CN117337341A

CN117337341A - High-strength steel sheet and method for producing same

Info

Publication number: CN117337341A
Application number: CN202280030541.1A
Authority: CN
Inventors: 荒尾亮; 竹下龙平; 橘俊一
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2021-07-02
Filing date: 2022-05-30
Publication date: 2024-01-02
Also published as: KR20230159568A; JP7173423B1; JPWO2023276516A1

Abstract

The invention provides a high-strength steel sheet and a method for producing the same. The high-strength steel material of the present invention has a specific composition, and the metal structure of the high-strength steel sheet at the position of 1/4 of the sheet thickness is composed of a soft phase as a main phase and a hard phase as a balance, the soft phase being composed of ferrite, the hard phase containing 1 or more of pearlite, bainite, and martensite, the fraction of the processed ferrite in the whole metal structure being 50% or more in terms of volume fraction, the aspect ratio of the processed ferrite being 1.5 or more, the average grain size of the processed ferrite being 50 [ mu ] m or less, the yield stress of the base material being 235MPa or more, the Charpy impact absorption energy at-60 ℃ of the base material being 200J or more, and the Charpy impact absorption energy at-60 ℃ of the weld heat affected zone after large-wire energy welding being 80J or more.

Description

High-strength steel sheet and method for producing same

Technical Field

The present invention relates to a steel sheet (thick steel sheet) used for ships, marine structures, middle-to-high buildings, bridges, tanks, and the like, and more particularly, to a high-strength steel sheet having high toughness even in a weld heat affected zone after welding the steel sheet, and a method for manufacturing the same.

Background

In recent years, the severity desired for the material properties of welding steels (welding steel sheets) used for structures such as ships, marine structures, middle-to-high buildings, bridges, and tanks has been increasing. In addition, in order to manufacture such a structure in a short period of time, it is desirable to use a large linear energy welding method represented by submerged arc welding, electric welding, electroslag welding, or the like. Therefore, as with the toughness of the steel itself, the severity required for the toughness of the weld heat affected zone (hereinafter also referred to as "HAZ") is also increasing.

However, in general, when the weld line energy increases, the structure of the HAZ coarsens, and the toughness of the HAZ decreases. For such a decrease in toughness of the HAZ due to the high heat input welding, for example, as in the techniques described in patent documents 1 to 5, a large number of countermeasures have been proposed so far.

Patent documents 1 and 2 disclose techniques for improving toughness of a HAZ (hereinafter, also referred to as "high heat input HAZ") caused by high heat input welding. Specifically, a method of suppressing coarsening of austenite grains by utilizing the pinning effect of TiN, al oxide, or the like has been proposed.

Patent documents 3, 4 and 5 disclose techniques for achieving refinement of the intra-grain structure by allowing a large number of ferrite transformation nuclei to exist in austenite grains. Specifically, refinement of the intra-grain structure is achieved by using TiN, mnS, ti oxide or the like as ferrite transformation nuclei, thereby achieving improvement of low-temperature toughness of the HAZ.

Prior art literature

Patent literature

Patent document 1: japanese patent laid-open No. 2002-256379

Patent document 2: japanese patent No. 2950076

Patent document 3: japanese patent publication No. 07-068577

Patent document 4: japanese patent publication No. 05-017300

Patent document 5: japanese patent No. 3733898

Disclosure of Invention

Problems to be solved by the invention

However, even if the techniques for refining the HAZ using the above-described precipitates disclosed in patent documents 1 to 5 are applied, coarsening of the HAZ structure is unavoidable when large-heat welding is performed, and deterioration of low-temperature toughness of the HAZ occurs, for example, in an environment below-60 ℃.

In recent years, in the field of ships, tanks, and the like, the use in an environment at a lower temperature than before has been studied. Therefore, there is a need for a steel product having significantly improved low-temperature toughness in the weld heat affected zone as compared with the steel product (steel sheet) to be subjected to the technique described in each of the above patent documents.

The present invention has been made in view of the above-described circumstances, and particularly aims to provide a high-strength steel sheet excellent in low-temperature toughness of a base material and a HAZ (high-heat-input HAZ) of a steel sheet used for the above-described applications, and a method for producing the same.

Here, "high strength" in the present invention means that the yield stress (YP) indicating the base material strength of the high-strength steel sheet is 235MPa or more. The term "excellent low-temperature toughness of the base material" in the present invention means that the absorption energy of the base material in the Charpy impact test at-60℃is 200J or more. In the present invention, "excellent low-temperature toughness of the HAZ (high heat input HAZ)" means that the absorption energy in the Charpy impact test at-60 ℃ of the HAZ after high heat input welding, that is, the HAZ of the single-sided 1-pass welded joint, is 80J or more. In particular, stable low-temperature toughness can be obtained for a welded joint having a weld line energy of 4kJ/mm or more, which is brought about by single-sided 1-pass welding such as submerged arc welding.

The yield stress and the absorption energy of the Charpy impact test can be measured by the methods described in examples described later.

Means for solving the problems

The present inventors have conducted intensive studies on a method for improving low-temperature toughness of a high heat input HAZ in order to solve the above-mentioned problems, and as a result, have obtained the following findings.

First, the present inventors focused on a coarse bainitic structure generated by high heat input welding. Coarse bainite has a coarse structure compared with a structure such as ferrite or pearlite. The brittle fracture of the coarse structure has a low critical stress, and therefore becomes a factor of decrease in toughness. Therefore, the present inventors considered to improve the low-temperature toughness of the high heat input HAZ by suppressing the generation of coarse bainite.

Further, as a result of intensive studies, the present inventors have found the following findings. Ferrite transformation is promoted by the composition designed to satisfy the condition of the following formula (1), and coarse bainite grain refinement can be achieved. Further, by designing the ratio of Ti to N (Ti/N) to be in the range of 1.5 to 4.0, coarsening of the prior austenite grain diameter due to welding can be suppressed, and suppression of coarse bainite can be achieved.

However, when the condition of formula (1) is satisfied, the fraction of ferrite as the soft phase increases, and thus it becomes difficult to secure the strength of the base material (steel sheet).

