KR101819356B1 - Ultra thick steel having superior brittle crack arrestability and method for manufacturing the steel - Google Patents

Abstract

The present invention provides high-strength ultra-thick steel having superior brittle crack propagation resistance and a manufacturing method thereof. According to the present invention, the ultra-thick steel having superior brittle crack propagation resistance comprises 0.03-0.09 wt% of C, 1.4-2.2 wt% of Mn, 0.2-0.9 wt% of Ni, 0.005-0.05 wt% of Nb, 0.005-0.04 wt% of Ti, 0.1-0.5 wt% of Cu, 0.05-0.5 wt% of Si, 0.01-0.05 wt% of Al, 100 ppm or lower of P, 40 ppm or lower of S, and the remainder consisting of Fe and inevitable impurities. A surface portion has a mixed structure of polygonal ferrite and bainite. A 1/2t ~ 1/4t (t is the steel thickness) thickness portion consists of 50 volume% or higher of acicular ferrite and 50 volume% or lower of bainite. A fraction of an area having a fully bainite structure in the entire steel thickness is 20% or lower. According to the present invention, high-strength ultra-thick steel having superior brittle crack propagation resistance can be provided at high productivity.

Description

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a steel material having excellent brittle crack propagation resistance and a method of manufacturing the steel material.

TECHNICAL FIELD The present invention relates to a superfine steel material having excellent brittle crack propagation resistance and a method of manufacturing the same. More particularly, the present invention relates to a superfine steel material having excellent brittle crack propagation resistance and productivity, and a method of manufacturing the same.

Recently, it is required to develop ultra high strength steel and high strength steel for the design of structures such as domestic and overseas ships.

When high-strength steel is used in the design of the structure, it is economically advantageous due to the weight reduction of the structure, and the thickness of the plate can be made thin, thus facilitating the processing and welding work.

Generally, when a high-strength steel is manufactured from a superficial material, the structure is not sufficiently deformed due to a decrease in the total reduction ratio, so that the structure becomes coarse. In the rapid cooling for securing strength, There is a difference in cooling rate between the sub-center portions, which causes a coarse low-temperature transformation phase such as bainite on the surface portion, which makes it difficult to secure toughness.

Especially, the application of high strength steel reinforced materials to major structures such as ships is increasing the demand for guaranteeing the brittle crack propagation resistance which indicates the stability of the structure.

However, as described above, when a high-strength steel is produced from a post-mortar material, a brittle crack propagation resistance is significantly lowered when a coarse low-temperature transformation phase is generated. Therefore, brittle crack propagation of the high- It is a very difficult situation to do.

Further, in the production of high-strength steel after-feed materials, since the finish rolling is carried out at a very low temperature in order to improve toughness, it is required to wait for a long period of time in an air-cooled state at a high temperature before completion of rough rolling, .

BACKGROUND ART [0002] Techniques for finely reducing the particle size of the surface portion in order to improve brittle crack propagation resistance in the production of high strength and high strength materials having a yield strength of 500 MPa or more are known.

As described above, there is known a technique for finely reducing the particle size of the surface portion by applying surface cooling during finishing rolling or adjusting the grain size by applying bending stress during rolling.

However, although the above-mentioned conventional techniques are helpful in refining the surface texture, it is impossible to solve the brittle crack propagation resistance because it can not solve the degradation of the impact toughness due to coarsening of the remaining tissues. There is a problem that the productivity is expected to deteriorate considerably and the productivity deterioration due to the long period of air cooling standby between the rough rolling and the finishing rolling can not be prevented.

In order to improve brittle crack propagation resistance, a technique of adding a large amount of elements such as Ni to improve toughness is known.

However, when a large amount of elements such as Ni is added, resistance to brittle crack propagation can be improved, but it is difficult to commercialize it in view of manufacturing cost because it is an expensive element.

Japanese Patent Publication No. 06-004903 Japanese Patent Publication No. Hei 06-006741

An aspect of the present invention is to provide a high strength ultra high strength steel having excellent brittle crack propagation resistance.

Another aspect of the present invention is to provide a method for producing a high strength superalloy steel excellent in brittle crack propagation resistance with high productivity.

