US3707409A - Nickel base alloy - Google Patents

Nickel base alloy Download PDF

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Publication number
US3707409A
US3707409A US55796A US3707409DA US3707409A US 3707409 A US3707409 A US 3707409A US 55796 A US55796 A US 55796A US 3707409D A US3707409D A US 3707409DA US 3707409 A US3707409 A US 3707409A
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strength
columbium
titanium
aluminum
ductility
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William J Boesch
Falih N Darmara
Francis J Warmuth
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ALLEGHENY INTERNATIONAL ACCEPTANCE Corp
Special Metals Corp
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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
    • C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%

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  • the present invention invention relates to a nickel base alloy and more particularly to a nickel base alloy with a superb combination of strength and ductility.
  • Nickel base alloys have been known and used at elevated temperatures for quite some time. These alloys have been disclosed with varying amounts of columbium, iron, cobalt, aluminum, titanium, boron, carbon, zirconium, chromium, and molybdenum. Particular nickel base alloys are discussed in United States Patent No. 3,107,167 which issued on Oct. 15, 1963.
  • the present invention teaches a nickel base alloy which has a superb combination of both strength and ductility at room temperature and at elevated temperatures. It is particularly useful as a turbine disc, especially in the heavy hub section and is also useful in other turbine applications as well as in sheet material and fastener stock applications. It contains especially critical ranges for columbium, cobalt, aluminum, titanium, iron, boron, zirconium, chromium, molybdenum, carbon, cerium, vanadium, hafnium, and tantalum.
  • FIG. 1 is a plot of alloy strength at varying columbium levels
  • FIG. 2 is a plot of alloy ductility at varying columbium levels
  • FIG. 3 is a photomicrograph at 1000 of a nickel base alloy having 'a columbium content of 052%.
  • FIG. 4 is a photomicrograph at 1000 of a nickel base alloy having a columbium content of 1.70%.
  • FIG. 5 is a plot of alloy strength, expressed in rupture time, at varying cobalt levels
  • FIG. 6 is a plot of alloy ductility at varying cobalt levels
  • FIG. 7 is a plot of alloy strength, expressed in rupture time, at varying iron levels.
  • FIG. 8 is a plot of alloy ductility at varying iron levels.
  • the alloy of the present invention has a composition consisting essentially of, in weight percent, from 0.2 to 1.2% columbium, from 10 to 15% cobalt, from 1.2 to 2.0% aluminum, from 2 to 3.5% titanium, from 0.10 to 2% iron, from 0.0025 to 0.125% boron, from 0.05 to 0.2% zirconium, from 15 to 19% chromium, up to 0.15% carbon, from 3 to 5% molybdenum, up to 0.1% cerium, up to 0.5% vanadium, up to 0.02% of at least one metal from Group IIA of the Periodic Table; e.g., magnesium, barium or calcium, balance essentially nickel.
  • Aluminum and titanium are present in amounts such that their sum is at least 3.6% and such that the titanium to aluminum ratio is less than 2.0.
  • the amount of columbium can vary between 0.2 and 1.2%. Columbium is necessary in amounts of at least 0.2% to stabilize randomly dispersed metallic carbides.
  • the carbides are believed to have the general composition MC, wherein the M portion is regarded as consisting mainly of titanium and molybdenum and is considered to have the approximate atomic proportions: titanium, 15 molybdenum. Without columbium, the carbides decompose during heating for working and/or during high temperature heat treatment and reprecipitate during cooling as a fine grain boundary MC. This fine MC is believed to cause a loss of ductility and strength. Amounts of columbium greater than 1.2% promote the formation of an acicular phase which detrimentally affects ductility and strength. This phase is believed to be a nickel-columbium intermetallic of the approximate composition, Ni Cb. A preferred range for columbium is between 0.4 and 0.8%.
  • Cobalt is present in the range from 10% to 15%.
  • a minimum of 10% cobalt is necessary for high temperature strength. Higher levels maintain the strength achieved with 10% additions and simultaneously increase the alloys ductility. Levels in excess of 15 are too high as they render the alloy susceptible to the formation of deleterious acicular phases, e.g., Ni Cb, which detrimentally affect ductility and strength.
  • Excessive cobalt causes a decrease in the nickel balance and initiates the formation of Ni Cb.
  • Ni Cb forms when there is not enough nickel to hold the residual columbium that is not tied up in carbides.
  • a preferred cobalt range is between 12 and 14%.
  • the amounts of aluminum and titanium are respectively between 1.2 and 2.0% and between 2 and 3.5%. In addition, the minimum amount of aluminum and titanium is at least 3.6%. A minimum of 1.2% aluminum is necessary to impart sufiicient strength to the alloy. Alloys with more than 2.0% aluminum lack sulficient strength at low temperatures and at elevated temperatures under high stress levels. A minimum of 2% titanium is necessary to insure the formation of sufficient gamma prime, which imparts strength to the alloy. Gamma prime is believed to have the general composition M (Al, Ti).
  • the M portion of the gamma prime composition is regarded as consisting mainly of nickel with some substitution of chromium and molybdenum and is considered to have the approximate proportions, nickel, 3 chromium, and 2 molybdenum.
  • Alloys with more than 3.5 titanium are susceptible to the formation of Ni Ti, which reduces ductility.
  • the ratio of titanium to aluminum must be less than 2.0 as the likelihood of forming Ni Ti increases as the ratio of titanium to aluminum increases.
  • the minimum amount of aluminum and titanium should be at least 3.6% to insure that the alloy consistently possesses high strength.
