KR102045641B1 - High strength steel for arctic environment having excellent resistance to fracture in low temperature, and method for manufacturing the same - Google Patents

Abstract

Provided are high strength steels for polar environments having excellent fracture resistance at low temperatures, and methods for manufacturing the same.
The high-strength steel having excellent fracture resistance at low temperatures of the present invention 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: 0.001-0.5%, Ca: 0.0002-0.005%, N: 0.001-0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% It contains below (excluding 0%), balance Fe and inevitable impurities, satisfies relations 1 and 2, and its microstructure contains more than 70 area% of total polygonal ferrite and acicular ferrite, (Martensite-austenite composite phase) is contained in 3.5 area% or less.

Description

High-strength steel for polar environments with excellent fracture resistance at low temperatures and manufacturing method thereof

The present invention relates to a high-strength steel for polar environments having excellent fracture resistance at low temperatures that can be preferably applied to shipbuilding, marine structural steel and a method of manufacturing the same.

Due to global warming, ice in the Arctic region is gradually decreasing, and as a result, interest in the Arctic sea route connecting Europe and East Asia is greatly increased. In recent years, cargo ships have been tested only during the summer months, and time and cost are reported to be reduced by up to 30% or more compared to existing routes through Southeast Asia. Furthermore, within 20-30 years, if the Arctic ice is completely gone, even straight routes across the North Pole are expected. Accordingly, the need for ships passing through the Arctic region is gradually becoming a reality, and the need for design of safe ships and required steel materials for polar environments in such polar environments is gradually increasing.

Existing structural steels are vulnerable to destruction in polar environments, that is, environments exposed to impacts of low temperatures up to -60 ° C and drift ice, and thus require new high-strength steels for polar environments with excellent fracture resistance at low temperatures. Do.

In general, the thick, high-strength steel used in large ships or petroleum mining platform is vulnerable to breakdown at low temperatures. High-strength ultra-thick steels have a high amount of alloying elements such as Mn and Mo to secure strength, and are coarse due to low rolling reduction rate and slow accelerated cooling rate in the production of ultra-thick steels. Granular bainite This is because hard phase tissues such as B and MA are easily formed. Due to this microstructure, the steel has a very weak characteristic of resistance to fracture at low temperatures. Therefore, in order to have high strength and excellent low temperature fracture characteristics of the ultra-thick material, it is necessary to minimize the microstructure of the tissue and to harden hard tissues such as granular bainite or M-A.

In order to solve the above-mentioned problems, ① the slab reheat temperature is drastically lowered and control rolling is carried out at low temperature, or ② the Cu is tempered at low temperature by adding 1% or more to improve the strength with fine Cu precipitates, or ③ In order to improve the low temperature toughness of the hard granular bainite and the like, a large amount of Ni is added, and ④ minimizing cognate elements such as C is used to dramatically reduce the ④M-A structure. However, structures such as ships are gradually enlarged, and as the use environment is changed to a polar environment, there is a problem that it is difficult to sufficiently secure the initiation of destruction and propagation resistance at low temperatures by simply applying the above-described conventional method.

Therefore, there is a demand for development of a high-strength steel and a method of manufacturing the same having improved fracture start and propagation resistance at low temperatures.

Korean Unexamined Patent Publication No. 2002-0028203

Accordingly, an object of the present invention is to provide a high-strength steel having excellent breakdown resistance and propagation resistance at low temperatures, and a method of manufacturing the same.

In addition, the subject of this invention is not limited to the content mentioned above. The problem of the present invention will be understood from the general contents of the present specification, those skilled in the art will have no difficulty understanding the additional problem of the present invention.

The present invention for achieving the above object,

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: 0.001-0.02%, Cu: 0.01-1.0%, Ni: 0.01-2.0%, Cr: 0.01-0.5%, Mo: 0.001-0.5%, Ca: 0.0002-0.005%, N: 0.001% to 0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% or less (except 0%), balance Fe and inevitable impurities Satisfying the following Expressions 1 and 2,

The microstructure contains a total of 70 area% or more of polygonal ferrite and acicular ferrite, and has high fracture resistance at low temperatures including MA phase (martensite-austenite composite phase) of 3.5 area% or less. It's about steel. .

[Relationship 1]

Mn + 0.5 x (Ni + Cu) ≥ 2.5wt%

[Relationship 2]

Equation 2: Mo + Cr + 1.5 x Si + 10 x Nb ≤ 0.5wt%

(However, in the above formula 1-2, each element is a value expressed in weight%.)

In addition, the present invention,

Providing a steel slab that satisfies the alloy composition described above;

Heating the steel slab to 1000-1200 ° C .;

Finishing hot rolling the heated slab in a temperature range of 650 ° C. or higher so that the total reduction ratio in the non-recrystallization temperature range is 30% or more (excluding the recrystallization reduction ratio); And

It relates to a method of manufacturing a high strength steel having excellent fracture resistance at low temperatures, including the step of cooling the finished hot-rolled hot rolled steel sheet to a cooling end temperature of 200 ~ 550 ℃ at a cooling rate of 2 ~ 30 ℃ / s. .

