Patent Application
20230392229
The present disclosure relates to a steel strip and a method for manufacturing the same, and more particularly, to a steel strip having high strength characteristics and excellent workability and a method for manufacturing the same.
As a structural member of a commercial vehicle and heavy equipment according to the related art, a plate having a thickness of 12 to 14 mm and a tensile strength of 440 MPa or more and manufactured by a thick plate process has been mainly used, but recently, a technique using a high-strength steel material having a tensile strength of 550 MPa or more has been developed for weight reductions and high strength. In particular, an ultra-thick steel material having a thickness of 15 to 25 mm that is applied to a large commercial vehicle, a special vehicle, and a heavy equipment part has been manufactured by a thick plate process, but a measure to apply a hot rolling process has been required to secure price competitiveness.
However, when a high-strength ultra-thick steel material is manufactured in a hot rolling process, it is difficult to form a uniform microstructure due to difficulties under high pressure during rolling, such that it is difficult to secure a stable yield strength, cracks are likely to occur during manufacturing of parts, and durability lifespan may be deteriorated due to a local stress concentration during use.
In this regard, as for a steel material in the related art, Patent Document 1 has proposed a technique in which an austenite region is subjected to general hot rolling and then coiling is performed at a high temperature to form a ferrite phase as a matrix structure and a fine precipitate so as to secure strength and ductility, or Patent Document 2 has proposed a technique in which a coiling temperature is cooled to a temperature at which a bainite phase is formed as a matrix structure so as not to generate a coarse pearlite structure, and then coiling is performed. In addition, Patent Document 3 has proposed a technique of refining austenite grains through reductions of two or more times at 20 to 40% in a non-recrystallized region during hot rolling by using Ti, Nb, and the like.
However, alloy components such as Si, Mn, Al, Mo, and Cr, which are mainly used in the above techniques for manufacturing high-strength thick steel, are effective in improving strength, but when a large amount of alloy components are added, segregation and non-uniformity of the microstructure occur, resulting in deterioration of workability, and microcracks generated on a shear surface easily propagate in a fatigue environment, resulting in breakage of parts. In particular, as a thickness of the steel is increased, microstructure non-uniformity between a thickness surface portion and a central portion may be increased, such that a local stress concentration is increased and a propagation speed of cracks in a fatigue environment is also increased, resulting in deterioration of durability.
In addition, it is effective to use precipitate-forming elements such as Ti, Nb, and V in order to refine grains of the thick steel material and obtain a precipitation strengthening effect, and when a cooling rate is not controlled during cooling after coiling at a high temperature of 500 to 700° C. at which precipitates are easily formed or hot rolling, coarse carbides are formed in the thickness central portion of the thick steel material, such that the quality of the shear surface is deteriorated.
In addition, the application of a reduction amount of 20 to 40% twice or more in the non-recrystallized region during hot rolling may be easily applied to a thin product, but it is difficult to apply when manufacturing a thick product with a small total rolling reduction compared to the thin product.
An aspect of the present disclosure is to provide a high-strength steel strip having excellent yield strength and elongation and prevents cracks when formed because a uniform microstructure is secured during a hot-rolling process of a steel material, and a method for manufacturing the same.
An object of the present disclosure is not limited to the above description. The object of the present disclosure will be understood from the general contents of the present specification, and those skilled in the art to which the present disclosure pertains will have no difficulties in understanding the additional objects of the present disclosure.
According to an aspect of the present disclosure, a steel strip contains, by wt %, 0.05 to 0.15% of C, 0.01 to 1.0% of Si, 1.0 to 2.0% of Mn, 0.005 to 1.0% of Cr, 0.01 to 0.1% of Al, 0.001 to 0.02% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.005 to 0.11% of Ti, 0.005 to 0.07% of Nb, and a balance of Fe and unavoidable impurities,
wherein an R value defined in the following Relational Expression 1 satisfies 0.3 to 1.0,
a surface portion (where t represents a thickness of the steel strip) in a range of 0 to t/4 and a central portion (not including t/4) in a range of t/4 to t/2 based on a cross section each contain, by area %, 90% or more of ferrite and bainite in total, less than 5% of pearlite and carbides having a diameter of 0.5 μm or more, and less than 5% of a martensite and austenite (MA) phase, as a microstructure,
R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]−0.9×[Si]−1.5×[Nb]
[C]*=[C]−[C]×Q
Q=([Nb]/93+[Ti]/48)/([C]/12) [Relational Expression 1]
The thickness of the steel strip may be 15 mm or more.
