KR20020040210A - High strength Steel plate to be precipitating TiN and complex oxide of Mg-Ti for welded structure, method for manufacturing the same - Google Patents

Abstract

PURPOSE: A method for manufacturing a steel plate for welded structure is provided which not only minimizes a toughness difference between a base metal and a weld heat-affected zone in the welding heat input ranging from medium heat input to ultra high heat input, but also accomplishes high strength of the base metal using TiN precipitate and Mg-Ti complex oxide. CONSTITUTION: The high strength steel plate for welded structure having TiN precipitate and Mg-Ti complex oxide comprises 0.03 to 0.17 wt.% of C, 0.01 to 0.5 wt.% of Si, 0.4 to 2.0 wt.% of Mn, 0.005 to 0.2 wt.% of Ti, 0.0005 to 0.1 wt.% of Al, 0.008 to 0.030 wt.% of N, 0.0003 to 0.01 wt.% of B, 0.001 to 0.2 wt.% of W, 0.03 wt.% or less of P, 0.03 wt.% or less of S, 0.002 to 0.03 wt.% of O and 0.001 to 0.005 wt.% of Mg satisfying 1.2<=Ti/N<=2.5, 10<=N/B<=40, 2.5<=Al/N<=7, 6.5<=(Ti+2Al+4B)/N<=14, 4<=Ti/O<= 10, 0.2<=Mg/O<=3 and 3<=(Ti+Mg)/O<=12, 30 to 80 wt.% of bainite, and a balance of ferrite which is 20 microns or less. The method for manufacturing the high strength steel plate for welded structure having TiN precipitate and Mg-Ti complex oxide comprises the processes of heating a steel slab comprising 0.03 to 0.17 wt.% of C, 0.01 to 0.5 wt.% of Si, 0.4 to 2.0 wt.% of Mn, 0.005 to 0.2 wt.% of Ti, 0.0005 to 0.1 wt.% of Al, 0.008 to 0.030 wt.% of N, 0.0003 to 0.01 wt.% of B, 0.001 to 0.2 wt.% of W, 0.03 wt.% or less of P, 0.03 wt.% or less of S, 0.002 to 0.03 wt.% of O and 0.001 to 0.005 wt.% of Mg satisfying 1.2<=Ti/N<=2.5, 10<=N/B<=40, 2.5<=Al/N<=7, 6.5<=(Ti+2Al+4B)/N<= 14, 4<=Ti/O<=10, 0.2<=Mg/O<=3 and 3<=(Ti+Mg)/O<=12 to the temperature range of 1100 to 1250 deg.C for 60 to 180 minutes; hot rolling the heated steel slab to a rolling ratio of 40% or more in the Austenite recrystallization zone; and cooling the hot rolled steel slab to the bainite transformation finishing temperature±10 deg.C in a cooling rate of 5 to 20 deg.C/sec.

Description

High strength steel plate to be precipitating TiN and complex oxide of Mg-Ti for welded structure, method for manufacturing the same

The present invention relates to structural steel used in welded structures, such as construction, bridges, shipbuilding, offshore structures, steel pipes, line pipes. More specifically, the present invention relates to a method for producing a welded structural steel material in which the base material is bainite + ferrite and the base material strength is improved by improving the toughness of the weld heat affected zone by using fine TiN precipitate and Ti-Mg composite oxide.

Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel. In order to weld such thick steels, high-efficiency welding is inevitable. As a technique for welding thickened steels, a high-pass heat submerged welding method and an electro-welding method capable of one-pass welding are widely used. In addition, in the field of shipbuilding and bridges, even when welding a steel plate having a plate thickness of 25 mm or more, the above-described high heat input welding method capable of one-pass welding is applied.

In general, in welding, the larger the amount of heat input, the larger the amount of welding, so that the number of welding passes decreases. Therefore, it is advantageous to enable high heat input welding in consideration of welding productivity. In other words, increasing the amount of heat input in the welding will be able to widen the range of use. The range of high heat input currently in use is about 100-200kJ / cm, but in order to weld steel materials with a thicker plate thickness of 50mm or more, it is possible to have a super heat input range of 200-500kJ / cm.

When the heat input is applied to the steel, the heat affected zone formed during welding, particularly the heat affected zone near the fusion boundary, is heated to a temperature close to the melting point by the amount of heat input. Accordingly, since the grains of the weld heat affected zone grow and coarse, and microstructures that are vulnerable to toughness such as upper bainite and martensite are formed during the cooling process, the weld heat affected zone is the site where the toughness of the weld deteriorates most.

Therefore, in order to secure the stability of the welded structure, it is necessary to suppress the growth of the austenite grains in the weld heat affected zone and to keep it fine. As a means to solve this problem, there is disclosed a technique for delaying grain growth of the weld heat affected zone during welding by appropriately distributing an oxide or Ti-based carbonitride, which is stable at a high temperature, to steel materials. For example, Japanese Patent Application Laid-Open No. 11-140582, No. 10-298708, No. 10-298706, No. 9-194990, No. 9-324238, No. 8-60292, (S) 60-245768, (Pyeong) 5-186848, (S) 58-31065, (S) 61-79745, Journal of the Japan Welding Society, Vol. 52, No. 2, 49 and Japanese Patent Laid-Open And 64-15320.

Japanese Patent Laid-Open No. 11-140582 is a representative technique using TiN precipitates, and has a toughness of about 200J at 0 ° C when 100 J / cm of heat input (maximum heating temperature of 1400 ° C) is applied. Is about 300J). In this prior art, Ti / N is substantially managed at 4-12 so that TiN precipitates of 0.05 µm or less are 5.8 × 10 ³ pieces / mm 2 to 8.1 × 10 ⁴ pieces / mm 2, and TiN precipitates of 0.03 to 0.2 μm are 3.9 ×. The toughness of the welded portion is secured by making the ferrite fine by depositing 10 ³ pieces / mm 2 to 6.2 × 10 ⁴ pieces / mm 2. This steel is a microstructure of ferrite and pearlite, which has a mechanical property of 581 MPa in tensile strength and 405 MPa in yield strength.

However, according to this prior art, when the 100 kJ / cm high heat input welding is applied, the toughness of the base material and the heat affected zone is generally low (the base material: 320J, the heat affected zone: 220J at the highest impact toughness of 0 ° C), and also the base material. Since the toughness difference between the and the heat affected zone is about 100J, there is a limit in securing the reliability of the steel structure due to superheated welding of the thickened steel. In addition, as a method for securing the desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or higher and quenched, followed by a reheating process for hot rolling. Becomes In addition, the microstructure of the steel is a composite structure of ferrite and pearlite has a mechanical property of tensile strength up to 581MPa, yield strength up to 405MPa.

Japanese Patent Laid-Open No. 9-194990 manages the ratio of Al and O to 0.3 ≦ Al / O ≦ 1.5 in low nitrogen steel (N ≦ 0.005%), and uses a composite oxide of Al, Mn, and Si. However, when the high heat input welding of about 100 kJ / cm is applied, the weld heat affected zone transition temperature is -50, which is not good at toughness.

In addition, Japanese Patent Application Laid-Open No. 10-298708 uses MgO having a particle size of 0.005-0.1 μm as a nucleus, and has TiN around it, and has a size of 0.05-0.5 μm of MgO-TiN composite precipitate of 1 square mm. Steels containing 1.0 × 10 ^{5 to} 1.0 × 10 ⁷ sugars have been proposed. In this prior art, a steel component system is managed with 0.005-0.025% of Ti, 0.0002-0.005% of Mg, and 0.0005-0.008% of O in steel containing 0.002-0.0008% of N to obtain MgO-TiN composite precipitates. . The steel provided from this prior art improves to some extent the toughness of the weld heat affected zone of the steel by using a composite precipitate of MgO-TiN, but the impact of the weld heat affected zone 0 ° C. when a high heat input welding of about 200 kJ / cm is applied. Toughness is 130J, toughness is not very good. In addition, the tensile strength of the controlled rolled and controlled cooled hot rolled steel sheet (base material) 500MPa, yield strength of 400MPa level of the mechanical properties of the base material is not very good.

