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

Abstract

The invention discloses a nickel-based alloy suitable for a large-section forge piece and a preparation method of 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 nickel-based alloy comprises the following components in percentage by mass: 17-19% of Cr; 8%-10% of Mo; co: 10.9%-13%; b: 0.003% to 0.006%; 1.2% to 1.4% of Al; 1.2% to 1.4% of Ti; 0.03% to 0.05% of C; and the balance of Ni and other inevitable impurities. The nickel-based alloy 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 nickel-based alloy suitable for large-section forgings and a preparation method thereof. Background Art

[0002] With the increasingly urgent demand for clean and efficient thermal power generation, as well as the trend of large-scale high-temperature nuclear power equipment and petrochemical equipment, 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 for steam temperatures exceeding 650°C. The diameter is generally more than 800mm, the medium-pressure rotor is more than 1100mm, the tonnage is more than 10 tons, the demand for ingots is more than 20 tons, and even the ingot is expected to reach 40-60 tons, and the room temperature yield strength is required to exceed 450MPa, and it is required to stably serve 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 nickel-based alloy suitable for large-section forgings and a preparation method thereof, 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 nickel-based alloy suitable for large-section forgings, wherein the components of the nickel-based alloy include, by mass percentage: Cr: 17% to 19%; Mo: 8% to 10%; Co: 10.9% 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.

[0006] Furthermore, in the composition of the nickel-based alloy, Al / Ti is controlled to be 0.9 to 1.05.

[0007] Furthermore, in the composition of the nickel-based alloy, Cr / C is controlled within the range of 350-500.

[0008] Furthermore, the components of the nickel-based alloy include, by mass percentage, Cr: 17.3% to 19%; Mo: 8.3% to 10%; Co: 10.9% to 12%; B: 0.004% to 0.006%; Al: 1.2% to 1.35%; Ti: 1.2% to 1.35%; C: 0.04% to 0.05%; the remainder is Ni and some other inevitable impurities.

[0009] Furthermore, nickel-based alloys can be used to prepare forgings with a diameter of more than 500 mm.

[0010] The present invention also provides a method for preparing the above-mentioned nickel-based alloy suitable for large-section forgings, comprising:

[0011] Step 1, smelting and pouring to obtain an ingot;

[0012] Step 2: Homogenize the ingot:

[0013] 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;

[0014] Step 4: Performing solid solution and aging treatment on the forging billet to obtain a nickel-based alloy.

[0015] Furthermore, in step 3, the grain size of the forging billet is finer than level 4.

[0016] Furthermore, in step 4, the solution treatment temperature is set according to the target grain size, and the solution treatment temperature satisfies: (1080-1100)+25×(4-a), where a is the required grain size level.

[0017] Furthermore, in step 4, the aging treatment temperature is 720-780°C.

[0018] Furthermore, in step 4, the grain size of the obtained nickel-based alloy is 1 to 4 levels, the grains are uniform, and the grain size difference from the center to the edge is less than 2 levels.

[0019] Compared with the prior art, the present invention has the following beneficial effects:

[0020] The nickel-based alloy suitable for large-section (e.g., above Φ500 mm) forgings of the present invention ensures a light degree of segregation by precisely controlling the types and contents of elements in the alloy and by synergistically controlling the synergistic quantitative relationship of different elements, combined with a precisely controlled preparation method, and is suitable for smelting large ingots, thereby manufacturing large-tonnage and large-section forgings.

[0021] The preparation method of the nickel-based alloy of the present invention ensures the uniformity of the grains by accurately 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 2 levels (for example, less than 1 level); at the same time, it also reduces the cracking tendency and improves the yield of the preparation of large-section forgings (for example, more than Φ500mm).

[0022] The nickel-based alloy of the present invention has an overall room temperature tensile strength of >950MPa (e.g. 960-1030MPa), a yield strength of >550MPa (e.g. 570-620MPa), an impact AKV of >35J (e.g. 50-80J), a 700°C tensile strength of >750MPa (e.g. 810-870MPa), a yield strength of >480MPa (e.g. 515-560MPa), an elongation and a cross-sectional expansion ratio at room temperature and 700°C of more than 25%, and a 100,000-hour endurance strength of more than 150MPa at 700°C. The alloy can meet the requirements of long-term stable service at 700°C.

