US3649379A - Co-precipitation-strengthened nickel base alloys and method for producing same - Google Patents

Co-precipitation-strengthened nickel base alloys and method for producing same Download PDF

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US3649379A
US3649379A US834942A US3649379DA US3649379A US 3649379 A US3649379 A US 3649379A US 834942 A US834942 A US 834942A US 3649379D A US3649379D A US 3649379DA US 3649379 A US3649379 A US 3649379A
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alloy
nickel
phase
precipitation
nickel base
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Peshotan Sohrab Kotval
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Haynes International Inc
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Cabot 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
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%

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  • This invention relates to precipitation strengthened nickel base alloys and more particularly to alloys of this type which are strengthend by the co-precipitation of monocarbides and an intermetallic phase.
  • the invention also relates to a method for producing co-precipitation strengthened nickel base alloys.
  • Precipitation strengthened nickel base alloys have heretofore been based upon the formation of the well known aluminum-titanium rich phase known as gamma prime. This strengthening mechanism requires the presence of aluminium and titanium in an amount of as much as 8% by weight.
  • Another object is to provide a precipitation strengthened nickel base alloy which does not require the presence of aluminum and titanium therein.
  • a further object is to provide a nickel base alloy which has been strengthened by the co-precipitation of an A M phase and an MC monocarbide phase wherein A consists essentially of nickel and M consists essentially of one or more elements selected from the group consisting of Ta, Cb and V, wherein C is carbon.
  • Yet another object is to provide a method for producing a co-precipitation strengthened nickel base alloy having an A M phase and a MC monocarbide phase as defined above.
  • FIGS. 1 and 2 are schematic representations of the microstructure of the alloys of the present invention.
  • FIG. 3 is a diagrammatic process flow chart outlining the steps required for making the alloys of the invention.
  • the present invention is based upon two discoveries which are utilized-in combination to produce new nickel base alloys as well as a method for producing same.
  • the first discovery resides in the fact that it is possible to precipitate an intermetalic strengthening phase of the .A M type as defined above from a nickel base matrix. It should be understood that this precipitation phase is distinct from the above mentioned gramma prime phase achieved through the use of aluminum and titanium in combination. It is known that the gamma prime phase has 3,649,379 Patented Mar. 14, 1972 a crystal structure of the ordered face centered cubic (FCC) type. In marked contrast to this, the A M phase produced within the alloys of the present invention is noncubic in crystal structure.
  • FCC ordered face centered cubic
  • the second discovery resides in the fact that it is possible to precipitate a well dispersed distriubtion of monocarbides having particle sizes of less than 250 angstroms in the alloy.
  • Most prior art wrought nicket base alloys contain sufiicient carbon so as to form monocarbides of a simple face centered cubic structure whenever monocarbide forming elements such as V, Ta and Cb are present. It is important to note that these monocarbides are usually in the form of undissolved primary particles present in the alloy matrix but which do not effectively strengthen the same.
  • primary monocarbide particles have particle diameters of between 10 and 20 microns in nickel base alloys.
  • This invention contemplates the co-precipitation of the abovementioned intermetallic and the monocarbide phases concurrently.
  • nickel-base alloy compositions having a new microstructure.
  • This microstructure is comprised of intermetallic-cum-monocarbide co-precipitated phases.
  • the interrnetallic phase may be defined as A M wherein A consists essentially of Nickel and M consists essentially of one or more elements selected from the group consisting of Ta, Cb and V.
  • the monocarbide phase may be defined as MC wherein M consists essentially of one or more elements selected from the group consisting of Ta, Cb and V, and C is carbon, the MC monocarbides being in the form of fine particles of less than 250 angstroms size.
  • the structure would also contain primary monocarbides of 5 microns or less in size.
  • the chemical composition of the abovementioned alloys of the invention consists essentially of (by weight): 15- 22% Cr, 312% M0, at least one of the group consisting of 7-10% Ta, 6-9% Cb and 69% V, the total Ta-i-Cb-i-V content not exceeding 10%.
