Preparation method of nickel-based alloy suitable for large-section forge piece and nickel-based alloy

Abstract

The invention discloses a preparation method of a nickel-based alloy suitable for a large-section forge piece and the nickel-based alloy, belongs to the technical field of nickel-based alloys, and solves the problem that the nickel-based alloy in the prior art is difficult to meet the comprehensive performance requirements of the large-section forge piece. The preparation method of the nickel-based alloy suitable for the large-section forge piece comprises the steps of preparing a forging stock through four-upsetting and four-drawing and carrying out solution treatment on the forging stock. Multi-section cooling is conducted in the four-upsetting and four-drawing process; the solution treatment temperature is (1080-1100) + 25 * (4-a), and a is the required grade number of the grain size. The nickel-based alloy prepared through the method can be used for manufacturing large-tonnage large-section forgings.

Description

Technical Field

[0001] The invention relates to the technical field of nickel-based alloys, and in particular to a preparation method of a nickel-based alloy suitable for large-section forgings and the nickel-based alloy. Background Art

[0002] At present, the demand for large-scale high-temperature alloy forgings is becoming more and more obvious. Taking ultra-supercritical steam turbine rotor forgings as an example, nickel-based alloys are required to be used when the steam temperature exceeds 650°C. The diameter is generally more than 800mm, the medium-pressure rotor is more than 1100mm, the tonnage is more than 10 tons, and the demand for ingots is more than 20 tons, and even the ingots are expected to reach 40-60 tons. The room temperature yield strength is required to exceed 450MPa, and it is required to be in stable service for 30 years at 650°C or even 700°C and above. Among the mature nickel-based alloys currently in use, there is no alloy that can meet the required comprehensive performance and manufacture large-section forgings. Summary of the invention

[0003] In view of the above situation, the present invention aims to provide a method for preparing a nickel-based alloy suitable for large-section forgings and a nickel-based alloy, so as to solve the problem that existing nickel-based alloys are difficult to meet the comprehensive performance requirements of large-section forgings.

[0004] The purpose of the present invention is mainly achieved through the following technical solutions:

[0005] On the one hand, the present invention provides a method for preparing a nickel-based alloy suitable for large-section forgings, comprising preparing a forging blank by four-upsetting and four-drawing and performing a solid solution treatment on the forging blank;

[0006] Multiple cooling stages during the four-upsetting and four-drawing process;

[0007] The temperature of the solution treatment is (1080~1100)+25×(4-a), where a is the required grain size level.

[0008] Further, the four-upsetting and four-drawing method for preparing the forging blank includes:

[0009] S301, first fire: upsetting: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%;

[0010] S302, second fire: drawing: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%;

[0011] S303, third fire: upsetting: heating temperature 1145 ~ 1150 ℃, deformation greater than 40%;

[0012] S304, the fourth fire: drawing, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%;

[0013] S305, the fifth fire: upsetting, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%;

[0014] S306, the sixth fire: drawing, heating temperature X + 5 ~ X + 15 ℃, deformation greater than 30%; S307, the seventh fire: upsetting, heating temperature within the range of X ± 5 ℃, deformation greater than 40%;

[0015] S308, the eighth fire: drawing, heating temperature X-5 ~ X-30 ℃, deformation greater than 30%, pressing rate 20 ~ 200mm / s;

[0016] The value of X is determined using the following formula: X=1115+6000C / Cr.

[0017] Furthermore, the preparation method also includes: smelting and pouring to obtain an ingot; and homogenizing the ingot.

[0018] Furthermore, the process of homogenizing the ingot is: keeping the ingot at a temperature range of 1180 to 1220° C. for 40 to 150 hours.

[0019] Furthermore, the components of the nickel-based alloy include, by mass percentage, Cr: 17% to 19%; Mo: 8% to 10%; Co: 11% to 13%; B: 0.003% to 0.006%; Al: 1.2% to 1.4%; Ti: 1.2% to 1.4%; C: 0.03% to 0.05%; the remainder is Ni and some other inevitable impurities.

[0020] Furthermore, in the composition of the nickel-based alloy, Al / Ti is controlled to be 0.9 to 1.05.

[0021] Furthermore, in the composition of the nickel-based alloy, Cr / C is controlled within the range of 350-500.

[0022] Furthermore, the grain size of the forging blank prepared by four-upsetting and four-drawing is finer than grade 4.

[0023] Furthermore, large-section forgings include forgings with a diameter of more than 500 mm.

[0024] The present invention also provides a nickel-based alloy, which is prepared by adopting the above-mentioned preparation method.

[0025] Compared with the prior art, the present invention has the following beneficial effects:

[0026] The preparation method of the nickel-based alloy suitable for large-section forgings of the present invention prepares a forging blank through four-upsetting and four-drawing, and the multi-stage cooling during the four-upsetting and four-drawing process can refine the grains, and at the same time, the large block carbides in the cast structure are broken as much as possible, so that the distribution of grains and carbides is more uniform; by accurately controlling the temperature of the solution treatment, suitable required grains can be obtained.

[0027] The preparation method of the present invention ensures the uniformity of the grains by precisely controlling the order and process parameters of each step, ensuring that the grain size difference from the core to the edge of the large-section forging is less than 1.5 levels (for example, less than 1 level); at the same time, it also reduces the cracking tendency and improves the yield of large-section forgings (for example, more than Φ500mm).

[0028] The preparation method of the present invention combines precise control of the types and contents of elements in the alloy, and coordinated control of the coordinated quantitative relationship of different elements, combined with a precise control preparation method, to ensure a light degree of segregation, suitable for smelting large ingots, thereby manufacturing large-tonnage and large-section forgings.

[0029] The nickel-based alloy of the present invention has an overall room temperature tensile strength of >960MPa (e.g. 970-1030MPa), a yield strength of >560MPa (e.g. 570-620MPa), an impact AKV of >40J (e.g. 55-80J), a 700°C tensile strength of >810MPa (e.g. 815-880MPa), a yield strength of >510MPa (e.g. 520-570MPa), an elongation and a cross-sectional expansion ratio at room temperature and 700°C of more than 25% (e.g. more than 30%), and a 100,000-hour endurance strength of more than 165MPa (e.g. more than 170MPa), which can meet the requirements of long-term stable service at 700°C.

