JP2010209433A - Steel material superior in toughness of weld heat-affected zone and fatigue characteristics of base metal, and method for manufacturing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a steel material which shows superior HAZ toughness even when having been welded at such a high heat input as an amount of 50 kJ/mm or more, and besides, has improved fatigue characteristics of a base metal itself. <P>SOLUTION: The steel material has following characteristics of: (a) including an oxide containing Zr, a REM and Ca; (b) having the composition satisfying 5-50% ZrO<SB>2</SB>, 10-50% oxide of a REM and 5.0-50% CaO, when the compositions of all oxides have been measured and converted into the composition of single oxide; (c) containing 120 pieces per 1 mm<SP>2</SP>or more of oxides with circle equivalent diameters of 0.1-2.0 μm, and 5 pieces per 1 mm<SP>2</SP>or less of oxides with circle equivalent diameters of more than 5.0 μm; (d) having a metal structure in which an average circle equivalent diameter D of crystal grains that are surrounded by a large-angle grain boundary having a crystal orientation difference of 15 degrees or more is 25 μm or less, and a ratio R of random grain boundaries occupying in the large-angle grain boundaries is 70 area% or less, when observed with an EBSP method; and (e) having an average hardness of 170 Hv or more. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention relates to a steel material used for a bridge, a high-rise building, a ship, and the like, and in particular, a part that is affected by heat when welded (hereinafter referred to as “welding heat affected zone” or “HAZ”). The steel material is excellent in toughness and the fatigue characteristics of the steel material itself, and a method for producing the steel material.

  The properties required for steel materials used in bridges, high-rise buildings, ships and the like have become increasingly severe in recent years, and particularly good toughness is required. Generally, these steel materials are often joined by welding, but among the welded joints, particularly HAZ has a problem that the toughness is likely to deteriorate due to thermal influence during welding. This toughness deterioration becomes more prominent as the heat input during welding increases, and the cause is that the larger the heat input during welding, the lower the cooling rate of the HAZ and the lower the hardenability, resulting in coarse island martensite. It is thought to be in the generation. Therefore, in order to improve the toughness of the HAZ, it is considered that the heat input during welding should be suppressed as much as possible. However, on the other hand, in order to increase the welding work efficiency, it is desired to employ a high heat input welding method in which the heat input of welding is 50 kJ / mm or more, such as electrogas welding, electroslag welding, submerged welding, and the like.

Therefore, the present applicant has proposed steel materials that suppress the HAZ toughness deterioration when the high heat input welding method is adopted in Patent Documents 1 to 3. These steel materials are characterized in that they contain REM oxide and / or CaO and ZrO ₂ as oxides. Since the oxide exists in a liquid state in molten steel, it is finely dispersed in the steel. In addition, the oxide is thermally stable, and, for example, it does not dissolve and disappear even when exposed to a high temperature of 1400 ° C. for a long time, which greatly contributes to the improvement of HAZ toughness.

  By the way, since structures such as bridges, high-rise buildings, and ships are subject to fatigue failure due to repeated stress, steel materials used as materials may have good fatigue characteristics from the viewpoint of ensuring the safety of the structure. It is desired. The fatigue process of a steel material is roughly classified into two processes: the generation of a crack in a stress concentration portion and the occurrence of the crack that propagates in the steel material.

  Regarding the occurrence of cracks, since fatigue cracks in structures often occur from the start end of welded joints, it has been studied to improve the shape of the start of welded joints in order to suppress the occurrence of fatigue cracks.

  In contrast, fatigue cracks that occur in steel materials propagate through the steel material, but when the cracks progress, the steel material itself eventually breaks. Alternatively, it is desirable to make the crack difficult to progress by making the crack propagation rate as small as possible. Non-patent documents 1 to 3 have been proposed as techniques for reducing the growth rate of fatigue cracks. Non-Patent Document 1 describes the progress of fatigue cracks by bypassing fatigue cracks by making the metal structure a mixed structure of a soft layer (ferrite) and a hard layer (martensite) and flattening the hard layer (martensite). A technique for reducing the speed is disclosed. Non-Patent Documents 2 and 3 disclose a technique in which a hard layer is dispersed in a soft layer (ferrite) to bypass the fatigue crack and reduce the fatigue crack growth rate. In Non-Patent Document 2, bainite is generated as a hard layer, and in Non-Patent Document 3, pearlite is generated as a hard layer.

Japanese Patent Laid-Open No. 2007-1001000 JP 2007-247004 A JP 2007-247005 A

Nakajima Kiyotaka et al., "Effect of hard second phase on fatigue crack growth and weld joint fatigue properties of thick steel plates", Proceedings of Welding Structure Symposium 2004, 2004, page 335 Noboru Honda et al., “New Functional Thick Steel Plate and Its Joint Properties that Self-Suppress Fatigue Crack Propagation”, Journal of Welding Society, Vol. 74, No. 4, 2005, p. 25 Satoshi Iki et al., “High-performance steel for shipbuilding-Life cycle cost reduction technology of JFE steel”, JFE Technical Report No. 5, JFE Steel, August 2004, p.13

  The present invention has been made in view of such circumstances, and its purpose is excellent in HAZ toughness even when large heat input welding with a heat input of 50 kJ / mm or more is performed, and the mother It is to provide a steel material with improved fatigue characteristics of the material itself. Moreover, the other objective of this invention is to provide the manufacturing method of the said steel material.

The steel material excellent in toughness and base metal fatigue characteristics of the weld heat-affected zone according to the present invention that has solved the above-mentioned problems is C: 0.02 to 0.12% (meaning “mass%”. The same applies to chemical components and oxides.), Si: 0.5% or less (including 0%), Mn: 1-2%, Zr: 0.0002-0.050%, REM: 0.0002-0. 050%, Ca: 0.0005 to 0.010%, Ti: 0.005 to 0.02%, N: 0.0040 to 0.01%, O: 0.0005 to 0.010%, P: 0 0.02% or less (excluding 0%), S: 0.015% or less (not including 0%), Al: 0.05% or less (including 0%), the balance being iron and inevitable impurities Steel material consisting of
(A) The steel material includes an oxide containing Zr, REM, and Ca,
(B) When the composition of all oxides contained in the steel is measured and converted as a single oxide, ZrO ₂ : 5 to 50%, REM oxide (M ₂ O _{3 when} REM is represented by the symbol M) ): 10 to 50%, and CaO: 5.0 to 50%,
(C) of the total oxides contained in the steel material, the circle equivalent diameter oxides 0.1~2.0μm is 1 mm ² per 120 or more, circles oxides 5.0μm greater in equivalent diameter 1 mm ² Less than 5,
(D) When the metal structure of the steel material is observed by a backscattered electron diffraction method (EBSP method), the following expressions (1) and (2) are satisfied:
(E) It has a gist in that the average hardness of the steel material is 170 Hv or more.
D ≦ 25 μm (1)
R ≦ 70 area% (2)
[However, in the formula (1), D is a measurement of the orientation difference between two adjacent crystals by the EBSP method, and the average equivalent circle diameter of a crystal grain surrounded by a large-angle grain boundary with a crystal orientation difference of 15 ° or more ( μm).
In the formula (2), R means the ratio (area%) of the random grain boundary in the large angle grain boundary. ]

The steel material, as another element,
(1) Ni: 0.4% or less (not including 0%), Cu: 0.3% or less (not including 0%), Cr: 1.5% or less (not including 0%), and Mo At least one element selected from the group consisting of 1% or less (excluding 0%),
(2) Nb: 0.1% or less (not including 0%), V: 0.1% or less (not including 0%), and B: 0.005% or less (not including 0%) At least one element selected from the group,
Etc. may be contained.

The steel material of the present invention adjusts the dissolved oxygen content of the molten steel to a range of 0.0010 to 0.0060%, and then stirs the molten steel to float and separate oxides in the molten steel, thereby reducing the total oxygen content to 0.000. After adjusting to the range of 0010 to 0.0070%, after adding Ti, then Zr, REM, and Ca to adjust the component composition, casting is performed, and at the time of hot rolling, the steel slab is changed to 1100 to 1250. After heating to ℃, the cumulative reduction ratio is 40% or more while cooling so that the average cooling rate per pass is 1.5 ° C / second or more in the temperature range of 1050 ° C or less and Ar ₃ point + 100 ° C. First rolling is performed, and then the second rolling is performed so that the maximum rolling reduction per pass is 12% or less and the cumulative rolling reduction is 50% or more in a temperature range exceeding Ar ₃ point + 100 ° C. and exceeding the Ar ₃ point. Do After it is prepared by cooling at an average cooling rate of 5 ° C. / sec or more to a temperature range where the surface temperature is 500 ° C. or less.

After adding Zr, REM, and Ca, it is preferable to perform casting after stirring the molten steel within a range not exceeding 40 minutes. The first rolling is preferably performed in the austenite recrystallization temperature range. After cooling to a temperature range where the surface temperature of the rolled material is 500 ° C. or less at an average cooling rate of 5 ° C./second or more, tempering treatment may be performed in a temperature range of 500 ° C. or more and less than Ac ₁ point.

  According to the present invention, a predetermined amount of oxide (an oxide containing Zr, REM, and Ca) serving as a nucleus of intragranular ferrite transformation is generated, and the size and number of oxides present in the steel ( Since the particle size distribution is also appropriately controlled, a steel material having excellent HAZ toughness during high heat input welding can be provided. In particular, in the steel material of the present invention, it has been clarified that not only a predetermined amount or more of a fine oxide having an equivalent circle diameter of 0.1 to 2 μm useful for improving HAZ toughness is present, but also adversely affecting HAZ toughness improvement. Since the number of coarse oxides having an equivalent circle diameter of more than 5.0 μm is significantly suppressed, even if welding is performed with a larger heat input than the HAZ toughness evaluation method disclosed in the example of Patent Document 1 above. HAZ toughness can be increased.

  Further, according to the present invention, when the metal structure is observed by the EBSP method, the metal structure is surrounded by a large-angle grain boundary having a crystal orientation difference of 15 ° or more (hereinafter sometimes simply referred to as “large-angle grain boundary”). The average equivalent circle diameter D of the crystal grains is controlled to 25 μm or less, the ratio R of random grain boundaries occupying the large angle grain boundaries is controlled to 70% or less, and the average hardness of the base material is controlled to 170 Hv or more. In addition, the fatigue characteristics of the base metal itself can be improved.

