KR20190131408A - Precipitation hardening and grain refined hot-rolled high strength abnormal steel sheet with a tensile strength of at least 600 MPa and a method of manufacturing - Google Patents

Abstract

The present invention is a method for producing an abnormal steel sheet, in weight percent, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al A process for producing a liquid steel comprising a chemical composition consisting of 0.01 to 0.1, N -0.007 max, Nb: 0.005-0.035 and V-0.06 max, a process of continuously casting the liquid steel to slabs, a finish rolling temperature (FRT ) Hot rolling the slab into a hot rolled sheet at 840 ± 30 ° C, hot rolled on a runout table with a cooling rate of 40-70 ° C / s to achieve an intermediate temperature (T _INT ) (720 ≤ T _INT ≤ 650) Cooling the steel sheet, naturally cooling the hot rolled steel sheet for a duration of 5 to 7 seconds, and quenching the hot rolled steel sheet at a cooling rate of 40-70 ° C./s to martens in austenite where residual carbon is concentrated. Transforming to a site to achieve a coiling temperature of less than 400 ° C.

Description

Precipitation hardening and grain refined hot-rolled high strength abnormal steel sheet with a tensile strength of at least 600 MPa and a method of manufacturing

The present disclosure relates to a process for manufacturing hot rolled high strength dual phase steel. The present disclosure also relates to hot-rolled high strength anomalous tissue steels having a tensile strength of greater than 600 MPa and 25% total elongation.

Automobile fuel consumption and its emissions are one of the main causes of air pollution. Lightweight, environmentally friendly vehicle designs are needed to solve the pollution problem. Successful lightweight cars will need to use advanced high strength steel (AHSS) sheets. However, AHHS steel sheet is not easily applicable to various automotive parts because of poor formability. Therefore, there is an increasing demand for the ductility and formability required for AHSS steel sheet. Therefore, to solve the current scenario, high combinations of excellent uniform elongation, working hardening rate and total elongation for automotive components such as wheel-web applications There is a need for development of hot rolled steel sheets having tensile strength.

Therefore, it is necessary to develop hot rolled steel sheets having not only a minimum tensile strength of 600 MPa but also good formability and surface quality to replace existing grade steels used for automobile structural and web-wheel applications.

European patent EP1398392A1 and US patent US8337643 disclose a method for producing hot rolled bistructure (ferrite + martensitic) steel with a minimum tensile strength of 590 MPa. The disclosed steels contain high content of Si (at least 0.5 wt% for European patents and 0.2 wt% for US patents), although the required strength levels have been achieved. The presence of Si will result in a surface scale, commonly referred to as a tiger mark.

European patent EP2053139B1 discloses a method of heat-treating a hot rolled steel sheet after the forming process to achieve varying tensile strength in the range of 440 to 640 MPa.

However, the heat treatment after the molding process, which is an essential part of the present disclosure, is likely to increase the processing cost and thus is not suitable for mass production.

European patent EP2578714A1 discloses a method for producing a hot rolled steel sheet having a minimum tensile strength of 590 MPa with good bake hardenability and stretch-flangeability. According to the disclosed method, the steel described above should contain 1.7 to 2.5 wt% Mn. When such a large amount is added, Mn tends to cause segregation in the center of the thickness direction, which not only induces cracking during the press forming process but also causes inconsistency in achieving the desired elongation flangeability.

In developing steel, it is also important to understand the wheels of automobiles. The car wheel consists of a disc and a rim. While the disk is press molded, the rim is flared and then roll formed after flash butt welding. Thus, the material needed to form the disk needs to have good deep drawability, stretch formability and stretchability, while the material needed to form the rim has good formability after welding. Need to have. After the wheel disk and the rim are formed by the respective processes, they are assembled by spot welding or arc welding. Thus, materials for both rim and disk use must have good spot weldability. In terms of application, the most important functional requirement for automobile wheels is durability, which can be increased by increasing the fatigue strength of the wheel material.