Until now, many techniques have been known to improve the toughness of the weld heat affected zone during high heat input welding using TiN precipitates and Al-based or MgO oxides. There have been no reports of achieving a high strength of the base metal at the same time with significant improvements. In particular, there are few technologies in which the toughness of the weld heat affected zone shows an equivalent level to that of the base metal. Therefore, if the above problems can be solved, super heat input welding of the thickened steel is possible, so that it is possible to achieve high efficiency of the welding operation as well as to secure the high-rise of the steel structure and the reliability of the steel structure.

The present invention, using the TiN and Ti-Mg composite oxide in the welding heat input range from the middle heat input to super heat input to minimize the difference in the toughness of the base material and the heat affected zone, and at the same time welding can achieve high strength of the base material It is an object of the present invention to provide a method for manufacturing structural steel.

Welded structural steel materials of the present invention for achieving the above object, by weight% C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1 %, N: 0.008-0.030%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 1.2 ≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10, 0.2≤Mg / O It satisfies <3, 3 <(Ti + Mg) / O <= 12, and consists of 30-80% of bainite and the remainder of 20 micrometers or less of ferrite.

In addition, the manufacturing method of the welded structural steel material of this invention is C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1% by weight. , N: 0.008-0.030%, B: 0.0003-0.01%, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 1.2≤ Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10, 0.2≤Mg / O≤ After slab satisfying 3, 3≤ (Ti + Mg) / O≤12 was heated in the range of 1100-1250 ℃ for 60-180 minutes, hot-rolled in the austenitic recrystallization zone with a rolling ratio of 40% or more, and then the bainite transformation was completed. It is comprised including cooling by the speed | rate of 5-20 degreeC / sec to temperature +/- 10 degreeC.

Hereinafter, the present invention will be described in detail.

In the present invention, the term "prior austenite" refers to austenite formed in the weld heat affected zone when the heat input welding is applied to the steel, and the austenite formed in the manufacturing process of the steel (hot rolling process). Use for convenience to distinguish from.

The present inventors have studied ways to improve the strength of the steel (base material) and at the same time improve the toughness of the weld heat affected zone, and as a result, when the microstructure of the base material is a composite structure of bainite and ferrite, When the TiN precipitate and Mg-Ti composite oxide are used together with the substrate having such a microstructure, the grain size of the austenite of the weld heat affected zone is reduced to about 80 µm or less. The main action of the composite oxide of Mg-Ti promotes the nucleation of needle-like ferrite from fine austenite, and it has been found that the toughness of the weld heat affected zone is remarkably improved.

Starting from this point of view, the inventors of the present invention have been able to derive the following method of improving the toughness of the weld heat affected zone by managing the size of the austenite grains below the critical value in the composite structure of bainite + ferrite.

[1] using TiN precipitate and Ti-Mg composite oxide together,

[2] The initial ferrite grain size of the steel is less than the critical level, and the prior austenite of the weld heat affected zone is refined to 80 µm or less. Also,

[3] Effective use of Ti-Mg complex oxides, BN and AlN precipitates to increase ferrite fraction in weld heat affected zones, and especially to improve the toughness improvement effect by promoting the transformation of polygonal or needle-like ferrite in the austenite . And,

[4] The base material strength is improved by securing the proper fraction of bainite microstructure in the base material through accelerated cooling in the rolling process. These [1] [2] [3] [4] are demonstrated in more detail.

[1] TiN precipitate and Ti-Mg complex oxide management

When heat input welding is applied to structural steels, the weld heat affected zone near the melting line is heated to a high temperature of about 1400 ° C or more, and TiN precipitates precipitated in the base metal are partially dissolved by welding heat or Ostwald ripening. Small precipitates are decomposed and diffuse into larger precipitates, resulting in larger precipitates). Only some precipitates are coarse, and the number of TiN precipitates is significantly reduced, thereby suppressing the inhibitory effect of the growth of austenite grains. .

The inventors have found that this phenomenon is caused by the diffusion of TiN precipitates dispersed in the base metal by the dissolution of the dissolved Ti atoms by the heat of welding, and the characteristics of TiN precipitates according to the ratio of Ti / N are as follows. / N ratio is low), it is found that the dissolved Ti concentration and diffusion rate of the dissolved Ti atoms and the high temperature stability of the TiN precipitates are improved. That is, when the ratio of Ti and N (Ti / N) is in the range of 1.2 to 2.5, the amount of solid solution Ti is extremely reduced and the high temperature stability of the TiN precipitate is increased, so that the fine TiN precipitate of 0.01-0.1 μm or less is 0.5 μm or less. A result of distribution of 1.0 × 10 ⁷ pieces / mm 2 or more at intervals was obtained. This increases the content of nitrogen at the same Ti content, so that all the Ti atoms dissolved are easily combined with the nitrogen atom, and in a high nitrogen environment, the amount of solid solution Ti decreases, making TiN precipitates more stable at higher temperatures than in the case of low nitrogen content. It is analyzed that this is because the solubility product is lowered. In view of the fact that the present invention can promote the aging due to the presence of solid solution N in a high nitrogen environment, the ratio of N / B, Al / N, V / N, and N and Ti + Al + B + (V) Is managed as a whole to precipitate N into BN, AlN, and VN.

Furthermore, in the present invention, the Ti-Mg composite oxide, which is stable at 1400 ° C. or more, together with the fine and uniformly distributed TiN precipitates, has another feature. The Ti-Mg composite oxide is dispersed in the base material to suppress the growth of the austenitic grains of the weld heat affected zone during welding and to promote the needle-like ferrite transformation to improve the toughness of the weld heat affected zone.

[2] ferrite grain size management

According to the study of the present invention, in order to make the size of the old austenite about 80 µm, it is important to make the size of the ferrite 20 µm or less while the microstructure of the base material is a ferrite + bainite composite. . At this time, the refinement of the ferrite is to suppress the growth of the ferrite grains generated in the cold process by using carbides (Fe ₃ C, WC, VC) as well as austenite grain refining by hot working during hot rolling.

[3] microstructure of weld heat affected zone

The facts of the present invention reveal that the toughness of the weld heat affected zone is not only the size of the former austenite grains when the base material is heated above about 1400 ° C., but also the amount and size of the ferrite (more than 70%) precipitated at the old austenite grain boundaries. (Less than 20㎛) And its shape has an important effect. In the present invention, the Mg / O ratio is 0.2-3, the Ti / O ratio is 4-10, and the (Ti + Mg) / O ratio is 3-12. In addition to optimizing the ratio of Mg, using a precipitate of AlN, BN, etc., a large amount of fine ferrite is produced in the mouth of the austenite. Greatly improves.

[4] bainite tissue fraction control

The present inventors can easily control the bainite texture fraction which can improve the strength of the base material when the slab is hot rolled and then the accelerated cooling rate is controlled (5 to 20 ° C./sec). It was confirmed that the physical properties are not related to the microstructure change of the base metal.

Hereinafter, the present invention will be described in detail by dividing the steel component and its manufacturing method.

[Welding Structural Steels]

It is preferable to make content of carbon (C) into 0.03 to 0.17%.

When the content of carbon (C) is less than 0.03%, securing strength as a structural steel is insufficient. In addition, when C exceeds 0.17%, microstructures susceptible to toughness, such as upper bainite, martensite and degenerate pearlite, are transformed during cooling to lower the low temperature impact toughness of structural steel, and also the hardness of the welded portion. Or increasing the strength resulting in deterioration of toughness and generation of weld cracks.

The content of silicon (Si) is preferably limited to 0.01-0.5%.

If the content of silicon is less than 0.01%, the deoxidation effect of molten steel is insufficient during steelmaking and the corrosion resistance of steel is reduced. If the content is more than 0.5%, the effect is saturated, It promotes the transformation of martensite and lowers the low temperature impact toughness.

The content of manganese (Mn) is preferably limited to 0.4-2.0%.

Mn is an effective element that improves deoxidation, weldability, hot workability and strength in steel, and precipitates in the form of MnS around Ti-based oxides, which affects the formation of acicular and polygonal ferrites effective for improving the toughness of weld heat affected zones. Element. Such Mn forms a solid solution to form a solid solution in the matrix to strengthen the matrix to secure the strength and toughness, for this purpose it is preferred to add more than 0.4%. However, when the Mn content exceeds 2.0%, tissue heterogeneity due to Mn segregation has a detrimental effect on the toughness of the weld heat affected zone, rather than the effect of strengthening the solid solution. In addition, when the Mn content is added more than 2.0%, macro segregation and micro segregation occurs depending on the segregation mechanism during steel coagulation, which promotes the formation of a central segregation zone in the center part during rolling, thereby creating a low temperature transformation structure in the center of the base material. .

