CN117363955A - Multi-type precipitated phase cooperative strengthening heat-resistant alloy and preparation method thereof - Google Patents

Abstract

The invention discloses a multi-type precipitation phase synergistically strengthened heat-resistant alloy and a preparation method thereof, and belongs to the technical field of heat-resistant alloys. The chemical composition mass percentage of the alloy is: C 0.01~0.04%, Si 0.1~0.4%, Mn 0.2~0.6%, Mo 0.8~1.6%, Cr 12~17%, Ni 28~33%, Ti 1.5~2.5% , Al 0.5～2.0%, Nb 0.2～0.8%, N＜0.004%, P＜0.005%, S＜0.004%, the balance is Fe. The advantage of the heat-resistant alloy of the present invention is to achieve synergistic strengthening through multiple types of nano-strengthening phases (γ' phase, (Ti, Nb)C and Fe ₂ Ti type Laves phase) with a volume fraction of more than 30%; the prepared heat-resistant alloy has high temperature tensile strength The strength is not less than 486Mpa, which meets the requirements for the use of heat-resistant alloy materials such as automobile engine exhaust valves and fasteners. It is also suitable for manufacturing aircraft engine load-bearing parts and gas turbine heat-resistant parts.

Description

A multi-type precipitation phase synergistically strengthened heat-resistant alloy and its preparation method

Technical field

The invention belongs to the technical field of heat-resistant alloys and relates to a multi-type precipitation phase synergistically strengthened heat-resistant alloy and a preparation method thereof, which is suitable for use in automobile engine exhaust valves.

Background technique

The automobile industry is one of my country's most important pillar industries, and its total automobile production and sales have ranked first in the world for 14 consecutive years. Fuel vehicles are still the mainstream of my country's automobile market, and their sales volume and proportion are still much higher than those of pure electric vehicles. Automobile carbon emissions account for about 7.5% of the total carbon emissions of society. The large number of fuel vehicles, fossil fuel combustion and low combustion efficiency during use are the main factors causing high carbon emissions from vehicles. In recent years, with the continuous improvement of the requirements for energy conservation and emission reduction, higher requirements have been put forward for the combustion efficiency of automobile engines. Valve steel is a key material for automobile engines. my country's annual valve output has exceeded 500 million (equivalent to about 48,000 tons of raw materials). The automobile engine exhaust valve discharges high-temperature corrosive exhaust gas, so the valve is subject to high temperature and high pressure. Exhaust valve materials are required to have excellent high-temperature strength, toughness, hardness, wear resistance, oxidation resistance and corrosion resistance at working temperature, as well as structural stability and dimensional stability under the alternating working conditions of hot and cold engines. At the same time, the valve material should have good hot and cold processing and welding properties during processing.

At present, the widely used valve alloy grades include high alloy steel 21-4N and 21-4NWNb, nickel-based high-temperature alloy GH4751 and Nimonic 80A, with operating temperatures between 680 and 820°C. However, it is difficult for the above-mentioned valve alloy materials to achieve a good match between high service temperature and high high-temperature strength, and they do not have the advantage of low cost. With the widespread application of technologies to improve combustion efficiency such as direct injection and turbocharging in automobile internal combustion engines, exhaust valve materials are required to have higher resistance to high-temperature exhaust gas corrosion, oxidation resistance and high-temperature strength. The current automobile exhaust valve steel materials 21-4N and 21-4NWNb cannot meet the temperature requirements of the engine combustion chamber of 700°C and above. The cost of using nickel-based valve alloy is very high. my country's high-performance valve alloys for automobile internal combustion engines that serve at temperatures between 680 and 760 degrees Celsius all rely on imports.

Therefore, there is an urgent need to develop new valve alloy materials that combine high performance and low cost. On the one hand, this is of great strategic significance to solve the problem of material dependence on imports, and at the same time, it also provides guarantee for meeting the demand for valve alloys with higher emission standards in the future.

Contents of the invention

In order to solve the above problems, the technical solution of the present invention provides a multi-type precipitation phase synergistically strengthened heat-resistant alloy and a preparation method thereof. By optimizing the design of alloy components and formulating a reasonable production process, a high-strength heat-resistant alloy can be obtained. After electroslag remelting, homogenization treatment, forging, high-temperature solid solution and aging heat treatment, the heat-resistant alloy prepared has a grain size of 5-7, and fine and dispersed multi-type strengthening phases (γ' phase, ( Ti, Nb)C and Fe ₂ Ti type Laves phase). Among them, the γ' phase strengthening phase has a spherical morphology with a volume fraction of 25% to 35%, a volume fraction of small block (Ti, Nb)C is 5% to 8%, and the Fe ₂ Ti type Laves phase is intermittently distributed along the grain boundaries. The volume fraction is 2% to 5%.