[0023] 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

[0024] 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.

[0025] Figure 1 The core grain size after the treatment in step 3 of Example 1;

[0026] Figure 2 The edge grain size after the treatment in step 3 of Example 1;

[0027] Figure 3 The core grain size after the treatment in step 4 of Example 1;

[0028] Figure 4 The edge grain size after the treatment in step 4 of Example 1;

[0029] Figure 5 The core grain size after the step 3 of Example 2;

[0030] Figure 6 The edge grain size after the treatment in step 3 of Example 2;

[0031] Figure 7 The core grain size after the treatment in step 4 of Example 2;

[0032] Figure 8 The edge grain size after the treatment in step 4 of Example 2;

[0033] Fig. 9 The core grain size after the treatment in step 3 of Example 3;

[0034] Fig.10The edge grain size after the treatment in step 3 of Example 3;

[0035] Fig.11 The core grain size after the treatment in step 4 of Example 3;

[0036] Fig.12 The edge grain size after the treatment in step 4 of Example 3;

[0037] Fig.13 The figure shows the intracrystalline γ′ phase of the present invention;

[0038] Fig.14 The fine particles M in the grain boundary and grain of the present invention are shown. 23 C6 type carbide. DETAILED DESCRIPTION

[0039] 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.

[0040] The invention provides a nickel-based alloy suitable for large-section forgings. The components of the nickel-based alloy include, by mass percentage, Cr: 17%-19%; Mo: 8%-10%; Co: 10.9%-13%; B: 0.003%-0.006%; Al: 1.2%-1.4%; Ti: 1.2%-1.4%; C: 0.03%-0.05%; the balance is Ni and some other inevitable impurities.

[0041] Specifically, the content ratio of Al to Ti (ie, Al / Ti) is controlled to be 0.9 to 1.05.

[0042] Specifically, the content ratio of Cr to C (ie, Cr / C) is controlled within the range of 350-500.

[0043] The following is a detailed description of the effects and dosage of the components in the present invention:

[0044] 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 23The 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%.

[0045] 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 23 Stability 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.

[0046] 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%.

[0047] 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.

[0048] 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.

[0049] 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.

[0050] After in-depth research, the inventors have controlled the C and Cr contents in 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 lower level; 2. The nickel-based alloy of the present invention is aimed at the product - ultra-supercritical turbine rotor, which has very high requirements for endurance performance, and larger grains are beneficial to the endurance performance of forgings. Generally, it will be controlled at levels 1 to 3, so its grain boundary area is small. If the C and Cr contents are too high, the M6C carbides precipitated at the grain boundaries will be 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.

[0051] 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.

[0052] Other elements: Elements such as Si, S, O, etc. must be controlled at the lowest level possible.

[0053] 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: 17.3% to 19%; Mo: 8.3% to 10%; Co: 10.9% to 12%; B: 0.004% to 0.006%; Al: 1.2% to 1.35%; Ti: 1.2% to 1.35%; C: 0.04% to 0.05%; the balance is Ni and some other inevitable impurities.

[0054] Specifically, the content ratio of Al to Ti (ie, Al / Ti) is controlled to be between 0.92 and 1.05.

[0055] Specifically, the content ratio of Cr to C (ie, Cr / C) is controlled at 370-450.

[0056] Specifically, the nickel-based alloy suitable for large-section forgings can be used to prepare forgings with a diameter of 500 mm or more (eg, 750 mm or more).

[0057] The present invention also provides a method for preparing the above nickel-based alloy, comprising:

[0058] Step 1, smelting and pouring to obtain an ingot;

[0059] Step 2: homogenize the ingot: keep it at a temperature of 1180-1220°C for 40-150h;

[0060] 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;

[0061] Step 4: Performing solid solution and aging treatment on the forging billet to obtain a nickel-based alloy.

