Patent Application
20250207215
The present invention relates to a high-strength steel sheet (specifically, hot-rolled steel sheet) that is suitable as an automotive member and improved particularly in terms of strength and fatigue resistance and a method for producing the high-strength steel sheet.
Reductions in CO2 emissions have been anticipated in a worldwide framework from the viewpoints of global environment conservation. In particular, there has been a strong demand for increases in the mileage of automobiles. Reductions in the weights of automotive bodies have been anticipated. It is effective to increase the strengths of steel sheets used as materials for automotive members and reduce the thicknesses of the steel sheets for reducing the weights of automotive bodies without reducing the strengths of the automotive bodies. In particular, a steel sheet having a tensile strength of 980 MPa or more has been considered as a promising material that enables the reductions in the thicknesses of the steel sheets and thereby markedly increases the mileage of automobiles.
In order to maintain durability, which is likely to degrade with reductions in the thicknesses of automotive parts, it is necessary to enhance the fatigue resistance of steel sheets. Automotive parts, that is, specifically, undercarriage parts, such as suspension parts, are subjected to cyclic loading through tires. Therefore, if the fatigue strengths of the above parts are low, the durability of the parts may fall below the designed durability with an increase in the distance travelled. Commonly, an increase in the strength of a steel sheet does not always result in an increase in the fatigue strength of the steel sheet.
For enhancing the fatigue strength of a steel sheet while increasing the tensile strength of the steel sheet, various studies (Patent Literatures 1 to 3) have been conducted in the related art.
The techniques known in the related art, such as those described in Patent Literatures 1 to 3, have the following issues.
In the technique described in Patent Literature 1, a tensile strength of 980 MPa or more cannot be achieved.
In the technique described in Patent Literature 2, a fatigue strength at which the steel sheet is practically used as an automotive part is not studied sufficiently.
In the technique described in Patent Literature 3, a high-strength steel sheet having excellent fatigue resistance is reportedly produced. However, fatigue resistance is not described specifically.
Thus, a technique related to a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent fatigue resistance has not been established in the related art.
Accordingly, aspects of the present invention were developed in light of the above-described circumstances. An object of aspects of the present invention is to provide a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and a markedly high fatigue strength and a method for producing the high-strength hot-rolled steel sheet.
In order to achieve the above object, the inventors of the present invention conducted extensive studies of a technique for enhancing the fatigue resistance of a hot-rolled steel sheet while maintaining a tensile strength of 980 MPa or more and consequently found the following facts. Specifically, a microstructure that includes upper bainite as a primary phase and an adequate amount of a martensite and/or retained austenite phase, which serves as a hard secondary phase, is formed. Furthermore, the dislocation density in all the phases included in the surface layer region which extends from the surface of the steel sheet to the position 1/10 of the thickness of the steel sheet increases. Moreover, the grain sizes of all the phases are controlled. This enables a steel sheet having a high strength of 980 MPa or more and a markedly high fatigue strength to be formed subsequent to a heat treatment that corresponds to baking coating.
Note that the term “upper bainite phase” refers to a microstructure phase that is an assembly of lath-shaped ferrite grains having a misorientation of less than 15° and includes a Fe-based carbide and/or a retained austenite phase interposed between the lath-shaped ferrite grains. Note that the above microstructure phase does not always include Fe-based carbide and/or retained austenite interposed between the lath-shaped ferrite grains.
Since lath-shaped ferrite grains have a lath-like shape and the inside of the grains has a relatively high dislocation density unlike lamellar (laminar) ferrite or polygonal ferrite in pearlite, they can be distinguished from each other using a SEM (scanning electron microscope) or a TEM (transmission electron microscope).
In the case where retained austenite is present between lath grains, only the lath-shaped ferrite parts are considered as upper bainite and distinguished from retained austenite.
A martensite and/or retained austenite phase appears brighter in a SEM image than an upper bainite phase, a lower bainite phase, or a polygonal ferrite phase. Thus, a martensite phase and/or retained austenite phase can be distinguished from the above microstructure phase using a SEM.
Although martensite and retained austenite phases appear to have the same degrees of brightness with a SEM, they can be distinguished from each other using an electron backscatter diffraction patterns (EBSD) method.
