CN117363955A

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

Info

Publication number: CN117363955A
Application number: CN202311480801.8A
Authority: CN
Inventors: 史成斌; 张淮; 王世舟
Original assignee: University of Science and Technology Beijing USTB
Current assignee: University of Science and Technology Beijing USTB
Priority date: 2023-11-08
Filing date: 2023-11-08
Publication date: 2024-01-09

Abstract

The invention disclosesA multi-type precipitated phase cooperative strengthening heat-resistant alloy and a preparation method thereof belong to the technical field of heat-resistant alloys. The alloy comprises the following chemical components in percentage by mass: 0.01 to 0.04 percent of C, 0.1 to 0.4 percent of Si, 0.2 to 0.6 percent of Mn, 0.8 to 1.6 percent of Mo, 12 to 17 percent of Cr, 28 to 33 percent of Ni, 1.5 to 2.5 percent of Ti, 0.5 to 2.0 percent of Al, 0.2 to 0.8 percent of Nb, less than 0.004 percent of N, less than 0.005 percent of P, less than 0.004 percent of S and the balance of Fe. The heat-resistant alloy has the advantages that the heat-resistant alloy is prepared by the multi-type nanometer strengthening phases (gamma' phase, (Ti, nb) C and Fe with the volume fraction of more than 30 percent ₂ Ti type Laves phase) to realize cooperative reinforcement; the prepared heat-resistant alloy has high-temperature tensile strength not lower than 486Mpa, meets the use requirements of heat-resistant alloy materials such as an exhaust valve and a fastener of an automobile engine, and is also suitable for manufacturing aeroengine bearing parts and heat-resistant parts of a gas turbine.

Description

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

Technical Field

The invention belongs to the technical field of heat-resistant alloy, and relates to a multi-type precipitated phase cooperative reinforcement heat-resistant alloy and a preparation method thereof, which are suitable for exhaust valves of automobile engines.

Background

The automobile industry is one of the most important pillar industries in China, and the total quantity of automobile production and marketing is stable and the first worldwide for 14 years. The fuel oil automobile is still the main stream of the automobile market in China, and the sales volume and the duty ratio of the fuel oil automobile are still far higher than those of the pure electric automobile. The carbon emission of the automobile accounts for about 7.5% of the carbon emission of the whole society. The fuel oil automobile has large holding quantity, low burning efficiency of fossil fuel in the using stage, and is a main factor causing high carbon emission of the automobile. In recent years, with the continuous increase of energy conservation and emission reduction requirements, higher requirements are being put on the combustion efficiency of automobile engines. Valve steel is a key material for automotive engines. The annual output of China exceeds 5 hundred million (about 4.8 ten thousand tons of raw materials are converted). Exhaust valves of automotive engines are subjected to high temperatures and pressures due to the exhaust of high temperature corrosive exhaust gases. The exhaust valve material is required to have excellent high temperature strength, toughness, hardness, wear resistance, oxidation resistance and corrosion resistance at the working temperature, and tissue stability and dimensional stability under the cold-hot alternating working condition of the engine. Meanwhile, the air valve material should have good cold and hot processing and welding performance during processing.

Currently, the widely used gas valve alloys are high alloy steels 21-4N and 21-4NWNb, nickel-based superalloy GH4751 and Nimonic 80A, and the working temperature is 680-820 ℃. However, the valve alloy material is difficult to achieve good matching of high service temperature and high temperature strength, and does not have the advantage of low cost. With the wide application of technologies for improving combustion efficiency such as direct injection, vortex supercharging and the like in an automobile internal combustion engine cylinder, the exhaust valve material is required to have higher high-temperature exhaust gas corrosion resistance, oxidation resistance and high-temperature strength. The existing automobile exhaust valve steel materials 21-4N and 21-4NWNb can not meet the temperature requirements of 700 ℃ and above of the engine combustion chamber. The nickel-based air valve alloy is adopted, and the cost is very high. The high-performance valve alloy for the internal combustion engine 680-760 ℃ of the automobile in China is all dependent on import.

Therefore, there is a need to develop new valve alloy materials that combine high performance with low cost. The method has important strategic significance in solving the problem of material dependence import, and simultaneously provides guarantee for meeting the requirements of future higher emission standard gas valve alloys.

Disclosure of Invention

In order to solve the problems, the technical scheme of the invention provides a multi-type precipitated phase cooperative strengthening heat-resistant alloy and a preparation method thereof, and the high-strength heat-resistant alloy is obtained by optimally designing alloy components and formulating a reasonable production process. After electroslag remelting, homogenization treatment, forging, high-temperature solid solution and aging heat treatment, the prepared heat-resistant alloy has a grain size of 5-7 grade and precipitates fine dispersed multi-type strengthening phases (gamma' phase, (Ti, nb) C and Fe ₂ Ti-type Laves phase). Wherein the gamma' phase strengthening phase is spherical, the volume fraction is 25-35%, the volume fraction of small block (Ti, nb) C is 5-8%, and Fe is contained in the mixture ₂ Ti-type Laves phases are intermittently distributed along grain boundaries, and the volume fraction is 2% -5%.

