EP1308529A1 - Intermetallische Verbindung aus Titan-Aluminium-Basislegierung und Verfahren zur Herstellung eines Produktes aus dieser Legierung - Google Patents

Intermetallische Verbindung aus Titan-Aluminium-Basislegierung und Verfahren zur Herstellung eines Produktes aus dieser Legierung Download PDF

Info

Publication number
EP1308529A1
EP1308529A1 EP02024552A EP02024552A EP1308529A1 EP 1308529 A1 EP1308529 A1 EP 1308529A1 EP 02024552 A EP02024552 A EP 02024552A EP 02024552 A EP02024552 A EP 02024552A EP 1308529 A1 EP1308529 A1 EP 1308529A1
Authority
EP
European Patent Office
Prior art keywords
based alloy
intermetallic compound
compound based
titanium aluminum
product
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Granted
Application number
EP02024552A
Other languages
English (en)
French (fr)
Other versions
EP1308529B1 (de
Inventor
Toshimitsu Nagasaki Res. & Dev. Center Tetsui
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Mitsubishi Heavy Industries Ltd
Original Assignee
Mitsubishi Heavy Industries Ltd
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Mitsubishi Heavy Industries Ltd filed Critical Mitsubishi Heavy Industries Ltd
Publication of EP1308529A1 publication Critical patent/EP1308529A1/de
Application granted granted Critical
Publication of EP1308529B1 publication Critical patent/EP1308529B1/de
Anticipated expiration legal-status Critical
Expired - Lifetime legal-status Critical Current

Links

Images

Classifications

    • 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
    • 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
    • C22F1/16—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of other metals or alloys based thereon
    • C22F1/18—High-melting or refractory metals or alloys based thereon
    • C22F1/183—High-melting or refractory metals or alloys based thereon of titanium or alloys based thereon

