WO2018030737A1 - Ultra-thick steel material having excellent resistance to brittle crack propagation and preparing method therefor - Google Patents

Abstract

The present invention provides a high-strength ultra-thick steel material having excellent resistance to brittle crack propagation and a preparing method therefor. The present invention provides an ultra-thick steel material having excellent resistance to brittle crack propagation and a preparing method therefor, wherein the steel material comprises, in terms of weight%, 0.03-0.09% of C, 1.4-2.2% of Mn, 0.2-0.9% of Ni, 0.005-0.05% of Nb, 0.005-0.04% of Ti, 0.1-0.5 % of Cu, 0.05-0.5% of Si, 0.01-0.05% of Al, 100 ppm or less of P, 40 ppm or less of S, and the balance Fe and unavoidable impurities, and wherein a surface portion is composed of a mixed phase of polygonal ferrite and bainite; a portion corresponding to a thickness of 1/2t-1/4t (here, t is the thickness of steel material) is composed of 50 vol% or more of acicular ferrite and 50 vol% or less of bainite; and a fraction of a region having a bainite single-phase structure in the overall thickness of the steel material is 20% or less. According to the present invention, a high-strength ultra-thick steel material having excellent resistance to brittle crack propagation can be produced at high productivity.

Description

Extremely thick steel with excellent brittle crack propagation resistance and manufacturing method

The present invention relates to an extremely thick steel material and a method for manufacturing the same having excellent brittle crack propagation resistance, and more particularly, to a structural ultra thick steel having excellent brittle crack propagation resistance and productivity.

Recently, in the design of structures such as domestic and overseas ships, development of ultra thick and high strength steels is required.

The use of high-strength steel in the design of the structure makes it possible to reduce the thickness of the structure and to reduce the thickness of the plate, thereby facilitating the ease of machining and welding.

In general, when the high strength steel is manufactured from the ultra-thick material, the structure becomes coarse because sufficient deformation is not made throughout the tissue due to the decrease in the total reduction ratio. There is a difference in the cooling rate between the sub-center portion, resulting in a coarse low-temperature transformation phase such as bainite in the surface portion is difficult to secure the toughness.

In particular, there is an increasing number of cases requiring the guarantee of brittle crack propagation resistance that indicates the stability of the structure when the high strength steel material is applied to the main structures such as ships.

However, as described above, when coarse low temperature transformation phase is produced when the high strength steel is made of the ultra thick material, the brittle crack propagation resistance is greatly lowered, thereby improving the brittle crack propagating resistance phase of the ultra thick high strength steel. It is a very difficult situation.

In addition, in the manufacture of high-strength steel ultra-thick material, because the filament rolling is carried out at a very low temperature to improve the toughness, after the rough rolling is completed, the air must be kept in an air-cooled state at a high temperature for a long time before finishing the rough rolling. Problem occurs.

In order to improve the brittle crack propagation resistance when manufacturing high strength steel material having a yield strength of 500 MPa or more, a technique for miniaturizing the surface portion is known.

As a conventional technique for miniaturizing the surface particle size as described above, a technique for controlling the particle size by applying surface cooling during finishing rolling or by applying a bending stress during rolling is known.

However, the above-described prior arts are helpful for miniaturizing the surface tissue, but the impact toughness due to coarsening of the remaining tissues cannot be solved, and thus, they cannot be a fundamental countermeasure against brittle crack propagation resistance. In order to apply, a great fall in productivity is expected, and there is a problem in that the productivity decrease due to the long time air-cooling atmosphere between rough rolling and finishing rolling cannot be prevented.

In addition, in order to improve the brittle crack propagation resistance, a technique of adding a large amount of elements such as Ni, which is helpful for improving toughness, is known.

However, in the case of adding a large amount of elements such as Ni, the brittle crack propagation resistance can be improved, but since it is an expensive element, commercial application is difficult in terms of manufacturing cost.

One aspect of the present invention is to provide a high strength ultra-thick steel material excellent in brittle crack propagation resistance.