According to a first aspect of the technical scheme of the invention, a preparation method of a multi-type precipitated phase cooperative reinforcement heat-resistant alloy is provided, which comprises the following steps:

(1) Smelting: smelting in a vacuum induction furnace, casting an electrode rod, protecting electroslag remelting in an inert atmosphere, or smelting by adopting an arc furnace+LF+VD+electroslag remelting method to produce an electroslag ingot, and then carrying out hot feed annealing on the electroslag ingot;

(2) Homogenizing: the homogenization treatment adopts a two-stage heat preservation process, wherein the homogenization treatment temperature in the first stage is 900-1050 ℃, and the heat preservation time is 2-10 hours, so that the gamma-gamma' eutectic phase and the Laves phase are fully dissolved back; the homogenization treatment temperature in the second stage is 1100-1150 ℃ and the heat preservation time is 2-16 h, so that (Ti, nb) C is fully dissolved, ti and Nb elements are uniformly diffused, and element segregation is eliminated; furnace cooling to 1030 ℃ and furnace discharging and air cooling after homogenization treatment is finished;

(3) Forging: the electroslag ingot is preserved at 1150-1180 ℃ and then forged; the forging temperature is 1150-1180 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 3-4;

(4) Solid solution and aging: carrying out high-temperature solid solution treatment and aging treatment on the forged alloy blank, wherein the high-temperature solid solution temperature is 950-1100 ℃, and the heat preservation time is 0.5-6 hours; the aging treatment temperature is 580-780 ℃ and the heat preservation time is 4-32 hours.

Further, electroslag remelting is carried out in an argon protective atmosphere, in order to prevent Ti from burning out in the electroslag process and ensure good surface quality of an electroslag ingot, a special slag system is adopted, and the slag system comprises the following components in percentage by mass: caF (CaF) ₂ ：50～55％，CaO：15～25％，Al ₂ O ₃ ：15～25％，MgO：1～4％，TiO ₂ ：2～5％，FeO≤0.5％，SiO ₂ ≤0.8％。

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

And further, demoulding after electroslag remelting is finished, covering the electroslag ingot by using a stainless steel heat-insulating cover, and slowly cooling for more than 5 hours, so that the surface cracking of the electroslag ingot is effectively prevented.

Further, the alloy structure after aging heat treatment is characterized in that: the grain size is 5-7 grade, spherical gamma' phase (the size is not more than 100nm and the volume fraction is 25% -35%) is precipitated in the crystal, and small block (Ti, nb) C (the size is not more than 200nm and the volume fraction is 5% -8%) and quasi-spherical or small block Fe are precipitated in the crystal boundary ₂ Ti type Laves phase (size is not more than 200nm, volume fraction is 2% -5%, laves phase is composed of Fe, ti, nb, ni, si, mo elements); wherein, the intra-crystal nano-scale gamma ' phase strengthening phase interacts with dislocation, and the mechanism of strengthening by coherent strain, strengthening by an ' Orowan ' bypassing mechanism, strengthening by a dislocation cutting mechanism and strengthening by a dislocation climbing mechanism are all adoptedThe number of gamma prime strengthening phases is a fundamental contributor, and whatever mechanism is active; nano-scale (Ti, nb) C and Fe precipitated at grain boundary ₂ Ti type Laves phase blocks dislocation movement, inhibits grain boundary migration and grain growth, and can improve strength and toughness of the area near the grain boundary. At the same time, nanoscale (Ti, nb) C and Fe ₂ The existence of Ti type Laves phase refines alloy crystal grains and plays a role in fine grain strengthening. Aging to separate out fine dispersed nanometer gamma' phase, (Ti, nb) C and Fe ₂ Ti type Laves phase realizes the cooperative reinforcement of multiple types of precipitation phases.

Further, the mechanical properties of the alloy after aging heat treatment satisfy the following conditions: high temperature mechanical properties at 760 ℃): yield strength R _p0.2 Not less than 432Mpa; tensile strength R _m ≥486Mpa。

According to a second aspect of the technical solution of the present invention, there is provided a multi-type precipitation-phase cooperative strengthening heat-resistant alloy, which is prepared by the preparation method according to any one of the above aspects,

wherein, the chemical components of the multi-type precipitated phase cooperative strengthening heat-resistant alloy comprise the following components in percentage by mass: 0.01 to 0.04 percent of C, 0.1 to 0.4 percent of Si, 0.2 to 0.6 percent of Mn, 0.8 to 1.6 percent of Mo, 12 to 17 percent of Cr, 28 to 33 percent of Ni, 1.5 to 2.5 percent of Ti, 0.5 to 2.0 percent of Al, 0.2 to 0.8 percent of Nb, less than 0.004 percent of N, less than 0.005 percent of P, less than 0.004 percent of S and the balance of Fe.

