CA2069557A1 - Cr-bearing gamma titanium aluminides and method of making same - Google Patents
Cr-bearing gamma titanium aluminides and method of making sameInfo
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- CA2069557A1 CA2069557A1 CA002069557A CA2069557A CA2069557A1 CA 2069557 A1 CA2069557 A1 CA 2069557A1 CA 002069557 A CA002069557 A CA 002069557A CA 2069557 A CA2069557 A CA 2069557A CA 2069557 A1 CA2069557 A1 CA 2069557A1
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- 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
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C14/00—Alloys based on titanium
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- 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
- C22C32/0047—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 carbides, nitrides, borides or silicides as the main non-metallic constituents
- C22C32/0073—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 carbides, nitrides, borides or silicides as the main non-metallic constituents only borides
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- Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
- Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
- Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
- Y10T428/00—Stock material or miscellaneous articles
- Y10T428/12—All metal or with adjacent metals
- Y10T428/12014—All metal or with adjacent metals having metal particles
- Y10T428/12028—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, etc.]
- Y10T428/12049—Nonmetal component
- Y10T428/12056—Entirely inorganic
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Abstract
Cr-BEARING GAMMA TITANIUM ALUMINIDES
AND METHOD OF MAKING SAME
Abstract of the Disclosure An article comprises a Cr-bearing, predominantly gamma titanium aluminide matrix including second phase dispersoids, such as TiB2, in an amount effective to increase both the strength and the ductility of the matrix.
AND METHOD OF MAKING SAME
Abstract of the Disclosure An article comprises a Cr-bearing, predominantly gamma titanium aluminide matrix including second phase dispersoids, such as TiB2, in an amount effective to increase both the strength and the ductility of the matrix.
Description
Cr-BEARING GAMMA TITANIUM ALUMINIDES
AND METHOD OF MAKING SAME
Field of the Invention The present invention relates to alloys of titanium and aluminum and, more particularly, to cr-bearing, predominantly gamma titanium aluminides thatexhibit an increase in both strength and ductility upon inclusion of second phase dispersoids therein.
Backaround of the Invention For the past several years, extensive research has been devoted to the development of intermetallic materials, such as titanium aluminides, for use in the manufacture of light weight structural components capable of withstanding high temperatures/stresses. Such components are represented, for example, by blades, vanes, disks, shafts, casings, and other components of the turbine section of modern gas turbine engines where higher gas and resultant component temperatures are desired to increase engine thrust/efficiency or other applications requiring lightweight high temperature materials .
AND METHOD OF MAKING SAME
Field of the Invention The present invention relates to alloys of titanium and aluminum and, more particularly, to cr-bearing, predominantly gamma titanium aluminides thatexhibit an increase in both strength and ductility upon inclusion of second phase dispersoids therein.
Backaround of the Invention For the past several years, extensive research has been devoted to the development of intermetallic materials, such as titanium aluminides, for use in the manufacture of light weight structural components capable of withstanding high temperatures/stresses. Such components are represented, for example, by blades, vanes, disks, shafts, casings, and other components of the turbine section of modern gas turbine engines where higher gas and resultant component temperatures are desired to increase engine thrust/efficiency or other applications requiring lightweight high temperature materials .
- 2 ~ 7 - P-311 Howmet 2 Intermetallic materials, such as gamma titanium aluminide, exhibit improved high temperature mechanical properties, including high strength-to-weight ratios, and oxidation resistance relative to conventional high temperature titanium alloys. However, general exploitation of these intermetallic materials has been limited by the lack of strength, room temperature ductility and toughness, as well as the technical challenges associated with processing and fabricating the matelrial into the complex end-use shapes that are exemplified, for example, by the aforementioned turbine components.
The Kampe et al U.S. Patent 4,915,905 issued April 10, 1990 describes in detail the development of various metallurgical processing techniques for improving the low (room) temperature ductility and toughness of intermetallic materials and increasing their high temperature strength. The Kampe et al '905 patent relates to the rapid solidification of metallic matrix composites. In particular, in this patent, an intermetallic-second phase composite is formed; for example, by reacting second phase-forming constituents P-311 Howmet 3 '~
ln the presence of a solvent metal, to form in-situ precipitated second phase particles, such as boride dispersoids, within an intermetallic-containing matrix, such as titanium aluminide. The intermetallic-second phase composite is then subjected to rapid solidification to produce a rapidly solidified composite. Thus, for example, a composite comprising in-situ precipitated TiB2 particles within a titanium aluminide matrix may be formed and then rapidly solidified to produce a rapidly solidified powder of the composite. The powder is then consolidated by such consolidation techniques as hot isostatic pressing, hot extrusion and superplastic forging to provide near-final (i.e., near-net) shapes.
U.S. Patent 4,836,982 to Brupbacher et al also relates to the rapid solidification of metal matrix composites wherein second phase-forming constituents are reacted in the presence of a solvent metal to form in-situ precipitated second phase particles, such as TiB2 or TiC, within the solvent metal, such as aluminum.
P-311 Howmet 4 2 ~
U.S. Patents 4,774,052 and 4,916,029 to Nagle et al are specifically directed toward the production of metal matrix-second phase composites in which the metallic matrix comprises an intermetallic material, such as titanium aluminide. In one embodiment, a first composite is formed which comprises a dispersion of second phase particles, such as TiB2, within a metal or alloy matrix, such as Al.
This composite is then introduced into an additional metal which is reactive with the matrix to form an intermetallic matrix. For example, a first composite comprising a dispersion of TiB2 particles within an Al matrix may be introduced into molten titanium to form a final composite comprising TiB2 dispersed within a titanium aluminide matrix. U.S. Patent 4,915,903 to Brupbacher et al describes a modification of the methods taught in the aforementioned Nagle et al patents.
U.S. Patents 4,751,048 and 4,916,030 to Christodalou et al relate to the production of metal matrix-second phase composites wherein a first composite which comprises second phase particles ~ P-311 Howmet 5 2~3~ ~
dispersed in a metal matrix is diluted in an additional amount of metal to form a final composite of lower second phase loading. For example, a first composite comprising a dispersion of TiB2 particles within an Al matrix may be introduced into molten titanium to form a final composite comprising TiB2 dispersed within a titanium aluminide matrix.
U.S. Patent 3,203,794 to Jaffee et al relates to gamma TiAl alloys which are said to maintain hardness and resistance to oxidation at elevated temperatures. The use of alloying additions such as In, Bi, Pb, Sn, Sb, Ag, C, 0, Mo, V, Nb, Ta, Zr, Mn, Cr, Fe, W, Co, Ni, Cu, Si, Be, B, Ce, As, S, Te and P is disclosed. However, such additions are said to lower the ductility of the TiAl binary alloys.