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

According to an aspect of the present invention, in weight%, C: 0.03-0.09%, Mn: 1.4-2.2%, Ni: 0.2-0.9%, Nb: 0.005-0.05%, Ti: 0.005-0.04%, Cu: 0.1 ~ 0.5%, Si: 0.05 ~ 0.5%, Al: 0.01 ~ 0.05%, P: 100ppm or less, S: 40ppm or less, including the remaining Fe and other unavoidable impurities, and the surface portion is a mixed phase of polygonal ferrite and bainite. And a thickness of 1 / 2t to 1 / 4t (where t: steel thickness) is composed of acicular ferrite of 50 vol% or more and bainite of 50 vol% or less, and bainite single phase in the entire steel thickness. An extremely thick steel material having excellent brittle crack propagation resistance having a fraction of 20% or less of a structured region is provided.

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

The steel may preferably have a yield strength of at least 500 MPa.

The steel material may preferably have a central impact transition temperature of −40 ° C. or less.

The steel material may preferably have a thickness of 50 mm or more.

According to another aspect of the present invention, in weight%, C: 0.03-0.09%, Mn: 1.4-2.2%, Ni: 0.2-0.9%, Nb: 0.005-0.05%, Ti: 0.005-0.04%, Cu: 0.1 ~ 0.5%, Si: 0.05 ~ 0.5%, Al: 0.01 ~ 0.05%, P: 100ppm or less, S: 40ppm or less, and reheat the steel slab containing the remaining Fe and other unavoidable impurities to a temperature of 1150 ~ 1000 ℃. step;

Rough rolling the reheated slab at a temperature of 1150 to 900 ° C;

Cooling the crude rolled bar by using cooling means;

Reheating the cooled bar above an Ac3 temperature on a surface basis;

Finishing rolling the regenerated bar at a temperature of Ar3 or higher on a 1 / 4t basis; And after finishing rolling, cooling to a temperature of 600 ° C. or less at a cooling rate of 3 ° C./s or more, wherein the cooling of the bar has a temperature of less than Ac 3 at a surface portion of the bar, where t is Bar thickness) region is provided a method for producing a very thick steel material excellent in brittle crack propagation resistance is carried out to have a temperature higher than 50 ℃ higher than the finish rolling start temperature.

According to the present invention, it is possible to provide a high-strength ultra-thick steel having excellent brittle crack propagation resistance with high productivity.

Hereinafter, the preferable example of this invention is demonstrated in detail.

According to an aspect of the present invention, the ultra-thick steel having excellent brittle crack propagation resistance is% by weight, C: 0.03 to 0.09%, Mn: 1.4 to 2.2%, Ni: 0.2 to 0.9%, Nb: 0.005 to 0.05%, and Ti. : 0.005 ~ 0.04%, Cu: 0.1 ~ 0.5%, Si: 0.05 ~ 0.5%, Al: 0.01 ~ 0.05%, P: 100ppm or less, S: 40ppm or less, remaining Fe and other unavoidable impurities It consists of a mixed phase of gonal ferrite and bainite, and the thickness of 1 / 2t to 1 / 4t (where t: steel thickness) is composed of acicular ferrite of 50 vol% or more and bainite of 50 vol% or less. The fraction of areas having bainite single phase tissue in the total thickness is less than 20%.

Hereinafter, the component of steel and its content are demonstrated.

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

C is the most important element for securing the basic strength in the present invention, it needs to be contained in steel within an appropriate range. However, when the content of C exceeds 0.09%, the toughness decreases due to the formation of a large amount of phase martensite and low temperature transformation phase in the weld heat affected zone, and when the content of C is less than 0.03%, the strength decreases. The content is limited to 0.03 to 0.09%. The content of C is preferably limited to 0.04% to 0.09%, and more preferably 0.05% to 0.08%.