Definitions

  • the present invention is related to a titanium aluminum intermetallic compound based alloy (hereinafter, referred to as a TiAl based alloy) and a method of fabricating a product from the alloy.
  • a TiAl based alloy titanium aluminum intermetallic compound based alloy
  • TiAl based alloys exhibit superior strength at high temperatures. Therefore, TiAl based alloys are expected to use for rotating components, such as turbine wheels of turbochargers and turbine blades of gas turbines.
  • Jp-A-Heisei 9-143599 discloses a TiAl based alloy.
  • the alloy consists of, in atomic percent, 45-48% aluminum, 0.5-2.0% nickel, 1.0-3.0% niobium, 0.2-1.0% tungsten, and 1.0-2.0% manganese, and a remainder including titanium plus impurities.
  • the alloy has superior internal friction properties because of the appropriate nickel concentration.
  • Fabrication of products from TiAl based alloys is often achieved by casting.
  • An issue of casting is that many defects, such as voids, are included in the cast products. Defects are desirably removed to increase mechanical properties of the cast products.
  • HIP hot isostatic pressing
  • HIP enlarges the lamellar spacing of lamellar structure of the TiAl based alloys.
  • the enlargement of the lamellar spacing is undesirable, because the narrow lamellar spacing of the lamellar structure is an origin of superior mechanical properties of the TiAl based alloy.
  • the increase in the lamella spacing degrades the creep strength of the TiAl based alloy.
  • the low cycle fatigue strength of TiAl based alloys is inferior to conventional metallic material because of its poor ductility. Especially in case of TiAl alloys made by casting, the low cycle fatigue strength is very small due to its large grain size. The poor low cycle fatigue strength is an issue, because gas turbines components are often exposed to severe thermal stress that causes low cycle fatigue.
  • An object of the present invention is to provide a TiAl based alloy made by casting that is superior in creep strength.
  • Another object of the present invention is to provide a TiAl based alloy made by casting that is superior in low cycle fatigue strength.
  • Still another object of the present invention is to provide an TiAl based alloy made by casting that is superior in both of creep strength and low cycle fatigue strength.
  • a TiAl based alloy has properties as follows.
  • the TiAl based alloy has lamellar structure, and a volume ratio of non-lamellar structure of the alloy is equal to or less than 3 volume percent. Diameters of lamellar grains included in the alloy are equal to or less than 200 ⁇ m. Lamellar spacing of lamella structure included in the alloy is equal to or less than 2 ⁇ m.
  • the alloy preferably includes boron, a concentration of which is advantageously in the range of 0.2 to 1.2 atomic percent.
  • an TiAl based alloy includes, in atomic percent, 43-48% aluminum, 2.0-5.0% niobium, 0.2-1.2% tungsten, 0.1-1.0% nickel, 0.2-1.2% boron, and a remainder including titanium plus impurities.
  • the TiAl based alloy preferably includes manganese, a concentration of which is advantageously in the range of 0.2 to 1.2 atomic percent.
  • the TiAl based alloy preferably includes chromium, a concentration of which is advantageously in the range of 0.2 to 1.2 atomic percent.
  • the TiAl based alloy preferably includes silicon, a concentration of which is advantageously in the range of 0.1 to 1.0 atomic percent.
  • the TiAl based alloy preferably includes carbon, a concentration of which is advantageously in the range of 0.1 to 0.5 atomic percent.
  • a method of fabricating a product from a TiAl based alloy is composed of
  • the product is preferably heated at 1320° to 1370° C. during the heat-treatment.
  • the product is preferably cooled to 1000 °C. at a cooling rate of 30° to 100° C./min. during the rapid cooling.
  • the fabricating method preferably includes:
  • the product is preferably heated at 900° to 1050° C. during the second heat-treatment.
  • the product is cooled at a cooling rate equal to or less than 10° C./min. during the slow cooling.
  • the fabricating method is advantageous when the TiAl based alloy includes boron, a concentration of which is preferably in the in the range of 0.2 to 1.2 atomic percent.
  • a TiAl based alloy has properties as follows.
  • the TiAl based alloy has a lamellar structure, and a volume ratio of non-lamellar structure thereof is equal to or less than 3 volume percent. Diameters of lamellar grains are equal to or less than 200 ⁇ m, and lamellar spacing is equal to or less than 2 ⁇ m.
  • the small volume ratio of the non-lamellar structure, and the small lamellar spacing improves the creep strength of the TiAl based alloy, while the small diameters of the lamellar grains improves the low cycle fatigue strength.
  • the TiAl based alloy is composed of, in atomic percent, 43-48% aluminum (Al), 2.0-5.0% niobium (Nb), 0.2-1.2% tungsten (W), 0.1-1.0 nickel (Ni), 0.2-1.2% manganese (Mn), and 0.2-1.2% boron (B), and a remainder including titanium plus impurities.
