KR20020041080A - Method for manufacturing steel plate having superior toughness in weld heat-affected zone by recrystallization controlled rolling - Google Patents

Abstract

PURPOSE: A method for manufacturing high strength steel plate having superior impact toughness in weld heat-affected zone by recrystallization controlled rolling is provided, in which microstructured TiN that is stable at even elevated high temperatures is diffused in steel homogeneously. CONSTITUTION: The method for manufacturing the high strength steel plate for welded structure 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 or 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, 0.005 wt.% or less of O to the temperature range of Ar3 to (Ar3+50°C), wherein each metal components satisfies the follow inequalities: 1.2<=Ti/N<=2.5, 10<=N/B<=40, 2.5<=Al/N<=7, 6.5<=(Ti+2Al+4B)/N<= 14; cooling the heated steel slab to the temperature range of Ar3+25°C to Ar3-25°C at a cooling rate of less than 25°C/sec; hot rolling at a reduction ratio of higher than 60%; and cooling to ambient temperature at a cooling rate of 5 to 30°C/sec.

Description

Method for manufacturing steel plate having superior toughness in weld heat-affected zone by recrystallization controlled rolling}

The present invention relates to a method for manufacturing 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 manufacturing a welded structural steel that can improve the toughness of the weld heat affected zone by uniformly distributing the fine TiN precipitate while increasing the fine ferrite phase fraction in the base material.

Recently, steel materials used in accordance with the trend of high-rise buildings, structures are being replaced by thick materials. In order to weld such thick materials, high-efficiency welding is inevitable. As a technique for welding thickened steel materials, a high-pass heat submerged welding method and an electro-welding method capable of 1-pass welding are widely used. In addition, in the field of shipbuilding and bridges, even in the case of 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, and in order to weld steel with a thicker 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 becomes the site where the toughness of the weld is most degraded.

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 known 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 high temperature, to steel. Such techniques include Japanese Patent Laid-Open No. 11-140582, Hei 10-298708 Hei 10-298706 Hei 9-194990 Hei 9-324238 Hei 8-60292, (S) 60-245768, (Square) 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 being precipitated at × 10 ³ pieces / mm 2 to 6.2 × 10 ⁴ pieces / mm 2.

However, according to this prior art, when 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), Since the difference is about 100J, there is a limit in securing the reliability of the steel structure according to superheat input welding of thickened steel. In addition, as a method for securing a desired TiN precipitate, the slab is heated at a temperature of 1050 ° C. or more and quenched and then reheated for hot rolling, thereby increasing the manufacturing cost due to two heat treatments. In addition, the base material has a poor ferrite + pearlite structure, yield strength of 400 MPa, tensile strength of 530 MPa, and elongation of 20% or less.

Until now, many techniques have been known to improve the toughness of the weld heat affected zone during high heat input welding. However, the case of remarkably improving the toughness of the weld heat affected zone during superheated heat welding, which is maintained for more than 1350 ° C for a long time while securing the mechanical properties of the base metal, It has not been announced yet. 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 high-rise of the steel structure and the reliability of the steel structure.

The present invention improves the mechanical properties of the base material by increasing the number of TiN precipitates having excellent high temperature stability while increasing the fraction of fine ferrite in the base material, as well as effectively inhibiting the growth of the austenite grains of the base material when the high heat input welding is applied. It is an object of the present invention to provide a method for manufacturing a welded structural steel that can minimize the difference in toughness of the base material and the heat affected zone.

The present invention for achieving the above object, in the 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: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤N / B Steel slabs satisfying ≤40, 2.5≤Al / N≤7, 6.5≤ (Ti + 2Al + 4B) / N≤14 are heated to a temperature range of Ar ₃ to (Ar ₃ + 50 ° C),

The heated slab is cooled to a temperature range of (Ar ₃ + 25 ° C) to (Ar ₃ -25 ° C) at a rate of 25 ° C / sec or less, hot rolled at a reduction ratio of 60% or more, and cooled at 5 to 30 ° C / sec. It comprises a cooling to room temperature at a speed.

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 studied a method of simultaneously improving the strength and toughness of steel (base metal) and the toughness of the weld heat affected zone. When uniformly distributed, it was confirmed that the crystal grain size of the austenite of the weld heat affected zone was reduced to about 80 μm or less, so that the toughness of the weld heat affected zone was not a problem.

Starting from this point of view, the present invention was able to derive the following method to improve the characteristics of the base material and the weld heat affected zone.

[1] While minimizing the distribution of the fine TiN precipitates, the solubility product showing stability of the precipitates at high temperatures is reduced.