Accordingly, the present inventors studied the metal structure of the base material. As a result, it was found that: at the position 1/4 of the plate thickness of the steel sheet, the fraction of the processed ferrite in the entire structure of the base material is 50% or more by volume fraction, the aspect ratio of the processed ferrite is 1.5 or more, and the average crystal grain size of the processed ferrite is 50 μm or less, whereby excellent base material strength can be obtained when the condition of formula (1) is satisfied. C (Mn)/6 (Cu (Ni)) 15 (Cr (Mo) V) 5 (0.35) … (1) with the concentration of C (Mn) and 6 (V) being 0.25 to 0

Wherein C, mn, cu, ni, cr, mo, V is the content (mass%) of each element, and the content of the element not contained is set to zero.

The present invention has been completed based on the above findings, and the gist thereof is as follows.

[1] A high-strength steel sheet comprising, in mass%, C: 0.010-0.070%, si:0.01 to 0.50 percent of Mn: 1.00-2.00%, P: less than 0.020%, S:0.0005 to 0.0100 percent of Al:0.035 to 0.100 percent of Ti:0.010 to 0.030% and N:0.0035 to 0.0100%, a carbon equivalent Ceq (IIW) represented by the following formula (1) satisfying 0.25 to 0.35% by mass, a Ti/N satisfying 1.5 to 4.0, and the balance being Fe and unavoidable impurities,

the metal structure of the high-strength steel sheet at the 1/4 position of the sheet thickness is composed of a soft phase as a main phase and a hard phase as a balance, the soft phase is composed of ferrite, the hard phase contains 1 or more of pearlite, bainite and martensite,

the fraction of the processed ferrite in the whole metallic structure is 50% or more in terms of volume fraction,

the aspect ratio of the processed ferrite is 1.5 or more,

the processed ferrite has an average crystal grain size of 50 μm or less,

the base material has a yield stress of 235MPa or more and a Charpy impact absorption energy of 200J or more at-60 ℃,

The Charpy impact absorption energy of the welding heat affected zone after large heat input welding at minus 60 ℃ is more than 80J,

Ceq.(IIW)＝[C]+[Mn]/6+([Cu]+[Ni])/15+([Cr]+[Mo]+[V])/5…(1)

here, the [ ] in the formula (1) is the content (mass%) of each element in the brackets, and the content of the element not contained is set to zero.

[2] The high-strength steel sheet according to the above [1], wherein the composition further comprises, in mass%, a composition selected from the group consisting of: less than 0.0030%, cu: less than 0.50%, ni: less than 1.50%, V: less than 0.100%, cr: less than 0.50%, mo: less than 0.50%, ca: less than 0.0030 percent of Mg:0.0050% or less and REM:0.1000% or less of 1 or 2 or more of the group consisting of.

[3] The high-strength steel sheet according to the item [1] or [2], wherein the size of TiN in the metal structure of the high-strength steel sheet at the 1/4 position of the sheet thickness is 5 to 200nm.

[4] A method for producing a high-strength steel sheet according to any one of the above [1] to [3], wherein,

the steel raw material with the above composition is heated to a temperature of 1050 ℃ or more and 1200 ℃ or less,

then, the following hot rolling was performed: rolling is started in a temperature range of 950 ℃ or higher, i.e., in a recrystallized gamma region, is started under conditions in which the rolling reduction in a temperature range of 850 ℃ or lower, i.e., in a non-recrystallized gamma region, is 30% or higher and the rolling reduction in a ferrite-austenite two-phase region having a (Ar 3 transformation point-80) DEG C to Ar3 transformation point is 30% or higher, the finishing rolling temperature is 650 ℃ or higher,

Then, the cooling is performed at an average cooling rate of 5 ℃ per second or more from a cooling start temperature of 650 ℃ or more to a cooling stop temperature in a temperature range of 600 ℃ or less and 300 ℃ or more.

[5] The method for producing a high-strength steel sheet according to [4], wherein the steel material is cast at an average casting speed of 0.3 to 1.0 m/min.

Effects of the invention

According to the present invention, even when the high-strength steel sheet of the present invention is subjected to high heat input welding, the high-strength steel sheet can have excellent low-temperature toughness in the base metal and the weld heat affected zone. Therefore, the high-strength steel sheet of the present invention can be suitably used as a steel sheet (steel material) for welding a structure such as a liquefied gas low-temperature storage tank or a ship operated in a low-temperature environment, which is constructed by a high-heat welding method such as electric gas welding, submerged arc welding or electroslag welding.

Detailed Description

The present invention will be described below. The present invention is not limited to the following embodiments.

First, the reason why the composition of the components of the high-strength steel sheet is limited in the present invention will be described. In the present invention, unless otherwise specified, "%" indicating the composition of the components means "% by mass".

C：0.010～0.070％

In the present invention, it is necessary to contain 0.010% or more of C in order to obtain the strength of the target base material (high-strength steel sheet). However, if C is contained in an amount exceeding 0.070%, island-like martensite increases and the low-temperature toughness of a weld Heat Affected Zone (HAZ) decreases, and therefore, the C content is set to 0.070% or less. The C content is preferably 0.020% or more, more preferably 0.030% or more, and still more preferably 0.050% or more. The C content is preferably 0.065% or less, more preferably 0.060% or less, and still more preferably 0.055% or less.

Si：0.01～0.50％

Si is a component necessary for ensuring the strength of the base material, deoxidizing, and the like, and in the present invention, si is contained in an amount of 0.01% or more. On the other hand, if the Si content exceeds 0.50%, the HAZ is hardened and the low-temperature toughness of the HAZ is reduced, and therefore the Si content is set to 0.50% or less. The Si content is preferably 0.1% or more, more preferably 0.15% or more. The Si content is preferably 0.40% or less, more preferably 0.3% or less.