According to the present invention, it is possible to effectively provide a steel material in which breakdown initiation and propagation resistance at low temperatures are dramatically improved.

1 is a graph showing the measurement of the Kca value for the steel in Example 1 in the present embodiment.
Figure 2 is a microstructure photograph of the steel of Inventive Example 3 in the present embodiment.

Hereinafter, preferred embodiments of the present invention will be described. However, embodiments of the present invention may be modified in various other forms, and the scope of the present invention is not limited to the embodiments described below. In addition, the embodiments of the present invention are provided to more completely explain the present invention to those skilled in the art.

The present inventors have conducted research and experiments to further improve fracture initiation and propagation resistance at low temperatures, and as a result, minimize the addition of alloying elements such as alloying elements, in particular, C, Mo, Cr, and Nb, which produce carbide At the same time, it was confirmed that it is necessary to precisely control in the direction of increasing the amount of the alloying element, which has the effect of simultaneously improving the strength and toughness of the ferrite matrix. And by controlling in this way, the steel microstructure includes more than 70 area% of polygonal ferrite and acicular ferrite in total, and contains less than 3.5 area% of MA phase (martensite-austenite composite phase). Accordingly, the present invention has been found to significantly improve the initiation of breakdown and propagation resistance at low temperatures.

In other words, the steel material having excellent fracture resistance at low temperatures of the present invention, in 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: 0.001-0.02%, Cu: 0.01-1.0%, Ni: 0.01-2.0%, Cr: 0.01- 0.5%, Mo: 0.001-0.5%, Ca: 0.0002-0.005%, N: 0.001-0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% or less (excluding 0%), the balance Fe and unavoidable impurities, and satisfy the above-described Equation 1-2. The steel microstructure includes a total of 70 area% or more of polygonal ferrite and acicular ferrite, and the MA phase (martensite-austenite composite phase) of 3.5 area% or less.

Hereinafter, first, the alloy composition of the steel of the present invention and the content limitation reason will be described in detail. Hereinafter, the unit of each element content is weight%.

C: 0.01 ~ 0.07%

C is an element that promotes formation of acicular ferrite or lath bainite, and plays an important role in securing strength by producing cementite or pearlite.

If the C content is less than 0.01%, there is almost no diffusion of C, so that transformation occurs relatively quickly, and thus, transformation into a coarse ferrite structure may reduce the strength and toughness of the steel. On the other hand, when the C content is more than 0.07%, not only the cementite or MA phase is excessively formed, but also coarse to be formed, which greatly deteriorates the resistance to breakdown at low temperatures.

Si: 0.005-0.3%

Si is generally an element added for the purpose of solid solution strengthening along with deoxidation and desulfurization effects. However, the effect of increasing yield and tensile strength is insignificant, while the stability of austenite in welding heat affected zones is increased to increase the fraction of MA phase, which significantly degrades the initiation resistance at low temperatures. In the invention, it is preferable to limit the amount to 0.3% or less. On the other hand, in order to control the Si content to less than 0.005%, the processing time in the steelmaking process is greatly increased, the production cost increases, and there is a problem that the productivity is lowered, so the lower limit of the Si content is preferably 0.005%.

Mn: 1.7-3.0%

Mn has a large effect of increasing strength due to solid solution strengthening and does not have a large decrease in toughness at low temperatures. Therefore, Mn is added at least 1.7% to secure sufficient high strength.

However, when Mn is excessively added, the segregation becomes severe at the center of the thickness direction of the steel sheet, and at the same time, it promotes the formation of the non-metallic inclusion MnS together with the segregated S. The MnS inclusions formed in the center are stretched by the subsequent rolling, and the segregation sites easily form high-temperature, low-temperature structures due to high hardenability, and consequently greatly lower the initiation of breakage and propagation resistance at low temperatures, thereby increasing the upper limit of the Mn content. Is preferably 3.0%.

Therefore, the Mn content is preferably 1.7 to 3.0%.

Sol.Al: 0.005 ~ 0.035%

Sol.Al, together with Si and Mn, is used as a strong deoxidizer in the steelmaking process, and at least 0.005% must be added at the time of single or complex deoxidation to obtain this effect.

However, when the Sol.Al content is more than 0.035%, the above-mentioned effect is saturated, and the fraction of Al ₂ O ₃ in the oxidative inclusions resulting from the deoxidation increases more than necessary, so that the size of the inclusions becomes coarse and during refining. It is difficult to remove the problem, which greatly reduces the low-temperature toughness of the steel. In addition, similar to Si, the formation of the MA phase in the weld heat affected zone can be promoted, thereby significantly reducing the onset of breakdown and propagation resistance at low temperatures.

Therefore, the content of Sol.Al is preferably 0.005 to 0.035%.

Nb: 0.02% or less (excluding 0%)

Nb is dissolved in austenite during slab reheating to increase the hardenability of austenite, and precipitated as fine carbonitrides (Nb, Ti) (C, N) during hot rolling to suppress recrystallization during rolling or cooling, resulting in final microstructure It is an element that is very effective in making fine. However, when Nb is added in an excessively large amount, since the hardenability in the weld heat affected zone is excessively increased, the formation of the MA phase is accelerated, and the breakdown at low temperatures and the propagation resistance are greatly reduced. Therefore, the Nb content in the present invention is 0.02% or less ( Limit to 0%).