According to the first aspect of the technical solution of the present invention, a method for preparing a heat-resistant alloy synergistically strengthened by multiple types of precipitation phases is provided, which includes the following steps:

(1) Smelting: vacuum induction furnace smelting → casting electrode rods → inert atmosphere protected electroslag remelting, or smelting using electric arc furnace + LF + VD + electroslag remelting method to produce electroslag ingots, which are then hot-sent for annealing;

(2) Homogenization treatment: The homogenization treatment adopts a two-stage heat preservation process. The first stage of homogenization treatment temperature is 900~1050℃, and the heat preservation time is 2~10h, so that the γ-γ' eutectic phase and Laves phase can be fully redissolved. ; The second stage of homogenization treatment temperature is 1100 ~ 1150 ℃, and the holding time is 2 ~ 16 hours, so that (Ti, Nb) C can be fully dissolved, and Ti and Nb elements can be evenly diffused to eliminate element segregation; after the homogenization treatment is completed, furnace After cooling to 1030℃, it is released from the furnace and air-cooled;

(3) Forging: The electroslag ingot is kept at 1150~1180℃ and then forged; the starting forging temperature is 1150~1180℃, the final forging temperature is not less than 950℃, and the forging ratio is 3~4;

(4) Solid solution and aging: The forged alloy billet undergoes high-temperature solution treatment + aging treatment. The high-temperature solution temperature is 950~1100℃, and the holding time is 0.5~6 hours; the aging treatment temperature is 580~780℃, and the heat preservation time is Time 4~32 hours.

Furthermore, electroslag remelting is carried out in an argon protective atmosphere. In order to prevent Ti burning during the electroslag process and ensure good surface quality of the electroslag ingot, a special slag system is used. The composition of the slag system in terms of mass percentage is: CaF ₂ : 50 to 55%, CaO: 15 to 25%, Al ₂ O ₃ : 15 to 25%, MgO: 1 to 4%, TiO ₂ : 2 to 5%, FeO ≤ 0.5%, SiO ₂ ≤ 0.8% .

Further, for the 50-200kg ingot type: the voltage is 25-30V and the current is 1500-3000A when the remelting is stable; for the 200-500kg ingot type: the voltage is 30-35V and the current is 3000-4000A when the remelting is stable; 500kg～1t ingot type: when remelting is stable, the voltage is 35～40V and the current is 4000～5000A.

Furthermore, after the electroslag remelting is completed, the mold is demoulded, and a stainless steel insulation cover is used to cover the electroslag ingot and slowly cool it. The cooling time is more than 5 hours, which effectively prevents surface cracking of the electroslag ingot.

Furthermore, the structural characteristics of the alloy after aging heat treatment are as follows: the grain size is 5-7, spherical γ' phase precipitates in the grains (size is not more than 100nm, volume fraction is 25% to 35%), and small blocks precipitate at the grain boundaries ( Ti, Nb)C (size is not larger than 200nm, volume fraction is 5% to 8%) and spherical or small block Fe ₂ Ti type Laves phase (size is not larger than 200nm, volume fraction is 2% to 5%, Laves phase Composed of Fe, Ti, Nb, Ni, Si, Mo elements); among them, the intragranular nanoscale γ' phase strengthening phase interacts with dislocations, through coherent strain strengthening, "Orowan" bypass mechanism strengthening, and dislocation cutting Mechanism strengthening and dislocation climbing mechanism strengthening, the above mechanisms all work, and no matter which mechanism works, the number of γ' phase strengthening phases is the fundamental influencing factor; nanoscale (Ti, Nb) C precipitated at the grain boundary And the Fe ₂ Ti type Laves phase hinders dislocation movement, inhibits grain boundary migration and grain growth, and can improve the strength and toughness of the area near the grain boundary. At the same time, the existence of nanoscale (Ti, Nb)C and Fe ₂ Ti type Laves phases refines the alloy grains and achieves the effect of grain refinement and strengthening. After aging treatment, finely dispersed nanometer γ' phase, (Ti, Nb)C and Fe ₂ Ti type Laves phase are precipitated to achieve synergistic strengthening of multiple types of precipitated phases.

Furthermore, the mechanical properties of the alloy after aging heat treatment meet: 760°C high temperature mechanical properties: yield strength R _p0.2 ≥ 432Mpa; tensile strength R _m ≥ 486Mpa.

According to the second aspect of the technical solution of the present invention, a multi-type precipitate phase synergistically strengthened heat-resistant alloy is provided, and the multi-type precipitate phase synergistically strengthened heat-resistant alloy is prepared by the preparation method according to any of the above aspects,

Among them, the chemical composition of the multi-type precipitation phase synergistically strengthened heat-resistant alloy includes, in terms of mass percentage: C0.01~0.04%, Si 0.1~0.4%, Mn 0.2~0.6%, Mo 0.8~1.6%, Cr 12 ~17%, Ni28~33%, Ti 1.5~2.5%, Al 0.5~2.0%, Nb 0.2~0.8%, N<0.004%, P<0.005%, S<0.004%, the balance is Fe.

