KR101568551B1 - Steel sheet for linepipe with excellent deformability and low-temperature toughness and method for manufacturing the same - Google Patents

Abstract

A steel plate for a line pipe excellent in deformability and low-temperature toughness and a method of manufacturing the same. The steel sheet for a line pipe excellent in deformability and low temperature toughness, which is one aspect of the present invention, contains 0.03 to 0.05% of C, 0.01 to 0.2% of Si, 0.5 to 2.0% of Mn, 0.01 to 0.3% : 0.01 to 0.3% of Nb, 0.01 to 0.05% of Ti, 0.01 to 0.02% of Ti, 0.05% of Al or less and the balance of Fe and unavoidable impurities. The microstructure is a polygonal ferrite of 60 to 80% 20 to 35 area% acicular ferrite, and not more than 5 area% of the second phase.

Description

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a steel sheet for a line pipe excellent in flexibility and low temperature toughness,

The present invention relates to a steel sheet for a line pipe excellent in deformability and low-temperature toughness and a method for producing the same.

Recently constructed line pipes are required to have excellent deformability with excellent strength and toughness. In particular, there is a growing demand for excellent deformability with excellent low temperature toughness in order to increase the fracture resistance against gradual or rapid deformation accompanied by ground movements occurring in harsh environments such as frozen soil, load of the structure itself, and earthquake . The steel for the line pipe is buried after corrosion-resistant coating, and the uniform elongation is greatly reduced and the deformability is deteriorated. On the other hand, it is known that a Cottrell atmosphere is formed at a coating temperature of 200 ° C in the coating and the carbon atoms in the ferrite are fixed to the dislocation to cause a yield point phenomenon and a uniform elongation is lowered.

There are known a method using a Ti-based or Nb-based carbide to reduce the amount of carbon atoms in the ferrite that promotes a phenomenon that the uniform elongation is lowered, and a method of precipitating cementite to remove carbon atoms in the ferrite. Among these, the Ti-based or Nb-based carbide is produced at a high temperature of 600 ° C or higher, and therefore it is stable to precipitate the cementite in order to reduce carbon atoms even at a temperature lower than 600 ° C. Therefore, it is desirable to reduce the content of elements such as Si that inhibits the formation of cementite. However, since silicon serves to deoxidize molten steel by assisting aluminum, it also has an effect as a solid solution strengthening element. Therefore, it is important to find an optimal silicon content by adding 0.01 wt% or more of silicon.

On the other hand, conventional high-strain steel for line pipe has polygonal ferrite (PF) as a main structure. However, such a high-strength steel has a disadvantage in that the low-temperature transformed phase having a high dislocation density and the fraction of the second phase are too low to exhibit discontinuous yielding behavior in a tensile test. Therefore, there have been many attempts to form the low-temperature transformation phase and the second phase at a certain level or more and to distribute the low-temperature transformation phase and the second phase more finely and evenly. However, no effective method has been proposed so far.

An aspect of the present invention is to provide a steel sheet for a line pipe having excellent low-temperature toughness and deformability, and a method of manufacturing the same.

In order to achieve the above object, one aspect of the present invention provides a ferritic stainless steel comprising 0.03 to 0.05% of C, 0.01 to 0.2% of Si, 0.5 to 2.0% of Mn, 0.01 to 0.3% of Ni, 0.01 to 0.3% of Cr, 0.01 to 0.3% of Cr, , The balance being Fe and inevitable impurities, the microstructure is composed of 60 to 80% by area of polygonal ferrite, 20 to 20% by area of Nb, 0.01 to 0.05% of Ti, 0.01 to 0.02% A steel sheet for a line pipe excellent in deformability and low temperature toughness comprising acicular ferrite having an area of 35% by area and a second phase of 5% by area or less is provided.

