CN108467997B - Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof - Google Patents
Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof Download PDFInfo
- Publication number
- CN108467997B CN108467997B CN201810224259.2A CN201810224259A CN108467997B CN 108467997 B CN108467997 B CN 108467997B CN 201810224259 A CN201810224259 A CN 201810224259A CN 108467997 B CN108467997 B CN 108467997B
- Authority
- CN
- China
- Prior art keywords
- temperature
- steel
- cooling
- yield strength
- less
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Active
Links
Images
Classifications
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/74—Methods of treatment in inert gas, controlled atmosphere, vacuum or pulverulent material
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0273—Final recrystallisation annealing
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
Landscapes
- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Heat Treatment Of Sheet Steel (AREA)
- Heat Treatment Of Steel (AREA)
Abstract
A precipitation strengthening ferritic steel with yield strength of 1100MPa grade comprises the following components in percentage by weight: c: 0.133-0.162%, Si: 0.01 to 0.09%, Mn: 1.94-2.32%, P: less than or equal to 0.008 percent, S: less than or equal to 0.003 percent, Cr: 0.21-0.32%, V: 0.21-0.34%, Ti: 0.162 to 0.196%, Mo: 0.173 to 0.328%, Nb: 0.091-0.132%, N: less than or equal to 0.005 percent. The production method comprises the following steps: heating the casting blank after conventional smelting and casting to form a blank; two-stage hot rolling; laminar cooling; coiling; naturally cooling to room temperature and then pickling; warm rolling; annealing; and (6) cooling. The invention can ensure that the yield strength of the steel is 1102-1143 MPa, the tensile strength is 1151-1189 MPa, the elongation is more than or equal to 17%, the average grain size of ferrite is 0.7-1.2 mu m, the precipitation strengthening contribution amount is 382-417 MPa, and the steel has excellent plasticity, forming performance and welding performance.
Description
Technical Field
The invention relates to a low-carbon ultrahigh-strength ferrite steel and a production method thereof, which exactly belong to precipitation strengthening type ultrafine grained ferrite steel with the yield strength of 1100MPa and the production method thereof, and are suitable for the fields of automobiles, transportation, engineering machinery and the like.
Background
In recent years, high-performance automobile steel is rapidly developed, is regarded as a key for automobile weight reduction, and is one of important means for realizing energy conservation and emission reduction of automobiles. The technology for researching and developing high-performance automobile steel has important significance for the development of steel and automobile industry in China. The ultrahigh-strength ferritic steel is an important development direction of high-performance automobile steel, has excellent plasticity, formability, welding performance and the like, is particularly suitable for manufacturing automobile parts with complex structures, and has simple production process and low cost. However, the strength of ferritic steel is low, and how to improve the strength becomes a key point for developing ultra-high strength ferritic automobile steel. At present, the ultra-high strength steel mainly comprises martensite steel and bainite steel, and the steel has ultra-high strength but low elongation and cannot meet the requirement of manufacturing automobile parts with complex structures. In addition, the production of the steel needs to add a large amount of alloy elements and needs to go through a complicated heat treatment process, the manufacturing flow is long, and the production cost is high, as searched:
the Chinese patent application No. 201210117567.8 discloses a non-quenched and tempered hot rolled strip steel with yield strength higher than 900MPa and a preparation method thereof, which adopts low-carbon component design, and the component range is as follows: 0.06-0.12% of C, 0.10-0.30% of Si, 0.80-1.20% of Mn, 0.00-0.04% of Nb0.00, 0.00-0.04% of V, 0.02-0.10% of Ti, 0.8-1.20% of Cr0.10-0.30% of Mo0.10, 0.001-0.003% of B, less than 0.012% of P and less than 0.01% of S, and the steel strip with yield strength of not less than 900MPa, tensile strength of not less than 940MPa and elongation after fracture of 12-16% is produced by adding microalloy elements and adopting the technologies of rolling and cooling control, rolling and cooling control and tempering. The metallographic structure is a structure of bainite/martensite or tempered bainite/tempered martensite with fine grains and a small amount of residual austenite. The yield strength of the steel reaches 900MPa, but the elongation is only 12-16%, the main strengthening mechanism is bainite/martensite phase transformation strengthening, and the steel has high strength but relatively low elongation. The use requirements of higher-level components are still not met.