V: 0.01% or less (except 0%)

When V reheats the slab, almost all of them are re-used and precipitated during cooling after rolling to improve the strength. However, in the heat affected zone, the V melts at a high temperature to greatly increase the hardenability to promote the formation of the MA phase. Therefore, the V content in the present invention is limited to 0.01% or less (excluding 0%).

Ti: 0.001-0.02%

Ti is present as a fine TiN-shaped hexagonal precipitate mainly at high temperatures, or when added together with Nb to form (Ti, Nb) (C, N) precipitates, thereby suppressing grain growth of the base metal and the weld heat affected zone. .

In order to sufficiently secure the above-mentioned effect, it is preferable to add Ti to 0.001% or more, and to maximize the effect, it is preferable to increase it in accordance with the amount of N added. On the other hand, when the Ti content is more than 0.02%, coarse carbonitride is produced more than necessary to act as a starting point of fracture cracking, rather it can greatly reduce the impact characteristics of the weld heat affected zone. Therefore, the Ti content is preferably 0.001 to 0.02%.

Cu: 0.01 ~ 1.0%

Cu is an element that can greatly improve the strength by solid solution and precipitation without significantly deteriorating the breakdown initiation and propagation resistance.

When the Cu content is less than 0.01%, the above effects are insufficient. On the other hand, if the Cu content is more than 1.0% may cause cracks on the surface of the steel sheet, Cu is an expensive element, the problem of cost increase occurs.

Ni: 0.01 ~ 2.0%

Ni has little effect of increasing strength, but is effective in initiating breakdown and propagation resistance at low temperatures, and in particular, when Cu is added, it has an effect of suppressing surface cracks due to selective oxidation generated when the slab is reheated. In addition, even if the weld heat affected zone is formed of coarse hard tissue by high temperature and fast cooling rate by Ni addition, there is an effect of improving toughness at low temperature.

When the Ni content is less than 0.01%, the above effects are insufficient. On the other hand, Ni is an expensive element and when the content is more than 2.0%, there is a problem of cost increase.

Cr: 0.01 ~ 0.5%

Cr has a small effect of increasing yield and tensile strength by solid solution, but due to high hardenability, the material is made to have a fine structure even at a slow cooling rate, thereby improving strength and toughness.

When the Cr content is less than 0.01%, the above effects are insufficient. On the other hand, when the Cr content is more than 0.5%, not only the cost increases but also the low temperature toughness of the weld heat affected zone may be inferior.

Mo: 0.01 ~ 0.65%

Mo is an element having the effect of delaying the phase transformation in the accelerated cooling process, resulting in a large increase in strength, and preventing the fall of toughness due to grain boundary segregation of impurities such as P.

If the Mo content is less than 0.01%, the above effects are insufficient. On the other hand, when the Mo content is more than 0.65%, due to the high hardenability, it is possible to promote the formation of the MA phase in the weld heat affected zone, thereby greatly reducing the onset of breakdown and propagation resistance at low temperatures.

Ca: 0.0002-0.005%

When Al is deoxidized and added to molten steel during steelmaking, it is combined with S, which is mainly present as MnS, to suppress MnS formation and to form spherical CaS, thereby exhibiting an effect of suppressing cracks in the center of steel. Therefore, in the present invention, Ca must be added in an amount of 0.0002% or more in order to sufficiently form the added S into CaS.

However, when the amount of Ca added is excessive, excess Ca combines with O to form coarse and hard oxidative inclusions, which are stretched and fractured in subsequent rolling to act as a crack initiation point at low temperatures. Therefore, the upper limit of the Ca content is preferably 0.005%.

N: 0.001-0.006%

N is an element which forms a precipitate together with the added Nb, Ti, and Al to refine the grains of the steel to improve the strength and toughness of the base metal. However, when excessively added, it is known as the most representative element that exists in an excess of atomic state, causing aging after cold deformation, thereby reducing low-temperature toughness. It is also known to promote surface cracks due to embrittlement at high temperatures in slab production by continuous casting.

Therefore, in the present invention, considering that the Ti content is 0.001% to 0.02%, the amount of N is limited to 0.001% to 0.006%.

P: 0.02% or less (except 0%)

P serves to increase strength but is inferior to low temperature toughness. In particular, there is a problem of greatly inferior low-temperature toughness due to grain boundary segregation in heat-treated steel. Therefore, it is desirable to control P as low as possible. However, the excessive removal of P in the steelmaking process is expensive, so it is limited to 0.02% or less.

S: 0.003% or less (except 0%)

S is a major cause of inferior low temperature toughness by combining with Mn to form MnS inclusions mainly in the thickness direction center of the steel sheet. Therefore, in order to secure the strain aging impact characteristics at low temperatures, it is desirable to remove S as much as possible in the steelmaking process. However, it may be excessive cost, so limit to 0.003% or less.