FIG. 1 is a schematic explanatory diagram showing the shape of a compact test piece (CT test piece) used in the examples. FIG. 2 is a graph showing the relationship between the average equivalent circle diameter D of a crystal grain surrounded by a large-angle grain boundary and the impact property (vE- ₆₀ ). FIG. 3 is a graph showing the relationship between the ratio R of random grain boundaries in the large-angle grain boundaries and the fatigue crack growth rate (da / dN). FIG. 4 is a graph showing the relationship between the average hardness of the rolled material and the fatigue crack growth rate (da / dN). FIG. 5 is a graph showing the relationship between the cumulative rolling reduction at Ar ₃ point + 100 ° C. or less and the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries. FIG. 6 is a graph showing the relationship between the γ grain size and the average equivalent circle diameter D of crystal grains surrounded by a large-angle grain boundary. FIG. 7 is a graph showing the relationship between the cumulative rolling reduction and γ particle size at 1050 ° C. or lower. FIG. 8 is a graph showing the relationship between the average equivalent circle diameter D of crystal grains surrounded by the large angle grain boundaries and the ratio R of random grain boundaries in the large angle grain boundaries.

  The present invention improves the intragranular ferrite transformation technique disclosed in Patent Document 1 above, and obtains a steel material that has good HAZ toughness and excellent fatigue characteristics of the base metal itself even when welding is performed with a larger heat input. For technology.

  That is, the present invention is based on the technique disclosed in Patent Document 1 and has been studied to further improve the HAZ toughness. As a result, all the oxides in the steel (the oxides serving as nuclei of intragranular ferrite transformation) The size and number of the oxides are not limited, and the number and size of the oxides are deeply involved in the improvement of HAZ toughness. In particular, the number of coarse oxides having a circle equivalent diameter of more than 5.0 μm is 5 or less. It was found that a steel material excellent in HAZ toughness can be obtained even if high heat input welding with a heat input of about 50 kJ / mm is performed if suppressed. Furthermore, the present inventors have repeatedly studied to improve the fatigue characteristics of the base metal itself after improving the HAZ toughness as described above. As a result, the inventors have found that fatigue characteristics can be improved by controlling the metal structure and hardness of the steel material in a well-balanced manner, thereby completing the present invention. Specifically, when the orientation difference between two adjacent crystals is measured by the EBSP method for the metallographic structure of the steel material, the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundary having a crystal orientation difference of 15 ° or more. It has been clarified that the fatigue characteristics of the base metal itself can be improved if the ratio R of random grain boundaries in the large-angle grain boundaries is 70 area% or less and the hardness of the steel is 170 Hv or more.

  That is, in order to improve the fatigue properties by improving the toughness of the base metal itself, it is effective to refine the metal structure, and in particular, crystal grains surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more. The average equivalent circle diameter D may be 25 μm or less. However, since the grain boundaries increase when the structure is refined, the growth rate of cracks generated in the steel material increases. This is because cracks tend to propagate through the grain boundaries, and the growth path of cracks increases as the grain boundaries increase. Thus, as a result of studies aimed at improving the fatigue characteristics of the base metal itself while miniaturizing the structure, it has been found that it is sufficient to reduce random grain boundaries where the grain boundary energy is high and fatigue cracks are likely to progress. That is, even in a large-angle grain boundary having a crystal orientation difference of 15 ° or more, energy is not high at all shift angles, and at a specific shift angle, a grain called “corresponding grain boundary” having extremely low grain boundary energy. There is a world (for example, “Materials Histology”: Seto Takagi, Kanzaki Tsuzaki, Asakura Shoten, page 45). In other words, large-angle grain boundaries are broadly divided into “corresponding grain boundaries” with low grain boundary energy and “random grain boundaries” with high grain boundary energy. Fatigue cracks tend to develop at random grain boundaries, but the corresponding grain boundaries are It is known that progress is difficult. Therefore, in the present invention, based on the viewpoint of reducing the ratio R of random grain boundaries in the large-angle grain boundaries having a crystal orientation difference of 15 ° or more, the ratio R is reduced to 70 area% or less to improve the fatigue characteristics of the base material. It is premised on improvement.

  However, as a result of further studies by the present inventors, the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries having a crystal orientation difference of 15 ° or more as described above is set to 25 μm or less, and the large-angle grain boundaries It has been found that simply setting the proportion R of the random grain boundaries to occupy 70% by area or less is not sufficient for improving the fatigue characteristics of the base material, and that it is extremely important that the hardness of the steel material be 170Hv or more.

  That is, if the ratio R of the random grain boundaries in the large-angle grain boundaries is reduced, the paths where fatigue cracks are likely to occur are reduced, so the fatigue cracks propagate through the steel while bypassing the crystal grain boundaries so as to select random grain boundaries. To do. However, fatigue cracks that can no longer propagate through the grain boundaries may propagate through the crystal grains, not the grain boundaries. Accordingly, as described above, even if the rate R of random grain boundaries in the large-angle grain boundaries is reduced to reduce the fatigue crack growth rate, if the fatigue crack growth rate in the crystal grains is not reduced at the same time, eventually, However, the fatigue characteristics of the base metal itself cannot be improved. And in this invention, the steel material hardness useful for prevention of the progress of the fatigue crack in a crystal grain was determined as 170 Hv or more based on the viewpoint that the speed of the fatigue crack which propagates in the crystal grain becomes larger as the crystal grain becomes softer. It depends on you.

  In the present specification, for the convenience of explanation, the former is used for distinguishing between an oxide that is a nucleus of intragranular ferrite transformation, that is, an oxide containing Zr, REM, and Ca, and all oxides contained in the steel material. Is particularly referred to as “Zr · REM · Ca-based oxide”, and the latter is particularly referred to as “total oxide”.

  In addition, the essential components (Zr, REM, and Ca) that constitute the “Zr / REM / Ca-based oxide” may be particularly referred to as intragranular ferrite transformation nucleation elements.

  Moreover, in this specification, among all the oxides contained in steel materials, an oxide with an equivalent circle diameter of 0.1 to 2.0 μm is referred to as a “fine oxide”, and an equivalent circle diameter of more than 5.0 μm. Oxides may be referred to as “coarse oxides”.

  According to the present invention, since the number of coarse oxides is remarkably suppressed, the HAZ toughness evaluation method disclosed in the example of Patent Document 1 [from 800 ° C. after holding at a heating temperature of 1400 ° C. for 5 seconds] A heat cycle that cools the temperature up to 500 ° C. in 300 seconds (heat input condition: 1400 ° C. × 5 seconds, cooling time Tc = 300 seconds) and measures absorbed energy at −40 ° C.] The HAZ toughness can be improved even if welding is performed. Specifically, the steel material of the present invention has an absorption energy when a heat cycle (heat input condition: 1400 ° C. × 30 seconds, cooling time Tc = 300 seconds) in which the holding time at 1400 ° C. is increased to 30 seconds is given. Even if it evaluates by the method measured similarly to the above, favorable HAZ toughness is shown (refer the example mentioned below).

  Hereinafter, the requirements (a) to (e) constituting the present invention will be described in detail.

[(A) Zr / REM / Ca oxide]
First, the Zr / REM / Ca-based oxide that is the starting point of the intragranular ferrite transformation will be described. The above-mentioned Zr / REM / Ca-based oxide means an oxide containing a Zr oxide, a REM oxide, and a Ca oxide. The elements (intragranular ferrite transformation nucleation elements) constituting the Zr / REM / Ca-based oxide are Zr, REM, and Ca. In addition to the above, for example, oxides such as Ti, Mn, Si, and Al It may contain forming elements and other steel components.

The existence form of the Zr / REM / Ca-based oxide is not particularly limited, and may be present as a single oxide containing an intragranular ferrite transformation nucleation element alone, or may be present as an intragranular ferrite transformation nucleation element. You may exist as complex oxide containing 2 or more types. Examples of the single oxide include ZrO _{2 for} Zr; CaO for Ca; and REM for REM represented by the symbol “M”, such as M ₂ O ₃ , M ₃ O ₅ , and MO ₂ . These oxides may exist in an aggregated state, or may exist in a form in which other compounds such as sulfides and nitrides are complex-deposited on the oxides.

The Zr / REM / Ca-based oxide preferably further contains an oxide of Ti. When Ti oxide further exists, intragranular ferrite transformation is promoted, and the improvement of HAZ toughness is further enhanced. The oxide of Ti may exist as a single oxide (for example, Ti ₂ O ₃ , Ti ₃ O ₅ , TiO ₂ ), or a composite oxide with the Zr / REM / Ca-based oxide. May exist.

[(B) Average composition of oxide]
When the steel material of the present invention measures the composition of all oxides contained in the steel material and is converted into mass as a single oxide (total is 100%), ZrO ₂ is 5 to 50%, REM oxide (REM When represented by the symbol M, M ₂ O ₃ ) satisfies 10 to 50% and CaO satisfies 5.0 to 50%, and this effectively acts as a nucleus of intragranular ferrite transformation. If the lower limit value of each oxide is not reached, the amount of oxide that forms intragranular ferrite nuclei at the time of welding is insufficient, and the effect of improving HAZ toughness is not exhibited. On the other hand, when the upper limit value of each oxide is exceeded, the oxide becomes coarse, the number of fine oxides that effectively act as the nuclei of intragranular ferrite decreases, and the HAZ toughness improving effect is not exhibited effectively.

The ZrO ₂ is 5% or more, preferably 10% or more, more preferably 13% or more, and further preferably 15% or more. On the other hand, the upper limit is 50%, preferably 45%, more preferably 40%.

The oxide of the REM is 10% or more, preferably 15% or more, more preferably 20% or more, and further preferably 30% or more. On the other hand, the upper limit is 50%, preferably 45%, more preferably 40%. In addition, when the REM is represented by the symbol M, the REM oxide is present in the steel material in the form of M ₂ O ₃ , M ₃ O ₅ , MO _2, etc., but all the REM oxide is converted to M ₂ O ₃ . It means the amount when converted.

  The CaO is 5.0% or more, preferably 10% or more, more preferably 15% or more, and still more preferably 18% or more. On the other hand, the upper limit is 50%, preferably 45%, more preferably 40%, and particularly preferably 30%.

The remaining components of the total oxide composition are not particularly limited, and examples thereof include oxides of oxide-forming elements contained in the steel material of the present invention (for example, SiO ₂ , Al ₂ O ₃ , MnO, etc.).