Various recent studies have shown that both precipitation hardened steels and dual phase (DP) steels are suitable for application to wheel discs. In view of the fatigue strength, the upper limit of tensile strength of steel for use in wheels is ~ 600 MPa (or 85 ksi) [T. Irie, K. Tsunoyama, M. Shinozaki and T. Kato: SAE Paper No. 880695, 1988]. This is because an increase in tensile strength exceeding 600 MPa results in notch sensitivity, which results in a drop in fatigue strength. Hot rolled DP steel with a tensile strength of 600 MPa (or HR-DP 600) is very widely used for wheel discs because of its excellent strength and formability as well as its excellent stretchability (high n value) and spot weldability. have. However, it is difficult to manufacture HR-DP 600 steel in any mill, because in order to obtain the desired microstructure characteristics which in turn determine the final mechanical properties, numerous process parameters, eg final rolling temperature This is because the cooling rate, etc., need to be optimized and fine tuned in consideration of the mill configuration, for example the length of the run out table, the volume of water available and the like. All existing patents and literature contain significant amounts of Si to increase ferrite strength to extend the fatigue life of steel.

In order to overcome the aforementioned limitations inherent in the prior art, an object of the present disclosure is to provide a hot rolled precipitation strengthened high strength having a tensile strength exceeding 600 MPa with a lower percentage of Si. It is to provide a method for manufacturing a dual phase steel sheet).

It is still another object of the present disclosure to provide a method of manufacturing a hot rolled precipitate reinforced high strength abnormal tissue steel sheet having a lower percentage of Si.

It is yet another object of the present disclosure to provide a hot rolled precipitate reinforced high strength abnormal steel sheet having a tensile strength exceeding 600 MPa with a lower percentage of Si.

It is yet another object of the present disclosure to provide a hot rolled precipitate reinforced high strength abnormal tissue steel sheet having a lower percentage of Si.

The present disclosure provides a method for manufacturing the abnormal tissue steel sheet. The method is wt%, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to 0.1, N -0.007 preparing a liquid steel comprising a chemical composition consisting of max, Nb: 0.005-0.035 and V-0.06 max; Continuous casting of the liquid steel into slabs; Hot rolling the slab with a hot rolled steel sheet at a finish rolling temperature (FRT) 840 ± 30 ° C .; The process of cooling the hot-rolled steel sheet on the run-out table with 70 ℃ / s - a medium temperature (T _INT) the cooling rate of 40 to achieve (720≤T _INT ≤ 650); Naturally cooling the hot rolled steel sheet for 5-7 seconds; And quenching the hot rolled steel sheet at a cooling rate of 40-70 ° C./s to transform the austenite containing residual carbon into martensite to achieve a coiling temperature of less than 400 ° C. do.

1 illustrates various processes for manufacturing a high strength abnormal tissue steel according to one embodiment of the present disclosure.
2 illustrates a schematic diagram of a cooling profile for obtaining high strength abnormal tissue steel according to one embodiment of the disclosure.
3 illustrates a tensile stress-strain graph of strip 1 according to one embodiment of the disclosure.
4 illustrates an optical micrograph of strip 1 (Nital etched) according to one embodiment of the disclosure.
5 is a Lepera etched sample white phase, martensite (α), according to one embodiment of the present disclosure; Dark phase: illustrates an optical image of ferrite α '.
6 illustrates an optical image of a Lepera etched sample in which particulates as small as 2 μm can be found, according to one embodiment of the disclosure.
7 shows a scanning electron microscope image of strip 1 according to an embodiment of the present disclosure.
Figure 8 (a) is a bright field TEM photograph of one of the precipitates (precipitates) in the ferrite matrix; (B) is a dark field image of FIG. 8 (a); 8 (c) shows a selected area diffraction pattern from Nb (C, N) precipitates; 8 (d) is a dark system image showing Nb (C, N) precipitates; 8 (e) is an EDS spectrum of the precipitate; And (f) are diagrams illustrating the composition of the precipitate.