  • the alloy must also contain 0.030.15% carbon with the balance being nickel and residual impurities.
  • the impurities should have not more than 7% Fe and not more than 8% Co.
  • the ability to produce the abovementioned nickel base alloy results from the discovery that by controlling the level of certain impurities in the nickel base alloy, (i.e. Fe) it is possible to solution heat treat the alloy at temperatures of about 2280 F. without causing liquation and melting. Although, ideally, higher temperatures would normally be required to dissolve monocaribdes, it is possible to achieve partial dissolution thereof by progressively subjecting the material to an alternating sequence of high temperature solution heat treatments at about 2280 F. and quenching together with controlled cold working steps.
  • certain impurities in the nickel base alloy i.e. Fe
  • Chromium is required in the alloy within the range disclosed above to provide strengthening and corrosion resistance. Less than 15% Cr yields an alloy with minimal corrosion resistance; over 22% Cr yields an alloy with reduced ductility. Molybdenum is present in the alloy within the ranges shown above to provide further solution strengthening and corrosion resistance as required. Molybdenum is preferred in the alloy, although tungsten may replace molybdenum in whole or in part. Tungsten may be present in greater quantities than molybdenum, i.e., up to a maximum of 16 percent, but preferably about 12 percent nominally. Carbon must be present in the alloy within the range about 0.03 to 0.15% by weight to promote the formation of carbides in the alloy.
  • the alloy system of this invention must also contain at least one of the group including tantalum, colnmbium, and vanadium. At least one of these elements must be present in the alloy together with carbon to provide the metal monocarbides that are precipitated in the nickel-chromium (molybdenum) matrix. The presence of these precipitated carbides in addition to the aforementioned A M intermetallic phase together with the critical proccessing steps that promote the controlled precipitation are the heart of the present invention.
  • each element in this group must be present within the ranges stated above when that element is the principal carbide and intermetallic former.
  • the total content of tantalum-lcolumbium+vanadium in the alloy system must not exceed by weight.
  • Iron may be present up to a maximum of 7%. This requirement results from the fact that the higher iron content tends to lower the melting point of the alloy thereby prohibiting the higher solution heat treatment temperature required to cause the dissolution of the primary monocarbides so that subsequent reprecipitation can be achieved as will be hereinafter described.
  • Cobalt may be present up to a maximum of 8%. Higher cobalt contents tend to lower the matrix stacking fault energy and thereby cause impediments in the nucleation of the dislocation reaction which is a necessary precondition for the desired morphology of monocarbide precipitation.
  • the use of aluminum and titanium in the alloy compositions of the invention is not required. It is desirable that the aluminum plus titanium content should be as low as possible, and preferably less than 1.4% If the aluminum plus titanium level is low, it will be possible to produce a good quality alloy with an air melting practice.
  • Boron, silicon, manganese, magnesium, and copper up to a total of about 2.5% may be present in the alloy within the ranges known in the art to be effective to enhance certain characteristics associated with these elements; i.e. the deoxidation step, casting fluidity, ductility and the like.
  • the balance of the alloy is nickel and adventitious impurities generally known to be present in this class of alloys.
  • FIG. 1 is a threedimensional representation of the microstructure as would be observed in a thin foil sample using transmission elec- -prirnary MC monocarbides P.
  • the actual size of the primary MC monocarbides P shown is about l-2 microns.
  • the unique structure further has dispersed within it, fine precipitates of MC monocarbides K which are associated with sheets of planar lattice defects.
  • the precipitates K- are of about 250 angstroms or less in size and being coherent with the matrix strengthen it.
  • the unique structure still further has, uniformly dispersed within the matrix, an A M intermetallic phase which is coherent therewith and which in combination with the precipitated MC monocarbide phase greatly strengthens the entire alloy system.
  • the alloys of the invention can be made by providing an alloy material with 15-22% Cr, 312% M0, at least one member selected from the group consisting of 710% Ta, 6-9% Cb and 69% V, the total of Ta+Cb+V not exceeding 10%, and 0.03-O.15% C, the balance being nickel and residual impurities said impurities consisting of not more than 7% Fe and not more than 8% Co, said alloy being in a form capable of being cold-worked.
