EP2799584A1 - Hochfeste dicke stahlplatte zur konstruktion mit hervorragenden magnetischen eigenschaften zur verhinderung der diffusion von spröden rissen und herstellungsverfahren dafür - Google Patents

Hochfeste dicke stahlplatte zur konstruktion mit hervorragenden magnetischen eigenschaften zur verhinderung der diffusion von spröden rissen und herstellungsverfahren dafür Download PDF

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EP2799584A1
EP2799584A1 EP12863408.6A EP12863408A EP2799584A1 EP 2799584 A1 EP2799584 A1 EP 2799584A1 EP 12863408 A EP12863408 A EP 12863408A EP 2799584 A1 EP2799584 A1 EP 2799584A1
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Prior art keywords
steel plate
less
temperature
central portion
thickness direction
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French (fr)
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EP2799584B1 (de
EP2799584A4 (de
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Yoshiko TAKEUCHI
Kazukuni Hase
Shinji Mitao
Yoshiaki Murakami
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JFE Steel Corp
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JFE Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2201/00—Treatment for obtaining particular effects
    • C21D2201/05—Grain orientation
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite

Definitions

  • the present invention relates to a high-strength thick steel plate for structural use excellent in terms of brittle crack arrestability and a method for manufacturing the steel plate, and in particular to a steel plate having a thickness of 50 mm or more which can be preferably used for ships.
  • Ni content As a method for improving the brittle crack arrestability of a steel material, a method in which Ni content is increased has been known in the past. 9%-Ni steel is commercially used for the storage tanks of liquefied natural gases.
  • Patent Literature 1 proposes a steel material having an ultra fine crystallization microstructure in the surface portion in order to improve brittle crack arrestability without an increase in alloy cost.
  • the steel material having excellent brittle crack arrestability according to Patent Literature 1 is characterized in that, focusing on the fact that shear-lips (plastic deformation areas), which are formed in the surface portion of a steel material when a brittle crack propagates, are effective for improving brittle crack arrestability, the crystal grain size in the shear-lip portions is decreased in order to absorb the propagation energy of a propagating brittle crack.
  • an ultra fine ferrite structure or bainite structure is formed in the surface portion of the steel material by repeating once or more a process, in which the surface portion of a hot-rolled steel plate is cooled down to a temperature equal to or lower than the A r3 transformation point by performing controlled cooling and then the controlled cooling is stopped in order to allow the surface portion to recuperate to have a temperature equal to or higher than the transformation point, while the steel material is rolled in order for transformation or recrystallization due to deformation to repeatedly occur.
  • Patent Literature 2 it is disclosed that, in order to improve the brittle crack arrestability of a steel material having a microstructure mainly including a ferrite-pearlite phase, it is important to form a layer, in either of the surface portions of the steel material, including 50% or more of a ferrite structure having ferrite grains with a circle-equivalent average grain size of 5 ⁇ m or less and an aspect ratio of the grains of 2 or more, and to prevent the variation of a ferrite grain size, and that, as a method for preventing the variation, the maximum rolling reduction per pass of finishing rolling is controlled to be 12% or less in order to prevent local recrystallization.
  • Patent Literature 3 discloses a technique which is a modification of TMCP and in which, focusing not only on a decrease in ferrite crystal grain size but also on a subgrain formed in a ferrite grain, brittle crack arrestability is improved.
  • brittle crack arrestability is improved by controlling (a) rolling conditions such that fine ferrite crystal grains are achieved, (b) rolling conditions such that a fine ferrite structure is formed in a portion constituting 5% or more of the thickness of the steel material, (c) rolling conditions such that subgrains are formed by growing a texture in the fine ferrite structure and by rearranging dislocations introduced by applying deformation (rolling) using thermal energy and (d) cooling conditions such that an increase in the grain size of the formed fine ferrite crystal grains and an increase in the grain size of the formed fine subgrains are prevented.
  • a method in which brittle crack arrestability is improved by applying reduction forth of rolling to a transformed ferrite phase in order to grow a texture, is also known.
  • resistance to brittle fracture is increased by forming a separation parallel to the plate surface on the fracture surface of a steel material in order to reduce stress at the brittle crack tip.