In another aspect of the present invention, there is provided a steel sheet comprising, by weight, 0.03 to 0.05% of C, 0.01 to 0.2% of Si, 0.01 to 2.0% of Mn, 0.01 to 0.3% of Ni, 0.01 to 0.3% of Cr, A slab reheating step of reheating a steel slab containing 0.01 to 0.05%, Ti: 0.01 to 0.02%, Al: not more than 0.05%, the remainder Fe and unavoidable impurities, and then extracting the steel slab at 1100 to 1140 ° C; A recrystallization reverse rolling step of starting rolling on the reheated steel slab to finish rolling at Tnr 占 폚 to Tnr + 100 占 폚; A non-recrystallization reverse rolling step of starting rolling the recrystallized reverse-rolled steel sheet at a temperature between Tnr-200 ° C and Tnr-150 ° C, and finishing the rolling at a temperature of Ar3-20 ° C to Ar3 ° C; And a cooling step of cooling the non-recrystallized back-rolled steel sheet at a rate of 10-20 ° C / sec to a temperature of Ms-80 ° C to Ms ° C. The present invention also provides a method of manufacturing a steel sheet for a line pipe excellent in deformability and low-temperature toughness.

In addition, the solution of the above-mentioned problems does not list all the features of the present invention. BRIEF DESCRIPTION OF THE DRAWINGS The various features and advantages and effects of the present invention will become more fully understood with reference to the following specific embodiments.

The steel sheet for a line pipe according to the present invention preferably has a DWTT elongation wavefront ratio of 85% or more at -30 캜, a uniform elongation after heat treatment of 11% or more after preliminary deformation, and is suitable for a steel sheet for a line pipe used in cold climates Lt; / RTI >

Fig. 1 shows microscopic microstructures of Inventive Example 1 and Comparative Example 1 observed under an optical microscope.
2 is a graph showing FE-TEM observation of TiNb composite carbonitrides of Inventive Example 1 and Comparative Example 1 of the present invention.

Hereinafter, a steel sheet for a line pipe having excellent deformability and low temperature toughness as one aspect of the present invention will be described in detail.

First, the alloy composition of a steel sheet for a line pipe having excellent deformability and low temperature toughness according to the present invention will be described in detail. Herein, unless otherwise stated, each component content means weight%.

Carbon (C): 0.03 to 0.05%

C is an effective element for strengthening steel by solid solution strengthening and precipitation strengthening, but when it is added in a large amount, the toughness of steel is lowered. If the content of carbon is less than 0.03%, it is difficult to secure a desired strength because it is difficult to obtain a second phase such as martensite, pearlite and cementite. On the other hand, when the carbon content exceeds 0.05% There is a problem. Therefore, in the present invention, the carbon content is preferably limited to 0.03 to 0.05%.

Silicon (Si): 0.01 to 0.2%

Si not only serves to deoxidize molten steel, but also serves as a solid solution strengthening element to improve the strength of steel. In order to obtain such an effect in the present invention, the content of silicon is preferably 0.01% or more. However, when the content of silicon exceeds 0.2%, the formation of two phases such as cementite is inhibited and the deformability is lowered, and toughness and weldability are deteriorated. In addition, the TiNb complex carbonitride is coarsened and low temperature toughness of the base material is deteriorated. Therefore, in the present invention, the content of silicon is preferably limited to 0.01 to 0.2%.

Manganese (Mn): 0.5 to 2.0%

Mn is a solid solution strengthening element and serves to improve the strength of the steel. In addition, in the present invention having a low carbon content, Mn plays a role of promoting the formation of martensite structure by compensating the hardenability reduced by the low carbon content. When the content of manganese is less than 0.5%, it is difficult to secure the desired strength. On the other hand, when the content of manganese exceeds 2.0%, center segregation is promoted during slab casting, and the toughness and weldability of steel are deteriorated. Therefore, in the present invention, the content of manganese is preferably limited to 0.5 to 2.0%.

Nickel (Ni): 0.01 to 0.3%

Ni is an effective element for improving the strength of steel in the present invention having a small carbon content without deteriorating the weldability in the field. If the content of nickel is less than 0.01%, there is a problem that strength and toughness of steel are lowered. On the other hand, if the content exceeds 0.3%, the manufacturing cost increases and it is economically disadvantageous. Therefore, in the present invention, the nickel content is preferably limited to 0.01-0.3%.

Cr (Cr): 0.01 to 0.3%

Cr is an element effective in securing sufficient curing ability upon cooling in the present invention having a low carbon content and forming a second phase such as cementite. In order to obtain such effects in the present invention, the content of chromium is preferably 0.01% or more. On the other hand, if the content exceeds 0.3%, the manufacturing cost is increased and economically disadvantageous. Therefore, in the present invention, the content of chromium is preferably limited to 0.01 to 0.3%.