  The composition of the total oxide contained in the steel material is determined by observing the surface of the steel material with, for example, an electron probe X-ray micro analyzer (EPMA) and quantitatively analyzing the oxides observed in the observation field. And measure. Details of the measurement conditions will be described in the column of Examples described later.

[(C) Particle size distribution of all oxides]
Next, the number and size of all oxides that characterize the present invention (all oxides present in the steel material, not limited to the above-described Zr / REM / Ca-based oxides) will be described.

The total oxide contained in the steel material of the present invention includes 120 or more fine oxides having an equivalent circle diameter of 0.1 to 2.0 μm per 1 mm ^{2 of the} observation visual field area, and 5.0 μm in equivalent circle diameter. The exceeding coarse oxide satisfies 5 or less per 1 mm ^{2 of the} observation visual field area.

  When the present inventors examined in detail the relationship between the particle size distribution of all oxides and the HAZ toughness, in particular, a fine oxide having an equivalent circle diameter of 0.1 to 2.0 μm and a coarse oxidation of more than 5.0 μm. Are closely related to the HAZ toughness of high heat input welding, and it is the number of fine oxides that greatly contributes to the improvement of HAZ toughness. Coarse oxides are the starting point of brittle fracture. It has been found that HAZ toughness is reduced. It has also been clarified that fine oxides having an equivalent circle diameter of less than 0.1 μm hardly contribute to the HAZ toughness improving effect by oxide dispersion. Therefore, in order to increase the HAZ toughness, it is preferable that the number of fine oxides is as large as possible. However, since the number of coarse oxides tends to increase as the number of fine oxides increases, It is necessary to appropriately control the number of oxides and coarse oxides.

The preferred number of fine oxides is 200 or more per 1 mm ^{2 of} the observation visual field area, more preferably 500 or more, and still more preferably 1000 or more.

The smaller the amount of coarse oxide, the better. The number is preferably 3 or less, more preferably 1 or less, and most preferably 0 per 1 mm ² of the viewing field area. The present invention is not limited to the number of oxides having sizes other than those described above, and it has been confirmed by experiments that the desired HAZ toughness can be obtained as long as the oxides having the above sizes are controlled.

  The “circle equivalent diameter” is a diameter of a circle assumed to have the same oxide area, and is recognized on a transmission electron microscope (TEM) observation surface.

[(D) Regarding the average equivalent circular diameter D of crystal grains surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more, and the ratio R of random grain boundaries in the large-angle grain boundaries]
The steel material of the present invention needs to satisfy the following formulas (1) and (2) when the metal structure is observed by a backscattered electron diffraction image method (EBSP method). Satisfying both equations improves the fatigue properties of the base metal itself.
D ≦ 25 μm (1)
R ≦ 70 area% (2)
In the above formula (1), D measures the orientation difference between two adjacent crystals by the EBSP method, and the average equivalent circle diameter (μm) of the crystal grains surrounded by a large-angle grain boundary with a crystal orientation difference of 15 ° or more. Means.

  In the present invention, the D value is 25 μm or less. By setting the D value to 25 μm or less, it is possible to ensure the toughness of the base material itself. Further, it is generally known that fatigue cracks bend, detour, or stop at large-angle grain boundaries having a crystal orientation difference of 15 ° or more. Therefore, by refining crystal grains surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more, the positions where fatigue cracks bend, detour, and stop increase, and as a result, the impact properties of the base material increase. In addition, the fatigue characteristics of the base metal itself are improved (see FIG. 2 below). Further, as will be described later, it is effective to refine the crystal grains surrounded by the large-angle grain boundaries in order to reduce the random grain boundary ratio R to a predetermined value or less (see FIG. 8 described later). This is because the ratio R of the random grain boundaries is easily reduced as the crystal grains surrounded by the large-angle grain boundaries are refined. The smaller the D value, the better, preferably 23 μm or less, and more preferably 20 μm or less. Note that the lower limit of the D value is about 13 μm.

  The D value is measured by observing the metal structure at the t / 2 position in the thickness direction when the thickness of the steel material is t (mm). The reason why the D value at the t / 2 position is used as a reference is that, as the plate thickness increases, it becomes more difficult to apply strain, and thus the metal structure is most difficult to control. A specific measurement procedure will be described in the section of Examples described later.

  In the above formula (2), R is the ratio (area%) of the random grain boundary occupying the large angle grain boundary where the crystal orientation difference is 15 ° or more when the orientation difference between two adjacent crystals is measured by the EBSP method. I mean.

  In the present invention, the R value is 70 area% or less. As described above, large-angle grain boundaries are divided into corresponding grain boundaries and random grain boundaries, but by reducing random grain boundaries where fatigue cracks tend to progress, the path for cracks to bypass becomes longer and the crack growth rate is increased. Since it can be reduced, the fatigue characteristics can be improved. In addition, the crack energy is attenuated in the process of propagation, so by increasing the propagation path of the crack, the crack is stopped by reducing the energy, and the fatigue characteristics of the base metal itself can be improved ( (See FIG. 3 below).

  The smaller the R value, the better, preferably 65 area% or less, more preferably 60 area% or less. Note that the lower limit of the R value is about 40 area%.

  The R value is measured by observing the metal structure at the surface in the thickness direction, at the t / 4 position, and at the t / 2 position, where the thickness of the steel material is t (mm). The reason for observing the surface of the steel material is that fatigue cracks are likely to occur on the surface of the steel material, and the reason for observing the t / 4 position is that it is a general reference for evaluating the characteristics of the entire steel material. The reason for observing the t / 2 position is that, as the plate thickness increases, it becomes difficult to impart strain, and therefore the metal structure is most difficult to control. A specific measurement procedure will be described in the section of Examples described later.

[(E) Average hardness]
In the present invention, the average hardness of the steel material is 170 Hv or more. By stiffening the steel material to some extent, even when a crack that has become difficult to propagate through the grain boundary propagates within the grain, the speed of propagation within the grain can be reduced. As a result, it is possible to reduce the growth rate of cracks that propagate in the steel (see FIG. 4 below).

  The higher the average hardness, the better, preferably 180 Hv or higher, more preferably 190 Hv or higher, and still more preferably 200 Hv or higher.

  In order to set the average hardness of the steel material to 170 Hv or more, it is recommended that the metal structure is mainly composed of bainite. This is because the hardness of the steel material tends to decrease as the ferrite fraction in the metal structure increases. By bainite-based, it means that the bainite fraction is about 60% by area or more when the metal structure is observed with an electron microscope. In addition, the upper limit of average hardness is about 260Hv in general. This is because if it becomes too hard, cracks do not propagate in the grains and it is difficult to reduce the crack growth rate. The upper limit of this average hardness is a value substantially equal to the average hardness of the bainite structure.

  The hardness of the steel material is measured at the t / 2 position in the thickness direction when the thickness of the steel material is t (mm). The reason for using the hardness at the t / 2 position as a reference is that as the plate thickness increases, cooling to the center of the plate thickness becomes difficult, and it is difficult to control the metal structure appropriately. This is because the hardness of steel generally depends on the form of the metal structure, and the form of the metal structure depends on the cooling rate. A specific measurement procedure will be described in the section of Examples described later.

  Next, the component composition in the steel material (base material) of the present invention will be described. In the steel material of the present invention, C: 0.02 to 0.12%, Si: 0.5% or less (including 0%), Mn: 1 to 2%, Zr: 0.0002 to 0.00 as basic components. 050%, REM: 0.0002 to 0.050%, Ca: 0.0005 to 0.010%, Ti: 0.005 to 0.02%, N: 0.0040 to 0.01%, O: 0 .0005 to 0.010%, P: 0.02% or less (not including 0%), S: 0.015% or less (not including 0%), Al: 0.05% or less (including 0%) ) Is satisfied. The reasons for setting these ranges are as follows.

  C is an element indispensable for securing the strength of the steel material (base material). In order to exert such effects, it is necessary to contain 0.02% or more. C is preferably contained in an amount of 0.04% or more, more preferably 0.05% or more. However, if C exceeds 0.12%, a lot of island martensite (MA) is generated in the HAZ at the time of welding to cause deterioration of the toughness of the HAZ, and also adversely affect the weldability. Therefore, C is 0.12% or less, preferably 0.1% or less, more preferably 0.08% or less.

  Si is an element that contributes to securing the strength of a steel material by solid solution strengthening. However, if Si exceeds 0.5%, not only does martensite (MA) form in the HAZ during welding to cause deterioration of the HAZ toughness, but also adversely affects weldability. Therefore, Si is 0.5% or less. Preferably it is 0.3% or less, More preferably, it is 0.2% or less, More preferably, it is 0.18% or less. In order to secure the strength of the steel material by adding Si, it is preferable to contain 0.02% or more, more preferably 0.05% or more, and still more preferably 0.1% or more. .

  Mn is an element that contributes to improving the strength of the steel material (base material). In order to exhibit such an effect effectively, it is necessary to contain 1% or more. Mn is preferably contained at 1.2% or more, more preferably 1.4% or more. However, if it exceeds 2%, the weldability of the steel material (base material) is deteriorated. Therefore, Mn needs to be suppressed to 2% or less. Preferably it is 1.8% or less, More preferably, you may be 1.6% or less.

Zr is an element necessary for generating ZrO ₂ . By containing ZrO ₂ , the oxide is easily finely dispersed, and the finely dispersed oxide serves as a nucleus for formation of intragranular ferrite, which contributes to improvement of HAZ toughness. However, if Zr is less than 0.0002%, the amount of fine oxides that form the nuclei of the intragranular ferrite decreases, so the HAZ toughness cannot be improved. Therefore, Zr needs to be contained in an amount of 0.0002% or more. Preferably it is 0.0005% or more, More preferably, it is 0.0010% or more. However, when Zr is added excessively, a large amount of coarse oxide is generated and the HAZ toughness deteriorates. In addition, fine Zr carbide that causes precipitation strengthening is formed, leading to a decrease in toughness of the base material itself. Therefore, Zr is suppressed to 0.050% or less. Preferably it is 0.04% or less, More preferably, it is 0.01% or less, More preferably, it is 0.008% or less.

  REM (rare earth element) and Ca are elements necessary to form respective oxides. By containing these oxides, the oxides are easily finely dispersed, and the finely dispersed oxides serve as nuclei for forming intragranular ferrite, thereby contributing to the improvement of HAZ toughness.

  REM should be contained in an amount of 0.0002% or more, preferably 0.0005% or more, more preferably 0.0010% or more, and further preferably 0.0015% or more. However, when REM is added excessively, a coarse oxide is formed, and the HAZ toughness deteriorates. Therefore, REM should be suppressed to 0.050% or less. Preferably it is 0.04% or less, More preferably, it is 0.01% or less.

