Method for improving high-temperature toughness of GH4099 alloy

Abstract

The invention relates to the technical field of nickel-based high-temperature alloy treatment, in particular to a method for improving the high-temperature toughness of a GH4099 alloy. According to the method, the GH4099 alloy in the solid solution state is subjected to ultrasonic rolling to form a rolling deformation layer, and then aging treatment is carried out; the aging treatment temperature ranges from 780 DEG C to 820 DEG C, a double-heterostructure (a gradient heterostructure from the surface layer to the core part and a heterostructure of particle precipitation and no particle precipitation isometric crystals in the surface layer) can be constructed in the GH4099 alloy, meanwhile, the high-temperature strength and toughness of the GH4099 alloy are improved, and the strength-plasticity inversion relation of a traditional gradient material is overcome.

Description

Technical Field

[0001] The present invention relates to the technical field of nickel-based superalloy treatment, and particularly to a method for improving the high-temperature strength and toughness of GH4099 alloy. Background Art

[0002] Nickel-based superalloys have excellent mechanical properties and thermal stability at high temperatures. These advantages make them one of the irreplaceable materials in extreme environments and occupy a very important position. Taking domestic GH4099 superalloy as an example, the main strengthening mechanisms of the alloy are solid-solution strengthening brought by Cr, W, Mo and Co elements, and second-phase precipitation strengthening brought by forming γ′ phase (Ni3(Al,Ti), FCC) in the austenite (A) matrix. GH4099 alloy has a wide range of applications, covering various aviation components, including engine combustion chambers and structural reinforcements, and is also commonly used in the nuclear power industry. With the deterioration of the service environment, the requirements for alloy performance are more stringent. In alloy design, improving mechanical properties through plastic deformation has always been a concern. Metallic materials that have undergone plastic deformation have excellent mechanical properties, especially strength. However, increasing strength through plastic deformation is often accompanied by an inevitable sharp decrease in plasticity. Strength and plasticity are the two most important mechanical properties of metallic materials. How to simultaneously improve the high-temperature strength and toughness of GH4099 alloy is still an important problem to be solved. Summary of the Invention

[0003] In view of this, the purpose of the present invention is to provide a method for improving the high-temperature strength and toughness of GH4099 alloy, which can simultaneously improve the high-temperature strength and toughness of GH4099 alloy.

[0004] In order to achieve the above-mentioned invention purpose, the present invention provides the following technical solutions:

[0005] The present invention provides a method for improving the high-temperature strength and toughness of GH4099 alloy, including the following steps: subjecting the solution-treated GH4099 alloy to ultrasonic rolling to form a rolled deformation layer and then performing aging treatment; the temperature of the aging treatment is 780 - 820 °C.

[0006] Preferably, the temperature of the aging treatment is 800 °C.

[0007] Preferably, the time of the aging treatment is 4 h.

[0008] Preferably, the cooling method of the aging treatment is air cooling.

[0009] Preferably, the aging treatment is carried out under inert gas protection or in a vacuum condition.

[0010] Preferably, the conditions for ultrasonic rolling include: the lathe speed is 100 - 200 r / min, the shot peening air pressure is 0.3 - 0.7 MPa, the feed rate F z of the rolling head perpendicular to the surface of the component = 0.05 - 0.2 mm / min, and the feed rate F x of the rolling head parallel to the surface of the component = 0.8 - 2 mm / min, and the ultrasonic vibration frequency is 27 - 42 kHz.

[0011] Preferably, the ultrasonic rolling is multi - pass ultrasonic rolling. After the ultrasonic rolling, the actual total reduction of the GH4099 alloy does not exceed 0.1 mm.

[0012] Preferably, the ultrasonic rolling is divided into 8 passes. The downward pressure of the ultrasonic rolling equipment is set to 0.8 mm and rolled 4 passes first, then the downward pressure is adjusted to 1 mm and rolled 3 passes continuously, and finally the downward pressure is adjusted to 1.2 mm and rolled 1 pass.

[0013] Preferably, the thickness of the rolled deformation layer is 100 - 200 μm.

[0014] Preferably, in terms of mass percentage, the elemental composition of the GH4099 alloy is: Cr 19.5%, W 6.8%, Mo 4.2%, Al 2.2%, Co 6.8%, Ti 1.3%, C 0.04%, and the balance is Ni.

