阅读说明：本技术 耐冲击性优异的热轧钢板、钢管、部件及其制造方法 (Hot-rolled steel sheet, steel pipe, and component having excellent impact resistance, and method for producing same ) 是由 成焕球 赵悦来 裵成范 于 2018-12-14 设计创作，主要内容包括：本发明的优选的方面提供一种耐冲击性优异的热轧钢板、利用该热轧钢板的钢管和部件以及它们的制造方法,以重量％计,所述热轧钢板包含：C：0.35-0.55％；Mn：0.7-1.5％；Si：0.3％以下(0％除外)；P：0.03％以下(包括0％)；S：0.004％以下(包括0％)；Al：0.04％以下(0％除外)；Cr：0.3％以下(0％除外)；Mo：0.3％以下(0％除外)；Ni：0.1-1.0％和Cu：0.1-1.0％中的一种或两种,Cu+Ni：0.4％以上；N：0.006％以下(0％除外)；余量的Fe和其它杂质,所述合金元素满足下述关系式1至关系式3,以体积％计,微细组织包含10％以上的铁素体和90％以下的珠光体。[关系式1](Mn/Si)≥3(重量比)[关系式2](Ni+Cu)/(C+Mn)≥0.2(重量比)[关系式3](Ni/Si)≥1(重量比)。(A preferred aspect of the present invention provides a hot-rolled steel sheet excellent in impact resistance, a steel pipe and a part using the same, and methods for producing the same, the hot-rolled steel sheet comprising, in wt%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); and Fe and other impurities in balance, the alloying elements satisfying the following relational expressions 1 to 3, and the fine structure containing 10% or more of ferrite and 90% or less of pearlite in volume%. [ equation 1] (Mn/Si) is not less than 3 (weight ratio) [ equation 2] (Ni + Cu)/(C + Mn) is not less than 0.2 (weight ratio) [ equation 3] (Ni/Si) is not less than 1 (weight ratio).)

1. A hot-rolled steel sheet excellent in impact resistance, comprising in wt.%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; fe and other impurities in balance, the alloying elements satisfying the following relational expressions 1 to 3, the fine structure containing 10 to 30% of ferrite and 70 to 90% of pearlite in volume%,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio).

2. The hot-rolled steel sheet excellent in impact resistance according to claim 1, characterized in that the hot-rolled steel sheet further comprises a component selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

3. The hot rolled steel sheet excellent in impact resistance according to claim 1, characterized in that the hot rolled steel sheet has a tensile strength of 600-1000 MPa.

4. A method for manufacturing a hot rolled steel sheet excellent in impact resistance, comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling; and

and carrying out laminar flow cooling on the hot-rolled steel plate, and carrying out winding at the temperature of 550-750 ℃.

5. The method of manufacturing a hot-rolled steel sheet excellent in impact resistance according to claim 4, characterized in that the steel slab further comprises a component selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

6. The method of manufacturing a hot rolled steel sheet excellent in impact resistance according to claim 4, characterized by further comprising a step of subjecting the hot rolled steel sheet to pickling treatment to obtain a hot-rolled pickled steel sheet.

7. A steel pipe comprising, in weight%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; fe and other impurities in balance, the alloying elements satisfying the following relational expressions 1 to 3, the fine structure containing 10 to 60% of ferrite and 40 to 90% of pearlite in volume%,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio).

8. The steel pipe of claim 7, further comprising a metal selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

9. A method of manufacturing a steel pipe comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling;

carrying out laminar cooling on the hot-rolled steel plate, and rolling at the temperature of 550-750 ℃;

welding the hot rolled steel sheets to obtain a steel pipe; and

and carrying out annealing heat treatment on the steel pipe.

10. The method of making a steel pipe of claim 9 wherein said steel slab further comprises a material selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

11. The method of manufacturing a steel pipe as claimed in claim 9 further comprising a step of drawing after the annealing heat treatment step.

12. The method of manufacturing a steel pipe as claimed in claim 9 or 11 wherein the annealing heat treatment of the steel pipe is Ac₁-50 ℃ to Ac₃At +150 ℃ for 3-60 minutes.

13. A part, comprising in weight percent: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; fe and other impurities in balance, the alloying elements satisfying the following relational expressions 1 to 3, and the microstructure containing 90% or more of one or both of martensite and tempered martensite and 10% or less of retained austenite,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio).

14. The component of claim 13, further comprising an additive selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

15. The component of claim 13, wherein the component has a yield strength of 1400MPa or greater and a tensile strength of 1800MPa or greater.

16. A method of manufacturing a component, comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less and 0% or less is excluded; p: less than 0.03% and including 0%; s: less than 0.004% and including 0%; al: 0.04% or less except 0%; cr: 0.3% or less and 0% or less is excluded; mo: 0.3% or less and 0% or less is excluded; ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% except 0%; the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling;

carrying out laminar cooling on the hot-rolled steel plate, and rolling at the temperature of 550-750 ℃;

welding the hot rolled steel sheets to obtain a steel pipe;

carrying out annealing heat treatment and drawing on the steel pipe;

hot forming the drawn steel pipe to obtain a part; and

and carrying out quenching treatment or quenching and tempering treatment on the component.

17. The method of manufacturing a component of claim 16, wherein said steel blank further comprises a material selected from the group consisting of Ti: 0.04% or less and 0% or less except, B: 0.005% or less and 0% or less excluding Sb: 0.03% or less and 0% or more of the other than the above.

18. According to claimThe method for producing a member according to claim 16, wherein the annealing heat treatment of the steel pipe is performed at Ac₁-50 ℃ to Ac₃At +150 ℃ for 3-60 minutes.

19. The method of manufacturing a component according to claim 16, wherein a cooling rate at the time of the quenching treatment is 10 to 70 ℃/sec.

20. The method of claim 16, wherein the tempering is performed at a tempering temperature of 150-230 ℃ for 120-3600 seconds.

Technical Field

The present invention relates to a hot-rolled steel sheet used for automobile body members such as suspension members of automobiles, steel pipes and parts using the hot-rolled steel sheet, and methods for producing the same, and more particularly, to a hot-rolled steel sheet excellent in impact resistance and rust resistance and exhibiting ultra high strength after heat treatment, a steel pipe and parts using the hot-rolled steel sheet, and methods for producing the same.

Background

Among automotive body members, suspension members are one of members that require high strength-high toughness, corrosion resistance, fatigue durability, and the like, and hot-rolled steel sheets are mainly used.

In addition, such a suspension member is manufactured by subjecting a tubular member to hot forming or cold forming and heat treatment, and in many cases, it is known that premature breakage occurs in the manufacturing process of the member or the use environment of the member. This is known to be caused by various causes, but is thought to be basically caused by quench cracking (quenchcrack) generated in the process of manufacturing a steel pipe using the produced steel sheet, or hydrogen-induced delayed fracture caused by hydrogen atoms and/or molecules mixed into the inside of the steel pipe in the manufacturing process or use environment. Among them, the Hydrogen Induced Delayed fracture includes technical terms such as Hydrogen embrittlement (Hydrogen fracture), Hydrogen Delayed fracture (Hydrogen Delayed Cracking), and Hydrogen Induced Cracking (Hydrogen Induced Cracking). This is indicated to have a significant effect on ultra-high strength steel sheets or pipes having a tensile strength after heat treatment of 1800MPa or more.