  In the present invention, REM means a lanthanoid element (15 elements from La to Ln), Sc (scandium) and Y (yttrium). Among these elements, it is preferable to contain at least one element selected from the group consisting of La, Ce and Y, more preferably La and / or Ce.

  Ca should be contained in an amount of 0.0005% or more, preferably 0.001% or more, and more preferably 0.0015% or more. However, when Ca is added excessively, a coarse oxide is formed and the HAZ toughness is deteriorated. Therefore, Ca needs to be suppressed to 0.010% or less. Preferably it is 0.008% or less, More preferably, it is 0.005% or less.

  Ti is an element that generates nitrides such as TiN and Ti oxide in the steel material and contributes to the improvement of HAZ toughness. In order to exert such effects, it is necessary to contain Ti by 0.005% or more. Preferably it is 0.007% or more, more preferably 0.010% or more. However, if Ti is added excessively, the toughness of the steel (base material) is deteriorated, so Ti should be suppressed to 0.02% or less. Preferably it is 0.018% or less, More preferably, it is 0.017% or less.

  N is an element that precipitates nitrides (for example, ZrN and TiN), and the nitrides prevent the austenite grains formed in the HAZ during welding and promote ferrite transformation by the pinning effect. , Contributing to the improvement of HAZ toughness. In order to exhibit such an effect effectively, it is necessary to contain 0.0040% or more. Preferably it is 0.005% or more, More preferably, it is 0.006% or more. As N increases, Ti-containing nitrides are formed and austenite grain refinement is promoted, which effectively improves the toughness of HAZ. However, when N exceeds 0.01%, the amount of solute N increases, the toughness of the base metal itself deteriorates, and the HAZ toughness also decreases. Therefore, N must be suppressed to 0.01% or less. Preferably it is 0.0095% or less, More preferably, it is 0.009% or less.

  O (oxygen) is an element necessary for generating a fine oxide that serves as a nucleus for formation of intragranular ferrite that contributes to the improvement of HAZ toughness. However, if O is less than 0.0005%, the amount of the fine oxide is insufficient, and the HAZ toughness cannot be improved. Therefore, it is necessary to contain O by 0.0005% or more. Preferably it is 0.0010% or more, More preferably, it is 0.0015% or more. However, if added excessively, the oxide becomes coarse and deteriorates instead of the HAZ toughness. Therefore, O should be suppressed to 0.010% or less. Preferably it is 0.008% or less, More preferably, it is 0.005% or less.

  P is an element that easily segregates, and particularly segregates at a grain boundary in the steel material to deteriorate the toughness of the base material. Therefore, P must be suppressed to 0.02% or less. Preferably it is 0.018% or less, More preferably, you may be 0.015% or less.

  S is a harmful element that combines with Mn to produce sulfide (MnS) and degrades the toughness of the base material and the ductility in the thickness direction. Therefore, S must be suppressed to 0.015% or less. Preferably it is 0.012% or less, More preferably, it is 0.008% or less, More preferably, it is 0.006% or less.

  Al is an element having a strong deoxidizing power and acting as a deoxidizing agent. However, if added excessively, the oxide is reduced, making it difficult to produce the Zr / REM / Ca oxide that contributes to the improvement of HAZ toughness. Therefore, Al needs to be suppressed to 0.05% or less. Al is preferably 0.04% or less, more preferably 0.035% or less. Al may be 0%.

  The steel material of the present invention contains the above elements as essential components, and the balance may be iron and inevitable impurities (for example, Mg, As, Se, etc.).

It is also effective for the steel material of the present invention to contain, as other elements, elements (Ni, Cu, Cr, Mo) for improving the strength of the steel material, elements (Nb, V, B) for further improving the HAZ toughness, etc. It is. In particular,
(1) Ni: 0.4% or less (not including 0%), Cu: 0.3% or less (not including 0%), Cr: 1.5% or less (not including 0%), and Mo At least one element selected from the group consisting of 1% or less (excluding 0%),
(2) Nb: 0.1% or less (not including 0%), V: 0.1% or less (not including 0%), and B: 0.005% or less (not including 0%) At least one element selected from the group,
Etc. are preferably contained. The reasons for setting these ranges are as follows.

[(1) Ni, Cu, Cr, Mo]
Ni, Cu, Cr, and Mo are all elements that contribute to increasing the strength of the steel material, and can be added alone or in combination.

  In particular, Ni is an element that contributes to increasing the strength of the steel material and improving the toughness of the steel material itself. In order to exhibit such an action effectively, it is preferable to contain 0.05% or more. More preferably, it is 0.1% or more, More preferably, it is 0.2% or more. Ni is preferably contained as much as possible, but since it is an expensive element, if it is contained excessively, the cost becomes high. Therefore, the upper limit is preferably 0.4% for economic reasons. More preferably, it is 0.38% or less, More preferably, it is 0.35% or less.

  Cu is an element that enhances the strength of the steel by solid solution strengthening. In order to exhibit such an action effectively, it is preferable to contain 0.05% or more. More preferably, it is 0.1% or more, More preferably, it is 0.2% or more. However, if the content exceeds 0.3%, the toughness of the steel material deteriorates, so Cu is preferably 0.3% or less. More preferably, it is 0.28% or less, More preferably, it is 0.25% or less.

  It is preferable to contain 0.1% or more of Cr. More preferably, it is 0.2% or more, More preferably, it is 0.5% or more. However, if the Cr content exceeds 1.5%, the strength of the steel (base material) is remarkably increased and the base material toughness deteriorates, so that the HAZ toughness decreases. Therefore, Cr is preferably 1.5% or less. More preferably, it is 1.2% or less, More preferably, it is 1% or less.

  It is preferable to contain Mo 0.1% or more. More preferably, it is 0.2% or more, More preferably, it is 0.3% or more. However, if it exceeds 1%, the strength of the steel material (base material) becomes extremely high and the base material toughness deteriorates instead, so that the HAZ toughness also decreases. Therefore, Mo is preferably 1% or less. More preferably, it is 0.7% or less, More preferably, it is 0.5% or less.

[(2) Nb, V, B]
Nb, V, and B are all elements that improve the HAZ toughness, and can be added alone or in combination.

  In particular, Nb and V suppress the coarsening of austenite grains during slab heating before rolling by the solid drag effect by solid solution and the pinning effect by precipitation of carbonitride, and refine the structure to improve the HAZ toughness. It is an element having an action.

  In order to effectively exhibit such action by adding Nb, it is preferable to contain 0.002% or more. More preferably, it is 0.005% or more, More preferably, it is 0.008% or more. However, if Nb exceeds 0.1%, the precipitated carbonitrides become coarse and the toughness of the base material deteriorates. Accordingly, Nb is preferably 0.1% or less. More preferably, it is 0.08% or less, More preferably, it is 0.05% or less.

  If V is added in an amount exceeding 0.1%, as with Nb, the precipitated carbonitrides become coarse and the toughness of the base material deteriorates. Therefore, V is preferably 0.1% or less. More preferably, it is 0.08% or less, More preferably, it is 0.05% or less. In order to effectively exhibit the effect of V addition, it is preferable to contain 0.002% or more. More preferably it is 0.005% or more, and still more preferably 0.01% or more.

  On the other hand, B is an element that suppresses the formation of grain boundary ferrite and improves the HAZ toughness. However, if B exceeds 0.005%, it precipitates as BN at the austenite grain boundary and causes a reduction in HAZ toughness. Therefore, B is preferably 0.005% or less. More preferably, it is 0.004% or less. In addition, in order to exhibit the effect | action by B addition effectively, it is preferable to make it contain 0.001% or more. More preferably, it is 0.0015% or more.

  Next, a production method that can be suitably employed in producing the steel material of the present invention will be described.

  First, at the time of casting, after adjusting the dissolved oxygen content of the molten steel to a range of 0.0010 to 0.0060%, the molten steel is stirred and the oxide in the molten steel is levitated and separated to reduce the total oxygen content to 0.0010. After adjusting to a range of ˜0.0070%, Ti, then Zr, REM, and Ca are added to adjust the component composition. Thereby, the requirements (a) to (c) can be satisfied.

Next, in hot rolling, after heating the steel slab to 1100 to 1250 ° C., the average cooling rate per pass is 1.5 ° C./second or more in a temperature range exceeding 1050 ° C. and Ar ₃ point + 100 ° C. 1st rolling (corresponding to rough rolling) is performed so that the cumulative rolling reduction is 40% or more while cooling, and then the maximum rolling per pass in the temperature range of Ar ₃ point + 100 ° C or lower and above Ar ₃ point. After performing the second rolling (corresponding to finish rolling) so that the rate is 12% or less and the cumulative reduction rate is 50% or more, the average cooling rate is 5 ° C / second or more until the surface temperature becomes 500 ° C or less. Cool with. Thereby, the requirements (d) and (e) can be satisfied.

  The outline of the production method of the present invention will be described as follows.

  First, by adding a predetermined alloy element in a predetermined order to molten steel in which the dissolved oxygen amount and the total oxygen amount are adjusted, a desired oxide serving as a nucleus for formation of intragranular ferrite can be generated. Particularly in the present invention, it is extremely important to adjust the total oxygen amount after adjusting the dissolved oxygen amount so that coarse oxides are not formed.

  Dissolved oxygen means oxygen in a free state that does not form an oxide and exists in molten steel. The total oxygen means the sum of all oxygen contained in the molten steel, that is, free oxygen and oxygen forming oxides.

Next, the steel slab obtained by casting is hot-rolled. The average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries is 25 μm or less, and the ratio R of the random grain boundaries occupying the large-angle grain boundaries is to a 70% by area or less so as to define a second rolling 12% maximum reduction rate per pass in a temperature range of Ar ₃ point from + 100 ° C. until Ar ₃ point or less, the cumulative reduction of 50% or more It is important to control the hot rolling. In order to satisfy D ≦ 25 μm while achieving R ≦ 70 area% by hot rolling in this temperature range, 1 in the temperature range from 1050 ° C. to Ar ₃ point + 100 ° C. as defined as the first rolling. Rolling is performed so that the cumulative reduction ratio is 40% or more while cooling so that the average cooling rate per pass is 1.5 ° C./second or more. By performing the first rolling in this manner and setting the austenite (γ) grain size to 50 μm or less, the metal structure is appropriately controlled in the second rolling after the first rolling, and the requirements specified in the present invention are satisfied. Adjusted to Hereinafter, the production method of the present invention will be described in detail step by step.