  • the aluminum concentration of the TiAl based alloy is advantageously in the range of 43 to 48 atomic percent.
  • An aluminum concentration less than 43 atomic percent degrades toughness of the alloy.
  • aluminum concentration more than 48 atomic percent decreases the ratio of ⁇ 2 phase in the lamellar structure, and thus degrades the strength at high temperature.
  • the doping of niobium improves anti-oxidation resistance.
  • the niobium concentration of the TiAl based alloy is advantageously in the range of 2.0 to 5.0 atomic percent. Niobium concentration less than 2.0 atomic percent is not effective to improve the anti-oxidation resistance. On the other hand, niobium concentration more than 5.0 atomic percent is not preferable because of increases in the specific gravity and the cost of the alloy.
  • the doping of tungsten improves the strength of the TiAl based alloy at high temperature.
  • the tungsten concentration is advantageously in the range from 0.2 to 1.2 atomic percent. Tungsten concentration less than 0.2 atomic percent is not effective to improvement of the mechanical strength at high temperature. On the other hand, tungsten concentration more than 1.2 atomic percent is not preferable because of increases in the specific gravity and the cost of the alloy.
  • the doping of nickel increases the internal friction of the TiAl based alloy.
  • the nickel concentration of the alloy is advantageously in the range of 0.1 to 1.0 atomic percent. Nickel concentration less than 0.1 atomic percent is not effective to increase in the internal friction. On the other hand, nickel concentration more than 1.0 atomic percent decreases the ductility at room temperature because of the production of harmful phases, such as Laves phases, in the alloy.
  • the doping of manganese improves low cycle fatigue strength.
  • the manganese concentration of the TiAl based alloy is advantageously in the range of 0.2 to 1.2 atomic percent. Manganese concentration less than 0.2 atomic percent is not effective to the improvement of low cycle fatigue strength. On the other hand, manganese concentration more than 1.2 atomic percent degrades anti-oxidation resistance.
  • the doping of boron decreases the size of lamellar grains of the TiAl based alloy.
  • the decrease in the grain size improves the low cycle fatigue strength of the alloy.
  • the boron concentration of the alloy is advantageously in the range of 0.2 to 1.2 atomic percent. Boron concentration less than 0.2 atomic percent is not effective to the decrease in the grain size. On the other hand, Boron concentration more than 1.2 atomic percent degrades the brittleness of the alloy because of the precipitation of boride.
  • the doping of chromium is preferable to further improve the low cycle fatigue strength.
  • the chromium concentration of the TiAl based alloy is advantageously in the range of 0.2 to 1.2 atomic percent. Chromium concentration less than 0.2 atomic percent is not effective to the improvement of the low cycle fatigue strength. On the other hand, chromium concentration more than 1.2 atomic percent degrades the creep strength because of the formation of ⁇ phase in the alloy.
  • the doping of silicon is preferable to further improve the creep strength of the TiAl based alloy.
  • the silicon concentration of the alloy is advantageously in the range of 0.1 to 1.0 atomic percent. Silicon concentration less than 0.1 atomic percent is not effective to the improvement of the creep strength. Silicon concentration more than 1.0 atomic percent degrades low cycle fatigue strength.
  • the doping of carbon is also preferable to further improve the creep strength of the alloy.
  • the carbon concentration of the alloy is advantageously in the range of 0.1 to 0.5 atomic percent. Carbon concentration less than 0.1 atomic percent is not effective to the improvement of the creep strength. Carbon concentration more than 0.5 atomic percent degrades low cycle fatigue strength.
  • the fabrication process starts with melting and casting the TiAl based alloy into a cast product.
  • the boron doping into the TiAl based alloy effectively improves the low cycle fatigue strength of the cast product through miniaturizing the size of the lamellar grains of the product.
  • the cast product is then exposed to HIP.
  • the conditions of the HIP are as follows.
  • the cast product is heated and pressed in an environment of substantially pure argon at a temperature of 1200° to 1300° C. for 1 to 3 hours.
  • the applied pressure is equal to or more than 100 MPa.
  • the argon environment effectively prevents the oxidation of the cast product.
