KR20020007468A - Steel plate having superior toughness in weld heat-affected zone and method for manufacturing the same, welding fabric using the same - Google Patents

Abstract

PURPOSE: A method of manufacturing a steel plate for welding fabric is provided to give superior toughness in weld heat-affected zone of the steel plate when welded by promoting transformation of acicular structure ferrite. CONSTITUTION: The steel plate is manufactured by the following step: (i) heat treatment of a steel slab that is composed of C 0.03-0.17wt.%, Si 0.01-0.5wt.%, Mn 0.4-2.0wt.%, Ti 0.005-0.2wt.%, Al 0.03-0.3wt.%, N 0.008-0.030wt.%, B 0.0003-0.01wt.%, W 0.001-0.2wt.%, P 0.03wt.% or less, O 0.00-0.03wt.%, the balance of Fe and inevitable impurities and that further satisfies below constituents ratio: 1.2<=Ti/N<=2.5, 10<=N/B<=40, 2.5<=Al/N<=7, 6.5<=(Ti+2Al+4B)/N<=14, 8<=Al/O<=22, 250<=Mn/O<=530 in the temperature range 1100 to 1250deg.C for 60-180min; (ii) hot rolling the steel slab at a reduction ratio higher than 40% during austenitic recrystallization; (iii) cooling the hot rolled steel slab at a cooling rate higher than 1deg.C/min until temperature of the steel slab reach a temperature of ferrite transformation temperature±10deg.C. In addition, the steel slab is manufactured by casting molten steel at 0.9-1.1m/min and then slowly cooling in 2nd cooling zone by spraying water 0.3-0.35L/kg. When the steel plate is welded, prior austenite smaller than 80μm is formed in heat affected zone, and acicular/polygonal structure ferrite smaller than 20μm occupies at least 70% in the heat affected zone when quenched.

Description

Steel plate having superior toughness in weld heat-affected zone and method for manufacturing the same, welding fabric using 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 welded structural steel material and a method of manufacturing the same, by using fine TiN precipitates and Al ₂ O ₃ · MnO composite oxides to improve the toughness of the weld heat affected zone.

Recently, the steel used in accordance with the trend of high-rise building, structure has been replaced by thick steel.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 used corresponds to approximately 100-200 kJ / cm, but in order to weld more thickened steel, that is, steel with a plate thickness of 50 mm or more, it is possible to have a super heat input range of 200-500 kJ / 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.

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

Japanese Patent Laid-Open No. 9-194990 discloses a composite oxide composed of Al, Mn, and Si by controlling the ratio of Al and O to 0.3 ≦ Al / O ≦ 1.5 in low nitrogen steel (N ≦ 0.005%). The ferrite nucleation function in the austenite mouth in the heat affected zone and the transition temperature of the heat affected zone (maximum heating temperature of 1450 ℃, cooling cycle of 800 to 500 ℃ for 60 seconds) is -50 to -60. Welded structural steels are presented. This steel has a rather low transition temperature. In this prior art, high nitrogen steels use less than 0.005% of low nitrogen steels because they deteriorate toughness.

To date, many techniques for improving the toughness of the weld heat affected zone during high heat input welding have been known. However, there have been no reports of a significant improvement in the toughness of the weld heat affected zone during long heat input welding maintained at 1350 ° C or longer. In particular, there are few technologies in which the toughness of the weld heat affected zone shows the same level as 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 structure of the steel structure and the reliability of the steel structure.

The present invention provides a welded structural steel having excellent toughness in the weld heat affected zone and a method of manufacturing the same, and the heat effect of the structural steel while the toughness of the base material and the heat affected zone is minimized in the weld heat input range from medium heat to super heat input. It is an object of the present invention to provide a welded structure using subsidiary properties.

Welded structural steels 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: 0.03-0.3 %, 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%, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 8≤Al / O≤22, 250≤Mn / O≤530, and the rest It is composed of Fe and other impurities, and microstructure is composed of a composite structure of ferrite and pearlite of 20 μm or less.

