WO2019125075A1 - High-strength steel with excellent toughness of welding heat affected zone and manufacturing method thereof - Google Patents

Abstract

The present invention relates to structural steel used as a material of storage tanks, pressure vessels, building structures, ship structures, etc. and, more particularly, to steel with excellent toughness of a welding heat affected zone and a manufacturing method thereof.

Description

High Strength Steel with Excellent Toughness of Welded Heat Affected Zone and Its Manufacturing Method

The present invention relates to a structural steel used as a material for a storage tank, a pressure vessel, an architectural structure, a ship structure, and the like. More particularly, the present invention relates to a high strength steel having excellent weld heat affected zone toughness and a method of manufacturing the same.

When constructing structures such as storage tank, pressure vessel, building structure mall, and ship structure using structural steel, it involves a lot of welding. Therefore, not only the performance of the base material but also the efficiency of welding and the stability of the welded structure should be secured. To this end, it is necessary to keep the austenite grain growth of the heat affected zone (HAZ) as much as possible and to keep the final transformed structure finely.

As a means for solving this problem, there is proposed a technique of appropriately distributing a Ti-based carbonitride or the like, which is stable at a high temperature, on a steel material to delay the growth of grain growth in the weld heat affected zone during welding.

For example, Patent Document 1 discloses a structural steel having a impact toughness of about 200 J (base material is about 300 J) at 0 캜 when a heat input of 100 J / cm 2 (maximum heating temperature of 1400 캜) is applied as a typical technique using a precipitate of TiN In the above technique, the TiN precipitates having a TiN content of substantially 4-12 are deposited at a ratio of 5.8 × 10 ³ / mm ² to 8.1 × 10 ⁴ / mm ² with a TiN precipitate of 0.05 μm or less, and 0.03 to 0.2 μm, Is precipitated at 3.9 × 10 ³ / mm ² to 6.2 × 10 ⁴ / mm ² to make the ferrite finer to secure the toughness of the welded portion.

However, in Patent Document 1, there is a problem that cracks are generated on the surface of the slab during performance by forming excess carbon and nitride. When a slab with many surface cracks is used to produce a heavy plate product, There is a problem in that there is a possibility that defective products such as surface repair or the like can not be manufactured due to the occurrence of problems such as edo cracks.

(Patent Document 1) Japanese Laid-Open Patent Publication No. 1999-140582

One aspect of the present invention is to provide a steel material having excellent strength and toughness of a base material and excellent welding heat affected zone (HAZ) after welding and stress annealing, and a method of manufacturing the steel material.

The object of the present invention is not limited to the above-mentioned matters. Additional objects and advantages of the invention will be set forth in part in the description which follows, and in part will become apparent to those having ordinary skill in the art upon examination of the following or may be learned from practice of the invention.

An aspect of the present invention relates to a method of manufacturing a semiconductor device, comprising: 0.16 to 0.20% of C, 1.0 to 1.5% of Mn, 0.3% or less of Si (excluding 0), 0.005 to 0.5% of Al, 0.01% or less, 0.005 to 0.02% of Ti, 0.01 to 0.1% of Nb, and 0.006 to 0.01% of N,

 At least one selected from the group consisting of Ca: not more than 0.006%, V: not more than 0.03%, Ni: not more than 2.0%, Cu: not more than 1.0%, Cr: not more than 1.0%, and Mo: not more than 1.0% Containing impurities,

The microstructure is composed of a ferrite-pearlite composite structure,

After the welding and stress relieving heat treatment, the microstructure had a precipitate of 100 nm or less of 1.27 x 10 < ⁶ > And more than 900 of them are distributed in a single crystal grain.

