KR100723201B1 - High strength and toughness steel having superior toughness in multi-pass welded region and method for manufacturing the same - Google Patents

Abstract

The present invention relates to a steel for welding structures used in ships, offshore structures, line pipes and the like.

In the present invention, by weight%, C: 0.06 to 0.1%, Si: 0.05 to 0.35%, Mn: 1.3 to 1.7%, Cu: 0.1 to 0.6%, Ni: 0.1 to 1.0%, Sol.Al: 0.005 to 0.06% , Nb: 0.005 to 0.03%, Ti: 0.005 to 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities, Ceq = C [%] + Mn [%] / 6+ (Cr [%] High strength toughness steel with excellent toughness in multi-layer welds where the carbon equivalent (Ceq) value defined by + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 satisfies the range of 0.36 to 0.42; It relates to a manufacturing method.

The present invention is to provide a high strength high toughness welded structural steel with excellent weld heat affected zone toughness in yield heat and middle heat welded multi-layer welding portion, yield strength of 420MPa or more, tensile strength of 520MPa or more, ductile-brittle impact energy transition temperature of -70 ℃ or less. Can be.

Multi-layer welding part, heat of quenching, heat of medium input, high strength, high toughness

Description

High-strength, high-toughness steel with excellent multi-layer weldability and its manufacturing method {HIGH STRENGTH AND TOUGHNESS STEEL HAVING SUPERIOR TOUGHNESS IN MULTI-PASS WELDED REGION AND METHOD FOR MANUFACTURING THE SAME}

1 is a graph showing the DBTT value of the weld heat affected zone according to Ceq in the heat of quenching of 0.7 kJ / mm and the middle heat of 4.5kJ / mm.

The present invention relates to a welded structural steel used in ships, offshore structures, line pipes, and the like, and more particularly, to a high strength high toughness steel excellent in multi-layer welded part toughness and a method of manufacturing the same.

Recently, as the size of the structure is accelerated due to economic reasons, the demand for high strength and high toughness steel is increasing. In order to secure the strength of the steel, it is necessary to increase the addition amount of the alloying element, but when the addition amount of the alloying element is excessive, there is a problem that the low temperature toughness of the weld portion is greatly reduced.

In order to solve this problem, Japanese Patent Laid-Open No. 4-032536 has defined an alloy element range of steel and an upper limit of carbon equivalent (Ceq) in order to improve low temperature toughness of a welded part. In addition, in order to prevent coarsening of austenite, which is a major cause of the decrease in weld toughness, Japanese Laid-Open Patent Publication No. 55-026164 uses Ti nitride to reduce weld weld toughness by miniaturizing weld structure. Ti nitride is hardly employed because it is near the Fusion Line where the maximum heating temperature of the weld heat affected zone exceeds 1400 ℃. In order to solve this problem, Japanese Patent Laid-Open No. 61-079745 or 62-103344 uses Ti oxide instead of Ti nitride to improve weld toughness.

The prior arts improve the toughness of a weld by suppressing coarsening of the austenite of the weld heat affected zone under high heat input. However, in areas where stability of structures is very important, such as offshore structures, instead of high heat input welding with a large heat affected zone (HAZ), small heat input and low heat input multi-layer welding are applied. Unlike the high heat input welds, the austenitic multi-column welds do not have large coarsening of austenite due to the small heat input. In addition, the toughness of the multi-layer weld is largely dependent on the toughness of the Local Brittle Zone (LBZ), known as the intercrically reheated coarse-grain HAZ (IRCGHAZ). Therefore, it is more important to improve the toughness of the local embrittlement region than to suppress the coarsening of the austenite structure in order to improve the toughness of the multi-layer weld under the heat input.

The present invention is to solve the problems of the prior art, while improving the low-temperature toughness in the most vulnerable local embrittlement area (IRCGHAZ) of the welding heat affected zone while ensuring excellent multi-layer weld toughness, at the same time yield strength of 420MPa or more To provide a high strength high toughness steel having a tensile strength of 520 MPa and a ductile-brittle impact energy transition temperature of -70 ° C. or less, and a method of manufacturing the same.

