JP7045459B2 - High-strength steel materials for polar environments with excellent fracture resistance at low temperatures and their manufacturing methods - Google Patents

Description

The present invention relates to a high-strength steel material for polar environment and a method for manufacturing the same, which has excellent fracture resistance at low temperature, and more specifically, to provide fracture resistance at low temperature, which can be suitably applied to steel materials for shipbuilding and offshore structures. Regarding excellent high-strength steel materials for polar environment and their manufacturing methods.

Due to global warming, the amount of ice in the Arctic region is gradually decreasing, and as a result, interest in the Arctic Sea Route connecting Europe and East Asia is increasing. In recent years, cargo ships may be operated on a trial basis only in the summer, and it has been reported that time and cost can be saved by up to 30% or more compared to the conventional route through Southeast Asia. In addition, if the Arctic ice disappears completely within 20 to 30 years, a straight route across the Arctic point is expected. Therefore, the need for ships passing through the Arctic region is becoming more and more real, and the need for safe ship design and polar environmental steel materials to be used in such polar environments is increasing.

Conventional structural steel materials are vulnerable to fracture in a polar environment, that is, an environment exposed to impacts such as low temperatures up to -60 degrees and drift ice, so a new polar environment with excellent fracture resistance at low temperatures can overcome this. High-strength steel material is required.
The reasons why thick high-strength steel materials used in large vessels and oil mining platforms are generally vulnerable to fracture at low temperatures are as follows. For high-strength extra-thick steel materials, the amount of alloying elements such as Mn and Mo must be increased in order to secure the strength, and low rolling reduction and slow acceleration during the production of extra-thick steel materials. This is because a hard phase structure such as coarse granular bainite or MA is likely to be formed depending on the cooling rate. Due to such a fine structure, the steel material has a characteristic that the resistance property against fracture at a low temperature is extremely weak. Therefore, in order to have high strength and excellent fracture resistance at low temperature of ultra-thick materials, it is necessary to miniaturize the structure and extremely reduce hard structures such as granular bainite and MA.

In order to solve the above problems, (1) the reheating temperature of the slab is lowered to the utmost and the structure is refined by controlled rolling at a low temperature, or (2) Cu is added at 1% or more and the temperature is low. By tempering with, the strength is improved with fine Cu precipitates, or (3) a large amount of Ni is added to improve the low temperature toughness of granular bainite, which is a hard phase, or (4) MA structure. In order to reduce the temperature to the utmost, a method such as minimizing the promoting element such as C is used. However, since structures such as ships are gradually becoming larger and the environment in which they are used is changing to polar environments, simply applying the conventional method as described above will result in fracture initiation and propagation resistance at low temperatures. There is a problem that it is difficult to secure enough.
Therefore, there is a demand for the development of high-strength steel materials having improved fracture initiation and propagation resistance at low temperatures and methods for producing the same.

Korean Published Patent No. 2002-0028203

The present invention has been made to solve the above-mentioned problems of the prior art, and an object thereof is to provide a high-strength steel material having excellent fracture initiation and propagation resistance at a low temperature and a method for producing the same. To do.
The subject of the present invention is not limited to the above contents, but can be understood from the entire contents of the present specification, and any person who has ordinary knowledge in the technical field to which the present invention belongs can add to the present invention. There is no difficulty in understanding the challenges.

The high-strength steel material having excellent fracture resistance at low temperature of the present invention made to achieve the above object is
By weight%, C: 0.005 to 0.07%, Si: 0.005 to 0.3%, Mn: 1.7 to 3.0%, Sol. Al: 0.001 to 0.035%, Nb: 0.02% or less (excluding 0%), V: 0.01% or less (excluding 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01 to 2.0%, Cr: 0.01 to 0.5%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001 to 0.008%, P: 0.02% or less (excluding 0%), S: 0.003% or less (excluding 0%), O: 0.003% It consists of the following (excluding 0%), the remaining Fe, and unavoidable impurities, and satisfies the following relational expression 1 and the following relational expression 2.
The microstructure is characterized by containing 70 area% or more of polygonal ferrite and acicular ferrite in total, and 3.5 area% or less of MA phase (martensite-austenite composite phase).

[Relational expression 1]
Mn + 0.5x (Ni + Cu) ≧ 2.5 wt%
[Relational expression 2]
Mo + Cr + 1.5xSi + 10xNb ≦ 0.5wt%
(However, in the above relational expressions 1 and 2, each element is a value shown in% by weight.)

Further, the method for producing a high-strength steel material having excellent fracture resistance at low temperatures according to the present invention is
At the stage of preparing a steel slab that meets the above alloy composition,
The step of heating the steel slab to 1000-1200 ° C.
The step of finishing and hot rolling the heated slab in a temperature range of 650 ° C. or higher so that the total reduction rate in the unrecrystallized region temperature section is 30% or more (excluding the reduction rate in the recrystallization region). When,
It is characterized by including a step of cooling the hot-rolled hot-rolled steel sheet to a cooling end temperature of 200 to 550 ° C. at a cooling rate of 2 to 30 ° C./s.

According to the present invention, it is possible to effectively provide a steel material having significantly improved fracture initiation and propagation resistance at a low temperature.