[0030] Other features and advantages of the present invention will be described in the following description, and part of them will become obvious from the description, or will be understood by practicing the present invention. The purpose and other advantages of the present invention can be realized and obtained by the structures particularly pointed out in the written description, claims, and drawings. BRIEF DESCRIPTION OF THE DRAWINGS

[0031] The drawings are only for the purpose of illustrating particular embodiments and are not to be considered limiting of the present invention. Like reference symbols denote like components throughout the drawings.

[0032] Figure 1 The core grain size after the treatment in step 3 of Example 1;

[0033] Figure 2 The edge grain size after the treatment in step 3 of Example 1;

[0034] Figure 3The longitudinal carbide distribution after the treatment in step 4 of Example 1;

[0035] Figure 4 The core grain size after the treatment in step 3 of comparative example 3;

[0036] Figure 5 The edge grain size after the treatment in step 3 of comparative example 3;

[0037] Figure 6 The longitudinal carbide distribution after the treatment in step 4 of comparative example 3;

[0038] Figure 7 The longitudinal carbide distribution after the treatment in step 4 of comparative example 4;

[0039] Figure 8 The figure shows the intracrystalline γ′ phase of the present invention;

[0040] Fig. 9 The fine particles M in the grain boundary and grain of the present invention are shown. 23 C6 type carbide. DETAILED DESCRIPTION

[0041] Preferred embodiments of the present invention will be described in detail below in conjunction with the accompanying drawings, wherein the accompanying drawings constitute a part of the present invention and are used to explain the principles of the present invention together with the embodiments of the present invention.

[0042] The present invention provides a method for preparing a nickel-based alloy suitable for large-section forgings. The method is used to prepare a nickel-based alloy for large-section forgings with a diameter of more than 500 mm (for example, more than 750 mm).

[0043] The preparation method of the invention comprises four-upsetting and four-drawing steps to prepare a forging blank, wherein the temperature is reduced in multiple stages during the four-upsetting and four-drawing steps, and the initial forging temperature of the last tempering step is 1100-1125°C.

[0044] Specifically, the preparation method of the present invention comprises:

[0045] Step 1, smelting and pouring to obtain an ingot;

[0046] Step 2: homogenize the ingot: keep it at a temperature of 1180-1220°C for 40-150h;

[0047] Step 3: Prepare forging blanks by multi-fire forging; multi-stage cooling between fires, the initial forging temperature of the last fire is 1100-1130°C, and the final forging temperature is 895-905°C;

[0048] Step 4: Performing solid solution and aging treatment on the forging billet to obtain a nickel-based alloy.

[0049] Specifically, the components of the nickel-based alloy suitable for large-section forgings of the present invention include, by mass percentage, Cr: 17% to 19%; Mo: 8% to 10%; Co: 11% to 13%; B: 0.003% to 0.006%; Al: 1.2% to 1.4%; Ti: 1.2% to 1.4%; C: 0.03% to 0.05%; the remainder is Ni and some other inevitable impurities.

[0050] Specifically, the content ratio of Al to Ti (ie, Al / Ti) is controlled to be 0.9 to 1.05.

[0051] Specifically, the content ratio of Cr to C (ie, Cr / C) is controlled within the range of 350-500.

[0052] The following is a specific description of the functions and dosages of the components contained in the alloy of the present invention:

[0053] Cr: It mainly plays a strengthening and anti-oxidation role. The strengthening effect is divided into two aspects. On the one hand, Cr itself is a solid solution strengthening element, which can play a strengthening role and improve the strength of the alloy. On the other hand, under appropriate heat treatment process, Cr can form fine particles of M 23 C6 carbides are precipitated at the grain boundaries, increasing the grain boundary strength of the material, thereby improving the durability of the alloy. Cr will oxidize in a high temperature environment and form a dense oxide film on the surface of the forging, protecting the inside of the forging from oxidation and improving the oxidation resistance of the alloy. If the Cr content is too low, the oxidation resistance will decrease, and if it is too high, M 23 The amount of C6 carbide precipitation increases and is easy to grow, losing the effect of strengthening the grain boundary. At the same time, it will promote the precipitation of σ phase, which is a TCP (topological dense packing) phase that is easy to grow into needles, destroying the continuity of the material and reducing M 23 The C6 content causes the performance of the forging to deteriorate or even fail, and too high a Cr content will lead to a significant increase in material costs. Therefore, the present invention controls the Cr content to 17% to 19%.

[0054] Mo: It mainly plays a role in solid solution strengthening. At the same time, since Mo itself is a high melting point element, it can significantly increase the thermal strength of the material, that is, increase the high temperature strength of the material. Mo can also partially replace M 23 The Cr element in C6 carbide increases M 23Stability of C6 carbides. If the Mo content is too low, the solid solution strengthening effect is not obvious. If it is too high, it is easy to promote the precipitation of M6C carbides. During long-term service, it will also promote the precipitation of μ phase, which is a TCP (topological dense packing) phase. It is easy to grow into long strips, destroying the continuity of the material, causing the performance of the forging to decline or even fail. At the same time, Mo will produce a certain segregation between dendrites during smelting and solidification, promoting the precipitation of MC carbides. Due to its large atomic diameter, it is difficult to diffuse. It is difficult to achieve a completely uniform distribution after homogenization, especially for large ingots. Its uneven distribution will cause the precipitation of unevenly distributed M6C carbides in the subsequent hot working process, resulting in mixed crystals of forgings. At the same time, the precipitation of M6C carbides will also affect M 23 The precipitation of C6 carbides affects the durability of the material. Therefore, the Mo content is controlled within 8% to 10% in the present invention.

[0055] Co: It mainly plays a role in solid solution strengthening. At the same time, since Co itself is a high melting point element, it can significantly increase the thermal strength of the material, that is, increase the high temperature strength of the material. Co can also partially replace M 23 The Cr element in C6 carbide increases M 23 The stability of C6 type carbide. However, too high Co will promote the precipitation of σ phase and μ phase, while too low Co will weaken the solid solution strengthening effect. Therefore, the present invention controls the Co content to 10.9% to 13%.

