CN111492085B

CN111492085B - High-strength steel material for polar environment having excellent fracture resistance at low temperature and method for producing same

Info

Publication number: CN111492085B
Application number: CN201880081799.8A
Authority: CN
Inventors: 严庆根; 李学哲; 金佑谦
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2017-12-22
Filing date: 2018-12-18
Publication date: 2021-10-29
Anticipated expiration: 2038-12-18
Also published as: EP3730659A1; JP2021507989A; JP7045459B2; WO2019124945A1; EP3730659A4; KR20190076758A; CN111492085A; KR102045641B1

Abstract

The invention provides a high-strength steel for polar environment with excellent fracture resistance at low temperature and a manufacturing method thereof. The high-strength steel material having excellent fracture resistance at low temperatures according to the present invention comprises, in wt%: 0.005-0.07%, Si: 0.005-0.3%, Mn: 1.7-3.0%, Sol.Al: 0.001-0.035%, Nb: 0.02% or less (except 0%), V: 0.01% or less (except 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01-2.0%, Cr: 0.01 to 0.5%, Mo: 0.001-0.5%, Ca: 0.0002 to 0.005%, N: 0.001-0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% or less (excluding 0%) and the balance of Fe and other inevitable impurities, and satisfies the following relational expressions 1 and 2, and the microstructure includes polygonal ferrite and acicular ferrite of 70 area% or more in total, and includes 3.5 area% or less of MA phase (martensite-austenite composite phase).

Description

High-strength steel material for polar environment having excellent fracture resistance at low temperature and method for producing same

Technical Field

The present invention relates to a high-strength steel material for polar environments excellent in fracture resistance at low temperatures, which is preferably used as a steel material for shipbuilding and offshore structures, and a method for producing the same.

Background

As global warming progresses, the ice layer in arctic regions decreases, and thus there is an increasing interest in arctic airlines connecting europe and east asia. In recent years, the test sailing of cargo ships is limited to summer. It is reported that time and cost can be reduced by more than 30% at most, compared to existing airlines via southeast asia. Further, if the ice layer in the arctic region completely disappears within 20-30 years, the straight line passing through the arctic is expected to be opened. Therefore, the necessity of ships crossing the arctic region is becoming more and more realistic, and the design of ships safe in such arctic environment and the demand for steel materials for the arctic environment are increasing.

Since conventional structural steel materials are likely to break in polar environments, that is, environments exposed to low temperatures of up to-60 degrees and impacts due to ice floes and the like, there is a need for a high-strength steel material for polar environments having excellent fracture resistance at low temperatures, which can solve the above-mentioned problems.

In general, the reason why high-strength steel materials with relatively thick thickness for large ships or oil platforms are easily broken at low temperature is as follows: in order to secure the strength of the high-strength super-thick steel, a large amount of alloying elements such as Mn, Mo, etc. can be added, and when the super-thick steel is manufactured, coarse granular bainite or hard phase structures such as M-a are easily formed due to low rolling reduction and slow accelerated cooling rate. Due to this microstructure, the steel material has extremely weak fracture resistance at low temperatures. Therefore, in order to obtain a super-thick material having high strength and excellent fracture characteristics at low temperatures, it is necessary to refine the structure and to greatly reduce the hard structure such as granular bainite or M-a.

In order to solve the above problems, the following methods are used: firstly, reducing the reheating temperature of a plate blank to be extremely low, and performing controlled rolling at low temperature to refine the structure; or ② adding more than 1% Cu, tempering at low temperature to improve the strength by fine Cu precipitate; or thirdly, a large amount of Ni is added to improve the low-temperature toughness of hard phase granular bainite and the like; or fourthly, promoting elements such as C and the like are reduced as much as possible so as to reduce the M-A structure to a minimum. However, as structures such as ships become larger and the use environment thereof becomes polar, the conventional method simply applied cannot sufficiently ensure fracture initiation and propagation resistance at low temperatures.

Therefore, it is required to develop a high-strength steel material with improved fracture initiation and propagation resistance at low temperatures and a method for manufacturing the same.

[ Prior Art document ]

(reference 1) Korean laid-open patent publication No. 2002-0028203

Disclosure of Invention

Technical problem

The present invention is made to solve the above-described problems of the prior art, and an object of the present invention is to provide a high-strength steel material excellent in fracture initiation and propagation resistance at low temperatures, and a method for producing the same.

In addition, the technical problem of the present invention is not limited to the above. Technical problems of the present invention can be understood through the entire contents of the present specification, and other technical problems of the present invention will not be understood with any difficulty to those of ordinary skill in the art.

Technical scheme

To achieve the above object, the present invention provides a high strength steel material excellent in fracture resistance at low temperature, comprising, in wt%: 0.005-0.07%, Si: 0.005-0.3%, Mn: 1.7-3.0%, Sol.Al: 0.001-0.035%, Nb: 0.02% or less (except 0%), V: 0.01% or less (except 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01-2.0%, Cr: 0.01 to 0.5%, Mo: 0.001-0.5%, Ca: 0.0002 to 0.005%, N: 0.001-0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), 0: 0.003% or less (excluding 0%) and the balance of Fe and other inevitable impurities, and satisfies the following relational expressions 1 and 2,

the microstructure includes polygonal ferrite and acicular ferrite of 70 area% or more in total, and includes 3.5 area% or less of MA phase (martensite-austenite composite phase).

