JP5234921B2 - High-strength thick steel plate with excellent strain aging characteristics and manufacturing method thereof - Google Patents

Description

  The present invention relates to a thick steel plate that is mainly used as a material for ship structural materials, and particularly has a high strength thickness that can reduce as much as possible a decrease in toughness after imparting strain while ensuring high strength. It relates to a meat steel plate and a useful method for producing such a steel plate.

  In recent years, ultra-large container ships have been promoted, and accordingly, the structural members of ships are becoming thicker. For example, from the 2002 maximum loading capacity of 6000 TEU, a plan to change it to 10,000 TEU is now underway, and it is necessary to further increase the thickness and strength of the steel sheet. In particular, in the hatch corner portion of a container ship, in order to reduce the corner curvature r from the viewpoint of increasing stress and increasing the number of containers loaded, a high-strength steel plate (for example, a thickness of 60 mm or more) is required. It has been demanded. Also, from the viewpoint of ensuring safety, it is also necessary to ensure toughness after being subjected to strain.

  Strain aging occurs when the steel plate is strained. This strain aging increases the yield strength of the steel plate because C and N in the steel fix the dislocations generated by applying the strain. This is a phenomenon in which material toughness (steel plate toughness) decreases. Further, the deterioration of toughness due to strain aging is more likely to occur as the steel sheet has higher strength. For this reason, the high-strength thick steel plates used for the above applications have good properties that do not cause toughness deterioration due to strain aging (this property is referred to as “strain aging properties” in the present invention). It is also required to be.

In order to prevent toughness deterioration due to strain aging, it is known that it is useful to reduce the amount of free C and N in the steel sheet. As a steel plate with improved strain aging characteristics, for example, the technique of Patent Document 1 has been proposed. In this technique, by controlling the content of Ti and Nb in the steel sheet, free C and N are fixed as precipitates, and by rolling in the non-recrystallization temperature range during finish rolling, The crystal grains at the grain boundaries are refined, thereby trapping free C and N and suppressing toughness deterioration due to strain aging. And in this technique, the manufacturing method which implemented rolling in the low recrystallization temperature range by the side of low temperature is implemented. However, just by refining the crystal grains at the large-angle grain boundaries, the strain aging characteristics are not necessarily improved, and the low temperature toughness required for shipbuilding E grade in the NK class (absorbed energy at −40 ° C. vE ₋₄₀ 100J or more) may not be secured.
Japanese Patent No. 3848091

  The present invention has been made paying attention to the circumstances as described above, and its purpose is to appropriately define each crystal orientation relationship, thereby ensuring high strength and high strain aging characteristics. It is to provide a strong thick steel plate and a useful method for producing such a steel plate.

  The high-strength thick steel plate of the present invention that has achieved the above object is C: 0.10 to 0.16% (meaning “mass%”, the same applies to the chemical composition), Si: 0.15 ˜0.30%, Mn: 1.30 to 1.60%, Al: 0.015 to 0.05%, Cu: 0.15 to 0.35, Ni: 0.10 to 0.30%, Mo : 0.10 to 0.25%, V: 0.030 to 0.05%, Nb: 0.005 to 0.015%, Ca: 0.005% or less (excluding 0%) and N: 0 0.002 to 0.008% of each steel sheet, the balance being iron and inevitable impurities, the average equivalent circle diameter of the crystal grains surrounded by large-angle grain boundaries where the orientation difference between the two crystals is 15 ° or more It is necessary that D is 35 μm or less and the random grain boundary fraction R measured from the crystal orientation distribution difference is 50 area% or more. And it has a.

  In the steel sheet of the present invention, it is effective to further contain (a) Ti: 0.005 to 0.020%, (b) Cr: 0.10% or less (not including 0%), etc., if necessary. In addition, the characteristics are further improved depending on the type of element contained.

In the high-strength thick steel plate as described above, the yield point is 480 MPa or more, the tensile strength is 590 MPa or more, and after applying 10% strain, at -40 ° C. when subjected to aging treatment at 250 ° C. for 1 hour. Thus, the characteristic that the average impact absorption energy vE- ₄₀ is 100 J or more is exhibited.

