CN113840933A - Thick steel plate and method for producing same - Google Patents

Thick steel plate and method for producing same Download PDF

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CN113840933A
CN113840933A CN202080029745.4A CN202080029745A CN113840933A CN 113840933 A CN113840933 A CN 113840933A CN 202080029745 A CN202080029745 A CN 202080029745A CN 113840933 A CN113840933 A CN 113840933A
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CN113840933B (en
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名古秀德
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Kobe Steel Ltd
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    • 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
    • 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
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0231—Warm 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
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • 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
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • 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
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • 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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • 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
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • 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
    • 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
    • C21D2201/00—Treatment for obtaining particular effects
    • C21D2201/05—Grain orientation
    • 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

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Abstract

A thick steel sheet having a composition of C, Si, Mn, P, S, Al, Nb, Ti, N and Ca satisfying the requirements, the balance being Fe and unavoidable impurities, and satisfying Di +10Nb as determined by the following formula (1): 1.20 to 2.50, in the crystal grains surrounded by the high angle crystal boundary with the crystal grain orientation difference of more than 15 degrees,the total area fraction SA of crystal grains having a circle equivalent diameter of 7.5 μm or less is 34% or more at the 1/4 position of the plate thickness and 27% or more at the 1/2 position of the plate thickness. Di is 1.16 × ([ C ]]/10)0.5×(0.7×[Si]+1)×(5.1×([Mn]－1.2)+5)×(0.35×[Cu]+1)×(0.36×[Ni]+1)×(2.16×[Cr]+1)×(3×[Mo]+1)×(1.75×[V]+1)×(200×[B]+1) … (1) formula (1) wherein [ element(s) ]]Each represents the content of each element expressed by mass%, and the element not included is 0.

Description

Thick steel plate and method for producing same
Technical Field
The present invention relates to a thick steel plate and a method for manufacturing the same. In particular, the present invention relates to a high-strength thick steel plate having excellent base material toughness, and a method for producing the thick steel plate.
Background
For example, as LPG tanks and the like are becoming larger, there is an increasing demand for thick steel plates having high strength and both low-temperature toughness of the base material and HAZ low-temperature toughness.
As a high-strength steel excellent in HAZ toughness, for example, patent document 1 discloses a steel satisfying predetermined components, satisfying that the C concentration of the average chemical analysis value of the center segregation portion of the steel sheet is 1.2 times or less the average C concentration of the steel sheet, the cleanliness of inclusions measured in JIS standard is 0.03% or less, and the number of oxide inclusions having an average diameter of 10 μm or more observed in the cross section of the steel sheet is 1 or less/1 mm2The number of precipitates of oxide and nitride of 0.05 to 5 μm is 100 or more per 1mm2。
In patent document 2, a low-yield-ratio high-tension steel sheet having excellent base material low-temperature toughness and HAZ low-temperature toughness can be obtained by satisfying predetermined chemical components, setting the acicular ferrite structure fraction to 50% or more, and setting the island Martensite (MA) structure fraction of 1 to 5 μm in terms of average equivalent circle diameter to 3 to 10%.
Patent document 3 discloses a composition satisfying a predetermined requirement, wherein the microstructure is a bainite structure, and the yield strength is 500N/mm2Above and tensile strength of 610N/mm2The above steel material. Further, it is disclosed that the steel material is not required to be subjected to annealing heat treatment for removing residual stress after welding, and is suitable for production of a tank for an LPG or ammonia cargo ship.
Patent document 4 discloses a method for producing a high-strength steel sheet having excellent sr (stress relief) resistance, characterized in that a steel satisfying a predetermined composition and having parameters of 9 × Ceq +4 × P ≥ 4.8 and [ C ]/([ Mo ] + [ Ti ] + [ Nb ] + [ V ]) of 0.6 to 1.7 is heated to a temperature of 1100 to 1300 ℃, hot-rolled at a rolling completion temperature of 750 ℃ or higher, then accelerated-cooled at a cooling rate of 20 ℃/s or higher to a temperature of less than 400 ℃, and immediately thereafter reheated at a temperature rise rate of 0.5 ℃/s or higher to 550 to 700 ℃.
Documents of the prior art
Patent document
Patent document 1: japanese laid-open patent publication No. 8-158006
Patent document 2: japanese laid-open patent publication No. 2009-
Patent document 3: japanese laid-open patent application No. 2008-025014
Patent document 4: japanese laid-open patent publication No. 2007-270194
In patent document 1, although a good balance between strength and toughness can be obtained, in the examples, only a plate thickness of 40mm or less is considered, and a technique considering a thicker steel plate is not proposed. In patent document 2, the low temperature toughness and strength of the base material and the HAZ are attempted to be combined, but the low temperature toughness is evaluated at-60 ℃. In patent document 3, the mechanical properties are evaluated only at the position of t/4 of the sheet thickness, and the mechanical properties up to the inside of the steel sheet are not further considered. Patent document 4 discloses a method for producing a thick steel plate having good mechanical properties even after SR, but the evaluation temperature of toughness is only-10 ℃, and toughness at lower temperatures is not studied.
Disclosure of Invention
The present invention has been made in view of the above circumstances, and an object thereof is to provide a thick steel plate which exhibits an excellent strength-toughness balance in the inner region of the steel plate even if the plate is thick, and particularly exhibits high strength and excellent toughness at a lower temperature than conventional ones, and a method for producing the thick steel plate.
Mode 1 of the present invention is a thick steel plate having a composition satisfying the following requirements
C: 0.020 to 0.070 mass percent,
Si: more than 0 mass% and not more than 0.40 mass%,
Mn: 1.30-1.95% by mass,
P: more than 0 mass% and not more than 0.015 mass%,
S: more than 0 mass% and not more than 0.005 mass%,
Al: 0.005-0.070 wt%,
Nb: 0.015-0.048 wt%,
Ti: 0.005-0.024 mass%,
N: 0.0030 to 0.0080 mass%, and
ca: higher than 0 mass% and not more than 0.0040 mass%,
the balance being Fe and unavoidable impurities,
and satisfies Di +10Nb as determined by the following formula (1): 1.20 to 2.50 of the total weight of the alloy,
the total area fraction SA of crystal grains having a circle-equivalent diameter of 7.5 [ mu ] m or less among crystal grains surrounded by high-angle grain boundaries having a grain orientation difference of 15 DEG or more is 34% or more at the 1/4 position of the plate thickness and 27% or more at the 1/2 position of the plate thickness.
