High-strength steel sheet and method for manufacturing the same

Abstract

The present invention provides a high-strength steel sheet and a method for manufacturing the same. The high-strength steel sheet of the present invention has a specific composition, a metal structure of the high-strength steel sheet at a 1 / 4-thickness position is composed of a soft phase as a main phase and a hard phase as a balance, the soft phase is composed of ferrite, the hard phase contains one or two or more of pearlite, bainite and martensite, a ratio of processed ferrite in the entire metal structure is 50% or more in volume fraction, a length-diameter ratio of the processed ferrite is 1.5 or more, an average crystal grain size of the processed ferrite is 50 μm or less, a yield stress of a base material is 235 MPa or more, a Charpy impact absorbed energy of the base material at -60°C is 200 J or more, and a Charpy impact absorbed energy of a weld heat-affected zone after a large heat input welding at -60°C is 80 J or more.

Description

Technical Field

[0001] This invention relates to steel plates (thick steel plates) used in, for example, ships, marine structures, mid- to high-rise buildings, bridges, and storage tanks, and particularly to high-strength steel plates that exhibit high toughness even in the weld heat-affected zone after welding, and a method for manufacturing the same. Background Technology

[0002] In recent years, the stringent requirements for the material properties of welded steel (welded steel plates) used in structures such as ships, marine structures, mid- to high-rise buildings, bridges, and storage tanks have been increasing. Furthermore, for such structures, high-heat welding methods, such as submerged arc welding, electric arc welding, and electroslag welding, are preferred for rapid manufacturing. Therefore, similar to the inherent toughness of the steel, the stringent requirements for the toughness of the weld heat-affected zone (HAZ) are also increasing.

[0003] However, generally speaking, if the welding line energy increases, the microstructure of the HAZ becomes coarser, and the toughness of the HAZ decreases. Numerous countermeasures have been proposed to date to address this decrease in HAZ toughness caused by high line energy welding, such as those described in Patent Documents 1-5.

[0004] Patent Documents 1 and 2 disclose techniques for improving the toughness of the high heat input zone (HAZ) caused by high heat input welding (hereinafter sometimes referred to as "high heat input HAZ"). Specifically, they propose methods to suppress the coarsening of austenite grains by utilizing the pinning effect of TiN, Al oxides, etc.

[0005] Furthermore, Patent Documents 3, 4, and 5 disclose techniques for refining the microstructure within austenite grains by introducing a large number of ferrite phase transformation nuclei. Specifically, by utilizing TiN, MnS, Ti oxides, etc., as ferrite phase transformation nuclei, the microstructure within the grains is refined, thereby improving the low-temperature toughness of the HAZ.

[0006] Existing technical documents

[0007] Patent documents

[0008] Patent Document 1: Japanese Patent Application Publication No. 2002-256379

[0009] Patent Document 2: Japanese Patent No. 2950076

[0010] Patent Document 3: Japanese Patent Publication No. 07-068577

[0011] Patent Document 4: Japanese Patent Publication No. 05-017300

[0012] Patent Document 5: Japanese Patent No. 3733898 Summary of the Invention

[0013] The problem that the invention aims to solve

[0014] However, even when applying the techniques disclosed in Patent Documents 1 to 5 for miniaturizing the HAZ using the precipitates described above, coarsening of the HAZ structure is unavoidable when performing high heat input welding. For example, the low-temperature toughness of the HAZ will deteriorate in environments below -60°C.

[0015] In recent years, research has been conducted on applications in ships, storage tanks, and other fields at even lower temperatures than before. Therefore, there is a need for steel that offers a significant improvement in the low-temperature toughness of the weld heat-affected zone compared to the target steels (steel plates) described in the aforementioned patent documents.

[0016] The present invention was made in view of the above-mentioned actual situation, and in particular, aims to provide a base material for steel plates used in the above-mentioned applications and a high-strength steel plate with excellent low-temperature toughness in HAZ (high heat input zone), and a method for manufacturing the same.

[0017] Here, "high strength" in this invention refers to a yield stress (YP) of 235 MPa or higher, representing the strength of the base material of the high-strength steel plate. "Excellent low-temperature toughness of the base material" in this invention refers to an absorbed energy of 200 J or higher in the Charpy impact test at -60°C. "Excellent low-temperature toughness of the HAZ (high heat input HAZ)" in this invention refers to an absorbed energy of 80 J or higher in the Charpy impact test at -60°C for the HAZ after high heat input welding, i.e., the HAZ of a single-pass weld joint on one side. Especially for weld joints with a welding heat input of 4 kJ / mm or higher resulting from single-pass welding such as submerged arc welding, stable low-temperature toughness can be obtained.

[0018] It should be noted that the yield stress and the absorbed energy of the Charpy impact test can be determined by the methods described in the examples below.

[0019] Methods for solving problems

[0020] In order to solve the above problems, the inventors conducted in-depth research on methods for improving the low-temperature toughness of high line energy HAZ, and obtained the following insights.

[0021] First, the inventors focused on the coarse bainitic structure generated by high heat input welding (HAZ). Compared to structures like ferrite and pearlite, coarse bainite is a coarse microstructure. The critical stress for brittle fracture of a coarse microstructure is low, thus becoming a major cause of reduced toughness. Therefore, the inventors considered improving the low-temperature toughness of HAZ by suppressing the formation of coarse bainite.

[0022] Furthermore, the inventors' in-depth research has yielded the following insights. By designing a composition that satisfies the conditions of equation (1), the ferrite phase transformation is promoted, enabling the refinement of coarse bainite. Moreover, by designing the ratio of Ti to N (Ti / N) to be in the range of 1.5 to 4.0, the coarsening of the original austenite grain size caused by welding can be suppressed, thereby achieving the suppression of coarse bainite.

[0023] However, due to the condition of (1), the fraction of ferrite as a soft phase increases, making it difficult to ensure the strength of the base material (steel plate).

[0024] Therefore, the inventors studied the microstructure of the base material. The results showed that at the 1 / 4 position of the steel plate thickness, when the proportion of processed ferrite in the total microstructure of the base material is 50% or more by volume, the aspect ratio of the processed ferrite is 1.5 or more, and the average grain size of the processed ferrite is 50 μm or less, excellent base material strength can be obtained while satisfying the conditions of equation (1). 0.25≤[C]+[Mn] / 6+([Cu]+[Ni]) / 15+([Cr]+[Mo]+[V]) / 5≤0.35…(1)

[0025] Wherein, C, Mn, Cu, Ni, Cr, Mo, and V represent the content (mass%) of each element, and the content of elements that are not present is set to zero.