The present invention for achieving the above object, in weight%, C: 0.06 ~ 0.1%, Si: 0.05 ~ 0.35%, Mn: 1.3 ~ 1.7%, Cu: 0.1 ~ 0.6%, Ni: 0.1 ~ 1.0%, Sol Al: 0.005 to 0.06%, Nb: 0.005 to 0.03%, Ti: 0.005 to 0.03%, N: 0.008% or less, and the remaining Fe and other unavoidable impurities.

Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 The present invention relates to a high strength high toughness steel having excellent multilayer weld zone toughness satisfying the range of 0.36 to 0.42.

In addition, the present invention is a weight%, C: 0.06 ~ 0.1%, Si: 0.05 ~ 0.35%, Mn: 1.3 ~ 1.7%, Cu: 0.1 ~ 0.6%, Ni: 0.1 ~ 1.0%, Sol.Al: 0.005 ~ 0.06%, Nb: 0.005 to 0.03%, Ti: 0.005 to 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities,

Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 Initiating a final hot rolling of the steel slab satisfying the range of 0.36 to 0.42 at an Ar ₃ to austenite recrystallization stop temperature to finish finishing hot rolling at an Ar ₃ or more with a reduction ratio of 40% or more; And

And a step of cooling the finished hot rolled steel sheet to 350 ° C. to 550 ° C. at a cooling rate of 3 ° C./sec or more.

In addition, the present invention is a weight%, C: 0.06 ~ 0.1%, Si: 0.05 ~ 0.35%, Mn: 1.3 ~ 1.7%, Cu: 0.1 ~ 0.6%, Ni: 0.1 ~ 1.0%, Sol.Al: 0.005 ~ 0.06%, Nb: 0.005 to 0.03%, Ti: 0.005 to 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities,

Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 Initiating a final hot rolling of the steel slab satisfying the range of 0.36 to 0.42 at an Ar ₃ to austenite recrystallization stop temperature to finish hot rolling at an Ar ₃ or more with a reduction ratio of 40% or more;

Cooling the finished hot rolled steel sheet to less than 150 to 350 ° C. at a cooling rate of 3 ° C./sec or more; And

It comprises a step of tempering the cooled steel sheet at 500 ~ 600 ℃; relates to a method for producing a high strength high toughness steel excellent in multi-layer weld portion toughness.

Hereinafter, the present invention will be described in detail by dividing it into a steel component and a manufacturing process.

[Steel Ingredients]

C: 0.06 to 0.1% by weight (hereinafter referred to as '%' only)

The C is a very effective element to secure the strength of the steel, it is difficult to secure the strength less than 0.06%, and when added in excess of more than 0.1% reduces the toughness of the base material and greatly deteriorates the toughness of the heat input multilayer welding, It is preferable to limit the content to 0.06 to 0.1%.

Si: 0.05 ~ 0.35%

The Si plays a role of deoxidation of molten steel with Al and needs to contain 0.05% or more in order to secure tensile strength. However, when Si exceeds 0.35% by weight, the structure of the middle heat input multi-layer weld is composed of the acicular ferrite and the upper bainite composite. Since the toughness is greatly reduced by forming a large amount of martensite Austenite Constituent (MA structure) in the heat affected zone at the upper bainite monostructure, it is preferable to limit the content to 0.05 to 0.35%.

Mn: 1.3-1.7%

The Mn is a useful element to secure the strength, if the content is less than 1.3% by weight of the annealed multi-layer welded tissue is changed from the composite structure of the upper bainite and lower bainite to the upper bainite monostructure to reduce the toughness, 1.7 In case of excessive input in excess of%, toughness of the heat input welding part is inhibited. Therefore, the content of Mn is preferably limited to 1.3 ~ 1.7%.

Cu: 0.1 ~ 0.6%

The Cu is an effective component for increasing the strength of the steel without significantly deteriorating the weld toughness. To obtain the effect of the addition, Cu should be added at least 0.1%, but when added in excess of 0.6%, the effect of addition is saturated. Therefore, the Cu content is preferably limited to 0.1 ~ 0.6%.

Ni: 0.1 ~ 1.0%

Ni is an element useful for improving the strength of the base material. If the content is less than 0.1%, the toughness of the heat-induced multi-layer welding heat affected zone decreases, and if it exceeds 1.0%, the Ni increases the manufacturing cost and secures the strength. On the side, the effect is saturated. Therefore, the content of Ni is preferably limited to 0.1 ~ 1.0%.