Di＝1.16×([C]/10)0.5×(0.7×[Si]+1)×(5.1×([Mn]－1.2)+5)×(0.35×[Cu]+1)×(0.36×[Ni]+1)×(2.16×[Cr]+1)×(3×[Mo]+1)×(1.75×[V]+1)×(200×[B]+1)…(1)
In the formula (1), [ C ], [ Si ], [ Mn ], [ Cu ], [ Ni ], [ Cr ], [ Mo ], [ V ] and [ B ] respectively represent the contents of C, Si, Mn, Cu, Ni, Cr, Mo, V and B in mass%, and the element not included is 0.
Embodiment 2 of the present invention further includes the thick steel plate according to embodiment 1
Cu: more than 0 mass% and not more than 0.75 mass%, and
ni: more than 0 mass% and not more than 1.4 mass%
One or more elements selected from the group consisting of the elements.
Embodiment 3 of the present invention further includes the thick steel plate according to embodiment 1 or 2
Mo: more than 0 mass% and not more than 0.50 mass%,
V: more than 0 mass% and not more than 0.060 mass%,
Cr: more than 0 mass% and not more than 0.8 mass%, and
b: more than 0 mass% and not more than 0.0007 mass%
One or more elements selected from the group consisting of the elements.
Embodiment 4 of the present invention is the thick steel sheet according to any one of embodiments 1 to 3, further comprising a steel sheet obtained from REM: more than 0 mass% and not more than 0.0060 mass%, and
zr: higher than 0 mass% and not more than 0.0050 mass%
One or more elements selected from the group consisting of the elements.
Embodiment 5 of the present invention is a method for producing a thick steel sheet according to any one of embodiments 1 to 4, including a step of heating a steel sheet having the composition according to any one of embodiments 1 to 4 to 1020 ℃ or higher and 1200 ℃ or lower, and a hot rolling step after the heating,
in the hot rolling step, the number of passes is set to 3 or more, and the hot rolling and the cooling after the hot rolling are performed so as to satisfy all of the following conditions (a) to (d).
(a) A cumulative reduction ratio of 40% or more in a temperature range of 850 ℃ or lower
(b) The average reduction rate of the final 3 passes of rolling is more than 5.5 percent
(c) The final rolling temperature is 720-830 DEG C
(d) After hot rolling, the steel sheet is cooled from a cooling start temperature of a finish rolling temperature to 690 ℃ to a cooling stop temperature of 320 to 550 ℃ at an average cooling rate of 0.5 to 20 ℃/s.
According to the present invention, even if the steel plate is thick, a thick steel plate exhibiting an excellent strength-toughness balance in the interior of the steel plate, and particularly exhibiting high strength and excellent toughness at a lower temperature than conventional steel plates, and a method for producing the same can be provided.
Drawings
FIG. 1 is a graph showing the relationship between the total area fraction SA and the value Y of crystal grains having an equivalent circle diameter of 7.5 μm or less among crystal grains surrounded by high angle grain boundaries having a grain orientation difference of 15 ° or more.
Detailed Description
In the present invention, intensive studies have been made to obtain a thick steel sheet having an improved balance between strength and toughness at a lower temperature than in the past in a state before SR, that is, in a state immediately after hot rolling. In particular, the following steel sheets have been intensively studied: even when the sheet thickness is large, the strength-toughness balance is excellent at 1/4 and 1/2 positions of the sheet thickness, and specifically, the parameter Y for the strength-toughness balance defined in the present invention is very low, i.e., 20 × vTrs-7 × YP.
As a result, it has been found that the above-mentioned excellent properties can be obtained if the composition of the components is controlled and a predetermined amount of fine acicular ferrite structure, more specifically, acicular ferrite having an equivalent circle diameter of 7.5 μm or less as measured by the method described later is secured at the 1/4 position and 1/2 position of the sheet thickness. Hereinafter, the 1/4 position of the plate thickness is referred to as "t/4 position", and the 1/2 position of the plate thickness is referred to as "t/2 position".
Further, it was found that it is effective to carry out all of the following (a) to (C) in order to form the fine acicular ferrite structure in a predetermined amount. The acicular ferrite is hereinafter referred to as "AF".
(A) Before transformation from the austenite phase to the ferrite phase, sufficient working strain is applied to the austenite phase by hot rolling. The AF crystal grains are generated using the dislocation structure and the deformation band introduced by the processing as nuclei, thereby realizing a fine structure.
(B) Solid solution of Nb is ensured before rolling in a hot non-recrystallized zone. In this way, the processing strain is easily obtained before the phase transformation, and the fine AF crystal grains are easily generated. As described later, it is effective to ensure the solid-solution Nb by increasing the heating temperature before rolling to 1020 ℃ or higher and reducing the reduction rate in rolling at 850 ℃ or higher.
(C) The phase transition temperature to the ferrite phase is appropriately controlled. When the transformation temperature into the ferrite phase is high, a grain boundary ferrite structure is formed before AF generation, and the AF amount decreases. On the other hand, when the transformation temperature into the ferrite phase is low, the martensite structure is generated in a state where the AF structure is not formed. In order to properly control the transformation temperature into the ferrite phase, the C content, Mn content, and the content of at least one of Cu and Ni when contained in the composition, and Di +10Nb may be controlled in respective ranges, and the average cooling rate in a predetermined temperature range after hot rolling is set to 0.5 ℃/s or more as described later.
Hereinafter, the steel structure, the composition of components, the properties and the manufacturing method of the steel plate according to the present invention will be described in order.
