Method for improving mechanical properties of alloys

Abstract

The application provides a method for improving mechanical properties of an alloy, comprising the following steps: step (1): uniformly treating an alloy ingot and then cooling to obtain a uniform alloy; step (2): performing hot working deformation on the uniform alloy in step (1) and then cooling to obtain a deformed alloy; step (3): performing solid solution treatment on the deformed alloy in step (2) and then cooling to obtain a solid solution alloy; and step (4): performing heat preservation on the solid solution alloy in step (3) at 650-700 DEG C for 10-55 h and then cooling to induce grain boundary precipitation of granular or short rod-shaped M23C6 type carbide. The method for improving mechanical properties of the alloy can effectively improve the mechanical properties.

Description

TECHNICAL FIELD

[0001] The present application belongs to the technical field of alloys, and particularly relates to a method for improving the mechanical properties of alloys. BACKGROUND

[0002] At present, as the fourth generation nuclear power plant technology, the high-temperature gas-cooled micro reactor has the advantages of high efficiency, flexibility, safety and the like. The control rod assembly is a key component in the gas-cooled micro reactor, controls the start, stop and operating power of the reactor, and serves for a long time at 450-650 DEG C, while facing the high temperature conditions of about 750 DEG C peak operating temperature and up to 900 DEG C transient accident conditions, which puts higher requirements on the high-temperature mechanical properties, microstructure stability, heat aging resistance and high-temperature long-time creep properties of the control rod material.

[0003] However, 800H alloy, as a Cr and Ni content high solid solution strengthening type single-phase austenitic alloy, has good high-temperature heat aging resistance and long-time creep properties, and is used as a candidate material for the control rod of the high-temperature gas-cooled micro reactor. Influenced by the service temperature conditions, the 800H alloy will age and precipitate during use, so that secondary carbides are precipitated at the grain boundaries, so the 800H alloy is generally not subjected to aging treatment in advance, and is usually directly used in the solid solution state. However, the mechanical properties of the 800H alloy in the solid solution state are insufficient and unstable during actual service.

[0004] Therefore, how to provide a method for improving the mechanical properties of alloys which can effectively improve the mechanical properties has become a problem to be solved by those skilled in the art. SUMMARY

[0005] Therefore, the technical problem to be solved by the present application is to provide a method for improving the mechanical properties of alloys, which can effectively improve the mechanical properties.

[0006] In order to solve the above problems, the present application provides a method for improving the mechanical properties of alloys, comprising the following steps:

[0007] Step (1): uniformly treating and cooling the alloy ingot to obtain a homogenized alloy;

[0008] Step (2): deforming and cooling the homogenized alloy in step (1) to obtain a deformed alloy;

[0009] Step (3): solid solution treating and cooling the deformed alloy in step (2) to obtain a solid solution alloy;

[0010] Step (4): after heat preservation at 650-700 DEG C for 10-55 h, cooling the solid solution alloy in step (3) to induce grain boundary precipitation of granular or short rod-shaped M23C6 type carbide.

[0011] Further, the M23C6 type carbide is semi-continuously and uniformly distributed;

[0012] Further, the size of the carbide is not greater than 1 μm.

[0013] Further, in step (1), the alloy ingot is cooled after homogenization treatment to obtain a homogenized alloy, comprising the following steps:

[0014] The alloy ingot is heated to 1160-1200 ℃ and held for 8-12 h.

[0015] Further, heating the alloy ingot to 1160-1200 ℃ comprises the following steps:

[0016] The alloy ingot is placed in a box-type resistance furnace and heated to 1160-1200 ℃;

[0017] Further, the heating rate is 5-10 ℃ / min.

[0018] Further, in step (2), the hot working deformation comprises a forging + hot rolling process or a hot extrusion process.

[0019] Further, in step (3), the deformed alloy in step (2) is cooled after solution treatment, comprising the following steps:

[0020] The deformed alloy is heated to a preset temperature and held for a preset time to eliminate secondary carbides formed during the hot working deformation and to make the grain size of the alloy coarser than 5 levels.