  • the temperature range of 1200° to 1300° C. is advantageous because of the following reasons. HIP at a temperature lower than 1200° C. is not effective to the elimination of defects because of the poor deformablilty of the alloy. On the other hand, HIP at a temperature higher than 1300° C. is not preferable, because the argon environment inevitably includes a small amount of oxygen gas as an impurity and the small amount of oxygen gas may oxidize the cast product, because its partial pressure is increased by highpressure such as 100 MPa.
  • HIP eliminates the defects of the cast product.
  • the creep strength of the cast product is degraded by HIP.
  • HIP increases lamellar spacing of the lamellar structure, and produces non-lamellar structure at grain boundaries. The non-lamellar structure degrades the conformability of the lamellar grains. The increase in the lamellar spacing and the degraded conformability degrade the creep strength of the alloy.
  • the cast product is subjected to the following process. Firstly, the cast product is heat-treated in vacuum at a temperature from 1320° to 1370° C. for 20 to 60 minutes, and then rapidly cooled. This process may be referred to as a VHRCP (vacuum heating and rapid cooling process), hereinafter.
  • the rapid cooling of the cast product is achieved by blowing argon gas having room temperature into the furnace used for the heat-treatment.
  • the cast product is preferably cooled at a cooling rate of 30° to 100° C./min. till the metal product is cooled to 1000° C.
  • the VHRCP restructures the lamellar structure of the alloy, and effectively improves the creep strength.
  • the VHRCP decreases the lamellar spacing and eliminates the non-lamella structure at the grain boundaries. The decrease in the lamellar spacing and the elimination of the non-lamella structure effectively improve the creep strength of the alloy.
  • the temperature of the heat-treatment in the range of 1320° to 1370° C. is advantageous because of the following reasons.
  • a heat treatment at a temperature lower than 1320° C. does not cause the restructuring of the lamella structure.
  • the VHRCP which is effective to the improvement of the creep strength, produces remaining stress in the cast product.
  • the remaining stress causes dynamic recrystallization, and thereby produces non-lamellar structure at the grain boundaries when an external mechanical stress is exerted on the cast product.
  • the non-lamellar structure degrades the creep strength of the cast product.
  • the cast product is annealed at a temperature from 900° to 1050° C. for 3 to 50 hours in air.
  • the annealing effectively removes the remaining stress.
  • the annealing temperature from 900° to 1050° C. is advantageous.
  • An annealing at a temperature lower than 900° C. is not effective to the removal of the remaining stress.
  • An annealing at a temperature higher than 1050° C. destroys the lamellar structure restructured by the VHRCP.
  • the treatment in inert gas atmosphere is not required, because oxidation of the TiAl alloy in the air is small in this temperature range.
  • the structure and compositions of the TiAl based alloy and the fabricating process effectively improves both the creep strength and the low cycle fatigue strength of the cast product.
  • Seven turbine wheels which have a shape shown in Fig. 1, were fabricated by casting. The turbine wheels were then subjected to the aforementioned HIP, VHRCP, and/or annealing. HIP is executed for 3 hours under the conditions of a temperature of 1250° C. and a pressure of 150 MPa. The fabricated seven turbine wheels are respectively referred to as Examples 1 to 7.
  • Fig. 2 shows alloys and fabrication process of Examples 1 to 7.
  • Examples 1 to 7 are made from one of Alloys 1 to 5 that have different compositions as shown in Fig. 3.
  • Alloy 1 has a comparative composition, which includes none of boron, chromium, silicon and carbon.
  • Alloys 2 to 5 are embodiments according to the present invention.
  • Alloy 2 is standard composition.
  • Alloy 3 includes chromium, which improves the low cycle fatigue strength.
  • Alloys 4 and 5 respectively include silicon and carbon, which improve creep strength.
  • Example 1 was fabricated from Alloy 1, and subjected to HIP, VHRCP, and annealing for removing remaining stress.
  • Examples 2 to 4 were fabricated from Alloy 2, while subjected to different processes.
  • Example 2 was exposed to HIP, not exposed to VHRCP and annealing.
  • Example 3 was exposed to HIP and VHRCP, not exposed to annealing.
  • Example 4 was exposed to HIP, VHRCP, and annealing.
  • Examples 5, 6, and 7 were respectively fabricated from Alloy 3, 4, and 5. All of Examples 5, 6, and 7 were exposed to HIP, VHRCP and annealing.
  • Example 2 was not subjected to VHRCP and annealing, and Example 3 was not subjected to the annealing. Examples 2 and 3 were utilized for showing the effects of the VHRCP and annealing.
  • Test pieces for a low cycle fatigue test and a creep rupture test were machined from the center portions of Examples 1 to 7, respectively.
  • the test pieces for the low cycle fatigue test were cylindrical rods measuring 3 mm in diameter and 8 mm in length.