  • Preferred ranges for aluminum and titanium are respectively between 1.3 and 1.8% and between 2.4 and 2.8%.
  • Alloys of the present invention have between 0.10 and 2% iron.
  • the alloys ductility increases as the iron level increases from 0.10 to 2%. Iron levels in excess of 2% are disadvantageous as they are accompanied by a loss in high temperature strength and not accompanied by further increase in ductility.
  • Boron is present within the range of 0.0025 to 0.0125%
  • a boron level of 0.0025 is required to impart sufiicient I strength to the alloy. Boron levels in excess of 0.0125 are detrimental as they increase the likelihood of forming boride clusters during solidification, which are not eliminated during subsequent heat treatment. Boride clusters cause a localized restriction of grains, i.e., they impede grain growth, which results in the formation of non-uniform grain structure. Non-uniform grain structures are harmful to tensile properties perpendicular to the direction of forming.
  • a preferred boron range is between 0.0050 and 0.0100%.
  • Zirconium is present in amounts from 0.05 to 0.2%. At least 0.05% zirconium is necessary for workability and hot ductility. Levels in excess of 0.2% are, however, detrimental as they promote the formation of nickel-zirconium intermetallics which decrease hot ductility. The intermetallics form since the solubility of zirconium in nickel is limited. A preferred zirconium range is from 0.06 to 0.09%.
  • Chromium is present within the range of 15-19%. A chromium level of 15% is necessary to impart sufficient oxidation resistance to the alloy. Chromium should not, however, exceed 19% as the alloys strength decreases with increasing chromium levels. Chromium increases the gamma prime solvus which in turn decrease the amount of gamma prime. A preferred chromium range is between 16 and 18%.
  • the carbon level should not exceed 0.15%. Higher carbon levels generate excessive MC which could decompose during heating for working and/or during high temperature heat treatments and reprecipitate during cooling as the detrimental fine, grain boundary MC which is believed to cause a loss in ductility and strength. A small amount of carbon is, however, often desirable as carbon increases the alloys yield strength. A preferred range for carbon is between 0.03 to 0.07%.
  • Molybdenum should be present within the range of from 3 to 5%. At least 3% molybdenum is necessary for high temperature strength. Molybdenum levels in excess of 5% are, however, detrimental as molybdenum ceases to act as a solid solution element at higher levels and precipitates as molybdenum rich carbides. The formation of molybdenum rich carbides increases the likelihood of forming detrimental acicular phases such as Ni Ti as it shifts the solvus for the intermetallics. As stated above, the acicular phases are believed to detrimentally affect both ductility and strength. A preferred range for molybdenum is between 3.5 and 4.5%.
  • the alloy can have up to 0.1% cerium which improves hot ductility, up to 0.5% vanadium as small amounts of vanadium have a tendency to increase hot workability, up to 0.02% of at least one metal from Group II-A of the periodic table as Group II-A metals are believed to improve grain boundary ductility and thus decrease the tendency towards grain boundary cracking in creep during long time service, and in substitution the alloy can have tantalum and/or hafnium in place of columbium. Two parts of hafnium and/or tantalum are necessary to replace one part of columbium. Columbium is, however, preferred over tantalum and hafnium.
  • the samples were (1) rolled to a thickness between 0.46 and 0.48 inch, soaked for 4 hours at 2200 F., equalized at 2150 F., and given a 15% reduction to a final thickness of about 0.40 inch; (2) given a three stage heat treatment which comprised the steps of heating at a temperature of 1825 F. for 4 hours and oil quenching, heating at a temperature of 1550 F. for 4 hours and air cooling and heating at a temperature of 1400 F. for 16 hours and air cooling; and (3) tested for strength and ductility. The results of the tests are given below in Table II.
  • FIGS. 1 and 2 show how columbium additions affect the alloys strength and FIG. 2 shows how columbium additions affect the alloys ductility.
  • the figures reveal how the strength and ductility areimproved with columbium additions within the range of 0.2 to 1.2% and how strength and ductility are optimized with columbium additions between 0.4 and 0.8%.
  • Photomicrographs at 1000 were taken of samples B and D to show the effect of excessive amounts of columbium.
  • Sample B shown in FIG. 3 had a columbium content of 0.52%, a columbium content within the range of this invention, and sample D shown in FIG. 4 had a columbium content of 1.70%, a columbium range outside the range of this invention.
  • Examination of FIGS. 3 and 4 reveals that sample D in FIG. 4 has an acicular phase which detrimentally affects ductility and strength whereas the presence of an acicular phase is not evident in sample B of FIG. 3.
  • the acicular phase shown in FIG. 4 is believed to be a nickel-columbium intermetallic of the approximate composition, Ni Cb.
  • Columbium is, however, necessary as working operations, e.g., rolling and forging, cannot always be tailored to particular practices, e.g., the practice used to obtain the data for Table III, and since columbium obviates the need to so tailor the working operation.
  • the sample B properties given in Table II are equal to or better than the sample A properties given in Table III.
  • sample E had a cobalt content of 8.2%
  • sample F had a cobalt content of 10.1%
  • sample G had a cobalt content of 12.1%
  • sample H had a cobalt content of 14.1%.
  • the composition of the samples is shown below in Table IV.