O: 0.003% or less (except 0%)

O is removed and made into an oxidative inclusion by addition of deoxidizers such as Si, Mn, and Al during steelmaking. If the addition amount of the deoxidizer and the inclusion removal process are insufficient, the amount of oxidative inclusions remaining in the molten steel increases, and at the same time, the size of the inclusions is greatly increased. The coarse oxidative inclusions not removed in this way remain in the form of spherical or spherical form during the rolling process in the steel manufacturing process, and serve as a starting point of breakdown at low temperatures or a propagation path of cracks. Therefore, in order to secure impact characteristics and CTOD characteristics at low temperatures, coarse oxidative inclusions should be suppressed as much as possible, and for this purpose, the O content is limited to 0.003% or less.

The remaining component of the present invention is iron (Fe). However, in the usual manufacturing process, elements or impurities which are not intended from the raw materials or the surrounding environment may be inevitably mixed, and thus may not be excluded. For example, boron (B) etc. may be contained in 5 ppm or less. Since these impurities are known to those skilled in the art, all of them are not specifically mentioned in the present specification.

In addition, the alloy composition of the present invention is required to contain Mn, Ni, Cu, Cr and Nb so as to not only satisfy the above-described respective element content but also satisfy the following relation 1-2.

[Relationship 1]

Mn + 0.5 x (Ni + Cu) ≥ 2.5wt%

[Relationship 2]

Equation 2: Mo + Cr + 1.5 x Si + 10 x Nb ≤ 0.5wt%

(However, in the above formula 1-2, each element is a value expressed in weight%.)

Mn, Ni, and Cu forming the above relation 1 are representative face-centered cubic metals, which are elements which not only increase the strength by solid solution strengthening when added to steel materials, but also do not significantly deteriorate toughness even at low temperatures. The present inventors designed relation 1 in consideration of the influence on the steel strength and toughness of the elements, and as the value of relation 1 increases, the solid solution strengthening effect increases, and thus the strength of the steel and the weld heat affected zone increases. Therefore, in order to obtain sufficient strength, it is preferable to control the value of the said relational expression to 2.5 or more.

Equation 2 is designed in consideration of the influence of the element to promote the formation of the MA phase, which is a representative structure that greatly degrades the toughness of the steel and welded heat affected zone, the fraction of the MA phase increases significantly with the increase of the relationship 2 value in the end of the steel The ductile-brittle transition temperature, which is a low temperature impact characteristic, increases. In other words, as the relation 2 increases, low temperature toughness tends to decrease. Therefore, in order to sufficiently secure the low-temperature impact characteristics, in particular the CTOD value of the steel, it is preferable to control the value of the above relation 2 to 0.5 or less.

Also, SC-HAZ (Sub-Critically Reheated Heat Affected Zone), which is the most important position to guarantee the low temperature CTOD value of welded part, has a microstructure almost similar to the microstructure of the base material because the temperature is lower than the ideal zone temperature during welding. By controlling the value of the relational expression 2 below 0.5, the low temperature impact characteristics and the CTOD value of the welded portion can also be sufficiently secured.

In the present invention, the microstructure of the steel of the present invention contains not less than 70 area% of polygonal ferrite and acicular ferrite in total, and less than 3.5 area% of MA phase (martensite-austenite composite phase).

Needle-shaped ferrite not only increases the strength due to the fine grain size effect, but also is the most important and basic microstructure to prevent the propagation of cracks occurring at low temperatures. Polygonal ferrite is a microstructure that contributes to the suppression of propagation at low temperature because it is relatively smaller than the needle-like ferrite, the contribution to the increase in strength is relatively small, but has a low dislocation density and a high angle grain boundary.

When the sum of the polygonal ferrite and the needle-like ferrite is less than 70 area%, it is difficult to suppress the initiation and propagation of cracks at low temperatures, and it is difficult to secure high strength. Therefore, it is preferable that the sum total of polygonal ferrite and acicular ferrite is 70 area% or more, More preferably, it is 85 area% or more, More preferably, it is 90 area% or more.

In the present invention, the polygonal ferrite and the needle-like ferrite have a ratio of grain boundaries of large diameter at which the grain orientation difference between grains is 15 ° or more is 40% or more of the total grain boundaries, and the length of the large diameter grain boundaries per unit area is 300 mm. It is preferable that it is / mm <^2> or more.

In addition, the MA phase does not accept deformation due to its high hardness, thereby concentrating deformation of the soft ferrite matrix around it, and above the threshold, the interface with the surrounding ferrite matrix is separated or the MA phase itself is destroyed to start the crack initiation point. Acts as Therefore, since it is the most important cause of deterioration of low temperature fracture characteristics of the steel, the MA phase should be controlled as low as possible, and it is preferable to control it to 3.5 area% or less.

At this time, in the present invention, the MA phase may have an average size of 2.5 μm or less as measured by a circular equivalent diameter. This is because when the average size of the MA phase is more than 2.5 μm, the stress is more concentrated, and thus the MA phase is easily broken and acts as a starting point of cracking.

In addition, in the present invention, the polygonal ferrite and the needle-like ferrite may not be work hardened by hot rolling. That is, the polygonal ferrite and acicular ferrite may be not drawn by hot rolling, and the polygonal ferrite and acicular ferrite may be produced after hot rolling.

