EP3730644B1

EP3730644B1 - High-strength steel with excellent toughness of welding heat affected zone and manufacturing method thereof

Info

Publication number: EP3730644B1
Application number: EP18891918.7A
Authority: EP
Inventors: Jae-Yong CHAE; Sang-Deok Kang
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2017-12-24
Filing date: 2018-12-21
Publication date: 2024-06-26
Anticipated expiration: 2038-12-21
Also published as: US20200325562A1; KR20190077186A; CN111542632A; KR101999016B1; EP3730644A4; EP3730644C0; EP3730644A1; WO2019125075A1

Description

[Technical Field]

The present invention relates to a structural steel used as a material for storage tanks, pressure vessels, building structures, ship structures, or the like, and more particularly, to a high strength steel having excellent toughness in a heat-affected zone and a method of manufacturing the same.

[Background Art]

When constructing structures such as storage tanks, pressure vessels, building structures, and ship structures using structural steel, a large amount of welding is involved. For this reason, not only the performance of the base material, but also the efficiency of welding and the stability of the welding structure must be secured. To this end, it is necessary to suppress the growth of austenite grains in the heat-affected zone (HAZ) as much as possible to keep a final transformation structure fine.
As a means to solve this, a technique has been proposed to delay the grain growth of the heat-affected zone during welding, by appropriately distributing Ti-based carbon and nitrides or the like, which are stable at high temperature, in steel.
As an example, Patent Document 1 is a representative technique using a precipitate of TiN and relates to a structural steel material having an impact toughness of about 200J at 0°C (about 300J in a base material) when a heat input of 100J/cm (a highest heating temperature of 1400°C) is applied. In the above technique, Ti/N is practically managed to be 4-12 and, thus, TiN precipitates of 0.05 µm or less are 5.8×10³pieces/mm² to 8.1×10⁴pieces/mm², and in addition, TiN precipitates of 0.03 to 0.2 µm are 3.9×10³pieces/mm² to 6.2×10⁴pieces/mm², in refining ferrite to secure toughness of a weld portion.
However, in Patent Document 1 a problem, in that cracks may be severely generated on the slab surface during continuous casting by forming excessive carbon and nitride, is caused. When a thick plate product is produced using the slabs having a large number of surface cracks as above, there is also a problem in which cracks or the like also occur in the surface of the final product. Therefore, there is a great possibility that problems such as surface repair or the like may occur, or defective products incapable of being repaired may be manufactured.
KR 2012 0014413 A discloses a high strength steel plate for welding structure with superior toughness for high heat input welding, the steel plate including by weight % C: 0.03~0.2; Si: 0.01~0.5; Mn: 0.6~2.0; Ti: 0.005~0.2; Al: 0.2~1.0; B: 0.0005~0.01; N: 0.008~0.030; P: less than 0.03; S: less than 0.03; O: 0.03; and the remainder Fe and incidental impurities. The steel material has microstructure consisting of 70~75 weight% ferrite, 20~25 weight% pearlite, and the remainder bainite and martensite.

(Patent Document 1) Japanese Patent Laid-open Publication No. 1999-140582
(Patent Document 2) KR 2012 0014413 A

[Disclosure]

[Technical Problem]

An aspect of the present invention is to provide a steel material capable of securing an excellent heat-affected zone (HAZ) while having excellent strength and toughness of a base material even after welding and a stress relief heat treatment, and a method of manufacturing the same.

[Technical Solution]

The invention is disclosed in the appended claims.
According to aspects of the invention, there is provided a high-strength steel having excellent toughness in a heat-affected zone as recited in Claim 1, and a method of manufacturing a high-strength steel having excellent toughness in a heat-affected zone as recited in Claim 6. Optional features of the invention are recited in the dependent claims.

[Advantageous Effects]

According to the invention, a steel material is provided having excellent toughness in a heat-affected zone resulting from large heat input welding without lowering strength and toughness of a base material even after a stress relief heat treatment after welding. In addition, since the strength of the base material is maintained even when stress annealing is performed, the steel may be suitably used in a storage tank, a pressure vessel, a structure, and the like. In addition, since the steel according to the present invention has no defects such as surface cracking, the steel may be suitably used as a structural steel material.