1. Steel structure
The steel structure of the steel plate of the present invention will be described in detail below. In the following description of the steel structure, a mechanism that has such a structure and can improve various properties is sometimes described. These are mechanisms considered by the present inventors based on the findings obtained at present, but it should be noted that they do not limit the technical scope of the present invention.
In the present invention, among the crystal grains surrounded by the high angle grain boundaries having a grain orientation difference of 15 ° or more, the crystal grains having an equivalent circle diameter of 7.5 μm or less are defined as a fine Acicular Ferrite (AF) structure. In the present invention, "forming a predetermined amount of fine Acicular Ferrite (AF) structure" means that crystal grains surrounded by high-angle grain boundaries having a grain orientation difference of 15 ° or more, that is, crystal grains having an equivalent circle diameter of 7.5 μm or less are secured by 34% or more at the t/4 position and 27% or more at the t/2 position, as described below.
In the present invention, in order to obtain a thick steel sheet exhibiting an excellent strength-toughness balance even when the sheet thickness is thick, the total area fraction SA of crystal grains having an equivalent circle diameter of 7.5 μm or less, that is, the area fraction of a fine Acicular Ferrite (AF) structure, among crystal grains surrounded by high-angle grain boundaries having a grain orientation difference of 15 DEG or more are defined at both the t/4 position and the t/2 position. Specifically, the total area fraction SA satisfies 34% or more at the t/4 position and 27% or more at the t/2 position. the total area ratio SA at the t/4 position is preferably 35% or more, more preferably 36% or more. The total area ratio SA at the t/2 position is preferably 28% or more, and more preferably 30% or more. From the viewpoint of obtaining an excellent strength-toughness balance, the upper limit of the total area fraction SA at each of the t/4 position and the t/2 position is not particularly limited. Considering the manufacturing conditions of the thick steel sheet according to the embodiment of the present invention, the upper limit of the total area fraction SA at the t/4 position is about 80%, and the upper limit of the total area fraction SA at the t/2 position is about 70%.
The microstructure other than the fine acicular ferrite includes bainite, ferrite, cementite, retained austenite, martensite, and the like. The acicular ferrite having an equivalent circle diameter of more than 7.5 μm may be present as long as the total area fraction SA falls within the range defined in the embodiment of the present invention.
2. Composition of
The composition of the steel plate of the present invention will be described below.
C: 0.020% to 0.070% by mass
C has the effect of properly controlling the ferrite transformation temperature and suppressing the generation of intergranular ferrite, which acts as a brittle fracture starting point and causes deterioration of the strength-toughness balance, before AF generation. From the viewpoint of exerting this effect, the C content is 0.020% by mass or more, preferably 0.023% by mass or more, and more preferably 0.030% by mass or more. On the other hand, if the amount of C is excessive, a hard martensite structure is generated, and acts as a brittle fracture origin, thereby deteriorating the strength-toughness balance. Therefore, the C content is 0.070 mass% or less. The amount of C is preferably 0.065 mass% or less, and more preferably 0.060 mass% or less.
Si: more than 0 mass% and not more than 0.40 mass%
Si is a deoxidizing element and is contained in an amount of more than 0 mass%. The Si content may be 0.05 mass% or more, and further 0.10 mass% or more. On the other hand, if the amount of Si is excessive, a hard martensite structure is generated, and the hard martensite structure acts as a brittle fracture origin, and the strength-toughness balance is deteriorated. Therefore, the Si content is 0.40 mass% or less, preferably 0.38 mass% or less, and more preferably 0.35 mass% or less.
Mn: 1.30 to 1.95% by mass
Mn has the effect of appropriately controlling the ferrite transformation temperature, and suppressing the generation of intergranular ferrite, which acts as a brittle fracture starting point and causes deterioration of the strength-toughness balance, before AF generation. From the viewpoint of exerting this effect, the Mn content is 1.30 mass% or more, preferably 1.40 mass% or more, and more preferably 1.45 mass% or more. On the other hand, if the Mn content is excessive, a hard martensite structure is generated, and the hard martensite structure acts as a brittle fracture origin, and the strength-toughness balance is deteriorated. Therefore, the Mn content is 1.95 mass% or less, preferably 1.90 mass% or less, and more preferably 1.80 mass% or less.
P: more than 0 mass% and not more than 0.015 mass%
P is an impurity element, and if it is contained excessively, grain boundaries are embrittled, and the strength-toughness balance is deteriorated. Therefore, the P content is 0.015 mass% or less. The P amount is preferably 0.008 mass% or less, and more preferably 0.007 mass% or less. On the other hand, in the industry, it is difficult to achieve a P amount of 0 mass%, so the lower limit of the P amount is higher than 0 mass%.
S: more than 0 mass% and not more than 0.005 mass%
S is an impurity element, and if it is contained excessively, grain boundaries are embrittled, and the strength-toughness balance is deteriorated. Therefore, the S content is 0.005 mass% or less. The amount of S is preferably 0.004 mass% or less, more preferably 0.003 mass% or less. On the other hand, in the industrial field, it is difficult to make the S amount to 0 mass%, so the lower limit of the S amount is higher than 0 mass%.
Al: 0.005 to 0.070 mass%
Al is a deoxidizing element. Since oxygen in the steel is reduced by sufficient deoxidation and deterioration of the strength-toughness balance by oxides is suppressed, the Al content is 0.005 mass% or more. The amount of Al is preferably 0.010 mass% or more, and more preferably 0.015 mass% or more. On the other hand, if the amount of Al is excessive, coarse oxides are formed, and the strength-toughness balance is deteriorated. Therefore, the Al amount is 0.070 mass% or less, preferably 0.050 mass% or less, and more preferably 0.045 mass% or less.