According to one aspect of the present invention, there is provided a steel sheet comprising, by weight%, 0.03 to 0.09% of C, 1.4 to 2.2% of Mn, 0.2 to 0.9% of Ni, 0.005 to 0.05% of Nb, 0.005 to 0.04% And the balance of Fe and other unavoidable impurities, and the surface portion is composed of a mixture of polygonal ferrite and bainite in a ratio of 0.5 to 0.5%, Si: 0.05 to 0.5%, Al: 0.01 to 0.05%, P: Wherein the portion of thickness 1 / 2t to 1 / 4t (where t is the steel thickness) is composed of 50% by volume or more of acicular ferrite and 50% by volume or less of bainite, and the bainite single phase And a fraction of the region having the structure of 20% or less is excellent in brittle crack propagation resistance.

The steel may preferably have an average grain size of 20 micrometers or less with a high-angle boundary of the central microstructure.

The steel may preferably have a yield strength of 500 MPa or more.

The steel may preferably have a center impact transient temperature of less than -40 占 폚.

The steel may preferably have a thickness of at least 50 mm.

According to another aspect of the present invention, there is provided a method of manufacturing a copper-clad laminate, comprising: 0.03 to 0.09% of C, 1.4 to 2.2% of Mn, 0.2 to 0.9% of Ni, 0.005 to 0.05% of Nb, 0.005 to 0.04% The steel slab containing 0.1 to 0.5% of Si, 0.05 to 0.5% of Si, 0.01 to 0.05% of Al, 100 ppm or less of P, 40 ppm or less of S and the balance of Fe and other unavoidable impurities is reheated to a temperature of 1150 to 1000 ° C step;

Subjecting the reheated slab to rough rolling at a temperature of 1150 to 900 캜;

Cooling the roughly rolled bar using a cooling means;

Recovering the cooled bar to an Ac3 temperature or higher on a surface basis;

Finishing the reheated bar at a temperature of Ar3 or higher on a 1/4 t basis; And cooling to a temperature of 600 DEG C or less at a cooling rate of 3 DEG C / s or more after finish rolling, wherein the cooling of the bar has a temperature of less than Ac3 at the surface portion of the bar, Bar thickness) region has a brittle crack propagation resistance resistance which is carried out so as to have a temperature higher by 50 DEG C or more than the finish rolling starting temperature.

According to the present invention, it is possible to provide a high strength superalloy having excellent brittle crack propagation resistance and high productivity.

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

The steel of the present invention having excellent brittle crack propagation resistance according to one aspect of the present invention comprises 0.03 to 0.09% of C, 1.4 to 2.2% of Mn, 0.2 to 0.9% of Ni, 0.005 to 0.05% of Nb, 0.005 to 0.05% of Nb, : 0.005 to 0.04%, Cu: 0.1 to 0.5%, Si: 0.05 to 0.5%, Al: 0.01 to 0.05%, P: 100 ppm or less, S: 40 ppm or less and the balance Fe and other unavoidable impurities, And a thickness of 1 / 2t to 1 / 4t (here, t: steel material thickness) is composed of 50% by volume or more of acicular ferrite and 50% by volume or less of bainite, The fraction of the region having a bainite single phase structure is 20% or less.

Hereinafter, the components of the steel and its content will be described.

C: 0.03 to 0.09% (hereinafter, the content of each component means weight%)

C is the most important element in securing the basic strength in the present invention, and therefore it is necessary to be contained in the steel in an appropriate range. However, when the content of C exceeds 0.09%, toughness is lowered due to generation of a large amount of ground martensite and promotion of generation of low-temperature transformation phase in the weld heat affected portion, and when the content of C is less than 0.03% The content is limited to 0.03 to 0.09%. The content of C is preferably limited to 0.04 to 0.09%, more preferably 0.05 to 0.08%.

Mn: 1.4 to 2.2%

Mn is a useful element for improving the hardenability so as to improve the strength by solid solution strengthening and to produce a low temperature transformation phase. In order to satisfy the strength of 500 MPa or more, Mn should be added in an amount of 1.4% or more. However, the addition of more than 2.2% accelerates the formation of upper bainite and martensite due to an increase in the excess hardening ability, which greatly reduces the impact toughness and brittle crack propagation resistance, so that the Mn content is 1.4 to 2.2% It limits. The Mn content is preferably limited to 1.5 to 2.1%, and more preferably to 1.6 to 2.0%.