In addition, the manufacturing method of the welded structural steel material of this invention is C: 0.03-0.17% by weight, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.03-0.3% , 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%, 0.5≤Ti / N≤3.0, 5 Steel slab satisfying ≤N / B≤50, 2≤Al / N≤10, 6.5≤ (Ti + 2Al + 4B) / N≤14 and composed of the remaining Fe and other impurities is 60 in the range of 1100-1250 ° C. After heating for -180 minutes, hot rolling is carried out in the austenite recrystallization zone with a rolling ratio of 40% or more, and then cooling to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min or more.

In addition, the welded structure of the present invention has the above-described composition and microstructure, wherein TiN precipitates of 0.01-0.1 μm are distributed at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less, and Al ₂ of 0.5-2.0 μm. The high heat input welding is applied to steel materials in which O ₃ · MnO oxide is distributed 1 × 10 ² -1 × 10 ³ pcs / mm ² to produce prior austenite of 80 μm or less in the weld heat affected zone. Subsequently, it is quenched and the microstructure of the weld heat affected zone is made of a ferrite of 20 µm or less with a phase fraction of 70% or more.

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 deeply investigated the growth behavior of prior austenite in the welding heat affected zone and the phase transformation during cooling when the heat input welding is applied to steel materials (base metal), and thus the critical old austenite grain size (about 80 µm) On the basis of this, it was confirmed that the toughness of the welded heat affected zone was remarkably changed and that the formation of fine ferrite in the austenite mouth did not cause a problem of the welded heat affected zone.

Based on these findings, in the present invention,

[1] While increasing the TiN high temperature stability while finely and uniformly distributing the precipitate of TiN,

[2] With the initial ferrite grain size of the steel below the critical level, when the high heat input welding is applied, the former austenite of the heat affected zone is refined to 80 μm or less. Also,

[3] The complex distribution of Al ₂ O ₃ · MnO is finely precipitated to improve the toughness of the heat affected zone by producing polygonal ferrite and acicular ferrite in the mouth of the austenite during cooling of the heat affected zone. These [1] [2] [3] are demonstrated in more detail.

[1] TiN precipitate 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 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. Increasing the nitrogen content at the same Ti content, all of the Ti atoms dissolved are easily combined with the nitrogen atom, and in the high nitrogen environment, the amount of solid solution Ti decreases, so that the TiN precipitates are stabilized at a higher temperature than when the nitrogen content is low. 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.

[2] ferrite grain size management

According to the present invention, it can be seen that it is important to make the size of the ferrite to 20 μm or less while the microstructure of the base material is a composite structure of ferrite + pearlite in order to make the size of the old austenite to 80 μm. At this time, the refinement of the ferrite may be obtained by controlling the growth of the ferrite grains generated in the cooling process after hot rolling as well as the austenite grain refinement by the steel working during hot rolling. For this purpose, it was confirmed that it is very effective to properly deposit and distribute carbides (VC, WC) effective for ferrite grain growth.

[3] oxide management

In the present invention, the Al / O ratio is 8 or more and the Mn / O ratio is controlled to 250 to 530 to precipitate a stable composite oxide (Al ₂ O ₃ · MnO) at high temperature, which is fine ferrite in the mouth of the austenite. It plays a major role in generating. Most of these ferrites are polygonal ferrites and acicular ferrites, which greatly improve the toughness of the heat affected zone.

Japanese Unexamined Patent Publication (Kokai) No. 9-194990 describes that when the Al / O ratio exceeds 1.5, the ratio of oxides of Mn and Si decreases so that the oxide does not act as a ferrite precipitation nucleus in the mouth. This difference is considered to be due to the fact that the complex oxide (Al-Mn-Si) of the prior art is different from that of the present invention (Al ₂ O ₃ · MnO).