An attempt to improve room temperature ductility by alloying intermetallic materials with one or more metals in combination with certain plastic forming techniques is disclosed in the Blackburn U.S.
Patent 4,294,615 wherein vanadium was added to a TiAl composition to yield a modified composition of Ti-31 P-311 Howmet to 36% Al-0 to 4% V (percentages by weight). The modified composition was melted and isothermally forged to shape in a heated die at a slow deformation rate necessitated by the dependency of ductility of the intermetallic material on strain rate. The isothermal forging process is carried out at above 1000C such that special die materials (e.g., a Mo alloy known as TZM) must be used. Generally, it is extremely difficult to process TiAl intermetallic materials in this way as a result of their high temperature properties and the dependence of their ductility on strain rate.
A series of U.S. patents comprising U.S.
Patents 4,836,983; 4,842,817; 4,842,819; 4,842,820;
4,857,268; 4,879,092; 4,897,127; 4,902,474; and 4,916,028, have described attempts to make gamma TiAl intermetallic materials having both a modified stoichiometric ratio of Ti/A1 and one or more alloyant additions to improve room temperature strength and ductility. The addition of Cr alone or with Nb, or with Nb and C, is described in the '819; '092 and '028 patents. In making cylindrical shapes from these P-311 Howmet 7 modified compositions, the alloy was typically first made into an ingot by electro-arc melting. The ingot was melted and melt spun to form rapidly solidified ribbon. The ribbon was placed in a suitable container and hot isostatically pressed (HIP'ped) to form a consolidated cylindrical plug. The plug was placed axially into a central opening of a billet and sealed therein. The billet was heated to 975C for 3 hours and extruded through a die to provide a reduction of about 7 to 1. Samples from the extruded plug were removed from the billet and heat treated and aged.
U.S. Patent 4,916,028 (included in the series of patents listed above) also refers to processing the TiAl base alloys as modified to include C, Cr and Nb additions by ingot metallurgy to achieve desirable combinations of ductility, strength and other properties at a lower processing cost than the aforementioned rapid solidification approach. In particular, the ingot metallurgy approach described in the '028 patent involves melting the modified alloy and solidifying it into a hockey puck-shaped ingot Oe simple ~eometry and small size (e.g., 2 inches in P-311 Howmet 8 C~ r~
diameter and .5 inch thick), homogenizing the ingot at 1250C for 2 hours, enclosing the ingot in a steel annulus, and then hot forging the annulus/ring assembly to provide a 50% reduction in ingot thickness. Tensile specimens cut from the ingot were annealed at various temperatures above 1225C prior to tensile testing. Tensile specimens prepared by this ingot metallurgy approach exhibited lower yield strengths but greater ductility than specimens prepared by the rapid solidification approach.
Despite the attempts described hereabove to improve the ductility and strength of intermetallic materials, there is a continuing desire and need in the high performance material-using industries, especially in the gas turbine engine industry, for intermetallic materials which have improved properties or combinationæ of properties and which are amenable to fabrication into usable, complex engineered end-use shapes on a relatively high volume basis at a relatively low cost. It is an object of the present invention to satisfy these desires and needs.
Summarv of the Invention In one embodiment, the present invention e~ ~ ~
P-311 Howmet 9 involves a titanium aluminide article, as well as method of making the article, wherein both the strength and ductility thereof can be increased by virtue of the inclusion of second phase dispersoids in a Cr-bearing, predominantly gamma titanium aluminide matrix. To this end, second phase dispersoids, such as, for example, TiB2, in an amount of about 0.5 to about 20.0 volume %, preferably about 0.5 to about 7.0 volume %, are included in a predominantly gamma titanium aluminide matrix including from about 0.5 to about 5.0 atomic % Cr, preferably from about 1.0 to about 3.0 atomic % Cr.
In another embodiment, the invention involves a titanium aluminum alloy consisting essentially of (in atomic %) about 40 to about 52% Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn, and about 0.5 about 5.0% Cr. A preferred alloy consists essentially of (in atomic %) about 41 to about 50% Ti, about 46% to 49% Al, about 1% to about 3% Mn, about 1% to about 3% Cr, up to abcut 3% V and up to about 3% Nb. Second phase dispersoids may be included in the alloy in an amount of about 0.5 to about 20.0 volume % to increase strength.
Unexpectedly, the titanium aluminide alloy exhibits an increase in ductility as well as strength upon the 2 ~
P-311 Howmet 10 inclusion of tne second phase dispersoids therein.
Brief Descri~tion of the Drawings Figures la and lb are bar graphs illustrating the change in strength and ductility of Cr-bearing, predominantly gamma titanium aluminide alloys of the invention upon the inclusion of titanium borides. Similar data is presented for a Ti-48Al-2V-2Mn alloy ~reference alloy) to illustrate the increase in strength but the decrease in ductility observed upon inclusion of the same boride levels therein.
Figures 2a, 2b, and 2c illustrate the microstructure of the Ti-48Al-2V-2Mn reference alloy after hot isostatic pressing and heat treatment at 1650F (900C) for 16 hours.
Figure~ 3a, 3b and 3c illustrate the microstructure of the Ti-48Al-2Mn-2Cr alloy of thP
invention after the same hot isostatic pressing and heat treatment as used in Figs. 2a-2c.
Figures 4a, 4b and 4c illustrate the microstructure of the Ti-48Al-2V-2Mn-2Cr alloy of the P-311 Howmet 11 invention after the same hot isostatic pressing and heat treatment as used in Figs. 2a-2c.
Figures 5a,5b and 6a,6b illustrate the change in strength and ductility of the aforementioned alloys of Fig. 1 after different heat treatments.
Figures 7a, 7b and 7c, 7d illustrate the effect of heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, on microstructure of the Ti-48al-2Mn-2Cr alloy of the invention devoid of TiB2 dispersoids.
Figures 8a, 8b and 8c, 8d illustrate the effect of heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, on microstructure of the Ti-48al-2Mn-2Cr alloy of the invention including 7 volume % TiB2 dispersoids.
Figure 9 illustrates the change in yield strength of the aforementioned alloys of Fig. 1 with the volume % of TiB2 dispersoids.
Figure 10 illustrates the measured grain size as a function of TiB2 volume % for the aforementioned alloys.
