KR20190076758A - High strength steel for arctic environment having excellent resistance to fracture in low temperature, and method for manufacturing the same - Google Patents

Abstract

Provided are high strength steel for an arctic environment having excellent resistance to fracture at low temperature, and a manufacturing method thereof. The high strength steel for the arctic environment having excellent resistance to fracture at low temperature contains, by weight percentages: 0.005-0.07% of C; 0.005-0.3% of Si; 1.7-3.0% of Mn; 0.001-0.035% of Sol.Al; 0.02 or less percent of Nb (excluding 0%); 0.01 less percent of V (excluding 0%); 0.001-0.02% of Ti; 0.01-1.0% of Cu; 0.01-2.0% of Ni; 0.01-0.5% of Cr; 0.001-0.5% of Mo; 0.0002-0.005% of Ca; 0.001-0.008% of N; 0.02 less percent of P (excluding 0%); 0.003 less percent of S (excluding 0%); 0.003 or less percent of O (excluding 0%); and remnant Fe and other inevitable impurities. The comprising elements satisfy the following relational equations: Mn + 0.5 x (Ni + Cu) >= 2.5wt% and Mo + Cr + 1.5 x Si + 10 x Nb >= 0.5wt%, where each element symbol is a value showing the content of each element by weight percentage. The microstructure of the present invention contains 70 area percentage of polygonal ferrite and acicular ferrite in total, and contains 3.5 or less area percentage a martensite-austenite compound phase (MA phase).

Description

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a high strength steel for use in a polar environment having excellent resistance to breakage at low temperatures,

The present invention relates to a high-strength steel material for polar environment having excellent resistance to breakdown at low temperatures, which can be preferably applied to steel for shipbuilding and marine structure, and a manufacturing method thereof.

As global warming gradually reduces the Arctic ice, there is a growing interest in Arctic routes connecting Europe and East Asia. In recent years, the cargo ship has been piloted only in summer, and the time and cost are reduced by more than 30% more than the existing route through Southeast Asia. Furthermore, within 20 to 30 years, even if the Arctic ice completely disappeared, straight-line routes crossing the North Pole are expected. As a result, the need for vessels passing through the Arctic region is becoming more and more important. In this polar environment, the need for the design of safe vessels and the required steel for the polar environment is gradually increasing.

Existing structural steels are required to have high-strength steels for polar environment with excellent resistance to breakdown at low temperatures, which can overcome them in an extreme environment, that is, at temperatures as low as -60 ° C and exposed to shocks caused by drift ice. Do.

In general, the reason why high-strength steel with high thickness used for large ships or oil mining platforms is vulnerable to destruction at low temperatures is as follows. In order to secure the strength of the high strength superfine steel material, the addition amount of alloying elements such as Mn and Mo is inevitably large, and due to the low rolling reduction rate and the slow accelerated cooling rate during the production of the superfine steel material, And a hard phase texture such as MA is likely to be generated. Due to such microstructure, the steel has a characteristic that the resistance to fracture is very weak at low temperature. Therefore, it is necessary to miniaturize the structure and to reduce the hard tissues such as granulabainite and M-A extremely in order to have the high strength of the extreme post material and the fracture property at the excellent low temperature.

In order to solve the above-mentioned problems, (1) the structure is finely refined by controlling rolling at a low temperature by reducing the slab reheating temperature to an extreme level; (2) the strength is improved by fine Cu precipitates by tempering at low temperature by adding at least 1% In order to improve low-temperature toughness such as granular bainite, which is a hard phase, a large amount of Ni is added, and (4) minimization of promoting elements such as C is used in order to extremely reduce the M-A structure. However, as structures of ships and the like are gradually becoming larger and the use environment is changed to an polar environment, there is a problem that it is difficult to sufficiently secure breakdown initiation and propagation resistance at low temperatures by simply applying the above-described conventional method.

Therefore, it is required to develop a high-strength steel having improved fracture initiation and propagation resistance at a low temperature and a manufacturing method thereof.

Korean Patent Publication No. 2002-0028203

Accordingly, it is an object of the present invention to provide a high-strength steel having excellent fracture initiation and propagation resistance at low temperatures and a method of manufacturing the same.

On the other hand, the object of the present invention is not limited to the above description. It will be understood by those of ordinary skill in the art that there is no difficulty in understanding the additional problems of the present invention.

According to an aspect of the present invention,

0.001 to 0.07% of Si, 0.005 to 0.3% of Si, 1.7 to 3.0% of Mn, 0.001 to 0.035% of Sol.Al, 0.02% or less of Nb (excluding 0% 0.001 to 0.02% of Ti, 0.01 to 1.0% of Cu, 0.01 to 2.0% of Ni, 0.01 to 0.5% of Cr, 0.001 to 0.5% of Mo, 0.0002 to 0.005% of Ca, (Excluding 0%), O: not more than 0.003% (excluding 0%), the remainder Fe and unavoidable impurities are contained in an amount of N: 0.001 to 0.008%, P: 0.02% And satisfy the following relational expression 1 and relational expression 2,

The microstructure includes 70% by area or more of polygonal ferrite and needle-like ferrite in total, and has a high strength with excellent fracture toughness at low temperature including 3.5% by area or less of MA phase (martensite-austenite composite phase) Steel. .