Further, the chemical components of the multi-type precipitated phase cooperative strengthening heat-resistant alloy comprise the following components in percentage by mass: 0.02% of C, 0.2% of Si, 0.4% of Mn, 1.2% of Mo, 15% of Cr, 30% of Ni, 1.9% of Ti, 0.6% of Al, 0.5% of Nb, less than 0.004% of N, less than 0.005% of P, less than 0.004% of S and the balance of Fe.

Further, the chemical components of the multi-type precipitated phase cooperative strengthening heat-resistant alloy comprise the following components in percentage by mass: 0.02% of C, 0.2% of Si, 0.4% of Mn, 1.2% of Mo, 15% of Cr, 30% of Ni, 2.0% of Ti, 1.4% of Al, 0.5% of Nb, less than 0.004% of N, less than 0.005% of P, less than 0.004% of S and the balance of Fe.

Further, the chemical components of the multi-type precipitated phase cooperative strengthening heat-resistant alloy comprise the following components in percentage by mass: 0.02% of C, 0.2% of Si, 0.4% of Mn, 1.2% of Mo, 15% of Cr, 30% of Ni, 2.4% of Ti, 0.6% of Al, 0.5% of Nb, less than 0.004% of N, less than 0.005% of P, less than 0.004% of S and the balance of Fe.

The invention has the beneficial effects that:

(1) And an ingot homogenization dynamics model is established to guide and formulate a high-temperature homogenization process. The ingot structure before homogenization treatment is distributed with brittle gamma-gamma' eutectic phase, brittle Laves phase and large size (Ti, nb) C.

The electroslag ingot is fully dissolved back into the brittle gamma-gamma' eutectic phase and Laves phase through two-stage homogenization treatment, so that (Ti, nb) C is fully dissolved, the defect of void in the alloy is effectively avoided, and the grain size is controlled to be excessively large. After the two-stage homogenization treatment, the residual segregation coefficient delta of Ti and Nb elements is less than 0.2, and the Ti and Nb elements are fully and uniformly diffused, so that the hot working plasticity of the alloy is improved. The element residual segregation coefficient is calculated as follows:

wherein the method comprises the steps ofAnd->Minimum and maximum concentrations of elements in the ingot, respectively, ">And->Respectively the minimum concentration and the maximum concentration of elements in the steel ingot after high-temperature diffusion annealing.

(2) The sum of the Ti+Al+Nb contents and the Ti/Al ratio are regulated to control the volume fraction, the size, the inversion domain boundary energy and the degree of mismatch with the substrate of the gamma' phase. The sum of the contents of Ti, al and Nb is 3 to 5 percent

Ensure the volume fraction of the gamma 'phase to be 25-35 percent, and 0.5 percent of Nb element is added to strengthen the gamma' phase-

Ni ₃ Thermal stability of (Ti, al, nb). The diffusion speed of Ti and Al elements is slowed down by adding 1.2% of Mo, the Ti/Al ratio is regulated as far as possible, and the generation of harmful phase eta phase after long-time aging is avoided. The intra-crystalline nanoscale gamma' -phase strengthening phase is strengthened by coherent strain strengthening, bypass mechanism strengthening, dislocation cutting mechanism strengthening and dislocation climbing mechanism, and nanoscale (Ti, nb) C and Fe are separated out from grain boundary ₂ The Ti type Laves phase improves the strength and toughness of the area near the grain boundary, and realizes the synergistic strengthening effect of multiple types of precipitated phases.

(3) The content of C, ti, nb, mo, si is strictly controlled, and (Ti, nb) C and Fe are avoided ₂ The Ti type Laves phase is continuously precipitated at the grain boundary, thereby securing the strength and toughness of the grain boundary. (Ti, nb) C and Fe in the alloy of the present invention ₂ The Ti-type Laves phase is intermittently distributed along the grain boundary, the volume fraction of (Ti, nb) C is 5-8%, fe ₂ The volume fraction of Ti-type Laves phase is 2-5%.

Drawings

In order to more clearly illustrate the embodiments of the present invention or the technical solutions in the prior art, the drawings that are required in the embodiments or the description of the prior art will be briefly described, and it is obvious that the drawings in the following description are only some embodiments of the present invention, and other drawings may be obtained according to the structures shown in these drawings without inventive effort for a person skilled in the art.