The Chinese patent application No. 201610713632.1 discloses a thin hot-formed steel with tensile strength more than or equal to 1100MPa directly rolled by a thin slab and a production method thereof, and the steel comprises the following components: c: 0.12-0.16%, Si: 0.15 to 0.20%, Mn: 0.7-1.0%, P is less than or equal to 0.02%, S is less than or equal to 0.008%, Als: 0.015 to 0.060%, Cr: 0.15-0.20%, Ti: 0.005-0.02% or Nb: 0.005-0.02% or V: 0.005-0.02% or a mixture of two or more thereof at an arbitrary ratio, B: 0.0005-0.0020% and N is less than or equal to 0.005%. The method adopts hot rolling raw materials produced by continuous casting and rolling of sheet billets to heat and austenitize; and then, stamping and forming by using a die, and then quenching to finally obtain the hot forming steel with the tensile strength of more than or equal to 1100MPa, wherein the elongation rate is not more than 9%, and the quenched finished product structure is martensite which is mainly strengthened by martensite phase transformation, so that the strength is high, but the elongation rate is relatively low.
However, the existing ferritic steel can meet the requirement on elongation, but the strength is lower, and the yield strength of the existing ferritic steel is generally less than 700 MPa. How to increase the strength grade of ferritic steel becomes the key to the development of ultra-high strength ferritic steel. The ferritic steel has excellent plasticity, formability, welding performance and the like, is particularly suitable for manufacturing automobile parts with complex structures, and has simple production process and low cost. Ultra-high strength ferritic steels are an important development direction for high performance automotive steels.
Disclosure of Invention
The invention aims to overcome the defects of the existing ferritic steel and provides the precipitation strengthening ferritic steel which can ensure that the yield strength of the steel is 1102-1143 MPa, the tensile strength is 1151-1189 MPa, the elongation is more than or equal to 17%, the average grain size of ferrite is 0.7-1.2 mu m, and the precipitation strengthening ferritic steel has excellent plasticity, formability, welding performance and the like and a production method thereof.
The measures for realizing the aim are as follows:
a precipitation strengthening ferritic steel with yield strength of 1100MPa comprises the following components in percentage by weight: c: 0.133-0.162%, Si: 0.01 to 0.09%, Mn: 1.94-2.32%, P: less than or equal to 0.008 percent, S: less than or equal to 0.003 percent, Cr: 0.21-0.32%, V: 0.21-0.34%, Ti: 0.162 to 0.196%, Mo: 0.173 to 0.328%, Nb: 0.091-0.132%, N: less than or equal to 0.005 percent, and the balance of Fe and inevitable impurities; the metallographic structure is full ferrite; mechanical properties: the yield strength is 1102-1143 MPa, the tensile strength is 1151-1189 MPa, the elongation is more than or equal to 17%, and the average grain size of ferrite is 0.7-1.2 mu m.
Preferably: the content of C is 0.135-0.160 wt%.
Preferably: the Mn content is 1.97-2.2% by weight.
Preferably: the Nb content is 0.099-0.127 wt%.
Preferably: the weight percentage content of V is 0.21-0.33%.
Preferably: the Ti content is 0.165-0.193% by weight.
Preferably: the weight percentage content of Mo is 0.185-0.297%.
A method for producing precipitation-strengthened ferritic steel having a yield strength of 1100MPa, comprising the steps of:
1) heating the casting blank after conventional smelting and casting to form a blank, wherein the temperature of the heated casting blank is controlled to be 1331-1352 ℃, and the heating time is 139-166 min;
2) carrying out two-stage hot rolling: wherein the rough rolling finishing temperature is controlled to be 1109-1121 ℃, and the cumulative reduction rate is 88-92%; controlling the finish rolling temperature to be 781-803 ℃, and controlling the accumulated reduction rate to be 87-91%;
3) carrying out laminar cooling, and cooling to the coiling temperature at the cooling speed of 79-87 ℃/s;
4) coiling, wherein the coiling temperature is controlled to be 583-607 ℃;
5) naturally cooling to room temperature and then carrying out conventional acid washing;
6) carrying out warm rolling, controlling the warm rolling temperature to be 243-332 ℃, and controlling the accumulated reduction rate to be 78-92%;
7) annealing under the protection of a total hydrogen atmosphere, controlling the annealing temperature to be 573-627 ℃, and preserving heat for 41-57 min at the temperature;
8) cooling to room temperature at a cooling rate of not less than 107 ℃/s.