P-311 Howmet 12 2 Q ~ 7 Detailed Descri~tion of the Invention The present invention contemplates a titanium aluminide article including second phase dispersoids (e.g., TiB2) in a Cr-bearing, predominantly gamma TiAl matrix in effective concentrations that result in an increase in both strength and ductility. In one embodiment of the invention, the alloy matrix consists essentially of, in atomic %, about 40 to about 52 % Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn and about 0.5 to about 5.0 % Cr to this end. Preferably, the alloy matrix consists essentially of, in atomic %, about 41 to about 50% Ti, about 46 to about 49% Al, about 1 to about 3% Mn, about 1 to about 3~ Cr, up to about 3% V, and up to about 3% Nb. The alloy matrix includes second phase dispersoids, such as preferably TiB2, in an amount not exceeding about 20.0 volume %.
Preferably, the second phase dispersoids are present in an amount of about 0.5 to about 12.0 volume %, more preferably from about 0.5 to about 7.0 volume %.
P-311 Howmet 13 The matrix is considered predominantly gamma in that a majority of the matrix microstructure in the as-cast or the cast/hot isostatically pressed/heat treated condition described hereafter comprises gamma phase. Alpha 2 and beta phases can also be present in minor proportions of the matrix microstructure; e.g., from about 2 to about 15 volume % of alpha 2 phase and up to about 5 volume ~ beta phase can be present.
The following Table I lists nominal and measured Cr-bearing titanium-aluminum ingot compositions produced in accordance with exemplary embodiments of the present invention. Also listed are the nominal and measured ingot composition of a Ti-48Al-2V-2Mn alloy used as a reference alloy for comparison purposes.
14 ~ , f~ ~
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P-311 Howmet 15 The dispersoids of TiB2 were provided in the ingots using a master sponge material comprising 70 weight % TiB2 in an Al matrix and available from Martin Marietta Corp., Bethesda, Md. and its licensees. The master sponge material was introduced into a titanium aluminum melt of the appropriate composition prior to casting into an investment mold in accordance with U.S. Patents 4,751,048 and 4,916,030, the teachings of which are incorporated herein by reference.
Segments of each ingot were sliced, remelted by a conventional vacuum arc remelting, to a superheat of +50F above the alloy melting temperature, and investment cast into preheated ceramic molds (600F) to form cast test bars having a diameter of 0.625 inch and a length of 6.0 inches. Each mold included a Zr203 facecoat and a plurality of Al203/Zr203 backup coats.
Following casting and removal from the investment molds, all test bars were hot isostatically pressed (HIP'ed) at 25 ksi and 2300F for 4 hours in an inert atmosphere (Ar).
Baseline mechanical tensile data were obtained using the investment cast test bars which had P-311 Howmet 16 2~9~7 been heat treated at 1650F (900C) for 16 hours following the aforementioned hot isostatic pressing operation. The TiB2 dispersoids present in the cast/HIP'ed/heat treated test bars typically had particle sizes (i.e., diameters) in the range of 0.3 to 5 microns.
The results of the tensile tests are shown in Fig. la plotted as a function of matrix alloy composition for 0, 7, and 12 volume % TiB2. From Fig.
la, it is apparent that the yield strength of all the alloys increases with the addition of 7 and 12 volume % TiB2.
However, from Fig. lb, the room temperature ductility of the Ti-48Al-2V-2Mn alloy was observed to decrease sùbstantially with the addition of these levels of TiB2 to the matrix alloy. Surprisingly, the ductility of the Cr-bearing alloys (i.e., Ti-48Al-2Mn-2Cr, Ti-48Al-2V-2Mn-2Cr and Ti-47Al-2Mn lNb-lCr) was observed to increase with the addition of these levels of TiB2, especially upon the addition of 7 volume % TiB2. Thus, for the TiAl alloys including chromium as an additional alloyant and TiB2 dispersoids, both the strength and the ductility were P-311 Howmet 17 2~3~ ~
found to increase unexpectedly.
Representative optlcal microstructures of these alloys after casting, hot isostatic pressing, and heat treatment are shown in Figs. 2a, 2b, 2c; 3a, 3b, and 3c; and 4a, 4b, and 4c. The photomicrographs illustrate that the microstructures of the alloys are predominantly lamellar (i.e., alternating lathes of gamma phase and alpha 2 phase) with some equiaxed grains residing at colony boundaries. Generally, there was little or no evidence of microstructural coarsening or other morphological transformations upon hot isostatic pressing and/or heat treatment.
The effect of longer time or higher temperature heat treatments on alloy strength and ductility are illustrated in Figs. 5a,5b and 6a,6b for heat treatments at 900C (1650F) for 50 hours (Figs.
5a,5b) and 1100C (2012F) for 16 hours (Figs. 6a,6b).
Yield strength is shown to increase with increasing percent TiB2. Moreover, increases in ductility were again noted for the Cr-bearing test bars having 7 volume % TiB2 in the matrix. In general, the 900oC
(1650F) heat treatments resulted in maximum ductility in all of the alloys shown. In the alloys of the P-311 Howmet 18 invention containing 7 and 12 volume % TiB2, maximum ductility occurred following heat treatment at 1650F
for 50 hours. In general, strength was relatively insensitive to heat treatment.
Figs. 7a,7b and 7c,7d illustrate the microstructures of alloy matrices following heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, for the Ti-48Al-2Mn-2Cr devoid of TiB2. Figs. 8a,8b and 8c,8d illustrate the alloy matrix microstructure for the same alloy with 7 volume % TiB2 after the same heat treatments. In the boride-free alloy, transformation of the matrix to a primarily equiaxed microstructure was observed after these heat treatments. On the other hand, the matrix microstructure including 7 volume % TiB2 exhibited very little change after these heat treatments, retaining a primarily lamellar microstructure.
Fig. 9 illustrates tensile yield strength as a function of dispersoid (TiB2) loading for the aforementioned alloys heat treated at 1650F for 16 hours. All alloys exhibit approximately linear increases in strength with increasing dispersoid loading (volume %). The Ti-48Al-2V-2Mn alloy P-311 Howmet 19 exhibited the strongest dependence~
Grain size analyses were performed on the alloys that had been heat treated at 1650F for 16 hours to determine the effect of dispersoid loading on grain size. Fig. 10 depicts large reductions in grain size due to the inoculative effect of the TiB2 dispersoids. A reduced sensitivity of grain size on dispersoid loading is apparent at higher volume fractions of dispersoids. The large variations in alloy grain size when no dispersoids are present appears to be a consequence primarily of the size and scale of the smaller, equiaxed grains that reside between large columnar, lamellar colonies.