Moreover, in manufacturing the steel plate of the present invention as described above, the steel slab is heated to a temperature of Ac ₃ transformation point or higher to 1200 ° C., and the cumulative reduction is performed in the austenite recrystallization temperature range where the average temperature of the steel plate is 900 ° C. or higher. The steel sheet is subjected to rolling at a rate of 10% or more, and then the average cooling rate between the passes of the entire steel sheet is 0.3 ° C./second or more in the non-recrystallization temperature range where the average temperature of the steel sheet is 800 ° C. or more and 890 ° C. or less. The steel sheet is rolled so that the cumulative reduction ratio is 25% or more and less than 50% while being cooled so that the average temperature of the steel sheet is (Ar ₃ transformation point + 10 ° C.) or more and (Ar ₃ transformation point + 90 ° C.) or less. From the temperature range to the temperature range where the steel sheet surface temperature is 500 ° C. or lower, the average cooling rate is cooled at a cooling rate of 5 ° C./second or more, and the tempering treatment is performed in the temperature range of 500 ° C. or higher and less than the Ac ₁ transformation. You can do it.

  In the steel sheet of the present invention, along with the chemical composition, by appropriately defining the grain size of crystal grains having a crystal orientation relationship and a specific crystal orientation difference, a steel sheet having excellent strain aging characteristics can be realized, Such a steel plate is useful as a material for a hatch corner portion of a super-large container ship.

  The inventor has studied means for improving the strain aging characteristics of a steel sheet from various angles. As a result, the following knowledge was obtained. In other words, the structure of the steel sheet is generated with some orientation relation, but the orientation relation of each crystal lattice selected by the chemical composition of the steel sheet, the formation temperature of the structure, other conditions, etc. If the crystal grains having a certain crystal orientation difference are refined and the fraction (area%) of the random grain boundaries among the large angle grain boundaries is increased to 50 area% or more, The present inventors have found that strain aging characteristics are improved and completed the present invention. Hereinafter, the effects of the present invention will be described along the background of the completion of the present invention.

  Strain aging is a phenomenon in which the yield strength increases and the base material toughness (steel plate toughness) decreases as a result of free C and N in the steel fixing dislocations generated by applying strain. Therefore, in order to improve the strain aging characteristics (to prevent characteristic deterioration due to strain aging), it is important to trap free C and N so that dislocations are not fixed.

  In order to achieve the effect as described above, it is necessary to satisfy the requirement that the average equivalent circle diameter D of the crystal grains surrounded by the large-angle grain boundary where the orientation difference between the two crystals is 15 ° or more is 35 μm or less. . A large-angle grain boundary having a crystal misorientation of 15 ° or more is a high-energy grain boundary, so that an interstitial element such as C and N enters and stabilizes, so that it becomes an effective trap site. In order to improve the strain aging characteristics, the average equivalent circle diameter D (hereinafter sometimes referred to as “average large angle particle diameter D”) of the crystal grains surrounded by the large angle grain boundaries is 35 μm or less. is necessary. The “crystal orientation difference” is also called “shift angle” or “tilt angle”. Moreover, in order to measure such a crystal orientation difference, an EBSP method (Electron Backscattering Pattern Method) may be employed.

  It has already been known that refinement of the average large-angle particle diameter D is effective in improving the strain aging characteristics (Patent Document 1). However, it has also been found that the strain aging characteristics may not always be improved by simply refining the average large angle particle diameter D.

  Therefore, the present inventor further studied to eliminate such inconvenience. As a result, when the fraction (area%) of the random grain boundary is increased to 50% or more among the crystal grains surrounded by the large-angle grain boundary (the orientation difference between the two crystals is 15 ° or more), it is good. Strain aging characteristics were obtained.

  Even at a large-angle grain boundary having a crystal orientation difference of 15 ° or more, energy is not high at all shift angles, and there exists a grain boundary called a corresponding grain boundary at which a grain boundary energy is extremely low at a specific shift angle ( For example, “Materials Histology”: Nobuo Takagi, Kanaki Tsuzaki, Asakura Shoten, page 45). Large-angle grain boundaries can be divided into corresponding grain boundaries and random grain boundaries, and there is a low possibility that free interstitial elements such as C and N enter the corresponding grain boundaries. For these reasons, it is important not only to refine the average large-angle grain diameter D but also to increase the fraction of random grain boundaries in order to improve the strain aging characteristics of the steel sheet.

  By the way, in order to reduce the average large-angle grain size D, the growth of austenite grains (γ grains) can be suppressed by rolling in the low-temperature-side non-recrystallization temperature region, and strain can be efficiently transformed before transformation. The structure after transformation can be refined. However, in general, the more the processing strain is introduced at a low temperature, the selection of the variant after the transformation (crystallographically equivalent orientation relationship existing in the cubic crystal) is determined, so the random grain boundary tends to decrease. (For example, “Recrystallization and Material Structure”: Eiichi Furubayashi, Uchida Otsukuru, page 45). This is considered to be the reason why the good strain aging characteristics were not exhibited in the conventional technique (Patent Document 1).