In another aspect of the present invention, there is provided a steel sheet comprising, by weight, 0.16 to 0.20% of C, 1.0 to 1.5% of Mn, 0.3% or less of Si (excluding 0), 0.005 to 0.5% of Al, S: 0.01% or less, Ti: 0.005 to 0.02%, Nb: 0.01 to 0.1%, and N: 0.006 to 0.01%

 At least one selected from the group consisting of Ca: not more than 0.006%, V: not more than 0.03%, Ni: not more than 2.0%, Cu: not more than 1.0%, Cr: not more than 1.0%, and Mo: not more than 1.0% Preparing a steel slab containing impurities;

Heating the steel slab in a temperature range of 1050 to 1250 占 폚;

Hot rolling the heated steel slab at a hot finish rolling temperature of 910 占 폚 or lower; And

And cooling the steel sheet at a cooling rate of 20 ° C / Hr or less after the hot rolling, to a method of manufacturing a high strength steel having excellent weld heat affected zone toughness.

According to the present invention, it is possible to provide a steel material which does not deteriorate the strength and toughness of the base material even when subjected to the post-weld stress relieving heat treatment and is excellent in the weld heat affected zone toughness at the time of heat welding. Also, since the strength of the base material is maintained even by the stress relaxation heat treatment, it can be suitably used for storage tanks, pressure vessels, structures, and the like. In addition, since the steel material of the present invention has no defects such as surface cracking, it can be suitably applied as a structural steel material.

BRIEF DESCRIPTION OF THE DRAWINGS Fig. 1 is a photograph of microstructure observed in Inventive Example 1 and Comparative Example 1 of the present invention. Fig.

2 is a TEM (transmission electron microscope) observation of the size and shape of the NbC precipitates observed in Example 1 of the present invention.

FIG. 3 is a TEM (transmission electron microscope) observation of the size and shape of Fe ₃ C precipitates observed in Comparative Example 6 of the present invention.

The inventors of the present invention have conducted extensive research to fundamentally solve the problem of cracks and other defects on the surface of steel during the manufacture of a steel material for use in conventional structural steel. As a result, The present invention has been accomplished on the basis of confirming that it is possible to secure not only the toughness but also the weld heat affected zone having excellent toughness by controlling the microstructure of the weld heat affected zone at the time of welding.

Particularly, since the toughness of the weld heat affected zone (HAZ) can be ensured well during the large heat welding such as submerged arc welding, the present invention can be suitably applied as structural steel.

On the other hand, as a result of the inventors of the present invention, a method for preventing a decrease in the strength of the base material, which may occur after the stress annealing process, which is generally performed for the purpose of stabilizing the material of the welded portion hardened structure in the production of the storage tank or the pressure vessel, It was confirmed that the strength after heat treatment can be ensured by the formation of fine precipitates when the alloy component is added, and the present invention has been accomplished. Therefore, the present invention can be suitably applied not only to conventional structural steel but also to storage tanks and pressure vessels.

Hereinafter, the present invention will be described in detail.

First, the alloy composition of the steel of the present invention will be described in detail. The steel of the present invention is characterized in that the steel contains 0.16 to 0.20% by weight of carbon (C), 1.0 to 1.5% of manganese (Mn), 0.3% (S): 0.01% or less, titanium (Ti): 0.005 to 0.02%, niobium (Nb): 0.01 to 0.1%, nitrogen (N) : 0.006 to 0.01%

 (Ni): not more than 1.0%, copper (Cu): not more than 1.0%, chromium (Cr): not more than 1.0%, and molybdenum (Mo): 1.0% or less.

Carbon (C): 0.16 to 0.20%

Since C is the element having the greatest influence on the slab solidification behavior, it is necessary to be contained in the steel within an appropriate range. If the content of C is less than 0.16%, there is a problem that the solidification layer strength becomes large at the time of phase transformation at the time of solidification of the slab, causing shrinkage and forming a heterogeneous solidification layer to facilitate cracking on the slab surface, , There is a problem that the carbon equivalent is too large and the hardenability of the welded portion is greatly increased, thereby deteriorating the toughness of the welded portion. Therefore, the content of C in the present invention is preferably 0.16 to 0.20%.

Manganese (Mn): 1.0 to 1.5%

The Mn is an element which is effective for increasing the hardenability of the steel and securing the strength of the steel sheet. In the present invention, however, it is necessary to appropriately limit the content of the Mn to the side for securing toughness of the weld heat affected zone (HAZ). In general, Mn tends to be segregated at the center of the thickness of the steel sheet, although the toughness of the weld heat affected zone is not greatly deteriorated. Since the Mn content is much higher than the average content in the region where Mn is segregated, There is a problem in that a brittle tissue that greatly damages the surface is easily generated. In consideration of this, in the present invention, it is preferable to contain 1.5% or less of high Mn. However, if the content is too low, there is a problem that it becomes difficult to secure the strength of the steel material. Therefore, the lower limit is preferably 1.0%.