In this embodiment, it is a graph which measured and showed the Kca value of the steel material of invention example 1. In this Example, it is a microstructure photograph of the steel material of Invention Example 3.

Hereinafter, preferred embodiments of the present invention will be described. However, embodiments of the present invention can be transformed into various other embodiments, and the scope of the invention is not limited to the embodiments described below. Also, embodiments of the invention are provided to more fully explain the invention to those with average knowledge in the art.

As a result of diligent research to further improve the initiation of fracture and propagation resistance at low temperatures, the present inventors have made alloying elements such as C, Mo, Cr, and Nb, which are alloying elements that generate carbide. It was confirmed that it is necessary to precisely control the amount of alloying elements added to the maximum, which has the effect of simultaneously improving the strength and toughness of the ferrite matrix while minimizing the amount of alloy added. By controlling in this way, the microstructure of the steel material contains 70 area% or more of polygonal ferrite and acicular ferrite in total, and 3.5 area% or less of MA phase (martensite-austenite composite phase). By doing so, it has been found that the initiation of fracture at low temperature and the propagation resistance can be remarkably improved, and the present invention has been presented.

That is, the steel material having excellent fracture resistance at low temperature of the present invention is C: 0.005 to 0.07%, Si: 0.005 to 0.3%, Mn: 1.7 to 3 in weight%. .0%, Sol. Al: 0.001 to 0.035%, Nb: 0.02% or less (excluding 0%), V: 0.01% or less (excluding 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01 to 2.0%, Cr: 0.01 to 0.5%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001 to 0.008%, P: 0.02% or less (excluding 0%), S: 0.003% or less (excluding 0%), O: 0.003% It is composed of the following (excluding 0%), the remaining Fe, and unavoidable impurities, and satisfies the above relational expressions 1 and 2. The microstructure of the steel material contains 70 area% or more of polygonal ferrite and acicular ferrite in total, and 3.5 area% or less of MA phase (martensite-austenite composite phase).

First, the alloy composition of the steel material of the present invention and the reasons for limiting the content thereof will be described in detail. Hereinafter, the unit of the content of each element is% by weight.

C: 0.01-0.07%
C is an element that promotes the formation of acicular ferrite or lath bainite and plays an important role in producing cementite, pearlite, etc. to ensure strength. When the C content is less than 0.01%, there is almost no diffusion of C and the transformation occurs relatively quickly, so that the transformation into a coarse ferrite structure may occur and the strength and toughness of the steel material may decrease. be. On the other hand, when the C content exceeds 0.07%, not only cementite and MA phase are excessively formed, but also they are formed coarsely, so that the fracture initiation resistance at low temperature may be significantly deteriorated. There's a problem. Therefore, the content of C is preferably in the range of 0.01 to 0.07%. The content of C is more preferably 0.01 to 0.06%, further preferably 0.01 to 0.05%.

Si: 0.005 to 0.3%
Si is an element that is generally added for the purpose of strengthening solid solution as well as deoxidizing and desulfurizing effects. However, while the effect of increasing the yield strength and tensile strength is small, the stability of austenite is greatly increased in the weld heat-affected zone and the fraction of the MA phase is increased, so that the fracture initiation resistance at low temperature is low. In the present invention, it is preferable to limit it to 0.3% or less because there is a problem that it can be significantly deteriorated. On the other hand, in order to control the Si content to less than 0.005%, there is a problem that the processing time in the steelmaking process increases significantly, the production cost increases, and the productivity is inferior. Therefore, the lower limit of the Si content is set. It is preferably 0.005%. Therefore, the Si content preferably ranges from 0.005 to 0.3%. The Si content is more preferably 0.005 to 0.25%, and even more preferably 0.005 to 0.2%.

Mn: 1.7-3.0%
Since Mn has a large effect of increasing strength by strengthening solid solution and does not significantly decrease toughness at low temperatures, it is added in an amount of 1.7% or more in order to secure sufficiently high strength. However, when Mn is added in an excessive amount, segregation at the center in the thickness direction of the steel sheet becomes severe, and together with the segregated S, the formation of MnS, which is a non-metal inclusion, is promoted. The MnS inclusions formed in the central part are stretched by the subsequent rolling, and the segregated site is likely to form a high-hardness low-temperature structure due to its high hardening ability, and as a result, the fracture initiation and propagation resistance at low temperatures are greatly reduced. Therefore, the upper limit of the Mn content is preferably 3.0%. Therefore, the Mn content is preferably 1.7 to 3.0%. Further, the content of Mn is more preferably 1.7 to 2.8%.

Sol. Al: 0.005 to 0.035%
Sol. Al is used as a powerful deoxidizing agent in the steelmaking process together with Si and Mn, and such an effect can be sufficiently obtained when at least 0.005% or more is added at the time of single or combined deoxidation. can. However, Sol. When the Al content exceeds 0.035%, the above effects are saturated and the fraction of Al ₂ O ₃ in the oxidizing inclusions produced as a result of deoxidation increases more than necessary. Since the size of the inclusions becomes coarse and is difficult to be removed during refining, there arises a problem that the low temperature toughness of the steel material is greatly reduced. Further, as with Si, the formation of the MA phase can be promoted at the weld heat-affected zone, and the fracture initiation and propagation resistance at low temperatures can be significantly reduced. Therefore, Sol. The Al content is preferably 0.005 to 0.035%. Sol. The Al content is more preferably 0.005 to 0.03%, and even more preferably 0.005 to 0.02%.