The dislocation density can be determined by irradiating a steel material with an X-ray and analyzing the resulting strength curve (line profile) with respect to the angle or energy of the diffracted X-ray. The analysis of line profile is conducted in accordance with “Method for Evaluating Dislocation Density Using X-ray Diffraction”, Material and Process, vol. 17 (2004) No. 3, p. 396-399. In accordance with aspects of the present invention, the dislocation density is calculated from the half-widths of (110), (211), and (220).
The inventors of the present invention conducted further studies on the basis of the above findings and consequently devised aspects of the present invention. The summary of aspects of the present invention is as follows.
According to aspects of the present invention, a high-strength steel sheet having a tensile strength of 980 MPa or more and excellent fatigue resistance and a method for producing the high-strength steel sheet can be provided.
Applying the high-strength steel sheet according to aspects of the present invention to automotive undercarriage parts, such as a suspension, structural parts, framework parts, or truck frame parts enhances safety, allows reductions in the weights of the automotive bodies, and therefore produces markedly advantageous effects from the viewpoint of industry.
In accordance with aspects of the present invention, the expression “excellent fatigue resistance” means that the ratio (fatigue limit ratio) of the fatigue strength at 2×106 cycles of plane bending performed in an alternating plane bending fatigue test relative to the tensile strength is 0.50 or more.
The FIGURE is a schematic diagram illustrating the shape of a test specimen used for a plane bending fatigue test in accordance with aspects of the present invention.
Embodiments of the present invention are described below. Note that the following description is intended to be illustrative of preferable embodiments of the present invention, and the present invention is not limited by the following embodiments.
The steel sheet has the chemical composition described below. In the following description, the symbol “%” used as the unit of the content of an element in the chemical composition means “% by mass” unless otherwise specified.
C is an element that effectively facilitates the formation of bainite and increases strength by enhancing hardenability. If the C content is less than 0.03%, the above advantageous effects are not produced to a sufficient degree, and a tensile strength of 980 MPa or more cannot be achieved. Accordingly, the C content is 0.03% or more, is preferably 0.04% or more, and is more preferably 0.05% or more. On the other hand, if the C content is more than 0.15%, the amount of martensite and retained austenite increases and, consequently, sufficiently high fatigue resistance cannot be achieved. Accordingly, the C content is 0.15% or less, is preferably 0.14% or less, and is more preferably 0.13% or less.
Si contributes to an increase in steel strength by solid solution strengthening of steel. Accordingly, the Si content is 0.1% or more, is preferably 0.3% or more, and is more preferably 0.5% or more. However, since Si is an element that facilitates the formation of ferrite, a Si content exceeding 3.0% causes the formation of ferrite and degrades fatigue resistance. Accordingly, the Si content is 3.0% or less, is preferably 2.5% or less, and is more preferably 2.0% or less.
Mn is an element that stabilizes austenite. Mn is also an element effective for suppressing the formation of ferrite and increasing strength. If the Mn content is less than 0.8%, the above advantageous effects are not produced to a sufficient degree and ferrite, etc. are formed. Consequently, a tensile strength of 980 MPa or more cannot be achieved. Accordingly, the Mn content is 0.8% or more, is preferably 1.0% or more, and is more preferably 1.2% or more. On the other hand, if the Mn content is more than 3.0%, the amount of martensite and retained austenite increases and, consequently, sufficiently high fatigue resistance cannot be achieved. Accordingly, the Mn content is 3.0% or less, is preferably 2.8% or less, and is more preferably 2.5% or less.
Since P degrades weldability, it is desirable to minimize the P content. The maximum P content allowable in accordance with aspects of the present invention is 0.1%. Thus, the P content is 0.1% or less. Since a P content of less than 0.001% reduces production efficiency, the lower limit is set to 0.001% or more.
Since S degrades weldability, it is desirable to minimize the S content. The maximum S content allowable in accordance with aspects of the present invention is 0.03%. Thus, the S content is 0.03% or less. Since a S content of less than 0.0001% reduces production efficiency, the lower limit is set to 0.0001% or more.