Various embodiments of the present disclosure provide a method for manufacturing a dual phase steel sheet, the method comprising:

in wt%, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to 0.1, N-0.007 max, Nb: Preparing a liquid steel comprising a chemical composition consisting of 0.005-0.035 and V-0.06 max;

Continuous casting of the liquid steel into slabs; Hot rolling the slab with a hot rolled steel sheet at a finish rolling temperature (FRT) 840 ± 30 ° C .; Cooling the hot rolled steel sheet on a Run Out table at a cooling rate of 40-70 ° C./s to achieve an intermediate temperature T _INT (720 ≦ T _INT ≦ 650); Naturally cooling the hot rolled steel sheet for 5 to 7 seconds; And quenching the hot rolled steel sheet at a cooling rate of 40-70 ° C./s to transform martensite from austenite with residual carbon concentrated to martensite to achieve a coiling temperature of less than 400 ° C. Process.

Another embodiment of the present disclosure, in wt%, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to An abnormal tissue steel sheet comprising a chemical composition consisting of 0.1, N-0.007 max, Nb: 0.005-0.035 and V-0.06 max is provided.

1, a method 100 of manufacturing an abnormal tissue steel sheet is shown. In step 104 a liquid steel is produced. The composition of the liquid steel is wt%, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to 0.1, N 0.007 max, Nb: 0.005-0.035 and V-0.06 max.

The addition of each alloying element and the limits imposed on its respective components are essential to achieving the desired microstructure and properties.

C : 0.03-0.12%: Carbon is one of the most effective and economic reinforcing elements. Carbon combines with Nb or V to form carbides or carbonitrides that lead to precipitation strengthening. This requires at least 0.03% C in steel. However, to have good weldability, the carbon content should be limited to less than 0.12%.

Mn : 0.8-1.5%: Manganese not only gives solid solution strengthening for ferrite, but also lowers the austenite to ferrite transformation temperature to reduce the grain size of ferrite. Purify. However, the Mn level cannot increase above 1.5%, because at this high level it increases center line segregation during continuous casting.

Si <0.1 wt.%: Like Mn, silicon is a very efficient solid solution strengthening element. However, Si causes surface scale problems in hot rolling and should be limited to less than 0.1% to prevent the formation of surface scales.

Nb : up to 0.035%: Niobium is the strongest microalloying element for grain refinement even when added in very small amounts. When in solid solution, this lowers the transformation temperature from austenite to ferrite to not only fine the ferrite particle size but also promote the formation of lower transformation products such as bainite. However, to ensure the effectiveness of Nb, no precipitation should be allowed before the transformation temperature is reached. The maximum Nb content is limited to 0.035% to ensure that the total Nb content remains in solution and added alone before rolling starts.

V : up to 0.06%: microalloying with vanadium leads to not only grain refinement but also precipitation strengthening. Since the solubility of vanadium in austenite is higher than that of other trace alloy elements, it is likely that it remains in solution prior to transformation. During phase transformation, vanadium precipitates as carbides and / or nitrides at grain boundaries depending on relative carbon and nitrogen content, resulting in precipitation refinement as well as grain refinement. It is necessary to add Nb or V to achieve the desired strengthening. Both may also be added. If V is added alone, up to 0.06 wt% is required.

P : max. 0.025%: The higher the phosphorus content, the lower the toughness and weldability due to segregation of P into grain boundaries, so the phosphorus content should be limited to a maximum of 0.025%.

S : 0.008% at maximum: The sulfur content needs to be limited, otherwise it will result in very high levels of inclusions that degrade moldability.

N <0.007: Too high an N content results in high melting temperatures of Nb (C, N), reducing the effect of Nb. Lowering the nitrogen level is also positive for resistance to inter-crystalline stress-corrosion cracking, as well as for ageing stability and toughness at the heat affected zone of the weld seam. Affects. Therefore, the N level should preferably be kept below 0.007.

