CN110205565B

CN110205565B - Dispersion nanometer strengthened 690 steel and manufacturing method thereof

Info

Publication number: CN110205565B
Application number: CN201910473681.6A
Authority: CN
Inventors: 周梓荣; 王艳林
Original assignee: Dongguan Dongguan Institute Of Science And Technology Innovation
Current assignee: Dongguan Dongguan Institute Of Science And Technology Innovation
Priority date: 2019-06-01
Filing date: 2019-06-01
Publication date: 2021-08-24
Anticipated expiration: 2039-06-01
Also published as: CN110205565A

Abstract

The invention discloses dispersion nano strengthened 690 steel and a manufacturing method thereof, wherein the manufacturing method comprises the working procedures of casting, heating, rolling and heat treatment, and the steel comprises the following chemical components in percentage by mass: 0.02-0.30% of C, 0.50-2.0% of Mn, 0.04-0.10% of Al, 0.002-0.013% of V, 0.01-0.3% of Nb, 0.005-0.3% of Ti, 0.001-0.010% of N, and the balance of Fe and inevitable impurities; the manufacturing method comprises fusion casting and TMCP processes. The method improves the hardenability of the steel by adjusting and optimizing the proportion of alloy elements in the steel, forms carbonitride nano-scale second phase particles at the same time by using trace alloy elements, and controls the precipitation behavior of the trace alloy elements and the precipitation amount, size and distribution of the particles by adopting a TMCP (thermal mechanical control processing) process so as to refine the particles, improve the obdurability of the steel and obtain good formability and corrosion resistance.

Description

Dispersion nanometer strengthened 690 steel and manufacturing method thereof

Technical Field

The invention relates to application of 690 steel in the field of ocean engineering equipment, in particular to dispersion nano strengthened 690 steel and a manufacturing method thereof.

Background

The steel is used as a key structural material of ocean engineering equipment and widely applied to drilling platforms, production platforms, submarine pipelines and the like. Because the service time is long, the steel plate is required to resist severe wind and wave conditions for a long time, the underwater repair and maintenance cost is extremely high, and the adopted steel plate has high strength, high toughness, fatigue resistance, lamellar tearing resistance, good weldability and cold processing performance, and particularly important marine atmosphere and seawater corrosion resistance.

Taking high-strength steel for ocean platforms as an example, the main grades are

yield strength

355, 420, 460, 500, 550, 620 and 690MPa, while the requirements of pile legs, cantilever beams, lifting electric gears, gear mechanisms and the like of the self-elevating ocean platform are more strict, high-strength or extra-thick special steel with the grade of 690MPa or above is needed, and the high-strength steel has excellent low-temperature toughness, welding performance and corrosion cracking resistance.

However, the higher the strength of steel, the greater the environmental susceptibility to fracture and the lower the resistance to stress corrosion (hydrogen embrittlement), which limits its application in complex marine environments.

The low-alloy high-strength steel is based on common carbon steel, utilizes a nano phase as a reinforcing second phase particle, improves the alloy strength, and keeps the toughness and the elongation of the alloy basically unchanged. The mechanism is that one or more small amount of alloy elements (the total amount of the alloy elements is less than 3%) are added, and the low-alloy high-strength steel meets the fracture sensitivity resistance and the stress corrosion resistance (hydrogen embrittlement) performance while ensuring the strength in strengthening modes such as solid solution strengthening, fine grain strengthening, dispersion strengthening, desolventizing strengthening, phase change strengthening and the like. The precipitation of microalloy carbide is one of the most important problems in the physical metallurgy process of microalloy steel, and mainly comprises the following steps: the hot rolling process is precipitated in austenite, and the cooling process is precipitated in ferrite or bainite. The carbides precipitated in austenite mainly play a role in fine-grain strengthening by inhibiting recrystallization and grain growth, the size of the carbides is mostly 20-50nm, the precipitation strengthening effect on a matrix is small, and the precipitation strengthening effect in ferrite or bainite is obvious. The size of carbide precipitated in ferrite is generally 10nm or less, and the precipitation strengthening effect is excellent. Through analysis, the carbide precipitation in the ferrite has two modes of interphase precipitation (interface precipitation) and supersaturated precipitation (supersaturated precipitation). Wherein, a large amount of fine interphase precipitation carbide can be obtained by a reasonable control means, and the strength of the steel is greatly improved, so that the interphase precipitation of ferrite is a strengthening way with great development prospect. The bainite ferrite matrix has relatively uniform high-density dislocation, and carbide is almost completely precipitated in the matrix in a dislocation nucleation mode and combined with a bainite structure, so that the strengthening effect can be further enhanced. The precipitation behavior (precipitation amount, shape, size and distribution of precipitation particles) of the microalloy carbonitride in bainite is controlled, and the structure and performance of the steel are effectively improved.

