US3642543A - Thermomechanical strengthening of the superalloys - Google Patents

Thermomechanical strengthening of the superalloys Download PDF

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US3642543A
US3642543A US864268A US3642543DA US3642543A US 3642543 A US3642543 A US 3642543A US 864268 A US864268 A US 864268A US 3642543D A US3642543D A US 3642543DA US 3642543 A US3642543 A US 3642543A
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temperature
alloy
aging
gamma
phase
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William A Owczaraki
John M Oblak
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RTX Corp
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United Aircraft Corp
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    • 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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  • the superalloys are strengthened in a process involving both thermal and defonnational treatments under controlled conditions.
  • the method is particularly effective for the nickel-base superalloys of the 'yy'-type having a volume fraction of the 7' phase exceeding about 25 percent at room temperature, and for the superalloys precipitating the topologically closepacked phases such as the sigma phase. it relies on the establishment of a microstructure wherein the 7 phase is precipitated in a uniformly distributed array having an interparticle spacing not exceeding about 5 microns; warm working the alloy to effect an area reduction of at least 15 percent; and subjecting the alloy to a stabilization heat treatment.
  • the strength increase is attributable to a particular thermally and mechanically stable array of microcrystalline imperfections thus established and, in those alloys precipitating the sigma phase, also by an altered sigma phase morphology.
  • the present invention is most conveniently characterized as a metal processing technique and it is particularly adapted to improving the mechanical properties of the nickel-base and cobalt-base superalloys.
  • the superalloys are, in general, those alloys which display very high strengths at very high temperatures and, thus, which have significant utility in the fabrication of gas turbine engine components.
  • the typical nickel-base superalloy of this type for example, is essentially a nickel-chromium solid solution (7 phase) hardened by the additions of elements such as aluminum and titanium to precipitate an intermetallic compound (7' phase).
  • the usual intermetallic compound which is represented by the formula Ni (Al, Ti), is an ordered facecentered-cubic structure with aluminum and titanium at the corners of the unit cell and nickel at the face centers.
  • These alloys also normally contain cobalt to raise the solvus temperature of the 7' phase, refractory metal additions for solution and carbon, boron and zirconium to promote ductility and fabricability. In the monocrystal form, these alloys may have reduced quantities of carbon to prevent crackproneness associated with the formulation of MC-type carbides.
  • alloys to which the present invention has particular applicability are those nickel-base superalloys having a quantity of the 'y' precipitate exceeding about 25 volume percent at the hot working temperatures and which is stable within the matrix at this same temperature.
  • alloys of this type are those identified in the industry as follows:
  • thermomechanical strengthening sequence of the present invention is the outcome of one such program.
  • the phase known as the sigma phase (0') is a particularly well known example of a precipitate of this general type and many of the commercially available superalloys, such as [NCO 901 and [NCO 718, are specifically formulated to avoid the formation of these undesirable phases.
  • the sigma forms as hard, brittle platelets which provide the natural sites for mechanical weakness and, in addition, compete with the 7 matrix phase for the solid solution strengthening elements.
  • a detailed discussion of the sigma phase may be found in an article by E. 0. Hall et al. The Institute for Metals, Vol. 1 1 (1966) p. 61.
  • Sigma belongs to a class of intermetallic phases identified as topologically close-packed (TCP) phases which typically form in a platelike morphology in a size comparable to the alloy grain size.
  • TCP phases such as the sigma, chi or mu phases, occur in both the nickel-base and cobalt-base alloys.
  • the present invention provides a method for improving the yield and tensile strengths and improving the creep and fatigue resistance of the advanced nickel-base superalloys by thermomechanical means.
  • the improved properties are provided by a stable array of microcrystalline imperfections established by controlled heat treatment and deformation, and preserved and stabilized by subsequent heat treatment causing further 'y' precipitation and in the case of the alloys prone to sigma formation, a finely dispersed precipitate of equiaxed sigma.
  • An initial microstructure is established by heat treatment involving solutioning plus aging to provide a microstructure consisting of a uniformly distributed y phase having an effective interparticle spacing not exceeding about 5 microns and a volume fraction not less than about 25 percent.
  • the material so heat treated is deformably processed at elevated temperatures subject to the following conditions:
  • the deformation temperature is selected to preserve the volume fraction and distribution previously established and usually corresponds to the aging temperature of step l b.
  • the deformation temperature must not exceed either the solvus temperature of the y phase or the gross 3 ,642,543 3 4 recrystallization temperature and must not be less then D RIP I N F THE PREFERRED EMBODIMENTS the minimum recovery temperature;
  • c. total deformation must be equivalent to at least about a percent reduction (but is not likely to exceed a reduction of about 60 percent).
  • a signl icant measure 0 strengt emng may e ac ieve m [M900 2300 2200 1850 18504050 either of two microcrystallme forms, the generation of which Udimet 700 2220 2100 1800 1800-1975 in the alloys of the y-y type is dependent upon the particular temperature within the aging temperature range at which the 'Coincidentwith the uppertemperaturclimit for planarslip aging and deformation is accomplished. The more significant strengthening is achieved by working the alloy nearer the lower end of the aging temperature range, i.e., nearer the Preferred, in context minimum recovery temperature.