[2] The recrystallization-controlled rolling refines the ferrite grains of the base material to manage the size of the old austenite of the weld heat affected zone to about 80 μm or less. By refining the ferrite grains of the steel material, the strength and toughness of the base material are simultaneously improved. Also,

[3] By reducing the Ti / N ratio, BN and AlN precipitates can be effectively precipitated to increase ferrite production in the welding heat direction, and in particular, to induce ferrite shapes into needles or polygons effective for toughness improvement. . [1] [2] [3] will be described in detail below.

[1] management of TiN precipitates

When heat input welding is applied to steel materials, the heat affected zone near the melting line is heated to about 1400 ° C or more, and the TiN precipitates deposited in the base materials are partially dissolved or coarsened to suppress the austenite grain growth. The effect disappears.

The inventors noticed that this phenomenon is caused by the diffusion of TiN precipitates distributed in the base metal by the dissolution of the dissolved Ti atoms by the heat of welding. As a result of examining the characteristics of TiN precipitates according to the ratio of Ti / N, the high nitrogen environment (Ti / 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.

According to the present invention, when the ratio of Ti and N (Ti / N) is heated to a 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, resulting in the fine TiN precipitate having a size of 0.01-0.1 μm. A result of distributing more than 1.0 × 10 ⁷ / mm ² at intervals of 0.5 μm or less can be obtained. Increasing the nitrogen content at the same Ti content causes all Ti atoms to be easily combined with nitrogen atoms, and the amount of solid solution Ti decreases in a high nitrogen environment. The solubility was lowered.

Furthermore, the present inventors are interested in the process of refining steel, which is important to solidify most of Ti without forming Ti as an oxide in molten steel in order to uniformly distribute a large amount of fine TiN precipitates. As a result, it was confirmed that Ti was not formed as an oxide and dissolved in steel when the amount of dissolved oxygen, which greatly affects the formation behavior of oxide, was controlled to 30 ppm or less.

According to the present invention, by managing the deoxidation conditions so that Ti is dissolved in the steel as much as possible, and when the ratio of Ti / N is controlled, fine TiN precipitates of 0.01-0.1 μm may be 1.0 × ^{10 7} / mm 2 or more at intervals of 0.5 μm or less. It can be distributed.

[2] recrystallization-controlled rolling (fine grains in ferrite grains)

According to the present invention, it is important to refine the size of the ferrite in the base material structure of ferrite + pearlite together with the management of the precipitates in order to make the size of the old austenite at 80 μm while improving the strength and toughness of the base material. . In the present invention, the refinement of the ferrite is achieved by using recrystallization control rolling. Recrystallization control rolling uses the precipitate (TiN) during high temperature rolling to suppress grain growth, grain refining by static recrystallization during rolling, and precipitation of carbides during austenite-ferrite transformation (WC, Fe ₃ C, VC, etc.) Ferrite grains are refined.

[3] microstructure of weld heat affected zone

The facts of the present invention revealed that the toughness of the weld heat affected zone has a significant influence not only on the size of the austenite grains but also on the amount and shape of the ferrite precipitated at the former austenite grain boundaries when the base material is heated above 1400 ° C. Will be. In particular, it is important to induce the transformation of polygonal ferrite and acicular ferrite in the austenite mouth. In the present invention, AlN, Fe ₂₃ (B, C) ₆ , BN precipitates are used for this purpose.

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

[Composition of Welding Structural Steels]

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

If the content of 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 during transformation are transformed to lower the low temperature impact toughness of structural steel, and Increasing hardness or strength results in toughness degradation and generation of weld cracks.

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

If the Si content is less than 0.01%, the deoxidation effect of the molten steel is insufficient during steelmaking and the corrosion resistance of the steel is deteriorated. If the content of Si is more than 0.5%, the effect is saturated, and the degree of quenchability during cooling after rolling is increased. 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 which deoxidizes in steel and improves weldability, hot workability, and strength. Mn forms a substituted solid solution in the matrix to strengthen the matrix to secure the strength and toughness. For this purpose, Mn is preferably added at least 0.4%. However, when the Mn content is more than 2.0%, it has a detrimental effect on the toughness of the weld heat affected zone due to tissue heterogeneity due to manganese segregation rather than a solid solution strengthening effect. In addition, when the steel solidifies, macro segregation and micro segregation occur depending on the segregation mechanism, which promotes the formation of a central segregation zone in the center of rolling, thereby causing a low temperature transformation structure in the center of the base metal. In particular, manganese is an element that precipitates in the form of MnS around the Ti-based oxide and affects the formation of acicular and polygonal ferrites, which is effective for improving weld heat affected toughness.

