Method for prolonging high-temperature creep life of beta-gamma type TiAl alloy

Abstract

The invention discloses a method for prolonging the high-temperature creep life of a beta-gamma type TiAl alloy. The method comprises the following steps that 1, an alloy ingot with the nominal composition of Ti-42A1-5Mn (at%) is smelted through vacuum induction; forging the alloy ingot, cooling to obtain a forged blank, and carrying out homogenization solution treatment; and 2, carrying out isothermal annealing treatment on the sample subjected to homogenization solution treatment in the step 1. According to the method, homogenization and isothermal two-step heat treatment are carried out on the beta-gamma type TiAl alloy, so that a beta o phase unfavorable for high-temperature mechanical properties in the alloy is subjected to solid-state phase change and completely disappears, and a newly-separated hard Laves phase coherent with a matrix remarkably inhibits dislocation proliferation and slippage, so that the high-temperature creep life of the alloy is remarkably prolonged.

Description

TECHNICAL FIELD

[0001] The present application relates to the technical field of metal materials, in particular to a method for improving the high-temperature creep life of a beta-gamma type TiAl alloy. BACKGROUND

[0002] Titanium-aluminum alloys are expected to replace heavy nickel-based high-temperature alloys and be applied to turbine blades of an aero-engine, so as to greatly improve the thermal efficiency and ecological compatibility of an aircraft energy conversion system, due to their low density, high specific strength, excellent high-temperature creep and oxidation resistance and other characteristics, which meet the design concept of "higher", "stronger" and "lighter" of aerospace.

[0003] In particular, beta-solidification gamma-TiAl alloys show a significantly better hot working capacity than traditional gamma-TiAl alloys, thereby attracting more and more attention from researchers. However, the addition of a stable beta phase element will inevitably cause the phase boundary to move, compress the alpha single-phase region to some extent, and even make it completely disappear. Therefore, the beta to alpha phase transition during solidification of the alloy is more or less inhibited, resulting in partial high-temperature beta phase ordering into room-temperature beta o phase. Generally, the ordered beta o phase has intrinsic brittleness and a strong decomposition tendency in a medium temperature range, such as the precipitation of omega phase, which is not conducive to the overall performance of the alloy, especially significantly deteriorating the high-temperature mechanical properties of the alloy. Therefore, it is urgent to find an effective method to improve the high-temperature mechanical properties of the alloy. SUMMARY

[0004] The present application aims to solve the problem that the ordered beta o phase of the TiAl alloy has intrinsic brittleness and a strong decomposition tendency in a medium temperature range, such as the precipitation of omega phase, which is not conducive to the overall performance of the alloy, especially significantly deteriorating the high-temperature mechanical properties of the alloy, and proposes a method for improving the high-temperature creep life of a beta-gamma type TiAl alloy.

[0005] In order to achieve the above technical purposes, the technical scheme provided by the present application is as follows: A method for improving the high-temperature creep life of a beta-gamma type TiAl alloy, comprising the following steps: Step 1: vacuum induction smelting an alloy ingot with a nominal composition of Ti-42A1-5Mn; after forging and cooling the alloy ingot, a forged blank is obtained, and then homogenization and solid solution treatment is performed; Step 2: isothermally annealing the sample subjected to the homogenization and solid solution treatment in step 1.

[0006] Further, the homogenization and solid solution treatment temperature is 1000℃-1400℃.

[0007] Further, the method for determining the homogenization solution treatment temperature comprises the following steps: firstly, heat treating the alloy at an initial temperature, then characterizing the microstructure of the alloy, if the β o phase in the alloy does not have lath-shaped precipitates, then the initial temperature is determined as the homogenization solution treatment temperature; otherwise, if the β o phase in the alloy has lath-shaped precipitates, then the homogenization treatment temperature of the alloy needs to be gradually increased by a gradient temperature until the β o phase in the alloy disappears after the homogenization treatment, thus determining the homogenization solution treatment temperature.

[0008] Further, the gradient temperature is 25-50°C.

[0009] Specifically, firstly, heat treating the alloy at 1000°C, then characterizing the microstructure of the alloy, if the β o phase in the alloy does not have lath-shaped precipitates, then the 1000°C is determined as the homogenization solution treatment temperature; otherwise, if the β o phase in the alloy has lath-shaped precipitates, then the homogenization treatment temperature of the alloy needs to be gradually increased by 50°C until the β o phase in the alloy disappears after the homogenization treatment, thus determining the homogenization solution treatment temperature.

[0010] Further, the homogenization solution treatment time is 0.5-10 h.

[0011] Further, the homogenization solution treatment time is 0.5-5 h.