Al is an element that serves as a deoxidizing agent and effectively enhances the cleanliness of steel. If the Al content is excessively low, the advantageous effects cannot always be produced to a sufficient degree. Accordingly, the Al content is 0.001% or more, is preferably 0.01% or more, and is more preferably 0.02% or more. However, since Al is an element that facilitates the formation of ferrite, an Al content exceeding 2.0% causes the formation of ferrite and reduces a fatigue strength. Accordingly, the Al content is 2.0% or less, is preferably 1.8% or less, and is more preferably 1.6% or less.
N binds to an element that forms a nitride to precipitate in the form of a nitride and thereby contributes to a reduction in grain size. For producing the above advantageous effects, the N content needs to be 0.001% or more. However, since N is likely to bind to Ti at high temperatures to form coarse nitride particles, an excessively high N content degrades fatigue resistance. Accordingly, the N content is 0.01% or less, is preferably 0.008% or less, and is more preferably 0.006% or less.
B is an element that effectively facilitates the formation of upper bainite and increases the strength of the steel sheet by segregating at prior-austenite grain boundaries to suppress the formation of ferrite. For producing the above advantageous effects, the B content needs to be 0.0002% or more. Therefore, the B content is 0.0002% or more, is preferably 0.0005% or more, and is more preferably 0.0007% or more. However, if the B content is more than 0.010%, the above advantageous effects become saturated. Thus, the B content is 0.010% or less, is preferably 0.009% or less, and is more preferably 0.008% or less.
Ti and Nb are elements that form a carbide and effectively increase strength by precipitation strengthening. Therefore, one or more elements selected from Ti and Nb need to be included in the chemical composition. The lower limits for the Ti and Nb contents are set to Ti: 0.01% or more and Nb: 0.001% or more. The Ti and Nb contents are preferably Ti: 0.02% or more and Nb: 0.002% or more and are more preferably Ti: 0.03% or more and Nb: 0.003% or more. However, if the Ti and Nb contents are more than Ti: 0.30% and Nb: 0.10%, carbide particles become coarsened to degrade hardenability and, consequently, it may become impossible to form the steel microstructure intended in accordance with aspects of the present invention. Accordingly, the upper limits for the Ti and Nb contents are set to Ti: 0.30% or less and Nb: 0.10% or less. The Ti and Nb contents are preferably Ti: 0.25% or less and Nb: 0.08% or less and are more preferably Ti: 0.20% or less and Nb: 0.05% or less.
The balance includes Fe and incidental impurities.
The above are constituents of the fundamental chemical composition of the high-strength steel sheet according to aspects of the present invention. The chemical composition of the high-strength steel sheet may optionally contain the following elements as needed.
Cr, Ni, Cu, V, and Mo are elements that stabilize austenite and are also elements effective for suppressing the formation of ferrite and increasing strength. In order to produce the above advantageous effects, one or more elements selected from the above elements are preferably included in the chemical composition. In the case where one or more elements selected from Cr, Ni, Cu, V, and Mo are included in the chemical composition, the contents of the above elements are preferably Cu: 0.005% to 2.0%, Ni: 0.005% to 2.0%, Cr: 0.005% to 2.5%, V: 0.001% to 0.5%, and Mo: 0.005% to 1.0%. If the Cr, Ni, Cu, V, and Mo contents are more than the respective upper limits described above, martensite and retained austenite are likely to remain and, consequently, it may become impossible to form the steel microstructure intended in accordance with aspects of the present invention. The lower limit for the Cr content is more preferably 0.1% or more. The upper limit for the Cu content is more preferably 0.6% or less. The lower limit for the Ni content is more preferably 0.1% or more. The upper limit for the Ni content is more preferably 0.6% or less. The lower limit for the Cu content is more preferably 0.1% or more. The upper limit for the Cu content is more preferably 0.6% or less. The lower limit for the V content is more preferably 0.005% or more. The upper limit for the V content is more preferably 0.3% or less. The lower limit for the Mo content is more preferably 0.1% or more. The upper limit for the Mo content is more preferably 0.5% or less.