[0056] B: can be distributed at the grain boundary, purify the grain boundary, and at the same time 23 C6 carbides are distributed at the edge, inhibiting M 23 If the B content is too low, the above effect will not be achieved. If the B content is too high, the precipitation of precipitated phases such as M3B2 will be promoted, thus deteriorating the performance of the material. Therefore, the B content is controlled within 0.003% to 0.006% in the present invention.

[0057] Al and Ti: Promote the precipitation of γ′ phase, increase the content of γ′ phase, and significantly improve the yield strength and endurance strength of the alloy. At the same time, a reasonable ratio of Ti and Al content can keep the mismatch between γ′ phase and matrix at a low level, thereby improving the stability of γ′ phase in long-term high-temperature service, thereby significantly improving the performance and stability of the material in high-temperature and long-term service. At the same time, it can also ensure that the precipitation temperature of γ′ phase is kept below 950°C during hot working, ensuring a wide hot working temperature range, and not too fast precipitation to cause excessive forging deformation resistance and cracking. Too high or too low of either of the two will affect the mismatch between γ′ phase and matrix in the alloy, and if both are too low, the strength improvement effect is not obvious, and too high will affect the hot working performance. At the same time, with a reasonable heat treatment process, the overall performance of large-section forgings can be uniform. Therefore, the content of Al and Ti should be controlled within 1.2% to 1.4%, and specifically, the value of Al / Ti should be controlled within the range of 0.9 to 1.05.

[0058] C: C in nickel-based alloys mainly precipitates carbides, and the desired carbide type is M 23 C6 type, but for large-tonnage nickel-based alloy ingots, due to the serious segregation during smelting, too high C will promote the precipitation of MC and M6C carbides, and also promote the M 23 The coarsening of C6 carbides should not be too high. Too low will cause M 23 The amount of C6 carbide precipitation is insufficient. Therefore, the present invention controls the C content to 0.03% to 0.05%, and with the reasonable grain size distribution after heat treatment, a large number of fine particles of M can be precipitated at the grain boundary. 23 C6 type carbide.

[0059] After in-depth research, the inventors have controlled the C and Cr contents in the alloy of the present invention to be relatively low. The main reasons are: 1. The present invention is aimed at large-section and large-tonnage ingots, and large-section and large-tonnage nickel-based alloy ingots are prone to serious segregation during the solidification process. Too high a C content will promote the precipitation of MC-type and M6C-type carbides. MC-type and M6C-type carbides mainly contain Mo, Co, Ti, etc. These elements are important strengthening elements. If too many MC-type and M6C-type carbides are precipitated, the strength of the material will be significantly reduced. Therefore, the C content must be controlled at a relatively low level; 2. The nickel-based alloy of the present invention is aimed at products - ultra-supercritical turbine rotors, which have very high requirements for endurance performance, and larger grains are beneficial to the endurance performance of forgings. Generally, they are controlled at levels 1 to 3, so their grain boundary area is small. If the C and Cr contents are too high, the M6C carbides precipitated at the grain boundaries will be reduced. 23 There will be more C6 carbides, resulting in grain boundary M 23C6 carbides are coarsened and continuous, and are prone to coarsening during subsequent service, thus affecting the continuity of the material and reducing its performance; 3.M 23 The precipitation of C6 carbides in the form of particles at the grain boundaries can significantly improve the durability of the material. 23 The precipitation amount of C6 type carbide, the C content cannot be too low; 4. Under the condition of lower C content, if the Cr content is not controlled within a reasonable range, too much Cr will promote the precipitation of σ phase, which is an undesirable precipitation phase, so the Cr content cannot be too high; 5. The service temperature of the material of the present invention reaches 700°C, and in a steam environment, therefore, in order to ensure sufficient steam corrosion resistance, the Cr content cannot be too low. Taking the above factors into consideration, the C content and Cr need to have a corresponding reasonable range, the C content should be controlled at 0.03% to 0.05%, the Cr content should be controlled at 17% to 19%, and the Cr / C should be controlled at 350 to 500, so that it can not only ensure the grain boundary M 23 The precipitation amount and precipitation morphology of C6 type carbides can also ensure sufficient Cr to play an antioxidant role and can also effectively inhibit the precipitation of σ phase.

[0060] The alloy of the present invention does not contain Nb, Ta, and W. This is because the inventors found in the process of in-depth research that such elements can play a good strengthening role in small forgings, but in large nickel-based alloy forgings, they are easy to cause segregation and form too much MC type carbide, thereby inhibiting the beneficial M at the grain boundary. 23 The precipitation of C6 carbides causes local mixed crystals, which affects the uniformity of the forgings and reduces the key properties of the forgings, such as fatigue and durability.

[0061] Other elements: Elements such as Si, S, O, etc. must be controlled at the lowest level possible.

[0062] Specifically, in the above step 1, the smelting process route is a triple smelting process of vacuum induction melting + electroslag remelting + vacuum consumables or a double smelting process of vacuum induction melting + electroslag remelting or vacuum induction melting + vacuum consumables.

[0063] Specifically, in the above step 2, since the alloy of the present invention contains a certain amount of Ti and a relatively high content of Mo, and the alloy is aimed at large forgings, the ingot diameter is generally large, exceeding 500mm or even exceeding 1000mm, and segregation is easy to occur, among which Ti and Mo are the main ones. The atomic diameter of Ti is relatively small, but the degree of segregation is serious. Although Mo has a relatively small tendency to segregate, its atomic diameter is relatively large and it is difficult to diffuse. Therefore, the present invention adopts a relatively high temperature to make such segregated elements diffuse evenly, so as to provide a guarantee for the final uniformity of the organization. After in-depth research, the temperature was controlled to be kept in the range of 1180-1220℃ for 40-150h. Long-term heat preservation in the temperature range of 1180-1220℃ can effectively eliminate the microscopic segregation formed during the smelting and solidification process; it can also make the M precipitated during the solidification process 23 All C6 and M6C carbides are dissolved back. In the alloy system of the present invention, the dissolution temperature of these two carbides is between 1120 and 1180°C, so the lowest homogenization temperature is selected to reach 1180°C; at the same time, some of the MC carbides that can be dissolved back are dissolved back as much as possible, and the edges and corners of the remaining MC carbides are blunted to avoid forming stress concentration points in the subsequent forging process, thereby forming forging cracks. In this alloy system, when the temperature is higher than 1220°C, it is easy to form initial melting at places such as grain boundaries, causing material continuity damage. Therefore, the control temperature cannot be higher than 1220°C.