The pearlite and the carbides having a diameter of 0.5 μm or more may be 3% or less and the MA phase may be 3% or less, in terms of area % in the central portion of the steel strip.
The bainite may be 20% or less, the pearlite and the carbides having a diameter of 0.5 μm or more may be less than 2%, and the MA phase may be 3% or less, in terms of area % in the surface portion of the steel strip.
A difference between an average hardness value and a maximum hardness value of hardness values measured at intervals of 0.5 mm from a point located at 0.5 mm directly below a surface of a specimen to a point located at 0.5 mm directly below a back surface based on an arbitrary line perpendicular to a thickness cross section of the steel strip may be 20 Hv or less.
According to another aspect of the present disclosure, a method for manufacturing a steel strip includes: reheating a steel slab containing, by wt %, 0.05 to 0.15% of C, 0.01 to 1.0% of Si, 1.0 to 2.0% of Mn, 0.005 to 1.0% of Cr, 0.01 to 0.1% of Al, 0.001 to 0.02% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.005 to 0.11% of Ti, 0.005 to 0.07% of Nb, and a balance of Fe and unavoidable impurities, and satisfying an R value defined in the following Relational Expression 1 of 0.3 to 1.0;
R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]−0.9×[Si]−1.5×[Nb]
[C]*=[C]−[C]×Q
Q=([Nb]/93+[Ti]/48)/([C]/12) [Relational Expression 1]
Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]80×[Si]−1.4×(t−8) [Relational Expression 2]
CRMin=76.6−157×[C]−25.2×[Si]−14.1×[Mn]−27.3×[Cr]+61×[Ti]+448×[Nb] [Relational Expression 3]
The reheating may be performed in a temperature range of 1,200 to 1,350° C.
During the first cooling, the cooling rate may be 80° C./sec or less.
During the second cooling, air cooling or water cooling may be performed to a temperature range of room temperature to 200° C.
As set forth above, according to an aspect of the present disclosure, it is possible to provide a steel strip having high strength characteristics and excellent workability because it has excellent tensile strength, yield strength, and elongation, and a method for manufacturing the same.
According to another aspect of the present disclosure, it is possible to provide a high-strength steel strip used for structural members of a large commercial vehicle, such as a wheel rim, a disk, members, and a frame, and a method for manufacturing the same.
Hereinafter, preferred exemplary embodiments in the present disclosure will be described. The exemplary embodiments in the present disclosure may be modified in various forms, and the scope of the present disclosure should not be interpreted to be limited to the exemplary embodiments set forth below. The exemplary embodiments are provided in order to describe the present disclosure in more detail to those skilled in the art to which the present disclosure pertains.
In order to solve the above problems, the inventors of the present disclosure have investigated a distribution of microstructures and detailed material changes for each thickness direction according to components and hot rolling and cooling conditions for ultra-thick rolled steel materials having various components.
As a result, the inventors of the present disclosure have found that a measure of imparting excellent yield strength and ductility to a thick hot-rolled steel strip, and in particular, have found that in a microstructure of a steel strip having a certain thickness or more, uniformity is secured and thus a hardness distribution in a thickness direction may be constant, thereby completing the present disclosure.
Hereinafter, the present disclosure will be described in detail.
Hereinafter, a steel composition of the present disclosure will be described in detail.
In the present disclosure, unless otherwise specified, % indicating a content of each element is based on weight.
A steel strip according to an aspect of the present disclosure may contain, by wt %, 0.05 to 0.15% of C, 0.01 to 1.0% of Si, 1.0 to 2.0% of Mn, 0.005 to 1.0% of Cr, 0.01 to 0.1% of Al, 0.001 to 0.02% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.005 to 0.11% of Ti, 0.005 to 0.07% of Nb, and a balance of Fe and unavoidable impurities.
Carbon (C): 0.05 to 0.15%
Carbon (C) is the most economical and effective element for strengthening steel, and when the amount of C added is increased, a precipitation strengthening effect or a bainite phase fraction is increased, resulting in an increase in tensile strength. When a thickness of a hot-rolled steel strip is increased, a cooling rate in a thickness central portion becomes slow during cooling after hot rolling, and thus, when a content of carbon (C) is large, coarse carbides or pearlite is easily formed. When the content of carbon (C) is less than 0.05%, it is difficult to obtain a sufficient strengthening effect, and when the content thereof exceeds 0.15%, coarse carbides or a pearlite phase and a band structure are formed in the thickness central portion, such that workability is deteriorated, durability is deteriorated, and weldability is also deteriorated.