Sol.Al: 0.005 ~ 0.06%

Sol.Al is an element necessary for deoxidation of steel and is an effective component for increasing the integrity of steel. If the content is less than 0.005%, the above effect cannot be obtained. If the content is more than 0.06%, the toughness of the base material is inhibited and the weld metal is mixed with the weld metal during welding, thereby reducing the toughness of the weld. Therefore, the content of the Sol.Al is preferably limited to 0.005 ~ 0.06%.

Nb: 0.005-0.03%

The Nb is a very effective element for improving the strength by increasing the hardenability of the steel as well as to improve the strength as well as to refine the steel to improve the toughness of the base material, in order to obtain this effect should be added 0.005% or more, but excess addition exceeding 0.03% If so, it is preferable to limit the content to 0.005% to 0.03%, since it forms islands martensite or precipitate hardening in the middle heat input multilayer welding heat affected zone, thereby greatly lowering the heat input toughness.

Ti: 0.005-0.03%

Ti is an effective component for miniaturizing austenite grains by forming TiN in combination with nitrogen and improving toughness. To obtain such effects, it is necessary to add 0.005% or more. However, when the addition exceeds 0.03%, the effect of the addition is saturated, so it is preferable to limit the content to 0.005 to 0.03%.

N: 0.008% or less

N serves to refine the tissue by combining with Ti or Al to form a nitride. However, when contained in excess of 0.008%, the toughness of the base material and the welded portion is greatly reduced by the solid solution N, and the surface quality of the continuously cast slab is reduced. Therefore, it is desirable to limit the content to 0.008% or less.

In addition, the steel of the present invention may include impurities such as P and S that are not completely removed in the steelmaking process, and they are more preferably limited to the following conditions in order to further improve the physical properties of the steel.

P: 0.012% or less, S: 0.005% or less

The P and S are elements that greatly reduce the toughness of the steel, but it is necessary to reduce as much as possible, but it causes a lot of load and cost in the steelmaking process to reduce the content to a very low level. Therefore, P and S do not cause a big problem at 0.012% and 0.005% or less, respectively, so the upper limits are set to 0.012% and 0.005%, respectively.

In the present invention, in addition to the above components, it is possible to add one or more elements selected from Cr and Mo as needed.

Cr: 0.2% or less

The Cr is an element that increases the strength and corrosion resistance of the base material, but when excessively added in excess of 0.2%, Cr may significantly reduce the toughness of the weld heat-affected zone, so the upper limit thereof is preferably limited to 0.2%.

Mo: 0.2% or less

The Mo contributes to improving the strength by increasing the hardenability of the steel, but when added in excess of 0.2%, there is a problem that greatly reduces the toughness of the middle heat input welding heat affected zone, it is preferable to limit the upper limit to 0.2%.

In the present invention, in addition to the composition range, the carbon equivalent (Ceq) of the steel represented by the following formula should be in the range of 0.36 ~ 0.42. The said carbon equivalent is one of the factors which show the weldability of steel.

Ceq = C [%] + Mn [%] / 6 + (Cr [%] + Mo [%]) / 5 + (Cu [%] + Ni [%]) / 15

Here, C [%], Mn [%], Cr [%], Mo [%], Cu [%], and Ni [%] represent the weight percentage of each element.

The results of the inventors of the present invention showed a general tendency that the toughness of the localized embrittlement zone decreases as the carbon equivalent increases in the heat input multilayer welding of 4 kJ / mm or more. However, in the quenched heat multi-layer welding having a heat input amount of 2 kJ / mm or less, it was found that the toughness of the local embrittlement region improved as the carbon equivalent increased.

The test results are shown in FIG. 1. FIG. 1 compares HAZ transition temperatures according to carbon equivalents in quench heat (0.7 kJ / mm) and medium heat (4.5 kJ / mm) using steels satisfying the component range of the present invention. Here, the HAZ transition temperature is the result of the impact test by processing the impact specimen after performing a thermal reproducing test to simulate the IRCGHAZ. In order to simulate IRCGHAZ, it was heated to the maximum heating temperature of 1350 ℃ in the first heat cycle, and then cooled at a cooling rate corresponding to the heat input amount in the temperature range of 800 ℃ to 500 ℃, followed by heating to 800 ℃, the ideal station temperature. Then cooled. After the thermal reproduction specimen was processed into an impact specimen, an impact test was performed at various temperatures.