  • Patent Literature 4 discloses that brittle fracture resistance is improved by performing controlled rolling in order to form a microstructure having an X-ray plane intensity ratio in the (110) plane showing a texture developing degree of 2 or more and including large-size grains having a diameter equivalent to a circle in the crystal grains of 20 ⁇ m or more in an amount of 10% or less.
  • Patent Literature 5 discloses, as a steel for welded structural use having excellent brittle crack arrestability in the joint part, a steel plate having an X-ray plane intensity ratio in the (100) plane showing a texture developing degree on a plane inside the plate parallel to the rolling surface of the plate of 1.5 or more. It is disclosed that the steel plate has excellent brittle crack arrestability owing to the difference in angle between the direction of applied stress and the direction of crack propagation as a result of the growth of the texture mentioned above.
  • Non Patent Literature 1 it is reported that, from the results of the evaluation of the brittle crack arrestability of a steel plate having a thickness of 65 mm, a brittle crack was not arrested in a large brittle crack arrestability test on a base metal.
  • the steel plates having excellent brittle crack arrestability according to Patent Literatures 1 through 5 described above are mainly intended for a steel plate having a thickness of about 50 mm or less as indicated by the manufacturing conditions and the disclosed experimental data, and therefore it is not clear whether specified properties can be obtained in the case where the disclosed techniques are applied to a steel plate having a thickness of more than 50 mm, and the properties regarding crack propagation in the thickness direction which are required for ship's hull structures have never been tested at all.
  • an object of the present invention is to provide a high-strength thick steel plate excellent in terms of brittle crack arrestability which can be stably manufactured using a very simple industrial process in which rolling conditions are optimized in order to control a texture in the thickness direction and a method for manufacturing the steel plate.
  • the present inventors diligently conducted investigations in order to solve the problems described above and found the following knowledge regarding a high-strength thick steel plate having excellent crack arrestability despite the steel plate having a heavy thickness.
  • the present invention is as follows.
  • a high-strength thick steel plate having a thickness of 50 mm or more excellent in terms of brittle crack arrestability, in which a texture is appropriately controlled in the thickness direction, and a method for manufacturing the steel plate can be provided, and it is effective to apply the present invention to a steel plate having a thickness of preferably more than 50 mm, more preferably 55 mm or more.
  • the present invention contributes to the improvement of the safety of ships by being applied to hutch side coamings and deck part materials in high-strength deck structures of large container carriers and bulk carriers, which results in a large advantage in industry.
  • Fig. 1 is a schematic diagram illustrating the fracture surface shape of an ESSO test compliant with WES 3003 of a thick steel plate having a thickness of more than 50 mm, where (a) is a diagram illustrating a plane view of a test piece and (b) is a diagram illustrating the fracture surface of the test piece.
  • toughness and the integration degree I of the RD//(110) plane in the central portion in the thickness direction are appropriately specified in accordance with desired brittle crack arrestability.
  • a Charpy fracture appearance transition temperature in the surface portion and the central portion in the thickness direction be -40°C or lower. It is preferable that the Charpy fracture appearance transition temperature in the central portion in the thickness direction be -50°C or lower.
  • cleavage planes are integrated diagonally to the main direction of a crack so as to form fine branched cracks, which brings a stress relaxation effect at a brittle crack tip and results in an increase in brittle crack arrestability.
  • the integration degree I of the RD//(110) plane in the central portion in the thickness direction be 1.5 or more, preferably 1.7 or more.
  • the integration degree I of the RD//(110) plane in the central portion in the thickness direction is defined in the following way. Firstly, by performing mechanical polishing and electrolytic polishing on a surface, being parallel to the steel plate surface, of a sample having a thickness of 1 mm cut out of the central portion in the thickness direction, a test piece for X-ray diffractometry is prepared. By performing X-ray diffraction measurement using a Mo X-ray source on this test piece, the pole figures of (200), (110) and (211) planes are obtained, and then three dimensional orientation distribution function is calculated from the obtained pole figures by using a Bunge method.
  • the Charpy fracture appearance transition temperature and the integration degree I of the RD//(110) plane in the central portion in the thickness direction satisfy the relational expression (1) below.