[ relational expression 1]

Mn+0.5x(Ni+Cu)≥2.5wt％

[ relational expression 2]

Relation 2: mo + Cr +1.5xSi +10xNb is less than or equal to 0.5wt percent

In the relational expressions 1 and 2, each element is a value expressed in weight%.

In another aspect, the present invention provides a method for producing a high-strength steel material having excellent fracture resistance at low temperatures, including the steps of:

preparing a steel billet satisfying the alloy components;

heating the steel billet to 1000-1200 ℃;

finish hot rolling the heated slab at a temperature of 650 ℃ or higher so that the total rolling reduction in the non-recrystallization zone temperature zone becomes 30% or higher (except for the rolling reduction in the recrystallization zone); and

and cooling the hot-rolled steel sheet after the hot finish rolling to a cooling finish temperature of 200-550 ℃ at a cooling rate of 2-30 ℃/sec.

Effects of the invention

According to the invention, a steel material with remarkably improved fracture initiation and expansion resistance at low temperature can be effectively provided.

Drawings

FIG. 1 is a graph showing the Kca values measured for the steel material of invention example 1 in this example.

FIG. 2 is a photograph showing the microstructure of the steel material of example 3 of the present invention.

Detailed Description

Hereinafter, preferred embodiments of the present invention will be described. However, the embodiments of the present invention may be modified in various different ways, and the scope of the present invention is not limited to the embodiments described below. In addition, embodiments of the present invention are provided to more fully describe the present invention to those of ordinary skill in the art.

As a result of repeated studies and experiments to further improve fracture initiation and propagation resistance at low temperatures, the present inventors have found that it is necessary to minimize the amount of alloying elements, particularly Carbide (Carbide) -forming alloying elements such as C, Mo, Cr, and Nb, while maximizing the amount of alloying elements that improve the strength and toughness of the ferritic matrix, by precise control. The present inventors have also found that by controlling in this way, the microstructure of a steel material can be made to include polygonal ferrite and acicular ferrite of 70 area% or more in total, and include MA phase (martensite-austenite composite phase) of 3.5 area% or less, whereby fracture initiation and propagation resistance at low temperatures can be significantly improved, and thus have made the present invention.

That is, the steel material excellent in fracture resistance at low temperatures according to the present invention comprises, in wt%: 0.005-0.07%, Si: 0.005-0.3%, Mn: 1.7-3.0%, Sol.Al: 0.001-0.035%, Nb: 0.02% or less (except 0%), V: 0.01% or less (except 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01-2.0%, Cr: 0.01 to 0.5%, Mo: 0.001-0.5%, Ca: 0.0002 to 0.005%, N: 0.001-0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), 0: 0.003% or less (excluding 0%) and the balance of Fe and other inevitable impurities, and satisfies the relational expressions 1 and 2. In addition, the steel microstructure includes polygonal ferrite and acicular ferrite of 70 area% or more in total, and includes an MA phase (martensite-austenite composite phase) of 3.5 area% or less.

Next, the alloy composition of the steel of the present invention and the reason for limiting the content thereof will be described in detail. Hereinafter, the content unit of each element is weight%.

C：0.01～0.07％

C promotes the formation of acicular ferrite or lath (lath) bainite, and forms cementite, pearlite, and the like, and is an important element for securing strength. When the content of C is less than 0.01%, since C hardly diffuses, transformation occurs relatively quickly, and thus transformation into a coarse ferrite structure occurs, and there is a problem that the strength and toughness of the steel may be greatly reduced. On the contrary, when the content of C exceeds 0.07%, not only cementite or MA phase is excessively formed but also coarse, and the fracture initiation resistance at low temperature may be greatly reduced. Therefore, the content of C is preferably 0.01 to 0.07%. The content of C is more preferably 0.01 to 0.06%, and still more preferably 0.01 to 0.05%.

Si：0.005～0.3％

Si is an element generally added for decarburization, desulfurization, and solid solution strengthening. Although the effect of improving the yield strength and tensile strength is not significant, the stability of austenite in the weld heat affected zone is greatly improved, so that the fraction of MA phase is increased, resulting in the possibility that the fracture initiation resistance at low temperature may be greatly reduced. Therefore, in the present invention, the content is preferably limited to 0.3% or less. In addition, in order to control the Si content to less than 0.005%, the treatment time in the steel making process is greatly increased, which leads to problems of increased production cost and decreased productivity, so the lower limit of the Si content is preferably 0.005%. Therefore, the content of Si is preferably 0.005 to 0.3%. The content of Si is more preferably 0.005 to 0.25%, and still more preferably 0.005 to 0.2%.

Mn：1.7～3.0％

Mn has a significant effect of improving strength by solid solution strengthening, and does not cause a significant decrease in toughness at low temperatures, and therefore is added in an amount of 1.7% or more in order to ensure sufficient high strength. However, when Mn is excessively added, segregation increases in the central portion in the thickness direction of the steel sheet, and formation of a nonmetallic inclusion MnS with the segregated S is promoted. Since the MnS inclusion formed at the central portion is stretched due to the subsequent rolling and the segregated portion has high hardenability, a high-hardness low-temperature structure is easily formed, and finally fracture initiation and expansion resistance at low temperature are greatly reduced, the upper limit of the Mn content is preferably 3.0%. Therefore, the content of Mn is preferably 1.7 to 3.0%. The Mn content is more preferably 1.7-2.8%.