  On the other hand, as the rolling temperature region becomes higher, the amount of processing strain introduced into the initial γ grains decreases, so that the variant selectivity is lost and the fraction of random grain boundaries tends to increase (for example, “recrystallization and "Material organization": Eiichi Furubayashi, Uchida Otsukaku, page 45). Further, since the cooling start temperature is increased, the hardenability effect is increased, and the introduction of random grain boundaries can be promoted. However, since the rolling temperature range is high, it is easily affected by temperature rise due to processing heat generation, and γ grains before transformation grow, and there is a concern about increase in average large-angle grain diameter D after transformation (that is, coarsening of crystal grains). Will be.

  In order to make the above two technologies compatible, it is necessary to strictly define the manufacturing conditions. Detailed conditions for producing the steel sheet of the present invention will be described later, but the points in the production method of the present invention are as follows.

  When rolling in the high-temperature non-recrystallized region as usual, it acts in the direction of increasing the fraction of random grain boundaries, but growth of γ grains occurs before transformation, and the large-angle grain boundary size after transformation is coarse. There is a concern about the change. Therefore, in the present invention, by performing rolling while actively cooling in the non-recrystallization temperature region on the high temperature side, the temperature rise of the steel sheet due to heat generated by plastic deformation is suppressed, and the increase of random grain boundaries and γ grains before transformation He tried to balance growth control. The detailed conditions will be described later.

  In the steel plate of the present invention, it is necessary to appropriately control the chemical component composition, but the reasons for limiting the ranges of these components are as follows.

[C: 0.10 to 0.16%]
C is an element necessary for ensuring the strength of the steel sheet. In order to obtain the minimum strength as a thick steel plate for shipbuilding, that is, about 590 MPa (depending on the thickness of the steel material to be used), it is necessary to contain 0.10% or more. However, if the content exceeds 0.16%, the weldability and the base metal toughness are adversely affected. For these reasons, the C content is set to 0.10 to 0.16%. In addition, the minimum with preferable C content is 0.11%, and a preferable upper limit is 0.14%.

[Si: 0.15 to 0.30%]
Si is an element useful for improving the strength of the base material and deoxidizing the molten steel. In order to exhibit the effect effectively, it is necessary to contain 0.15% or more. However, if the content exceeds 0.30%, weldability and base metal toughness deteriorate. In addition, the minimum with preferable Si content is 0.17%, and a preferable upper limit is 0.25%.

[Mn: 1.30 to 1.60%]
Mn is useful as an element for improving the strength of a steel sheet, and in order to exert such effects, it is necessary to contain Mn in an amount of 1.30% or more. However, if it is contained excessively, weldability and base metal toughness are deteriorated, so it is necessary to be 1.60% or less. In addition, the minimum with preferable Mn content is 1.4%, and a preferable upper limit is 1.5%.

[Al: 0.015 to 0.05%]
Al is an element that is useful as deoxidation and contributes to the refinement of the base material structure by forming nitride (AlN). In order to exhibit such an effect, Al needs to be contained 0.015% or more. However, if the Al content is excessive, the toughness of the steel sheet is coarsened, so it is necessary to make it 0.05% or less. In addition, the minimum with preferable Al content is 0.020%, and a preferable upper limit is 0.04%.

[Cu: 0.15-0.35%]
Cu is an element effective for miniaturization of austenite crystal grains. In order to exhibit such an effect, it is necessary to contain 0.15% or more of Cu. However, if the Cu content is excessive, the weldability of the base material is deteriorated, so it is necessary to make it 0.35% or less. In addition, since it will become easy to generate | occur | produce a hot crack when Cu is added independently, it is necessary to contain the following Ni simultaneously and to prevent a hot crack.

[Ni: 0.10 to 0.30%]
Ni is an element effective for improving low-temperature toughness. In order to exert such effects, Ni needs to be contained by 0.10% or more. However, if the Ni content is excessive, the cost is increased, so it is necessary to set it to 0.30% or less. In addition, the minimum with preferable Ni content is 0.15%, and a preferable upper limit is 0.25%.

[Mo: 0.10 to 0.25%]
Mo is an element effective for precipitating carbonitrides and increasing the strength of the steel sheet. In order to exhibit such an effect, it is necessary to contain 0.10% or more. However, if the Mo content is excessive, weldability and base metal toughness deteriorate, so it is necessary to set the content to 0.25% or less.

[V: 0.030 to 0.05% and Nb: 0.005 to 0.015%]
V and Nb are elements effective for refining austenite grains during rolling and suppressing recrystallization due to the formation of carbonitride, and for refining the structure after transformation (for example, ferrite structure). In order to exert such an effect, it is necessary to contain 0.030% or more in V and 0.005% or more in Nb. However, if these contents are excessive, the weldability of the steel sheet is hindered, so it is necessary to set V to 0.05% or less and Nb to 0.015% or less.