Silicon (Si): 0.3% or less (excluding 0)

The Si promotes the morphology of the martensite (MA) by increasing the strength of the steel sheet and suppressing the formation of cementite when the elements necessary for deoxidation of molten steel or the unstable austenite are decomposed, HAZ) is greatly deteriorated. In consideration of this, the content of Si in the present invention is preferably 0.3% or less, and if it exceeds 0.3%, a coarse Si oxide is formed and brittle fracture may occur starting from such inclusions.

Aluminum (Al)): 0.005 to 0.5%

The above-mentioned Al is an element capable of inexpensively deoxidizing molten steel, and is preferably added in an amount of 0.005% or more. However, if the content exceeds 0.5%, there is a problem that nozzle clogging occurs during continuous casting, and the solidified Al can form the martensite on the welded portion, which may result in deterioration of the toughness of the welded portion. The content of Al is preferably 0.005 to 0.5%.

Phosphorus (P): not more than 0.02%

P is an element favoring strength improvement and corrosion resistance, but it is an element which greatly hinders impact toughness. Therefore, it is advantageous to keep P as low as possible, so that the upper limit is preferably 0.02%.

Sulfur (S): not more than 0.01%

Since S is an element that significantly inhibits impact toughness by forming MnS or the like, it is advantageous to keep it as low as possible, and therefore, the upper limit is preferably 0.01%.

Titanium (Ti): 0.005 to 0.02%

The Ti bonds with nitrogen (N) to form fine nitrides, thereby alleviating the coarsening of crystal grains that may occur in the vicinity of the welding fusion wire, thereby suppressing the deterioration of toughness. At this time, if the content of Ti is too low, the effect of suppressing the coarsening can not be sufficiently exerted due to the insufficient number of the Ti nitrides. Therefore, it is preferable that the Ti content is 0.005% or more. However, if the content is excessively large, there is a problem that the effect of the grain boundary fixing is deteriorated due to the formation of a coarse Ti nitride. Therefore, the upper limit is preferably 0.02%.

Niobium (Nb): 0.01 to 0.1%

The Nb precipitates in the form of NbC or Nb (C, N), thereby greatly improving the strength of the base material and the welded portion. In addition, Nb solidified at the time of reheating at a high temperature suppresses transformation of austenite and transformation of ferrite or bainite, thereby exhibiting an effect of making the structure finer. Therefore, in order to secure the strength of the base material even in the case of a post-weld stress relaxation heat treatment such as a storage container, it is preferable that the strength is 0.01% or more. However, when the content exceeds 0.1%, excessive brittle cracks may appear on the edge of the steel material, and the toughness of the welded heat affected zone is greatly lowered. Therefore, it is preferable that the content does not exceed 0.1%.

Nitrogen (N): 0.006 to 0.01%

The N bonds with the above-mentioned Ti to form a fine nitride, thereby alleviating the coarsening of crystal grains that may occur near the welding fusion wire, thereby preventing a decrease in toughness. In order to obtain the above effect, N should be contained in an amount of 0.006% or more. However, if the content is excessively large, there is a problem that the toughness is greatly reduced. Therefore, it is preferable that the content is not more than 0.01%.

The steel sheet of the present invention may further contain, in addition to the above-described alloy composition, elements capable of ensuring favorable physical properties in the present invention. As a preferable example, it is preferable to use an alloy containing not more than 0.006% of Ca, not more than 0.03% of vanadium, not more than 2.0% of nickel, not more than 1.0% of copper, not more than 1.0% of chromium, (Mo): 1.0% or less, and the like. Hereinafter, these will be described in detail.