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

Al is not only an element necessary for deoxidation, but also an indispensable element for forming fine AlN precipitates in steel. In addition, Al is an element which reacts with oxygen to form an Al oxide, and thus is necessary for Ti to form fine TiN precipitates without reacting with oxygen. In order to form a fine AlN precipitate, the Al should be added 0.0005% or more, but if the content exceeds 0.1%, Weedman Stetten is vulnerable to toughness during the cooling of the welding heat affected zone of AlN precipitated and the remaining solid Al It promotes the production of ferrite (Widmanstatten ferrite) and phase martensite to lower the toughness of the high heat input welding heat affected zone.

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 fine TiN precipitates that are stable at high temperatures. In order to obtain such a fine TiN precipitation effect, Ti should be added more than 0.005%, but when the content exceeds 0.2%, coarse TiN precipitates and Ti oxides are formed in molten steel, which does not inhibit austenite grain growth of the weld heat affected zone. It is not preferable because it is.

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

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%, it is not preferable because the hardenability increases due to hardening and low temperature cracking of the weld heat affected zone.

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 suppresses austenite grain growth of welding heat affected zone at the time of high heat input welding and precipitates such as TiN, AlN, BN, VN, NbN Increase the amount. In particular, the content of TiN and AlN precipitates and the spacing of precipitates, the distribution of precipitates, the frequency of complex precipitation with oxides, the high temperature stability of the precipitates themselves, etc. have a significant influence, so the content is preferably set to 0.008% or more. However, when the N content exceeds 0.03%, the effect is saturated, toughness is reduced due to the increase in the amount of solid solution nitrogen distributed in the weld heat affected zone, and it is mixed in the weld metal due to dilution during welding to reduce the toughness of the weld metal. It is preferable to limit the upper limit to 0.03% because of this.

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

W is an element that uniformly precipitates in the base material as tungsten carbide (WC) after hot rolling to effectively suppress ferrite grain growth after ferrite transformation, and also suppress growth of austenite grains in the initial heating of the weld heat affected zone. If the content is less than 0.001%, the tungsten carbide for inhibiting ferrite grain growth is less distributed during the hot rolling and it is not preferable because 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. However, in the case of sulfur, it precipitates in the form of MnS around the Ti-based oxide and thus affects the formation of needle-shaped and polygonal ferrites effective for improving the toughness of the weld heat affected zone. It is desirable to limit the amount from 0.003% to 0.03% or less.

Oxygen (O) is preferably at most 0.005%.

When the oxygen content exceeds 0.005%, the Ti element is not preferable because Ti element is formed of Ti oxide in molten steel, and thus TiN precipitate is not formed. It is not desirable because it is crazy.

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 distribute uniformly while increasing the number of TiN precipitates. In other words, if the nitrogen content is increased at the same Ti content, all Ti atoms that are employed in the cooling process during the playing process are combined with the nitrogen atoms to distribute fine TiN precipitates at narrow intervals. Secondly, the solubility product showing stability at high temperature is small, thereby preventing re-use of the Ti. That is, in a high nitrogen environment, Ti has a stronger tendency to bind with N rather than solid solution, so that the TiN precipitate becomes stable at high temperature. However, when the Ti / N ratio is less than 1.2, it is not preferable because the amount of solid solution nitrogen in the base material increases, which is detrimental to the toughness of the weld heat-directed portion. 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. It is not preferable because it adversely affects the impact toughness.

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 which promotes the ferrite transformation of polygons at the austenite grain boundary during the cooling process after welding is insufficient, and when the N / B ratio exceeds 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 in the grain and the polygonal ferrite The size and number of distributions of TiN, AlN, BN, and VN precipitates for formation and control of the tissue fraction are insufficient, and the effect is saturated when (Ti + 2Al + 4B) / N exceeds 14. On the other hand, when V is added, it is preferable to set the ratio of (Ti + 2Al + 4B + V) / N to 7-17.

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, and the content of this 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. On the other hand, when the V / N ratio exceeds 9, the weld heat influences because the size of the VN precipitates deposited at the TiN + MnS precipitate boundary is coarsened, and the number of VN precipitations deposited at the TiN + MnS composite precipitate boundary is reduced. Reduce the ferrite phase fraction effective for negative toughness.

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 for improving the strength and toughness of the base material by solid solution strengthening. To achieve this effect, Ni should be added more than 0.1%, but if the content exceeds 3.0%, the hardenability increases to decrease the toughness of the weld heat affected zone, and there is a possibility of high temperature crack in the weld heat affected zone and the weld metal. Because it is not desirable.