[Ca: 0.005% or less (excluding 0%)]
Ca is an element effective for improving the base material toughness. Such an effect increases as the content thereof increases. However, even if Ca is excessively contained, the effect is saturated, so the Ca content is preferably 0.005% or less. In addition, in order to exhibit said effect effectively, it is more preferable to contain 0.0005% or more with Ca.

[N: 0.002 to 0.008%]
N is an element that exhibits the effect of forming nitrides with elements such as Al, Nb, Ti, Nb, and V contained in steel and making the base material structure finer. In order to exhibit such an effect, N needs to be contained by 0.002% or more. However, if the N content is excessive, the amount of solute N is increased, and particularly the toughness of the welded portion is deteriorated.

  The basic components in the steel sheet of the present invention are as described above, and the balance is composed of iron and inevitable impurities (for example, P, S, B, O, etc.), but if necessary, (a) Ti: 0.005 It is also effective to contain ~ 0.020%, (b) Cr: 0.10% or less (excluding 0%), etc., and the characteristics are further improved depending on the type of element contained. . The reasons for limiting the range when these elements are contained are as follows.

[Ti: 0.005 to 0.020%]
Ti is an element effective for finely dispersing TiN in steel to prevent coarsening of austenite grains. In order to exhibit such an effect, it is preferable to contain Ti 0.005% or more. However, if the Ti content is excessive, the toughness of the base material is lowered, so that the content is preferably 0.020% or less.

[Cr: 0.10% or less (excluding 0%)]
Cr is an effective element for precipitating carbonitride and increasing the strength of the steel sheet. Such an effect increases as the content increases. However, if the content is excessive, the weldability and the base metal toughness deteriorate. Therefore, the content is preferably 0.10% or less. In addition, the more preferable minimum for exhibiting the said effect by Cr is 0.03%.

In producing the steel sheet of the present invention, the steel slab is heated to a temperature of not less than the Ac ₃ transformation point to 1200 ° C., and the cumulative rolling reduction is not less than 10% in the austenite recrystallization temperature range where the average temperature of the steel sheet is 900 ° C. or more. After that, in the non-recrystallization temperature range where the average temperature of the steel sheet is 800 ° C. or higher and 890 ° C. or lower, cooling is performed such that the average cooling rate between passes of the entire steel sheet is 0.3 ° C./second or higher. From the temperature range in which the cumulative rolling reduction is 25% or more and less than 50%, and the average temperature of the steel sheet is (Ar ₃ transformation point + 10 ° C.) or more and (Ar ₃ transformation point + 90 ° C.) or less, The steel sheet may be cooled to a temperature range where the steel sheet surface temperature is 500 ° C. or less at an average cooling rate of 5 ° C./second or more, and tempered in a temperature range of 500 ° C. or more and less than the Ac ₁ transformation. Hereinafter, these conditions will be described in order.

In order to set the heating temperature of the steel slab to austenite, it is necessary to set the Ac ₃ transformation point. 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 heating temperature is preferably 1200 ° C. or lower.

  The reduction ratio is not particularly limited in the recrystallization temperature range of austenite where the average temperature of the target part [t / 4 part (t: plate thickness)] is 900 ° C. or higher, but the viewpoint of productivity and the upper limit of the finish rolling reduction ratio Therefore, it is preferable that the cumulative rolling reduction is 10% or more. The cumulative rolling reduction at this time is a value obtained by the following equation (1).

Cumulative rolling reduction = [(t ₀ −t ₁ ) / t ₀ ] × 100 (%)
However, t ₀ : rolling start thickness (mm) when the steel sheet average temperature is in the target temperature range
t ₁ : Finishing thickness (mm) when the average steel sheet temperature is in the target temperature range

  Random grain boundaries can be introduced / increased in rolling in a non-recrystallization temperature region by rolling at a relatively high temperature of 800 ° C. or higher and 890 ° C. or lower in average temperature of the target part. In order to obtain such an effect, it is necessary to reduce the cumulative rolling rate by less than 50% (see FIG. 4 below). If the reduction ratio at this time is 50% or more, selection of variants due to excessive strain occurs and random grain boundaries decrease.

  In rolling in the high temperature side non-recrystallization temperature range where the average temperature of the target part is 800 ° C. or more and 890 ° C. or less, the average large angle after transformation is achieved even in the high temperature side non-recrystallization temperature range by carrying out active cooling. The particle diameter D can be refined. In order to exert such an effect, it is necessary to perform cooling so that the reduction at that time is 25% or more in terms of cumulative reduction, and the average cooling rate between passes of the entire steel sheet is 0.3 ° C./second or more. (See FIG. 5 below).