Mn：1.00～2.00％

Mn is required to be contained in an amount of 1.00% or more in order to secure the strength of the base material. On the other hand, when Mn is contained in an amount exceeding 2.00%, not only weldability is deteriorated but also the cost of the steel sheet increases. Therefore, the Mn content is set to 1.00 to 2.00%. The Mn content is preferably 1.20% or more, more preferably 1.40% or more, and still more preferably 1.50% or more. The Mn content is preferably 1.90% or less, more preferably 1.75% or less, and still more preferably 1.60% or less.

P: less than 0.020%

P is an unavoidable impurity, and if the content of P exceeds 0.020%, the low-temperature toughness of the base material and the welded portion is lowered, and therefore the upper limit is set to 0.020%. Therefore, the P content is set to 0.020% or less. The P content is preferably 0.015% or less. In order to obtain good low-temperature toughness, the content of P is more preferably 0.010% or less, and still more preferably 0.007% or less. The lower limit of the P content is not particularly limited, but the P content is preferably 0.001% or more because the extremely low P treatment causes an increase in cost.

S：0.0005～0.0100％

In order to produce CaS or MnS, which is a core of a composite inclusion required for ferrite core production, S is required to be contained in an amount of 0.0005% or more. On the other hand, if the S content exceeds 0.0100%, the low-temperature toughness of the base material is deteriorated. Therefore, the S content is set to 0.0100% or less. The S content is preferably 0.0090% or less, more preferably 0.0030% or less. The S content is preferably 0.0010% or more, more preferably 0.0015% or more.

Al：0.035～0.100％

Al is required to be contained in an amount of 0.035% or more in view of deoxidation of steel. On the other hand, if Al is contained in an amount exceeding 0.100%, the low-temperature toughness of the base material is lowered, and the low-temperature toughness of the weld metal is deteriorated. Therefore, the Al content is set to 0.100% or less. The Al content is preferably 0.095% or less, more preferably 0.090% or less, and still more preferably 0.080% or less. The Al content is preferably 0.040% or more, more preferably 0.050% or more.

Ti：0.010～0.030％

Ti becomes TiN and precipitates during solidification of the steel, and contributes to suppression of coarsening of austenite in the HAZ and formation of ferrite transformation nuclei, thereby contributing to high toughness. If the Ti content is less than 0.010%, the effect is small, and if it exceeds 0.030%, the desired effect cannot be obtained due to coarsening of TiN particles. Therefore, the Ti content is set to be in the range of 0.010 to 0.030%. The Ti content is preferably 0.011% or more, more preferably 0.013% or more, and still more preferably 0.015% or more. The Ti content is preferably 0.028% or less, more preferably 0.025% or less, and still more preferably 0.020% or less.

N：0.0035～0.0100％

N is contained in an amount of 0.0035% or more in order to form TiN for bonding with Ti. When the content of N increases, the solid solution N increases, which results in a decrease in low-temperature toughness of the HAZ, and therefore the upper limit of the N content is set to 0.0100%. Therefore, the N content is set to 0.0100% or less. The N content is preferably 0.0040% or more, more preferably 0.0045% or more, and even more preferably 0.0052% or more. The N content is preferably 0.0095% or less, more preferably 0.0090% or less, and even more preferably 0.0075% or less.

Carbon equivalent ceq. (IIW): 0.25 to 0.35 mass%

The composition of the high-strength steel sheet of the present invention is adjusted so that the carbon equivalent ceq (IIW) represented by the following formula (1) satisfies the range of 0.25 to 0.35 mass%. Ceq. (IIW) = [ C ] + [ Mn ]/6+ ([ Cu ] + [ Ni ])/[ 15+ ([ Cr ] + [ Mo ] + [ V ])/[ 5 … (1)

Regarding the carbon equivalent Ceq. (IIW), 0.35 mass% or less is set in order to ensure that vTrs (fracture transition temperature) is-60 ℃ or less in the weld heat affected zone. On the other hand, if the carbon equivalent ceq. (IIW) is less than 0.25 mass%, the yield stress in the base material cannot be ensured to be 235MPa or more. Accordingly, the carbon equivalent ceq. (IIW) is set to 0.25 to 0.35 mass%. The carbon equivalent ceq. (IIW) is preferably 0.27 mass% or more, more preferably 0.28 mass% or more. The carbon equivalent ceq. (IIW) is preferably 0.33 mass% or less, more preferably 0.32 mass% or less.

The reason why the vTrs is set to-60 ℃ or lower in the weld heat affected zone is that it is estimated that the design temperature of the liquefied gas storage tank, which is expected to be increased in the future, is-60 ℃ or lower.

Ti/N：1.5～4.0

In the present invention, ti is added and the Ti content is adjusted so as to satisfy 1.5.ltoreq.Ti/N.ltoreq.4.0 in consideration of the relation with the N content (here, ti and N are defined as the contents (mass%) of the respective elements). By adjusting the Ti content, an optimal size and amount of TiN can be ensured, and as a result, coarsening of austenite can be suppressed. When Ti/N is less than 1.5, tiN is refined, and thus TiN is solid-dissolved in the weld heat affected zone. As a result, the amount of TiN required for improving the low-temperature toughness of the welded portion cannot be ensured. On the other hand, if the Ti/N ratio exceeds 4.0, the low-temperature toughness of the base material and the low-temperature toughness of the welded portion are lowered due to the formation of TiC particles and the coarsening of TiN. Therefore, the ratio (T/Ni) of the Ti content to the N content is set to 1.5 or more and 4.0 or less. T/Ni is preferably 2.0 or more, more preferably 2.5 or more. T/Ni is preferably 3.4 or less, more preferably 3.2 or less.

In the present invention, the size of TiN is preferably adjusted to 5nm or more and 200nm or less. This is because, when the size is outside this range, a sufficient effect of suppressing coarsening of austenite grains cannot be obtained. In the present invention, the "size of TiN" means the diagonal length of TiN, which is a rectangle, and can be measured by the method described in examples described below.

In the high-strength steel sheet of the present invention, the balance other than the above-mentioned components is iron (Fe) and unavoidable impurities.

In the present invention, the above elements are set to a basic component composition, and the target characteristics of the present invention can be obtained by the basic component composition. In the present invention, in order to further improve the characteristics, the following elements may be contained in addition to the above basic component composition, if necessary. Note that, since each component B, cu, ni, V, cr, mo, ca, mg and REM described below may be contained as necessary, these components may be 0%.

