WO2016055013A1 - TiAl金属间化合物单晶材料及其制备方法 - Google Patents

TiAl金属间化合物单晶材料及其制备方法 Download PDF

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WO2016055013A1
WO2016055013A1 PCT/CN2015/091508 CN2015091508W WO2016055013A1 WO 2016055013 A1 WO2016055013 A1 WO 2016055013A1 CN 2015091508 W CN2015091508 W CN 2015091508W WO 2016055013 A1 WO2016055013 A1 WO 2016055013A1
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alloy
bar
single crystal
tial
intermetallic compound
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PCT/CN2015/091508
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English (en)
French (fr)
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陈�光
彭英博
郑功
祁志祥
王敏智
李沛
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南京理工大学
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Priority claimed from CN201410528019.3A external-priority patent/CN104328501B/zh
Priority claimed from CN201410529844.5A external-priority patent/CN104278173B/zh
Priority claimed from CN201510244611.5A external-priority patent/CN104878452A/zh
Application filed by 南京理工大学 filed Critical 南京理工大学
Priority to US15/517,165 priority Critical patent/US10570531B2/en
Priority to RU2017115945A priority patent/RU2701438C2/ru
Priority to JP2017538285A priority patent/JP6944874B2/ja
Priority to EP15849516.8A priority patent/EP3205753B1/en
Publication of WO2016055013A1 publication Critical patent/WO2016055013A1/zh

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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C1/00—Making non-ferrous alloys
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    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C22C33/06—Making ferrous alloys by melting using master alloys
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    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B11/00—Single-crystal growth by normal freezing or freezing under temperature gradient, e.g. Bridgman-Stockbarger method
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    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B11/00—Single-crystal growth by normal freezing or freezing under temperature gradient, e.g. Bridgman-Stockbarger method
    • C30B11/001—Continuous growth
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B11/00—Single-crystal growth by normal freezing or freezing under temperature gradient, e.g. Bridgman-Stockbarger method
    • C30B11/002—Crucibles or containers for supporting the melt
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B11/00—Single-crystal growth by normal freezing or freezing under temperature gradient, e.g. Bridgman-Stockbarger method
    • C30B11/003—Heating or cooling of the melt or the crystallised material
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B13/00—Single-crystal growth by zone-melting; Refining by zone-melting
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B21/00—Unidirectional solidification of eutectic materials
    • C30B21/04—Unidirectional solidification of eutectic materials by zone-melting
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B29/00—Single crystals or homogeneous polycrystalline material with defined structure characterised by the material or by their shape
    • C30B29/10—Inorganic compounds or compositions
    • C30B29/52—Alloys
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B33/00—After-treatment of single crystals or homogeneous polycrystalline material with defined structure
    • C30B33/02—Heat treatment
    • C—CHEMISTRY; METALLURGY
    • C30—CRYSTAL GROWTH
    • C30B—SINGLE-CRYSTAL GROWTH; UNIDIRECTIONAL SOLIDIFICATION OF EUTECTIC MATERIAL OR UNIDIRECTIONAL DEMIXING OF EUTECTOID MATERIAL; REFINING BY ZONE-MELTING OF MATERIAL; PRODUCTION OF A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; SINGLE CRYSTALS OR HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; AFTER-TREATMENT OF SINGLE CRYSTALS OR A HOMOGENEOUS POLYCRYSTALLINE MATERIAL WITH DEFINED STRUCTURE; APPARATUS THEREFOR
    • C30B35/00—Apparatus not otherwise provided for, specially adapted for the growth, production or after-treatment of single crystals or of a homogeneous polycrystalline material with defined structure
    • C30B35/007—Apparatus for preparing, pre-treating the source material to be used for crystal growth