Therefore, the content of carbon (C) may be 0.05 to 0.15%. More preferably, the content of carbon (C) may be 0.06% or more and may be 0.12% or less.
Silicon (Si): 0.01 to 1.0%
Silicon (Si) has an effect of deoxidizing molten steel and a solid solution strengthening effect, and is only an element advantageous for improving the workability by delaying formation of coarse carbides. When a content of silicon (Si) is less than 0.01%, the solid solution strengthening effect is insufficient, and the effect of delaying formation of carbides is also insufficient, such that it is difficult to improve the workability. When the content thereof exceeds 1.0%, a phase transformation temperature is increased, such that during hot rolling of a low-temperature region of an ultra-thick steel material, coarse grains are easily formed due to rolling of a local ferrite region in a surface portion, a red scale by silicon (Si) is formed on a surface of the steel strip, such that the surface quality of the steel strip is significantly deteriorated and ductility and weldability are also deteriorated.
Therefore, the content of silicon (Si) may be 0.01 to 1.0%. More preferably, the content of silicon (Si) may be 0.1% or more and may be 0.9% or less.
Manganese (Mn): 1.0 to 2.0%
Similar to Si, manganese (Mn) is an element that is effective in strengthening solid solution of steel, and facilitates formation of a bainite phase during cooling after hot rolling by increasing hardenability of steel. When a content of manganese (Mn) is less than 1.0%, the above effect according to the addition may not be obtained, and when the content thereof exceeds 2.0%, the hardenability is significantly increased, such that martensite phase transformation is likely to occur, and formation of pearlite is promoted during high-temperature coiling. In addition, in a casting process, a segregation portion is significantly developed in a thickness central portion when a slab is cast. When cooling is performed after hot rolling, a microstructure is non-uniformly formed in a thickness direction, such that workability and durability are deteriorated.
Therefore, the content of manganese (Mn) may be 1.0 to 2.0%. More preferably, the content of manganese (Mn) may be 1.1% or more.
Chromium (Cr): 0.005 to 1.0%
Chromium (Cr) is an element that strengthens solid solution of steel, and serves to help formation of bainite by delaying ferrite phase transformation during cooling. When a content of chromium (Cr) is less than 0.005%, the above effect according to the addition may not be obtained, and when the content thereof exceeds 1.0%, the ferrite transformation is excessively delayed to form a martensite phase, resulting in deterioration of an elongation. In addition, similar to Mn, the segregation portion is significantly developed in the thickness central portion, and a microstructure in the thickness direction is non-uniform, resulting in deterioration of the workability and durability.
Therefore, the content of chromium (Cr) may be 0.005 to 1.0%. More preferably, the content of chromium (Cr) may be 0.1% or more and may be 0.9% or less. [80]
Aluminum (Al): 0.01 to 0.1%
Aluminum (Al) is an element mainly added for deoxidation. When a content of aluminum (Al) is less than 0.01%, the addition effect is insufficient, and when the content thereof exceeds 0.1%, Al is combined with N to form AlN, such that corner cracks are likely to occur in the slab during continuous casting, and defects due to inclusion formation are likely to occur.
Therefore, the content of aluminum (Al) may be 0.01 to 0.1%.
Phosphorus (P): 0.001 to 0.02%
Similar to Si, phosphorus (P) is an element having solid solution strengthening and ferrite transformation promoting effects at the same time. When a content of phosphorus (P) exceeds 0.02%, brittleness occurs due to grain boundary segregation, and microcracks are likely to occur when formed, and the workability and durability are significantly deteriorated. On the other hand, in order to manufacture the steel strip with a content of phosphorus (P) of less than 0.001%, a lot of manufacturing cost is required, which is economically unfavorable and is insufficient to obtain strength.
Therefore, the content of phosphorus (P) may be 0.001 to 0.02%.
Sulfur (S): 0.001 to 0.01%
Sulfur (S) is an impurity present in steel. When a content of sulfur (S) exceeds 0.01%, S is combined with Mn and the like to form a non-metallic inclusion. Accordingly, microcracks are likely to occur during steel cutting processing, and the workability and durability are deteriorated. On the other hand, in order to manufacture the steel strip with a content of sulfur (S) of less than 0.001%, it takes a significant amount of time to perform a steelmaking process, and thus, productivity is reduced.
Therefore, the content of sulfur (S) may be 0.001 to 0.01%.