Nb: 0.015 to 0.048 mass%
Nb is an element that promotes AF generation. In order to sufficiently generate fine AF structure and obtain a good strength-toughness balance, the Nb content is 0.015 mass% or more, preferably 0.016 mass% or more, and more preferably 0.018 mass% or more. On the other hand, if the amount of Nb is excessive, a hard martensite structure is generated, and this structure acts as a brittle fracture origin and deteriorates the strength-toughness balance. Therefore, the Nb content is 0.048 mass% or less, preferably 0.045 mass% or less, and more preferably 0.040 mass% or less.
Ti: 0.005 to 0.024 mass%
Ti is an element contributing to improvement of HAZ toughness by forming TiN. From the viewpoint of exerting this effect, the Ti content is 0.005 mass% or more, preferably 0.007 mass% or more, and more preferably 0.009 mass% or more. On the other hand, if the amount of Ti is excessive, coarse crystalline TiN is generated, and the balance between strength and toughness deteriorates. Therefore, the Ti content is 0.024% by mass or less, preferably 0.022% by mass or less, and more preferably 0.020% by mass or less.
N: 0.0030 to 0.0080 wt%
N is an element contributing to improvement of HAZ toughness by forming TiN. From the viewpoint of exerting this effect, the N amount is 0.0030 mass% or more, preferably 0.0032 mass% or more, and more preferably 0.0035 mass% or more. On the other hand, if the amount of N is excessive, the amount of solid solution N increases, and the strength-toughness balance deteriorates. Therefore, the N content is 0.0080% by mass or less, preferably 0.0075% by mass or less, and more preferably 0.0070% by mass or less.
Ca: more than 0 mass% and not more than 0.0040 mass%
Ca is a deoxidizing element, and the content thereof is higher than 0 mass%. In addition, when the amount of Mn is large in the steel, coarse MnS is likely to be formed at the t/2 position due to the thickening of Mn during casting, and it is considered that the toughness at the t/2 position is likely to be lowered. In order to suppress formation of MnS, the Ca content is preferably higher than 0 mass%, more preferably 0.0008 mass% or more, and still more preferably 0.0010 mass% or more. On the other hand, if the amount of Ca is excessive, coarse oxides are formed, and the strength-toughness balance is deteriorated. Therefore, the Ca amount is 0.0040 mass% or less, preferably 0.0028 mass% or less, and more preferably 0.0025 mass% or less.
The balance being Fe and unavoidable impurities. As inevitable impurities, trace elements such As, Sb, and Sn are allowed to be mixed in depending on the conditions of raw materials, manufacturing facilities, and the like. Further, for example, as P and S are used, the smaller the content is, the more preferable the content is, and therefore, the impurities are inevitable, but there are elements separately defined in the above-described manner with respect to the composition range thereof. Therefore, in the present specification, the "inevitable impurities" constituting the balance are a concept excluding elements whose composition ranges are separately defined.
The thick steel sheet according to the embodiment of the present invention may contain the above-described elements in the component composition. The optional elements described below may not be included, but if necessary, they may be included together with the above elements, so that high strength and the like can be more easily achieved. In addition, a desired structure can be more easily secured, and the strength-toughness balance required in the embodiment of the present invention can be more easily achieved. Hereinafter, the selection element is explained.
From Cu: more than 0 mass% and below 0.75 mass% and Ni: more than 0 mass% and not more than 1.4 mass% of one or more elements selected from the group consisting of
These elements have the effect of properly controlling the ferrite transformation temperature and suppressing the generation of intergranular ferrite before AF generation, which acts as a brittle fracture starting point and causes deterioration of the strength-toughness balance. From the viewpoint of exerting this effect, Cu is contained in an amount of preferably more than 0 mass%, more preferably 0.05 mass% or more, further preferably 0.10 mass% or more, and further preferably 0.15 mass% or more. When Ni is contained, it is preferably higher than 0 mass%, more preferably 0.10 mass% or more, further preferably 0.15 mass% or more, and further preferably 0.20 mass% or more. On the other hand, if these elements are excessive, a hard martensite structure is formed, and this structure acts as a brittle fracture initiation point, and causes deterioration of the strength-toughness balance. Therefore, the Cu content is preferably 0.75 mass% or less, more preferably 0.70 mass% or less, and further preferably 0.68 mass% or less. The amount of Ni is preferably 1.4 mass% or less, more preferably 1.2 mass% or less, and further preferably 1.0 mass% or less.
From Mo: more than 0 mass% and less than 0.50 mass%, V: more than 0 mass% and not more than 0.060 mass% of Cr: above 0 mass% and below 0.8 mass% and B: more than 0 mass% and not more than 0.0007 mass% of one or more elements selected from the group consisting of
These elements are effective for improving strength. From the viewpoint of exerting this effect, Mo is contained in an amount of preferably more than 0 mass%, more preferably 0.05 mass% or more, and further preferably 0.10 mass% or more. When V is contained, it is preferably higher than 0% by mass, more preferably 0.01% by mass or more, and further preferably 0.02% by mass or more. When Cr is contained, it is preferably higher than 0 mass%, more preferably 0.10 mass% or more, and further preferably 0.20 mass% or more. When B is contained, it is preferably higher than 0% by mass, and more preferably 0.0003% by mass or more.
On the other hand, if the content of these elements is excessive, a hard martensite structure is formed, and this structure acts as a brittle fracture initiation point, and the strength-toughness balance is deteriorated. Therefore, the Mo amount is preferably 0.50 mass% or less, more preferably 0.45 mass% or less, and further preferably 0.40 mass% or less. The V amount is preferably 0.060 mass% or less, more preferably 0.050 mass% or less, and further preferably 0.045 mass% or less. The amount of Cr is preferably 0.8 mass% or less, more preferably 0.70 mass% or less, and still more preferably 0.60 mass% or less. The amount of B is preferably 0.0007% by mass or less, more preferably 0.0006% by mass or less.