[0026] This invention is based on the above insights, and its main points are as follows.

[0027] [1] A high-strength steel plate having the following composition by mass%: C: 0.010-0.070%, Si: 0.01-0.50%, Mn: 1.00-2.00%, P: less than 0.020%, S: 0.0005-0.0100%, Al: 0.035-0.100%, Ti: 0.010-0.030%, and N: 0.0035-0.0100%, wherein the carbon equivalent Ceq.(IIW) shown in the following formula (1) satisfies 0.25-0.35 by mass%, the Ti / N ratio satisfies 1.5-4.0, and the balance is composed of Fe and unavoidable impurities.

[0028] The microstructure of the high-strength steel plate at 1 / 4 of its thickness consists of a soft phase as the main phase and a hard phase as the balance. The soft phase is composed of ferrite, and the hard phase includes one or more of pearlite, bainite, and martensite.

[0029] The proportion of processed ferrite in the overall metal microstructure, expressed as a volume fraction, is more than 50%.

[0030] The aspect ratio of the aforementioned processed ferrite is 1.5 or higher.

[0031] The average grain size of the processed ferrite is less than 50 μm.

[0032] The yield stress of the base material is above 235 MPa, and the Charpy impact absorption energy of the base material at -60℃ is above 200 J.

[0033] The Charpy impact absorption energy of the weld heat-affected zone at -60℃ after high heat input welding is over 80J.

[0034] Ceq.(IIW)=[C]+[Mn] / 6+([Cu]+[Ni]) / 15+([Cr]+[Mo]+[V]) / 5…(1)

[0035] Here, in equation (1), [] represents the content (mass%) of each element within the brackets, and the content of elements that are not present is set to zero.

[0036] [2] The high-strength steel plate described in [1] above, wherein the above composition, by mass %, further contains one or more of the following: B: less than 0.0030%, Cu: less than 0.50%, Ni: less than 1.50%, V: less than 0.100%, Cr: less than 0.50%, Mo: less than 0.50%, Ca: less than 0.0030%, Mg: less than 0.0050%, and REM: less than 0.1000%.

[0037] [3] The high-strength steel plate described in [1] or [2] above, wherein the size of TiN in the metal structure of the high-strength steel plate at the 1 / 4 position of the plate thickness is 5 to 200 nm.

[0038] [4] A method for manufacturing a high-strength steel plate, which is the method for manufacturing a high-strength steel plate as described in any one of [1] to [3] above, wherein,

[0039] The steel raw material with the above composition is heated to a temperature above 1050°C and below 1200°C.

[0040] Then, hot rolling is performed as follows: rolling begins in the temperature range above 950°C, i.e., the recrystallization γ region; rolling is carried out with a reduction rate of 30% or more in the temperature range below 850°C, i.e., the non-recrystallization γ region; and rolling is performed with a reduction rate of 30% or more in the ferrite-austenite dual-phase region from (Ar3 transformation point -80)°C to the Ar3 transformation point; the finishing rolling temperature is above 650°C.

[0041] Then, the temperature is cooled from the cooling start temperature above 650°C to the cooling stop temperature range of below 600°C and above 300°C at an average cooling rate of 5°C / s or higher.

[0042] [5] The method for manufacturing high-strength steel plates described in [4] above, wherein the steel raw material is cast at an average casting speed of 0.3 to 1.0 m / min.

[0043] Invention Effects

[0044] According to the present invention, even when the high-strength steel plate of the present invention is subjected to high heat input welding, it can still exhibit excellent low-temperature toughness in both the base material and the weld heat-affected zone. Therefore, the high-strength steel plate of the present invention can be appropriately used as a welding steel plate (steel) for structures constructed by high heat input welding methods such as electric welding, submerged arc welding, and electroslag welding, such as cryogenic storage tanks for liquefied gas and ships used in cryogenic environments. Detailed Implementation

[0045] The present invention will now be described. It should be noted that the present invention is not limited to the following embodiments.

[0046] First, the reasons for defining the composition of the high-strength steel plate in this invention will be explained. It should be noted that, unless otherwise specified, the "%" designation for the composition in this invention refers to "mass %".

[0047] C: 0.010~0.070%

[0048] Regarding carbon (C), in order to obtain the strength of the target base material (high-strength steel plate), it is necessary to contain 0.010% or more. However, if the C content exceeds 0.070%, island martensite increases, thereby reducing the low-temperature toughness of the weld heat-affected zone (HAZ). Therefore, the C content is set to 0.070% or less. The C content is preferably 0.020% or more, more preferably 0.030% or more, and even more preferably 0.050% or more. The C content is preferably 0.065% or less, more preferably 0.060% or less, and even more preferably 0.055% or less.

[0049] Si: 0.01~0.50%

[0050] Si is a component required for ensuring the strength of the base material and for deoxidation, etc. In this invention, it contains 0.01% or more Si. On the other hand, if the Si content exceeds 0.50%, HAZ hardening occurs, thereby reducing the low-temperature toughness of the HAZ. Therefore, the Si content is set to 0.50% or less. The Si content is preferably 0.1% or more, more preferably 0.15% or more. The Si content is preferably 0.40% or less, more preferably 0.3% or less.

[0051] Mn: 1.00~2.00%

[0052] Regarding Mn, to ensure the strength of the base material, it needs to contain 1.00% or more. On the other hand, if the Mn content exceeds 2.00%, not only will the weldability deteriorate, but the cost of the steel plate will also increase. Therefore, the Mn content is set in the range of 1.00 to 2.00%. The Mn content is preferably 1.20% or more, more preferably 1.40% or more, and even more preferably 1.50% or more. The Mn content is preferably 1.90% or less, more preferably 1.75% or less, and even more preferably 1.60% or less.

[0053] P: below 0.020%

[0054] Phosphorus (P) is an unavoidable impurity. If the P content exceeds 0.020%, it reduces the low-temperature toughness of the base material and the weld. Therefore, the upper limit is set to 0.020%. Thus, the P content is set to 0.020% or less. The P content is preferably set to 0.015% or less. To obtain good low-temperature toughness, the P content is more preferably 0.010% or less, and even more preferably 0.007% or less. It should be noted that the lower limit of the P content does not need to be specifically limited, but since implementing extremely low P treatment will increase costs, the P content is preferably set to 0.001% or more.