Hereinafter, the present invention will be described in detail by dividing the steel 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 increase in 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 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 element that reacts with Al oxide to form Al-Mn complex oxide (Al ₂ O ₃ · MnO). Al-Mn complex oxide (Al ₂ O ₃ · MnO) plays an important role in promoting intragranular needle-like ferrite transformation to improve toughness in the heat affected zone. Mn is also an effective element for improving deoxidation, weldability, hot workability and strength in steel. Mn forms a substituted solid solution in the matrix structure to strengthen the matrix to secure strength and toughness. If the amount of Mn is less than 0.4%, the number of Al-Mn complex oxides (Al ₂ O ₃ · MnO) is small, so that Mn content of 0.4% or more is required. If Mn content exceeds 2.0%, coarse Mn oxide is used. It is not preferable because it forms. In addition, when the Mn content is more than 2.0%, tissue heterogeneity due to segregation adversely affects the weld heat affected zone toughness.

The content of aluminum (Al) is preferably limited to 0.03-0.3%.

Al is an indispensable element for forming fine AlN precipitates in steel as well as the necessary element because it forms Al composite oxide by reacting with oxygen as a deoxidizer. For this purpose, it is preferable to contain Al more than 0.03%, but when it exceeds 0.3%, coarse Al oxide is formed, and solid solution Al is Wiedmanstatten ferrite and the phase which is vulnerable to toughness during cooling of the weld heat affected zone. It promotes the production of martensite and lowers the toughness of the heat input welding heat affected zone.

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

Ti is an indispensable element in the present invention because it combines with N to form a fine TiN precipitate that is stable at high temperatures. It is preferable to add more than 0.005% of Ti in order to obtain such a fine TiN precipitation effect, but when it exceeds 0.2%, coarse TiN precipitates and Ti oxides are formed in molten steel, and thus it is impossible to suppress the growth of the austenite grains of the weld heat affected zone. Because it is not desirable.

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 precipitates. In particular, since the TiN and AlN precipitates have a remarkable effect on the size, precipitate spacing, precipitate distribution, complex precipitation frequency with oxide, and high temperature stability of the precipitate itself, the content is preferably set at 0.008% or more. 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 in the base material as tungsten carbide (WC) after hot rolling, effectively inhibiting ferrite grain growth after ferrite transformation, and also suppressing the growth of the initial austenite grains during 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. Particularly, in the case of sulfur, it precipitates 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 the amount from 0.003% to 0.03% or less.

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

O is an indispensable element for reacting with Al and Mn in molten steel to form a composite oxide of Al and Mn. Al ₂ O ₃ · MnO composite oxides promote the transformation of acicular ferrite during ferrite transformation in the austenite in the heat affected zone. If the O content is less than 0.0020%, it is not preferable because the number of Al ₂ O ₃ · MnO composite oxides formed is small, and if it exceeds 0.03%, it is not preferable because it forms coarse Al oxide and other oxides such as FeO.

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 the Ti atoms dissolved in the cooling process during the playing process combine with the nitrogen atom, thereby increasing 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.

It is preferable to make Al / O ratio into 8 <= Al / O <= 22.

If the Al / O ratio is less than 8, the number of Al ₂ O ₃ · MnO (Galaxite) oxides required for suppressing the growth of the austenite grains is insufficient, and the Al ratio in the oxides becomes small, resulting in loss of function as a ferrite nucleation site in the mouth. The intragranular acicular ferrite phase fraction effective for improving the toughness of the weld heat affected zone is lowered. In addition, when the Al / O ratio is more than 22, the former austenite grain growth inhibition effect is saturated, and the ratio of Al component contained in the oxide becomes smaller, which is undesirable because it loses the function of nucleation sites of the ferrite in the mouth.

It is preferable to make Mn / O ratio into 250 <= Mn / O <= 530 range.