The content of aluminum (Al) is preferably limited to 0.0005-0.1%.

Al is an element necessary as a deoxidizer and forms an Al oxide by reacting with oxygen to prevent Ti from reacting with oxygen, which not only helps Ti to form fine TiN precipitates, but also is effective for forming fine AlN precipitates in steel. . In order to form fine AlN precipitates, the Al content is preferably made 0.0005% or more. However, if it exceeds 0.1%, AlN is precipitated and the remaining solid Al promotes the formation of Widmanstatten ferrite and phase martensite, which are vulnerable to toughness during the cooling process of the weld heat affected zone. Lowers.

The content of titanium (Ti) is preferably limited to 0.005-0.2%.

Ti is indispensable in the present invention because it combines with N to form a fine TiN precipitate that is stable at high temperature and also reacts with oxygen to form a Ti-Mg composite oxide. In order to obtain such a fine TiN precipitation effect and Ti-Mg complex oxide, it is preferable to add more than 0.005% of Ti, but when it exceeds 0.2%, coarse TiN crystals and coarse Ti-Mg oxides are formed in molten steel. It is not preferable because the crystallization of the crystal is present in the base metal and does not suppress the growth of the former austenite grains in the weld heat affected zone during welding of the steel.

The content of boron (boron, B) is preferably limited to 0.0003-0.01%.

B is a very effective element for producing polygonal ferrite at grain boundaries as well as acicular ferrite having excellent toughness in grains. B forms a BN precipitate, which hinders the growth of the old austenite grains and forms Fe carbide in the grain boundary and in the mouth to promote ferrite transformation of acicular and polygons having excellent toughness. If the B content is less than 0.0003%, such an effect cannot be expected, and if it exceeds 0.01%, the hardenability increases, which may cause hardening of the weld heat affected zone and low temperature cracking.

The content of nitrogen (N) is preferably limited to 0.008-0.03%.

N is an indispensable element for forming TiN, AlN, BN, VN, NbN, etc., and it is possible to minimize the growth of the old austenite grains in the weld heat affected zone during the high heat input welding and to increase TiN, AlN, BN, VN, NbN, etc. Increase the amount of precipitate; Particularly, the content of TiN and AlN precipitates and the spacing of precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, and the high temperature stability of the precipitates themselves are significantly affected. However, when the nitrogen content exceeds 0.03%, the effect is saturated, and toughness decreases due to an increase in the amount of solid solution nitrogen distributed in the weld heat affected zone, and it may be incorporated into the weld metal due to dilution during welding, resulting in a decrease in the toughness of the weld metal. Can be.

The content of tungsten (W) is preferably limited to 0.001-0.2%.

Tungsten is an element that uniformly precipitates tungsten carbide (WC) in the base material after hot rolling, thereby effectively suppressing ferrite grain growth after ferrite transformation, and also suppressing the growth of the initial austenite grains in the initial heating of the weld heat affected zone. If the content is less than 0.001%, there is less distribution of tungsten carbide for suppressing ferrite grain growth upon cooling after hot rolling, and the effect is saturated when more than 0.2% is added.

The content of phosphorus (P) and sulfur (S) is preferably limited to 0.030% or less.

P is preferably as low as possible because it is an impurity element that promotes central segregation during rolling and hot cracking during welding. In order to improve the toughness of the base metal, the toughness of the weld heat affected zone, and to reduce the center segregation, it is recommended to manage it to 0.03% or less.

Since S forms a low melting point compound such as FeS when present in a large amount, it is preferable to manage S as low as possible. In order to reduce the base material toughness, weld heat affected zone toughness and central segregation, it is recommended that the S content be 0.03% or less. In particular, sulfur is precipitated in the form of MnS around Ti-based oxides, which affects the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of welding heat affected zones. It is desirable to limit it to 0.03% or less.

The content of oxygen (O) is preferably limited to 0.0020-0.03%.

O is an indispensable element for forming Ti-Mg complex oxide by reacting with Ti in molten steel. Ti-Mg complex oxide promotes the transformation of acicular ferrite during the ferrite transformation from the former austenite in the weld heat affected zone. If the O content is less than 0.0020%, the number of Ti-Mg composite oxides is not preferable, and if it is more than 0.03%, coarse Ti-Mg composite oxides and other oxides such as FeO are formed, which is preferable for the toughness of the weld heat affected zone. Not.

The content of magnesium (Mg) is preferably limited to 0.001-0.005%.

Mg is a useful element that combines with Ti and O in molten steel to form Ti-Mg composite oxides to promote needle ferrite transformation in the weld heat affected zone. Mg content of 0.001% or more is required to form fine Ti-Mg composite oxides. It is preferable to add. However, when the Mg content exceeds 0.005%, coarse Mg oxide is formed in molten steel, which adversely affects the mechanical properties of the base metal.

The ratio of Ti / N is preferably 1.2 to 2.5.

In the present invention, the Ti / N ratio is lowered to 2.5 or less, which has two advantages. First, it is possible to increase the amount of TiN, that is, the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all Ti atoms that are dissolved in the cooling process during the playing process are combined with the nitrogen atoms to increase the fine TiN precipitation. Second, TiN is stable at high temperatures. That is, since the solubility product which shows the stability of the precipitate at high temperature such as the weld heat affected zone becomes smaller, the precipitate such as high nitrogen TiN is more stable than the case where the nitrogen content is low. On the other hand, if the Ti / N ratio is higher than 2.5, coarse TiN is crystallized in molten steel, which is a steelmaking process, and a uniform distribution of TiN is not obtained. Also, excess Ti remaining without precipitation as TiN remains in a solid solution to weld heat. Affects bad toughness. If the Ti / N ratio is less than 1.2, the amount of solid solution nitrogen in the base metal increases, which is detrimental to the toughness of the weld heat-oriented part.

It is preferable to make ratio of N / B into 10-40.

In the present invention, if the N / B ratio is less than 10, the precipitation amount of BN that promotes the ferrite transformation of polygons at the old austenite grain boundary during the post-weld cooling process is insufficient, and when the N / B ratio is over 40, the effect is saturated and dissolved. This is because the amount of nitrogen is increased to lower the toughness of the weld heat affected zone.

It is preferable to make Al / N ratio into 2.5-7.

In the present invention, when the Al / N ratio is less than 2.5, AlN precipitates for inducing needle-like ferrite transformation are insufficient, and the amount of solid solution nitrogen in the weld heat affected zone may increase, resulting in a weld crack, and an Al / N ratio of more than 7 In the case the effect is saturated.

It is preferable that ratio of (Ti + 2Al + 4B) / N is 6.5-14.

In the present invention, when the ratio of (Ti + 2Al + 4B) / N is less than 6.5, the growth inhibition of the austenite grain growth of the weld heat affected zone, the generation of fine polygonal ferrite at the grain boundary, the amount of solid solution nitrogen, the needle shape and the polygonal shape in the grain boundary Insufficient size and number of distribution of TiN, AlN, BN, and VN precipitates for ferrite formation and control of tissue fraction, and the effect is saturated when (Ti + 2Al + 4B) / N is more than 14 seconds. If V is added, the ratio of (Ti + 2Al + 4B + V) / N is preferably 7-17.

The ratio of Ti / O is preferably set to 4-10.

If the Ti / O ratio is less than 4, the number of TiO oxides required to inhibit austenite grain growth is insufficient, and the Ti ratio in TiO oxides becomes small, thus losing the function of acicular ferrite nucleation sites and improving the toughness of the weld heat affected zone. The effective acicular ferrite phase fraction is lowered. If the ratio of Ti / O is more than 10, the effect of inhibiting the weld heat affected zone austenite grain size crystallization is saturated, and the ratio of components such as Mn contained in the oxide becomes rather small, thus losing the function of the nucleation site of the ferrite in the mouth.

The Mg / O ratio is preferably limited to 0.2-3.