  • the test pieces were subjected to the low cycle fatigue test under the conditions of 0.7% strain and a temperature of 750° C. to determine cycles to low cycle fatigue failure.
  • the creep test pieces were cylindrical bars measuring 6 mm in diameter and 30 mm in length.
  • the creep test pieces are subjected to stress of 100 MPa at 900° C. to determine creep rupture lives.
  • Example 1 which does not include boron and was subjected to HIP, VHRCP, and annealing, exhibits superior creep strength. The superior creep strength arises from the VHRCP and annealing. Example 1, however, is poorest in low cycle fatigue strength.
  • Fig. 4 shows an optical micrograph of a cross section of Example 1. The grains of Example 1, which is not doped with boron, have a diameter larger than 200 ⁇ m. The large grain size causes the poor low cycle fatigue strength of Example 1.
  • Example 2 which includes boron, exhibits superior low cycle fatigue strength.
  • Fig. 5 shows an optical micrograph of a cross section of Example 2.
  • the grain size of Example 2 (that is, size of lamellar grains) is smaller than 200 ⁇ m, that is, much smaller than that of Example 1.
  • the small grain size of Example 2 effectively improves the low cycle fatigue strength.
  • Example 2 however, exhibits poor creep strength.
  • Fig. 6 shows a back scattered electron micrograph of a cross section of Example 2 in the vicinity of a grain boundary.
  • the lamellar spacing of Example 2 is larger than 2 ⁇ m.
  • Example 2 includes non-lamellar structure.
  • the non-lamellar structure has a volume percent more than 3%.
  • the large lamellar spacing and the non-lamellar structure of Example 2 causes the poor creep strength.
  • Example 3 which includes boron and were subjected to HIP and VHRCP, exhibits superior low cycle fatigue strength. As discussed above, this results from the small grain size caused by the boron doping.
  • Example 3 is superior to Example 2 in creep strength.
  • Fig. 7 shows a back scattered electron micrograph of a cross section of Example 3 before the creep rupture test.
  • the lamellar spacing is small and less than 2 ⁇ m.
  • non-lamellar structure is not observed in the vicinity of grain boundary.
  • the backscattered electron micrograph proves that Example 3 has a desirable structure for improving the creep strength.
  • Example 3 however, suffers dynamic recrystallization during the creep test.
  • Fig. 8 shows an optical micrograph of a cross section of Example 3 after the failure of the creep test. A large number of small grains formed by the recrystallization are observed on the boundaries of the lamellar grains. Creep voids are found where the recrystallized grains exist.
  • the recrystallization is considered as being caused by remaining stress.
  • the rapid cooling during the VHRCP produces large remaining stress at the grain boundaries, and the remaining stress causes dynamic recrystallization during the creep test. Removal of the remaining stress is considered as being effective to improvement of creep strength.
  • Example 4 which includes boron and is exposed to HIP, VHRCP, and annealing, exhibits superior low cycle fatigue strength and creep strength.
  • the superior low cycle fatigue is achieved by the doping of boron, while the improvement of the creep strength is achieved by the VHRCP and annealing.
  • the annealing removes remaining stress, and thus suppresses the dynamic recrystallization.
  • Example 5 which was fabricated from Alloy 3, exhibits superior low cycle fatigue strength, while it is inferior to Example 4 in creep strength. This results from the doping of chromium that improves ductility while deteriorating strength. In this invented alloy, the doping of chromium is effective for the application requesting superior low cycle fatigue strength.
  • Example 6 which is fabricated from Alloy 4, exhibits superior creep strength, while it is inferior to Example 4 in low cycle fatigue strength. This results from the doping of silicon that improves mechanical strength while deteriorating ductility. In this invented alloy, the doping of silicon is effective for the application requesting superior creep strength.
  • Example 7 which is fabricated from Alloy 5, also exhibits superior creep strength, while it is inferior to Example 4 in low cycle fatigue strength. This results from the doping of carbon that improves strength while deteriorating ductility. In this invented alloy, the doping of carbon is effective for the application requesting superior creep strength.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Turbine Rotor Nozzle Sealing (AREA)
  • Powder Metallurgy (AREA)
EP02024552A 2001-11-05 2002-11-04 Intermetallische Verbindung aus Titan-Aluminium-Basislegierung und Verfahren zur Herstellung eines Produktes aus dieser Legierung Expired - Lifetime EP1308529B1 (de)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
JP2001339980 2001-11-05
JP2001339980A JP4107830B2 (ja) 2001-11-05 2001-11-05 ＴｉＡｌ金属間化合物基合金、及び鋳造部品の製造方法