  • No'rE.Times are those utilized for convenience and represent no limitation on the maximum or minimum allowable times.
  • the conditions are defects along these boundaries. normally selected to dissolve the maximum quantity of the y
  • the material at one temperaprecipitate into solid solution.
  • all of the 7 phase can be solutioned. With condition, or by processing at intermediate temperatures, it is MAR-M200 most, but perhaps not all, of the 7' phase is dispossible to produce both kinds of dislocation arrays in the solved.
  • FIG. 1 is a photomicrograph taken utilizing electron as MAR-M200 r where the solvus terhpel'mlh'e approaches microscopy techniques showing the microstructure of a Sohdhs temperature solhhonihg y actually be done Slightly Udimet 700 alloy sample subjected to a conventional below the solvus temperature- Strengthening h treatment M 24,0QOX) Aging of the alloy results in precipitation of the 7' phase.
  • MAR-M200 r where the solvus terhpel'mlh'e approaches microscopy techniques showing the microstructure of a Sohdhs temperature solhhonihg y actually be done Slightly Udimet 700 alloy sample subjected to a conventional below the solvus temperature- Strengthening h treatment M 24,0QOX) Aging of the alloy results in precipitation of the 7' phase.
  • FIG 2 i a photomicrograph f a Udi 700 alloy
  • aging may be undertaken processed to produce the polygonal substructure of the at any temperature above the mihhhhm recover) temperature present invention (24,Q()() or that temperature above which nonplanar slip occurs, and
  • FIG. 3 is a photomicrograph of the same alloy processed to below the SOh/hs temperature of the 7' p Aging and Workproduce the warm-worked substructure of the present invenihg hear the upper end of this range Promotes formation of 3 tion, (24,000X) polygonal substructure, while aging and working near the FIG. 4 is a photomicrograph f a sigma prone ll lower end ofthe range results in a warm-worked microcrystalprocessed conventionally. (500x) line array.
  • FIG. 5 is a photomicrograph f h ll f FIG 4 processed established, particularly in a temperature-time relationship, to according to the present invention.
  • NUS. 0 anti I are graphs computing the hardness oi the D l'orlmlilon is ur 'mhly undertaken M numlr-ml rm treatments prone alloys as conventionally proc ssed and n lure at which the alloy has been aged. However. it may he rm processed according to the present invention. dertaken at any elevated temperature within the aging range provided the microstructure established in aging is essentially maintained. The total deformation must exceed that equivalent to about a 15 percent reduction of area, and the strength increase which finally results is usually achieved in the range of a 15-60 percent reduction. This is not to say that further reduction cannot be made without deleterious effect, but rather that the maximum advantages in physical properties will have been achieved at the point where a 60 percent reduction has been achieved.
  • Deformation is normally accomplished with a 5-10 percent reduction per pass with reheat between passes to reestablish the temperature. With greater degrees of deformation, particularly when imposed in a limited number of passes, the synergistic effects of working plus external heat are necessarily considered It is desirable to prevent precipitation simultaneously with the working process, hence, the desirability of working the alloy at the same temperature at which it was aged is established. And, as previously mentioned, the selection of the aging-working temperatures is additionally influenced by the particular form of strengthening desired, i.e., the polygonal or warm-worked microstructure or combination thereof.
  • the final postwork heat treatment comprises a normal stabilization and precipitation aging.
  • the temperature in the postwork heat treatment sequence must not exceed the temperature of deformation and final aging. Its purpose is to promote the final arrangement of the microcrystal imperfections into a stable array in the matrix through the further precipitation of the 7' phase and, in alloys containing more than about 0.05 weight percent carbon, to precipitate intragranular carbides. Both of these subsequent precipitation events further stabilize the microdefect array and additionally strengthen by normal precipitation hardening.
  • EXAMPLE 2 A monocrystal specimen formed from low carbon Udimet 700 was thermomechanically treated to produce a polygonal subslitture as follows: 7
  • the 0.2 percent yield strength was 152 k.s.i.
  • the 0.2 percent yield strength is 121 k.s.i.
  • the comparable lifetime of an untreated monocrystal of the same composition is about 900 cycles, that of an untreated polycrystalline sample of the same composition about 200 cycles.
  • the 0.2 percent yield strength of this sample was 173 k.s.i. which compares to a k.s.i. yield strength for the untreated alloy. in low cycle fatigue at room temperature with a 1.5 percent strain amplitude, life was 3.038 cycles as compared to a lifetime of about 700 cycles for the untreated monocrystal.
  • the physical property improvements verified by the foregoing data were achieved without change in alloy composition by the provision of a specific substructure by thermomechanical treatment of the alloy.
  • the substructure size is of the order of 5 microns or less and was achieved because the original particle size and distribution of the 'y' precipitate, as established by heat treatment, provided a network of particulate barriers with a spacing of 5 microns or less.
  • the strength increase established is maintained as long as the substructure is maintained, and, hence, is evident up to at least the minimum recovery temperature for the particular alloy involved.
  • Bar A was then aged at l,975 F. for 4 hours to precipitate the y phase. It was then swaged to a 60 percent reduction in area at 1.975 F. using a reduction of 6 percent per pass with a minute reheat between passes. No difficulty was experienced during the working operation.