Niobium (Nb): 0.01 to 0.05%

Nb is an element that plays a role in simultaneously improving the strength and toughness of steel through grain refinement. The Nb carbonitride produced during hot rolling reduces the austenite recrystallization and grain growth, thereby finer austenite grains. In order to obtain such an effect in the present invention, the content of niobium is preferably 0.01% or more. On the other hand, when the content exceeds 0.05%, not only the effect is saturated but also Nb carbonitride is excessively precipitated to excessively increase the austenite non-recrystallization temperature, thereby increasing the material anisotropy and increasing the cost, There is a problem that the toughness of the weld heat affected zone (HAZ) deteriorates. Therefore, in the present invention, the content of niobium is preferably limited to 0.01 to 0.05%.

Titanium (Ti): 0.01 to 0.02%

Ti forms precipitates in the solidification process of steel to suppress the growth of austenite grains during the slab heating and hot rolling process, thereby improving the strength of the steel by making the grain size of the final structure finer. In order to obtain such effects in the present invention, the content of titanium is preferably 0.01% or more. On the other hand, when the content exceeds 0.02%, there is a problem that the precipitates are coarse and the toughness of the steel is lowered.

Aluminum (Al): 0.05% or less (excluding 0)

Al, like silicon, serves to deoxidize molten steel. If the content exceeds 0.05%, Al ₂ O ₃ , which is a nonmetal oxide, is formed and toughness of the base material and welded part is lowered. Therefore, The content of aluminum is preferably limited to 0.05% or less.

The remainder Fe and unavoidable impurities. The inevitable impurities include phosphorus (P), sulfur (S), nitrogen (N) and the like, and the content of these elements is preferably minimized. On the other hand, addition of an effective component other than the above-mentioned composition is not excluded.

Hereinafter, the microstructure and precipitates of the steel sheet for a line pipe according to the present invention will be described in detail.

The steel sheet for a line pipe according to the present invention not only satisfies the above-mentioned composition but also has a microstructure of polyhedral ferrite having 60 to 80 area%, acicular ferrite having 20 to 35 area% and less than 5 area% And a second phase of the second phase. According to an embodiment of the present invention, the second phase may be at least one selected from the group consisting of pearlite and cementite.

On the other hand, the average effective grain size of the composite structure is more preferably 10 탆 or less. Here, the average effective grain size is defined as the grain boundary in which the misorientation between neighboring grain orientations is 15 degrees or more irrespective of the type of microstructure, Mean the average size. In the present invention, excellent DWTT ductile waveguide ratio can be secured by controlling the mean effective grain size of the composite structure to 10 탆 or less as described above.

On the other hand, that according to one embodiment of the invention, at least the present invention as a line pipe steel sheet comprises a TiNb composite carbonitride, the TiNb an average particle diameter of the complex carbonitride is 20nm or less, number per unit area is 30 / ㎛ ² desirable. When the average particle diameter of the TiNb composite carbonitride is more than 20 nm or the number per unit area is less than 30 particles / 탆 ² , the effect of fineness is insignificant. Therefore, the average particle diameter and the unit of the TiNb composite carbonitride It is desirable to control the number per area.

The steel sheet satisfying the above-mentioned component system and internal structure has a DWTT elongation wavefront ratio of 85% or more at -30 캜 and a uniform elongation after annealing after preliminary deformation of 11% or more, It is a steel plate.

Hereinafter, a method of manufacturing a steel sheet for a line pipe having excellent deformability and low temperature toughness, which is another aspect of the present invention, will be described in detail as a preferred example for producing the steel sheet for a line pipe.

Slab reheat step

First, the steel slab having the above composition is reheated, and then the steel slab is extracted at 1100 to 1140 ° C. The reheating step of the slab is a step of heating the steel so as to smoothly perform the subsequent rolling process and sufficiently obtain the physical properties of the target steel sheet, and therefore, the heating must be performed within an appropriate temperature range according to the purpose. During the reheating step, it is necessary to uniformly heat the precipitation-type elements in the steel sheet to such an extent that the precipitation-type elements can be sufficiently solidified, and to prevent excessive coarsening of the crystal grains by heating at an excessively high temperature.