If the Mg / O ratio is less than 0.2, the number of Ti-Mg complex oxides required for austenite grain growth inhibition is insufficient, and the Mg ratio contained in the oxide becomes small, and thus loses its function as a needle-like ferrite nucleation site. The needle ferrite phase fraction effective for improvement is lowered. In addition, when the Mg / O ratio exceeds 3, the austenite grain growth inhibitory effect is saturated, and the ratio of components such as Mn contained in the oxide becomes smaller, thus losing the function of nucleation sites of the ferrite in the mouth.

The (Ti + Mg) / O ratio is preferably limited to 3-12.

When the (Ti + Mg) / O ratio is less than 3, the number of Ti-Mg complex oxides required for the austenite grain growth inhibition is insufficient, and thus the toughness of the weld heat affected zone is improved by losing its function as a needle ferrite nucleation site in the austenite grain. The effective acicular ferrite phase fraction decreases. In addition, when the (Ti-Mg) / O ratio exceeds 12, the austenite grain growth inhibitory effect is saturated, and the ratio of the Mn component contained in the oxide becomes rather small, thus losing the function of the nucleation site of the ferrite in the mouth.

In order to further improve the toughness of the steel material (base material) and the heat-affected portion formed as described above, V is further added.

The content of vanadium (V) is preferably limited to 0.01-0.2%.

V is an element that combines with N to form VN to promote ferrite formation in the weld heat affected zone, and VN precipitates alone or precipitates in TIN precipitates to promote ferrite transformation. In addition, V combines with C to form VC, which acts to inhibit ferrite grain growth after ferrite transformation. When the V content is less than 0.01%, it is difficult to obtain the ferrite transformation promoting effect in the weld heat affected zone because the VN deposition amount is small. On the other hand, exceeding 0.2% is not preferable because it causes toughness of the base metal and the weld heat affected zone (HAZ) and improves the weld hardenability, which may cause the low temperature crack of the weld.

Moreover, it is preferable to make ratio of V / N into 0.3-9. In the present invention, when the V / N ratio is less than 0.3, it is difficult to secure an appropriate number and size of VN precipitates deposited and distributed at the TiN + MnS composite precipitate boundary for improving the toughness of the weld heat affected zone. If the V / N ratio exceeds 9, the size of the VN precipitates deposited at the TiN + MnS precipitate boundary is coarsened, and the frequency of VN precipitation precipitated at the TiN + MnS composite precipitate boundary is reduced, which is effective for the toughness of the weld heat affected zone. Reduce ferrite phase percentage.

In the present invention, in order to further improve the mechanical properties in the steel composition as described above, one or more selected from the group of Ni, Cu, Nb, Mo, Cr is further added.

The content of nickel (Ni) is preferably limited to 0.1-3.0%.

Ni is an effective element which improves the strength and toughness of the base material by solid solution strengthening. In order to achieve this effect, the Ni content is preferably 0.1% or more, but when the content exceeds 3.0%, the hardenability is increased to reduce the toughness of the weld heat affected zone and the possibility of high temperature cracking in the weld heat affected zone and the weld metal. This is not desirable because there is.

The content of copper (Cu) is preferably limited to 0.1-1.5%.

Cu is an element which is effective to secure the base material strength and toughness due to solid solution at the base. For this purpose, Cu content should be contained more than 0.1%, but if it exceeds 1.5%, it is not preferable because it increases the hardenability by increasing the hardenability in the weld heat affected zone and promotes high temperature crack in the weld heat affected zone and the weld metal. . In particular, Cu is an element that affects the formation of acicular and polygonal ferrites, which are effective in improving the toughness of the welded heat affected zone by depositing CuS around Ti-based oxides with sulfur, so that the content is 0.3-1.5%. desirable.

In addition, in the case of complex addition of Cu and Ni, the total sum thereof is preferably less than 3.5%. The reason is that less than 3.5% of the hardenability increases, which adversely affects the weld heat affected zone toughness and weldability.

The content of niobium (Nb) is preferably limited to 0.01-0.10%.

Nb is an effective element from the viewpoint of securing the strength of the base material. For this purpose, Nb is added in an amount of 0.01% or more. However, Nb is undesirable because it causes coarse precipitation of coarse NbC, which is detrimental to the toughness of the base material.

Chromium (Cr) is preferably made 0.05 to 1.0%.

Cr increases the hardenability and also improves the strength. If the content is less than 0.05%, the strength cannot be obtained and when the content exceeds 1.0%, the base metal and the HAZ toughness deteriorate.

Molybdenum (Mo) is preferably 0.05-1.0%.

Mo is also an element that increases the hardenability and improves the strength. The content thereof is 0.05% or more for securing the strength, but the upper limit is set to 1.0% like Cr for suppressing the HAZ hardening and the welding low temperature crack.

In addition, in the present invention, one or two kinds of Ca and REM are further added to suppress the grain growth of the austenite at the time of heating.

Ca and REM form an oxide having excellent high temperature stability, thereby suppressing the growth of the austenite grains when heated in the base metal and improving the toughness of the weld heat affected zone. In addition, Ca has the effect of controlling the coarse MnS shape during steelmaking. To this end, it is preferable to add more than 0.0005% of calcium (Ca) and more than 0.005% of REM, but if Ca exceeds 0.005% of REM of more than 0.05%, large inclusions and clusters are generated to harm the cleanliness of the steel. As REM, 1 type, or 2 or more types, such as Ce, La, Y, and Hf, may be used, and any of the above effects can be obtained.

· Microstructure of steel

In the present invention, the steel is a composite structure of ferrite + bainite, the tissue fraction of bainite is preferably in the range of 30-80%. The reason is that less than 30% is difficult to secure the appropriate base material strength for showing the effect of the present invention, and if it is more than 80% it is difficult to secure the base material toughness.

In the composite structure of ferrite + bainite, the ferrite grain size is preferably 20 µm or less. This is because when the grain size of the ferrite is larger than 20 μm, the austenite grain size of the weld heat affected zone becomes 80 μm or more during high heat input welding, which is detrimental to the weld heat affected zone toughness.

Precipitate

The former austenite grains of the weld heat affected zone are greatly influenced by the size, number and distribution of oxides or nitrides distributed in the base material when the austenite grain size of the base material is constant. In addition, since 30-40% of the nitrides distributed in the base material are welded to the base material at the time of high heat input welding (above the heating temperature of 1400 ℃ or more), the effect of inhibiting the growth of the austenite grains in the weld heat affected zone is reduced. There is a need for a uniform distribution of further nitrides taking into account the re-used nitrides. In order to suppress the growth of the old austenite in the weld heat affected zone, it is important to uniformly distribute the fine TiN precipitate to suppress the Ostwald ripening phenomenon in which some precipitates are coarsened. For this purpose, the TiN precipitates should be controlled to 0.5 μm or less to make the TiN distribution uniform.

In addition, it is preferable to limit the particle diameter and the critical number of TiN to 0.01-0.1 μm and 1.0 × 10 ⁷ or more per 1 mm ² . The reason for this is that less than 0.01 μm is easily re-used to the base metal during high heat input welding, and the effect of inhibiting the growth of the old austenite grains is insufficient. When the thickness exceeds 0.1 μm, pinning of the old austenite grains This is because the growth inhibition effect is reduced and behaves like coarse non-metallic inclusions, which adversely affects mechanical properties. In addition, when the number of precipitates is less than 1.0 × 10 ⁷ per 1 mm ² , it is difficult to control the size of the old austenite grains of the weld heat affected zone at the time of welding higher than the heat input to be 80 μm or less, which is a threshold value.

·oxide

In the present invention, the particle diameter and the critical number of the Ti-Mg oxide are preferably in the range of 0.5-2.0 μm and 1.0 × ^{10 2} -1.0 × 10 ³ per 1 mm ² . The reason for this is that when the oxide size is less than 0.5 µm, the needle ferrite nucleation promoting effect is insufficient in the weld heat-affected grains. This is because the amount of transformation of the acicular ferrite is reduced.

The reason for limiting the number of oxide having an adverse effect on toughness because less acicular ferrite amount because there is less number of the oxide of the acicular ferrite nucleation in the welding heat affected portion below 1.0x10 ² per 1mm ² also 1mm ² per 1.0x10 Above ³ , the number of Ti-Mg complex oxides is not preferable because the amount of acicular ferrite in the grains decreases because the primary acicular ferrites nucleating in the oxides grow and collide with each other to limit the formation of the acicular acicular ferrites.