[0055] S: 0.0005~0.0100%

[0056] Regarding sulfur (S), it needs to contain at least 0.0005% S to facilitate the formation of CaS or MnS, which are required for the formation of composite inclusions necessary for ferrite nuclei. On the other hand, if the S content exceeds 0.0100%, the low-temperature toughness of the base material deteriorates. Therefore, the S content is set to 0.0100% or less. The S content is preferably 0.0090% or less, more preferably 0.0030% or less. The S content is preferably 0.0010% or more, more preferably 0.0015% or more.

[0057] Al: 0.035~0.100%

[0058] Regarding Al, from the perspective of steel deoxidation, it needs to contain 0.035% or more. On the other hand, if the Al content exceeds 0.100%, the low-temperature toughness of the base metal decreases, and the low-temperature toughness of the weld metal deteriorates. Therefore, the Al content is set to 0.100% or less. The Al content is preferably 0.095% or less, more preferably 0.090% or less, and even more preferably 0.080% or less. The Al content is preferably 0.040% or more, more preferably 0.050% or more.

[0059] Ti: 0.010~0.030%

[0060] Ti precipitates as TiN during steel solidification, which helps suppress the coarsening of austenite in the hardened austenite zone (HAZ) and acts as a nucleus for ferrite transformation, thus contributing to high toughness. If the Ti content is less than 0.010%, this effect is minimal; if it exceeds 0.030%, the desired effect cannot be achieved due to the coarsening of TiN particles. Therefore, the Ti content is set in the range of 0.010 to 0.030%. The Ti content is preferably 0.011% or more, more preferably 0.013% or more, and even more preferably 0.015% or more. The Ti content is preferably 0.028% or less, more preferably 0.025% or less, and even more preferably 0.020% or less.

[0061] N: 0.0035~0.0100%

[0062] Regarding nitrogen (N), it is contained at least 0.0035% to combine with Ti and form TiN. If the N content increases, the amount of N dissolved in the solid solution increases, leading to a decrease in the low-temperature toughness of the HAZ (high-temperature zone). Therefore, the upper limit of the N content is set at 0.0100%. Thus, the N content is set at 0.0100% or less. The N content is preferably 0.0040% or more, more preferably 0.0045% or more, and even more preferably 0.0052% or more. The N content is preferably 0.0095% or less, more preferably 0.0090% or less, and even more preferably 0.0075% or less.

[0063] Carbon equivalent (Ceq.(IIW): 0.25–0.35% by mass)

[0064] For the high-strength steel plate of the present invention, its composition is adjusted so that the carbon equivalent Ceq.(IIW) shown in the following formula (1) is in the range of 0.25 to 0.35% by mass. Ceq.(IIW)=[C]+[Mn] / 6+([Cu]+[Ni]) / 15+([Cr]+[Mo]+[V]) / 5…(1)

[0065] Here, in equation (1), [] represents the content (mass%) of each element within the brackets, and the content of elements that are not present is set to zero.

[0066] Regarding the carbon equivalent Ceq.(IIW), in order to ensure that the vTrs (fracture transition temperature) is below -60°C in the weld heat-affected zone, it is set to 0.35% by mass or less. On the other hand, if the carbon equivalent Ceq.(IIW) is less than 0.25% by mass, it is impossible to ensure that the yield stress in the base material is above 235 MPa. Therefore, the carbon equivalent Ceq.(IIW) is set to 0.25 to 0.35% by mass. The carbon equivalent Ceq.(IIW) is preferably 0.27% by mass or more, more preferably 0.28% by mass or more. The carbon equivalent Ceq.(IIW) is preferably 0.33% by mass or less, more preferably 0.32% by mass or less.

[0067] It should be noted that the reason for setting the above vTrs to below -60°C in the welding heat-affected zone is that it is assumed that the design temperature of the liquefied gas storage tanks that are expected to be added in the future will be below -60°C.

[0068] Ti / N ratio: 1.5–4.0

[0069] Furthermore, in this invention, considering the relationship with N content, Ti is added and the Ti content is adjusted to satisfy 1.5 ≤ Ti / N ≤ 4.0 (here, Ti and N are the contents of each element (mass%)). By adjusting the Ti content, optimal size and amount of TiN can be ensured, resulting in the suppression of austenite coarsening. When Ti / N is less than 1.5, TiN becomes finer, thereby causing TiN to dissolve in the weld heat-affected zone. As a result, the amount of TiN required to improve the low-temperature toughness of the weld cannot be ensured. On the other hand, when Ti / N exceeds 4.0, the low-temperature toughness of the base material and the low-temperature toughness of the weld decrease due to the formation of TiC particles and the coarsening of TiN. Therefore, the ratio of Ti content to N content (T / Ni) is set to 1.5 or more and 4.0 or less. T / Ni is preferably 2.0 or more, more preferably 2.5 or more. T / Ni is preferably 3.4 or less, more preferably 3.2 or less.

[0070] It should be noted that, in this invention, the size of TiN is preferably adjusted to be 5 nm or more and 200 nm or less. The reason for this is that if the size is outside this range, a sufficient effect on suppressing austenite grain coarsening cannot be obtained. In this invention, "the size of TiN" refers to the diagonal length of the rectangular TiN, which can be measured using the method described in the embodiments below.

[0071] In the high-strength steel plate of the present invention, the balance other than the above-mentioned components is iron (Fe) and unavoidable impurities.

[0072] In this invention, the aforementioned elements are set as the basic components, and the target characteristics of this invention can be obtained through this basic component composition. In order to further improve the characteristics, the following elements can be added as needed to the above-mentioned basic component composition. It should be noted that the components B, Cu, Ni, V, Cr, Mo, Ca, Mg, and REM listed below can be included as needed, and therefore these components can also be 0%.

[0073] Choose one or more from the group consisting of B: less than 0.0030%, Cu: less than 0.50%, Ni: less than 1.50%, V: less than 0.100%, Cr: less than 0.50%, Mo: less than 0.50%, Ca: less than 0.0030%, Mg: less than 0.0050%, and REM: less than 0.1000%.