From REM: above 0 mass% and below 0.0060 mass% and Zr: more than 0 mass% and not more than 0.0050 mass% of one or more elements selected from the group consisting of
These elements are deoxidizing elements. In order to exert this effect, when REM is contained, it is preferably higher than 0 mass%, more preferably 0.0010 mass% or more, and further preferably 0.0015 mass% or more. When Zr is contained, it is preferably higher than 0 mass%, more preferably 0.0010 mass% or more, and further preferably 0.0012 mass% or more. On the other hand, if these elements are excessive, coarse oxides are formed, and the strength-toughness balance is deteriorated. Therefore, the REM amount is preferably 0.0060 mass% or less, more preferably 0.0050 mass% or less, and still more preferably 0.0045 mass% or less. The Zr content is preferably 0.0050 mass% or less, more preferably 0.0045 mass% or less, and further preferably 0.0040 mass% or less. The REM is meant to include lanthanoids (15 elements from La to Lu), Sc (scandium), and Y (yttrium).
Di +10 Nb: 1.20 to 2.50(Di is obtained by the following formula (1))
Di＝1.16×([C]/10)0.5×(0.7×[Si]+1)×(5.1×([Mn]－1.2)+5)×(0.35×[Cu]+1)×(0.36×[Ni]+1)×(2.16×[Cr]+1)×(3×[Mo]+1)×(1.75×[V]+1)×(200×[B]+1)…(1)
In the formula (1), [ C ], [ Si ], [ Mn ], [ Cu ], [ Ni ], [ Cr ], [ Mo ], [ V ] and [ B ] respectively represent the contents of C, Si, Mn, Cu, Ni, Cr, Mo, V and B in mass%, and the element not included is 0.
Di +10Nb is a parameter which affects the ferrite transformation temperature. By appropriately controlling the ferrite transformation temperature, generation of grain boundary ferrite before AF generation can be suppressed, and excessive generation of pearlite structure and martensite structure other than AF can be controlled to promote AF generation. From these viewpoints, in the embodiment of the present invention, Di +10Nb is set to be in the range of 1.20 to 2.50. If the Di +10Nb value is too small, a pearlite structure is formed together with grain boundary ferrite, and the strength-toughness balance is deteriorated. Therefore, Di +10Nb is 1.20 or more. Di +10Nb is preferably 1.25 or more, more preferably 1.30 or more. On the other hand, if the value of Di +10Nb is too large, a hard martensite structure is formed, and the hard martensite structure acts as a brittle fracture origin, and the strength-toughness balance is deteriorated. Therefore, Di +10Nb is 2.50 or less. Di +10Nb is preferably 2.20 or less, more preferably 2.00 or less.
The thickness of the thick steel sheet according to the embodiment of the present invention is preferably 16mm or more, more preferably 30mm or more, further preferably 40mm or more, and particularly preferably higher than 40 mm. The upper limit of the plate thickness is not particularly limited, but is preferably 80mm or less, for example.
The thick steel plate of the present invention has the following characteristics, and is suitable for manufacturing large LPG tanks, ships, and the like.
3. Characteristics of
In the present invention, the value Y, which is a parameter represented by the following formula (2), is used for the evaluation of the balance between the strength and the toughness at a lower temperature than the conventional one. The Y value is represented by the following formula (2) and includes a yield strength YP and a ductile-brittle transition temperature vTrs. As the yield strength YP, SS curve (also referred to as "stress-strain curve") in the following formula (2), in the case of a circular arc whose yield point is unknown, 0.2% yield strength (0.2YS) is used, and in the case of having a yield point, YP is used.
The thick steel sheet according to the embodiment of the present invention is evaluated to have an excellent strength-toughness balance when Y at the t/4 position is less than-5200 and Y at the t/2 position is less than-4700 in the C direction perpendicular to the rolling direction in a hot-rolled state.
Y＝20×vTrs－7×YP…(2)
The value of Y at the t/4 position is preferably-5300 or less, more preferably-5400 or less, and the value of Y at the t/2 position is preferably-4800 or less, more preferably-5000 or less. Lower values of these values mean a better strength-toughness balance.
In the embodiment of the present invention, the above parameters may be achieved as characteristics. The yield strength YP and ductile-brittle transition temperature vTrs are respectively provided on the premise that the above parameters are achieved, and for example, the yield strength YP is preferably in the range of 350 to 550MPa, and the ductile-brittle transition temperature vTrs is preferably in the range of less than-70 ℃, and more preferably in the range of less than-80 ℃. Further, depending on the measurement position, in the t/4 position, on the premise that the above parameters are achieved, for example, the yield strength YP is preferably 410MPa or more, more preferably 430MPa or more, further preferably 460MPa or more, further preferably 480MPa or more, particularly 500MPa or more, and the ductile-brittle transition temperature vTrs is preferably-90 ℃ or less, more preferably-95 ℃ or less, further preferably-100 ℃ or less, and further preferably-110 ℃ or less. In addition, in the t/2 position, on the premise of achieving the above parameters, for example, the yield strength YP is preferably 400MPa or more, more preferably 430MPa or more, further preferably 460MPa or more, still more preferably 480MPa or more, particularly 500MPa or more, and the ductile-to-brittle transition temperature vTrs is preferably-80 ℃ or less, more preferably-90 ℃ or less, further preferably-95 ℃ or less, still more preferably-100 ℃ or less.
4. Manufacturing method
The method for producing a thick steel sheet according to the present invention comprises a step of heating a steel sheet having the above-described composition to a temperature higher than 1020 ℃ and lower than 1200 ℃, and a hot rolling step after the heating, wherein the hot rolling step is performed so that the number of passes is 3 or more and the conditions (a) to (d) below are satisfied.
(a) A cumulative reduction ratio of 40% or more in a temperature range of 850 ℃ or lower
(b) The average reduction rate of the final 3 passes of rolling is more than 5.5 percent
(c) The final rolling temperature is 720-830 DEG C
(d) After hot rolling, the steel sheet is cooled from a cooling start temperature of a finish rolling temperature to 690 ℃ to a cooling stop temperature of 320 to 550 ℃ at an average cooling rate of 0.5 to 20 ℃/s.
Hereinafter, each production condition will be described in detail.