[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 can adopt three upsetting and three drawing, and the specific steps of three upsetting and three 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 (X + 5) ~ (X + 15) ℃, deformation greater than 40%, single hammer reduction greater than 20%, reduction rate 10 ~ 200mm / s; wherein the value of X is determined by the following formula: X = 1115 + 6000C / Cr;

[0069] S305, fifth fire: upsetting, heating temperature X ± 5 ° C, deformation greater than 40%, pressing rate 10 ~ 200mm / s;

[0070] S306, the sixth fire: drawing, heating temperature X-5 ~ X-30 ℃, deformation greater than 30%, single hammer pressing amount to be controlled between 15% ~ 25%, pressing rate 20 ~ 200mm / s.

[0071] 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.

[0072] 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.

[0073] 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.

[0074] 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.

[0075] 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%.

[0076] 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℃.

[0077] 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.

[0078] Specifically, through the operation of S304, the grains inside the forging can be further refined than the third fire, and a good foundation is laid for the continued grain refinement of the next fire. Specifically, in S304, the heating temperature can be 1135-1150°C. Specifically, in S304, the pressing rate is 10-80 mm / s.

[0079] Specifically, in S305, 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.

[0080] Specifically, in S305, 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.

[0081] Specifically, in S306, the heating temperature can be 1110-1125° C. In S306, the temperature is further reduced, and the deformation is selected to be greater than 30% in order to further refine the grains and ensure that the overall grain size of the final forging is finer than level 4. The single hammer reduction is controlled between 15% and 25% to ensure that the critical deformation is reached and the grains do not grow abnormally.

[0082] Specifically, in S306, due to the low forging heating temperature, the product is generally more complex in shape and may not reach the final size in this firing. A seventh firing is required, and the forging heating temperature of the seventh firing should be reduced by 5 to 10°C based on the heating temperature of the sixth firing.

[0083] Specifically, after S306 treatment, the grain size of the forging is finer than grade 4.

[0084] 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 satisfies: (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), ensuring that the 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.

[0085] 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.

[0086] 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.

[0087] 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 2 levels (for example, less than 1 level); the microstructure is austenite structure + fine granular M 23 C6 carbide + individual M6C and MC carbides at grain boundaries + nano-scale γ′ phase evenly distributed in the grains.

[0088] Specifically, the overall room temperature tensile strength of the nickel-based alloy after the treatment in step 4 is greater than 950 MPa (e.g., 960-1030 MPa), the yield strength is greater than 550 MPa (e.g., 570-620 MPa), the impact AKV is greater than 35 J (e.g., 50-80 J), the tensile strength at 700°C is greater than 750 MPa (e.g., 810-870 MPa), the yield strength is greater than 480 MPa (e.g., 515-560 MPa), the elongation and cross-sectional expansion ratio at room temperature and 700°C are both greater than 25%, and the endurance strength at 700°C and 100,000 hours is greater than 150 MPa.

[0089] Compared with the prior art, the nickel-based alloy suitable for large-section forgings of the present invention ensures a light degree of segregation by precisely controlling the types and contents of elements in the alloy and by synergistically controlling the synergistic quantitative relationship of different elements, combined with a precisely controlled preparation method, and is suitable for smelting large ingots, thereby manufacturing large-tonnage and large-section forgings.

[0090] The preparation method of the nickel-based alloy of the present invention ensures the uniformity of the grains by accurately 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 2 levels (for example, less than 1 level); at the same time, it also reduces the cracking tendency and improves the yield of the preparation of large-section forgings (for example, more than Φ500mm).

[0091] The nickel-based alloy of the present invention has an overall room temperature tensile strength of >950MPa (e.g. 960-1030MPa), a yield strength of >550MPa (e.g. 570-620MPa), an impact AKV of >35J (e.g. 50-80J), a 700°C tensile strength of >750MPa (e.g. 810-870MPa), a yield strength of >480MPa (e.g. 515-560MPa), an elongation and a cross-sectional expansion ratio at room temperature and 700°C of more than 25%, and a 100,000-hour endurance strength of more than 150MPa at 700°C. The alloy can meet the requirements of long-term stable service at 700°C.