[0021] Further, in step (3), the deformed alloy in step (2) is cooled after solution treatment, comprising the following steps: the muffle furnace is heated to a preset temperature, the deformed alloy is put into the furnace, and cooled after holding for a preset time.

[0022] Further, the preset temperature is 1120-1200 ℃;

[0023] Further, the preset time is 2-3 h.

[0024] Further, the cooling mode in step (1) is air cooling.

[0025] Further, the cooling mode in step (2) is air cooling.

[0026] Further, the cooling mode in step (3) is water cooling.

[0027] Further, the cooling mode in step (4) is water cooling.

[0028] Further, the alloy is 800H alloy

[0029] The method for improving the mechanical properties of the alloy can effectively improve the mechanical properties; the low-temperature aging at 650 DEG C-700 DEG C can improve the nucleation rate of M23C6 type grain boundary carbide and reduce the size of carbide precipitation. The aging time is selected as 10h-55h, and the purpose is to obtain grain boundary carbide with suitable morphology, size and distribution. The morphology of the carbide should be granular or short rod, and the average size is less than 1um, so as to reduce the stress concentration of the alloy during stress process. The carbide is semi-continuous and uniformly distributed in the grain boundary, which can effectively hinder the crack propagation and improve the strength of the alloy. BRIEF DESCRIPTION OF DRAWINGS

[0030] Figure 1 The grain boundary carbide precipitation state under different heat treatment systems;

[0031] Figure 2 The SEM microstructure morphology photo of the sample obtained in Example 5;

[0032] Figure 3 The SEM microstructure morphology photo of the sample obtained in Comparative Example 1.

[0033] In Figure 1 , the method comprises the following steps:

[0034] (a) solid solution state; (b) 650 DEG C for 2h; (c) 650 DEG C for 5h; (d) 650 DEG C for 10h; (e) 650 DEG C for 20h; (f) 650 DEG C for 55h; (g) 650 DEG C for 100h; (h) 650 DEG C for 210h. DETAILED DESCRIPTION

[0035] For the defects of insufficient and unstable mechanical properties of the 800H alloy in the actual service process in the solid solution state, the inventors of the present application found that the 800H alloy in the solid solution state has the problem of insufficient mechanical properties in the initial stage of service because the grain boundary carbide has not been precipitated in time or the amount of precipitation is small, and the grain boundary strengthening effect is weak. And because the service temperature of the control rod fluctuates, it may lead to uncontrollable precipitation state of the grain boundary carbide of the 800H alloy, so there is a risk of unstable mechanical properties.

[0036] In combination with the method for improving the mechanical properties of the alloy provided in the present application, the method comprises the following steps: Figures 1-3

[0037] Step (1): uniformly heat treating the alloy ingot and cooling to obtain a uniform alloy;

[0038] Step (2): deforming the uniform alloy in step (1) by hot working and cooling to obtain a deformed alloy;

[0039] ​Step (3): solid solution treatment and cooling of the wrought alloy in step (2) to obtain a solid solution alloy;

[0040] Step (4): aging of the solid solution alloy in step (3) at 650-700℃ for 10-55h and then cooling to induce fine and uniformly distributed granular or short rod-shaped M23C6 carbides at grain boundaries. In step (4), low-temperature aging at 650-700℃ can increase the nucleation rate of M23C6 grain boundary carbides and reduce the size of the carbide precipitates. In the 650-700℃ low-temperature range, the aging time is selected to be 10-55h to obtain grain boundary carbides with suitable morphology, size and distribution. The carbides are granular or short rod-shaped with an average size of less than 1μm, which can reduce stress concentration during alloy loading, and the carbides are semi-continuously and uniformly distributed at grain boundaries, which can effectively hinder crack propagation and improve alloy strength. The average size of the carbides refers to the average diameter for granular carbides and the average length for short rod-shaped carbides. M in M23C6 mainly refers to Cr, i.e. Cr23C6, and 800H alloy has a high Cr content, which can precipitate. The present application achieves the technical effect of inducing the precipitation of granular or short rod-shaped M23C6 carbides at grain boundaries. Each step has requirements based on the cooperation of the steps, and each step can be implemented. In the cooperative process, the step of aging at 650-700℃ for 10-55h is mainly relied on. The precipitation temperature mainly determines whether the carbides can be precipitated, and the precipitation time mainly adjusts the amount and size of the precipitates, but the two cannot be separated and need to be cooperated.