  In rolling in the non-recrystallization temperature range, since the deformation resistance is high, the steel sheet temperature rises due to heat generated by plastic deformation, γ grains grow before transformation, and there is a concern that the average large-angle grain size D after transformation increases. . Therefore, even if the cumulative rolling reduction is 25% or more, when the inter-pass cooling rate is less than 0.3 ° C./second, the steel plate temperature rises due to the plastic deformation heat generation, and the γ grains grow before transformation. Cannot be suppressed (see FIG. 6 below). In addition, when the inter-pass cooling rate (cooling rate during rolling) exceeds 0.4 ° C./second, an increase in processing strain due to a rapid decrease in temperature, and rolling may not be completed at a temperature of 800 ° C. or higher. Since the random grain boundary fraction is reduced, the cooling rate is preferably about 0.3 to 0.4 ° C./second (see Test No. 14 described later).

  However, if the rolling temperature is higher than 890 ° C., even if the above-described positive cooling is performed, the growth coarsening of γ grains before transformation cannot be suppressed, and the average large-angle particle diameter D may increase (described later). FIG. 9) discloses rolling at 890 ° C. or lower. When the rolling end temperature is lower than 800 ° C., variants are selected by low temperature rolling, and random grain boundaries are reduced (see FIG. 10 described later).

After finishing rolling, the average cooling rate (DQ cooling rate) of the steel plate is 5 ° C./second or more from the temperature range where the average temperature of the steel plate is (Ar ₃ transformation point + 10 ° C.) or more and (Ar ₃ transformation point + 90 ° C.) or less. To 500 ° C. or lower. When the temperature at the start of cooling exceeds (Ar ₃ transformation point + 90 ° C.), the strain introduced by rolling disappears during the waiting time from the end of rolling to the start of cooling, and the average large-angle grain boundary diameter D may increase. Therefore, it is necessary to start cooling from a temperature not higher than (Ar ₃ transformation point + 90 ° C.). By starting the cooling from this temperature range, it is possible to ensure the hardenability necessary to obtain an increase in random grain boundaries. However, when the cooling start temperature is lower than (Ar ₃ transformation point + 10 ° C.), it is difficult to increase the random grain boundary fraction due to insufficient hardenability (see FIG. 8 below). In addition, the cooling stop temperature is set to 500 ° C. or lower in order to completely complete the transformation.

After the positive cooling (accelerated cooling) as described above, tempering is performed in a temperature range of 500 ° C. or higher and lower than the Ac ₁ transformation. This treatment is for eliminating the island-like martensite phase (MA phase) generated in the structure. If such a phase remains, it may become a starting point of destruction.

The steel sheet of the present invention is manufactured by the manufacturing method as described above, and satisfies the above-described requirements. In such a steel sheet, the yield point YP is 480 MPa or more, the tensile strength TS is 590 MPa or more, and a strain of 10% is obtained. After the application, the characteristics are such that the average impact absorption energy vE- ₄₀ at −40 ° C. when subjected to an aging treatment at 250 ° C. for 1 hour is 100 J or more. The high-strength steel sheet of the present invention assumes a case where the steel sheet has a thickness of 60 mm or more, and the effect is particularly remarkable in the case of such a thick steel sheet.

  EXAMPLES Hereinafter, the present invention will be described more specifically with reference to examples. However, the present invention is not limited by the following examples, but may be appropriately modified within a range that can meet the purpose described above and below. Of course, it is possible to implement them, and they are all included in the technical scope of the present invention.

Using steel slabs of each steel type (steel types A to D) having the chemical composition shown in Table 1 below, under the manufacturing conditions shown in Table 2 below, [plate thickness (product thickness), slab heating temperature, recrystallization temperature range reduction rate (Cumulative rolling reduction), steel sheet cooling rate, rolling start temperature, non-recrystallization temperature range rolling reduction (cumulative rolling reduction), interpass cooling rate, rolling end temperature, post rolling rolling start temperature, post rolling rolling cooling rate (DQ cooling rate) ), Cooling stop temperature, tempering temperature]. Regarding the temperature at this time, the slab heating temperature, the cooling stop temperature, and the tempering temperature are controlled by the surface temperature of the steel sheet, and the others are controlled by the average temperature. The detailed temperature management procedure is as follows. It is as follows. Further, the Ar ₃ transformation point and Ac ₁ transformation point shown in Table 1 are calculated by the following equations (2) and (3), respectively.