[0012] Further, the homogenization solution treatment time is 1-2 h.

[0013] Further, the isothermal annealing temperature is 700-900°C.

[0014] Further, the method for determining the isothermal annealing temperature comprises the following steps: heat treating the sample after the homogenization solution treatment at an initial temperature, then characterizing the microstructure of the alloy, if the β o phase in the alloy is completely transformed into the mixed microstructure of Laves+γ+α2, then the initial temperature is the isothermal annealing temperature; otherwise, if the β o phase still exists in the alloy, then the isothermal annealing temperature of the alloy needs to be gradually decreased by a gradient temperature until the β o phase in the alloy disappears after the isothermal annealing treatment, thus determining the isothermal annealing temperature.

[0015] Further, the gradient temperature is 50°C.

[0016] Specifically, the homogenization and solid solution treated sample is heat treated at 900℃, and then the microstructure of the alloy is characterized, if the initial β o phase is completely transformed into a mixed structure of Laves+γ+α2, then 900℃ is the isothermal annealing temperature; otherwise, if the β o phase still exists in the alloy, it is necessary to gradually reduce the isothermal annealing temperature of the alloy according to the gradient temperature until the β o phase completely disappears, so as to determine the isothermal annealing temperature.

[0017] Further, the isothermal annealing heat treatment time is 9h-480h.

[0018] Further, the isothermal annealing heat treatment time is 24h-480h.

[0019] Further, the isothermal annealing heat treatment time is 360h-480h.

[0020] Further, the cooling mode of the homogenization and solid solution treatment and the isothermal annealing treatment is water quenching.

[0021] Specifically, the phase structure of the TiAl alloy obtained by the above two-step heat treatment is: lamellar structure composed of γ / α2+ primary blocky γ phase+α2 phase+secondary γ+ellipsoidal Laves phase, wherein the Laves phase is precipitated from the β phase during isothermal annealing.

[0022] The present application has the following beneficial effects: 1. The present application carries out two-step heat treatment of homogenization and isothermal annealing on the TiAl alloy, so that the β o phase in the alloy which is not conducive to high-temperature mechanical properties is completely eliminated by solid-state phase transition, and the newly precipitated Laves phase which is coherent with the matrix significantly suppresses the multiplication and slip of dislocations, thereby significantly improving the high-temperature creep life of the alloy.

[0023] 2. The present application significantly prolongs the creep rupture life of the heat treated sample under the condition of 300MPa or 350MPa uniaxial stress at 750℃: under the condition of constant stress of 350MPa, the creep life is increased from 18.20h of the untreated state to 28.57h after heat treatment; under the condition of constant stress of 300MPa, the creep life is increased from 116.48h of the untreated state to 183.88h after heat treatment. BRIEF DESCRIPTION OF DRAWINGS

[0024] Figure 1 X-ray diffraction patterns of Ti-42Al-5Mn alloy before and after isothermal annealing after homogenization treatment are shown.

[0025] Figure 2 The β phase region of the Ti-42Al-5Mn alloy after isothermal annealing at 800℃ for 480h is shown. o The high-angle annular dark-field image and the corresponding element mapping of the phase region scanning transmission electron microscopy are shown.

[0026] Figure 3 The creep curves of the Ti-42Al-5Mn alloy before and after heat treatment at 750℃ under different stress conditions are shown.

[0027] Figure 4 The interaction between the Laves phase and dislocations near the creep fracture of the heat-treated sample under a constant stress of 300 MPa at 750℃ is shown. DETAILED DESCRIPTION

[0028] The technical solutions of the present application will be described below in a clear and complete manner. Obviously, the described embodiments are part of the embodiments of the present application, rather than all the embodiments. Based on the embodiments in the present application, all other embodiments obtained by those skilled in the art without creative labor fall within the scope of protection of the present application.

[0029] In addition, the technical features involved in different embodiments of the present application described below can be combined with each other as long as they do not conflict with each other.

[0030] Embodiment 1 The alloy composition selected in this embodiment is Ti-42Al-5Mn (atomic percentage), and the specific implementation steps are as follows: a Ti-42Al-5Mn alloy in an annealed state after forging is selected.

[0031] (1) A plurality of φ10 mm x 12 mm round bar samples are obtained by electric spark cutting.

[0032] (2) After the sample is cleaned with alcohol, it is packaged in a quartz glass tube filled with argon to prevent surface oxidation of the sample during heat treatment.