Ni: 0.2 to 0.9%

Ni is an important element that improves the impact strength and improves the hardenability by facilitating cross slip at dislocations at low temperatures. It is an impact element in high-strength steels having a yield strength of 500 MPa or more and brittle crack propagation In order to improve the resistivity, it is preferable to add at least 0.2%, but when it is added in excess of 0.9%, the hardenability is excessively increased to cause a low temperature transformation phase to lower the toughness and increase the manufacturing cost. And is preferably limited to 0.9%. The Ni content is preferably limited to 0.3 to 0.9%, more preferably 0.4 to 0.8%.

Nb: 0.005 to 0.05%

Nb improves the hardenability and precipitates in the form of NbC or NbCN to improve the strength of the base material. In addition, the Nb solidified at the time of reheating at a high temperature is extremely finely precipitated in the form of NbC at the time of rolling, thereby suppressing the recrystallization of austenite, and thereby making the structure finer. Therefore, it is preferable that Nb is added in an amount of 0.005% or more in order to obtain such an additive effect. However, when excessive amount of Nb is added, there is a possibility of causing a brittle crack at the edge of the steel, so the upper limit of the Nb content is limited to 0.05%. The Nb content is preferably limited to 0.01 to 0.04%, more preferably 0.015 to 0.03%.

Ti: 0.005 to 0.04%

Ti is an element that precipitates as TiN during reheating and greatly improves the low-temperature toughness by suppressing grain growth of the base material and weld heat affected zone. Ti should be added in an amount of 0.005% or more for effective precipitation of TiN. However, excessive addition of more than 0.04% causes problems such as clogging of performance nozzles, precipitation of coarse TiN, coarse (TiNb), and (C, N) forms to reduce toughness. Therefore, the Ti content is 0.005 to 0.04 %. The Ti content is preferably limited to 0.01 to 0.03%, more preferably 0.012 to 0.025%.

Cu: 0.1 to 0.5%

Cu is a major element which improves the hardenability, improves the strength of the steel by strengthening the solid solution, and increases the yield strength through the formation of the epsilon Cu precipitate when the tempering is applied. In order to obtain such an additive effect, Is preferably added. However, the slab may crack due to hot shortness in the steelmaking process when added in a large amount, so the upper limit of the Cu content is preferably limited to 0.5%. The Cu content is preferably limited to 0.1 to 0.4%, more preferably 0.2 to 0.4%.

Si: 0.05 to 0.5% and Al: 0.01 to 0.05%

Si and Al are alloying elements which are necessary for deoxidizing work by precipitating dissolved oxygen in molten steel in steelmaking and casting process, and Si and Al are contained more than 0.05% and 0.01% It is essential. However, when added in a large amount, Si and Al complex oxides can be produced coarsely or a large amount of martensite can be produced in a large amount in the microstructure. Therefore, Si is preferably added in an amount of 0.5% or less and Al is added in an amount of 0.05% or less.

P: not more than 100 ppm and S: not more than 40 ppm

P and S are elements which induce brittleness in grain boundaries or cause coarse inclusions to induce brittleness. In order to improve brittle crack propagation resistance, it is preferable to limit P to not more than 100 ppm and S to not more than 40 ppm.

Hereinafter, the microstructure and physical properties of the steel will be described.

The microstructure of the surface portion of the steel material of the present invention is composed of a mixed phase of polygonal ferrite and bainite and is composed of fine grains of 1 / 4t portion [thickness 1 / 2t to 1 / 4t (here, t: The structure consists of 50 volume% or more acicular type ferrite and 50 volume% or less of bainite.

The surface portion of the steel material can be defined, for example, from the surface directly under the surface to the area of 10 mm from the surface.

For example, the microstructure of the steel surface portion includes 70 to 90% by volume of polygonal ferrite and 10 to 30% by volume of bainite in the case of 2 mm portion directly under the surface, and 20 to 30% by volume of polygonal And a mixed phase comprising ferrite and 70 to 80% by volume of bainite.

The fraction of the region having a bainite single phase structure in the entire steel sheet is 20% or less.

In the present invention, the bar immediately before the finish rolling has an austenite structure. The surface portion of the bar is subjected to a cooling and a reheating process under a suitable condition to form a fine structure, for example, a bainite phase, And a mixed phase thereof and the like have a reverse-transformed fine austenite.