Selected from the group consisting of: less than 0.0030%, cu: less than 0.50%, ni: less than 1.50%, V: less than 0.100%, cr: less than 0.50%, mo: less than 0.50%, ca: less than 0.0030 percent of Mg:0.0050% or less and REM:0.1000% or less of 1 or 2 or more of the group consisting of.

B: less than 0.0030 percent

B is an element that effectively contributes to the enhancement of strength of the steel sheet (base material). Such an effect becomes remarkable when the content of B is 0.0002% or more. On the other hand, if B is excessively contained, the low-temperature toughness of the HAZ of the welded portion is adversely affected, and therefore, the B content is preferably 0.0030% or less. Therefore, when B is contained, the B content is preferably 0.0030% or less. The B content is preferably 0.0002% or more, more preferably 0.0007% or more. The B content is more preferably 0.0012% or less.

Cu: less than 0.50%

Cu is an element for improving hardenability of steel, and contributes to improvement of not only strength of a base material after rolling but also functions such as high-temperature strength and weather resistance. These effects are exhibited by containing 0.01% or more of Cu. On the other hand, excessive Cu content adversely deteriorates low-temperature toughness and weldability of the HAZ of the welded portion. The Cu content is preferably 0.50% or less. Therefore, when Cu is contained, the Cu content is preferably 0.50% or less. The Cu content is preferably 0.01% or more, more preferably 0.04% or more. The Cu content is more preferably 0.10% or less.

Ni: less than 1.50 percent

Ni is an element for improving the hardenability of steel, and contributes to the improvement of the strength of the base material after rolling, as well as the improvement of the functions such as low-temperature toughness, high-temperature strength, weather resistance, and the like of the base material. These effects are exhibited by the inclusion of 0.01% or more of Ni. On the other hand, excessive Ni content adversely deteriorates low-temperature toughness and weldability of the HAZ of the welded portion, and increases the cost of the alloy. The Ni content is preferably 1.50% or less. Therefore, when Ni is contained, the Ni content is preferably 1.50% or less. The Ni content is preferably 0.01% or more, more preferably 0.02% or more. The Ni content is more preferably 0.50% or less.

V: less than 0.100%

V is an element effective for improving the strength and low-temperature toughness of the base material, and is an element that acts as a ferrite generation nucleus of the VN form. Such an effect is exhibited by containing 0.005% or more of V. On the other hand, when V is contained in an amount exceeding 0.100%, the low-temperature toughness of the HAZ of the welded portion is lowered. The V content is preferably 0.100% or less. Therefore, when V is contained, the V content is preferably 0.100% or less. The V content is preferably 0.005% or more, more preferably 0.009% or more. The V content is more preferably 0.080% or less.

Cr: less than 0.50%

Like Cu, cr is an element for improving hardenability of steel, and contributes to improvement of strength of a base material after rolling, as well as improvement of functions such as high-temperature strength and weather resistance. These effects are exhibited by the inclusion of 0.01% or more of Cr. On the other hand, excessive Cr content adversely deteriorates the low-temperature toughness and weldability of the HAZ of the welded portion. The Cr content is preferably 0.50% or less. Therefore, when Cr is contained, the Cr content is preferably 0.50% or less. The Cr content is preferably 0.01% or more, more preferably 0.02% or more. The Cr content is more preferably 0.10% or less.

Mo: less than 0.50%

Like Cu and Cr, mo is an element for improving hardenability of steel, and contributes to improvement of strength of a base material after rolling, high-temperature strength, weather resistance, and other functions. These effects are exhibited by the Mo content of 0.01% or more. On the other hand, excessive Mo content adversely deteriorates low-temperature toughness and weldability of the HAZ of the welded portion. The Mo content is preferably 0.50% or less. Therefore, when Mo is contained, the Mo content is preferably 0.50% or less. The Mo content is preferably 0.01% or more, more preferably 0.02% or more. The Mo content is more preferably 0.10% or less.

Ca: less than 0.0030 percent

Ca is an element useful for improving the low-temperature toughness of the base material and the HAZ due to the fixation of S. If the content of Ca exceeds 0.0030%, the effect is saturated, and therefore, ca is set to be 0.0030% or less. On the other hand, if the content of Ca is less than 0.0005%, the fixation of S becomes insufficient. Ca is contained in an amount of 0.0005% or more. Therefore, when Ca is contained, the Ca content is preferably 0.0030% or less. The Ca content is preferably 0.0005% or more, more preferably 0.0010% or more. The Ca content is more preferably 0.0025% or less.

Mg: less than 0.0050%, REM: less than 0.1000%

Mg and REM (rare earth metal) both have strong deoxidizing power in molten steel and have an effect of assisting formation of fine oxides, and thus are added as needed. Regarding the content each showing the deoxidizing effect, mg was 0.0002% or more, REM was 0.0010% or more. On the other hand, if the amount is large, coarse inclusions are generated and the properties of the base material are impaired, so that the respective contents are preferably set to 0.0050% or less of Mg and 0.1000% or less of REM. Therefore, when Mg and REM are contained, mg is preferably 0.0050% or less and REM is preferably 0.1000% or less. The Mg content is preferably 0.0002% or more. The REM content is preferably 0.0010% or more.

The high-strength steel sheet of the present invention can achieve an improvement in low-temperature toughness by satisfying the above-described composition. On the other hand, as described above, it is difficult to secure the strength of the base material. Therefore, in order to secure the strength, it is also important to define the metallic structure of the high-strength steel sheet as follows in the present invention.

Hereinafter, the metal structure of the high-strength steel sheet of the present invention will be described.

The high-strength steel sheet of the present invention has a metallic structure comprising a soft phase as a main phase and a hard phase as a balance at 1/4 of the sheet thickness, wherein the soft phase is composed of ferrite, and the hard phase contains at least 1 or 2 of pearlite, bainite, and martensite. In addition, regarding the above-mentioned processed ferrite, the fraction of the processed ferrite of the high-strength steel sheet in the entire metal structure at the plate thickness 1/4 position is 50% or more in terms of volume fraction, the aspect ratio of the processed ferrite is 1.5 or more, and the average crystal grain size of the processed ferrite is 50 μm or less.