[0064] Specifically, in the above step 3, the multi-fire forging adopts four upsetting and four drawing, and the specific steps of four upsetting and four drawing include:

[0065] S301, first fire: upsetting: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%, pressing rate 1 ~ 100mm / s;

[0066] S302, second fire: drawing: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%, single hammer reduction greater than 20%, reduction rate 1 ~ 100mm / s;

[0067] S303, third fire: upsetting: heating temperature 1145 ~ 1150 ℃, deformation greater than 40%, pressing rate 1 ~ 100mm / s;

[0068] S304, the fourth fire: drawing, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%, single hammer pressing greater than 20%, pressing rate 20 ~ 200mm / s;

[0069] S305, the fifth fire: upsetting, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%, pressing rate 1 ~ 100mm / s;

[0070] S306, the sixth fire: drawing, heating temperature X + 5 ~ X + 15 ℃, deformation greater than 30%, single hammer pressing amount controlled between 15% ~ 45%, pressing rate 10 ~ 100mm / s;

[0071] S307, seventh fire: upsetting, heating temperature X ± 5 ℃ range, deformation greater than 40%, pressing rate 5 ~ 100mm / s;

[0072] S308, the eighth fire: drawing, heating temperature X-5 ~ X-30 ℃, deformation greater than 30%, single hammer pressing amount controlled between 15% ~ 25%, pressing rate 20 ~ 200mm / s;

[0073] The value of X is determined using the following formula: X=1115+6000C / Cr.

[0074] It should be noted that in S301, the purpose of the first fire upsetting is to break the coarse dendrites in the as-cast structure and undergo dynamic recrystallization to refine the grains through high-temperature and large-deformation upsetting deformation, while breaking up the large block carbides in the as-cast structure as much as possible; in the ingot state, due to the coarse as-cast grain size, the presence of columnar crystals, etc., the forging plasticity is poor, so in order to avoid cracking, the reduction rate needs to be controlled. At the same time, a lower reduction rate is conducive to more dynamic recrystallization nucleation inside the ingot, increasing the recrystallization ratio, rather than local dynamic recrystallization and softening to form an adiabatic shear band, resulting in uneven deformation resistance inside the forging, causing cracking and mixed crystals.

[0075] It should be noted that in S301, the selection of 1150-1200°C is based on: according to the alloy system of the present invention, it has very good forging plasticity within the range of 1000-1150°C, so forging is generally carried out within this temperature range. Specifically for the first fire, its main purpose is to maximize the fragmentation of the cast structure. The higher the forging heating temperature, the greater the deformation, the higher the degree of recrystallization of the ingot, and the more obvious the fragmentation of the cast structure. However, when the temperature is higher than 1200°C, the deformation heat during the deformation process will cause the internal temperature of the forging to be higher than the heating temperature, which is likely to cause overheating and cracking. In order to maximize the fragmentation of the cast structure and give full play to the forging plasticity, 1150-1200°C is generally selected as the forging heating temperature. The selection of a deformation amount greater than 40% is to maximize the fragmentation of the cast structure.

[0076] Specifically, in S301, at the beginning of upsetting, the pressing rate is 70-100 mm / s, and as the pressing amount increases, it slowly decreases to 3-10 mm / s. The change in the pressing rate from faster at the beginning to slower later is related to the press capacity and the actual process. At the beginning of pressing, the temperature of the billet is high, the cross-section is small, the overall deformation resistance is low, and the pressing rate of the press is fast. As the deformation proceeds, the temperature of the billet drops to a certain extent, and the deformation cross-sectional area increases, the pressure demand on the press increases, and therefore the pressing rate decreases. In the actual operation process, a relatively fast pressing rate at the beginning can effectively break the cast structure without causing local deformation, thereby causing failure and cracking. The subsequent slowing of the pressing rate is more conducive to breaking the recrystallization of the cast structure.

[0077] It should be noted that in S302, the purpose of the drawing is to continue to break the coarse dendrites in the as-cast structure through high-temperature and large deformation drawing deformation on the basis of the first fire, and to cause dynamic recrystallization, refine the grains, and break the large block carbides in the as-cast structure as much as possible. A lower pressing rate is conducive to more dynamic recrystallization nucleation inside the ingot, increasing the recrystallization ratio, rather than local dynamic recrystallization and softening, forming an adiabatic shear band, causing uneven deformation resistance inside the forging, cracking and mixed crystals. Therefore, the heating temperature is controlled at 1150-1200°C, and the pressing rate is 1-100mm / s. For example, at the beginning of the drawing, the pressing rate is 70-100mm / s, and as the pressing amount increases, it slowly decreases to 3-10mm / s.

[0078] In S302, in order to maximize the fragmentation of the as-cast structure and reserve enough deformation for the next fire, if the deformation of the current fire is small, it is difficult to achieve a sufficient deformation in the next fire; therefore, the deformation is controlled to be greater than 40%. Considering that the original as-cast structure is coarse due to the high forging heating temperature, although recrystallization has occurred in most parts of the ingot in step S301, the grains are still coarse, and the critical deformation at this time is large, so the single hammer reduction should be relatively large, so that the coarse grains can be fully recrystallized. If it is too small, recrystallization will not occur or only a small area will recrystallize, and the grains in most areas will grow instead. Therefore, the single hammer reduction is controlled to be greater than 20%.