The austenite refinement by the inverse transformation of the surface portion raises the air-cooled ferrite transformation temperature, so that at least a part of the fine austenite is transformed into ferrite before the cooling process after the finishing rolling, and the austenite not transformed into ferrite It is transformed into bainite by cooling.

Therefore, the microstructure of the steel surface portion has a mixed phase of ferrite and bainite.

In this way, by making the microstructure of the steel surface portion to be a mixed phase of ferrite and bainite, it is possible to achieve a fraction of the area having a bainite single-phase structure of 20% or less in the entire steel material thickness.

If the fraction of the region having a bainite single-phase structure in the total thickness of the steel exceeds 20%, the brittle crack propagation resistance is lowered.

As the content of C, Mn and Ni increases, the fraction of bainite as a whole increases, and accordingly the strength also increases.

 The steel may preferably have an average grain size of 20 micrometers or less with a high-angle boundary of the central microstructure.

When the grain size having a high-angle boundary exceeds 20 micrometers on average, the brittle crack propagation resistance can be lowered.

The steel may preferably have a yield strength of 500 MPa or more.

The steel may preferably have a center impact transient temperature of less than -40 占 폚.

The steel may preferably have a thickness of at least 50 mm.

Hereinafter, a method of manufacturing the steel material of the present invention will be described.

The method for producing a steel material of the present invention includes a process of slab reheating-roughing-bar cooling-double heat-finishing rolling-cooling.

Slab reheating temperature: 1150 ~ 1000 ℃

The slab is reheated to a temperature of 1150 to 1000 占 폚 before rough rolling the slab.

The heating temperature of the slab is preferably 1000 ° C or higher, in order to solidify the carbonitride of Ti and / or Nb formed during the casting.

In order to sufficiently solidify the carbonitride of Ti and / or Nb, it is more preferable to heat it to 1050 占 폚 or more. However, if the slab is reheated at an excessively high temperature, there is a possibility that the austenite is coarsened. Therefore, the slab reheating temperature is preferably 1150 DEG C or lower.

Rough rolling temperature: 1150 to 900 ° C

The reheated slab is subjected to rough rolling after heating to adjust its shape.

The rolling temperature should be at least the temperature (Tnr) at which the austenite recrystallization stops. It is possible to obtain the effect of reducing the grain size through recrystallization of coarse austenite along with the destruction of the cast structure such as dendrites formed during casting by rolling.

In order to obtain such an effect, the rough rolling temperature is preferably limited to 1150 to 900 ° C.

In order to make the structure fine by causing sufficient recrystallization, the total cumulative rolling reduction during rough rolling is preferably 40% or more

Bar Cooling:

After the rough rolling, the bar is rapidly cooled to the finish rolling temperature or more by using a cooling means. By cooling, a fine structure is formed on the bar surface portion. For example, by cooling, a bainite phase, an acicular ferrite phase, a mixed phase thereof, or the like may be generated on the bar surface portion.

The cooling end temperature is preferably 50 ° C or higher than the finish rolling starting temperature on a 1/4 t basis, and the cooling rate is preferably 0.5 ° C / s or more on a 1/4 t basis.

When the cooling end temperature is lower than 50 캜 from the finishing rolling start temperature, the double surface of the surface portion is not sufficiently generated, so that fine structures such as a bainite phase, an acicular ferrite phase, There is a possibility that the toughness is lowered because the steel is not reversely transformed into austenite again. Therefore, it is preferable that the cooling end temperature is limited to a temperature of 50 DEG C or higher than the finish rolling starting temperature.

On the other hand, when the cooling end temperature exceeds 100 deg. C higher than the finish rolling starting temperature, the austenite grows due to the high temperature after the double refining, and the grain size becomes large. May be deteriorated. Therefore, it is preferable that the cooling end temperature is limited to a temperature of 100 DEG C or lower than the finish rolling starting temperature.

When the cooling rate is less than 0.5 占 폚 / s (sec) on the basis of 1 / 4t, the recrystallized austenite structure in the center of the bar is coarsened, and the grain size having the high-angle boundary of the central microstructure after cooling of the finished rolled steel The cooling rate is preferably 0.5 deg. C / s (sec) or more, more preferably 1 to 10 deg. C / s (sec) on the basis of 1 / 4t, , Still more preferably 2 to 5 占 폚 / s (second).