According to an embodiment of the present invention, the reheating of the steel slab is preferably performed at a temperature of 1150 to 1250 ° C. If the reheating temperature is less than 1150 占 폚, the slab may not be heated sufficiently, and alloying elements such as Nb may not be sufficiently employed. On the other hand, when the reheating temperature exceeds 1250 占 폚, austenite coarsening due to coarse TiN precipitation or Mixed austenite structure can be produced, and such coarse austenite is difficult to recrystallize during rough rolling and remains mainly in an elongated state. Such elongated coarse austenite generally has a high hardenability and tends to be transformed into coarse bainite after final cooling. This leads to deterioration of the low temperature DWTT characteristic.

Meanwhile, it is preferable that the reheated slab is extracted at a temperature of 1100 to 1140 ° C before extraction. When the extraction temperature is less than 1100 ° C, the subsequent process is difficult to hot-roll, while when it is more than 1140 ° C, it is difficult to miniaturize the structure.

Recrystallization station  Rolling step

Thereafter, the reheated slab is subjected to recrystallization back-rolling. Recrystallization reverse rolling is intended to homogenize the austenite grain size.

At this time, it is preferable that the rolling finish temperature is Tnr 占 폚 to Tnr + 100 占 폚. If the rolling finish temperature exceeds Tnr + 100 ° C, the grain size of the recrystallized austenite may be coarsened. On the other hand, when the rolling finish temperature is lower than Tnr ° C, there is a fear that coarse unrecrystallized austenite is generated have. Tnr = 887 + (464xC) + (6445xNb) - (644xNb) Theoretically, the Tnr temperature means a boundary temperature between a temperature region where austenite is recrystallized and a temperature region where austenite is not recrystallized. ) + ((732 x V) - (230 x V)) + (890 x Ti) + (363 x Al) - (357 x Si). However, the above formula is for convenience of application, and the Tnr can be confirmed by various test methods commonly used in the technical field of the present invention. However, if there are differences in the results of various methods, the above formula shall be followed.

Meanwhile, according to an embodiment of the present invention, it is more preferable that the average rolling reduction of the final pass during the recrystallization reverse rolling is 10% or more. If the average rolling reduction of the final pass is less than 10%, coarse unrecrystallized austenite may remain and the drop weight tearing test (DWTT) characteristic may be degraded.

Non recrystallization station  Rolling step

Thereafter, rolling is started on the recrystallization-back-rolled steel sheet at a temperature of Tnr-200 ° C to Tnr-150 ° C, and non-recrystallization reverse-rolling is performed to finish rolling at Ar 3 ° C to Ar 3 + 20 ° C. The purpose of the non-recrystallized reverse rolling is to form a microstructure and a bainite by forming a deformation zone in the steel sheet while stretching the austenite structure in the rolling direction.

When the rolling start temperature exceeds Tnr-150 캜 at the time of the non-recrystallization rolling, coarse austenite is formed due to occurrence of partial recrystallization, which leads to formation of a coarse low temperature transformation texture after cooling, . ≪ / RTI > On the other hand, in order to apply the target rolling finish temperature to be described later, it is preferable that the rolling starting temperature is performed at Tnr-200 DEG C or higher.

On the other hand, when the rolling finish temperature is higher than Ar3 DEG C during the non-recrystallization rolling, pancaking of the crystal grains is not sufficient and it is difficult to miniaturize the grain for improving the toughness of the steel. The Ar3 temperature means a temperature at which austenite is transformed into ferrite. In theory, Ar3 = 910- (273xC) - (74xMn) - (57xNi) - (16xCr) - - (5 x Cu).

Meanwhile, according to one embodiment of the present invention, the cumulative rolling reduction RRR4 in the non-recrystallization rolling is more preferably 75% or more. If the cumulative rolling reduction is less than 75%, there is a fear that a coarse bainite is generated after cooling.

Cooling step

The non-recrystallized reverse-rolled steel sheet is cooled to a temperature of Ms-80 ° C to Ms ° C at a rate of 10 to 20 ° C / sec. When the cooling rate is lower than 10 ° C / sec, the grain size becomes coarse and the toughness of the steel decreases. On the other hand, when the cooling rate exceeds 20 ° C / sec, There is a problem that the toughness of the steel is deteriorated. When the cooling end temperature exceeds Ms 占 폚 at the time of cooling, the effective grain size increases due to the formation of coarse bainite structure and low temperature tear deterioration occurs. Here, Ms means a martensitic transformation starting temperature, which is theoretically derived by Ms = 561- (474 x C) - (33 x Mn) - (17 x Ni) - (17 x Cr) - (21 x Mo) I can do it.