[ procedure of heating the steel sheet having the composition to a temperature higher than 1020 ℃ and lower than 1200 ]
When the heating temperature is 1020 ℃ or lower in the heating in hot rolling, NbC generated in casting cannot be sufficiently dissolved in a solid solution, and an AF generation promoting effect by the dissolved Nb cannot be obtained. Therefore, the heating temperature is higher than 1020 ℃, preferably 1040 ℃ or higher, more preferably 1050 ℃ or higher, and still more preferably 1060 ℃ or higher. On the other hand, when the heating temperature is 1200 ℃ or higher, austenite grains are coarsened and the entire structure is coarsened. Therefore, the heating temperature is less than 1200 ℃, preferably 1180 ℃ or less, and more preferably 1150 ℃ or less.
[ Hot Rolling Process after heating ]
As will be described later, the present invention is characterized in that the average reduction ratio of the final 3-pass rolling is controlled, after the heating, the hot rolling is performed with 3 or more passes, and the total number of passes (the number of passes) is not limited as long as the structure and the properties are not affected. The rolling pass is more preferably 7 passes or more, and still more preferably 10 passes or more, and may be 60 passes or less from the viewpoint of productivity.
In an embodiment of the present invention, the hot rolling is performed so as to satisfy the following conditions (a) to (d) and cooling after the hot rolling is performed.
(a) Cumulative reduction at a temperature range of 850 ℃ or lower: more than 40 percent
In order to obtain a sufficient AF structure, it is necessary to set the cumulative reduction (also referred to as "total reduction") in the hot rolling at a temperature range of 850 ℃. However, when the reduction ratio in the rolling at 850 ℃ or higher is increased, it is considered that NbC is precipitated and solid-solution Nb is decreased in the rolling. From the viewpoint of suppressing the cumulative reduction in rolling at 850 ℃ or higher, the cumulative reduction in the temperature range of 850 ℃ or lower is 40% or higher. The cumulative rolling reduction is preferably 50% or more, more preferably 55% or more. The upper limit of the cumulative rolling reduction is about 80% from the viewpoint of productivity.
(b) Average reduction of final 3 passes of rolling: 5.5% or more
The dislocation structure introduced at the final 3 passes of reduction is transferred to the cooling step in a state in which the dislocation structure is relatively recovered, and therefore, the AF promoting effect is large. In the embodiment of the present invention, sufficient working strain is introduced into the austenite phase in order to introduce the dislocation structure, and therefore the average reduction ratio in the final 3 passes of rolling is set to 5.5% or more. In the embodiment of the present invention, by controlling the average reduction ratio of the final 3-pass rolling, it is possible to obtain a desired structure having an AF structure at a certain level or more even in the steel sheet, particularly at the t/2 position, unlike the conventional method in which the average reduction ratio of the final 3-pass rolling is not controlled. The average reduction ratio is preferably 5.8% or more, and more preferably 6.0% or more. On the other hand, from the viewpoint of rolling mill load, the upper limit of the average reduction ratio is about 20%.
(c) Finishing Rolling Temperature (FRT): 720-830 DEG C
If the temperature of the steel material is higher than 850 ℃, sufficient working strain cannot be introduced into the austenite phase even if the final 3 passes of rolling are performed at the above average reduction ratio, and the AF microstructure amount is insufficient. In the embodiment of the present invention, the finishing temperature is set to 830 ℃ or lower in order to sufficiently secure the AF structure amount. The finish rolling temperature is preferably 820 ℃ or lower, more preferably 810 ℃ or lower. On the other hand, if the finish rolling temperature is lower than 720 ℃, coarse ferrite is formed during rolling, and the toughness is deteriorated. Therefore, the finishing temperature is 720 ℃ or higher, preferably 750 ℃ or higher, and more preferably 760 ℃ or higher.
(d) After hot rolling, the steel sheet is cooled from a cooling start temperature of a finish rolling temperature to 690 ℃ to a cooling stop temperature of 320 to 550 ℃ at an average cooling rate of 0.5 to 20 ℃/s.
By cooling a constant temperature region at an average cooling rate of 0.5 to 20 ℃/s, AF tissue can be sufficiently secured. If the average cooling rate is higher than 20 ℃/s, martensite transformation occurs in a state where the AF microstructure cannot be sufficiently formed, which is not preferable. The average cooling rate is preferably 15 ℃/s or less, and more preferably 12 ℃/s or less. As described above, in order to properly control the transformation temperature into the ferrite phase and sufficiently secure the AF structure, the C content, the Mn content, and the content of at least one of Cu and Ni when contained, and the respective ranges of Di +10Nb in the component composition are controlled so that the average cooling rate is 0.5 ℃/s or more. If the average cooling rate is less than 0.5 ℃/s, coarse grain boundary ferrite is generated during cooling, and the AF microstructure amount is insufficient. The average cooling rate is preferably 2.0 ℃/s or more, and more preferably 3.0 ℃/s or more. As a cooling method of the average cooling rate, for example, water cooling is exemplified.
The start temperature of the cooling is an arbitrary temperature from the finishing temperature to 690 ℃. If the temperature of the steel material is less than 710 ℃, there is a problem that grain boundary ferrite is easily generated before cooling starts, or work strain introduced into an austenite phase is easily recovered. As a result, the AF tissue amount tends to be insufficient. In the embodiment of the present invention, in order to secure a sufficient amount of AF structure, the Start Cooling Temperature (SCT) at the above average Cooling rate is set to 690 ℃. The cooling start temperature is preferably 710 ℃ or higher, and more preferably 720 ℃ or higher.
The Temperature at which the Cooling is completed (FCT) is set to any Temperature of 320 to 550 ℃. If cooling at the average cooling rate, for example, water cooling, is stopped in a temperature range higher than 550 ℃, grain boundary ferrite is generated during slow cooling after the water cooling is stopped, and it is difficult to sufficiently secure the AF structure. Therefore, the end temperature is 550 ℃ or less. The end temperature is preferably 500 ℃ or less, more preferably 480 ℃ or less. On the other hand, if the cooling at the average cooling rate is performed to a temperature range of, for example, less than 320 ℃, martensite transformation occurs in a state where the AF structure is not sufficiently formed. Therefore, the end temperature is 320 ℃ or higher. The end temperature is preferably 340 ℃ or higher, more preferably 360 ℃ or higher.