Calcium (Ca): Not more than 0.006%

The Ca is mainly used as an element which controls the shape of the MnS inclusions and improves the low-temperature toughness. However, excessive Ca addition causes formation of coarse inclusions due to formation and bonding of a large amount of CaO-CaS, thereby deteriorating the weldability of the steel as well as lowering the cleanliness of the steel. Therefore, Ca is preferably 0.006% or less.

Vanadium (V): not more than 0.03%

The above V is low in the temperature to be employed as compared with other alloying elements and is excellent in the effect of precipitating on the weld heat affected zone (HAZ) to prevent the strength from lowering. However, if V is too high, , The content thereof is preferably 0.03% or less.

Nickel (Ni): not more than 2.0%

The Ni is an almost unique element capable of simultaneously improving the strength and toughness of the base material. However, since it is an expensive element, exceeding 2.0% is not only disadvantageous in terms of economy but also deteriorates the weldability. Therefore, it is preferable that the Ni content does not exceed 2.0%.

Copper (Cu): not more than 1.0%

The Cu is an element capable of minimizing the toughness of the base material and improving the strength of the steel. However, when it is added too much, there is a problem that the surface quality of the product is significantly lowered, so it is preferable that the content is 1.0% or less.

Cr (Cr): Not more than 1.0%

The Cr increases the hardenability and has a great effect on the strength improvement. However, if it is added too much, there is a problem that the weldability is significantly lowered. Therefore, it is preferable that the content does not exceed 1.0%.

Molybdenum (Mo): not more than 1.0%

The addition of only a small amount of Mo significantly improves the hardenability and has an effect of suppressing the formation of a ferrite phase and is an element capable of greatly improving the strength. However, if it is added too much, there is a problem that the hardness of the welded portion is greatly increased and the toughness is deteriorated. Therefore, the content thereof is preferably not more than 1.0%.

The steel sheet of the present invention is an iron (Fe) component in addition to the above-mentioned alloying elements. However, the impurities which are not intended from the raw material or the environment of caution can be inevitably incorporated in a normal manufacturing process, and therefore this can not be excluded. Since these impurities can be known to any ordinary technician, they do not fully exploit all of them.

The steel material of the present invention preferably has a surface crack sensitivity index (Cs) defined by the following relational expression 1 of not more than 0.3.

[Relation 1]

Cs = (71.4 x [C] ² ) - (30.3 x [C]) + 3.32

Here, [C] means the weight% value of the content of C.

As mentioned above, C has the greatest influence on the slab solidification behavior. In the present invention, when the C content is less than 0.16%, the surface crack sensitivity index (Cs) of the relational expression 1 exceeds 0.3. That is, when the slab is solidified, the solidification layer strength is large at the time of occurrence of the phase transformation, causing shrinkage, and forming a non-uniform solidification layer, thereby facilitating cracks on the surface of the slab. Therefore, in order to provide a steel material free from surface cracking, it is preferable that the surface crack sensitivity index (Cs) of the above-mentioned relational expression 1 is 0.3 or less. The Cs value in the above-mentioned relational expression 1 is preferably as small as possible, but C is present in the shell, and therefore, the Cs value is preferably more than zero.

On the other hand, the steel of the present invention preferably has a value of Free-N defined by the following relational expression (1)

[Relation 2]

Free-N = [N] - {([Ti] /47.887) x 14.01} - {([B] /10.81) x 14.01}

Here, [N], [Ti], and [B] refer to the weight percent values of N, Ti and B, respectively.

In the present invention, for example, NbC and Nb (C) N type precipitates, which are produced by adding Nb, play a major role in ensuring strength after stress relaxation heat treatment. At this time, N may be combined with Ti, Al, B, and the like to preferentially form another type of precipitate such as TiN or BN, thereby negatively affecting securing the intended Nb precipitate. When free-N is less than 0, Ti and B, which do not sufficiently form a nitrogen-based precipitate, can bond with C to form a coarse precipitate. Therefore, the free- The value of N is preferably greater than zero. The upper limit of the Free-N is not particularly limited, but is preferably 0.008148 or lower.