[0092] The advantages of precise control of the composition and process parameters of the nickel-based alloy of the present invention are demonstrated below with specific examples and comparative examples. The embodiments of the present invention all provide a nickel-based alloy and a preparation method thereof.

[0093] Embodiment 1:

[0094] This embodiment provides a nickel-based alloy suitable for large-section forgings and a preparation method thereof. The main chemical components are shown in Table 1. The ingot diameter of this embodiment is Φ900 mm, and the final nickel-based alloy forging diameter is Φ800 mm.

[0095] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0096] (2) Step 2, homogenizing the ingot: keeping the ingot at a temperature of 1200° C. for 100 h;

[0097] (3) Step 3, forging: three upsetting and three drawing, forging temperature range 1170~900℃; details are as follows:

[0098] 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;

[0099] Second fire: drawing temperature 1170℃, drawing deformation 60%, drawing reduction rate: controlled at 5-50mm / s, single hammer reduction during drawing controlled at ≥30%;

[0100] 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;

[0101] Fourth fire: drawing temperature 1140℃, drawing deformation 60%, drawing reduction rate: controlled at 15-50mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0102] Fifth fire: upsetting temperature 1130℃, upsetting deformation 60%, upsetting reduction rate: 100mm / s at the beginning, and then slowly reduced to 10mm / s as the reduction increases;

[0103] The sixth fire: the drawing temperature is 1125℃, the drawing deformation is 65%, the drawing reduction rate is controlled at 15-50mm / s, and the single hammer reduction is controlled at 20% during drawing;

[0104] (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;

[0105] 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 core grain size after step 4 treatment. Figure 4The figure shows the edge grain size after step 4. It can be seen that 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. After step 4, the grain size of the core is level 2.5 and the grain size of the edge is level 3, with a level difference of 0.5. The structure of the forging of this embodiment is a single austenite structure + fine grained M at the grain boundary. 23 C6 type carbide + a small amount of M6C / MC type carbide + uniformly dispersed γ′ phase. Fig.13 The γ′ phase in the crystal is shown; Fig.14 The fine particles M in the grain boundary and grain 23 C6 type carbide.

[0106] The overall room temperature tensile strength of the forging of this embodiment is >960MPa, the yield strength is >570MPa, the impact AKV is >45J, the 700℃ tensile strength is >810MPa, the yield strength is >520MPa, the elongation and cross-sectional expansion ratio at room temperature and 700℃ are both above 30%, and the extrapolated endurance strength at 700℃ and 100,000 hours is higher than 160MPa.

[0107] Embodiment 2:

[0108] This embodiment provides a nickel-based alloy suitable for large-section forgings and a preparation method thereof. The main chemical components are shown in Table 1. The ingot diameter of this embodiment is Φ1100 mm, and the final nickel-based alloy forging diameter is Φ1000 mm.

[0109] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0110] (2) Step 2, homogenizing the ingot: keeping the temperature at 1195°C for 120h;

[0111] (3) Step 3, forging: three upsetting and three drawing, forging temperature range 1180~900℃; details are as follows:

[0112] First fire: upsetting temperature 1180℃, upsetting deformation 50%, upsetting reduction rate: 70mm / s at the beginning, and slowly reduced to 3mm / s as the reduction increases;

[0113] Second fire: drawing temperature 1180℃, drawing deformation 50%, drawing reduction rate: controlled at 3-40mm / s, single hammer reduction during drawing controlled at ≥25%;

[0114] 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;

[0115] Fourth fire: drawing temperature 1145℃, drawing deformation 45%, drawing reduction rate: controlled at 10-40mm / s, and the single hammer reduction during drawing is controlled at ≥25%;

[0116] Fifth 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;

[0117] The sixth fire: the drawing temperature is 1120℃, the drawing deformation is 65%, the drawing reduction rate is controlled at 15-40mm / s, and the single hammer reduction is controlled at 20% during drawing;

[0118] (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.