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

Cu is an element that is effective at securing the strength and toughness of the base material due to solid solution at the base. In order to achieve this effect, Cu should be added more than 0.1%, but if the content exceeds 1.5%, it is desirable because it increases the hardenability in the weld heat affected zone and lowers the toughness and promotes high temperature crack in the weld heat affected zone and the weld metal. I can't. In particular, the Cu precipitates in the form of CuS around the Ti-based oxide together with sulfur, thereby affecting the formation of needle-shaped and polygonal ferrites, which are effective for improving the toughness of the weld heat affected zone, and thus the content is 0.3-1.5%. desirable.

In addition, when adding Cu and Ni compositely, it is preferable to make these sum total less than 3.5%. The reason for this is that when it exceeds 3.5%, the hardenability becomes large, 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 base material strength, and this effect cannot be obtained when the Nb content is less than 0.01%. On the other hand, if it exceeds 0.1%, it causes undesired precipitation of coarse NbC, which is detrimental to the toughness of the base metal.

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

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

Molybdenum (Mo) is preferably 0.05-1.0%.

Mo is an element that has an effect of increasing hardenability and improving strength. The content of Mo is preferably set to 0.05% or more for securing strength, but in order to suppress HAZ hardening and welding low temperature cracking, the upper limit thereof is 1.0%. It is preferable to set it as.

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

Ca and REM form an oxide having excellent high temperature stability, thereby suppressing austenite grain growth 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 impair 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.

[Method of manufacturing welded structural steel]

Refining (Deoxidation, Degassing) Process

In general, the steel refining process consists of an out-of-furnace refining process after the first refining of the converter and the second refining of the molten steel of the converter by ladle. ). Usually deoxidation takes place between primary and secondary refining.

A feature of the present invention is that in such a deoxidation process, the dissolved oxygen is adjusted to an appropriate level or lower, and then Ti is added, so that Ti is mostly dissolved in molten steel without forming Ti as an oxide. For this purpose, it is preferable to inject and deoxidize an element having a greater deoxidizing power than Ti before adding Ti. The deoxidizing power of the deoxidizer is as follows.

Cr <Mn <Si <Ti <Al <REM <Zr <Ca ≒ Mg

The amount of dissolved oxygen is greatly influenced by the formation behavior of the oxide, and the deoxidizer having a high affinity with oxygen has a very high rate of bonding with oxygen in the molten steel. Therefore, if deoxidation is performed using an element having a greater deoxidizing power before adding Ti, it is possible to prevent Ti from forming an oxide as much as possible. Of course, before the addition of the element (Al), which has a greater deoxidizing power than Ti, by adding and deoxidizing Mn, Si and the like, which are the five major elements of steel, and then deoxidizing by adding Al, the amount of deoxidizer added is preferable.

On the other hand, floating separation of inclusions in molten steel is generally known to proceed in the following order. (Dissolution of deoxygenation element in molten steel) → (Involvement of nucleation) → (Growth of inclusions) → (Continuous growth and injury due to collisions between inclusions) → (Removal of slag from slag on molten steel surface) Since the speed of each step varies depending on the type, deoxidation using a strong deoxidation element can lower the dissolved oxygen amount more easily.

In the present invention, the amount of dissolved oxygen is added as low as possible by injecting a strong deoxidation element before the introduction of Ti, and in order to maximize the amount of Ti dissolved in molten steel, it is preferable to set it as at least 30 ppm or less. The reason for this is that when the dissolved oxygen amount exceeds 30 ppm, the oxygen in the molten steel and Ti are combined to form Ti oxide easily when Ti is added, and the amount of solid solution Ti decreases.

On the other hand, according to the 'Stoke Law' widely used in steelmaking, the higher the density of inclusions, the more difficult the inclusions are. Injuries are difficult As the inclusions increase in the steel, the addition of a deoxidation element forming a high density inclusion may be used as an additional advantage according to the distribution of oxides. It does not affect the effect of.

After adjusting the amount of dissolved oxygen according to the present invention, it is preferable to add Ti within 10 minutes so that the content is 0.005-0.2%. If less than 0.005% of Ti is contained in the molten steel after deoxidation, it is difficult to form a large amount of fine TiN in the casting process, and if it contains more than 0.2%, the effect is saturated and the TiN is coarsened to expect the austenite grain suppression effect. Difficult to do The reason for adding Ti within 10 minutes is that Ti oxide is generated and the amount of solid solution Ti decreases as the time after Ti is added. In the case where the vacuum degassing treatment ('RH') is performed in the refining process, the addition of Ti can be performed either before or after the vacuum degassing treatment.

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 set to 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.

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 for showing the effect of the present invention by coarsening 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 (recrystallization controlled rolling)

In the present invention, it is characterized by setting the rolling conditions that can make the ferrite grains fine by recrystallization control rolling.