If the Mn / O ratio is less than 250, the number of Al ₂ O ₃ · MnO (Galaxite) oxides required for suppressing the growth of the old austenite grains is insufficient, and the Mn ratio contained in the oxides becomes small, thus losing the function of the ferrite nucleation site in the mouth. As a result, the intragranular acicular ferrite phase fraction effective for improving the toughness of the weld heat affected zone is lowered. In addition, when the Mn / O ratio exceeds 530, the former austenite grain growth inhibitory effect is saturated, and the ratio of Mn component contained in the oxide becomes smaller, which is not preferable because it loses the function of 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 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 thus the VN precipitation frequency deposited at the TiN + MnS complex 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 when 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, after the hot rolling, the microstructure of the steel is ferrite + pearlite, and the size of the ferrite grains is preferably 20 μm or less. The reason is that when the grain size of the ferrite is larger than 20 μm, the size of the old austenite grains in 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.

In addition, the higher the percentage of ferrite in the composite structure of ferrite + perlite, the higher the toughness and elongation of the base metal, and the ferrite is 20% or more and most preferably 70% or more.

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 less than or equal to the threshold of 80 μm.

·oxide

According to the present invention, when the ratio of Al / O and Mn / O is controlled, an Al-based composite oxide is formed, and when they have a range of 1.0x10 ² -1.0x10 ³ per 1 mm ^{2 with} a particle diameter in the range of 0.5-2.0 μm, the heat affected zone As the fine ferrite phase fraction increases, acicular ferrite can be obtained. When the Al-based oxide is less than 0.5 μm, the needle ferrite nucleation promoting effect is insufficient in the weld heat-affected grains. When the Al-based oxide exceeds 2.0 μm, the needle-ferrite grain growth inhibition effect is insufficient and nucleated. The amount of transformation is reduced. In addition, when the number of composite oxides is 1.0x10 ² or less per 1 mm ² , the needle ferrite is insufficient in transformation amount, and when the number of composite oxides is 1.0x10 ³ or more, there is a risk of causing cracks due to dislocation accumulation at the interface between the composite oxide and the matrix. This is undesirable because it reduces the toughness of the base metal.

[Method of manufacturing welded structural steel]

Casting process

In the present invention, the molten steel composed of the chemical component system is continuously cast through a conventional refining process to make a slab. 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 1.1 m / min, more preferably about 0.9 to 1.1 m / min, which is lower than the usual casting speed of about 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. When it is faster than 1.1m / min, cast surface cracks are more likely to occur.

In the secondary cooling zone, the specific water amount is preferably as low as possible, that is, 0.3 to 0.35 l / kg. When the specific amount is less than 0.3 L / kg, it is difficult to control the proper size and number of TiN to show the effect of the present invention due to coarsening of TiN precipitates. In addition, when the specific water content exceeds 0.35L / kg, the precipitation frequency of the TiN precipitates is small, and it is difficult to control the number, size, and the like of TiN precipitates for showing the effects of the present invention.

Hot rolling process

In the present invention, the slab is heated at 1100-1250 ° C. for 60-180 minutes. It is a problem because the number of TiN precipitates is small because the rate of diffusion of solute atoms is less than 1100 ℃, Ti precipitates, etc. are coarse or decomposed when it exceeds 1250 ℃, the precipitates decrease the number of precipitates Not desirable 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.

After rolling as above, it is cooled at a rate of 1 ° C./min or more to the ferrite transformation end temperature ± 10 ° C., preferably at a rate of 1 ° C./min to the ferrite transformation end temperature, and thereafter air-cooled. Of course, cooling at a rate of 1 ° C./min to room temperature is not a problem in terms of ferrite miniaturization, but it is not economical, and is not preferable. have. When the cooling rate is less than 1 ° C / min, it causes a grain growth of the recrystallized fine ferrite, it is difficult to secure the ferrite grain size of the steel sheet 20㎛ or less.

[Welding Structure]

The welded steel material provided according to the present invention is a composite structure of ferrite (70% or more) + pearlite, and the size of the ferrite grains is 20 µm or less. In addition, TiN precipitates have a size of 0.01-0.1 μm, 1.0 × 10 ⁷ or more per 1 mm ² , and an interval of 0.5 μm or less, with Al-based composite oxides being finely distributed.