Mn: 1.4-2.2%

Mn is a useful element that improves the strength by solid solution strengthening and improves the hardenability so that low-temperature transformation phase is generated. It is required to add 1.4% or more to satisfy the strength of 500 MPa or more. However, addition of more than 2.2% promotes the formation of upper bainite and martensite due to excessive increase in hardenability, which greatly reduces the impact toughness and brittle crack propagation resistance, so the Mn content is 1.4-2.2%. It is limited. The Mn content is preferably limited to 1.5 to 2.1%, more preferably 1.6 to 2.0%.

Ni: 0.2 ~ 0.9%

Ni is an important element that facilitates cross slip of dislocation at low temperature, improves impact toughness, and improves hardening ability to improve strength. Impact toughness and brittle crack propagation in high strength steel with yield strength of 500 MPa or more In order to improve the resistance, it is preferable to add 0.2% or more, but when it is added in excess of 0.9%, there is a problem of excessively increasing the hardenability to form low-temperature transformation phase, lowering toughness and increasing manufacturing cost, so the upper limit of Ni content is It is preferable to limit to 0.9%. The Ni content is preferably limited to 0.3 to 0.9%, more preferably 0.4 to 0.8%.

Nb: 0.005-0.05%

Nb improves hardenability and precipitates in the form of NbC or NbCN to improve the base metal strength. In addition, Nb dissolved in reheating at a high temperature precipitates very finely in the form of NbC during rolling, thereby suppressing recrystallization of austenite, thereby miniaturizing the structure. Therefore, in order to obtain such an addition effect, Nb is preferably added at 0.005% or more. However, when excessively added, there is a possibility of causing brittle cracks at the corners 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: 0.005-0.04%

Ti is reconstituted with TiN during reheating to suppress grain growth of the base metal and the weld heat affected zone, thereby greatly improving low temperature toughness. For effective TiN precipitation, Ti must be added at least 0.005%. However, excessive addition of more than 0.04% may cause clogging of the playing nozzle, precipitation of coarse TiN or precipitation of coarse (TiNb), (C, N), resulting in a decrease in toughness. Therefore, the Ti content is 0.005 to 0.04. It is limited to%. The Ti content is preferably limited to 0.01 to 0.03%, more preferably 0.012 to 0.025%.

Cu: 0.1 ~ 0.5%

Cu is the major element that improves the hardenability, strengthens the solid solution, improves the strength of the steel, and increases the yield strength through the generation of epsilon Cu precipitates when tempering is applied. It is preferred to be added. However, when a large amount is added, the slab may be cracked due to hot shortness in the steelmaking process, 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: 0.05-0.5% and Al: 0.01-0.05%

Si and Al are essential alloy elements for deoxidation by precipitating dissolved oxygen in molten steel in the form of slag during steelmaking and casting process.Since Si and Al are used, Si is more than 0.05% and Al is 0.01% or more. It is essential. However, when a large amount is added, Si and Al composite oxides may be coarsely generated or coarse amounts of island martensite may be coarsely generated in the microstructure. Therefore, Si is preferably 0.5% or less and Al is preferably added at 0.05% or less.

P: 100 ppm or less and S: 40 ppm or less

P, S is an element that causes brittleness or forms coarse inclusions at grain boundaries, and is preferably limited to P: 100 ppm or less and S: 40 ppm or less in order to improve brittle crack propagation resistance.

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

The microstructure of the steel surface portion of the present invention is composed of a mixed phase of polygonal ferrite and bainite, and is fine in a 1 / 4t portion (thickness 1 / 2t to 1 / 4t (where t: steel thickness) portion) at the center of the steel. The tissue consists of at least 50% by volume of acicular ferrite and up to 50% by volume of bainite.

The surface portion of the steel material may be defined, for example, from directly below the surface to an area of the surface ˜10 mm.

For example, the microstructure of the steel surface part contains 70 to 90% by volume of polygonal ferrite and 10 to 30% by volume of bainite at 2mm below the surface, and 20 to 30% by volume of polygonal at 10mm below the surface. It is preferred to consist of a mixed phase comprising ferrite and 70 to 80% by volume of bainite.