In addition, as one of methods for improving the fatigue durability of steel pipe members, various studies have been made to identify the cause of Hydrogen Delayed Fracture (Hydrogen Delayed Fracture or Hydrogen induced Cracking) and to derive an improvement method in order to suppress premature breakage or premature Fracture of steel pipe members.

Patent document 1 discloses that delayed fracture is suppressed for at least about 24 hours under the conditions of U-bend and immersion in HCl (pH 1) in a cold-rolled steel sheet quenched after annealing heat treatment or a quenched-tempered cold-rolled steel sheet obtained by cold-rolling a steel sheet obtained by adding a large amount of Nb element less than 0.1% to a steel used for cold-rolled steel sheets so as to control the Prior Austenite Grain Size (PAGS) of the steel sheet to be less than 20 μm, preferably less than 15 μm.

Similarly to what is proposed in patent document 2, it is mentioned therein that hydrogen in steel is trapped on grain boundaries refined by Nb or Ti precipitates, and thus the delayed fracture resistance is improved by dispersing the effect of hydrogen of the critical amount causing delayed fracture.

In addition, patent document 1 mentions that in steel in which the amount of Si added is as high as 0.5% or more, it is confirmed that Ni element deteriorates delayed fracture resistance, and therefore Ni element is added in an amount of less than 0.5%, preferably, an impurity level controlled to 0.03% as much as possible. The reason why the delayed fracture characteristics are poor is judged from the experimental results confirmed by bending a steel sheet test piece quenched by rapid cooling at 100 ℃/sec or more (rapid cooling in water) in a U-shape and immersing the steel sheet test piece in HCl acid, or by quenching-tempering the heat-treated steel sheet test piece: the hydrogen-induced delayed fracture of the steel is promoted in a form of a critical stress required to reduce the generation or propagation of cracks, because the stress concentration portion is formed by diffusion of hydrogen, which is caused by cracks remaining in the quenched steel sheet having a martensite phase structure, or has entered or is entering the steel, to a plurality of defect sites containing dislocations (dislocations) formed by rapid cooling in water.

Further, a method of improving delayed fracture resistance of steel is proposed, in which local corrosion (pitting) of steel is suppressed, or penetration of hydrogen atoms into the interior of steel is minimized, or various defect sites (sites) including dislocation/grain boundary/precipitate interfaces are formed in the interior of steel to trap the penetrated hydrogen atoms so as not to exceed a critical content. In particular, patent document 2 proposes that a steel containing Si at a level as high as 1 to 3% is subjected to a continuous annealing process and a heating-rapid cooling-tempering process to produce a cold-rolled steel sheet, and when cold forming is performed using the cold-rolled steel sheet, the shape of retained austenite in the formed microstructure composed of bainitic ferrite + martensite + retained austenite is controlled so that the axial ratio (major axis/minor axis) of the retained austenite is 5 or more, thereby suppressing cleavage fracture during observation of the sheet cross section after a tensile test of the steel member, and thus improving the hydrogen embrittlement property. Further, this is a steel sheet having a tensile strength characteristic after heat treatment of less than 1500Mpa, and it is considered that the steel sheet is relatively less sensitive to hydrogen embrittlement than martensite or tempered martensite single-phase structure steel. Further, a method for improving the fatigue life of a wire member by the delayed fracture property of the martensite single-phase structure is proposed, and patent document 3 proposes a method for suppressing hydrogen penetration into the member by controlling the content ratio of B/Cr to less than 0.04 in a steel containing a high content of Si + Cr to form a boron (B) -concentrated layer on the surface layer of the steel member.

In addition, the tempering heat treatment at a relatively high temperature in the range of 350-.

As a result of studies on the production processes of the steel sheets and steel members proposed in the above patent documents, there has been no proposal of a hot-rolled steel sheet and a steel pipe which have impact resistance, have tensile strength of 1800MPa or more in heat treatment of heating-rapid cooling or heating-rapid cooling-tempering, and are excellent in impact resistance and rust resistance, and which do not undergo premature breakage or premature fracture in drawing quenched steel, and a method for producing the same.

(patent document 1) Korean laid-open patent No. 10-2016-0086877

(patent document 2) Korean laid-open patent No. 10-2006-0076741

(patent document 3) Korean laid-open patent No. 10-2007-0068665

Disclosure of Invention

Technical problem to be solved

It is an object of a preferred aspect of the present invention to provide a hot rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture at the time of a tensile test, and exhibiting ultra high strength after heat treatment.

It is another object of a preferred aspect of the present invention to provide a method of manufacturing a hot-rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture in a tensile test, and exhibiting ultra high strength after heat treatment.

It is another object of a preferred aspect of the present invention to provide a steel pipe manufactured using a hot rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture at the time of a tensile test, and exhibiting ultra high strength after heat treatment.

A preferred another aspect of the present invention provides a method of manufacturing a steel pipe using a hot rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture in a tensile test, and exhibiting ultra high strength after heat treatment.

A preferred another aspect of the present invention provides a component using a steel pipe manufactured using a hot-rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture at the time of a tensile test, and exhibiting ultra high strength after heat treatment.

A preferred another aspect of the present invention provides a method of manufacturing a part using a steel pipe manufactured using a hot rolled steel sheet having excellent impact resistance and rust resistance even under a short natural aging time, not causing premature breakage and abnormal fracture at the time of a tensile test, and exhibiting ultra high strength after heat treatment.

Technical scheme

According to a preferred aspect of the present invention, there is provided a hot rolled steel sheet excellent in impact resistance, comprising in weight%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3, and the fine structure comprising 10 to 30% of ferrite and 70 to 90% of pearlite in volume%.

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio)

According to another preferred aspect of the present invention, there is provided a method of manufacturing a hot rolled steel sheet excellent in impact resistance, the method comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is more than or equal to 3 (weight ratio),

[ relational expression 2]

(Ni + Cu)/(C + Mn) is not less than 0.2 (weight ratio),

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling; and

and carrying out laminar flow cooling on the hot-rolled steel plate, and carrying out winding at the temperature of 550-750 ℃.

The method of manufacturing a hot rolled steel sheet excellent in impact resistance may further include a step of subjecting the hot rolled steel sheet to pickling treatment to obtain a hot-rolled pickled steel sheet.

According to another preferred aspect of the present invention, there is provided a steel pipe comprising, in weight%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3, and the fine structure comprising 10 to 60% of ferrite and 40 to 90% of pearlite in volume%.

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio)

According to another preferred aspect of the present invention, there is provided a method of manufacturing a steel pipe, the method comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is more than or equal to 3 (weight ratio),

[ relational expression 2]

(Ni + Cu)/(C + Mn) is not less than 0.2 (weight ratio),

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling;

carrying out laminar cooling on the hot-rolled steel plate, and rolling at the temperature of 550-750 ℃;

welding the hot rolled steel sheets to obtain a steel pipe; and

and carrying out annealing heat treatment on the steel pipe.