According to the invention the slab is heated to a temperature range of Ar ₃ ~ (Ar ₃ + 50 ℃). It is because below Ar ₃ becomes ferrite + austenite biphasic structure and coarse structure in which coarse ferrite remains in post process is not desirable to secure a homogeneous microstructure, and above Ar ₃ + 50 ° C, austenitic particles are too coarse This is because the ferrite refining effect in the later step is less, which degrades the mechanical properties of the steel.

The heated slab is cooled to a temperature range of (Ar ₃ +25) to (Ar ₃ -25 ° C) at a rate of 25 ° C / sec or less, and then to a temperature of (Ar ₃ +25) to (Ar ₃ -25 ° C). Rolling is started. This is because when the cooling rate is more than 25 ° C / sec, it is difficult to secure the austenite recrystallization region and thus it is difficult to secure fine ferrite grains in the abnormal reverse rolling.

After cooling to a temperature range of (Ar ₃ + 25 ° C)-(Ar ₃ -25 ° C), it is rolled in this temperature range. If the rolling temperature is lower than Ar ₃ -25 ℃ has the ferrite to precipitate a large amount of steel in the crude residue in a coarse rolling and then stretching the shape before rolling has an adverse effect on the finely divided ferrite. When the rolling temperature is higher than Ar ₃ + 25 ° C., deformation organic transformation does not occur, so it is difficult to obtain fine ferrite, and even if deformation organic transformation occurs, the particle size of ferrite is not preferable. It is preferable to make the reduction ratio at this time into 60% or more. If the reduction ratio is less than 60%, the ferrite grains formed at the grain boundaries of the austenite may increase because the austenite grains are coarsened during recrystallization.

After hot rolling as above, it is cooled to room temperature at a cooling rate of 5 to 30 ° C / sec. If the cooling rate is less than 5 ℃ / sec, it is difficult to suppress the growth of ferrite grains, and if the cooling rate is faster than 30 ℃ / sec is not preferable because residual austenite or martensite is formed, which adversely affects the toughness of the base material .

[Organization of Steel]

· Microstructure

Steel produced according to the present invention has a microstructure of ferrite + pearlite, the grain size of the ferrite is preferably to have 6㎛ or less. As the ferrite becomes finer, the toughness of the weld heat affected zone can be improved by atomizing the old austenite grains of the weld heat affected zone to 80 μm or less during high heat input welding.

In addition, the higher the percentage of ferrite in the composite structure of ferrite + perlite increases the toughness and elongation of the base material. In consideration of this, it is most preferable that the phase fraction of the ferrite is 80% or more.

Distribution of precipitates

The former austenite grains of the weld heat affected zone are greatly influenced by the size and number of nitrides distributed in the base metal and their distribution when the austenite grain size of the base material is constant. In addition, 30-40% of the nitrides distributed in the base material when the heat input welding is higher than the heat input temperature (above the heating temperature of 1400 ℃) are re-used as the base material, thereby reducing the austenite grain growth inhibiting effect of the welding heat affected zone. There is a need for a uniform distribution of further nitrides taking into account the nitrides that become. In order to suppress the growth of 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 the high heat input welding, so that the effect of inhibiting the growth of the austenite grains is insufficient. If the thickness exceeds 0.1 μm, pinning of the austenite grains occurs. This is because the growth inhibition effect is reduced and the same behavior as coarse nonmetallic inclusions has a detrimental effect on the 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 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.

In the hot rolling process of the present invention, hot charge rolling and direct rolling, which are well-known according to a user's use, may be applied, and various techniques such as known 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 according to the conditions of Tables 2 and 4 by melting in a converter, and the composition ratios between the alloying elements of the steel types according to the conditions of the present invention are shown in Table 3 to show the effects of the present invention. .

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

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 weld heat affected zone, it is heated to the maximum heating temperature (1200 ~ 1400 ℃) to 140 ℃ / sec condition by using a simulated welding simulator (1). Hold for a second and then 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).