Main phase: ferrite body

From the viewpoint of improving the strength of the base material, the high-strength steel sheet of the present invention has ferrite as a main phase at the 1/4 position of the sheet thickness. In the present invention, the "main phase" means 50% or more by volume. The volume ratio of ferrite is preferably 70% or more, more preferably 75% or more, still more preferably 80% or more, and still more preferably 90% or more.

In the present invention, as described above, in order to obtain excellent base material strength when the condition of formula (1) is satisfied, the fraction of processed ferrite at the 1/4 position of the plate thickness of the steel sheet is also specified in addition to ferrite as the soft phase. Details of the ferrite processing will be described later.

The balance: hard phase containing 1 or 2 or more of pearlite, bainite and martensite

The remaining structure other than ferrite is a hard phase containing 1 or 2 or more of pearlite, bainite, and martensite at the 1/4 position of the plate thickness from the viewpoint of securing strength. The remaining tissues are preferably set to 25% or less in total of the volume fractions of the respective tissues. The total volume ratio of the remaining tissues is more preferably 15% or less, and still more preferably 10% or less.

Volume fraction of processed ferrite: more than 50 percent

In the present invention, when manufacturing a high strength steel sheet, according to hot rolling conditions described below, dislocation is added to ferrite having a temperature range in a dual phase region by dual phase region rolling, thereby improving strength. In order to obtain such an effect, it is necessary to have processed ferrite into which dislocations are introduced by two-phase region rolling in a ratio of a certain or more. In order to secure the strength of the base material, the processed ferrite needs to be 50% or more in terms of volume fraction of the entire metal structure of the high-strength steel sheet at the position of 1/4 of the plate thickness. The processed ferrite is preferably 60% or more in terms of volume fraction.

The processed ferrite occupies 55% or more of the ferrite. Preferably 70% or more.

The upper limit of the processed ferrite is not particularly limited, but is preferably set to 90% or less in terms of volume fraction from the viewpoint of the load of the rolling mill and the reason of preventing the reduction in absorption energy due to the occurrence of peeling. The processed ferrite is more preferably 80% or less in terms of volume fraction.

The processed ferrite occupies 96% or less of the ferrite. Preferably 93% or less.

Herein, "processed ferrite" in the present invention means that the value of the dislocation density ρ obtained by X-ray diffraction (XRD) is 1.0x10 ¹⁴ m ^-2 The above ferrite. By making the value of the dislocation density ρ 1.0X10 ¹⁴ m ^-2 As described above, work hardening works, and strength increases. The value of the dislocation density ρ is preferably set to 2.0X10 ¹⁴ m ^-2 . However, if the two-phase region pressure is excessively applied, dislocations are excessively introduced, and the dislocations are difficult to move, with the result that the low-temperature toughness of the steel sheet is lowered. Therefore, the upper limit of the value of the dislocation density ρ is set to 2.5X10 ¹⁵ m ^-2 The following is given. In the present invention, the dislocation density can be measured by the method described in examples described below.

Aspect ratio of processed ferrite: 1.5 or more

If the aspect ratio of the processed ferrite at the plate thickness 1/4 position is less than 1.5, the specific texture may not be sufficiently developed, and ductile cracks may occur. In addition, the dislocation strengthening and grain refining effects are not exhibited, and the toughness of the base material at low temperatures is reduced. Therefore, the aspect ratio of the processed ferrite is set to 1.5 or more. The aspect ratio of the processed ferrite is preferably 2.0 or more, more preferably 2.5 or more. The upper limit value of the aspect ratio of the processed ferrite is not particularly limited. From the viewpoint of the performance of the rolling mill, the aspect ratio of the processed ferrite is preferably 4.0 or less, more preferably 3.8 or less.

Average grain size of processed ferrite: 50 μm or less

When the average grain size of the processed ferrite at the position 1/4 of the plate thickness is 50 μm or less, grain boundaries per unit volume increase, and as a result, dislocation is less likely to move, and strength of the steel sheet is improved. The average grain size of the processed ferrite is preferably 45 μm or less, more preferably 40 μm or less, and even more preferably 30 μm or less. The upper limit of the average crystal grain size of the processed ferrite is not particularly limited. From the viewpoint of applying the steel sheet to the above-mentioned applications, it is preferably 5 μm or more, more preferably 15 μm or more.

Here, regarding the grain size of the processed ferrite of the present invention, the difference in orientation of adjacent grains is obtained, and the boundary where the difference in orientation is 15 ° or more is defined as the grain boundary and measured. The average crystal grain size was obtained by obtaining an arithmetic average of grain sizes from the obtained crystal grain boundaries, and obtaining an average circle equivalent diameter.

The volume fraction of ferrite, pearlite, bainite, martensite, processed ferrite, the aspect ratio of processed ferrite, and the average crystal grain size of processed ferrite can be measured by the methods described in examples described later.

Next, an embodiment of the method for producing a high-strength steel sheet according to the present invention will be described.

First, molten steel having the above-described composition is melted by a melting method such as a converter or an electric furnace. In addition, refining may be performed 2 times using a vacuum degassing furnace. Thereafter, a steel material such as a billet having a predetermined size is preferably obtained by a casting method such as a continuous casting method or an ingot-cogging method.

In the present invention, a billet may be used as the steel stock. In the case of producing the steel material by the continuous casting method, the casting conditions thereof preferably satisfy the following conditions.

Specifically, the average casting speed at the time of billet casting is preferably set to 0.3 m/min or more and 1.0 m/min or less. The cooling of the steel billet (steel stock) can be controlled by the casting speed. If the average casting speed is less than 0.3 m/min, the size of TiN of the base material (high-strength steel sheet) becomes large. When the size of TiN increases, the TiN density of the base material (high strength steel sheet) decreases, and thus the pinning effect decreases. As a result, ferrite refinement cannot be sufficiently achieved in the HAZ of the welded portion, and the low-temperature toughness of the HAZ may be deteriorated. As described above, the size of TiN serving as a core is preferably 5nm to 200 nm. The casting speed is an average speed of the whole casting process.