The fraction of the region having bainite single phase structure in the total thickness of the steel is 20% or less.

In the present invention, the bar just before finishing rolling has an austenite structure, and the surface portion of the bar has a fine structure such as bainite phase, acicular ferrite phase, or the like by cooling and recuperating the roughly rolled bar under appropriate conditions. These mixed phases have fine austenite reversely transformed.

Due to the austenite miniaturization due to reverse transformation of the surface portion, the air-cooled ferrite transformation temperature is increased, and thus, at least a part of the fine austenite is transformed into ferrite before the cooling process after finishing rolling, and the austenite not transformed into ferrite is By cooling, it transforms into bainite.

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 consist of a mixed phase of ferrite and bainite, it is possible to achieve a fraction of the region having bainite single phase structure in the overall thickness of the steel to 20% or less.

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

As the C, Mn and Ni content increases, the fraction of bainite generally increases, and thus the strength also increases.

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

Brittle crack propagation resistance may be degraded when the particle size having a high angle boundary exceeds 20 micrometers on average.

The steel may preferably have a yield strength of at least 500 MPa.

The steel material may preferably have a central impact transition temperature of −40 ° C. or less.

The steel material may preferably have a thickness of 50 mm or more.

Hereinafter, the method of manufacturing the steel material of this invention is demonstrated.

The steel manufacturing method of the present invention includes the process of slab reheating-rough rolling-bar cooling-reheating-finish rolling-cooling.

Slab reheating temperature: 1150 ~ 1000 ℃

Before rough-rolling the slab, reheat the slab to a temperature between 1150 and 1000 ° C.

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

In order to sufficiently solidify the carbonitride of Ti and / or Nb, heating to 1050 ° C or more is more preferable. However, when the slab is reheated at an excessively high temperature, austenite may be coarsened, so the slab reheating temperature is preferably 1150 ° C or lower.

Rough rolling temperature: 1150 ~ 900 ℃

The reheated slabs are subjected to rough rolling after heating to adjust their shape.

The rolling temperature is at least the temperature Tnr at which recrystallization of austenite stops. The effect of reducing the particle size can also be obtained through the recrystallization of coarse austenite with the destruction of the casting structure such as the dendrite formed during casting by rolling.

In order to acquire such an effect, it is preferable to limit rough rolling temperature to 1150-900 degreeC.

In order to cause sufficient recrystallization and refine the tissue, the total cumulative reduction rate during rough rolling is preferably 40% or more.

Bar Cooling:

After the rough rolling bar is quickly cooled to the finish rolling temperature or more using a cooling means. By cooling, fine structure is produced in the bar surface part. For example, by cooling, the bar surface portion may generate a bainite phase, acicular ferrite phase, a mixed phase thereof, or the like.

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

When the cooling end temperature is less than 50 ° C. below the finish rolling start temperature, recuperation of the surface portion does not occur sufficiently, such that a fine structure generated on the surface portion during cooling, for example, a bainite phase, a needle-like ferrite phase, or a mixed phase thereof There is a possibility that the toughness is lowered because it is not reversely transformed into austenite. Therefore, the cooling end temperature is preferably limited to a temperature of 50 ° C. or more than the finish rolling start temperature.

On the other hand, when the cooling end temperature exceeds 100 ℃ than the finish rolling start temperature, because the temperature after the recuperation is high because the austenite grows to increase the grain size, or after the recuperation after a long temperature due to the large temperature difference until finish rolling productivity This may fall. Therefore, the cooling end temperature is preferably limited to a temperature of 100 ° C or less than the finish rolling start temperature.

When the cooling rate is less than 0.5 ° C / s (seconds) on the basis of 1 / 4t, coarsening of the recrystallized austenite structure in the center of the bar occurs, and the particle size having the high angle boundary of the central microstructure after cooling of the finished rolled steel is averaged. Since it may exceed 20 micrometers, the cooling rate is preferably 0.5 ° C / s (sec) or more on a 1 / 4t basis, and more preferably 1 to 10 ° C / s (sec) on a 1 / 4t basis. More preferably, it is 2-5 degree-C / s (second).