FIG. 1 is a flow chart of a preparation method of a multi-type precipitated phase cooperative strengthening heat-resistant alloy according to the technical scheme of the invention;

FIG. 2 shows the result of equilibrium phase precipitation in the heat resistant alloy of the present invention;

FIG. 3 shows the nano-scale multi-type precipitated phases (gamma' phase, (Ti, nb) C and Fe after aging at-770 ℃ for 4 hours ₂ Ti type Laves phase) and matrix structure;

FIG. 4 shows the nano-scale multi-type precipitated phases (gamma' -phase, (Ti, nb) C and Fe after aging at-710 ℃ for 28 hours ₂ Ti type Laves phase) and matrix groupWeaving;

FIG. 5 shows the nano-scale multi-type precipitated phases (gamma' -phase, (Ti, nb) C and Fe after aging at-740 ℃ for 4 hours ₂ Ti-type Laves phase) and matrix structure.

Detailed Description

Embodiments of the present invention will be described in detail below in conjunction with examples of the present invention to provide a better understanding of the advantages and features of the present invention to those skilled in the art. It will be apparent that the embodiments described below are only some, but not all, embodiments of the invention. All other embodiments, which can be made by those skilled in the art based on the embodiments of the invention without making any inventive effort, are intended to be within the scope of the invention.

The invention discloses a multi-type precipitated phase cooperative strengthening heat-resistant alloy and a preparation method thereof, and belongs to the technical field of heat-resistant alloys. The alloy comprises the following chemical components in percentage by mass: 0.01 to 0.04 percent of C, 0.1 to 0.4 percent of Si, 0.2 to 0.6 percent of Mn, 0.8 to 1.6 percent of Mo, 12 to 17 percent of Cr, 28 to 33 percent of Ni, 1.5 to 2.5 percent of Ti, 0.5 to 2.0 percent of Al, 0.2 to 0.8 percent of Nb, less than 0.004 percent of N, less than 0.005 percent of P, less than 0.004 percent of S and the balance of Fe. Adopting a vacuum induction smelting and atmosphere protection electroslag remelting duplex process for smelting, carrying out two-stage homogenization treatment on an electroslag ingot at 900-1150 ℃, and discharging and air cooling after the homogenization treatment is finished and the furnace is cooled to 1030 ℃. The electroslag ingot after homogenization treatment is subjected to heat preservation at 1150-1180 ℃ and then is forged; the forging temperature is 1150-1180 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 3-4; the blank after forging is subjected to solution treatment for 0.5 to 6 hours at 950 to 1100 ℃; and aging treatment is carried out for 4 to 32 hours at 580 to 780 ℃. Thus, by volume fraction greater than 30% of the multiple types of nano-reinforcement phases (gamma' -phase, (Ti, nb) C and Fe ₂ Ti type Laves phase) to realize cooperative reinforcement; the prepared heat-resistant alloy has high-temperature tensile strength not lower than 486Mpa, meets the use requirements of heat-resistant alloy materials such as an exhaust valve and a fastener of an automobile engine, and is also suitable for manufacturing aeroengine bearing parts and heat-resistant parts of a gas turbine.

Specifically, the technical scheme of the invention firstly provides a preparation method of a multi-type precipitated phase cooperative reinforcement heat-resistant alloy, as shown in fig. 1, the method comprises the following steps:

(S101) smelting process: smelting by adopting a vacuum induction smelting and atmosphere protection electroslag remelting duplex process; electroslag remelting is carried out in an argon protective atmosphere, and the remelting process is carried out under the condition of low electrode melting rate, so that the cleanliness and homogeneity of the heat-resistant alloy cast ingot are ensured, and the mechanical property of the alloy is improved.

(S102) homogenization treatment: the element segregation phenomenon inevitably exists in the electroslag ingot, and brittle gamma-gamma' eutectic phase, brittle Laves phase and large-size (Ti, nb) C are precipitated. The presence of these low melting brittle phases and large size (Ti, nb) C can reduce the hot working plasticity of the alloy, resulting in forging cracking. Therefore, the electroslag ingot must undergo homogenization treatment prior to forging to dissolve the large-size eutectic phase and eliminate element segregation.

For this purpose, a two-stage homogenization treatment process is determined based on thermodynamic equilibrium calculations (fig. 2) in combination with production practices. The homogenization treatment temperature in the first stage is 900-1050 ℃, and the heat preservation time is 2-10 hours, so that the gamma-gamma' eutectic phase and the Laves phase are fully dissolved back; the homogenization treatment temperature in the second stage is 1100-1150 ℃ and the heat preservation time is 2-16 h, so that (Ti, nb) C is fully dissolved, ti and Nb elements are uniformly diffused, and element segregation is eliminated. And after the high-temperature diffusion annealing is finished, furnace cooling is carried out firstly to 1030 ℃, and then furnace discharging and air cooling are carried out.

It should be noted here that if a one-stage homogenization treatment is used, the temperature is higher than the initial melting temperature of the Laves phase and the (Ti, nb) C, and the precipitated phase will melt, resulting in void defects in the alloy.