On the other hand, if the average casting speed exceeds 1.0 m/min, the TiN density increases, but the size of TiN becomes smaller than the above range. Thus, tiN is solid-dissolved by large line energy at the time of welding of the base material (high-strength steel sheet). As a result, the austenite grain size coarsens, and the low-temperature toughness of the HAZ may be degraded.

The reasons for limiting the production conditions of the high-strength steel sheet excellent in low-temperature toughness and low-temperature toughness of the high-heat input HAZ for producing the steel material as a base material will be described in detail.

In the present invention, the steel material is heated to a temperature of 1050 ℃ or more and 1200 ℃ or less, and then hot rolled as follows: the rolling is started in a recrystallized gamma region at 950 ℃ or higher, and is performed under conditions such that the rolling reduction in an unrecrystallized gamma region at 850 ℃ or lower is 30% or higher and the rolling reduction in a ferrite-austenite two-phase region from (Ar 3 transformation point-80) DEG C to Ar3 transformation point is 30% or higher, and the finishing rolling temperature is 650 ℃ or higher. After the hot rolling, the hot rolled steel sheet is cooled at an average cooling rate of 5 ℃ per second or more from a cooling start temperature of 650 ℃ or more to a cooling stop temperature in a temperature range of 600 ℃ or less and 300 ℃ or more.

In the following description of the production method, unless otherwise specified, the reference symbol "°c" concerning the temperature refers to the surface temperature of each steel material or steel sheet. The surface temperature may be measured by, for example, a radiation thermometer or the like. The temperature at the center of the thickness of the steel slab or the steel sheet may be obtained by, for example, installing a thermocouple at the center of the thickness of the steel sheet, measuring the temperature, calculating the temperature distribution in the cross section of the steel sheet by thermal conduction analysis, and correcting the result by the surface temperature of the steel sheet.

[ heating temperature of Steel Material ]

The heating temperature of the steel material (for example, billet) needs to be 1050 ℃ or higher and 1200 ℃ or lower. This is because, when the heating temperature is less than 1050 ℃, coarse inclusions generated during solidification of the billet and adversely affecting the low-temperature toughness may not be melted and remain. On the other hand, if heating is performed at a high temperature, there is a possibility that precipitates produced by cooling at a controlled rate during casting may be remelted. Based on this point, a heating temperature of 1200 ℃ or less is sufficient as a heating temperature in the meaning of ending the phase transition. It is considered that coarsening of crystal grains occurring during the heat retention can be prevented in advance by the above-described pinning effect of TiN. For the above reasons, the heating temperature is limited to 1050 ℃ or higher and 1200 ℃ or lower. The heating temperature is preferably 1180 ℃ or lower, more preferably 1100 ℃ or lower.

[ Hot Rolling conditions ]

Rolling start temperature: 950 ℃ or higher

The heated steel material was rolled in a recrystallization temperature range (recrystallization γ zone) which is a temperature range of 950 ℃ or higher. In this temperature range, austenite grains are recrystallized by rolling. As a result, the structure can be made finer. If the rolling is started from a temperature lower than 950 ℃, recrystallization of austenite grains cannot be sufficiently performed, and thus the microstructure is insufficiently miniaturized. As a result, the base material strength is reduced. The rolling start temperature is preferably 970 ℃ or higher, more preferably 1000 ℃ or higher. The upper limit of the rolling start temperature is not particularly limited, but is preferably 1100 ℃ or lower from the viewpoint of suppressing the rolling roll loss due to contact with the high-temperature steel material.

The reduction rate of unrecrystallized gamma region below 850℃: more than 30 percent

Hot rolling is performed at a reduction ratio of 30% or more in a temperature range of 850 ℃ or less, that is, in a non-recrystallized temperature range (non-recrystallized gamma region). The reason for this is as follows. In this temperature range, recrystallization of austenite grains does not occur, the austenite grains deform flatly, and defects such as deformation bands are introduced into the austenite grains. This accumulated internal energy is applied to the driving force of the ferrite transformation thereafter. If the reduction ratio of the unrecrystallized gamma region is less than 30%, the internal energy accumulated is insufficient, and therefore ferrite is not sufficiently refined, and the strength of the base material cannot be ensured. The reduction ratio of the unrecrystallized γ region is preferably 35% or more, and more preferably 40% or more. The upper limit of the reduction ratio of the unrecrystallized gamma region is not particularly specified. If the unrecrystallized gamma region reduction ratio is excessively increased, the production efficiency is reduced, and therefore, the reduction ratio is preferably 55% or less, more preferably 50% or less.

(Ar 3 transformation point-80). Degree.C to Ar3 transformation point ferrite-austenite two-phase region reduction ratio: more than 30 percent

It is necessary to perform hot rolling with a reduction ratio of 30% or more in a ferrite-austenite two-phase temperature range, which is a temperature range from (Ar 3 transformation point-80) DEG C to Ar3 transformation point. The reason for this is as follows. The increase in the amount of depression in the above two-phase temperature range has the following effect: the improvement in strength associated with dislocation strengthening by ferrite processing during rolling and the improvement in low-temperature toughness due to the effect of grain refining by the formation of secondary grains by processing. Further, by setting the rolling reduction in the ferrite-austenite two-phase temperature range to 30%, the rolling texture of ferrite is developed, which contributes to improvement of low-temperature toughness. For this reason, the rolling reduction in the ferrite-austenite two-phase temperature range is limited to 30% or more. The reduction ratio in the two-phase temperature range is preferably 35% or more, and more preferably 40% or more. The upper limit of the reduction ratio in the two-phase temperature range is not particularly limited. From the viewpoint of ensuring flatness of the steel sheet, it is preferably 50% or less.

Here, the Ar3 transformation point can be obtained by the following equation.