[Description of Drawings]

FIGS. 1A and 1B are images of microstructures of Example 1 and Comparative Example 1 of the present disclosure.
FIG. 2 illustrates the size and shape of NbC precipitates observed with a transmission electron microscope (TEM) in Inventive Example 1 of the present disclosure.
FIG. 3 illustrates the size and shape of Fe₃C precipitates observed with a transmission electron microscope (TEM) in Comparative Example 6 of the present disclosure observed.

[Best Mode for Invention]

The inventors have studied in depth to fundamentally solve the problem of defects such as cracks on the steel surface when manufacturing thick steel materials for use as existing structural steel materials, and have confirmed that the heat-affected zone having excellent toughness could be secured by controlling the microstructure of the heat-affected zone during welding, as well as securing base material strength and toughness, when optimizing the steel composition and manufacturing conditions, in completing the present invention.
In detail, since excellent toughness of the heat-affected zone (HAZ) may be secured during large heat input welding such as submerged arc welding, a steel according to the present disclosure may have an effect in which the steel may be suitably applied as a structural steel.
On the other hand, as a result of finding a method to prevent the strength reduction of the base material that may occur after the stress relief heat treatment, which is generally performed to stabilize the material of the welded hardened structure when manufacturing a storage tank, a pressure vessel or the like, it was confirmed that the strength may be secured after heat treatment when some alloy components are added to generate fine precipitates; and thus the present invention has been completed. Therefore, there is an effect that may be suitably applied not only to existing structural steels, but also to storage tanks, pressure vessels and the like.
Hereinafter, the present invention will be described in detail.
First, an alloy composition of a steel according to the present invention will be described in detail. The steel according to the present invention includes, by weight% (hereinafter, %), carbon (C): 0.16 to 0.20%, manganese (Mn): 1.0 to 1.5%, silicon (Si): 0.3% or less (excluding 0), aluminum ( Al): 0.005 to 0.5%, phosphorus (P): 0.02% or less, sulfur (S): 0.01% or less, titanium (Ti): 0.005 to 0.02%, niobium (Nb): 0.01 to 0.1%, and nitrogen (N): 0.006 to 0.01%.
The steel optionally includes one or more selected from the group consisting of calcium (Ca): 0.006% or less, vanadium (V): 0.03% or less, nickel (Ni): 2.0% or less, copper (Cu): 1.0% or less, chromium (Cr): 1.0% or less, and molybdenum (Mo): 1.0% or less.

Carbon (C): 0.16 to 0.20%

Since C is an element having a greatest influence on the slab solidification behavior, it needs to be contained in the steel within an appropriate range. If the content of C is less than 0.16%, the strength of a solidified layer increases when the phase transformation occurs during slab solidification. Therefore, there is a problem in which the occurrence of cracking on the slab surface may be facilitated by causing shrinkage and forming a non-uniform solidification layer. On the other hand, if the content thereof exceeds 0.20%, the carbon equivalent becomes too large. Therefore, in this case, there is a problem in that the toughness of the weld portion deteriorates as the hardenability of the weld portion is greatly increased. Therefore, in the present invention, the content of C is 0.16 to 0.20%.

Manganese (Mn): 1.0 to 1.5%

The Mn is an element useful for securing the strength of the steel sheet by increasing the hardenability of the steel, but in the present invention, it is necessary to appropriately limit the content thereof to secure toughness of the heat-affected zone (HAZ). In general, Mn does not significantly deteriorate the toughness of the heat-affected zone, but tends to be segregated in the center of the thickness of the steel sheet. The Mn segregated portion as described above has a very high Mn content, compared to the average content, and thus, there is a problem of easily generates a brittle structure that greatly harms the toughness of the weld-heat-affected zone. In consideration thereof, in the present invention, Mn is in amount of 1.5% or less. However, if the content is too low, there is a problem that it is difficult to secure the strength of the steel, and thus, the lower limit is set to 1.0%.

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

The Si increases the strength of the steel sheet and is an element necessary for deoxidation of molten steel, but Si inhibits the formation of cementite when unstable austenite is decomposed, and thus promotes a martensite austenite constituent (MA), in which there is a problem in which toughness of the heat-affected zone (HAZ) is significantly lowered. Considering this, in the present invention, the Si content is 0.3% or less, and if it exceeds 0.3%, coarse Si oxide is formed, and unpreferably, brittle fracture may occur around such inclusions.