Further, the chemical composition of the multi-type precipitation phase synergistically strengthened heat-resistant alloy includes, in mass percentage: C 0.02%, Si 0.2%, Mn 0.4%, Mo 1.2%, Cr 15%, Ni 30%, Ti 1.9 %, Al0.6%, Nb0.5%, N＜0.004%, P＜0.005%, S＜0.004%, and the balance is Fe.

Further, the chemical composition of the multi-type precipitation phase synergistically strengthened heat-resistant alloy includes, in mass percentage: C 0.02%, Si 0.2%, Mn 0.4%, Mo 1.2%, Cr 15%, Ni 30%, Ti 2.0 %, Al1.4%, Nb0.5%, N＜0.004%, P＜0.005%, S＜0.004%, and the balance is Fe.

Further, the chemical composition of the multi-type precipitation phase synergistically strengthened heat-resistant alloy includes, in mass percentage: C 0.02%, Si 0.2%, Mn 0.4%, Mo 1.2%, Cr 15%, Ni 30%, Ti 2.4 %, Al0.6%, Nb0.5%, N＜0.004%, P＜0.005%, S＜0.004%, and the balance is Fe.

Beneficial effects of the present invention:

(1) A kinetic model of ingot homogenization was established to guide the formulation of high-temperature homogenization process. The structure of the ingot before homogenization treatment includes brittle γ-γ' eutectic phase, brittle Laves phase and large-sized (Ti, Nb)C.

The electroslag ingot undergoes a two-stage homogenization treatment to fully dissolve the brittle γ-γ' eutectic phase and Laves phase, and fully dissolve (Ti, Nb)C, effectively avoiding void defects inside the alloy and controlling the grain size. Not overgrown. After the two-stage homogenization treatment, the residual segregation coefficient δ of Ti and Nb elements is <0.2. The Ti and Nb elements have been fully diffused and evenly distributed, which improves the hot working plasticity of the alloy. The calculation formula of element residual segregation coefficient is as follows:

in and/> are respectively the minimum and maximum concentrations of elements in the ingot,/> and/> are the minimum and maximum concentrations of elements in the steel ingot after high-temperature diffusion annealing, respectively.

(2) Adjust the sum of the Ti+Al+Nb content and the Ti/Al ratio to control the volume fraction, size, reverse phase domain boundary energy and mismatch degree of the γ' phase with the substrate. The sum of Ti+Al+Nb content is 3% to 5%

Ensure that the volume fraction of the γ' phase is 25% to 35%, and add 0.5% Nb element to enhance the γ' phase -

Thermal stability of Ni ₃ (Ti, Al, Nb). Add 1.2% Mo to slow down the diffusion rate of Ti and Al elements, and adjust the Ti/Al ratio as much as possible to avoid the formation of harmful eta phase after long-term aging. The intragranular nanoscale γ' phase strengthening phase is strengthened by coherent strain strengthening, bypass mechanism strengthening, dislocation cutting mechanism strengthening and dislocation climbing mechanism. Nanoscale (Ti, Nb) C and Fe ₂ Ti types precipitated at the grain boundary The Laves phase improves the strength and toughness of the area near the grain boundary and achieves a synergistic strengthening effect of multiple types of precipitates.

(3) Strictly control the contents of C, Ti, Nb, Mo, and Si to avoid the continuous precipitation of (Ti, Nb)C and Fe ₂ Ti type Laves phases at the grain boundaries, thereby ensuring the strength and toughness of the grain boundaries. In the alloy of the present invention, the (Ti, Nb) C and Fe ₂ Ti type Laves phases are intermittently distributed along the grain boundaries. The volume fraction of (Ti, Nb) C is 5% to 8%, and the volume fraction of the Fe ₂ Ti type Laves phase is 2% to 2%. 5%.

Description of the drawings

In order to explain the embodiments of the present invention or the technical solutions in the prior art more clearly, the drawings needed to be used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings in the following description are only These are some embodiments of the present invention. For those of ordinary skill in the art, other drawings can be obtained based on the structures shown in these drawings without exerting creative efforts.

Figure 1 is a flow chart of a method for preparing a heat-resistant alloy synergistically strengthened by multiple types of precipitation phases according to the technical solution of the present invention;

Figure 2 shows the equilibrium phase precipitation results in the heat-resistant alloy of the present invention;

Figure 3 shows the nanoscale multi-type precipitated phases (γ' phase, (Ti, Nb)C and Fe ₂ Ti type Laves phase) and matrix structure of the heat-resistant alloy of the present invention after solid solution aging at -770°C for 4 hours;

Figure 4 shows the nanoscale multi-type precipitated phases (γ' phase, (Ti, Nb)C and Fe ₂ Ti type Laves phase) and matrix structure of the heat-resistant alloy of the present invention after solid solution aging at -710°C for 28 hours;

Figure 5 shows the nanoscale multi-type precipitated phases (γ' phase, (Ti, Nb)C and Fe ₂ Ti type Laves phase) and matrix structure of the heat-resistant alloy of the present invention after solid solution aging at -740°C for 4 hours.