[Relation 1]

Mn + 0.5 x (Ni + Cu)? 2.5 wt%

[Relation 2]

Relation 2: Mo + Cr + 1.5 x Si + 10 x Nb? 0.5 wt%

(In the above relational expression 1-2, each element is a value expressed in weight%).

Further, according to the present invention,

Providing a steel slab satisfying the alloy composition described above;

Heating the steel slab to a temperature of 1000 to 1200 캜;

Subjecting the heated slab to a finish hot rolling such that the total rolling reduction in the temperature range of not lower than 650 ° C is not less than 30% (excluding the recrystallization reverse rolling reduction rate) in the non-recrystallization reverse temperature range; And

And cooling the finished hot-rolled steel sheet at a cooling rate of 2 to 30 占 폚 / s to a cooling end temperature of 200 to 550 占 폚, and a method of manufacturing a high-strength steel material excellent in low- .

INDUSTRIAL APPLICABILITY According to the present invention, it is possible to effectively provide a steel material having remarkably improved fracture initiation and propagation resistance at low temperatures.

1 is a graph showing the Kca value of a steel material in Inventive Example 1 measured in this embodiment.
2 is a microstructure photograph of the steel material of Inventive Example 3 in this embodiment.

Hereinafter, preferred embodiments of the present invention will be described. However, the embodiments of the present invention can be modified into various other forms, and the scope of the present invention is not limited to the embodiments described below. Further, the embodiments of the present invention are provided to more fully explain the present invention to those skilled in the art.

As a result of researches and experiments conducted to further improve breakdown initiation and propagation resistance at low temperatures, the present inventors have found that the addition amount of alloying elements such as C, Mo, Cr, and Nb, which are alloying elements that generate carbide, It is necessary to precisely control the addition amount of the alloying element, which has the effect of simultaneously improving the strength and toughness of the ferrite base, in the direction of maximizing the increase. By controlling in this manner, the microstructure of the steel material can contain 70% or more of polygonal ferrite and acicular ferrite in total, and can contain 3.5% or less of the MA phase (martensite-austenite composite phase) , Thereby enabling breakthrough initiation and propagation resistance at low temperatures to be dramatically improved.

That is, according to the present invention, the steel material having excellent resistance against fracture at low temperature is 0.005 to 0.07% of C, 0.005 to 0.3% of Si, 1.7 to 3.0% of Mn, 0.001 to 0.035% of Sol.Al, , Ni: not more than 0.02% (excluding 0%), V: not more than 0.01% (excluding 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0% 0.001 to 0.5% of Mo, 0.0002 to 0.005% of Ca, 0.001 to 0.008% of N, 0.02% or less of P (excluding 0%), S of 0.003% or less 0.003% or less (excluding 0%), the remainder Fe and unavoidable impurities, and satisfies the above relational expression 1-2. The steel microstructure contains 70% by area or more of polygonal ferrite and acicular ferrite in total, and contains 3.5% by area or less of MA phase (martensite-austenite composite phase).

Hereinafter, 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 unit of each element content is% by weight.

C: 0.01 to 0.07%

C is an element that plays an important role in promoting the formation of needle-like ferrite or lath bainite and securing strength by generating cementite or pearlite.

If the C content is less than 0.01%, the diffusion of C is hardly occurred and the transformation occurs relatively quickly, so that the steel is transformed into a coarse ferrite structure and the strength and toughness of the steel may be deteriorated. On the other hand, when the C content is more than 0.07%, cementite or MA phase is not only excessively formed, but also has a problem in that it is formed to a great extent and can significantly deteriorate the fracture initiation resistance at low temperature.

Si: 0.005 to 0.3%

Si is an element that is generally added for the purpose of strengthening employment together with deoxidation and desulfurization effect. However, the effect of increasing the yield and tensile strength is negligible, while the stability of the austenite in the weld heat affected zone is greatly increased and the fraction of the MA phase is increased. In the invention, it is preferable to limit it to 0.3% or less. On the other hand, in order to control the Si content to less than 0.005%, the treatment time in the steelmaking process is greatly increased, resulting in an increase in production cost and a decrease in productivity. Therefore, the lower limit of the Si content is preferably 0.005%.

Mn: 1.7 to 3.0%

Mn has a large effect of increasing the strength by solid solution strengthening, and toughness reduction at low temperature is not large. Therefore, Mn is added in an amount of 1.7% or more to ensure a sufficient strength.

However, if Mn is added excessively, segregation becomes serious in the thickness direction center of the steel sheet, and at the same time, it promotes the formation of MnS, which is a non-metallic inclusion together with segregated S. The MnS inclusions generated in the center portion are stretched by the subsequent rolling, and the segregation site has a high hardening ability, so that the low-temperature structure of high hardness is easily generated, and as a result, the breakdown initiation and propagation resistance at low temperatures are greatly lowered, Is preferably 3.0%.

Therefore, the Mn content is preferably 1.7 to 3.0%.

Sol.Al: 0.005 to 0.035%

Sol.Al is used as a strong deoxidizer in the steelmaking process together with Si and Mn, and at least 0.005% should be added at the time of single or multiple deoxidation to obtain sufficient effect.