[0015] The present invention provides a method for improving the high - temperature strength and toughness of GH4099 alloy, which includes the following steps: subjecting the solution - treated GH4099 alloy to ultrasonic rolling to form a rolled deformation layer and then performing aging treatment; the temperature of the aging treatment is 780 - 820 °C. The present invention first performs ultrasonic rolling on the solution - treated GH4099 alloy, introducing a rolled deformation layer on the surface of the GH4099 alloy, forming a heterogeneous structure of surface and core grains. This structure can effectively achieve back - stress strengthening in the GH4099 alloy, significantly improving the strength at room temperature and high temperature, but the plasticity shows a downward trend, and the strength - plasticity inversion relationship has not been solved. Then, the present invention performs aging treatment at 780 - 820 °C. By controlling the aging temperature, a γˊ - rich deformed grain region and a γˊ - poor recrystallized equiaxed grain region are formed in the deformation layer after aging treatment, and the two form a heterogeneous structure in the deformation layer. Under 700 °C high - temperature tension, the dual heterogeneous structure inside the GH4099 alloy increases the strain gradient during the deformation process. In order to adapt to this strain gradient, a high - density of geometrically necessary dislocations (GND) accumulates near the soft - hard interface, triggering the heterogeneous deformation - induced (HDI) effect at high temperature. Finally, when the GH4099 alloy is under 700 °C high - temperature tension, it shows a significant strength - plasticity synergistic effect.

[0016] The results of the examples show that after ultrasonic rolling and aging treatment at 800 °C, the yield strength of the U800 sample is 816 Mpa, which is 180 MPa higher than that of the S800 sample treated only by aging. The total elongation rate is 26.5%, an increase of 3.1%, overcoming the strength-ductility inversion relationship of traditional gradient materials. Description of the Drawings

[0017] Figure 1 It is a schematic diagram of the size of the tensile specimen;

[0018] Figure 2 It is an optical micrograph of the surface layer of the samples after solution treatment and aging at different temperatures. Among them, (a) solution-treated GH4099, (b) U750, (c) U800, (d) U850;

[0019] Figure 3 It is the precipitation of γ' phase in the core of the rolled specimens at different aging temperatures. Among them, (a) U750, (b) U800, (c) U850, (d) statistics of the γ' phase size and content in the core of different specimens;

[0020] Figure 4 It is a micrograph of the rolled layer at a distance of 45 μm from the surface at different aging temperatures. Among them, (a) U750, (b) U800, (c) U850;

[0021] Figure 5 It is the microstructure (a) of the U750 specimen at a distance of 45 μm from the surface layer and the distribution maps of Al and Ti elements (b, c);

[0022] Figure 6 It is the microstructure (a) of the U800 specimen at a distance of 45 μm from the surface layer and the distribution maps of Al and Ti elements (b, c);

[0023] Figure 7 It is the microstructure (a) of the U850 specimen at a distance of 45 μm from the surface layer and the distribution maps of Al and Ti elements (b, c);

[0024] Figure 8 It is the microhardness distribution map of the gradient layer in the true stress-strain curve and work hardening rate curve of the U750, U800, and U850 specimens;

[0025] Figure 9 It is the uniaxial tensile property test results of the S750, S800, S850, U750, U800, and U850 specimens at 700 °C. Among them, (a) S750 and U750, (b) S800 and U800, (c) S850 and U850;

[0026] Figure 10For the strain hardening rate curves of S750, S800, S850, U750, U800, and U850;

[0027] Figure 11 Schematic diagram of the microstructure evolution of rolled GH4099 at different aging temperatures;

[0028] Figure 12 TEM image at a distance of 45 μm from the surface layer after tensile deformation of the U750 specimen. Among them, (a) is the low magnification image, and (b) is the high magnification image;

[0029] Figure 13 EBSD maps at different positions from the tensile fracture of U750. Among them, (a, a1) is the far fracture of U750, (b, b1) is the near fracture of U750, (a, b) is the inverse pole figure, and (a1, b1) is the KAM map;

[0030] Figure 14 TEM image at a distance of 45 μm from the surface layer after tensile deformation of the U800 specimen. Among them, (a) and (b) are diagrams showing equiaxed recrystallized grains, and (c) is a schematic diagram of the stress source between the recrystallized grains and the particle precipitation zone;

[0031] Figure 15 EBSD maps at different positions from the tensile fracture of U800. Among them, (a, a1) is the far fracture of U800, (b, b1) is the near fracture of U800, (a, b) is the inverse pole figure, and (a1, b1) is the KAM map;

[0032] Figure 16 Fracture analysis of GH4099 under aging treatment at different temperatures. Among them, (a1 - a3) is U750, (b1 - b3) is U800, and (c1 - c3) is U850. Specific implementation method

[0033] The present invention provides a method for improving the high - temperature strength and toughness of GH4099 alloy, including the following steps: subjecting the solution - treated GH4099 alloy to ultrasonic rolling to form a rolled deformation layer and then performing aging treatment; the temperature of the aging treatment is 780 - 820 °C.

[0034] The present invention has no special requirements for the composition of the GH4099 alloy, and any well - known GH4099 alloy in the art can be used. In the embodiments of the present invention, by mass percentage, the elemental composition of the GH4099 alloy is: Cr 19.5%, W 6.8%, Mo 4.2%, Al 2.2%, Co 6.8%, Ti 1.3%, C 0.04%, and the balance is Ni.