[Manufacturing method]

Refining (Deoxidation, Degassing) Process

In general, the steel refining process is first refined in the refining furnace (electric furnace, electric furnace), and then the molten steel of the refining furnace is pulled out by the ladle (LF) to the secondary refining (external refining). After out-refining, degassing (RH process) is performed. Normal deoxidation takes place in primary and secondary refining.

The present inventors focused on the fact that the amount of dissolved oxygen in the deoxidation process greatly affects the formation behavior of the oxide to uniformly distribute a large amount of Ti into an appropriate amount of Ti-Mg composite oxide and a large amount of fine TiN precipitate. The following facts could be found.

(1) The amount of dissolved oxygen greatly affects the formation behavior of oxides, and in order to generate a large amount of oxides, an appropriate amount of dissolved oxygen exists, which is about 50-200 ppm,

(2) When Mg is first introduced into molten steel having such dissolved oxygen, and then Ti is added, Mg oxide has a significantly higher affinity with oxygen than other oxides, and thus a finer oxide distribution can be obtained. Some form Ti-Mg complex oxides to increase the number of oxides, and these Ti-Mg complex oxides increase the precipitation frequency of precipitates such as TiN, MnS, CuS, VN, and the rest of the added Ti in molten steel. The solid solution forms a fine Ti precipitate upon solidification, which will be described in detail.

In the present invention, before the addition of Mg and Ti, molten steel is added to the molten steel with a higher deoxidizing element than Ti to adjust the dissolved oxygen amount to 50-200 ppm. The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

The deoxidizer with greater affinity with oxygen has a higher rate of binding to oxygen in molten steel and a stronger binding force. Therefore, before adding Zr and Ti, deoxidizer with greater deoxidizing power than Ti should be added to secure proper dissolved oxygen (50-200ppm) to make Ti into TiN precipitate while forming the target TiO oxide while reducing the progress time of deoxidation process. Can be. That is, since the deoxidizer with greater deoxidizing power than Ti is pre-injected, the ratio of Ti remaining in solid solution rather than coarse primary oxide in molten steel increases considerably, thus securing fine amounts of Ti oxide and TiN nitride during solidification. It can be.

In the present invention, before adding an element (Al, REM, Zr, Ca, Mg) having a higher deoxidizing power than Ti, by deoxidizing by adding elements such as Mn, Si, and then deoxidizer having a higher deoxidizing power than Ti such as Al It is also preferable to inject. This method has the advantage of adjusting the amount of oxygen of molten steel through the separation of coarse primary oxide by adjusting Al and adding Al and adding Al by adding weak Al.

As an example of the deoxidation operation pattern in the present invention, Mn, Si, Al are added in the order of primary refining, followed by Mg in the secondary refining, and finally Ti is added. In the present invention, the final Al content is limited to the range of 0.005-0.1%. The reason is that the oxygen content in the molten steel is controlled by floating separation by reacting with the oxygen in the molten steel by Al injecting after the first deoxidation in the above content range. Because it is effective to. As a result of experiment to derive the proper range of Al, when Al is less than 0.005%, the influence on the control of the proper dissolved oxygen amount is not sufficient, and when it exceeds 0.1%, the amount of oxygen in the molten steel is insufficient to form Ti-Mg complex oxide. This is because the amount decreases.

As described above, in the present invention, it is preferable to deoxidize the dissolved oxygen amount to 50-200 ppm before the introduction of Ti, for the following reason. If the amount of dissolved oxygen is less than 50 ppm, the amount of oxygen in the molten steel is too small to form an appropriate amount of Ti-Mg-based oxide for showing the effect of the present invention, and in the case of more than 200 ppm, the effect is saturated and not only the target element This is because there is a problem of oxidizing another alloying element to obtain an alloy component system composed of accurate chemical components.

As described above, the amount of dissolved oxygen is controlled by using Si, Mn, and Al elements, and then Mg and Ti are added as deoxidation elements. Since Mg has a high affinity with oxygen in molten steel, Mg reacts with oxygen present in molten steel to form Mg oxide. In addition, the Mg oxide has a high affinity with Ti, which has a direct effect on the formation of Ti-Mg composite oxide, which has a significant effect on increasing the number of Ti-based oxides distributed in the steel during solidification. When Ti and Mg are added at the same time, Mg has a greater affinity for oxygen than Ti and preferentially combines with oxygen in molten steel to form Mg oxides rather than Ti oxides. Therefore, since there is a problem of greatly reducing the number of Ti oxides acting as a polygonal or needle-like ferrite transformation nucleus, the ferrite nucleation is influenced by first distributing the Mg oxide in the deoxidation process by Mg and then deoxidizing by Ti. It is advantageous to increase the number of Ti oxides which affect the

In the present invention, Mg is added in the range of 0.001-0.005% as the tertiary deoxidation element, because the Mg content is less than 0.001%, the number of Mg oxides formed in molten steel is too small to form Ti-Mg complex oxides. This is because if the affinity between Mg and oxygen is greater than Ti oxide, the amount of oxygen that can be combined with Ti is insufficient at 0.005% or more, so that the number of TiO oxides as well as Ti-Mg composite oxides for showing the effect of the present invention is reduced. .

As described above, 0.005-0.2% Ti is added in the molten steel state after the addition of Mg deoxidation element. Rather than forming Ti oxide in the molten steel state, the solid solution in most molten steel and forming Ti oxide during solidification has an effect of increasing the number and finer Ti oxide. Therefore, when Ti is added after the addition of Mg deoxidation element, the ratio remaining as solid solution Ti is considerably higher than that produced as an oxide, and thus a fine amount of Ti oxide and TiN precipitate can be secured during solidification. At this time, the preferred amount of Ti is in the range of 0.005-0.2%, which is in consideration of MnS, CuS, BN, and VN precipitation behaviors effective for improving the toughness of the weld heat affected zone as well as the proper distribution of Ti oxide and TiN precipitates for showing the effects of the present invention. to be.

As described above, Mg is added followed by Ti, followed by refining, for example, degassing, wherein the holding time of the molten steel until continuous casting is preferably limited to within 0.5-20 minutes. If the retention time of molten steel is less than 0.5 minutes, there is no time for coarse primary oxides to float and remain. After coagulation, the average size of Ti-based oxides is increased. Because this is very high. Therefore, when Ti element is added within 0.5-20 minutes after Mg element is added, a fine Ti-Mg composite oxide can be obtained during solidification.

Casting process

In the present invention, the molten steel refined as described above is continuously cast into slabs. Continuous casting is preferable from the viewpoint of productivity improvement by casting at low speed and giving a weak cooling condition in the secondary cooling zone in consideration of the high possibility of occurrence of cast surface cracks in high nitrogen steel. Cooling conditions in the secondary cooling zone are important factors affecting the refinement and uniform distribution of TiN precipitates.

According to the study of the present invention, the continuous casting speed is less than 1.2 m / min at low speed, more preferably about 0.9 to 1.2 m / min. The reason is that when the casting speed is less than 0.9m / min, it is advantageous for cast surface cracks, but the productivity is lowered.

Hot rolling process

In the present invention, the slab is heated at 1100-1250 ° C. for 60-180 minutes. Below 1100 ° C., the number of TiN precipitates is small because the rate of diffusion of solute atoms is small. If it exceeds 1250 ° C., Ti-based precipitates are coarsened or decomposed, and thus the number of precipitates decreases. I can't. On the other hand, if the heating time is less than 60 minutes, it is not preferable because there is no segregation reduction effect of the solute atoms and there is not enough time for the solute atoms to diffuse to form precipitates. In addition, when the heating time exceeds 180 minutes, coarsening of austenite grain size occurs, which is not preferable in terms of work productivity.

After heating as above, it is preferable to hot-roll at a rolling ratio of 40% or more at the austenite recrystallization zone temperature. The austenite recrystallization zone temperature is affected by the emphasis and the previous reduction amount, and the austenite recrystallization zone temperature is in the range of about 1050 to 850 ° C. in consideration of the usual reduction amount in the emphasis of the present invention. In this section, a rolling ratio of at least 40% should be given. If the rolling ratio is less than 40%, the ferrite nucleation site in the austenite grain is insufficient and the effect of refining the ferrite grains due to austenite recrystallization is insufficient. It affects the precipitate behavior which effectively affects the toughness of the affected zone.