The precipitation "nose tip" temperature of microalloy carbonitride in austenite ranges from about 850 to 900 ℃, and the precipitation "nose tip" temperature in ferrite ranges from about 650 to 750 ℃. If ultra-rapid cooling is performed in this section, the strain-induced precipitation behavior in the austenite region and interphase precipitation during cooling are suppressed, and a large amount of carbonitride may diffuse out at the dynamic transformation point.

Compared with the traditional TMCP, the NG-TMCP can inhibit strain induced precipitation in the hot rolling process in the control of precipitates in steel, so that more microalloy elements are reserved in a ferrite or bainite phase transformation region, the precipitation size is small, the strength of steel is obviously improved, precipitation of carbide in the conventional cooling process during the process of passing through an austenite region and a high-temperature ferrite region is avoided, the growth of the precipitates in the cooling process is inhibited, and the optimal carbide precipitation technological parameters can be obtained through precise cooling path control.

Disclosure of Invention

Based on the design, the invention designs the ultra-low carbon multi-element alloy, improves the hardenability of the steel by adding V, Nb and Ti, enhances the precipitation strengthening effect by forming fine second phase particles by Nb and Ti with carbon and nitride, and forms a mode of coupling strengthening of multiple alloy trace elements among V, Nb and Ti. In-situ synthesis technology and fusion casting smelting technology are adopted, the reinforcement is synthesized in situ through reaction in the smelting process, and the precipitation behavior of microalloy elements, such as the precipitation amount of crystal grains, the morphology, the size and the distribution of precipitated phases, is controlled by combining TMCP (thermo Mechanical Control process) technology, so that the crystal grains are refined, the toughness of the steel is improved, and good formability, weldability and corrosion resistance are obtained.

The technical scheme of the invention is as follows: a dispersion nanometer strengthened 690 steel and a manufacturing method thereof are as follows: the composite material is characterized by comprising the following chemical components in percentage by mass: 0.02-0.30% of C, 0.50-2.0% of Mn, 0.04-0.10% of Al, 0.002-0.013% of V, 0.01-0.3% of Nb, 0.005-0.3% of Ti, 0.01-0.20% of N, and the balance of Fe and inevitable impurities. The manufacturing method comprises the steps of smelting process as a fusion casting process and TMCP process as a heating, rolling and heat treatment process.

Preferably, the manufacturing method of the dispersion nano-reinforced 690 steel is characterized in that the fusion casting process comprises the following steps: the smelting process is adopted, a 150Kg vacuum induction furnace is adopted for smelting and pouring, the shrinkage cavity is cut off, and the billet with the diameter of 100mm multiplied by 120mm is forged. And reheating the steel billet to 1250 ℃, preserving the temperature for 2h, and carrying out 7-pass rolling on a 450nm two-roll reversible hot rolling experimental rolling mill in a laboratory, wherein the final plate thickness is about 12 mm. And (3) placing a 12mm thick steel plate in a box type resistance furnace at 1250 ℃ for heat preservation for 72h to ensure that carbide is completely dissolved in the matrix as far as possible, and then quenching to room temperature.

Preferably, the manufacturing method of the dispersion nano-reinforced 690 steel is characterized in that the heating, rolling and heat treatment processes are TMCP processes, the sample is heated to 1200 ℃, the temperature is kept for 3min, and then the following scheme is carried out: (1) cooling to 900 deg.C at cooling rate of 10 deg.C/s for single compression, wherein the deformation amount is 60%, and the deformation rate is 1.0s^-1Then cooling to 500 ℃, 580 ℃, 620 ℃ and 660 ℃ respectively at a cooling rate of 80 ℃/s, and preserving heat for 5min, 10min, 20min and 60min respectively; (2) cooling to 900 deg.C at cooling rate of 10 deg.C/s for single compression, wherein the deformation amount is 60%, and the deformation rate is 1.0s^-1Then cooled to 500 deg.C, 580 deg.C, 620 deg.C and 660 deg.C at a cooling rate of 20 deg.C/s, respectively, and kept at the temperature for 5min and 20min, respectively.

Preferably, the TMCP process involved in the method for manufacturing dispersion nano-reinforced 690 steel is characterized in that the sample steel is heated to 1200 ℃ for 3min, then cooled to 900 ℃ at a cooling rate of 10 ℃/s, and subjected to 60% deformation, and then kept at 20s to eliminate the temperature gradient inside the sample, and then cooled to room temperature at a cooling rate of v ═ 0.5,1,2,5,10,15,20,25,30,40,60,80 ℃/s, respectively, in order to measure the static continuous cooling transformation curve of the sample steel.