[0035] The present invention has no special requirements for the source of the solution-treated GH4099 alloy. Any solution-treated GH4099 alloy with a well-known source in the art can be used. In the embodiments of the present invention, specifically, it is obtained by solution-treating commercially forged GH4099 alloy purchased. The temperature of the solution treatment is 1100 °C, the time is 1 h, and the cooling method is air cooling.

[0036] In the present invention, the conditions of the ultrasonic rolling preferably include: the lathe speed is 100 - 200 r / min, the shot peening air pressure is 0.3 - 0.7 MPa, the feed rate F of the rolling head perpendicular to the surface of the component z = 0.05 - 0.2 mm / min, and the feed rate F of the rolling head parallel to the surface of the component x = 0.8 - 2 mm / min, and the ultrasonic vibration frequency is 27 - 42 kHz; in specific embodiments, the lathe speed can be 100 r / min, 150 r / min or 200 r / min, the shot peening air pressure can be 0.3 MPa, 0.4 MPa, 0.5 MPa, 0.6 MPa or 0.7 MPa, and F z can be 0.05 mm / min, 0.1 mm / min, 0.15 mm / min or 0.2 mm / min, and F x can be 0.8 mm / min, 1.0 mm / min, 1.5 mm / min or 2.0 mm / min, and the ultrasonic vibration frequency can be 27 kHz, 30 kHz, 35 kHz, 38 kHz or 42 kHz. By controlling the conditions of the ultrasonic rolling within the above range, the present invention is beneficial to the formation of a rolled deformation layer.

[0037] In the present invention, the ultrasonic rolling is multi-pass ultrasonic rolling. After the ultrasonic rolling, the actual total reduction of the GH4099 alloy preferably does not exceed 0.1 mm.

[0038] In specific embodiments, the ultrasonic rolling is divided into 8 passes. The set reduction of the ultrasonic rolling equipment is 0.8 mm and it is rolled 4 passes first, then the reduction is adjusted to 1 mm and it is rolled 3 passes continuously, and finally the reduction is adjusted to 1.2 mm and it is rolled 1 pass. In the present invention, since the GH4099 alloy will produce springback after pressing down, the set reduction of the ultrasonic rolling equipment is significantly less than the actual reduction. After the above 8 passes of ultrasonic rolling, the actual total reduction of the GH4099 alloy does not exceed 0.1 mm.

[0039] In the present invention, the ultrasonic rolling is carried out at room temperature.

[0040] After ultrasonic rolling, a rolled deformation layer is introduced on the surface of the GH4099 alloy of the present invention. The thickness of the rolled deformation layer is preferably 100-200 μm. The rolled deformation layer has high residual stress and high hardness, and the grains are significantly refined compared with the core matrix, forming a heterogeneous structure of surface and core grains. This structure can effectively achieve back stress strengthening in the GH4099 alloy, significantly improving the strength at room temperature and high temperature, but the plasticity shows a downward trend, and the strong-plastic inversion relationship has not been resolved.

[0041] After completing ultrasonic rolling, the present invention subjects the ultrasonically rolled GH4099 alloy to aging treatment.

[0042] In the present invention, the temperature of the aging treatment is 780-820 °C, and in specific embodiments, it can be 780 °C, 790 °C, 800 °C, 810 °C or 820 °C, and most preferably 800 °C.

[0043] In the present invention, the time of the aging treatment is preferably 4 h; the cooling method of the aging treatment is preferably air cooling. The present invention preferably performs the aging treatment under inert gas protection or vacuum conditions. The present invention has no special requirements for the type of the inert gas, and any well-known inert gas in the art can be used, such as nitrogen and argon.

[0044] Through the aging treatment of the present invention, partial recrystallization occurs in the rolled deformation layer. A considerable part of the grains undergo recrystallization, forming equiaxed grains with a low dislocation density, and a relatively high density of dislocations still exists locally. At the same time, there is also an uneven distribution of γ' phase. There is almost no γ' phase in the low dislocation density region, and the γ' phase is mainly distributed in the high dislocation density region. Thus, regions of γ' phase enrichment and depletion are formed in the rolled layer, and finally, a deformed crystal region rich in γ' phase and a recrystallized equiaxed grain region poor in γ' phase are formed in the deformed layer.

[0045] By introducing a rolled deformation layer, the present invention forms a heterogeneous structure of surface and core grains. Through aging treatment at a specific temperature, a heterogeneous structure is formed in the deformed layer, specifically, a deformed crystal region rich in γ' phase and a recrystallized equiaxed grain region poor in γ' phase are formed in the deformed layer. Under 700 °C high-temperature tension, the dual heterogeneous structure inside the GH4099 alloy increases the strain gradient during the deformation process. To adapt to this strain gradient, a high density of GND accumulates near the soft-hard interface, inducing the HDI effect at high temperature. Finally, when the GH4099 alloy is under 700 °C high-temperature tension, it exhibits a significant strong-plastic synergy effect.