Nb: 0.02% or less (excluding 0%)
Nb is dissolved in austenite during reheating of the slab to increase the curing ability of austenite, and is precipitated as fine carbonitrides (Nb, Ti) (C, N) during hot rolling during rolling and cooling. It is an element that has a great effect of making the final microstructure finer by suppressing recrystallization. However, when Nb is added in an excessive amount, the curing ability in the weld heat-affected zone is excessively increased to promote the formation of the MA phase, and the fracture initiation at low temperature and the propagation resistance are greatly reduced. In the present invention, the content of Nb is limited to 0.02% or less (excluding 0%). The content of Nb is more preferably 0.015% or less, and further preferably 0.012% or less.

V: 0.01% or less (excluding 0%)
Almost all of V is re-dissolved when the slab is reheated, and most of it is precipitated during cooling after rolling to improve the strength. Promotes the production of. Therefore, in the present invention, the content of V is limited to 0.01% or less (excluding 0%). The content of V is more preferably 0.008% or less, and further preferably 0.005% or less.

Ti: 0.001 to 0.02%
Ti mainly exists as a precipitate of a hexagonal surface in the form of a fine TiN at high temperature, or when added together with Nb or the like, a (Ti, Nb) (C, N) precipitate is formed, and the base metal and the weld heat-affected zone are formed. It has the effect of suppressing the growth of crystal grains. In order to sufficiently secure the above effect, it is preferable to add 0.001% or more of Ti, and in order to maximize the effect, it is preferable to increase it according to the content of N added. On the other hand, when the Ti content exceeds 0.02%, an unnecessarily coarse carbonitride is generated and acts as a starting point of fracture cracks, so that the impact characteristics of the weld heat affected zone are significantly reduced. obtain. Therefore, the Ti content is preferably 0.001 to 0.02%. The Ti content is more preferably 0.001 to 0.017%, even more preferably 0.001 to 0.015%.

Cu: 0.01-1.0%
Cu is an element that can greatly improve the strength by solid solution and precipitation without significantly impairing the fracture initiation and propagation resistance. If the Cu content is less than 0.01%, the above effect is insufficient. On the other hand, if the Cu content exceeds 1.0%, cracks may be induced on the surface of the steel sheet, and Cu is an expensive element, which causes a problem of cost increase. Therefore, the Cu content preferably ranges from 0.01 to 1.0%. The Cu content is more preferably 0.01 to 0.6%, even more preferably 0.01 to 0.4%.

Ni: 0.01-2.0%
Ni has almost no effect of increasing strength, but is effective in starting fracture at low temperature and improving propagation resistance, and in particular, when Cu is added, the surface due to selective oxidation generated when the slab is reheated. It has the effect of suppressing cracks. Further, the addition of Ni has the effect of improving the toughness at a low temperature even if a coarse hard structure is generated in the weld heat affected zone due to a high temperature and a high cooling rate. When the Ni content is less than 0.01%, the above-mentioned effect is insufficient. On the other hand, since Ni is an expensive element, there is a problem of cost increase when its content exceeds 2.0%. Therefore, the Ni content preferably ranges from 0.01 to 2.0%. The Ni content is more preferably 0.2 to 1.8%, and even more preferably 0.3 to 1.2%.

Cr: 0.01-0.5%
Cr has a small effect of increasing yield strength and tensile strength due to solid dissolution, but has a high hardening ability, so that a fine structure is generated in a thick material even at a slow cooling rate, and an effect of improving strength and toughness is achieved. There is. When the Cr content is less than 0.01%, the above effect is insufficient. On the other hand, when the Cr content exceeds 0.5%, not only the cost increases but also the low temperature toughness of the weld heat affected zone can be deteriorated. Therefore, the Cr content preferably ranges from 0.01 to 0.5%. The Cr content is more preferably 0.01 to 0.4%, and even more preferably 0.01 to 0.25%.

Mo: 0.01-0.65%
Mo is an element that has the effect of delaying the phase transformation in the accelerated cooling process and, as a result, greatly increasing the strength, and has the effect of preventing a decrease in toughness due to grain boundary segregation of impurities such as P. When the Mo content is less than 0.01%, the above effect is insufficient. On the other hand, when the Mo content exceeds 0.65%, the high curing ability can promote the formation of the MA phase in the weld heat-affected zone, and can greatly reduce the initiation of fracture at low temperature and the propagation resistance. Therefore, the Mo content preferably ranges from 0.01 to 0.65%. The content of Mo is more preferably 0.01 to 0.5%, further preferably 0.01 to 0.4%.

Ca: 0.0002 to 0.005%
After Al deoxidation, when Ca is added to the molten steel in steelmaking, it binds to S mainly existing as MnS, suppresses the formation of MnS, and forms spherical CaS to form cracks in the center of the steel material. It exerts the effect of suppressing. Therefore, in the present invention, in order to sufficiently form the added S as CaS, 0.0002% or more of Ca should be added. However, when the amount of Ca added is excessive, excess Ca combines with O to form coarse and hard oxidizing inclusions, which are stretched and fractured in the subsequent rolling and act as a crack starting point at low temperature. Will come to do. Therefore, the upper limit of the Ca content is preferably 0.005%. Therefore, the Ca content preferably ranges from 0.0002 to 0.005%. The Ca content is more preferably 0.0005 to 0.003%, even more preferably 0.0005 to 0.0025%.