Preferably, the TMCP process involved in the method for manufacturing dispersion nano-strengthened 690 steel is characterized in that the dynamic continuous cooling transformation curve of the sample is measured by: heating the sample to 1200 deg.C at a speed of 10 deg.C/s, holding for 3min, cooling to 900 deg.C at a cooling rate of 10 deg.C/s, holding for 20s to eliminate the temperature gradient in the sample, and holding for 5s^-1Is applied to the specimen at a true strain rate of 0.5, and then the specimen is subjected to a constant strain at a v of 0.5,cooling to room temperature at a cooling rate of 1,2,5,10,15,20,25,30,40 ℃/s.

Preferably, the TMCP process involved in the manufacturing method of the dispersion nano-strengthened 690 steel is characterized in that the two-stage cooling process of slow cooling under the ultra-fast cooling condition: heating the sample to 1200 ℃, preserving heat for 3min, cooling to 900 ℃ at a cooling rate of 10 ℃/s, preserving heat for 5s, cooling to 620 ℃ and 660 ℃ at a cooling rate of 80 ℃/s, cooling to 400 ℃ at the speed of 0.1, 0.25 and 0.5 ℃/s, and rapidly cooling to room temperature.

Advantageous effects

The method improves the hardenability of the steel by adjusting and optimizing the proportion of alloy elements in the steel, forms carbonitride nano-scale second-phase particles at the same time by using trace alloy elements, and controls the precipitation behavior of the micro-alloy elements by adopting a TMCP (thermal mechanical control processing) process so as to refine crystal grains, improve the obdurability of the steel and obtain good formability and corrosion resistance. Because the TMCP process produces the steel with high strength and high toughness under the conditions of not adding excessive alloy elements and not needing complex subsequent heat treatment, the TMCP process has outstanding economic advantages while ensuring the good comprehensive performance of the low-carbon alloy steel.

The study of binding (1), (2) and (3) resulted in:

1) the precipitated phases of the sample steel at all temperatures have the characteristics of taking certain binary precipitation as a main part and compositely precipitating and coexisting other microalloy elements. With the temperature reduction, the Nb-Ti steel takes TiN, NbC and VC precipitated phases as main phases in sequence, and simultaneously has (Nb, V, Ti) (CN) composite precipitated phases.

2) The trace alloy elements Nb, V and Ti in the sample steel enable the CCT curve to integrally move towards the lower right, and the hardenability of the experimental steel is improved.

3) In the isothermal quenching process under the ultra-fast cooling condition and the conventional cooling speed, a ferrite and martensite dual-phase structure is obtained by isothermal quenching in a high-temperature ferrite region, the martensite obtained by the isothermal quenching at 660 ℃ is relatively uniform in distribution, the size distribution is about 4.1nm, and the number of precipitated particles is relatively large; the hardness of isothermal quenching in 60min reaches HV285, and the converted tensile strength is 980 MPa. Isothermal quenching at 580 deg.C for more than 20min to increase average size of precipitated particles from 4.5nm to 4.7nm, with stable size, hardness of HV245 after isothermal quenching for 60min, and converted tensile strength of 850 MPa.

4) According to the ultrafast cooling and slow cooling two-stage cooling process, the average diameters of precipitated phases obtained by slow cooling at the cooling speed of 0.1 ℃/s at 660 ℃ and 620 ℃ are 5.89nm and 3.9nm respectively, the average size of the precipitated phases is smaller as the slow cooling temperature is lower, the higher the dislocation density in the obtained ferrite matrix is, and the larger the number of the precipitated phases is when the slow cooling temperature is 620 ℃.

Drawings

FIG. 1 sample steel static CCT experimental scheme

FIG. 2 dynamic CCT experimental scheme of sample steel

FIG. 3 is a schematic diagram of an isothermal quenching process test scheme

FIG. 4 is a schematic view of a simulation scheme of a slow cooling process

FIG. 5 sample Steel equilibrium Compound carbonitride precipitated phase chemistry

FIG. 6 shows the chemical composition of equilibrium state of compound carbonitride precipitate phase of sample steel

FIG. 7 transition curve of static continuous cooling of sample steel

FIG. 8 metallographic structure of sample steel at different cooling rates without deformation

FIG. 9 dynamic continuous cooling transition curve of experimental steel

FIG. 10 metallographic structure of sample steels at different cooling rates for deformation

Figure 11620 ℃ precipitate morphology and size distribution: (a) (c)20 min; (b) (d)60min

Graph 12660 ℃ precipitate morphology and size distribution: (a) (c)20 min; (b) (d)60min

Figure 13660 ℃ isothermal precipitate morphology and energy spectrum analysis for 60min

FIG. 14580 deg.C isothermal 60min precipitate morphology and energy spectrum analysis

FIG. 15580 ℃ isothermal precipitation particle size distribution: (a)20 min; (b)60min