[0046] The following combines examples to detail a method for improving the high-temperature strength and toughness of the GH4099 alloy provided by the present invention, but they cannot be understood as limiting the protection scope of the present invention.

[0047] Examples 1 - 3 and Comparative Examples 1 - 3

[0048] The material used in this experiment is commercially forged GH4099 superalloy, and its typical elemental composition is shown in Table 1. Before surface ultrasonic rolling process (USRP), round bars with a size of Φ10mm×97mm were solution-treated at 1100°C for 1h and then air-cooled to dissolve all the precipitates and impurities into the matrix to obtain a uniform supersaturated solid solution. Then, the solution-treated bar samples were subjected to surface ultrasonic rolling process at room temperature. The ultrasonic rolling parameters used are as follows: spindle speed V1 = 100 r / min, atmospheric pressure P = 0.3 MPa, feed rate F z = 0.1 mm / min, F x = 1 mm / min, initial reduction X1 = 0.8 mm, and under this reduction, it was rolled for 4 passes. Subsequently, the reduction was adjusted to X2 = 1 mm and rolled for 3 more passes. Finally, the reduction was adjusted to X3 = 1.2 mm and rolled for 1 pass, for a total of 8 passes. Through repeated plastic deformation, the surface grains were refined, so that a grain-refined layer with high residual stress and high hardness could be formed on the surface of GH4099. Then, the samples after USRP were respectively solution-treated (750°C / 4h, 800°C / 4h, 850°C / 4h, air-cooled) under nitrogen protection, and were named U750, U800, and U850 respectively. For comparison, the samples that were not rolled but underwent the same solution heat treatment were marked as standard samples and named S750, S800, and S850 respectively. The solution treatment temperature and duration settings for each sample are shown in Table 2.

[0049] Table 1 Chemical composition of GH4099 (wt%)

[0050] Element Cr W Mo Al Co Ti C Ni Wt.% 19.5 6.8 4.2 2.2 6.8 1.3 0.04 Remainder

[0051] Table 2 Treatment conditions for each sample

[0052]

[0053] Structure characterization and property testing:

[0054] Mechanical property testing method

[0055] The Vickers microhardness of the ultrasonic rolled specimens was measured using an HVS-1000B Vickers hardness tester. A load of 100 g was used, and the average value from 5 parallel measurements was taken as the final hardness value. A uniaxial tensile test was carried out at high temperature using an MTS Landmark test system, and the dimensions of the tensile specimens are as Figure 1 shown, with the unit of mm. The tensile parameters are as follows: tensile temperature 700°C, holding time 15 min, constant strain rate 1 s -1To ensure the accuracy of the results, three reproducible tensile tests were performed on the specimens for each treatment condition.

[0056] Microstructure Characterization Methods

[0057] The microstructure inside the plastic deformation layer on the surface of GH4099 nickel-based alloy was studied. First, the processed rod-shaped sample was cut into blocks of about 3 mm along the cross-sectional direction using an electric spark wire cutting machine. Then, the microstructure of the surface of the ultrasonic rolling sample was characterized by an optical microscope (OM, SDPTOP ICX4IM). The OM sample was mechanically polished and then corroded in a high-temperature alloy corrosion solution prepared by 0.5g CuCl2+100mL CH3CH2OH+100mL HCl for 5 minutes and 30 seconds. Backscattered electron diffraction (EBSD, EDAX Hikari Plus) was used to analyze the misorientation angle and geometric necessary dislocation (GND) density after tensile testing. The specific experimental parameters are as follows: electrolytic polishing, acceleration voltage 20kV, step size 1200nm. Transmission electron microscopy (TEM, Philips CM12) was used to characterize the microstructure before and after the tensile test, and the shooting points were all 45μm away from the surface of the sample (inside the rolling deformation layer). The TEM sample was first ground to a thickness of 50 μm, and then the disc was ground into a thin foil using an ion polisher. The TEM images taken were analyzed using Image-Pro Plus software to count the size and content of the γ′ phase. In order to eliminate experimental errors, three images were selected for each sample for measurement. A scanning electron microscope (SEM, ZEISS SUPRA40) was used to characterize the fracture morphology after the tensile test, and the strengthening mechanism brought by the dual heterogeneous structure inside the rolling deformation layer was analyzed in combination with the fracture form.