N: 0.001 to 0.006%
N is an element that forms a precipitate together with the added Nb, Ti, and Al to refine the crystal grains of the steel and improve the strength and toughness of the base metal. However, when excessively added, the excess N exists in an atomic state and is known as the most representative element that causes an aging phenomenon after cold deformation and lowers low temperature toughness. It is also known that when a slab is manufactured by continuous casting, embrittlement at a high temperature promotes cracks on the surface portion. Therefore, in the present invention, considering that the Ti content is 0.001 to 0.02%, the amount of N added is limited to the range of 0.001 to 0.006%. The content of N is more preferably 0.001 to 0.005%, further preferably 0.001 to 0.0045%.

P: 0.02% or less (excluding 0%)
P is an element that plays a role in increasing strength but deteriorates low temperature toughness. In particular, in heat-treated steel, there is a problem that low-temperature toughness is significantly deteriorated due to grain boundary segregation. Therefore, it is preferable to control P as low as possible. However, since it costs a lot to remove P extremely in the steelmaking process, it is limited to 0.02% or less. The content of P is more preferably 0.015% or less, and further preferably 0.012% or less.

S: 0.003% or less (excluding 0%)
S combines with Mn to form MnS inclusions mainly in the central portion of the steel sheet in the thickness direction, which is a main cause of deterioration of low temperature toughness. Therefore, in order to secure the deformation aging impact characteristics at low temperatures, it is preferable to remove S as much as possible in the steelmaking process. However, since it costs too much, it is limited to the range of 0.003% or less. The content of S is more preferably 0.002% or less, and further preferably 0.0015% or less.

O: 0.003% or less (excluding 0%)
O is removed as an oxidizing inclusion by adding a deoxidizing agent such as Si, Mn, and Al in the steelmaking process. If the amount of the deoxidizing agent added and the step of removing the inclusions are not sufficient, the amount of the oxidizing inclusions remaining in the molten steel increases, and the size of the inclusions also greatly increases. The coarse oxidative inclusions not removed in this way will remain inside in the crushed or spherical form during the rolling process in the subsequent steel manufacturing process, and will be the starting point of fracture at low temperature. Or it acts as a propagation path for cracks. Therefore, in order to secure the impact characteristics and CTOD characteristics at low temperature, coarse oxidizing inclusions should be suppressed as much as possible, and therefore, the content of O is limited to 0.003% or less. The content of O is more preferably 0.0025% or less, further preferably 0.0022% or less.

In the present invention, the remaining component is iron (Fe). However, in the normal manufacturing process, unintended elements or impurities may be unavoidably mixed from the raw materials or the surrounding environment, and therefore this cannot be excluded. For example, it may contain 5 ppm or less of boron (B) or the like. All of these impurities are not specifically mentioned herein, as they are well known to any technician in the normal manufacturing process.

Further, the alloy composition of the present invention is required to contain Mn, Ni, Cu, Cr, and Nb so as to satisfy the contents of each of the above elements and the following relational formulas 1 and 2.
[Relational expression 1]
Mn + 0.5x (Ni + Cu) ≧ 2.5 wt%
[Relational expression 2]
Mo + Cr + 1.5xSi + 10xNb ≦ 0.5wt%
(However, in the above relational expressions 1 and 2, each element is a value shown in% by weight.)

Mn, Ni, and Cu forming the above relational expression 1 are typical face-centered cubic metals, and when added to steel materials, they not only increase the strength by solid solution strengthening but also do not significantly impair the toughness even at low temperatures. It is an element. The present inventors designed the relational expression 1 in consideration of the degree of influence of the metal on the strength and toughness of the steel material. At this time, as the value of the relational expression 1 increases, the effect of solid solution strengthening increases, and as a result, the strength of the steel material and the weld heat-affected zone increases. Therefore, in order to obtain sufficient strength, it is preferable to control the value of the relational expression 1 to 2.5 or more.

The relational expression 2 is a formula designed in consideration of the degree of influence of an element that promotes the formation of the MA phase, which is a typical structure that greatly impairs the toughness of the steel material and the weld heat affected zone, and is the relational expression 2. As the value increases, the MA phase fraction increases significantly, resulting in an increase in the ductile-brittle transition temperature, which is the low temperature impact characteristic of steel. That is, as the value of the relational expression 2 increases, the low temperature toughness tends to decrease. Therefore, in order to sufficiently secure the low temperature impact characteristics of the steel material, particularly the CTOD value, it is preferable to control the value of the relational expression 2 to 0.5 or less. The welded part, especially the SC-HAZ (Sub-Critically reheated Heat Affected Zone), which is the most important position for guaranteeing the low temperature CTOD value, has a base material because the temperature at the time of welding is lower than the two-phase temperature. It has a microstructure that is almost similar to the microstructure. Therefore, by controlling the value of the relational expression 2 to 0.5 or less, the low temperature impact characteristic and the CTOD value of the welded portion can be sufficiently secured. The value of the relational expression 2 is more preferably 0.48 or less, and further preferably 0.45 or less.