FIG. 16620 ℃ and 660 ℃ hardness at isothermal temperature as a function of isothermal time

FIG. 17500 ℃ and 580 ℃ changes in isothermal hardness with isothermal time

FIG. 18 is a graph of precipitate morphology and size distribution from 660 ℃ and 620 ℃ with 0.1 ℃/s slow cool

FIG. 19 sample Steel hardness as a function of Slow Cooling Rate

Detailed Description

The technical solutions of the embodiments of the present invention will be described clearly and completely with reference to the accompanying drawings, and it is to be understood that the described embodiments are only a part of the embodiments of the present invention, and not all of the embodiments; all other embodiments, which can be derived by a person skilled in the art from the embodiments of the present invention without any inventive step, are within the scope of the present invention.

A 690 low-carbon microalloyed steel, which comprises the following components: c: 0.02-0.30%; mn: 0.50-2.0%;

c, carbon is an element required for improving the yield property, tensile strength and welding property of the steel, and the mechanism is that the carbon forms carbide with other elements to play the roles of structure strengthening and precipitation strengthening. Meet the performance requirement of 690 steel structural members for ocean engineering, and when the content is lower than 0.02 percent, the solid solution strengthening effect is insufficient. The content thereof is preferably 0.07% or more, more preferably 0.10% or more. However, since an excessive amount of C affects the weldability, toughness and corrosion resistance of the steel, the maximum amount of C should be 0.3% or less. More preferably 0.24% or less, and still more preferably 0.20% or less.

Mn and Mn are elements required for improving the toughness, tensile strength and hardness of steel and improving hardenability, and play a role in solid solution strengthening to make up for the strength loss caused by the reduction of the C content in the low-carbon steel. Meets the performance requirement of 690 steel structure member for ocean engineering, and when the content is less than 0.5%, the strength compensation effect cannot be fully exerted, and the content is preferably more than 0.8%, and more preferably more than 1.15%. However, too high Mn content increases the carbon equivalent thereof, and has a high tendency to segregate, thereby deteriorating the weldability of the material, and reducing the corrosion resistance of the material and the impact resistance of the steel sheet core. Therefore, the maximum Mn content should be 2.00% or less. More preferably 1.60% or less, and still more preferably 1.45% or less.

And further contains V: 0.01-0.3%; nb: 0.005-0.3%; ti: 0.01-0.20% of one or more than one;

v and V are excellent deoxidizing agents for steel, and are elements that improve the impact toughness of steel materials and have a strong secondary hardening effect. The improvement mechanism is that vanadium reacts with carbon and nitrogen to form V (C, N), the vanadium-carbon compound after heat treatment is dispersed and distributed in ferrite of austenite crystal boundary in a fine particle mode to precipitate and separate out, recrystallization of austenite can be inhibited and grain growth can be prevented in the rolling process, and ferrite crystal refining and precipitation strengthening effects are achieved, so that creep and endurance strength and hydrogen corrosion resistance of the steel are improved. The content thereof is preferably 0.01% or more, more preferably 0.06% or more. However, if the vanadium content is too high, the vanadium is locally segregated, which causes pitting corrosion and deteriorates the corrosion resistance. Therefore, the maximum amount of V should be 0.30% or less. More preferably 0.20% or less, and still more preferably 0.12% or less.

Nb and niobium have obvious fine-grain strengthening and medium-level precipitation strengthening effects, and are beneficial to improving the yield strength of the steel plate. Niobium mainly exists in 3 forms in steel: 1) undissolved niobium, wherein undissolved particles prevent the coarsening of the soaking austenite grains and play a role in grain refinement; 2) the solid-solution state of niobium reduces the mobility of grain boundaries and subgrain boundaries. When the grain boundary moves, solute atoms exert resistance, niobium in a solid solution state at high temperature interacts with dislocation in austenite, austenite grains are prevented from being recrystallized, austenite recrystallization temperature is increased, and Nb is subjected to strain induction to precipitate carbonitride on the dislocation, the subgrain boundary and the grain boundary along with the reduction of temperature and the progression of deformation, so that the nucleation rate of recrystallized grains is reduced. 3) Nb which forms carbon, nitrogen or carbon nitride is precipitated and precipitated, the compound is fine and dispersed, recrystallization and structure grain growth are inhibited together with solid-solution niobium when the Nb is precipitated below 900 ℃ in the rolling process, and the Nb is precipitated at low temperature to play a role of precipitation strengthening. In order to effectively exert such an effect, the content is preferably 0.005% or more, and more preferably 0.015% or more. Since niobium is scarce and expensive, the maximum amount of Nb should be 0.07% or less. More preferably 0.05% or less, and still more preferably 0.035% or less.