The method of manufacturing a steel pipe may further include a drawing step after the annealing heat treatment step.

According to another preferred aspect of the present invention, there is provided a component comprising, in weight%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); and Fe and other impurities in balance, the alloy elements satisfying the following relational expressions 1 to 3, and the microstructure including 90% or more of one or both of martensite and tempered martensite and 10% or less of retained austenite.

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio)

According to another preferred aspect of the present invention, there is provided a method of manufacturing a component, the method comprising the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is more than or equal to 3 (weight ratio),

[ relational expression 2]

(Ni + Cu)/(C + Mn) is not less than 0.2 (weight ratio),

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling;

carrying out laminar cooling on the hot-rolled steel plate, and rolling at the temperature of 550-750 ℃;

welding the hot rolled steel sheets to obtain a steel pipe;

carrying out annealing heat treatment and drawing on the steel pipe;

hot forming the drawn steel pipe to obtain a part; and

and carrying out quenching treatment or quenching and tempering treatment on the component.

Advantageous effects

According to a preferred aspect of the present invention, it is possible to provide a hot rolled steel sheet and a steel pipe having excellent impact toughness and rust resistance without early breakage in a tensile test, and having an effect that hydrogen embrittlement that may occur during a manufacturing process of the steel pipe or use (in-service) of a steel pipe member can be reduced.

Drawings

Fig. 1 is a tensile curve showing the fracture morphology of the inventive materials (4, 6, 15) and the comparative material (3) in the examples.

FIG. 2 shows the distribution of copper (Cu) element present in the surface layers of the hot rolled steel sheets of the inventive materials (4) and (12) in the examples.

FIG. 3 shows the distribution of nickel (Ni) elements present in the surface layers of the hot rolled steel sheets of inventive materials (4) and (12) in the examples.

Fig. 4 shows the optical microstructures of the drawn pipe of the invention material (4) before and after the heat treatment in the example, (a) shows the microstructure of the drawn pipe before the heat treatment, and (b) shows the microstructure of the drawn pipe after the heat treatment.

Best mode for carrying out the invention

The present invention will be explained below.

First, a hot-rolled steel sheet excellent in impact resistance, which is a preferred aspect of the present invention, will be described.

A hot-rolled steel sheet excellent in impact resistance according to a preferred aspect of the present invention includes, in wt%: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3.

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio)

C: 0.35 to 0.55 wt% (hereinafter, also referred to as "%")

The carbon (C) is an effective element for improving the strength of the steel, and increases the strength after the quenching heat treatment. When the content of carbon is less than 0.35%, it is difficult to secure sufficient strength of 1800Mpa or more after the tempering heat treatment, and on the other hand, when the content of carbon exceeds 0.55%, martensite having excessively high hardness is formed, so that cracks are generated in the steel plate material or the steel pipe member, possibly resulting in deterioration of fatigue durability. Therefore, the content of carbon (C) is preferably limited to 0.35 to 0.55%.

Mn：0.7-1.5％

The manganese (Mn) is an essential element for improving the strength of the steel, and increases the strength after the quenching heat treatment of the steel. When the content of manganese is less than 0.7%, it is difficult to secure sufficient strength of 1800Mpa or more after the tempering heat treatment, and on the other hand, when the content of manganese exceeds 1.5%, segregation zones may be formed inside and/or outside the continuous casting slab and the hot rolled steel sheet, and high frequency machining defects may be caused when manufacturing the steel pipe. Also, after the tempering heat treatment, the strength may be excessively increased, and the fatigue durability may be deteriorated. Therefore, the content of manganese (Mn) is preferably limited to 0.7 to 1.5%.

Si: below 0.3% (except 0%)

The silicon (Si) is an element added to improve strength or ductility, and is added within a range in which the problem of surface scale of the hot-rolled steel sheet and the hot-rolled pickled steel sheet does not occur. When the content of silicon exceeds 0.3% or more, silicon oxide is formed and surface defects occur, and it is difficult to remove them by acid washing, so that the content of silicon is limited to 0.3% or less (except 0%).

P: less than 0.03% (including 0%)

The phosphorus (P) segregates at austenite grain boundaries and/or phase boundaries and may induce brittleness. Therefore, the content of phosphorus (P) is kept as low as possible, with the upper limit limited to 0.03%. The content of phosphorus (P) is preferably 0.02% or less. In the present invention, the presence of S element is confirmed at the position of the steel where the quench crack is generated at the time of quenching, not the P content, and therefore the control of P is relatively less strict, but phosphate (H) of the pipe is performed for removing scale in the drawing manufacturing process of the pipe₃PO₄) When the pickling treatment is not performed properly after the treatment, the content of the P element is preferably controlled to a low level because defects in the inner wall of the steel pipe may be caused by the remaining P element.

S: less than 0.004% (including 0%)

The sulfur (S) forms MnS nonmetallic inclusions in the steel or segregates during continuous casting solidification to possibly induce high-temperature cracks. Further, the impact toughness of the heat-treated steel sheet or pipe may be deteriorated, and therefore, it is necessary to control the sulfur content to as low a level as possible. Therefore, in the present invention, the content of sulfur (S) is kept as low as possible, and the upper limit of the content of sulfur (S) is preferably limited to 0.004%.

Al: below 0.04% (except 0%)

The aluminum (Al) is an element added as a deoxidizer. In addition, the aluminum reacts with nitrogen (N) in steel to precipitate AlN, and slab cracking is caused under the slab cooling condition in which these precipitates precipitate when manufacturing a thin slab, and thus the quality of the cast slab or hot-rolled steel sheet may be degraded. Therefore, the content of aluminum (Al) is preferably limited to 0.04% or less (except for 0%).

Cr: below 0.3% (except 0%)

The chromium (Cr) is an element that increases the hardenability at the quenching heat treatment of steel and increases the heat treatment strength by delaying the transformation of austenite to ferrite. When more than 0.3% of chromium (Cr) is added to a steel containing 0.35% or more of carbon (C), excessive hardenability of the steel may be induced, and therefore the content of chromium (Cr) is limited to 0.3% or less (except for 0%).

Mo: below 0.3% (except 0%)

The molybdenum (Mo) increases the hardenability of the steel and forms fine precipitates, thereby making it possible to refine the austenite crystal grains. Further, it is effective to improve the strength and toughness of the steel after heat treatment, but when the content of molybdenum exceeds 0.3%, the manufacturing cost of the steel may be increased, so the content of molybdenum (Mo) is limited to 0.3% or less (except 0%).

In the present invention, one or both of Ni and Cu are contained.

Ni：0.1-1.0％

The nickel (Ni) is an element that improves both hardenability and toughness of steel. In addition, in the present invention, when tensile properties are evaluated after heat treatment of a steel sheet or steel pipe in which the content of nickel (Ni) is increased in the base component, the strength after heat treatment decreases as the content of Ni increases, which is considered to be because the movement of dislocations introduced into martensite is promoted by the nickel (Ni) element. When the content of nickel is less than 0.1%, the effect of increasing the hardenability and toughness is insufficient, and on the other hand, when the content of nickel exceeds 1.0%, the manufacturing cost of the steel plate is sharply increased and the weldability for manufacturing the steel pipe may be deteriorated in spite of the above-described advantages. In addition, the increase in the Ni content suppresses diffusion of hydrogen that is concentrated on the surface of the heat-treated member and flows into the interior of the member, and/or suppresses penetration of hydrogen by forming a dense corrosion product (Cu — Ni-rich FeOOH) in a corrosive environment, thereby having a beneficial effect of increasing resistance to stress corrosion cracking. Therefore, the content of nickel (Ni) is limited to the range of 0.1 to 1.0%.