The microstructure of the steel of the present invention may include bainitic ferrite, cementite, etc. in addition to the above-described polygonal ferrite, acicular ferrite, and MA phase.

Bainitic ferrite is a metamorphic structure at low temperature and has many dislocations inside, but has a relatively coarse feature compared to various ferrites, and also contains a MA phase inside, so it has high strength but is vulnerable to crack initiation and propagation. Should be controlled to a minimum.

In addition, the steel of the present invention may include an inclusion having a size of 10 μm or more in the range of 11 pieces / cm ² or less. The size is the size measured by the equivalent diameter. If the size of the inclusions having a size of 10 µm or more is more than 11 / cm ² , there occurs a problem of acting as a crack initiation point at low temperature. In order to control such coarse inclusions, it is preferable to add Ca or a Ca alloy at the final stage of the secondary refining, followed by bubbling and refluxing with Ar gas for at least 3 minutes.

In addition, the steel of the present invention may have a yield strength of 460 MPa or more, an impact energy value of 279J or more at -60 ° C, and a CTOD value of 0.2mm or more at -20 ° C. In addition, the steel material of the present invention may have a tensile strength of 570 MPa or more. And the steel of the present invention may have a DBTT (ductile-brittle transition temperature) is less than -80 ℃.

Next, a method for producing a high strength steel having excellent fracture initiation and propagation resistance at low temperatures of the present invention will be described.

Steel manufacturing method of the present invention comprises the steps of providing a steel slab that satisfies the above-described alloy composition; Heating the steel slab to 1000-1200 ° C .; Finishing hot rolling the heated slab in a temperature range of 650 ° C. or higher; And cooling the finished 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.

Steps to raise the river slab

A steel slab that satisfies the alloy composition as described above is provided.

At this time, in the present invention in preparing a steel slab, the step of injecting Ca or Ca alloy into the molten steel in the final stage of the secondary refining of molten steel; And after the Ca or Ca alloy is added, bubbling and refluxing with Ar gas for at least 3 minutes. This is to control coarse inclusions.

Steel slab heating stage

The steel slab is heated to 1000-1200 ° C.

When the slab heating temperature is less than 1000 ° C., it is difficult to re-use carbides and the like generated in the slab during the performance, and the homogenization treatment of segregated elements is insufficient. Therefore, it is preferable to heat the temperature to 1000 ° C. or more, which is a temperature at which 50% or more of the added Nb can be reused.

On the other hand, when the slab heating temperature exceeds 1200 ° C, the austenite grain size may grow too coarsely, and micronization may be insufficient due to subsequent rolling, and mechanical properties such as tensile strength and low temperature toughness of the steel sheet may be greatly reduced. .

Hot rolling stage

The heated slab is finished hot-rolled at 650 ° C. or more, which is bainite formation starting temperature, to obtain a hot rolled steel sheet.

When the finish hot rolling temperature is less than 650 ° C, coarse bainite is produced and processed and hardened during rolling, the strength is excessively increased more than necessary, and the impact toughness at low temperature is greatly reduced. It is preferable to limit to. In other words, when the hot rolling temperature is low, coarse salt-bearing ferrite is formed before the hot rolling finish, which is then stretched by rolling to form work hardening, and the remaining austenite remains in a band form and has a high density of MA hardening phase. This is because the transformation lowers the low-temperature toughness.

In addition, the present invention is carried out at a total reduction ratio of 30% or more (excluding the recrystallization reduction ratio) in the non-recrystallization temperature range to accumulate sufficient strain energy in austenite, and thus polygonal and needle type for favorable low temperature toughness during subsequent transformation It is desirable to produce enough ferrite and to ensure the ratio and density of the large-diameter grain boundary.

Cooling stage

Next, in the present invention, the finish hot rolled hot rolled steel sheet is cooled.

At this time, it is preferable to cool the hot-rolled steel sheet to the cooling end temperature of 200 ~ 550 ℃ at a cooling rate of 2 ~ 30 ℃ / s. If the cooling rate is less than 2 ℃ / s, the cooling rate is too slow to avoid coarse ferrite, pearlite and bainite transformation zones can cause heat and strength toughness, and if it is above 30 ℃ / s granules La bainite or martensite are formed to increase strength, but low temperature toughness can be very inferior.

When the cooling end temperature is higher than 550 ° C., microstructures such as acicular ferrite may be hardly generated, and coarse bainite or pearlite may be generated. On the other hand, if the temperature is less than 200 ° C., there is no disadvantage in the microstructure, but there is a problem that productivity is greatly reduced due to excessive time required for cooling.

On the other hand, in the present invention, after heating the cooled hot-rolled steel sheet to 450 ~ 650 ℃, if necessary, after maintaining the (1.3 x t + 5) to (1.3 x t + 200) minutes for further adding a tempering step It may include as (wherein t is a value measured in mm unit of the hot rolled steel sheet). When MA or martensite is excessively produced, it decomposes MA or martensite, removes high dislocation density inside, and deposits traced but solid solution of Nb as carbonitride to improve yield strength or low temperature toughness. For sake.