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

In addition, it is possible to prevent the generation of a long air-cooled atmosphere until the finish rolling to obtain the effect of improving productivity.

Reheat: Above Ac3 Temperature at Surface

After the rough rolling, the bar cooled by the cooling means is cooled by air for a predetermined time to restore the excessively cooled surface temperature. When the bar is cooled, the surface temperature at the time of recuperation is changed to Ac3 temperature in order to transform the microstructures formed on the surface portion, for example, bainite phase, needle-like ferrite phase, or mixed phase thereof, into austenite, that is, reverse transformation. It is preferable to recuperate until it becomes abnormal. The more preferable surface part temperature at reheating is Ac3 degreeC-Ac3 + 100 degreeC, and even more preferable surface part temperature is Ac3 + 20 degreeC-Ac3 + 70 degreeC.

The surface portion of the bar may be defined, for example, from directly below the surface to an area of the surface ˜10 mm.

As the recuperation as described above, the microstructure of the bar surface portion generated during the cooling, for example, the bainite phase, the needle-like ferrite phase, or a mixed phase thereof is inversely transformed into austenite, so that the surface austenite is minutely generated. As a result, the air-cooled ferrite transformation temperature is increased to reduce the generation of bainite single phase structure in the steel.

The particle size of the austenite reversely transformed in the microstructure, for example, bainite phase, acicular ferrite phase, or a mixed phase thereof may be, for example, 50 micrometers (μm) or less.

Finish rolling temperature: Ar3 or higher at 1 / 4t

The rough rolled bar is subjected to finish rolling in the unrecrystallized region. The finish rolling finish temperature is at least the ferrite generation temperature (Ar3). When rolling is performed at a temperature below Ar3, a large amount of air-cooled ferrite is generated in the overall microstructure in the thickness direction, which may make it difficult to secure yield strength of 500 MPa or more.

Cooling condition after finish rolling: Cooled down to 600 ℃ or lower with cooling rate over 3 ℃ / s

Due to the austenite miniaturization due to reverse transformation of the surface portion, the air-cooled ferrite transformation temperature is increased, and thus, at least a part of the fine austenite is transformed into ferrite before the cooling process after finishing rolling, and the austenite not transformed into ferrite is By cooling, it transforms into bainite.

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

The finished rolled steel is cooled to 600 ° C. or less at a cooling rate of 3 ° C./s or more.

When cooling after finishing rolling, if cooling rate becomes lower than 3 degree-C / s or cooling ends at the temperature higher than 600 degreeC, microstructure will not be formed suitably and yield strength may become less than 500 Mpa.

The steel material may preferably have a thickness of 50 mm or more.

Through the above manufacturing method, the surface microstructure of the steel is made of a mixed phase of polygonal ferrite and bainite, 1 / 4t in the center of the steel is made of more than 50% needle-like ferrite and less than 50% bainite , Steel can be produced in which the fraction of the region having bainite single phase structure in the total thickness of the steel is 20% or less.

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

The steel may preferably have a yield strength of at least 500 MPa.

The steel material may preferably have a central impact transition temperature of −40 ° C. or less.

As such, the steel composition and manufacturing conditions are controlled to secure the microstructure of the product having excellent brittle crack propagation resistance, and the air cooling air time from rough rolling to finishing rolling is shortened through bar cooling and reheating to improve productivity and refine the particle size. It is possible to provide a very thick high strength steel having a yield strength of 500MPa or more and a central impact transition temperature of -40 ° C or less.

In particular, in the present invention, since the bar is cooled by using a cooling means, the air cooling waiting time is reduced and austenite is prevented from growing, thereby improving productivity, and increasing the high angle of the steel core microstructure. A boundary particle size can be maintained at an average of 20 micrometers or less.