Detailed ways

The implementation of the present invention will be described in detail below with reference to the embodiments of the present invention, so that those in the art can better understand the advantages and features of the present invention. Obviously, the embodiments described below are only some of the embodiments of the present invention, rather than all of the embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those of ordinary skill in the art without any creative work fall within the scope of protection of the present invention.

The invention discloses a multi-type precipitation phase synergistically strengthened heat-resistant alloy and a preparation method thereof, and belongs to the technical field of heat-resistant alloys. The chemical composition mass percentage of the alloy is: C 0.01~0.04%, Si 0.1~0.4%, Mn 0.2~0.6%, Mo 0.8~1.6%, Cr 12~17%, Ni 28~33%, Ti 1.5~2.5% , Al 0.5～2.0%, Nb0.2～0.8%, N＜0.004%, P＜0.005%, S＜0.004%, the balance is Fe. Vacuum induction melting + atmosphere-protected electroslag remelting double process is used for smelting. The electroslag ingot undergoes two-stage homogenization treatment at 900-1150°C. After the homogenization treatment is completed, the furnace is cooled to 1030°C and then air-cooled out of the furnace. The electroslag ingot after homogenization treatment is kept at 1150~1180℃ and then forged; the starting forging temperature is 1150~1180℃, the final forging temperature is not less than 950℃, the forging ratio is 3~4; the forged blank is at 950~1100 Solution treatment at ℃ for 0.5~6 hours; and aging treatment at 580~780℃ for 4~32 hours. As a result, synergistic strengthening is achieved through multiple types of nano-strengthening phases (γ' phase, (Ti, Nb) C and Fe ₂ Ti type Laves phase) with a volume fraction greater than 30%; the high-temperature tensile strength of the prepared heat-resistant alloy is not less than 486Mpa, which meets the requirements for the use of heat-resistant alloy materials such as automobile engine exhaust valves and fasteners. It is also suitable for manufacturing aircraft engine load-bearing parts and gas turbine heat-resistant parts.

Specifically, the technical solution of the present invention first provides a method for preparing a heat-resistant alloy that is synergistically strengthened by multiple types of precipitation phases. As shown in Figure 1, the method includes the following steps:

(S101) Smelting process: Vacuum induction melting + atmosphere protected electroslag remelting dual process smelting is used; electroslag remelting is carried out in an argon protective atmosphere, and the remelting process is ensured to be carried out under the condition of low electrode melting rate. To ensure the cleanliness and homogeneity of heat-resistant alloy ingots and improve the mechanical properties of the alloy.

(S102) Homogenization treatment: Element segregation is inevitable in electroslag ingots, and brittle γ-γ' eutectic phase, brittle Laves phase and large-sized (Ti, Nb) C are precipitated. The presence of these low melting point brittle phases and large size (Ti, Nb)C will reduce the hot working plasticity of the alloy and lead to forging cracking. Therefore, electroslag ingots must undergo homogenization treatment before forging to dissolve large-sized eutectic phases and eliminate element segregation.

To this end, based on thermodynamic equilibrium calculation (Figure 2) and combined with production practice, a two-stage homogenization process was determined. The first stage of homogenization treatment temperature is 900 ~ 1050 ℃, the holding time is 2 ~ 10 hours, so that the γ-γ' eutectic phase and Laves phase can be fully redissolved; the second stage homogenization treatment temperature is 1100 ~ 1150 ℃, the holding time is 2 ~16h, (Ti, Nb)C is fully dissolved, Ti and Nb elements are evenly diffused, and element segregation is eliminated. After high-temperature diffusion annealing, the furnace is first cooled to 1030°C and then air-cooled out of the furnace.

Here, it should be noted that if a one-stage homogenization process is used and the temperature is higher than the initial melting temperature of the Laves phase and (Ti, Nb)C, the precipitated phase will melt and void defects will occur inside the alloy.

(S103) Forging: The electroslag ingot is kept at 1150~1180°C and then forged. The starting forging temperature is 1150~1180℃, the final forging temperature is not lower than 950℃, and the forging ratio is 3~4.

(S104) Solid solution and aging: The solid solution treatment is mainly used to control the redissolution of the precipitated phase, the uniformity of the structure and the grain size. The aging treatment is mainly used to control the precipitation of the strengthening phase γ', (Ti, Nb) C and Laves phase. Quantity, size and distribution. The high temperature solid solution temperature is 950~1100℃, and the holding time is 0.5~6 hours; the aging treatment temperature is 580~780℃, and the holding time is 4~32 hours.