[0079] Specifically, in S303, on the basis of the first two fires fully crushing the as-cast structure, the recrystallization ratio inside the forging is further increased to obtain a complete recrystallized structure. On the basis of S301 and S302, the structure at most positions has become an equiaxed crystal structure. At this time, the heating temperature is appropriately reduced. The main reason is that after the deformation of the first two fires, the overall grain size of the forging has been refined to a certain extent, the recrystallization nucleation points have increased, the recrystallization difficulty has been reduced, and sufficient recrystallization can occur at a lower temperature. At this time, reducing the heating temperature can not only allow sufficient recrystallization to occur, but also keep the recrystallized grains at a finer level. Especially for the large-section forging of the present invention, the core temperature will have a certain temperature rise due to the deformation heat during the forging process. Within a certain period of time after forging, due to its small thermal conductivity, the temperature will drop slowly, and the grains will grow to a certain extent. The higher the temperature, the greater the grain growth. In order to control the final grain size, this fire begins to reduce the heating temperature in a targeted manner on the basis of the first two fires. Considering that if the temperature is too low, the dynamic recrystallization cannot be completely carried out inside the forging, and the mixed crystal is serious; if the temperature is too high, the grains obtained by recrystallization will grow too large, which is not conducive to the grain size control in the subsequent forging process. Therefore, the temperature is controlled at 1145~1150℃.

[0080] Specifically, in S303, at the beginning of upsetting, the pressing rate is 70-100 mm / s, and as the pressing amount increases, it slowly decreases to 3-10 mm / s.

[0081] Specifically, through the operation of S304, the grains inside the forging can be further recrystallized than the third fire, and the carbides inside the forging can be more evenly distributed, and a good foundation can be laid for the next fire to continue to refine the grains. Specifically, in S304, the heating temperature can be 1145-1150°C. Specifically, in S304, the pressing rate is 20-200 mm / s.

[0082] Specifically, in S305, the heating temperature may be 1145-1150° C. The recrystallization ratio is further increased by deformation, so that the distribution of grains and carbides is more uniform.

[0083] Specifically, in S305, at the beginning of upsetting, the pressing rate is 80-100 mm / s, and as the pressing amount increases, it slowly decreases to 3-10 mm / s.

[0084] Specifically, through the operation of S306, the grains inside the forging can be further refined than the fifth fire, the carbide distribution is more uniform, and a good foundation is laid for the next fire to continue to refine the grains. Specifically, in S306, the heating temperature can be 1135-1150°C. Specifically, in S306, the pressing rate is 10-80 mm / s.

[0085] Specifically, in S307, the heating temperature may be 1125-1140° C. The grains are further refined by deformation. Since the heating temperature is relatively low, the reduction rate may be appropriately increased to ensure the smooth progress of the forging process. This operation may lay a good foundation for the grain size of the final forming fire.

[0086] Specifically, in S307, at the beginning of upsetting, the pressing rate is 90-100 mm / s, and as the pressing amount increases, it slowly decreases to 5-10 mm / s.

[0087] Specifically, in S308, the heating temperature can be 1110-1125°C. In S308, the temperature is continued to be lowered, and the deformation amount is selected to be greater than 30% in order to continue to refine the grains and ensure that the overall grain size of the final forging is finer than level 4. The single hammer pressing amount should be controlled between 15% and 25% to ensure that the critical deformation amount is reached and the grains will not grow abnormally.

[0088] Specifically, in S308, due to the low forging heating temperature, the product is generally more complex in shape and may not reach the final size in this fire. A ninth fire of forging is required. The forging heating temperature of the ninth fire should be reduced by 5 to 10°C based on the heating temperature of the eighth fire, and the deformation in the final fire should be controlled above 20%.

[0089] Specifically, after S308 treatment, the grain size of the forging billet is finer than level 4.

[0090] Specifically, in the above step 4, the solution treatment temperature is 1100-1150°C, and is set according to the target grain size. The solution treatment temperature is: (1080-1100)+25×(4-a), where a is the required grain size level, the solution time is: 3-10h, and the solution holding time (unit h) is determined according to the cross-sectional diameter d (holding time = (2-3) + (d / 100-d / 300). The solution treatment ensures that all precipitations except MC carbides are completely dissolved, and the grain size is uniform and controlled within the required range. After the insulation, water cooling is used to reduce the precipitation of M6C carbides and γ′ phase as much as possible.

[0091] Specifically, in the above step 4, the aging treatment temperature is 720-780°C, and the aging time is 1-30h. The aging treatment temperature and time need to match and are determined according to the cross section (holding time = 0-20 + d / 100-d / 300), and ultimately the core and edge performance can be consistent.

[0092] Specifically, in the above step 4, if the surface quality after forging is good without defects such as cracks, it can be directly put into the heat treatment furnace for heat treatment to reduce energy consumption; if there are defects on the surface, it needs to be air-cooled or air-cooled before surface machining, and heat treatment is performed after the surface defects are removed.

[0093] Specifically, the grain size of the nickel-based alloy after the treatment in step 4 is 1 to 4, the grains are uniform, and the grain size difference from the center to the edge is less than 1.5 levels (for example, less than 1 level); the microstructure is austenite structure + uniformly distributed fine granular M 23 C6 carbide + individual M6C and MC carbides at grain boundaries + nano-scale γ′ phase evenly distributed in the grains.

[0094] The present invention also provides a nickel-based alloy, which is prepared by the above-mentioned preparation method.

[0095] In order to further improve the comprehensive properties of the above-mentioned nickel-based alloy, the components of the above-mentioned nickel-based alloy include, by mass percentage: Cr: 17.1% to 18.9%; Mo: 8.3% to 9.9%; Co: 11% to 11.8%; B: 0.004% to 0.0055%; Al: 1.2% to 1.32%; Ti: 1.2% to 1.33%; C: 0.04% to 0.05%; the remainder is Ni and some other inevitable impurities.

[0096] Specifically, the content ratio of Al to Ti (ie, Al / Ti) is controlled to be between 0.92 and 1.05.

[0097] Specifically, the content ratio of Cr to C (ie, Cr / C) is controlled at 375-450.