As described above, by controlling the steel composition and performing a cooling and recuperation process under suitable conditions before the manufacturing process, in particular, after the rough rolling and finishing rolling process, high strength ultra-thick steel having excellent brittle crack propagation resistance can be provided with high productivity. .

Hereinafter, the present invention will be described in more detail with reference to Examples. However, it is necessary to note that the following examples are provided only to illustrate the present invention by way of example and not to limit the scope of the present invention. This is because the scope of the present invention is determined by the matters described in the claims and the matters reasonably inferred therefrom.

Hereinafter, the present invention will be described through examples.

(Example)

After reheating a 400 mm thick steel slab having the composition shown in Table 1 to a temperature of 1070 ° C., rough rolling was performed at a temperature of 1025 ° C. to prepare a bar. The cumulative rolling reduction rate was roughly 50% for rough rolling.

The rough rolled bar had a thickness of 200 mm, and after the rough rolling, the bar was cooled and then recuperated. The surface recuperation temperature of Table 2 was 1 / 4t and 1 / 2t of the bar after considering the bar thickness. It means the surface temperature measured value at the time when a temperature difference becomes less than 20 degreeC. The cooling of the bar was carried out such that the surface of the bar had a temperature of less than Ac3, and the 1 / 4t (where t is the bar thickness) region had a temperature of at least 50 ° C. above the finish rolling start temperature. At this time, the cooling rate at the time of cooling the bar was 1 ~ 5 ℃ / sec.

At the completion of the recuperation, finish rolling was performed immediately to obtain a steel sheet having the thickness shown in Table 2, and then cooled to a temperature in the range of 500 to 300 ° C. at a cooling rate of 3.5 to 5 ° C./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 Inventive Steel 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 Inventive Steel 5 0.066 1.82 0.27 0.15 0.018 0.028 0.31 0.037 46 24 Comparative Steel 1 0.12 2.01 0.52 0.21 0.021 0.019 0.20 0.041 49 9 Comparative Steel 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. Steel grade Steel plate thickness (mm) Whether to apply bar cooling Surface recuperation temperature after bar cooling (℃) Particle size of inversely transformed austenite after recuperation (μm) Inventive Example 1 Inventive Steel 1 80 ○ Ac3 + 45 50 or less Inventive Example 2 Inventive Steel 2 90 ○ Ac3 + 22 50 or less Inventive Example 3 Invention Steel 3 95 ○ Ac3 + 38 50 or less Inventive Example 4 Inventive Steel 4 100 ○ Ac3 + 52 50 or less Inventive Example 5 Inventive Steel 5 80 ○ Ac3 + 39 50 or less Comparative Example 1 Inventive Steel 2 90 ○ Ac3-52 Measurement impossible due to partial transformation Comparative Example 2 Invention Steel 3 95 ○ Ac3-76 Measurement impossible due to partial transformation Comparative Example 3 Inventive Steel 1 80 X - - Comparative Example 4 Inventive Steel 4 100 X - - Comparative Example 5 Comparative Steel 1 80 ○ Ac3 + 33 50 or less Comparative Example 6 Comparative Steel 2 85 ○ Ac3 + 59 50 or less Comparative Example 7 Comparative Steel 3 90 ○ Ac3 + 65 50 or less Comparative Example 8 Comparative Steel 4 90 ○ Ac3 + 29 50 or less Comparative Example 9 Comparative Steel 5 95 ○ Ac3 + 45 50 or less

The microstructure analysis results and yield strength / central impact transition temperature results for the steels prepared according to Tables 1 and 2 are shown in Table 3 below.

In addition, the CAT (Crack Arrest Test) evaluation for the steel material was carried out at -10 ℃ by using the ESSO equipment, crack propagation (Propagate) / stop (Arrest) is shown in Table 3 below.

The core particle size of the following Table 3 was measured by the EBSD method, and means the value measured by calculating the grain boundary having a high boundary angle of 15 degrees or more by using the measurement results.