As described above, it is possible to prevent the recrystallized austenite structure from coarsening during air cooling through the bar cooling, thereby obtaining an effect of finely finishing the final microstructure.

In addition, it is possible to prevent an air-cooling atmosphere from being generated for a long period of time before finishing rolling, thereby improving the productivity.

Double heat: Surface part Ac3 temperature or higher

After the rough rolling, the bar cooled by the cooling means is air-cooled for a certain period of time to recover the temperature of the excessively cooled surface portion. In order to transform the fine structure formed on the surface portion during cooling of the bar into austenite again such as a bainite phase, an acicular ferrite phase, or a mixed phase thereof, that is, in order to reverse- It is preferable to heat it until it is more than the above. More preferably, the surface portion temperature is Ac 3 ° C to Ac 3 + 100 ° C, and more preferably, the surface portion temperature is Ac 3 + 20 ° C to Ac 3 + 70 ° C.

The surface portion of the bar can be defined, for example, from the surface directly underneath to the area of 10 mm from the surface.

As described above, the fine structure of the surface portion of the bar produced during the cooling, for example, a bainite phase, an acicular ferrite phase, a mixed phase thereof, or the like is reversely transformed into austenite, Whereby the air-cooled ferrite transformation temperature rises and an effect of lowering the bainite single phase texture formation in the steel can be obtained.

The particle size of the austenite that has been inversely transformed in the fine structure, for example, a bainite phase, an acicular ferrite phase, a mixed phase thereof, or the like may be, for example, 50 micrometers (占 퐉) or less.

Finishing rolling temperature: Ar3 or more based on 1 / 4t

The roughly rolled bars are subjected to finish rolling in the non-recrystallized region. The finishing rolling finishing temperature is set to be equal to or higher than the ferrite forming temperature (Ar3). When rolling is performed at a temperature less than Ar3, a large amount of air-cooled ferrite is generated in the entire microstructure in the thickness direction, which may make it difficult to secure a yield strength of 500 MPa or more.

Cooling condition after finishing rolling: cooling to 600 ° C or less at a cooling rate of 3 ° C / s or more

The austenite refinement by the inverse transformation of the surface portion raises the air-cooled ferrite transformation temperature, so that at least a part of the fine austenite is transformed into ferrite before the cooling process after the finishing rolling, and the austenite not transformed into ferrite It is transformed into bainite by cooling.

Therefore, the microstructure of the steel surface portion has a mixed phase of ferrite and bainite.

The finish-rolled steel is cooled to 600 캜 or lower at a cooling rate of 3 캜 / s or higher.

If the cooling rate after finishing rolling is lower than 3 ° C / s or when the cooling is terminated at a temperature higher than 600 ° C, the microstructure may not be properly formed, and the yield strength may become less than 500 MPa.

The steel may preferably have a thickness of at least 50 mm.

The surface microstructure of the steel is composed of a mixed phase of polygonal ferrite and bainite. In the center portion of the steel, 1 / 4t is composed of 50% or more acicular ferrite and 50% or less of bainite , A steel material having a fraction of a region having a bainite single-phase structure of not more than 20% at a total thickness of steel can be produced.

The steel may preferably have an average grain size of 20 micrometers or less with a high-angle boundary of the central microstructure.

The steel may preferably have a yield strength of 500 MPa or more.

The steel may preferably have a center impact transient temperature of less than -40 占 폚.

As described above, by controlling the steel composition and the manufacturing conditions, it is possible to secure the microstructure of the product having excellent brittle crack propagation resistance, to shorten the air-cooling standby time occurring until rough rolling after rough rolling, through bar cooling and double heat, Strength steel having a yield strength of 500 MPa or more and a core impact transition temperature of -40 占 폚 or less.

In particular, in the present invention, since the roughly rolled bar is cooled using the cooling means, the productivity can be improved by reducing the air-cooling waiting time and preventing the growth of austenite, The grain size with boundaries can be maintained below 20 micrometers on average.

As described above, by controlling the steel composition and performing the cooling and the double refining process under the proper conditions before the manufacturing process, particularly the rough rolling process after the rough rolling, it is possible to provide the high strength superalloy steel excellent in brittle crack propagation resistance with high productivity .

Hereinafter, the present invention will be described more specifically by way of examples. It should be noted, however, that the following examples are intended to illustrate the present invention by way of illustration and not to limit the scope of the present invention. The scope of the present invention is determined by the matters set forth in the claims and the matters reasonably inferred therefrom.