However, when the content of Sol.Al is more than 0.035%, the above-mentioned effect is saturated and the fraction of Al ₂ O ₃ in the oxidative inclusions produced as a result of deoxidation increases more than necessary, the size of the inclusions becomes large, There is a problem that the low temperature toughness of the steel material is greatly reduced. Also, similar to Si, the generation of the MA phase in the weld heat affected zone is promoted, and the breakdown initiation and propagation resistance at low temperatures can be greatly reduced.

Therefore, the content of Sol.Al is preferably 0.005 to 0.035%.

Nb: 0.02% or less (excluding 0%)

Nb is dissolved in the austenite during the reheating of the slab to increase the hardenability of the austenite and precipitates into fine carbonitrides (Nb, Ti) (C, N) during hot rolling to inhibit recrystallization during rolling and cooling, Is a very large element. However, when Nb is added in an excessively large amount, the hardenability of the weld heat affected zone is excessively increased to promote the generation of the MA phase, which significantly lowers destruction initiation and propagation resistance at low temperatures. Therefore, the Nb content in the present invention is 0.02% Excluding 0%).

V: 0.01% or less (excluding 0%)

V is almost completely re-heated at the time of reheating of the slab, and it is mostly precipitated during cooling after rolling to improve strength. In the heat affected zone of welding, it dissolves at high temperature to greatly increase hardenability, thereby promoting the formation of MA phase. Therefore, the V content in the present invention is limited to 0.01% or less (excluding 0%).

Ti: 0.001 to 0.02%

Ti has an effect of suppressing crystal grain growth of the base material and the weld heat affected zone by forming precipitates of (Ti, Nb) (C, N) precipitates mainly in the form of fine hexagonal TiN type precipitates at high temperatures or by adding them such as Nb .

In order to sufficiently secure the above-mentioned effect, it is preferable to add Ti in an amount of 0.001% or more, and in order to maximize the effect, it is preferable to increase it in accordance with the content of N added. On the other hand, when the Ti content is more than 0.02%, coarse carbonitride is produced more than necessary, which acts as a starting point of the fracture crack, which can greatly reduce the impact characteristics of the weld heat affected zone. Therefore, the Ti content is preferably 0.001 to 0.02%.

Cu: 0.01 to 1.0%

Cu is an element capable of significantly improving the strength by solubilization and precipitation without greatly deteriorating breakdown initiation and propagation resistance.

When the Cu content is less than 0.01%, the above-mentioned effect is insufficient. On the other hand, when the Cu content exceeds 1.0%, cracks may be generated on the surface of the steel sheet, and Cu is an expensive element, causing a problem of cost increase.

Ni: 0.01 to 2.0%

Ni has almost no effect of increasing the strength, but is effective in improving fracture initiation and propagation resistance at low temperatures. In particular, when Cu is added, it has an effect of suppressing surface cracking due to selective oxidation occurring at reheating of the slab. Also, the addition of Ni has the effect of improving the toughness at low temperature even if a coarse hard tissue is produced due to the high temperature and rapid cooling rate of the weld heat affected zone.

When the Ni content is less than 0.01%, the above-mentioned effect is insufficient. On the other hand, Ni is an expensive element, and when the content thereof exceeds 2.0%, there is a problem of cost increase.

Cr: 0.01 to 0.5%

Cr has a small effect of increasing the yield and tensile strength due to employment, but it has an effect of improving strength and toughness by allowing fine materials to be formed at a slow cooling rate of a post-material because of its high hardenability.

When the Cr content is less than 0.01%, the above-mentioned effect is insufficient. On the other hand, when the Cr content exceeds 0.5%, not only the cost increases, but also the low temperature toughness of the weld heat affected zone can be lowered.

Mo: 0.01 to 0.65%

Mo has the effect of delaying the phase transformation in the accelerated cooling process and consequently increasing the strength, and is an element having an effect of preventing the deterioration of toughness due to grain boundary segregation of impurities such as P and the like.

When the Mo content is less than 0.01%, the above-mentioned effect is insufficient. On the other hand, when the Mo content exceeds 0.65%, the generation of the MA phase in the weld heat affected zone is promoted due to the high hardenability, and the breakdown initiation and propagation resistance at low temperatures can be greatly reduced.

Ca: 0.0002 to 0.005%

When Ca is Al-deoxidized and added to molten steel during steelmaking, it is combined with S existing mainly in MnS, thereby suppressing MnS formation and forming spherical CaS, thereby suppressing cracks in the center of the steel. Therefore, Ca should be added in an amount of 0.0002% or more in order to sufficiently form added S in CaS.

However, when Ca is excessively added, excess Ca is combined with O to form a coarse hard, oxidative inclusion, which is then stretched and fractured at the subsequent rolling and acts as a crack initiation point at a low temperature. Therefore, the upper limit of the Ca content is preferably 0.005%.

N: 0.001 to 0.006%

N is an element that forms a precipitate together with added Nb, Ti and Al to improve the strength and toughness of the base material by refining the crystal grains of the steel. However, it is known as the most representative element to reduce the low-temperature toughness due to aging phenomenon after the cold deformation when it is present in excess atomic state in the excessive addition. It is also known that slabs produced by continuous casting promote surface cracking due to embrittlement at high temperatures.

Therefore, in the present invention, the addition amount of N is limited to the range of 0.001 to 0.006% considering that the Ti content is 0.001 to 0.02%.