  • Bar B was given an 1,875 F. age for 4 hours. Swaging at 1,875 F. resulted in severe cracking during the first swaging pass. The l,875 F. age was found to promote the precipitation of both 7' and the platelike sigma.
  • the hardness of the warm-worked material was significantly higher than that of conventional nickel-base superalloys. As shown in FIG. 6, hardness values in the R 50s were obtained after aging at temperatures below 1,700 F. For comparison, the hardness of fully heated treated Udimet 700 is W R 39. In fact, the hardness of the sigma containing material is nearly equivalent to the low R 60 values of high-speed tool steel. For this alloy, a maximum hardness of R, 57.5 was obtained after a 24 hour age at 1,400 F. The Vickess hardness curves of FIG. 7 indicate that the peak hardness after a 4 hour age is nearer l,500 F. The shift in peak hardness to l,400 F. with increasing aging time indicates that any softening mechanisms, such as recovery or coarsening, occurring at l,400 F., are more then compensated, presumably by continued precipitation of the sigma phase.
  • the yield strength of the warmworked material in compression was significantly higher than that of the conventionally processed nickel-base superalloys, as shown in Table IV.
  • Table IV At room temperature, the ductility of the hardest warm-worked specimen is above that of tool steel but the yield strength is 100,000 p.s.i. lower.
  • the strength of the tool steel has dropped significantly while the warm-worked superalloy has lost very little strength.
  • the yield strengths of the two alloys are about equal but the ductility of the tool steel is inferior.
  • the present invention has revealed that not only may the detrimental effects of sigma be alleviated in the superalloys, but also significant improvements may be provided therefrom.
  • the presence of substantial quantities of dispersed, equiaxed sigma in the superalloys may be utilized to enhance strength and to provide high hardness without embrittlement. This suggests that these modified alloys may have utility as high temperature bearing alloys for, while high speed tool steels are rarely used in bearing applications above 600 F. and are limited to temperatures below l,000 F., the equiaxed sigma superalloys are metallurgically stable to temperatures well over 1,000 F.
  • heat treating the alloy including the step of aging above the minimum recovery temperature ofthe alloy, to establish a microstructure having the y precipitate in a stable homogenous distribution having an effective interparticle spacing not exceeding about 5 microns and a volume fraction of the precipitate not less than about 25 percent at the working temperature; warm-working the alloy to effect a reduction in area of at least 15 percent while maintaining essentially the same 7' phase morphology established in the prior heat treatment, providing a microdefect array of regular geometry; and
  • the initial heat treatment includes aging at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.
  • the initial heat treatment includes aging at a temperature sufficiently close to the minimum recovery temperature of the alloy to promote formation upon warm-working of the microdefect array comprising a warm-worked metallurgical substructure consisting of a randomly nonoriented homogenous dislocation distribution.
  • the aging in the initial heat treatment is also above the precipitation temperature for the topologically closepacked phases.
  • the method of processing the nickel-base superalloys of the 7-7 type having a quantity of the 7' phase at room temperature exceeding about 25 volume percent and, optionally, a sigma phase precipitate which comprises the steps of: heat treating the alloy to solution the precipitate; aging the alloy to precipitate the 7' phase to a minimum of about 25 volume percent in a stable homogenous dis tribution having an effective interparticle spacing not exceeding about 5 microns at a temperature sufficiently high to prevent substantial sigma phase precipitation;
  • the alloy is worked about the aging temperature.
  • the alloy is worked to effect an area reduction of l560 percent.
  • the alloy is aged at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon warm-working of the regular microdefect array comprising a polygonal subcell structure.
  • aging of the alloy at a temperature between the solvus temperature of the 7/ phase and the minimum recovery temperature to reprecipitate at least 25 volume percent of the 7 phase, based on the overall alloy composition, in a stable uniform distribution at an effective interparticle spacing not exceeding about 5 microns; working the aged alloy at about the aging temperature to effect a deformation corresponding to at least a 15 percent area reduction while maintaining essentially the same volume percent and distribution of the 7' phase established in the aging process, providing a regular array of microcrystalline imperfections; and heat treating the alloy at a temperature not exceeding the temperature of aging and working to cause precipitation of that portion of the y phase remaining in solution after aging and to cause, in those alloys containing more than 0.05 weight percent carbon, precipitation of intragranular carbides, to provide a thermally and mechanically stable array of microcrystalling imperfections.
  • aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation upon working of a polygonal metallurgical substructure having the dis ocations aligned at subcell bounthe 'y'y' type enriched in those elements promoting the precipitation of topologically close-packed phases which comprises the steps of:
  • aging is performed at a temperature sufficiently above the minimum recovery temperature of the alloy to promote formation of a polygonal metallurgical substructure having the dislocations aligned at subcell boundaries to provide a regular array of defects along these boundaries.
  • aging is performed at a temperature sufficiently close to the minimum recovery temperature for the alloy to promote formation of warm-worked metallurgical substructure comprising a randomly nonoriented homogenous dislocation distribution.
  • the final heat treatment includes a sequential heat treatment, one heat treatment being selected to precipitate the topologically close-packed phases and another being selected to precipitate the 7' phase.