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

(Example)

A steel slab having a thickness of 400 mm having the composition shown in the following Table 1 was reheated to a temperature of 1070 캜, followed by rough rolling at a temperature of 1025 캜 to prepare bars. The cumulative rolling reduction rate during rough rolling was 50%.

The thickness of the rough-rolled bar was 200 mm. After the rough rolling, the bar was cooled and then subjected to a double heat. The surface recuperation temperature in Table 2 was set to 1 / 4t and 1 / 2t Means a surface temperature measurement value at a time point when the temperature difference becomes less than 20 占 폚. The cooling of the bars was carried out so that the surface portion of the bars had a temperature of less than Ac3 and the area of 1/4 t (where t is the bar thickness) had a temperature of 50 DEG C or more higher than the finish rolling start temperature. At this time, the cooling rate at the time of cooling the bar was 1 to 5 ° C / sec.

The steel sheet having the thickness shown in the following Table 2 was obtained by performing the finish rolling at the time of completing the heat treatment, and then cooled to a temperature in the range of 500 to 300 占 폚 at a cooling rate of 3.5 to 5 占 폚 / sec.

Steel grade
Steel composition (% by weight) C Mn Ni Cu Ti Nb Si Al P (ppm) S (ppm) Inventive Steel 1 0.061 1.53 0.63 0.21 0.023 0.018 0.30 0.031 55 17 Invention river 2 0.071 1.65 0.52 0.3 0.012 0.012 0.32 0.021 65 11 Invention steel 3 0.039 2.11 0.45 0.26 0.017 0.025 0.23 0.040 79 23 Inventive Steel 4 0.077 1.78 0.62 0.29 0.022 0.023 0.35 0.023 81 22 Invention steel 5 0.066 1.82 0.27 0.15 0.018 0.028 0.31 0.037 46 24 Comparative River 1 0.12 2.01 0.52 0.21 0.021 0.019 0.20 0.041 49 9 Comparative River 2 0.071 2.35 0.71 0.29 0.013 0.021 0.45 0.020 78 28 Comparative Steel 3 0.026 1.32 0.39 0.18 0.019 0.018 0.21 0.045 59 12 Comparative Steel 4 0.075 1.93 1.24 0.41 0.021 0.015 0.31 0.033 65 16 Comparative Steel 5 0.062 1.69 0.44 0.24 0.042 0.051 0.33 0.025 57 12

Example No. 2. Steel grade Steel plate thickness (mm) Whether bar cooling is applied Surface cooling temperature after bar cooling (℃) Particle size (탆) of reversely transformed austenite after double heat treatment Inventory 1 Inventive Steel 1 80 ○ Ac3 + 45 Less than 50 Inventory 2 Invention river 2 90 ○ Ac3 + 22 Less than 50 Inventory 3 Invention steel 3 95 ○ Ac3 + 38 Less than 50 Honorable 4 Inventive Steel 4 100 ○ Ac3 + 52 Less than 50 Inventory 5 Invention steel 5 80 ○ Ac3 + 39 Less than 50 Comparative Example 1 Invention river 2 90 ○ Ac3-52 Can not be measured with partial transformation Comparative Example 2 Invention steel 3 95 ○ Ac3 - 76 Can not be measured with partial transformation Comparative Example 3 Inventive Steel 1 80 X - - Comparative Example 4 Inventive Steel 4 100 X - - Comparative Example 5 Comparative River 1 80 ○ Ac3 + 33 Less than 50 Comparative Example 6 Comparative River 2 85 ○ Ac3 + 59 Less than 50 Comparative Example 7 Comparative Steel 3 90 ○ Ac3 + 65 Less than 50 Comparative Example 8 Comparative Steel 4 90 ○ Ac3 + 29 Less than 50 Comparative Example 9 Comparative Steel 5 95 ○ Ac3 + 45 Less than 50

Table 3 shows the microstructure analysis results and the yield strength / core impact transition temperature results for the steels produced according to Tables 1 and 2 above.

In addition, the steel material is subjected to a CAT (Crack Arrest Test) evaluation at -10 ° C using ESSO equipment, and whether or not crack propagation / arresting is performed is shown in Table 3 below.