[0074] B: Below 0.0030%

[0075] Boron (B) is an element that effectively contributes to the high strength of the steel plate (base material). This effect becomes significant when the B content is 0.0002% or more. On the other hand, if B is present in excess, it will adversely affect the low-temperature toughness of the HAZ in the welded area; therefore, the B content is preferably set to 0.0030% or less. Thus, when B is present, the B content is preferably set to 0.0030% or less. The B content is preferably set to 0.0002% or more, more preferably 0.0007% or more. The B content is more preferably set to 0.0012% or less.

[0076] Cu: less than 0.50%

[0077] Cu is an element that improves the hardenability of steel. Besides contributing to increased strength of the rolled base material, it also contributes to improved high-temperature strength and weather resistance. These effects are achieved by containing 0.01% or more Cu. On the other hand, excessive Cu content can degrade the low-temperature toughness and weldability of the welded area (HAZ). The Cu content is preferably 0.50% or less. Therefore, when Cu is present, the Cu content is preferably 0.50% or less. The Cu content is preferably 0.01% or more, more preferably 0.04% or more. The Cu content is even more preferably 0.10% or less.

[0078] Ni: below 1.50%

[0079] Ni is an element that improves the hardenability of steel. Besides contributing to increased strength of the rolled base material, it also improves its low-temperature toughness, high-temperature strength, and weather resistance. These effects are achieved by containing 0.01% or more Ni. On the other hand, excessive Ni content deteriorates the low-temperature toughness and weldability of the weld zone (HAZ) and increases the cost of the alloy. The Ni content is preferably 1.50% or less. Therefore, when Ni is present, the Ni content is preferably 1.50% or less. The Ni content is preferably 0.01% or more, more preferably 0.02% or more. The Ni content is even more preferably 0.50% or less.

[0080] V: Below 0.100%

[0081] V is an effective element for improving the strength and low-temperature toughness of the base material, and it functions as a ferrite nucleus in the form of VN. This effect is achieved by containing 0.005% or more of V. On the other hand, if V is contained in amounts exceeding 0.100%, it leads to a decrease in the low-temperature toughness of the HAZ in the weld. The V content is preferably set to 0.100% or less. Therefore, when V is present, the V content is preferably set to 0.100% or less. The V content is preferably set to 0.005% or more, more preferably 0.009% or more. The V content is more preferably set to 0.080% or less.

[0082] Cr: less than 0.50%

[0083] Similar to Cu, Cr is an element that improves the hardenability of steel. Besides contributing to increased strength of the rolled base material, it also contributes to improved high-temperature strength and weather resistance. These effects are achieved by containing 0.01% or more Cr. On the other hand, excessive Cr content deteriorates the low-temperature toughness and weldability of the welded area (HAZ). The Cr content is preferably 0.50% or less. Therefore, when Cr is present, the Cr content is preferably 0.50% or less. The Cr content is preferably 0.01% or more, more preferably 0.02% or more. The Cr content is even more preferably 0.10% or less.

[0084] Mo: 0.50% or less

[0085] Similar to Cu and Cr, Mo is an element that improves the hardenability of steel. Besides contributing to increased strength of the rolled base material, it also contributes to improved high-temperature strength and weather resistance. These effects are achieved by containing 0.01% or more Mo. On the other hand, excessive Mo content deteriorates the low-temperature toughness and weldability of the welded area (HAZ). The Mo content is preferably 0.50% or less. Therefore, when Mo is present, the Mo content is preferably 0.50% or less. The Mo content is preferably 0.01% or more, more preferably 0.02% or more. The Mo content is even more preferably 0.10% or less.

[0086] Ca: below 0.0030%

[0087] Ca is a useful element for improving the low-temperature toughness of the base material and HAZ due to the fixation of S. If the Ca content exceeds 0.0030%, the effect saturates, so the Ca content is set to be 0.0030% or less. On the other hand, if the Ca content is less than 0.0005%, the fixation of S becomes insufficient. The Ca content is set to be 0.0005% or more. Therefore, when Ca is present, the Ca content is preferably set to 0.0030% or less. The Ca content is preferably set to 0.0005% or more, more preferably 0.0010% or more. The Ca content is more preferably set to 0.0025% or less.

[0088] Mg: less than 0.0050%, REM: less than 0.1000%

[0089] Both Mg and REM (rare earth metals) have strong deoxidizing power in molten steel and play a role in assisting the formation of fine oxides; therefore, they are added as needed. Regarding the content at which they exhibit deoxidizing effects, Mg is 0.0002% or more, and REM is 0.0010% or more. On the other hand, if added in large quantities, coarse inclusions are generated, damaging the properties of the base material; therefore, the respective contents are preferably set at Mg 0.0050% or less and REM 0.1000% or less. Therefore, when Mg and REM are present, it is preferable to set Mg 0.0050% or less and REM 0.1000% or less. The Mg content is preferably 0.0002% or more. The REM content is preferably 0.0010% or more.

[0090] The high-strength steel sheet of the present invention achieves improved low-temperature toughness by satisfying the above-described composition. On the other hand, as mentioned above, it is difficult to ensure the strength of the base material. Therefore, in order to ensure this strength, it is also important in the present invention to specify the metal structure of the high-strength steel sheet as follows.

[0091] The following describes the metal structure of the high-strength steel plate of the present invention.

[0092] The high-strength steel sheet of the present invention has a metallic structure at a position of 1 / 4 of the sheet thickness, consisting of a soft phase as the main phase and a hard phase as the balance. The soft phase is composed of ferrite, and the hard phase includes one or more of pearlite, bainite, and martensite. Furthermore, regarding the processed ferrite, at the 1 / 4 of the sheet thickness, the processed ferrite accounts for 50% or more of the overall metallic structure by volume fraction, the aspect ratio of the processed ferrite is 1.5 or more, and the average grain size of the processed ferrite is 50 μm or less.

[0093] Principal phase: Ferrite

[0094] From the viewpoint of improving the strength of the base material, the high-strength steel sheet of the present invention has ferrite as the main phase at the 1 / 4 position of the sheet thickness. In the present invention, "main phase" refers to a volume fraction of 50% or more. The volume fraction of ferrite is preferably 70% or more, more preferably 75% or more, even more preferably 80% or more, and even more preferably 90% or more.