[0058] Test Results

[0059] Evolution of microstructure

[0060] The microstructure of GH4099 after solution treatment (1100℃ for 1h) is as follows: Figure 2 As shown in (a), the average grain size of the matrix is ​​about 58 μm, and there are a small amount of twins inside the grains. Figure 2 (b, c, d) show the grain structure of the USRP cold-deformed specimens after aging at different temperatures. During the ultrasonic rolling process, the rapidly rotating metal rolling head transmits compressive stress and ultrasonic vibration to the parallel segment surface of the rotating rod-shaped tensile specimen, which produces an impact and extrusion effect on it, causing the GH4099 material to undergo a large-scale plastic deformation and form a rolling deformation layer (DL) with a thickness of about 150μm. In addition, due to the combined effect of ultrasonic impact vibration and shear stress, the near-surface grains are deflected and elongated along the rotational shear direction of the specimen, showing plastic flow characteristics.

[0061] The size, morphology and distribution of the strengthening phase γ' phase in the equiaxed grain region of the core after different aging treatments are as follows Figure 3 shown. Among them, (a) U750, (b) U800, (c) U850, (d) statistics of the size and content of the γ' phase in the core of different specimens. After aging at different temperatures, there are obvious differences in the content and size of the γ' phase in the core of the specimens. Statistical analysis of the content and size of the γ' phase in different specimens shows the results in Figure 3 (d) in. As can be seen from Figure 3 , with the increase of the aging treatment temperature, the size of the γ' phase in the specimen coarsens and the quantity shows a downward trend.

[0062] The microstructure in the rolling layer under different aging temperatures has significant changes. The microstructure of the rolling layer 45 μm away from the surface is as follows Figure 4 shown. Among them, (a) U750, (b) U800, (c) U850. In the U750 specimen, there are still obvious directional deformed grains in the rolling deformation layer, and there are high-density dislocation pile-ups in the grains (see Figure 4 (a) in); after aging treatment at 800 °C, partial recrystallization occurs in the rolling deformation layer, and a considerable part of the grains undergo recrystallization to form equiaxed grains with a low dislocation density, and there are still high-density dislocations locally (see Figure 4 (b) in). After aging at 850 °C, the recrystallization of the rolling layer is relatively complete, and the rolling layer is composed of relatively uniform equiaxed grains, the dislocation density decreases significantly, and twins exist in local areas (see Figure 4 (c) in).

[0063] Perform TEM element scanning on the rolling layer of specimens at different aging temperatures, and observe the distribution state of the γ' phase in the rolling layer through the distribution of the main forming elements of the γ' phase. Figure 5 It is the microstructure (a) of the U750 specimen at 45 μm from the surface layer and the distribution maps of Al and Ti elements (b, c). Its grains still maintain the directionality caused by rolling, and the distribution of the γ' phase is relatively fine and uniform (see Figure 5 (a) in), and (b, c) are the distribution maps of the main forming elements Al and Ti of the γ' phase in this area, and it can be observed that the γ' phase is evenly and finely distributed in the rolling layer with grain distortion.

[0064] There are equiaxed recrystallized grains in the rolling layer of the U800 specimen, and at the same time, there is also an uneven distribution of the γ' phase. Figure 6In (a), there are a recrystallized region with a low dislocation density and an unrecrystallized region with a high dislocation density. In (b) and (c), the composition distributions of Al and Ti elements are respectively shown to reflect the distribution of γ' phase. There is almost no γ' phase in the region with a low dislocation density, and the γ' phase is mainly distributed in the region with a high dislocation density. Thus, regions with γ' phase enrichment and depletion are formed in the rolled layer. Forming such a structure on the surface layer is crucial for improving the strength and toughness. The particles are the precipitated strengthening phase γ' phase, and this region is the strengthening region with a high hardness. The region enveloped by the red dotted line is an equiaxed recrystallized region without precipitation, where there are no precipitated γ' phase particles and it has a low hardness.

[0065] When the aging temperature is increased to 850 °C, the rolled layer recrystallizes sufficiently and the dislocation density decreases significantly (see (a) in Figure 7 ), as shown by the size and distribution of the γ' phase reflected by the distribution of Al and Ti elements in (b, c) in Figure 7 . The γ' phase precipitated by strain induction in the deformed layer has significantly aggregated and grown and has the characteristic of grain boundary distribution, and there is no γ' phase particle precipitation region.

[0066] After the rolled specimens are treated at different aging temperatures, there are significant differences in the microstructural characteristics. The surface hardness distribution of the specimens is as shown in Figure 8 . Among them, the average hardness of the surface and the core of the U750 specimen is the highest, which is related to its fine γ' phase and the residual high dislocation density of the deformed grains. The U800 specimen has the microstructural characteristics of nano-scale equiaxed recrystallized grains in the rolled layer. The size of the γ' phase in its rolled layer is larger than that of the U750 specimen, so its surface microhardness is reduced by about 20 Hv and its core microhardness is reduced by about 15 Hv. Due to the more sufficient recrystallization in the surface rolled layer of the U850 specimen and the further coarsening of the γ' phase at a higher aging temperature, its surface microhardness is reduced by about 100 Hv and its core microhardness is reduced by about 50 Hv.