Nitrogen (N): 0.001 to 0.01%
Nitrogen (N) is a typical solid solution strengthening element together with C, and forms coarse precipitates together with Ti, Al, and the like. In general, although the solid solution strengthening effect of nitrogen (N) is superior to that of C, toughness is significantly decreased as the amount of nitrogen (N) in steel is increased. Therefore, an upper limit thereof is set to 0.01%. On the other hand, in order to manufacture the steel strip with a content of nitrogen (N) of less than 0.001%, it takes a significant amount of time to perform a steelmaking process, and thus, productivity is reduced.
Therefore, the content of nitrogen (N) may be 0.001 to 0.01%.
Titanium (Ti): 0.005 to 0.11%
Titanium (Ti) is a typical precipitation strengthening element, and forms coarse TiN in steel with strong affinity with N. TiN has an effect of suppressing a growth of grains during a heat process for hot rolling. In addition, titanium (Ti) remaining after reacting with N is solid-dissolved in steel and combined with C to form TiC precipitates, and thus, Ti is useful for improving the strength of steel. When a content of titanium (Ti) is less than 0.005%, the above effect is not obtained, and when the content thereof exceeds 0.11%, a local stress concentration occurs when formed due to generation of coarse TiN and coarsening of the precipitates, such that cracks are likely to occur.
Therefore, the content of titanium (Ti) may be 0.005 to 0.11%. More preferably, the content of titanium (Ti) may be 0.01% or more and may be 0.1% or less.
Niobium (Nb): 0.005 to 0.07%
Niobium (Nb) is a typical precipitation strengthening element together with Ti, and precipitates during hot rolling and thus is effective in improving the strength and impact toughness of steel due to a grain refinement effect caused by recrystallization delay. When a content of niobium (Nb) is less than 0.005%, the above effect is not obtained, and when the content thereof exceeds 0.07%, the workability and durability are deteriorated due to formation of elongated grains and formation of coarse composite precipitates caused by excessive recrystallization delay during hot rolling.
Therefore, the content of niobium (Nb) may be 0.005 to 0.07%. More preferably, the content of niobium (Nb) may be 0.01% or more.
The steel material of the present disclosure may contain a balance of iron (Fe) and unavoidable impurities in addition to the composition described above. Since the unavoidable impurities may be unintentionally incorporated in a general manufacturing process, the unavoidable impurities may not be excluded. Since these impurities are known to those skilled in a general steel manufacturing field, all the contents thereof are not particularly described in the present specification.
In the steel of the present disclosure, an R value defined in the following Relational Expression 1 may be 0.3 to 1.0.
When R in Relational Expression 1 is controlled, segregation of C, Mn, P, S, and the like and formation of MnS, which occur during solidification of the steel and cooling of the slab in the casting process, are minimized, such that uniformity of the microstructure may be improved. It is commonly known that the segregation of alloying elements such as C and Mn occurs in a cast structure formed during solidification, and P is mainly segregated at grain boundaries when the steel strip is hot-rolled and cooled and then maintained at a high temperature, which causes grain boundary embrittlement. Such segregation is highly dependent on the content of alloying elements. In particular, C and Mn form coarse carbides and a pearlite structure during cooling after hot rolling, which causes deterioration of the quality of the shear surface. In addition, Mn forms MnS, which is a non-metallic inclusion, together with Sn, and MnS is elongated during rolling, which causes significant deterioration of the workability of the final product. On the other hand, Si suppresses formation of coarse carbides and has a large solid solution strengthening effect even with a small amount of alloy, and Nb and Ti form fine precipitates and have an effect of refining a grain size, which is effective in solving the segregation and grain boundary embrittlement problems.
In the present disclosure, as a result of measuring the hardness at the cross section after manufacturing a thick steel material by hot rolling steel having various alloy components, it is confirmed that the uniformity of the microstructure, the alloy components, and the contents thereof are correlated, thereby deriving Relational Expression 1.
When the R value defined in the following Relational Expression 1 is less than 0.3, it is difficult to secure the physical properties targeted in the present disclosure, whereas when the value exceeds 1.0, the non-uniformity of the microstructure is increased, and thus, the hardness value at the cross section fluctuates greatly. More preferably, a lower limit of the R value may be 0.5 and an upper limit of the R value may be 0.8.