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US864268A 1969-09-26 1969-09-26 Thermomechanical strengthening of the superalloys Expired - Lifetime US3642543A (en)

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Cited By (15)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3816920A (en) * 1972-11-30 1974-06-18 Gillette Co Novel cutting edges and processes for making them
WO1982000477A1 (en) * 1980-08-11 1982-02-18 United Technologies Corp Heat treated single crystal articles and process
US4514360A (en) * 1982-12-06 1985-04-30 United Technologies Corporation Wrought single crystal nickel base superalloy
US4528048A (en) * 1982-12-06 1985-07-09 United Technologies Corporation Mechanically worked single crystal article
US4908069A (en) * 1987-10-19 1990-03-13 Sps Technologies, Inc. Alloys containing gamma prime phase and process for forming same
US5074925A (en) * 1990-06-25 1991-12-24 The United States Of America As Represented By The Secretary Of The Air Force Thermomechanical fabrication of net shape single crystal airfoils
US5169463A (en) * 1987-10-19 1992-12-08 Sps Technologies, Inc. Alloys containing gamma prime phase and particles and process for forming same
US5470371A (en) * 1992-03-12 1995-11-28 General Electric Company Dispersion strengthened alloy containing in-situ-formed dispersoids and articles and methods of manufacture
US5534085A (en) * 1994-04-26 1996-07-09 United Technologies Corporation Low temperature forging process for Fe-Ni-Co low expansion alloys and product thereof
US5725692A (en) * 1995-10-02 1998-03-10 United Technologies Corporation Nickel base superalloy articles with improved resistance to crack propagation
RU2119842C1 (ru) * 1996-06-21 1998-10-10 Институт проблем сверхпластичности металлов РАН Способ изготовления осесимметричных деталей и способ получения заготовок для его осуществления (варианты)
US5820700A (en) * 1993-06-10 1998-10-13 United Technologies Corporation Nickel base superalloy columnar grain and equiaxed materials with improved performance in hydrogen and air
US6328827B1 (en) * 1994-07-13 2001-12-11 Societe Nationale d'Etude et de Construction de Moteurs d'Aviation “SNECMA” Method of manufacturing sheets made of alloy 718 for the superplastic forming of parts therefrom
CN114214583A (zh) * 2021-12-16 2022-03-22 西北工业大学 一种高效强化镍基高温合金的时效热处理工艺
FR3117506A1 (fr) * 2020-12-16 2022-06-17 Safran Aircraft Engines Procede de fabrication d'une piece en superalliage monocristallin