Sol.Al：0.005～0.035％

Al is used as a strong deoxidizer together with Si and Mn in a steelmaking process, and the effect can be fully obtained only by adding more than 0.005 percent when the deoxidizer is used singly or compositely. However, when the content of sol.al exceeds 0.035%, the above effect is saturated and Al in oxide inclusions generated as deoxidation products is saturated₂O₃The fraction (c) is too large, the size of inclusions becomes coarse, and the inclusions are not easily removed during refining, resulting in a significant decrease in the low-temperature toughness of the steel. In addition, similar to Si, the formation of MA phase is promoted in the welding heat affected zone, possibly causing a great reduction in fracture initiation and propagation resistance at low temperatures. Therefore, the content of the sol.al is preferably 0.005% to 0.035%. The content of the sol.al is more preferably 0.005% to 0.03%, and still more preferably 0.005% to 0.02%.

Nb: below 0.02% (except 0%)

Nb has a significant effect in that Nb is dissolved in austenite in solid solution at slab reheating to improve hardenability of austenite, and is precipitated as fine carbides (Nb, Ti) (C, N) at hot rolling to suppress recrystallization during rolling or cooling, thereby refining the final microstructure. However, when Nb is excessively added, the hardenability of the weld heat-affected zone will be excessively increased to promote the formation of MA phase, resulting in a great decrease in fracture initiation and propagation resistance at low temperatures. Therefore, in the present invention, the content of Nb is limited to 0.02% or less (excluding 0%). The content of Nb is preferably 0.015% or less, and more preferably 0.012% or less.

V: below 0.01% (except 0%)

Most of V re-dissolves when the slab is reheated, mostly dissolves during cooling after rolling to improve strength, but dissolves at high temperature in the weld heat affected zone to greatly improve hardenability, thereby promoting the formation of MA phase. Therefore, in the present invention, the content of V is limited to 0.01% or less (except for 0%). The content of V is preferably 0.008% or less, and more preferably 0.005% or less.

Ti：0.001～0.02％

Ti exists mainly as a hexahedral precipitate in the form of fine TiN at high temperature, or when added together with Nb, forms a (Ti, Nb) (C, N) precipitate, thereby having an effect of suppressing the grain growth of the substrate and the weld heat affected zone. In order to sufficiently secure the above effect, it is preferable to add 0.001% or more of Ti, and in order to maximize the effect, it is preferable to increase the addition amount according to the content of N added. On the contrary, when the content of Ti exceeds 0.02%, too coarse carbonitride is formed as a crack initiation point, thereby greatly lowering the impact characteristics of the weld heat affected zone. Therefore, the content of Ti is preferably 0.001 to 0.02%. The content of Ti is more preferably 0.001-0.017%, and still more preferably 0.001-0.015%.

Cu：0.01～1.0％

The strength of Cu is greatly improved through solid solution and precipitation, and fracture initiation and expansion resistance are not influenced obviously. When the content of Cu is less than 0.01%, the above effects are insufficient. In contrast, when the content of Cu exceeds 1.0%, cracks are caused on the surface of the steel sheet, and Cu is an expensive element, there is a problem that the cost increases. Therefore, the content of Cu is preferably 0.01 to 1.0%. The Cu content is more preferably 0.01 to 0.6%, and still more preferably 0.01 to 0.4%.

Ni：0.01～2.0％

Ni has little strength-increasing effect, but has an effect in improving fracture initiation and propagation resistance at low temperatures. In particular, when Cu is added, it has an effect of suppressing surface cracks caused by selective oxidation generated when the slab is reheated. In addition, even if a coarse hard structure is generated in the weld heat affected zone due to high temperature and rapid cooling rate, the toughness at low temperature can be improved by adding Ni. When the content of Ni is less than 0.01%, the above effects are insufficient. In contrast, when the content of Ni exceeds 2.0%, since Ni is an expensive element, a problem of cost increase will be caused. Therefore, the content of Ni is preferably 0.01 to 2.0%. The Ni content is preferably 0.2 to 1.8%, and more preferably 0.3 to 1.2%.

Cr：0.01～0.5％

Cr is not much effective in increasing yield strength and tensile strength by solid solution, but has high hardenability, so that the thick material also forms a microstructure at a slow cooling rate, thereby having an effect of improving strength and toughness. When the content of Cr is less than 0.01%, the above effect is insufficient. In contrast, when the content of Cr exceeds 0.5%, not only the cost increases, but also the low-temperature toughness of the welding heat affected zone becomes poor. Therefore, the content of Cr is preferably 0.01 to 0.5%. The content of Cr is preferably 0.01 to 0.4%, more preferably 0.01 to 0.25%.

Mo：0.01～0.65％

Mo finally greatly improves the strength by delaying phase change in the accelerated cooling process, and has the effect of preventing the toughness from being reduced due to grain boundary segregation of impurities such as P and the like. When the content of Mo is less than 0.01%, the above effects are insufficient. In contrast, when the content of Mo exceeds 0.65%, the formation of MA phase is promoted in the welding heat affected zone due to high hardenability, possibly resulting in a great decrease in fracture initiation and propagation resistance at low temperatures. Therefore, the content of Mo is preferably 0.01 to 0.65%. The content of Mo is more preferably 0.01 to 0.5%, and still more preferably 0.01 to 0.4%.