Definitions

  • the invention belongs to the technical field of lightweight high-strength structural materials, and particularly relates to a TiAl intermetallic compound single crystal material and a preparation method thereof.
  • TiAl intermetallic compound is a new type of lightweight high-temperature structural material with a specific gravity less than 50% of nickel-based superalloy. It has high specific strength, high specific ratio, corrosion resistance, wear resistance, high temperature resistance, high elastic modulus and excellent Its oxidation resistance, creep resistance and good high temperature strength, its use temperature can reach 750 ⁇ 900 ° C, similar to Ni-based superalloy; but its density is only half of the high temperature alloy, it is the ideal Ni-based high temperature Alternative materials for alloys can be used in a wide range of high temperature components such as blades, turbine discs and exhaust valves for automotive or aerospace engines. For example, TiAl alloy is aerospace high temperature material in grams as a weight loss unit, especially the best candidate material for engines.
  • the United States GE company successfully used the Ti-48Al-2Cr-2Nb (4822) alloy to develop the two-stage low-pressure turbine blades of the Boeing aircraft, which reduced the aircraft weight by about 200Kg.
  • the high-temperature mechanical properties, creep resistance and oxidation performance of high-Nb-TiAl alloys are significantly higher than those of ordinary TiAl alloys, and the use temperature is about 60-100 °C, which is the most suitable TiAl alloy for engineering application.
  • the poor brittleness of the TiAl alloy at room temperature has become a major cause of hindering its industrial application.
  • the current service temperature of the 4822 alloy is only 650 ° C, and its high temperature performance needs to be further improved. Therefore, a large amount of research has focused on regulating the microstructure of TiAl alloys to improve room temperature brittleness and increase service temperature. Due to the obvious anisotropy of strength and plasticity in the TiAl alloy PST crystal, the TiAl alloy was prepared by directional solidification to produce a full-sheet PST crystal, and its sheet orientation was parallel to the growth direction of the crystal in the directional solidification. It can improve the mechanical properties of TiAl alloy.
  • the mechanical properties of the full-lamellar TiAl alloy are closely related to its sheet orientation.
  • PST polycrystalline twinned crystal
  • its strength and plasticity showed significant anisotropy. Due to this anisotropy of the full-sheet structure, when the sheet orientation is appropriate, it is more suitable for aeronautical engine blades such as those that require high temperature resistance and are only subjected to one-dimensional load.
  • the TiAl alloy can be produced by directional solidification to produce a full-sheet structure of the engine blade, and the orientation of the layer is parallel to the axial direction of the blade (the direction of growth of the crystal in the directional solidification), it is undoubtedly extremely advantageous. Yamaguchi et al.
  • the control methods of TiAl alloy sheet orientation at home and abroad mainly include seed crystal method and non-seed crystal method which changes the solidification path.
  • Yamaguchi, Johnson et al. obtained the single crystal PST whose sheet orientation was completely parallel to the growth direction by the ⁇ phase solidification seed crystal method and the Ti-Al-Si alloy as the seed crystal.
  • the seed crystal composition usually differs from the parent alloy composition, resulting in uneven composition and properties of the directionally solidified alloy, and the preparation process of the seed crystal is complicated. Therefore, the seed crystal method has obvious deficiencies.
  • This method requires two times of directional solidification of the same process, one more solidification process than the ordinary non-seed method, which aggravates the contamination of the alloy by the tantalum material, which is unfavorable for the industrialization of the directionally solidified TiAl alloy.
  • the preferred growth direction is ⁇ 001>, and its phase relationship is: ⁇ 110 ⁇ ⁇ // ⁇ 0001 ⁇ ⁇ // ⁇ 111 ⁇ ⁇ [25] , and 12 variables of ⁇ 110 ⁇ ⁇ Four of them are parallel to the growth direction, and eight are inclined to 45° with the growth direction [16 , 26] , and only one-third of the habits in the lamellar structure formed after the solid phase transformation are oriented parallel to the growth direction.
  • the orientation of the final layer structure of the TiAl alloy depends not only on the growth direction of the primary ⁇ phase but also on the subsequent solid phase transformation process. Therefore, the ⁇ solid phase transition process is also the key to controlling the orientation of the sheet. So far, research on TiAl sheet orientation control has focused on the solidification process, while neglecting the solid phase transformation process after solidification.
  • Intermetallic compound single crystal material The material has an ideal sheet orientation and uniformity and no pollution.
  • the high strength (729 MPa) is maintained while the room temperature tensile plasticity reaches 6.9%, the yield strength at 900 ° C is 637 MPa, and the ductile-brittle transition temperature reaches 900 ° C or higher.
  • Another object of the present invention is to provide a method for producing the above TiAl intermetallic compound single crystal material.
  • the object of the invention can be achieved by the following measures:
  • the TiAl intermetallic compound single crystal material of the invention can be prepared by a non-seed method optically floating zone directional solidification method, and the method comprises the following steps:
  • the mother alloy bar is cut into two parts, the upper and lower bars, respectively, as the raw material rod and the seed crystal rod of the optical floating zone directional solidification furnace, and the distance between the upper raw material bar and the lower seed bar is controlled to be 1 ⁇ 5mm; the distance between the upper and lower bars is 1 ⁇ 5mm; firstly, the raw material rod and the seed crystal rod are coaxial and perpendicular to the horizontal plane, and the inert gas is used for protection during the direction solidification, and the upper and lower bars are adjusted to the opposite Direction rotation, relative rotation speed is 10 ⁇ 40rpm, start heating, make the opposite ends of the upper and lower bars melt first, adjust the position of the upper and lower bars, make the opposite end gradually approach and then join, adjust the power of the equipment and keep warm 5 - After 10 minutes, when the surface of the floating zone is smooth and evenly melted, the growth rate is adjusted to 2.5-30 mm/h, and directional solidification is started; after the solidification is finished, the power is slowly reduced, and the solidified sample is slowly separated from the remaining feedstick
  • the prepared TiAl alloy single crystal bar is subjected to vacuum heat treatment using a heat treatment method of 1250 ° C to 1350 ° C ⁇ 12 h to 24 h + 900 ° C ⁇ 30 min / furnace cooling or air cooling.
  • the electromagnetic induction suspension melting is performed by using a water-cooled copper crucible, and the number of times of the mother alloy is not less than 3 times, further preferably not less than 4 times.
  • the size of the mother alloy bar is ⁇ (4-8) mm ⁇ 120 mm; the suction casting method adopts differential pressure suction casting, and the pressure difference is maintained at 3 MPa.
  • the protective gas pressure is at two-thirds of the standard atmospheric pressure.
  • the prepared mother alloy round bar may have a size of ⁇ (4 to 8) mm.
  • the purity of the Al, Ti, C or Si raw material is 99.999% or more, and the purity of the Nb pure metal raw material is 99.9% or more.
  • the length of the lower seed rod is 20-30 mm, and the length of the upper material rod is less than 190 mm.
  • the inert gas is argon or nitrogen, and the flow rate of the inert gas introduced during the directional solidification is 3 to 5 L/min.
  • step (2) the position of the upper and lower bars is adjusted so as to be gradually approached and joined to the opposite end, and then the power of the apparatus is adjusted for heat preservation and melting.
  • the total power of the device is 4.0 KW, its power is adjusted to 55-70% of the total power.
  • the single crystal rod is subjected to a de-segregation vacuum heat treatment of "1250 ° C ⁇ 24 h + 900 ° C ⁇ 30 min + air cooling".
  • the present invention further provides a method for preparing the above-mentioned single TiAl intermetallic compound single crystal material, the method comprising the following steps:
  • the first step selecting Ti, Al, Nb pure metal raw materials with a purity of 99.999% or more, according to the alloy composition expression, smelting the mother alloy in a cold crucible suspension melting furnace with a vacuum of less than 10 -3 Pa, via 3 ⁇ 4 times of melting to homogenize the alloy composition and suck-cast into a directional solidified bar;
  • the second step the TiAl alloy test bar is placed in a high-purity yttrium oxide coated corundum crucible for directional solidification, vacuuming to 5 ⁇ 10 -3 Pa, and then charging the system with high-purity argon shielding gas;
  • the third step adjusting the power of the induction power supply to heat the sample, the holding temperature is 1450 ⁇ 1650K, the holding time is 15 ⁇ 30min, the directional solidification begins, and the controlled solidification drawing rate is 5 ⁇ 20 ⁇ m / s; continuous growth to the sample At a length of 50 mm, the rapid quenching was started to rapidly quench the directional solidified sample, and the solid-liquid interface was retained.
  • the size of the directionally solidified bar is ⁇ (4 to 6 mm) x 100 mm.
  • the high-purity cerium oxide coating has a corundum size of ⁇ (7 to 9 mm) ⁇ 100 mm; and the high purity argon shielding gas is charged in an amount of 0.04 to 0.06 MPa.
  • the principle of the method is to control the orientation of the TiAl alloy sheet by the Bridgman directional solidification method.
  • the primary phase is ensured to be the whole ⁇ phase
  • the single crystal is obtained by the grain competition in the solidification process, and There is a critical temperature during solidification corresponding to a specific draw rate at which the draw rate is
  • the final phase orientation and the ⁇ phase with a growth direction of 45° are eliminated by phase boundary migration, and only the ⁇ phase of the final layer orientation and the growth direction is retained among the 12 ⁇ variables obtained in the ⁇ phase transition, thereby completing Control of sheet orientation.
  • the invention has the following advantages:
  • the preparation method of the TiAl alloy material proposed by the invention can greatly improve the room temperature mechanical properties of the alloy, and particularly improve the room temperature brittleness.
  • the invention can effectively improve the high temperature mechanical properties of the alloy by adjusting the content of Nb element and adding a small amount of C and Si strengthening elements.
  • Non-seed crystal optical floating zone directional solidification technology is used to prevent alloy contamination while avoiding the disadvantages of complex processing and uneven composition of the seed crystal method, which can avoid the problem of alloy contamination caused by directional solidification of traditional Bridgman.
  • a TiAl-Nb single crystal is obtained efficiently.