The surprising increase in both strength and ductility of the Cr-Bearing, predominantly gamma titanium aluminides of Fig. 1 is also observed at elevated temperatures as illustrated in Table II
wherein investment cast, HIP'd, and heat treated (900C for 50 hours) specimens were tensile tested at ~16C.
, , P-311 Howmet 20 ~9 TABLE II
Tensile Testing a 816C
a (ksi) o (k~ elong yield ult Ti-48Al-2Mn-2Cr 49.5 56.2 18.1 Ti-48Al-2Mn-2Cr + 7 v% Ti82 45-0 52.4 22.8 Ti-48Al-2Mn-2Cr + 12 v~ TiB2 47.5 55.3 20.3 Ti-47Al-2Mn-lNb-lCr 51.9 68.0 4.9 Ti-47Al-2Mn-lNb-lCr + 7~ v% TiB2 51.2 76.5 ~12.3 The creep resistance of the Ti-47Al-2Mn-lNb-lCr alloy without and with 7 volume % TiB2 dispe_soids was evaluated at 1500F and 20.0 ksi load. The specimens were investment cast, HIP'ed, and heat treated at 900C for 50 hours. As indicated in Table III, the boride-free and boride-bearing specimens exhibited generally comparable rupture lives. The creep resistance of the Ti-47Al-2Mn-lNb-lCr alloy thus was not adversely affected by the inclusion of 7 volume % TiB2 dispersoids.
TABLE III
Creep Data at 1500F/20.0 ksi RuDture Llfe ~hrs) Ti-47Al-2Mn-lNb-lCr 96.3/111.7 Tl-47Al-2Mn-lNb-lCr + 7 v~ TL8z 102.~3/110.7 P-311 Howmet 21 ~ 7 In practicing the present invention, the concentration of Cr should not exceed about 5.0 atomic % of the TiAl alloy composition in order to provide the aforementioned predominantly gamma titanium aluminide matrix microstructure. For example, a TiAl ingot nominally comprising Ti-48Al-2V-2Mn-6Cr (measured composition, in atomic %, 44.1 Ti-45.8Al-20Mn-6.2Cr-1.9V) was prepared and investment cast, HIP'ed, and heat treated as described hereinabove for the alloys of Fig. 1. The ingot included about 7.0 volume % TiB2. Examination of the microstructure of the ingot before and after a 1650~F/16 hour heat treatment revealed volume fractions of beta phase well in excess of 5 volume %, primarily at grain (colony) boundaries and along lamellar interfaces. The heat treatment resulted in spherodization and a relatively homogeneous distribution of the beta phase in the microstructure. The heat treated alloy exhibited a tensile yield strength of about 90 ksi but a substantially reduced ductility at room temperature of only 0.15 %.
Thus, in practicing the invention the upper limit of the Cr concentration should not exceed about 5.0 atomic % of the alloy composition. On the other P-311 Howmet 22 2 ~
hand, the lower limit of the Cr concentration should be sufficient to result in an increase in both strength and ductility when appropriate amounts of dispersoids are included in the matrix. To this end, in accordance with the present invention, the Cr concentration is preferably from about 0.5 to about 5.0 atomic % of the alloy matrix, more preferably from about 1.0 to about 3.0 atomic % of the alloy matrix.
While the invention has been described in terms of specific embodiments thereof, it is not intended to be limited thereto but rather only to the extent set forth in the following claims.
The Kampe et al U.S. Patent 4,915,905 issued April 10, 1990 describes in detail the development of various metallurgical processing techniques for improving the low (room) temperature ductility and toughness of intermetallic materials and increasing their high temperature strength. The Kampe et al '905 patent relates to the rapid solidification of metallic matrix composites. In particular, in this patent, an intermetallic-second phase composite is formed; for example, by reacting second phase-forming constituents P-311 Howmet 3 '~
ln the presence of a solvent metal, to form in-situ precipitated second phase particles, such as boride dispersoids, within an intermetallic-containing matrix, such as titanium aluminide. The intermetallic-second phase composite is then subjected to rapid solidification to produce a rapidly solidified composite. Thus, for example, a composite comprising in-situ precipitated TiB2 particles within a titanium aluminide matrix may be formed and then rapidly solidified to produce a rapidly solidified powder of the composite. The powder is then consolidated by such consolidation techniques as hot isostatic pressing, hot extrusion and superplastic forging to provide near-final (i.e., near-net) shapes.
U.S. Patent 4,836,982 to Brupbacher et al also relates to the rapid solidification of metal matrix composites wherein second phase-forming constituents are reacted in the presence of a solvent metal to form in-situ precipitated second phase particles, such as TiB2 or TiC, within the solvent metal, such as aluminum.
P-311 Howmet 4 2 ~
U.S. Patents 4,774,052 and 4,916,029 to Nagle et al are specifically directed toward the production of metal matrix-second phase composites in which the metallic matrix comprises an intermetallic material, such as titanium aluminide. In one embodiment, a first composite is formed which comprises a dispersion of second phase particles, such as TiB2, within a metal or alloy matrix, such as Al.
This composite is then introduced into an additional metal which is reactive with the matrix to form an intermetallic matrix. For example, a first composite comprising a dispersion of TiB2 particles within an Al matrix may be introduced into molten titanium to form a final composite comprising TiB2 dispersed within a titanium aluminide matrix. U.S. Patent 4,915,903 to Brupbacher et al describes a modification of the methods taught in the aforementioned Nagle et al patents.
U.S. Patents 4,751,048 and 4,916,030 to Christodalou et al relate to the production of metal matrix-second phase composites wherein a first composite which comprises second phase particles ~ P-311 Howmet 5 2~3~ ~
dispersed in a metal matrix is diluted in an additional amount of metal to form a final composite of lower second phase loading. For example, a first composite comprising a dispersion of TiB2 particles within an Al matrix may be introduced into molten titanium to form a final composite comprising TiB2 dispersed within a titanium aluminide matrix.
U.S. Patent 3,203,794 to Jaffee et al relates to gamma TiAl alloys which are said to maintain hardness and resistance to oxidation at elevated temperatures. The use of alloying additions such as In, Bi, Pb, Sn, Sb, Ag, C, 0, Mo, V, Nb, Ta, Zr, Mn, Cr, Fe, W, Co, Ni, Cu, Si, Be, B, Ce, As, S, Te and P is disclosed. However, such additions are said to lower the ductility of the TiAl binary alloys.