Cu：0.1-1.0％

The copper (Cu) is an alloy element that increases corrosion resistance of steel and can effectively increase quenching and quenching-tempering strength after heat treatment. When the copper content is less than 0.1%, it is difficult to secure the above effects, while when the copper content exceeds 1.0%, cracks are generated in the hot-rolled steel sheet, and thus the steel sheet production yield is reduced, the strength rapidly increases after heat treatment to generate cracks, or the strength rapidly increases after heat treatment to reduce the toughness. Therefore, the content of copper (Cu) is limited to the range of 0.1 to 1.0%. In addition, copper (Cu) itself may cause surface cracks of the hot-rolled steel sheet, and therefore, it is more preferable to use it together with nickel (Ni) element than to use copper (Cu) alone.

Cu + Ni: more than 0.4 percent

The sum of Cu + Ni is important to increase the rust resistance and toughness of the steel plate and the steel pipe.

In the present invention, when the sum of the contents of Cu + Ni added to the steel containing 0.35% or more of carbon (C) is less than 0.4%, it is difficult to simultaneously secure the above-described effects, and therefore the sum of Cu + Ni is set to 0.4% or more. Further, when a steel sheet or a steel pipe member added with the sum of Cu + Ni of 0.4% or more in a steel containing carbon (C) and manganese (Mn) in appropriate contents is subjected to heat treatment, it was confirmed that the advantageous effects of a reduced depth of a decarburized layer generated in the surface layer of the steel sheet or the steel pipe member, an improved impact toughness, rust resistance, and the like are exhibited. In particular, an increase in the depth of the decarburized layer becomes a factor of deteriorating fatigue durability of the steel pipe member. Therefore, the sum of the Cu + Ni contents is limited to 0.4% or more.

N: less than 0.006% (except 0%)

The nitrogen (N) is an austenite stabilizing element and a nitride forming element. If the content of nitrogen (N) exceeds 0.006%, coarse AlN nitrides are formed, which may cause fatigue cracks to occur when evaluating the fatigue durability of heat-treated steel sheets or steel pipe members, and may deteriorate the fatigue durability. Therefore, the content of nitrogen (N) is limited to 0.006% or less (except for 0%).

Further, when nitrogen (N) is added together with boron (B) element, in order to increase the effective boron (B) content, it is necessary to control the nitrogen (N) content to be as low as possible.

The Mn and Si need to satisfy the following relational formula 1.

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

The ratio of Mn to Si is an important parameter for determining the quality of a welded portion of a steel pipe. When the ratio of Mn/Si is less than 3, the content of Si is relatively high, silicon oxide is formed in the molten metal in the welded portion, and if forced discharge is not performed, defects are formed in the welded portion, which may cause defects in the production of the steel pipe, and therefore, the ratio of Mn/Si is limited to 3 or more.

The C, Mn, Ni and Cu need to satisfy the following relational expression 2.

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

The ratio of (Ni + Cu)/(C + Mn) is a necessary condition for securing strength after quenching or quenching-tempering heat treatment and securing a satisfactory level of impact toughness and hydrogen embrittlement resistance. When the ratio of (Ni + Cu)/(C + Mn) is less than 0.2, quench cracks may occur at the time of water quenching or water quenching and oil quenching or oil quenching, or hydrogen-induced delayed fracture of the steel pipe or steel pipe member may occur without long-term natural aging after quenching. Further, when the ratio of (Ni + Cu)/(C + Mn) exceeds 0.2 or more, there is an advantage that hydrogen-induced delayed fracture can be effectively suppressed by natural aging for a short time at the time of quenching of steel.

The Ni and Si need to satisfy the following relational expression 3.

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio)

The ratio of Ni/Si is an important parameter affecting the quenched strength by the quenching heat treatment or the tempered strength by the quenching-tempering heat treatment of the steel. The present invention is characterized in that a relatively higher content of nickel (Ni) element is added as compared with silicon (Si) element. When the ratio of Ni/Si is less than 1, the content of silicon (Si) in the steel is relatively high, and the strength of the hot rolled steel sheet is relatively high, so that the deformation resistance of the material to hot rolling increases, and it is difficult to manufacture a thin hot rolled steel sheet having a thickness of less than 3mm, for example. On the other hand, when the ratio of Ni/Si is 1 or more, the content of Ni is relatively high, the strength of the hot-rolled steel sheet is relatively low, and the quenching strength and the quenching-tempering strength are relatively low, so that it is advantageous in securing the toughness of the hot-rolled steel sheet or steel pipe member, and the fraction of retained austenite remaining inside the martensite or tempered martensite phase by the quenching or quenching-tempering heat treatment is relatively small, so that the critical content of diffusible hydrogen trapped at the austenite/base iron interface may be high, but hydrogen permeating into the interior of the heat-treated steel sheet or steel pipe member may be blocked at a relatively high amount, so that it is considered that the hydrogen embrittlement resistance may be further improved. In addition, an increase in the content of retained austenite in martensite or tempered martensite may become a factor for reducing the durability of steel. Therefore, the ratio of Ni/Si is limited to 1 or more.

In the present invention, the balance is composed of Fe and other impurities in addition to the above components.

In addition, in order to improve other characteristics, other alloying elements may be further added to the steel having the above-described composition.

In the present invention, the composition may further comprise a compound selected from Ti: 0.04% or less (except 0%), B: 0.005% or less (except 0%) and Sb: 0.03% or less (except 0%).

Ti: below 0.04% (except 0%)

The titanium (Ti) is an element forming precipitates (TiC, TiCN, TiNbCN) in the hot-rolled steel sheet, and increases the strength of the hot-rolled steel sheet by inhibiting the growth of austenite grains.

When the titanium content exceeds 0.04%, the strength of the quench-temper heat-treated steel is increased and it is effective for trapping diffusible hydrogen at the TiN interface, but when it exists in the form of coarse crystals rather than in the form of fine precipitates in the hot-rolled steel sheet, the toughness is deteriorated or it becomes a starting point for the generation of fatigue cracks, and therefore the fatigue durability of the heat-treated steel sheet or steel pipe member may be lowered. Therefore, the content of titanium (Ti) is limited to 0.04% or less (except for 0%).

B: less than 0.005% (except 0%)

The boron (B) is an advantageous element that can greatly increase the hardenability of the steel even at a low content. When the boron is added in an appropriate amount, the formation of ferrite is suppressed, and thus it is effective for improving hardenability, but when the boron is excessively contained, the austenite recrystallization temperature is increased, and weldability is deteriorated. When the content of the boron (B) exceeds 0.005%, the above-mentioned effects are saturated or it is difficult to secure appropriate strength and toughness. Therefore, the content of boron (B) is limited to 0.005% or less. More preferably, the content of boron (B) is limited to 0.003% or less, which is more effective for ensuring both strength and toughness of the heat-treated steel.