[Temperature measurement method during rolling]
1. Using the process computer, the heating temperature of the steel slab is calculated based on the ambient temperature from the start of heating to the end of heating and the in-furnace time.
2. Using the calculated heating temperature, based on the rolling pass schedule during rolling and the data of the cooling method (water cooling or air cooling) between passes, a method suitable for calculation such as the difference method is used to calculate the rolling temperature at any position in the plate thickness direction. Rolling is carried out while calculating using this.
3. The surface temperature of the steel sheet is measured using a radiation type thermometer installed on the rolling line. However, the theoretical value is also calculated in the process computer.
4). The surface temperature of the steel sheet measured at the start of rough rolling, at the end of rough rolling, and at the start of finish rolling is collated with a calculated temperature calculated from a process computer.
5. If the difference between the calculated temperature and the measured temperature is ± 30 ° C or more, recalculate the calculated temperature so that it matches the measured temperature to obtain the calculated temperature on the process computer. If the calculated temperature is less than ± 30 ° C, calculate from the process computer. The calculated temperature is used as it is.
6). Using the calculated temperature calculated above, the rolling temperature in the region to be controlled is managed.

Ar ₃ transformation point (° C.) = 910−310 × [C] −80 × [Mn] −20 × [Cu] −15 × [Cr] −55 × [Ni] −80 × [Mo] +0.35 (t ₂ -8) ... (2)
However, [C], [Mn], [Cu], [Cr], [Ni] and [Mo] indicate the contents (mass%) of C, Mn, Cu, Cr, Ni and Mo, respectively, and t ₂ indicates the plate thickness (finished product thickness: mm).

Ac ₁ transformation point (° C.) = 723-14 × [Mn] + 22 × [Si] −14.4 × [Ni] + 23.3 × [Cr] (3)
However, [Mn], [Si], [Ni], and [Cr] indicate the contents (mass%) of Mn, Si, Ni, and Cr, respectively.

  For each steel plate obtained, the average large-angle grain size D (average diameter D of crystal grains surrounded by large-angle grain boundaries with a crystal orientation difference of 15 ° or more) and crystal orientation difference distribution (random grain boundary fraction R) are shown below. The base material tensile properties, base material impact properties, and strain aging properties were measured by the following methods. These results are collectively shown in Table 3 below.

[Measurement Method of Average Large Angle Particle Size D]
(A) A sample including the front and back surfaces of the plate thickness cut in parallel with the rolling direction of the steel plate was prepared.
(B) A cross-section is polished using a wet emery polishing paper of # 150 to # 1000 or a polishing method having a function equivalent to that, and mirror-finished using an abrasive such as diamond slurry.
(C) In the above cross section, using an EBSP apparatus (trade name: “OIM”) manufactured by Tex SEM Laboratories, in a thickness direction t / 4 (t: thickness), measurement area: 200 × 200 (μm), Measurement pitch: Measured at intervals of 0.5 μm, and a large-angle grain boundary was measured using a boundary having a crystal orientation difference of 15 ° or more as a grain boundary. At this time, measurement points having a confidence index indicating the reliability of the measurement direction smaller than 0.1 were excluded from the analysis target.
(D) As an analysis method of the text data, those having a crystal grain size of 2.5 μm or less are judged to be measurement noise and deleted, and the average grain diameter on the observation surface is calculated (equivalent circle diameter). D.

[Method of measuring crystal orientation difference distribution (random grain boundary fraction R)]
(A) A sample including the front and back surfaces of the plate thickness cut in parallel with the rolling direction of the steel plate was prepared.
(B) A cross-section is polished using a wet emery polishing paper of # 150 to # 1000 or a polishing method having a function equivalent to that, and mirror-finished using an abrasive such as diamond slurry.
(C) In the above cross section, using an EBSP apparatus (trade name: “OIM”) manufactured by Tex SEM Laboratories, in a thickness direction t / 4 (t: thickness), measurement area: 200 × 200 (μm), Measurement pitch: measured at intervals of 0.5 μm. At this time, measurement points having a confidence index indicating the reliability of the measurement direction smaller than 0.1 were excluded from the analysis target.
(D) The crystal orientation difference of less than 5.5 ° was determined as measurement noise, and the distribution at each orientation difference up to the crystal orientation difference of 62.5 ° was obtained.
(E) The random grain boundary fraction R was calculated by associating the crystal orientation difference of (d) with the corresponding map. 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 [random grain boundaries other than corresponding grain boundaries (> Σ49)] and random grain boundary fraction R (average value) were measured.

[Evaluation of base material tensile properties]
From the t / 4 (t: plate thickness) part, a U4 test piece determined by the NK (Japan Maritime Association) classification was collected, and a tensile test was performed according to JIS Z2241. The judgment criteria were a yield point YP: 480 MPa or more and a tensile strength TS: 590 MPa or more.