  • the material is solution heat-treated at a temperature within the range 22502300 F. for a period of about 24 hours to homogenize the matrix and to at least partially dissolve the primary monocarbide particles and thereby cause the monocarbide-formingelement and the carbon to be put into solid solution in the alloy matrix.
  • the material is water-quenched and cold worked.
  • the abovementioned sequence of solution-heat-treatment followed by quenching and coldworking is repeated. Thereafter, the cold-worked material is annealed for a period not exceeding one hour within the solution heat-treatment temperature range of 2250-2300 F.
  • the material After annealing, the material is water-quenched to about ambient temperature and is thereafter aged at a temperature within the range of 1100-1450 F. for a period of 24 to hours.
  • both the A M intermetallic phase and the fine (less than 250 angstroms) monocarbide phase are precipitated uniformly throughout the matrix.
  • the occurrence of these two coherent phases in combination is the desired strengthening-mechanism of this invention.
  • EXAMPLE I A 5 lb. heat of material consisting of (by wt.): 20.34% Cr, 8.90% Mo, 0.11% C, 8.58% Ta, 4.8% Fe, 0.08% Cb, 0.23% Al, 0.01% Ti, the balance being nickel was electron-beam melted and cast into a 1.5" round bar. This bar was solution heat treated at a temperature of 2282 F. for a period of 24 hours and water-quenched. Thereafter, the material was sequentially cold-worked, solution-heattreated at 2282 F. and quenched. After the process was repeated five times, the resultant 0.004" sheet produced was solution heat treated (annealed) for a period of 1 hour at 2282 F. and quenched.
  • the material was cut into various specimens, each of which was aged at a different temperature and time.
  • the range of aging temperatures was ll001450 F. and the range of aging times was 0.5 hours to 1500 hours.
  • thin foil transmission electron microscopy anddiffraction revealed that the requisite structure, as depicted in FIGS. 1 and 2, and consisting of a uniform intragranular distribution of the two co-precipitated coherent phases i.e. an A M intermetallic phase and a MC monocarbide ph'asewas obtained.
  • the examination further showed a complete absence of a face centered cubic gamma prime phase.
  • EXAMPLE II A 50 1b. heat of material consisting of (by weight): 20.47% Cr, 5.41% Mo, 0.089% C, 9.82% Ta, 5.59% Fe and 5.02% Co, the balance being nickel was vacuum induction melted and cast. After homogenization at a temperature of about 2282 F. over a period of 24 hours the ingot was forged to one inch square bar. .This asforged material was thereafter solution heat treated..(ho mogenized) at 2282 F. for a period of 24 hours and water-quenched. Thereafter, the material was sequentially cold-worked, solution heat-treated at 2282". F. and quenched until a form of 0.004" sheet was attained. The material was thereafter further processed as in Example 1. Electron microscopy and. diffraction analysis revealed that the structure shown in FIGS. 1 and 2 and described in Example I was attained. As in Example I, there was a complete absence of any gamma prime formation present in the alloy produced.
  • a nickelbase alloy consisting essentially by weight of 15 to 22% Cr, 3 to 12% Mo, 0.03 to 0.15% C and at least one of the group consisting of 7 to 10% Ta, 6 to 9% Cb and 6 to 9% V, the total of said group not exceeding about 10%, and with the balance being nickel and residual impurities, provided that the alloy is essentially free of aluminum and titanium and contains not more than 7% Fe and not more than 8% Co, said alloy having both a coherent, non-cubic intermetallic phase represented by general formula A M and a coherent facecentered cubic monocarbide phase represented by the general formula MC wherein A consists essentially of Ni, M is a metal from the group consisting of Ta, Cb and V and C is carbon, said monocarbide phase being present both as primary particles of up to about microns in size and as fine particles of less than 250 angstroms in size which are co-precipitated with the said intermetallic A M phase.
  • a nickel-base alloy as claimed in claim 1 which is essentially free of conventional precipitated gamma prime phase.
  • a method for producing a nickel-base alloy strengthened by co-precipitation of two coherent phases namely a non-cubic intermetallic phase A M and a FCC monocarbide phase MC, wherein A consists essentially of Ni, C is carbon, and M is a metal from the group consisting of Ta, Cb and V, said MC monocarbide phase being dispersed as fine particles of less than 250 angstroms in size in addition to primary particles of up to about 5 microns in size, comprising the steps of:
  • steps (d), (e) and (f) are repeated as necessary in order to attain the desired form before aging step (g) is carried out.
  • step (g) 8. A method as claimed in claim 6 wherein the material is aged in step (g) for a period of about 30 to 100 hours.