[0098] Specifically, the overall room temperature tensile strength of the above-mentioned nickel-based alloy is >960MPa (for example, 970-1030MPa), the yield strength is >560MPa (for example, 570-620MPa), the impact AKV is >40J (for example, 55-80J), the 700°C tensile strength is >810MPa (for example, 815-880MPa), the yield strength is >510MPa (for example, 520-570MPa), the elongation and cross-sectional expansion ratio at room temperature and 700°C are both above 25% (for example, above 30%), and the endurance strength at 700°C and 100,000 hours is above 165MPa (for example, above 170MPa).

[0099] Compared with the prior art, the preparation method of the nickel-based alloy suitable for large-section forgings of the present invention ensures the uniformity of the grains by precisely controlling the order and process parameters of each step, and ensures that the grain size difference from the core to the edge of the large-section forging is less than 1.5 levels (for example, less than 1 level); at the same time, it also reduces the cracking tendency and improves the yield of large-section forgings (for example, above Φ500mm).

[0100] The preparation method of the present invention combines precise control of the types and contents of elements in the alloy, and coordinated control of the coordinated quantitative relationship of different elements, combined with a precise control preparation method, to ensure a light degree of segregation, suitable for smelting large ingots, thereby manufacturing large-tonnage and large-section forgings.

[0101] The nickel-based alloy of the present invention has an overall room temperature tensile strength of >960MPa (e.g. 970-1030MPa), a yield strength of >560MPa (e.g. 570-620MPa), an impact AKV of >40J (e.g. 55-80J), a 700°C tensile strength of >810MPa (e.g. 815-880MPa), a yield strength of >510MPa (e.g. 520-570MPa), an elongation and a cross-sectional expansion ratio at room temperature and 700°C of more than 25% (e.g. more than 30%), and a 100,000-hour endurance strength of more than 165MPa (e.g. more than 170MPa), which can meet the requirements of long-term stable service at 700°C.

[0102] The advantages of precise control of the composition and process parameters of the nickel-based alloy in the preparation method of the present invention are demonstrated below with specific examples and comparative examples. The embodiments of the present invention all provide a preparation method and a nickel-based alloy suitable for large-section forgings.

[0103] Embodiment 1:

[0104] This embodiment provides a preparation method of a nickel-based alloy suitable for large-section forgings and a nickel-based alloy. The main chemical composition of the nickel-based alloy is shown in Table 1. The ingot diameter of this embodiment is Φ900 mm, and the final nickel-based alloy forging diameter is Φ800 mm.

[0105] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0106] (2) Step 2, homogenizing the ingot: keeping the ingot at a temperature of 1200° C. for 100 h;

[0107] (3) Step 3, forging: four upsetting and four drawing, forging temperature range 1170~900℃; details are as follows:

[0108] First fire: upsetting temperature 1170℃, upsetting deformation 60%, upsetting reduction rate: 80mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0109] Second fire: drawing temperature 1170℃, drawing deformation 60%, drawing reduction rate: controlled at 5-50mm / s, single hammer reduction during drawing controlled at ≥30%;

[0110] The third fire: upsetting temperature 1150℃, upsetting deformation 60%, upsetting reduction rate: 80mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0111] Fourth fire: drawing temperature 1150℃, drawing deformation 60%, drawing reduction rate: controlled at 15-50mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0112] Fifth fire: upsetting temperature 1145℃, upsetting deformation 60%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0113] The sixth fire: the drawing temperature is 1140℃, the drawing deformation is 60%, the drawing reduction rate is controlled at 15-50mm / s, and the single hammer reduction is controlled at 40% during drawing;

[0114] The seventh fire: upsetting temperature 1130℃, upsetting deformation 60%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0115] Eighth fire: drawing temperature 1125℃, drawing deformation 65%, drawing reduction rate: controlled at 15-50mm / s, and single hammer reduction during drawing controlled at 20%;

[0116] (4) Step 4, heat treatment: one solution treatment + one aging treatment, solution treatment temperature: 1130°C, solution treatment time: 7h, cooling by water cooling; aging treatment temperature: 730°C, aging time: 10h;

[0117] Figure 1 The figure shows the core grain size after step 3 of this embodiment. Figure 2 The edge grain size after step 3 is shown. Figure 3 The figure shows the distribution of carbides in the longitudinal section after the step 4 treatment. It can be seen that after the step 3 treatment of this embodiment, the grain size of the core of the forging is 4.5, and the grain size of the edge is 5; after the step 4 treatment, the grain size of the core is 2.5, the grain size of the edge is 3, and the grade difference is 0.5; the structure of the forging of this embodiment is a single austenite structure + fine grained M at the grain boundary 23 C6 carbide + a small amount of M6C / MC carbide + uniformly dispersed γ′ phase, Figure 3 From the longitudinal section, it can be seen that the overall carbide fragmentation effect is good and evenly distributed, and there is basically no strip carbide. Figure 8 The γ′ phase in the crystal is shown; Fig. 9 The fine particles M in the grain boundary and grain 23 C6 type carbide.

[0118] The overall room temperature tensile strength of the nickel-based alloy of this embodiment is >990MPa (e.g., 995-1010MPa), the yield strength is >580MPa (e.g., 585-600MPa), the impact AKV is >55J (e.g., 58J or more), the 700°C tensile strength is >860MPa (e.g., 865-880MPa), the yield strength is >530MPa (e.g., 533-540MPa), the elongation and cross-sectional expansion ratio at room temperature and 700°C are both above 30%, and the extrapolated endurance strength at 700°C and 100,000 hours is above 170MPa.

[0119] Embodiment 2:

[0120] This embodiment provides a preparation method of a nickel-based alloy suitable for large-section forgings and a nickel-based alloy. The main chemical components are shown in Table 1. The ingot diameter of this embodiment is Φ1300 mm, and the final nickel-based alloy forging diameter is Φ1300 mm.