  • the relational expression (1) it is possible to achieve better brittle crack arrestability.
  • vTrs (1/2t) fracture appearance transition temperature (°C) in the central portion in the thickness direction
  • I RD//(110) [1/2t] integration degree of the RD//(110) plane in the central portion in the thickness direction
  • t thickness (mm).
  • a metallographic structure mainly includes a bainite phase.
  • a metallographic structure mainly includes a bainite phase means that the area fraction of a bainite phase is 80% or more with respect to the whole metallographic structure.
  • the area fraction of the remainder consisting of, for example, a ferrite phase, a martensite phase (including martensite islands) and a pearlite phase is 20% or less.
  • the microstructure formed in the rolling performed in the austenite non-recrystallization temperature range is transformed into a bainite phase
  • a texture having a specified orientation is preferentially formed, that is, so-called variant selection occurs, which results in the integration degree I of the RD//(110) plane becoming 1.5 or more, preferably 1.7 or more. Therefore, the metallographic structure, which is obtained after rolling and cooling have been performed, mainly includes a bainite phase.
  • C is a chemical element which increases the strength of steel and it is necessary that the C content be 0.03% or more in order to achieve the desired strength in the present invention
  • the C content in the case where the C content is more than 0.20%, there is not only a decrease in weldability but also a negative influence on toughness. Therefore, it is preferable that the C content be in the range of 0.03% to 0.20%, more preferably 0.05% to 0.15%.
  • Si is effective as a deoxidizing chemical element and as a chemical element for increasing the strength of steel, the effect cannot be realized in the case where the Si content is less than 0.03%.
  • the Si content is more than 0.5%, there is not only the deterioration of the surface quality of steel but also a significant decrease in toughness. Therefore, it is preferable that the Si content be 0.03% or more and 0.5% or less.
  • Mn is added as a chemical element for increasing strength. Since the effect is insufficient in the case where the Mn content is less than 0.5%, and since there is a decrease in weldability and an increase in steel material cost in the case where the Mn content is more than 2.5%, it is preferable that the Mn content be 0.5% or more and 2.5% or less.
  • Al is effective as a deoxidizing agent, and it is necessary that the Al content be 0.005% or more in order to realize this effect, but, in the case where the Al content is more than 0.08%, there is not only a decrease in toughness but also a decrease in the toughness of a weld metal when welding is performed. Therefore, it is preferable that the Al content be in the range of 0.005% to 0.08%, more preferably 0.02% to 0.04%.
  • N increases the strength of steel by controlling a crystal grain size as a result of combining with Al in steel to form AlN when rolling is performed, but, since there is a decrease in toughness in the case where the N content is more than 0.0050%, it is preferable that the N content be 0.0050% or less.
  • the content of P and S be respectively 0.03% or less and 0.01% or less, more preferably 0.02% or less and 0.005% or less respectively.
  • a small content of Ti is effective for increasing the toughness of a base metal by decreasing a crystal grain size as a result of forming a nitride, carbide or carbonitride. This effect is realized in the case where the Ti content is 0.005% or more, but, since there is a decrease in the toughness of a base metal and a welded heat affected zone in the case where the Ti content is more than 0.03%, it is preferable that the Ti content be in the range of 0.005% to 0.03%.
  • the chemical composition described above is the base chemical composition in the present invention
  • one or more of Nb, Cu, Ni, Cr, Mo, V, B, Ca and REM may be added in order to further improve the properties.
  • Nb contributes to an increase in strength as a result of precipitating in the form of NbC when ferrite transformation occurs or reheating is performed.
  • Nb since Nb is effective for expanding a temperature range in which recrystallization does not occur when rolling is performed under conditions for forming an austenite phase, which results in a decrease in bainite packet grain size, Nb contributes to an increase in toughness. This effect is realized in the case where the Nb content is 0.005% or more, but, since there is conversely a decrease in toughness as a result of the precipitation of large-size NbC in the case where the Nb content is more than 0.05%, it is preferable that the upper limit of the Nb content be 0.05%.
