Patent Application
20180363107
The present disclosure relates to a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance, and to a method of manufacturing the same.
Recently, there has been demand for the development of ultra-thick steel sheets having high strength properties in consideration of the design requirements of structures to be used in the shipping, maritime, architectural, and civil engineering fields, domestically and internationally.
In a case in which high-strength steel is included in the design of a structure, economic benefits may be obtained due to reductions in the weight of structures while processing and welding operations may be easily undertaken using a steel sheet having a relatively reduced thickness.
In general, in the case of high-strength steel, due to a reduction in a reduction ratio when thick steel plates are manufactured, sufficient deformation may not be performed, as compared with a thin steel sheet. Thus, microstructures of thick steel plates may be coarse, so that low temperature properties on which grain sizes have the most significant effect may be degraded.
In detail, in a case in which brittle crack arrestability representing stability of a structure is applied to a main structure, such as a ship's hull, the number of cases of assurances being demanded has increased. However, in a case in which microstructures become coarse, a phenomenon in which brittle crack arrestability is significantly degraded may occur. Thus, it may be difficult to improve brittle crack arrestability of an ultra-thick high-strength steel material.
In the meantime, in the case of high-strength steel having yield strength of 390 MPa or greater, various technologies, such as adjustment of grain size, by applying surface cooling during finishing milling and applying bending stress during rolling to refine the grain size of a surface portion, in order to improve brittle crack arrestability, have been introduced.
However, such technologies may contribute to refining a structure of a surface portion, but may not solve a problem in which impact toughness is degraded due to coarsening of structures other than that of the surface portion. Thus, such technologies may not be fundamental countermeasures to brittle crack arrestability.
In addition, recently, a design concept to improve safety of a ship by controlling brittle crack initiation of a steel material applied to large container ships has been introduced. Thus, in general, the number of cases of guaranteeing brittle crack initiation of a heat affected zone (HAZ), the most vulnerable portion in terms of brittle crack initiation, has increased.
In general, since, in the case of high-strength steel, the microstructure in a HAZ includes low temperature transformation ferrite having high strength, such as bainite, there is a limitation in which HAZ properties, in detail, toughness, are significantly reduced.
In detail, in the case of brittle crack initiation resistance generally evaluated through a crack tip opening displacement (CTOD) test to evaluate the stability of the structure, martensite-austenite generated from untransformed austenite, when low temperature transformation ferrite is generated, becomes an active nucleation site of brittle crack occurrence. Thus, it maybe difficult to improve brittle crack initiation resistance of a high-strength steel material.
In the case of high-strength steel of the related art having a yield strength of 400 MPa or greater, in order to improve welding zone brittle crack initiation resistance, an effort to refine a microstructure in a HAZ using TiN or to form ferrite in a HAZ using oxide metallurgy has been made. However, the effort partially contributes to forming impact toughness through refining a structure, but does not have a great effect on reducing a fraction of martensite-austenite having a significant influence on reducing brittle crack initiation resistance.
In addition, in the case of brittle crack initiation resistance of a base material, martensite-austenite may be transformed to have a different phase through tempering, or the like, to secure physical properties. However, in the case of a HAZ, in which an effect of tempering disappears due to thermal history, it is impossible to apply brittle crack initiation resistance.
In the meantime, in order to minimize the formation of martensite-austenite, the amount of elements, such as carbon (C) and niobium (Nb), should be reduced. However, in this case, it may be difficult to secure a specific level of strength. To this end, a relatively large amount of high-priced elements, such as molybdenum (Mo) and nickel (Ni), should be added. Thus, there is a limitation in which economic efficiency is deteriorated.
An aspect of the present disclosure may provide a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance.
Another aspect of the present disclosure may provide a method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance.