Al 0.01 to 0.1: Al is used to remove undesirable oxygen from the molten steel, so the steel contains some amount of Al, which may be up to 0.05 wt%. Excessive (high amounts) Al is an important issue in steel fabrication because it reduces not only nozzle clogging during casting but also high temperature deformation of the cast slab. Therefore, Al needs to be limited to 0.1 wt%.

In process 108 the liquid steel is continuously cast into slabs.

The liquid steel of the specified composition is first subjected to continuous casting in a conventional continuous caster or thin slab caster. When casting in a thin slab caster, the temperature of the casting slab should not drop to a temperature below 950 ° C. This is because Nb precipitation occurs when the thin slab temperature drops below 950 ° C. Subsequently, it becomes difficult to completely dissolve the precipitates in the subsequent reheating process, making them ineffective for precipitation strengthening.

Reheating : After casting a slab having the specific composition described above, the slab is reheated to a temperature of 1100 to 1200 ° C. for 20 minutes to 2 hours. The reheat temperature must be higher than 1100 ° C. to completely dissolve the precipitates of Nb and / or V which may have formed in the previous processes. Higher reheat temperatures above 1200 ° C. are undesirable because they lead to grain coarsening and / or excessive scale loss of austenite.

In process 112, the slab is hot rolled into a hot rolled steel sheet at a finish rolling temperature (FRT) 840 ± 30 ° C.

Hot rolling consists of a roughing step at a temperature above the recrystallization temperature and a finishing rolling step at a temperature below the recrystallization temperature, wherein the rolling is a conventional hot strip rolling mill. hot strip mill). When continuous strip processing is used to produce such steel in the absence of a separate roughing mill, the deformation schedule results in the casting structure being destroyed at the initial stand. It is designed in such a way that the finishing takes place at a temperature lower than the recrystallization temperature. More specifically, finish rolling in either setting must be performed at a temperature T _FRT given at 840 +/- 30 ° C.

Laminar cooling at run-out table (ROT) : The hot rolled sheet in process 116 is cooled at the run-out table at a cooling rate of 40-70 ° C./s. The cooling rate is maintained to achieve an intermediate temperature (T _INT ) 720 ≦ T _INT ≦ 650.

The cooling rate will have to be higher than 40 ° C./sec to prevent pearlite formation. The formation of pearlite or degenerate pearlite deteriorates both tensile strength and stretch-flangeability. High cooling rates also lead to a drop in the ferrite starting temperature which results in a refinement of the ferrite particle size (particle size). This also prevents the growth of ferrite. By increasing the cooling rate and controlling the rolling schedule, the desired particle size of 2-6 μm can be achieved. The cooling rate should not exceed 70 ° C./s since no desired amount of ferrite will be formed. This rapid cooling continues to an intermediate temperature, the intermediate temperature (T _INT ) should be 650 <T _INT <720 ° C.

In process 120, the strip is allowed to cool naturally while being transported on the ROT. The duration of air cooling is very important and is 5 to 7 seconds. If the strip is cooled for less than 5 seconds, a sufficient amount of ferrite will not form. On the other hand, if the strip is allowed to air cool for more than seven seconds, the amount of martensite will be insufficient.

During this time, austenite transforms into ferrite. However, because there is not enough time for complete metamorphosis, the entire austenite will not transform into ferrite. As a result, the ferrite will not be able to accommodate the average carbon content in the steel, so the residual austenite will be concentrated to carbon at the end of natural cooling.

The strip is naturally cooled in step 120 and then rapidly cooled further in step 124. This ensures the transformation of the austenitic austenitic martensite with residual carbon. The cooling rate during this period is 40 to 70 ° C / s to achieve a winding temperature of less than 400 ° C. The winding temperature may be as low as about 100 ° C.

Contributions to reinforcement from solid solution elements and trace alloy elements are limited. In addition, the extent of possible grain refinement by controlled rolling and cooling is limited to 2 μm, thereby obtaining high strength abnormal tissue steels.