  • the vacuum heat treatment completely eliminates a large amount of brittle B2 phase and Nb-rich brittle segregation phase remaining in the microstructure after directional solidification of TiAl alloy, thereby obtaining an alloy material with uniform microstructure and excellent room temperature performance, and avoiding coarsening of the sheet.
  • the method can also adopt the common Bridgman directional solidification method to control the continuous orientation liquid-solid phase transformation-oriented solid phase transformation by adjusting the solidification parameter, ensuring the full ⁇ phase growth and controlling the final sheet orientation by controlling the solid phase transformation. And a TiAl alloy single crystal structure in which the sheet orientation is completely parallel to the growth direction is obtained.
  • the invention effectively avoids the disadvantage of uneven performance of the seed crystal component, and at the same time obtains an ideal lamellar oriented single crystal structure in a single directional solidification process, which simplifies the process.
  • the single-wafer layer orientation can be completely controlled under a certain range of solidification parameters.
  • the invention provides a theoretical basis for the industrial application of directionally solidified TiAl alloy.
  • the preparation method has the advantages of simple preparation process, low cost, remarkable effect of improving room temperature brittleness, universal applicability and promotion value.
  • Figure 1 is a phase diagram of a prior art portion of a Ti-Al binary alloy.
  • FIG. 2 is a microstructural view of the longitudinal section (a) and the sheet orientation (b) of the directionally solidified sample of the present invention.
  • Fig. 3 is a longitudinal sectional view of the competition section of the directionally solidified sample of the present invention.
  • FIG. 4 is a microstructural view of the longitudinal section (a) and the sheet orientation (b) of the directionally solidified sample of the present invention.
  • Fig. 5 is a longitudinal sectional view of the competition section of the directionally solidified sample of the present invention.
  • Figure 6 is a microstructural view of the longitudinal section (a) and the sheet orientation (b) of the directionally solidified sample of the present invention.
  • Figure 7 is a quenching solid-liquid interface of a directional solidification sample of the present invention.
  • the growth direction of the microstructure in Figures 2-7 is from right to left.
  • Figure 8 is a flow chart for preparing a high strength and high plasticity TiAl alloy material.
  • Figure 9 is a TiAl alloy directionally solidified single crystal (a) and sheet orientation (b) microstructure.
  • Fig. 10 is a microscopic structure of segregation before and after the different heat treatment processes of the TiAl alloy single crystal (a is before heat treatment, and b is after heat treatment).
  • Figure 11 is an XRD diffraction pattern of a TiAl alloy single crystal before and after different heat treatment processes.
  • Figure 12 is a graph showing the tensile mechanical properties of a TiAl alloy at room temperature.
  • Fig. 13 is a schematic view showing the solid-liquid interface morphology (a) and the neck-removing principle (b) for solidification of a TiAl-Nb single crystal.
  • Figure 14 is an optical morphology obtained after directional solidification of a TiAl-Nb alloy.
  • Fig. 15 is a scanning electron micrograph of segregation before (a) and (b) of the TiAl-Nb single crystal.
  • Figure 16 is a picture of the interlaminar spacing of (a) and (b) before the heat treatment of the TiAl-Nb single crystal.
  • Figure 17 is a displacement intensity curve of TiAl-Nb stretched at 900 °C.
  • the TiAl intermetallic compound single crystal with fully controlled lamellar orientation is prepared in accordance with the accompanying drawings and a Bridgman directional solidification method.
  • the specific embodiment is as follows:
  • a Ti-Al-Nb ternary alloy in which the primary phase is a total ⁇ phase is selected.
  • the precipitated phase is all ⁇ phase. Specifically, the content of Nb is increased, and the relative proportion of Al is lowered to form a wider ⁇ phase region.
  • the high-purity metal component is used, and under the protection of high-purity Ar gas, the master alloy is melted by a cold-blow electromagnetic suspension melting equipment. The master alloy is smelted multiple times to obtain a uniform master alloy ingot and sucked into a master alloy bar.
  • the holding temperature is 1450 ⁇ 1650K
  • the holding time is 15 ⁇ 30min
  • the directional solidification begins, and the directional solidification growth rate is controlled to 5 ⁇ 20 ⁇ m / s;
  • the alloy composition used in the experiment is Ti 47 Al 45 Nb 8 (atomic percent at%), and the purity of the metal component is 99.999%.
  • the cold cathode electromagnetic is used under the vacuum degree of 5 ⁇ 10-3Pa.
  • the suspension smelting equipment melts the master alloy.
  • a uniform mother alloy ingot was obtained by 4 times of smelting, and was drawn into a ⁇ 4 ⁇ 100 mm mother alloy bar.
  • the TiAl alloy test bar was placed in a corundum crucible coated with high-purity yttrium oxide on the inner wall to conduct a directional solidification test, vacuumed to 5 ⁇ 10-3 Pa, and then charged with 0.05 MPa high purity argon shielding gas into the system.
  • the holding temperature is 1550K
  • the holding time is 25min
  • the directional solidification is started
  • the directional solidification growth rate is controlled to be 5 ⁇ m/s; when the drawing length is 50mm to the length of the sample, the rapid quenching is started.
  • the sample is subjected to a rapid quenching treatment to retain a solid-liquid interface.
  • the maximum longitudinal section of the cylindrical specimen was characterized by microstructure. The precipitating phase, grain size and lamellar orientation of the solidification at the drawing rate were observed and analyzed, as shown in Fig. 2(a) and Fig. 2(b). As shown, it was found that a TiAl alloy single crystal having a sheet orientation parallel to the growth direction was obtained.
  • the enrichment of the solute can be sufficiently diffused, the growth can be stably performed, and the crystal grains have a sufficient time to grow, so that the obtained crystal grains are coarsened until single crystal growth is obtained.
  • Figure 3 shows the microstructure of the directional solidification competition segment at 5 ⁇ m/s. Since in the ⁇ solid-state phase transition, the difference in mismatch between the two interfaces forming the 0° and 45° sheets results in different phase boundary mobility, so there is a critical pull rate of 5 ⁇ m/s. Below this drawing rate, the 0° and 45° sheet-oriented ⁇ -grain nucleation forms a 0° grain growth driving force, and the 45° grain is finally eliminated, and a single crystal whose sheet orientation is parallel to the growth direction is obtained. .
  • the holding temperature was 1550 K
  • the holding time was 25 min
  • the directional solidification was started
  • the directional solidification growth rate was controlled to be 15 ⁇ m/s; as shown in Fig. 4(a) and Fig. 4(b) It is shown that at this drawing rate, the ⁇ solid phase transition remains under the 45° sheet-oriented ⁇ phase, so the final structure is a single crystal with a sheet orientation of 45°.
  • Figure 5 shows the microstructure of the directional solidification competition segment at 15 ⁇ m/s.
  • the driving force of the 45° grain solid phase deformation nucleus is greater than 0° grain, so that the 0° grain cannot grow, and a TiAl alloy single crystal with a sheet orientation and a growth direction of 45° is obtained.
  • the holding temperature was 1550 K, the holding time was 25 min, the directional solidification was started, and the directional solidification growth rate was controlled to be 20 ⁇ m/s; as shown in Fig. 6(a) and Fig. 6(b) It is shown that a single crystal having a sheet orientation and a growth direction of 45° was obtained.
  • Figure 7 shows the solid-liquid interface retained by the quenching treatment.
  • the dendrite growth morphology is 4-fold symmetrical, with obvious secondary dendrites and a 90° vertical relationship with the primary dendrites. It can be inferred that the directional solidification process In the middle, the ⁇ phase of the cubic system is the primary phase.
  • the alloy composition used was Ti 55 Al 43 Nb 2 , the holding temperature was 1650 K, the holding time was 30 min, the directional solidification growth rate was 5 ⁇ m/s, and TiAl parallel to the growth direction of the sheet orientation was obtained. Alloy single crystal.
  • the alloy composition used was Ti 48 Al 43 Nb 9 , the holding temperature was 1450 K, the holding time was 30 min, the directional solidification growth rate was 10 ⁇ m/s, and the sheet orientation and growth direction were 45°. Single crystal of TiAl alloy.
  • the alloy composition used was Ti 51 Al 45 Nb 6 , the holding temperature was 1650 K, the holding time was 15 min, the directional solidification growth rate was 5 ⁇ m/s, and TiAl which was parallel to the growth direction of the sheet orientation was obtained. Alloy single crystal.
  • the alloy composition used was Ti 42 Al 49 Nb 9 , the holding temperature was 1550 K, the holding time was 25 min, the directional solidification growth rate was 5 ⁇ m/s, and TiAl which was parallel to the growth direction of the sheet orientation was obtained. Alloy single crystal.
  • the atomic percentage of the alloy composition of the high-strength high-plastic TiAl alloy material is: (44 to 51) Ti-(43 to 47) Al-(6 to 9) Nb.
  • the precipitated phase is all ⁇ phase.
  • the prepared TiAl alloy single crystal bar is subjected to vacuum heat treatment; after heating for a certain period of time in the ⁇ single-phase region, the film is annealed after heat preservation; the brittle B2 phase and residual stress are completely eliminated, and a high-strength and high-plastic TiAl alloy material is obtained.
  • the alloy composition selected for the preparation of the master alloy ingot of the present invention is Ti 47 Al 45 Nb 8 (atomic percent), and the purity of each metal component is 99.999%, and the Nb is 99.95%.
  • the mother alloy ingot is melted in a water-cooled copper crucible electromagnetic induction suspension melting furnace: after the surface of the metal raw material is mechanically polished to remove the surface oxide scale, according to the designed distribution ratio material preparation; according to each ingot
  • the weight of about 70g is put into the water-cooled copper crucible in the melting furnace, and vacuumed to 5 ⁇ 10 -3 Pa; the furnace is filled with a certain amount of high-purity argon (99.999%), argon gas.
  • the pressure range is from 0.8 to 1 MPa. Multi-pass smelting 3 to 4 times to obtain a uniformly mixed master alloy ingot.
  • the master alloy ingot was then suction cast into a ⁇ 6 x 120 mm bar.
  • the mother alloy bar is cut into two parts, the upper and lower bars, respectively, as the raw material rod and the seed crystal rod of the optical floating zone directional solidification furnace; the lower end is a seed crystal rod with a length of 30 mm, and the upper end is a feeding rod with a length of less than 100 mm;
  • the raw material rod is firstly arranged coaxially with the seed crystal rod and perpendicular to the horizontal plane.
  • the distance between the upper and lower bars is 5 mm and the interval is at the center of the four filament focusing; high purity argon is introduced at 5 L/min.