The production method of the present invention is not particularly limited except for the hot rolling step, and may be carried out under ordinary conditions.
If a person skilled in the art who has come into contact with the above-described method for producing a high-strength steel sheet according to the embodiment of the present invention, then, through trial and error, there is a possibility that the high-strength steel sheet according to the embodiment of the present invention may be obtained by a production method different from the above-described production method.
[ examples ] A method for producing a compound
The embodiments of the present invention will be described in more detail below with reference to examples. The present invention is not limited to the following examples, and can be suitably modified and implemented within the scope that can meet the aforementioned and later-described objectives, and all of them are included in the technical scope of the present invention.
1. Sample preparation
Steels having the compositions shown in Table 1 were melted in a 150kg VIF (Vacuum Induction Furnace) or an actual converter, and slabs obtained by casting were hot-rolled under various conditions shown in Table 2 to obtain steel sheets having the thicknesses shown in Table 2. In the hot rolling, the total pass number (the number of rolling passes) is higher than 20. In Table 1, "-" indicates an unintentional addition. In table 2, FRT represents the finish rolling temperature, SCT represents the cooling start temperature, and FCT represents the cooling end temperature. The cooling stop temperature FCT is measured at 1 to 3 points along the longitudinal direction of the surface of the steel sheet by a radiation thermometer, and the average value thereof is calculated. Further, FRT and SCT were obtained by measuring 1 point on the surface of the steel material with a radiation thermometer.
[ TABLE 1 ]
Figure BDA0003309859100000151
[ TABLE 2 ]
Figure BDA0003309859100000161
2. Steel structure
EBSD (Electron Back Scatter diffraction) measurements were carried out at the t/4 (t: sheet thickness) and t/2 positions of the cross section perpendicular to the rolling width direction of the hot rolled material. The measurement conditions are as follows.
EBSD measurement conditions
An apparatus: JEOL-5410 or JSM-IT 100 manufactured by Japan Electron
Observation magnification: 400 times of
Area of measurement: 200 μm × 200 μm
Step size (pixel): 0.4 μm
The phases considered: ferrite, austenite
The obtained EBSD data were analyzed by OIM Analysis software manufactured by TSLSolutions, inc. In the obtained data, the high angle grain boundary was defined as a grain boundary having a crystal orientation of 15 ° or more with respect to the adjacent pixel, except for the point where the Confidence Index was 0.100 or less. Among the cells surrounded by the large angle grain boundaries, the cells having a pixel size of 10 or more are regarded as large angle grains. In addition, the high angle grains at the end of the measurement field were excluded from the analysis. The average equivalent circle diameter of the high-angle crystal grains is obtained, and the total area fraction SA of the high-angle crystal grains having an average equivalent circle diameter of 7.5 μm or less in the total area of the high-angle crystal grains is calculated.
3. Mechanical properties
(yield strength YS)
In the t/4 position and the t/2 position of the hot-rolled steel sheet, an ASTM round bar tensile test piece was taken out in parallel to the sheet width direction (C direction), and a tensile test was carried out in accordance with the ASTM procedure to measure the yield strength YS.
(Low temperature toughness of parent Material)
V-notch Charpy impact test pieces were taken out in parallel to the sheet width direction (C direction) at the t/4 position and t/2 position of the hot-rolled steel sheet, and Charpy impact tests were carried out in accordance with the ASTM procedure. Then, the temperature vTrs at which the brittle fracture rate was 50% was evaluated.
The yield strength YP and ductile-brittle transition temperature vTrs obtained by the above measurement are substituted into the following formula (2), and Y values at the t/4 position and the t/2 position are obtained. The results are also shown in Table 2. Also in Table 2, values for Y are calculated as vTrs ═ 30 ℃ at the time of calculation at vTrs > -30 ℃. Further, when vTrs < -130 ℃, the Y value is calculated as vTrs ═ 130 ℃ at the time of calculation.
Y＝20×vTrs－7×YP…(2)
The following is clear from tables 1 and 2. In experiments nos. 1 to 11, since the thick steel plate was produced under the predetermined conditions while satisfying the composition specified in the embodiment of the present invention, the obtained thick steel plate had an excellent balance between strength and toughness in the steel plate even when the plate thickness was large. In particular, the steel sheet has high strength and excellent toughness at lower temperatures than conventional steel sheets. In contrast, in experiments No.12 to 17, at least any one of the component composition and the production conditions was out of the range defined in the embodiment of the present invention, and therefore the obtained thick steel plate was inferior in the balance between strength and toughness.
In experiment No.12, although the composition of the components was within the range of the embodiment of the present invention, the cumulative reduction ratio in the temperature range of 850 ℃ or lower was insufficient under the production conditions, and therefore the AF structure was insufficient at any position of t/4 and t/2, resulting in poor balance between strength and toughness.
In experiments No.13 and 14, although the composition of the composition was within the range of the embodiment of the present invention, in the production conditions, the AF structure was insufficient at any position of t/4 and t/2 because the finish rolling temperature was high, resulting in poor strength-toughness balance.
In experiment No.15, Di +10Nb is less than the specified range, and this resulted in poor balance between strength and toughness. In experiment No.15, it is considered that Di +10Nb is less than the predetermined range, and therefore, a pearlite structure is formed together with grain boundary ferrite. As a result, it was found that the AF structure was insufficient and the balance between strength and toughness was poor
In experiments No.16 and 17, Di +10Nb was higher than the specified range, and therefore, the balance between strength and toughness was poor. In experiment nos. 16 and 17, it is considered that the above-mentioned Di +10Nb is higher than the predetermined range, and a hard martensite structure is generated. As a result, it is considered that the hard martensite structure acts as a brittle fracture origin and the strength-toughness balance is poor.