The steel material of the present invention preferably has a microstructure and a ferrite-pearlite composite structure as a main structure. In addition to the ferrite and pearlite composite structure, it is preferable that the second phase such as bainite or martensite is not generated as much as possible. When the bainite or martensite structure is formed, the physical properties, the physical properties of the weld heat affected zone, and the like are completely different, and it is difficult to realize the steel properties intended in the present invention. It is preferable that the ferrite-pearlite composite structure has an area fraction of 50 to 75% of pearlite and the remainder is ferrite.

Steel material of the present invention is 1.27 × 10 ⁶ per precipitates having a size of stress after annealing heat treatment, the diameter or less 100㎚ performed after substituting hot welding 1㎟ And more than 900 precipitates are distributed in the single crystal grains. Through the distribution of the precipitates, the strength and toughness of the base material can be prevented from lowering even after the stress relaxation heat treatment.

The welded heat affected zone is rapidly heated to a high temperature close to the melting point and then quickly cooled down to room temperature depending on the degree of proximity from the welding point. In this case, a low temperature phase such as bainite or martensite may be generated, Even if ferrite is produced, a kind of microstructure having a high stress is generated in the interior like needle-like ferrite. The microstructure of the welded heat affected zone has a problem that brittleness is generated and it is easily broken in the processing and use environment of the steel. Therefore, in the manufacturing process of the storage tank, the pressure vessel, the building structure and the ship structure, Heat treatment is performed, which is intended to reduce the stress of the welded portion and the heat affected portion to reduce embrittlement, thereby lowering the possibility of fracture in a use environment. The stress relieving heat treatment conditions vary depending on the welding conditions and the thickness of the steel material. For example, A516-70, a medium-temperature pressure vessel steel, is heat-treated at a temperature of 620 ° C for 120 minutes.

The stress relieving heat treatment may have a negative influence on the base material itself, not the welded portion or the heat affected portion. Ferrite, and pearlite, the formation and coarsening of precipitates including carbide can be actively performed by performing the stress relaxation heat treatment at a temperature of 400 to 800 ° C. Such a carbide is proportional to time So that the coarsening of the carbide occurs and the concentration of carbonization in the base structure is reduced, so that the overall strength can be lowered. Therefore, it is necessary to appropriately manage the formation of precipitates containing carbide in order to prevent the strength of the base material from lowering by the welding and stress annealing heat treatment.

On the other hand, when the steel material is broken, propagation generally proceeds along the grain boundary, the soft phase, or the segregation zone. The fine precipitate having a size of 100 nm or less obstructs the propagation of the fracture when the steel material is broken Thereby improving the strength and toughness of the steel material. Since the base structure of the steel material of the present invention has a ferrite-pearlite structure, the relatively soft ferrite structure is vulnerable to fracture, but in many cases, fracture progresses along the pearlite band. Therefore, the fine precipitates are uniformly distributed .

However, even if precipitates are formed in a coarse form such as Fe ₃ C, VC, MoC, and Ce ₂₃ C ₆ or formed into a fine size, coarsening of the precipitates may not contribute to the bad effect of the breaking interference Rather, it acts as a starting point of fracture to reduce strength and toughness, so it is important that the size of the precipitate is fine and appropriately distributed.

Particularly, it is preferable to distribute evenly more than 900 pieces in a single crystal grain, rather than being concentrated at a specific position in a single ferrite or a pearlite, in order to improve strength and impact toughness.

The precipitate of the present invention is preferably Nb-based carbide. More preferably, it is NbC. The Nb-based carbide is mainly produced and grown at a relatively low temperature of 600 to 700 ° C (austenite-to-ferrite transformation point direct lower temperature zone), and in the process, it acts to suppress the strength and ferrite grain growth .

On the other hand, the steel material of the present invention is a steel material having improved sinterability compared to the conventional steel material, and a desired structure can be formed in the steel material without rapid water cooling or the like. However, when the hardness of the steel material is improved and the hard texture is easily formed, the low temperature toughness is often deteriorated. In the present invention, by specifying the preferable structure of the steel material, There is an effect of preventing deterioration of toughness characteristics.