An attempt to improve room temperature ductility by alloying intermetallic materials with one or more metals in combination with certain plastic forming techniques is disclosed in the Blackburn U.S.
Patent 4,294,615 wherein vanadium was added to a TiAl composition to yield a modified composition of Ti-31 P-311 Howmet to 36% Al-0 to 4% V (percentages by weight). The modified composition was melted and isothermally forged to shape in a heated die at a slow deformation rate necessitated by the dependency of ductility of the intermetallic material on strain rate. The isothermal forging process is carried out at above 1000C such that special die materials (e.g., a Mo alloy known as TZM) must be used. Generally, it is extremely difficult to process TiAl intermetallic materials in this way as a result of their high temperature properties and the dependence of their ductility on strain rate.
A series of U.S. patents comprising U.S.
Patents 4,836,983; 4,842,817; 4,842,819; 4,842,820;
4,857,268; 4,879,092; 4,897,127; 4,902,474; and 4,916,028, have described attempts to make gamma TiAl intermetallic materials having both a modified stoichiometric ratio of Ti/A1 and one or more alloyant additions to improve room temperature strength and ductility. The addition of Cr alone or with Nb, or with Nb and C, is described in the '819; '092 and '028 patents. In making cylindrical shapes from these P-311 Howmet 7 modified compositions, the alloy was typically first made into an ingot by electro-arc melting. The ingot was melted and melt spun to form rapidly solidified ribbon. The ribbon was placed in a suitable container and hot isostatically pressed (HIP'ped) to form a consolidated cylindrical plug. The plug was placed axially into a central opening of a billet and sealed therein. The billet was heated to 975C for 3 hours and extruded through a die to provide a reduction of about 7 to 1. Samples from the extruded plug were removed from the billet and heat treated and aged.
U.S. Patent 4,916,028 (included in the series of patents listed above) also refers to processing the TiAl base alloys as modified to include C, Cr and Nb additions by ingot metallurgy to achieve desirable combinations of ductility, strength and other properties at a lower processing cost than the aforementioned rapid solidification approach. In particular, the ingot metallurgy approach described in the '028 patent involves melting the modified alloy and solidifying it into a hockey puck-shaped ingot Oe simple ~eometry and small size (e.g., 2 inches in P-311 Howmet 8 C~ r~
diameter and .5 inch thick), homogenizing the ingot at 1250C for 2 hours, enclosing the ingot in a steel annulus, and then hot forging the annulus/ring assembly to provide a 50% reduction in ingot thickness. Tensile specimens cut from the ingot were annealed at various temperatures above 1225C prior to tensile testing. Tensile specimens prepared by this ingot metallurgy approach exhibited lower yield strengths but greater ductility than specimens prepared by the rapid solidification approach.
Despite the attempts described hereabove to improve the ductility and strength of intermetallic materials, there is a continuing desire and need in the high performance material-using industries, especially in the gas turbine engine industry, for intermetallic materials which have improved properties or combinationæ of properties and which are amenable to fabrication into usable, complex engineered end-use shapes on a relatively high volume basis at a relatively low cost. It is an object of the present invention to satisfy these desires and needs.
Summarv of the Invention In one embodiment, the present invention e~ ~ ~
P-311 Howmet 9 involves a titanium aluminide article, as well as method of making the article, wherein both the strength and ductility thereof can be increased by virtue of the inclusion of second phase dispersoids in a Cr-bearing, predominantly gamma titanium aluminide matrix. To this end, second phase dispersoids, such as, for example, TiB2, in an amount of about 0.5 to about 20.0 volume %, preferably about 0.5 to about 7.0 volume %, are included in a predominantly gamma titanium aluminide matrix including from about 0.5 to about 5.0 atomic % Cr, preferably from about 1.0 to about 3.0 atomic % Cr.
In another embodiment, the invention involves a titanium aluminum alloy consisting essentially of (in atomic %) about 40 to about 52% Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn, and about 0.5 about 5.0% Cr. A preferred alloy consists essentially of (in atomic %) about 41 to about 50% Ti, about 46% to 49% Al, about 1% to about 3% Mn, about 1% to about 3% Cr, up to abcut 3% V and up to about 3% Nb. Second phase dispersoids may be included in the alloy in an amount of about 0.5 to about 20.0 volume % to increase strength.
Unexpectedly, the titanium aluminide alloy exhibits an increase in ductility as well as strength upon the 2 ~
P-311 Howmet 10 inclusion of tne second phase dispersoids therein.
Brief Descri~tion of the Drawings Figures la and lb are bar graphs illustrating the change in strength and ductility of Cr-bearing, predominantly gamma titanium aluminide alloys of the invention upon the inclusion of titanium borides. Similar data is presented for a Ti-48Al-2V-2Mn alloy ~reference alloy) to illustrate the increase in strength but the decrease in ductility observed upon inclusion of the same boride levels therein.
Figures 2a, 2b, and 2c illustrate the microstructure of the Ti-48Al-2V-2Mn reference alloy after hot isostatic pressing and heat treatment at 1650F (900C) for 16 hours.
Figure~ 3a, 3b and 3c illustrate the microstructure of the Ti-48Al-2Mn-2Cr alloy of thP
invention after the same hot isostatic pressing and heat treatment as used in Figs. 2a-2c.
Figures 4a, 4b and 4c illustrate the microstructure of the Ti-48Al-2V-2Mn-2Cr alloy of the P-311 Howmet 11 invention after the same hot isostatic pressing and heat treatment as used in Figs. 2a-2c.
Figures 5a,5b and 6a,6b illustrate the change in strength and ductility of the aforementioned alloys of Fig. 1 after different heat treatments.
Figures 7a, 7b and 7c, 7d illustrate the effect of heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, on microstructure of the Ti-48al-2Mn-2Cr alloy of the invention devoid of TiB2 dispersoids.
Figures 8a, 8b and 8c, 8d illustrate the effect of heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, on microstructure of the Ti-48al-2Mn-2Cr alloy of the invention including 7 volume % TiB2 dispersoids.
Figure 9 illustrates the change in yield strength of the aforementioned alloys of Fig. 1 with the volume % of TiB2 dispersoids.
Figure 10 illustrates the measured grain size as a function of TiB2 volume % for the aforementioned alloys.