The technical solution of the present invention also provides a multi-type precipitation phase synergistically strengthened heat-resistant alloy, which is characterized in that the chemical composition of the heat-resistant alloy contains: C 0.01~0.04%, Si 0.1~0.4%, Mn0. 2～0.6%, Mo 0.8～1.6%, Cr 12～17%, Ni 28～33%, Ti 1.5～2.5%, Al 0.5～2.0%, Nb 0.2～0.8%, N＜0.004%, P＜0.005% , S<0.004%, the balance is Fe.

Among them, the optimized mass percentage of Al+Ti is: 2.5% ≤ Al + Ti ≤ 3.4%, and the Ti/Al ratio is 1.4 ≤ Ti/Al ≤ 4.

Here, the preferred alloy component Al content is controlled at 0.5% to 1.4%, and the Ti content is controlled at 1.9% to 2.4%. Increased Al content promotes the precipitation of γ' phase, and Al is an important element for improving the oxidation resistance of the alloy. Ti is a strong carbide-forming element. In addition to forming the γ' phase, the Ti added to the alloy is also used to fix carbon, form stable and difficult-to-decompose carbides, eliminate the depletion of Cr at the grain boundaries, and thereby eliminate the grain boundaries of the alloy. Interval corrosion. At the same time, the generated nanoscale carbides can play a role in precipitation strengthening. The Ti/Al ratio determines the antiphase domain boundary energy of the γ' phase. As the Ti/Al ratio of the alloy is increased, the antiphase domain boundary energy of the γ' phase increases. By increasing the Ti content and the Ti/Al ratio, the lattice constant of the γ' phase increases, and the mismatch between the γ' phase and the substrate increases. Add 0.5% Nb to enhance the thermal stability of the γ' phase.

It should be noted that, in addition to the matrix iron, the specific selection reasons for the above main chemical components are as follows (the alloy components in this specification are all based on mass percentage):

·Carbon (C): C can form and stabilize austenite and form carbides with other elements. On the one hand, if the C content in the alloy is too high, large-sized eutectic carbides, such as M ₂₃ C ₆ , MC and M ₆ C, will continuously precipitate on the grain boundaries, reducing the grain boundary strength and adversely affecting the toughness of the alloy. On the other hand, excessive carbides distributed in a network will be formed, which will increase the sensitivity of liquefaction cracks in the welding heat-affected zone and reduce the welding performance of the alloy. Controlling an appropriate amount of C content in the alloy and aging treatment will cause carbides to precipitate discontinuously on the grain boundaries, which will help hinder grain boundary slip and crack expansion, and improve the durability of the alloy. However, too low C content will reduce the nucleation driving force of carbides, make it difficult for carbides to precipitate, and reduce the strength and toughness of grain boundaries. At the same time, too low C content leads to the weakening of the solid solution strengthening effect of C element. Therefore, the C content is strictly controlled within the range of 0.01 to 0.04% in the present invention.

·Silicon (Si): Adding Si to heat-resistant alloys is beneficial to improving the gamma matrix strength, steam corrosion resistance and high-temperature oxidation resistance. However, too high Si content promotes the precipitation of intermetallic σ phase and reduces the grain boundary strength, which is not conducive to the impact toughness and lasting life of the alloy. Considering that this heat-resistant alloy serves in environments with high-temperature steam corrosion and high-temperature exhaust gas corrosion, at least 0.1% Si needs to be added to enhance the corrosion resistance and high-temperature oxidation resistance of the alloy. Therefore, the Si content in the alloy of the present invention is strictly controlled at 0.1 to 0.4%.

Manganese (Mn): Mn can form and stabilize austenite. Adding Mn to the alloy can increase the strength and improve hot processing, corrosion resistance and welding properties. However, excessive Mn will form MnS with S, reducing the cleanliness of the alloy. Mn tends to segregate at grain boundaries, causing the grain boundary strength to be weakened and the alloy's lasting strength to be reduced. Considering that forging is an essential step in the production of this alloy, at least 0.2% Mn needs to be added to enhance the hot workability of the alloy. Therefore, the Mn content in the alloy of the present invention is strictly controlled at 0.2 to 0.6%.

·Chromium (Cr): Cr added to heat-resistant alloys can play a solid solution strengthening role. Cr in the matrix causes the crystal lattice to be distorted and generate an elastic stress field that interacts with dislocations, thus increasing the strength of the γ solid solution. Cr forms a dense Cr ₂ O ₃ type oxide film during the service process of heat-resistant alloys, which improves resistance to high-temperature oxidation and hot corrosion. However, excessive Cr content will promote the precipitation of the intermetallic σ phase, destroy the structural stability, and damage the mechanical properties of the alloy. At least 12% Cr needs to be added to this heat-resistant alloy to form a Cr ₂ O ₃ type oxide film. Based on the above considerations, the Cr content in the alloy of the present invention ranges from 12 to 17%.