According to an aspect of the present disclosure, a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises, by wt %, carbon (C): 0.05% to 0.09%, manganese (Mn): 1.5% to 2.0%, nickel (Ni): 0.3% to 0.8%, niobium (Nb): 0.005% to 0.04%, titanium (Ti): 0.005% to 0.04%, copper (Cu): 0.1% to 0.5%, silicon (Si): 0.05% to 0.3%, aluminum (Al): 0.005% to 0.05%, phosphorus (P): 100 ppm or less, sulfur (S): 40 ppm or less, iron (Fe) as a residual component thereof, and inevitable impurities, wherein a microstructure of a central portion includes, by area %, acicular ferrite in an amount of 70% or greater, pearlite in an amount of 10% or less, and one or more selected from a group consisting of ferrite, bainite, and martensite-austenite (MA), as residual components; a circle-equivalent diameter of pearlite being 15 μm or less; a surface portion microstructure in a region at a depth of 2 mm or less, directly below a surface, includes, by area %, ferrite in an amount of 30% or greater and one or more of bainite, martensite, and pearlite as residual components; and a heat affected zone (HAZ) formed during welding includes, by area %, martensite-austenite (MA) in an amount of 5% or less.
A weight ratio of Cu and Ni (a Cu/Ni weight ratio) may be set to be 0.8 or less, and in more detail, 0.6 or less.
The high-strength steel material may have yield strength of 390 MPa or greater.
The high-strength steel material may have a Charpy fracture transition temperature of −40° C. or lower in a 1/2t position in a steel material thickness direction, where t is a steel sheet thickness.
According to another aspect of the present disclosure, a method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises rough rolling a slab at a temperature of 900° C. to 1100° C. after reheating the slab at 1000° C. to 1100° C., including, by wt %, C: 0.05% to 0.09%, Mn: 1.5% to 2.0%, Ni: 0.3% to 0.8%, Nb: 0.005% to 0.04%, titanium (Ti): 0.005% to 0.04%, copper (Cu): 0.1% to 0.5%, silicon (Si): 0.1% to 0.3%, aluminum (Al): 0.005% to 0.05%, phosphorus (P): 100 ppm or less, sulfur (S): 40 ppm or less, iron (Fe) as a residual component thereof, and inevitable impurities; obtaining a steel sheet by finish rolling a bar obtained from the rough rolling a slab, at a temperature in a range of Ar3+60° C. to Ar3° C., based on a temperature of a central portion; and cooling the steel sheet to 700° C. or lower.
A reduction ratio per pass of three final passes during the rough rolling a slab may be 5% or greater, and a total cumulative reduction ratio may be 40% or greater.
A strain rate of three final passes during the rough rolling a slab may be 2/sec or lower.
A grain size of a central portion in a bar thickness direction before finish rolling after the rough rolling a slab may be 150 μm or less, in detail, 100 μm or less, and more specifically, 80 μm or less.
A reduction ratio during the finish rolling may be set such that a ratio of a slab thickness (mm) to a steel sheet thickness (mm) after the finish rolling may be 3.5 or greater, and in more detail, 4 or greater.
A cumulative reduction ratio during the finish rolling may be maintained to be 40% or greater, while the reduction ratio per pass, not including skin pass rolling, may be maintained to be 4% or greater.
The skin pass rolling is performed to secure a shape of a sheet (to secure a flat sheet) at a relatively low reduction rate, less than 5% in 1 to 2 passes of finish rolling.
The cooling the steel sheet may be performed at a cooling rate of the central portion of 1.5° C./s or higher.
The cooling the steel sheet may be performed at an average cooling rate of 2° C./s to 300° C./s.
In addition, the present inventive concept may be exemplified in many different forms and should not be construed as being limited to the specific embodiments set forth herein. Rather, these embodiments are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the disclosure to those skilled in the art.
According to an aspect of the present disclosure, a high-strength steel material having a relatively high level of yield strength, as well as excellent brittle crack arrestability and welding zone brittle crack initiation resistance may be provided.
The inventors of the present disclosure conducted research and experiments to improve yield strength, brittle crack arrestability, and welding zone brittle crack initiation resistance of a thick steel material and proposed the present disclosure based on results thereof.
In an exemplary embodiment, a steel composition, a structure, and manufacturing conditions of a steel material may be controlled, thereby improving yield strength, brittle crack arrestability, and welding zone brittle crack initiation resistance of the thick steel material.
A main concept of an exemplary embodiment is as follows.
1) The steel composition is appropriately controlled to improve strength through solid solution strengthening. In detail, contents of manganese (Mn), nickel (Ni), copper (Cu), and silicon (Si) are optimized for solid solution strengthening.