[0041] The above alloy is 800H alloy.

[0042] The present application can promote the fine and uniform precipitation of M23C6 carbides at the grain boundaries of 800H alloy, hinder crack propagation, improve grain boundary strength, and thus optimize the mechanical properties of the alloy.

[0043] The present application also discloses some embodiments in which the M23C6 carbides are semi-continuously and uniformly distributed.

[0044] The present application also discloses some embodiments in which the size of the carbides is less than 1μm to reduce stress concentration during alloy loading, and the carbides are semi-continuously and uniformly distributed at grain boundaries, which can effectively hinder crack propagation and improve alloy strength. The larger the size of the carbides, the more likely it is to cause stress concentration and cracking during alloy loading. Therefore, the small size of the carbides (limited to an average size of less than 1μm in the present application) can reduce stress concentration. The semi-continuous and uniform distribution of the carbides at grain boundaries can effectively hinder crack propagation and improve alloy strength, which is the essence of second phase grain boundary strengthening, i.e. the hindering effect of second phase particles at grain boundaries. Too few precipitates have small hindering effect, and semi-continuous precipitation has better hindering effect.

[0045] The application also discloses some embodiments, wherein the alloy ingot is cooled after the homogenization treatment in step (1) to obtain a homogenized alloy, and the method comprises the following steps:

[0046] The alloy ingot is heated to 1160-1200 DEG C and kept for 8-12 hours. In step (1) of the application, the alloy ingot is subjected to high-temperature homogenization treatment at 1160-1200 DEG C / 8-12h before hot deformation, which aims to dissolve large-size carbides precipitated during slow cooling of the ingot, promote element homogenization diffusion, and reduce re-precipitation of the dissolved carbides in the subsequent hot deformation process. After the carbides are dissolved, the local position of the solute elements is prone to re-precipitation, and the homogenization refers to eliminating the local enrichment of the solute elements to make the elements uniformly distributed.

[0047] The application also discloses some embodiments, wherein the alloy ingot is heated to 1160-1200 DEG C, and the method comprises the following steps:

[0048] The alloy ingot is placed into a box-type resistance furnace, and is heated to 1160-1200 DEG C.

[0049] The application also discloses some embodiments, wherein the heating rate is 5-10 DEG C / min.

[0050] The application also discloses some embodiments, wherein the hot working deformation in step (2) comprises a forging+hot rolling process or a hot extrusion process.

[0051] The application also discloses some embodiments, wherein the deformed alloy in step (2) is subjected to solution treatment and then cooled in step (3), and the method comprises the following steps:

[0052] The deformed alloy is heated to a preset temperature and kept for a preset time, so as to eliminate secondary carbides formed in the hot working deformation process and make the grain size of the alloy coarser than 5. In the solution treatment in step (3), the solution temperature and time should be selected to make the secondary carbides formed in the hot working deformation process sufficiently dissolved, and the solute elements sufficiently diffused, so as to eliminate the influence of the unsolved carbides and the solute concentration gradient on the carbide aging re-precipitation morphology and distribution in advance; the solution system should also be selected to make the grain size coarser than 5, so as to improve the high-temperature long-time creep performance of the alloy during service.

[0053] The application also discloses some embodiments, wherein the deformed alloy in step (2) is subjected to solution treatment and then cooled in step (3), and the method comprises the following steps: the muffle furnace is heated to a preset temperature, the deformed alloy is put into the furnace, and the deformed alloy is cooled after being kept for a preset time.

[0054] The application also discloses some embodiments, wherein the preset temperature is 1120-1200 DEG C.

[0055] This application also discloses some embodiments, with a preset time of 2-3 hours. The preset temperature for solution treatment is generally between 1120-1200℃, which is related to the smelting composition and the microstructure after hot deformation. The selection of the preset temperature should be sufficient to eliminate secondary carbides formed during the hot working deformation of the alloy, while ensuring that the alloy grain size is coarser than level 5.