P-311 Howmet 12 2 Q ~ 7 Detailed Descri~tion of the Invention The present invention contemplates a titanium aluminide article including second phase dispersoids (e.g., TiB2) in a Cr-bearing, predominantly gamma TiAl matrix in effective concentrations that result in an increase in both strength and ductility. In one embodiment of the invention, the alloy matrix consists essentially of, in atomic %, about 40 to about 52 % Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn and about 0.5 to about 5.0 % Cr to this end. Preferably, the alloy matrix consists essentially of, in atomic %, about 41 to about 50% Ti, about 46 to about 49% Al, about 1 to about 3% Mn, about 1 to about 3~ Cr, up to about 3% V, and up to about 3% Nb. The alloy matrix includes second phase dispersoids, such as preferably TiB2, in an amount not exceeding about 20.0 volume %.
Preferably, the second phase dispersoids are present in an amount of about 0.5 to about 12.0 volume %, more preferably from about 0.5 to about 7.0 volume %.
P-311 Howmet 13 The matrix is considered predominantly gamma in that a majority of the matrix microstructure in the as-cast or the cast/hot isostatically pressed/heat treated condition described hereafter comprises gamma phase. Alpha 2 and beta phases can also be present in minor proportions of the matrix microstructure; e.g., from about 2 to about 15 volume % of alpha 2 phase and up to about 5 volume ~ beta phase can be present.
The following Table I lists nominal and measured Cr-bearing titanium-aluminum ingot compositions produced in accordance with exemplary embodiments of the present invention. Also listed are the nominal and measured ingot composition of a Ti-48Al-2V-2Mn alloy used as a reference alloy for comparison purposes.
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ooo ooo ooo oo Z O O O O O O O O O O O
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0 ,~ ~ to o~ o o~ o h _~ N ~ ~ ~ O ~1 O` ~ ~D O~ O ~ O ~ O ~ O
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r o~ o r ~ o r r o ~o ~~ ~ o ~ ,~ ~o ~ ~ :
r r ~ o o~ ~ ~D R
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P-311 Howmet 15 The dispersoids of TiB2 were provided in the ingots using a master sponge material comprising 70 weight % TiB2 in an Al matrix and available from Martin Marietta Corp., Bethesda, Md. and its licensees. The master sponge material was introduced into a titanium aluminum melt of the appropriate composition prior to casting into an investment mold in accordance with U.S. Patents 4,751,048 and 4,916,030, the teachings of which are incorporated herein by reference.
Segments of each ingot were sliced, remelted by a conventional vacuum arc remelting, to a superheat of +50F above the alloy melting temperature, and investment cast into preheated ceramic molds (600F) to form cast test bars having a diameter of 0.625 inch and a length of 6.0 inches. Each mold included a Zr203 facecoat and a plurality of Al203/Zr203 backup coats.
Following casting and removal from the investment molds, all test bars were hot isostatically pressed (HIP'ed) at 25 ksi and 2300F for 4 hours in an inert atmosphere (Ar).
Baseline mechanical tensile data were obtained using the investment cast test bars which had P-311 Howmet 16 2~9~7 been heat treated at 1650F (900C) for 16 hours following the aforementioned hot isostatic pressing operation. The TiB2 dispersoids present in the cast/HIP'ed/heat treated test bars typically had particle sizes (i.e., diameters) in the range of 0.3 to 5 microns.
The results of the tensile tests are shown in Fig. la plotted as a function of matrix alloy composition for 0, 7, and 12 volume % TiB2. From Fig.
la, it is apparent that the yield strength of all the alloys increases with the addition of 7 and 12 volume % TiB2.
However, from Fig. lb, the room temperature ductility of the Ti-48Al-2V-2Mn alloy was observed to decrease sùbstantially with the addition of these levels of TiB2 to the matrix alloy. Surprisingly, the ductility of the Cr-bearing alloys (i.e., Ti-48Al-2Mn-2Cr, Ti-48Al-2V-2Mn-2Cr and Ti-47Al-2Mn lNb-lCr) was observed to increase with the addition of these levels of TiB2, especially upon the addition of 7 volume % TiB2. Thus, for the TiAl alloys including chromium as an additional alloyant and TiB2 dispersoids, both the strength and the ductility were P-311 Howmet 17 2~3~ ~
found to increase unexpectedly.
Representative optlcal microstructures of these alloys after casting, hot isostatic pressing, and heat treatment are shown in Figs. 2a, 2b, 2c; 3a, 3b, and 3c; and 4a, 4b, and 4c. The photomicrographs illustrate that the microstructures of the alloys are predominantly lamellar (i.e., alternating lathes of gamma phase and alpha 2 phase) with some equiaxed grains residing at colony boundaries. Generally, there was little or no evidence of microstructural coarsening or other morphological transformations upon hot isostatic pressing and/or heat treatment.
The effect of longer time or higher temperature heat treatments on alloy strength and ductility are illustrated in Figs. 5a,5b and 6a,6b for heat treatments at 900C (1650F) for 50 hours (Figs.
5a,5b) and 1100C (2012F) for 16 hours (Figs. 6a,6b).
Yield strength is shown to increase with increasing percent TiB2. Moreover, increases in ductility were again noted for the Cr-bearing test bars having 7 volume % TiB2 in the matrix. In general, the 900oC
(1650F) heat treatments resulted in maximum ductility in all of the alloys shown. In the alloys of the P-311 Howmet 18 invention containing 7 and 12 volume % TiB2, maximum ductility occurred following heat treatment at 1650F
for 50 hours. In general, strength was relatively insensitive to heat treatment.
Figs. 7a,7b and 7c,7d illustrate the microstructures of alloy matrices following heat treatment at 1650F for 50 hours and 2012F for 16 hours, respectively, for the Ti-48Al-2Mn-2Cr devoid of TiB2. Figs. 8a,8b and 8c,8d illustrate the alloy matrix microstructure for the same alloy with 7 volume % TiB2 after the same heat treatments. In the boride-free alloy, transformation of the matrix to a primarily equiaxed microstructure was observed after these heat treatments. On the other hand, the matrix microstructure including 7 volume % TiB2 exhibited very little change after these heat treatments, retaining a primarily lamellar microstructure.
Fig. 9 illustrates tensile yield strength as a function of dispersoid (TiB2) loading for the aforementioned alloys heat treated at 1650F for 16 hours. All alloys exhibit approximately linear increases in strength with increasing dispersoid loading (volume %). The Ti-48Al-2V-2Mn alloy P-311 Howmet 19 exhibited the strongest dependence~
Grain size analyses were performed on the alloys that had been heat treated at 1650F for 16 hours to determine the effect of dispersoid loading on grain size. Fig. 10 depicts large reductions in grain size due to the inoculative effect of the TiB2 dispersoids. A reduced sensitivity of grain size on dispersoid loading is apparent at higher volume fractions of dispersoids. The large variations in alloy grain size when no dispersoids are present appears to be a consequence primarily of the size and scale of the smaller, equiaxed grains that reside between large columnar, lamellar colonies.