The microstructure obtained comprises martensite particles / phases in the ferrite matrix. The microstructures described above are uniform or, in other words, the martensite phase is uniformly distributed throughout the ferrite matrix. In addition, bainite or degenerate pearlite / pearlite and grain boundary cementite are avoided, and high strength anomalous steel sheets achieve good work hardening rates, low yield points and continuous yield. do. The contribution of each microstructural component is described below:

 a) Ferrite: The hot rolled steel sheet according to the present disclosure has a ferrite (volume basis) of 75-90%. Ferrite is enhanced by solid solutions that enhance the contribution of Mn. By using suitable processing conditions the particle size is limited to 2-5 μm. This particle tabletting of ferrite leads to the strengthening of the ferrite, the amount of ferrite being determined by the Hall-Petch relationship. It is also precipitate strengthened by the formation of fine Nb, V (CN) precipitates.

b) Martensite: The amount of martensite in the microstructure is 10-25% (by volume). Reinforcement from martensite results from its structure, carbon content and higher dislocation density.

c) Bainite: The amount of martensite in the microstructure is less than 5% (by volume).

The high-strength abnormal tissue steel sheet obtained an improved fatigue life due to the presence of fine precipitates in the ferrite matrix bonded with martensite as the second phase.

The yield stress of the obtained high strength abnormal steel sheet is 350 to 500 MPa. The tensile strength obtained is at least 600 MPa. The minimum uniform elongation is 16% and the minimum total elongation is 22%.

In addition, the strain hardening exponent ("n") of the high strength abnormal steel sheet is 0.15-0.16. Yield to Tensile strength ratio of ideal tissue steel is 0.6-0.8, and hole expansion ratio in punched condition is about 40%.

Experimental analysis

For illustrative purposes only, slabs of the composition according to Method 100 (strip 1) (given in Table 1) were continuously cast in a CSP mill. The slab was hot rolled. ROT cooling was performed in accordance with the present disclosure, the cooling profile of which is provided in FIG. 2. The mechanical properties of the steel sheet are listed in Tables 2, 3 and 4. The microstructures of the steel are shown in FIGS. 4, 5, 6 and 7. From the mechanical properties and microstructure achieved, it is clear that the targeted properties can be achieved when the chemical properties and ROT cooling parameters meet the requirements of the present disclosure.

Optical (both Nital and Le pera etch) and SEM microstructures are shown in FIGS. 4, 5, 6 and 7 consisting of ferrite and martensite. Tensile test samples of 50 mm gauge length were prepared according to ASTM E8 standard. A graph of a typical tensile test is provided in FIG. 3. The newly developed steel has a minimum 600MPa tensile strength, 16% uniform elongation and at least 22% total elongation, and the strip has a high strain hardening coefficient of 0.15 and a yield ratio between 0.6 and 0.8 (yield strength versus tensile strength). Will be apparent from the above figures and tables. The steel described above has a dispersion of fine precipitates in the ferrite matrix. The stagnation of these deposits is confirmed using Energy Dispersive Spectroscopy (EDS) and Selective Area Diffraction (SAD) techniques in the TEM. The precipitate is mainly Nb (C, N) as described in FIGS. 8A-8F. The steel also has a very fine average particle size of less than 3 μm.

Composition of test samples (% by weight) Strip 1 C Mn Si Cr Nb Al P S N 0.058 1.18 0.02 0.59 0.02 0.036 0.02 0.003 <0.005

Tensile Properties of Test Samples Sample ID Explanation YS ( MPa ) TS ( MPa ) TEL ( % ) YS / TS N Strip 1 Rolling direction (RD) 441 652 25 0.68 0.16 457 666 25 0.69 0.16 437 653 24 0.67 0.16 Average RD 445 657 25 0.68 0.16 Transverse direction (TD) 447 658 22 0.68 0.15 458 655 23 0.70 0.15 444 655 22 0.68 0.16 Average TD 450 656 22 0.69 0.15

Hole Expansion of Test Samples Sample ID Sample No.
Punched  Dia. (D1) Extended Dia (D2) HER = (D2-D1) / D1 HER * 100 ( % ) Remarks
Strip 1 One 10.04 13.48 0.34 34.0 Multiple extensions 2 10.04 14.85 0.48 48.3 Too big crack 3 10.03 14.88 0.49 49.2 Too big crack 4 10.04 14.14 0.40 40.0 One small crack 5 10.03 13.97 0.39 38.5 One small crack