Sb: below 0.03% (except 0%)

The antimony (Sb) element is an advantageous element that can suppress decarburization of the surface layer of the high carbon hot-rolled steel sheet. When the appropriate amount of antimony is added, it is effective to suppress decarburization of the surface layer of the steel sheet by enriching the antimony in the surface layer of the hot-rolled steel sheet, but when the amount of antimony is excessively increased, high-temperature ductility of the steel is lowered during cooling of the slab, and cracks are generated at the edge portion of the slab, thereby deteriorating the surface quality of the slab. When the content of antimony (Sb) exceeds 0.03%, the effect of suppressing decarburization is saturated or the surface quality of the slab is deteriorated, and defects are generated on the surface of the hot-rolled steel sheet, so that the yield of the hot-rolled coil may be reduced. Therefore, the content of antimony (Sb) is limited to 0.03% or less. More preferably, the content of antimony (Sb) is limited to 0.02% or less, which is more effective for ensuring both the effect of suppressing surface decarburization and the surface quality of the slab or hot-rolled steel sheet.

A hot rolled steel sheet excellent in impact resistance and rust resistance of a preferred aspect of the present invention has a fine structure containing 10 to 30% by volume of ferrite and 70 to 90% by volume of pearlite. When the fraction of ferrite is less than 10%, since the content of pearlite excessively increases, the strength becomes high, and thus it may be difficult to manufacture a steel sheet having a thickness of, for example, 3mm or less. Therefore, the fraction of ferrite is preferably limited to 10% or more. The fraction of ferrite is preferably 10 to 30%.

The hot rolled steel sheet may have a thickness of 2 to 7 mm.

The hot rolled steel sheet may have a tensile strength of 600 and 1000 Mpa.

Hereinafter, a method for producing a hot-rolled steel sheet excellent in impact resistance and rust resistance according to a preferred aspect of the present invention will be described.

A method of manufacturing a hot rolled steel sheet excellent in impact resistance and rust resistance according to a preferred aspect of the present invention includes the steps of:

heating a steel slab to a temperature in the range of 1150-: c: 0.35-0.55%; mn: 0.7-1.5%; si: 0.3% or less (except 0%); p: less than 0.03% (including 0%); s: less than 0.004% (including 0%); al: 0.04% or less (except 0%); cr: 0.3% or less (except 0%); mo: 0.3% or less (except 0%); ni: 0.1-1.0% and Cu: 0.1-1.0%, one or two of Cu + Ni: more than 0.4 percent; n: less than 0.006% (excluding 0%); the balance of Fe and other impurities, the alloying elements satisfying the following relational expressions 1 to 3,

[ relational expression 1]

(Mn/Si) is not less than 3 (weight ratio)

[ relational expression 2]

(Ni + Cu)/(C + Mn) ≥ 0.2 (weight ratio)

[ relational expression 3]

(Ni/Si) is more than or equal to 1 (weight ratio);

hot rolling the heated slab at an Ar3 or higher temperature to obtain a hot rolled steel sheet, the hot rolling including rough rolling and finish rolling; and

and carrying out laminar flow cooling on the hot-rolled steel plate, and carrying out winding at the temperature of 550-750 ℃.

Heating step of billet

The steel billet with the composition is heated to the temperature range of 1150-1300 ℃.

The reason for heating the slab to the temperature range of 1150-1300 ℃ is to have a uniform structure and composition distribution in the slab, and when the heating temperature of the slab is as low as less than 1150 ℃, precipitates formed in the continuously cast slab cannot be solid-dissolved and uniformity of composition cannot be ensured.

In addition, when the heating temperature of the slab exceeds 1300 ℃, the decarburizing depth excessively increases and crystal grains grow, so that it is difficult to secure the target material and surface quality of the hot-rolled steel sheet. Therefore, the heating temperature of the billet is limited to the range of 1150-1300 ℃.

Step of obtaining Hot rolled Steel sheet

At Ar₃And (4) carrying out hot rolling on the heated steel billet to obtain a hot rolled steel plate, wherein the hot rolling comprises rough rolling and finish rolling.

The hot rolling is preferably performed at Ar₃The hot finish rolling was performed as described above. When the hot rolling is below Ar₃When the temperature of (3) is increased, a part of austenite is transformed into ferrite, and the material is unevenly deformed in hot rolling, and the passability including the flatness of the steel sheet is deteriorated, so that there is a high possibility of occurrence of an operation failure such as a sheet break. In particular, when the finish rolling temperature exceeds 950 ℃, scale defects and the like are generated, so the finish rolling temperature is preferably limited to 950 ℃ or less.

Winding step

The hot rolled steel sheet obtained by hot rolling as described above was subjected to laminar cooling and was wound at a temperature of 550-.

The reason why the hot rolling is followed by the laminar cooling and the rolling is performed in the temperature range of 550-750 ℃ is to ensure a uniform material of the hot rolled steel sheet, and when the rolling temperature is excessively low below 550 ℃, a low-temperature phase change phase such as bainite or martensite is introduced into an edge portion in the width direction of the steel sheet, so that the strength of the steel sheet may be sharply increased and a deviation of the hot rolled strength in the width direction is increased.

In addition, when the coiling temperature exceeds 750 ℃, internal oxidation of the surface layer portion of the steel sheet is promoted, and surface defects such as cracks or surface irregularities may be generated on the surface after hot rolling pickling. In addition, coarsening of pearlite may cause variation in surface hardness of the steel sheet. Therefore, the coiling temperature of the hot rolled steel sheet after cooling is limited to 550-750 ℃.

In the present invention, the hot-rolled steel sheet produced as described above may be further subjected to pickling treatment to produce a hot-rolled pickled steel sheet. The pickling method is not limited to a specific method, and any pickling method may be used as long as it is a pickling method generally used in a hot rolling pickling process.

According to the method for producing a hot-rolled steel sheet excellent in impact resistance and rust resistance according to a preferred aspect of the present invention, a hot-rolled steel sheet having a fine structure containing 10% by volume or more of ferrite and 90% by volume or less of pearlite can be produced.

The hot rolled steel sheet may have a thickness of 2 to 7 mm.

The hot rolled steel sheet may have a tensile strength of 600 and 1000 Mpa.

Hereinafter, a steel pipe and a method for manufacturing the same according to another preferred aspect of the present invention will be described.

A steel pipe of another preferred aspect of the present invention is produced using the hot rolled steel sheet of the present invention described above, and has the alloy composition of the hot rolled steel sheet of the present invention described above and a fine structure containing 10 to 60% ferrite and 40 to 90% pearlite in volume%. Preferably, the fine structure of the steel pipe may include 20 to 60% ferrite by volume%.

A method of manufacturing a steel pipe according to another preferred aspect of the present invention is a method of manufacturing a steel pipe using a hot-rolled steel sheet manufactured by the method of manufacturing a hot-rolled steel sheet according to the present invention described above.