  • the gas is used as the atmosphere protection, and the axial relative rotational speed of the upper and lower bars is adjusted to 30 rpm, and heating is started to melt the opposite ends of the upper and lower bars first, and the positions of the upper and lower bars are adjusted so that the opposite ends are gradually approached and then joined. Adjusting the power to 68% of the total power, keeping the surface of the floating zone smooth and evenly uniform (ie, when the floating zone has no obvious jitter), adjusting the growth rate to 5 mm/h, and starting directional solidification; stopping the directional solidification when growing to 80 mm, Slowly reduce the power while slowly separating the solidified sample from the remaining feed bar sample;
  • the single crystal portion of the directionally solidified bar was placed in a corundum tube, vacuumed to 10-3 Pa, and the tube was sealed, and placed in a heat treatment furnace, and subjected to a heat treatment process of 1300 ° C ⁇ 24 h + 900 ° C ⁇ 30 min / furnace cooling.
  • Fig. 9a is a macroscopic photograph of the test rod after directional solidification of the optical floating zone. It can be seen that the sample rapidly becomes single crystal growth after undergoing short competition in directional solidification, and Fig. 9b shows that the orientation of the single wafer layer is parallel to the growth direction.
  • Fig. 10(a) and Fig. 10(b) are microstructure diagrams before and after heat treatment. In combination with the XRD pattern of Fig. 11, it can be seen that a large amount of B2 phase is distributed inside the microstructure before heat treatment, and B2 is completely eliminated after heat treatment for 24 hours.
  • Figure 12 is a room temperature drawing of the prepared high strength and high plasticity TiAl alloy. The stress-strain curve has a yield strength of 729 MPa and a plastic strain of 6.9%, and has excellent room temperature mechanical properties.
  • the alloy composition was Ti 44 Al 47 Nb 9 (atomic percent), and the optical floating zone directional solidification process was a relative rotational speed of 20 rpm, a heating power of 55%, a growth rate of 2.5 mm/h, and a vacuum heat treatment.
  • the process is 1250 ° C ⁇ 12 h + 900 ° C ⁇ 30 min / furnace cooling, B2 phase is completely eliminated, the tensile yield strength of the TiAl alloy material obtained at room temperature is 550 MPa, and the plastic strain is 6.0%.
  • the alloy composition was Ti 51 Al 40 Nb 9 (atomic percent), and the optical floating zone directional solidification process was a relative rotational speed of 25 rpm, a heating power of 70%, a growth rate of 10 mm/h, and a vacuum heat treatment process.
  • the B2 phase is completely eliminated, and the tensile yield strength of the TiAl alloy material is 628 MPa and the plastic strain is 6.5%.
  • the alloy composition was Ti 48 Al 43 Nb 9 (atomic percent), and the optical floating zone directional solidification process was a relative rotational speed of 20 rpm, a heating power of 68%, a growth rate of 15 mm/h, and a vacuum heat treatment process.
  • the B2 phase is completely eliminated, and the tensile yield strength of the TiAl alloy material is 660 MPa and the plastic strain is 6.2%.
  • the alloy composition was Ti 48 Al 43 Nb 9 (atomic percent), and the optical floating zone directional solidification process was a relative rotational speed of 20 rpm, a heating power of 70%, a growth rate of 15 mm/h, and a vacuum heat treatment process.
  • the B2 phase is completely eliminated, and the tensile yield strength at room temperature of the TiAl alloy material is 593 MPa, and the plastic strain is 6.8%.
  • the alloy composition was Ti 48 Al 46 Nb 6 (atomic percent), and the optical floating zone directional solidification process was a relative rotational speed of 30 rpm, a heating power of 60%, a growth rate of 20 mm/h, and a vacuum heat treatment process.
  • the optical floating zone directional solidification process was a relative rotational speed of 30 rpm, a heating power of 60%, a growth rate of 20 mm/h, and a vacuum heat treatment process.
  • B2 phase is not completely eliminated, as shown in the XRD pattern of Figure 10b, a small amount of B2 phase remains in the 12h heat treatment, and the tensile yield strength of the TiAl alloy material is 656MPa, and the plastic strain is 3.0%.
  • the alloy composition was Ti 44 Al 45 Nb 8 (atomic percent), and the optical floating zone directional solidification process was a relative rotation speed of 25 rpm, a heating power of 55%, and a growth rate of 30 mm/h, to obtain a sheet layer.
  • the TiAl alloy single crystal with orientation and growth direction is 45°
  • the vacuum heat treatment process is 1250°C ⁇ 12h+900°C ⁇ 30min/furnace cooling
  • the B2 phase is completely eliminated.
  • the tensile yield strength of the TiAl alloy material is 430MPa and the plastic strain is 7.8%
  • the atomic percentage of the alloy composition is: Ti-45Al-8Nb-0.3C-0.2Si, the balance is Ti, the initial raw material is 99.999% high purity Al, Ti, C and Si and 99.95% high purity Nb,
  • the TiAl-Nb master alloy ingot was obtained by repeated smelting in a cold heading electromagnetic induction suspension melting furnace at a vacuum of 5 ⁇ 10-3 MPa;
  • the shielding gas with a flow rate of 4L/min was introduced, and the seeding rod and the feeding rod were adjusted to rotate in the opposite direction at 30r/min, and the heating power was increased to 68% within 10 minutes to melt the alloy.
  • the temperature was 15 mm/
  • the growth rate of h is directional solidification. Due to the heating characteristics of the optical floating zone, the solid-liquid interface is a convex interface as shown in Fig. 13(a).
  • the schematic diagram is as shown in Fig. 13(b), the intermediate crystal grains will grow along the growth direction, and the crystal grains on both sides will It grows obliquely to both sides.
  • Figure 14 shows that the power is slowly reduced after the directional solidification is completed, and the solidified sample is slowly separated from the remaining feed rod sample;
  • the prepared TiAl-Nb single crystal is subjected to vacuum de-segregation heat treatment.
  • the segregation morphology before heat treatment is as shown in Fig. 15(a).
  • the segregation phase can be eliminated.
  • the final single crystal was obtained by air cooling, and the heat treatment was completely eliminated as shown in Fig. 15 (b).
  • Fig. 16 shows the variation of the lamellar spacing before and after the heat treatment. Since the air cooling rate is fast, the lamellar layer cannot be roughened.
  • the heat-treated single crystal is processed into a tensile specimen with a gauge length of ⁇ 3 mm ⁇ 20 mm, and the tensile curve is at a stretching temperature of 1 ⁇ 10 -3 S -1 and a tensile temperature of 900 ° C as shown in Fig. 17, indicating
  • the TiAl-Nb single crystal has a yield strength of 637 MPa at 900 ° C, an elongation of 8.1%, and a ductile-brittle transition temperature of more than 900 ° C, which is much higher than that of the general TiAl alloy.
  • Example 15 Ti-45Al-8Nb-0.4C-0.5Si was prepared, and the balance was Ti alloy, but the gravity casting method was used to obtain a round bar sample of ⁇ 8 mm, which was treated by necking and crystallizing.
  • the alloy of this diameter can also quickly obtain a single crystal sample, and the tensile strength after the same de-segregation heat treatment is 618 MPa, and the elongation is 9.2%.
  • the alloy composition was changed to Ti-45Al-8Nb-0.4Si-0.6C, and the balance was Ti (atomic percent), and the same heat treatment process was also employed, since a small amount of C and Si were not large.
  • the amplitude changes the phase transition temperature point, but brings about a high temperature strengthening effect, so that the material has a tensile yield strength of 650 MPa at 900 ° C and a plastic strain of 7.6%.
  • Example 15 The same preparation method as in Example 15 was employed, the alloy composition was Ti-45Al-8Nb-0.5Si, and the balance was Ti, and the drawing rate was changed to 40 mm/h, even if the temperature gradient was small, due to the neck-removing treatment, The single crystal was still obtained at a faster growth rate, and the tensile yield strength at 900 ° C after the heat treatment was 595 MPa, and the elongation was 8.7%.
  • Example 15 In the same preparation method as in Example 15, the alloy composition was changed to Ti-43Al-10Nb-0.3C-0.3Si, and the balance was Ti. Although the Nb element brought about strengthening, the segregation increased accordingly, but the heat treatment was carried out. The process still eliminates the brittle segregation phase, and the 900 °C tensile test results in a yield strength of 668 MPa and an elongation of 6%.
  • the alloy composition is Ti-45Al-8Nb-0.4C, and the balance is Ti, and the growth rate of the optical floating zone directional solidification process is changed to 5 mm/h, and the lower growth rate is more advantageous for single use.
  • the formation of crystals is manifested in the shorter distance of the elimination section.
  • the single crystal alloy material After vacuum heat treatment to segregation heat treatment, the single crystal alloy material has a tensile yield strength of 602 MPa and a plastic strain of 7.6% at 900 °C.
  • the alloy composition is Ti-45Al-8Nb, the balance is Ti, and the optical floating zone directional solidification process is a relative rotation speed of 20 rpm, and the temperature is more uneven due to the slower rotation speed, thereby intermediate grain growth.
  • the single crystal was obtained faster, the tensile yield strength at 900 ° C was 620 MPa, and the plastic strain was 7%.
  • the alloy composition was Ti-45Al-8Nb-0.4Si-0.6C, the balance was Ti, and the optical floating zone heating power was 65%, although the low heating temperature reduced the temperature gradient. Conducive to the formation of single crystal, but the method of neck-removing successfully succeeded in making the heating power still able to form a single crystal, and the yield strength after stretching at 900 ° C was 639 MPa, and the elongation was 7.2%.
  • the performance of the TiAl single crystal alloy prepared by the optical floating zone method was tested by conventional room temperature and high temperature tensile tests, and the alloy was found to have significantly higher room temperature and high temperature performance than other similar alloys (see the table for specific properties).
  • the room temperature brittleness of TiAl intermetallic compounds has been a major problem limiting its application.
  • the room temperature elongation of the TiAl alloy is 2 to 3%, and the TiAl alloy obtained in this patent has a room temperature elongation of 6.9% while maintaining a high strength (729 MPa).
  • the room temperature high plasticity solves the inherent room temperature processing of the TiAl alloy.
  • the problem is that the alloy can be easily machined to the desired part shape and its room temperature brittleness is improved. See Table 1 for comparison with the performance of some TiAl alloy single crystals.
  • Excellent yield strength at high temperature (900°C/637MPa): The yield strength of this patented alloy reaches 637MPa at 900°C, which is 30-50% higher than that of other TiAl alloys at 900°C. It is expected that the alloy will be used at temperatures from the current 650°C- 700 ° C increased to 900 ° C (currently, the United States GE company successfully applied Ti-48Al-2Cr-2Nb alloy in the Boeing 787 passenger aircraft low-pressure turbine six or seven-stage blades, its operating temperature is 650 ° C). The performance of 900 ° C with other TiAl alloys is shown in Table 2.
  • the excellent room temperature and high temperature performance of the TiAl single crystal is expected to expand the range of use in the engine blades of Boeing and Airbus aircraft to replace the engine blades at 650-900 ° C, which will bring energy saving and emission reduction. Great gains.
  • it has important application prospects in automotive turbocharger and exhaust valves, space momentum interceptor engine skirts, satellite engine nozzles, and reversible turbine rotors for spacecraft.