2) The steel composition is appropriately controlled to improve strength by increasing hardenability. In detail, the contents of Mn, Ni, and Cu, as well as a carbon (C) content are optimized to increase hardenability.
A fine structure is secured in a central portion of the thick steel material even at a relatively slow cooling rate.
3) In detail, a weight ratio of Cu to Ni may be controlled.
In a case in which the weight ratio of Cu to Ni is controlled as described above, surface quality may be improved.
4) A composition is appropriately controlled to control a fraction of martensite-austenite in a heat affected zone (HAZ) formed during welding. In detail, contents of C, Si, and niobium (Nb), affecting generation of martensite-austenite, are optimized.
As such, the steel composition may be optimized, thereby securing excellent brittle crack initiation resistance even in the HAZ.
5) More specifically, a structure of the steel material may be controlled to improve strength and brittle crack arrestability. In detail, a structure of the central portion and a surface layer region is controlled in a direction of a steel material thickness.
As such, a microstructure may be controlled, thereby securing strength required in the steel material, while the microstructure facilitating generation of a crack may be excluded, thereby improving brittle crack arrestability.
6) In detail, rough rolling conditions may be controlled to refine the structure of the steel material. In detail, the fine structure is secured in the central portion by controlling rolling conditions during rough rolling. Using a process described above, the generation of acicular ferrite is also facilitated.
7) Finish rolling conditions may be controlled to further refine the structure of the steel material. In detail, the fine structure is secured in the central portion by controlling rolling conditions during rough rolling. As such, the generation of acicular ferrite is also facilitated.
Hereinafter, the high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance according to an aspect of the present disclosure will be described in detail.
According to an aspect of the present disclosure, a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises, by wt %, carbon (C): 0.05% to 0.09%, manganese (Mn): 1.5% to 2.0%, nickel (Ni): 0.3% to 0.8%, niobium (Nb): 0.005% to 0.04%, titanium (Ti): 0.005% to 0.04%, copper (Cu): 0.1% to 0.5%, silicon (Si): 0.05% to 0.3%, aluminum (Al): 0.005% to 0.05%, phosphorus (P): 100 ppm or less, sulfur (S): 40 ppm or less, iron (Fe) as a residual component thereof, and inevitable impurities, wherein a microstructure of a central portion includes, by area %, acicular ferrite in an amount of 70% or greater, pearlite in an amount of 10% or less, and one or more selected from a group consisting of ferrite, bainite, and martensite-austenite (MA), as residual components; a circle-equivalent diameter of pearlite being 15 μm or less; a surface portion microstructure in a region at a depth of 2 mm or less, directly below a surface, includes, by area %, ferrite in an amount of 30% or greater and one or more of bainite, martensite, and pearlite as residual components; and a heat affected zone (HAZ) formed during welding includes, by area %, martensite-austenite (MA) in an amount of 5% or less.
Hereinafter, a steel component and a component range of an exemplary embodiment will be described.
Carbon (C): 0.05 wt % to 0.09 wt % (hereinafter, referred to as “%”)
Since C is the most significant element used in securing basic strength, C is required to be contained in steel within an appropriate range. In order to obtain an effect of addition, C may be added in an amount of 0.05% or greater.
However, in a case in which a C content exceeds 0.09%, a large amount of martensite-austenite is generated in the HAZ to degrade brittle crack initiation resistance. Low temperature toughness is degraded due to a relatively high level of strength of ferrite of a base material and the generation of a relatively large amount of low temperature transformation ferrite. Thus, the C content maybe limited to 0.05% to 0.09%. In detail, the C content may be limited to 0.061% to 0.085%, and more specifically, to 0.065% to 0.075%.
Manganese (Mn): 1.5% to 2.0%
Mn is a useful element improving strength through solid solution strengthening and increasing hardenability to generate low temperature transformation ferrite. In addition, since Mn may generate low temperature transformation ferrite even at a relatively low cooling rate due to improved hardenability, Mn is a main element used to secure strength of a central portion of a thick steel plate.
Therefore, in order to obtain an effect described above, Mn may be added in an amount of 1.5% or greater.
However, in a case in which a Mn content exceeds 2.0%, generation of upper bainite and martensite may be facilitated due to an increase in excessive hardenability, thereby degrading impact toughness and brittle crack arrestability and toughness of the HAZ.