The center particle size in Table 3 below is measured by EBSD method and means the value obtained by calculating the grain boundaries having a high boundary angle of 15 degrees or more using the measurement results.

Example
No. Steel grade Surface portion
Microstructure Bainite single phase tissue area (%) Center ~ 1 / 4t Microstructure fraction (%) surrender
burglar
(Mpa) center
Average particle size (탆) Central impact transition temperature (캜) CAT
(@ -10 ° C)
Inventory 1 Inventive Steel 1 PF + B 15 AF (83) B (17) 512 17.3 -53 Arrest Inventory 2 Invention river 2 PF + B 17 AF (74) B (26) 532 18.5 -49 Arrest Inventory 3 Invention steel 3 PF + B 8 AF (61) B (39) 575 16.1 -65 Arrest Honorable 4 Inventive Steel 4 PF + B 13 AF (89) B (11) 506 17.7 -78 Arrest Inventory 5 Invention steel 5 PF + B 12 AF (59) B (41) 568 13.1 -71 Arrest Comparative Example 1 Invention river 2 B 28 AF (68) B (32) 544 14.6 -59 Propagate Comparative Example 2 Invention steel 3 B 26 AF (66) B (34) 559 19.5 -53 Propagate Comparative Example 3 Inventive Steel 1 PF + B 16 AF (72) B (28) 529 25.3 -36 Propagate Comparative Example 4 Inventive Steel 4 PF + B 14 AF (86) B (14) 516 23.6 -31 Propagate Comparative Example 5 Comparative River 1 PF + B 49 AF (18) B (82) 678 28.7 -32 Propagate Comparative Example 6 Comparative River 2 PF + B 38 AF (22) B (78) 659 24.8 -28 Propagate Comparative Example 7 Comparative Steel 3 PF + B 6 PF (45) AF (49) P (6) 387 15.4 -64 Arrest Comparative Example 8 Comparative Steel 4 PF + B 37 AF (26) B (74) 595 26.8 -37 Propagate Comparative Example 9 Comparative Steel 5 PF + B 26 AF (51) B (49) 577 17.6 -45 Propagate

[PF: Polygonal Ferrite, P: Pearlite, AF: Acicular Ferrite, B: Bainite]

As shown in Table 3, in the case of Comparative Examples 1 and 2, since the double heat temperature after the Bar cooling was below the Ac3 temperature, the surface sub-bainite structure formed during the Bar cooling did not transform into the austenite again and remained coarse , The bainite single phase structure area exceeded 20%. Therefore, it can be seen that the crack was not stopped and propagated in the -10 ° C CAT test showing brittle crack propagation resistance.

In the case of Comparative Examples 1 and 2, since it was reheated to a temperature below the Ac3 temperature, only a part of the austenite did not occur in the bainite produced during the cooling and partly occurred, so that it was impossible to measure the grain size of the austenite after the recrystallization.

In the case of the comparative steels 3 and 4, since the bar cooling described in the present invention is not applied, the microstructure fraction is included in the range suggested in the present invention, Therefore, it can be seen that the center average particle size is more than 20 microns and the center impact transitional temperature is more than -40 ° C. In the -10 ° C CAT test showing the brittle crack propagation resistance, the crack is not stopped and propagated.

 In the case of Comparative Example 5, by having a value higher than the upper limit of the C specified in the present invention, a large amount of bainite structure was produced due to excessive curing ability, thereby causing a central impact transit temperature of -40 ° C or higher and brittle crack propagation resistance , It can be seen that the crack did not stop and propagated in the -10 ° C CAT test.

In the case of Comparative Example 6, by having a value higher than the upper limit of Mn shown in the present invention, a large amount of bainite structure was generated due to excessive curing ability, and thus, the central impact transition temperature was -40 ° C or more, In the -10 ° C CAT test showing resistance, it can be seen that the crack was not stopped and propagated.

In the case of Comparative Example 7, since the lower limit of the C and Mn lower limit values suggested in the present invention was used, a large amount of polygonal ferrite and pearlite structure were produced due to insufficient curing ability, and thus the yield strength was 500 MPa or less.

In the case of Comparative Example 8, a higher value than the upper limit of Ni suggested in the present invention resulted in the formation of a large amount of bainite structure due to excessive curing ability, thereby causing a central impact transition temperature of -40 ° C or higher, In the -10 ° C CAT test showing resistance, it can be seen that the crack was not stopped and propagated.