Sb is an element that reduces the likelihood of elements being removed from the surface of a steel material when the steel material is heated and thereby effectively limits a reduction in steel strength. Accordingly, in the case where the chemical composition contains Sb, the Sb content is preferably 0.005% to 0.2%. If the Sb content is more than the above upper limit, embrittlement of the steel sheet may occur. The lower limit for the Sb content is more preferably 0.01% or more. The upper limit for the Sb content is more preferably 0.050% or less.
Sn is an element that suppresses the formation of pearlite and thereby effectively limits a reduction in steel strength. In order to produce the above advantageous effects, in the case where the chemical composition contains Sn, the Sn content is preferably 0.001% to 0.05%. If the Sn content is more than the above upper limit, embrittlement of the steel sheet may occur. The lower limit for the Sn content is more preferably 0.005% or more. The upper limit for the Sn content is more preferably 0.03% or less.
Ca, Mg, and REMs are elements that effectively enhance workability by shape control of inclusions. In order to produce the above advantageous effects, one or more elements selected from the above elements are preferably included in the chemical composition. In the case where the chemical composition includes one or more elements selected from Ca, Mg, and REMs, the contents of the above elements are preferably Ca: 0.0005% to 0.01%, Mg: 0.0005% to 0.01%, and REM: 0.0005% to 0.01%. However, if the Ca, Mg, and REM contents are more than the respective upper limits, the amount of inclusion increases and workability may become degraded consequently. The lower limit for the Ca content is more preferably 0.001% or more. The upper limit for the Ca content is more preferably 0.005% or less. The lower limit for the Mg content is more preferably 0.001% or more. The upper limit for the Mg content is more preferably 0.005% or less. The lower limit for the REM content is more preferably 0.001% or more. The upper limit for the REM content is more preferably 0.005% or less. Note that the term “REM (rare-earth element)” used herein refers collectively to Sc, Y, and the 15 elements from lanthanum (La) with an atomic number of 57 to lutetium (Lu) with an atomic number of 71. The term “REM content” used herein refers to the total content of these elements.
The advantageous effects according to aspects of the present invention are not impaired even when the Mo, V, Cr, Ni, Cu, Sb, Sn, Ca, Mg, and REM contents are less than the respective lower limits described above. Thus, when the contents of the above constituents are less than the respective lower limits described above, it is considered that the chemical composition contains the above elements as incidental impurities.
The microstructure of the high-strength steel sheet according to aspects of the present invention is described below.
A surface layer region of the high-strength steel sheet according to aspects of the present invention which extends from the surface of the steel sheet to the position 1/10 of the thickness of the steel sheet has the microstructure described below. Specifically, an upper bainite phase, the area fraction of which is 75% or more and less than 98.5%, is the primary phase, and a microstructure phase consisting of a martensite phase and/or a retained austenite phase, the area fraction of which is 1.5% or more and less than 25%, is the secondary phase. Furthermore, the average grain size of all the phases included in the surface layer region is 6.0 μm or less, and the dislocation density in all the phases is 8.0×1014/m2 or more.
<Upper Bainite Phase: 75% or More and Less than 98.5% by Area>
The microstructure of the high-strength steel sheet according to aspects of the present invention includes upper bainite as a primary phase. If the area fraction of upper bainite is less than 75%, a markedly high fatigue strength cannot be achieved. Accordingly, the lower limit for the area fraction of upper bainite is set to 75% or more and is preferably 85% or more. However, if the upper bainite phase is 98.5% or more, the intended dislocation density cannot be achieved. Accordingly, the upper limit for the area fraction of upper bainite is set to be less than 98.5% and is preferably 97% or less.
<Martensite Phase and/or Retained Austenite Phase: 1.5% or More and Less Than 25% by Area>
The microstructure of the high-strength steel sheet according to aspects of the present invention includes a martensite phase and/or a retained austenite phase. If the martensite phase and/or retained austenite phase is less than 1.5%, it becomes impossible to achieve a tensile strength of 980 MPa or more and excellent fatigue resistance. On the other hand, if the area fraction of martensite and/or retained austenite is 25% or more, the amount of interfaces between martensite and/or retained austenite and upper bainite, which may act as the origins of fatigue cracks, increases and, consequently, fatigue resistance may become degraded. For the above reasons, it is necessary to limit the total area fraction of martensite and/or retained austenite to be less than 25%. The above total area fraction is preferably 20% or less and is more preferably 15% or less. Note that, in accordance with aspects of the present invention, the term “martensite” refers to as-quenched martensite.