[0033] (3) Take one sample and first perform homogenization heat treatment at a temperature of 1000℃ in a high-temperature resistance furnace. In order to minimize the temperature drop during quenching, the sample is placed in a self-designed semi-automatic isothermal quenching furnace. When the set holding time and temperature are reached, one-key quenching can be realized. The time interval from heating temperature to quenching is less than 1s, which ensures a small temperature drop. The entire heat treatment process is as follows: the sample is placed in the high-temperature resistance furnace after the target set temperature is reached, and the heat treatment time is 1h.

[0034] (4) The sample obtained in step (3) is subjected to general metallographic preparation (mainly including sample grinding, polishing and etching) to obtain a sample for observing the microstructure of the alloy, and the microstructure of the alloy is observed by using a scanning electron microscope.

[0035] (5) Microstructure characterization shows that the β0 phase of the alloy after homogenization heat treatment at 1000°C has residual γ phase in the form of lath.

[0036] (6) Another sample is taken, the homogenization heat treatment temperature is increased by 50°C, that is, heat treatment at 1050°C, steps (3) to (5) are repeated, and the results show that the β o phase of the alloy has residual γ phase in the form of lath.

[0037] (7) Another sample is taken, the homogenization heat treatment temperature is increased by 50°C, that is, heat treatment at 1100°C, steps (3) to (5) are repeated, and the results show that the β o phase of the alloy has no residual phase.

[0038] (8) Several samples are taken for homogenization heat treatment at 1100°C, and thus the homogenization heat treatment is completed.

[0039] (9) The sample obtained in step (8) is subjected to subsequent isothermal annealing treatment. One sample is first subjected to 900°C isothermal heat treatment in a high-temperature resistance furnace, in order to minimize the temperature drop during quenching, the sample is placed in a self-designed semi-automatic isothermal quenching furnace, and one-key quenching can be realized after the set holding time and temperature are reached, the time interval from heating temperature to quenching of the sample is less than 1 s, which ensures a small temperature drop. The whole heat treatment process is that the sample is put into the high-temperature resistance furnace after the target set temperature is reached, and the heat treatment time is 360 h.

[0040] (10) The sample obtained in step (9) is subjected to general metallographic preparation (mainly including sample grinding, polishing and etching) to obtain a sample for observing the microstructure of the alloy, and the microstructure of the alloy is observed by using a scanning electron microscope.

[0041] (11) Microstructure characterization shows that the β o phase is residual in the alloy after 900°C isothermal heat treatment.

[0042] (12) Another sample is taken, the isothermal heat treatment temperature is decreased by 50°C, that is, heat treatment at 850°C, steps (9) to (11) are repeated, and the results show that the β o phase is residual in the alloy.

[0043] (13) Another sample was taken, and the isothermal heat treatment temperature was reduced by 50°C, i.e., heat treatment was performed at 800°C, and steps (9) to (11) were repeated. The results show that the β o phase disappears.

[0044] Through the above steps, the heat treatment parameters of homogenization heat treatment and isothermal annealing treatment were determined.

[0045] The TiAl alloy obtained through the above two steps of heat treatment was subjected to creep loading at 700°C under a condition of 300 MPa or 350 MPa, and the TiAl alloy without heat treatment was subjected to creep loading under the same condition.

[0046] The X-ray diffraction patterns of the samples before and after isothermal heat treatment are shown in Figure 1 Before annealing treatment, only the diffraction peaks of the α2 phase, the γ phase and the β o phase were found in the sample, and no diffraction peak of the Mn compound was found, i.e., the Mn element was completely solid-solved in the above-mentioned constituent phases. However, in the sample isothermally annealed for 480 h (A-480h), in addition to the corresponding diffraction peaks of the aforementioned three constituent phases (α2, γ, β o ), the diffraction peaks of the Laves phase (TiMn2) at 2θ = 42.48° and 43.19° additionally appeared, as shown in the enlarged area on the left side of the box in Figure 1 .

[0047] Figure 2 The scanning transmission electron microscopy high-angle annular dark-field image and the corresponding element face scanning graph of the β o phase region of the sample after isothermal annealing for 480 h are shown in FIG. 6. From the morphology, the ellipsoidal Laves grains (Mn element enrichment region) and the γ grains are adjacent to each other and embedded in the matrix region of the β o phase in a network, and the β o phase has completely transformed into the α2 phase.