Ca：0.0002～0.005％

When Ca is added to molten steel during steel making after Al deoxidation, Ca bonds with S mainly present as MnS to suppress formation of MnS and forms spherical CaS, thereby having an effect of suppressing generation of cracks in the central portion of the steel. Therefore, in the present invention, Ca needs to be added in an amount of 0.0002% or more in order to sufficiently form the added S into CaS. However, when Ca is excessively added, the excess Ca bonds with 0 to form coarse hard oxide inclusions, which are then stretched and broken in the subsequent rolling, thereby forming crack initiation sites at low temperatures. Therefore, the upper limit of the content of Ca is preferably 0.005%. Therefore, the content of Ca is preferably 0.0002 to 0.005%. The content of Ca is more preferably 0.0005 to 0.003%, and still more preferably 0.0005 to 0.0025%.

N：0.001～0.006％

N forms precipitates with the added Nb, Ti, and Al, making the crystal grains of the steel fine, thereby improving the strength and toughness of the base material. However, when N is excessively added, it is present in a residual atomic state, thereby causing an aging phenomenon after cold deformation to lower the low-temperature toughness, and is the most representative element in this respect. In addition, it is known that when a slab is manufactured by a continuous casting process, N promotes surface cracking due to embrittlement at high temperatures. Therefore, in the present invention, considering that the content of Ti is 0.001% to 0.02%, the range of the addition amount of N is limited to 0.001% to 0.006%. The content of N is more preferably 0.001 to 0.005%, and still more preferably 0.001 to 0.0045%.

P: below 0.02% (except 0%)

P acts to improve strength, but causes deterioration of low-temperature toughness. In particular, heat-treated steel has a problem that low-temperature toughness is greatly reduced due to grain boundary segregation. Therefore, it is preferable that the content of P is controlled as low as possible. However, in the steel making process, excessive removal of P requires a large cost, and thus the content of P is limited to 0.02% or less. The content of P is preferably 0.015% or less, and more preferably 0.012% or less.

S: below 0.003% (except 0%)

S is a main cause of deterioration in low-temperature toughness, and forms MnS inclusions mainly in the center portion in the thickness direction of the steel sheet in combination with Mn. Therefore, in order to ensure strain aging impact characteristics at low temperatures, it is preferable to remove S as much as possible in the steel making process. However, it may require an excessively high cost, and thus is limited to 0.003% or less. The content of S is preferably 0.002% or less, and more preferably 0.0015% or less.

0: below 0.003% (except 0%)

0 is to remove oxide inclusions by adding deoxidizer such as Si, Mn, Al, etc. in the steel-making process. When the amount of addition of the deoxidizer and the inclusion removal process are insufficient, the amount of oxide inclusions remaining in molten steel increases and the size of inclusions also significantly increases. The coarse oxide inclusions which are not removed in this way remain inside in the form of cracks or spheres in the rolling process of the steel manufacturing process, and become fracture initiation points or crack propagation paths at low temperatures. Therefore, in order to ensure impact characteristics at low temperatures and CTOD (crack tip opening displacement) performance, it is necessary to suppress coarse oxide inclusions as much as possible. For this reason, the content of 0 is limited to 0.003% or less. The content of 0 is preferably 0.0025% or less, more preferably 0.0022% or less.

The balance of the present invention is iron (Fe). However, the conventional manufacturing process inevitably involves mixing of unexpected impurities from raw materials or the surrounding environment, and thus the mixing of impurities cannot be excluded. For example, boron (B) may be contained in an amount of 5ppm or less. These impurities are known to anyone skilled in the art of conventional manufacturing processes and therefore all relevant matters are not described in this specification.

In addition, for the alloy composition of the present invention, not only the above-described contents of the respective elements are required to be satisfied, but also Mn, Ni, Cu, Cr and Nb are required to satisfy the following relational expressions 1 to 2.

[ relational expression 1]

Mn+0.5x(Ni+Cu)≥2.5wt％

[ relational expression 2]

Relation 2: mo + Cr +1.5xSi +10xNb is less than or equal to 0.5wt percent

Mn, Ni and Cu of the above-described relational expression 1 are representative face-centered cubic metals, and when added to a steel material, not only improve strength by solid solution strengthening, but also do not significantly reduce toughness even at low temperature. The present inventors designed relational expression 1 in consideration of the degree of influence of the elements on the strength and toughness of the steel, and as the value of relational expression 1 increases, the solid solution strengthening effect increases, and finally the strength of the steel and the weld heat affected zone increases. Therefore, in order to sufficiently obtain the strength, the value of the relational expression 1 is preferably controlled to be 2.5 or more.

The relation 2 is designed in consideration of the influence degree of elements that promote the formation of the MA phase, which is a representative structure that significantly reduces the toughness of the steel and the weld heat affected zone, the fraction of the MA phase greatly increases as the value of the relation 2 increases, and the low temperature impact characteristic, i.e., the ductile-brittle transition temperature, of the final steel material increases. That is, as the value of relation 2 increases, the low temperature toughness tends to decrease. Therefore, in order to sufficiently ensure the low-temperature impact characteristics, particularly the CTOD value, of the steel material, the value of the relational expression 2 is preferably controlled to 0.5 or less. In the welded portion, particularly, the Sub-critical regenerated Heat Affected Zone (SC-HAZ) which is an important position for securing the low-temperature CTOD value, since the temperature at the time of welding is equal to or lower than the temperature of the two-phase region, the welded portion has a microstructure almost similar to the microstructure of the base material, and therefore, the value of the relational expression 2 is controlled to 0.5 or lower, whereby the low-temperature impact characteristics and the CTOD value of the welded portion can be sufficiently secured. The value of the relational expression 2 is more preferably 0.48, and still more preferably 0.45 or less.