Ti, titanium produces strong precipitation strengthening and moderate fine grain strengthening. Since Ti is chemically very active, it easily forms a compound with C, N and the like in steel. The micro Ti is adopted to treat the fixed nitrogen and generate TiN, so that the content of solid-solution nitrogen in steel is reduced, the TiN precipitated in a solidified state is fine, austenite grains can be effectively prevented from growing in a heating stage, and meanwhile, the preferential precipitation of Ti and N is favorable for improving the solid solubility of Nb in austenite. The carbide formed by Ti and C has strong binding force, is extremely stable and is not easy to decompose, and Ti can also generate insoluble carbide particles with Fe and C, and the insoluble carbide particles are enriched at the crystal boundary of steel, so that the coarsening of the steel is prevented, the hardenability and the tempering stability are improved, and the secondary hardening effect is realized. The content thereof is preferably 0.01% or more, more preferably 0.05% or more. However, too high titanium content may reduce the toughness of the steel and seriously affect the casting quality of the steel. Therefore, the maximum amount of Ti should be 0.20% or less. More preferably 0.15% or less, and still more preferably 0.10% or less.

Further, when the total of V + Nb + Ti is less than 0.22%, the cost increases and the weld workability is not good, and the plasticity is also affected if the total of V + Nb + Ti exceeds 0.22%.

Or the composition can further contain 0.002-0.013% of N; al: 0.04-0.10% of one or more;

n and nitrogen can enhance the yield strength and the tensile strength of the material through the comprehensive action of fine-grain strengthening and precipitation strengthening, the increase of the nitrogen content enables the steel to have a smaller bainite ferrite lath substructure, the size of a precipitated phase is reduced to below 15nm, and the volume fraction is increased. The content thereof is preferably 0.002% or more, more preferably 0.007% or more; the maximum amount should be 0.013% or less, more preferably 0.01% or less.

Al and Al have the functions of deoxidation and nitrogen fixation in molten steel, and simultaneously an inner rust layer can form an alkaline PH value environment, and a passivation film is formed by passivating an anode to improve the corrosion resistance. Al can refine grains and improve impact toughness, and the content of the Al is preferably more than 0.015 percent, and more preferably more than 0.045 percent; too high an Al content causes casting difficulties and also results in the formation of a large number of dispersed needle-like Al2O3 inclusions in the steel, impairing the integrity of the steel sheet, the plasticity and fatigue creep resistance and the weldability, with the maximum amount being 0.10% or less, more preferably 0.07% or less.

The balance of Fe and inevitable impurities;

TABLE 1690 Steel chemical composition (Unit,%)

C

Mn

Al

V

Nb

Ti

N

0.02～0.30

0.50～2.0

0.04～0.10

0.002～0.013

0.01～0.3

0.005～0.3

0.01～0.20

The specific implementation manner of the patent is as follows:

the casting and smelting process of the sample steel comprises the following steps: the chemical composition of the sample steels is shown in table 1. Smelting and pouring in a 150Kg vacuum induction furnace, cutting off shrinkage cavities, and forging into billets of 100mm multiplied by 120 mm. And reheating the steel billet to 1250 ℃, preserving the temperature for 2h, and carrying out 7-pass rolling on a 450nm two-roll reversible hot rolling experimental rolling mill in a laboratory, wherein the final plate thickness is about 12 mm. And (3) placing a 12mm thick steel plate in a box type resistance furnace at 1250 ℃ for heat preservation for 72h to ensure that carbide is completely dissolved in the matrix as far as possible, and then quenching to room temperature. A standard thermal simulation sample having a diameter of 8mm and a length of 15mm was cut out from the treated steel sheet.

Determination of the static continuous cooling transformation curve of the test steel (static CCT): the method is carried out on an MMS-300 thermal simulation testing machine, as shown in figure 1, a sample is heated to 1200 ℃, is kept for 3min, is cooled to 900 ℃ at a cooling speed of 10 ℃/s, is applied with 60% of deformation, is kept for 20s to eliminate the temperature gradient in the sample, and is then cooled to room temperature at a cooling rate of v ═ 0.5,1,2,5,10,15,20,25,30,40,60 and 80 ℃/s.

Determination of dynamic continuous cooling transformation curve of sample (dynamic CCT): applying MMS-300 thermal simulation testing machine, in order to prevent sample oxidation during heating and deformation, adopting high purity nitrogen protection, heating sample to 1200 deg.C at speed of 10 deg.C/s as shown in FIG. 2, holding for 3min, cooling to 900 deg.C at cooling rate of 10 deg.C/s, holding for 20s to eliminate temperature gradient inside sample, and holding for 5s^-1Is applied to the sample at a true strain of 0.5 and subsequently cooled to room temperature at a cooling rate of v 0.5,1,2,5,10,15,20,25,30,40 ℃/s, respectively.