[0095] It should be noted that, in this invention, as described above, in order to obtain excellent base material strength while satisfying the condition of equation (1), in addition to specifying ferrite as a soft phase, the fraction of processed ferrite at the 1 / 4 position of the steel plate thickness is also specified. Details regarding processed ferrite are described below.

[0096] Balance: Contains one or more hard phases selected from pearlite, bainite, and martensite.

[0097] Regarding the microstructure of the remainder other than ferrite, from the viewpoint of ensuring strength, a hard phase comprising one or more of pearlite, bainite, and martensite is provided at the 1 / 4 position of the plate thickness. The total volume fraction of the remainder is preferably 25% or less. More preferably, the total volume fraction of the remainder microstructure is 15% or less, and even more preferably, 10% or less.

[0098] The volume fraction of processed ferrite is 50% or more.

[0099] In this invention, during the manufacture of high-strength steel plates, dislocations are introduced into ferrite within the two-phase region through rolling under the hot-rolling conditions described later, thereby increasing strength. To achieve this effect, processed ferrite with dislocations introduced through two-phase rolling needs to be present in a certain proportion or higher. To ensure the strength of the base material, the processed ferrite needs to be set at 50% or more in volume fraction relative to the overall metal structure of the high-strength steel plate at the 1 / 4 position of the plate thickness. Preferably, the processed ferrite is set at 60% or more in volume fraction.

[0100] The processed ferrite constitutes more than 55% of the aforementioned ferrite. Preferably, it is more than 70%.

[0101] It should be noted that there is no particular limit to the upper limit of processed ferrite, but from the perspective of rolling mill load and to prevent the reduction of absorbed energy due to peeling, it is preferable to set it to 90% or less in terms of volume fraction. More preferably, the processed ferrite is set to 80% or less in terms of volume fraction.

[0102] The processed ferrite comprises 96% or less of the aforementioned ferrite. Preferably, it comprises 93% or less.

[0103] Here, "processed ferrite" in this invention refers to a dislocation density ρ determined by X-ray diffraction (XRD) with a value of 1.0 × 10⁻⁶. 14 m -2 The above ferrite. By making the dislocation density ρ 1.0 × 10 14 m -2 As described above, work hardening takes effect, increasing strength. The dislocation density ρ is preferably set to 2.0 × 10⁻⁶. 14 m -2 However, excessive two-phase compression introduces an excessive number of dislocations, making them difficult to move, resulting in reduced low-temperature toughness of the steel plate. Therefore, the upper limit of the dislocation density ρ is set at 2.5 × 10⁻⁶. 15 m -2 Hereinafter, in this invention, the dislocation density described above can be determined by the method described in the embodiments below.

[0104] The aspect ratio of the machined ferrite is 1.5 or higher.

[0105] If the aspect ratio of the processed ferrite at the 1 / 4 position of the plate thickness is less than 1.5, the specific texture cannot be sufficiently developed, potentially leading to ductile cracking. Furthermore, dislocation strengthening and grain refinement effects will not be observed, resulting in reduced toughness of the base material at low temperatures. Therefore, the aspect ratio of the processed ferrite is set to 1.5 or higher. Preferably, the aspect ratio of the processed ferrite is 2.0 or higher, more preferably 2.5 or higher. It should be noted that the upper limit of the aspect ratio of the processed ferrite does not need to be specifically specified. From the viewpoint of rolling mill performance, the aspect ratio of the processed ferrite is preferably 4.0 or lower, more preferably 3.8 or lower.

[0106] Average grain size of processed ferrite: less than 50 μm

[0107] By setting the average grain size of the processed ferrite at one-quarter of the plate thickness to 50 μm or less, the number of grain boundaries per unit volume increases, resulting in greater difficulty for dislocations to move, thereby increasing the strength of the steel plate. The average grain size of the processed ferrite is preferably 45 μm or less, more preferably 40 μm or less, and even more preferably 30 μm or less. It should be noted that there is no particular upper limit to the average grain size of the processed ferrite. From the viewpoint of applying the steel plate to the above-described applications, it is preferable to set it to 5 μm or more, more preferably 15 μm or more.

[0108] Here, for the grain size of the processed ferrite of the present invention, the orientation difference between adjacent grains is calculated, and the boundary with an orientation difference of 15° or more is defined as a grain boundary and measured. The arithmetic mean of the grain size obtained from the obtained grain boundaries is used to obtain the average circle equivalent diameter, which is defined as the above-mentioned average grain size.

[0109] It should be noted that the volume fractions of ferrite, pearlite, bainite, martensite, and processed ferrite, the aspect ratio of processed ferrite, and the average grain size of processed ferrite can be determined by the methods described in the examples below.

[0110] Next, one embodiment of the method for manufacturing the high-strength steel plate of the present invention will be described.

[0111] First, molten steel with the above-mentioned composition is smelted using a converter, electric furnace, or other smelting methods. Additionally, a second refining process can be performed using a vacuum degassing furnace. Afterward, steel billets and other raw materials of specified dimensions are preferably obtained through casting methods such as continuous casting or ingot-rolling.

[0112] In this invention, steel billets can be used as the aforementioned steel raw material. When manufacturing this steel raw material using a continuous casting method, the casting conditions preferably satisfy the following conditions.

[0113] Specifically, the average casting speed during billet casting is preferably set to 0.3 m / min or more and 1.0 m / min or less. The cooling of the billet (raw steel) can be controlled by the casting speed. If the average casting speed is less than 0.3 m / min, the size of TiN in the base material (high-strength steel plate) becomes larger. If the TiN size increases, the TiN density in the base material (high-strength steel plate) decreases, thereby reducing the pinning effect. As a result, sufficient ferrite refinement cannot be achieved in the HAZ of the welded section, and the low-temperature toughness of the HAZ may deteriorate. It should be noted that, as mentioned above, the size of the TiN forming the nucleation site is preferably 5 nm or more and 200 nm or less. The casting speed mentioned above is the average speed of the entire casting process.

[0114] On the other hand, if the average casting speed exceeds 1.0 m / min, the density of TiN increases, but the size of TiN decreases relative to the aforementioned range. Consequently, TiN undergoes solid solution due to the high heat input during welding of the base material (high-strength steel plate). As a result, the austenite grain size coarsens, potentially leading to a deterioration in the low-temperature toughness of the HAZ.