Therefore, the Mn content may be limited to 1.5% to 2.0%. In detail, the Mn content may be limited to 1.61% to 1.92%, and more specifically, to 1.7% to 1.9%.
Nickel (Ni): 0.3% to 0.8%
Ni is a significant element used in improving impact toughness by facilitating a dislocation cross slip at a relatively low temperature and increasing strength by improving hardenability. In order to obtain an effect described above, Ni may be added in an amount of 0.8% or greater. However, in a case in which Ni is added in an amount of 1.2% or greater, hardenability may be excessively increased to generate low temperature transformation ferrite, thereby degrading toughness, and manufacturing costs may be increased due to a relatively high cost of Ni, as compared with other hardenability elements. Thus, an upper limit value of the Ni content may be limited to 0.8%.
In detail, the Ni content may be limited to 0.37% to 0.71%, and more specifically, to 0.4% to 0.6%.
Niobium (Nb): 0.005% to 0.04%
Nb is educed to have a form of NbC or NbCN to improve strength of a base material.
In addition, Nb solidified when being reheated at a relatively high temperature is significantly finely educed to have the form of NbC during rolling to suppress recrystallization of austenite, thereby having an effect of refining a structure.
Therefore, Nb may be added in an amount of 0.005% or greater. However, in a case in which Nb is added excessively, generation of martensite-austenite in the HAZ may be facilitated to degrade brittle crack initiation resistance and cause a brittle crack in an edge of the steel material. Thus, an upper limit value of an Nb content maybe limited to 0.04%.
In detail, the Nb content may be limited to 0.012% to 0.031%, and more specifically, to 0.017% to 0.03%.
Titanium (Ti): 0.005% to 0.04%
Ti is a component educed to be TiN when being reheated and inhibiting growth of the base material and a grain in the HAZ to greatly improve low temperature toughness. In order to obtain an effect of addition, Ti may be added in an amount of 0.005% or greater.
However, in a case in which Ti is added excessively, low temperature toughness may be degraded due to clogging of a continuous casting nozzle or crystallization of the central portion. Thus, a Ti content maybe limited to 0.005% to 0.04%.
In detail, the Ti content may be limited to 0.012% to 0.023%, and more specifically, to 0.014% to 0.018%.
Silicon (Si): 0.05% to 0.3%
Si is a substitutional element improving strength of the steel material through solid solution strengthening and having a strong deoxidation effect, so that Si maybe an element essential in manufacturing clean steel. Thus, Si may be added in an amount of 0.05% or greater. However, when a relatively large amount of Si is added, a coarse martensite-austenite phase may be formed to degrade brittle crack arrestability and welding zone brittle crack initiation resistance. Thus, an upper limit value of a Si content may be limited to 0.3%.
In detail, the Si content may be limited to 0.1% to 0.27%, and more specifically, to 0.19% to 0.25%.
Copper (Cu): 0.1% to 0.5%
Cu is a main element used in improving hardenability and causing solid solution strengthening to enhance strength of the steel material. In addition, Cu is a main element used in increasing yield strength through the generation of an epsilon Cu precipitate when tempering is applied. Thus, Cu may be added in an amount of 0.1% or greater. However, when a relatively large amount of Cu is added, a slab crack may be generated by hot shortness in a steelmaking process. Thus, an upper limit value of a Cu content may be limited to 0.5%.
In detail, the Cu content may be limited to 0.15% to 0.31%, and more specifically, to 0.2% to 0.3%.
Contents of Cu and Ni may be set such that the weight ratio of Cu to Ni may be 0.8 or less, and in more detail, 0.6 or less. More specifically, the weight ratio of Cu to Ni may be limited to 0.5 or less.
In a case in which the weight ratio of Cu to Ni is set as described above, surface quality may be improved.
Aluminum (Al): 0.005% to 0.05%
Al is a component functioning as a deoxidizer. In a case in which an excessive amount of Al is contained, an inclusion may be formed to degrade toughness. Thus, an Al content may be limited to 0.005% to 0.05%.
Phosphorus (P): 100 ppm or less, Sulfur (S): 40 ppm or less
P and S are elements causing brittleness in a grain boundary or forming a coarse inclusion to cause brittleness. In order to improve brittle crack arrestability, a P content may be limited to 100 ppm or less, while an S content may be limited to 40 ppm or less.