(S103) forging: and preserving the temperature of the electroslag ingot at 1150-1180 ℃, and forging. The 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 solution treatment is mainly used for controlling the back dissolution of the precipitated phases, the uniformity of the structure and the grain size, and the aging treatment is mainly used for controlling the quantity, the size and the distribution of the precipitation of the strengthening phases gamma', (Ti, nb) C and Laves. The high-temperature solid solution temperature is 950-1100 ℃, and the heat preservation time is 0.5-6 hours; the aging treatment temperature is 580-780 ℃ and the heat preservation time is 4-32 hours.

The technical scheme of the invention also provides a multi-type precipitated phase cooperative reinforcement heat-resistant alloy, which is characterized by comprising the following chemical components in percentage by mass: 0.01 to 0.04 percent of C, 0.1 to 0.4 percent of Si, 0.2 to 0.6 percent of Mn, 0.8 to 1.6 percent of Mo, 12 to 17 percent of Cr, 28 to 33 percent of Ni, 1.5 to 2.5 percent of Ti, 0.5 to 2.0 percent of Al, 0.2 to 0.8 percent of Nb, less than 0.004 percent of N, less than 0.005 percent of P, less than 0.004 percent of S and the balance of Fe.

Wherein, the mass percentage of Al+Ti which is optimally regulated is as follows: al+Ti is more than or equal to 2.5% and less than or equal to 3.4%, and the ratio of Ti to Al is more than or equal to 1.4 and less than or equal to 4.

Here, the Al content of the preferred alloy composition is controlled to be 0.5% to 1.4%, the Ti content is controlled to be 1.9% to 2.4%, the Al content is increased to promote the precipitation of gamma' -phase, and Al is an important element for improving the oxidation resistance of the alloy. Ti is a strong carbide forming element, and Ti added into the alloy is used for fixing carbon besides forming gamma' phase, forming stable carbide which is not easy to decompose, eliminating depletion of Cr at grain boundary, and thus eliminating intergranular corrosion of the alloy. Meanwhile, the generated nano-scale carbide can play a role in precipitation strengthening. The Ti/Al ratio determines the phase inversion domain boundary energy of the gamma 'phase, improves the Ti/Al ratio of the alloy, and increases the phase inversion domain boundary energy of the gamma' phase. Increasing the Ti content and increasing the Ti/Al ratio increases the lattice constant of the gamma 'phase and increases the degree of mismatch between the gamma' phase and the substrate. The addition of 0.5% Nb enhances the thermal stability of the gamma' phase.

It should be noted that the above main chemical components are specifically selected from the following reasons (in the present specification, the alloy components are all in mass percent) except for the matrix iron:

carbon (C): c can form and stabilize austenite and form carbides with other elements. An excessive C content in the alloy can lead to continuous precipitation of large-size eutectic carbides such as M on grain boundaries ₂₃ C ₆ MC and M ₆ C, etc., the grain boundary strength is lowered, and the alloy toughness is adversely affected. On the other hand, excessive carbide distributed in a net shape can be formed, so that the liquidation crack sensitivity of a welding heat affected zone is improved, and the welding performance of the alloy is reduced. The content of proper C in the alloy is controlled, and carbide is discontinuously precipitated on the grain boundary by aging treatment, which is favorable for resistancePreventing grain boundary sliding and crack propagation, and improving the durability life of the alloy. However, too low a content of C lowers the nucleation driving force of the carbide, which makes precipitation difficult and lowers the strength and toughness of the grain boundary. Meanwhile, too low a content of C results in a decrease in the solid solution strengthening effect of the C element. Therefore, the C content is strictly controlled to be 0.01-0.04%.

Si is added into the silicon (Si) heat-resistant alloy to improve the strength, the steam corrosion resistance and the high-temperature oxidation resistance of the gamma matrix. However, the Si content is too high to promote the precipitation of intermetallic phase sigma phase, reduce the grain boundary strength and are unfavorable for the impact toughness and the lasting life of the alloy. Considering that the heat-resistant alloy is in service in the environment of high-temperature steam corrosion and high-temperature exhaust gas corrosion, at least 0.1% of 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 to be 0.1 to 0.4%.

Mn, mn can form and stabilize austenite. Mn is added into the alloy to improve the strength and the hot working, corrosion resistance and welding performance. However, excessive Mn and S form MnS, and the cleanliness of the alloy is reduced. Mn tends to be biased to grain boundaries, so that the grain boundary strength is weakened, and the alloy endurance strength is reduced. Considering that forging is an essential element for producing the present alloy, it is necessary to add at least 0.2% Mn to enhance the hot workability of the alloy. Therefore, the Mn content in the alloy of the present invention is strictly controlled to be 0.2-0.6%.