In hot rolling, the austenite grain size is affected by the temperature, time in the reheating furnace, and the rolling amount. The grain size of the austenite affects the quenchability, so that the desired bainite fraction can be obtained by controlling it. In order to increase the bainite fraction, it is recommended to set the austenite grain size to 10 μm or more.However, if the austenite grain size becomes larger than 50 μm, the hardenability becomes too large during transformation, which may cause martensite transformation. .

In the present invention, the reason for limiting the cooling rate in the range of 5-20 ° C / sec to the bainite transformation end temperature ± 10 ° C after hot rolling is as follows. The phase transformation of the present invention steel occurs in a section within the bainite transformation end temperature ± 10 ℃, the cooling rate must be controlled up to this section. If the accelerated cooling rate is less than 5 ° C / sec it is difficult to secure the bainite phase fraction for showing the effect of the present invention, and in the case of more than 20 ° C / sec martensite phase ratio increases to be harmful to the base material toughness.

Casting of the steel in the present invention can be produced by slab by continuous casting or mold casting. In this case, if the cooling rate is high, it is advantageous to finely disperse the precipitate, so continuous casting having a high cooling rate is preferable. For the same reason, slabs are advantageously thinner. In addition, the slab may be immersed according to the present invention, and then hot charge rolling and direct rolling may be applied according to the user's use in the hot rolling process. Technology can be applied. In addition, heat treatment may be applied to improve the mechanical properties of the hot rolled sheet produced according to the present invention. However, even if the well-known techniques are applied to the present invention, it is natural that they are interpreted to be substantially within the technical scope of the present invention as a simple change of the present invention.

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

EXAMPLE

Steel grades having the composition as shown in Table 1 were prepared in the slab by the continuous casting method by dissolving them in a converter, and the composition ratio between alloying elements by steel type to show the effect of the present invention is shown in Table 3. The solidification rate, slab heating temperature, heating time, rolling start temperature and end temperature, rolling reduction, and cooling rate of the rolled material having a thickness of 25 to 40 mm in the rolling process are shown in Tables 2 and 4. At this time, the rolling reduction ratio of all the steel grades was 60% or more.

The test pieces for evaluating the mechanical properties of the base metal from the hot rolled plates as described above were taken at the center of the plate thickness of the rolled material, the tensile test piece was taken in the rolling direction, and the Charpy impact specimen was taken in the direction perpendicular to the rolling direction. It was.

Tensile test piece was used KS standard (KS B 0801) No. 4 test piece and the tensile test was tested at the cross head speed (5mm / mim). The impact test piece was manufactured according to KS (KS B 0809) No. 3 test piece, and the notch direction was processed on the side of the rolling direction (L-T) in the case of the base material and in the welding line direction on the welding material. In addition, in order to investigate the austenite grain size according to the maximum heating temperature of the welding heat affected zone, it is heated to 140 ℃ / sec condition for 1 second after the heating up to the maximum heating temperature (1200 ~ 1400 ℃) by using the simulation welding simulator (simulator). After holding, it was quenched with He gas. The quenched specimens were ground and corroded to determine the austenite grain size at the highest heating temperature condition by KS (KS D 0205).

The size, number, and spacing of precipitates and oxides, which have a significant effect on the microstructure analysis and the toughness of the weld affected zone after cooling, were measured by the point counting method using an image analyzer and an electron microscope. At this time, the test surface was evaluated based on 100 mm ² .

Impact toughness evaluation of the weld heat affected zone is 800-500 ℃ cooling after heating the welding conditions equivalent to about 80 kJ / cm, 150 kJ / cm, 250 kJ / cm, that is, the maximum heating temperature to 1400 ℃ After the welding heat cycles of 60 seconds, 120 seconds, and 180 seconds were applied, the surface of the test piece was polished, processed into an impact test piece, and evaluated through a Charpy impact test at -40 ° C.

Chemical composition (% by weight) C Si Mn P S Al Ti B (ppm) N (ppm) W Cu Ni Cr Mo Nb V Ca Mg REM O (ppm) Inventive Steel 1 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7 115 0.005 - - - - - 0.01 - 0.003 - 32 Inventive Steel 2 0.07 0.12 1.50 0.006 0.005 0.07 0.05 10 275 0.002 - 0.2 - - - 0.01 - 0.002 - 54 Invention Steel 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015 3 112 0.003 0.1 - - - - 0.02 - 0.002 - 28 Inventive Steel 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02 5 80 0.001 - - - - - 0.05 - 0.001 - 31 Inventive Steel 5 0.08 0.15 1.52 0.006 0.004 0.09 0.05 15 300 0.002 0.1 - 0.1 - - 0.05 - 0.002 - 52 Inventive Steel 6 0.10 0.14 1.50 0.007 0.005 0.025 0.02 10 100 0.004 - - - 0.1 - 0.09 - 0.001 - 32 Inventive Steel 7 0.13 0.14 1.48 0.007 0.005 0.04 0.015 8 115 0.15 0.1 - - - - 0.02 - 0.001 - 30 Inventive Steel 8 0.11 0.15 1.52 0.007 0.005 0.06 0.018 10 120 0.001 - - - - 0.015 0.01 - 0.002 - 44 Inventive Steel 9 0.13 0.21 1.50 0.007 0.005 0.025 0.02 4 90 0.002 - - 0.1 - - 0.02 0.001 0.002 - 43 Inventive Steel 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025 6 100 0.05 - 0.3 - - 0.01 0.02 - 0.002 0.01 52 Inventive Steel 11 0.09 0.12 1.48 0.006 0.003 0.048 0.019 10 130 0.01 - 0.2 - - - - - 0.001 - 32 Conventional Steel 1 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 - - - - - - - - - - 22 Conventional Steel 2 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 - - - - - - - - - - 32 Conventional Steel 3 0.13 0.24 1.44 0.012 0.003 0.0044 0.010 1.2 127 - 0.3 - - - 0.05 - - - - 138 Conventional Steel 4 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 - - - 0.14 0.15 - 0.028 - 0.0022 - 25 Conventional Steel 5 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 - 0.75 0.58 0.24 0.14 0.015 0.037 - 0.0024 - 27 Conventional Steel 6 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 - 0.35 1.15 0.53 0.49 0.001 0.045 - 0.0019 - 25 Conventional Steel 7 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 - 0.3 - - - 0.036 - - - - - Conventional Steel 8 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 - 0.32 0.35 - - 0.013 - - - - - Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 - - - 0.21 0.19 0.025 0.035 - - - - Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 - - 1.09 0.51 0.36 0.021 0.021 - - - - Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 - - - - - - 0.069 - - - - Conventional steels (1, 2, 3) are invention steels (5, 32, 55) of Japanese Patent Application Laid-Open No. 9-194990. Conventional steels (4, 5, 6) are Japanese Patent Application Laid-open Nos. Invented steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invented steels (48, 58, 60, 61) of Japanese Patent Application Laid-Open No. 8-60292. Is invention steel F of JP-A-11-140582

Steel grade used division Primary deoxidation sequence Amount of dissolved oxygen after Al deoxidation (ppm) Mg addition after Al deoxidation (%) Ti content after Mg addition (%) Molten steel holding time (min) Oxygen content after coagulation (ppm) Casting condition Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 68 0.003 0.014 15 30 1.1 0.35 Invention 2 Mn → Si 72 0.003 0.014 15 32 1.1 0.35 Invention 3 Mn → Si 56 0.003 0.014 15 34 1.1 0.35 Comparative Material 1 Mn → Si 35 0.003 0.014 15 9 1.1 0.35 Comparative Material 2 Mn → Si 264 0.003 0.014 15 134 1.1 0.35 Inventive Steel 2 Invention 4 Mn → Si 64 0.002 0.05 15 54 1.2 0.30 Invention Steel 3 Invention 5 Mn → Si 62 0.002 0.015 13 28 1.2 0.30 Inventive Steel 4 Invention 6 Mn → Si 62 0.001 0.02 14 31 1.2 0.35 Inventive Steel 5 Invention 7 Mn → Si 59 0.002 0.05 18 52 1.2 0.30 Inventive Steel 6 Invention Material 8 Mn → Si 58 0.001 0.02 15 32 1.2 0.35 Inventive Steel 7 Invention Material 9 Mn → Si 62 0.001 0.015 14 30 1.2 0.30 Inventive Steel 8 Invention 10 Mn → Si 73 0.002 0.018 16 44 1.1 0.32 Inventive Steel 9 Invention 11 Mn → Si 82 0.002 0.02 15 43 1.1 0.35 Inventive Steel 10 Invention Material12 Mn → Si 69 0.002 0.025 15 52 1.2 0.35 Inventive Steel 11 Invention Material 13 Mn → Si 62 0.001 0.019 15 32 1.1 0.30 Conventional steel (1-11) is not specifically described its manufacturing conditions