Preferably: the warm rolling temperature is 264-311 ℃, and the cumulative reduction rate is 77-86%.
Preferably: the annealing temperature is 585-620 ℃.
The action and mechanism of the main strengthening element and process in the invention
C: the low carbon design is selected to reduce the amount of cementite in the steel microstructure and inhibit the formation of pearlite. When the C content is less than 0.13%, if the carbon content is too low, it is difficult to form a nano-scale precipitate by bonding with the microalloying elements Ti, Nb, V, etc., and thus it is difficult to perform the precipitation strengthening function. When the C content is more than 0.162%, the welding and low temperature toughness are rapidly deteriorated, so that the C content is controlled to 0.13 to 0.162%, preferably 0.135 to 0.160% by weight.
Si: while the steel serves to strengthen the steel by solid solution and contains a deoxidizing element, the content of Si should be controlled to 0.01% or more, but if the content of Si is more than 0.09%, the formation of an inner rust layer is promoted, which makes descaling difficult during rolling, thereby deteriorating the surface quality of the steel strip, and if the content of Si is too high, the weldability of the steel is lowered, and thus, the content of Si is controlled to 0.01 to 0.09%.
Mn: the Mn-Mn alloy is an important toughening element in steel, improves the Mn content in the steel, can expand a gamma region, reduces the transformation temperature, expands the rolling range and promotes the grain refinement, thereby increasing the toughness of the steel, and the impact transformation temperature is hardly changed, so the Mn content is more than 1.94 percent, in addition, when the Mn content is higher, casting blank cracks are easily generated in the continuous casting process, the welding performance of the steel can be reduced when the Mn content is too high, and the Mn content is controlled to be 1.94-2.32 percent.
P: p in steel deteriorates the toughness of steel, and in particular drastically lowers the low-temperature impact toughness of steel, so that the P content is controlled to 0.008% or less.
S: MnS inclusions generated by excessively high S content in steel can cause obvious difference in longitudinal and transverse properties of steel and deteriorate low-temperature toughness. The S content should be controlled below 0.003%.
Cr: since Cr element in steel can improve the strength and hardness of steel, the Cr content should be controlled to 0.21% or more, and 0.32% or less in consideration of the economy of the composition.
Ti: the Ti can be preferentially combined with N to form TiN when the Ti content is low, the TiN particles are relatively large in size and cannot be dissolved under the high-temperature conditions of heating and welding, the welding performance of the steel is obviously improved, in addition, the TiN can also effectively pin austenite crystal boundaries and help to prevent the austenite crystal grains from growing large, when the Ti content is high, besides the TiN can be formed, the residual Ti in the steel can be combined with C in the steel to form TiC particles with small sizes and can play a role in precipitation strengthening, the Ti content in the steel is too low, the Ti cannot play a role in fine grain strengthening and precipitation strengthening, the Ti content is too high, the Ti cannot be completely dissolved, and unnecessary alloy element loss is caused. Comprehensively considering the solid solubility product of Ti and related elements in austenite and ferrite and the contents of Nb, N, S, C and other elements in steel, the control range of the Ti content in the invention is as follows: 0.162-0.196%, preferably 0.165-0.193% by weight of Ti.
Mo: the Mo content is greater than 0.173%, because the Mo is a precious metal, the Mo content is controlled to be less than 0.328% in view of production cost, and preferably the weight percentage content of the Mo is 0.185-0.297%.
Nb: is a strong carbonitride forming element, and trace Nb in the steel can inhibit the recrystallization of deformed austenite, prevent the growth of austenite grains, improve the recrystallization temperature of the austenite, refine the grains and improve the strength and toughness of the steel. In addition, the precipitation of Nb (CN) in the cooling process can play a role of precipitation strengthening, the mechanical property of the steel is improved, the solid solubility product of Nb and related elements in austenite and ferrite and the contents of elements such as Ti, N, S, C and the like in the steel are comprehensively considered, so the control range of the Nb content is 0.091-0.132%, if the Nb content is less than 0.091%, the Nb cannot play a role in fine grain strengthening and precipitation strengthening, and if the Nb content is more than 0.132%, the Nb element cannot be completely dissolved, and unnecessary alloy element loss is caused. The content of Nb is preferably 0.099-0.127% by weight.