Aluminum (Al): 0.005 to 0.5%

The Al is an element capable of deoxidizing molten steel inexpensively, and for this use, it is added at 0.005% or more. However, if the content exceeds 0.5%, there is a problem of causing nozzle clogging during continuous casting, and the solidified Al may form the martensite austenite constituent in the weld portion and may result in a decrease in toughness of the welding portion. The Al content is 0.005 to 0.5%.

Phosphorus (P): 0.02% or less

The P is an element that is advantageous for strength improvement and corrosion resistance, but since it is an element that greatly inhibits impact toughness, the content thereof may be advantageously managed to be as low as possible, and thus, the upper limit is 0.02%.

Sulfur (S): 0.01% or less

Since S is an element that greatly inhibits impact toughness by forming MnS or the like, the content thereof may be advantageously managed to be as low as possible, and thus, the upper limit is set to 0.01%.

Titanium (Ti): 0.005 to 0.02%

The Ti is combined with nitrogen (N) to form a fine nitride, thereby reducing grain coarsening that may occur near the welding melting line, to suppress a decrease in toughness. In this case, if the content of Ti is too low, the number of Ti nitrides is insufficient, and thus, the effect of suppressing coarsening is not sufficiently exhibited, and thus, Ti is in an amount of 0.005% or more. However, if the content is too excessive, there is a problem that the grain boundary fixation effect is lowered due to the generation of coarse Ti nitride, and thus, the upper limit thereof is 0.02%.

Niobium (Nb): 0.01 to 0.1%

The Nb is precipitated in the form of NbC or Nb(C, N) to greatly improve the strength of the base material and the weld portion. In addition, solidified Nb during reheating at a high temperature suppresses recrystallization of austenite and transformation of ferrite or bainite, thereby exhibiting an effect that the structure is refined. Therefore, NB is 0.01% or more to secure the strength of the base material even after undergoing stress relief heat treatment after welding, such as for a storage container. However, if the content exceeds 0.1% and is excessively added, brittle cracks may appear at the corners of the steel and greatly reduce the toughness of the heat-affected zone, and thus, NB is limited to 0.1%.

Nitrogen (N): 0.006 to 0.01%

The N is combined with the above-described Ti to form a fine nitride to alleviate grain coarsening that may occur near the weld melting line to prevent toughness from deteriorating. To obtain the above effects, it is necessary to contain N in an amount of 0.006% or more. However, if the content is too excessive, there is a problem of significantly reducing toughness, and thus, N does not exceed 0.01%.
In addition to the alloy composition described above, the steel sheet according to the present invention may further include elements capable of securing advantageous physical properties in the present invention. As a detailed example, the steel sheet may further include calcium (Ca): 0.006% or less, vanadium (V): 0.03% or less, nickel (Ni): 2.0% or less, copper (Cu): 1.0% or less, chromium (Cr): 1.0% or less, molybdenum (Mo): 1.0% or less.

Calcium (Ca): 0.006% or less

The Ca is mainly used as an element that controls the shape of the MnS inclusion and improves low-temperature toughness. However, excessive Ca addition causes a large amount of CaO-CaS to form and combine to form coarse inclusions, thus impairing the cleanliness of the steel and spoiling weldability in the field. Therefore, Ca may be 0.006% or less.

Vanadium (V): 0.03% or less

The V has a solid-solution temperature lower than other alloying elements, and has an excellent effect of preventing a decrease in strength by depositing in a heat-affected zone (HAZ). However, if the content is too excessive, there is a problem of rather decreasing toughness. Therefore, V may be 0.03% or less.

Nickel (Ni): 2.0% or less

Ni is almost the only element capable of simultaneously improving the strength and toughness of the base material, but since it is an expensive element, exceeding 2.0% is not only very disadvantageous in terms of economy, but also has a problem of deterioration of weldability. Therefore, it may not exceed 2.0% when the Ni is added.

Copper (Cu): 1.0% or less

The Cu is an element capable of improving the strength of steel while significantly reducing a decrease in toughness of the base material. However, if it is added excessively, there is a problem of significantly deteriorating the surface quality of the product, and therefore, copper may be in an amount of 1.0% or less.

Chromium (Cr): 1.0% or less

The Cr has a great effect on strength improvement by increasing hardenability. However, if it is added excessively, there is a problem that the weldability is greatly deteriorated, and thus, Cr may not exceed 1.0%.