When the high heat input welding is applied to such steels, the old austenite grain size is 80 μm or less. If the austenite grain size is greater than 80 µm, low-temperature structure (martensite or upper bainite) is easily formed due to the increase in hardenability, which is detrimental to the weld heat affected zone toughness. In addition, even if ferrites having different nucleation sites are produced at the old austenite grain boundaries, ferrite coalesces during grain growth, which adversely affects toughness. Therefore, the austenite grain size of the weld heat affected zone for showing the effect of the present invention must be controlled to be 80 μm or less.

When the high heat input welding is applied and quenched as described above, the ferrite of 20 micrometers or less of the microstructure size of a heat-affected part has a phase ratio of 70% or more. If the grain size of the ferrite is larger than 20 μm, the ferrite fraction of the side plate or allotriomorphs harmful to the weld heat affected zone toughness increases. When the ferrite of the present invention has the characteristics of polygonal ferrite and acicular ferrite, it is more advantageous for toughness, which is an Al-based composite oxide according to the present invention.

In the hot rolling process of the present invention, well-known hot charge rolling and direct rolling may be applied according to a user's use, and various techniques such as control rolling and control cooling may be applied. In addition, in order to improve the mechanical properties of the hot rolled sheet produced according to the present invention, heat treatment may be applied. Thus, even if the known techniques are applied to the present invention, it is natural to interpret that this is a simple change of the present invention substantially within the technical idea 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 as slabs by a continuous casting method by dissolving them in a converter, and the composition ratio between the alloying elements for each steel type to show the effect of the present invention is shown in Table 2. Table 3 shows the solidification rate of slab, slab heating temperature, heating time, rolling start temperature and end temperature, rolling reduction, and rolling rate of 25 ~ 40mm thickness in rolling process. At this time, the rolling reduction ratio of all steel grades was made 75% 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. In addition, the test for evaluating the mechanical properties (hardness, toughness impact toughness) of the weld heat affected zone was performed by the hardness test in the cross section direction of the weld zone, and the impact test piece was evaluated by welding test pieces taken from the rolling direction in the vertical direction.

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.

Alloy element composition ratio for showing the effect of the present invention Al / O Mn / O Ti / N Al / N N / B V / N (Ti + 2Al + 4B + V) / N Inventive Steel 1 18.8 481 1.2 5.2 17.1 0.8 8.9 Inventive Steel 2 13.0 277 1.82 2.5 28.0 0.4 7.3 Invention Steel 3 21.4 529 1.3 5.4 36.7 1.8 14.2 Inventive Steel 4 9.7 477 2.5 3.8 16.0 6.3 14.0 Inventive Steel 5 17.3 292 1.6 3.0 20.0 1.7 9.5 Inventive Steel 6 12.5 469 2 4.0 10.0 9.0 16.4 Inventive Steel 7 16.7 493 1.3 4.3 14.4 1.7 10.3 Inventive Steel 8 13.6 345 1.5 5.0 12.0 0.8 12.7 Inventive Steel 9 9.3 349 2.2 4.4 22.5 2.2 10.2 Inventive Steel 10 8.6 279 2.5 4.5 16.7 2.0 13.7 Inventive Steel 11 10.9 322 1.4 3.8 12.0 - 9.3 Conventional Steel 1 0.6 595 4.1 0.6 13.8 - 5.7 Conventional Steel 2 1.1 418 2.5 0.8 96.0 - 4.0 Conventional Steel 3 0.3 103 0.8 0.4 105.8 - 1.5 Conventional Steel 4 1.1 540 4.1 0.8 4.0 8.8 15.5 Conventional Steel 5 0.8 325 6.5 1.1 4.0 18.5 28.1 Conventional Steel 6 0.4 392 3.2 0.4 2.6 16.1 21.6 Conventional Steel 7 - - 1.0 2.5 9.9 - 6.5 Conventional Steel 8 - - 1.2 0.4 14.3 - 2.2 Conventional Steel 9 - - 0.8 2.1 9.1 3.9 9.2 Conventional Steel 10 - - 0.6 3.2 9.5 1.5 8.9 Conventional Steel 11 - - 5.5 3.4 12.7 7.8 20.3