[0115] Then, the reasons for limiting the manufacturing conditions of high-strength steel plates with excellent low-temperature toughness and high heat input energy (HAZ) for manufacturing the above-mentioned steel raw materials as base materials are explained in detail.

[0116] In this invention, the aforementioned steel raw material is heated to a temperature of 1050°C or higher and 1200°C or lower, and then hot-rolled as follows: rolling begins in the recrystallization γ-zone at 950°C or higher, rolling is performed under the following conditions: a reduction rate of 30% or higher in the non-recrystallization γ-zone below 850°C, and a reduction rate of 30% or higher in the ferrite-austenite dual-phase zone from (Ar3 transformation point -80)°C to the Ar3 transformation point; the finishing rolling temperature is 650°C or higher. After this hot rolling, the hot-rolled steel sheet is cooled from a cooling start temperature of 650°C or higher to a cooling stop temperature in the range of 600°C or lower and 300°C or higher at an average cooling rate of 5°C / s or higher.

[0117] It should be noted that, unless otherwise specified, the temperature reference "°C" in the following manufacturing method description refers to the surface temperature of each steel raw material or steel plate. Surface temperature can be measured using, for example, a radiation thermometer. Furthermore, the temperature at the center of the steel billet or steel plate thickness can be measured, for example, by installing a thermocouple at the center of the steel plate thickness, or by calculating the temperature distribution within the steel plate cross-section through thermal conductivity analysis, and then correcting the result using the surface temperature of the steel plate.

[0118] [Heating temperature of steel raw materials]

[0119] The heating temperature of the steel raw material (e.g., steel billet) needs to be set to 1050°C or higher and 1200°C or lower. This is because when heating to a temperature lower than 1050°C, there is a possibility that large inclusions generated during the solidification of the steel billet, which adversely affect low-temperature toughness, may remain unmelted. On the other hand, if heating is performed at high temperatures, there is a possibility that precipitates created by controlling the cooling rate during casting may remelt. Based on this, a heating temperature of 1200°C or lower is sufficient, in the sense of ending the phase transformation. It should be noted that the grain coarsening that is thought to occur during heat holding can also be prevented in advance by the aforementioned pinning effect of TiN. For the above reasons, the heating temperature is limited to 1050°C or higher and 1200°C or lower. The heating temperature is preferably 1180°C or lower, and more preferably 1100°C or lower.

[0120] [Hot Rolling Conditions]

[0121] Rolling start temperature: above 950℃

[0122] For heated steel raw materials, rolling should begin within a temperature range of 950°C or higher, i.e., the recrystallization temperature range (recrystallization γ-region). Within this temperature range, austenite grains recrystallize due to rolling. As a result, the microstructure can be refined. If rolling begins at a temperature below 950°C, sufficient recrystallization of austenite grains cannot occur, resulting in insufficient microstructure refinement. Consequently, the strength of the base material decreases. The rolling start temperature is preferably set to 970°C or higher, more preferably 1000°C or higher. While there is no particular upper limit to the rolling start temperature, from the viewpoint of suppressing roll wear caused by contact with the high-temperature steel raw material, it is preferably set to 1100°C or lower.

[0123] The reduction rate in the non-recrystallized γ-region below 850℃ is 30% or more.

[0124] Hot rolling with a reduction rate of 30% or more is performed in the temperature range below 850°C, i.e., the non-recrystallization temperature range (non-recrystallization γ-region). The reasons are as follows: In this temperature range, austenite grain recrystallization does not occur; the austenite grains deform flatten, and defects such as deformation bands are introduced into the interior of the austenite grains. This accumulated internal energy is applied to the driving force of the subsequent ferrite phase transformation. If the reduction rate in the non-recrystallization γ-region is less than 30%, the accumulated internal energy is insufficient, and therefore the ferrite is not sufficiently refined, thus failing to ensure the strength of the base material. It should be noted that the reduction rate in the non-recrystallization γ-region is preferably 35% or more, more preferably 40% or more. There is no specific upper limit for the reduction rate in the non-recrystallization γ-region. If the reduction rate in the non-recrystallization γ-region is excessively increased, manufacturing efficiency decreases; therefore, it is preferably 55% or less, more preferably 50% or less.

[0125] The reduction rate of the ferrite-austenite two-phase region from the Ar3 phase transformation point (-80)℃ to the Ar3 phase transformation point is over 30%.

[0126] Hot rolling with a reduction rate of 30% or more is required within the temperature range of (Ar3 transformation point - 80)℃ to the Ar3 transformation point, i.e., the ferrite-austenite two-phase temperature range. The reasons are as follows: Increasing the reduction rate within the aforementioned two-phase temperature range has the following effects: increased strength associated with dislocation strengthening due to ferrite processing during rolling, and improved low-temperature toughness resulting from the grain refinement effect achieved through the formation of secondary grains caused by processing. Furthermore, by setting the reduction rate within the ferrite-austenite two-phase temperature range to 30%, the ferrite rolling texture is developed, which contributes to improved low-temperature toughness. Based on these reasons, the reduction rate within the aforementioned ferrite-austenite two-phase temperature range is limited to 30% or more. It should be noted that the reduction rate within the aforementioned two-phase temperature range is preferably 35% or more, and more preferably 40% or more. There is no particular upper limit for the reduction rate within the aforementioned two-phase temperature range. From the viewpoint of ensuring the flatness of the steel sheet, it is preferable to set it to 50% or less.

[0127] Here, the Ar3 phase transition point can be obtained by the following formula.

[0128] Ar3(℃)=910-273×C-74×Mn-57×Ni-16×Cr-9×Mo-5×Cu

[0129] In this formula, each element represents its content (mass%), and the content of elements that are not present is set to zero.

[0130] Finishing rolling temperature: above 650℃

[0131] The finishing rolling temperature (finishing end temperature) in hot rolling is set to 650°C or higher. That is, hot rolling is completed (finished) at a temperature above 650°C because if finishing rolling is performed at a temperature below 650°C, the ferrite formed due to phase transformation is subjected to strain exceeding the necessary limits. As a result, the low-temperature toughness of the base material decreases. It should be noted that the finishing rolling temperature is preferably set to 670°C or higher, and more preferably 680°C or higher. There is no specific upper limit for the finishing rolling temperature. To ensure rolling within the ferrite-austenite two-phase temperature range, it is preferably set to 710°C or lower.