  • Cu, Ni, Cr and Mo are all chemical elements which increase the hardenability of steel. Since these chemical elements directly contribute to an increase in strength after rolling has been performed and may be added in order to improve functional properties such as toughness, high temperature strength or weather resistance, and since these effects are realized in the case where the contents of these chemical elements are respectively 0.01% or more, it is preferable that the contents of these chemical elements be respectively 0.01% or more in the case where these chemical elements are added. However, since there is a decrease in toughness and weldability in the case where the contents of these chemical elements are excessively large, it is preferable that the upper limits of the contents of Cu, Ni, Cr and Mo be respectively 0.5%, 1.0%, 0.5% and 0.5% in the case where these chemical elements are added.
  • V 0.001% to 0.10%
  • V is a chemical element which increases the strength of steel by precipitation strengthening as a result of precipitating in the form of V(C,N).
  • the V may be contained in the amount of 0.001% or more in order to realize this effect, but there is a decrease in toughness in the case where the V content is more than 0.10%. Therefore, in the case where V is added, it is preferable that the V content be in the range of 0.001% to 0.10%.
  • a small amount of B may be added as a chemical element which increases the hardenability of steel.
  • the B content is more than 0.0030%, since there is a decrease in the toughness of a weld zone, it is preferable that the B content be 0.0030% or less in the case where B is added.
  • Ca and REM increase toughness as a result of decreasing a grain size in a microstructure in a welded heat affected zone and there is no decrease in the effect of the present invention even in the case where these chemical elements are added, these chemical elements may be added as needed.
  • the upper limit of the contents of Ca and REM be respectively 0.0050% and 0.010% in the case where these are added.
  • manufacturing conditions such as the heating temperature of a slab as a steel material, hot rolling conditions and cooling conditions be specified.
  • hot rolling it is preferable to specify, in addition to total cumulative rolling reduction in the austenite recrystallization temperature range and austenite non-recrystallization temperature range, cumulative rolling reduction for each of the cases where the central portion in the thickness direction has a temperature in the austenite recrystallization temperature range and where the central portion in the thickness direction has a temperature in the austenite non-recrystallization temperature range and to specify a temperature condition in rolling while the central portion in the thickness direction has a temperature in the austenite non-recrystallization temperature range.
  • molten steel having the chemical composition described above is produced using, for example, a converter furnace and made into a slab using, for example, a continuous casting method. Subsequently, the slab is heated at a temperature of 1000°C to 1200°C and then hot-rolled.
  • the heating temperature it is difficult to secure sufficient time for performing rolling in the austenite recrystallization temperature range in the case where the heating temperature is lower than 1000°C.
  • the heating temperature in the case where the heating temperature is higher than 1200°C, since there is not only a decrease in toughness due to an increase in austenite grain size but also a decrease in yield due to a significant loss caused by oxidation, it is preferable that the heating temperature be 1000°C to 1200°C, more preferably in the range of 1000°C to 1150°C from the viewpoint of toughness.
  • the cumulative rolling reduction is 20% or more.
  • the cumulative rolling reduction is less than 20%, since an austenite grain size does not become sufficiently small, there is not an increase in the toughness of the microstructure which is finally obtained.
  • the central portion in the thickness direction has a temperature in the austenite non-recrystallization temperature range, under the conditions that the cumulative rolling reduction is 40% or more.
  • the cumulative rolling reduction in this temperature range is 40% or more, since a texture in the central portion in the thickness direction can be sufficiently grown, it is possible to control the integration degree I of the RD//(110) plane in the central portion in the thickness direction to be 1.5 or more, preferably 1.7 or more.
  • the difference in rolling temperature between the first and the last passes in rolling performed while the central portion in the thickness direction has a temperature in the austenite non-recrystallization temperature range be 40°C or less.
  • rolling temperature means the temperature of the central portion in the thickness direction of a steel material immediately before the steel material is rolled.
  • the temperature of the central portion in the thickness direction can be derived from, for example, the thickness, the surface temperature and the cooling conditions by, for example, simulation calculation. For example, by calculating the temperature distribution in the thickness direction using a difference method, the temperature of the central portion in the thickness direction of the steel plate can be derived.