Families Citing this family (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4116723A (en) * 1976-11-17 1978-09-26 United Technologies Corporation Heat treated superalloy single crystal article and process
CH654593A5 (de) * 1983-09-28 1986-02-28 Bbc Brown Boveri & Cie Verfahren zur herstellung eines feinkoernigen werkstuecks aus einer nickelbasis-superlegierung.
FR2557145B1 (fr) * 1983-12-21 1986-05-23 Snecma Procede de traitements thermomecaniques pour superalliages en vue d'obtenir des structures a hautes caracteristiques mecaniques
US4579602A (en) * 1983-12-27 1986-04-01 United Technologies Corporation Forging process for superalloys
US4574015A (en) * 1983-12-27 1986-03-04 United Technologies Corporation Nickle base superalloy articles and method for making
JPS6153181U (enrdf_load_stackoverflow) * 1984-09-13 1986-04-10
US4957567A (en) * 1988-12-13 1990-09-18 General Electric Company Fatigue crack growth resistant nickel-base article and alloy and method for making

Citations (1)

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Publication number Priority date Publication date Assignee Title
US3147155A (en) * 1961-08-02 1964-09-01 Int Nickel Co Hot-working process

Patent Citations (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3147155A (en) * 1961-08-02 1964-09-01 Int Nickel Co Hot-working process

Cited By (18)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3816920A (en) * 1972-11-30 1974-06-18 Gillette Co Novel cutting edges and processes for making them
US4328045A (en) * 1978-12-26 1982-05-04 United Technologies Corporation Heat treated single crystal articles and process
WO1982000477A1 (en) * 1980-08-11 1982-02-18 United Technologies Corp Heat treated single crystal articles and process
US4392894A (en) * 1980-08-11 1983-07-12 United Technologies Corporation Superalloy properties through stress modified gamma prime morphology
US4514360A (en) * 1982-12-06 1985-04-30 United Technologies Corporation Wrought single crystal nickel base superalloy
US4528048A (en) * 1982-12-06 1985-07-09 United Technologies Corporation Mechanically worked single crystal article
US5169463A (en) * 1987-10-19 1992-12-08 Sps Technologies, Inc. Alloys containing gamma prime phase and particles and process for forming same
US4908069A (en) * 1987-10-19 1990-03-13 Sps Technologies, Inc. Alloys containing gamma prime phase and process for forming same
US5074925A (en) * 1990-06-25 1991-12-24 The United States Of America As Represented By The Secretary Of The Air Force Thermomechanical fabrication of net shape single crystal airfoils
US5470371A (en) * 1992-03-12 1995-11-28 General Electric Company Dispersion strengthened alloy containing in-situ-formed dispersoids and articles and methods of manufacture
US5820700A (en) * 1993-06-10 1998-10-13 United Technologies Corporation Nickel base superalloy columnar grain and equiaxed materials with improved performance in hydrogen and air
US5534085A (en) * 1994-04-26 1996-07-09 United Technologies Corporation Low temperature forging process for Fe-Ni-Co low expansion alloys and product thereof
US6328827B1 (en) * 1994-07-13 2001-12-11 Societe Nationale d'Etude et de Construction de Moteurs d'Aviation “SNECMA” Method of manufacturing sheets made of alloy 718 for the superplastic forming of parts therefrom
US5725692A (en) * 1995-10-02 1998-03-10 United Technologies Corporation Nickel base superalloy articles with improved resistance to crack propagation
US5788785A (en) * 1995-10-02 1998-08-04 United Technology Corporation Method for making a nickel base alloy having improved resistance to hydrogen embittlement
RU2119842C1 (ru) * 1996-06-21 1998-10-10 Институт проблем сверхпластичности металлов РАН Способ изготовления осесимметричных деталей и способ получения заготовок для его осуществления (варианты)
FR3117506A1 (fr) * 2020-12-16 2022-06-17 Safran Aircraft Engines Procede de fabrication d'une piece en superalliage monocristallin
CN114214583A (zh) * 2021-12-16 2022-03-22 西北工业大学 一种高效强化镍基高温合金的时效热处理工艺

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DE2046409A1 (de) 1971-04-01
BE756653A (fr) 1971-03-01
FR2062845A5 (enrdf_load_stackoverflow) 1971-06-25
JPS4916012B1 (enrdf_load_stackoverflow) 1974-04-19
NL7014192A (enrdf_load_stackoverflow) 1971-03-30
CH576525A5 (enrdf_load_stackoverflow) 1976-06-15

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