The steel material of the present invention is excellent not only in tensile strength of the base material of 500 MPa or more and a Charpy impact energy at 0 캜 of 150 J or more even after a stress annealing process (for example, at 620 캜 for 120 minutes) The microstructure of the welded heat affected zone (HAZ) has an impact martensite fraction of 3% or less and excellent impact toughness at a temperature of 0 ° C and a impact energy of 100 J or more.

Hereinafter, a method of manufacturing a steel material according to the present invention will be described in detail. The following production method is a preferred example for producing the steel sheet of the present invention, but the present invention is not limited thereto.

The manufacturing method of the present invention includes preparing a steel slab satisfying the above-described alloy composition, heating, and hot rolling and cooling the steel slab. Hereinafter, each process will be described in detail.

First, a steel slab having the above-described alloy composition is prepared, and then the steel slab is heated. At this time, it is preferable to heat in a temperature range of 1050 to 1250 캜. The heating is preferably carried out at a temperature of 1050 DEG C or higher, in order to solidify Ti and / or Nb carbonitride formed during casting. That is, in order to sufficiently solidify the Ti and / or Nb carbonitride formed during the casting, it is necessary to heat at 1050 ° C or higher. However, in case of heating to an excessively high temperature, the austenite may be coarsened, and therefore, it is preferable to limit the reheating temperature to 1250 ° C or less.

The heated steel slab is hot-rolled. The hot-rolled steel sheet is preferably subjected to rough rolling of the heated steel slab under normal conditions, and then hot-rolled at a predetermined temperature to produce a hot-rolled steel sheet. At this time, the hot finish rolling is performed at 910 占 폚 or lower. The hot finish rolling is for transforming the austenite structure into a nonuniform microstructure. When the hot finish rolling temperature exceeds 910 占 폚, coarse texture is formed and the impact toughness is deteriorated. More preferably in a temperature range of 850 to 910 < 0 > C. If the rolling finish temperature is lowered to less than 850 占 폚, there is a problem that it becomes difficult to control the shape of the plate material.

Preferably, the hot-rolled steel sheet obtained by the hot-rolling is cooled, and the cooling is preferably performed at a low-speed cooling lower than a normal air-cooling level. Particularly, it is preferable to cool at a cooling rate of 20 ° C / Hr or less in a temperature range of 800 to 435 ° C. This makes it possible to stably secure the optimum strength and toughness of the steel material of the present invention. The temperature interval is a main temperature interval in which precipitates are generated and grown. The cooling rate is preferably 1 ° C / Hr or more. In order to ensure the strength and toughness of the steel material before and after the stress relaxation annealing process, the minimum cooling power can be ensured through the cooling process in order to secure the target fraction and distribution of precipitates. As a method of implementing the slow cooling as described above, there is a method of separately sending equipment, and a method of stacking steel plates of similar dimensions after completion of hot rolling without a separate warming equipment can be achieved.

Hereinafter, the present invention will be described in more detail by way of examples. It should be noted, however, that the embodiments described below are for illustrating and embodying the present invention, and not for limiting the scope of the present invention. And the scope of the present invention is determined by the matters described in the claims and the matters reasonably deduced therefrom.

(Example)

Steel slabs having the composition shown in the following Table 1 were prepared, and then each steel slab was rolled under the conditions shown in Table 2 below and then cooled to produce a hot-rolled steel sheet.

division	C	Mn	Si	Al	P	S	Ti	Nb	N	Ni	Cu	Cr	Mo	V	Ca	Equation (1)	Equation (2)
Steel A	0.166	1.41	0.18	0.021	0.011	0.001	0.016	0.014	0.0077	-	-	-	-	0.0145	0.0019	0.26	0.0030
Grade B	0.185	1.38	0.14	0.038	0.009	0.002	0.018	0.016	0.0079	-	-	-	-	-	-	0.16	0.0026
Grade C	0.171	1.38	0.167	0.024	0.01	0.002	0.018	0.014	0.0065	-	-	-	-	0.016	-	0.23	0.0012
Grade D	0.17	1.3	0.15	0.035	0.013	0.002	0.017	0.007	0.0075	0.4	-	0.3	-	-	-	0.23	0.0025
Steel E	0.19	1.45	0.17	0.013	0.012	0.002	0.02	0.004	0.008	0.3	0.2	-	0.1	-	-	0.14	0.0021
Grade F	0.18	1.42	0.26	0.031	0.005	0.0004	0.012	0.026	0.0027	0.23	0.17	0.052	0.089	0.015	0.0017	0.18	-0.0008
Grade G	0.172	1.3	0.4	0.024	0.013	0.005	0.015	0.04	0.007	-	-	0.5	-	-	-	0.22	0.0026

(The components in Table 1 are% by weight, and the remainder is composed of Fe and unavoidable impurities.) In Table 1, the expressions (1) and (2) mean relational expressions 1 and 2, respectively.