The advantageous effects according to aspects of the present invention are not impaired when the area fraction of the remaining microstructure phase, which is a phase other than any of upper bainite and martensite and/or retained austenite, is 2.0% or less at the maximum. The remaining microstructure phase includes known microstructure phases, such as ferrite and pearlite.
It is considered that fatigue cracks are formed as a result of slip deformation occurring in the crystal grains included in the surface layer. Grain boundaries reduce the likelihood of the slip deformation being propagated to adjacent crystal grains and consequently delay the occurrence of cracking. That is, fatigue strength can be increased by reducing the size of crystal grains. Reducing the grain size also contributes to an increase in strength. Accordingly, the average grain size is 6.0 μm or less and is preferably 5.0 μm or less. However, if the average grain size is excessively small, elongation may be reduced with an increase in strength. Accordingly, the average grain size is preferably 2.0 μm or more. Note that the term “average grain size” used herein refers to the average grain size of all the phases included in the surface layer region of the steel sheet which extends from the surface of the steel sheet to the position 1/10 of the thickness of the steel sheet. In the case where the surface layer region, which extends to the position 1/10 of the thickness, includes the remaining microstructure phase, the remaining microstructure phase is also included in the “all the phases”.
<Dislocation Density: 8.0×1014/m2 or More>
Most of the fatigue cracks occur at the surface of the steel sheet. After the fatigue cracks have grown to several tens of micrometers in length, a fatigue crack propagation stage starts. In high-cycle fatigue, fatigue life is primarily affected by the number of the cycles performed until the occurrence of cracking. Therefore, for increasing a fatigue strength at 2×106 cycles, it is necessary to reduce the formation of cracks. It is important to control the dislocation behavior of the surface layer region, which extends from the surface of the steel sheet to the position 1/10 of the thickness. In the high-strength steel sheet according to aspects of the present invention, the dislocations introduced to the microstructure are pinned as a result of the heat treatment being performed in the subsequent step to obstruct the movement of the dislocations. This prevents the movement and rearrangement of the dislocations, delays cyclic softening, and consequently enhances fatigue resistance. In order to produce the above advantageous effects, the dislocation density is limited to 8.0×1014/m2 or more. The dislocation density is preferably 1.0×1015/m2 or more and is more preferably 1.2×1015/m2 or more. The upper limit for the dislocation density is not set but preferably 4.0×1015/m2 or less. Although it is most important to control the dislocation density in the primary phase included in the surface layer region, which extends from the surface of the steel sheet to the position 1/10 of the thickness, it is difficult to measure the dislocation density in the primary phase only. Therefore, the dislocation density measured in accordance with aspects of the present invention is the dislocation density in all the phases included in the surface layer region, which extends to the position 1/10 of the thickness. In the case where the surface layer region, which extends to the position 1/10 of the thickness, includes the remaining microstructure phase, the remaining microstructure phase is also included in the “all the phases”.
The high-strength steel sheet according to aspects of the present invention has a tensile strength of 980 MPa or more and a fatigue limit ratio of 0.50 or more. The term “fatigue limit ratio” used herein refers to the ratio of the fatigue strength at 2×106 cycles of plane bending to the tensile strength. Thus, the high-strength steel sheet according to aspects of the present invention has a high tensile strength and is capable of maintaining safety even when the thickness of the steel sheet is reduced. The high-strength steel sheet according to aspects of the present invention can be applied to members for trucks or automobiles.
Note that, in accordance with aspects of the present invention, the area fractions and mechanical properties of the above microstructure phases are determined by the methods described in Examples below.
A method for producing the high-strength steel sheet according to an embodiment of the present invention is described below. In the following description, the symbol “° C.” used for describing temperature refers to the surface temperature of the object (steel material or steel sheet) unless otherwise specified.
The high-strength steel sheet according to aspects of the present invention can be produced by subjecting a steel material to the treatments (1) to (6) below in order. Each of the steps is described below.