The isothermal quenching process under the ultra-fast cooling condition and the conventional cooling speed comprises the following steps: the MMS-300 thermal simulation experiment machine is used for researching the influence of the isothermal temperature and the isothermal time in the isothermal quenching process of the sample steel on the precipitation (dispersion) of the carbonitride precipitate of the tissue and the microalloy. The samples were heated to 1200 ℃ as in FIG. 3, incubated for 3min, and then run according to the following protocol: (1) cooling to 900 deg.C at cooling rate of 10 deg.C/s for single compression, wherein the deformation amount is 60%, and the deformation rate is 1.0s^-1Then cooling to 500 ℃, 580 ℃, 620 ℃ and 660 ℃ respectively at a cooling rate of 80 ℃/s, and preserving heat for 5min, 10min, 20min and 60min respectively; (2) cooling to 900 deg.C at cooling rate of 10 deg.C/s for single compression, wherein the deformation amount is 60%, and the deformation rate is 1.0s^-1Then cooled to 500 deg.C, 580 deg.C, 620 deg.C and 660 deg.C at a cooling rate of 20 deg.C/s, respectively, and kept at the temperature for 5min and 20min, respectively.

And (3) a slow-cooling two-stage cooling process under an ultra-fast cooling condition: by using a Formaster-FII full-automatic phase change instrument, as shown in figure 4, a sample is heated to 1200 ℃, is kept warm for 3min, is cooled to 900 ℃ at a cooling rate of 10 ℃/s, is kept warm for 5s, is cooled to 620 ℃ and 660 ℃ at a cooling rate of 80 ℃/s, is cooled to 400 ℃ at a speed of 0.1, 0.25 and 0.5 ℃/s, and is then rapidly cooled to room temperature.

And (3) testing the tissue performance: the thermal simulation sample wire is cut at a position about 1mm below the thermocouple, and is corroded for 15s by 4 percent (volume fraction) nitric acid alcohol solution after mechanical grinding and polishing, and the metallographic structure of the sample wire is observed by an Optical Microscope (OM). In order to observe the precipitation behavior of the experimental steel in detail, a wafer with the thickness of about 300um is cut from a thermal simulation sample, the wafer is mechanically thinned to 50um, then a double-spraying thinning instrument is adopted to carry out double-spraying thinning in a perchloric acid alcohol solution with the volume fraction of 9%, the double-spraying voltage is 30-35V, the temperature is-20 ℃, a Tecnai G2F 20 field emission Transmission Electron Microscope (TEM) is used for observing the size, the morphology and the distribution rule of the precipitated particles, and the chemical components of the precipitated particles are determined by an energy spectrum (EDS). And (3) counting the sizes of the nanometer-sized precipitated particles in the TEM Image by using Image ProPlus professional Image processing software, wherein the number of the precipitated particles counted in each Image is 200-600, 5-10 images are counted in each sample, and the average value is determined as the sizes of the precipitated particles of the samples.

As can be seen from fig. 5 and 6, the sample steel has three types of microalloy face-centered cubic (fcc) phases, and the precipitated phases are mainly TiN, VC, and NbC by composition calculation. Wherein in the high-temperature stage (900-1200 ℃): the fcc #1 precipitated phase of the sample steel contained a small amount of Nb and C, indicating that there is a (VNb, Ti) (CN) composite precipitated phase; in the medium temperature stage (700-900 ℃): the fcc #1 precipitated phase of the sample steel contains a small amount of V, C element, indicating that the temperature range has a (V, Nb, Ti) (CN) composite precipitated phase; in the low-temperature stage (400-700 ℃): the fcc #2 precipitated phase of the sample steel contained a high content of C element, indicating that Ti was mainly precipitated as TiC and Ti (C, N) at low temperature.

The transformation curve of the static continuous cooling of the sample steel is shown in figure 7. The static continuous cooling transition curve of the sample steel was obtained from fig. 7. The temperature at which austenite transforms is lower as the cooling rate increases, and the curve gradually decreases from the upper right to the lower left. The dotted line in the figure is the static continuous cooling transition curve of the Ti steel sample. It can be seen that the ferrite phase transformation area of the microalloyed steel is obviously smaller than that of Ti steel, and the bainite and martensite phase transformation areas of the microalloyed steel are larger at each cooling speed, so that the addition of V and Nb promotes the improvement of the hardenability of the low-carbon steel. At low cooling speed (<5 ℃/s), the trace elements Nb and V not only reduce ferrite phase transformation, but also improve the phase transformation starting temperature of AF/B, so that Nb-containing V steel can obtain more acicular ferrite structures at lower cooling speed.