  First, the dissolved oxygen content of the molten steel is adjusted to a range of 0.0010 to 0.0060%. If the amount of dissolved oxygen in the molten steel is less than 0.0010%, the amount of dissolved oxygen in the molten steel is insufficient, so it is not possible to secure a predetermined amount of Zr, REM, and Ca-based oxides that become the core of intragranular ferrite transformation. Excellent HAZ toughness cannot be obtained. Moreover, when the amount of dissolved oxygen is insufficient, Zr, which could not form an oxide, forms a carbide, or REM or Ca forms a sulfide, which causes the toughness of the base material itself to deteriorate. Therefore, the amount of dissolved oxygen is set to 0.0010% or more. The dissolved oxygen is preferably 0.0013% or more, more preferably 0.0020% or more.

  On the other hand, if the amount of dissolved oxygen exceeds 0.0060%, the amount of oxygen in the molten steel is too large, so that the reaction between the oxygen in the molten steel and the above elements becomes violent, which is not preferable for melting work, and coarse oxidation. A product is produced and the HAZ toughness is deteriorated. Therefore, the amount of dissolved oxygen should be suppressed to 0.0060% or less. The amount of dissolved oxygen is preferably 0.0055% or less, more preferably 0.0053% or less.

  By the way, the amount of dissolved oxygen in molten steel primarily refined in a converter or electric furnace usually exceeds 0.010%. Therefore, in the production method of the present invention, it is necessary to adjust the amount of dissolved oxygen in the molten steel to the above range by some method.

  Examples of the method for adjusting the amount of dissolved oxygen in the molten steel include a method of vacuum C deoxidation using an RH type degassing refining device, a method of adding a deacidifying element such as Si, Mn, Ti, and Al. The amount of dissolved oxygen may be adjusted by appropriately combining these methods. Moreover, you may adjust the amount of dissolved oxygen using a ladle heating type refining apparatus, a simple molten steel processing facility, etc. instead of the RH type degassing refining apparatus. In this case, since the amount of dissolved oxygen cannot be adjusted by vacuum C deoxidation, a method of adding a deacidifying element such as Si may be employed to adjust the amount of dissolved oxygen. When employing a method of adding a deoxidizing element such as Si, the deoxidizing element may be added when steel is removed from the converter to the ladle.

  After adjusting the dissolved oxygen content of the molten steel to the range of 0.0010 to 0.0060%, the molten steel is stirred, and the oxide in the molten steel is floated and separated, whereby the total oxygen content in the molten steel is 0.0010 to 0.00. Adjust to 0070%. As described above, in the present invention, the molten steel in which the amount of dissolved oxygen is appropriately controlled is stirred, and unnecessary oxides are removed, and then an intragranular ferrite transformation nucleation element such as Zr is added. Oxide generation can be prevented. In Patent Document 1 described above, since this floating separation step is not performed, a coarse oxide is generated, and good HAZ toughness cannot be ensured (see Examples described later).

  If the total oxygen amount is less than 0.0010%, the desired oxide amount is insufficient, and therefore, it is not possible to ensure the amount of oxide serving as a nucleus for formation of intragranular ferrite contributing to the improvement of HAZ toughness. Therefore, the total oxygen amount is set to 0.0010% or more. The total oxygen amount is preferably 0.0015% or more, more preferably 0.0018% or more.

  On the other hand, if the total oxygen amount exceeds 0.0070%, the amount of oxide in the molten steel becomes excessive, and a coarse oxide is generated to deteriorate the HAZ toughness. Therefore, the total oxygen amount should be suppressed to 0.0070% or less. The total oxygen amount is preferably 0.0060% or less, more preferably 0.0050% or less.

  Since the total amount of oxygen in the molten steel changes in correlation with the stirring time of the molten steel, it can be controlled by adjusting the stirring time. Specifically, the total amount of oxygen in the molten steel is appropriately controlled while appropriately measuring the total amount of oxygen in the molten steel after stirring the molten steel and removing the floating oxide.

  After adjusting the total amount of oxygen in the molten steel to the above range, Ti is added, and then Zr, REM, and Ca are added before casting. By adding the above elements to the molten steel with the total oxygen content adjusted, a Zr · Ca · REM oxide that becomes the nucleus of the desired intragranular ferrite transformation is obtained. Since Ti oxide has a smaller interfacial energy with molten steel than Zr / REM / Ca-based oxides, if alloy elements are added in this order, Ti oxide is refined, resulting in HAZ. Fine oxides that contribute to toughness can be generated.

  The form of REM, Ca, Zr, Ti added to the molten steel is not particularly limited. For example, as REM, pure La, pure Ce, pure Y, or pure Ca, pure Zr, pure Ti, and further Fe-Si- A La alloy, Fe—Si—Ce alloy, Fe—Si—Ca alloy, Fe—Si—La—Ce alloy, Fe—Ca alloy, Ni—Ca alloy, or the like may be added. Moreover, you may add misch metal to molten steel. Misch metal is a mixture of cerium group rare earth elements, and specifically contains about 40 to 50% of Ce and about 20 to 40% of La. However, since misch metal often contains Ca as an impurity, when the misch metal contains Ca, the range specified in the present invention must be satisfied.

  In the present invention, for the purpose of promoting the removal of coarse oxides, it is preferable to stir the molten steel within a range not exceeding 40 minutes after adding Zr, REM, and Ca. When the stirring time exceeds 40 minutes, fine oxides aggregate and coalesce in the molten steel, so that the oxides become coarse and the HAZ toughness deteriorates. Therefore, the stirring time is preferably within 40 minutes. The stirring time is more preferably within 35 minutes, and even more preferably within 30 minutes. The lower limit of the stirring time of the molten steel is not particularly limited, but if the stirring time is too short, the concentration of the additive element becomes non-uniform, and the desired effect cannot be obtained as a whole steel material. Accordingly, a desired stirring time corresponding to the container size is required.

  By adding Zr, REM, and Ca as described above, molten steel with an adjusted component composition can be obtained. It casts using the obtained molten steel, and obtains a steel piece.

  Although the obtained slab is hot-rolled, the heating temperature at the time of hot rolling shall be 1100-1250 degreeC. The lower limit of the heating temperature is 1100 ° C. because the metal structure of the steel slab is austenite. However, if the heating temperature is too high, the initial austenite structure becomes too coarse, and it becomes difficult to sufficiently refine the structure after transformation. Therefore, the upper limit of the heating temperature is 1250 ° C.

After heating, the average cooling rate per pass is 1.5 ° C./temperature in the temperature range where the temperature at the t / 2 position of the steel slab (t is the thickness of the steel slab) is 1050 ° C. or less and Ar ₃ point + 100 ° C. The first rolling is performed so that the cumulative rolling reduction is 40% or more while cooling to be at least 2 seconds.

By performing rolling while cooling in a temperature range of 1050 ° C. or lower and Ar ₃ point + 100 ° C., austenite grain size can be prevented and the austenite grain size can be reduced to 50 μm or less. By rolling 2 times, the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries can be controlled to 25 μm or less (see FIG. 6 described later).

  However, when the average cooling rate when rolling in the above temperature range is less than 1.5 ° C./second or the cumulative rolling reduction is less than 40%, the austenite becomes coarse and the austenite grain size exceeds 50 μm, as described later. The average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries cannot be controlled to 25 μm or less even if the second rolling is performed under conditions suitable for the above (see FIG. 7 described later). Therefore, the fatigue characteristics of the base material cannot be improved. Therefore, the average cooling rate when rolling in the above temperature range is 1.5 ° C./second or more. Preferably it is 2 degreeC / second or more, More preferably, it is 2.5 degreeC or more.

  Further, the cumulative rolling reduction is 40% or more. Preferably it is 50% or more, More preferably, it is 60% or more. The upper limit of the cumulative rolling reduction is not particularly limited, but is generally about 65%.

The temperature of the Ar ₃ point can be calculated from the following equation (3). In the formula, [] represents the content (% by mass) of each element, and t represents the finished thickness (mm) of the product.
Ar ₃ (° C.) = 910−310 × [C] −80 × [Mn] −20 × [Cu] −15 × [Cr] −55 × [Ni] −80 × [Mo] + 0.35 × (t− 8) ... (3)

The cumulative rolling reduction is a value obtained by the following equation (4). t ₀ is the rolling start thickness (mm) when the average temperature of the steel slab is in the target temperature range, and t ₁ is the rolling end thickness (mm) when the average temperature of the steel slab is in the target temperature range. means.
Cumulative rolling reduction = [(t ₀ −t ₁ ) / t ₀ ] × 100 (%) (4)

After performing the first rolling in a temperature range of 1050 ° C. or lower and Ar ₃ point + 100 ° C., the second rolling described later may be performed continuously, but the first rolling is performed at 1050 ° C. or lower and Ar ₃ point. Of the temperature range exceeding + 100 ° C., it is preferable to perform in a temperature range where austenite recrystallizes. In order to roll at a temperature exceeding the Ar ₃ point + 100 ° C. from the austenite recrystallization temperature (this temperature range may be referred to as the austenite non-recrystallization temperature range hereinafter), it is necessary to increase the rolling reduction and the equipment load is increased. Take it.

  The upper limit value of the austenite recrystallization temperature varies somewhat depending on the chemical composition of the steel slab, but is generally 1060 ° C. in the steel types targeted by the present invention. By finishing the first rolling in the austenite recrystallization temperature range, the austenite grain size can be reduced.

  The temperature at which the first rolling is completed is preferably set to the temperature of the lower limit value (the temperature immediately before reaching the austenite non-recrystallization region) in the austenite recrystallization temperature region. This is because austenite is recrystallized and coarsened when held for a long time in the austenite recrystallization temperature range after the first rolling.

After the first rolling is finished, the maximum reduction rate per pass in the temperature range where the temperature at the t / 2 position of the steel slab (t is the thickness of the steel slab) is Ar ₃ point + 100 ° C. or less and exceeds the Ar ₃ point. The second rolling is performed so that the cumulative reduction ratio is 12% or less and the cumulative reduction ratio is 50% or more.

When the second rolling is performed at a temperature exceeding the Ar ₃ point + 100 ° C., the crystal grains surrounded by the large-angle grain boundaries having a crystal orientation difference of 15 ° or more cannot be refined. On the other hand, even if the second rolling is performed at a temperature not higher than the Ar ₃ point, the crystal grains surrounded by the large-angle grain boundaries cannot be refined, and the rolling load is too low, resulting in an increased equipment load.