The surprising increase in both strength and ductility of the Cr-Bearing, predominantly gamma titanium aluminides of Fig. 1 is also observed at elevated temperatures as illustrated in Table II
wherein investment cast, HIP'd, and heat treated (900C for 50 hours) specimens were tensile tested at ~16C.
, , P-311 Howmet 20 ~9 TABLE II
Tensile Testing a 816C
a (ksi) o (k~ elong yield ult Ti-48Al-2Mn-2Cr 49.5 56.2 18.1 Ti-48Al-2Mn-2Cr + 7 v% Ti82 45-0 52.4 22.8 Ti-48Al-2Mn-2Cr + 12 v~ TiB2 47.5 55.3 20.3 Ti-47Al-2Mn-lNb-lCr 51.9 68.0 4.9 Ti-47Al-2Mn-lNb-lCr + 7~ v% TiB2 51.2 76.5 ~12.3 The creep resistance of the Ti-47Al-2Mn-lNb-lCr alloy without and with 7 volume % TiB2 dispe_soids was evaluated at 1500F and 20.0 ksi load. The specimens were investment cast, HIP'ed, and heat treated at 900C for 50 hours. As indicated in Table III, the boride-free and boride-bearing specimens exhibited generally comparable rupture lives. The creep resistance of the Ti-47Al-2Mn-lNb-lCr alloy thus was not adversely affected by the inclusion of 7 volume % TiB2 dispersoids.
TABLE III
Creep Data at 1500F/20.0 ksi RuDture Llfe ~hrs) Ti-47Al-2Mn-lNb-lCr 96.3/111.7 Tl-47Al-2Mn-lNb-lCr + 7 v~ TL8z 102.~3/110.7 P-311 Howmet 21 ~ 7 In practicing the present invention, the concentration of Cr should not exceed about 5.0 atomic % of the TiAl alloy composition in order to provide the aforementioned predominantly gamma titanium aluminide matrix microstructure. For example, a TiAl ingot nominally comprising Ti-48Al-2V-2Mn-6Cr (measured composition, in atomic %, 44.1 Ti-45.8Al-20Mn-6.2Cr-1.9V) was prepared and investment cast, HIP'ed, and heat treated as described hereinabove for the alloys of Fig. 1. The ingot included about 7.0 volume % TiB2. Examination of the microstructure of the ingot before and after a 1650~F/16 hour heat treatment revealed volume fractions of beta phase well in excess of 5 volume %, primarily at grain (colony) boundaries and along lamellar interfaces. The heat treatment resulted in spherodization and a relatively homogeneous distribution of the beta phase in the microstructure. The heat treated alloy exhibited a tensile yield strength of about 90 ksi but a substantially reduced ductility at room temperature of only 0.15 %.
Thus, in practicing the invention the upper limit of the Cr concentration should not exceed about 5.0 atomic % of the alloy composition. On the other P-311 Howmet 22 2 ~
hand, the lower limit of the Cr concentration should be sufficient to result in an increase in both strength and ductility when appropriate amounts of dispersoids are included in the matrix. To this end, in accordance with the present invention, the Cr concentration is preferably from about 0.5 to about 5.0 atomic % of the alloy matrix, more preferably from about 1.0 to about 3.0 atomic % of the alloy matrix.
While the invention has been described in terms of specific embodiments thereof, it is not intended to be limited thereto but rather only to the extent set forth in the following claims.
Claims (26)
1. An article comprising a Cr-bearing, predominantly gamma titanium aluminide matrix having second phase dispersoids present in the matrix in an amount effective to increase both the strength and the ductility thereof as compared to the strength and ductility of the matrix devoid of the dispersoids.
2. The article of claim 1 wherein Cr is present in the matrix in an amount of about 0.5 to about 5.0 atomic % of the matrix.
3. The article of claim 2 wherein Cr is present in an amount of about 1.0 to about 3.0 atomic %.
4. The article of claim 1 wherein the second phase dispersoids are present in the matrix in an amount of about 0.5 to about 20.0 volume %.
5. The article of claim 1 wherein the second phase dispersoids are present in an amount of about 0.5 to about 12.0 volume %.
P-311 Howmet 24
P-311 Howmet 24
6. The article of claim 5 wherein the second phase dispersoids are present in an amount of about 0.5 to about 7.0 volume %.
7. The article of claim 1 wherein the second phase dispersoids comprise a boride of titanium.
8. An article comprising a Cr-bearing, predominantly gamma titanium aluminide matrix consisting essentially of, in atomic %, about 40 to about 52% Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn, and about 0.5 to about 5.0% Cr, and second phase dispersoids present in the matrix in an amount effective to increase both the strength and the ductility thereof as compared to the strength and ductility of the matrix devoid of the dispersoids.
9. The article of claim 8 wherein the second phase dispersoids are present in the matrix in an amount of about 0.5 to about 12.0 volume %.
10. The article of claim 8 wherein the second phase dispersoids comprise a boride of titanium.
P-311 Howmet 25
P-311 Howmet 25
11. An article comprising a Cr-bearing, predominantly gamma titanium aluminide matrix consisting essentially of, in atomic %, about 41 to about 50% Ti, about 46 to about 49% Al, about 1 to 3%
Mn, about 1 to about 3% Cr up to about 3% V, and up to about 3% Nb, and second phase dispersoids present in the matrix in an amount effective to increase both the strength and the ductility thereof as compared to the strength and ductility of the matrix devoid of the dispersoids.
Mn, about 1 to about 3% Cr up to about 3% V, and up to about 3% Nb, and second phase dispersoids present in the matrix in an amount effective to increase both the strength and the ductility thereof as compared to the strength and ductility of the matrix devoid of the dispersoids.
12. The article of claim 11 wherein the second phase dispersoids are present in the matrix in an amount of about 0.5 to about 12.0 volume %.
13. The article of claim 11 wherein the second phase dispersoids comprise a boride of titanium.
14. A titanium aluminum alloy consisting essentially of, in atomic %, about 40 to about 52% Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn and about 0.5 to about 5.0% Cr, said alloy being amenable to an increase in both strength and ductility by virtue of the inclusion of second phase dispersoids P-311 Howmet 26 therein.