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Abstract

一种TiAl金属间化合物单晶材料及其制备方法。该材料的合金成分为Ti aAl bNb c(C,Si) d，其中43≤b≤49，2≤c≤10，a+b+c＝100，0≤d≤1（at.%）。

Description

TiAl金属间化合物单晶材料及其制备方法 技术领域
本发明属于轻质高强结构材料技术领域，具体涉及一种TiAl金属间化合物单晶材料及其制备方法。
背景技术
TiAl金属间化合物是一种新型轻质高温结构材料，其比重不到镍基高温合金的50％，具有高比强、高比刚、耐蚀、耐磨、耐高温、高弹性模量以及优异的抗氧化性、抗蠕变和良好的高温强度等优点，其使用温度可达750～900℃，与Ni基高温合金相近；但其密度仅为高温合金的一半，因而是理想的Ni基高温合金的替代材料，可广泛应用于汽车或航空发动机的高温部件如叶片、涡轮盘和排气阀等。例如TiAl合金是以克为减重单位的航空航天高温用材料，特别是发动机用最佳候选材料。美国GE公司成功利用Ti-48Al-2Cr-2Nb(4822)合金研制了波音飞机后两级低压涡轮叶片，使飞机减重200Kg左右。高Nb-TiAl合金其高温力学性能、抗蠕变性能及氧化性能显著高于普通TiAl合金，提高使用温度约60～100℃，是最有工程应用前景的TiAl合金。
但由于金属间化合物的本征脆性，TiAl合金室温脆性差成为阻碍其工业化应用的主要原因。同时，目前应用的4822合金服役温度仅为650℃，其高温性能有待进一步提高。所以，大量研究集中在调控TiAl合金组织结构以改善室温脆性并提高服役温度。由于TiAl合金PST晶体中强度与塑性呈现的明显的各向异性，将TiAl合金采用定向凝固的方法制作出全片层组织PST晶体，并使其片层组织取向平行于定向凝固中晶体的生长方向，可以提高TiAl合金的力学性能。
全片层组织TiAl合金的力学性能与其片层取向有着密切的关系。通过对具有单一取向的全片层多孪晶晶体PST(Polysynthetic twinned crystal)的研究，发现其强度与塑性呈现出明显的各向异性。由于全片层组织的这种各向异性，当片层取向合适时，使其更适合于航空发动机叶片这样一些要求耐高温，而又只受一维方向载荷的服役条件。如果能将TiAl合金采用定向凝固的方法制作出全片层组织的发动机叶片，并使其片层组织取向平行于叶片的轴向(定向凝固中晶体的生长方向)，无疑是极其有利的。Yamaguchi等人系统研究了TiAl合金片层取向对力学性能的影响，发现当载荷方向与片层取向平行时，屈服强度 和延伸率达到最佳组合。因此，要进一步提高TiAl合金的使用性能，就必须对最终组织的片层取向进行控制，以获得取向与载荷方向一致的TiAl金属间化合物单晶全片层组织。
目前，国内外TiAl合金片层取向的控制方法主要包括籽晶法和改变凝固路径的非籽晶法。Yamaguchi、Johnson等人通过α相凝固籽晶法，选用Ti-Al-Si系合金作为籽晶，通过缩颈选晶法得到了片层取向完全平行于生长方向的单晶PST。籽晶成分通常与母合金成分存在差异导致定向凝固合金的成分和性能不均匀，而且籽晶的制备工艺复杂。因此，籽晶法具有明显的不足。
非籽晶法中，目前国内外没有研究得到与生长方向平行的全片层TiAl单晶组织。林均品等在较低的G/V条件下，对Ti-46Al-5Nb合金采用“自籽晶法”(double directional solidification)的方法得到了与生长方向平行的全片层单晶组织。他们认为，较低的G/V工艺使β相枝晶间距在合适的条件下，可以通过全包晶反应得到与生长方向平行的单一取向的α相，而不会发生固态相变β→α生成不同位相的α变体，从而完成对片层取向的控制。这种方法需要进行两次相同工艺的定向凝固，较普通非籽晶法多一次凝固过程，加重了坩埚材料对合金的污染，对定向凝固TiAl合金的工业化不利。
之前国内外关于非籽晶法控制片层取向的研究均为改变凝固路径，不能控制单晶片层取向，而且没有得到与生长方向完全平行的单晶片层组织。为解决此技术难题，TiAl合金固态定向相变过程成为控制片层取向的关键。由相图可知，全片层组织TiAl合金凝固后还须经历β→α和α→α2+γ的固态相变。当初生相为β相时，择优生长方向为<001>，其位相关系为：{110}β//{0001}α//{111}γ [25]，且{110}β的12个变量中4个平行于生长方向，8个与生长方向倾斜成45°[16,26]，经历固态相变后形成的片层组织中仅有1/3惯习面的取向平行于生长方向。显然，TiAl合金最终片层组织的取向，不仅取决于初生β相的生长方向，还取决于之后的固态相变过程。所以，β→α固态相变过程也是控制片层取向的关键。而迄今为止关于TiAl片层取向控制的研究，均集中于凝固过程，而忽视了凝固之后的固态相变过程。
因此，不但要控制凝固过程，使定向凝固初生相为β相，而且要控制TiAl合金定向固态相变过程中新相的形核长大、相界定向迁移过程，使其在定向固态相变中只保留下来与生长方向呈0°的片层取向，完成在连续定向液固相变—固态相变条件下对TiAl合金片层取向的控制。
发明内容
本发明的目的是提供一种片层取向完全可控的具有优异室温及高温性能的TiAl金属 间化合物单晶材料。该材料具有理想片层取向且组织均匀无污染，保持较高的强度(729MPa)的同时室温拉伸塑性达到6.9％，900℃高温下屈服强度为637MPa，且韧脆转变温度达到900℃以上。
本发明的另一目的是提供一种上述TiAl金属间化合物单晶材料的制备方法。
本发明的目的可以通过以下措施达到：
一种TiAl金属间化合物单晶材料，以原子百分比计，该材料的合金成分表达式为TiaAlbNbc(C,Si)d，其中43≤b≤49，2≤c≤10，a+b+c＝100，0≤d≤1。
在一种方案中，该材料的合金成分表达式为TiaAlbNbc(C,Si)d，其中42≤a≤55，43≤b≤49，2≤c≤9，d＝0。
在另一种方案中，该材料的合金成分表达式为TiaAlbNbc(C,Si)d，其中44≤a≤51，43≤b≤47，6≤c≤9，d＝0。
在另一种方案中，该材料的合金成分表达式为TiaAlbNbc(C,Si)d，其中43≤b≤47，6≤c≤10，a+b+c＝100，0.1≤d≤1。
本发明TiAl金属间化合物单晶材料可以通过非籽晶法光学浮区定向凝固的方法制备得到，该方法包括如下步骤：
(1)分别选取纯度在99.9％以上的各纯物质原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚电磁感应悬浮熔炼炉中熔炼母合金锭，再采用重力铸造法或吸铸法获得母合金棒材；
(2)将母合金棒材切割为上、下棒料两部分，分别作为光学浮区定向凝固炉的原料棒与籽晶棒，控制上端原料棒与下端籽晶棒之间的距离为1～5mm；上、下棒料之间的距离为1～5mm；首先将原料棒与籽晶棒同轴且垂直于水平面设置，定向凝固时通入惰性气体进行保护，调节上、下棒料以相反方向旋转，相对转速为10～40rpm，启动加热，使上、下棒料的相对一端先熔化，调整上、下棒料的位置，使其相对一端逐渐接近后接合，调节设备功率并保温5-10min后当浮区表面光滑且熔化均匀时，调节生长速率为2.5～30mm/h，开始定向凝固；凝固结束后缓慢降低功率，同时并将已凝固试样与剩余送料棒试样进行缓慢分离；
(3)将制备的TiAl合金单晶棒材进行真空热处理，其采用1250℃～1350℃×12h～24h+900℃×30min/炉冷或空冷的热处理方式。
在一种优选方案中，在步骤(1)中，电磁感应悬浮熔炼采用水冷铜坩埚，母合金熔炼次数不少于3次，进一步优选不少于4次。
在一种优选方案中，在步骤(1)中，所述母合金棒材的尺寸为Φ(4～8)mm×120mm；所述吸铸法采用差压吸铸，其压强差保持在3MPa；采用重力铸造法时，保护气体压强在三分之二的标准大气压。所制备的母合金圆棒材的尺寸可以为Φ(4～8)mm。
在一种优选方案中，在步骤(1)中，Al、Ti、C或Si原材料的纯度在99.999％以上，Nb纯金属原材料的纯度在99.9％以上。
在一种优选方案中，在步骤(1)中，下端籽晶棒的长度为20-30mm，上端原料棒的长度小于190mm。
在一种优选方案中，在步骤(2)中，所述惰性气体为氩气或者氮气，定向凝固时通入的惰性气体流量为3～5L/min。
在一种优选方案中，在步骤(2)中，调整上下棒料的位置使其相对一端逐渐接近后接合，然后调节设备功率进行保温熔化。当设备总功率为4.0KW时，调节其功率至总功率的55～70％。
在一种优选方案中，在步骤(3)中，对单晶棒进行工艺为“1250℃×24h+900℃×30min+空冷”的去偏析真空热处理。
本发明进一步提供了另一种上述TiAl金属间化合物单晶材料的制备方法，该方法包括如下步骤：
第一步：选取纯度在99.999％以上的Ti、Al、Nb纯金属原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚悬浮熔炼炉中熔炼母合金，经3～4次熔炼使合金成分均匀化，并吸铸成定向凝固棒材；
第二步：将TiAl合金试棒放入高纯氧化钇涂层的刚玉坩埚中进行定向凝固，抽真空至5×10-3Pa，再向系统中充入高纯氩保护气；
第三步：调节感应电源功率对试样进行加热，保温温度为1450～1650K，保温时间为15～30min，开始定向凝固，控制定向凝固抽拉速率为5～20μm/s；持续生长至试样长度50mm处，启动快淬对定向凝固试样进行快淬处理，保留固液界面。
在一种优选方案中，在第一步中，定向凝固棒材的尺寸为Φ(4～6mm)×100mm。
在一种优选方案中，在第二步中，高纯氧化钇涂层的刚玉坩埚尺寸为Φ(7～9mm)×100mm；高纯氩保护气充入量为0.04～0.06MPa。
该方法的原理为采用Bridgman定向凝固方法控制TiAl合金片层取向，通过改变凝固参量温度梯度和生长速率，首先保证初生相为全β相，其次通过凝固过程中晶粒竞争淘汰获得单晶，并且在凝固过程中存在一个临界温度对应特定的抽拉速率，在此抽拉速率下 最终片层取向与生长方向为45°的α相通过相界迁移而消除，使β→α相变中得到的12个α变量中只保留最终片层取向与生长方向平行的α相，从而完成对片层取向的控制。
本发明与现有技术相比，具有如下优点：
(1)本发明提出的TiAl合金材料制备工艺方法能大幅度提高该合金的室温力学性能，尤其改善了室温脆性。本发明通过调节Nb元素含量，并添加少量C和Si强化元素，可有效改善合金高温力学性能。