Ar3(℃)＝910-273×C-74×Mn-57×Ni-16×Cr-9×Mo-5×Cu

In the formula, the content (mass%) of each element is zero, and the content of the element not contained is zero.

Finish rolling temperature: 650 ℃ above

The finishing temperature (finishing temperature) during hot rolling is 650 ℃ or higher. That is, the hot rolling is completed (ended) at a temperature of 650 ℃ or higher because: when finish rolling is performed at a temperature lower than 650 ℃, a necessary strain or more is applied to ferrite generated by transformation. As a result, the low-temperature toughness of the base material is reduced. The finish rolling temperature is preferably 670 ℃ or higher, more preferably 680 ℃ or higher. The upper limit of the finish rolling temperature is not particularly limited. In order to perform rolling in the ferrite-austenite two-phase temperature range, it is preferable to set the temperature to 710 ℃ or lower.

[ Cooling condition after Hot Rolling ]

In the present invention, after the hot rolling, the hot rolled steel sheet is cooled under the following conditions. By this cooling, the strength of the base material can be improved.

Cooling start temperature: 650 ℃ above

The reason for starting cooling from a temperature above 650 ℃ is because: if cooling is started from a temperature lower than 650 ℃, it is disadvantageous from the viewpoint of hardenability, and there is a possibility that the desired base material strength cannot be obtained. The cooling start temperature is preferably 670 ℃ or higher, more preferably 680 ℃ or higher. The upper limit of the cooling start temperature is not particularly limited. In order to start cooling at a temperature of not more than the Ar3 transformation point, it is preferably not more than 710 ℃.

Cooling stop temperature: a temperature in a temperature range of 600 ℃ or less and 300 ℃ or more

The hot-rolled steel sheet after hot rolling is cooled from the cooling start temperature to a temperature (cooling stop temperature) in a temperature range of 300 ℃ to 600 ℃. This is because, when cooling is stopped at a temperature exceeding 600 ℃, it is difficult to secure sufficient strength from the standpoint of hardenability. Further, since the cooling stop at a temperature lower than 300 ℃ does not impart a significant change in the properties of the steel sheet, only the operational load becomes large. The cooling stop temperature is preferably 570 ℃ or lower, more preferably 520 ℃ or lower.

Average cooling rate: 5 ℃/s or more

When the average cooling rate in the above temperature range is less than 5 ℃/s, it is difficult to obtain a steel having a uniform metal structure, and the strength and low-temperature toughness of the base material cannot be ensured. The average cooling rate is preferably 7 ℃/s or more, more preferably 10 ℃/s or more. The average cooling rate is preferably 100 ℃/s or less, more preferably 80 ℃/s or less, and even more preferably 60 ℃/s or less.

For this reason, after finishing hot rolling at a finish rolling temperature of 650 ℃ or higher, the steel sheet is cooled at an average cooling rate of 5 ℃/s or higher from a cooling start temperature of 650 ℃ or higher to a cooling stop temperature of 300 ℃ or higher and 600 ℃ or lower.

The high-strength steel sheet produced under the above production conditions has the above metal structure in addition to the above component composition. That is, the ferrite structure is composed of a ferrite structure as a main phase and a structure having 1 or more than 2 kinds of pearlite, bainite, and martensite as the balance, and has the above processed ferrite structure.

As described above, the high-strength steel sheet of the present invention has characteristics suitable as a steel sheet for the above-described applications (for example, a low-temperature storage tank for liquefied gas, a structure for a ship or the like to be operated in a low-temperature environment), and the like. In particular, it is possible to obtain a base material having a yield stress of 235MPa or more and a high low-temperature toughness, specifically a Charpy impact absorption energy of 200J or more at-60 ℃. In addition, when the high-strength steel sheet of the present invention is used as a joint for welding and the steel sheets are welded by high heat input, the Charpy impact absorption energy at-60 ℃ in the weld heat affected zone is 80J or more, and the low-temperature toughness in the weld heat affected zone is also excellent.

Examples

The present invention will be specifically described below based on examples. The following examples illustrate preferred examples of the present invention, but the present invention is not limited to these examples.

Billets (steel materials) having the composition shown in table 1 were produced by a converter-ladle refining-continuous casting method. The blank column in table 1 indicates that an element is not intentionally added, and means that the element is not only not contained (0%) but also inevitably contained.

The billets thus obtained were heated, cooled, hot-rolled, and then cooled under the various conditions shown in tables 2-1 and 2-2, to obtain high-strength steel sheets (thick steel sheets) having a sheet thickness (finish rolling thickness) of 8 to 25 mm.

First, using each of the obtained high-strength steel sheets, the metal structure was evaluated by the following method.

(1) Evaluation of metallic Structure of high-Strength Steel sheet

[ Metal Structure ]

Samples were cut from the obtained high-strength steel sheet so that the 1/4 position of the sheet thickness of the steel sheet was set as the center and the surface perpendicular to the sheet width direction was set as the observation surface. The surface of the sample was mirror polished, etched with an aqueous solution of nitric acid and then observed with an optical microscope (magnification: 200 times). The structure was photographed, and the ferrite structure fraction (% by area) was calculated using an image analysis device. The 10 fields of view were observed, and the average value of the structure fraction (area%) of each ferrite was calculated. Since the area ratio corresponds to the volume ratio when the anisotropy of the microstructure is small, the average ferrite area ratio is referred to as the volume ratio in this patent.

In tables 3-1 and 3-2, ferrite is represented by F, pearlite is represented by P, bainite is represented by B, and martensite is represented by M.

[ fraction of processed ferrite ]

Regarding the fraction of processed ferrite, a sample was cut from the 1/4 position of the plate thickness of the steel plate in the same manner as described above, the surface was mirror polished, fine polished with colloidal silica, and then EBSD (electron beam back scattering diffraction) measurement was performed using SEM (scanning electron microscope). The crystal orientation was measured at 500 times magnification, and a boundary at which the orientation difference between adjacent measurement points was 15 ° or more was defined as a crystal grain boundary based on the obtained data. Then, a ferrite region in which the GAM (Grain Average Misorientation, average grain orientation difference) value of each grain surrounded by the grain boundaries was 1.0 or more was defined as processed ferrite, and the area fraction was obtained. 10 fields of view were measured, and the average area ratio of processed ferrite was used as the volume ratio in the same manner as described above.