[0119] Figure 5 The figure shows the core grain size after step 3 of this embodiment. Figure 6 The edge grain size after step 3 is shown. Figure 7 The figure shows the core grain size after step 4 treatment. Figure 8 The figure shows the edge grain size after step 4. It can be seen that 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 6. After step 4, the grain size of the core is level 1.5 and the grain size of the edge is level 1.5, with a grade difference of 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.

[0120] The overall room temperature tensile strength of the forging is >950MPa, the yield strength is >565MPa, the impact AKV is >45J, the 700℃ tensile strength is >760MPa, the yield strength is >490MPa, the elongation and cross-sectional expansion ratio at room temperature and 700℃ are both above 30%, and the extrapolated endurance strength at 700℃ and 100,000 hours is higher than 150MPa.

[0121] Embodiment 3:

[0122] This embodiment provides a nickel-based alloy suitable for large-section forgings and a preparation method thereof. The main chemical components are shown in Table 1. The ingot diameter of this embodiment is Φ810 mm, and the final nickel-based alloy forging diameter is Φ780 mm.

[0123] (1) smelting and pouring according to the alloy composition to obtain an ingot;

[0124] (2) Step 2, homogenizing the ingot: keeping the temperature at 1190° C. for 80 h;

[0125] (3) Step 3, forging: three upsetting and three drawing, forging temperature range 1160~920℃; details are as follows:

[0126] First fire: upsetting temperature 1155℃, upsetting deformation 50%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0127] Second fire: drawing temperature 1155℃, drawing deformation 50%, drawing reduction rate: controlled at 10~70mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0128] The third fire: upsetting temperature 1145℃, upsetting deformation 55%, upsetting reduction rate: 100mm / s at the beginning, and slowly reduced to 10mm / s as the reduction increases;

[0129] Fourth fire: drawing temperature 1140℃, drawing deformation 55%, drawing reduction rate: controlled at 20~80mm / s, and single hammer reduction during drawing controlled at ≥30%;

[0130] Fifth fire: upsetting temperature 1130℃, upsetting deformation 60%, upsetting reduction rate: 100mm / s at the beginning, and then slowly reduced to 10mm / s as the reduction increases;

[0131] The sixth fire: the drawing temperature is 1110℃, the drawing deformation is 60%, the drawing reduction rate is controlled at 30-50mm / s, and the single hammer reduction is controlled at 25% during drawing;

[0132] (4) Step 4, heat treatment: one solution treatment + one aging treatment, solution treatment temperature: 1120°C, solution treatment time: 4.5h, cooling by water cooling. Aging treatment temperature: 765°C, aging time: 8h.

[0133] Fig. 9 The figure shows the core grain size after step 3 of this embodiment. Fig.10 The edge grain size after step 3 is shown. Fig.11 The figure shows the core grain size after step 4 treatment. Fig.12 The figure shows the edge grain size after step 4. It can be seen that 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 6.5. After step 4, the grain size of the core is level 2.5, and the grain size of the edge is level 3.5, with a level difference of 1. 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 forging is >950MPa, the yield strength is >565MPa, the impact AKV is >45J, the 700℃ tensile strength is >760MPa, the yield strength is >490MPa, the elongation and cross-sectional expansion ratio at room temperature and 700℃ are both above 30%, and the extrapolated endurance strength at 700℃ and 100,000 hours is higher than 150MPa.

[0135] 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.

[0136] Comparative Example 1

[0137] 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.

[0138] 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.

[0139] Comparative Example 2

[0140] 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.

[0141] 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.

[0142] Comparative Example 3

[0143] This comparative example provides a nickel-based alloy, the components are the same as those in Example 1. The preparation method of this comparative example is compared with Example 1, but the step-down forging method is not used, but the forging heating temperature is always fixed at 1150°C, and the other steps are the same. After forging, the core grain size is level 2, and the edge grain size is level 6.5. After solution treatment at 1130°C, the core grain size is level 1, and the edge grain size is level 3.5. There is an obvious mixed crystal phenomenon, which will have a certain impact on the durability and fatigue performance of the entire forging, as shown in Table 2.