TiN precipitate size, number, and spacing, 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 welding heat affected zone is 800-500 ℃ cooling after heating the welding conditions corresponding 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 REM O (ppm) Invention 1 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7 120 0.005 - - - - - 0.01 - - 11 Invention 2 0.09 0.12 1.50 0.006 0.005 0.07 0.05 10 280 0.002 - 0.2 - - - 0.01 - - 12 Invention 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015 3 110 0.003 0.1 - - - - 0.02 - - 10 Invention 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02 5 80 0.001 - - - - - 0.05 - - 9 Invention 5 0.11 0.15 1.52 0.006 0.004 0.09 0.05 15 300 0.002 0.1 - 0.1 - - 0.05 - - 12 Invention 6 0.12 0.14 1.50 0.007 0.005 0.025 0.02 10 100 0.004 - - - 0.1 - 0.09 - - 9 Invention 7 0.13 0.14 1.48 0.007 0.005 0.04 0.015 8 115 0.15 0.1 - - - - 0.02 - - 11 Invention Material 8 0.11 0.15 1.52 0.007 0.005 0.06 0.018 10 120 0.001 - - - - 0.015 0.01 - - 10 Invention Material 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 - 12 Invention 10 0.12 0.16 1.45 0.008 0.006 0.045 0.025 6 100 0.05 - 0.3 - - 0.01 0.02 - 0.01 8 Inventive Steel 11 0.09 0.12 1.48 0.006 0.003 0.048 0.019 10 130 0.01 - 0.1 - - - - - - 14 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 - - 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 - - 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 - - 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, and conventional steels (4, 5, 6) are inventions of Japanese Patent Application Laid-Open No. 10-298708. Steels (14, 24, 28) and conventional steels (7, 8, 9, 10) are invention steels (48, 58, 60, 61) of JP-A-8-60292. Inventive Steel F of Patent Publication No. 11-140582

Steel grade used division Primary deoxidation sequence Dissolved oxygen after Al addition (ppm) Ti addition after finishing deoxidation (%) Casting speed (m / min) Specific quantity (ℓ / kg) Inventive Steel 1 Invention 1 Mn → Si 19 0.014 1.1 0.35 Invention 2 Mn → Si 18 0.014 1.1 0.35 Invention 3 Mn → Si 18 0.014 1.1 0.35 Comparative Material 1 Mn → Si 79 0.014 1.1 0.35 Comparative Material 2 Mn → Si 118 0.014 1.1 0.35 Inventive Steel 2 Invention 4 Mn → Si 16 0.05 1.2 0.35 Invention Steel 3 Invention 5 Mn → Si 15 0.015 1.2 0.35 Inventive Steel 4 Invention 6 Mn → Si 15 0.02 1.2 0.35 Inventive Steel 5 Invention 7 Mn → Si 12 0.05 1.2 0.35 Inventive Steel 6 Invention Material 8 Mn → Si 17 0.02 1.2 0.35 Inventive Steel 7 Invention Material 9 Mn → Si 18 0.015 1.1 0.35 Inventive Steel 8 Invention 10 Mn → Si 14 0.018 1.1 0.35 Inventive Steel 9 Invention 11 Mn → Si 19 0.02 1.2 0.35 Inventive Steel 10 Invention Material12 Mn → Si 23 0.025 1.1 0.35 Inventive Steel 11 Invention Material 13 Mn → Si 23 0.019 1.1 0.35 The manufacturing conditions of the conventional steel (1-11) are not specifically described.

Alloy element composition ratio for showing the effect of the present invention Ti / N N / B Al / N V / N (Ti + 2Al + 4B + V) / N Invention 1 1.2 17.1 3.3 0.8 8.9 Invention 2 1.2 17.1 3.3 0.8 8.9 Invention 3 1.2 17.1 3.3 0.8 8.9 Comparative Material 1 1.2 17.1 3.3 0.8 8.9 Comparative Material 2 1.2 17.1 3.3 0.8 8.9 Invention 4 1.8 28.0 2.5 0.4 7.3 Invention 5 1.4 36.7 5.5 1.8 14.2 Invention 6 2.5 16.0 2.5 6.3 14.0 Invention 7 1.7 20.0 3.0 1.7 9.5 Invention Material 8 2.0 10.0 2.5 9.0 16.4 Invention Material 9 1.3 14.4 3.5 1.7 10.3 Invention 10 1.5 12.0 5.0 0.8 12.7 Invention 11 2.2 22.5 2.8 2.2 10.2 Invention Material12 2.5 16.7 4.5 2.0 13.7 Invention Material 13 1.5 13.0 3.7 - 9.2 Conventional Steel 1 4.1 13.8 0.6 - 5.7 Conventional Steel 2 2.5 96.0 0.8 - 4.0 Conventional Steel 3 0.8 105.8 0.4 - 1.5 Conventional Steel 4 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 6.5 4.0 1.1 18.5 28.1 Conventional Steel 6 3.2 2.6 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) Cooling rate (℃ / sec) Hot Rolling Condition Cooling rate (℃ / sec) Rolling Start Temperature (℃) Rolling end temperature (℃) Cumulative pressure drop (%) Invention 2 Inventive Example 1 930 150 5 920 780 80 11 Inventive Example 2 930 150 5 920 780 80 12 Inventive Example 3 930 150 5 920 780 80 11 Comparative Example 1 1200 50 5 1050 950 40 11 Comparative Example 2 1100 310 5 1020 940 50 11 Invention 1 Inventive Example 4 930 120 8 920 790 80 12 Invention 3 Inventive Example 5 940 110 8 920 790 80 12 Comparative Material 1 Comparative Example 3 1200 130 8 930 890 70 0.5 Comparative Material 2 Comparative Example 4 1250 130 7 1050 890 70 12 Invention 4 Inventive Example 6 1050 120 5 1000 780 80 11 Invention 5 Inventive Example 7 920 110 5 920 780 80 11 Invention 6 Inventive Example 8 930 120 6 920 770 80 11 Invention 7 Inventive Example 9 930 140 5 910 770 80 10 Invention Material 8 Inventive Example 10 940 150 5 910 780 80 11 Invention Material 9 Inventive Example 11 920 130 5 910 780 80 10 Invention 10 Inventive Example 12 940 150 8 920 780 80 10 Invention 11 Inventive Example 13 940 160 8 920 790 80 11 Invention Material12 Inventive Example 14 940 140 7 920 790 80 10 Invention Material 13 Inventive Example 15 940 150 5 920 790 80 10 Conventional Steel 11 1200 - Ar ₃ or higher Cooling Manufacturing conditions of conventional steel (1-10) are not specifically presented