[ aspect ratio of processed ferrite ]

Regarding the aspect ratio of the processed ferrite, a sample was cut from a 1/4 position of the plate thickness of the steel plate in the same manner as described above, and the surface was mirror polished and etched to expose the processed ferrite grain boundaries. Then, 10 to 20 fields of view were photographed at a magnification of 200 times by an optical microscope, and the maximum length in the rolling direction was divided by the maximum length in the plate thickness direction for each processed ferrite grain in each field of view, to calculate the average value of all grains. This was used as the aspect ratio of the processed ferrite.

[ average grain size of processed ferrite ]

The average grain size of the processed ferrite was measured by cutting a sample from a position 1/4 of the plate thickness of the steel plate in the same manner as described above, mirror-polishing the surface, finely polishing the surface with colloidal silica, and then performing EBSD measurement using SEM. The crystal orientation was measured at 500 times magnification, and a boundary at which the orientation difference between adjacent measurement points was 15 ° or more was defined as a crystal grain boundary based on the obtained data. Then, the ferrite region having a GAM value of 1.0 or more in each crystal grain surrounded by the grain boundary is defined as processed ferrite, and the area of the processed ferrite is calculated. The circle equivalent diameter equal to the area of the processed ferrite was set as the respective crystal grain size, and the average value of the obtained crystal grain sizes was set as the average crystal grain size of the processed ferrite.

[ size of TiN ]

Regarding the size of TiN, a thin film sample was cut from a position 1/4 of the plate thickness of the steel plate, and the deposition was measured by using a TEM (transmission electron microscope). The diagonal length of TiN, which is a rectangle, was measured by observing 10 fields of view of 1 μm×1 μm. The average value of the diagonal length of all TiNs was calculated as the size of TiN.

Then, using each of the obtained high-strength steel sheets, the characteristics of the base material and the weld heat affected zone after the high heat input welding were evaluated by the following methods.

(2) Evaluation of characteristics of base Material

From the obtained high-strength steel sheets, a tensile test piece according to JIS Z2241 (2011) was cut at a position 1/4 of the sheet thickness. Then, a tensile test based on JIS Z2241 (2011) was performed to measure the yield stress (YP) of the base material. In this example, it was determined that the base material had excellent strength (high strength) with a yield stress of 235MPa or more.

Further, test pieces according to JIS Z2242 (2018) were cut from 1/4 of the plate thickness of each high-strength steel plate. Then, a Charpy impact test was performed in accordance with JIS Z2242 (2018), and Charpy impact absorption energy (vE) at-60℃of the base material was measured _-60 ). In this example, it was determined that the base material had excellent low-temperature toughness when the average value of the absorption energy of 3 pieces at-60℃was 200J or more.

(3) Evaluation of characteristics of weld Heat affected zone

Then, a test piece for producing a welded joint was cut from 1/4 of the plate thickness of each of the obtained high-strength steel plates, and the single-sided 1-pass welding was performed to produce a high heat input welded joint. The welding conditions were set to the weld line energies shown in tables 3-1 and 3-2. JIS No. 4 impact test pieces having the notch position as the joint were cut from these welded joints, and the Charpy impact test was performed to measure Charpy impact absorption energy (vE) at-60℃of the HAZ _-60 ). In this example, it was determined that the average value of the absorption energy of 3 roots at-60℃was 80J or more and that the low-temperature toughness of the HAZ was excellent.

The measurement results are shown in tables 3-1 and 3-2.

As shown in tables 3-1 and 3-2, the high-strength steel sheet of the present invention has a composition satisfying the condition of formula (1) and a metallic structure satisfying the base material at the 1/4 position of the sheet thickness as described above. This confirmed that the base material had both high strength and excellent low-temperature toughness. It was also confirmed that the absorption energy (vE) of the Charpy impact test was measured at the location where the high-strength steel sheet of the present invention was subjected to high heat input welding _-60 ) Meets the requirement of more than 80J and has excellent low-temperature toughness of HAZ.

In contrast, the comparative examples which deviate from the scope of the present invention do not satisfy the above characteristics.

Claims

1. A high-strength steel sheet comprising, in mass%, C: 0.010-0.070%, si:0.01 to 0.50 percent of Mn: 1.00-2.00%, P: less than 0.020%, S:0.0005 to 0.0100 percent of Al:0.035 to 0.100 percent of Ti:0.010 to 0.030% and N:0.0035 to 0.0100%, a carbon equivalent Ceq (IIW) represented by the following formula (1) satisfying 0.25 to 0.35% by mass, a Ti/N satisfying 1.5 to 4.0, and the balance being Fe and unavoidable impurities,

the aspect ratio of the processed ferrite is more than 1.5,

the processed ferrite has an average crystal grain size of 50 μm or less,

Ceq.(IIW)＝[C]+[Mn]/6+([Cu]+[Ni])/15+([Cr]+[Mo]+[V])/5…(1)

2. The high-strength steel sheet according to claim 1, wherein the composition of the components further contains, in mass%, a composition selected from the group consisting of B: less than 0.0030%, cu: less than 0.50%, ni: less than 1.50%, V: less than 0.100%, cr: less than 0.50%, mo: less than 0.50%, ca: less than 0.0030 percent of Mg:0.0050% or less and REM:0.1000% or less of 1 or 2 or more of the group consisting of.

3. The high-strength steel sheet according to claim 1 or 2, wherein TiN has a size of 5 to 200nm in a metal structure of the high-strength steel sheet at a plate thickness 1/4 position.

4. A method for producing a high-strength steel sheet according to any one of claims 1 to 3, wherein,

heating the steel stock having the composition of the components to a temperature above 1050 ℃ and below 1200 ℃,

5. The method for producing a high-strength steel sheet according to claim 4, wherein the steel stock is cast at an average casting speed of 0.3 to 1.0 m/min.