As can be seen in Figure 1, under the heat input amount of 4.5 kJ / mm, the heat input, the toughness of the local embrittlement area deteriorates as the carbon equivalent increases. However, it can be seen that under 0.7 kJ / mm quenching heat, the toughness of the localized embrittlement region is rather improved as the carbon equivalent is increased.

This phenomenon seems to be due to the change in the microstructure of IRCGHAZ, which is a local embrittlement region of the multilayer weld, according to the heat input. In the heat input multi-layer weld, the multi-wall weld toughness decreases because the microstructure of the local embrittlement zone is changed from the complex structure of the acicular ferrite and the upper bainite to the upper bainite structure containing a lot of island martensite. However, in quench heat multi-layer welding, the toughness is improved because the microstructure of the local embrittlement zone changes from the upper bainite structure to the lower bainite structure with excellent toughness as the carbon equivalent increases.

In the present invention, in order to ensure excellent multi-weld weld toughness at a heat input amount of less than the heat input including hardening heat, it is desirable to ensure that the impact energy transition temperature of the local embrittlement zone is -20 ° C or lower under heat and heat input. Therefore, as shown in FIG. 1, the carbon equivalent should be 0.42 or less for securing the toughness of the heat input multilayer weld, and the carbon equivalent should be 0.36 or more for securing the low temperature toughness under the heat of quenching. It is preferable to limit it to 0.42 or less.

Hereinafter, the manufacturing process of this invention is demonstrated in detail.

[Manufacture process]

In the present invention, a high-strength high toughness steel is manufactured by hot rolling the steel slabs formed as described above and then cooling them.

In the hot rolling, the finish rolling start temperature is preferably limited to Ar ₃ to austenite recrystallization stop temperature (Tnr), because when the finish rolling start temperature is less than Ar _3, the temperature is reached before the start of finish rolling. There is a problem that the air-cooling waiting time greatly increases until the rolling productivity decreases, and when the austenite recrystallization stop temperature is exceeded, sufficient austenite grain refining effect cannot be secured, thereby making it difficult to secure strength and toughness. Therefore, it is preferable that the finish rolling start temperature is set in a range of not less than the abnormal reverse transformation start temperature or less than the austenite recrystallization stop temperature.

In addition, the finish rolling temperature is preferably limited to Ar ₃ or more, because the rolling productivity may be lowered if the finishing rolling temperature is less than Ar ₃ .

When the rolling reduction is less than 40%, the grain size is not sufficiently refined, so that the rough grains generated during reheating do not disappear and the toughness of the base material remains largely impaired. Therefore, it is preferable to limit the amount of reduction during finishing rolling to 40% or more.

Preferably, the cooling rate is limited to 3 ° C./sec or more, because when the cooling rate is less than 3 ° C./sec, polygonal ferrite may be formed at the center of thickness, thereby making it difficult to secure strength.

In addition, when the cooling end temperature of the steel material during the cooling exceeds 550 ℃, it is difficult to secure the strength of the base material, and thus the cooling end temperature is preferably limited to 550 ℃ or less. More preferably, the cooling end temperature is limited to 350 to 550 ° C. This is because, when the cooling end temperature is less than 350 ℃, it is possible to secure the tensile strength, but because the low-temperature structure is excessively generated and yield yield decreases, it is difficult to secure the yield strength. In addition, when the cooling end temperature is less than 350 ℃, more preferably less than 150 ~ 350 ℃ it is possible to ensure the yield strength by tempering treatment in the section of 500 ~ 600 ℃ after the cooling end.

Hereinafter, the present invention will be described in more detail with reference to examples, which are only illustrative of preferred embodiments of the present invention, but the present invention is not limited by the description of these examples.

Example 1

After roughly rolling the steel slabs having the composition shown in Table 1, finish rolling was started at a temperature of Ar ₃ + 30 ° C. for each composition, and finish rolling was finished at a temperature of Ar ₃ + 15 ° C. At this time, the amount of finish rolling reduction was 45%. After the finish rolling was cooled to 450 ℃ at a cooling rate of about 4 ℃ / second, the thickness of the final steel sheet was 76mm. The content of each element described in Table 1 is weight percent.