Note that the steel material may be any steel material having the above-described chemical composition. The high-strength steel sheet that is to be produced finally has the same chemical composition as the steel material used. The steel material may be, for example, a steel slab. The method for producing the steel material is not limited. For example, a molten steel having the above-described chemical composition is prepared using a known method, such as a converter, and the molten steel is formed into a steel material by a casting method, such as continuous casting. A method other than continuous casting, such as ingot casting-blooming rolling, may also be used. Alternatively, steel scrap may also be used as a raw material. After the steel material has been produced by continuous casting or the like, it may be directly subjected to the subsequent heating step. In another case, the steel material may be cooled to prepare warm or cold steel pieces, which are subjected to the heating step.
First, the steel material is heated to a heating temperature of 1150° C. or more. In the steel material cooled to low temperatures, most of the carbonitride-forming elements, such as Ti, are present in the form of coarse carbonitride particles in a nonuniform manner. The presence of the coarse and nonuniform precipitates degrades various properties (e.g., strength and fatigue resistance) commonly required for high-strength steel sheets used for producing parts for trucks or automobiles. Therefore, it is necessary to heat the steel material prior to hot rolling to dissolve the coarse precipitates. Accordingly, the temperature to which the steel material is heated is 1150° C. or more. The above heating temperature is preferably 1180° C. or more and is more preferably 1200° C. or more. However, if the temperature to which the steel material is heated is excessively high, slab flaws may occur. Furthermore, yields may be reduced due to descaling. Accordingly, the temperature to which the steel material is heated is preferably 1350° C. or less, is more preferably 1300° C. or less, and is further preferably 1280° C. or less.
In the heating step, it is preferable to hold the temperature of the steel material at the above heating temperature after the temperature of the steel material has been increased to the heating temperature in order to make the temperature of the steel material uniform. Although the amount of time (holding time) during which the temperature of the steel material is held at the heating temperature is not limited, the holding time is preferably 1800 seconds or more in order to increase the uniformity in the temperature of the steel material. However, if the holding time is more than 10000 seconds, the amount of the scale generated may be increased. This increases, for example, the possibility of entanglement of the scale in the subsequent hot rolling step and consequently may reduce yields due to surface flaw defects. Accordingly, the holding time is preferably 10000 seconds or less and is more preferably 8000 seconds or less. Alternatively, subsequent to casting, the steel material that has not been hot-rolled may be directly subjected to hot rolling (hot direct rolling) while the temperature of the steel material is high (i.e., while the temperature of the steel material is held to fall within the above heating temperature range).
The heated steel material (or, the as-cast steel material having a high temperature) is subjected to a hot rolling step in which rough rolling and finish rolling are performed. The conditions under which the rough rolling is performed are not limited and may be any conditions with which the resulting sheet bar has intended dimensions. The steel material is rough-rolled to form a rough-rolled steel sheet bar. Prior to finish rolling, the rough-rolled steel sheet bar may be subjected to descaling (high-pressure water descaling) in which high-pressure water is sprayed at the entry side of the finish rolling mill.
In accordance with aspects of the present invention, when temperatures RC1 and RC2 are defined by Formulae (1) and (2) below, finish rolling is performed such that the total of the rolling reductions achieved in the temperature range of equal to or more than (RC1−150)° C. and equal to or less than RC1° C. is 35% or more. The amount of residence time during which the temperature is held in the temperature range is not limited and may be 3 seconds or more and 20 seconds or less. Moreover, the finish rolling delivery temperature is set to a temperature equal to or more than (RC2−100)° C. and equal to or less than (RC2+50)° C. RC1 is an austenite 50%-recrystallization temperature estimated from the chemical composition and RC2 is the lower limit for the austenite recrystallization temperature which is estimated from the chemical composition. If the above total rolling reduction achieved in the temperature range of equal to or more than (RC1−150)° C. and equal to or less than RC1° C. is less than 35%, the average grain size increases and, consequently, it becomes impossible to enhance fatigue resistance. Accordingly, the total rolling reduction achieved in the temperature range of equal to or more than (RC1−150)° C. and equal to or less than RC1° C. is 35% or more. The above total rolling reduction is preferably 45% or more and is more preferably 60% or more.