[Evaluation of impact characteristics of base material]
From the t / 4 (t: thickness) part, a U4 test piece defined by the NK (Japan Maritime Association) classification was collected and a V-notch Charpy test was performed (test method based on JIS Z 2242). In shipbuilding E grade in the NK (Nippon Kaiji Kyokai) class, the impact properties of the base material are determined at a test temperature of −40 ° C., so the average absorbed energy (vE ₋₆₀ ) at the test temperature of −60 ° C. was measured. And it evaluated that the thing of the value of vE- ₆₀ was 100J or more was excellent in toughness.

[Strain aging characteristics of base material]
Strain aging was imparted by the method prescribed in the NK (Japan Maritime Association) classification. Specifically, after 10% strain was applied to the tensile test piece, an aging treatment was performed at 250 ° C. for 1 hour. The shape of the test piece (TP) at this time is shown in FIG. Thereafter, a V-notch Charpy test piece was collected from a location where the amount of strain applied was in the range of 9.6 to 10.4%, and the test was conducted according to JIS Z 2242. Test temperature: average absorbed energy at −40 ° C. ( vE- ₄₀ ) was measured. And the thing whose average value of vE- ₄₀ is 100J or more was evaluated as being excellent in strain aging characteristics. Note that the amount of strain imparted at this time is determined by the following equation (4).

Strain imparting amount = (L−L ₀ ) / L ₀ × 100 (%) (4)
However, L ₀ : Gage distance before giving strain (mm)
L: Gage distance after distortion (mm)

  From these results, it can be considered as follows. Test No. 1 to 5, 15, 16, 18, 19, and 20 satisfy the requirements defined in the present invention, and have excellent strain aging characteristics while ensuring the tensile characteristics and impact characteristics of the base material. It can be seen that a certain high-strength thick steel plate can be realized.

  In contrast, test no. Those of 6 to 14 and 17 lack any of the requirements defined in the present invention, and at least one of the characteristics is deteriorated. Specifically, Test No. 6 to 9 are those in which the rolling reduction in the non-recrystallization temperature region deviates from the range defined by the present invention (25% or more and less than 50%), and the increase in average large-angle particle diameter D or random grain boundary content. The rate R is insufficient, and at least one of the base material toughness and strain aging characteristics is deteriorated.

Test No. Nos. 10 and 11 are those in which the cooling start temperature after rolling is outside the range defined by the present invention [(Ar ₃ transformation point + 10 ° C.) or more and (Ar ₃ transformation point + 90 ° C.) or less], and the average large-angle grains An increase in the diameter D or an insufficient random grain boundary fraction occurs, and at least one of the base material toughness and strain aging characteristics is deteriorated.

  Test No. Nos. 12 and 13 are those in which the rolling start temperature and the rolling end temperature deviate from the ranges specified in the present invention (800 ° C. or more and 890 ° C. or less), respectively, and an increase in average large-angle particle size D or random grain boundary fraction. Insufficiency has occurred, and at least one of the base material toughness and strain aging characteristics has deteriorated.

  Test No. Nos. 14 and 17 are those in which the inter-pass cooling rate (cooling rate during rolling) deviates from the preferred range (0.3 to 0.4 ° C./second) defined in the present invention, and the average large-angle particle size D is increased. Alternatively, the random grain boundary fraction is insufficient, and at least one of the base material toughness and strain aging characteristics is deteriorated.

Based on these results, the relationship between the average large-angle particle diameter D and the strain aging characteristics (10% strain imparted toughness vE _-40 ) is shown in FIG. 2. In order to achieve the strain aging characteristics, the average large-angle particle diameter D It can be seen that it is useful to set the thickness to 35 μm or less. FIG. 3 shows the relationship between the random grain boundary fraction R and strain aging characteristics (10% strain imparted toughness vE- ₄₀ ). In order to achieve the strain aging characteristics, the random grain boundary fraction R is 50 areas. It turns out that it is useful to set it as% or more.

  FIG. 4 shows the relationship between the reduction rate in the non-recrystallization temperature range and the random grain boundary fraction, and FIG. 5 shows the relationship between the reduction rate in the non-recrystallization temperature range and the average large-angle grain size D. It can be seen that by setting the rolling reduction in the temperature range to 25% or more and less than 50%, the random grain boundary fraction R is ensured and the average large-angle particle diameter D is refined. FIG. 6 shows the relationship between the cooling rate between passes (cooling rate during rolling) and the average large-angle particle size D. By making the cooling rate during rolling 0.3 ° C./second or more, the average large-angle particle size D is refined. It can be seen that is achieved.