P: 0.02% or less (excluding 0%)

P acts to increase the strength, but it is an element that lowers the low temperature toughness. Particularly, there is a problem that low-temperature toughness is largely deviated due to grain boundary segregation in the heat-treated steel. Therefore, it is preferable to control P as low as possible. However, excessively removing P from the steelmaking process is expensive, so it is limited to 0.02% or less.

S: 0.003% or less (excluding 0%)

S is a main cause of MnS inclusions mainly in the thickness direction center of the steel sheet by binding with Mn, thereby lowering the low temperature toughness. Therefore, it is desirable to remove S as much as possible in the steelmaking process in order to secure the deformation aging property at low temperature. However, it may be excessive cost, so it should be limited to less than 0.003%.

O: 0.003% or less (excluding 0%)

O is made into an oxidative inclusion by adding a deoxidizing agent such as Si, Mn, Al in the steel making process. If the amount of the deoxidizing agent and the process for removing inclusions are insufficient, the amount of the oxidative inclusions remaining in the molten steel increases, and the size of the inclusions increases greatly. The coarse oxidative inclusions which have not been removed in this way are then left in a crushed form or spherical form during the rolling process in the steel making process and serve as a starting point of fracture at low temperature or as propagation paths of cracks. Therefore, in order to secure impact characteristics and CTOD characteristics at low temperatures, it is necessary to suppress coarse oxidative inclusions as much as possible and limit the O content to 0.003% or less.

The remainder of the present invention is iron (Fe). However, it is not possible to exclude elements or impurities which are not intended from the raw material or the surrounding environment in a conventional manufacturing process, since they may be inevitably incorporated. For example, 5 ppm or less of boron (B) or the like. These impurities are not specifically mentioned in this specification, as they are known to any person skilled in the art of manufacturing.

Further, the alloy composition of the present invention is required to contain Mn, Ni, Cu, Cr, and Nb so as to satisfy not only the above-described respective element content but also the following relational expression 1-2.

[Relation 1]

Mn + 0.5 x (Ni + Cu)? 2.5 wt%

[Relation 2]

Relation 2: Mo + Cr + 1.5 x Si + 10 x Nb? 0.5 wt%

(In the above relational expression 1-2, each element is a value expressed in weight%).

Mn, Ni, and Cu constituting the above-mentioned relational expression 1 are typical face-centered cubic metals. These elements are elements that not only increase the strength by solid solution strengthening when added to steel materials, but also do not significantly affect toughness even at low temperatures. The inventors of the present invention designed Equation 1 in consideration of the influence of the elements on the steel strength and toughness. As the value of Equation 1 increases, the solid solution strengthening effect increases and the strength of the steel material and weld heat affected portion increases. Therefore, in order to obtain sufficient strength, it is preferable to control the value of the relational expression 1 to 2.5 or more.

(2) is a formula designed in consideration of the influence of an element promoting the formation of an MA phase, which is a representative structure that largely affects the toughness of the steel material and the weld heat affected zone. As the value of the relational expression 2 increases, the MA phase fraction increases greatly, The ductile-brittle transition temperature, which is a low-temperature impact characteristic, increases. That is, as the value of the relational expression 2 increases, the low temperature toughness tends to decrease. Therefore, it is preferable to control the value of the relational expression (2) to 0.5 or less in order to sufficiently secure the low temperature impact property of the steel material, particularly the CTOD value.

In addition, the SC-HAZ (Sub-Critically Reheated Heat Affected Zone), which is the most important position for guaranteeing the welded part, especially the low temperature CTOD value, has microstructure almost similar to the microstructure of the base material , And by controlling the value of the relational expression 2 to 0.5 or less, the low-temperature impact characteristics and the CTOD value of the welded portion can be sufficiently secured.

The microstructure of the steel according to the present invention contains 70% by area or more of polygonal ferrite and needle-like ferrite in total, and contains 3.5% by area or less of MA phase (martensite-austenite composite phase).

The acicular ferrite is the most important and basic microstructure to not only increase the strength due to the fine grain size effect but also to prevent crack propagation at low temperatures. Since polygonal ferrite is relatively large compared to acicular ferrite, it contributes relatively little to the increase in strength, but it has a low dislocation density and high grain boundaries and is a microstructure that contributes greatly to suppressing propagation at low temperatures.

When the sum of the polygonal ferrite and the needle-shaped ferrite is less than 70% by area, it is difficult to suppress crack initiation and propagation at low temperatures, and it is difficult to ensure high strength. Accordingly, the total of the polygonal ferrite and the needle-shaped ferrite is preferably 70% by area or more, more preferably 85% by area or more, still more preferably 90% by area or more.

In the present invention, the polygonal ferrite and the needle-shaped ferrite have a ratio of the grain size of the large diameter angle in which the difference in crystal orientation between the crystal grains is defined as not less than 15 degrees is not less than 40% in the total grain boundaries and the length of the large- / mm < 2 > or more.

Since the MA phase does not accept deformation due to its high hardness, it not only concentrates the deformation of the soft ferrite base around it, but also separates the interface with the surrounding ferrite base or destroys the MA phase, Lt; / RTI > Therefore, it is the most important cause for deteriorating the low-temperature fracture characteristics of the steel, so the MA phase should be controlled as low as possible, and it is preferable to control the MA phase to 3.5% or less.