However, if the heating temperature is less than 450 ° C., the ferrite matrix is not softened sufficiently, and embrittlement phenomenon due to P segregation or the like appears, which may deteriorate toughness. On the other hand, when the heating temperature is higher than 650 ° C., the recovery and growth of the grains occur rapidly, and when the temperature is higher, some reverse transformation into austenite may cause the yield strength to be lowered, and at the same time, the low-temperature toughness may deteriorate.

When the holding time is less than (1.3 × t + 5) minutes, the homogenization of the tissue is not sufficiently performed, and when it is more than (1.3 × t + 200) minutes, there is a problem in that productivity is lowered.

Hereinafter, the present invention will be described in more detail with reference to Examples.

(Example)

The slabs having the composition shown in Table 1 below were heated, hot rolled, and cooled under the conditions shown in Table 2 to prepare steel materials.

The microstructure of the prepared steel was observed, and the physical properties thereof were listed in Table 3 below.

In addition, after welding the prepared steel material, the CTOD value (-20 ° C.) of the welding heat affected zone (SCHAZ) was measured and described in Table 3 below. Since the CTOD value (-20 ° C) of the steel is higher than the weld heat affected zone, the CTOD value (-20 ° C) for the steel was not separately measured.

At this time, the microstructure of the steel was polished to the mirror surface of the prepared steel and then etched with Nital or LePera according to the purpose, and the image was measured at a magnification of 100 to 5000 times with an optical or scanning electron microscope. The fraction of phases was measured from the measured images using an image analyzer. In order to obtain statistically significant values, the same specimens were repeatedly measured at different positions and their average values were obtained.

In addition, in order to observe the characteristics of the fabricated structure in detail, the Nital-etched specimen was subjected to Electron Back Scatter Diffraction (EBSD) measurement using a scanning electron microscope to quantitatively measure grain boundary properties of the manufactured steel.

The physical properties of the steels were measured and described from the nominal strain-nominal stress curves obtained by conventional tensile tests.

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 to machine the specimen in the size of B (thickness) x B (width) x 5B (length) perpendicular to the rolling direction in accordance with the BS 7448 standard, so that the fatigue crack length is approximately 50% of the specimen width. CTOD tests were performed at −20 ° C. after the fatigue crack was inserted. Where B is the thickness of the produced steel.

The Kca value was tested three times by ESSO test method, and the graph of the propagation stop temperature and K value of the cracks measured in each test was obtained, and the Kca value was obtained from the K value (Kca: crack arrest K) when the temperature was -10 degrees. In addition, CAT (Crack arrest temperature) was measured from the value calculated from the conversion equation of equation 1 by measuring the nil-ductility transition temperature (NDTT) from the NRL test. Here, B represents the thickness of the steel.

division
Steel grade
Alloy composition (% by weight) Relationship 1 Relationship 2 C Si Mn P S Sol.Al Cu Ni Cr Mo Ti Nb V N Ca foot
persons
River a 0.035 0.18 2.2 0.005 0.0011 0.005 0.28 1.3 0.05 0.02 0.012 0.005 0.003 0.0035 0.002 3.0 0.39 b 0.046 0.01 1.9 0.005 0.0008 0.02 0.31 0.99 0.03 0.29 0.012 0.011 0.003 0.0042 0.0025 2.6 0.45 c 0.015 0.008 2.5 0.005 0.001 0.005 0.09 0.61 0.02 0.24 0.012 0.003 0.005 0.0026 0.0016 2.9 0.30 ratio
School
River


d 0.081 0.015 1.9 0.005 0.0012 0.022 0.31 1.13 0.03 0.2 0.012 0.012 0.003 0.0035 0.0028 2.6 0.37 e 0.021 0.011 3.2 0.005 0.0013 0.015 0.03 0.25 0.03 0.02 0.012 0.007 0.003 0.0035 0.0016 3.3 0.14 f 0.029 0.01 2.1 0.005 0.0009 0.011 0.29 0.75 0.12 0.15 0.012 0.038 0.003 0.0035 0.0011 2.6 0.67 g 0.025 0.015 2.4 0.005 0.0012 0.005 0.12 0.65 0.23 0.35 0.012 0.018 0.003 0.0035 0.0021 2.8 0.78 h 0.061 0.051 1.7 0.005 0.0018 0.028 0.15 0.25 0.03 0.2 0.012 0.005 0.003 0.0035 0.0018 1.9 0.36 i 0.034 0.017 1.9 00.057 0.0008 0.008 0.22 0.53 0.33 0.19 0.007 0.012 0.004 0.0039 0.0015 2.3 0.67

In Table 1, relationship 1 is Mn + 0.5 x (Ni + Cu), relationship 2 is Mo + Cr + 1.5 x Si + 10 x Nb.