Metallographic structure plots of the experimental steels of fig. 8 cooled to room temperature at 900 c at different cooling rates. It was found that the structure of the sample steel was an acicular ferrite + pearlite + a small amount of granular bainite structure and a large amount of granular bainite at a cooling rate of 0.5 to 2 ℃/s. At a cooling rate of 5-15 ℃/s, the structure of the sample steel is an all-bainite structure, and the morphology of bainite is complex. When the cooling rate is more than 15 ℃/s, lath martensite starts to appear, and the structure is a mixed structure of bainite and martensite. When the cooling rate reaches 80 ℃/s, all the martensite laths are obtained.

As can be seen from the dynamic continuous cooling transformation curve of the sample of fig. 9, the transformation process is divided into three regions, i.e., ferrite + pearlite transformation (F + P), acicular ferrite/granular bainite/lath bainite transformation (AF/GB/LB), and martensite transformation (M). At cooling rates below 5 ℃/s, the F + P + GB transition occurs. When the cooling rate exceeds 5 ℃/s, the phase transformation completely avoids the occurrence of high-temperature transformation, and the transformation is completely bainite phase transformation within the cooling rate range of 5-25 ℃/s. The martensitic transformation occurs when the cooling rate exceeds 25 c/s. The structure of the sample steel after deformation at different cooling rates is shown in fig. 10. Ferrite transformation can be avoided when the cooling rate is more than 5 ℃/s, martensite transformation can be started when the cooling rate is more than 25 ℃, but bainite transformation still continues to be 40 ℃/s later.

Fig. 11 and 12 show the morphology and size distribution of precipitates after the test steel is cooled to the ferrite zone for austempering.

As can be seen from fig. 11(a) (b) and 12(a) (b): when the ferrite region is isothermal at high temperature, microalloy precipitates in the ferrite are mainly precipitated in the ferrite matrix, except on dislocation lines. Since the energy at the crystal defect such as dislocation line is high, nucleation starts mainly at these positions at the time of precipitation.

From the statistical distribution results of sizes of precipitated particles with isothermal temperatures of 620 ℃ and 660 ℃ in FIGS. 11(c) (d) and 12(c) (d), respectively, the average sizes of the precipitated particles at the isothermal temperature of 620 ℃ for 20min and 60min are 2.2nm and 2.7nm, respectively, and the number of the precipitated particles is similar; the average sizes of the precipitated particles at 660 ℃ were 3.5nm and 4.1nm at 20min and 60min, respectively. The number of the precipitated particles is about 430/um²And 370 pieces/um²And the number of precipitated particles isothermal for 60min is large, which shows that the precipitation equilibrium is not reached even isothermal for 20min at 660 ℃, the temperature of 660 ℃ is closer to the maximum nucleation rate temperature, and the temperature of 620 ℃ is closer to the nose point temperature of a precipitation curve.

FIG. 13 is a graph showing the energy spectrum analysis of precipitated particles from the steel sample. Obtaining a precipitate through energy spectrum analysis; the particles are (Nb, V, Ti) (CN) composite precipitates.

As can be seen from FIG. 14, fine spherical precipitated phases can be observed at different isothermal times after bainite is rapidly cooled to 580 ℃, the size is about 1-10 nm, and the distribution is uniform. The spectrum analysis revealed that the precipitates were Nb complex-precipitated (Nb, V, Ti) (CN), Ti complex-precipitated, and Ti (N, C) precipitated particles, where Nb/(Nb + Ti) ═ 0.135.

As can be seen from FIG. 15, the growth of precipitated particles occurred after 20min isothermal treatment at 580 deg.C, and the average size of the precipitated particles increased from 4.5nm to 4.7nm at a temperature of 20min isothermal treatment, with almost no change. Nanoscale precipitates were also observed in the bainitic ferrite matrix laths and along the prior austenite grain boundaries.

FIGS. 16 and 17 show the hardness of the austempered steel at 620/660 ℃ and the austempered steel at 500/580 ℃ as a function of the isothermal time. Generally, when the temperature is constant at 620 ℃, the temperature is lower than 660 ℃ or the like. The solid solution strengthening and the precipitation strengthening are combined. The tissues obtained by the experimental steel at the temperature of 620 ℃ and the temperature of 660 ℃ in isothermal different time are ferrite and martensite, so that the influence of the phase change strengthening on the isothermal hardness change at the two temperatures is small. The lower the temperature, the lower the solid solubility of the solid solution atoms, the more supersaturated solid solution atoms will desolventize in the subsequent isothermal process, so the isothermal solid solution strengthening effect at 660 ℃ is stronger. On the other hand, the number of precipitated particles at 660 ℃ isothermal temperature is significantly increased as compared with 620 ℃ isothermal temperature, and the contribution of precipitation strengthening action is large, so that the hardness of 660 ℃ isothermal temperature is higher than that of 620 ℃ isothermal temperature. The hardness at 500 ℃ and 580 ℃ has approximately the same trend along with the change of time, the hardness is reduced along with the prolonging of the isothermal time, the hardness is also the result of the interaction of solid solution strengthening and precipitation strengthening, the humidity temperature of a bainite area and the like is relatively low, the atomic diffusion is slow, and particularly the temperature of 500 ℃ is close to the Ms point. In addition, the bainite region has different isothermal transformation types, lath bainite transformation occurs isothermally at 500 ℃, granular bainite transformation occurs isothermally at 580 ℃, and the hardness of lath bainite is higher than that of granular bainite. Because more nano-scale microalloy carbonitride is precipitated at the temperature of 580 ℃, the precipitation strengthening contribution is larger, and the hardness is less reduced, the hardness of the 580 ℃ isothermal steel is higher along with the increase of the isothermal time.