At this time, in the present invention, the MA phase may have an average size measured at a circle-equivalent diameter of 2.5 mu m or less. When the average size of the MA phase is more than 2.5 탆, the MA is more likely to be broken due to more concentrated stress, and acts as a starting point of crack initiation.

Further, in the present invention, the polygonal ferrite and the needle-like ferrite may not be work-hardened by hot rolling. That is, the polygonal ferrite and the needle-like ferrite may not be drawn by hot rolling, and the polygonal ferrite and the needle-shaped ferrite may be produced after hot rolling.

The microstructure of the steel of the present invention may include bainitic ferrite, cementite, etc. in addition to the polygonal ferrite, acicular ferrite and MA phase described above.

The bainitic ferrite is a transformed structure at low temperature and has many internal potentials. However, bainitic ferrite has a relatively stronger characteristic than ferrites and also contains an MA phase therein, so that its strength is high. However, And therefore should be controlled to a minimum.

Further, the steel material of the present invention may contain inclusions having a size of 10 m or more in a range of 11 pieces / cm ² or less. The size is the size measured in the circle equivalent diameter. When the inclusions having a size of 10 탆 or more are more than 11 pieces / cm ² , there arises a problem of acting as a crack initiation point at a low temperature. In order to control the coarse inclusions in this way, it is preferable to bubbling and refluxing with Ar gas for 3 minutes or more after charging Ca or Ca alloy in the last stage of secondary refining.

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

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

A method of manufacturing a steel material according to the present invention includes the steps of: preparing a steel slab satisfying the above-described alloy composition; Heating the steel slab to a temperature of 1000 to 1200 캜; Finishing hot-rolling the heated slab in a temperature range of 650 ° C or higher; And cooling the finish hot-rolled hot-rolled steel sheet to a cooling end temperature of 200 to 550 ° C at a cooling rate of 2 to 30 ° C / s.

Step of preparing steel slab

A steel slab satisfying the alloy composition as described above is provided.

At this time, in the present invention, in preparing the steel slab, a step of injecting Ca or a Ca alloy into molten steel at a second refining end of the molten steel; And bubbling and refluxing with Ar gas for 3 minutes or more after the Ca or Ca alloy is charged. This is to control coarse inclusions.

Steel slab heating stage

The steel slab is heated to 1000 to 1200 ° C.

When the heating temperature of the slab is less than 1000 ° C., it is difficult to reuse the carbides generated in the slab during the performance, and the homogenization of the segregated elements becomes insufficient. Therefore, it is preferable to heat to at least 1000 ° C, at which 50% or more of the added Nb can be reused.

On the other hand, if the slab heating temperature exceeds 1200 ° C, the austenite grain size may grow excessively large, and further fineness may be insufficient due to subsequent rolling, and the mechanical properties such as tensile strength and low temperature toughness of the steel sheet may be greatly reduced .

Hot rolling step

The heated slab is subjected to finish hot rolling at 650 DEG C or higher, which is the bainite formation initiation temperature, to obtain a hot-rolled steel sheet.

When the finish hot rolling temperature is lower than 650 ° C, coarse bainite is produced and the work hardens during rolling, and the strength is excessively increased excessively. On the contrary, the impact toughness at low temperature is greatly reduced, . That is, when the hot rolling temperature is low, coarse erosion ferrite is produced before the hot rolling finish, and after that, it is stretched by rolling and work hardening is performed, and the remaining austenite remains in a band form, And the low-temperature toughness is lowered.

In the present invention, sufficient strain energy is accumulated in the austenite by performing a total reduction ratio of 30% or more (except for the recrystallization reverse reduction ratio) in the non-recrystallization inverse temperature region, so that polygonal and needle- It is preferable to sufficiently generate ferrite and ensure the ratio and density of the large-diameter grain boundaries.

Cooling step

Next, in the present invention, the finished hot-rolled steel sheet is cooled.

At this time, it is preferable to cool the hot-rolled steel sheet to a cooling end temperature of 200 to 550 ° C at a cooling rate of 2 to 30 ° C / s. If the cooling rate is less than 2 DEG C / s, the cooling rate is too slow to coarse the ferrite, pearlite and bainite transformation sections, and the strength and low-temperature toughness may suffer from thermal degradation. Rabynite or martensite is formed and the strength is increased, but the low-temperature toughness may be extremely dull.

When the cooling end temperature is higher than 550 ° C, fine structures such as needle-like ferrite are hardly generated and there is a high possibility that coarse bainite or pearlite is produced. On the other hand, when the temperature is lower than 200 ° C, there is no disadvantage in terms of microstructure, but there is a problem that the time required for cooling is excessive and productivity is greatly reduced.

Meanwhile, in the present invention, if necessary, the cooled hot-rolled steel sheet is heated to 450 to 650 ° C., maintained for (1.3 × t + 5) to (1.3 × t + 200) (Where t is the thickness of the hot-rolled steel sheet measured in mm). This is because when MA or martensite is excessively produced, MA or martensite is decomposed to remove the high dislocation density therein, and precipitated Nb or the like, which is a small amount, is solidified with carbonitrides to further improve the yield strength or low temperature toughness It is for this reason.