Example No. Steel grade Surface microstructure Bainite Single Phase Area (%) Central to 1 / 4t microstructure fraction (%) Yield strength (Mpa) Center Average Particle Size (㎛) Impact Transition Temperature (℃) CAT (@ -10 ℃) Inventive Example 1 Inventive Steel 1 PF + B 15 AF (83) B (17) 512 17.3 -53 Arrest Inventive Example 2 Inventive Steel 2 PF + B 17 AF (74) B (26) 532 18.5 -49 Arrest Inventive Example 3 Invention Steel 3 PF + B 8 AF (61) B (39) 575 16.1 -65 Arrest Inventive Example 4 Inventive Steel 4 PF + B 13 AF (89) B (11) 506 17.7 -78 Arrest Inventive Example 5 Inventive Steel 5 PF + B 12 AF (59) B (41) 568 13.1 -71 Arrest Comparative Example 1 Inventive Steel 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 Steel 1 PF + B 49 AF (18) B (82) 678 28.7 -32 Propagate Comparative Example 6 Comparative Steel 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

[Table 3] PF: Polygonal Ferrite, P: Pearlite, AF: Acicular Ferrite, B: Bainite

As shown in Table 3, in Comparative Examples 1 and 2, since the recuperation temperature after cooling the bar is lower than Ac3 temperature, the surface bainite structure generated during cooling of the bar remains coarse without being transformed into austenite again. As the bainite single-phase tissue area is more than 20%, it can be seen that the crack did not stop and propagated in the -10 ° C CAT test indicating brittle crack propagation resistance.

In the case of Comparative Examples 1 and 2 it was re-reduced below the Ac3 temperature, so that only a part of the austenite produced during the cooling did not occur again with austenite.

In the case of Comparative Steels 3 and 4, the bar cooling proposed in the present invention is not applied, so that the microstructure fraction is coarse during the air cooling after the rough rolling even though the microstructure fraction is included in the present invention. Therefore, it can be seen that the average particle size of the core is 20 microns or more, and thus the center impact transition temperature is -40 ° C or higher, and the crack does not stop in the -10 ° C CAT test indicating brittle crack propagation resistance.

 In the case of Comparative Example 5 has a value higher than the upper limit of the C proposed in the present invention, a large amount of bainite structure is generated due to excessive hardening ability, the center impact transition temperature is -40 ℃ or more, brittle crack propagation resistance It can be seen that the crack did not stop in the CAT test of -10 ° C and propagated.

In the case of Comparative Example 6 by having a value higher than the Mn upper limit proposed in the present invention, a large amount of bainite structure was generated due to excessive hardening ability, the center impact transition temperature is -40 ℃ or more, brittle crack propagation The resistivity of the -10 ° C CAT test indicates that the crack did not stop and propagated.

In the case of Comparative Example 7 has a lower value than the lower limit of C, Mn presented in the present invention, the hardenability is insufficient to produce a large amount of polygonal ferrite and pearlite structure, it can be seen that the yield strength is 500MPa or less.

In the case of Comparative Example 8, by having a value higher than the upper limit of Ni presented in the present invention, a large amount of bainite structure was generated due to excessive hardening ability, which resulted in a central impact transition temperature of -40 ° C or higher, and brittle crack propagation. The resistivity of the -10 ° C CAT test indicates that the crack did not stop and propagated.

In the case of Comparative Example 9, by having a value higher than the upper limit of Ti and Nb proposed in the present invention, a large amount of bainite structure is generated due to excessive hardening ability, and coarse TiN or (TiNb), (C, N) precipitates. Thus, it can be seen that the crack propagated without stopping in the CAT test at -10 ° C showing brittle crack propagation resistance.