[0056] The preset solution treatment temperature is related to the smelting composition and the microstructure after hot deformation. This is influenced by many factors. For example, a higher nitrogen content in the alloy will refine the grains, requiring a higher solution temperature to grow them to a grain size coarser than grade 5. Furthermore, different hot deformation processes (such as hot extrusion or hot rolling) will produce hot-deformed grains of different sizes. The finer the initial hot-deformed grain size, the higher the solution temperature required to grow them to a grain size coarser than grade 5. Therefore, this application stipulates that the solution treatment regime should be selected to ensure a grain size coarser than grade 5.

[0057] Generally, the composition limits for each element in an 800H alloy are defined as a range. All alloys within this range meet the 800H composition requirements. However, the solution temperatures for different components within this range vary. Each component requires adjustment to its corresponding solution temperature.

[0058] This application also discloses some embodiments in which the cooling method in step (1) is air cooling;

[0059] This application also discloses some embodiments in which the cooling method in step (2) is air cooling;

[0060] This application also discloses some embodiments in which the cooling method in step (3) is water cooling;

[0061] This application also discloses some embodiments in which the cooling method in step (4) is water cooling.

[0062] To effectively prevent the precipitation of the second phase during the slow cooling process, these cooling methods employ rapid cooling.

[0063] from Figure 1 As can be seen from the grain boundary carbide precipitation states and corresponding tensile properties under different heat treatment regimes in Table 1, the alloy in the solution state has no grain boundary carbides. Figure 1 a) Grain boundaries are weak points under high-temperature stress, resulting in low alloy strength. Grain boundary carbides appear after holding at 650℃ for 2 hours. Grain boundary precipitation strengthens the alloy, increasing its strength compared to the solid solution state. However, some grain boundaries still lack precipitated carbides. Figure 1 b). After holding at the temperature for 5 hours, grain boundary carbides precipitated at most grain boundaries, further improving the alloy strength. However, the density of grain boundary carbides was low and their distribution along the grain boundaries was sparse. Figure 1 c). When the heat preservation time is extended to 10-55 hours ( Figure 1d-f), the grain boundary carbide is semi-continuous along the grain boundary, the carbide maintains a granular or short rod shape, the average size is less than 1 μm, at this time the grain boundary precipitation strengthening effect is optimal, and the alloy strength is significantly improved. When the holding time is further extended to within 100 h, the grain boundary carbide gradually grows, but the size growth is not significant, and the performance change is not obvious. When the holding time is 100-210 h, Figure 1 g-h), the grain boundary carbide size is large, and is distributed in a film shape along the grain boundary, and the mechanical properties of the alloy show a downward trend.

[0064] Table 1 800H alloy tensile property data at 650 ℃ after aging for different times

[0065]

[0066] From the above test analysis, it can be seen that: the water cooling treatment at 650 ℃ for 10-55 h can make the alloy obtain fine and uniform granular or short rod-shaped M23C6 type grain boundary carbide in size before service, so as to solve the problem of insufficient mechanical properties of the alloy in the initial service stage, and avoid the uncontrollable precipitation state of the grain boundary carbide due to the service temperature fluctuation. Different from the coarse dendritic, flaky and long rod-shaped carbide precipitated in the high-temperature slow cooling process, due to the high nucleation rate of the carbide, the large number of precipitates, the small solute concentration gradient, the stable carbide interface, the fine granular or short rod-shaped carbide is mainly precipitated, and is semi-continuous and uniformly distributed along the grain boundary. In the high-temperature tensile process, the grain boundary position is the main weak link, which seriously affects the high-temperature strength of the alloy, and the small size grain boundary carbide can hinder crack propagation and improve the grain boundary strength, thereby optimizing the mechanical properties of the alloy.

[0067] The application has the following beneficial effects:

[0068] The process method is aimed at the 800H alloy of the gas-cooled micro reactor control rod material, and the process method promotes the fine and uniform precipitation of M23C6 type carbide at the grain boundary of the 800H alloy through heat treatment, the grain boundary carbide hinders crack propagation, improves the grain boundary strength, and thus improves the mechanical properties of the alloy.