·Nickel (Ni): Ni can stabilize and expand the austenite phase area to obtain a single-phase austenite structure. The addition of Ni can improve the composition and properties of the Cr ₂ O ₃ type oxide film, and the alloy's resistance to high-temperature oxidation is improved. The addition of Ni improves the corrosion resistance and plastic toughness of the alloy. However, excessive Ni content will increase the coarsening rate of NbNi ₃ , reduce the thermal strength, and increase the cost of the alloy. Considering that this heat-resistant alloy serves in a high-temperature oxidation environment, at least 28% Ni needs to be added to meet the performance requirements of the alloy against high-temperature oxidation. Therefore, the Ni content in the alloy of the present invention ranges from 28 to 33%.

·Molybdenum (Mo): Mo mainly plays a solid solution strengthening role in heat-resistant alloys. Mo can slow down the diffusion rate of Cr, Al and Ti at high temperatures, improve the interatomic bonding force of γ solid solution, and significantly improve the thermal strength of the alloy. The hardness of the alloy can be improved by precipitating fine Mo-rich intermetallic compounds (Laves phase) during aging. The segregation coefficient of Mo element is less than 1, and it segregates in the interdendritic area during solidification. If the Mo content is too high, the segregation will be serious. On the one hand, it will promote the precipitation of large-sized M ₆ C-type carbides. On the other hand, it is easy to generate TCP harmful phases, such as μ phase. This heat-resistant alloy needs to add at least 0.8% Mo to slow down the diffusion rate of Al and Ti elements at high temperatures and suppress the coarsening of the γ' strengthening phase. Therefore, the Mo content in the alloy of the present invention is strictly controlled at 0.8 to 1.6%.

·Titanium (Ti): About 90% of the Ti added to the heat-resistant alloy forms γ'-Ni ₃ (Ti, Al), and about 10% enters the γ solid solution for solid solution strengthening. Under the condition of a certain Al content, the increase of Ti content promotes the precipitation of γ' phase and improves the high temperature strength of the alloy. Ti is also a key element to enhance the alloy's resistance to hot corrosion and improve the stability of the surface structure. However, if the Ti/Al ratio is too high, the tendency of the γ' phase to transform into the eta-Ni ₃ Ti phase increases. After the alloy is aged for a long time, a needle-like eta phase forms at the grain boundary, destroying the structural stability and reducing the impact toughness of the alloy. The heat-resistant alloy of the present invention needs to add at least 1.5% Ti to form γ' phase, Fe ₂ Ti type Laves phase and (Ti, Nb) C strengthening phase. Therefore, the Ti content in the alloy of the present invention is strictly controlled at 1.5 to 2.5%.

·Aluminum (Al): About 80% of the Al added to the heat-resistant alloy forms γ'-Ni ₃ (Ti, Al), and about 20% enters the γ solid solution for solid solution strengthening. The increase in Al content promotes the precipitation of γ' phase, and Al is an important element to improve the oxidation resistance of the alloy and increases the stability of the alloy surface structure. However, if the Al content is too high, harmful β-NiAl phase may precipitate. In order to ensure the high-temperature strength of the alloy, the present invention requires adding at least 0.5% Al to precipitate a γ' strengthening phase with a volume fraction of at least 15%. At the same time, 0.5% Al cooperates with Cr element to further enhance the alloy's high-temperature oxidation resistance. Therefore, the Al content in the alloy of the present invention is strictly controlled at 0.5 to 2.0%.

·Niobium (Nb): The atomic radius of Nb is larger than that of Mo, and the solid solution strengthening effect is good. Nb can replace Ti and Al in the γ' phase, increase the mismatch between the γ' phase and the γ matrix, and improve the strengthening ability of the γ' phase. Adding Nb is beneficial to increasing the thermal stability and volume fraction of the γ' phase, and Nb and C can form fine dispersed nanoscale MC carbides, improving the structural stability and creep strength of the alloy. However, if the Nb content is too high, a large number of micron-sized eutectic carbides will be formed, which is detrimental to the plastic toughness, welding performance and corrosion resistance of the alloy. In addition, too high Nb content will reduce the oxidation resistance of the alloy, especially the cyclic oxidation resistance. The present invention requires the addition of at least 0.2% Nb to enhance the thermal stability of the γ' phase. At the same time, Nb combines with C elements to form fine dispersed nanoscale carbides to strengthen grain boundaries. Therefore, the Nb content in the alloy of the present invention is strictly controlled at 0.2 to 0.8%.

In addition, in order to ensure the performance of the alloy, the content of five harmful elements and other impurity elements should be as low as possible.