  • the total cumulative rolling reduction in the austenite recrystallization temperature range and in the austenite non-recrystallization temperature range be 65% or more. This is because sufficient reduction cannot be applied to a microstructure in the case where the total reduction is small, which results in the target toughness and strength not being achieved, and because, by controlling the total cumulative rolling reduction to be 65% or more, it is possible to apply sufficient reduction to a microstructure, which results in the target toughness and strength being achieved.
  • the austenite recrystallization temperature range and the austenite non-recrystallization temperature range are determined by performing preliminary experiments using steel having the chemical composition described above in which the steel is subjected to heating and processing history under various conditions.
  • finishing temperature of hot rolling it is preferable that the finishing temperature be in the austenite non-recrystallization temperature range from the view point of rolling efficiency.
  • the rolled steel plate be cooled down to a temperature of 450°C or lower at a cooling rate of 4°C/s or more.
  • a cooling rate of 4°C/s or more.
  • the cooling rate is less than 4°C/s, there is an increase in grain size in a microstructure and the progression of ferrite transformation throughout the thickness, which results not only in the desired microstructure not being achieved but also in a decrease in strength of the steel plate.
  • the cooling stop temperature By controlling the cooling stop temperature to be 450°C or lower, since bainite transformation can be sufficiently progressed, it is possible to achieve a metallographic structure having the desired toughness and texture. In the case where the cooling stop temperature is higher than 450°C, since bainite transformation is not sufficiently progressed, structures such as a ferrite phase and a pearlite phase are also formed, which results in a microstructure mainly including a bainite phase which is a target microstructure in the present invention not being achieved.
  • the cooling rate and cooling stop temperature described above are determined by using the temperature of the central portion in the thickness direction of the steel plate.
  • the temperature of the central portion in the thickness direction can be derived from, for example, the thickness, the surface temperature and the cooling conditions by, for example, simulation calculation. For example, by calculating the temperature distribution in the thickness direction using a difference method, the temperature of the central portion in the thickness direction of the steel plate can be derived.
  • a temper treatment may be performed on the cooled steel plate. By performing a temper treatment, it is possible to further increase the toughness of a steel plate. By controlling a tempering temperature to be equal to or lower than the A C1 point in terms of the average temperature of the steel plate, it is possible to prevent the desired microstructure obtained through rolling and cooling from being lost.
  • the A C1 point (°C) is derived using the equation below.
  • a C1 point 751 - 26.6C + 17.6Si - 11.6Mn -169Al - 23Cu - 23Ni + 24.1Cr + 22.5Mo + 233Nb - 39.7V - 5.7Ti - 895B, where an atomic symbol in the equation above represents the content (mass%) of a chemical element in steel represented by the symbol, and where the symbol is assigned a value of 0 in the case where the chemical element is not contained.
  • the average temperature of the steel plate can also be derived from, for example, the thickness, the surface temperature and the cooling conditions by, for example, simulation calculation, as is the case with the temperature of the central portion in the thickness direction.
  • Kca value at a temperature of -10°C was determined by performing an ESSO test compliant with WES 3003.
  • the integration degree I of the RD//(110) plane in the central portion in the thickness direction was derived in the following way. Firstly, by performing mechanical polishing and electrolytic polishing on the surface parallel to the steel plate surface of a sample having a thickness of 1 mm cut out of the central portion in the thickness direction, a test piece for X-ray diffractometry was prepared. By performing X-ray diffraction measurement using a Mo X-ray source on this test piece, the pole figures of (200), (110) and (211) planes were obtained. Three dimensional orientation distribution function was calculated from the obtained pole figures by using a Bunge method.
  • Sample steel plates (serial Nos. 1 through 13 and 27 through 30), which had the fracture appearance transition temperature and textures in the central portion in the thickness direction which were within the range according to the present invention, had a Kca(-10°C) of 6000 N/mm 3/2 or more, which means these sample steel plates had excellent brittle crack arrestability.