A method of manufacturing a steel pipe according to another preferred aspect of the present invention includes the steps of: welding the hot rolled steel sheet manufactured by the method of manufacturing a hot rolled steel sheet of the present invention described above to obtain a steel pipe; and annealing the steel pipe.

Step of obtaining a Steel pipe

The hot-rolled steel sheet manufactured by the method of manufacturing a hot-rolled steel sheet of the present invention described above is welded to obtain a steel pipe.

The steel pipe is obtained by using the hot-rolled steel sheet or hot-rolled pickled steel sheet, and pipe-making by, for example, electric resistance welding, induction heating welding, or the like.

Annealing heat treatment step of steel pipe

The steel pipe obtained through pipe making as described above is subjected to annealing heat treatment.

The present invention may further comprise the step of drawing the steel pipe subjected to the annealing heat treatment. The steel pipe is cold drawn, so that the pipe diameter of the steel pipe can be reduced. The drawing method may be a cold drawing method.

In the present invention, it is possible to use the hot rolled steel sheet or the hot rolled pickled steel sheet and manufacture a small-diameter steel pipe using a conventional cold forming method including, for example, processes of manufacturing a steel pipe by electric resistance welding or induction heating welding, annealing heating, and cold drawing.

The annealing heat treatment of the steel pipe is preferably Ac₁-50 ℃ to Ac₃At +150 ℃ for 3-60 minutes. The annealing heat treatment may include furnace cooling and air cooling. When the temperature of the annealing heat treatment is too low or the time is short, a Pearlite (Pearlite) band structure is formed in the microstructure of the steel pipe, and the cold rolling reduction rate or the cross-sectional area reduction rate is lowered at the time of cold drawing of the steel pipe. On the other hand, when the annealing heat treatment temperature is too high or the annealing heat treatment is carried out for a long time, coarse spherical Fe is formed in the microstructure of the steel pipe₃C, or the surface layer or the inner wall layer of the steel sheet may be decarburized.

Hereinafter, a member and a method for manufacturing the member according to another preferred aspect of the present invention will be described.

A member according to another preferred aspect of the present invention is produced by the steel pipe according to the present invention, has the alloy composition of the steel pipe according to the present invention, and has a microstructure containing 90% or more of one or both of martensite and tempered martensite and 10% or less of retained austenite.

When the fraction of the martensite and tempered martensite is less than 90%, there is a problem that it is difficult to secure a desired yield strength of 1400MPa or more or a tensile strength of 1800MPa or more. When the content of the retained austenite exceeds 10%, although the hydrogen-induced delayed fracture resistance can be increased by trapping the diffusible hydrogen, it becomes a fatigue crack site, and thus the fatigue durability may be lowered.

The component of another preferred aspect of the invention may have a yield strength of 1400MPa or more and a tensile strength of 1800MPa or more.

The parts of another preferred aspect of the present invention have an ultra high strength after heat treatment, and have excellent impact resistance and rust resistance even at a short natural aging time of less than 45 hours, and do not undergo premature breakage or abnormal fracture in a tensile test.

A method of manufacturing a component of another preferred aspect of the present invention comprises the steps of: subjecting the steel pipe obtained by the above-described method for manufacturing a steel pipe of the present invention to annealing heat treatment and drawing; hot forming the drawn steel pipe as described above to obtain a part; and subjecting the component to a quenching treatment or a quenching and tempering treatment.

Step of obtaining a part

The steel pipe drawn as described above is shaped to obtain a part.

The steel pipe can be formed by, for example, heating the steel pipe to a high temperature and hot forming the steel pipe. As an example of the component, a suspension member can be cited.

In the hot forming of the steel pipe, the steel pipe with a specific length is heated to the temperature range of 900-980 ℃, is isothermally kept for 60-1000 seconds or less, is taken out, and is hot formed by a mold or the like to obtain a part.

The heating of the steel pipe to the temperature range of 900-980 ℃ is to austenitize the microstructure of the steel pipe member and to make the composition uniform, and when the heating temperature of the steel pipe is lower than 900 ℃, the temperature is greatly lowered in the hot forming and quenching heat treatment process, and ferrite is formed on the surface of the steel pipe, so that it is difficult to secure sufficient strength after heat treatment. On the other hand, when the heating temperature of the steel pipe exceeds 980 ℃, the austenite grain size of the steel pipe increases, or the inner/outer wall of the steel pipe is decarburized, and thus the fatigue strength of the final member may be reduced.

Further, when heated to the above temperature or higher, it is difficult to secure the target strength after the heat treatment of the final member. Therefore, the heating temperature of the steel pipe is preferably limited to the temperature range of 900-.

In addition, in order to ensure the above-mentioned sufficient heat treatment strength and prevent the occurrence of decarburization, heat treatment is performed for a time in the range of 60 to 1000 seconds. When the heating (holding) time is less than 60 seconds, it is difficult to ensure uniform composition distribution and structure, and when the heating and holding time exceeds 1000 seconds, it is difficult to prevent growth or decarburization of crystal grains.

Therefore, the time for holding at the above heating temperature is preferably limited to the range of 60 to 1000 seconds.

Quenching step or quenching and tempering step of component

The member obtained by the above hot forming is subjected to quenching treatment or quenching and tempering treatment.

The heating temperature during the quenching treatment can be 900-.

In the quenching treatment, for example, the hot-formed part may be directly immersed in cooling medium water or oil to be water-cooled or oil-cooled to 200 ℃ or less to form a martensite phase structure.

The quenching heat treatment is performed on the parts obtained by the hot forming by using water, a mixture of water and oil, or oil as a cooling medium, in order to make the structure of the hot formed parts (members) have a martensite phase, and the hot formed members are quenched (rapidly cooled) by immersing the hot formed members in the cooling medium so that the temperature of the parts (members) becomes 200 ℃ or lower. At this time, for example, in a temperature range from Ms (martensite start temperature) to Mf (martensite finish temperature), the cooling rate may be 10 to 70 ℃/sec.

In the Ms to Mf temperature range, when the cooling rate is less than 10 ℃/sec, the martensite phase is difficult to form, and when the cooling rate exceeds 70 ℃/sec, an excessive martensite phase is formed due to a rapid cooling deviation of the inner wall/outer wall of the steel pipe, so that dimensional defects such as a change in the shape of a part (component) or manufacturing defects of a component such as a quench crack are liable to occur. In particular, this remarkably appears in steel sheets or parts (members) exhibiting tensile physical properties after heat treatment of 1800MPa or more, and in order to minimize manufacturing defects of the members, the cooling rate of the parts is preferably limited to the range of 10-70 ℃/sec in the Ms to Mf temperature range.

Further, it is more preferable that the cooling rate is limited to a range of 20 to 60 ℃/sec in order to effectively secure the tensile strength of the member after the heat treatment. In order to ensure the cooling rate, the temperature of the cooling medium of water or oil and water or oil may be raised from normal temperature to high temperature and used.

In the present invention, the component may be subjected to only the quenching treatment as described above, but may be subjected to a tempering treatment after the quenching treatment as described above to impart toughness (toughnesss).

The tempering treatment may be performed by holding the quenched component (member) at a tempering temperature of 150-.