[0048] The creep properties of the Ti-42Al-5Mn alloy before and after heat treatment were characterized under the conditions of 750°C and two constant loads of 300 MPa and 350 MPa. The main purpose of selecting the above temperature and stress combination is to minimize the influence of phase transformation during creep on the creep properties, and at the same time, to be as close as possible to the actual working condition. The creep test was carried out until macroscopic fracture of the sample occurred. The curves of the creep strain versus time under different stress conditions are shown in Figure 3The alloy transits from the deceleration creep stage to the acceleration creep stage in a short time without obvious steady creep stage under the stress state of 350 MPa. The alloy experiences the three typical stages of the creep curve completely: the deceleration creep stage, the steady creep stage and the acceleration creep stage under the stress state of 300 MPa. According to the change of the instantaneous creep rate with time, it is found that the heat-treated sample enters the steady creep stage (131 h) from the deceleration creep stage (about 13.5 h) under the low stress condition, and the steady creep rate is 7.63x10-6s-1. -8 s -1 .

[0049] The creep rupture life of the heat-treated sample is significantly prolonged under the stress state of 300 MPa or 350 MPa: the creep life is increased from 18.20 h of the untreated state to 28.57 h after heat treatment under the constant stress of 350 MPa; the creep life is increased from 116.48 h of the untreated state to 183.88 h after heat treatment under the constant stress of 300 MPa.

[0050] The creep life of the heat-treated sample under the constant stress of 300 MPa at 750℃ is 183.88 h. The transmission electron microscope observation of the fracture of the sample shows that the Laves phase interacts with the dislocation, and the Laves phase significantly inhibits the proliferation and sliding of the dislocation, as shown in FIG. 4. Figure 4

[0051] Although the embodiments of the present application have been shown and described, it is to be understood that for the purpose of the present application, the embodiments can be changed, modified, replaced and varied in many ways without departing from the principles and spirit of the present application, and the scope of the present application is defined by the appended claims and their equivalents.

[0052] The above detailed description of the application is a detailed description of the application, and cannot be considered as limiting the specific embodiments of the application to these descriptions. For those skilled in the art, without departing from the concept of the present application, a number of simple deductions and substitutions can be made, which should be considered as falling within the protection scope of the present application.​

Claims

1. A method for improving the high-temperature creep life of β-γ type TiAl alloys, characterized in that, Includes the following steps: Step 1: Vacuum induction smelting of Ti-42Al-5Mn alloy ingots with nominal composition; the aforementioned alloy ingots are forged and cooled to obtain forging blanks, which are then subjected to homogenization solution treatment; Step 2: Perform isothermal annealing on the sample homogenized and solution-treated in Step 1.

2. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 1, characterized in that, The homogenization solution heat treatment temperature is 1000℃~1400℃.

3. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 2, characterized in that, The method for determining the homogenization solution heat treatment temperature includes the following steps: first, heat-treating the alloy at an initial temperature, then characterizing the microstructure of the alloy, and when the β of the alloy... o If there are no lamellar precipitates in the phase, then this initial temperature is determined to be the homogenization solution heat treatment temperature; otherwise, if the alloy's β... o If lamellar precipitates are present, the homogenization heat treatment temperature of the alloy needs to be gradually increased according to a gradient temperature until the β of the alloy after homogenization heat treatment is reached. o There were no lamellar precipitates in the phase, thus the homogenization solution treatment temperature was determined.

4. The method for improving the high-temperature creep life of TiAl alloy according to claim 3, characterized in that, The gradient temperature is 25℃-50℃.

5. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 1, characterized in that, The homogenization solution treatment time is 0.5 h to 10 h.

6. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 5, characterized in that, The homogenization solution treatment time is 0.5 h to 5 h.

7. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 6, characterized in that, The homogenization solution treatment time is 1 h to 2 h.

8. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 1, characterized in that, The isothermal annealing temperature is 700℃~900℃.

9. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 8, characterized in that, Determining the isothermal annealing temperature involves the following steps: heat-treating the homogenized solution-treated sample at the initial temperature, followed by characterizing the alloy's microstructure. If the initial β... o If the phase completely transforms into the precipitated phase, the initial temperature is the isothermal annealing temperature; conversely, if the alloy still contains β... o If the β phase exists, the isothermal annealing temperature of the alloy needs to be gradually decreased according to the gradient temperature until the β phase in the alloy is reduced after isothermal annealing heat treatment. o The phase completely disappears, thus determining the isothermal annealing temperature.

10. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 9, characterized in that, The gradient temperature is 50°C.

11. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 1, characterized in that, The isothermal annealing heat treatment time is 9 h to 480 h.

12. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 11, characterized in that, The isothermal annealing heat treatment time is 24 h to 480 h.

13. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 12, characterized in that, The isothermal annealing heat treatment time is 360 h ~ 480 h.

14. The method for improving the high-temperature creep life of β-γ type TiAl alloy according to claim 1, characterized in that, The cooling method for the homogenization solution treatment and isothermal annealing treatment is water quenching.

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