On the other hand, the microstructure of the steel material of the present invention contains 70 area% or more of polygonal ferrite and acicular ferrite in total, and 3.5 area% or less of MA phase (martensite-austenite composite phase).
Needle-shaped ferrite is the most important and basic microstructure that not only increases the strength due to the fine grain size effect, but also hinders the propagation of cracks generated at low temperatures. Polygonal ferrite is coarser than needle-shaped ferrite, so its contribution to increasing strength is relatively small, but it has a low dislocation density and high-inclined grain boundaries, which greatly contributes to suppressing propagation at low temperatures. It is a fine structure.
When the total of the polygonal ferrite and the needle-shaped ferrite is less than 70 area%, there is a problem that it is difficult to suppress the start and propagation of cracks at a low temperature and it is difficult to secure high strength. Therefore, the total of the polygonal ferrite and the needle-shaped ferrite is preferably 70 area% or more, more preferably 85 area% or more, and further preferably 90 area% or more.

Further, in the present invention, in the polygonal ferrite and the needle-like ferrite, the ratio of the large tilt angle grain boundaries defined that the crystal orientation difference between the crystal grains is 15 ° or more is 40% or more in the total grain boundaries. Further, it is preferable that the length of the large tilt angle crystal grain boundary per unit area is 300 mm / mm ² or more.
Since the MA phase does not accept deformation due to its high hardness, it not only concentrates on the deformation of the soft ferrite matrix around it, but also separates the interface with the surrounding ferrite matrix above the limit point. The MA phase itself is destroyed and acts as a crack starting point. Therefore, the MA phase should be controlled as low as possible because it is the most important cause of deterioration of the low temperature fracture characteristics of the steel material, and it is preferable to control it to 3.5 area% or less.

At this time, in the present invention, it is preferable that the MA phase has an average size of 2.5 μm or less measured with a diameter equivalent to a circle. This is because when the average size of the MA phase exceeds 2.5 μm, the stress is further concentrated, so that the MA phase is easily broken and acts as a crack starting point.
Further, in the present invention, the polygonal ferrite and the needle-like ferrite may not be work-hardened by hot rolling. That is, it is preferable that the polygonal ferrite and the needle-like ferrite are not stretched by hot rolling, but are produced after hot rolling.

The microstructure of the steel material according to the present invention may contain bainitic ferrite, cementite and the like in addition to the above-mentioned polygonal ferrite, acicular ferrite and MA phase.
Bainitic ferrite is a structure transformed at low temperature and has many dislocations inside, but it has the characteristic that it is relatively coarse compared to various ferrites, and it is also inside. It contains the MA phase and is strong, but it is vulnerable to crack initiation and propagation and should be controlled to a minimum.

Further, the steel material of the present invention can contain inclusions having a size of 10 μm or more in a range of 11 pieces / cm ² or less. The size is a size measured with a diameter equivalent to a circle. When the number of inclusions having a size of 10 μm or more exceeds 11 pieces / cm ² , the problem of acting as a crack starting point at a low temperature arises. In order to control such coarse inclusions, it is preferable to add Ca or a Ca alloy in the final stage of the secondary refining, and then bubbling and refluxing with Ar gas for 3 minutes or more.

Further, the steel material of the present invention can have a yield strength of 460 MPa or more, an impact energy value of 300 J or more at −60 ° C., and a CTOD value of 0.2 mm or more at −20 ° C. Further, the steel material of the present invention can have a tensile strength of 570 MPa or more. The steel material of the present invention can have a DBTT (ductility-brittle transition temperature) of −80 ° C. or lower.
Next, a method for producing a high-strength steel material having excellent fracture initiation and propagation resistance at a low temperature of the present invention will be described.

The method for producing a steel material of the present invention includes a step of preparing a steel slab satisfying the above alloy composition, a step of heating the steel slab to 1000 to 1200 ° C., and a step of heating the heated slab in a temperature range of 650 ° C. or higher. It includes a step of hot-rolling for finishing and a step of cooling the hot-rolled hot-rolled steel sheet to a cooling end temperature of 200 to 550 ° C. at a cooling rate of 2 to 30 ° C./s.

Steel slab preparation stage Prepare a steel slab that meets the above alloy composition.
At this time, in the present invention, when preparing the steel slab, in the final stage of the secondary refining of the molten steel, the step of adding Ca or Ca alloy to the molten steel and the step of adding the Ca or Ca alloy to the molten steel for 3 minutes or more. It is preferable to carry out a step including a step of bubbling and reflux treatment with Ar gas. This is to control coarse inclusions.

Steel slab heating stage The steel slab is heated to 1000-1200 ° C.
When the heating temperature of the slab is less than 1000 ° C., it is difficult to re-dissolve the carbides generated in the slab during continuous casting, and the homogenization treatment of the segregated elements becomes insufficient. Therefore, it is preferable to heat to 1000 ° C. or higher, which is a temperature at which 50% or more of the added Nb can be re-dissolved.

On the other hand, when the heating temperature of the slab exceeds 1200 ° C., the size of the austenite crystal grains may grow excessively coarsely, and the subsequent rolling cannot sufficiently miniaturize the austenite crystal grains. Mechanical properties such as low temperature toughness can be significantly reduced.
The heating temperature of the steel slab is more preferably 1000 to 1160 ° C, even more preferably 1000 to 1140 ° C.