V: the element is a strong carbonitride forming element, in the rolling process of a recrystallization zone, solid-solution V atoms can enable the migration speed of grain boundaries or subboundary to be slower through the interaction with dislocation, further the coarsening of austenite grains is inhibited to a certain extent, in the rolling process of a non-recrystallization zone, a large amount of V (C, N) grains are separated out on the austenite grain boundaries and the subboundary due to the strain-induced precipitation effect, after austenite is recrystallized, the grain boundaries and the subboundary are pinned by the separated grains, therefore, the grains are difficult to grow continuously, in the control rolling process of a non-recrystallization zone, ferrite is easy to nucleate the previously separated V (C, N) precipitates, and further the effect of obviously refining the grains is achieved. The control range of the content of V in the steel comprehensively considers the solid solubility product of V and related elements in austenite and ferrite and the content of elements such as Ti, Nb, N, S, C and the like in the steel is 0.21-0.34%, and the content of V in the steel is preferably 0.21-0.33% by weight.
N: nitrogen in the steel can be combined with elements such as Ti, Nb, V and the like at high temperature to form corresponding compounds, and the compounds can coarsen and grow at high temperature, which seriously damages the plasticity and toughness of the steel. In addition, such coarse carbonitride particles formed at high temperatures contribute little to precipitation strengthening and consume the effective Ti, Nb, and V contents in the steel, so the contents thereof are controlled to 0.005% or less.
The invention controls the warm rolling temperature to be 243-332 ℃ and the accumulated reduction rate to be 78-92% because the common warm rolling is only suitable for steel types which are difficult to be directly cold rolled at room temperature, such as high-carbon and high-alloy steel, and is used for improving the work hardening of the steel in the cold rolling process and improving the machinability of the steel in the cold rolling process. The warm rolling process mainly aims to further refine the ferrite grain size and improve the fine grain strengthening effect at a proper rolling temperature. The warm rolling temperature is higher than 332 ℃, so that ferrite is recrystallized, the grain size is overlarge, and the fine grain strengthening effect is weakened. In addition, the warm rolling temperature is too high, which causes coarsening of second phase particles generated during the hot rolling process, resulting in a reduction in precipitation strengthening effect. When the warm rolling temperature is lower than 243 ℃ and the cumulative deformation is higher than 79%, the deformation resistance is large in the rolling process, and the rolling is difficult. When the accumulated deformation is less than 63%, the accumulated deformation is too small, so that the ferrite grain size is difficult to be further refined, and the ferrite with ultrafine grains is difficult to obtain; preferably, the warm rolling temperature is 264-311 ℃, and the cumulative reduction rate is 77-86%.
The annealing temperature is controlled to 573-627 ℃, and the temperature is kept for 41-57 min; the annealing temperature is preferably 585-620 ℃, because the annealing temperature is higher than 627 ℃, and the holding time is longer than 57 minutes, the second phase particles grow and coarsen, so that the precipitation strengthening effect is weakened, and the annealing temperature is lower than 573 ℃, so that the thermodynamic and kinetic conditions for separating out the second phase particles are poor, the second phase particles are difficult to separate out, and the precipitation strengthening effect is weakened. When the cooling rate after annealing is lower than 107 ℃/s, the second phase particles grow and coarsen during the cooling process, thereby weakening the precipitation strengthening effect. Therefore, the annealing is followed by cooling to room temperature at a cooling rate of not less than 107 ℃/s.
The production method mainly adopts a Ti-Mo-Nb-V composite microalloying technology and adopts a conventional hot rolling, warm rolling and annealing mode to obtain the precipitation strengthening type ultrafine grained ferrite steel with the yield strength ranging from 1102 MPa to 1143MPa, the tensile strength ranging from 1151 MPa to 1189MPa, the elongation rate ranging from 17.3% to 19.6%, the average grain size of ferrite ranging from 0.7 μm to 1.2 μm and the precipitation strengthening contribution amount ranging from 382 MPa to 417 MPa. The invention is characterized in that the warm rolling is carried out after the hot rolling, the difficulty that the direct cold rolling is difficult due to the overhigh hot rolling strength of the steel is solved, the size of ferrite grains is fully refined through the setting of reasonable temperature and deformation in the warm rolling process, and the ferrite steel with the average grain size of 0.7-1.2 mu m ultrafine grains is obtained. Breaks through the limit value of the average grain size of 3.0 mu m of the ferritic steel produced by the traditional method. In addition, through reasonable formulation of hot rolling, warm rolling and annealing processes, fine and dispersed second phase particles are fully separated out while ferrite grains are fully refined, the precipitation strengthening contribution amount is 382-417 MPa, the limit value of 320MPa of the precipitation strengthening contribution amount of the traditional Ti-Nb-V composite micro-alloy steel is broken through, and ultrahigh strengthening of the ferrite steel is realized. The steel plate produced by the invention has high strength and good plasticity.