[0067] The results of the uniaxial tensile property tests of the S750, S800, S850, U750, U800, and U850 specimens at 700 °C are shown in Figure 9 , and the specific data are shown in Table 3.

[0068] Table 3 Mechanical properties of U750, U800, and U850 specimens at 700 °C

[0069] S750 U750 S800 U800 S850 U850 Yield strength MPa 656 825 636 800 606 663 Tensile strength MPa 1003 977 952 1003 959 946 Elongation % 25.5 13.4 23.4 26.5 37.8 39.8

[0070] As can be seen from Table 3, compared with the corresponding standard specimens, the mechanical properties of the specimens after rolling have changed significantly. Among them, the yield strength of the 750 aging specimen U750 has increased by about 170 MPa relative to the S750 specimen, but the tensile plasticity has decreased significantly, by about 12%; compared with S800, the rolled layer of the U800 specimen has a significantly uneven equiaxed grain structure, showing a simultaneous increase in strength and plasticity, with the yield strength increasing by about 170 MPa and the elongation increasing by 3.1%; the increase in the yield strength and plasticity of the U850 specimen at high temperature relative to the S850 specimen is relatively small, with the yield strength increasing by about 57 MPa and the elongation increasing by 2%.

[0071] USRP treatment can effectively promote the formation of differences in microstructure and properties from the surface layer to the core of the specimen, thereby effectively improving the strength index at 700 °C. Among them, the rolled layer of the specimen aged at 800 °C has a relatively special heterogeneous structure distribution characteristic, bringing the performance characteristics of simultaneous improvement of strength and plasticity.

[0072] The strain-hardening rate (θ) of the GH4099 alloy at different aging temperatures is as Figure 10 shown. In the early stage of plastic deformation, the strain-hardening rate of the U800 specimen is greater than that of other specimens. In addition, the U800 specimen has a two-stage strain-hardening behavior. First, the strain-hardening rate θ drops sharply, which is characteristic of the elastic-plastic transition stage. Then, θ experiences a continuous fluctuation stage of sharp increase and decrease, which represents the discontinuous yield of the U800 specimen. This is because there are few mobile dislocations at the beginning of plastic deformation, and a greater stress needs to be applied to promote the rapid sliding of dislocations to adapt to the applied constant strain rate. No θ rebound phenomenon similar to that of the U800 specimen is observed in other specimens, and the θ of other specimens shows a monotonically decreasing trend, indicating that only the U800 specimen with a dual heterogeneous structure will produce this unique behavior, and the back stress strengthening effect of the U800 specimen can be observed from the strain-hardening rate curve.

[0073] Gradient structure tissue characteristics:

[0074] During the ultrasonic rolling process, the grains in the rolled deformation layer of GH4099 are deformed, and the dislocation density on the surface layer is significantly increased. Subsequently, aging treatments are carried out at 750 °C, 800 °C, and 850 °C for 4 h respectively. In the U750 specimen, the deformed layer is composed of deformed grains, and there are high-density dislocations and γˊ phases distributed in the deformed grains. The size of the γˊ phase is 15 nm. In the U800 specimen, partial recrystallization occurs inside the deformed layer, forming grains with high-density dislocations and recrystallized equiaxed grains with a small amount of twins. During the recrystallization process of nickel-based alloys, the coordinated deformation between grains can promote the decomposition of some residual second phases rich in Al and Ti in the matrix. Under the condition of meeting the recrystallization temperature, as the recrystallized grain boundaries migrate and displace Al and Ti elements, when the migration rate of the recrystallized grain boundaries is high enough, the concentration in the non-recrystallized region (solute-rich region) in front of the grain boundaries is sufficient to cause constitutional supercooling, and γˊ phases precipitate in the non-recrystallized region at a suitable aging temperature. Therefore, 750 °C does not meet the recrystallization condition, and γˊ phases precipitate uniformly in the rolled layer with a high dislocation density. Aging at 800 °C meets the recrystallization temperature condition. The local recrystallized region cannot meet the precipitation of γˊ phases due to the depletion of solute atoms Al and Ti, so a γˊ phase depletion zone is formed, and γˊ phases with a size of 32 nm precipitate in the non-recrystallized region. The deformed layer inside shows a heterogeneous structure composed of a γˊ phase-rich deformed grain region and a γˊ phase-poor equiaxed recrystallized region, and the γˊ phases formed due to the enrichment of Al and Ti on the recrystallized grain boundaries are relatively coarse. At 850 °C aging, sufficient recrystallization occurs. The migration of the recrystallized grain boundaries promotes the enrichment of Al and Ti in front of them, and a large number of massive γˊ phases precipitate at the grain boundaries. The microstructure evolution of the rolled GH4099 at different aging temperatures is as Figure 11 shown.