division Precipitate properties Base material structure characteristics Base material mechanical properties Count (pcs / mm ² ) Average size (㎛) Average interval (㎛) AGS FGS Ferrite Percentage (%) Thickness (mm) Yield strength (MPa) Tensile Strength (MPa) Elongation (%) Inventive Example 1 2.3 X ^{10 8} 0.016 0.3 17 5 82 20 564 683 33 Inventive Example 2 3.1 X ^{10 8} 0.017 0.3 15 3 84 20 532 651 32 Inventive Example 3 2.4 X ^{10 8} 0.012 0.4 13 3 83 20 520 662 32 Comparative Example 1 4.2 X 10 ⁶ 0.154 2.5 38 24 68 20 383 524 28 Comparative Example 2 5.4 X 10 ⁶ 0.155 2.6 34 21 65 20 392 530 29 Inventive Example 4 3.2 X ^{10 8} 0.025 0.5 15 5 83 25 550 688 32 Inventive Example 5 2.3 X ^{10 8} 0.013 0.4 14 5 82 25 548 682 31 Comparative Example 3 2.3 X 10 ⁴ 0.264 2.4 38 21 48 25 384 524 20 Comparative Example 4 2.3 X 10 ⁴ 0.257 2.3 36 20 50 25 365 560 22 Inventive Example 6 3.2 X ^{10 8} 0.026 0.5 15 4 84 25 539 683 31 Inventive Example 7 4.1X10 ⁸ 0.024 0.4 16 5 83 30 538 684 33 Inventive Example 8 4.2 X ^{10 8} 0.014 0.4 15 5 82 30 542 682 34 Inventive Example 9 3.2 X ^{10 8} 0.028 0.5 15 5 81 30 535 686 32 Inventive Example 10 4.2 X ^{10 8} 0.021 0.6 15 4 82 30 557 686 33 Inventive Example 11 2.7 X ^{10 8} 0.029 0.4 14 4 84 25 550 696 33 Inventive Example 12 2.2 X ^{10 8} 0.025 0.4 16 5 83 25 542 682 33 Inventive Example 13 1.7 X ^{10 8} 0.024 0.5 17 6 84 25 564 695 30 Inventive Example 14 2.1X10 ⁸ 0.021 0.5 18 6 82 25 572 687 31 Inventive Example 15 2.3 X ^{10 8} 0.026 0.6 18 5 81 25 568 692 30 Conventional Steel 1 35 406 436 31 Conventional Steel 2 35 405 441 30 Conventional Steel 3 25 629 681 28 Conventional Steel 4 Precipitates of MgO-TiN 3.03 × 10 ⁶ pcs / mm ² 40 472 609 32 Conventional Steel 5 Precipitates of MgO-TiN 4.07 × 10 ⁶ pcs / mm ² 40 494 622 32 Conventional Steel 6 Precipitates of MgO-TiN 2.80 × 10 ⁶ pcs / mm ² 50 812 912 24 Conventional Steel 7 25 629 681 23 Conventional Steel 8 50 504 601 23 Conventional Steel 9 60 526 648 25 Conventional Steel 10 60 760 829 24 Conventional Steel 11 50 401 514 18.3