Hot-rolling step The heated slab is finished and hot-rolled at a bainite formation start temperature of 650 ° C. or higher to obtain a hot-rolled steel sheet.
When the temperature of the finish hot rolling is less than 650 ° C, coarse bainite is generated and work-hardened during rolling to increase the strength more than necessary, and conversely, the impact toughness at low temperature is greatly reduced. Therefore, it is preferable to limit the rolling end temperature to 650 ° C. or higher. That is, when the temperature of hot rolling is low, coarse proeutectoid ferrite is generated before hot rolling finish, and it is stretched and work-hardened by subsequent rolling, and the remaining austenite remains in a band shape. At the same time, the MA hardened phase is transformed into a structure with a high density, and the low temperature toughness is lowered.

Further, in the present invention, by setting the total reduction rate in the unrecrystallized region temperature section to 30% or more (excluding the reduction rate in the recrystallized region), sufficient deformation energy is accumulated in austenite, and the austenite is subsequently subjected to. It is preferable to sufficiently generate polygonal and acicular ferrite which are advantageous for low temperature toughness at the time of transformation, and to secure the ratio and density of large tilt angle grain boundaries.
The reduction rate is more preferably 40% or more, further preferably 45% or more.

Cooling stage Then, in the present invention, the hot-rolled hot-rolled steel sheet is cooled.
At this time, it is preferable to cool the hot-rolled steel sheet at a cooling rate of 2 to 30 ° C./s to a cooling end temperature of 200 to 550 ° C. When the cooling rate is less than 2 ° C / s, the cooling rate is too slow to avoid coarse ferrite, pearlite, and bainite transformation sections, and the strength and low temperature toughness deteriorate, resulting in 30 ° C / s. If it exceeds, granular bainite or martensite is formed and the strength is increased, but the low temperature toughness may be significantly deteriorated.

When the cooling end temperature exceeds 550 ° C., fine structures such as needle-like ferrite are unlikely to be formed, and coarse bainite or pearlite is likely to be formed. On the other hand, when the cooling end temperature is less than 200 ° C., there is no problem in terms of microstructure, but there is a problem that the cooling takes an excessive time and the productivity is significantly inferior.
The cooling end temperature is more preferably 200 to 500 ° C, further preferably 200 to 450 ° C.

On the other hand, in the present invention, if necessary, the cooled hot-rolled steel sheet is heated to 450 to 650 ° C., maintained for (1.3 × t + 5) to (1.3 × t + 200) minutes, and then cooled. Further tempering steps may be included [where t is a value measured in mm for the thickness of the hot-rolled steel sheet]. This decomposes MA and martensite when MA and martensite are excessively generated, removes the high dislocation density inside, and precipitates a small amount of solid-dissolved Nb and the like as carbonitride. This is to further improve the yield strength or low temperature toughness.

However, when the heating temperature is less than 450 ° C., the ferrite matrix is not sufficiently softened and embrittlement due to P segregation or the like occurs, so that the toughness may be deteriorated. On the other hand, when the heating temperature exceeds 650 ° C, the recovery and growth of crystal grains occur rapidly, and when the temperature becomes higher, austenite is partially reverse-transformed, the yield strength is rather lowered, and the low temperature toughness is reduced. May get worse.
When the maintenance time is less than (1.3 × t + 5) minutes, the tissue is not sufficiently homogenized, and when it exceeds (1.3 × t + 200) minutes, the productivity is lowered. There is a problem.

Hereinafter, the present invention will be described in more detail with reference to examples.
(Example)
A slab having the composition shown in Table 1 below was heated, hot-rolled, and cooled under the conditions shown in Table 2 below to produce a steel material.
The microstructure of the manufactured steel material was observed, the physical properties were measured, and the results are shown in Table 3 below.

Further, after welding the manufactured steel material, the CTOD value (-20 ° C.) of the weld heat-affected zone (SCHAZ) was measured and shown in Table 3 below. Since the CTOD value (-20 ° C) of the steel material is higher than that of the weld heat affected zone, the CTOD value (-20 ° C) of the steel material was not measured separately.
At this time, the microstructure of the steel material is mirror-polished on the cross section of the manufactured steel material and then etched with Nital or LePera depending on the purpose, and a certain area of the test piece is optically or scanned with an electron microscope. The image was measured at a magnification of 100 to 5000 times. The fraction of each phase was measured from the measured image using an image analyzer. In order to obtain a statistically significant value, the same test piece was repeatedly measured while changing its position, and the average value was calculated.

Further, in order to observe the characteristics of the manufactured structure in more detail, EBSD (Electron Backscatter Diffraction) measurement was performed on the test piece etched by Nital with a scanning electron microscope, and the grain boundaries of the manufactured steel material were measured. The characteristics of the were quantitatively measured.
The physical properties of the steel material are described by measuring from the nominal strain-nominal stress curve obtained by a normal tensile test.
The impact energy value (-60 ° C.) of the weld heat-affected zone was measured by performing a Charpy V-notch impact test.

The CTOD value (-20 ° C) is based on the BS7448 standard, and the test piece is processed to a size of B (thickness) x B (width) x 5B (length) perpendicular to the rolling direction, and the length of the fatigue crack is tested. After fatigue cracking was made to be about 50% of the width of the piece, a CTOD test was performed at −20 ° C. Here, B is the thickness of the manufactured steel material.