Molybdenum (Mo): 1.0% or less

The Mo has an effect of inhibiting the formation of a ferrite phase by greatly improving the hardenability even with a relatively small amount, and is an element capable of greatly improving the strength. However, if it is added excessively, there is a problem of significantly increasing hardness of the weld portion and inhibiting the toughness, and thus, Mo may not exceed 1.0%.
The steel sheet according to the invention has as balance iron (Fe). However, in the normal manufacturing process, unavoidable impurities may be inevitably mixed from the raw material or a surrounding environment, and cannot be excluded. These impurities are known to those skilled in the art, and descriptions thereof are not all provided in detail.
The steel according to the invention may have a surface crack sensitivity index (Cs) of 0.3 or less, which is defined by the following relational expression 1. $Cs = (71.4 \times {[C]}^{2}) - (30.3 \times [C]) + 3.32,$
where [C] indicates the weight percent value that is the content of C.
As described above, C is an element that has a greatest influence on the slab solidification behavior, and if the C content is less than 0.16% in the present invention, the surface crack sensitivity index (Cs) of the relational expression 1 exceeds 0.3. For example, when the solidification of the slab occurs, the strength of the solidification layer is relatively great at the time of occurrence of phase transformation, causing shrinkage, and forming a non-uniform solidification layer to facilitate crack generation on the slab surface. Therefore, to provide a steel material without surface cracking, the surface crack sensitivity index (Cs) of the relational expression 1 may be 0.3 or less. The Cs value of the relational expression 1 is preferably as low as possible, but since C is present in the steel, the Cs value may be preferably greater than 0.
On the other hand, in the case of the steel according to the invention, the value of Free-N defined by the following relational expression 2 may be greater than zero. Free-N = [N] - {([Ti]/47.887)×14.01} - {([B]/10.81)×14.01}
In this case, the [N], [Ti], and [B] respectively indicate the content weight percent value of each of N, Ti and B.
In the present invention, as an example of Nb precipitates generated by the addition of Nb, NbC, Nb(C)N-type precipitates, etc., play a major role in securing strength after a stress relief heat treatment. At this time, N is combined with Ti, Al, B, etc., to preferentially form another type of precipitate, such as TiN, BN, etc., thereby negatively affecting securing the intended Nb precipitate. Accordingly, if the free-N is less than 0, Ti and B that do not form sufficient nitrogen-based precipitates may be combined with C to form coarse precipitates. Therefore, the value of free N defined by the foregoing relational expression 2 may be greater than zero. The upper limit of the Free-N is not particularly limited, but may be preferably 0.008148 or less.
The steel according to the present invention has a ferrite-pearlite composite structure as a microstructure, as a main structure. Secondary phases such as bainite, martensite, etc are not produced. When the bainite or martensite structure is formed, physical properties, heat-affected zone properties, and the like are completely differently changed, and thus it is difficult to implement steel properties intended in the present invention. The ferrite-pearlite composite structure may have pearlite of 50 to 75% in area fraction, and the remainder thereof is ferrite.
In the steel according to the present invention after the stress relief heat treatment performed after high-input heat welding, precipitates of at least 1.27×10⁶ precipitates per 1 mm² having a diameter of 100 nm or less, and 900 or more precipitates in a single crystal grain, are distributed. Through the distribution of the precipitates, the strength and toughness of the base material may be prevented from being degraded even after a stress relief heat treatment.
In the case of large heat input welding, in the heat-affected zone, depending on the degree of proximity from the welding point, a most neighboring part is rapidly heated to a high temperature close to the melting point, and then rapidly cooled to room temperature. At this time, a low-temperature phase such as bainite or martensite may be generated, and even when ferrite is generated, a microstructure type having a high stress therein, such as acicular ferrite, is generated. The microstructure of the heat-affected zone has a problem of easily breaking in the processing or use environment of steel due to embrittlement occurrence. Therefore, in the manufacturing process of storage tanks, pressure vessels, building structures, ship structures, etc., the stress relief heat treatment of the weld portion is performed, which relieves stress of the weld portion and the heat-affected zone to reduce embrittlement and reduce the possibility of breakage that may occur in the use environment. The stress relief heat-treatment conditions are diverse depending on the welding conditions and the thickness of the steel. For example, in the case of A516-70, a pressure vessel steel material for medium and normal temperature, heat treatment is performed at a temperature of 620°C for 120 minutes.
The stress relief heat treatment may have a negative effect on the base material itself, not the weld portion or the heat-affected zone. In the case of a steel material composed of microstructures such as ferrite and pearlite, when stress relief heat treatment at a level of 400 to 800°C is performed, generation and coarsening of precipitates containing carbides may occur actively. In the case of these carbides, carbide coarsening occurs in proportion to time, and a decrease in carbonization concentration in the matrix structure occurs, thereby causing a decrease in overall strength. Therefore, it is necessary to appropriately manage the formation of precipitates containing carbides to prevent the strength from being degraded by the welding and stress relief heat treatment.