R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]−0.9×[Si]−1.5×[Nb]
[C]*=[C]−[C]×Q
Q=([Nb]/93+[Ti]/48)/([C]/12) [Relational Expression 1]
([C], [Mn], [P], [S], [Si], [Nb], and [Ti] in Relational Expression 1 represent wt % of the corresponding alloying elements, respectively)
Hereinafter, a microstructure of the steel of the present disclosure will be described in detail.
In the present disclosure, unless otherwise specified, % indicating a fraction of a microstructure is based on area.
In the steel satisfying the alloy composition according to an aspect of the present disclosure, a surface portion (where t represents a thickness of the steel strip) in a range of 0 to t/4 and a central portion (not including t/4) in a range of t/4 to t/2 based on a cross section each contain, by area %, 90% or more of ferrite and bainite in total, less than 5% of pearlite and carbides having a diameter of 0.5 μm or more, and less than 5% of a martensite and austenite (MA) phase, as a microstructure.
A microstructure of high-strength steel is determined during cooling. Bainite and a martensite and austenite (MA) phase are easily formed in a surface portion where a cooling rate is fast, whereas coarse carbides and pearlite are easily formed in a central portion where the cooling rate is slow.
In general, the MA phase formed in the surface portion is a hard phase and exhibits a higher hardness than that of the surrounding microstructure, such that a non-uniform hardness distribution occurs, and microcracks occur due to a difference in hardness between the MA phase and the matrix structure when formed. In addition, the coarse carbides and pearlite formed in the central portion exhibit a higher hardness than that of the surrounding microstructure and are simultaneously brittle, which causes microcracks during shear formation.
Therefore, in the present disclosure, in order to simultaneously solve the problems of the surface portion and the central portion, a fraction of the pearlite and carbides having a diameter of 0.5 μm or more is limited to less than 5%, and a fraction of the MA phase is limited to less than 5%. In this case, the fractions of the pearlite and carbides having a diameter of 0.5 μm or more and the MA phase may be equally applied to each of the surface portion and the central portion.
In the present disclosure, containing 90% or more of ferrite and bainite is to suppress formation of unnecessary coarse carbides and pearlite so as to have a uniform hardness distribution for each thickness position and to secure excellent yield strength and elongation, and when less than 90% of ferrite and bainite are contained, it is difficult to secure the product value (YSxT-El) of the yield strength and the elongation targeted in the present disclosure. Therefore, in the present disclosure, 90% or more of ferrite and bainite in total may be contained.
In terms of securing the physical properties targeted the present disclosure, more preferably, in the central portion, the pearlite and carbides having a diameter of 0.5 μm or more may be 3% or less and the MA phase may be 3% or less, and in the surface portion, the bainite may be 20% or less, the pearlite and carbides having a diameter of 0.5 μm or more may be less than 2%, and the MA phase may be 3% or less.
In the present disclosure, the microstructure has the same characteristics in the surface portion and the central portion of the steel, and the microstructure proposed in the present disclosure is equally applied to the entire steel. In addition, in the present disclosure, the surface portion means a region in a range of 0 to t/4 (t represents a thickness of the steel strip) based on the cross section, and the central portion means a region in a range of t/4 to t/2 (not including t/4).
Hereinafter, a method for manufacturing the steel of the present disclosure will be described in detail.
Steel according to an aspect of the present disclosure may be manufactured by subjecting a steel slab satisfying the alloy composition described above to reheating, hot rolling, first cooling, coiling, and second cooling.
Slab Reheating
A steel slab satisfying the alloy composition described above may be reheated in a temperature range of 1,200 to 1,350° C.
When the reheating temperature is lower than 1,200° C., precipitates are not sufficiently solid-dissolved again, such that formation of precipitates in a subsequent process is reduced, and coarse TiN remains. On the other hand, when the temperature exceeds 1,350° C., the strength of the steel is reduced due to an abnormal grain growth of austenite grains.
Hot Rolling
The reheated steel slab may be hot-rolled in a temperature range of 800 to 1,150° C. at a reduction ratio of 20 to 50%, and the rolling may be finished in a temperature range of Tn−50 to Tn defined in the following Relational Expression 2.
When the hot rolling temperature exceeds 1,150° C., the temperature of the steel strip is excessively increased, such that the grain size is coarsened and the surface quality of the hot-rolled steel strip is deteriorated. On the other hand, when the temperature is lower than 800° C., elongated grains are developed due to excessive recrystallization delay, resulting in severe anisotropy and deterioration of the workability, and when rolling is performed at a temperature below an austenite temperature range, a non-uniform microstructure is developed more severely. Therefore, microcracks are likely to occur in non-uniform portions when formed, which also causes deterioration of the ductility.