[0132] Cooling conditions after hot rolling

[0133] In this invention, after the above-described hot rolling process, the hot-rolled steel sheet is cooled under the following conditions. This cooling process improves the strength of the base material.

[0134] Cooling start temperature: above 650℃

[0135] The reason for starting cooling at a temperature above 650°C is that starting cooling at a temperature below 650°C is disadvantageous from a hardenability point of view, and there is a possibility that the desired base material strength cannot be obtained. The cooling start temperature is preferably set to 670°C or higher, more preferably 680°C or higher. There is no specific upper limit for the cooling start temperature. To begin cooling below the Ar3 phase transformation point, it is preferably set to 710°C or lower.

[0136] Cooling stop temperature: Temperature range between 300°C and 600°C

[0137] The hot-rolled steel sheet is cooled from the aforementioned cooling start temperature to a temperature range of 300°C to 600°C (cooling stop temperature). This is because, from a hardenability point of view, it is difficult to ensure sufficient strength when cooling stops above 600°C. Furthermore, since cooling stops below 300°C does not impart significant changes to the steel sheet's properties, only the operational load increases. The cooling stop temperature is preferably set to 570°C or below, and more preferably to 520°C or below.

[0138] Average cooling rate: 5℃ / s or more

[0139] When the average cooling rate within the aforementioned temperature range is less than 5°C / s, it is difficult to obtain steel with a uniform metallic structure, and the strength and low-temperature toughness of the base material cannot be guaranteed. The average cooling rate is preferably 7°C / s or higher, more preferably 10°C / s or higher. The average cooling rate is preferably 100°C / s or lower, more preferably 80°C / s or lower, and even more preferably 60°C / s or lower.

[0140] For this reason, after the steel sheet is hot rolled at a finishing rolling temperature of 650°C or higher, it is cooled from the cooling start temperature of 650°C or higher to the cooling stop temperature of 300°C or higher and 600°C or lower at an average cooling rate of 5°C / s or higher.

[0141] The high-strength steel sheet manufactured under the above manufacturing conditions has the above-described composition and also the above-described metallic structure. That is, it is a structure consisting of ferrite as the main phase and one or more of pearlite, bainite, and martensite as the balance, and has the above-described processed ferrite.

[0142] As explained above, the high-strength steel plate of the present invention possesses characteristics suitable for use as a raw material for the aforementioned applications (e.g., cryogenic storage tanks for liquefied gas, structures such as ships used in cryogenic environments). In particular, it exhibits characteristics such as a yield stress of 235 MPa or higher in the base material and high low-temperature toughness, specifically a Charpy impact absorption energy of 200 J or higher at -60°C. Furthermore, joints constructed using the high-strength steel plate of the present invention as welding steel plates and through high heat input welding exhibit a Charpy impact absorption energy of 80 J or higher at -60°C in the weld heat-affected zone, and also demonstrate excellent low-temperature toughness in the weld heat-affected zone.

[0143] Example

[0144] Next, the present invention will be described in detail based on embodiments. It should be noted that the following embodiments illustrate a preferred example of the present invention, and the present invention is not limited to this embodiment.

[0145] Steel billets (steel raw materials) with the composition shown in Table 1 were produced using the converter-ladle refining-continuous casting method. It should be noted that blank columns in Table 1 indicate unintentional addition of elements, meaning that this includes not only cases where the element is not present (0%), but also cases where the element is unavoidably present.

[0146] The obtained steel billet is heated and then cooled according to the various conditions shown in Tables 2-1 and 2-2, followed by hot rolling and then cooling treatment to obtain a high-strength steel plate (thick steel plate) with a thickness (finished rolling thickness) of 8 to 25 mm.

[0147] First, the microstructure of the obtained high-strength steel plates is evaluated using the method shown below.

[0148] (1) Evaluation of the metal structure of high-strength steel plates

[0149] [Metal Structure]

[0150] A sample was cut from the obtained high-strength steel plate, centered at 1 / 4 of the plate thickness and with the surface perpendicular to the plate width as the observation plane. The surface of the sample was mirror-polished, then etched with nitric acid-ethanol solution, and observed using an optical microscope (magnification: 200x). The microstructure was photographed, and the ferrite fraction (area %) was calculated using an image analysis device. Ten fields of view were observed, and the average ferrite fraction (area %) was calculated for each field. Given the low anisotropy of the microstructure, the area fraction is equivalent to the volume fraction; therefore, in this patent, the average ferrite area fraction is used as the volume fraction.

[0151] It should be noted that in Tables 3-1 and 3-2, ferrite is represented by F, pearlite by P, bainite by B, and martensite by M.

[0152] [Ferrite fraction]

[0153] Regarding the fraction of processed ferrite, samples were cut from 1 / 4 of the steel plate thickness, mirror-polished, and then finely polished with colloidal silica. EBSD (electron beam backscattering diffraction) measurements were then performed using SEM (scanning electron microscope). Crystal orientation was measured at 500x magnification. Based on the obtained data, boundaries with an orientation difference of 15° or more between adjacent measurement points were defined as grain boundaries. Then, ferrite regions within each grain surrounded by these grain boundaries, where the GAM (Grain Average Misorientation) value reached 1.0 or more, were defined as processed ferrite, and their area fraction was calculated. Ten fields of view were measured, and the average processed ferrite area fraction was used as the volume fraction, similar to the method described above.

[0154] [Aspect Ratio of Machining Ferrite]

[0155] Regarding the aspect ratio of the processed ferrite, samples were cut from the steel plate at 1 / 4 of its thickness, as described above. The surface was mirror-polished and etched to expose the grain boundaries of the processed ferrite. Then, 10–20 fields of view were photographed using an optical microscope at 200x magnification. For each processed ferrite grain in each field of view, the average value of all grains was calculated by dividing the maximum length in the rolling direction by the maximum length in the thickness direction. This average value was used as the aspect ratio of the processed ferrite.

[0156] Average grain size of processed ferrite

[0157] Regarding the average grain size of the processed ferrite, samples were cut from 1 / 4 of the steel plate thickness, mirror-polished, and then finely polished with colloidal silica, followed by EBSD analysis using SEM. The crystal orientation was measured at 500x magnification. Based on the obtained data, boundaries with an orientation difference of 15° or more between adjacent measurement points were defined as grain boundaries. Then, ferrite regions with a GAM value of 1.0 or higher within each grain enclosed by these grain boundaries were defined as processed ferrite, and the area of ​​the processed ferrite was calculated. The equivalent diameter of a circle with the same area as these processed ferrites was taken as their respective grain size, and the average value of the obtained grain sizes was taken as the average grain size of the processed ferrite.