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  • Chemical & Material Sciences (AREA)
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US834942A 1969-06-20 1969-06-20 Co-precipitation-strengthened nickel base alloys and method for producing same Expired - Lifetime US3649379A (en)

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JP (1) JPS5013734B1 (fr)
AT (1) AT304890B (fr)
BE (1) BE752287A (fr)
CA (1) CA930574A (fr)
DE (1) DE2029963A1 (fr)
FR (1) FR2047000B1 (fr)
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Cited By (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3985582A (en) * 1973-07-30 1976-10-12 Office National D'etudes Et De Recherches Aerospatiales (O.N.E.R.A.) Process for the improvement of refractory composite materials comprising a matrix consisting of a superalloy and reinforcing fibers consisting of a metal carbide
US4574015A (en) * 1983-12-27 1986-03-04 United Technologies Corporation Nickle base superalloy articles and method for making
US4579602A (en) * 1983-12-27 1986-04-01 United Technologies Corporation Forging process for superalloys
US4769087A (en) * 1986-06-02 1988-09-06 United Technologies Corporation Nickel base superalloy articles and method for making
US9828658B2 (en) 2013-08-13 2017-11-28 Rolls-Royce Corporation Composite niobium-bearing superalloys
US9938610B2 (en) 2013-09-20 2018-04-10 Rolls-Royce Corporation High temperature niobium-bearing superalloys
CN117051342A (zh) * 2023-07-27 2023-11-14 北京科技大学 一种优质gh4738合金大尺寸铸锭的预处理工艺及设计方法
CN118650392A (zh) * 2024-08-21 2024-09-17 贵州航飞精密制造有限公司 一种狭窄空间用扩口螺栓的加工方法

Families Citing this family (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS57207145A (en) * 1981-06-15 1982-12-18 Toshiba Corp Wear resistant alloy

Cited By (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3985582A (en) * 1973-07-30 1976-10-12 Office National D'etudes Et De Recherches Aerospatiales (O.N.E.R.A.) Process for the improvement of refractory composite materials comprising a matrix consisting of a superalloy and reinforcing fibers consisting of a metal carbide
US4574015A (en) * 1983-12-27 1986-03-04 United Technologies Corporation Nickle base superalloy articles and method for making
US4579602A (en) * 1983-12-27 1986-04-01 United Technologies Corporation Forging process for superalloys
US4769087A (en) * 1986-06-02 1988-09-06 United Technologies Corporation Nickel base superalloy articles and method for making
US9828658B2 (en) 2013-08-13 2017-11-28 Rolls-Royce Corporation Composite niobium-bearing superalloys
US9938610B2 (en) 2013-09-20 2018-04-10 Rolls-Royce Corporation High temperature niobium-bearing superalloys
CN117051342A (zh) * 2023-07-27 2023-11-14 北京科技大学 一种优质gh4738合金大尺寸铸锭的预处理工艺及设计方法
CN118650392A (zh) * 2024-08-21 2024-09-17 贵州航飞精密制造有限公司 一种狭窄空间用扩口螺栓的加工方法

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NL7009044A (fr) 1970-12-22
FR2047000B1 (fr) 1973-01-12
CA930574A (en) 1973-07-24
FR2047000A1 (fr) 1971-03-12
DE2029963A1 (de) 1970-12-23
BE752287A (fr) 1970-12-01
JPS5013734B1 (fr) 1975-05-22
GB1320775A (en) 1973-06-20
AT304890B (de) 1973-01-25

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