In the case of Comparative Example 9, a large amount of bainite structure was produced due to excessive curing ability, and coarse TiN or (TiNb) and (C, N) precipitated , It can be seen that the crack was not stopped and propagated in the -10 ° C CAT test showing brittle crack propagation resistance.

On the other hand, in the case of Inventive Examples 1 to 5 in which the surface range of the surface portion is recovered at a temperature higher than the Ac3 temperature after satisfying the composition range proposed by the present invention and performing bar cooling after rough rolling, the center particle size is fine to be 20 micrometers or less, It can be understood that the bainite single phase region is 20% or less with respect to the entire thickness region, the acute portion of the surface portion to 1/4 t portion is composed of the acicular ferrite having 50% or more of the microstructure and 50% or less of the bainite. As a result, the inventive examples 1 to 5 show excellent brittle crack propagation resistance with a yield strength of 500 MPa or more, a core impact transition temperature of -40 ° C or less, and a crack stopping at -10 ° C CAT test.

Claims (13)

The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.03 to 0.09% of C, 1.4 to 2.2% of Mn, 0.2 to 0.9% of Ni, 0.005 to 0.05% of Nb, 0.005 to 0.04% of Ti, 0.5 to 0.5%, Al: 0.01 to 0.05%, P: not more than 100 ppm, S: not more than 40 ppm, and the balance of Fe and other unavoidable impurities. The surface portion is a mixed phase of polygonal ferrite and bainite, 1 / 4t (where, t: steel material thickness) portion is composed of 50% by volume or more acicular type ferrite and 50% by volume or less of bainite, and the fraction of the region having a bainite single- Or less and an average grain size of 20 micrometers or less with a high-hardness boundary of the center microstructure, is excellent in brittle crack propagation resistance.
delete The superfine steel material according to claim 1, wherein the yield strength of the steel is 500 MPa or more and is excellent in brittle crack propagation resistance.
The superfine steel material according to claim 1, wherein the steel material has a brittle crack propagation resistance of not more than -40 캜 at a central portion shock transition temperature.
The steel material according to claim 1, wherein the steel has a thickness of 50 mm or more and is excellent in brittle crack propagation resistance.
The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.03 to 0.09% of C, 1.4 to 2.2% of Mn, 0.2 to 0.9% of Ni, 0.005 to 0.05% of Nb, 0.005 to 0.04% of Ti, Reheating the steel slab containing Fe, Fe, and other unavoidable impurities to a temperature of 1150 to 1000 캜; 0.5 to 0.5%; Al: 0.01 to 0.05%; P: not more than 100 ppm;
Subjecting the reheated slab to rough rolling at a temperature of 1150 to 900 캜;
Cooling the roughly rolled bar using a cooling means;
Recovering the cooled bar to an Ac3 temperature or higher on a surface basis;
Finishing the reheated bar at a temperature of Ar3 or higher on a 1/4 t basis; And cooling the rolled bar to a temperature of not more than 600 DEG C at a cooling rate of not less than 3 DEG C / s after the finish rolling, wherein the surface portion of the bar has a temperature of less than Ac3, , and t is the bar thickness) region is 50 ° C or more higher than the finish rolling starting temperature.
The method of claim 6, wherein the steel has a thickness of 50 mm or more and is excellent in brittle crack propagation resistance.
[7] The method of claim 6, wherein the surface temperature of the rebar in the reheating step is Ac3 + 20 deg. C to Ac3 + 70 deg. C, wherein the brittle crack propagation resistance is excellent.
The method of claim 6, wherein in the step of cooling the bar, a bainite phase, an acicular ferrite phase, or a mixed phase thereof is formed on the surface portion of the bar. Way.
The method according to claim 9, wherein the bainite phase of the surface portion, the acicular ferrite phase, or the mixed phase thereof is reversely transformed into austenite in the step of reheating the cooled bar, .
The method of claim 10, wherein the grain size of the austenite is less than 50 micrometers (탆).
The method of claim 6, wherein the cooling rate of the roughly rolled bars is 1 to 10 占 폚 / s (sec) on the basis of 1 / 4t.
The method of claim 6, wherein the cooling rate of the roughly rolled bar is 2 to 5 占 폚 / s (sec) on a 1/4 t basis.