The Kca value was tested three times by the ESSO test method, and the graph of the crack propagation stop temperature and the K value measured in each test was obtained, and the K value (Kca: arrest K) when the temperature was -10 degrees. ). Further, CAT (Crack arrest temperature) measured NDTT (Nil-ductility transition temperature) from the NRL test, and obtained this from the value calculated by the conversion formula of Equation 1. Here, B indicates the thickness of the steel material.

＊前記表３において、フェライト系は、ポリゴナルフェライトと針状フェライトの合計を意味する。

* In Table 3 above, the ferrite system means the total of polygonal ferrite and acicular ferrite.

As shown in Tables 1 to 3, Invention Examples 1 to 4 satisfying both the alloy composition and the production conditions presented in the present invention have low temperatures when the yield strength, tensile strength, impact energy value, Kca, CAT and the like are taken into consideration. It can be confirmed that the fracture resistance is excellent and the CTOD value at the weld heat affected zone is also high. In particular, as shown in FIG. 1, the Kca value measured in Example 1 of the present invention is much higher than the required value of 8000. It can be confirmed that such excellent strength and low temperature toughness properties are also the result obtained from the sufficiently produced fine polygonal and acicular ferrite structures, as shown in FIG.

On the other hand, Comparative Example 1 is a case where the content of C exceeds the range of the present invention, and the added C is the most potent element that promotes granular bainite and MA. Therefore, due to the excessive addition of C, the fraction of ferrite, which is advantageous for toughness, is greatly reduced, and the strength of the base metal is high, but the low temperature toughness such as the impact energy value deteriorates, and in particular, the CTOD value of the weld heat affected zone becomes high. It dropped significantly.
Comparative Example 2 is a case where the content of the added Mn exceeds the range of the present invention. In this case, since the Mn content is high, the probability of segregation of the central portion of the steel material is greatly increased, the impact energy at the central portion in the thickness direction of the steel material is greatly deteriorated, and the segregation of the central portion is also performed at the weld heat affected zone. In the band, a hardened structure with extremely high hardness was partially generated, an early fracture phenomenon (pop-in) appeared, and the CTOD value was greatly reduced.

Comparative Example 3 is a case where the content of Nb, which is generally widely used for improving the strength and miniaturizing the structure, exceeds the range of the present invention. In general, the addition of Nb is advantageous for refining the structure and increasing both strength and toughness, but when added more than necessary, it suppresses the formation of polygonal and acicular ferrite which are advantageous for toughness, and granular bainite or the like. Acts to promote the tissue of. Therefore, the ratio and density of the large tilt angle grain boundaries of 15 ° or more, which is advantageous for suppressing crack propagation, are greatly reduced, and crack propagation is relatively facilitated. As a result, as shown in Table 3, the Kca value measured in Comparative Example 3 is at the level of 5860, which is far below the required value of 8000. In addition, the weld heat-affected zone greatly promoted the formation of MA structure, which has a particularly disadvantageous effect on low-temperature toughness, and as a result, had the effect of greatly reducing CTOD.

In Comparative Examples 4, 5, and 6, the range of the content of each element satisfies the range of the present invention, but the values of the relational expression 1 and the relational expression 2 are out of the range of the invention, and the strength is low. It can be seen that the low temperature toughness is greatly reduced.
Specifically, Comparative Example 4 is a case where the relational expression 1 composed of components advantageous for improving low temperature toughness is satisfied, but the relational expression 2 composed of components impairing low temperature toughness exceeds the scope of the present invention. .. As a result, although the strength is sufficiently high, the impact energy value in the base metal and the CTOD value in the weld heat affected zone are deteriorated.

Further, in Comparative Example 5, the relational expression 2 satisfies the scope of the invention, but the relational expression 1 is out of the scope of the present invention, and the amount of the component added is required to secure the strength of the steel material as a whole. There was a shortage, and the strength of the base metal was greatly reduced.
Comparative Example 6 is a case where both the relational expression 1 and the relational expression 2 are out of the scope of the invention. That is, there is a case where the components advantageous for low temperature toughness are insufficient, the components disadvantageous for low temperature toughness are exceeded, and the characteristic values of low temperature toughness are all deteriorated.

Comparative Example 7 is a case where all the components of the steel material satisfy the scope of the invention, but the total rolling reduction amount of unrecrystallized area rolling does not satisfy the scope of the invention in the manufacturing process of the steel material. That is, the amount of reduction in the unrecrystallized region is insufficient, the fraction of ferrite that hinders the propagation of cracks in the fine structure of the steel material is low, and the ratio and density of the large grain boundaries are greatly reduced. The characteristic value of low temperature toughness was not good.
Further, in Comparative Example 8, all the components of the steel material satisfy the scope of the invention, but in the steel material manufacturing process, the cooling rate is slow because it is manufactured by air cooling without applying accelerated cooling after controlled rolling. As a result, ferrite, which is advantageous for low-temperature toughness, is sufficiently produced but coarse, and the strength is greatly reduced.

Although the embodiments of the present invention have been described in detail above, the technical scope of the present invention is not limited to this, and various modifications are made within the range not deviating from the technical idea of the present invention described in the claims. And the possibility of modification is obvious to those with ordinary knowledge in the art.