When a rolling end temperature exceeds Tn, the microstructure of the steel is coarsened and non-uniform. When the temperature is lower than Tn−50, in a high-strength ultra-thick steel strip having a thickness of 15 to 25 mm, a fraction of a fine ferrite phase is increased due to the promotion of ferrite phase transformation in the surface portion where the temperature is relatively low, but an elongated grain shape is formed, which causes cracks to propagate quickly, and a non-uniform microstructure may remain in the central portion, resulting in unfavorable durability.
The rolling end temperature determined by Relational Expression 2 of the present disclosure means a temperature of the hot-rolled steel strip at the end of hot rolling.
Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]80×[Si]−1.4×(t−8) [Relational Expression 2]
A reduction amount in the hot-rolling temperature range may be 20 to 50%.
When the reduction amount is less than 20%, it is difficult to obtain the recrystallization delay effect and thus non-uniform coarse grains are easily formed, and when the reduction amount exceeds 50%, an excessively elongated microstructure is formed and carbides are formed along the grain boundaries, such that cracks are likely to occur along the grain boundaries when formed. In addition, fine precipitates are also reduced and the precipitation strengthening effect is also reduced.
First Cooling and Coiling
The hot-rolled steel strip may be subjected to first cooling to a temperature range of 450 to 550° C. at a cooling rate equal to or higher than CRMin defined in the following Relational Expression 3, and then the cooled hot-rolled steel strip may be coiled.
A temperature range from immediately after the hot rolling to the cooling end temperature corresponds to a temperature section in which the ferrite phase transformation occurs during cooling. Since a cooling rate in the thickness central portion is slower than that in the thickness surface portion of the rolled strip, a coarse ferrite phase and coarse carbides are formed in the thickness central portion, and thus the steel strip has a non-uniform microstructure. Therefore, in order to suppress this, in the present disclosure, it is required to perform cooling faster than a specific cooling rate (CRMin). However, when an average cooling rate in the above temperature region exceeds 80° C./sec, a difference in cooling rate between the surface portion and the central portion is excessively increased, and a difference in hardness between the surface portion and the central portion is excessively increased, such that the workability and durability are deteriorated.
The cooling rate determined by Relational Expression 3 of the present disclosure means a cooling rate of the hot-rolled steel strip after hot rolling.
CRMin=76.6−157×[C]−25.2×[Si]−14.1×[Mn]−27.3×[Cr]+61×[Ti]+448×[Nb] [Relational Expression 3]
When the cooling end temperature and the coiling temperature exceed 550° C., a pearlite phase is formed as a band structure or a large amount of coarse carbides are formed, which causes deterioration of the workability and durability, and when the temperatures are lower than 450° C., a martensite phase and an MA phase are excessively formed, which causes deterioration of the workability and durability.
Second Cooling
The coiled steel strip may be subjected to second cooling to a temperature range of room temperature to 200° C., and the second cooling may be air cooling or water cooling.
In the present disclosure, air cooling means cooling performed in the air at room temperature and a cooling rate of 0.001 to 10° C./hour. Even in a case where the cooling rate exceeds 10° C./hour, when it complies with the coiling temperature and the first cooling conditions, transformation of some of the untransformed phases in the steel into an MA phase may be suppressed. Therefore, water cooling may be performed. In the present disclosure, water cooling means cooling performed by charging a coil into a water bath at room temperature. However, in order to control the cooling rate to less than 0.001° C./hour, a separate heating and heat preservation facility and the like are required, which is economically unfavorable. Therefore, a lower limit of the cooling rate may be 0.001° C./hour.
The steel strip of the present disclosure manufactured as described above is a steel strip having a thickness of 10 mm or more, more preferably, may have a thickness of 15 mm or more, and may be a steel strip having an upper limit of a thickness of 25 mm. A difference between an average hardness value and a maximum hardness value of hardness values measured at intervals of 0.5 mm from a point located at 0.5 mm directly below a surface of a specimen to a point located at 0.5 mm directly below a back surface based on an arbitrary line perpendicular to a thickness cross section of the steel strip may be 20 Hv or less, and more preferably, the average hardness value may be 160 to 300 Hv. In addition, since the product (YSxT-E) value of the yield strength and the elongation is 16,000 MPa·% or more, high strength and excellent workability may be provided.
Hereinafter, the present disclosure will be described in more detail with reference to Examples. However, the following Examples are provided to illustrate and describe the present disclosure in detail, but are not intended to limit the scope of the present disclosure.