  The maximum rolling reduction per pass when rolling in the above temperature range is 12% or less. When the maximum rolling reduction exceeds 12%, excessive strain accumulates to cause recrystallization, and the ratio R of random grain boundaries to the large angle grain boundaries increases. Therefore, the maximum rolling reduction is 12% or less. Preferably it is 11.5% or less, More preferably, it is 11% or less.

  The cumulative rolling reduction when rolling in the above temperature range is 50% or more. When the cumulative rolling reduction is less than 50%, the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries having a crystal orientation difference of 15 ° or more becomes large (see FIG. 5 to be described later). Since the grain boundary ratio R increases, the fatigue crack growth rate increases and the fatigue characteristics of the base material cannot be improved. Therefore, the cumulative rolling reduction is 50% or more. Preferably it is 55% or more, More preferably, it is 60% or more.

When the first rolling is finished in the austenite recrystallization temperature range, the second rolling is started after cooling to a temperature range in which the temperature at the t / 2 position of the steel slab becomes Ar ₃ point + 100 ° C. or less.

  In addition, if it takes time from the end of the first rolling to the start of the second rolling for the convenience of production or facilities, after the first rolling is completed, it is quickly cooled to the austenite non-recrystallization temperature range. In this temperature range, the second rolling may be started when ready. This is to prevent the refined austenite from coarsening again by setting the standby temperature of the billet to the austenite amorphous temperature range.

After the completion of the second rolling, the steel slab is actively cooled (accelerated cooling) at an average cooling rate of 5 ° C./second or higher from a temperature exceeding the Ar ₃ point to 500 ° C. or lower from the temperature exceeding the Ar ₃ point. If the cooling start temperature is lower than the Ar ₃ point, a large amount of ferrite is generated and the hardness cannot be ensured. Therefore, the cooling start temperature is set to a temperature exceeding the Ar ₃ point. The cooling stop temperature is set to 500 ° C. or lower in order to completely complete the transformation.

The average cooling rate from the temperature exceeding the Ar ₃ point to 500 ° C. or less is 5 ° C./second or more. When the average cooling rate is less than 5 ° C./second, a large amount of ferrite is generated and the hardness cannot be ensured. Therefore, the average cooling rate is 5 ° C./second or more. It is preferably 7 ° C./second or more, more preferably 9 ° C./second or more.

After cooling to a temperature range where the surface temperature of the steel slab is 500 ° C. or less, the tempering treatment may be performed in a temperature range of 500 ° C. or more and less than Ac ₁ point. By performing the tempering treatment, the strain introduced by rolling or transformation disappears, so that the base material toughness can be improved.

The temperature at the Ac ₁ point can be calculated from the following equation (5). In the formula, [] indicates the content (% by mass) of each element.
Ac ₁ (° C.) = 723-14 × [Mn] + 22 × [Si] −14.4 × [Ni] + 23.3 × [Cr] (5)

  The steel material according to the present invention can be used, for example, as a material for structures such as bridges, high-rise buildings, ships, and the like. The toughness deterioration of the affected part can be prevented.

  The steel material of the present invention is intended for a steel material such as a thick steel plate having a thickness of about 3.0 mm or more. The excellent HAZ toughness improving effect according to the present invention is effectively exhibited when a thick steel plate having a plate thickness of 20 mm or more (particularly 40 mm or more) is used and large heat input welding with a heat input of 50 kJ / mm or more is performed.

  Hereinafter, the present invention will be described in more detail with reference to examples. However, the following examples are not intended to limit the present invention, and may be implemented with appropriate modifications within a range that can meet the purpose described above and below. These are all possible and are within the scope of the present invention.

  Using a vacuum melting furnace (capacity 150 kg), the test steel containing the chemical components (mass%) shown in Table 2 below under the conditions shown in Table 1 below (steel types a to w, the balance being iron and inevitable impurities) Was melted, cast into a 150 kg ingot, and cooled. After heating the obtained ingot to the heating temperature shown in Table 3 or Table 4 below, the slab was cooled until the average temperature at the t / 2 position (t is the thickness of the slab, the same applies hereinafter) was 1050 ° C. or less. Subsequently, hot rolling was performed to produce a rolled material having a product thickness of 10 to 80 mm. Hereinafter, the manufacturing conditions will be specifically described.

  In melting the test steel in the vacuum melting furnace, the components other than Ti, Zr, Al, REM, and Ca are adjusted, and at least one element selected from C, Si, and Mn is used. The amount of dissolved oxygen in the molten steel was adjusted by deoxidation. The amount of dissolved oxygen (ppm) after adjustment is shown in Table 1 below.

  The total amount of oxygen in the molten steel was adjusted by stirring the molten steel in which the amount of dissolved oxygen was adjusted for about 1 to 10 minutes to float and separate oxides in the molten steel. The total oxygen amount (ppm) after adjustment is shown in Table 1 below.

  After adding Ti to the molten steel in which the total oxygen content was adjusted, Zr, REM, and Ca were added to obtain a molten steel whose components were adjusted. Ti is in the form of an Fe-Ti alloy, Zr is in the form of an Fe-Zr alloy, REM is in the form of a misch metal containing about 25% La and about 50% Ce, Ca is a Ni-Ca alloy, Alternatively, each was added in the form of a Ca—Si alloy or a Fe—Ca green compact.

  However, no. In No. 22, FeO was added to adjust the total oxygen amount. In Table 1, No. In No. 7, Ti was immediately added without stirring the molten steel in which the amount of dissolved oxygen was adjusted.

  The molten steel in which the elements (Zr, REM, and Ca) were added and the components were adjusted was stirred for the time shown in Table 1 below, then cast into an ingot and cooled.

  The obtained ingot was heated to the heating temperature shown in Table 3 or 4 below, and then hot-rolled. In this example, the first rolling was performed using a roughing mill, and the second rolling was performed using a finish rolling mill.

  The rolling start temperature of rough rolling is shown in Table 3 or Table 4 below. The rolling start temperature was controlled by the average temperature at the t / 2 position of the steel slab.

Rough rolling was performed while cooling so that the average cooling rate per pass would be the rate shown in Table 3 or 4 below. The end temperature of rough rolling was Ar ₃ + 120 ° C. or higher. Table 3 or Table 4 below shows the cumulative rolling reduction in rough rolling.

After the rough rolling, the steel slab was cooled until the average temperature at the t / 2 position became Ar ₃ point + 100 ° C. or less, and finish rolling was performed. The finish rolling start temperature was Ar ₃ + 100 ° C. or lower, and the finish rolling end temperature was Ar ₃ + 40 ° C. or higher. Table 3 or Table 4 below shows the maximum reduction rate and cumulative reduction rate per pass in finish rolling.

  After the finish rolling, the hot rolled material was cooled until the surface temperature became 500 ° C. or lower. The cooling start temperature after finishing rolling is shown in Table 3 or Table 4 below. The average cooling rate from the cooling start temperature to 500 ° C. is shown in Table 3 or Table 4 below.

  In Table 3, No. 1-2 to 1-8 are the steel types a shown in Table 2 and No. 1 in Table 3. 2-2 is the steel type b shown in Table 2 and No. 2 in Table 3. 3-2 is an example in which the steel type c in Table 2 was used, and the hot rolling conditions and the like were changed.

  In Table 3, No. 1-2 and No.1. 1-3 is an example in which after the surface temperature of the hot-rolled material was cooled to a temperature range of 500 ° C. or lower, the tempering treatment was performed at the tempering temperature (managed by the surface temperature) shown in Table 3 below.

  The temperature at the t / 2 position of the billet was calculated according to the following procedure.

<Calculation method of average temperature>
(1) Using a process computer, the heating temperature at an arbitrary position in the thickness direction from the front surface to the back surface of the steel slab is calculated based on the ambient temperature from the start of heating to extraction and the in-furnace time.
(2) Using the calculated heating temperature, based on the rolling pass schedule during rolling and the cooling method (water cooling or air cooling) data between passes, the rolling temperature at any position in the plate thickness direction can be calculated using the difference method, etc. Roll while calculating using a suitable method.
(3) The surface temperature of the steel slab is measured using a radial thermometer installed on the rolling line (however, it is also calculated on the process computer).
(4) The surface temperature of the steel slab measured at the start of rough rolling, at the end of rough rolling, and at the start of finish rolling is collated with the calculated surface temperature on the process computer.
(5) When the difference between the calculated surface temperature and the measured surface temperature of the steel slab is ± 30 ° C. or more, the measured surface temperature of the steel slab is replaced with the calculated surface temperature to obtain the calculated surface temperature on the process computer.
(6) Using the corrected calculated surface temperature, obtain the temperature at the t / 2 position.

  On the other hand, the surface temperature of the hot-rolled material was measured using a radial thermometer installed on the rolling line.

Table 3 or Table 4 below also shows the product thickness (mm) of the rolled material obtained by cooling. Table 3 or Table 4 below also shows the values of the Ar ₃ point and Ac ₁ point calculated using the formulas (3) and (5) based on the chemical composition shown in Table 2 above. .

  A sample was cut out from the cross section at the t / 4 position (t is the thickness of the rolled material) of the obtained rolled material, the composition of all oxides contained in the sample was measured, and the oxide was converted into mass as a single oxide. The average composition of was calculated.

The cut sample surface was observed 600 times using “EPMA-8705 (device name)” manufactured by Shimadzu Corporation, and the component composition was quantitatively analyzed for particles having a maximum diameter of 0.2 μm or more. The observation conditions are an acceleration voltage of 20 kV, a sample current of 0.01 μA, an observation visual field area of 1 to 5 cm ² , and an analysis number of 100, and quantitative analysis of the component composition at the center of the particle by wavelength dispersion spectroscopy of characteristic X-rays. did. The analysis target elements are Mn, Ti, Zr, La, Ce, Ca, Si, Al, and O (oxygen). Using a known substance, the relationship between the electron beam intensity and the element concentration of each element is obtained in advance as a calibration curve. Then, the element concentration of the particles was quantified from the electron beam intensity obtained from the particles and the calibration curve in advance.