  • the samples were worked, given a three stage heat treatment which comprised the steps of heating at a temperature of 1975 F. for 4 hours and air cooling, heating at a temperature of 1550 F. for 24 hours and air cooling, and heating at a temperature of 1400" F. for 16 hours and air cooling and subsequently stress-rupture tested at 1500 F. under a stress of 47,500 p.s.i.
  • the results of the tests are given below in Table V.
  • FIGS. 5 and 6 show how cobalt additions affect the alloys high temperature strength and FIG. 6 shows how cobalt additions aifect the alloys ductility. The figures reveal how high temperature strength is attained with a minimum cobalt level of 10% and how ductility improves with increasing amounts of cobalt.
  • Samples for the first four tests were given a three-stage heat treatment which comprised the steps of heating at a temperature of 1825 F. for four hours and oil quenching, heating at a temperature of 1550 F. for 4 hours and air cooling and heating at a temperature of 1400 F. for 16 hours and air cooling.
  • Samples for the fifth test were given a three-stage heat treatment which comprised the steps of heating at a temperature of 1975 F. for 4 hours and air cooling, heating at a temperature of 1550 F. for 24 hours and air cooling, and heating at a temperature of 1400 F. for 16 hours and air cooling.
  • Test 1 was a stress rupture test at 1350 F.
  • test 2 under a stress of 70,000 p.s.i.
  • test 3 was a stress rupture test at 1500 F. under a stress of 47,500 p.s.i.
  • test 4 was a room temperature tensile test and test 5 was a room temperature tensile test.
  • the results of tests 1 through 5 are respectively given in Tables VII through XI.
  • Tables VII, X, and XI show how the alloys strength decreases with increasing aluminum contents above 1.60% and frequently becomes too low at levels in excess of 2%.
  • Table VII was based upon a test conducted at an elevated temperature under a high stress level and Tables X and XI were based upon room temperature tensile tests.
  • Tables VIII and IX show how aluminum contents up to and in excess of 2% increase the strength of alloys tested at elevated temperatures under moderate stresses. In the past this data could have led to the choice of an improper aluminum content for alloys requiring good room temperature strength and/r good strength at elevated temperatures under high stress levels, as do the alloys of this invention.
  • sample N had an iron content of 0.14%
  • sample 0 had an iron content of 1.82%
  • sample P had an iron content 'of 3.99%
  • sample Q had an iron content of 5.85%.
  • the composition of the samples is shown below in Table XII.
  • the samples were worked, given a three-stage heat treatment which comprised the steps of heating at a temperature of 1975" F. for 4 hours and air cooling, heating at a temperature of 1550 F. for 24 hours and air cooling, and heating at a temperature of 1400 F. for 16 hours and air cooling, and subsequently stress-rupture tested at 1500 F. under a stress of 47,500 p.s.i.
  • the results of the tests are given below in Table XI'II.
  • FIGS. 7 and 8 show how iron additions affect the alloys high temperature strength and FIG. 8 shows how iron additions affect the alloys ductility. Note the increase in ductility; i.e., higher reduction in area, which accompanies iron level increases frqrn 0.10 to 0.25% and up up 2.0%,
  • samples S and T were hot worked from 5.75 inch diameter ingots to 2.25 inch square bars;
  • samples R, S and T were given a three stage heat treatment which comprised the steps of heating at a temperature of 1850 F. for 4 hours and oil quenching, heating at a temperature of 1550 F. for 4 hours and air cooling, and heating at a temperature of l400 for 16 hours and air cooling;
  • samples R, S and T were tensile tested at room temperature and at 1000 F. in the longitudinal direction; i.e., the direction of working, stress rupture tested at 1500 F. under a stress of 42,500 p.s.i.
  • Table XV gives the room temperature tensile test results in the longitudinal direction
  • Table XVI gives the 1000 F. tensile test results in the longitudinal direction
  • Table XVII gives the stress rupture test results
  • Table XVIII gives the room temperature tensile test results in the transverse direction.
  • the tested alloy had an elongation of over 24% and a 26 to 27% reduction in area, in the transverse test direction.
  • These figures are quite impressive when contrasted With prior art alloys which have displayed an elongation of 9 to 15% and an 11 to 16% reduction in area, in the transverse test direction.
  • a nickel base alloy having stabilized metallic carbides and being substantially free of acicular phases, consisting essentially of, in weight percent from 0.4 to 0.8% columbium, from to cobalt, from 1.2 to 2.0% aluminum, from 2 to 3.5% titanium, from 0.10 to 2% iron, between 0.005 and 0.010% boron, from 0.05 to 0.2% zirconium, from 15 to 19% chromium, up to 0.15% carbon, from 3 to 5% molybdenum, up to 0.1% cerium, up to 0.5% vanadium, up to 0.02% of at least one metal from Group II-A of the periodic table, balance essentially nickel; said percentages of aluminum and titanium being at least 3.6% and said titanium and aluminum being present in a titanium to aluminum ratio of less than 2.0; said stabilized metallic carbides being MC carbides wherein M consists essentially of titanium, molybdenum and columbium.
  • a nickel base alloy according to from 12 to 14% cobalt.
  • a nickel base alloy according from 2.4 to 2.8% titanium 4.
  • a nickel base alloy according from 0.06 to 0.09% zirconium.
  • a nickel base alloy according from 16 to 18% chromium.
  • a nickel base alloy according from 0.03 to 0.07% carbon 7.
  • claim 1 having claim 1 having to claim 1 having to claim 1 having to claim 1 having to claim 1 having 8.
  • a nickel base alloy according to claim 1 having from 0.4 to 0.8% columbium, from 12 to 14% cobalt, from 1.3 to 1.8% aluminum, from 2.4 to 2.8% titanium, from 0.0050 to 0.0100% boron, from 0.06 to 0.09% zorconium, from 16 to 18% chromium, from 0.03 to 0.07% carbon and from 3.5 to 4.5 molybdenum.