However, if the heating temperature is lower than 450 占 폚, the softening of the ferrite base is not sufficient, and the embrittlement phenomenon due to the P segregation or the like appears, which may deteriorate toughness. On the other hand, when the heating temperature is higher than 650 ° C, the recovery and growth of the crystal grains occur rapidly, and when the temperature is higher, the steel is partially transformed into austenite and the yield strength is lowered and the low temperature toughness may be deteriorated.

When the holding time is less than (1.3 x t + 5) minutes, the homogenization of the tissue is not sufficiently performed, and when the holding time is more than (1.3 x t + 200) minutes, the productivity is lowered.

Hereinafter, the present invention will be described in more detail by way of examples.

(Example)

The slabs having the composition shown in the following Table 1 were heated, hot-rolled and cooled under the conditions shown in Table 2 to produce steels.

The microstructure of the steel material thus prepared was observed, and physical properties thereof were measured and are shown in Table 3 below.

The CTOD value (-20 DEG C) of the weld heat affected zone (SCHAZ) was measured after welding the above-prepared steel material, and the results are shown in Table 3 below. The CTOD value (-20 ° C) of the steel is higher than that of the weld heat affected part, so the CTOD value (-20 ° C) for the steel is not separately measured.

At this time, the microstructure of the steel material was polished to a specular surface after polishing the cross section of the steel material, and etched with Nital or LePera according to the purpose, and a certain area of the specimen was measured with an optical or scanning electron microscope at a magnification of 100 to 5000 times, The fraction of the phase was measured from the measured image using an image analyzer. In order to obtain a statistically significant value, the same specimen was repeatedly measured by changing its position, and the average value thereof was determined.

Also, in order to observe the characteristics of the fabricated structure more precisely, the Nittal etched specimens were subjected to EBSD (Electron Back Scatter Diffraction) measurement with a scanning electron microscope to quantitatively measure the grain boundary characteristics of the produced steel

The physical properties of the steel are described by measuring from the nominal strain-nominal stress curve obtained by ordinary tensile tests.

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

The CTOD value (-20 ° C) shall be determined by machining the specimen in the size of B (thickness) x B (width) x 5B (length) perpendicular to the rolling direction according to BS 7448 standard and to make the fatigue crack length approximately 50% After the fatigue crack was inserted, the CTOD test was performed at -20 ° C. Where B is the thickness of the steel produced.

The value of Kca was tested three times by the ESSO test method, and the graph of the propagation stop temperature and the K value of the crack measured in each test was obtained. The K value was obtained from the K value (Kca: crack arrest K) at the temperature of -10 degrees. Also, the crack arrest temperature (CAT) was measured from the NRL-ductility transition temperature (NDTT) and calculated from the equation (1). Here, B represents the thickness of the steel material.

division
Steel grade
Alloy composition (% by weight) Relationship 1 Relation 2 C Si Mn P S Sol.Al Cu Ni Cr Mo Ti Nb V N Ca foot
persons
River a 0.035 0.18 2.2 0.005 0.0011 0.005 0.28 1.3 0.05 0.02 0.012 0.005 0.003 0.0035 0.002 3.0 0.39 b 0.046 0.01 1.9 0.005 0.0008 0.02 0.31 0.99 0.03 0.29 0.012 0.011 0.003 0.0042 0.0025 2.6 0.45 c 0.015 0.008 2.5 0.005 0.001 0.005 0.09 0.61 0.02 0.24 0.012 0.003 0.005 0.0026 0.0016 2.9 0.30 ratio
School
River


d 0.081 0.015 1.9 0.005 0.0012 0.022 0.31 1.13 0.03 0.2 0.012 0.012 0.003 0.0035 0.0028 2.6 0.37 e 0.021 0.011 3.2 0.005 0.0013 0.015 0.03 0.25 0.03 0.02 0.012 0.007 0.003 0.0035 0.0016 3.3 0.14 f 0.029 0.01 2.1 0.005 0.0009 0.011 0.29 0.75 0.12 0.15 0.012 0.038 0.003 0.0035 0.0011 2.6 0.67 g 0.025 0.015 2.4 0.005 0.0012 0.005 0.12 0.65 0.23 0.35 0.012 0.018 0.003 0.0035 0.0021 2.8 0.78 h 0.061 0.051 1.7 0.005 0.0018 0.028 0.15 0.25 0.03 0.2 0.012 0.005 0.003 0.0035 0.0018 1.9 0.36 i 0.034 0.017 1.9 00.057 0.0008 0.008 0.22 0.53 0.33 0.19 0.007 0.012 0.004 0.0039 0.0015 2.3 0.67

* In Table 1, the relation 1 is Mn + 0.5 x (Ni + Cu), and the relation 2 is Mo + Cr + 1.5 x Si + 10 x Nb.