FIGS. 7 and 8 show the influence of the cooling start temperature after rolling (cooling start temperature: position from the Ar ₃ point) on the average large-angle grain size D and the random grain boundary fraction R, respectively (however, the non-recrystallization temperature) When the rolling reduction temperature in the region is 25 to 30%) and the cooling start temperature is (Ar ₃ transformation point + 10 ° C.) or more and (Ar ₃ transformation point + 90 ° C.) or less, the average large angle grain boundary diameter D is fine. It can be seen that the formation of a random grain boundary fraction R is achieved.

  FIG. 9 shows the relationship between the rolling start temperature and the average large-angle grain size D, and FIG. 10 shows the relationship between the rolling end temperature and the random grain boundary fraction R (however, the unrecrystallization temperature range reduction ratio: 25-30%). thing). As apparent from these results, the average large-angle particle diameter D can be made 35 μm or less by setting the rolling start temperature to 890 ° C. or less, and the random grain boundary fraction R is made to be by rolling the end temperature of 800 ° C. or more. It turns out that it can be 50 area% or more.

It is explanatory drawing which shows the shape of the tension test piece which provided the distortion. It is a graph which shows the relationship between average large angle particle diameter D and a strain aging characteristic (10% strain imparted toughness vE- ₄₀ ). It is a graph which shows the relationship between the random grain boundary fraction R and a strain aging characteristic (10% strain imparted toughness vE- ₄₀ ). It is a graph which shows the relationship between the unrecrystallized temperature range reduction rate and the random grain boundary fraction R. 4 is a graph showing the relationship between the unrecrystallized temperature range reduction rate and the average large-angle particle size D. It is a graph which shows the relationship between the cooling rate during rolling, and the average large angle particle size D. After rolling cooling start temperature (cooling start temperature position from Ar ₃ point) is a graph showing the effect on the average high angle grain size D. Rolling After cooling start temperature (cooling start temperature position from Ar ₃ point) is a graph showing the effect of random grain boundaries fraction R. It is a graph which shows the relationship between rolling start temperature and average large angle particle size D. It is a graph which shows the relationship between rolling completion temperature and the random grain boundary fraction R.

Claims (3)

C: 0.10 to 0.16% (meaning “mass%”, the same applies to the chemical composition), Si: 0.15 to 0.30%, Mn: 1.30 to 1.60%, Al: 0.015~0.05%, Cu: 0.15~0.35%, Ni: 0.10~0.30%, Mo: 0.10~0.25%, V: 0.030~0. 05%, Nb: 0.005 to 0.015%, Ca: 0.005% or less (excluding 0%) , N: 0.002 to 0.008% , Ti: 0.005 to 0.020 % And / or Cr: steel plates each containing 0.10% or less (excluding 0%) , the balance being iron and unavoidable impurities, at a large angle grain boundary where the orientation difference between the two crystals is 15 ° or more The average equivalent circle diameter D of the enclosed crystal grains is 35 μm or less, and the random grain boundary fraction measured from the crystal orientation distribution difference High strength thick steel plate but excellent strain aging characteristic, characterized in that it is 50 area% or more. The yield point is 480 MPa or more, the tensile strength is 590 MPa or more, and the average impact absorption energy vE- ₄₀ at −40 ° C. is 100 J or more when aging treatment is performed at 250 ° C. for 1 hour after applying 10% strain. The high-strength thick steel plate according to claim 1 . In producing the steel sheet according to claim 1 or 2 , the steel slab is heated to a temperature of not less than the Ac ₃ transformation point to 1200 ° C, and the cumulative reduction is performed in an austenite recrystallization temperature range where the average temperature of the steel sheet is 900 ° C or more. The steel sheet is subjected to rolling at a rate of 10% or more, and then the average cooling rate between the passes of the entire steel sheet is 0.3 ° C./second or more in the non-recrystallization temperature range where the average temperature of the steel sheet is 800 ° C. or more and 890 ° C. or less. The steel sheet is rolled so that the cumulative reduction ratio is 25% or more and less than 50% while cooling is performed, and the average temperature of the steel sheet is subsequently (Ar ₃ transformation point + 10 ° C.) or more, (Ar ₃ transformation point + 90 ° C.) The average cooling rate from the following temperature range to the temperature range where the steel sheet surface temperature is 500 ° C. or lower is cooled at a cooling rate of 5 ° C./second or more, and tempering is performed at a temperature range of 500 ° C. or higher and less than the Ac ₁ transformation point. Strain aging characteristics characterized by A method for producing excellent high-strength thick steel plates.

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