Category Casting speed (m / min) Heating temperature (℃) Heating time (min) Rolling Start Temperature (℃) Rolling end temperature (℃) Rolling amount at recrystallization station (%) Cumulative pressure drop (%) Cooling rate (℃ / min) Inventive Steel 1 Invention 1 1.1 1150 160 980 810 55 75 5 Invention 2 1.0 1200 120 1010 820 55 75 5 Invention 3 1.0 1250 80 1020 820 55 75 5 Comparative Material 1 0.6 1100 60 1020 820 55 75 5 Comparative Material 2 1.4 1300 180 1020 820 55 75 5 Inventive Steel 2 Invention 4 0.95 1180 160 1030 830 50 75 7 Invention Steel 3 Invention 5 0.9 1180 150 1000 810 50 75 8 Inventive Steel 4 Invention 6 0.95 1210 130 990 800 50 75 10 Inventive Steel 5 Invention 7 0.96 1220 120 980 790 50 75 5 Inventive Steel 6 Invention Material 8 1.0 1210 120 990 790 50 75 6 Inventive Steel 7 Invention Material 9 1.1 1200 130 1010 800 50 75 5 Inventive Steel 8 Invention 10 1.05 1220 110 1010 810 50 75 5 Inventive Steel 9 Invention 11 1.02 1220 110 1020 820 50 75 5 Inventive Steel 10 Invention Material12 1.05 1190 140 1020 820 50 75 5 Inventive Steel 11 Invention Material 13 1.06 1180 150 1010 820 50 75 5 Conventional Materials 11 - 1200 - Ar ₃ or higher 960 50 80 5 Conventional material (1-10) does not specifically describe its manufacturing conditions

As shown in Table 4, the number of precipitates (Ti-based nitrides) of the hot rolled material produced by the present invention has a range of 2.4 × ^{10 8} / mm ² or more. On the contrary, in the case of the conventional material 11, the range of 11.1 X 10 ³ / mm ² or less, it can be seen that the number of the invention material significantly increased while having a fairly uniform and fine precipitate size compared to the conventional material. In addition, the number of oxides (Al-Mn complex oxide) is about 1.8x10 ² -5.6x10 ³ ranges and the average size is about 1.1-1.3㎛ range. Meanwhile, it can be seen that the ferrite grain size (FGS) of the inventive steel in the base steel structure of the inventive steel is very fine compared to that of the comparative material in the range of about 10-12 μm, and the base ferrite percentage of the inventive steel is all higher than 80%. It consists of fractions.

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 (℃) Invention 1 23 34 56 74 15 372 -74 332 -67 293 -63 Invention 2 22 35 55 77 13 384 -76 350 -69 302 -64 Invention 3 23 35 56 75 13 366 -72 330 -67 295 -63 Comparative Material 1 54 86 182 38 24 124 -43 43 -34 28 -28 Comparative Material 2 65 92 198 36 26 102 -41 30 -32 17 -25 Invention 4 25 38 63 76 14 353 -71 328 -68 284 -65 Invention 5 26 41 57 78 15 365 -71 334 -67 295 -62 Invention 6 25 32 53 75 14 383 -73 354 -69 303 -63 Invention 7 24 35 55 77 14 365 -71 337 -67 292 -63 Invention Material 8 27 37 53 74 13 362 -71 339 -67 296 -62 Invention Material 9 24 36 52 78 15 368 -72 330 -67 284 -63 Invention 10 22 34 53 75 14 383 -72 345 -66 293 -63 Invention 11 26 35 64 75 14 356 -71 328 -68 282 -64 Invention Material12 27 39 62 76 16 363 -72 333 -67 286 -65 Invention Material 13 28 38 61 77 15 364 -72 336 -68 284 -64 Conventional Materials 1 -58 Conventional material 2 -55 Conventional Materials 3 -54 Conventional Materials 4 230 132 (0 ℃) Conventional material 5 180 129 (0 ℃) Conventional Materials 6 250 60 (0 degrees Celsius) Conventional Materials 7 -61 Conventional Materials 8 -48 Conventional Materials 9 -42 Conventional Materials 10 -45 Conventional Materials 11 219J (0 ℃)