Steel grade C Si Mn P S Cu Ni Sol.Al Nb Ti N Cr Mo Ceq Inventive Steel A 0.082 0.13 1.57 0.011 0.003 0.27 0.53 0.031 0.012 0.011 0.004 - - 0.397 Inventive Steel B 0.062 0.15 1.51 0.011 0.002 0.27 0.51 0.025 0.013 0.012 0.003 - - 0.366 Invention Steel C 0.096 0.13 1.54 0.011 0.003 0.25 0.58 0.033 0.011 0.009 0.003 - - 0.408 Inventive Steel D 0.084 0.31 1.54 0.008 0.003 0.31 0.59 0.039 0.008 0.013 0.004 - - 0.401 Inventive Steel E 0.083 0.15 1.37 0.009 0.003 0.35 0.58 0.034 0.013 0.015 0.005 - - 0.373 Inventive Steel F 0.079 0.13 1.64 0.011 0.003 0.32 0.54 0.032 0.011 0.005 0.003 - - 0.410 Invention Steel G 0.081 0.16 1.52 0.01 0.003 0.11 0.51 0.034 0.013 0.008 0.003 - - 0.376 Inventive Steel H 0.078 0.16 1.52 0.008 0.003 0.57 0.55 0.031 0.01 0.012 0.004 - - 0.406 Inventive Steel I 0.079 0.16 1.52 0.01 0.003 0.35 0.11 0.031 0.012 0.013 0.003 - - 0.363 Invention Steel J 0.08 0.14 1.53 0.011 0.003 0.3 0.91 0.032 0.011 0.011 0.003 - - 0.416 Inventive Steel K 0.082 0.16 1.54 0.011 0.002 0.27 0.52 0.031 0.01 0.009 0.004 0.06 0.08 0.419 Inventive Steel L 0.081 0.14 1.53 0.011 0.003 0.31 0.54 0.03 0.021 0.01 0.004 - - 0.393 Comparative Steel A 0.091 0.27 1.6 0.011 0.003 0.45 0.67 0.03 0.01 0.011 0.003 - - 0.432 Comparative Steel B 0.087 0.16 1.63 0.01 0.003 0.43 0.62 0.033 0.012 0.009 0.003 - - 0.429 Comparative Steel C 0.072 0.14 1.42 0.009 0.003 0.23 0.43 0.031 0.01 0.012 0.004 - - 0.353 Comparative Steel D 0.067 0.11 1.43 0.01 0.003 0.18 0.38 0.03 0.009 0.01 0.004 - - 0.343 Comparative Steel E 0.051 0.16 1.51 0.008 0.003 0.26 0.47 0.031 0.012 0.013 0.003 - - 0.351 Comparative Steel F 0.12 0.13 1.56 0.009 0.003 0.31 0.55 0.035 0.002 0.007 0.004 - - 0.437 Comparative Steel G 0.081 0.41 1.53 0.008 0.004 0.22 0.54 0.035 0.009 0.008 0.003 - - 0.387 Comparative Steel H 0.078 0.16 1.25 0.011 0.003 0.31 0.56 0.025 0.012 0.012 0.004 - - 0.344 Comparative Steel I 0.083 0.13 1.81 0.009 0.003 0.36 0.45 0.031 0.009 0.011 0.003 - - 0.439 Comparative Steel J 0.085 0.16 1.55 0.008 0.003 - - 0.031 0.012 0.011 0.004 - - 0.343 Comparative Steel K 0.078 0.13 1.5 0.009 0.003 0.68 1.12 0.034 0.011 0.014 0.004 - - 0.448 Comparative Steel L 0.077 0.17 1.46 0.01 0.003 0.31 0.5 0.029 0.011 0.012 0.003 0.24 0.25 0.472 Comparative Steel M 0.079 0.15 1.51 0.01 0.003 0.28 0.51 0.033 - 0.012 0.003 - - 0.383 Comparative Steel N 0.079 0.13 1.57 0.009 0.003 0.28 0.54 0.031 0.033 0.014 0.003 - - 0.395