When the tempering temperature is less than 150 ℃, although the strength after the heat treatment is high, the normal temperature impact toughness is very low, and when the tempering temperature exceeds 230 ℃, temper embrittlement (tempembrittlement) in which the total elongation or uniform elongation of the member is drastically reduced may occur, and it is difficult to secure a desired strength after the heat treatment, or alloy elements need to be added to secure sufficient hardenability in order to secure a desired strength after the heat treatment, but it is not preferable from an economical point of view. Further, it is difficult to ensure desired strength. Therefore, the tempering temperature is preferably defined as 150-230 ℃.

In order to ensure sufficient strength and impact toughness after heat treatment, it is preferable to maintain the temperature at the tempering temperature of 150-230 ℃ for 120-3600 seconds.

When the holding time is less than 120 seconds, the dislocation density introduced into the martensite phase of the part subjected to the quenching heat treatment is not greatly changed, and thus the yield strength is low, the tensile strength is very high, and thus the impact toughness is insufficient, and when the holding time exceeds 3600 seconds, although relatively satisfactory impact toughness may be secured, it may be difficult to secure the strength after the heat treatment. Therefore, the time of holding at the tempering temperature is preferably limited to the range of 120-3600 seconds.

According to the method for manufacturing a part of the present invention, it is possible to manufacture a part having an ultra-high strength after heat treatment, which has excellent impact resistance and rust resistance without premature breakage or abnormal fracture in a tensile test even at a short natural aging time of less than 45 hours.

Detailed Description

The present invention will be described in more detail below with reference to examples.

(examples)

Using steels having the compositions shown in tables 1 and 2 below, hot rolling was carried out under the conditions shown in Table 3 below to obtain hot rolled steel sheets having a thickness of 3mm, which were then subjected to pickling treatment. Prior to hot rolling, a slab manufactured on site or a steel slab manufactured in a laboratory is heated in the range of 1200. + -. 20 ℃ for 200 minutes to perform a homogenization treatment, and then each slab or slab is subjected to rough rolling and finish rolling and is wound at a temperature of 600 ℃ and 700 ℃ to prepare a hot rolled steel sheet having a thickness of 3 mm.

In tables 1 and 2 below, the inventive steels (1 to 14) satisfy relational expressions (1) to (3), and the sum of Cu + Ni satisfies 0.4 or more. The comparative steels (1 to 7) do not satisfy at least one of the relational expressions (1) to (3). The Ms temperature is calculated using the empirical formula Ms 539-423C-30.4Mn-12.1Cr-17.7Ni-7.5 Mo.

For the hot rolled steel sheet manufactured as described above, the microstructure, the Yield Strength (YS), the Tensile Strength (TS), and the Elongation (EL) were measured, and the results thereof are shown in table 3 below. The fine structure other than ferrite is pearlite.

The hot rolled steel sheet was pickled, and a steel pipe having a diameter of 28mm was partially produced by resistance welding, annealed and cold drawn to obtain a drawn steel pipe having a diameter of 23.5 mm. At this time, the annealing temperature was 721 ℃. The steel pipes were subjected to heating-hot forming-quenching heat treatment or heating-hot forming-quenching-tempering heat treatment under the conditions of table 4 below to prepare parts.

At this time, the part was heated to 930-950 ℃ and immersed in cooling medium oil for 200 seconds to cool the part to 200 ℃ or lower, and as completely as possible, the part was cooled to room temperature.

After the quenching heat treatment, whether or not cracks were generated in the component was confirmed, and the results thereof are shown in table 4 below. Whether or not cracks are generated is classified into: o, no generation: x, no generation: x (after natural aging time) etc. and shown.

For the parts manufactured as described above, Yield Strength (YS), Tensile Strength (TS), Elongation (EL), Yield Ratio (YR), and impact energy were measured, and the results thereof are shown in table 5 below.

Further, with respect to the parts manufactured as described above, corrosion resistance (rust), fine structure, and surface layer decarburization depth were measured, and the results thereof are shown in table 6 below.

The mechanical physical property values of the hot rolled steel sheet and the member are values obtained by taking and measuring JIS 5 test pieces in a direction parallel to the rolling direction at the w/4 position of the width.

The susceptibility to quench cracking and hydrogen embrittlement is the result of varying the natural aging time of test pieces subjected to a separate quenching heat treatment and performing a tensile test.

The room temperature impact test value was evaluated for a test piece obtained by processing a test piece subjected to quenching-tempering heat treatment into a small size (sub-size) and grinding-off the surface of both sides of the test piece to remove a decarburized layer, according to ASTM E23.

The rust evaluation result is a value obtained by spraying water on the surface of a steel pipe or a flat test piece before/after heat treatment of each steel type, exposing the steel pipe or flat test piece to the air, and measuring the time taken for rust to form on the surface of the test piece. The above results are considered to be indirect evidence that the degree of corrosion resistance of the steel grade can be judged.

The microstructure of the part is measured using a quantitative analysis device including an optical microscope, a scanning Electron microscope, a transmission Electron microscope, an Electron Back Scattering Diffractometer (EBSD).

The depth of the decarburized layer is divided into Ferrite decarburization (full decarburization) and total decarburization (total decarburization) and measured.

Further, the inventive materials (4, 6, and 15) and the comparative material (3) were subjected to a tensile test 45 hours after the natural aging treatment, and the results thereof are shown in fig. 1.

Further, with respect to the hot rolled steel sheets of the invention material (4) and the invention material (12), the distributions of the copper (Cu) element and the nickel (Ni) element in the surface layer portion were analyzed, and the results are shown in fig. 2 and fig. 3, respectively.

The microstructure of the drawn pipe of the inventive material (4) before and after the heat treatment was observed, and the results are shown in fig. 4. Fig. 4 (a) shows the microstructure of the drawn pipe before the heat treatment, and (b) shows the microstructure of the drawn pipe after the heat treatment.

[ Table 1]