Hereinafter, the present invention will be described more specifically by way of examples. It should be noted, however, that the following examples are intended to illustrate the invention in more detail and not to limit the scope of the invention. The scope of the present invention is determined by the matters set forth in the claims and the matters reasonably inferred from them.

( Example  One)

Steel slabs having the compositions shown in the following Table 1 were reheated, rolled and cooled under the manufacturing conditions shown in Table 2 below to produce a steel sheet for a line pipe having a thickness of 33 mm.

Subsequently, a part of the finally prepared steel sheet was taken and subjected to a DWTT test at -30 ° C to measure the DWTT ductile wave fracture ratio and tensile properties. The texture and precipitates of each steel sheet were analyzed, Respectively. Meanwhile, in Table 3, the average effective grain size was measured by EBSD (Electron Back-Scattered Diffraction) without discrimination between microstructures and the average effective grain size was defined as a misorientation angle of 15 ° or more.

Thereafter, each steel sheet was subjected to heat treatment at an atmospheric temperature of 200 DEG C for 5 minutes after 2% of preliminary strain was applied, and then tensile properties were evaluated. The results are shown in Table 4.

Steel grade Composition (% by weight) theory
Tnr (占 폚) theory
Ar3 (占 폚) theory
Ms (占 폚) C Si Mn Al Cr Ni Ti Nb A 0.047 0.13 1.5 0.029 0.153 0.097 0.014 0.039 1009.56 778.192 490.174 B 0.049 0.156 1.57 0.029 0.155 0.1 0.015 0.043 1021.513 772.263 486.899 C 0.05 0.17 1.52 0.026 0.13 0.1 0.013 0.039 994.6932 776.09 487.65 D 0.052 0.26 1.6 0.028 0.157 0.1 0.016 0.04 971.712 769.192 484.521 E 0.05 0.258 1.59 0.028 0.157 0.1 0.015 0.039 965.7832 770.478 485.799 F 0.038 0.254 1.45 0.031 - 0.3 0.018 0.04 970.227 775.226 490.038 G 0.053 0.245 1.67 0.031 0.157 - 0.016 0.042 988.3293 769.439 483.437 H 0.056 0.233 1.5 0.03 0.156 0.103 0.016 0.041 988.7779 775.345 485.857

Steel grade Reheating Recrystallization reverse rolling Non-recrystallized reverse rolling Cooling Remarks Extraction temperature
(° C) Termination temperature
(° C) Initiation temperature
(° C) Termination temperature
(° C) Cooling rate (° C / s) End temperature (캜) A 1100 1063 823 768 13 420 Inventory 1 B 1110 1055 822 768 10 480 Inventory 2 C 1110 1060 818 768 18 450 Inventory 3 D 1100 1065 821 773 11 480 Comparative Example 1 E 1110 1057 823 768 15 500 Comparative Example 2 F 1109 1057 822 770 21 480 Comparative Example 3 G 1110 1062 820 770 19 430 Comparative Example 4 H 1100 1061 823 769 19 390 Comparative Example 5

Steel grade Microstructure
(area%) Average effective grain size (占 퐉) TiNb precipitates Mechanical properties Remarks Average particle size (nm) Number (number / 탆 ² ) YS
(MPa) TS
(MPa) Hand
(%) U.El
(%) DWTT
at -30 ℃
(%) A PF75.3 + AF21.6 7.5 18.7 34 457 527 31 13.2 90 Inventory 1 B PF62.5 + AF30.1 9.1 17.6 39 473 533 30 12.7 88 Inventory 2 C PF74.0 + AF22.8 8.1 19.0 40 482 542 32 13.0 100 Inventory 3 D PF89.9 + AF8.0 11.0 20.1 21 483 544 35 14.8 81 Comparative Example 1 E PF83.8 + AF13.8 10.4 22.5 19 493 547 35 14.8 83 Comparative Example 2 F PF85.6 + AF11.6 12.4 25.1 20 474 537 26.9 12.2 84 Comparative Example 3 G PF84.7 + AF12.3 16.3 23.4 15 502 570 32.2 12.7 81 Comparative Example 4 H PF82.7 + AF11.8 18.3 22.7 14 481 547 32.4 13.5 75 Comparative Example 5

