches the germinative grain growth temperature.
  • the residence time of the alloy bar at the germinative grain growth temperature should be at least sufiicient to promote germinative grain growth. Tests have indicated that once the extruded alloy bar has reached the appropriate germinative grain growth temperature, the residence time should be at least about 5 minutes and, more preferably, in excess of about 10 minutes.
  • the temperature reached at the center of the bar being heated lags behind the temperature imposed by the travelling hot zone.
  • the maximum temperature of the bar being heated generally will be below the maximum temperature of the hot zone.
  • the desired temperature may not be reached in the bar being heated or the time at the germinative grain growth temperature may be too short.
  • EXAMPLE 2 An extruded alloy bar No. 5 was produced having a diameter of about five-eighths of an inch from mechanically alloyed powders of substantially saturation hardness and having the following composition by weight: about 20.5% chromium, about 2.9% titanium, about 1.5% aluminum, about 0.06% zirconium, about 0.007% boron, about 0.07% carbon, about 1.23% Y O having a particle size of about 350 A. and the balance essentially nickel.
  • the bar was extruded from a 3.5 inch diameter can of powder at a temperature of about 2000 R, an extrusion ratio of about 31.8, and a ram speed exceeding one inch per second.
  • the extruded alloy was subject to zone annealing at various zone speeds and the following results obtained:
  • the residence time at 2300 F. should preferably exceed about minutes.
  • EXAMPLE 3 This example illustrates the efiect of too slow a speed through the hot zone. Material from the same alloy bar as that of Example 2 was employed. The results obtained 10 the same as those of Example 1. Yttria of varying particle size in varying amounts was included in each of the charges. The resulting powders were packed in 3 /2 inch diameter metal cans and were hot extruded without evacnation of the cans. In each instance, the extrusion ratio was 22:1 and the heating temperature prior to extrusion was 2000" F. with the exception of Alloy 6 for which the heating temperature was 1950 F.
  • Alloy 6 contained 2.25% by volume of yttria having a particle size of 162 angstroms
  • Alloy 7 contained 2.25 volume percent of yttria having a particle size of 170 angstroms.
  • Alloys 8, 9 and 10 contained, respectively, 1%, 4% and 5.5% volume of yttria having a particle size of about 151 angstroms and Alloy 11 contained 5.5 volume percent of yttria having a particle size of 278 angstroms.
  • Test bars made from each of the extrusions were subjected to conventional heat treatment at high temperature followed by aging seven hours at 1975 F. and 16 hours at 1300 F., while corresponding test bars from each extrusion were also subjected to a high temperature zone anneal using a are s follows: 20
  • the rate of travel of the hot zone may range from ape 81 hours Elons RA above 1 inch to 17 inches per hour, preferably from about 2.325 q ax d- 16 0 0.8 0.8 7. 2,325 --do 15 25 1.3 1.6 2 lnches to 15 inches per hour. It W111 be appreclated that g 2425 15 0,1 34 the upper speed limit is a function of bar diameter and '2 gig? .5 .3 g furnace characteristics. The important thing is to bring 1 :3 11: 15 1 0,9 the entire bar cross-section, including the center, to the TABLE 7 Zone Reerystullization-1,900 F. Stress Rupture IOU-hour Zone Percent stress tempora- Grain Stress Life, improveture,F. shape k.s.i. hours Elong. RA. ment, k.s.i.
  • the residence time for germinative grain growth at temperatures ranging from 2150 F. to below the incipient melting point of the alloy should be in excess of 5 minutes, preferably at least about 10 minutes, and range 7 up to about 4 hours.
  • the zone annealing heat traetment may comprise at least a first step by zone annealing at an elevated annealing temperature to solution treat, homogenize and germinatively grow the grains and form large coarse grains in the direction of working of the hot worked product.
  • a resolution treatment may be required because of coarse gamma prime precipitation occurring during slow cooling from the high zone annealing temperatures.
  • Subsequent heat treatment to age the alloy to the desired hardness and strength may or may not be employed. Aging may not be required where the alloy is used at a temperature at which aging occurs in situ. Aging treatments comprise one or more steps and may be a series of aging sub-steps at successively lower temperatures where desirable.
  • a three-step heat treatment found particularly advantageous comprises: (1) subjecting the hot worked alloy to a grain growth temperature of about 2325 F. to 2400 F. by means of a moving hot zone, wherein a protective environment, e.g., argon, may be provided; (2) thereafter, heating at a resolution temperature of 1975 F. for a suitable time, e.g., 7 hours, in air followed by air cooling; and (3) finally aging the alloy at 1300 F. for a longer time, such as 24 hours, in air and then air cooling.