15. A titanium aluminum alloy consisting essentially of, in atomic %, about 41 to about 50% Ti, about 46 to about 49% Al, about 1 to about 3 % Mn, about 1 to about 3% Cr, up to about 3% V, and up to 3%
Nb, said alloy being amenable to an increase in both strength and ductility by virtue of the inclusion of second phase dispersoids therein.
Nb, said alloy being amenable to an increase in both strength and ductility by virtue of the inclusion of second phase dispersoids therein.
16. A method of making a titanium aluminide article, comprising including second phase dispersoids in a Cr-bearing, predominantly gamma titanium aluminide matrix in an amount effective to increase both the strength and ductility of the matrix as compared to the matrix devoid of the dispersoids.
17. The method of claim 16 wherein Cr is included in the matrix in an amount of about 0.5 to about 5 atomic % thereof.
18. The method of claim 16 wherein the second phase dispersoids comprise a boride of titanium present in an amount of about 0.5 to about 20.0 volume %.
P-311 Howmet 27
P-311 Howmet 27
19. The method of claim 16 wherein the second phase dispersoids are present in an amount of about 0.5 to about 12.0 volume %.
20. The method of claim 19 wherein the second phase dispersoids are present in an amount of about 0.5 to about 7.0 volume %.
21. The method of claim 16 wherein the dispersoids are included in the matrix by introducing preformed dispersoids into a Cr-bearing titanium-aluminum alloy melt and then solidifying the melt.
22. The method of claim 21 wherein the melt is investment cast to solidify it.
23. A method of making a titanium aluminide article, comprising including second phase dispersoids in a Cr-bearing, predominantly gamma titanium aluminide matrix consisting essentially of, in atomic %, about 40 to 52% Ti, about 44 to about 52% Al, about 0.5 to about 5.0% Mn, and about 0.5 to about 5.0% Cr, said dispersoids being included in an amount effective to increase both strength and ductility of the matrix as compared to the matrix devoid of the dispersoids.
P-311 Howmet 28
P-311 Howmet 28
24. The method of claim 23 wherein the second phase dispersoids comprise a boride of titanium present in an amount of about 0.5 to about 12.0 volume %.
25. A method of making a titanium aluminide article, comprising including second phase dispersoids in a Cr-bearing, predominantly gamma titanium aluminide matrix consisting essentially of, in atomic %, about 41 to about 50% Ti, about 46 to about 49% Al, about 1 to about 3% Mn, about 1 to about 3% Cr, up to about 3% V, and up to about 3% Nb, said dispersoids being included in an amount effective to increase both strength and ductility of the matrix as compared to the matrix devoid of the dispersoids.
26. The method of claim 25 wherein the second phase dispersoids comprise a boride of titanium present in an amount of about 0.5 to about 12.0 volume %.
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|---|---|---|---|
| US07/716,951 US5354351A (en) | 1991-06-18 | 1991-06-18 | Cr-bearing gamma titanium aluminides and method of making same |
| US716,951 | 1991-06-18 |
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| JP2679109B2 (en) * | 1988-05-27 | 1997-11-19 | 住友金属工業株式会社 | Intermetallic compound TiA-based light-weight heat-resistant alloy |
| US4879092A (en) * | 1988-06-03 | 1989-11-07 | General Electric Company | Titanium aluminum alloys modified by chromium and niobium and method of preparation |
| US4906430A (en) * | 1988-07-29 | 1990-03-06 | Dynamet Technology Inc. | Titanium diboride/titanium alloy metal matrix microcomposite material and process for powder metal cladding |
| US4897127A (en) * | 1988-10-03 | 1990-01-30 | General Electric Company | Rapidly solidified and heat-treated manganese and niobium-modified titanium aluminum alloys |
| US4902474A (en) * | 1989-01-03 | 1990-02-20 | General Electric Company | Gallium-modified titanium aluminum alloys and method of preparation |
| US4916028A (en) * | 1989-07-28 | 1990-04-10 | General Electric Company | Gamma titanium aluminum alloys modified by carbon, chromium and niobium |
| JP2749165B2 (en) * | 1989-12-25 | 1998-05-13 | 新日本製鐵株式会社 | TiA-based composite material and method for producing the same |
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| US5098653A (en) * | 1990-07-02 | 1992-03-24 | General Electric Company | Tantalum and chromium containing titanium aluminide rendered castable by boron inoculation |
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| US5284620A (en) * | 1990-12-11 | 1994-02-08 | Howmet Corporation | Investment casting a titanium aluminide article having net or near-net shape |
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| US5204058A (en) * | 1990-12-21 | 1993-04-20 | General Electric Company | Thermomechanically processed structural elements of titanium aluminides containing chromium, niobium, and boron |
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| US5228931A (en) * | 1991-12-20 | 1993-07-20 | General Electric Company | Cast and hipped gamma titanium aluminum alloys modified by chromium, boron, and tantalum |
-
1991
- 1991-06-18 US US07/716,951 patent/US5354351A/en not_active Expired - Lifetime
-
1992
- 1992-05-26 CA CA002069557A patent/CA2069557A1/en not_active Abandoned
- 1992-06-16 DE DE69217732T patent/DE69217732D1/en not_active Expired - Lifetime
- 1992-06-16 DE DE69229971T patent/DE69229971T2/en not_active Expired - Lifetime
- 1992-06-16 EP EP96111924A patent/EP0753593B1/en not_active Expired - Lifetime
- 1992-06-16 EP EP92420209A patent/EP0519849B1/en not_active Expired - Lifetime
- 1992-06-17 JP JP4181563A patent/JP2651975B2/en not_active Expired - Fee Related
-
1993
- 1993-12-02 US US08/161,324 patent/US5433799A/en not_active Expired - Lifetime
- 1993-12-02 US US08/161,323 patent/US5458701A/en not_active Expired - Lifetime
Also Published As
| Publication number | Publication date |
|---|---|
| JP2651975B2 (en) | 1997-09-10 |
| DE69217732D1 (en) | 1997-04-10 |
| EP0519849B1 (en) | 1997-03-05 |
| DE69229971D1 (en) | 1999-10-14 |
| US5458701A (en) | 1995-10-17 |
| DE69229971T2 (en) | 2000-03-30 |
| US5354351A (en) | 1994-10-11 |
| JPH06293928A (en) | 1994-10-21 |
| EP0753593A1 (en) | 1997-01-15 |
| EP0519849A3 (en) | 1993-06-09 |
| US5433799A (en) | 1995-07-18 |
| EP0753593B1 (en) | 1999-09-08 |
| EP0519849A2 (en) | 1992-12-23 |
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Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| FZDE | Discontinued |