[0121] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0122] (2) Step 2, homogenizing the ingot: keeping the temperature at 1220° C. for 120 h;

[0123] (3) Step 3, forging: four upsetting and four drawing, forging temperature range 1190~900℃; details are as follows:

[0124] First fire: upsetting temperature 1190℃, upsetting deformation 50%, upsetting reduction rate: 80mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0125] Second fire: drawing temperature 1180℃, drawing deformation 50%, drawing reduction rate: controlled at 5-50mm / s, single hammer reduction during drawing controlled at ≥25%;

[0126] The third fire: upsetting temperature 1150℃, upsetting deformation 45%, upsetting reduction rate: 70mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0127] Fourth fire: drawing temperature 1145℃, drawing deformation 45%, drawing reduction rate: controlled at 10-40mm / s, and single hammer reduction during drawing controlled at ≥25%;

[0128] Fifth fire: upsetting temperature 1145℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0129] The sixth fire: the drawing temperature is 1140℃, the drawing deformation is 50%, the drawing reduction rate is controlled at 15-40mm / s, and the single hammer reduction is controlled at 20% during drawing;

[0130] The seventh fire: upsetting temperature 1135℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 5mm / s as the reduction increases;

[0131] Eighth fire: drawing temperature 1120℃, drawing deformation 50%, drawing reduction rate: controlled at 15-40mm / s, and single hammer reduction during drawing controlled at 20%;

[0132] (4) Step 4, heat treatment: one solution treatment + one aging treatment, solution treatment temperature: 1145°C, solution treatment time: 9h, cooling by water cooling; aging treatment temperature: 750°C, aging time: 12h.

[0133] After step 3 of this embodiment, the grain size of the core of the forging is level 4, and the grain size of the edge is level 5.5; after step 4, the grain size of the core is level 1.5, the grain size of the edge is level 1.5, and the level difference is 0; the structure of the forging of this embodiment is a single austenite structure + grain boundary granular M 23 C6 type carbide + a small amount of M6C / MC type carbide + uniformly dispersed γ′ phase.

[0134] The overall room temperature tensile strength of the nickel-based alloy of this embodiment is >970MPa (e.g., 975-1010MPa), the yield strength is >580MPa (e.g., 585-610MPa), the impact AKV is >50J (e.g., above 55J), the 700°C tensile strength is >800MPa (e.g., 815-850MPa), the yield strength is >520MPa (e.g., 525-540MPa), the elongation and cross-sectional shrinkage at room temperature and 700°C are both above 30%, and the extrapolated endurance strength at 700°C and 100,000 hours is above 170MPa.

[0135] Embodiment 3:

[0136] This embodiment provides a preparation method of a nickel-based alloy suitable for large-section forgings and a nickel-based alloy. The main chemical components are shown in Table 1. The ingot diameter of this embodiment is Φ960 mm, and the final nickel-based alloy forging diameter is Φ800 mm.

[0137] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0138] (2) Step 2, homogenizing the ingot: keeping the temperature at 1190° C. for 90 h;

[0139] (3) Step 3, forging: four upsetting and four drawing, as follows:

[0140] First fire: upsetting temperature 1165℃, upsetting deformation 55%, upsetting reduction rate: 90mm / s at the beginning, and slowly reduced to 20mm / s as the reduction increases;

[0141] Second fire: drawing temperature 1165℃, drawing deformation 50%, drawing reduction rate: controlled at 10~70mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0142] The third fire: upsetting temperature 1150℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0143] Fourth fire: drawing temperature 1150℃, drawing deformation 55%, drawing reduction rate: controlled at 20~80mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0144] Fifth fire: upsetting temperature 1145℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0145] The sixth fire: the drawing temperature is 1140℃, the drawing deformation is 50%, the drawing reduction rate is controlled at 30-50mm / s, and the single hammer reduction is controlled at 25% during drawing;

[0146] The seventh fire: upsetting temperature 1130℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0147] Eighth fire: drawing temperature 1110℃, drawing deformation 45%, drawing reduction rate: controlled at 30-50mm / s, and single hammer reduction during drawing controlled at 25%;

[0148] Ninth fire: drawing temperature 1105℃, drawing deformation 20%, drawing reduction rate: controlled at 30-50mm / s, single hammer reduction during drawing controlled at 15%;

[0149] (4) Step 4, heat treatment: one solution treatment + one aging treatment, solution treatment temperature: 1120°C, solution treatment time: 6h, cooling by water cooling. Aging treatment temperature: 755°C, aging time: 9h.

[0150] After step 3 of this embodiment, the grain size of the core of the forging is 4.5, and the grain size of the edge is 6; after step 4, the grain size of the core is 3, and the grain size of the edge is 3.5, with a grade difference of 0.5; the structure of the forging of this embodiment is a single austenite structure + grain boundary granular M 23C6 type carbide + a small amount of M6C / MC type carbide + uniformly dispersed γ′ phase.

[0151] The overall room temperature tensile strength of the nickel-based alloy of this embodiment is >980MPa (e.g., 985-1020MPa), the yield strength is >590MPa (e.g., 592-615MPa), the impact AKV is >60J (e.g., 65J or more), the 700°C tensile strength is >830MPa (e.g., 835-870MPa), the yield strength is >530MPa (e.g., 535-570MPa), the elongation and cross-sectional expansion ratio at room temperature and 700°C are both above 30%, and the extrapolated endurance strength at 700°C and 100,000 hours is above 170MPa.

[0152] The components of Examples 1-3 are shown in Table 1 below, and the properties are shown in Table 2 below. The temperature, deformation amount, and deformation rate during the forging process of Examples 1-3 are all carried out in accordance with the requirements of the present invention, so no obvious cracking occurs except that the local temperature of the blank surface is too low due to the contact between the hammer head and the blank during the forging process, thereby forming unavoidable local surface cracking.

[0153] Comparative Example 1

[0154] This comparative example provides a nickel-based alloy, the components of which are shown in Table 1. The preparation method of this comparative example is the same as that of Example 1.

[0155] The Al content in this comparative example is slightly lower than that in Example 1, and the Ti content is much lower than that in the present invention. After being prepared by the same preparation method as in Example 1, its overall performance is lower than that in Example 1, especially in terms of tensile strength and endurance strength. The specific data are shown in Table 2.

[0156] Comparative Example 2

[0157] This comparative example provides a nickel-based alloy, the components of which are shown in Table 1. The preparation method of this comparative example is the same as that of Example 1.