  • sample steel plates (serial Nos. 1 through 13), which had the Charpy fracture appearance transition temperature and the integration degrees I of the RD//(110) plane in the surface portion and the central portion in the thickness direction satisfying the relational expression (1), had higher Kca(-10°C) than sample steel plates (serial Nos. 27 through 30), which had the Charpy fracture appearance transition temperature and the integration degree I of the RD//(110) plane in the surface portion and the central portion in the thickness direction not satisfying the relational expression (1).
  • sample steel plates (serial Nos. 21 through 26), which had chemical compositions of the steel plate within the preferable ranges according to the present invention and were prepared under manufacturing conditions regarding heating and rolling conditions for the steel plate out of the preferable range according to the present invention, had a Kca(-10°C) of less than 6000 N/mm 3/2 .
  • the textures of sample steel plates did not satisfy the specifications according to the present invention.
  • Sample steel plates (serial Nos. 14 through 20), which had chemical compositions of the steel plate out of the preferable ranges according to the present invention, had toughness not satisfying the specifications according to the present invention and a Kca(-10°C) of less than 6000 N/mm 3/2 .
  • sample steel plates (serial Nos. 14 through 26), which had at least one of the toughness and texture in the central portion in the thickness direction out of the ranges according to the present invention, had a Kca(-10°C) of less than 6000 N/mm 3/2 .

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  • Physics & Mathematics (AREA)
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  • Heat Treatment Of Steel (AREA)
EP12863408.6A 2011-12-27 2012-05-18 Herstellungsverfahren für eine hochfeste dicke stahlplatte zur konstruktion mit hervorragenden magnetischen eigenschaften zur verhinderung der diffusion von spröden rissen Active EP2799584B1 (de)

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JP2011285570 2011-12-27
JP2012111158A JP5304925B2 (ja) 2011-12-27 2012-05-15 脆性亀裂伝播停止特性に優れた構造用高強度厚鋼板およびその製造方法
PCT/JP2012/063409 WO2013099318A1 (ja) 2011-12-27 2012-05-18 脆性亀裂伝播停止特性に優れた構造用高強度厚鋼板およびその製造方法

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CN105980588B (zh) * 2013-12-12 2018-04-27 杰富意钢铁株式会社 钢板及其制造方法
EP3239332B1 (de) * 2014-12-24 2019-11-20 Posco Hochfester stahl mit hervorragender sprödbruchstablität und herstellungsverfahren dafür
CN107109590A (zh) * 2014-12-24 2017-08-29 Posco公司 耐脆性裂纹扩展性优异的高强度钢材及其制造方法
KR101657827B1 (ko) * 2014-12-24 2016-09-20 주식회사 포스코 취성균열전파 저항성이 우수한 구조용 극후물 강재 및 그 제조방법
CN107109597B (zh) * 2014-12-24 2020-01-31 Posco公司 耐脆性裂纹扩展性优异的高强度钢材及其制造方法
WO2017047088A1 (ja) 2015-09-18 2017-03-23 Ｊｆｅスチール株式会社 構造用高強度厚鋼板およびその製造方法
CN108779525A (zh) * 2016-02-24 2018-11-09 杰富意钢铁株式会社 脆性裂纹传播停止特性优良的高强度极厚钢板及其制造方法
KR102193527B1 (ko) * 2016-08-09 2020-12-21 제이에프이 스틸 가부시키가이샤 고강도 후강판 및 그의 제조 방법
WO2018030186A1 (ja) * 2016-08-09 2018-02-15 Ｊｆｅスチール株式会社 高強度厚鋼板およびその製造方法
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RU2630086C1 (ru) * 2016-06-14 2017-09-05 Федеральное государственное бюджетное образовательное учреждение высшего образования "Калининградский государственный технический университет" Малокремнистая судостроительная сталь

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JP2013151732A (ja) 2013-08-08
BR112014015779B1 (pt) 2019-04-09
EP2799584A4 (de) 2015-01-07
CN104024462A (zh) 2014-09-03
KR20140097463A (ko) 2014-08-06
CN104024462B (zh) 2016-03-23
KR101588258B1 (ko) 2016-01-25
BR112014015779A8 (pt) 2017-07-04
JP5304925B2 (ja) 2013-10-02
BR112014015779A2 (pt) 2017-06-13
WO2013099318A1 (ja) 2013-07-04

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