Example 1

Table 1 shows the alloy components (weight percentage) of the examples; Table 2 shows the slag components (weight percentage) used to smelt the heat-resistant alloy of the present invention; Table 3 shows the examples and comparative examples at 760°C. Comparison of high temperature mechanical properties.

According to the composition of heat-resistant alloy 1# shown in Table 1, the alloy ingot is smelted using a double process of vacuum induction melting + atmosphere-protected electroslag remelting. The electroslag ingot undergoes two-stage homogenization treatment. The first stage homogenization treatment temperature is 930°C and the holding time is 4 hours; the second stage homogenization treatment temperature is 1140°C and the holding time is 6 hours. The homogenized electroslag ingot is kept at 1160°C and then forged. The starting forging temperature is 1160℃, the final forging temperature is not less than 950℃, and the forging ratio is 4. The prepared alloy billet is solution treated at 1015°C, kept for 1 hour and then water-cooled to room temperature; then it is kept at an aging temperature of 770°C for 4-16 hours and then air-cooled to room temperature to obtain a multi-type precipitation-strengthened heat-resistant alloy.

Table 1

Table 2

As shown in Figure 3, the scanning electron microscope photo of the typical structural characteristics of heat-resistant alloy 1# after aging treatment at 770°C for 4 hours. The γ' phase has a spherical morphology, the volume fraction is about 29%, and (Ti, Nb)C is It has a small block morphology with a volume fraction of about 7%. The Fe ₂ Ti type Lvaes phase has a small block morphology with a volume fraction of about 3%.

Example 2

According to the composition of heat-resistant alloy 2# shown in Table 1, the alloy ingot is smelted using a double process of vacuum induction melting + atmosphere-protected electroslag remelting. The electroslag ingot undergoes a two-stage homogenization treatment. The first stage homogenization treatment temperature is 980°C and the holding time is 8 hours; the second stage homogenization treatment temperature is 1140°C and the holding time is 12 hours. The homogenized electroslag ingot is kept at 1160°C and then forged. The starting forging temperature is 1160℃, the final forging temperature is not less than 950℃, and the forging ratio is 4. The prepared alloy billet is solution treated at 1050°C, kept for 0.5 hours and then water-cooled to room temperature; then it is kept at an aging temperature of 710°C for 4-32 hours and then air-cooled to room temperature to obtain a multi-type precipitation-strengthened heat-resistant alloy.

As shown in Figure 4, the scanning electron microscope photo of the typical structural characteristics of heat-resistant alloy 2# after aging treatment at 710°C for 28 hours. The γ' phase has a spherical morphology with a volume fraction of about 34%. (Ti, Nb) C and The Fe ₂ Ti type Lvaes phase has a small block morphology.

Example 3

According to the composition of heat-resistant alloy 3# shown in Table 1, the alloy ingot is smelted using a double process of vacuum induction melting + atmosphere-protected electroslag remelting. The electroslag ingot undergoes a two-stage homogenization treatment. The first stage homogenization treatment temperature is 900°C and the holding time is 6 hours; the second stage homogenization treatment temperature is 1140°C and the holding time is 10 hours. The homogenized electroslag ingot is kept at 1160°C and then forged. The starting forging temperature is 1160℃, the final forging temperature is not less than 950℃, and the forging ratio is 4. The prepared alloy billet is solution treated at 1015°C, kept for 1 hour and then water-cooled to room temperature; then it is kept at an aging temperature of 740°C for 4-16 hours and then air-cooled to room temperature to obtain a multi-type precipitation-strengthened heat-resistant alloy.

As shown in Figure 5, the scanning electron microscope photo of the typical structural characteristics of heat-resistant alloy 3# after aging treatment at 740°C for 4 hours. The γ' phase has a spherical morphology with a volume fraction of about 30%, (Ti, Nb) C and The Fe ₂ Ti type Laves phase has a small block morphology.

Table 3 shows the high-temperature tensile properties and comparison of the alloys in comparative examples of the heat-resistant alloys of the present invention after different aging times. Table 3 adds the high-temperature tensile property data of A286 alloy and another 15Cr-30Ni-3.3Cu heat-resistant alloy developed by the inventor for comparison with the alloy of the present invention.

Table 3 Comparison of high temperature mechanical properties (760℃)

The performance data of the alloys in the comparative examples of the alloys of the present invention are shown in Table 3. By further optimizing the design of alloy components and rationally regulating electroslag smelting, diffusion annealing, thermal processing, solid solution and aging heat treatment processes, the alloy of the present invention can have excellent high-temperature strength and plasticity. Compared with A286 alloy and 15Cr-30Ni-3.3Cu alloy, the high-temperature strength and plasticity of the heat-resistant alloy of the present invention are significantly improved, which fully meets the requirements for the use of heat-resistant alloy materials such as automobile engine exhaust valves and fasteners, and can also be used to manufacture aviation products. Engine load-bearing parts and gas turbine heat-resistant parts have a wider application range. At the same time, it has obvious cost advantages compared with nickel-based alloys.