Of the obtained quantitative results, particles having an oxygen content of 5% or more were regarded as oxides, and the average composition of those converted in mass as a single oxide was defined as the average composition of oxides. The average composition of all oxides is shown in Table 3 or Table 4 below. The REM oxide, when the metal element is represented by M, exists in the form of M ₂ O ₃ , M ₃ O ₅ , and MO _{2 in} the steel material, but all oxides are converted to M ₂ O _3. The composition was measured. When a plurality of elements were observed from one inclusion, the composition of the oxide was calculated in terms of a single oxide of each element from the ratio of X-ray intensity indicating the presence of these elements.

As a result of observing the sample surface with EPMA, most of the observed oxides were composite oxides containing Ti, Zr, REM, and Ca, but Ti ₂ O ₃ , ZrO ₂ , REM as single oxides. An oxide of CaO was also produced.

In addition, the equivalent circle diameter of the oxide having an oxygen content of 5% or more is measured in the obtained quantitative result, and the number of oxides having an equivalent circle diameter (particle diameter) of 0.1 to 2.0 μm and the equivalent circle. The number of oxides having a diameter (particle diameter) exceeding 5.0 μm was calculated. Table 5 below shows the number of oxides converted per observation field of 1 mm ² .

  Next, in order to evaluate the toughness of HAZ which is affected by heat during welding, a welding reproduction test shown below was performed by simulating high heat input welding. In the welding reproduction test, a sample cut from the rolled material was heated to 1400 ° C., kept at this temperature for 30 seconds, and then given a heat cycle for cooling. The cooling rate was adjusted so that the cooling time from 800 ° C. to 500 ° C. was 300 seconds.

The impact characteristics of the sample after cooling were evaluated by measuring the absorbed energy (vE _-40 ) at -40 ° C by a V-notch Charpy test. A material having a vE- ₄₀ of 100 J or more is regarded as acceptable (haz toughness is good). The measurement results are shown in Table 5 below.

  Next, the metal structure of the obtained rolled material is observed by the following procedure, and the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries having a crystal orientation difference of 15 ° or more, and the random grains occupying the large-angle grain boundaries. The field ratio R was determined. The values of D (μm) and R (area%) are shown in Table 6 below.

<< Calculation method of D >>
(1) Prepare a sample cut parallel to the rolling direction (longitudinal direction) so as to include both the front and back surfaces of the rolled material.
(2) Polishing with a wet emery polishing paper of # 150 to # 1000 or a polishing method having the same function as that, and a mirror finish using a polishing agent such as diamond slurry.
(3) The mirror-polished surface is an EBSP (Electron Back Scattering Pattern) device manufactured by TexSEM Laboratories. The measurement range is 200 μm × 200 μm and the pitch is 0.5 μm at the t / 2 position in the plate thickness direction. The difference is measured, and a boundary having a crystal orientation difference of 15 ° or more is defined as a large-angle grain boundary. Note that measurement points whose confidence index indicating the reliability of the measurement direction is smaller than 0.1 are excluded from the analysis target.
(4) In the grain distribution map, the maximum width (usually along the plate thickness direction) and the maximum length (usually along the rolling direction) of the crystal grain surrounded by the large-angle grain boundary having a crystal orientation difference of 15 ° or more. Length), the area of the crystal grain is calculated, the equivalent circle diameter of the crystal grain is calculated, and the average value is obtained.

<< R calculation method >>
(1) The ratio R of the random grain boundaries in the large-angle grain boundaries is a EBSP apparatus manufactured by TexSEM Laboratories, using a sample that has been mirror-finished under the same conditions as when D is calculated. At the surface in the thickness direction, at the t / 2 position and at the t / 4 position, the measurement range is 200 μm × 200 μm, the pitch is 0.5 μm, and the orientation difference between the two crystals is measured. Note that measurement points whose confidence index indicating the reliability of the measurement direction is smaller than 0.1 are excluded from the analysis target.
(2) Among the measurement results, those having a crystal orientation difference of less than 5.5 ° are considered to be noise and deleted, and the distribution at each orientation difference up to 62.5 ° is obtained.
(3) By associating the crystal orientation difference distribution created in the step (2) with the corresponding grain boundary map (a table in which the number of each corresponding grain boundary is described), the large-angle grain boundary at each plate thickness position The ratio R of the occupied random grain boundary is calculated. Specifically, the distribution of each corresponding grain boundary is obtained by dividing each corresponding grain boundary (Σ1 to 49) by the number of large-angle grain boundaries having an orientation difference of 15 ° or more obtained from the crystal orientation distribution. By subtracting from 100%, the ratio R of random grain boundaries to the large angle grain boundaries at each plate thickness position is calculated. The maximum random grain boundary ratio R at each plate thickness position is defined as the random grain boundary ratio R in the rolled material [other than the corresponding grain boundary is set as a random grain boundary (> Σ49)].

  In addition, “TSL OIM Data Collection ver5.2” manufactured by TSL Co., Ltd. was used for the measurement of the corresponding grain boundary, and “TSL OIM Analysis ver5.0” manufactured by TSL Co., Ltd. was used for the analysis.

  Moreover, the metal structure of the obtained rolled material was observed according to the following procedure, and the austenite particle size (γ particle size) and hardness were determined. The values of γ particle size are shown in Table 6 below.

<Measuring method of γ particle size>
(1) The γ grain size of the rolled material was corroded using an extremely low carbon corrosive solution using a sample that had been mirror-finished under the same conditions as when calculating D, and the former γ grain boundary was revealed. Then, a region at the position t / 2 in the plate thickness direction was photographed with an optical microscope at a magnification of 100 times or 400 times to obtain a photograph of 6 cm × 8 cm. That is, it corresponds to 600 μm × 800 μm at an observation magnification of 100 × and 150 μm × 200 μm at an observation magnification of 400 ×.
(2) The 6 cm side of the photograph corresponds to the plate thickness direction, and the 8 cm side corresponds to the rolling direction. The photographed photograph was subjected to image analysis, and the average particle diameter of γ grains found in the observation field was determined.

<Measurement method of hardness>
The hardness of the rolled material was measured with a Vickers hardness tester using a sample that was mirror-finished under the same conditions as when D was calculated. The measurement position was a t / 2 position in the thickness direction of the rolled material, the load was 10 kgf, the number of measurements was 20 points, and the average value was the hardness of the rolled material.

  Next, the fatigue characteristics of the obtained rolled material were evaluated by the following procedure.

<Evaluation of impact characteristics>
The impact characteristics of the rolled material were evaluated by performing a V-notch Charpy test and measuring the impact characteristics of the rolled material by measuring the absorbed energy (vE- ₆₀ ) at -60 ° C. The measurement of vE- ₆₀ was carried out according to JIS Z2242 by collecting U4 test piece defined by NK (Japan Maritime Association) classification from t / 4 position. The measurement results are shown in Table 6 below. In the shipbuilding E grade in the NK class, the impact characteristics of the base material are evaluated at a test temperature of −40 ° C. In this experimental example, the conditions are more stringent and the test temperature is −60 ° C., and the absorbed energy (vE _-60 ) The average value of 100 J or more was determined to be acceptable (the low temperature toughness of the base material was good).

<Measurement of fatigue crack growth rate>
A compact test piece (CT test piece) having a thickness of 12.5 mm was taken from the t / 4 position of the rolled material. The shape of the collected CT test piece is shown in FIG. The fatigue crack growth rate was calculated | required by implementing a fatigue crack growth test based on ASTM E647 using this CT test piece. The test conditions are as follows.

Test method: An electro-hydraulic servo type ± 10 ton fatigue tester was used, and the crack length was measured by a computer-controlled compliance method. Compliance means the ratio (δ / P) of crack opening displacement δ to load P, and the crack length is automatically measured from this compliance.
Test environment: Room temperature, in air Control method: Load control Control waveform: Sine wave Stress ratio: R = 0.1
Test speed: 600-1200 cpm

Assuming that the crack occurred from the weld toe and progressed to the HAZ and the base metal, the evaluation was made with a value of ΔK = 20 (MPa · √m), which is the middle ΔK region. The ΔK region in this test was found to be a stable growth region where the Paris law defined by the following equation (6) is satisfied.
da / dN = C (ΔK) ^m (6)
Here, a: crack length, n: number of repetitions, C, m: constant determined by matters such as material and load, and ΔK: stress intensity factor range, respectively.

In addition, since it is known that the value of the fatigue crack growth rate by the CT test varies, the average value of the data of the conventional steel plate was evaluated as a reference. Since the average value of the data of the conventional steel sheet is about 5.4 × 10 ⁻⁵ mm / cycle, in the present invention, less than 5.4 × 10 ⁻⁵ mm / cycle was accepted. In addition, the average value of the data of the conventional steel sheet was determined based on “Metallic material fatigue crack growth resistance data” published by the Japan Society of Materials Science.

Based on these results, FIG. 2 shows the relationship between the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries and the impact characteristics (vE- ₆₀ ). From FIG. 2, it can be seen that it is useful to set the average equivalent circle diameter D to 25 μm or less in order to improve impact characteristics at −60 ° C.

FIG. 3 shows the relationship between the ratio R of random grain boundaries in the large-angle grain boundaries and the fatigue crack growth rate (da / dN). FIG. 3 shows that it is useful to set the ratio R of the random grain boundaries to 70 area% or less in order to make the fatigue crack growth rate less than 5.4 × 10 ⁻⁵ mm / cycle.

FIG. 4 shows the relationship between the average hardness of the rolled material and the fatigue crack growth rate (da / dN). FIG. 4 shows that it is useful to set the average hardness of the rolled material to 170 Hv or more in order to make the fatigue crack growth rate less than 5.4 × 10 ⁻⁵ mm / cycle.

FIG. 5 shows the relationship between the cumulative rolling reduction at Ar ₃ point + 100 ° C. or less and the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries. FIG. 5 shows that it is useful to set the cumulative rolling reduction at Ar ₃ point + 100 ° C. or less to 50% or more in order to set the average equivalent circle diameter D to 25 μm or less.

  FIG. 6 shows the relationship between the γ grain size and the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries. FIG. 6 shows that it is useful to set the γ particle size to 50 μm or less in order to make the average equivalent circle diameter D 25 μm or less.

  FIG. 7 shows the relationship between the cumulative rolling reduction at 1050 ° C. or less and the γ particle size. FIG. 7 shows that it is useful to set the cumulative rolling reduction at 1050 ° C. or less to 40% or more in order to make the γ particle size 50 μm or less.

  FIG. 8 shows the relationship between the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries and the ratio R of random grain boundaries in the large-angle grain boundaries. From FIG. 8, there is a correlation between the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundaries and the ratio R of the random grain boundaries in the large-angle grain boundaries, and the crystal grains surrounded by the large-angle grain boundaries are made smaller. It can be seen that random grain boundaries can be reduced.