Chromium (Cr): cr is added into heat-resistant alloy to play the role of solid solution strengthening. Cr in the matrix distorts the crystal lattice, and generates elastic stress field to interact with dislocation, so as to raise the strength of gamma solid solution. Cr forms compact Cr in the service process of the heat-resistant alloy ₂ O ₃ The oxide film can improve high temperature oxidation resistance and hot corrosion resistance. However, too high Cr content promotes precipitation of intermetallic phase sigma phase, damages the structural stability and damages the mechanical properties of the alloy. At least 12% Cr is added into the heat-resistant alloy to form Cr ₂ O ₃ And a type oxide film. In summary, the Cr content in the alloy of the present invention ranges from 12 to 17%.

Nickel (Ni): ni can stabilize and enlarge the austenite phase region, and a single-phase austenite structure is obtained. Ni addition can be changedGood Cr ₂ O ₃ The high-temperature oxidation resistance of the alloy is improved. The addition of Ni improves the corrosion resistance and the ductility and toughness of the alloy. However, too high Ni content increases NbNi ₃ Reduces the heat resistance and increases the cost of the alloy. Considering that the heat-resistant alloy is in service in a high-temperature oxidation environment, at least 28% of Ni is required to be added to meet the high-temperature oxidation resistance requirement of the alloy. 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 speed of Cr, al and Ti at high temperature, improve the interatomic bonding force of gamma solid solution and obviously improve the heat resistance of the alloy. Aging out of fine Mo-rich intermetallic compounds (Laves phases) can increase the hardness of the alloy. The segregation coefficient of Mo element is less than 1, and the Mo element is biased to the interdendritic region during solidification. Segregation is severe after the Mo content is too high, on one hand, large-size M can be promoted ₆ The C-type carbide precipitates, and on the other hand, TCP deleterious phases such as a mu phase are easily formed. The heat-resistant alloy needs to be added with at least 0.8% of Mo to slow down the diffusion speed of Al and Ti elements at high temperature and inhibit coarsening of gamma' strengthening phases. Therefore, the Mo content in the alloy is strictly controlled to be 0.8-1.6%.

Titanium (Ti) about 90% of the Ti added to the heat resistant alloy forms gamma' -Ni ₃ (Ti, al), about 10% enters into gamma solid solution to play a solid solution strengthening role. Under the condition of a certain Al content, the increase of the Ti content promotes the precipitation of gamma' phase and improves the high-temperature strength of the alloy. Ti is also a key element for enhancing the hot corrosion resistance of the alloy and improving the stability of a surface layer structure. However, when the Ti/Al ratio is too high, the gamma' phase is increased to eta-Ni ₃ The Ti phase tends to transform. After the alloy is aged for a long time, a needle-shaped eta phase is formed at the grain boundary, the tissue stability is destroyed, and the impact toughness of the alloy is reduced. The heat resistant alloy of the present invention requires at least 1.5% Ti to form gamma prime phase, fe ₂ Ti-type Laves phase and (Ti, nb) C strengthening phase. Therefore, the Ti content in the alloy of the invention is strictly controlled to be 1.5-2.5%.

Aluminum (Al) about 80% of the Al added to the heat resistant alloy forms gamma' -Ni ₃ About 20% of the (Ti, al) particles enter into a gamma solid solution to perform a solid solution strengthening function. Promotion of Al content increaseAnd gamma' phase is separated out, al is an important element for improving the oxidation resistance of the alloy, and the surface structure stability of the alloy is improved. However, too high an Al content may precipitate a detrimental beta-NiAl phase. In order to ensure the high temperature strength of the alloy, the invention at least needs to add 0.5 percent of Al to precipitate gamma' strengthening phase with the volume fraction of at least 15 percent. Meanwhile, 0.5% of Al cooperates with Cr element to further enhance the high-temperature oxidation resistance of the alloy. Therefore, the Al content in the alloy of the present invention is strictly controlled to be 0.5-2.0%.

Niobium (Nb) has a larger Nb atom radius than Mo, and has a good solid solution strengthening effect. Nb can replace Ti and Al in the gamma ' phase, increase the mismatching degree between the gamma ' phase and the gamma matrix, and improve the strengthening capability of the gamma ' phase. The addition of Nb is beneficial to increasing the thermal stability and volume fraction of gamma' phase, and Nb and C can form fine dispersion nanoscale MC carbide, thereby improving the structural stability and creep strength of the alloy. However, the Nb content is too high to form a large amount of micron-sized eutectic carbide, which is unfavorable for the toughness, welding performance and corrosion resistance of the alloy. Furthermore, an excessively high Nb content reduces the oxidation resistance, especially the cyclic oxidation resistance, of the alloy. The present invention requires the addition of at least 0.2% Nb to enhance the thermal stability of the gamma prime phase. Meanwhile, nb and C elements are combined to form fine dispersed nanoscale carbide, and grain boundaries are reinforced. Therefore, the Nb content in the alloy of the invention is strictly controlled to be 0.2-0.8%.