On the other hand, when a fracture occurs in the steel, propagation usually proceeds along a grain boundary, a soft phase, or a segregation zone, and the fine precipitate having a size of 100 nm or less interferes with the propagation of the fracture when the steel breaks, and thus has the effect of improving the strength and toughness of the steel. Since the matrix structure of the steel according to the present invention has ferrite-pearlite, a relatively soft ferrite structure is susceptible to fracture; but in many cases the fracture also proceeds along the pearlite band, and it may be preferable that the fine precipitates are evenly distributed regardless of the matrix structure.
However, in the case in which precipitates are produced in coarse form, such as Fe₃C, VC, MoC, Ce₂₃C₆, or the like, or even if the precipitates are formed in a fine size, in the case in which coarsening of the precipitate occurs, precipitates do not contribute significantly to the disturbance of propagation, and furthermore, rather may act as a starting point for fracture, to serve as reducing strength and toughness. Hence it is important that the size of the precipitate is fine and the precipitates are properly distributed.
In detail, 900 or more precipitates are evenly distributed in a single crystal grain, thereby improving the strength and impact toughness.
The precipitate according to the invention is an Nb-based carbide. In more detail this may be NbC. The Nb-based carbide is produced and grown in a relatively low temperature zone of 600 to 700°C (in a temperature zone directly below the ferrite transformation point in austenite), and serves to suppress a decrease in strength and ferrite grain growth in the process thereof.
On the other hand, the steel according to the present invention is a steel material with improved quenchability, as compared to a related art steel material, and a required structure may be formed inside the steel material without rapid water cooling or the like. However, in a case in which quenchability of a steel material is improved and a hard structure is easily formed therein, low-temperature toughness deteriorates in most cases. Therefore in the present invention, by defining the preferred structure shape of the steel material, even if the quenchability of the steel material is improved there is an effect of preventing deterioration of low-temperature toughness characteristics.
The steel according to the present invention may have excellent tensile strength of 500 MPa or more and Charpy impact energy at 0°C of 150 J or more, even after the stress relief heat treatment (for example, 120 minutes at 620°C) after fabrication of the welded structure. Furthermore, the steel has excellent impact toughness in which the fraction of martensite austenite constituent in the microstructure of the heat-affected zone (HAZ) is 3% or less and the Charpy impact energy at 0°C is 100 J or more.
Hereinafter, a method of manufacturing a steel according to the present invention will be described in detail.
The manufacturing method according to the invention includes preparing a steel slab that satisfies the above-described alloy composition, heating, hot rolling and cooling the steel slab. Hereinafter, respective processes 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 may be preferable to heat the steel slab in the temperature range of 1050 to 1250°C. The heating may be preferably performed at 1050°C or higher, to solidify Ti and/or Nb carbon/nitride formed during casting. For example, it is necessary to heat the steel slab at 1050°C or higher to sufficiently solidify Ti and/or Nb carbon and nitride formed during casting. However if heating to an excessively high temperature, austenite may be coarsened and thus it may be preferable to limit the reheating temperature to 1250°C or lower in consideration thereof.
The heated steel slab is hot rolled. The hot rolling is performed to produce a hot rolled steel sheet by performing hot finish rolling at a predetermined temperature after roughly rolling the heated steel slab under normal conditions. At this time, the hot finish rolling is performed at 910°C or lower. The hot finish rolling is for transforming the austenite structure into a non-uniform microstructure, and if the hot finish rolling temperature exceeds 910°C, a coarse structure is formed and impact toughness is deteriorated. More advantageously, the hot finish rolling may be more preferably performed at a temperature in a range of 850 to 910°C. If the rolling termination temperature is lowered to less than 850°C, there is a problem that it is difficult to control the shape of a plate material.
The hot rolled steel sheet obtained by the hot finish rolling is cooled at a low speed lower than a normal air cooling level respectively at a cooling rate of 20°C/Hr or less at a temperature in a range of 800 to 435°C. In this case, the steel having optimum strength and toughness according to the present invention is obtained. The temperature range is a main temperature section in which precipitates are generated and grown. The cooling rate may be preferably 1°C/Hr or more. To secure a target fraction and distribution of precipitates to secure the strength and toughness of the steel according to the present invention before and after stress relief heat treatment, a minimum thermal driving force may be secured through the cooling process. As a method of implementing the slow cooling as described above, there is also a method of using a separate cold storage facility, or a method of stacking, in multiple stages, steel sheets of similar dimensions after hot rolling without additional thermal insulation.