Publications (2)

Publication Number Publication Date
EP1308529A1 true EP1308529A1 (de) 2003-05-07
EP1308529B1 EP1308529B1 (de) 2005-01-26

Family

ID=19154248

Family Applications (1)

Application Number Title Priority Date Filing Date
EP02024552A Expired - Lifetime EP1308529B1 (de) 2001-11-05 2002-11-04 Intermetallische Verbindung aus Titan-Aluminium-Basislegierung und Verfahren zur Herstellung eines Produktes aus dieser Legierung

Country Status (4)

Country Link
US (1) US20030111141A1 (de)
EP (1) EP1308529B1 (de)
JP (1) JP4107830B2 (de)
DE (1) DE60202731T2 (de)

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP1507017A1 (de) * 2003-08-14 2005-02-16 Rolls-Royce Plc Verfahren zur Wärmebehandlung von TiAl-Legierungen
EP2808488A1 (de) * 2013-05-29 2014-12-03 MTU Aero Engines GmbH TiAl-Schaufel mit Oberflächenmodifizierung

Families Citing this family (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20070178383A1 (en) * 2006-01-31 2007-08-02 Viavattine Joseph J Current collector
WO2014070510A1 (en) * 2012-11-02 2014-05-08 Borgwarner Inc. Process for producing a turbine wheel
RU2502824C1 (ru) * 2012-11-13 2013-12-27 Федеральное государственное автономное образовательное учреждение высшего профессионального образования "Национальный исследовательский технологический университет "МИСиС" Способ термообработки отливок из сплавов на основе гамма алюминида титана
KR101551961B1 (ko) 2013-08-19 2015-09-09 현대자동차주식회사 크리프 저항성이 우수한 티타늄(Ti)-알루미늄(Al) 합금

Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4842820A (en) * 1987-12-28 1989-06-27 General Electric Company Boron-modified titanium aluminum alloys and method of preparation
US5350466A (en) * 1993-07-19 1994-09-27 Howmet Corporation Creep resistant titanium aluminide alloy
US5653828A (en) * 1995-10-26 1997-08-05 National Research Council Of Canada Method to procuce fine-grained lamellar microstructures in gamma titanium aluminides
JP2000199025A (ja) * 1999-01-05 2000-07-18 Mitsubishi Heavy Ind Ltd ＴｉＡｌ系金属間化合物基合金およびその製造方法、タ―ビン部材およびその製造方法
US6294132B1 (en) * 1996-10-28 2001-09-25 Mitsubishi Heavy Industries Ltd. TiAl intermetallic compound-based alloy
WO2001088214A1 (de) * 2000-05-17 2001-11-22 Gfe Metalle Und Materialien Gmbh Bauteil auf basis von gamma-tial-legierungen mit bereichen mit gradiertem gefüge

Family Cites Families (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
FR2579344B1 (fr) * 1985-03-21 1991-06-21 Dupont Daniel Stimulateur analogique de circuits de commande actionnes a partir d'un poste ou habitacle

Patent Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4842820A (en) * 1987-12-28 1989-06-27 General Electric Company Boron-modified titanium aluminum alloys and method of preparation
US4842820B1 (de) * 1987-12-28 1992-05-12 Gen Electric
US5350466A (en) * 1993-07-19 1994-09-27 Howmet Corporation Creep resistant titanium aluminide alloy
US5653828A (en) * 1995-10-26 1997-08-05 National Research Council Of Canada Method to procuce fine-grained lamellar microstructures in gamma titanium aluminides
US6294132B1 (en) * 1996-10-28 2001-09-25 Mitsubishi Heavy Industries Ltd. TiAl intermetallic compound-based alloy
JP2000199025A (ja) * 1999-01-05 2000-07-18 Mitsubishi Heavy Ind Ltd ＴｉＡｌ系金属間化合物基合金およびその製造方法、タ―ビン部材およびその製造方法
WO2001088214A1 (de) * 2000-05-17 2001-11-22 Gfe Metalle Und Materialien Gmbh Bauteil auf basis von gamma-tial-legierungen mit bereichen mit gradiertem gefüge

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
PATENT ABSTRACTS OF JAPAN vol. 2000, no. 10 17 November 2000 (2000-11-17) *

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP1507017A1 (de) * 2003-08-14 2005-02-16 Rolls-Royce Plc Verfahren zur Wärmebehandlung von TiAl-Legierungen
EP2808488A1 (de) * 2013-05-29 2014-12-03 MTU Aero Engines GmbH TiAl-Schaufel mit Oberflächenmodifizierung