In addition, in order to ensure the performance of the alloy, the lower the content of five harmful elements and other impurity elements is, the better.

Example 1

Table 1 shows the alloy compositions (weight percent) of the examples; table 2 shows the slag composition (weight percent) used in the smelting of the heat resistant alloy of the present invention; table 3 shows the high temperature mechanical properties of each example and comparative example at 760 ℃.

Alloy ingots were smelted by a duplex process of vacuum induction melting and atmosphere protection electroslag remelting according to the compositions shown in heat-resistant alloy 1# of table 1. Carrying out two-stage homogenization treatment on the electroslag ingot, wherein the homogenization treatment temperature in the first stage is 930 ℃, and the heat preservation time is 4h; the homogenization treatment temperature in the second stage is 1140 ℃ and the heat preservation time is 6h. The electroslag ingot after homogenization treatment is preserved at 1160 ℃ and then is forged. The forging temperature is 1160 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 4. The prepared alloy blank is subjected to solution treatment at 1015 ℃, and is cooled to room temperature after heat preservation for 1 hour; and then preserving heat for 4-16 hours at the aging temperature of 770 ℃, and then air-cooling to room temperature to obtain the multi-type precipitation-phase reinforced heat-resistant alloy.

TABLE 1

TABLE 2

As shown in FIG. 3, the heat-resistant alloy 1# is subjected to aging treatment at 770 ℃ for 4 hours and then subjected to Scanning Electron Microscope (SEM) photograph of typical structure characteristics, wherein the gamma' -phase is spherical, the volume fraction is about 29%, the (Ti, nb) C is small-block, the volume fraction is about 7%, and the Fe is ₂ The Ti-type lvae phase has a small block morphology, and the volume fraction is about 3%.

Example 2

Alloy ingots were smelted by a duplex process of vacuum induction melting and atmosphere protection electroslag remelting according to the compositions shown in heat-resistant alloy # 2 in table 1. Carrying out two-stage homogenization treatment on the electroslag ingot, wherein the homogenization treatment temperature in the first stage is 980 ℃, and the heat preservation time is 8 hours; the homogenization treatment temperature in the second stage is 1140 ℃ and the heat preservation time is 12h. The electroslag ingot after homogenization treatment is preserved at 1160 ℃ and then is forged. The forging temperature is 1160 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 4. The prepared alloy blank is subjected to solution treatment at 1050 ℃, is kept for 0.5 hour, and is cooled to room temperature; and then preserving the temperature at 710 ℃ for 4-32 hours, and then air-cooling to room temperature to obtain the multi-type precipitation-phase reinforced heat-resistant alloy.

As shown in FIG. 4, the heat-resistant alloy 2# is subjected to aging treatment at 710 ℃ for 28 hours and then has a typical structure, and the gamma' -phase is spherical, with a volume fraction of about 34%, (Ti, nb) C and Fe ₂ The Ti type Lvaes phase is in a small block shape.

Example 3

Alloy ingots were smelted by a duplex process of vacuum induction melting and atmosphere protection electroslag remelting according to the compositions shown in heat-resistant alloy 3# of table 1. Carrying out two-stage homogenization treatment on the electroslag ingot, wherein the homogenization treatment temperature in the first stage is 900 ℃, and the heat preservation time is 6h; the homogenization treatment temperature in the second stage is 1140 ℃ and the heat preservation time is 10 hours. The electroslag ingot after homogenization treatment is preserved at 1160 ℃ and then is forged. The forging temperature is 1160 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 4. The prepared alloy blank is subjected to solution treatment at 1015 ℃, and is cooled to room temperature after heat preservation for 1 hour; and then preserving heat for 4-16 hours at the aging temperature of 740 ℃, and then air-cooling to room temperature to obtain the multi-type precipitation-phase reinforced heat-resistant alloy.

As shown in FIG. 5, the heat-resistant alloy 3# is subjected to aging treatment at 740 ℃ for 4 hours and then has a typical structure characteristic, and the gamma' -phase is spherical, and has a volume fraction of about 30%, (Ti, nb) C and Fe ₂ Ti-type Laves phase is in a small block shape.

The high temperature tensile properties of the alloys of the comparative examples of the present invention after various time periods and the comparisons are shown in Table 3. The high temperature tensile property data of another 15Cr-30Ni-3.3Cu heat resistant alloy developed by the present inventors, to which the A286 alloy was added, was compared with the alloy of the present invention in Table 3.

TABLE 3 comparison of high temperature mechanical Properties (760 ℃ C.)