Steel grade YS (MPa) TS (MPa) El (%) Y.El (%) Remarks A 468 543 28.6 11.1 Inventory 1 B 472 514 31.7 12.8 Inventory 2 C 446 534 28.5 11.1 Inventory 3 D 480 555 27.3 10.2 Comparative Example 1 E 485 538 28.6 10.6 Comparative Example 2 F 447 509 29.5 10.4 Comparative Example 3 G 467 548 24.4 8.2 Comparative Example 4 H 454 521 26.5 8.7 Comparative Example 5

As can be seen from Tables 3 and 4, in Examples 1 to 3, which satisfy the alloy composition and manufacturing conditions proposed in the present invention, the DWTT elongation wavefront ratio at 30 ° C is 85% or more, It can be confirmed that the uniform elongation is 11% or more.

However, Comparative Examples 1 to 5 did not satisfy the alloy composition and / or manufacturing conditions proposed by the present invention, and DWTT elongation wavefront ratios at a temperature of 30 ° C were shown to be less than 85%, and the uniform elongation after annealing after pre- And 11% less.

1 (a) is a photograph of microstructure of Inventive Example 1, and Fig. 1 (b) is a micrograph of microstructure of Inventive Example 1 and Comparative Example 1, It is an organization photograph. 2 is a photograph of the TiNb composite carbonitride of Inventive Example 1 and Comparative Example 1 observed by FE-TEM. Fig. 2 (a) is a photograph of the TiNb composite carbonitride of Inventive Example 1, and Fig. 2 b) is a photograph of the TiNb composite carbonitrides of Comparative Example 1 observed.

Claims (9)

The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.03 to 0.05% of C, 0.01 to 0.2% of Si, 0.5 to 2.0% of Mn, 0.01 to 0.3% of Ni, 0.01 to 0.3% of Cr, 0.01 to 0.3% 0.02%, Al: not more than 0.05% (excluding 0%), the balance Fe and unavoidable impurities,
A microstructure having a composite structure comprising 60 to 80 area% of polygonal ferrite, 20 to 35 area% of acicular ferrite and 5% or less of a second phase,
Wherein the mean circle equivalent diameter of the effective crystal grains of the composite structure is 10 占 퐉 or less and excellent in deformability and low temperature toughness.
The method according to claim 1,
Wherein the second phase is at least one selected from the group consisting of pearlite and cementite and has excellent deformability and low temperature toughness.
delete The method according to claim 1,
Wherein the steel sheet comprises a TiNb composite carbonitride, the average particle diameter of the TiNb composite carbonitride is 20 nm or less, the number per unit area is 30 pieces / 占 퐉 ² or more, and the low temperature toughness is excellent.
The method according to claim 1,
Wherein the steel sheet has a DWTT elongation wave-surface ratio of 85% or more at -30 ° C, a strain of uniform elongation after heat treatment after pre-strain of 11% or more, and low temperature toughness.
The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.03 to 0.05% of C, 0.01 to 0.2% of Si, 0.5 to 2.0% of Mn, 0.01 to 0.3% of Ni, 0.01 to 0.3% of Cr, 0.01 to 0.3% 0.02%, Al: 0.05% or less (excluding 0%), the remainder Fe and inevitable impurities, and then extracting the steel slab at 1100 to 1140 ° C;
A recrystallization reverse rolling step of starting rolling on the reheated steel slab to finish rolling at Tnr 占 폚 to Tnr + 100 占 폚;
A non-recrystallization reverse rolling step of starting rolling the recrystallized reverse-rolled steel sheet at a temperature between Tnr-200 ° C and Tnr-150 ° C, and finishing the rolling at a temperature of Ar3-20 ° C to Ar3 ° C; And
And a cooling step of cooling the non-recrystallized reverse-rolled steel sheet at a rate of 10 to 20 ° C / sec to a temperature of Ms-80 ° C to Ms ° C, and having excellent deformability and low temperature toughness.
The method according to claim 6,
Wherein the reheating of the steel slab is carried out at a temperature of 1150 to 1250 占 폚 and is excellent in low temperature toughness.
The method according to claim 6,
Wherein the average rolling reduction of the final pass is 10 to 20% at the time of recrystallization reverse rolling, and is excellent in deformability and low temperature toughness.
The method according to claim 6,
Wherein the cumulative rolling reduction ratio of the non-recrystallized rolling is 75% or more and the low temperature toughness is excellent.

Images