Drawings
FIG. 1 is a metallographic structure representation of an example of the invention;
FIG. 2 is a graph of the morphology of precipitates in accordance with an embodiment of the present invention.
Detailed Description
The present invention is described in detail below:
table 1 is a list of chemical compositions for each example of the present invention and comparative example;
table 2 is a table of the main process parameters of each example of the present invention and comparative example;
table 3 is a table of the results of the performance tests of the examples of the present invention and the comparative examples.
The production of each embodiment of the invention is carried out according to the following steps:
1) heating the casting blank after conventional smelting and casting to form a blank, wherein the temperature of the heated casting blank is controlled to be 1331-1352 ℃, and the heating time is 139-166 min;
2) carrying out two-stage hot rolling: wherein the rough rolling finishing temperature is controlled to be 1109-1121 ℃, and the cumulative reduction rate is 88-92%; controlling the finish rolling temperature to be 781-803 ℃, and controlling the accumulated reduction rate to be 87-91%;
3) carrying out laminar cooling, and cooling to the coiling temperature at the cooling speed of 79-87 ℃/s;
4) coiling, wherein the coiling temperature is controlled to be 583-607 ℃;
5) naturally cooling to room temperature and then carrying out conventional acid washing;
6) carrying out warm rolling, controlling the warm rolling temperature to be 243-332 ℃, and controlling the accumulated reduction rate to be 78-92%;
7) annealing under the protection of a total hydrogen atmosphere, controlling the annealing temperature to be 573-627 ℃, and preserving heat for 41-57 min at the temperature;
8) cooling to room temperature at a cooling rate of not less than 107 ℃/s.
TABLE 1 tabulated (wt%) chemical compositions for inventive and comparative examples
| Examples | C | Si | Mn | Ti | Mo | V | Nb | N | Cr | S | P |
| 1 | 0.145 | 0.02 | 2.12 | 0.171 | 0.299 | 0.22 | 0.098 | 0.004 | 0.25 | 0.002 | 0.005 |
| 2 | 0.162 | 0.05 | 2.27 | 0.183 | 0.173 | 0.34 | 0.132 | 0.003 | 0.32 | 0.001 | 0.007 |
| 3 | 0.143 | 0.03 | 2.32 | 0.162 | 0.179 | 0.31 | 0.095 | 0.002 | 0.24 | 0.002 | 0.006 |
| 4 | 0.157 | 0.09 | 2.07 | 0.176 | 0.315 | 0.25 | 0.129 | 0.005 | 0.27 | 0.003 | 0.007 |
| 5 | 0.160 | 0.06 | 2.20 | 0.193 | 0.297 | 0.21 | 0.127 | 0.003 | 0.30 | 0.001 | 0.008 |
| 6 | 0.153 | 0.04 | 2.03 | 0.180 | 0.159 | 0.24 | 0.126 | 0.004 | 0.26 | 0.002 | 0.005 |
| 7 | 0.146 | 0.07 | 1.97 | 0.165 | 0.256 | 0.33 | 0.099 | 0.005 | 0.23 | 0.003 | 0.006 |
| 8 | 0.135 | 0.05 | 2.13 | 0.179 | 0.185 | 0.27 | 0.113 | 0.004 | 0.29 | 0.002 | 0.008 |
| 9 | 0.161 | 0.08 | 1.94 | 0.196 | 0.181 | 0.30 | 0.091 | 0.002 | 0.21 | 0.001 | 0.005 |
| 10 | 0.133 | 0.01 | 2.26 | 0.182 | 0.328 | 0.23 | 0.128 | 0.003 | 0.31 | 0.003 | 0.006 |
| Comparative example 1 | 0.021 | 1.33 | 0.22 | 0.011 | 0.010 | 0.02 | 0.001 | 0.03 | 0.005 | 0.102 | 0.037 |
| Comparative example 2 | 0.052 | 0.02 | 2.01 | 0.005 | 0.016 | 0.01 | 0.005 | 0.04 | 0.076 | 0.161 | 0.076 |
| Comparative example 3 | 0.457 | 1.01 | 1.33 | 0.017 | 0.033 | 0.03 | 0.012 | 0.07 | 0.005 | 0.156 | 0.351 |
TABLE 2 List of the main parameters of the processes of the examples of the invention and the comparative examples