The various performance data of the alloys of the present invention in the comparative examples are shown in Table 3. According to the invention, through further optimizing and designing alloy components, electroslag smelting, diffusion annealing, hot working process, solid solution and aging heat treatment process are reasonably regulated and controlled, so that the alloy has excellent high-temperature strength and plasticity. Compared with the A286 alloy and the 15Cr-30Ni-3.3Cu alloy, the heat-resistant alloy has obviously improved high-temperature strength and plasticity, completely meets the use requirements of heat-resistant alloy materials such as an exhaust valve and a fastener of an automobile engine, can also be used for manufacturing aeroengine bearing parts and heat-resistant parts of a gas turbine, and has wider application range. Meanwhile, compared with nickel base alloy, the nickel base alloy has obvious cost advantage.

The embodiments of the present invention have been described above with reference to the accompanying drawings, but the present invention is not limited to the above-described embodiments, which are merely illustrative and not restrictive, and many forms may be made by those having ordinary skill in the art without departing from the spirit of the present invention and the scope of the claims, which are to be protected by the present invention.

Claims

1. A preparation method of a multi-type precipitated phase cooperative strengthening heat-resistant alloy is characterized by comprising the following steps of: the method comprises the following steps:

(1) Smelting: vacuum induction furnace smelting, electrode rod casting, inert atmosphere protection electroslag remelting, or arc furnace +LF +VD +electroslag remelting method smelting to produce electroslag ingot, and then hot-feeding annealing of the electroslag ingot;

(2) Homogenizing: adopting a two-stage homogenization treatment process to treat the electroslag ingot, and discharging the electroslag ingot from the furnace for air cooling after the homogenization treatment is finished and cooling the electroslag ingot to a certain temperature;

(3) Forging: firstly preserving heat of the electroslag ingot treated in the step (2) and then forging to produce an alloy blank;

(4) Solid solution and aging: and carrying out high-temperature solution treatment and aging heat treatment on the alloy blank, thereby obtaining the multi-type precipitated phase cooperative strengthening heat-resistant alloy.

2. The method of manufacturing according to claim 1, wherein: in the step (1), electroslag remelting is carried out in an argon protective atmosphere, a special slag system is adopted, and the slag system comprises the following components in percentage by mass: caF (CaF) ₂ ：50～55％，CaO：15～25％，Al ₂ O ₃ ：15～25％，MgO：1～4％，TiO ₂ ：2～5％，FeO≤0.5％，SiO ₂ ≤0.8％。

3. The method of manufacturing as claimed in claim 2, wherein: aiming at 50-200 kg of ingot type: the voltage is 25-30V and the current is 1500-3000A when remelting and stabilizing; aiming at 200-500 kg of ingot type: the voltage is 30-35V and the current is 3000-4000A when remelting and stabilizing; aiming at 500 kg-1 t ingot: the voltage is 35-40V and the current is 4000-5000A when remelting and stabilizing.

4. The method of manufacturing as claimed in claim 2, wherein: in the step (1), demoulding is carried out after electroslag remelting is finished, and a stainless steel heat-insulating cover is used for covering the electroslag ingot for slow cooling, so that the cooling time is longer than 5 hours, and the surface of the electroslag ingot is effectively prevented from cracking.

5. The method of manufacturing according to claim 1, characterized in that: in the step (2), the two-stage homogenization treatment process specifically comprises:

the homogenization treatment temperature in the first stage is 900-1050 ℃, and the heat preservation time is 2-10 h;

the homogenization treatment temperature in the second stage is 1100-1150 ℃ and the heat preservation time is 2-16 h.

6. The method according to claim 1, wherein in the step (3), the forging temperature is 1150 to 1180 ℃, the final forging temperature is not lower than 950 ℃, and the forging ratio is 3 to 4.

7. The method according to claim 1, wherein in the step (4), the high-temperature solid solution temperature is 950-1100 ℃, and the heat preservation time is 0.5-6 hours; the aging treatment temperature is 580-780 ℃ and the heat preservation time is 4-32 hours.

8. The method of manufacturing according to claim 7, wherein: in the step (4), the alloy structure after aging heat treatment is characterized in that: the grain size is 5-7 grade, and the crystal is precipitated in the crystal:

the size of the spherical gamma' phase is not more than 100nm, and the volume fraction is 25% -35%;

small block (Ti, nb) C with the size not more than 200nm and the volume fraction of 5-8%;

spheroid or small block Fe ₂ Ti type Laves phase with size not more than 200nm and volume fraction of 2-5%.

9. The method of manufacturing according to claim 7, wherein: in the step (4), the mechanical properties of the alloy after aging heat treatment are as follows:

high temperature mechanical properties at 760 ℃): yield strength R _p0.2 Not less than 432Mpa; tensile strength R _m ≥486Mpa。

10. A multi-type precipitated phase cooperative strengthening heat-resistant alloy is characterized in that: the multi-type precipitated phase cooperative strengthening heat-resistant alloy is prepared by adopting the preparation method according to any one of claims 1 to 9,