(2)采用非籽晶光学浮区法定向凝固技术，防止合金污染的同时避免了籽晶法中加工复杂、成分组织不均匀的缺点，可避免传统Bridgman定向凝固带来合金污染问题，而且快速高效得到TiAl-Nb单晶。运用真空热处理完全消除了TiAl合金定向凝固后组织中残留的大量脆性B2相和富Nb脆性偏析相，从而获得组织均匀、室温性能优异的合金材料，而且避免了片层粗化。
(3)本方法还可采用普通的Bridgman定向凝固方法，通过调节凝固参量，控制连续定向液固相变—定向固态相变，保证全β相生长并且通过在固态相变控制最终片层取向，并得到片层取向完全平行于生长方向的TiAl合金单晶组织。本发明有效避免了籽晶法成分性能不均匀的缺点，同时在单次定向凝固过程中便得到了理想的片层取向的单晶组织，简化了工艺。
(4)本发明在制备TiAl合金单晶过程中，在一定范围凝固参量下能够完全控制其单晶片层取向。本发明为定向凝固TiAl合金的工业化应用提供了理论基础。
(5)本方法制备工艺简单，成本低，改善室温脆性效果显著，具有普遍适用性及推广价值。
附图说明
图1为现有技术部分Ti-Al二元合金相图。
图2为本发明定向凝固试样最大纵截面(a)及片层取向(b)的显微组织图。
图3为本发明定向凝固试样竞争段纵截面显微组织图。
图4为本发明定向凝固试样最大纵截面(a)及片层取向(b)的显微组织图。
图5为本发明定向凝固试样竞争段纵截面显微组织图。
图6为本发明定向凝固试样最大纵截面(a)及片层取向(b)的显微组织图。
图7为本发明定向凝固试样淬火固液界面。
注：附图2-7中显微组织生长方向为从右向左。
图8是为一种高强高塑TiAl合金材料的制备流程图。
图9是TiAl合金定向凝固单晶(a)及片层取向(b)显微组织。
图10是TiAl合金单晶不同热处理工艺前后偏析的显微结构(a为热处理前，b为热处理后)。
图11是不同热处理工艺前后TiAl合金单晶的XRD衍射图。
图12是TiAl合金室温拉伸力学性能曲线。
图13是为TiAl-Nb单晶凝固时固液界面形貌(a)以及缩颈选晶原理示意图(b)。
图14是TiAl-Nb合金定向凝固后得到的光学形貌。
图15是TiAl-Nb单晶热处理前(a)后(b)偏析扫描电镜图片。
图16是TiAl-Nb单晶热处理前(a)后(b)片层间距图片。
图17是TiAl-Nb在900℃下拉伸的位移强度曲线。
具体实施方式
以下结合附图和实施例，对本发明的内容做进一步说明。但本发明的保护范围并不局限于下述各例。
下面结合附图和一种Bridgman定向凝固方法制备片层取向完全可控的TiAl金属间化合物单晶，其具体实施方式如下：
(1)选择初生相为全β相的Ti-Al-Nb三元合金。根据多元合金相图及相选择原理，如图1，通过调整原子成分之间的配比关系，使其先析出相全部为β相。具体而言，提高Nb的含量，降低Al的相对比例，形成较宽的β相区。
(2)根据1)所得到的合金成分，采用高纯金属组元配置，并在高纯Ar气保护下，采用冷坩埚电磁悬浮熔炼设备熔制母合金。母合金多次熔炼得到均匀的母合金锭，并吸铸成母合金棒材。
(3)将TiAl合金棒材置入内壁烧结高纯氧化钇涂层的刚玉坩埚，刚玉坩埚尺寸为Φ(5～8mm)×100mm，放入Bridgman定向凝固炉，抽真空度置5×10-3Pa时，充入0.04～0.06MPa高纯氩保护气。
(4)调节感应电源功率对试样进行加热，保温温度为1450～1650K，保温时间为15～30min，开始定向凝固，控制定向凝固生长速率为5～20μm/s；
(5)在一定速率持续生长至试样长度50mm处，启动快淬对定向凝固试样进行快淬处理，保留固液界面。
下面结合具体实施例1-7和图1-7对本发明作进一步说明。
实施例1
实验所用合金成分为Ti47Al45Nb8(原子百分比at％)，其金属组元纯度均为99.999％，在高纯Ar气保护下，在真空度为5×10-3Pa下采用冷坩埚电磁悬浮熔炼设备熔制母合金。经4次熔炼得到均匀的母合金锭，并吸铸成Φ4×100mm母合金棒材。将TiAl合金试棒放入内壁涂有高纯氧化钇的刚玉坩埚中进行定向凝固实验，抽真空至5×10-3Pa，再向系统中充入0.05MPa高纯氩保护气。，调节感应电源功率对试样进行加热，保温温度为1550K，保温时间为25min，开始定向凝固，控制定向凝固生长速率为5μm/s；当抽拉长度至试样长度50mm处，启动快淬对试样进行快淬处理，保留固液界面。对该圆柱试样的最大纵截面进行显微组织表征，观察该抽拉速率下的凝固先析出相、晶粒大小和片层取向并进行分析，如图2(a)和图2(b)所示，发现得到片层取向平行于生长方向的TiAl合金单晶。生长速率较小为5μm/s时，溶质的富集能够充分扩散，生长能够稳定进行，晶粒有充分的时间长大，所以所得晶粒较为粗大直至获得单晶生长。
图3为5μm/s时定向凝固竞争段显微组织。由于在β→α固态相变中，由于形成0°和45°片层两种界面的错配度不同，导致不同的相界迁移率的不同，所以一个存在临界抽拉速率5μm/s，在此抽拉速率以下，形成0°与45°片层取向的α晶粒形核后0°晶粒生长驱动力较大，最终淘汰45°晶粒，得到片层取向平行于生长方向的单晶。
实施例2
采用与实施例1中相同的合金成分及方法，保温温度为1550K，保温时间为25min，开始定向凝固，控制定向凝固生长速率为15μm/s；如图4(a)和图4(b)所示，在此抽拉速率下β→α固态相变保留下的是45°片层取向的α相，所以最终组织为片层取向为45°的单晶。
图5为15μm/s时定向凝固竞争段显微组织。在此抽拉速率时，45°晶粒固态相变形核驱动力大于0°晶粒，以致0°晶粒不能生长，得到片层取向与生长方向呈45°的TiAl合金单晶。
实施例3
采用与实施例1中相同的合金成分及方法，保温温度为1550K，保温时间为25min，开始定向凝固，控制定向凝固生长速率为20μm/s；如图6(a)和图6(b)所示，获得片层取向与生长方向呈45°的单晶。
图7为快淬处理保留的固液界面，其枝晶生长形貌呈4重对称，具有较明显的二次枝晶且与一次枝晶干呈90°垂直关系，可以推断出在定向凝固过程中，立方晶系的β相是初生相。
实施例4
采用与实施例1相同的方法，所用合金成分为Ti55Al43Nb2，保温温度为1650K，保温时间为30min，定向凝固生长速率为5μm/s，获得与片层取向平行于生长方向的TiAl合金单晶。
实施例5
采用与实施例1相同的方法，所用合金成分为Ti48Al43Nb9，保温温度为1450K，保温时间为30min，定向凝固生长速率为10μm/s，获得片层取向与生长方向呈45°的TiAl合金单晶。
实施例6
采用与实施例1相同的方法，所用合金成分为Ti51Al45Nb6，保温温度为1650K，保温时间为15min，定向凝固生长速率为5μm/s，获得与片层取向平行于生长方向的TiAl合金单晶。
实施例7
采用与实施例1相同的方法，所用合金成分为Ti42Al49Nb9，保温温度为1550K，保温时间为25min，定向凝固生长速率为5μm/s，获得与片层取向平行于生长方向的TiAl合金单晶。
下面结合附图8和另一种高强高塑TiAl合金材料的制备方法，对本发明做进一步详细描述，其具体实施方式如下：
(1)结合图8，该高强高塑TiAl合金材料的合金成分原子百分比为：(44～51)Ti-(43～47)Al-(6～9)Nb。通过调整原子成分之间的配比关系，使其先析出相全部为β相。
(2)选取水冷铜坩埚电磁感应悬浮熔炼TiAl合金母合金纽扣锭，采用吸铸法获得母合金棒材。
(3)将母合金棒材切割为原料棒与籽晶棒两部分，通过光学浮区法进行定向凝固；通入高纯氩气作为气氛保护，调节上下段相对转速、加热功率以及生长速率控制TiAl合金片层取向并获得单晶生长。
(4)将制备的TiAl合金单晶棒材进行真空热处理；采用在α单相区加热一定时间后保温后退火；完全消除脆性B2相及残余应力，获得高强高塑TiAl合金材料。
(5)利用OM、XRD对制备的TiAl合金进行微观结构表征，并进一步对其进行力 学性能表征，以确定具有最佳综合力学性能的TiAl合金材料微观组织及其相应的制备工艺参数。
下面结合具体实施例8-13和图8-12对本发明作进一步说明。
实施例8
(1)原材料的选用：
本发明制备母合金锭选用的合金成分为Ti47Al45Nb8(原子百分比)，各金属组元的纯度Ti、Al为99.999％，Nb为99.95％。
(2)母合金锭的制备：
高纯氩气保护条件下，用水冷铜坩埚电磁感应悬浮熔炼炉熔制母合金锭：将金属原料的表面机械打磨去掉表面的氧化皮后，按照设计好的成分配比料备料；按照每锭70g左右的重量将配好的料放入熔炼炉内的水冷铜坩埚内，抽真空至5×10-3Pa；向炉内充入一定量压力的高纯氩气(99.999％)，氩气压力范围为0.8～1MPa。多道次熔炼3～4次得到混合均匀的母合金锭。随后将母合金锭吸铸成Φ6×120mm棒材。
(3)光学浮区定向凝固：
将母合金棒材切割为上、下棒料两部分，分别作为光学浮区定向凝固炉的原料棒与籽晶棒；下端为长度在30mm的籽晶棒，上端为长度小于100mm的送料棒；定向凝固时首先将原料棒与籽晶棒同轴且垂直于水平面设置，上、下棒料之间的距离为5mm且该间隔位于四个灯丝聚焦中心处；通入5L/min的高纯氩气作为气氛保护，调节上、下棒料的轴向相对转速30rpm，启动加热，使上、下棒料的相对一端先熔化，调整上、下棒料的位置，使其相对一端逐渐接近后接合，调节功率至总功率的68％，保温5min使浮区表面光滑熔化均匀时(即浮区无明显抖动时)，调节生长速率为5mm/h，开始定向凝固；生长至80mm时停止定向凝固，缓慢降低功率，同时并将已凝固试样与剩余送料棒试样进行缓慢分离；。
(4)真空热处理
将定向凝固棒材单晶部分放入刚玉管内，抽真空至10-3Pa后封管，放入热处理炉，采用1300℃×24h+900℃×30min/炉冷的热处理工艺。
(5)结构和性能表征
图9a是光学浮区定向凝固后的试棒宏观照片，可见试样在定向凝固中经历短暂竞争淘汰后迅速成为单晶生长，图9b表明单晶片层取向平行于生长方向。图10(a)与图10(b)是热处理前后显微组织图，结合图11的XRD图谱可以看出热处理前组织内部分布着大量的B2相，24h热处理后B2完全消除。图12为所制备的高强高塑TiAl合金的室温拉伸 应力-应变曲线，其屈服强度为729MPa，同时塑性应变达到6.9％，具有非常优异的室温力学性能。
实施例9
采用与实施例8相同的制备方法，合金成分为Ti44Al47Nb9(原子百分比)，光学浮区定向凝固工艺为相对转速20rpm，加热功率55％，生长速率为2.5mm/h，真空热处理工艺为1250℃×12h+900℃×30min/炉冷，B2相完全消除，获得TiAl合金材料室温拉伸屈服强度为550MPa，塑性应变为6.0％。