On the contrary, in the case of Inventive Examples 1 to 5, which satisfied the component range proposed in the present invention and performed the bar cooling after rough rolling, and then restored the surface portion to a temperature higher than Ac3 temperature, the particle size of the center part was fine and less than 20 microns, It can be seen that the bainite single phase region is 20% or less with respect to the entire thickness region, and the central portion of the surface portion to 1 / 4t submicrostructure is composed of 50% or more acicular ferrite and 50% or less bainite. Therefore, in the case of the invention examples 1 to 5 it can be seen that the yield strength 500MPa or more, the center impact transition temperature -40 ℃ or less, -10 ℃ exhibits excellent brittle crack propagation resistance that the crack is stopped in the CAT test.

Claims (13)

By weight%, C: 0.03-0.09%, Mn: 1.4-2.2%, Ni: 0.2-0.9%, Nb: 0.005-0.05%, Ti: 0.005-0.04%, Cu: 0.1-0.5%, Si: 0.05- 0.5%, Al: 0.01 ~ 0.05%, P: 100ppm or less, S: 40ppm or less, including the remaining Fe and other unavoidable impurities, the surface part is made of a mixed phase of polygonal ferrite and bainite, and the thickness is 1 / 2t to The 1 / 4t (where t: steel thickness) section consists of acicular ferrite of 50 vol% or more and bainite of 50 vol% or less, and has a 20% fraction of the area having bainite single phase structure in the overall thickness of the steel. Extremely thick steel with excellent brittle crack propagation resistance. The ultra-thin steel having excellent brittle crack propagation resistance according to claim 1, wherein the grain size having a high-angle boundary of the central microstructure of the steel has an average of 20 micrometers or less. The ultra-thick steel material of claim 1, wherein the steel has a brittle crack propagation resistance of 500 MPa or more. The ultra-thick steel material of claim 1, wherein the impact shock transition temperature of the steel is excellent in brittle crack propagation resistance of -40 ° C or less. The ultra-thick steel material of claim 1, wherein the steel has a brittle crack propagation resistance of 50 mm or more. By weight%, C: 0.03-0.09%, Mn: 1.4-2.2%, Ni: 0.2-0.9%, Nb: 0.005-0.05%, Ti: 0.005-0.04%, Cu: 0.1-0.5%, Si: 0.05- Reheating the steel slab containing 0.5%, Al: 0.01% to 0.05%, P: 100ppm or less, S: 40ppm or less, and remaining Fe and other unavoidable impurities to a temperature of 1150 to 1000 ° C; Rough rolling the reheated slab at a temperature of 1150 to 900 ° C; Cooling the crude rolled bar by using cooling means; Reheating the cooled bar above an Ac3 temperature on a surface basis; Finishing rolling the regenerated bar at a temperature of Ar3 or higher on a 1 / 4t basis; And after finishing rolling, cooling to a temperature of 600 ° C. or less at a cooling rate of 3 ° C./s or more, wherein the cooling of the bar has a temperature of less than Ac 3 at a surface portion of the bar, where t is Bar thickness) region is a method for producing an extremely thick steel material excellent in brittle crack propagation resistance is carried out to have a temperature higher than 50 ℃ above the finish rolling start temperature. The method of claim 6, wherein the thickness of the steel is 50mm or more. 7. The method according to claim 6, wherein the surface temperature of the bar reheated in the step of reheating is Ac3 + 20 ° C to Ac3 + 70 ° C. The method of claim 6, wherein in the step of cooling the bar, a bainite phase, a needle-like ferrite phase or a mixed phase thereof is formed on the surface of the bar, the preparation of the ultra-thick steel with excellent brittle crack propagation resistance. Way. 10. The method according to claim 9, wherein in the step of reheating the cooled bar, the bainite phase, acicular ferrite phase, or a mixed phase thereof in the surface portion is inversely transformed into austenite. . The method of claim 10, wherein the inversely transformed austenite has a particle size of 50 micrometers (µm) or less. The method of claim 6, wherein the cooling rate of the roughly rolled bar is about 1 to 10 ° C / s (sec) on a 1 / 4t basis, and has excellent brittle crack propagation resistance. The method of claim 6, wherein the cooling rate of the roughly rolled bar is about 2-5 ° C / s (sec) on a 1 / 4t basis.