division Steel grade Product thickness (mm) Reheating temperature of slab (℃) Rolling end temperature (캜) Non-recrystallization reverse total rate (%) Cooling rate (° C / s) Cooling end temperature (캜) Whether tempering is carried out Inventory 1 a 85 1060 790 45 4.0 390 x Inventory 2 b 100 1100 750 40 3.4 340 x Inventory 3 c 51 1120 800 57 6.9 280 x Honorable 4 b 76 1100 810 50 4.0 260 ○ Comparative Example 1 d 76 1110 780 48 4.5 380 x Comparative Example 2 e 51 1140 840 55 7.5 330 x Comparative Example 3 f 85 1120 805 35 3.1 280 x Comparative Example 4 g 80 1100 800 45 4.5 440 x Comparative Example 5 h 80 1100 780 40 4.3 360 x Comparative Example 6 i 100 1180 770 35 3.3 310 x Comparative Example 7 a 90 1170 820 20 3.5 270 x Comparative Example 8 c 76 1160 800 50 - - x

division Steel grade Ferrite phase fraction (%) Percentage of grain boundary grain boundaries (%) Confinement angle grain boundary length density (mm / mm ² ) Yield strength (MPa) Tensile Strength (MPa) Impact energy value (-60 ℃ base material): J Kca (-10 ° C) CAT (캜) HAZ
CTOD value
(-20 ° C): mm Inventory 1 a 78 53 390 480 620 315 9350 -16 0.82 Inventory 2 b 73 42 320 525 679 279 - -17 0.36 Inventory 3 c 83 63 490 488 627 332 - -20 1.21 Honorable 4 b 85 50 390 533 637 322 - -26 0.76 Comparative Example 1 d 56 43 380 596 766 75 - 11 0.18 Comparative Example 2 e 57 35 220 520 668 56 - -4 0.14 Comparative Example 3 f 55 33 190 540 696 90 5860 3 0.08 Comparative Example 4 g 46 42 210 587 754 103 - 2 0.12 Comparative Example 5 h 75 51 390 428 554 268 - -25 0.35 Comparative Example 6 i 55 38 220 471 610 85 - -9 0.08 Comparative Example 7 a 61 23 180 475 615 112 - -2 0.15 Comparative Example 8 c 86 56 480 435 545 185 - -36 0.29

* In Table 3, the ferrite system means the sum of polygonal ferrite and needle-shaped ferrite.

As shown in Table 1-3, Inventive Examples 1 to 4, which satisfy both the alloy composition and the manufacturing conditions proposed in the present invention, are excellent in fracture at low temperature in consideration of yield strength, tensile strength, impact energy value, Kca and CAT It can be confirmed that the toughness resistance is excellent and the CTOD value in the weld heat affected zone is also high. Particularly, as shown in FIG. 1, the value of Kca measured in Inventive Example 1 shows a value greatly exceeding the required value of 8000. These excellent strengths and low temperature toughness characteristics are also obtained from the fine polygonal and acicular ferrite structures sufficiently generated as shown in Fig.

On the other hand, in Comparative Example 1, when the C content exceeds the range of the present invention, added C is the most powerful element promoting granulobenite and MA. Therefore, the excessively high C content causes a significant reduction in the fraction of ferrite that is favorable to toughness, so that the strength in the base material is high, but the low temperature toughness such as impact energy value is poor.

Comparative Example 2 is a case where the added Mn content exceeds the range of the present invention. In this case, since the Mn content is high, the probability of segregation at the center of the steel greatly increases, impact energy at the center of the thickness direction of steel is largely damped, and the hardened structure, And the CTOD value dropped significantly due to a pop-in phenomenon.

Comparative Example 3 is a case where the content of Nb, which is widely used for strength improvement and texture refinement, generally exceeds the range of the present invention. In general, the addition of Nb is advantageous in simultaneously increasing the strength and toughness by refining the structure. However, if the Nb is added more than necessary, the formation of polygonal and needle-like ferrite favorable to toughness is suppressed and the structure such as granular bainite is promoted . Therefore, it is possible to relatively easily propagate cracks by greatly reducing the density and the ratio of the large-diameter grain boundaries of 15 DEG or more, which is advantageous for suppressing propagation of cracks. As a result, as shown in Table 3, the value of Kca measured in Comparative Example 3 was 5860, which is far below the required value of 8,000. In addition, in the heat-affected zone of the weld, the formation of M-A structure, which is particularly disadvantageous to low-temperature toughness, was greatly promoted and consequently the CTOD was significantly lowered.

 In Comparative Examples 4, 5 and 6, the ranges of the respective elemental contents satisfied the range of the present invention. However, when the values of relational expression 1 and relational expression 2 were outside the scope of the present invention, it was found that the strength was low or the low- have.

Specifically, in the comparative example 4, the relationship 1, which is composed of the components favorable for improving the low-temperature toughness, is satisfied, but the relationship 2 composed of the components that deteriorate the low-temperature toughness exceeds the scope of the present invention. As a result, the strength was sufficiently high, but the impact energy value in the base material and the CTOD value in the weld heat affected zone were inferior.

In Comparative Example 5, Relation 2 satisfied the scope of the invention. However, when Relation 1 deviates from the scope of the present invention, the addition amount of the component was insufficient to secure the strength of the steel as a whole, and the strength of the base material was greatly lowered.

Comparative Example 6 is a case where the equations (1) and (2) are outside the scope of the invention. That is, the component favorable to low-temperature toughness is insufficient, and the component disadvantageous to low-temperature toughness is exceeded, and all the low-temperature toughness properties are heated.

Comparative Example 7 is a case in which the components of the steel meet all of the inventions but do not fall within the range of the invention of the non-recrystallized reverse rolling total reduction in the manufacturing process of the steel. That is, the amount of ferrite which interferes with the propagation of cracks in the microstructure of the steel is low due to insufficient amount of reduction in the non-recrystallized zone, and the ratio and density of the large-angle grain boundaries are drastically decreased, .