  • the first step results in a marked increase in grain size by zone annealing having a preferred orientation relative to a working direction.
  • the coarse gra-ins are elongated and are disposed or exhibit a preferred orientation in the direction of extrusion, that is, the longitudinal axis of the elongated product.
  • the grains tend to be plate-like and to be disposed or show a preferred orientation in the direction of the major axis of working, that is, in the longitudinal direction and thus exhibit higher mechanical properties in the longitudinal direction.
  • the coarse grains generally exhibit aspect ratios of greater than about 3:1, in some cases greater than 10:1 or :1 or even higher.
  • the aspect ratio is that ratio that defines grain configuration correlated to the direction of interest, e.g., direction of applied stress.
  • the ratio is determined as the average dimension of the grain parallel to the direction of interest divided by its average dimension along a minor axis.
  • FIG. 2 illustrates photomacrographs (A to D) taken at 2 times magnification showing the improved grain structures (after zone annealing heat treatment at 2325 F.) obtained on an alloy comprising by weight 18.2% chromium, 1.3% total aluminum, 2.3% titanium, 0.004% boron, 0.06% zirconium, 1.34% Y O 0.055% carbon, 0.68% total oxygen, and the balance essentially nickel.
  • the alloy was produced from mechanically alloyed powder of substantially saturation hardness provided as described in Example 1, which powder was then placed into mild steel cans of about 3.5 inches in diameter. The cans were welded shut to form extrusion billets.
  • FIG. 2A depicts the kind of macrostructure obtained when the extruded bar is soaked conventionally at the germinative grain growth temperature of about 2325 F. for 2 hours, while FIGS. 2B, 2C and 2D show the microstructure obtained after zone annealing at the same temperature.
  • the hot zone traversed the bar at about 4.4 inches per hour. This specimen was pointed at the starting end to encourage the formation of a limited number of grains.
  • FIG. 3 illustrates a par produced similarly to those of FIGS. 2A through 2D which was removed from the furnace and cooled before complete passage of the hot zone which traveled at 18.3 inches per hour.
  • the fine grain of the untreated material is clearly evident as is the effect of advancing the hot zone upon the grain structure in the zone annealed portion of the bar.
  • the effect of the temperature lag in the center of the bar as compared to the bar surfaces is also shown.
  • the capability of producing parts which are fined-grained in one portion and 12 coarse-grained in another portion is clearly demonstrated by FIG. 3.
  • the heat treatment of hot worked alloy products produced from mechanically alloyed powder may vary over the following ranges:
  • First step Zone annealing from about 2200 F. to below the incipient melting point as described hereinabove;
  • Optional second step Solution heating at about 1750 F.
  • Optional third step And then aging at about 1150 F. to
  • the invention is particularly applicable to the following range of compositions: about 5% to 60% chromium, about 0.5% to 6.5% aluminum; about 0.5% to 6.5% titanium; up to about 15% molybdenum, up to about 20% tungsten; up to about 10% columbium; up to about 10% tantalum, up to about 3% vanadium; up to about 2% manganese; up to about 2% silicon; up to about 0.75% carbon; up to about 0.1% boron; up to about 1% zirconium; up to about 0.2% magnesium; up to about 6% hafnium; up to about 35% iron; up to 10 volume percent of a dispersoid, the balance essentially atleast about 40% nickel.
  • the alloys to which the invention is applicable have a melting point of at least about 2300 F.
  • the composition may comprise cobalt since nickel is generally considered an equivalent of cobalt.
  • the dispersoid may include those selected from the group consisting of ThO Y O La O and the rare earth mixtures didymia and Rare Earth Oxides, and other oxides having free energies of formation exceeding kilocalories per gram atom of oxygen at about 25 C.
  • the size of dispersoid found advantageous in producing dispersion-strengthened superalloys may range from about 50 angstroms to 5000 angstroms, and, more advantageously, from about angstroms to 1000 angstroms.
  • the alloys contain about 10% to 35% chromium, about 0.5% to 6% aluminum, about 1% to 5% titanium, up to about 5% molybdenum, up to about 10% tungsten, up to about 3% columbium, up to about 5% tantalum, up to about 15 cobalt, up to about 1% vanadium, up to about 2% manganese, up to about 1% silicon, up to about 0.2% carbon, up to about 0.1% boron, up to about 0.5% zirconium, up to 0.2% magnesium, up to 2% hafnium, up to about 10% iron, about 0.5 volume percent to 5 volume percent of a dispersoid, the balance essentially at least about 40% nickel.