[0075] Figure 12 Figure (a) and (b) are TEM characterizations of the U750 specimen after tensile deformation, where (a) is at low magnification and (b) is at high magnification. The rolled deformation layer of the U750 specimen is composed of high-density dislocations and fine γˊ phases. The dislocations generated by deformation are difficult to move, and the high-density dislocations accumulate inside the slender nanocrystals, which will cause premature stress concentration on the surface layer of the U750 specimen. As Figure 12 indicated by the arrows in (a) and (b) of

[0076] The strain distribution of the specimens at different positions from the fracture is characterized by EBSD, and the results are shown in Figure 13 , where (a, a1) is the far fracture of U750, (b, b1) is the near fracture of U750, (a, b) is the inverse pole figure, and (a1, b1) is the KAM map. Since a high-strain region will be formed on the surface before and after tension, the surface layer is red in the KAM map. This region is not favorable for coordinated strain, which will lead to poor plasticity, and the depth of the green region is relatively shallow during the tension process, indicating that the strain is concentrated on the surface layer. As Figure 13As shown in (a1), in the low-strain region of the U750 specimen, dislocations are mainly concentrated in the rolled surface layer. In the high-strain region near the fracture surface, the overall dislocation density in the rolled deformation layer increases, but the dislocation density in its core remains unchanged significantly. See Figure 13 in (b2), there is a large macroscopic deformation region on the surface.

[0077] Therefore, during high-temperature tensile testing, deformation is always concentrated in the surface layer. The low change in dislocation density in the core indicates that the core of the specimen undertakes less plastic deformation during the deformation process. A hard shell layer is formed on the surface of the U750 specimen, and the surface hard shell layer undertakes the main tensile load during the tensile deformation process. When the dislocations that the deformation layer can accommodate reach the limit, it is extremely easy to cause premature fracture of the U750 specimen surface due to stress concentration.

[0078] In the rolled deformation layer of the U800 specimen, equiaxed recrystallized grains with depleted γˊ phase and non-recrystallized regions with high-density γˊ phase are formed, thus having regions with different properties in the deformation layer. During high-temperature tensile testing, in addition to the HDI strengthening effect formed by the heterogeneous structure between the rolled layer and the core, there are also performance heterogeneous regions with uneven γˊ phase distribution within the rolled layer. When entering the elastic-plastic deformation stage of tension, the soft regions in the deformation layer undergo plastic deformation first, while the hard regions, due to their higher resistance to plastic deformation, remain in the elastic deformation stage. This deformation difference leads to the formation of inhomogeneous deformation and strain gradient of the material, and the flow stress shows a significant change between different soft and hard micro-regions, highlighting the HDI strengthening effect and work-hardening effect. To adapt to the strain gradient during the deformation process, the GND emitted from the Frank-Read source accumulates at the interface of the soft regions. As Figure 14 shown in (c), let the shear force acting on the slip system be τ a , and the number of GND be n, then the total stress generated by the GND accumulation at this boundary is nτ a . The accumulated GND causes the slip plane to bend and applies a long-range internal stress opposite to τ a called back stress. The back stress will prevent the Frank-Read source from emitting more GND, making the dislocations in the recrystallized soft regions not easy to move and accumulate at the interface, thus significantly enhancing the strength of the soft regions. Due to the interaction of forces, there will be a force at the interface of the hard regions with the same magnitude and opposite direction to the back stress called normal stress. During the deformation process, the soft grains yield first under low stress. Only when nτ a increases to its critical deformation stress, the hard grains will undergo deformation. Subsequently, the GND begins to slip into the hard regions, resulting in overall yielding. The combined effect of the back stress and the normal stress is defined as HDI strengthening. Therefore, the heterogeneous structure within the rolled deformation layer has a significant HDI strengthening effect on the mechanical properties of U800, which makes the yield strength and plasticity of the U800 specimen significantly improved compared with S800.Figure 14 In (a, b), significant accumulation of GND at the interface between hard and soft grains in the U800 specimen is shown, and this observation microscopically confirms the typical characteristics of HDI strengthening in the U800 specimen.

[0079] The EBSD characterization results of the U800 specimen are shown in Figure 15 , where (a, a1) is the far fracture surface of U800, (b, b1) is the near fracture surface of U800, (a, b) is the inverse pole figure, and (a1, b1) is the KAM map. As Figure 15 shown in (a1), compared with the U750 specimen ( Figure 13 ), the difference in dislocation density between the surface deformation layer and the core matrix in the low-strain region is relatively reduced. In the region with a higher strain, the strain in the surface deformation layer of the U800 specimen does not increase significantly, while the dislocation density in its core increases significantly. During the tensile deformation process of the U800 specimen with dual heterogeneous tissue characteristics, the recrystallized grains on the specimen surface can accommodate more dislocation movements, effectively coordinating local strain. At the same time, the core also undertakes plastic deformation, relieving the stress concentration during the tensile deformation process, which is beneficial to enhancing the work-hardening ability of the U800 specimen.