  Next, based on the results shown in Table 5, the HAZ toughness when welding will be considered.

No. 1-1, no. 2-1. 3-1. 23 is an example that satisfies the requirements defined in the present invention, and contains a predetermined amount of ZrO ₂ , REM oxide, and CaO when the composition of all oxides is measured and converted into mass as a single oxide. In order to prevent the generation of a coarse oxide having an equivalent circle diameter of more than 5.0 μm after adjustment, a large amount of fine oxide having an equivalent circle diameter of 0.1 to 2.0 μm is generated. A steel material having good HAZ toughness is obtained.

  On the other hand, no. Examples 4 to 22 are examples that do not meet any of the requirements defined in the present invention.

  No. 4 is an example in which the amount of dissolved oxygen in the molten steel is inadequate, and the amount of fine oxides that form the nucleus of intragranular ferrite contributing to the improvement of HAZ toughness cannot be ensured, improving the HAZ toughness. Not done. No. No. 5 is an example in which the amount of dissolved oxygen in the molten steel is excessive, and a coarse oxide is generated to deteriorate the HAZ toughness. No. 6 is an example in which the amount of dissolved oxygen in the molten steel is insufficient, and is an example in which the total amount of oxygen in the molten steel before adding Ti or the like is also insufficient. For this reason, the amount of fine oxides that serve as nuclei for intragranular ferrite contributing to the improvement of HAZ toughness cannot be ensured, and the HAZ toughness cannot be improved.

  No. 7 is an example of simulating the composition of the steel material disclosed in Patent Document 1. After adjusting the amount of dissolved oxygen in the molten steel, the total amount of oxygen before adding Ti exceeds the range specified in the present invention because the total amount of oxygen is not adjusted by stirring the molten steel. Therefore, the coarse oxide increases and the HAZ toughness is deteriorated.

  No. 8 and no. In No. 9, since Zr, REM, and Ca are added and the molten steel is stirred for a long time, oxides in the molten steel are aggregated together to generate a large amount of coarse oxide. Therefore, the HAZ toughness is deteriorated.

No. No. 10 is an example in which Al is too much, a coarse oxide is generated, and the HAZ toughness is deteriorated. No. No. 11 is an example in which there is too little Ti, and the HAZ toughness cannot be improved because the number of fine nitrides that form the intranuclear ferrite nuclei contributing to the improvement of the HAZ toughness decreases. No. No. 12 is an example in which there is too much Ti, and the oxide is coarsened and the HAZ toughness is deteriorated. No. 13 is an example in which there is too little REM, because the amount of oxide of REM when converted to a single oxide is small, and the amount of fine oxide that becomes the nucleus of formation of intragranular ferrite contributing to the improvement of HAZ toughness decreases. HAZ toughness has not been improved. No. No. 14 is an example in which there is too much REM, and the amount of REM oxide when converted to a single oxide is large, and a coarse oxide is produced to deteriorate the HAZ toughness. No. No. 15 is an example in which Zr is too small, since the amount of ZrO ₂ when converted to a single oxide is small, and the amount of fine oxides that become the formation nuclei of intragranular ferrite contributing to the improvement of HAZ toughness is reduced. The toughness has not been improved. No. No. 16 is an example in which there is too much Zr, and the amount of ZrO ₂ when converted to a single oxide is large, and a coarse oxide is generated to deteriorate the HAZ toughness. No. 17 is an example in which there is too little Ca, and the amount of CaO when converted to a single oxide is small, and the amount of fine oxides that form the intranuclear ferrite nuclei contributing to the improvement of HAZ toughness is reduced. Has not improved. No. No. 18 is an example in which there is too much Ca, and the amount of CaO when converted to a single oxide is large, and a coarse oxide is generated to deteriorate the HAZ toughness. No. No. 19 is an example in which there is too much N. The amount of solute N increases, the toughness of the base metal itself deteriorates, and the HAZ toughness also decreases. No. No. 20 is an example in which the amount of N is too small, and since the formation of Ti-containing nitrides is suppressed, coarsening of austenite grains due to the pinning effect cannot be prevented, and the HAZ toughness is deteriorated. No. No. 21 is an example in which there is too little O, and the amount of fine oxides that form the intranuclear ferrite formation nuclei is insufficient, and the HAZ toughness cannot be improved. No. No. 22 is an example in which O is too much, and the oxide is coarsened and the HAZ toughness is deteriorated.

  Next, based on the results shown in Table 6, the fatigue characteristics of the base metal itself will be considered.

  No. 1-1 to 1-6, no. 2-1. 2-2, No. 3-1. 3-2, no. No. 23 is an example using steel types a to c that satisfy the requirements specified in the present invention, and the metal structure of the steel material is appropriately controlled, so that the fatigue characteristics of the base material are excellent.

  No. 1-7 and No. 1 Although 1-8 is an example using the steel type a which satisfies the requirements prescribed | regulated by this invention, since the steel materials are too hard, the fatigue characteristics of a base material are inferior.

  No. Nos. 15 to 17 have inferior fatigue properties because the metal structure of the steel is not properly controlled. In particular, since the average equivalent circle diameter D of crystal grains surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more exceeds 25 μm, the impact characteristics are also inferior.

  No. Nos. 10, 12 to 14, 18, 21 and 22 have good impact characteristics among fatigue characteristics because the metallographic structure of the steel material is not properly controlled, but the fatigue crack growth rate cannot be reduced.

  In addition, No. Nos. 4-9, 11, 19, and 20 have good fatigue characteristics of the base metal because the metallographic structure of the steel is appropriately controlled, but the component composition of the steel satisfies the requirements defined in the present invention. Therefore, the HAZ toughness when welded as described above cannot be improved.

  Next, the results shown in Tables 5 to 6 will be considered together.

  No. in Table 6 1-1 to 1-6, no. 2-1. 2-2, No. 3-1. 3-2, no. 23 is an example using the steel types a to c and the steel type w that satisfy the requirements specified in the present invention. As is apparent from Table 6, the fatigue characteristics of the base material are also excellent.

  On the other hand, no. 1-7, No. 1 1-8, No. 1 Nos. 4 to 22 are examples that do not satisfy any of the requirements defined in the present invention, and are examples in which at least one of HAZ toughness or base metal fatigue characteristics is inferior.

Claims (7)

C: 0.02 to 0.12% (meaning “mass%”. The same applies to chemical components and oxides hereinafter),
Si: 0.5% or less (including 0%),
Mn: 1-2%
Zr: 0.0002 to 0.050%,
REM: 0.0002 to 0.050%,
Ca: 0.0005 to 0.010%,
Ti: 0.005 to 0.02%,
N: 0.0040 to 0.01%
O: 0.0005 to 0.010%,
P: 0.02% or less (excluding 0%),
S: 0.015% or less (excluding 0%),
Al: 0.05% or less (including 0%) is satisfied,
The balance is steel consisting of iron and inevitable impurities,
(A) The steel material includes an oxide containing Zr, REM, and Ca,
(B) When the composition of all oxides contained in the steel material is measured and converted as a single oxide,
ZrO ₂ : 5 to 50%,
REM oxide (REM represented by M ₂ O ₃ ): 10 to 50%, and CaO: 5.0 to 50%,
(C) Of all oxides contained in the steel material,
More than 120 oxides with a circle equivalent diameter of 0.1 to 2.0 μm per 1 mm ² ,
There are 5 or less oxides with a circle equivalent diameter of more than 5.0 μm per 1 mm ² ,
(D) When the metal structure of the steel material is observed by a backscattered electron diffraction method (EBSP method), the following expressions (1) and (2) are satisfied:
(E) A steel material excellent in toughness and base metal fatigue characteristics of the weld heat affected zone, wherein the steel material has an average hardness of 170 Hv or more.
D ≦ 25 μm (1)
R ≦ 70 area% (2)
[However, in the formula (1), D is a measurement of the orientation difference between two adjacent crystals by the EBSP method, and the average equivalent circle diameter of a crystal grain surrounded by a large-angle grain boundary with a crystal orientation difference of 15 ° or more ( μm).
In the formula (2), R means the ratio (area%) of the random grain boundary in the large angle grain boundary. ] The steel material is still another element,
Ni: 0.4% or less (excluding 0%),
Cu: 0.3% or less (excluding 0%),
The steel material according to claim 1, comprising at least one selected from the group consisting of Cr: 1.5% or less (not including 0%) and Mo: 1% or less (not including 0%). The steel material is still another element,
Nb: 0.1% or less (excluding 0%),
3. The composition according to claim 1, comprising at least one selected from the group consisting of V: 0.1% or less (not including 0%) and B: 0.005% or less (not including 0%). Steel material. A method for producing the steel material according to any one of claims 1 to 3,
After adjusting the amount of dissolved oxygen in the molten steel to a range of 0.0010 to 0.0060%, the total amount of oxygen is reduced to 0.0010 to 0.0070% by stirring the molten steel to float and separate oxides in the molten steel. After adjusting to the range, after adding Ti, then Zr, REM, and Ca to adjust the component composition according to any one of claims 1 to 3, performing casting,
At the time of hot rolling, after the steel slab is heated to 1100 to 1250 ° C., the average cooling rate per pass is 1.5 ° C./second or more in a temperature range of 1050 ° C. or less and Ar ₃ point + 100 ° C. The first rolling is performed so that the cumulative rolling reduction is 40% or more while cooling,
Next, after performing the second rolling so that the maximum reduction rate per pass is 12% or less and the cumulative reduction rate is 50% or more in a temperature range exceeding Ar ₃ point + 100 ° C. and Ar ₃ point,
A method for producing a steel material excellent in toughness and base metal fatigue characteristics of a weld heat-affected zone, characterized by cooling at an average cooling rate of 5 ° C / second or more to a temperature range where the surface temperature is 500 ° C or less. The manufacturing method of Claim 4 which casts, after adding the said Zr, REM, and Ca, stirring a molten steel in the range which does not exceed 40 minutes. The manufacturing method according to claim 4 or 5, wherein the first rolling is performed in an austenite recrystallization temperature range. 7. The tempering treatment is performed in a temperature range of 500 ° C. or more and less than Ac ₁ point after cooling at a mean cooling rate of 5 ° C./second or more to a temperature range where the surface temperature of the rolled material becomes 500 ° C. or less. The manufacturing method in any one of.

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