Steel grade	Slab heating temperature (℃)	Hot finish rolling temperature (℃)	Cooling rate (° C / Hr)	Remarks
Steel A	1137	905	20	Inventory 1
	1142	895	20	Inventory 2
	1146	931	20	Comparative Example 1
Grade B	1180	890	20	Inventory 3
	1024	889	20	Comparative Example 2
	1270	893	20	Comparative Example 3
Grade C	1180	867	20	Honorable 4
	1162	920	20	Comparative Example 4
	1154	880	60	Comparative Example 5
Grade D	1170	952	20	Comparative Example 6
	1175	909	20	Comparative Example 7
	1140	873	20	Comparative Example 8
Steel E	1100	873	20	Comparative Example 9
	1185	884	20	Comparative Example 10
Grade F	1146	878	20	Comparative Example 11
	1152	875	20	Comparative Example 12
	1137	867	20	Comparative Example 13
Grade G	1170	880	20	Comparative Example 14
	1157	878	20	Comparative Example 15

The steel material thus produced was welded at 200 kJ / cm and subjected to a stress relaxation heat treatment in which the steel material was held at 620 占 폚 for 120 minutes.

After the heat treatment, the microstructure of the base material, the distribution of precipitates of 100 nm or less and the number of precipitates present in the crystal grains were measured, and tensile strength and impact toughness were measured. The results are shown in Table 3. The impact toughness and the martensite fraction of the welded heat affected zone were measured, and the results are shown in Table 3. The impact toughness was measured by Charpy V-notch impact test at 0 캜. The above-mentioned on-road martensite analysis was carried out by Le-Pera etching and then the position and relative area fraction estimated by using point-counting method were measured.

In Table 3, F means ferrite and P means pearlite.

On the other hand, Fig. 2 shows the size (nm) of the precipitate of the precipitate of Inventive Example 1 observed by TEM. As shown in FIG. 2, Inventive Example 1 of the present invention shows that NbC precipitates of 100 nm or less are uniformly formed. On the other hand, FIG. 3 shows the precipitate of Comparative Example 6 observed by TEM and the size (nm) of the precipitate. In Comparative Example 6, coarse FeC precipitates were formed.

From the results of Table 3, it can be seen that not only the base material secures high strength and impact toughness but also high impact toughness of the weld heat affected zone (HAZ) even after the stress relaxation heat treatment after welding in the case of the invention example. That is, the steel material of the present invention can secure the toughness of the HAZ even when it is welded by large heat, and can produce a steel material free from defects such as surface cracks.

Comparative Examples 1 and 4 satisfied the alloy composition of the present invention, but the hot finish rolling temperature was too high to ensure sufficient toughness of the base material due to microstructure coarsening. Figs. 1 (a) and 1 (b) are photographs showing the microstructure of the base material of Inventive Example 1 and Comparative Example 1, respectively. All of the ferrite and pearlite were the same, but in the case of Comparative Example 1, it was considered that the impact strength was affected by the coarsening of the particle size.

Comparative Examples 2 and 3 also satisfy the alloy composition of the present invention, but the slab heating temperature is outside the scope of the present invention, so that the elements inhibiting austenite grain growth at high temperatures such as Nb are not sufficiently dissolved, So that the strength and impact toughness of the base material are lowered. In the case of Comparative Example 5, since the cooling rate of the steel after the hot rolling was out of the range proposed in the present invention, precipitates of the steel could not be secured and the strength in the present invention could not be obtained.