Physical properties measured from the steel sheet prepared as described above are shown in Table 2 below. In Table 2, the transition temperature refers to the impact energy transition temperature for ductile-brittle fracture. In addition, the reproducible HAZ transition temperature is the result obtained by using the reproducible HAZ test, and the practical welding CTOD value is obtained by performing multi-layer welding on each heat input, and after collecting the welded part, after performing CTOD test according to BS 7448 specified in API RP2Z. Obtained results. The CTOD test was performed at −10 ° C. for grain coarsening area (Coarse Grain HAZ, CGHAZ). In addition, FCAW (Flux Cored Arc Welding) welding was applied to a heat input amount of 0.7 kJ / mm, and Submerged Arc Welding (SAW) welding was applied to a heat input amount of 4.5 kJ / mm.

Steel grade Base material characteristics Reproduction HAZ Transition Temperature (℃) Practical Contact CTOD (mm) Yield strength (MPa) Tensile Strength (MPa) Transition temperature (℃) 0.7kJ / mm 4.5kJ / mm 0.7kJ / mm 4.5kJ / mm Inventive Steel A 443 550 -85 -56 -44 1.21 1.87 Inventive Steel B 425 533 -102 -23 -67 0.92 2.03 Invention Steel C 458 567 -81 -62 -33 1.46 1.74 Inventive Steel D 450 562 -75 -63 -43 - - Inventive Steel E 427 531 -89 -32 -65 - - Inventive Steel F 467 570 -77 -61 -29 - - Invention Steel G 433 535 -88 -33 -64 1.69 0.73 Inventive Steel H 469 573 -75 -68 -30 - - Inventive Steel I 449 575 -75 -21 -68 - - Invention Steel J 472 589 -88 -79 -22 - - Inventive Steel K 476 591 -73 -67 -28 1.31 0.62 Inventive Steel L 462 572 -76 -55 -45 - - Comparative Steel A 483 597 -54 -83 -3 1.24 0.21 Comparative Steel B 471 579 -64 -84 -4 - - Comparative Steel C 411 516 -85 -13 -70 - - Comparative Steel D 405 508 -83 -9 -67 0.32 1.89 Comparative Steel E 389 497 -81 -11 -78 - - Comparative Steel F 485 605 -45 -95 5 - - Comparative Steel G 461 564 -55 -55 -13 - - Comparative Steel H 379 480 -89 -19 -69 - - Comparative Steel I 461 585 -63 -64 2 - - Comparative Steel J 389 503 -67 -21 -77 - - Comparative Steel K 499 623 -58 -76 -11 - - Comparative Steel L 523 645 -19 -11 29 - - Comparative Steel M 404 492 -55 -26 -46 - - Comparative Steel N 469 578 -41 -40 -16 - -

Comparative steels E to N are those in which the carbon equivalent or alloy composition range of the steel is outside the scope of the present invention. When the carbon equivalent is low or when a specific alloy component is added below the lower limit of the present invention, the strength of the base material is insufficient or reproducible. The transition temperature was -20 ° C or higher, indicating that the toughness of the quenched heat multilayer weld was poor. On the contrary, when the carbon equivalent is very high or when the alloying component is added beyond the range of the present invention, the strength characteristics are excellent, but the toughness of the base material or the toughness of the multi-layered welding portion of the base material is inferior to those of the present invention steels. .

Comparative steels A to D are alloy compositions within the range defined by the present invention, but the carbon equivalent range is out of range. When the carbon equivalent exceeds the upper limit of the present invention, it can be seen that the inlet heat reproducing HAZ transition temperature increases. When the carbon equivalent was low, it was found that low temperature toughness was poor under heat of quenching. This trend can also be seen in the utility junction CTOD test. Practical contact CTOD of Comparative Steel A was 0.21 mm under medium heat of 4.5 kJ / mm, which is very low compared to the inventive steel. On the contrary, the practical weld CTOD of comparative steel D having a carbon equivalent of 0.343, which is lower than the lower limit of carbon equivalent of the present invention, showed poor results under heat of quenching.

However, the inventive steels A to L have a high base material strength and toughness characteristics because the chemical composition and the carbon equivalent satisfy the requirements of the present invention, and are compared to the comparative steels under the heat input amounts of 0.7 kJ / mm and 4.5 kJ / mm. Compared with the reproducible HAZ transition temperature and the utility junction CTOD.