The microstructure of the steel material of the present invention includes polygonal ferrite and acicular ferrite in a total amount of 70 area% or more, and also includes an MA phase (martensite-austenite composite phase) in an area% or less of 3.5 area%.

Acicular ferrite not only improves strength due to the effect of fine grain size, but also is the most important and fundamental microstructure to hinder crack propagation generated at low temperature. Polygonal ferrite is coarse as compared with acicular ferrite and thus contributes relatively less to the improvement of strength, but has a great contribution to the suppression of expansion at low temperatures due to the low dislocation density and high angle grain boundaries.

When the total of the polygonal ferrite and the acicular ferrite is less than 70 area%, it is difficult to suppress the initiation and propagation of cracks at low temperatures, and there is a problem in that it is difficult to secure high strength. Therefore, the total of polygonal ferrite and acicular ferrite is preferably 70 area% or more, more preferably 85 area% or more, and further preferably 90 area% or more.

Further, in the present invention, for the polygonal ferrite and the acicular ferrite, a crystal orientation difference between crystal grains is definedThe proportion of the large-angle grain boundaries of 15 DEG or more accounts for 40% or more of the entire grain boundaries, and the length of the large-angle grain boundaries per unit area is preferably 300mm/mm²The above.

Further, since the MA phase is high in hardness and does not deform, not only the deformation of the soft ferrite matrix around the MA phase is concentrated, but also, when the upper limit is exceeded, the interface with the surrounding ferrite matrix is separated, or the MA phase itself is broken to become a crack initiation origin. Therefore, the MA phase needs to be controlled as low as possible, preferably 3.5 area% or less, because this is the most important cause of deterioration of the low-temperature fracture characteristics of the steel material.

In the present invention, the average size of the MA phase measured as the equivalent circle diameter may be 2.5 μm or less. This is because, when the average size of the MA phase exceeds 2.5 μm, stress is more concentrated, so that the MA phase is likely to break to become a starting point of crack initiation.

In addition, in the present invention, the polygonal ferrite and the acicular ferrite may not be hot-roll hardened. That is, the polygonal ferrite and the acicular ferrite may be elongated without hot rolling, and the polygonal ferrite and the acicular ferrite may be formed after hot rolling.

As for the microstructure of the steel material of the present invention, bainitic ferrite, cementite, etc. may be included in addition to the polygonal ferrite, acicular ferrite, and MA phase.

Bainitic ferrite is a structure transformed at a low temperature, has many dislocations inside, but has a relatively coarse characteristic compared to various ferrites, and includes MA phase inside, thus having a high strength, but has a weak characteristic to crack initiation and propagation, and thus needs to be controlled to a minimum.

The number of the steel materials of the present invention may be 11/cm²The following range includes inclusions having a size of 10 μm or more. The dimension is a dimension measured as an equivalent circle diameter. When the number of inclusions with a size of 10 μm or more exceeds 11/cm²This may cause a problem of crack initiation at low temperature. In order to control the coarse inclusions, it is preferable to perform a secondary refining processAdding Ca or Ca alloy, and then carrying out bubbling and refluxing treatment for more than 3 minutes by using Ar gas.

In addition, the steel material of the present invention may have a yield strength of 460MPa or more, an impact energy value at-60 ℃ of 300J or more, and a CTOD value at-20 ℃ of 0.2mm or more. The tensile strength of the steel material of the present invention may be 570MPa or more. Further, the steel of the present invention may have a DBTT (ductile-brittle transition temperature) of-80 ℃ or lower.

Next, a method for producing a high-strength steel material excellent in fracture initiation and propagation resistance at low temperatures according to the present invention will be described.

The method for producing a steel material of the present invention comprises the steps of: preparing a steel billet satisfying the alloy components; heating the billet to 1000-1200 ℃; hot finish rolling the heated billet at a temperature of 650 ℃ or higher; and cooling the hot-rolled steel sheet after the hot finish rolling to a cooling finish temperature of 200 to 550 ℃ at a cooling rate of 2 to 30 ℃/sec.

Step of preparing billet

A steel slab satisfying the alloy composition described above was prepared.

In this case, in the present invention, the preparation of the billet is preferably carried out by a process including the steps of: adding Ca or Ca alloy into molten steel in the last step of secondary refining of the molten steel; after the addition of the Ca or Ca alloy, bubbling and refluxing treatment with Ar gas is carried out for at least 3 minutes or more. This is to control coarse inclusions.

Step of heating billet

And heating the steel billet to 1000-1200 ℃.

When the slab heating temperature is less than 1000 ℃, carbides and the like formed in the slab during continuous casting hardly re-dissolve, and the homogenization treatment of the segregation elements is insufficient. Therefore, it is preferable to heat the added Nb to 1000 ℃ or higher at which 50% or more of the added Nb can be re-dissolved.