  • the product provided in accordance with the invention is useful in the production of articles such as gas turbine blades and vanes and other articles subjected in use to the combined effects of elevated temperature and stress.
  • the heat resistant alloy has a composition ranging by weight from about 5% to 60% chromium, about 0.5 to 6.5% aluminum, about 0.5% to 6.5% titanium, up to about 15% molybdenum, up to about 20% tungsten, up to about columbium, up to about 10% tantalum, up to about 3% vanadium, up to about 2% manganese, up to about 2% silicon, up to about 0.75% carbon, up to about 0.1% boron, up to about 1% zirconium, up to about 0.2% magnesium, up to about 6% hafnium, up to about 35 iron, an eifective amount of up to about 10% by volume of a refractory dispersoid, and the balance essentially a metal from the group consisting of nickel and cobalt in an amount at least about 40% of the total composition, the alloy having been hot worked under conditions of temperature and reduction ratio lying within the closed area of FIG. 1, the residence time at the germinative grain growth temperature being substantially in excess of 5 minutes.
  • composition ranges from about 10% to 35% chromium, about 0.5% to 6% aluminum, about 1% to 5% titanium, up to about 5% molybdenum, up to about 10% tungsten, up to about 3% columbium, up to about 4% tantalum, up to about cobalt, up to about 1% vanadium, up to about 2% manganese, up to about 1% silicon, up to about 0.2% carbon, up to about 0.1% boron, up to about 0.5%
  • the alloy comprises about 19% chromium, about 2.4% titanium, about 1.2% aluminum, about 0.07% zirconium, about 0.007% boron, about 0.05% carbon, about 2.25% by volume of a dispersoid and the balance essentially nickel, and wherein said alloy is subjected to germinative grain growth by zone heating it to a temperature of at least about 2300" F. but below the incipient melting point of the alloy for up to about 4 hours.
  • a zone annealed high temperature heat resistant dispersion-strengthened alloy produced in accordance with the method of claim 1.
  • Line 75 for "-00.25%” read --o.25%-- Col 6, line 45', for "toneY' read --ton-- Col. 7, line 64, fer "33%” read dbdur i 33%- Col. 8, line 70, for. "1.23%” read "1.32%".

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Cited By (21)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3844847A (en) * 1973-09-11 1974-10-29 Int Nickel Co Thermomechanical processing of mechanically alloyed materials
US3847680A (en) * 1970-04-02 1974-11-12 Sherritt Gordon Mines Ltd Dispersion strengthened metals and alloys and process for producing same
US3850702A (en) * 1970-03-02 1974-11-26 Gen Electric Method of making superalloy bodies
US3909309A (en) * 1973-09-11 1975-09-30 Int Nickel Co Post working of mechanically alloyed products
DE3436878A1 (de) * 1983-10-26 1985-05-09 BBC Aktiengesellschaft Brown, Boveri & Cie., Baden, Aargau Vorrichtung zum zonengluehen eines aus einem hochtemperatur-werkstoff bestehenden werkstuecks und verfahren zum zonengluehen
EP0158844A1 (en) * 1984-03-19 1985-10-23 Inco Alloys International, Inc. Promoting directional grain growth in objects
US4588552A (en) * 1981-09-03 1986-05-13 Bbc Brown, Boveri & Co., Ltd. Process for the manufacture of a workpiece from a creep-resistant alloy
US4627959A (en) * 1985-06-18 1986-12-09 Inco Alloys International, Inc. Production of mechanically alloyed powder
US4665828A (en) * 1983-11-23 1987-05-19 Voest-Alpine Aktiengesellschaft Penetrator for a driving-cage projectile and the process of manufacturing the same
EP0232477A1 (de) * 1985-12-19 1987-08-19 BBC Brown Boveri AG Verfahren zum Zonenglühen eines metallischen Werkstücks
US4717435A (en) * 1985-10-26 1988-01-05 National Research Institute For Metals Gamma-prime precipitation hardening nickel-base yttria particle-dispersion-strengthened superalloy