[0158] The Cr content in this comparative example is significantly higher than that in Example 1, and the contents of other elements are basically the same as those in Example 1. After being prepared by the same preparation method as in Example 1, its short-term tensile properties are not significantly different from those in Example 1, but in terms of impact, due to excessive Cr, M 23 Excessive precipitation of C6 carbides at the grain boundaries reduces the impact properties, and needle-shaped σ phases are easily precipitated during long-term aging, making the material's extrapolated endurance strength at 700°C and 100,000 hours far lower than 160 MPa in Example 1, and can only reach about 80 MPa.

[0159] Comparative Example 3

[0160] This comparative example provides a nickel-based alloy, the components of which are the same as those of Example 1. The preparation method of this comparative example is compared with that of Example 1, and adopts a three-upsetting and three-drawing forging method. Except for the lack of the fourth and fifth tempers, the other processes and steps are the same. Figure 4 The figure shows the core grain size after step 3 of this embodiment. Figure 5 The edge grain size after step 3 is shown. Figure 6 The figure shows the carbide distribution in the longitudinal section after the treatment in step 4. It can be seen that after forging, the grain size in the center is level 4 and the grain size at the edge is level 5. After the treatment in step 4, the grain size in the center is level 2.5 and the grain size at the edge is level 3. From the axial sampling of the forging, it can be seen that there is still a certain distribution of strip carbides along the deformation direction, which has a certain influence on its endurance performance, as shown in Table 2.

[0161] Comparative Example 4

[0162] This comparative example provides a nickel-based alloy, the components are the same as those in Example 1, and the preparation method of this comparative example is compared with that in Example 1, and adopts two-upsetting and two-drawing forging, except that the third to sixth fires are missing, the other processes and steps are the same. After forging, the core grain size is level 3 and the edge grain size is level 5.5. After the treatment of step 4, the core grain size is level 2.5 and the edge grain size is level 4. Figure 7 The figure shows the carbide distribution in the longitudinal section after the treatment in step 4. It can be seen that there is a relatively obvious strip carbide distribution along the deformation direction, which has a certain influence on its durability performance. See Table 2 for details.

[0163] The chemical compositions of the embodiments and comparative examples are shown in Table 1 below.

[0164] Table 1 Main chemical components wt%

[0165] Steel Grade Cr Mo Co B Al Ti C Al / Ti Cr / C Example 1 18.7 9.3 11.3 0.0049 1.22 1.28 0.049 0.95 382 Example 2 18.5 9.8 11.3 0.0051 1.28 1.28 0.047 1.00 394 Example 3 17.1 8.8 11.1 0.0050 1.23 1.27 0.043 0.97 398 Comparative Example 1 18.8 9.1 11.0 0.0045 0.9 0.15 0.045 6 418 Comparative Example 2 27.0 9.0 11.5 0.0045 1.31 1.35 0.049 0.97 551 Comparative Example 3 18.9 9.1 11.5 0.005 1.25 1.31 0.050 0.95 375 Comparative Example 3 18.5 9.5 11.4 0.0046 1.21 1.26 0.049 0.96 378

[0166] The properties of the embodiments and comparative examples are shown in Table 2.

[0167] Table 2 Performance

[0168]

[0169]

[0170] The above description is only a preferred specific implementation manner of the present invention, but the protection scope of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by any technician familiar with the technical field within the technical scope disclosed by the present invention should be covered within the protection scope of the present invention.

Claims

1. A method for preparing a nickel-based alloy suitable for large-section forgings, characterized in that: It includes four-upsetting and four-drawing to prepare the forging blank and solution treatment of the forging blank; Multiple cooling stages during the four-upsetting and four-drawing process; The temperature of the solution treatment is (1080~1100)+25×(4-a), where a is the required grain size level.

2. The preparation method according to claim 1, characterized in that: The four-upsetting and four-drawing method for preparing a forging blank comprises: S301, first fire: upsetting: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%; S302, second fire: drawing: heating temperature 1150 ~ 1200 ℃, deformation greater than 40%; S303, third fire: upsetting: heating temperature 1145 ~ 1150 ℃, deformation greater than 40%; S304, the fourth fire: drawing, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%; S305, the fifth fire: upsetting, heating temperature 1145 ~ 1150 ℃, deformation greater than 40%; S306, the sixth fire: pulling, heating temperature X + 5 ~ X + 15 ℃, deformation greater than 30%; S307, seventh fire: upsetting, heating temperature is within the range of X±5℃, deformation is greater than 40%; S308, the eighth fire: drawing, heating temperature X-5 ~ X-30 ℃, deformation greater than 30%, pressing rate 20 ~ 200mm / s; The value of X is determined using the following formula: X=1115+6000C / Cr.

3. The preparation method according to claim 1, characterized in that: The preparation method further comprises: smelting and pouring to obtain an ingot; and homogenizing the ingot.

4. The preparation method according to claim 3, characterized in that: The process for homogenizing the ingot is: keeping the ingot at a temperature range of 1180 to 1220° C. for 40 to 150 hours.

5. The preparation method according to claim 1, characterized in that: The components of the nickel-based alloy are expressed in mass percentage Including: Cr: 17% ~ 19%; Mo: 8% ~ 10%; Co: 11%~13%; B: 0.003% ~ 0.006%; Al: 1.2% ~ 1.4%; Ti: 1.2%~1.4%; C: 0.03% to 0.05%; the remainder is Ni and some other inevitable impurities.

6. The preparation method according to claim 5, characterized in that: In the composition of the nickel-based alloy, Al / Ti is controlled at 0.9-1.

05.

7. The preparation method according to claim 5, characterized in that: In the composition of the nickel-based alloy, Cr / C is controlled within the range of 350-500.

8. The preparation method according to claim 1, characterized in that: The grain size of the forging billet prepared by four upsetting and four drawing is finer than grade 4.

9. The preparation method according to any one of claims 1 to 8, characterized in that: The large-section forgings include forgings with a diameter of 500 mm or more.

10. A nickel-based alloy, characterized in that: The nickel-based alloy is prepared by the preparation method according to any one of claims 1 to 9.

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