On the contrary, when the heating temperature of the steel sheet exceeds 1200 ℃, the grain size of austenite becomes too coarse and is not sufficiently refined by the subsequent rolling, which may cause a great decrease in mechanical properties such as tensile strength and low-temperature toughness of the steel sheet.

The heating temperature of the billet is more preferably 1000-1160 ℃, and further preferably 1000-1140 ℃.

Step of Hot Rolling

And performing finish hot rolling on the heated slab at a bainite formation start temperature of 650 ℃ or higher to obtain a hot-rolled steel sheet.

When the finish hot rolling temperature is less than 650 ℃, coarse bainite is formed and work hardened during rolling, resulting in an excessive increase in strength and a great decrease in impact toughness at low temperatures, and therefore the finish hot rolling temperature is preferably controlled to 650 ℃ or more. That is, when the hot rolling temperature is low, coarse pro-eutectoid ferrite is generated before the hot rolling is completed, and the pro-eutectoid ferrite is elongated to form work hardening in the subsequent rolling, and the remaining austenite remains in a band shape and is transformed into a structure having a high MA hardening phase density, thereby causing a decrease in low-temperature toughness.

In the present invention, it is preferable to perform the non-recrystallization zone temperature zone at a total rolling reduction of 30% or more (excluding the rolling reduction of the recrystallization zone) in order to accumulate sufficient strain energy in austenite so that polygonal and acicular ferrite favorable for low-temperature toughness is sufficiently generated at the time of subsequent transformation, while ensuring the proportion and density of large-angle grain boundaries.

The reduction ratio is preferably 40% or more, and more preferably 45% or more.

Step of Cooling

Next, in the present invention, the hot-rolled steel sheet after the hot finish rolling is cooled.

In this case, the hot-rolled steel sheet is preferably cooled to a cooling completion temperature of 200 to 550 ℃ at a cooling rate of 2 to 30 ℃/s. When the cooling rate is less than 2 ℃/s, coarse ferrite, pearlite and bainite transformation regions cannot be avoided due to too low a cooling rate, possibly resulting in deterioration of strength and low-temperature toughness. When it exceeds 30 ℃/s, strength may be increased due to the formation of granular bainite or martensite, but low-temperature toughness may be significantly deteriorated.

In addition, when the cooling completion temperature exceeds 550 ℃, the microstructure such as acicular ferrite is difficult to be generated, and coarse bainite or pearlite is highly likely to be generated. On the contrary, when it is less than 200 ℃, there is no adverse effect on the microstructure, but the time required for cooling is excessively long, so that there is a problem that the productivity is greatly lowered.

The cooling completion temperature is more preferably 200 to 500 ℃, and still more preferably 200 to 450 ℃.

In addition, the present invention may further include a tempering step of heating the cooled hot-rolled steel sheet to 450 to 650 ℃, holding the heated hot-rolled steel sheet for (1.3 × t +5) to (1.3 × t +200) minutes, and then cooling the hot-rolled steel sheet (wherein t is a value measured in mm as a thickness of the hot-rolled steel sheet), if necessary. This is because, when MA or martensite is excessively produced, MA or martensite is decomposed to eliminate the internal high dislocation density, and Nb or the like which is dissolved in solid solution (even a small amount) is precipitated as carbonitride, thereby further improving the yield strength and the low-temperature toughness.

However, when the heating temperature is less than 450 ℃, the ferrite matrix cannot be sufficiently softened, and embrittlement phenomenon due to P segregation or the like occurs, which may rather cause deterioration of toughness. On the contrary, when the heating temperature exceeds 650 ℃, recovery and growth of crystal grains occur rapidly, and when the temperature is higher, a part of the crystal grains are reversed to austenite, which rather causes a great decrease in yield strength, and low-temperature toughness may also be deteriorated.

Further, if the holding time is less than (1.3 × t +5) minutes, the uniformity of the structure is insufficient, and if it exceeds (1.3 × t +200) minutes, there is a problem that the productivity is lowered.

Modes for carrying out the invention

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

(examples)

The slabs having the compositions shown in table 1 below were heated, hot-rolled, and cooled under the conditions shown in table 2 below to manufacture steels.

The microstructure of the produced steel was observed, and the physical properties were measured and shown in Table 3.

The produced steel materials were welded, and the CTOD value (-20 ℃) of the weld heat affected zone (SCHAZ) was measured and shown in Table 3. Since the CTOD (-20 ℃ C.) of the steel material was higher than that of the weld heat affected zone, the CTOD (-20 ℃ C.) of the steel material was not separately measured.

In this case, the microstructure of the steel material is prepared by grinding a cross section of the produced steel material into a mirror surface, etching the mirror surface with Nital or LePera as necessary, measuring an image of a specific area of the sample by a factor of 100 to 5000 using an optical or scanning electron microscope, and measuring the fraction of each phase from the measured image using an image analyzer (image analyzer). In order to obtain statistically significant values, the same samples were repeatedly measured by changing the position, and the average value was determined.

In order to observe the characteristics of the produced structure in more detail, the grain boundary characteristics of the produced steel material were quantitatively measured by Electron Back Scattering Diffraction (EBSD) measurement using a scanning Electron microscope on a sample etched by Nital.

The physical properties of the steel are described after measurement from the engineering strain rate-engineering stress curve obtained by a conventional tensile test.