EP0274631A1 (de) * 1986-12-19 1988-07-20 BBC Brown Boveri AG Verfahren zur Erhöhung der Duktilität eines in groben Längsgerichteten stengelförmigen Kristalliten vorliegenden Werkstücks aus einer oxyddispersiongsgehärteten Nickelbasis-Superlegierung bei Raumtemperatur
EP0330081A1 (en) * 1988-02-22 1989-08-30 Inco Alloys International, Inc. Oxide dispersion-strengthened alloy having high strength at intermediate temperatures
EP0384608A1 (en) * 1989-02-08 1990-08-29 Inco Alloys International, Inc. Mechanically alloyed nickel-cobalt-chromium-iron composition of matter
EP0442545A1 (de) * 1990-02-14 1991-08-21 PM HOCHTEMPERATUR-METALL GmbH Verfahren zur Herstellung wärmebehandelter Profile und Formkörper
EP0447858A1 (de) * 1990-03-20 1991-09-25 Asea Brown Boveri Ag Verfahren zur Erzeugung von längsgerichteten grobkörnigen Stengelkristallen in einem aus einer oxyddispersionsgehärteten Nickelbasis-Superlegierung bestehenden Werkstück
US20040208775A1 (en) * 2003-04-16 2004-10-21 National Research Council Of Canada Process for agglomeration and densification of nanometer sized particles
US20070131318A1 (en) * 2005-12-12 2007-06-14 Accellent, Inc. Medical alloys with a non-alloyed dispersion and methods of making same
US20100021338A1 (en) * 2008-07-25 2010-01-28 Alstom Technology Ltd High-temperature alloy
US8961646B2 (en) 2010-11-10 2015-02-24 Honda Motor Co., Ltd. Nickel alloy
CN105238958A (zh) * 2015-10-28 2016-01-13 无棣向上机械设计服务有限公司 镍基高温合金

Families Citing this family (12)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
GB1435796A (en) * 1972-10-30 1976-05-12 Int Nickel Ltd High-strength corrosion-resistant nickel-base alloy
US4111685A (en) * 1976-11-04 1978-09-05 Special Metals Corporation Dispersion-strengthened cobalt-bearing metal
EP0045984B1 (de) * 1980-08-08 1984-03-14 BBC Aktiengesellschaft Brown, Boveri & Cie. Verfahren zur Herstellung eines Werkstückes aus einer warmfesten Legierung
US4668290A (en) * 1985-08-13 1987-05-26 Pfizer Hospital Products Group Inc. Dispersion strengthened cobalt-chromium-molybdenum alloy produced by gas atomization
US4732622A (en) * 1985-10-10 1988-03-22 United Kingdom Atomic Energy Authority Processing of high temperature alloys
GB2181454B (en) * 1985-10-10 1990-04-04 Atomic Energy Authority Uk Processing of high temperature alloys
DE3714239C2 (de) * 1987-04-29 1996-05-15 Krupp Ag Hoesch Krupp Verfahren zur Herstellung eines Werkstoffs mit einem Gefüge nanokristalliner Struktur
AT399165B (de) * 1992-05-14 1995-03-27 Plansee Metallwerk Legierung auf chrombasis
US5827377A (en) * 1996-10-31 1998-10-27 Inco Alloys International, Inc. Flexible alloy and components made therefrom
DE102018208736A1 (de) * 2018-06-04 2019-12-05 Siemens Aktiengesellschaft Y, Y' gehärtete Kobalt-Nickel-Basislegierung, Pulver, Komponente und Verfahren
DE102018208737A1 (de) * 2018-06-04 2019-12-05 Siemens Aktiengesellschaft Y, Y` gehärtete Kobalt-Nickel-Basislegierung, Pulver, Komponente und Verfahren
DE102019208666A1 (de) * 2019-06-14 2020-12-17 MTU Aero Engines AG Rotoren für hochdruckverdichter und niederdruckturbine eines getriebefantriebwerks sowie verfahren zu ihrer herstellung

Cited By (28)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3850702A (en) * 1970-03-02 1974-11-26 Gen Electric Method of making superalloy bodies
US3847680A (en) * 1970-04-02 1974-11-12 Sherritt Gordon Mines Ltd Dispersion strengthened metals and alloys and process for producing same
US3844847A (en) * 1973-09-11 1974-10-29 Int Nickel Co Thermomechanical processing of mechanically alloyed materials
US3909309A (en) * 1973-09-11 1975-09-30 Int Nickel Co Post working of mechanically alloyed products
US4588552A (en) * 1981-09-03 1986-05-13 Bbc Brown, Boveri & Co., Ltd. Process for the manufacture of a workpiece from a creep-resistant alloy
CH657151A5 (de) * 1983-10-26 1986-08-15 Bbc Brown Boveri & Cie Vorrichtung zum zonengluehen eines aus einem hochtemperatur-werkstoff bestehenden werkstuecks und verfahren zum zonengluehen.