[0080] The fracture surface characterization of the specimen is shown in Figure 16 , where (a1 - a3) is U750, (b1 - b3) is U800, (c1 - c3) is U850, and the orange dashed line represents the boundary between the sample rolling deformation layer and the sample substrate. There is an interface between the surface layer and the core of the U750 specimen (see Figure 16 in (a1, a2)), and the surface layer has a high proportion of cleavage fracture characteristics (see Figure 16 in (a3)). This is because the high residual dislocation density and high-density γˊ phase in the surface layer of the U750 sample limit the movement of dislocations during the tensile deformation process, resulting in a low work-hardening rate. The high GND density leads to high intergranular stress concentration. The fracture surface morphology of the U800 specimen is mainly ductile fracture, and the transition from the surface layer to the core is relatively gentle (see Figure 16 in (b1 - b2)), and its surface fracture has a high proportion of dimples and a small number of cleavage steps (see Figure 16 in (b3)). The recrystallized coarse grains in the rolling layer relieve the intergranular strain concentration to a certain extent, and the U800 shows different fracture surface characteristics from those of the U750. The tensile curve of the U800 specimen shows excellent strength-ductility synergy. In the U850 specimen, the fracture surface shows a large number of deep dimples (see Figure 16 in (c2 - c3)), and no brittle fracture characteristics such as cleavage fracture or intergranular fracture are observed. The fracture mode changes to complete ductile fracture, so it has good plasticity. The change in fracture characteristics is caused by the γˊ phase with a high-density heterogeneous distribution and heterogeneous grain sizes.

[0081] Based on the above results, after USRP, a rolling deformation layer with a thickness of about 150 μm was introduced on the surface of the specimens, forming a heterogeneous structure of surface and core grains. After aging at 750 °C, the γˊ phase was uniformly distributed within the undeformed grains. After aging treatment at 800 °C, a deformed grain region rich in γˊ phase and a recrystallized equiaxed grain region poor in γˊ phase were formed within the deformation layer, and the two formed a heterogeneous structure within the deformation layer of the U800 specimen. When the aging temperature was increased to 850 °C, the heterogeneous structure within the deformation layer transformed into a uniformly recrystallized coarse grain structure.

[0082] Under high-temperature tension at 700 °C, the dual heterogeneous structure within the U800 specimen increased the strain gradient during the deformation process. To adapt to this strain gradient, a high density of GND was accumulated near the soft-hard interface, indicating that the heterogeneous deformation of soft and hard micro-regions induced the HDI effect at high temperatures. Eventually, when under high-temperature tension at 700 °C, the U800 exhibited a significant strength-ductility synergy effect.

[0083] The above are only the preferred embodiments of the present invention. It should be noted that for those of ordinary skill in the art, without departing from the principle of the present invention, several improvements and refinements can be made, and these improvements and refinements should also be regarded as the protection scope of the present invention.

Claims

1. A method for improving the high temperature toughness of GH4099 alloy, characterized in that: The following steps are involved: The solid solution GH4099 alloy is subjected to ultrasonic rolling to form a rolling deformation layer and then subjected to aging treatment; the temperature of the aging treatment is 780-820°C.

2. The method according to claim 1, characterized in that The temperature of the aging treatment is 800°C.

3. The method according to claim 1 or 2, characterized in that: The aging treatment time is 4 hours.

4. The method according to claim 1 or 2, characterized in that: The cooling method of the aging treatment is air cooling.

5. The method according to claim 1 or 2, characterized in that: The aging treatment is carried out under inert gas protection or vacuum conditions.

6. The method according to claim 1, characterized in that The conditions of ultrasonic rolling include: the lathe speed is 100-200 r / min, the shot peening pressure is 0.3-0.7 MPa, and the rolling head feed rate F is vertical to the component surface. z =0.05~0.2mm / min, the feed rate F of the rolling head parallel to the component surface x =0.8~2mm / min, ultrasonic vibration frequency is 27~42kHz.

7. The method according to claim 1 or 6, characterized in that: The ultrasonic rolling is a multi-pass ultrasonic rolling, and the actual total reduction of the GH4099 alloy after the ultrasonic rolling does not exceed 0.1 mm.

8. The method according to claim 7, characterized in that The ultrasonic rolling is divided into 8 passes in total. The ultrasonic rolling equipment is set to a downward pressure of 0.8 mm for 4 passes, then the downward pressure is adjusted to 1 mm for 3 passes, and finally the downward pressure is adjusted to 1.2 mm for 1 pass.

9. The method according to claim 1 or 6, characterized in that: The thickness of the rolling deformation layer is 100-200 μm.

10. The method according to claim 1, characterized in that The element composition of the GH4099 alloy is, in terms of mass percentage, 19.5% Cr, 6.8% W, 4.2% Mo, 2.2% Al, 6.8% Co, 1.3% Ti, 0.04% C and the balance Ni.

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