A residual component of an exemplary embodiment is Fe.
However, since, in a manufacturing process of the related art, unintended impurities maybe inevitably mixed from a raw material or an external source, which may not be excluded.
Since the impurities are apparent to those skilled in the art, all the contents thereof are not specifically described in the present disclosure.
In the case of a steel material of an exemplary embodiment, a microstructure of a central portion includes, by area %, acicular ferrite in an amount of 70% or greater, pearlite in an amount of 10% or less, and one or more selected from a group consisting of ferrite, bainite, and martensite-austenite (MA), as residual components; a circle-equivalent diameter of pearlite being 15 μm or less; a surface portion microstructure in a region at a depth of 2 mm or less, directly below a surface, includes, by area %, ferrite in an amount of 30% or greater and one or more of bainite, martensite, and pearlite as residual components; and a heat affected zone (HAZ) formed during welding includes, by area %, martensite-austenite (MA) in an amount of 5% or less.
Ferrite refers to polygonal ferrite, while bainite refers to granular bainite and upper bainite.
In a case in which a fraction of acicular ferrite of the microstructure of the central portion is less than 70%, generation of coarse bainite may cause degradation of toughness.
In detail, the fraction of acicular ferrite may be 75% or greater, and more specifically, may be limited to 80% or greater.
In a case in which a fraction of pearlite in the central portion exceeds 10%, a microcrack may be generated in a front end of a crack during brittle crack propagation, thereby degrading brittle crack arrestability. Thus, the fraction of pearlite in the central portion may be 10% or less.
In detail, the fraction of pearlite may be limited to 8% or less, and more specifically, to 5% or less.
In a case in which the circle-equivalent diameter of pearlite in the central portion exceeds 15 μm, there is a problem in which a crack may be easily generated despite a relatively low fraction of pearlite being present in the central portion. Thus, the circle-equivalent diameter of pearlite in the central portion may be 15 μm or less.
In a case in which the surface portion microstructure in the region at a depth of 2 mm or less, directly below the surface, includes ferrite in an amount of 30% or greater, crack propagation may be effectively prevented on the surface at a time of brittle crack propagation, thereby improving brittle crack arrestability.
In detail, the fraction of ferrite may be limited to 40% or greater, and more specifically, to 50% or greater.
In a case in which a fraction of martensite-austenite in the HAZ formed when the steel material is welded exceeds 5%, martensite-austenite functions as a starting point of a crack, thereby degrading brittle crack initiation resistance. Thus, the fraction of martensite-austenite in the HAZ may be 5% or less.
Welding heat input during welding may be 0.5 kJ/mm to 10 kJ/mm.
A welding method during welding is not specifically limited and may include, for example, flux cored arc welding (FCAW), submerged arc welding (SAW), and the like.
The steel material may have yield strength of 390 MPa or greater.
The steel material may have a Charpy fracture transition temperature of −40° C. or lower in a 1/2t position in a steel material thickness direction, where t is a steel sheet thickness.
The steel material have a thickness of 50 mm or greater, in detail, a thickness of 60 mm to 100 mm, and more specifically, 80 mm to 100 mm.
Hereinafter, a method of manufacturing a high-strength steel material having excellent brittle crack arrestability according to another aspect of the present disclosure will be described in detail.
According to another aspect of the present disclosure, the method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises rough rolling a slab at a temperature of 900° C. to 1100° C. after reheating the slab at 1000° C. to 1100° C., including, by wt %, C: 0.05% to 0.09%, Mn: 1.5% to 2.0%, Ni: 0.3% to 0.8%, Nb: 0.005% to 0.04%, Ti: 0.005% to 0.04%, Cu: 0.1% to 0.5%, Si: 0.1% to 0.3%, Al: 0.005% to 0.05%, P: 100 ppm or less, S: 40 ppm or less, Fe as a residual component thereof, and inevitable impurities; obtaining a steel sheet by finish rolling a bar obtained from the rough rolling a slab, at a temperature in a range of Ar3+60° C. to Ar3° C., based on a temperature of a central portion; and cooling the steel sheet to 700° C. or lower.
Reheating a Slab
A slab is reheated before rough rolling.