DE3436878A1 (de) * 1983-10-26 1985-05-09 BBC Aktiengesellschaft Brown, Boveri & Cie., Baden, Aargau Vorrichtung zum zonengluehen eines aus einem hochtemperatur-werkstoff bestehenden werkstuecks und verfahren zum zonengluehen
US4665828A (en) * 1983-11-23 1987-05-19 Voest-Alpine Aktiengesellschaft Penetrator for a driving-cage projectile and the process of manufacturing the same
EP0158844A1 (en) * 1984-03-19 1985-10-23 Inco Alloys International, Inc. Promoting directional grain growth in objects
US4921549A (en) * 1984-03-19 1990-05-01 Inco Alloys International, Inc. Promoting directional grain growth in objects
US4627959A (en) * 1985-06-18 1986-12-09 Inco Alloys International, Inc. Production of mechanically alloyed powder
US4717435A (en) * 1985-10-26 1988-01-05 National Research Institute For Metals Gamma-prime precipitation hardening nickel-base yttria particle-dispersion-strengthened superalloy
EP0232477A1 (de) * 1985-12-19 1987-08-19 BBC Brown Boveri AG Verfahren zum Zonenglühen eines metallischen Werkstücks
US4743309A (en) * 1985-12-19 1988-05-10 Bbc Brown, Boveri & Company, Limited Method for zone heat treatment of a metallic workpiece
EP0274631A1 (de) * 1986-12-19 1988-07-20 BBC Brown Boveri AG Verfahren zur Erhöhung der Duktilität eines in groben Längsgerichteten stengelförmigen Kristalliten vorliegenden Werkstücks aus einer oxyddispersiongsgehärteten Nickelbasis-Superlegierung bei Raumtemperatur
EP0330081A1 (en) * 1988-02-22 1989-08-30 Inco Alloys International, Inc. Oxide dispersion-strengthened alloy having high strength at intermediate temperatures
EP0384608A1 (en) * 1989-02-08 1990-08-29 Inco Alloys International, Inc. Mechanically alloyed nickel-cobalt-chromium-iron composition of matter
EP0442545A1 (de) * 1990-02-14 1991-08-21 PM HOCHTEMPERATUR-METALL GmbH Verfahren zur Herstellung wärmebehandelter Profile und Formkörper
EP0447858A1 (de) * 1990-03-20 1991-09-25 Asea Brown Boveri Ag Verfahren zur Erzeugung von längsgerichteten grobkörnigen Stengelkristallen in einem aus einer oxyddispersionsgehärteten Nickelbasis-Superlegierung bestehenden Werkstück
US5180451A (en) * 1990-03-20 1993-01-19 Asea Brown Boveri Ltd. Process for the production of longitudinally-directed coarse-grained columnar crystals in a workpiece consisting of an oxide-dispersion-hardened nickel-based superalloy
US20040208775A1 (en) * 2003-04-16 2004-10-21 National Research Council Of Canada Process for agglomeration and densification of nanometer sized particles
US7235118B2 (en) 2003-04-16 2007-06-26 National Research Council Of Canada Process for agglomeration and densification of nanometer sized particles
US20070131318A1 (en) * 2005-12-12 2007-06-14 Accellent, Inc. Medical alloys with a non-alloyed dispersion and methods of making same
US20070131317A1 (en) * 2005-12-12 2007-06-14 Accellent Nickel-titanium alloy with a non-alloyed dispersion and methods of making same
US20100021338A1 (en) * 2008-07-25 2010-01-28 Alstom Technology Ltd High-temperature alloy
US8153054B2 (en) * 2008-07-25 2012-04-10 Alstom Technology Ltd High-temperature alloy
US8961646B2 (en) 2010-11-10 2015-02-24 Honda Motor Co., Ltd. Nickel alloy
CN105238958A (zh) * 2015-10-28 2016-01-13 无棣向上机械设计服务有限公司 镍基高温合金

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SE391737B (sv) 1977-02-28
AT321595B (de) 1975-04-10
ZA73505B (en) 1973-10-31
FR2170100B1 (US07652168-20100126-C00068.png) 1976-11-05
CH601466A5 (US07652168-20100126-C00068.png) 1978-07-14
AU474229B2 (en) 1976-07-15
CA967461A (en) 1975-05-13
JPS4886713A (US07652168-20100126-C00068.png) 1973-11-15
AU5150173A (en) 1974-08-01
GB1413762A (en) 1975-11-12
DE2303802C3 (de) 1978-06-29
DD104924A5 (US07652168-20100126-C00068.png) 1974-04-05
IT994044B (it) 1975-10-20
BE794801A (fr) 1973-07-31
DE2303802B2 (de) 1977-11-03
NL7301286A (US07652168-20100126-C00068.png) 1973-08-02
FR2170100A1 (US07652168-20100126-C00068.png) 1973-09-14
DE2303802A1 (de) 1973-08-16
JPS5641686B2 (US07652168-20100126-C00068.png) 1981-09-30

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