On the other hand, in Comparative Examples 6 to 10, since the content of Nb in the steel material is not sufficient, C is deposited on the coarse cementite particles or MoC or the like, so that sufficient strength can not be ensured and the toughness of HAZ can not be secured. In Comparative Examples 11 to 13, the N content of the steel material did not fall within the range of the present invention, and coarse precipitates such as MoC, Fe ₃ C and VC, etc., And the distribution of the precipitates is different from the present invention, so that the impact toughness of the HAZ is heated. In Comparative Examples 14 and 15, the Si content exceeded the range of the present invention, and the physical properties of the base material were consistent with the intended range of the present invention, but the fraction of the martensite on the surface of the HAZ was excessive and the impact toughness of the HAZ was weakened .

Claims (10)

1. Al: 0.005-0.5%, P: 0.02% or less, S: 0.01% or less, Ti: 0.1 to 0.20%, C: 0.16 to 0.20%, Mn: 1.0 to 1.5% 0.005 to 0.02%, Nb: 0.01 to 0.1%, and N: 0.006 to 0.01%

    At least one selected from the group consisting of Ca: not more than 0.006%, V: not more than 0.03%, Ni: not more than 2.0%, Cu: not more than 1.0%, Cr: not more than 1.0%, and Mo: not more than 1.0% Containing impurities,

   The microstructure is composed of a ferrite-pearlite composite structure,

   After the welding and stress relieving heat treatment, the microstructure had a precipitate of 100 nm or less of 1.27 x 10 < ⁶ > And more than 900 in a single grain. Welded heat affected zone High strength steel excellent in toughness.
2. The method according to claim 1,

   The precipitates are Nb-based carbides, and are excellent in toughness at the weld heat affected zone.
3. The method according to claim 1,

   Wherein the steel material has a surface crack sensitivity index (Cs) of 0.3 or less as defined by the following relational expression 1:

   [Relation 1]

   Cs = (71.4 x [C] ² ) - (30.3 x [C]) + 3.32

   ([C] is the weight% content of the component)
4. The method according to claim 1,

   The steel material is excellent in toughness of weld heat affected zone where Free-N exceeding 0, which is defined by the following relational expression 2, is high strength steel.

   [Relation 2]

   Free-N = [N] - {([Ti] /47.887) x 14.01} - {([B] /10.81) x 14.01}

   ([N], [Ti], and [B] are weight% contents of the corresponding components)
5. The method according to claim 1,

   The steel material has a tensile strength of 500 MPa or more and a Charpy impact absorption energy of 150 J or more at 0 캜 even after the stress relaxation heat treatment.
6. The method according to claim 1,

   The high-strength and high-strength steel material excellent in weld heat-affected portion toughness having an area fraction of martensite (MA) of not more than 3% as a weld heat affected portion of the steel.
7. The method according to claim 1,

   The steel material has a Charpy impact absorption energy of 100 J or more at 0 ° C after welding heat welding, and a welded heat-affected part has excellent toughness at weld heat-affected zone.
8. Al: 0.005-0.5%, P: 0.02% or less, S: 0.01% or less, Ti: 0.1 to 0.20%, C: 0.16 to 0.20%, Mn: 1.0 to 1.5% 0.005 to 0.02%, Nb: 0.01 to 0.1%, and N: 0.006 to 0.01%

    At least one selected from the group consisting of Ca: not more than 0.006%, V: not more than 0.03%, Ni: not more than 2.0%, Cu: not more than 1.0%, Cr: not more than 1.0%, and Mo: not more than 1.0% Preparing a steel slab containing impurities;

   Heating the steel slab in a temperature range of 1050 to 1250 占 폚;

   Hot rolling the heated steel slab at a hot finish rolling temperature of 910 占 폚 or lower; And

   After the hot rolling, cooling is performed at a cooling rate of 20 DEG C / Hr or less

   Wherein the welded heat affected zone toughness is excellent.
9. The method of claim 8,

   Wherein the hot rolling is performed in a temperature range of 850 to 910 占 폚.
10. The method of claim 8,

    Wherein the cooling is performed at a cooling rate of 20 ° C / Hr or less in a temperature range of 800 to 435 ° C.

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