A reheating temperature of the slab may be 1000° C. or higher so that a carbonitride of Ti and/or Nb, formed during casting, may be solidified.
However, in a case in which the slab is reheated at a significantly high temperature, austenite may become coarse. Thus, an upper limit value of the reheating temperature maybe 1100° C.
Rough Rolling
A reheated slab is rough rolled.
A rough rolling temperature may be a temperature Tnr at which recrystallization of austenite is halted, or higher. Due to rolling, a cast structure, such as a dendrite formed during casting, may be destroyed, and an effect of reducing a size of austenite may also be obtained. In order to obtain the effect, the rough rolling temperature may be limited to 900° C. to 1100° C.
In more detail, the rough rolling temperature may be 950° C. to 1050° C.
In an exemplary embodiment, in order to refine a structure of the central portion during rough rolling, a reduction ratio per pass of three final passes during rough rolling may be 5% or greater, and a total cumulative reduction ratio may be 40% or greater.
In the case of a structure recrystallized by initial rolling during rough rolling, grain growth occurs due to a relatively high temperature. However, when three final passes are performed, a bar is air cooled while waiting for a rolling process, so that grain growth speed may be decreased. Thus, during rough rolling, a reduction ratio of the three final passes has the greatest impact on a grain size of a final microstructure.
In addition, in a case in which the reduction ratio is reduced per pass of rough rolling, sufficient deformation may not be transmitted to the central portion, so that toughness may be degraded due to coarsening of the central portion. Therefore, the reduction ratio per pass of the three final passes may be limited to 5% or greater.
In detail, the reduction ratio per pass may be 7% to 20%.
In the meantime, in order to refine a structure of the central portion, the total cumulative reduction ratio during rough rolling may be set to be 40% or greater.
In detail, the total cumulative reduction ratio may be 45% or greater.
A strain rate of the three final passes during rough rolling may be 2/sec or lower.
In general, rolling is difficult at a relatively high reduction ratio due to a relatively great thickness of the bar during rough rolling. Thus, there is a limitation in which it is difficult to transmit a rolling reduction to the central portion of a thick steel plate, thereby allowing an austenite grain size in the central portion to be coarsened. However, as the strain rate is reduced, deformation is transmitted to the central portion even at a relatively low rolling reduction ratio. Thus, the grain size may be refined.
Therefore, in terms of the three final passes having the greatest impact on the final grain size during rough rolling, the strain rate may be limited to 2/sec or lower, thereby refining the grain size of the central portion. Thus, generation of acicular ferrite may be facilitated.
Finish Rolling
A rough rolled bar may be finish rolled at a temperature of Ar3 (a ferrite transformation initiation temperature)+60° C. to Ar3° C. to obtain a steel sheet so that a further refined microstructure may be obtained.
In a case in which rolling is performed at a temperature higher than Ar3, a relatively large amount of strain bands may be generated in austenite to secure a relatively large number of ferrite nucleation sites, thereby obtaining an effect of securing a fine structure in the central portion of a steel material.
In addition, in order to effectively generate a relatively large amount of strain bands in austenite, a cumulative reduction ratio during finish rolling may be maintained to be 40% or greater. The reduction ratio per pass, not including skin pass rolling, may be maintained to be 4% or greater.
In detail, the cumulative reduction ratio may be 40% to 80%.
In detail, the reduction ratio per pass may be 4.5% or greater.
In a case in which a finish rolling temperature is reduced to Ar3 or lower, coarse ferrite is generated before rolling and is elongated during rolling, thereby reducing impact toughness. In a case in which finish rolling is performed at a temperature of Ar3+60° C. or higher, the grain size is not effectively refined, so that the finish rolling temperature during finish rolling may be set to be a temperature of Ar3+60 ° C. to Ar3° C.
In an exemplary embodiment, a reduction ratio in an unrecrystallized region maybe limited to 40% to 80% during finish rolling.
As described above, since the reduction ratio in the unrecrystallized region is controlled, thereby increasing a number of nucleation sites of acicular ferrite, generation of structures described above may be facilitated.
In a case in which the reduction ratio in the unrecrystallized region is significantly low, acicular ferrite may not be sufficiently secured. In a case in which the reduction ratio in the unrecrystallized region is significantly high, strength may be reduced due to generation of pro-eutectoid ferrite caused by a relatively high reduction ratio.