[0069] Embodiment

[0070] Embodiment 1:

[0071] A 800H alloy ingot is melted by vacuum induction melting, and the alloy ingredient before melting is (wt. %): C 0.092 %, Cr 20.9 %, Ni 30.4 %, Al 0.58 %, Ti 0.59 %, Si 0.33 %, Mn 0.98 %, and the balance is Fe.

[0072] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1160°C x 12h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into a 40x40mm bar after 2h of holding at 1200°C, and the forging process is two upsetting and two drawing, with intermediate re-melting. The forged bar is hot rolled into a Φ16mm bar after 1h of holding at 1120°C, and then air cooled. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, with a solid solution system of 1160°C x 2h, water quenching, and a grain size of 4-5 levels. Finally, 10h of holding at 700°C, water cooling.

[0073] Example 2:

[0074] The 800H alloy ingot is smelted by vacuum induction melting, and the alloy ingredient before smelting is (wt. %): C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance is Fe.

[0075] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1180°C x 11h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into a 40x40mm bar after 2h of holding at 1200°C, and the forging process is two upsetting and two drawing, with intermediate re-melting. The forged bar is hot rolled into a Φ16mm bar after 1h of holding at 1120°C. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, with a solid solution system of 1160°C x 2h, water quenching, and a grain size of 4-5 levels. Finally, 10h of holding at 680°C, water cooling.

[0076] Example 3:

[0077] The 800H alloy ingot is smelted by vacuum induction melting, and the alloy ingredient before smelting is (wt. %): C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance is Fe.

[0078] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1200°C x 10h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into a 40x40mm bar after 2h of holding at 1200°C, and the forging process is two upsetting and two drawing, with intermediate re-melting. The forged bar is hot rolled into a Φ16mm bar after 1h of holding at 1120°C. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, with a solid solution system of 1160°C x 2h, water quenching, and a grain size of 4-5 levels. Finally, 10h of holding at 650°C, water cooling.

[0079] Example 4:

[0080] The 800H alloy ingot was vacuum induction melted, and the alloy ingredient before smelting was (wt. %): C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance was Fe.

[0081] After the alloy ingot smelting was completed, high-temperature homogenization treatment of 1200°C x 10h was carried out in a box-type resistance furnace, and after completion, air cooling was carried out. The alloy ingot was forged into a 40x40mm bar after 2h at 1200°C, and the forging process was two upsetting and two drawing, and the intermediate was returned to the furnace. The forged bar was hot-rolled into a Φ16mm bar at 1120°C for 1h. The hot-rolled bar was subjected to high-temperature solid solution treatment in a muffle furnace, and the solid solution system was 1160°C x 2h, water quenching, and the grain size was 4-5 grade. Finally, water cooling was carried out at 650°C for 20h.

[0082] Example 5:

[0083] The 800H alloy ingot was vacuum induction melted, and the alloy ingredient before smelting was (wt. %): C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance was Fe.

[0084] After the alloy ingot smelting was completed, high-temperature homogenization treatment of 1200°C x 8h was carried out in a box-type resistance furnace, and after completion, air cooling was carried out. The alloy ingot was forged into a 40x40mm bar after 2h at 1200°C, and the forging process was two upsetting and two drawing, and the intermediate was returned to the furnace. The forged bar was hot-rolled into a Φ16mm bar at 1120°C for 1h. The hot-rolled bar was subjected to high-temperature solid solution treatment in a muffle furnace, and the solid solution system was 1160°C x 2h, water quenching, and the grain size was 4-5 grade. Finally, water cooling was carried out at 650°C for 55h. The SEM microstructure of the obtained sample is shown in Figure 2 It can be seen that the fine granular M23C6 carbide is uniformly distributed along the grain boundary, and the average size of the carbide is less than 1μm.

[0085] Comparative Example 1:

[0086] The 800H alloy ingot was vacuum induction melted, and the alloy ingredient before smelting was (wt. %): C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance was Fe.