实施例10
采用与实施例8相同的制备方法，合金成分为Ti51Al40Nb9(原子百分比)，光学浮区定向凝固工艺为相对转速25rpm，加热功率70％，生长速率为10mm/h，真空热处理工艺为1300℃×20h+900℃×30min/炉冷，B2相完全消除，获得TiAl合金材料室温拉伸屈服强度为628MPa，塑性应变为6.5％。
实施例11
采用与实施例8相同的制备方法，合金成分为Ti48Al43Nb9(原子百分比)，光学浮区定向凝固工艺为相对转速20rpm，加热功率68％，生长速率为15mm/h，真空热处理工艺为1350℃×24h+900℃×30min/炉冷，B2相完全消除，获得TiAl合金材料室温拉伸屈服强度为660MPa，塑性应变为6.2％。
实施例12
采用与实施例8相同的制备方法，合金成分为Ti48Al43Nb9(原子百分比)，光学浮区定向凝固工艺为相对转速20rpm，加热功率70％，生长速率为15mm/h，真空热处理工艺为1350℃×12h+900℃×30min/炉冷，B2相完全消除，获得TiAl合金材料室温拉伸屈服强度为593MPa，塑性应变为6.8％。
实施例13
采用与实施例8相同的制备方法，合金成分为Ti48Al46Nb6(原子百分比)，光学浮区定向凝固工艺为相对转速30rpm，加热功率60％，生长速率为20mm/h，真空热处理工艺为1250℃×12h+900℃×30min/炉冷，B2相未完全消除，如图10b的XRD图谱中12h热处理发现少量B2相残留，获得TiAl合金材料室温拉伸屈服强度为656MPa，塑性应变为3.0％。
实施例14
采用与实施例8相同的制备方法，合金成分为Ti44Al45Nb8(原子百分比)，光学浮区 定向凝固工艺为相对转速25rpm，加热功率55％，生长速率为30mm/h，得到片层取向与生长方向呈45°的TiAl合金单晶，真空热处理工艺为1250℃×12h+900℃×30min/炉冷，B2相完全消除，获得TiAl合金材料室温拉伸屈服强度为430MPa，塑性应变为7.8％
实施例15
(1)选用合金成分原子百分比为：Ti-45Al-8Nb-0.3C-0.2Si，余量为Ti，初始原料为99.999％的高纯Al、Ti、C和Si以及99.95％的高纯Nb，在5×10-3MPa真空度下，采用冷坩埚电磁感应悬浮熔炼炉反复4次熔炼获得TiAl-Nb母合金铸锭；
(2)采用压强差为3MPa的差异吸铸法获得Φ6mm圆棒状合金。
(3)采用光学浮区法进行定向凝固，将吸铸获得的合金棒材切割成下端为20mm的籽晶棒，上端为150mm的送料棒，并对送料棒进行缩颈选晶处理。定向凝固时首先将送料棒与籽晶棒安装好，送料棒与籽晶棒必须保持同轴且垂直于水平面设置，送料棒与籽晶棒相距1-3mm，且其相对一端位于四个灯丝聚焦中心处；通入流速为4L/min的保护气体，并调节籽晶棒和送料棒分别以30r/min反方向旋转，加热功率在10min内升为68％使合金熔化，保温5min后以15mm/h的生长速率进行定向凝固。由于光学浮区加热特点致使固液界面为图13(a)所示的凸形界面，原理图为图13(b)所示，中间晶粒将沿着生长方向生长，而两侧晶粒将向两侧斜方向生长，因此对送料棒进行缩颈选晶处理后，将会使中间生长的晶粒淘汰两侧晶粒，快速长大为单晶，从竞争生长到最终稳定生长的过程过图14所示；定向凝固结束后缓慢降低功率，同时并将已凝固试样与剩余送料棒试样进行缓慢分离；
(4)将制备好的TiAl-Nb单晶进行真空去偏析热处理，热处理前的偏析形貌如图15(a)所示，在1250℃的α单相区加热24h后可使偏析相消除，再经过900℃的30min均匀化处理，以空冷方式得到最终单晶，图15(b)所示说明热处理已经完全消除。图16为热处理前后片层间距变化，由于空冷冷速较快，片层来不及发生粗化。
(5)将热处理单晶加工成标距为Φ3mm×20mm拉伸试样，在拉伸速率为1×10-3S-1和900℃拉伸温度下拉伸曲线如图17所示，表明该TiAl-Nb单晶在900℃时屈服强度为637MPa，延伸率为8.1％，韧脆转变温度大于900℃，远高于一般TiAl合金。
实施例16
采用与实施例15相同的方法制备Ti-45Al-8Nb-0.4C-0.5Si，余量为Ti的合金，但改变采用重力铸造方法得到Φ8mm的圆棒状试样，由于采用缩颈选晶处理，该直径的合金也能快速得到单晶试样，相同去偏析热处理后拉伸强度为618MPa，延伸率为9.2％。
实施例17
采用与实施例15相同的制备方法，改变合金成分为Ti-45Al-8Nb-0.4Si-0.6C,余量为Ti(原子百分比)，同样采用相同热处理工艺，由于少量的C和Si并没有大幅度改变相变温度点，但是带来了高温强化作用，使材料900℃拉伸屈服强度为650MPa，塑性应变为7.6％。
实施例18
采用与实施例15相同的制备方法，合金成分为Ti-45Al-8Nb-0.5Si，余量为Ti，改变抽拉速率为40mm/h，即使温度梯度较小，由于采用缩颈选晶处理，较快生长速度下依然得到了单晶，热处理后900℃拉伸屈服强度为595MPa，延伸率为8.7％。
实施例19
采用与实施例15相同的制备方法，改变合金成分为Ti-43Al-10Nb-0.3C-0.3Si，余量为Ti，虽然Nb元素带来强化作用，但来偏析也会相应增多，但该热处理工艺依然可消除脆性偏析相，900℃拉伸结果表明屈服强度达到668MPa，延伸率6％。
实施例20
采用与实施例15相同的制备方法，合金成分为Ti-45Al-8Nb-0.4C，余量为Ti，改变光学浮区定向凝固工艺生长速率为5mm/h，较低的生长速率更有利于单晶的形成，表现在淘汰段距离变短，真空热处理去偏析热处理后，单晶合金材料在900℃拉伸屈服强度为602MPa，塑性应变为7.6％。
实施例21
采用与实施例15相同的制备方法，合金成分为Ti-45Al-8Nb，余量为Ti，光学浮区定向凝固工艺为相对转速20rpm，由于转速变慢使得温度更加不均匀，从而中间晶粒生长更快而得到单晶，900℃拉伸屈服强度为620MPa，塑性应变为7％。
实施例22
采用与实施例15相同的制备方法，合金成分为Ti-45Al-8Nb-0.4Si-0.6C，余量为Ti，光学浮区加热功率为65％，虽然低的加热温度使温度梯度减小不利于单晶形成，但是缩颈选晶这一方法成功使得该加热功率依然能够形成单晶，900℃拉伸后屈服强度为639MPa，延伸率为7.2％。
实施例23、应用及性能对比
通过常规的室温和高温拉伸测试对光学浮区法制备的TiAl单晶合金进行性能检测，发现该合金具有显著优于其他类似合金的室温和高温性能(具体性能参见表格)。
TiAl金属间化合物的室温脆性一直是限制其应用的一大难题。通常TiAl合金室温延伸率在2～3％，而本专利得到的TiAl合金，其室温延伸率高达6.9％，同时保持了较高的强度(729MPa)，室温高塑性解决了TiAl合金固有的室温加工难题，从而使合金容易的加工出所需要的零件形状，并改善了其室温脆性。与部分TiAl合金单晶性能对比见表1。
高温下优良的屈服强度(900℃/637MPa)：本专利合金900℃时屈服强度达到637MPa，较其它TiAl合金900℃强度提高30～50％，有望使该合金的使用温度从目前的650℃-700℃提高到900℃(目前，美国GE公司成功应用Ti-48Al-2Cr-2Nb合金在波音787客机低压涡轮六、七级叶片，其工作温度为650℃)。与其他TiAl合金900℃性能对比如表2。该TiAl单晶优良的室温和高温性能，有望在波音飞机和空中客车飞机中的发动机叶片中扩大使用范围，以代替650-900℃使用温度下的发动机叶片，将在节能减排等方面带来巨大收益。另外，在汽车增压涡轮和排气阀，空间动量拦截器发动机尾裙，卫星发动机喷管，航天飞行器用可逆涡轮转子等部件上具有重要的应用前景。
表1.TiAl合金PST单晶室温拉伸性能对比
Figure PCTCN2015091508-appb-000001
表2.不同组织TiAl合金900℃高温拉伸性能对比
Table 2.Comparison of elevated temperature(900C)mechanical properties of some TiAl alloys with different microstructures,including fully lamellar(FL),nearly fully lamellar(NFL),near gamma(NG),nearly lamellar(NL),degraded fully lamellar(DFL),refined fully lamellar(RFL)and duplex(DP).
Figure PCTCN2015091508-appb-000002
1.D.R.Johnson,H.Inui,M.Yamaguchi,Acta Mater.44,2523-2535(1996).
2.D.R.Johnson,Y.Masuda,H.Inui and M.Yamaguchi,Acta Mater.45,2523-2533(1997).
3.I.S.Jung,H.S.Jang,M.H.Oh,J.H.Lee,D.M.Wee,Mater.Sci.Eng.A 329-331,13-18(2002).
4.T.Yamanaka,D.R.Johnson,H.Inui,M.Yamaguchi,Intermetallics7,779-784(1997).
5.D.R.Johnson,K.Chihara,H.Inui,M.Yamaguchi,Acta Mater.46,6529-6540(1998).
6.H.N.Lee,D.R.Johnson,H.Inui,M.H.Oh,D.M.Wee,M.Yamaguchi,Acta Mater.48,3221-3233(2000).
7.G.L.Chen,W.J.Zhang,Z.C.Liu,S.J.Li,Y.W.Kim,Gamma titanium aluminides,31-40(1999).
8.H.Clemens,I.Rumberg,P.Schretter.Intermetallics,2(3),179–184(1994).
9.T.Tetsui.Structural Intermetallics,489-493(1997).
10.Y.W.Kim.Mater.Sci.Eng.A.192-193,519-533(1995).
11.Z.C.Liu,J.P.Lin,S.J.Li.Intermetallics10(7),653-659(2002).