The steel composition and the thickness of the final hot-rolled strip of each steel type are shown in Table 1. Table 2 shows the values of the rolling end temperature (FDT), the total reduction amount (%), the coiling temperature (CT), the cooling rate (CR*) up to the coiling temperature, which is the cooling end temperature after hot rolling, Tn and Tn−50 defined in Relational Expression 2, and the minimum cooling rate (CRMin) defined in Relational Expression 3 of the steel types shown in Table 1. The reheating temperature, the hot rolling temperature, and the cooling rate of the steel strip after coiling not shown in Table 2 were equally applied as 1,250° C., 800 to 1,150° C., and 1° C./hour, respectively.
R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]−0.9×[Si]−1.5×[Nb]
[C]*=[C]−[C]×Q
Q=([Nb]/93+[Ti]/48)/([C]/12) [Relational Expression 1]
Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]80×[Si]−1.4×(t−8) [Relational Expression 2]
CRMin =76.6−157×[C]−25.2×[Si]−14.1×[Mn]−27.3×[Cr]+61×[Ti]+448×[Nb] [Relational Expression 3]
The microstructure characteristics and mechanical properties of the steel types are shown in Tables 3 and 4.
The microstructure shown in Table 3 is a result of analysis at a point located at 0.5 mm directly below a surface and a central portion of the hot-rolled strip. In the present disclosure, the surface portion means a range of 0 to t/4 based on the thickness (t), and the central portion means a range of t/4 to t/2 (not including t/4). In Table 3, the microstructure in the surface portion is a result of analysis at a point located at 0.5 mm directly below the surface, and the microstructure in the central portion is a result of analysis at t/2, which is a thickness central portion. The area fraction of the MA phase was measured after etching by Lepera etching method using an optical microscope and an image analyzer and was analyzed at 1,000 magnification. The area fractions of the martensite and austenite phase (MA), ferrite phase (F), bainite phase (B), and pearlite phase (P) were analyzed using a scanning electron microscope (SEM) at 3,000 to 5,000 magnification. Here, ferrite (F) is polygonal ferrite having an equiaxed crystal shape, and bainite (B) means a ferrite phase observed in a low-temperature region such as bainite, acicular ferrite, or bainitic ferrite. In addition, an area fraction of pearlite (P) means the sum of area fractions of pearlite and carbides having a size of 0.5 μm or more.
YS, TS, and T-El in Table 4 mean 0.2% off-set yield strength, tensile strength, elongation at break, respectively, which are test results obtained by taking a JIS No. 5 standard test piece parallel to a rolling direction. In addition, the hardness at a cross section of the specimen is measured and shown together. The hardness was measured at intervals of 0.5 mm from a point located at 0.5 mm directly below a surface of the specimen to a point located at 0.5 mm directly below a back surface based on an arbitrary line perpendicular to a thickness cross section of the specimen with a Micro-vickers tester, and a load of 500 g was applied. Table 4 shows a maximum hardness value and an average hardness value at the thickness cross section among the measured hardness values, and a difference between two hardness values is shown. Peak (number) means the number of portions where the difference between the hardness value and the average hardness value at the thickness point exceeds 20 Hv.
As shown in Table 4, in Inventive Steels 1 to 7 satisfying the alloy composition, manufacturing method, and Relational Expressions 1 to 3 proposed in the present disclosure, all the mechanical properties targeted in the present disclosure were secured.
In the cases of Comparative Steels 1 to 4, Relational Expression 1 proposed in the present disclosure was not satisfied, and in the case of Comparative Steel 1, the content of C satisfied the range of the present disclosure, but was out of the proposed range of Relational Expression 1 considering segregation. Therefore, excessive pearlite was formed over the central portion and the surface portion of the microstructure, and when the hardness was measured in the thickness direction, a locally high hardness difference was shown. The ductility was also insufficient, and thus, the result out of the range proposed in the present disclosure was shown. In the cases of Comparative Steels 2 and 3, the component range of Mn was out of the range proposed in the present disclosure, and in addition, Relational Expression 1 was not satisfied. In the case of Comparative Steel 2, since the content of Mn was small, segregation in the thickness direction of the rolled steel strip or coarse carbides and non-uniform pearlite were not formed, but the yield strength and the tensile strength were insufficient, and thus, the properties targeted in the present disclosure were not obtained. In the case of Comparative Steel 3, since the content of Mn was excessive, bainite was formed in the surface portion due to high hardenability, whereas pearlite was excessively formed in the central portion, and elongated MnS inclusions were also observed. In particular, when the hardness was measured in the thickness direction, a locally high hardness difference was exhibited, and the ductility was also insufficient. Comparative Steel 4 was a case in which the content of P was out of the range proposed in the present disclosure, and Relational Expression 1 was not satisfied at the same time. In the microstructure of Comparative Steel 4, the range proposed in the present disclosure was satisfied, and the strength and the elongation were also excellent, but when the hardness was measured, a local hardness difference was exhibited, which may cause a high possibility of brittleness when used after manufacturing parts.