division Steel grade Product thickness (mm) Slab reheating temperature (℃) Rolling end temperature (℃) Total unloading rate (%) Cooling rate (℃ / s) Cooling end temperature (℃) Whether to perform tempering Inventive Example 1 a 85 1060 790 45 4.0 390 x Inventive Example 2 b 100 1100 750 40 3.4 340 x Inventive Example 3 c 51 1120 800 57 6.9 280 x Inventive Example 4 b 76 1100 810 50 4.0 260 ○ Comparative Example 1 d 76 1110 780 48 4.5 380 x Comparative Example 2 e 51 1140 840 55 7.5 330 x Comparative Example 3 f 85 1120 805 35 3.1 280 x Comparative Example 4 g 80 1100 800 45 4.5 440 x Comparative Example 5 h 80 1100 780 40 4.3 360 x Comparative Example 6 i 100 1180 770 35 3.3 310 x Comparative Example 7 a 90 1170 820 20 3.5 270 x Comparative Example 8 c 76 1160 800 50 - - x

division Steel grade Ferritic phase fraction (%) Large-diameter grain boundary ratio (%) Confrontation angle grain boundary length density (mm / mm ² ) Yield strength (MPa) Tensile Strength (MPa) Impact energy value (-60 ° C base material): J Kca (-10 ℃) CAT (℃) HAZ part
CTOD value
(-20 ℃): mm Inventive Example 1 a 78 53 390 480 620 315 9350 -16 0.82 Inventive Example 2 b 73 42 320 525 679 279 - -17 0.36 Inventive Example 3 c 83 63 490 488 627 332 - -20 1.21 Inventive Example 4 b 85 50 390 533 637 322 - -26 0.76 Comparative Example 1 d 56 43 380 596 766 75 - 11 0.18 Comparative Example 2 e 57 35 220 520 668 56 - -4 0.14 Comparative Example 3 f 55 33 190 540 696 90 5860 3 0.08 Comparative Example 4 g 46 42 210 587 754 103 - 2 0.12 Comparative Example 5 h 75 51 390 428 554 268 - -25 0.35 Comparative Example 6 i 55 38 220 471 610 85 - -9 0.08 Comparative Example 7 a 61 23 180 475 615 112 - -2 0.15 Comparative Example 8 c 86 56 480 435 545 185 - -36 0.29

* In Table 3, the ferrite refers to the sum of polygonal ferrite and acicular ferrite.

As shown in Table 1-3, Inventive Examples 1 to 4 satisfying all of the alloy composition and manufacturing conditions presented in the present invention are fracture at low temperatures when considering yield strength, tensile strength, impact energy value, Kca and CAT, etc. It can be confirmed that the toughness resistance is excellent and the CTOD value in the weld heat affected zone is also high. In particular, as shown in Fig. 1, the Kca value measured in Inventive Example 1 shows a value that greatly exceeds the required value of 8000. These excellent strength and low temperature toughness characteristics can also be confirmed that the results obtained from fine polygonal and needle-like ferrite structure produced sufficiently as shown in FIG.

In contrast, Comparative Example 1 is the case where the C content exceeds the scope of the present invention, the added C is the strongest element to promote granular bainite and MA. Therefore, due to excessive addition of C, the fraction of ferrite, which is advantageous for toughness, is greatly reduced, and the strength in the base material is high, but low-temperature toughness, such as impact energy value, is inferior. In particular, the CTOD value of the weld heat affected zone is greatly reduced.

Comparative Example 2 is a case where the added Mn content exceeds the range of the present invention. In this case, due to the high Mn content, the probability of segregation in the center of the steel is greatly increased, so that the impact energy in the center of the steel thickness direction is inferior. In addition, the hardened structure having a large hardness is partially in the segregation zone of the center even in the weld heat affected zone. Premature pop-in, resulting in a significant drop in CTOD.

Comparative Example 3 is generally the case where the content of Nb widely used for strength improvement and microstructure of the tissue exceeds the scope of the present invention. In general, the addition of Nb is advantageous to increase the strength and toughness at the same time by miniaturizing the tissue, but if it is added more than necessary, it suppresses the production of polygonal and needle-like ferrite, which is beneficial to toughness, and promotes tissues such as granular bainite. Let's do it. Therefore, the ratio and density of the large-diameter grain boundary of 15 ° or more, which are advantageous for suppressing the propagation of the crack, are greatly reduced, thereby making the propagation of the crack relatively easy. As a result, as shown in Table 3, the Kca value measured in Comparative Example 3 was 5860, which was far below the required value of 8000. In addition, the weld heat affected zone greatly promoted the generation of M-A tissues that adversely affect the low-temperature toughness, and as a result, significantly lowered the CTOD.

 In Comparative Examples 4, 5, and 6, the range of each element content satisfies the scope of the present invention, but the relationship 1 and the value of relationship 2 are outside the scope of the invention, indicating that the strength is low or the low temperature toughness is greatly reduced. have.

Specifically, Comparative Example 4 satisfies the relational formula 1 composed of components that are advantageous for improving low-temperature toughness, but the relational formula 2 composed of components that impair low-temperature toughness exceeds the scope of the present invention. As a result, although the strength was high enough, the impact energy value in the base material and the CTOD value in the weld heat affected zone were inferior.

In addition, in Comparative Example 5, the relational formula 2 satisfies the scope of the invention, but the relational formula 1 is out of the scope of the present invention, the amount of the component is insufficient to secure the strength of the steel as a whole, the strength of the base material is greatly reduced.