[0087] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1160°C x 12h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into 40x40mm bar after 1200°C x 2h, and the forging process is two upsetting and two drawing, and the intermediate is returned to the furnace. The forged bar is hot rolled into Φ16mm bar at 1120°C x 1h. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, and the solid solution system is 1160°C x 2h, water quenching, and the grain size is 4-5 grade. Finally, water cooling is carried out at 900°C x 10h.

[0088] Comparative Example 2:

[0089] The 800H alloy ingot is smelted by vacuum induction melting, and the alloy ingredient before smelting is (wt.%) : C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance is Fe.

[0090] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1160°C x 12h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into 40x40mm bar after 1200°C x 2h, and the forging process is two upsetting and two drawing, and the intermediate is returned to the furnace. The forged bar is hot rolled into Φ16mm bar at 1120°C x 1h. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, and the solid solution system is 1160°C x 2h, water quenching, and the grain size is 4-5 grade. Finally, water cooling is carried out at 900°C x 10h.

[0091] Comparative Example 3:

[0092] The 800H alloy ingot is smelted by vacuum induction melting, and the alloy ingredient before smelting is (wt.%) : C 0.092%, Cr 20.9%, Ni 30.4%, Al 0.58%, Ti 0.59%, Si 0.33%, Mn 0.98%, and the balance is Fe.

[0093] After the alloy ingot smelting is completed, high temperature homogenization treatment of 1160°C x 12h is carried out in a box resistance furnace, and after completion, air cooling is carried out. The alloy ingot is forged into 40x40mm bar after 1200°C x 2h, and the forging process is two upsetting and two drawing, and the intermediate is returned to the furnace. The forged bar is hot rolled into Φ16mm bar at 1120°C x 1h. The hot rolled bar is subjected to high temperature solid solution treatment in a muffle furnace, and the solid solution system is 1160°C x 2h, water quenching, and the grain size is 4-5 grade. The SEM microstructure of the obtained sample is shown in Figure 3 It can be seen that no carbide precipitates at the grain boundary.

[0094] The samples obtained in the above examples and comparative examples are subjected to high temperature tensile property test at 650°C, and the average value of 2 parallel samples in each group is taken, and the results are shown in Table 2.

[0095] Table 2 Comparison of tensile properties of samples obtained in examples and comparative examples at 650℃

[0096]

[0097] The microstructure of examples 3, 4 and 5 corresponds to d, e and f in Table 1, respectively. Figure 1 The microstructure of example 1 is basically similar to that of f in Table 1. Figure 1 The microstructure of example 2 is basically similar to that of e in Table 1. Figure 1 The microstructure of example 3 is basically similar to that of d in Table 1. The microstructure of example 4 is basically similar to that of e in Table 1.

[0098] The microstructure of comparative example 1 is basically similar to that of b in Table 1. Figure 1 The microstructure of comparative example 2 is basically similar to that of b in Table 1.

[0099] Therefore, the grain boundary carbides precipitated in the examples are fine carbides with a size of less than 1 micron, and the number of carbides is large, which are semi-continuous and uniformly distributed, and can effectively hinder crack propagation. In comparative example 1, the number of grain boundary carbides precipitated is small, and the size is large, which can easily lead to stress concentration. In comparative example 2, the grain boundary carbides precipitated are small in size because they are not easy to grow due to low temperature, and therefore the effect of hindering cracks is not obvious. In comparative example 3, no grain boundary carbides are precipitated in the solid solution state, and the grain boundary is the weak link of cracking.

[0100] The present application is directed to a method for improving the performance of 800H alloy, which comprises the following steps:

[0101] (1) Put the 800H alloy ingot after smelting and cooling into a box-type resistance furnace, and heat it to the target temperature at a heating rate of 5-10℃ / min, and then heat it at the target temperature for 8-12h for high-temperature homogenization treatment and air cooling. The target temperature for high-temperature homogenization treatment is 1160-1200℃.

[0102] (2) Hot working deformation is performed on the 800H alloy after homogenization treatment, and a forging+hot rolling process or a hot extrusion process is used for hot working deformation, and air cooling is performed after hot deformation.