Alloy element composition ratio for showing the effect of the present invention Ti / O Mg / O (Ti + Mg) / O Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Invention 1 4.7 1.0 5.7 1.2 16.4 3.5 0.9 9.3 Invention 2 4.4 0.9 5.3 1.2 16.4 3.5 0.9 9.3 Invention 3 4.0 0.9 5.0 1.2 16.4 3.5 0.9 9.3 Comparative Material 1 15.6 3.3 18.9 1.2 16.4 3.5 0.9 9.3 Comparative Material 2 1.0 0.2 1.3 1.2 16.4 3.5 0.9 9.3 Invention 4 9.3 0.4 13.0 1.8 27.5 2.5 0.4 7.4 Invention 5 5.4 0.7 9.6 1.3 37.3 5.4 1.8 14.0 Invention 6 6.5 0.3 6.1 2.5 16.0 2.5 6.3 14.0 Invention 7 9.6 0.4 6.8 1.7 20.0 3.0 1.7 9.5 Invention Material 8 6.3 0.3 10.0 2.0 10.0 2.5 9.0 16.4 Invention Material 9 5.0 0.3 5.3 1.3 14.4 3.5 1.7 10.3 Invention 10 4.1 0.5 4.5 1.5 12.0 5.0 0.8 12.7 Invention 11 4.7 0.5 5.1 2.2 22.5 2.8 2.2 10.2 Invention Material12 4.8 0.4 5.2 2.5 16.7 4.5 2.0 13.7 Invention Material 13 5.9 0.3 6.3 1.5 13.0 3.6 - 9.0 Conventional Steel 1 0.4 - 0.4 4.1 13.8 0.6 - 5.7 Conventional Steel 2 3.8 - 3.8 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.7 - 0.7 0.8 105.8 0.4 - 1.5 Conventional Steel 4 5.2 0.9 6.08 4.1 2.5 0.8 8.8 24.5 Conventional Steel 5 4.8 0.9 5.7 6.5 0.25 10.5 18.5 47 Conventional Steel 6 3.6 0.8 4.36 3.2 0.39 0.4 16.1 21.6 Conventional Steel 7 - - - 1.0 9.9 2.5 - 6.5 Conventional Steel 8 - - - 1.2 14.3 0.4 - 2.2 Conventional Steel 9 - - - 0.8 9.1 2.1 3.9 9.2 Conventional Steel 10 - - - 0.6 9.5 3.2 1.5 8.9 Conventional Steel 11 - - - 5.5 12.7 3.4 7.8 20.3

Steel grade used division Heating temperature (℃) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount / accumulated loading amount at recrystallization area (%) Cooling rate (℃ / min) Cooling end temperature (℃) Invention 2 Inventive Example 1 1150 150 1030 790 65/80 15 550 Inventive Example 2 1200 130 1040 790 65/80 15 550 Inventive Example 3 1240 90 1040 790 65/80 15 550 Comparative Example 1 950 40 1040 790 65/80 15 550 Comparative Example 2 1350 250 1035 790 65/80 15 550 Invention 1 Inventive Example 4 1200 130 1020 790 65/80 16 550 Invention 3 Inventive Example 5 1200 130 1040 790 65/80 16 550 Comparative Material 1 Comparative Example 3 1210 120 1030 780 65/80 0.1 Room temperature Comparative Material 2 Comparative Example 4 1210 120 1030 790 65/80 35 Room temperature Invention 4 Inventive Example 6 1180 150 1020 780 60/80 17 550 Invention 5 Inventive Example 7 1190 140 1010 800 60/80 18 550 Invention 6 Inventive Example 8 1220 110 1010 810 60/75 17 550 Invention 7 Inventive Example 9 1220 110 1020 800 60/75 11 550 Invention Material 8 Inventive Example 10 1210 120 1010 790 60/75 10 550 Invention Material 9 Inventive Example 11 1240 100 1000 780 55/70 9 550 Invention 10 Inventive Example 12 1210 120 1010 790 55/70 19 550 Invention 11 Inventive Example 13 1190 100 1000 800 55/70 18 550 Invention Material12 Inventive Example 14 1220 110 1020 780 55/70 12 550 Invention Material 13 Inventive Example 15 1180 150 1020 780 65/75 12 550 Conventional Steel 11 1200 - Ar ₃ or higher 960 80 Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

division Precipitate properties Oxide properties Base material texture characteristics Count (pcs / mm ² ) Average size (㎛) Average interval (㎛) Count (pcs / mm ² ) Average size (㎛) Thickness (mm) Yield strength (MPa) Tensile Strength (MPa) Elongation (%) FGS (μm) Ferrite Percentage (%) Bainite fraction (%) Impact toughness at -40 ° C (J) Inventive Example 1 2.4 X ^{10 8} 0.016 0.25 2.3 X 10 ² 1.8 25 494 653 38 11 66 32 358 Inventive Example 2 3.2 X ^{10 8} 0.017 0.24 2.1X10 ² 1.4 25 495 651 39 9 63 35 362 Inventive Example 3 2.5 X ^{10 8} 0.012 0.26 2.3 X 10 ² 1.2 25 496 650 39 10 63 38 357 Comparative Example 1 2.3 X 10 ⁶ 0.174 1.6 2.7 X 10 ¹ 2.2 25 493 654 26 16 50 40 106 Comparative Example 2 3.4 X 10 ⁶ 0.165 1.8 2.6 X 10 ¹ 2.3 25 492 660 17 17 31 32 45 Inventive Example 4 3.2 X ^{10 8} 0.025 0.32 3.2 X 10 ² 1.2 30 496 658 38 11 63 32 349 Inventive Example 5 2.6 X ^{10 8} 0.013 0.34 2.6 X 10 ² 1.2 30 496 662 38 10 63 30 354 Comparative Example 3 1.3 X 10 ⁶ 0.182 1.2 1.5X10 ¹ 3.4 30 484 664 28 21 63 32 220 Comparative Example 4 4.3X10 ⁶ 0.177 1.4 1.2X10 ¹ 3.7 30 492 682 15 24 20 42 208 Inventive Example 6 3.3 X ^{10 8} 0.026 0.35 1.8 X 10 ² 1.3 30 490 663 38 10 62 32 364 Inventive Example 7 4.6 X ^{10 8} 0.024 0.32 3.2 X 10 ² 1.2 35 490 664 39 10 65 34 360 Inventive Example 8 4.3X10 ⁸ 0.014 0.40 2.1X10 ² 1.6 35 492 642 36 11 62 32 365 Inventive Example 9 5.6 X ^{10 8} 0.028 0.29 2.4 X 10 ² 1.5 35 491 636 37 10 64 31 359 Inventive Example 10 5.2 X ^{10 8} 0.021 0.28 2.3 X 10 ² 1.5 35 494 666 36 10 63 29 375 Inventive Example 11 3.7 X ^{10 8} 0.029 0.25 2.7 X 10 ² 1.7 40 490 666 37 12 63 28 364 Inventive Example 12 3.2 X ^{10 8} 0.025 0.31 1.7 X 10 ² 1.5 40 496 642 38 11 65 30 356 Inventive Example 13 3.3 X ^{10 8} 0.042 0.34 1.8 X 10 ² 1.5 40 406 664 38 12 62 26 348 Inventive Example 14 3.6X10 ⁸ 0032 0.28 2.2 X 10 ² 1.6 40 387 650 37 10 63 31 349 Inventive Example 15 4.2 X ^{10 8} 0.018 0.26 1.9 X 10 ² 1.7 30 489 649 39 9 66 28 368 Conventional Steel 1 35 406 436 Conventional Steel 2 35 405 441 Conventional Steel 3 25 629 681 Conventional Steel 4 Precipitates of MgO-TiN 3.03 × 10 ⁶ pcs / mm ² 40 472 609 Conventional Steel 5 Precipitates of MgO-TiN 4.07 × 10 ⁶ pcs / mm ² 40 494 622 Conventional Steel 6 Precipitates of MgO-TiN 2.80 × 10 ⁶ pcs / mm ² 50 812 912 Conventional Steel 7 25 629 681 Conventional Steel 8 50 504 601 Conventional Steel 9 60 526 648 Conventional Steel 10 60 760 829 Conventional Steel 11 50 401 514