[0158] [TiN dimensions]

[0159] Regarding the size of TiN, thin film samples were cut from the steel plate at 1 / 4 of its thickness, and the precipitates were measured using TEM (transmission electron microscopy). Ten 1μm × 1μm fields of view were observed, and the diagonal length of the TiN samples, which were arranged as rectangles, was measured. The average diagonal length of all TiN samples was calculated as the size of the TiN.

[0160] Then, using the obtained high-strength steel plates, the characteristics of the base material and the weld heat-affected zone after high heat input welding are evaluated by the method shown below.

[0161] (2) Evaluation of the properties of the base material

[0162] Tensile test pieces based on JIS Z2241 (2011) were cut from the 1 / 4 position of the thickness of each high-strength steel plate. Then, tensile tests based on JIS Z2241 (2011) were performed to determine the yield stress (YP) of the base material. In this embodiment, a yield stress of 235 MPa or higher was considered to indicate excellent strength (high strength) of the base material.

[0163] In addition, test pieces based on JIS Z2242 (2018) were cut from 1 / 4 of the thickness of each high-strength steel plate. Then, Charpy impact tests based on JIS Z2242 (2018) were performed to determine the Charpy impact absorption energy (vE) of the base material at -60°C. -60 In this embodiment, the average absorbed energy of the three roots at -60℃ being above 200J is considered to indicate that the base material has excellent low-temperature toughness.

[0164] (3) Evaluation of the characteristics of the heat-affected zone in welding

[0165] Then, test pieces for weld joint fabrication were cut from the 1 / 4 position of the thickness of each high-strength steel plate. High heat input weld joints were fabricated using the single-sided, single-pass welding method described above. The welding conditions were set to the welding heat inputs shown in Tables 3-1 and 3-2. JIS No. 4 impact test pieces were cut from these weld joints, with the notch position as the joint area, and Charpy impact tests were performed to determine the Charpy impact absorbed energy (vE) at -60°C in the HAZ. -60 In this embodiment, the average absorbed energy of the three roots at -60℃ is determined to be above 80J, indicating that the low-temperature toughness of the HAZ is excellent.

[0166] The results of the above measurements are shown in Tables 3-1 and 3-2.

[0167]

[0168]

[0169]

[0170]

[0171]

[0172] As shown in Tables 3-1 and 3-2, the high-strength steel sheet of the present invention has the composition that satisfies the condition of equation (1) as described above, and the metal structure of the base material is satisfied at the 1 / 4 position of the plate thickness. Thus, it is confirmed that the high strength and excellent low-temperature toughness of the base material are combined. Furthermore, it is confirmed that in the areas where high heat input welding is performed on the high-strength steel sheet of the present invention, the absorbed energy (vE) of the Charpy impact test is... -60 It meets the requirements of 80J and has excellent low-temperature toughness at HAZ.

[0173] In contrast, comparative examples that fall outside the scope of this invention fail to meet the above-mentioned characteristics.

Claims

1. A high-strength steel sheet having a composition consisting of, in mass%, C: 0.010 to 0.070%, Si: 0.01 to 0.50%, Mn: 1.00 to 2.00%, P: 0.020% or less, S: 0.0005 to 0.0100%, Al: 0.035 to 0.100%, Ti: 0.010 to 0.030%, and N: 0.0035 to 0.0100%, a carbon equivalent Ceq. (IIW) represented by the following formula (1) satisfying 0.25 to 0.35 mass%, Ti / N satisfying 1.5 to 4.0, and the balance consisting of Fe and unavoidable impurities, the metal structure of the high-strength steel sheet at a sheet thickness 1 / 4 position is composed of a soft phase as a main phase and a hard phase as a balance, the soft phase is composed of ferrite, and the hard phase contains one or two or more of pearlite, bainite, and martensite, a fraction of worked ferrite in the entire metal structure is 50% or more in volume fraction, an aspect ratio of the worked ferrite is 1.5 or more, an average crystal grain diameter of the worked ferrite is 50 μm or less, a yield stress of the base material is 235 MPa or more, and a Charpy impact absorbed energy at -60°C of the base material is 200 J or more, a Charpy impact absorbed energy at -60°C of a weld heat-affected zone after a large heat input welding is 80 J or more, Ceq. (IIW) = [C] + [Mn] / 6 + ([Cu] + [Ni]) / 15 + ([Cr] + [Mo] + [V]) / 5... (1) Here, [] in formula (1) is a content (mass%) of each element in the bracket, and a content of an element not contained is set to zero.

2. The high-strength steel sheet according to claim 1, wherein the composition further contains, in mass%, one or two or more selected from the group consisting of B: 0.0030% or less, Cu: 0.50% or less, Ni: 1.50% or less, V: 0.100% or less, Cr: 0.50% or less, Mo: 0.50% or less, Ca: 0.0030% or less, Mg: 0.0050% or less, and REM: 0.1000% or less.

3. The high-strength steel sheet according to claim 1 or 2, wherein, In the metal structure of the high-strength steel sheet at the sheet thickness 1 / 4 position, a size of TiN is 5 to 200 nm.

4. A method of producing a high-strength steel sheet according to any one of claims 1 to 3, wherein a steel raw material having the composition is heated to a temperature of 1050°C or higher and 1200°C or lower, then hot rolling is performed by starting rolling at a temperature range of 950°C or higher, that is, a recrystallization γ region, and rolling at a reduction ratio of 30% or more in a temperature range of 850°C or lower, that is, a non-recrystallization γ region, and a reduction ratio of 30% or more in a ferrite-austenite two-phase region of (Ar3 transformation point - 80) °C to the Ar3 transformation point, with a finish rolling temperature of 650°C or higher, then, cooling is performed at an average cooling rate of 5°C / s or more from a cooling start temperature of 650°C or higher to a cooling stop temperature in a temperature range of 600°C or lower and 300°C or higher.

5. The method of producing a high-strength steel sheet according to claim 4, wherein The steel material is cast at an average casting speed of 0.3 to 1.0 m / min. The steel material is cast at an average casting speed of 0.3 to 1.0 m / min.

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