Hot rolling is performed such that the finish rolling delivery temperature is equal to or more than (RC2−100)° C. and equal to or less than (RC2+50)° C. If the finish rolling delivery temperature is less than (RC2−100)° C., ferrite is formed and, consequently, a tensile strength of 980 MPa or more cannot be achieved. Accordingly, the finish rolling delivery temperature is equal to or more than (RC2−100)° C., is preferably equal to or more than (RC2−90)° C., and is more preferably equal to or more than (RC2−70)° C. On the other hand, if the finish rolling delivery temperature is more than (RC2+50)° C., austenite grains become coarsened and the average grain size of upper bainite increases consequently. This reduces strength. Accordingly, the finish rolling delivery temperature is equal to or less than (RC2+50)° C., is preferably equal to or less than (RC2+40)° C., and is more preferably equal to or less than (RC2+30)° C. RC1 and RC2 are defined by Formulae (1) and (2) below.
The hot-rolled steel sheet is cooled (first cooling). In the cooling step, the time interval between the end of hot rolling and the start of cooling (cooling start time) is limited to 2.0 s or less after the end of finish rolling. If the above cooling start time is less than 2.0 s, austenite grains grow disadvantageously and a tensile strength of 980 MPa or more cannot be achieved consequently. Accordingly, the cooling start time is 2.0 s or less, is preferably 1.5 s or less, and is more preferably 1.0 s or less.
In the cooling step, if the average cooling rate at which the temperature is reduced from the finish rolling delivery temperature to the cooling stop temperature is excessively low, ferrite transformation may disadvantageously occur prior to upper bainite transformation and the intended area fraction of upper bainite phase cannot be formed consequently. Accordingly, the average cooling rate is 20° C./s or more, is preferably 30° C./s or more, and is more preferably 50° C./s or more. Although the upper limit is not set, if the average cooling rate is excessively high, it becomes difficult to control the cooling stop temperature and, consequently, it may become difficult to form the intended microstructure. Therefore, the average cooling rate is preferably 500° C./s or less, is more preferably 300° C./s or less, and is further preferably 150° C./s or less. In the cooling step, forced cooling may be performed such that the above average cooling rate can be achieved. The cooling method is not limited. It is preferable to perform water cooling or the like.
The cooling stop temperature is set to a temperature equal to or more than Trs°C and equal to or less than (Trs+180)° C. If the cooling stop temperature is less than Trs°C, the microstructure includes lower bainite. Although lower bainite is a microstructure phase having a high strength, it may have low fatigue resistance after subjected to a heat treatment. Accordingly, the cooling stop temperature is set to a temperature equal to or more than Trs°C. On the other hand, if the cooling stop temperature is more than (Trs+180)° C., ferrite may be generated disadvantageously. This makes it impossible to achieve a tensile strength of 980 MPa or more. Accordingly, the cooling stop temperature is set to a temperature equal to or less than (Trs+180)° C. Trs is defined using Formula (3) below.
The cooled hot-rolled steel sheet is coiled at a coiling temperature equal to or more than Trs°C and equal to or less than (Trs+180)° C. If the coiling temperature is less than Trs°C, lower bainite transformation may occur subsequent to coiling and, consequently, intended martensite and/or retained austenite cannot be formed. Accordingly, the coiling temperature is limited to a temperature equal to or more than Trs°C, is preferably equal to or more than (Trs+10)° C., and is more preferably equal to or more than (Trs+30)° C. On the other hand, if the coiling temperature is more than (Trs+180)° C., ferrite may be generated disadvantageously. This makes it impossible to achieve a tensile strength of 980 MPa or more. Accordingly, the coiling temperature is limited to a temperature equal to or less than (Trs+180)° C., is preferably equal to or less than (Trs+150)° C., and is more preferably equal to or less than (Trs+120)° C.