Example 2

Rough rolling was carried out using a steel slab having a composition of Inventive Steel A to Inventive Steel L in Table 1, which is a composition according to the conditions of the present invention, followed by finishing rolling and cooling under the conditions described in Table 3 below, to a final thickness of 76 mm. The steel plate of was prepared. Some steel sheets were tempered at 550 ° C. after cooling. Table 4 shows the results of the evaluation of the physical properties of the specimens from the final steel sheet. In Table 3, Tnr indicates austenite recrystallization stop temperature, and Ar ₃ indicates an abnormal reverse transformation start temperature.

division Steel grade Tnr (℃) Ar ₃ (℃) Finish rolling start temperature (℃) Finish rolling end temperature (℃) Finish Rolling Rate (%) Cooling end temperature (℃) Cooling rate (℃ / sec) Tempering treatment Invention 1 Inventive Steel A 906 748 780 765 45 450 4.0 × Invention 2 906 748 880 865 45 445 4.1 × Invention 3 906 748 750 736 45 448 3.9 × Invention 4 906 748 780 767 40 463 3.8 × Invention 5 906 748 780 766 45 359 4.2 × Invention 6 906 748 780 765 45 545 4.1 × Invention Material7 906 748 780 766 45 453 3.2 × Invention Material 8 906 748 780 765 45 215 4.5 ○ Invention Material 9 Inventive Steel B 892 760 790 776 45 453 4.3 × Invention 10 Invention Steel C 908 744 775 759 45 455 4.1 × Invention Material 11 Inventive Steel D 835 746 775 761 45 442 4.2 × Invention Material12 Inventive Steel E 908 760 790 775 45 439 3.9 × Invention Material 13 Inventive Steel F 897 742 770 757 45 467 3.8 × Invention 14 Invention Steel G 897 757 785 772 45 456 4.1 × Invention 15 Inventive Steel H 888 746 775 761 45 436 4.2 × Invention 16 Inventive Steel I 896 775 805 789 45 441 4.1 × Invention 17 Invention Steel J 899 731 760 745 45 465 4.0 × Invention 18 Inventive Steel K 887 744 775 759 45 454 3.8 × Invention Material 19 Inventive Steel L 936 750 780 765 45 453 4.3 × Comparative Material 1 Inventive Steel A 906 748 924 905 45 434 3.9 × Comparative Material 2 906 748 770 754 30 445 4.0 × Comparative Material 3 906 748 787 762 45 577 4.2 × Comparative Material 4 906 748 801 784 45 471 2.1 ×

division Steel grade Yield strength (MPa) Tensile Strength (MPa) Transition temperature (℃) Invention 1 Inventive Steel A 443 550 -85 Invention 2 441 551 -78 Invention 3 447 556 -89 Invention 4 436 539 -77 Invention 5 426 573 -73 Invention 6 423 523 -89 Invention Material7 432 538 -75 Invention Material 8 446 547 -98 Invention Material 9 Inventive Steel B 421 527 -89 Invention 10 Invention Steel C 455 562 -81 Invention Material 11 Inventive Steel D 459 567 -73 Invention Material12 Inventive Steel E 429 543 -86 Invention Material 13 Inventive Steel F 456 573 -75 Invention 14 Invention Steel G 429 538 -88 Invention 15 Inventive Steel H 456 557 -76 Invention 16 Inventive Steel I 425 527 -79 Invention 17 Invention Steel J 462 576 -99 Invention 18 Inventive Steel K 469 582 -74 Invention Material 19 Inventive Steel L 459 577 -76 Comparative Material 1 Inventive Steel A 437 552 -45 Comparative Material 2 430 548 -47 Comparative Material 3 402 498 -74 Comparative Material 4 383 479 -64

Comparative material 1 is a steel material which starts finish rolling at a temperature higher than Tnr, the upper limit of finish rolling start temperature specified in the present invention, and does not secure sufficient austenite grain refining effect during finish rolling. It was found that the impact toughness of the base material was inferior as compared with the present invention at -45 ° C.

Comparative Material 2 is a case where the finish rolling is carried out at a reduction amount lower than 40%, which is the minimum finish rolling reduction prescribed in the present invention, and thus the austenite grains cannot be sufficiently refined, and thus the impact characteristics are not better than that of the present invention.