Claims (9)

By weight%, C: 0.005 to 0.07%, Si: 0.005 to 0.3%, Mn: 1.7 to 3.0%, Sol. Al: 0.001 to 0.035%, Nb: 0.02% or less (excluding 0%), V: 0.01% or less (excluding 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01 to 2.0%, Cr: 0.01 to 0.5%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001 to 0.008%, P: 0.02% or less (excluding 0%), S: 0.003% or less (excluding 0%), O: 0.003% It consists of the following (excluding 0%), the remaining Fe, and unavoidable impurities, and satisfies the following relational expression 1 and the following relational expression 2.
Its fine structure contains 70 area% or more of polygonal ferrite and acicular ferrite in total, and has a fracture resistance property at low temperature characterized by containing 3.5 area% or less of the MA phase (martensite-austenite composite phase). Excellent high-strength steel material.
[Relational formula 1] Mn + 0.5x (Ni + Cu) ≧ 2.5 wt%
[Relational formula 2] Mo + Cr + 1.5xSi + 10xNb ≦ 0.5 wt%
(However, in the above relational expressions 1 and 2, each element is a value shown in% by weight.) In the polygonal ferrite and the acicular ferrite, the ratio of the large tilt angle grain boundaries defined as the crystal orientation difference between the crystal grains is 15 ° or more is 40% or more in the total grain boundaries, and the unit area is The high-strength steel material having excellent fracture resistance at low temperatures according to claim 1, wherein the length of the grain boundaries with a large tilt angle per hit is 300 mm / mm ² or more. The yield strength is 460 MPa or more, the impact energy value at -60 ° C is 250 J or more, and the Kca value measured in the ESSO test is 8000 N / mm ^3/2 or more, or NDTT (Nil) measured in the NRL test. The high-strength steel material having excellent fracture resistance at low temperatures according to claim 1, wherein the CAT (crack arrest temperature) calculated from the ductility measurement temperature is less than −10 ° C. The high-strength steel material having excellent fracture resistance at low temperatures according to claim 1, wherein the tensile strength is 570 MPa or more and the DBTT (ductility-brittle transition temperature) is −80 ° C. or lower. The high-strength steel material having excellent fracture resistance at low temperatures according to claim 1, wherein inclusions having a size of 10 μm or more measured in a circle equivalent diameter are contained in a range of 11 pieces / cm ² or less. By weight%, C: 0.005 to 0.07%, Si: 0.005 to 0.3%, Mn: 1.7 to 3.0%, Sol. Al: 0.001 to 0.035%, Nb: 0.02% or less (excluding 0%), V: 0.01% or less (excluding 0%), Ti: 0.001 to 0.02%, Cu: 0.01 to 1.0%, Ni: 0.01 to 2.0%, Cr: 0.01 to 0.5%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001 to 0.008%, P: 0.02% or less (excluding 0%), S: 0.003% or less (excluding 0%), O: 0.003% Below (excluding 0%), the stage of preparing a steel slab consisting of the remaining Fe and unavoidable impurities and satisfying the following relational expressions 1 and 2 and
The step of heating the steel slab to 1000-1200 ° C.
The step of finishing and hot rolling the heated slab in a temperature range of 650 ° C. or higher so that the total reduction rate in the unrecrystallized region temperature section is 30% or more (excluding the reduction rate in the recrystallization region). When,
The step of producing a steel material by cooling the hot-rolled hot-rolled steel sheet to a cooling end temperature of 200 to 550 ° C. at a cooling rate of 2 to 30 ° C./s is included.
The cooled steel material is characterized in that its fine structure contains 70 area% or more of polygonal ferrite and acicular ferrite in total, and 3.5 area% or less of MA phase (martensite-austenite composite phase). A method for manufacturing high-strength steel with excellent fracture resistance at low temperatures.
[Relational formula 1] Mn + 0.5x (Ni + Cu) ≧ 2.5 wt%
[Relational formula 2] Mo + Cr + 1.5xSi + 10xNb ≦ 0.5 wt%
(However, in the above relational expressions 1 and 2, each element is a value shown in% by weight.) A claim comprising a tempering step in which the cooled hot-rolled steel sheet is heated to 450 to 650 ° C., maintained for (1.3 × t + 5) minutes to (1.3 × t + 200) minutes, and then cooled. Item 6. The method for producing a high-strength steel material having excellent fracture resistance at low temperatures. At the stage of preparing the steel slab,
At the final stage of secondary refining of molten steel, the stage of adding Ca or Ca alloy to the molten steel, and
The excellent fracture resistance at low temperature according to claim 6, wherein a step including a step of bubbling and reflux treatment with Ar gas is performed for at least 3 minutes after the Ca or Ca alloy is charged. A method for manufacturing high-strength steel materials. In the polygonal ferrite and the needle-like ferrite, the ratio of the large tilt angle grain boundaries defined as the crystal orientation difference between the crystal grains is 15 ° or more is 40% or more in the total grain boundaries, and the unit area. The method for producing a high-strength steel material having excellent fracture resistance at low temperatures according to claim 6, wherein the length of the grain boundaries with a large tilt angle per hit is 300 mm / mm ² or more.

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