[Mode for Invention]

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

(Example)

After preparing a steel slab having the composition of the components illustrated in Table 1 below, each steel slab was rolled under the conditions of Table 2 below and then cooled to prepare a hot rolled steel sheet.

[Table 1]

Cla ssi fic ati on	C	Mn	Si	Al	P	S	Ti	Nb	N	Ni	Cu	Cr	Mo	V	Ca	Exp res sio n (1)	Exp res sio n (2)
Ste el Gra de A	0.1 66	1.4 1	0.1 8	0.0 21	0.0 11	0.0 01	0.0 16	0.0 14	0.0 077	-	-	-	-	0.0 145	0.0 019	0.2 6	0.0 030
Ste el Gra de B	0.1 85	1.3 8	0.1 4	0.0 38	0.0 09	0.0 02	0.0 18	0.0 16	0.0 079	-	-	-	-	-	-	0.1 6	0.0 026
Ste el	0.1 71	1.3 8	0.1 67	0.0 24	0.0 1	0.0 02	0.0 18	0.0 14	0.0 065	-	-	-	-	0.0 16	0.0 -	0.2 3	0.0 012
Gra de C
Ste el Gra de D	0.1 7	1.3	0.1 5	0.0 35	0.0 13	0.0 02	0.0 17	0.0 07	0.0 075	0.4	-	0.3	0.3 -	-	-	0.2 3	0.0 025
Ste el Gra de E	0.1 9	1.4 5	0.1 7	0.0 13	0.0 12	0.0 02	0.0 2	0.0 04	0.0 08	0.3	0.2	0.2 -	0.1	0.1 -	-	0.1 4	0.0 021
Ste el Gra de F	0.1 8	1.4 2	0.2 6	0.0 31	0.0 05	0.0 004	0.0 12	0.0 26	0.0 027	0.2 3	0.1 7	0.0 52	0.0 89	0.0 15	0.0 017	0.1 8	- 0.0 008
Ste el Gra de G	0.1 72	1.3	0.4	0.0 24	0.0 13	0.0 05	0.0 15	0.0 4	0.0 07	-	-	0.5	0.5 -	-	-	0.2 2	0.0 026

(The components in Table 1 are weight%, and the rest are composed of Fe and unavoidable impurities. On the other hand, in Table 1, expressions (1) and (2) mean Relational Expressions 1 and 2, respectively.)

[Table 2]

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

After welding at 200 kJ/cm for the steel prepared as described above, a stress relief heat treatment was performed at 620°C for 120 minutes.

After the heat treatment, the microstructure of the base material and the distribution of precipitates of 100 nm or less and the number of precipitates in the crystal grains were measured, and the tensile strength and impact toughness were measured, and the results are illustrated in Table 3. In addition, the impact toughness of the heat-affected zone and the fraction of the martensite austenite constitute were measured and the results are illustrated in Table 3. On the other hand, the impact toughness was measured by performing a Charpy V-notch impact test at 0°C. In the above-described martensite austenite constitute analysis, after performing Le-Pera etching, a point-counting method was used to measure the estimated position and relative area fraction of the martensite austenite constituent.