The impact energy value (-60 ℃ C.) of the weld heat affected zone was measured by performing Charpy V-notch impact test.

For the CTOD value (-20 ℃ C.), specimens having a dimension perpendicular to the rolling direction of B (thickness) xB (width) x5B (length) were processed according to the BS 7448 standard, and fatigue cracks having a length of about 50% of the specimen width were inserted, followed by performing a CTOD test at-20 ℃. Wherein B is the thickness of the steel material to be produced.

The Kca values were measured by the ESSO test method 3 times to obtain graphs of the crack growth stopping temperature and the K value measured in each test, and the Kca values were obtained from the K value at a temperature of-10 degrees (Kca: crack arrest K). Further, CAT (Crack arrest temperature) is a value obtained by measuring NDTT (nickel-ductility transition temperature) from the NRL test and calculating it from the conversion formula of relational expression 1. Wherein B represents the thickness of the steel material.

[ relational expression 1]

[ TABLE 1]

In table 1, the relation 1 is Mn +0.5x (Ni + Cu), and the relation 2 is Mo + Cr +1.5xSi +10 xNb.

[ TABLE 2]

[ TABLE 3 ]

Ferrite in table 3 means the sum of polygonal ferrite and acicular ferrite.

As shown in tables 1 to 3, in consideration of yield strength, tensile strength, impact energy value, Kca, CAT, and the like, inventive examples 1 to 4, which satisfy the alloy components and manufacturing conditions proposed in the present invention, have excellent fracture toughness at low temperatures, and it can be seen that the CTOD value of the welding heat affected zone is also high. In particular, as shown in FIG. 1, the Kca value measured in inventive example 1 greatly exceeded required value 8000. As shown in fig. 2, it can also be seen that such excellent strength and low temperature toughness characteristics are obtained from the fine polygonal and acicular ferrite structures which are sufficiently formed.

Comparative example 1 is the case where the C content exceeds the range of the present invention, and the added C is the most effective element for promoting the formation of granular bainite and MA. Therefore, when C is excessively added, the ferrite fraction which is advantageous in toughness is greatly reduced, and although the strength in the base material is high, the low-temperature toughness such as impact energy value is poor, and in particular, the CTOD value in the weld heat affected zone is greatly reduced.

Comparative example 2 is a case where the amount of Mn added exceeds the range of the present invention. At this time, since the Mn content is high, the possibility of segregation in the central portion of the steel material is greatly increased, so that the impact energy in the central portion in the thickness direction of the steel material is greatly reduced, and a hardened structure having a very high hardness is formed in the segregation zone of the central portion also in the welding heat affected zone, so that a premature fracture phenomenon (pop-in) occurs to cause the CTOD value to be greatly reduced.

In comparative example 3, the content of Nb, which is widely used for improving the strength and refining the structure, is more than the range of the present invention. Generally, the addition of Nb facilitates structure refinement to improve both strength and toughness. However, when excessively added, the formation of polygonal and acicular ferrite, which is advantageous in toughness, is suppressed, and the formation of a structure such as granular bainite is promoted. Therefore, the proportion and density of large-angle grain boundaries of 15 ° or more, which are advantageous for suppressing crack propagation, will be greatly reduced, resulting in relatively easy crack propagation. As a result, as shown in table 3, the Kca value measured in comparative example 3 was 5860, which is significantly smaller than required value 8000. In addition, the formation of M-A structure, which has an adverse effect on particularly low-temperature toughness, is greatly promoted in the weld heat-affected zone, with the result that CTOD is greatly reduced.

For comparative examples 4, 5 and 6, although each element content range satisfies the range of the present invention, the values of relational expressions 1 and 2 are out of the range of the present invention. It can be seen that the strength is low or the low temperature toughness is greatly reduced.

Specifically, with comparative example 4, while relational expression 1 composed of a component advantageous for improving low-temperature toughness is satisfied, relational expression 2 composed of a component impairing low-temperature toughness is out of the scope of the present invention. As a result, although the strength is sufficiently high, the impact energy value in the base material or the CTOD value in the weld heat affected zone is poor.

In addition, with comparative example 5, although relational expression 2 satisfied the scope of the invention, relational expression 1 was out of the scope of the invention, and the component addition amount thereof was insufficient to secure the overall strength of the steel material, thereby greatly decreasing the strength of the base material.

Comparative example 6 is a case where both of the relational expression 1 and the relational expression 2 are out of the scope of the invention. That is, the components advantageous to low-temperature toughness are insufficient, and the components disadvantageous to low-temperature toughness are excessive, so that all low-temperature toughness characteristic values are inferior.

With comparative example 7, although the composition of the steel material satisfied all the ranges of the invention, the total rolling reduction in the unrecrystallized region in the manufacturing process of the steel material did not reach the ranges of the invention. That is, since the reduction amount in the non-recrystallized region is insufficient, the fraction of ferrite that acts to inhibit crack propagation in the microstructure of the steel material is low, and the proportion and density of large-angle grain boundaries are greatly reduced, so that the low-temperature toughness characteristic value is not good.

In addition, comparative example 8 is also a steel whose composition satisfies all the scope of the invention, but in the manufacturing process of the steel, the accelerated cooling is not performed after controlled rolling, but cooling is performed by air cooling, because the cooling rate is slow, ferrite favorable to low temperature toughness is sufficiently generated, but coarsening causes a great reduction in strength.