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  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
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CH (1) CH552677A (it)
DE (1) DE2135741A1 (it)
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GB (1) GB1302980A (it)
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Cited By (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4755240A (en) * 1986-05-12 1988-07-05 Exxon Production Research Company Nickel base precipitation hardened alloys having improved resistance stress corrosion cracking
US20090257865A1 (en) * 2008-03-31 2009-10-15 Kabushiki Kaisha Toshiba Ni-base alloy for turbine rotor of steam turbine and turbine rotor of steam turbine
US20090285692A1 (en) * 2008-03-17 2009-11-19 Kabushiki Kaisha Toshiba Ni-base alloy for turbine rotor of steam turbine and turbine rotor of steam turbine

Families Citing this family (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS61119640A (ja) * 1984-11-16 1986-06-06 Honda Motor Co Ltd 排気バルブ用合金
DE3638855A1 (de) * 1985-11-26 1987-05-27 United Technologies Corp Superlegierung auf nickelbasis
CN116875844B (zh) * 2023-07-28 2024-02-09 北京钢研高纳科技股份有限公司 一种盘轴一体涡轮盘及其制备方法

Cited By (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4755240A (en) * 1986-05-12 1988-07-05 Exxon Production Research Company Nickel base precipitation hardened alloys having improved resistance stress corrosion cracking
US20090285692A1 (en) * 2008-03-17 2009-11-19 Kabushiki Kaisha Toshiba Ni-base alloy for turbine rotor of steam turbine and turbine rotor of steam turbine
US8828313B2 (en) 2008-03-17 2014-09-09 Kabushiki Kaisha Toshiba Ni-base alloy for turbine rotor of steam turbine and turbine rotor of steam turbine
US20090257865A1 (en) * 2008-03-31 2009-10-15 Kabushiki Kaisha Toshiba Ni-base alloy for turbine rotor of steam turbine and turbine rotor of steam turbine

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AU3140471A (en) 1973-01-25
CH552677A (fr) 1974-08-15
CA940346A (en) 1974-01-22
SE376636B (it) 1975-06-02
IT945882B (it) 1973-05-10
FR2103033A5 (it) 1972-04-07
GB1302980A (it) 1973-01-10
DE2135741A1 (de) 1972-02-10

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