Table 5 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 52-65 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. In addition, the ferrite phase fraction of the steel of the present invention at a heat input of 100 kJ / cm is about 70% or more.

As described above, the present invention controls the high heat input welding heat affected zone austenite grains by using an Al-based composite oxide together with TiN precipitates and promotes acicular ferrite transformation in the crystal grains, thereby simultaneously providing excellent high heat input weld heat affected zone toughness. It is possible to provide a structural steel for welding that can be secured.

Claims (10)

By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.03-0.3%, 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%, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤ 7, 6.5≤ (Ti + 2Al + 4B) / N≤14, 8≤Al / O≤22, 250≤Mn / O≤530, satisfying the composition of Fe and other impurities, the microstructure is 20㎛ or less Welded structural steel with excellent toughness of welded heat affected zone consisting of a composite structure of ferrite and ferrite. 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. Welded structural steels with excellent toughness of weld heat affected zone, characterized in that the satisfactory. 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%. A welded structural steel with excellent toughness in the weld heat affected zone, characterized by containing one or two or more selected from the group of Ca: 0.0005-0.005% and REM: 0.005 to 0.05%. The Al ₂ O ₃ · MnO oxide according to claim 1 or 2, wherein TiN precipitates of 0.01-0.1 μm are distributed in the steel material at 1.0 × 10 ⁷ / mm 2 or more at intervals of 0.5 μm or less, and 0.5-2.0 μm. A welded structural steel with excellent toughness in the weld heat affected zone, characterized in that 1 × 10 ² -1x10 ³ / mm ^{2 is} distributed. By weight C: 0.03-0.17%, Si: 0.01-0.5%, Mn: 0.4-2.0%, Ti: 0.005-0.2%, Al: 0.03-0.3%, 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: 0.002-0.3%, 0.5≤Ti / N≤3.0, 5≤N / B≤50, 2≤Al / N≤ 10, 6.5≤ (Ti + 2Al + 4B) / N≤14, 40% in austenite recrystallization zone after heating steel slab composed of remaining Fe and other impurities for 60-180 minutes in the range of 1100-1250 ℃ Hot rolling at the above rolling ratio, and then cooling to a ferrite transformation end temperature ± 10 ° C at a rate of 1 ° C / min or more. 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. Method for producing a welded structural steel with excellent weld heat affected zone toughness 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%. 1 or 2 or more selected, and Ca: 0.0005-0.005%, REM: 0.005 to 0.05% of the method for producing a welded structural steel having excellent toughness of the weld heat affected zone characterized in that it contains one or two selected from the group. . [6] The welding heat affected zone according to claim 5, wherein the slab is continuously cooled by casting molten steel at a rate of 0.9 to 1.1 m / min and lightly cooled to a specific amount of 0.3 to 0.35 l / kg in a secondary cooling zone. Method for producing welded structural steel with good toughness. The welding is applied to the steel (base material) of claim 4 to form a prior austenite of 80 μm or less in the weld heat affected zone, and then quenched to 70% of the ferrite of 20 μm or less in the microstructure of the weld heat affected zone. Welding structure with excellent toughness of the weld heat affected zone composed of the above ordinary fraction. 10. The weld structure of claim 9, wherein the ferrite is polygonal ferrite and acicular ferrite.