Comparative Example 6 is a case where both formulas (1) and (2) are outside the scope of the invention. That is, a component that is advantageous for low temperature toughness is insufficient, and a component that is disadvantageous for low temperature toughness is exceeded, and all low temperature toughness characteristic values are inferior.

The comparative example 7 is a case where the components of steel materials satisfy | fill all the scope of invention, but fall short of the invention of the unrecrystallized reverse rolling total pressure drop in the steel manufacturing process. In other words, the fraction of ferrite, which prevents the propagation of cracks in the microstructure of the steel due to insufficient rolling amount in the unrecrystallized zone, was also low, and the low temperature toughness characteristic values were not good because the ratio and density of the large-angle grain boundary were greatly lowered. .

In Comparative Example 8, although the steel component satisfies all the scope of the invention, it is manufactured by air cooling without applying accelerated cooling after controlled rolling in the steel manufacturing process. Is sufficiently produced, but is coarse and the strength is greatly reduced.

Although the embodiments of the present invention have been described in detail above, the scope of the present invention is not limited thereto, and various modifications and variations can be made without departing from the technical spirit of the present invention described in the claims. It will be obvious to those who have ordinary knowledge of.

Claims (10)

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: 0.005% 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 0.03%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001% to 0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% or less (except 0%), balance Fe and inevitable impurities Satisfying the following Expressions 1 and 2,
The microstructure contains not less than 70 area% of total polygonal ferrite and acicular ferrite, less than 3.5 area% of MA phase (martensite-austenite composite phase), and the polygonal ferrite and needle shape Ferrite has a fracture resistance at low temperatures where the ratio of grain orientations of large-diameter grains defined by the difference of grain orientations between grains is greater than 40% among the grain boundaries and the length of the large-grain grain boundaries per unit area is 300 mm / mm ² or more. Excellent high strength steels.
[Relationship 1]
Mn + 0.5 x (Ni + Cu) ≥ 2.5wt%
[Relationship 2]
Equation 2: Mo + Cr + 1.5 x Si + 10 x Nb ≤ 0.5wt%
(However, in the above formula 1-2, each element is a value expressed in weight%.)
delete The method of claim 1, wherein the steel has a yield strength of 460 MPa or more, the impact energy value at -60 ℃ is 279J or more, the Kca value measured in ESSO test is 8000 N / mm ^3/2 or more NRL test High strength steel with excellent fracture resistance at low temperature, characterized in that the crack arrest temperature (CAT) calculated from the NDTT (Nil-ductility transition temperature) measured at less than -10 ℃.
The high strength steel of claim 1, wherein the steel has a tensile strength of 570 MPa or more and a DBTT (ductility-brittle transition temperature) of -80 ° C or less.
The high-strength steel having excellent fracture resistance at low temperatures according to claim 1, wherein the steel comprises 11 inclusions / cm ² or less having inclusions having a size of 10 µm or more as measured by a circular equivalent diameter.
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: 0.005% 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 0.03%, Mo: 0.001 to 0.5%, Ca: 0.0002 to 0.005%, N: 0.001% to 0.008%, P: 0.02% or less (except 0%), S: 0.003% or less (except 0%), O: 0.003% or less (except 0%), balance Fe and inevitable impurities Preparing a steel slab that satisfies the following relations 1 and 2;
Heating the steel slab to 1000-1200 ° C .;
Finishing hot rolling the heated slab in a temperature range of 650 ° C. or higher such that the total reduction ratio in the non-recrystallization temperature range is 30% or more (excluding the recrystallization reduction ratio); And
Cooling the finish hot rolled hot rolled steel sheet to a cooling end temperature of 200 ~ 550 ℃ at a cooling rate of 2 ~ 30 ℃ / s,
The cooled steel material has a microstructure of at least 70 area% of total polygonal ferrite and acicular ferrite, and includes less than 3.5 area% of MA phase (martensite-austenite composite phase), and the poly Gonal ferrite and needle-type ferrite have a ratio of grain boundaries of large diameter at which the grain orientation difference between grains is 15 ° or more is 40% or more of the total grain boundaries, and the length of large diameter grain boundaries per unit area is 300 mm / mm ² or more. A method for producing high strength steel with excellent fracture resistance at low temperatures.
[Relationship 1]
Mn + 0.5 x (Ni + Cu) ≥ 2.5wt%
[Relationship 2]
Equation 2: Mo + Cr + 1.5 x Si + 10 x Nb ≤ 0.5wt%
(However, each of the elements in the above formula 1-2 is a value expressed in weight%.)
The method of claim 6, further comprising a tempering step of heating the cooled hot-rolled steel sheet to 450 ~ 650 ℃, and then cooled after maintaining for (1.3 x t + 5) minutes (1.3 x t + 200) minutes Method for producing high strength steel with excellent fracture resistance at low temperatures.
The method of claim 6,
In providing the steel slab,
Injecting Ca or Ca alloy into the molten steel in the final stage of the secondary refining of molten steel; And bubbling and refluxing with Ar gas for at least 3 minutes after the Ca or Ca alloy is added. The method of manufacturing a high strength steel having excellent fracture resistance at low temperatures, comprising:
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