[Table 3]

Class ifica tion	Base material			Base material extract				HAZ
Class ifica tion	Tensi le stren gth (MPa)	Impact toughn ess (J, @0°C)	Mic ros tru ctu re	Type	Fracti on (area% )	Distrib ution (p cs/mm²)	Numbe r(pcs ) of preci pitat es in grain)	Impact toughn ess (J, @0°C)	MA Fract ion (area %)
Inven tive Examp le 1	522	194	F+P	NbC	1	1.7×10⁶	1214	125	1.7
Inven tive Examp le 2	519	183	F+P	NbC	1	1.6×10⁶	1139	111	2.2
Compa rativ e Examp le 1	550	87	F+P	NbC	1	0.8×10⁶	574	less than 50	4.0
Inven tive Examp le 3	504	193	F+P	NbC	1	1.6×10⁶	1166	104	2.5
Compa rativ e Examp le 2	473	91	F+P	NbC	1	0.7×10⁶	516	less than 50	3.5
Compa rativ e Examp le 3	499	76	F+P	NbC	1	0.6×10⁶	455	less than 50	3.5
Inven tive Examp le 4	512	216	F+P	NbC	1	1.9×10⁶	1326	168	2.6
Compa rativ e Examp le 4	507	79	F+P	NbC	1	0.7×10⁶	480	105	2.5
Compa rativ e Examp le 5	480	99	F+P	Unpro duced	0	0	0	less than 50	2.5
Compa rativ e	463	152	F+P	Fe₃C	6	6.0×10⁴	42	137	2.2
Examp le 6
Compa rativ e Examp le 7	442	302	F+P	Fe₃C	8	1.1×10⁵	80	117	2.4
Compa rativ e Examp le 8	487	287	F+P	Fe₃C	8	1.2×10⁵	84	151	2.5
Compa rativ e Examp le 9	433	217	F+P	Fe₃C	1	7.9×10⁴	56	129	2.3
Compa rativ e Examp le 10	445	222	F+P	MoC	1	8. 0×10⁴	59	103	2.2
Compa rativ e Examp le 11	573	187	F+P	MoC	3	9. 0×10⁴	64	19	2.1
Compa rativ e Examp le 12	544	166	F+P	Fe₃C, VC	5	8. 0×10⁴	54	55	1.8
Compa rativ e Examp le 13	535	201	F+P	Fe₃C, VC	5	9.0×10⁴	64	59	1.8
Compa rativ e Examp le 14	528	175	F+P	Cr Carbi de	4	8. 0×10⁴	55	89	4.1
Compa rativ e Examp le 15	518	178	F+P	Cr Carbi de	4	8. 0×10⁴	55	92	3.9

In Table 3, F refers to ferrite and P refers to pearlite.
On the other hand, FIG. 2 illustrates the size (nm) of the precipitate by observing the precipitate of Inventive example 1 by TEM. As illustrated in FIG. 2, it can be seen that in Inventive Example 1 of the present invention, NbC precipitates of 100 nm or less are evenly formed. On the other hand, FIG. 3 illustrates the size (nm) of the precipitate by observing the precipitates of Comparative Example 6 by the TEM, and it can be seen that in Comparative Example 6, a coarse FeC precipitate was formed.
From the results of Table 3 it can be seen that, in the case of the inventive example, even after the stress relief heat treatment after welding, the base material not only secures high strength and impact toughness but also the heat-affected zone (HAZ) may secure high impact toughness. For example, the steel according to the present invention ensures excellent toughness of HAZ even during large heat input welding, and may be produced as a steel material without defects such as surface cracks.
Comparative Examples 1 and 4 satisfy the alloy composition of the present invention, but the hot finish rolling temperature is too high. As a result sufficient toughness of the base material cannot be secured due to coarsening of the microstructure. FIGS. 1A and 1B are images of the base material microstructures of Inventive Example 1 and Comparative Example 1, respectively. Although all of the microstructures are formed of ferrite and pearlite in the same manner, in the case of Comparative Example 1 it is considered that the grain size is coarse and thus the impact toughness was lowered.
Comparative Examples 2 and 3 also satisfy the alloy composition of the present invention, but the slab heating temperature is outside the range of the present invention. Thus, an element that inhibits austenite grain growth at high temperature, such as Nb or the like, is not sufficiently solidified, or the austenite grain size is excessively coarse due to high temperature, resulting in a decrease in strength and impact toughness of the base material. In the case of Comparative Example 5, after hot rolling, the cooling rate of the steel was outside the range proposed in the present invention during cooling, and thus, the precipitation of the steel was not secured. Thus, the strength in the present invention was not obtained.
On the other hand, it can be seen that in Comparative Examples 6 to 10, the content of Nb in the steel was insufficient, and C was precipitated as coarse cementite particles, MoC, etc., and thus sufficient strength could not be secured and toughness of HAZ could not be secured. In Comparative Examples 11 to 13, it can be seen that the N content of the steel does not reach the scope of the present invention, Free-N defined by the relational expression 2 does not satisfy the conditions of the present invention and hence coarse precipitates such as MoC, Fe₃C, VC, etc are formed. It can be seen that the impact toughness of the HAZ is inferior because the distribution of precipitates is different from that in the present invention. In Comparative Examples 14 and 15, the Si content exceeded the range of the present invention, and the physical properties of the base material matched the range intended in the present invention, but the impact toughness of the HAZ was inferior due to excessive fraction of a martensite austenite constituent in the HAZ.