In experiment No.18, although the composition of the components was within the range of the embodiment of the present invention, the heating temperature before hot rolling was low and the finish rolling temperature was high under the production conditions, so that sufficient AF structure could not be secured, resulting in poor strength-toughness balance.
In experiment No.19, although the composition of the components was within the range of the embodiment of the present invention, the heating temperature before hot rolling was high under the production conditions, and therefore, a fine AF structure could not be secured to some extent or more, resulting in poor balance between strength and toughness.
In experiment No.20, since the B content was 0.0008% and higher than 0.0007% of the upper limit, hard martensite was formed, and SA was insufficient at the t/4 position and the t/2 position, resulting in poor properties.
Although the composition of experiment No.21 was within the range of the embodiment of the present invention, the average reduction ratio in the final 3 passes of hot rolling was too low under the production conditions, and therefore, a fine AF structure could not be secured to some extent or more, resulting in poor balance between strength and toughness.
In experiment No.22, the Nb content was insufficient, and the cumulative reduction ratio of で in the temperature range of 850 ℃ or lower and the average reduction ratio of the final 3 passes of rolling were also low under the production conditions, so that the fine AF structure could not be secured to some extent or more, resulting in poor balance between strength and toughness.
Fig. 1 is a graph showing the relationship between the total area fraction SA and the value Y of crystal grains having an equivalent circle diameter of 7.5 μm or less among crystal grains surrounded by high angle grain boundaries having a grain orientation difference of 15 ° or more, which are produced based on the above examples. Also, the downward arrow in fig. 1 means that the Y value is presumed to be a value lower than the plotted value because the measured vTrs is lower than-130 ℃, and the upward arrow means that the Y value is presumed to be a value higher than the plotted value because the measured vTrs is higher than-30 ℃.
As can be seen from fig. 1, the total area fraction SA and the Y value are correlated at both the t/4 position and the t/2 position of the steel sheet, and it is necessary to make the total area fraction SA 34% or more at the t/4 position in order to make the Y value lower than-5200, and to make the Y value lower than-4700 at the t/2 position in order to make the total area fraction SA 27% or more.
The present application is accompanied by the claims of priority based on japanese patent application No. 2019-22/4/2019-081271, and the claims of priority based on japanese patent application No. 2020-007626/21/2020/1/2020. Japanese patent application Nos. 2019 and 081271 and 2020 and 007626 are incorporated herein by reference.

Claims (3)

1. A thick steel plate having a composition satisfying the following requirements
C: 0.020 to 0.070 mass percent,
Si: more than 0 mass% and not more than 0.40 mass%,
Mn: 1.30-1.95% by mass,
P: more than 0 mass% and not more than 0.015 mass%,
S: more than 0 mass% and not more than 0.005 mass%,
Al: 0.005-0.070 wt%,
Nb: 0.015-0.048 wt%,
Ti: 0.005-0.024 mass%,
N: 0.0030 to 0.0080 mass%, and
ca: higher than 0 mass% and not more than 0.0040 mass%,
the balance being Fe and unavoidable impurities,
and satisfies Di +10Nb as determined by the following formula (1): 1.20 to 2.50 of the total weight of the alloy,
the total area fraction SA of crystal grains having an equivalent circle diameter of 7.5 μm or less among crystal grains surrounded by high-angle grain boundaries having a grain misorientation of 15 DEG or more is 34% or more at the 1/4 position of the plate thickness and 27% or more at the 1/2 position of the plate thickness,
Di＝1.16×([C]/10)0.5×(0.7×[Si]+1)×(5.1×([Mn]－1.2)+5)×(0.35×[Cu]+1)×(0.36×[Ni]+1)×(2.16×[Cr]+1)×(3×[Mo]+1)×(1.75×[V]+1)×(200×[B]+1)…(1)
in the formula (1), [ C ], [ Si ], [ Mn ], [ Cu ], [ Ni ], [ Cr ], [ Mo ], [ V ] and [ B ] respectively represent the contents of C, Si, Mn, Cu, Ni, Cr, Mo, V and B expressed in mass%, and the element not included is 0.
2. The high-strength steel sheet according to claim 1, wherein any 1 or more of the following (i) to (iii) are satisfied:
(i) further comprising a Cu: more than 0 mass% and less than 0.75 mass%, and Ni: more than 0 mass% and 1.4 mass% or less of one or more elements selected from the group consisting of;
(ii) further contains a compound selected from Mo: more than 0 mass% and less than 0.50 mass%, V: more than 0 mass% and not more than 0.060 mass% of Cr: more than 0 mass% and less than 0.8 mass%, and B: more than 0 mass% and 0.0007 mass% or less of one or more elements selected from the group consisting of;
(iii) also contains a secondary amino acid sequence selected from REM: more than 0 mass% and less than 0.0060 mass%, and Zr: more than 0 mass% and not more than 0.0050 mass%.
3. A method for manufacturing a thick steel plate according to claim 1 or 2, wherein,
comprising a step of heating a steel sheet having the composition of claim 1 or 2 to a temperature of more than 1020 ℃ and less than 1200 ℃ and a hot rolling step after the heating,
in the hot rolling step, the number of passes is set to 3 or more, and the hot rolling and the cooling after the hot rolling are performed so as to satisfy all of the following conditions (a) to (d),
(a) a cumulative reduction ratio at a temperature of 850 ℃ or lower of 40% or more;
(b) the average reduction rate of the final rolling of 3 passes is more than 5.5 percent;
(c) the finishing temperature is 720-830 ℃;
(d) after hot rolling, the steel sheet is cooled from a cooling start temperature of a finish rolling temperature of 690 ℃ to a cooling stop temperature of 320 to 550 ℃ at an average cooling rate of 0.5 to 20 ℃/s.
CN202080029745.4A 2019-04-22 2020-03-30 Thick steel plate and method for producing same Active CN113840933B (en)

Applications Claiming Priority (5)

Application Number Priority Date Filing Date Title
JP2019081271 2019-04-22
JP2019-081271 2019-04-22
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