The embodiments of the present invention have been described in detail above, but the scope of claims of the present invention is not limited to the above embodiments, and various modifications and changes can be made by those skilled in the art within the scope not exceeding the technical idea described in the claims.

Claims

1. A high-strength steel material excellent in fracture resistance at low temperatures, comprising, in% by weight, C: 0.005-0.07%, Si: 0.005-0.3%, Mn: 1.7-3.0%, Sol.Al: 0.001-0.035%, Nb: 0.02% or less and 0% or less, V: 0.01% or less and 0% or less except, Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01-2.0%, Cr: 0.01 to 0.5%, Mo: 0.001-0.5%, Ca: 0.0002 to 0.005%, N: 0.001-0.008%, P: 0.02% or less and 0% or less except, S: 0.003% or less and 0% or less except, O: 0.003% or less except 0%, and the balance of Fe and other inevitable impurities, and satisfies the following relational expressions 1 and 2,

the microstructure includes polygonal ferrite and acicular ferrite of 70 area% or more in total, and includes a martensite-austenite composite phase of 3.5 area% or less,

[ relational expression 1]

Mn+0.5x(Ni+Cu)≥2.5wt％

[ relational expression 2]

Relation 2: mo + Cr +1.5xSi +10xNb less than or equal to 0.5wt percent

2. The high-strength steel material excellent in fracture resistance at low temperatures according to claim 1,

the polygonal ferrite and the acicular ferrite are such that a large angle grain boundary proportion in which a crystal orientation difference between crystal grains is defined to be 15 DEG or more accounts for 40% or more of the entire grain boundary, and a large angle grain boundary length per unit area is 300mm/mm²The above.

3. The high-strength steel material excellent in fracture resistance at low temperatures according to claim 1,

the yield strength of the steel is more than 460MPa, the impact energy value at-60 ℃ is more than 250J, and the Kca value measured in an ESSO test is 8000N/mm^3/2Above or below-10 ℃ as calculated from NDTT measured in NRL test.

4. The high-strength steel material excellent in fracture resistance at low temperatures according to claim 1,

the steel has a tensile strength of 570MPa or more and a DBTT (ductile-brittle transition temperature) of-80 ℃ or less.

5. The high-strength steel material excellent in fracture resistance at low temperatures according to claim 1,

the steel is at 11 pieces/cm²The following range includes inclusions having a size of 10 μm or more as measured by the equivalent circle diameter.

6. A method for producing a high-strength steel material having excellent fracture resistance at low temperatures, comprising the steps of:

preparing a steel slab comprising, in weight percent, C: 0.005-0.07%, Si: 0.005-0.3%, Mn: 1.7-3.0%, Sol.Al: 0.001-0.035%, Nb: 0.02% or less and 0% or less, V: 0.01% or less and 0% or less except, Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01-2.0%, Cr: 0.01 to 0.5%, Mo: 0.001-0.5%, Ca: 0.0002 to 0.005%, N: 0.001-0.008%, P: 0.02% or less and 0% or less except, S: 0.003% or less and 0% or less except, O: 0.003% or less except 0%, and the balance of Fe and other inevitable impurities, and satisfies the following relational expressions 1 and 2;

heating the billet to 1000 ℃ to 1200 ℃;

finish hot rolling the heated slab at a temperature of 650 ℃ or higher so that the total rolling reduction in the temperature zone of the non-recrystallization zone is 30% or higher, excluding the rolling reduction in the recrystallization zone; and

cooling the hot-rolled steel sheet after the finish hot rolling at a cooling rate of 2 to 30 ℃/sec to a cooling completion temperature of 200 to 550 ℃,

[ relational expression 1]

Mn+0.5x(Ni+Cu)≥2.5wt％

[ relational expression 2]

Relation 2: mo + Cr +1.5xSi +10xNb less than or equal to 0.5wt percent

7. The method for producing a high-strength steel material excellent in fracture resistance at low temperatures according to claim 6, further comprising a tempering step,

the tempering step is to heat the cooled hot-rolled steel sheet to 450 to 650 ℃, hold the heated steel sheet for (1.3 × t +5) to (1.3 × t +200) minutes, and then cool the steel sheet.

8. The method of producing a high-strength steel material excellent in fracture resistance at low temperatures according to claim 6, characterized in that:

when preparing the steel billet, the process comprising the following steps is implemented:

adding Ca or Ca alloy into molten steel in the last step of the secondary refining process; after the addition of the Ca or Ca alloy, bubbling and refluxing treatment with Ar gas is carried out for at least 3 minutes or more.

9. The method of producing a high-strength steel material excellent in fracture resistance at low temperatures according to claim 6, characterized in that:

the microstructure of the cooled steel material includes polygonal ferrite and acicular ferrite of 70 area% or more in total, and includes a martensite-austenite composite phase of 3.5 area% or less.

10. The method for producing a high-strength steel material excellent in fracture resistance at low temperatures according to claim 9, characterized in that:

the polygonal ferrite and the acicular ferrite are such that a proportion of large-angle grain boundaries, in which a difference in crystal orientation between grains is defined to be 15 DEG or more, accounts for 40% or more of the entire grain boundaries, and the length of the large-angle grain boundaries per unit area is 300mm/mm²The above.