The grain size of the central portion of the bar in a thickness direction after rough rolling before finish rolling may be 150 μm or less, in detail, 100 μm or less, and more specifically, 80 μm or less.
The grain size of the central portion of the bar in a thickness direction after rough rolling before finish rolling may be controlled depending on a rough rolling condition, or the like.
As described above, in a case in which the grain size of the bar after rough rolling before finish rolling may be controlled, a final microstructure is refined due to refinement of an austenite grain. Thus, an advantage of improving low temperature impact toughness may be added.
The reduction ratio during finish rolling may be set such that a ratio of a slab thickness (mm) to a steel sheet thickness (mm) after finish rolling may be 3.5 or greater, and in detail, 4 or greater.
As described above, in a case in which the reduction ratio is controlled, as the rolling reduction is increased during rough rolling and finish rolling, an advantage of improving toughness of the central portion may be added by increasing yield strength/tensile strength, improving low temperature toughness, and decreasing the grain size of the central portion in the thickness direction through refinement of the final microstructure.
After finish rolling, the steel sheet may have a thickness of 50 mm or greater, in detail, 60 mm to 100 mm, and more specifically, 80 mm to 100 mm.
Cooling
The steel sheet is cooled to a temperature of 700° C., or lower, after finish rolling.
In a case in which a cooling end temperature exceeds 700° C., a microstructure may not be properly formed, so that sufficient yield strength may be difficult to secure. For example, yield strength of 390 MPa or greater may be difficult to secure.
The cooling end temperature may be 300° C. to 600° C.
In a case in which the cooling end temperature is lower than 300° C., an increase in a generation amount of bainite may degrade toughness.
The steel sheet may be cooled at a cooling rate of the central portion of 1.5° C/s or higher. In a case in which the cooling rate of the central portion of the steel sheet is lower than 1.5° C/s, the microstructure may not be properly formed, so that it maybe difficult to secure sufficient yield strength. For example, yield strength of 390 MPa or greater may be difficult.
In addition, the steel sheet may be cooled at an average cooling rate of 2° C/s to 300° C/s.
Hereinafter, the present disclosure will be described in more detail through exemplary embodiments. However, an exemplary embodiment below is intended to describe the present disclosure in more detail through illustration thereof, but not to limit right scope of the present disclosure, as the scope of rights thereof is determined by the contents of the appended claims and those able to be reasonably inferred therefrom.
A steel slab having a composition illustrated in Table 1 below, which is 400 mm in thickness, was reheated to a temperature of 1060° C., and then rough rolling was performed at a temperature of 1025° C., thereby manufacturing a bar. A cumulative reduction ratio of 50% during rough rolling was equally applied to an entirety of steel grades.
A thickness of a bar having been rough rolled was 200 mm, while a grain size of a central portion after rough rolling before finish rolling, as illustrated in Table 2, was 75 μm to 89 μm. A reduction ratio of three final passes during rough rolling was within a range of 7.2% to 14.3%. A strain rate during rolling was within a range of 1.29/s to 1.66/s.
After rough rolling, finish rolling was performed at a temperature equal to a difference between a finish rolling temperature and an Ar3 temperature, illustrated in Table 2 below to obtain a steel sheet having a thickness illustrated in Table 3 below, and then the steel sheet was cooled to a temperature of 412° C. to 496° C. at a cooling rate of 4.5° C./sec.
In terms of the steel sheet manufactured as illustrated above, a microstructure, yield strength, a Kca value (a brittle crack arrestability coefficient), and a crack tip opening displacement (CTOD) value (a brittle crack initiation resistance) were examined, and results thereof are illustrated in Tables 3 and 4 below.
The Kca value in Table 4 is a value derived by performing an ESSO test on the steel sheet.
A FCAW (0.7 kJ/mm) welding process was performed to carry out structure analysis and a CTOD test on the HAZ, and results thereof were illustrated in Tables 3 and 4 below.
Surface properties illustrated in Table 3 below were measured to determine whether a star crack in a surface portion was generated by hot shortness occurring depending on a Cu to N addition ratio.