The embodiments of the present invention have been described above in conjunction with the accompanying drawings. However, the present invention is not limited to the above-mentioned specific implementations. The above-mentioned specific implementations are only illustrative and not restrictive. Those of ordinary skill in the art will Under the inspiration of the present invention, many forms can be made without departing from the spirit of the present invention and the scope protected by the claims, and these all fall within the protection of the present invention.

Claims (10)

1. A method for preparing a heat-resistant alloy synergistically strengthened by multiple types of precipitation phases, which is characterized by: including the following steps: (1) Smelting: Vacuum induction furnace smelting → Casting electrode rods → Inert atmosphere protected electroslag remelting smelting, or using electric arc furnace + LF + VD + electroslag remelting method to produce electroslag ingots, which are then heated and annealed ; (2) Homogenization treatment: The electroslag ingot is processed using a two-stage homogenization treatment process. After the homogenization treatment is completed, it is first cooled to a certain temperature in the furnace and then air-cooled out of the furnace; (3) Forging: The electroslag ingot processed in step (2) is first kept warm and then forged to produce alloy billets; (4) Solid solution and aging: The alloy billet is subjected to high-temperature solution treatment + aging heat treatment, thereby obtaining a multi-type precipitated phase synergistically strengthened heat-resistant alloy. 2. The preparation method according to claim 1, characterized in that: in the step (1), electroslag remelting is carried out in an argon protective atmosphere, and a special slag system is used, and the composition of the slag system is based on mass percentage. The content is: CaF ₂ : 50 to 55%, CaO: 15 to 25%, Al ₂ O ₃ : 15 to 25%, MgO: 1 to 4%, TiO ₂ : 2 to 5%, FeO ≤ 0.5%, SiO ₂ ≤0.8%. 3. The preparation method according to claim 2, characterized in that: for the 50-200kg ingot type: when the remelting is stable, the voltage is 25-30V and the current is 1500-3000A; for the 200-500kg ingot type: when the remelting is stable The voltage is 30~35V and the current is 3000~4000A; for 500kg~1t ingot type: the voltage is 35~40V and the current is 4000~5000A when remelting is stable. 4. The preparation method according to claim 2, characterized in that: in the step (1), after the electroslag remelting is completed, the electroslag ingot is demoulded, and a stainless steel insulation cover is used to cover the electroslag ingot for slow cooling, and the cooling time is greater than 5h. , effectively prevent surface cracking of electroslag ingots. 5. The preparation method according to claim 1, characterized in that: in the step (2), the two-stage homogenization process specifically includes: The first stage of homogenization treatment temperature is 900~1050℃, and the holding time is 2~10h; The second stage of homogenization treatment temperature is 1100~1150℃, and the holding time is 2~16h. 6. The preparation method according to claim 1, characterized in that in the step (3), the starting forging temperature is 1150-1180°C, the final forging temperature is not lower than 950°C, and the forging ratio is 3-4. 7. The preparation method according to claim 1, characterized in that in the step (4), the high temperature solid solution temperature is 950-1100°C, the holding time is 0.5-6 hours; the aging treatment temperature is 580-780°C, Keeping time is 4 to 32 hours. 8. The preparation method as claimed in claim 7, characterized in that: in the step (4), the structural characteristics of the alloy after aging heat treatment are: the grain size is 5-7 levels, and intragranular precipitation: Globular γ' phase, the size is not larger than 100nm, and the volume fraction is 25% to 35%; Small lumps of (Ti, Nb)C, size no more than 200nm, volume fraction 5% to 8%; The spherical or small Fe ₂ Ti type Laves phase has a size of no more than 200nm and a volume fraction of 2% to 5%. 9. The preparation method as claimed in claim 7, characterized in that: in the step (4), the mechanical properties of the alloy after aging heat treatment satisfy: 760℃ high temperature mechanical properties: yield strength R _p0.2 ≥432Mpa; tensile strength R _m ≥486Mpa. 10. A multi-type precipitate phase synergistically strengthened heat-resistant alloy, characterized in that: the multi-type precipitate phase synergistically strengthened heat-resistant alloy is prepared by the preparation method according to any one of claims 1 to 9, Among them, the chemical composition of the multi-type precipitation phase synergistically strengthened heat-resistant alloy includes, in terms of mass percentage: C 0.01~0.04%, Si 0.1~0.4%, Mn 0.2~0.6%, Mo 0.8~1.6%, Cr 12~ 17%, Ni28～33%, Ti1.5～2.5%, Al 0.5～2.0%, Nb 0.2～0.8%, N＜0.004%, P＜0.005%, S＜0.004%, the balance is Fe.