As shown in Table 5, the number of precipitates (Ti-based nitride) of the hot rolled material produced by the present invention has a range of 2.4X10 ⁸ / mm ² or more, whereas in the case of conventional steel 4.07 X10 ⁶ / mm ² Since the following ranges show that the invention material has a fairly uniform and fine precipitate size compared to the conventional material, the number is also significantly increased. In addition, the present invention shows a number of oxides (Ti-Mg) of about 1.7x10 ² ~ 3.2x10 ³ ranges and the average size is about 1.2-1.8㎛. Meanwhile, in the base metal structure of the present invention steel, the present invention steel is composed of bainite and fine ferrite.

division Austenitic grain size of welding heat affected zone (㎛) Microstructure with welding heat effect of 100kJ / cm heat input Reproduction Weld Heat Affected Zone -40 ℃ Impact Toughness (J) (Maximum Heating Temperature: 1400 ℃) 1200 (℃) 1300 (℃) 1400 (℃) Ferrite Percentage (%) Ferrite Average Grain Size (㎛) Δt _800-500 = 60 seconds Δt _800-500 = 120 seconds Δt _800-500 = 180 seconds Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Impact Toughness (J) Transition temperature (℃) Inventive Example 1 23 34 56 74 15 372 -74 332 -67 293 -63 Inventive Example 2 22 35 55 77 13 384 -76 350 -69 302 -64 Inventive Example 3 23 35 56 75 13 366 -72 330 -67 295 -63 Comparative Example 1 54 86 182 38 24 124 -43 43 -34 28 -28 Comparative Example 2 65 92 198 36 26 102 -40 30 -32 17 -25 Inventive Example 4 25 38 63 76 14 353 -71 328 -68 284 -65 Inventive Example 5 26 41 57 78 15 365 -71 334 -67 295 -62 Comparative Example 3 56 80 178 40 26 108 -39 56 -32 24 -24 Comparative Example 4 63 88 184 39 28 64 -28 39 -30 10 -21 Inventive Example 6 25 32 53 75 14 383 -73 354 -69 303 -63 Inventive Example 7 24 35 55 77 14 365 -71 337 -67 292 -63 Inventive Example 8 27 37 53 74 13 362 -71 339 -67 296 -62 Inventive Example 9 24 36 52 78 15 368 -72 330 -67 284 -63 Inventive Example 10 22 34 53 75 14 383 -72 345 -66 293 -63 Inventive Example 11 26 35 64 75 14 356 -71 328 -68 282 -68 Inventive Example 12 27 39 64 74 15 353 -71 321 -67 276 -62 Inventive Example 13 23 38 68 74 14 354 -71 320 -67 254 -62 Inventive Example 14 25 35 64 70 15 342 -71 326 -67 248 -63 Inventive Example 15 23 36 53 76 16 349 -72 332 -68 293 -94 Conventional Steel 1 -58 Conventional Steel 2 -55 Conventional Steel 3 -54 Conventional Steel 4 230 93 132 (0 ℃) Conventional Steel 5 180 87 129 (0 ℃) Conventional Steel 6 250 47 60 (0 degrees Celsius) Conventional Steel 7 -60 -61 Conventional Steel 8 -59 -48 Conventional Steel 9 -54 -42 Conventional Steel 10 -57 -45 Conventional Steel 11 219 (0 ℃)

Table 6 shows the properties of the weld heat affected zone of the present invention steel and conventional steel. The austenitic grain size at the maximum heating temperature of 1400 ° C. such as the welding heat affected zone shows that the present invention has a range of 52-64 μm, while the conventional material has a very coarse range of about 180 μm or more. Can be. Therefore, in the present invention, it can be seen that the austenite grain suppression effect of the weld heat affected zone during welding is very excellent. Further, in the case of the present invention, the heat input portion of the heat-affected portion is composed of about 70% or more at a welding heat input of 100 kJ / cm.

As described above, the present invention can provide a structural steel for welding that can secure excellent high heat input welding heat affected zone toughness simultaneously by using Ti-Mg composite oxide together with TiN precipitates with high strength base material properties. will be.

Claims (9)

By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1%, N: 0.008-0.030%, B: 0.0003-0.01 %, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 1.2≤Ti / N≤2.5, 10≤N / B≤40 , 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10, 0.2≤Mg / O≤3, 3≤ (Ti + Mg) / O≤12 A high strength welded structural steel having a TiN precipitate composed of 30 to 80% of bainite and remaining 20 µm or less of ferrite and a composite oxide of Mg-Ti. The steel material according to claim 1, wherein V is contained in an amount of 0.01 to 0.2%, a ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B + V) / N≤17. High strength welded structural steel having a composite oxide of TiN precipitate and Mg-Ti characterized in that it satisfies. The steel material according to claim 1 or 2, wherein the steel comprises Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. High strength welding with a composite oxide of TiN precipitate and Mg-Ti, characterized in that one or two or more selected, and one or two selected from the group Ca: 0.0005-0.005%, Rem: 0.005 to 0.05% Structural steels. The TiN precipitate according to claim 1 or 2, wherein the TiN precipitate having a size of 0.01-0.1 µm is distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 µm or less, and the Ti-Mg composite having 0.5-2.0 µm. A steel for high strength welded structure having a TiN precipitate and a composite oxide of Mg-Ti, wherein the oxide is distributed 1 × 10 ² -1 × 10 ³ pcs / mm ² . By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.0005-0.1%, N: 0.008-0.030%, B: 0.0003-0.01 %, W: 0.001-0.2%, P: 0.03% or less, S: 0.03% or less, O: 0.002-0.03%, Mg: 0.001-0.005%, 1.2≤Ti / N≤2.5, 10≤N / B≤40 , 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 4≤Ti / O≤10, 0.2≤Mg / O≤3, 3≤ (Ti + Mg) / O≤12 After slab satisfying the heat treatment was heated for 60-180 minutes in the range of 1100-1250 ℃, hot-rolled at a rolling ratio of 40% or more in the austenitic recrystallization zone, and then the rate of 5-20 ℃ / sec until the bainite transformation temperature ± 10 ℃ A method for producing a high strength welded structural steel having a TiN precipitate and a composite oxide of Mg-Ti comprising cooling with a furnace. 6. The steel material according to claim 5, wherein V is contained in an amount of 0.01 to 0.2%, a ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B + V) / N≤17. A method for producing a high strength welded structural steel having a composite oxide of TiN precipitate and Mg-Ti, characterized in that it satisfies. The steel material according to claim 5 or 6, wherein the steel comprises Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%. High strength welding with a composite oxide of TiN precipitate and Mg-Ti, characterized in that one or two or more selected, and one or two selected from the group Ca: 0.0005-0.005%, REM: 0.005 to 0.05% Method of manufacturing structural steels. According to claim 5, The slab is deoxidized element having a greater deoxidizing power than Ti is added before the introduction of Ti to deoxidize the dissolved oxygen amount to 50 ~ 200ppm, Mg is added to make the amount of Mg 0.005-0.1% A method of manufacturing a high strength welded structural steel having a TiN precipitate and a composite oxide of Mg-Ti, characterized in that the casting is carried out by degassing after adding Ti to 0.005 to 0.2% by adding Ti. The method of claim 5 or 8, wherein the continuous casting of the slab is characterized in that the continuous cooling of molten steel at a rate of 0.9 ~ 1.2m / min is weakly cooled in a non-amount of 0.3 ~ 0.35ℓ / kg in the secondary cooling zone A method for producing a high strength welded steel having a composite oxide of TiN precipitate and Mg-Ti.