As illustrated in Tables 1 to 4, in the case of Comparative Example 1, a steel composition satisfies an exemplary embodiment, but the difference between the finish rolling temperature during finish rolling and the Ara temperature, proposed in an exemplary embodiment, was controlled to be 60° C. or higher. Since sufficient reduction was not applied to the central portion, a fraction of acicular ferrite (AF) in the central portion is less than 50%. In addition, cooling was started in an initial stage, so that ferrite of 30% or greater was not generated in a surface portion. Thus, it can be confirmed that the Kca value measured at a temperature of −10° C. may not exceed 6000 required in a steel material for shipbuilding of the related art.
In the case of Comparative Example 2, a C content had a value higher than an upper limit value of a C content of an exemplary embodiment. It can be confirmed that a relatively large amount of bainite was generated in the central portion during rough rolling, so an AF fraction of a final microstructure is less than 50%. Therefore, the Kca value measured at a temperature of −10° C. was 6000 or less. It can be confirmed that a relatively large amount of a martensite-austenite (MA) structure was also generated in the HAZ, so the CTOD value was 0.25 mm or less.
In the case of Comparative Example 3, a Sicontent had a value higher than an upper limit value of a Si content of an exemplary embodiment. It can be confirmed that a relatively large amount of Si was added to generate a relatively large amount of a coarse MA structure, so the microstructure in the central portion contains a relatively large amount of AF. However, the Kca value has a relatively low value similar to 6000 at a temperature of −10° C. It can be confirmed that a relatively large amount of MA is also generated in the HAZ, so the CTOD value is 0.25 mm or less.
In the case of Comparative Example 4, a Mn content has a value higher than an upper limit value of a Mn content of an exemplary embodiment. It can be confirmed that due to having a relatively high level of hardenability, a microstructure in a base material is provided as upper bainite, thereby allowing the fraction of AF to be less than 50%. Thus, the Kca value is 6000 or less at a temperature of −10° C.
In the case of Comparative Example 5, an Ni content had a value higher than an upper limit value of an Ni content of an exemplary embodiment. It can be confirmed that due to a relatively high level of hardenability, the microstructure of the base material is provided as granular bainite and upper bainite, and the fraction of acicular ferrite is less than 50%. Thus, the Kca value is 6000 or less at a temperature of −10° C.
In the case of Comparative Example 6, an Nb and Ti content has a value higher than an upper limit value of an Nb and Ti content of an exemplary embodiment. It can be confirmed that an entirety of other conditions satisfies a condition suggested in an exemplary embodiment, but due to a relatively high Nb and Ti content, a relatively large amount of the MA structure is generated in the HAZ, thereby allowing the CTOD value to be 0.25 mm or less.
Inventive Example 7 includes a component exceeding a ratio of Cu to Ni suggested in an aspect of the present disclosure. It can be confirmed that despite having other, significantly excellent physical properties, a star crack was generated, thereby causing a defect in surface quality.
In the case of Comparative Example 7, a C and Mn content has a value lower than a lower limit value of a C and Mn content of an exemplary embodiment. It can be confirmed that due to a relatively low level of hardenability, AF in the central portion is formed in an amount of less than 50%, and most structures have ferrite and a pearlite structure in an amount of 10% or greater. As pearlite has an average grain size of 15 μm or greater, the Kca value is 6000 or less at a temperature of −10° C.
On the other hand, it can be confirmed that, in the case of Inventive Examples 1 to 6, satisfying a composition range, a manufacturing range, and the Cu to Ni ratio of an exemplary embodiment, the fraction of AF of the microstructure in the central portion is 70% or greater, the fraction of pearlite in the central portion is 10% or less, a circle-equivalent diameter of pearlite in the central portion is 15 μm or less, and a fraction of MA phase in the HAZ is less than 5%.
It can be confirmed that, in Inventive Examples 1 to 6, yield strength satisfies 390 MPa or greater, the Kca value satisfies a value of 6000 or greater at a temperature of −10° C., and the CTOD value also represents a relatively high value of 0.25 mm or greater.
While exemplary embodiments have been shown and described above, it will be apparent to those skilled in the art that modifications and variations could be made without departing from the scope of the present invention as defined by the appended claims.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2015-0172687 | Dec 2015 | KR | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/KR2016/014124 | 12/2/2016 | WO | 00 |