As shown in Table 5, the number of precipitates (Ti-based nitrides) of the hot rolled material produced by the present invention had a range of 1.0 × ^{10 8} / mm ² or more, while the precipitates of conventional materials were 4.6X10 ⁵ / mm ^Since showing the range of ² or less, it can be seen that the number of the invention material is significantly increased while the invention material has a very fine precipitate size compared to the conventional material. In addition, it can be seen that the ferrite grain size (FGS) of the present invention steel is very fine in the structure of the base material, and the ferrite phase fraction is all composed of a high ferrite fraction of 80% or more.

division Austenitic grain size of welded 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 65 84 186 32 29 110 -40 50 -28 18 -20 Comparative Example 4 60 90 210 28 32 120 -38 54 -30 16 -15 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 24 34 52 72 16 340 -70 320 -66 245 -58 Inventive Example 14 26 30 50 73 15 342 -71 318 -64 240 -60 Inventive Example 15 22 29 48 72 16 351 -70 330 -66 278 -62 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 ℃)

As shown in Table 6, the austenitic grain size at the maximum heating temperature of 1400 ° C., such as the weld heat affected zone, ranges from 52 to 65 μm in the case of the present invention. You can see that we go to the range. Therefore, it can be seen that the steel of the present invention has a very good effect of suppressing the austenite grains in the weld heat affected zone during welding. In addition, the ferrite phase fraction of the steel of the present invention at the heat input of 100 kJ / cm was about 70% or more.

On the other hand, when comparing the impact toughness of the weld heat affected zone during high heat input welding, in the case of the present invention under the high heat input weld heat input condition of 250 kJ / cm (180-500 ° C. cooling time of 180 seconds), The impact toughness showed excellent toughness value of about 280J or more, and the transition temperature also showed a value of about -60 ° C. or less, indicating that the high heat input welding heat affected zone impact toughness was excellent. On the other hand, in the case of conventional steel, the impact toughness of 0 ° C. was very low as 60-132J. Therefore, it can be seen that the steels according to the present invention can significantly improve the impact toughness and transition temperature of the welded heat affected zone compared to the existing steels.

As described above, the present invention can further improve the impact toughness of the welding heat affected zone by further distributing a large amount of TiN precipitates that are stable at high temperature, and also increases the phase ratio of fine ferrite in the base material. There is a useful effect that can provide a welded structural steel with excellent strength and toughness of the base material).

Claims (7)

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: 0.005% or less, 1.2≤Ti / N≤2.5, 10≤N / B≤40, 2.5≤Al / N≤7 , Heating the steel slab satisfying 6.5≤ (Ti + 2Al + 4B) / N≤14 to a temperature range of Ar ₃ ~ (Ar ₃ +50 ℃), The heated slab is cooled to a temperature range of (Ar ₃ + 25 ° C) to (Ar ₃ -25 ° C) at a rate of 25 ° C / sec or less, hot rolled at a reduction ratio of 60% or more, and cooled at 5 to 30 ° C / sec. A method for producing a welded structural steel having excellent weld heat affected zone toughness comprising cooling to room temperature at a rate. The method of claim 1, wherein the steel material contains 0.01 to 0.2%, and the ratio of V and N (V / N) is 0.3 to 9, and 7≤ (Ti + 2Al + 4B) / N≤17. A method of manufacturing a welded structural steel having excellent weld heat affected zone toughness. 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 method for manufacturing a welded structural steel having excellent toughness in weld heat affected zone, characterized by containing one or two or more selected from the group selected from Ca: 0.0005-0.005% and REM: 0.005 to 0.05%. The method of claim 1, wherein the steel slab is added to the molten steel prior to Ti injecting a deoxidation element having a greater deoxidizing power than Ti to deoxidize dissolved oxygen of the molten steel to less than 30ppm, and then added within 10 minutes to Ti 0.005 ~ 0.2% Method for producing a welded structural steel with excellent toughness of the weld heat affected zone, characterized in that it is made by continuous casting of degassed molten steel. The method for manufacturing a welded structural steel according to claim 1 or 4, wherein the deoxidation is performed in the order of Mn, Si, and Al. The steel slab of claim 1 or 4, wherein the steel slab is continuously cooled at a rate of 0.9 to 1.2 m / min, and coldly cooled to a specific amount of 0.3 to 0.35 l / kg in a secondary cooling zone. Method for manufacturing welded structural steel with excellent weld heat affected zone toughness. The steel material according to claim 1, 2, or 4, wherein the steel is composed of a composite structure of ferrite and perlite having a microstructure of 8 µm or less, and a TiN precipitate of 0.01-0.1 µm at 1.0 intervals of 0.5 µm or less. Method for manufacturing welded steel with excellent toughness of weld heat affected zone distributed over x10 ⁷ pcs / mm2.