Steel grade	C	Si	Mn	P	S	S.Al	Cr	Mo	Ti	Cu	Ni	B	N
														Invention steel 1	0.405	0.247	1.290	0.0150	0.0020	0.033	0.147	0.148	0.038	0.103	0.306	0.0026	0.0040
Invention steel 2	0.405	0.255	1.300	0.0170	0.0022	0.031	0.147	0.147	0.040	0.106	0.870	0.0026	0.0036
														Invention steel 3	0.420	0.094	1.330	0.0100	0.0020	0.028	0.200	0.151	0.030	0.300	0.155	0.0021	0.0039
Invention steel 4	0.427	0.093	1.310	0.0095	0.0022	0.0333	0.199	0.149	0.030	0.299	0.310	0.0021	0.0044
														Invention steel 5	0.427	0.095	1.000	0.0096	0.0020	0.028	0.197	0.101	0.030	0.095	0.710	0.002	0.0036
Invention steel 6	0.420	0.095	1.000	0.0090	0.0018	0.022	0.198	0.102	0.028	0.096	0.924	0.0018	0.0032
														Invention steel 7	0.420	0.091	1.010	0.0100	0.0015	0.033	0.198	0.100	0.030	0.710	0.100	0.0021	0.0035
Invention steel 8	0.425	0.092	1.030	0.0100	0.0017	0.031	0.201	0.104	0.032	0.916	0.098	0.0021	0.0042
														Invention steel 9	0.416	0.089	1.010	0.0095	0.0017	0.022	0.198	0.100	0.001	0.105	0.905	0.0019	0.0033
Invention steel 10	0.425	0.092	1.020	0.0090	0.0021	0.033	0.197	0.102	0.031	0.101	0.923	0.0003	0.0044
														Invention steel 11	0.423	0.091	1.320	0.0095	0.002	0.033	0.200	0.149	0.030	0.299	0.910	0.0020	0.0037
Invention steel 12	0.412	0.092	1.310	0.0090	0.0026	0.025	0.199	0.150	0.029	0.300	0.903	0.0021	0.0043
														Invention steel 13	0.412	0.092	1.000	0.0095	0.0020	0.032	0.196	0.147	0.029	0.293	0.901	0.0021	0.0043
Inventive Steel 14	0.544	0.093	0.909	0.0090	0.0019	0.026	0.200	0.100	0.030	0.101	0.915	0.0019	0.0036
														Comparative Steel 1	0.402	0.098	1.300	0.0090	0.0022	0.030	0.200	0.148	0.029	0.000	0.000	0.0019	0.0053
Comparative Steel 2	0.450	0.360	0.809	0.0090	0.0019	0.031	0.195	0.001	0.030	0.300	0.310	0.0019	0.0041
														Comparative Steel 3	0.430	0.632	0.535	0.0110	0.0020	0.030	0.160	0.160	0.030	0.110	0.517	0.0022	0.0042
Comparative Steel 4	0.412	0.108	1.320	0.0095	0.0020	0.024	0.203	0.149	0.030	0.200	0.100	0.0021	0.0055
														Comparative Steel 5	0.410	0.260	1.340	0.0100	0.0023	0.007	0.15	0.153	0.042	0.110	0.103	0.0026	0.0047
Comparative Steel 6	0.420	0.095	1.320	0.0090	0.0020	0.025	0.199	0.150	0.029	0.001	0.000	0.0020	0.0034
														Comparative Steel 7	0.438	0.099	1.310	0.0100	0.0020	0.030	0.199	0.149	0.029	0.002	0.001	0.0020	0.0041

[ Table 2]

Steel grade	Relational expression (1) (Mn/Si)	Relation (2) (Cu + Ni)/(C + Mn)	Relational expression (3) (Ni/Si)
				Invention steel 1	5.2	0.24	1.24
Invention steel 2	5.1	0.57	3.41
				Invention steel 3	14.1	0.26	1.65
Invention steel 4	14.1	0.35	3.33
				Invention steel 5	10.5	0.56	7.47
Invention steel 6	10.5	0.72	9.73
				Invention steel 7	11.1	0.57	1.10
Invention steel 8	11.2	0.70	1.07
				Invention steel 9	11.3	0.71	10.17
Invention steel 10	11.1	0.71	10.03
				Invention steel 11	14.5	0.69	10.00
Invention steel 12	14.2	0.70	9.82
				Invention steel 13	10.9	0.85	9.79
Inventive Steel 14	9.8	0.70	9.84
				Comparative Steel 1	13.3	0.00	0.00
Comparative Steel 2	2.2	0.48	0.86
				Comparative Steel 3	0.8	0.65	0.82
Comparative Steel 4	12.2	0.17	0.93
				Comparative Steel 5	5.2	0.12	0.40
Comparative Steel 6	13.9	0.00	0.00
				Comparative Steel 7	13.2	0.00	0.01

[ Table 3]

[ Table 4]

Steel grade	Test piece numbering	Heating temperature (. degree.C.)	Cooling Rate (. degree. C./second)	Quenching crack	Tempering temperature (. degree.C.)
						Invention steel 1	Inventive Material 1	930	25	O → X (& gt 15 hours)	200
Invention steel 2	Inventive Material 2	930	25	X	200
						Invention steel 3	Invention Material 3	930	25	O → X (& gt 15 hours)	200
Invention steel 4	Inventive Material 4	930	20	O → X (& gt 15 hours)	200
						Invention steel 5	Inventive Material 5	930	50	X	220
Invention steel 6	Inventive Material 6	950	25	X	220
						Invention steel 7	Inventive Material 7	930	25	X	200
Invention steel 8	Inventive Material 8	900	25	X	220
						Invention steel 9	Inventive Material 9	930	20	X	220
Invention steel 10	Inventive Material 10	930	20	X	200
						Invention steel 11	Inventive Material 11	930	20	X	200
Invention steel 12	Inventive Material 12	930	20	X	200
						Invention steel 13	Invention Material 13	900	20	X	200
Invention steel 13	Inventive Material 14	950	50	O → X (& gt 15 hours)	-
						Inventive Steel 14	Inventive Material 15	930	20	X	200
Comparative Steel 1	Comparative Material 1	930	20	O	200
						Comparative Steel 1	Comparative Material 2	930	20	O	250
Comparative Steel 2	Comparative Material 3	930	20	O	200
						Comparative Steel 3	Comparative Material 4	930	25	O	200
Comparative Steel 4	Comparative Material 5	930	20	O	200
						Comparative Steel 5	Comparative Material 6	930	20	O	200
Comparative Steel 6	Comparative Material 7	930	20	O	200
						Comparative Steel 7	Comparative Material 8	930	20	O	200

[ Table 5]

[ Table 6]

As shown in tables 1 to 6, it is understood that the inventive materials (1 to 15) manufactured using the inventive steels (1 to 14) satisfying the relational expressions (1) to (3) do not generate quench cracks, or normal fracture (at the time of tensile test) occurs without abnormal fracture even after being held for a short time after quenching. On the other hand, comparative materials (1 to 8) manufactured using comparative steels (1 to 7) that do not satisfy at least one of relational expressions (1) to (3) generate quench cracks, or normal fracture occurs only after being held for a long time after the quenching heat treatment. Here, the abnormal fracture refers to a pre-failure (pre-fracture) in which the value of the total elongation in the stress-strain rate curve is very low in the tensile test.

Further, it was found that the inventive materials (1 to 15) exhibited yield strengths of 1400-.

Further, it is found that decarburized layers having a relatively shallow depth are generated in the inventive materials (1 to 15) as compared with the comparative materials (1 to 8).

As shown in fig. 1, it can be seen that the inventive materials (4, 6, 15) showed normal fracture, but the comparative material (3) showed premature fracture. That is, the comparative material (3) was broken before showing the maximum tensile stress value, and the elongation value was very low.

Further, as shown in FIGS. 2 and 3, it is known that the contents of copper and nickel are relatively higher than those of the concentrated layer in the interior of the steel sheet and the enrichment of nickel element is relatively high in the surface layer of the hot rolled steel sheets of the inventive materials (4) and (12).

As shown in fig. 4, it can be known that the drawn pipe before the quenching-tempering heat treatment [ fig. 4 (a) ] is composed of ferrite and pearlite phases, but the drawn pipe after the quenching-tempering heat treatment [ fig. 4 (b) ] has a typical tempered martensite phase.