As shown in FIG. 18, the morphology and the size distribution of precipitated phases obtained by slow cooling at 660 ℃ and 620 ℃ at a cooling rate of 0.1 ℃/s show that the average diameters of the precipitated phases gradually cooled from 660 ℃ and 620 ℃ are 5.89nm and 3.9nm, respectively, and the average size of the precipitated phase becomes smaller as the slow cooling temperature becomes lower, and the number of the precipitated phases becomes larger as the slow cooling temperature becomes 620 ℃.

As can be seen from fig. 19, the analysis in combination with table 2 shows that the pearlite transformation strengthening effect and the precipitation effect enhance the hardness value at 620 ℃ under slow cooling at a cooling rate of 0.1 ℃/s, the dispersion strengthening effect and the precipitation effect enhance the hardness value at 620 ℃ under slow cooling at a cooling rate of 0.25 ℃/s, and the strengthening effect of acicular ferrite and granular bainite enhance the hardness value at 620 ℃ under slow cooling at a cooling rate of 0.5 ℃/s.

TABLE 2 variation of pearlite amounts and grain sizes with slow cooling rates

The above description is only for the preferred embodiment of the present invention, but the scope of the present invention is not limited thereto, and any person skilled in the art should be considered to be within the technical scope of the present invention, and the technical solutions and the inventive concepts thereof according to the present invention should be equivalent or changed within the scope of the present invention.

Claims

1. The dispersion nano strengthened 690 steel is characterized by comprising the following chemical components in percentage by mass: 0.10-0.20% of C, 1.15-1.45% of Mn, 0.045-0.07% of Al, 0.06-0.12% of V, 0.015-0.035% of Nb, 0.05-0.10% of Ti, 0.01% of N, and the balance of Fe and inevitable impurities; and V + Nb + Ti is less than 0.22 percent;

the dispersion nano-strengthened 690 steel forms a ferrite matrix phase and a carbonitride nano-scale precipitated phase which is precipitated in the ferrite matrix phase in a dislocation nucleation mode and is uniformly distributed; the carbonitride nano precipitated phase mainly comprises TiN, NbC and VC precipitated phases, and simultaneously has a composite precipitated phase strengthened by coupling of a plurality of alloy trace elements formed among carbon nitride, V, Nb and Ti; the diameter of the carbonitride nano precipitated phase is 1-10 nm.

2. A method of manufacturing the dispersion nano-strengthened 690 steel of claim 1, wherein: the steel comprises a casting process and a TMCP process, wherein the steel comprises the following chemical components in percentage by mass: 0.10-0.20% of C, 1.15-1.45% of Mn, 0.045-0.07% of Al, 0.06-0.12% of V, 0.015-0.035% of Nb, 0.05-0.10% of Ti, 0.01% of N, and the balance of Fe and inevitable impurities; and V + Nb + Ti is less than 0.22 percent;

the TMCP process takes ultra-fast cooling as a core, and the specific process is as follows:

heating the sample to 1200 ℃, preserving heat for 3min, cooling to 900 ℃ at a cooling rate of 10 ℃/s, preserving heat for 5s, cooling to 620 ℃ and 660 ℃ at a cooling rate of 80 ℃/s, cooling to 400 ℃ at speeds of 0.1, 0.25 and 0.5 ℃/s, and rapidly cooling to room temperature to obtain dispersion nano strengthened 690 steel;

the casting process comprises the following steps: adopting a smelting process, adopting a 150Kg vacuum induction furnace to smelt and pour, cutting off a shrinkage cavity, and forging into a billet with the thickness of 100mm multiplied by 120 mm; reheating the steel billet to 1250 ℃, preserving heat for 2h, and carrying out 7-pass rolling on a 450nm two-roll reversible hot rolling experimental rolling mill in a laboratory, wherein the final plate thickness is 12 mm; and (3) placing a 12mm thick steel plate in a box type resistance furnace at 1250 ℃ for heat preservation for 72h to ensure that carbide is completely dissolved in the matrix as far as possible, and then quenching to room temperature.