Claims (20)

  1. 一种TiAl金属间化合物单晶材料，其特征在于以原子百分比计，该材料的合金成分表达式为TiaAlbNbc(C,Si)d，其中43≤b≤49，2≤c≤10，a+b+c＝100，0≤d≤1。
  2. 根据权利要求1所述的TiAl金属间化合物单晶材料，其特征在于其中42≤a≤55，43≤b≤49，2≤c≤9，d＝0。
  3. 根据权利要求1所述的TiAl金属间化合物单晶材料，其特征在于其中44≤a≤51，43≤b≤47，6≤c≤9，d＝0。
  4. 根据权利要求1所述的TiAl金属间化合物单晶材料，其特征在于其中43≤b≤47，6≤c≤10，a+b+c＝100，0.1≤d≤1。
  5. 根据权利要求1～4中任意一项所述的TiAl金属间化合物单晶材料，其特征在于该材料由包括如下步骤的方法制备：
    (1)分别选取纯度在99.9％以上的各纯物质原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚电磁感应悬浮熔炼炉中熔炼母合金锭，再采用重力铸造法或吸铸法获得母合金棒材；
    (2)将母合金棒材切割为上、下棒料两部分，分别作为光学浮区定向凝固炉的原料棒与籽晶棒；控制上端原料棒与下端籽晶棒之间的距离为1～5mm；首先将原料棒与籽晶棒同轴且垂直于水平面设置，定向凝固时通入惰性气体进行保护，调节上、下棒料以相反方向旋转，相对转速为10～40rpm，启动加热，使上、下棒料的相对一端先熔化，调整上、下棒料的位置，使其相对一端逐渐接近后接合，调节设备功率并保温5-10min后当浮区表面光滑且熔化均匀时，调节生长速率为2.5～30mm/h，开始定向凝固；凝固结束后缓慢降低功率，同时并将已凝固试样与剩余送料棒试样进行缓慢分离；
    (3)将制备的TiAl合金单晶棒材进行真空热处理，其采用1250℃～1350℃×12h～24h+900℃×30min/炉冷或空冷的热处理方式。
  6. 根据权利要求5所述的TiAl金属间化合物单晶材料，其特征在于在步骤(1)中，电磁感应悬浮熔炼采用水冷铜坩埚，母合金熔炼次数不少于3次。
  7. 根据权利要求5所述的TiAl金属间化合物单晶材料，其特征在于在步骤(1)中，所述母合金棒材的尺寸为Φ(4～8)mm×120mm；所述吸铸法采用差压吸铸，其压强差保持在3MPa；采用重力铸造法时，保护气体压强在三分之二的标准大气压。
  8. 根据权利要求5所述的TiAl金属间化合物单晶材料，其特征在于在步骤(2)中，所述惰性气体为氩气或者氮气，定向凝固时通入的惰性气体流量为3～5L/min。
  9. 根据权利要求5所述的TiAl金属间化合物单晶材料，其特征在于步骤(1)中，所述Al、Ti、C或Si原材料的纯度在99.999％以上，Nb纯金属原材料的纯度在99.9％以上。
  10. 根据权利要求5所述的TiAl金属间化合物单晶材料，其特征在于步骤(1)中，下端籽晶棒的长度为20-30mm，上端原料棒的长度小于190mm。
  11. 根据权利要求1～4中任意一项所述的TiAl金属间化合物单晶材料，其特征在于该材料由包括如下步骤的方法制备：
    第一步：选取纯度在99.9％以上的各纯物质原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚电磁感应悬浮熔炼炉中熔炼母合金锭，经3～4次熔炼使合金成分均匀化，并吸铸成定向凝固棒材；
    第二步：将TiAl合金试棒放入高纯氧化钇涂层的刚玉坩埚中进行定向凝固，抽真空至5×10-3Pa，再向系统中充入高纯氩保护气；
    第三步：调节感应电源功率对试样进行加热，保温温度为1450～1650K，保温时间为15～30min，开始定向凝固，控制定向凝固抽拉速率为5～20μm/s；持续生长至试样长度50mm处，启动快淬对定向凝固试样进行快淬处理，保留固液界面。
  12. 根据权利要求11所述的TiAl金属间化合物单晶材料，其特征在于在第一步中，定向凝固棒材的尺寸为Φ(4～6mm)×100mm；在第二步中，高纯氧化钇涂层的刚玉坩埚尺寸为Φ(7～9mm)×100mm；高纯氩保护气充入量为0.04～0.06MPa。
  13. 一种权利要求1～4中任意一项所述的TiAl金属间化合物单晶材料的制备方法，其特征在于其包括如下步骤：
    (1)分别选取纯度在99.9％以上的各纯物质原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚电磁感应悬浮熔炼炉中熔炼母合金锭，再采用重力铸造法或吸铸法获得母合金棒材；
    (2)将母合金棒材切割为上、下棒料两部分，分别作为光学浮区定向凝固炉的原料棒与籽晶棒；上端原料棒与下端籽晶棒之间的距离为1～5mm；首先将原料棒与籽晶棒同轴且垂直于水平面设置，定向凝固时通入惰性气体进行保护，调节上、下棒料以相反方向旋转，相对转速为10～40rpm，启动加热，使上、下棒料的相对一端先熔化，调整上、下棒料的位置，使其相对一端逐渐接近后接合，调节设备功率并保温5-10min后当浮区表面光滑且熔化均匀时，调节生长速率为2.5～30mm/h，开始定向凝固；凝固结束后缓慢降低功 率，同时并将已凝固试样与剩余送料棒试样进行缓慢分离；
    (3)将制备的TiAl合金单晶棒材进行真空热处理，其采用1250℃～1350℃×12h～24h+900℃×30min/炉冷或空冷的热处理方式。
  14. 根据权利要求13所述的方法，其特征在于在步骤(1)中，电磁感应悬浮熔炼采用水冷铜坩埚，合金熔炼次数不少于3次。
  15. 根据权利要求13所述的方法，其特征在于在步骤(1)中，所述母合金棒材的尺寸为Φ(4～8)mm×120mm；所述吸铸法采用差压吸铸，其压强差保持在3MPa；采用重力铸造法时，保护气体压强在三分之二的标准大气压。
  16. 根据权利要求13所述的方法，其特征在于步骤(1)中，所述Al、Ti、C或Si原材料的纯度在99.999％以上，Nb纯金属原材料的纯度在99.9％以上。
  17. 根据权利要求13所述的方法，其特征在于步骤(1)中，下端籽晶棒的长度为20-30mm，上端原料棒的长度小于190mm。
  18. 根据权利要求13所述的方法，其特征在于在步骤(2)中，所述惰性气体为氩气或者氮气，定向凝固时通入的惰性气体流量为3～5L/min。
  19. 一种权利要求1～4中任意一项所述的TiAl金属间化合物单晶材料的制备方法，其特征在于其包括如下步骤：
    第一步：选取纯度在99.9％以上的各纯物质原材料，按照合金成分表达式进行配比，在真空度小于10-3Pa的冷坩埚电磁感应悬浮熔炼炉中熔炼母合金锭，经3～4次熔炼使合金成分均匀化，并吸铸成定向凝固棒材；
    第二步：将TiAl合金试棒放入高纯氧化钇涂层的刚玉坩埚中进行定向凝固，抽真空至5×10-3Pa，再向系统中充入高纯氩保护气；
    第三步：调节感应电源功率对试样进行加热，保温温度为1450～1650K，保温时间为15～30min，开始定向凝固，控制定向凝固抽拉速率为5～20μm/s；持续生长至试样长度50mm处，启动快淬对定向凝固试样进行快淬处理，保留固液界面。
  20. 根据权利要求19所述的方法，其特征在于在第一步中，定向凝固棒材的尺寸为Φ(4～6mm)×100mm；在第二步中，高纯氧化钇涂层的刚玉坩埚尺寸为Φ(7～9mm)×100mm；高纯氩保护气充入量为0.04～0.06MPa。
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* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP3326746A1 (en) * 2016-11-25 2018-05-30 Helmholtz-Zentrum Geesthacht Zentrum für Material- und Küstenforschung GmbH Method for joining and/or repairing substrates of titanium aluminide alloys
RU2725229C1 (ru) * 2019-12-03 2020-06-30 Федеральное государственное бюджетное учреждение науки Институт физики твердого тела Российской академии наук (ИФТТ РАН) Способ получения Ti2MnAl
CN112746187B (zh) * 2020-12-21 2021-10-01 安徽工程大学 一种大尺寸的TiAl基合金籽晶的制备方法
CN112746329A (zh) * 2020-12-21 2021-05-04 安徽工程大学 一种用于TiAl基合金定向凝固的籽晶制备方法
CN114908422B (zh) * 2022-06-29 2024-06-14 合肥工业大学 锶掺杂六硼化镧单晶及其制备方法
CN115198356B (zh) * 2022-07-15 2023-07-21 郑州大学 一种特定取向的大规格金属单晶及其制备方法
CN115558811B (zh) * 2022-09-10 2023-06-16 哈尔滨工业大学 利用超声与电磁场制备TiAl半固态材料的装备与方法
CN116516213B (zh) * 2023-03-23 2024-06-07 北京科技大学 一种含Si高Nb-TiAl合金的制备方法

Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4461659A (en) * 1980-01-17 1984-07-24 Cannon-Muskegon Corporation High ductility nickel alloy directional casting of parts for high temperature and stress operation
CN1958817A (zh) * 2006-09-22 2007-05-09 北京科技大学 一种利用放电等离子烧结制备高铌钛铝合金材料的方法
CN102011195A (zh) * 2010-11-23 2011-04-13 北京科技大学 一种定向凝固高铌钛铝合金单晶的制备方法
CN102921929A (zh) * 2012-11-01 2013-02-13 哈尔滨工业大学 一种高铌钛铝金属间化合物的无污染定向凝固方法
CN104278173A (zh) * 2014-10-09 2015-01-14 南京理工大学 一种高强高塑TiAl合金材料及其制备方法
CN104328501A (zh) * 2014-10-09 2015-02-04 南京理工大学 一种片层取向完全可控的TiAl单晶合金及其制备方法
CN104878452A (zh) * 2015-05-13 2015-09-02 南京理工大学 一种高温高强TiAl-Nb单晶及其制备方法

Family Cites Families (12)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP0530968A1 (en) * 1991-08-29 1993-03-10 General Electric Company Method for directional solidification casting of a titanium aluminide
US6059015A (en) * 1997-06-26 2000-05-09 General Electric Company Method for directional solidification of a molten material and apparatus therefor
JP3054697B2 (ja) * 1998-06-30 2000-06-19 京都大学長 Ｔｉ−Ａｌ系合金の製造方法
US7418993B2 (en) * 1998-11-20 2008-09-02 Rolls-Royce Corporation Method and apparatus for production of a cast component
US6436208B1 (en) * 2001-04-19 2002-08-20 The United States Of America As Represented By The Secretary Of The Navy Process for preparing aligned in-situ two phase single crystal composites of titanium-niobium alloys
US6755239B2 (en) * 2001-06-11 2004-06-29 Santoku America, Inc. Centrifugal casting of titanium alloys with improved surface quality, structural integrity and mechanical properties in isotropic graphite molds under vacuum
TWI265198B (en) * 2002-12-02 2006-11-01 Univ Nat Taiwan The method and equipments for controlling the solidification of alloys in induction melting using cold crucible
KR100644880B1 (ko) * 2004-11-30 2006-11-15 한국과학기술원 층상조직의 열적안정성 및 기계적 성질이 매우 우수한 일방향응고 TiAlNbSiC계 합금
US20120152483A1 (en) * 2010-12-16 2012-06-21 General Electric Company Unidirectional solidification process and apparatus and single-crystal seed therefor
CN102400074A (zh) * 2011-12-02 2012-04-04 南京理工大学 一种减少定向凝固高铌钛铝合金显微偏析的热处理工艺
CN102672150A (zh) * 2012-05-14 2012-09-19 北京科技大学 一种钛铝铌合金片层组织方向控制方法
CN103789598B (zh) * 2014-02-28 2015-12-02 南京理工大学 一种定向TiAl基合金及其制备方法

Patent Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4461659A (en) * 1980-01-17 1984-07-24 Cannon-Muskegon Corporation High ductility nickel alloy directional casting of parts for high temperature and stress operation
CN1958817A (zh) * 2006-09-22 2007-05-09 北京科技大学 一种利用放电等离子烧结制备高铌钛铝合金材料的方法
CN102011195A (zh) * 2010-11-23 2011-04-13 北京科技大学 一种定向凝固高铌钛铝合金单晶的制备方法
CN102921929A (zh) * 2012-11-01 2013-02-13 哈尔滨工业大学 一种高铌钛铝金属间化合物的无污染定向凝固方法
CN104278173A (zh) * 2014-10-09 2015-01-14 南京理工大学 一种高强高塑TiAl合金材料及其制备方法
CN104328501A (zh) * 2014-10-09 2015-02-04 南京理工大学 一种片层取向完全可控的TiAl单晶合金及其制备方法
CN104878452A (zh) * 2015-05-13 2015-09-02 南京理工大学 一种高温高强TiAl-Nb单晶及其制备方法

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
See also references of EP3205753A4 *

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RU2701438C2 (ru) 2019-09-26
US20170268127A1 (en) 2017-09-21
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