TABLE 3 Table of mechanical Properties of each example and comparative example of the present invention
As can be seen from Table 3, the steel plate of the embodiment has the yield strength ranging from 1102 to 1143MPa, the tensile strength ranging from 1151 to 1189MPa, the elongation ranging from 17.3 to 19.6%, the ferrite average grain size ranging from 0.7 to 1.2 μm, the precipitation strengthening contribution amount ranging from 382 to 417MPa, and the 180 DEG cold bending performance is all qualified, while the yield strength of the comparison sample ranges from 311 to 708MPa, the tensile strength ranges from 437 to 752MPa, the elongation ranges from 4.3 to 5.3%, the ferrite average grain size ranging from 15.3 to 19.5 μm, the precipitation strengthening contribution amount ranging from 23 to 42MPa, and the 180 DEG cold bending performance is not all qualified. It can be seen that the performance indexes of the examples in this patent are all better than those of the comparative examples.
The present embodiments are merely preferred examples, and are not intended to limit the scope of the present invention.
Claims (4)
1. A precipitation strengthening ferritic steel with yield strength of 1100MPa comprises the following components in percentage by weight: c: 0.133-0.162%, Si: 0.01 to 0.09%, Mn: 2.12-2.32%, P: less than or equal to 0.008 percent, S: less than or equal to 0.003 percent, Cr: 0.21% or 0.31-0.32% of Cr, V: 0.21-0.34%, Ti: 0.162 to 0.196%, Mo: 0.173 to 0.256%, Nb: 0.113-0.132%, N: less than or equal to 0.005 percent, and the balance of Fe and inevitable impurities; the metallographic structure is full ferrite; mechanical properties: the yield strength is 1102-1143 MPa, the tensile strength is 1151-1189 MPa, the elongation is more than or equal to 17%, and the average grain size of ferrite is 0.7-1.2 mu m.
2. A precipitation-strengthened ferritic steel of 1100MPa class in yield strength as set forth in claim 1 wherein: the content of C is 0.135-0.160 wt%.
3. A method of producing a precipitation-strengthened ferritic steel having a yield strength of the 1100MPa class as claimed in claim 1, comprising the steps of:
1) heating the casting blank after conventional smelting and casting to form a blank, wherein the temperature of the heated casting blank is controlled to be 1331-1352 ℃, and the heating time is 139-166 min;
2) carrying out two-stage hot rolling: wherein the rough rolling finishing temperature is controlled to be 1109-1121 ℃, and the cumulative reduction rate is 88-92%; controlling the finish rolling temperature to be 781-803 ℃, and controlling the accumulated reduction rate to be 87-91%;
3) carrying out laminar cooling, and cooling to the coiling temperature at the cooling speed of 79-87 ℃/s;
4) coiling, wherein the coiling temperature is controlled to be 583-607 ℃;
5) naturally cooling to room temperature and then carrying out conventional acid washing;
6) carrying out warm rolling, controlling the warm rolling temperature to be 243-287 ℃, and controlling the accumulated reduction rate to be 78-92%;
7) annealing under the protection of a total hydrogen atmosphere, controlling the annealing temperature to be 573-627 ℃, and preserving heat for 41-57 min at the temperature;
8) cooling to room temperature at a cooling rate of not less than 107 ℃/s.