Quantification of microstructural constituents Sample ID Ferrite (%) Martensite ( % ) Bainite (%) Particle size (μm) Hardness( VHN  with 0.1kgf) Strip 1 80 17 <5 2-3 208 ± 10

Claims (21)

In the method of manufacturing the abnormal tissue steel sheet,
In weight percent, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to 0.1, N -0.007 max, Nb A process for producing a liquid steel comprising a chemical composition consisting of 0.005-0.035 and V-0.06 max;
Continuous casting of the liquid steel into slabs;
Hot rolling the slab with a hot rolled steel sheet at a finish rolling temperature (FRT) 840 ± 30 ° C .;
Cooling the hot rolled steel sheet on the runout table at a cooling rate of 40-70 ° C./s to achieve an intermediate temperature T _INT (720 ≦ T _INT ≦ 650);
Naturally cooling the hot rolled steel sheet for a duration of 5 to 7 seconds; And
Quenching the hot rolled steel sheet at a cooling rate of 40-70 ° C./s to transform martensite from austenite with residual carbon concentrated to achieve a coiling temperature of less than 400 ° C. How to manufacture.
The method of claim 1, wherein the abnormal tissue steel sheet is reheated to a temperature in the range of 1100-1200 ° C. for dissolution of the precipitate.
The method of claim 1, wherein the yield stress of the abnormal steel sheet is 350-500 MPa.
The method of claim 1, wherein the abnormal tissue steel sheet has a tensile strength of at least 600 MPa.
The method of claim 1, wherein the abnormal tissue steel sheet has a minimum uniform elongation of 16%.
The method of claim 1, wherein the abnormal tissue steel sheet has a minimum total length of 22%.
The method of claim 1, wherein the stress hardening index (“n”) of the abnormal tissue steel sheet is 0.15-0.16.
The method of claim 1, wherein the yield to tensile strength (ratio) of the abnormal tissue steel is 0.6-0.8.
The method of claim 1, wherein the hole expansion ratio in the punched condition of the abnormal tissue steel sheet is about 40%.
The method of claim 1, wherein the abnormal tissue steel sheet has, by volume, 75-90% ferrite, 10-25% martensite, and <5% bainite.
The method of claim 1, wherein the grain size of the high strength abnormal tissue steel is 2-5 μm.
In the abnormal tissue steel plate,
By weight, C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S-0.008 max, P-0.025 max, Al-0.01 to 0.1, N-0.007 max, Nb: An ideal tissue steel sheet comprising a chemical composition consisting of 0.005-0.035 and V-0.06 max.
The abnormal tissue steel sheet of claim 12, wherein the yield stress of the abnormal steel sheet is 350-500 MPa.
The abnormal tissue steel sheet of claim 12, wherein the abnormal steel sheet has a tensile strength of at least 600 MPa.
The abnormal tissue steel sheet of claim 12, wherein the abnormal tissue steel sheet has a minimum uniform elongation of 16%.
13. The abnormal tissue steel sheet of claim 12, wherein the abnormal tissue steel sheet has a minimum total growth rate of 22%.
The abnormal tissue steel sheet of claim 12, wherein the stress hardening index (“n”) of the abnormal steel sheet is 0.15 to 0.16.
13. The abnormal tissue steel sheet of claim 12 wherein the yield to tensile strength (ratio) is 0.6-0.8.
13. The abnormal tissue steel sheet of claim 12 wherein the hole expansion ratio under punching conditions is about 40%.
The abnormal tissue steel sheet according to claim 12, wherein the abnormal tissue steel sheet has a volume of 75-90% ferrite, 10-25% martensite, and <5% bainite.
The abnormal tissue steel sheet according to claim 12, wherein the grain size of the abnormal steel sheet is 2-5 탆.

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