[0144] Comparative Example 4

[0145] This comparative example provides a nickel-based alloy, the components are the same as those in Example 1. In the preparation method of this comparative example, compared with the example, except for the solution temperature and time selection, the other operations are the same. In this comparative example, the solution temperature is selected as 1080°C and the time is selected as 3h. The final grain size is 4 in the center and 6 at the edge, and the carbides in the grain boundaries are unevenly distributed and large in size, which has a certain impact on the performance of the entire forging, as shown in Table 2.

[0146] The chemical compositions of the embodiments and comparative examples are shown in Table 1 below.

[0147] Table 1 Main chemical components wt%

[0148] Steel Grade Cr Mo Co B Al Ti C Al / Ti Cr / C Example 1 18.9 9.1 11.5 0.005 1.25 1.31 0.05 0.95 375 Example 2 18.3 9.9 11.3 0.005 1.29 1.26 0.045 1.02 407 Example 3 17.3 8.6 10.9 0.0048 1.29 1.25 0.046 1.03 376 Comparative Example 1 19.0 9.1 11.1 0.0046 1.0 0.1 0.044 10 432 Comparative Example 2 25 9.2 11.4 0.0048 1.3 1.33 0.045 0.98 556 Comparative Example 3 18.6 9.0 11.2 0.005 1.26 1.25 0.047 1.01 395 Comparative Example 4 18.7 9.5 11.3 0.0049 1.27 1.28 0.048 0.99 398

[0149] The properties of the embodiments and comparative examples are shown in Table 2.

[0150] Table 2 Performance

[0151]

[0152]

[0153] 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 nickel-based alloy suitable for large-section forgings, characterized in that: The components of the nickel-based alloy are expressed in mass percentage Including: Cr: 17%~19%; Mo: 8%~10%; Co: 10.9%~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.

2. The nickel-based alloy suitable for large-section forgings according to claim 1, characterized in that: In the composition of the nickel-based alloy, Al / Ti is controlled at 0.9-1.

05.

3. The nickel-based alloy suitable for large-section forgings according to claim 1, characterized in that: In the composition of the nickel-based alloy, Cr / C is controlled within the range of 350-500.

4. The nickel-based alloy suitable for large-section forgings according to claim 1, characterized in that: The components of the nickel-based alloy include, by mass percentage, Cr: 17.3% to 19%; Mo: 8.3% to 10%; Co: 10.9% to 12%; B: 0.004% to 0.006%; Al: 1.2% to 1.35%; Ti: 1.2% to 1.35%; C: 0.04% to 0.05%; the remainder is Ni and some other inevitable impurities.

5. The nickel-based alloy suitable for large-section forgings according to any one of claims 1 to 4, characterized in that: The nickel-based alloy can be used to prepare forgings with a diameter of more than 500 mm.

6. A method for preparing a nickel-based alloy suitable for large-section forgings according to any one of claims 1 to 5, characterized in that: include: Step 1, smelting and pouring to obtain an ingot; Step 2: Homogenize the ingot: 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; Step 4: Performing solid solution and aging treatment on the forging billet to obtain a nickel-based alloy.

7. The preparation method according to claim 6, characterized in that: In the step 3, the grain size of the forging blank is finer than level 4.

8. The preparation method according to claim 6, characterized in that: In step 4, the solution treatment temperature is set according to the target grain size, and the solution treatment temperature satisfies: (1080-1100)+25×(4-a), where a is the required grain size level.

9. The preparation method according to claim 6, characterized in that: In the step 4, the aging treatment temperature is 720-780°C.

10. The preparation method according to any one of claims 6 to 9, characterized in that: In the step 4, the grain size of the obtained nickel-based alloy is 1 to 4 levels, the grains are uniform, and the grain size difference from the center to the edge is less than 2 levels.

Images