[0103] (3) The alloy after hot deformation is subjected to solid solution treatment, the muffle furnace is heated to the target temperature, and then put into the furnace, and then water-cooled after 2h of heat preservation. The target temperature for solid solution treatment should be selected to satisfy the elimination of secondary carbides formed during the hot working deformation of the alloy, and at the same time, the grain size of the alloy should be coarser than 5 levels.

[0104] (4) The alloy after solid solution treatment is heat treated at 650-700℃ for 10-55h and then water-cooled, to induce the precipitation of fine and uniformly distributed granular or short rod-shaped M23C6 type carbides at the grain boundary, and the average size of the grain boundary carbides is not greater than 1μm. The grain boundary carbides can hinder crack propagation and improve the grain boundary strength, thereby optimizing the mechanical properties of the alloy.

[0105] Those skilled in the art can easily understand that the above-mentioned advantageous modes can be freely combined and superimposed without conflict.

[0106] The above merely describes the preferred embodiments of the present application, but should not be used to limit the present application. Any modification, equivalent replacement, and improvement made within the spirit and principle of the present application shall be included in the protection scope of the present application. The above merely describes the preferred embodiments of the present application, but should not be used to limit the present application. Any modification, equivalent replacement, and improvement made within the spirit and principle of the present application shall be included in the protection scope of the present application.

Claims

1. A method of improving the mechanical properties of an alloy, characterized in that, The method comprises the following steps: Step (1): homogenizing the alloy ingot and cooling to obtain a homogenized alloy; Step (2): deforming the homogenized alloy in step (1) by hot working and cooling to obtain a deformed alloy; Step (3): solid-solutionizing the deformed alloy in step (2) and cooling to obtain a solid-solutionized alloy; Step (4): cooling the solid-solutionized alloy in step (3) after holding at 650-700℃ for 10-55h to induce grain boundary precipitation of granular or short rod-shaped M23C6 carbide; The alloy is 800H alloy; the M23C6 carbide is semi-continuously and uniformly distributed; and the size of the carbide is less than 1μm.

2. The method of improving the mechanical properties of an alloy according to claim 1, characterized in that, In step (1), the homogenized alloy is obtained by homogenizing the alloy ingot and cooling, comprising the following steps: The alloy ingot is heated to 1160-1200℃ and held for 8-12h.

3. The method of claim 2 wherein the alloy is subjected to a heat treatment comprising, The alloy ingot is heated to 1160-1200℃, comprising the following steps: The alloy ingot is placed in a box-type resistance furnace and heated to 1160-1200℃.

4. The method of improving the mechanical properties of an alloy according to claim 3, wherein The heating rate is 5-10℃ / min.

5. The method of claim 1 wherein the alloy is subjected to a heat treatment comprising, In step (2), the hot working deformation comprises a forging + hot rolling process or a hot extrusion process.

6. The method of claim 1 wherein the alloy is subjected to a heat treatment. In step (3), the solid-solutionizing of the deformed alloy in step (2) and cooling, comprising the following steps: The deformed alloy is heated to a preset temperature and held for a preset time to eliminate secondary carbide formed during the hot working deformation and make the grain size of the alloy coarser than 5 grade to improve the long-term creep performance of the alloy during service.

7. The method of improving the mechanical properties of an alloy of claim 1, wherein In step (3), the solid-solutionizing of the deformed alloy in step (2) and cooling, comprising the following steps: heating the muffle to a preset temperature, putting the deformed alloy into the muffle, holding for a preset time, and then cooling.

8. The method of improving the mechanical properties of an alloy according to any one of claims 6-7, wherein, The preset temperature is 1120-1200℃; And / or, the preset time is 2h-3h.

9. The method of claim 1 wherein the alloy is subjected to a mechanical property enhancement process comprising, The cooling mode in step (1) is air cooling; And / or, the cooling mode in step (2) is air cooling; And / or, the cooling mode in step (3) is water cooling; And / or, the cooling mode in step (4) is water cooling.

Images