Comparative material 3 is a steel material which is finished cooling at a temperature higher than the upper limit of the cooling end temperature defined in the present invention as 550 ℃, but the impact toughness of the base material is good, but the strength level of the steel material is very lower than the invention material of the present invention.

Comparative material 4 was not good in strength characteristics compared with the present invention because the cooling rate did not reach the lower limit of the cooling rate specified in the present invention 3 ℃ / sec.

However, inventive materials 1 to 19, since the manufacturing conditions are within the range specified in the present invention, the yield strength and tensile strength of the base material exceeded 420MPa and 520MPa, respectively, and the transition temperature was -70 ° C or less, which was very excellent. . In particular, Inventive Material 8 had a very low cooling end temperature, but the yield strength and impact toughness of the base metal were recovered by cooling after cooling, thereby showing excellent physical properties.

As described above, according to the present invention, the high strength and high toughness of the weld heat-affecting zone toughness, the yield strength of 420 MPa or more, the tensile strength of 520 MPa or more, and the ductile-brittle impact energy transition temperature of the quench heat and middle heat multi-layered welds are -70 ° C. Welded steel can be provided.

Claims (6)

By weight%, C: 0.06 to 0.1%, Si: 0.05 to 0.35%, Mn: 1.3 to 1.7%, Cu: 0.1 to 0.6%, Ni: 0.1 to 1.0%, Sol.Al: 0.005 to 0.06%, Nb: 0.005 ~ 0.03%, Ti: 0.005 ~ 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities, Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 Satisfies the range of 0.36 to 0.42, High strength, high toughness steel with excellent toughness in multi-layer welds with weld heat input in the range of 0.7 ~ 4.5kJ / mm. The high strength and high toughness steel of claim 1, wherein at least one of Cr: 0.2% or less and Mo: 0.2% or less is further added to the steel. By weight%, C: 0.06 to 0.1%, Si: 0.05 to 0.35%, Mn: 1.3 to 1.7%, Cu: 0.1 to 0.6%, Ni: 0.1 to 1.0%, Sol.Al: 0.005 to 0.06%, Nb: 0.005 ~ 0.03%, Ti: 0.005 ~ 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities, Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 Initiating a final hot rolling of the steel slab satisfying the range of 0.36 to 0.42 at an Ar ₃ to austenite recrystallization stop temperature to finish hot rolling at an Ar ₃ or more with a reduction ratio of 40% or more; And Cooling the finished hot-rolled steel sheet to 350 to 550 ° C. at a cooling rate of 3 ° C./sec or more; a high strength high toughness steel having excellent toughness of a multi-layer weld part in a range of 0.7 to 4.5 kJ / mm. Manufacturing method. The method of claim 3, wherein at least one of Cr: 0.2% or less and Mo: 0.2% or less is further added to the steel. By weight%, C: 0.06 to 0.1%, Si: 0.05 to 0.35%, Mn: 1.3 to 1.7%, Cu: 0.1 to 0.6%, Ni: 0.1 to 1.0%, Sol.Al: 0.005 to 0.06%, Nb: 0.005 ~ 0.03%, Ti: 0.005 ~ 0.03%, N: 0.008% or less, consisting of the remaining Fe and other unavoidable impurities, Ceq = Ceq value defined by C [%] + Mn [%] / 6+ (Cr [%] + Mo [%]) / 5+ (Cu [%] + Ni [%]) / 15 Initiating a final hot rolling of the steel slab satisfying the range of 0.36 to 0.42 at an Ar ₃ to austenite recrystallization stop temperature to finish hot rolling at an Ar ₃ or more with a reduction ratio of 40% or more; Cooling the finished hot rolled steel sheet to less than 150 to 350 ° C. at a cooling rate of 3 ° C./sec or more; And Tempering treatment of the cooled steel sheet to 500 ~ 600 ℃; The welded heat input comprising a multi-weld weld toughness excellent in toughness of the multi-layer weld in the range of 0.7 ~ 4.5kJ / mm. 6. The method of claim 5, wherein at least one of Cr: 0.2% or less and Mo: 0.2% or less is further added to the steel.

Images