Also Published As

Publication number Publication date
US20030111141A1 (en) 2003-06-19
DE60202731T2 (de) 2005-12-29
DE60202731D1 (de) 2005-03-03
JP2003147475A (ja) 2003-05-21
JP4107830B2 (ja) 2008-06-25
EP1308529B1 (de) 2005-01-26

Similar Documents

Publication Publication Date Title
JP6889418B2 (ja) Ｎｉ基超耐熱合金の製造方法およびＮｉ基超耐熱合金
RU2370561C2 (ru) Сплав на основе алюминидов титана
JP6839401B1 (ja) Ｎｉ基超耐熱合金及びＮｉ基超耐熱合金の製造方法
JP5792500B2 (ja) Ｎｉ基超合金材およびタービンロータ
JP6687118B2 (ja) ＴｉＡｌ合金及びその製造方法
JP6842316B2 (ja) Ｎｉ基合金、ガスタービン材およびクリープ特性に優れたＮｉ基合金の製造方法
JPH0823053B2 (ja) 加工性に優れた高強度チタン合金およびその合金材の製造方法ならびにその超塑性加工法
JP6315319B2 (ja) Ｆｅ−Ｎｉ基超耐熱合金の製造方法
JP2022037155A (ja) 高温チタン合金
WO2020189215A1 (ja) 熱間鍛造用のチタンアルミナイド合金材及びチタンアルミナイド合金材の鍛造方法並びに鍛造体
KR20230085948A (ko) 내크리프성 티타늄 합금
WO2000020652A1 (en) Creep resistant gamma titanium aluminide alloy
JP4756974B2 (ja) Ｎｉ３（Ｓｉ，Ｔｉ）系箔及びその製造方法
JP2009041065A (ja) 高温疲労強度および耐クリープ性に優れた耐熱部材用チタン合金
JP4264411B2 (ja) 高強度α＋β型チタン合金
EP1308529B1 (de) Intermetallische Verbindung aus Titan-Aluminium-Basislegierung und Verfahren zur Herstellung eines Produktes aus dieser Legierung
JP2020026568A (ja) チタン合金、その製造方法およびそれを用いたエンジン部品
JP7387139B2 (ja) チタン合金、その製造方法およびそれを用いたエンジン部品
JP4228166B2 (ja) 疲労強度の優れた継目無銅合金管
EL et al. Microstructure and ductility of TiAl alloys modified by Cr additions
JP5228708B2 (ja) 耐クリープ性および高温疲労強度に優れた耐熱部材用チタン合金
JP2008144202A (ja) 耐熱ばね及びその製造方法
JP2002097537A (ja) Ｃｏ−Ｎｉ基耐熱合金およびその製造方法
JP2003247039A (ja) 耐酸化性、高温強度及び熱間加工性に優れたＮｉ基合金
JP2728305B2 (ja) 金属間化合物ＴｉＡ▲ｌ▼基合金の熱間加工法

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

17P Request for examination filed

Effective date: 20021104

AK Designated contracting states

Designated state(s): AT BE BG CH CY CZ DE DK EE ES FI FR GB GR IE IT LI LU MC NL PT SE SK TR

AX Request for extension of the european patent

Extension state: AL LT LV MK RO SI

17Q First examination report despatched

Effective date: 20030714

AKX Designation fees paid

Designated state(s): DE NL SE

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): DE NL SE

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

REF Corresponds to:

Ref document number: 60202731

Country of ref document: DE

Date of ref document: 20050303

Kind code of ref document: P

REG Reference to a national code

Ref country code: SE

Ref legal event code: TRGR

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20051027

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: DE

Payment date: 20081103

Year of fee payment: 7

Ref country code: NL

Payment date: 20081116

Year of fee payment: 7

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: SE

Payment date: 20081107

Year of fee payment: 7

REG Reference to a national code

Ref country code: NL

Ref legal event code: V1

Effective date: 20100601

EUG Se: european patent has lapsed
PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: NL

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20100601

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: DE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20100601

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20091105