In the case of Comparative Steel 5 in which Relational Expression 1 was satisfied, but the content of Si of the present disclosure was not satisfied, it was confirmed that coarse ferrite was formed in the surface portion in the microstructure, and an MA phase was formed in the surface portion and the central portion. In addition, a slightly low hardness value was exhibited in the surface portion, and the product of the yield strength and the elongation was out of the targeted range of the present disclosure. This was because the phase transformation temperature was increased due to the excessive addition of Si, and ferrite was formed in the surface portion during hot rolling and subjected to two-phase rolling, and some untransformed ferrite was formed as an MA phase.
Comparative Steel 6 satisfied the alloy component range of the present disclosure, but did not satisfy Relational Expression 1. In this case, segregation of the components was not observed, an MA phase and coarse carbides were hardly formed in the microstructure, and only fine pearlite was observed around grain boundaries. Therefore, the hardness distribution in the thickness direction was also relatively uniform. However, the hardness value targeted in the present disclosure was not secured.
Comparative Steels 7 and 8 did not satisfy Relational Expression 2 and the reduction ratio. In the case of Comparative Steel 7, the rolling was terminated in a temperature range satisfying Relational Expression 2, but a non-uniform microstructure was formed during cooling due to an insufficient reduction ratio. Therefore, the compositional fraction of the microstructure satisfied the present disclosure, but coarse grains were mixed in the ferrite matrix structure, resulting in a low yield strength. The durability of the steel having such a microstructure may be deteriorated during use of parts. In the case of Comparative Steel 8 in which both Relational Expression 2 and reduction ratio conditions were not satisfied, it could be confirmed that an excessively elongated microstructure was formed in the surface portion due to delay of recrystallization during rolling caused by a large reduction amount, whereas equiaxed ferrite and pearlite were mainly formed in the central portion to form a non-uniform microstructure depending on the thickness portion, which caused the deterioration of the durability of the parts and the deterioration of the elongation.
In Comparative Steels 9 and 10, the coiling temperature conditions were not satisfied. In Comparative Steel 9, the cooling end temperature and the coiling temperature were higher than those of the temperature ranges proposed in the present disclosure, it was confirmed that pearlite was locally formed, and in particular, the pearlite band structure was observed in the central portion. Therefore, when the hardness in the thickness direction was measured, a locally high hardness difference was shown. In Comparative Steel 10, the cooling end temperature and the coiling temperature were lower than those of the ranges proposed in the present disclosure. In Comparative Steel 10, bainite was excessively formed in the microstructure in the surface portion, and the elongation was insufficient.
In Comparative Steel 11 in which the cooling rate conditions of the cooling rate relational expression 3 were not satisfied and the cooling rate during cooling after hot rolling was lower than the range of the present disclosure, pearlite and coarse carbides were formed in the thickness central portion, and thus a locally high hardness difference was shown.
In Comparative Steels 12 and 13, the reduction amount and cooling end temperature conditions were not satisfied. In Comparative Steel 12, the reduction amount was insufficient in a temperature range in which recrystallization was delayed during hot rolling, and the coiling temperature was low, such that the size of the ferrite grains was non-uniform and bainite in the microstructure in the surface portion was excessively formed. In addition, pearlite was also locally observed in the central portion, resulting in a low elongation. In Comparative Steel 13, the reduction amount was insufficient in the temperature region in which crystallization was delayed, the coiling temperature was high, and the cooling rate did not satisfy Relational Expression 3. Therefore, it was confirmed that the microstructure was non-uniform, pearlite was formed as a band structure, and the yield strength was low.
Hereinabove, the present disclosure has been described in detail by the exemplary embodiments, but other exemplary embodiments having different forms are possible. Therefore, the technical spirit and scope of the claims set forth below are not limited by the exemplary embodiments.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2020-0138314 | Oct 2020 | KR | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/KR2021/014695 | 10/20/2021 | WO |