In Comparative Example 8, the steel composition satisfies all the ranges of the invention. However, in the case where the steel is manufactured by air cooling without applying accelerated cooling after controlled rolling in the manufacturing process of steel, ferrite Is sufficiently generated, but the strength of the coarsening is greatly lowered.

While the present invention has been particularly shown and described with reference to exemplary embodiments thereof, it is to be understood that the scope of the present invention is not limited to the disclosed exemplary embodiments, but various modifications and changes may be made without departing from the scope of the invention. To those of ordinary skill in the art.

Claims (10)

0.001 to 0.07% of Si, 0.005 to 0.3% of Si, 1.7 to 3.0% of Mn, 0.001 to 0.035% of Sol.Al, 0.02% or less of Nb (excluding 0% 0.001 to 0.02% of Ti, 0.01 to 1.0% of Cu, 0.01 to 2.0% of Ni, 0.01 to 0.5% of Cr, 0.001 to 0.5% of Mo, 0.0002 to 0.005% of Ca, (Excluding 0%), O: not more than 0.003% (excluding 0%), the remainder Fe and unavoidable impurities are contained in an amount of N: 0.001 to 0.008%, P: 0.02% And satisfy the following relational expression 1 and relational expression 2,
The microstructure includes 70% by area or more of polygonal ferrite and needle-like ferrite in total, and has a high strength with excellent fracture toughness at low temperature including 3.5% by area or less of MA phase (martensite-austenite composite phase) Steel.
[Relation 1]
Mn + 0.5 x (Ni + Cu)? 2.5 wt%
[Relation 2]
Relation 2: Mo + Cr + 1.5 x Si + 10 x Nb? 0.5 wt%
(In the above relational expression 1-2, each element is a value expressed in weight%).
The polygonal ferrite according to claim 1, wherein the polygonal ferrite and the needle-like ferrite have a ratio of a large grain size grain boundary in which the difference in crystal orientation between crystal grains is defined as 15 degrees or more, is 40% or more in the total grain boundaries, Is not less than 300 mm / mm < 2 >.
The steel according to claim 1, wherein the steel has a yield strength of 460 MPa or more, an impact energy value at -60 캜 of 250 J or more, a Kca value measured by ESSO test of 8000 N / mm ^3/2 or more, Characterized in that the crack arrest temperature (CAT) calculated from the Nil-ductility transition temperature (NDTT) measured at the low temperature is less than -10 占 폚.
The high strength steel material according to claim 1, wherein the steel material has a tensile strength of 570 MPa or more and a DBTT (ductile-brittle transition temperature) of -80 占 폚 or less.
The high strength steel material according to claim 1, wherein the steel material includes inclusions having a size of 10 탆 or more measured at a circle equivalent diameter in a range of 11 pieces / cm ² or less.
0.001 to 0.07% of Si, 0.005 to 0.3% of Si, 1.7 to 3.0% of Mn, 0.001 to 0.035% of Sol.Al, 0.02% or less of Nb (excluding 0% 0.001 to 0.02% of Ti, 0.01 to 1.0% of Cu, 0.01 to 2.0% of Ni, 0.01 to 0.5% of Cr, 0.001 to 0.5% of Mo, 0.0002 to 0.005% of Ca, (Excluding 0%), O: not more than 0.003% (excluding 0%), the remainder Fe and unavoidable impurities are contained in an amount of N: 0.001 to 0.008%, P: 0.02% Providing a steel slab satisfying the following relational expression 1 and relational expression 2;
Heating the steel slab to a temperature of 1000 to 1200 캜;
Subjecting the heated slab to a finish hot rolling such that the total rolling reduction in the temperature range of not lower than 650 ° C is not less than 30% (excluding the recrystallization reverse rolling reduction rate) in the non-recrystallization reverse temperature range; And
And cooling the finished hot-rolled steel sheet at a cooling rate of 2 to 30 占 폚 / s to a cooling end temperature of 200 to 550 占 폚.
[Relation 1]
Mn + 0.5 x (Ni + Cu)? 2.5 wt%
[Relation 2]
Relation 2: Mo + Cr + 1.5 x Si + 10 x Nb? 0.5 wt%
(In the above relational expression 1-2, each element is a value expressed in weight%).
The method according to claim 6, further comprising a tempering step of heating the cooled hot-rolled steel sheet to 450 to 650 ° C, then maintaining the cooled hot-rolled steel sheet for (1.3 × t + 5) minutes to (1.3 × t + 200) Which is excellent in fracture resistance at low temperatures.
The method according to claim 6,
In providing the steel slab,
Introducing Ca or a Ca alloy into the molten steel at a final stage of secondary refining of the molten steel; And bubbling and refluxing with Ar gas for at least 3 minutes after the Ca or Ca alloy is charged.
7. The steel according to claim 6, wherein the microstructure of the cooled steel includes 70% or more of the total of polygonal ferrite and acicular ferrite, and 3.5% or less of the MA phase (martensite-austenite composite phase) The method of manufacturing a high strength steel material according to claim 1,
The method according to claim 9, wherein the polygonal ferrite and the needle-shaped ferrite have a ratio of a grain boundary of large-diameter angles in which the difference in crystal orientation between crystal grains is defined as not less than 15 degrees is not less than 40% Is not less than 300 mm / mm < 2 >.

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