4. A method of producing a precipitation-strengthened ferritic steel of the 1100MPa class yield strength as set forth in claim 3 wherein: the annealing temperature is 585-620 ℃.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| CN201810224259.2A CN108467997B (en) | 2018-03-19 | 2018-03-19 | Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| CN201810224259.2A CN108467997B (en) | 2018-03-19 | 2018-03-19 | Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| CN108467997A CN108467997A (en) | 2018-08-31 |
| CN108467997B true CN108467997B (en) | 2020-04-21 |
Family
ID=63264485
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CN201810224259.2A Active CN108467997B (en) | 2018-03-19 | 2018-03-19 | Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof |
Country Status (1)
| Country | Link |
|---|---|
| CN (1) | CN108467997B (en) |
Families Citing this family (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN110747322A (en) * | 2019-09-26 | 2020-02-04 | 中国电力科学研究院有限公司 | 1180MPa nanoparticle reinforced ferrite steel and preparation method thereof |
Family Cites Families (3)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN1223687C (en) * | 2002-08-30 | 2005-10-19 | 上海宝钢集团公司 | Submicron crystalline grain steel plate separated out nanometer and its manufacturing method |
| CN106636911B (en) * | 2017-03-17 | 2018-09-18 | 武汉科技大学 | With the 900MPa grade hot rolled steel sheets and its manufacturing method of sheet billet Direct Rolling |
| CN107190203B (en) * | 2017-05-31 | 2019-07-23 | 武汉钢铁有限公司 | With the yield strength >=800MPa hot rolled sheet and production method of sheet billet Direct Rolling |
-
2018
- 2018-03-19 CN CN201810224259.2A patent/CN108467997B/en active Active
Also Published As
| Publication number | Publication date |
|---|---|
| CN108467997A (en) | 2018-08-31 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| CN112048681B (en) | 980 MPa-grade high-formability cold-rolled DH steel and preparation method thereof | |
| CN108474081B (en) | Steel material for press forming, formed member thereof, and heat treatment method | |
| CN109023036B (en) | Ultrahigh-strength hot-rolled complex-phase steel plate and production method thereof | |
| CN111979489B (en) | 780 MPa-grade high-plasticity cold-rolled DH steel and preparation method thereof | |
| CN107619993B (en) | Cold-rolled martensite steel plate with yield strength of 750MPa and manufacturing method thereof | |
| CN111979490B (en) | High-ductility and high-formability cold-rolled DH590 steel and production method thereof | |
| CN112095046B (en) | Ultrahigh-strength cold-rolled DH1180 steel and preparation method thereof | |
| WO2018036348A1 (en) | Thin thermoformed steel directly rolled using thin slabs and having tensile strength of ≥1500 mpa, and method for producing same | |
| CN111254351B (en) | High-performance hot-rolled wear-resistant steel sheet and production method thereof | |
| CN113388779B (en) | 1.5 GPa-grade ultrahigh-strength high-plasticity high-hole-expansion DH steel plate and preparation method thereof | |
| WO2018036346A1 (en) | Thin thermoforming steel rolled directly from thin slab and having tensile strength greater than or equal to 1900 mpa and production method | |
| JP2015516511A (en) | High-strength cold-rolled steel sheet and method for producing such a steel sheet | |
| CN113549823B (en) | 900 MPa-grade hot-rolled and pickled multiphase steel with low yield ratio and high hole expansion rate and production method thereof | |
| CN110863138B (en) | 1800 MPa-grade hot forming steel and manufacturing method thereof | |
| CN106011644A (en) | High-elongation and high-strength cold-rolled steel plate and preparation method thereof | |
| JP2013057114A (en) | Medium carbon steel plate having excellent workability and hardenability and method for producing the same | |
| CN112877591B (en) | High-strength and high-toughness hardware tool and steel for chain and manufacturing method thereof | |
| CN111647801A (en) | 690 MPa-grade ferrite-martensite dual-phase steel, and preparation method and application thereof | |
| CN113481436A (en) | 800 MPa-grade hot-rolled complex phase steel and production method thereof | |
| CN108467997B (en) | Precipitation-strengthened ferritic steel with yield strength of 1100MPa and production method thereof | |
| CN109182700B (en) | Low-yield-ratio hot-rolled steel plate with excellent hole expansion performance for automobile and manufacturing method thereof | |
| CN108277434B (en) | Precipitation strengthening ferritic steel with yield strength of 900MPa and production method thereof | |
| CN115369324A (en) | 38MnVS5 rare earth free-cutting hot-rolled round steel for automobiles and preparation method thereof | |
| CN110172636A (en) | A kind of low-carbon hot forming steel and preparation method thereof | |
| JPH0425343B2 (en) |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| PB01 | Publication | ||
| PB01 | Publication | ||
| SE01 | Entry into force of request for substantive examination | ||
| SE01 | Entry into force of request for substantive examination | ||
| GR01 | Patent grant | ||
| GR01 | Patent grant |