Claims

A high-strength steel having excellent toughness in a heat-affected zone, comprising:
in weight%, C: 0.16 to 0.20%, Mn: 1.0 to 1.5%, Si: 0.3% or less excluding 0, Al: 0.005 to 0.5%, P: 0.02% or less, S: 0.01% or less, Ti: 0.005 to 0.02%, Nb: 0.01 to 0.1%, and N: 0.006 to 0.01%; and

at least one selected from the group consisting of Ca: 0.006% or less, V: 0.03% or less, Ni: 2.0% or less, Cu: 1.0% or less, Cr: 1.0% or less, and Mo: 1.0% or less, and a balance of Fe and unavoidable impurities,

wherein a microstructure of the high-strength steel is composed of a ferrite-pearlite composite structure; and

in the microstructure, 1.27×10⁶ or more of Nb-based carbides of 100 nm or less are present per 1mm², and 900 or more of Nb-based carbides are distributed in a single crystal grain, measured by a transmission electron microscope.
The high-strength steel having excellent toughness in a heat-affected zone of claim 1, wherein the steel has a surface crack sensitivity index (Cs) of 0.3 or less, defined by the following relational expression 1, $Cs = (71.4 \times {[C]}^{2}) - (30.3 \times [C]) + 3.32,$
where [C] is a weight% content of a corresponding component.
The high-strength steel having excellent toughness in a heat-affected zone of claim 1, wherein the steel has free-N defined by the following relational expression 2, exceeding 0, Free-N = [N] - {([Ti]/47.887)×14.01} - {([B]/10.81)×14.01}, where [N], [Ti] and [B] are weight% contents of corresponding components.
The high-strength steel having excellent toughness in a heat-affected zone of claim 1, wherein the steel has a tensile strength of 500 MPa or higher and a Charpy impact absorption energy of 150 J or higher.
The high-strength steel having excellent toughness in a heat-affected zone of claim 1, wherein in the heat-affected zone of the steel, a martensite austenite constituent (MA) is 3% or less in area fraction.
A method of manufacturing a high-strength steel according to Claim 1 having excellent toughness in a heat-affected zone, comprising:
preparing a steel slab,

the steel slab including,

in weight%, C: 0.16 to 0.20%, Mn: 1.0 to 1.5%, Si: 0.3% or less excluding 0, Al: 0.005 to 0.5%, P: 0.02% or less, S: 0.01% or less, Ti: 0.005 to 0.02%, Nb: 0.01 to 0.1%, and N: 0.006 to 0.01%, and

at least one selected from the group consisting of Ca: 0.006% or less, V: 0.03% or less, Ni: 2.0% or less, Cu: 1.0% or less, Cr: 1.0% or less and Mo: 1.0% or less, and a balance of Fe and unavoidable impurities;

heating the steel slab to a temperature in a range of 1050 to 1250°C;

hot rolling the heated steel slab at a temperature of finish hot rolling of 910°C or lower to obtain hot-rolled steel sheet;

performing cooling of the hot-rolled steel sheet at a temperature in a range of 800°C to 435°C at a cooling rate of 20°C/Hr or less; and

welding and stress-relief heat treating the cooled hot-rolled steel sheet.
The method of manufacturing a high-strength steel having excellent toughness in a heat-affected zone of claim 6, wherein the hot finish rolling is performed at a temperature in a range of 850 to 910°C.