Subsequently, the temperature is reduced to a temperature equal to or less than (Trs−250)° C. at an average cooling rate of 1° C./s or less (second cooling). If the average cooling rate at which the temperature is reduced from the coiling temperature to a temperature equal to or less than (Trs−250)° C. is more than 1° C./s, bainite transformation does not occur to a sufficient degree and the amount of martensite and retained austenite increases consequently. This makes it impossible to form the microstructure intended in accordance with aspects of the present invention. Accordingly, the average cooling rate at which the temperature is reduced from the coiling temperature to a temperature equal to or less than (Trs−250)° C. is limited to 1° C./s or less, is preferably 0.8° C./s or less, and is more preferably 0.5° C./s or less. Although cooling may be performed to any temperature equal to or less than (Trs−250)° C., it is preferable to reduce the temperature to about 10° C. to 30° C. The steel sheet may be cooled in any form. For example, the steel sheet may be cooled after it has been wound into a coil.
The cooled steel sheet is temper-rolled at a rolling reduction of 0.1% or more and 5.0% or less. If the rolling reduction is less than 0.1%, the dislocation density becomes insufficient and a markedly high fatigue strength cannot be achieved. Accordingly, the rolling reduction is limited to 0.1% or more, is preferably 0.3% or more, and is more preferably 0.5% or more. However, if temper rolling is performed at a rolling reduction of more than 5.0%, the amount of load applied to the rolls increases. This disadvantageously increases the number of times the rollers need to be replaced and the manufacturing costs. Accordingly, the rolling reduction is limited to 5.0% or less, is preferably 4.0% or less, and is more preferably 3.0% or less.
The high-strength steel sheet according to aspects of the present invention can be produced by the above-described steps. Optionally, for example, pickling may be performed after temper rolling in accordance with the conventional method in order to remove scales formed on the surface of the steel sheet.
Molten steels having the compositions described in Table 1 were prepared using a converter and formed into steel slabs by continuous casting, which were used as steel materials.
a
0.022
b
0.178
c
3.21
d
0.0001
e
0.68
f
3.21
g
0.123
h
Cr: 2.70
i
0.34
j
0.00
0.000
k
0.0140
a
b
c
d
e
f
g
h
i
j
k
The steel materials were heated to the respective heating temperatures described in Table 2. The heated steel materials were subjected to a hot-rolling process consisting of rough rolling and finish rolling to form hot-rolled steel sheets. Table 2 lists the finish rolling delivery temperatures in the hot rolling process.
The hot-rolled steel sheets were each cooled (first cooling) with the average cooling rate and the cooling stop temperature listed in Table 2. The cooled hot-rolled steel sheets were coiled at the respective coiling temperatures listed in Table 2. The coiled steel sheets were cooled (second cooling) with the respective average cooling rates listed in Table 2 to form high-strength steel sheets. Subsequent to cooling, temper rolling was performed at the rolling reductions listed in Table 2. Then, pickling was performed. Pickling was performed at a temperature of 85° C. using a 10-mass % aqueous solution of hydrochloric acid. Subsequently, the steel sheets were subjected to a heat treatment (170° C. and 20 minutes) that corresponded to baking coating. Hereby, high-strength hot-rolled steel sheets were prepared.
25
690
2.2
890
310
650
700
950
280
670
30
380
600
2
0.0
370
630
3
420
A test specimen was taken from each of the high-strength steel sheets, and the microstructure and mechanical properties of the specimen were determined by the following steps.
A test specimen for microstructure observation was taken from each of the high-strength steel sheets such that a cross section of the steel sheet which was taken in the thickness direction so as to be parallel to the rolling direction was exposed as an observation plane. The surface of the test specimen was ground and then corroded with an etchant (3% nital solution) in order to cause the microstructure to appear. Subsequently, an image of the surface layer that extended from the surface to the position 1/10 of the thickness of the steel sheet was taken using a scanning electron microscope (SEM) at a 5000-fold magnification in 10 fields of view in order to obtain SEM images of microstructure. The SEM images were analyzed by image processing in order to determine the area fractions of upper bainite (UB), polygonal ferrite (F), and lower bainite (LB). Since it is difficult to distinguish martensite (M) and retained austenite (γ) from each other with a SEM, the area fractions and average grain sizes thereof were determined by making identifications using an electron back scatter diffraction patterns (EBSD) method. Table 3 lists the area fractions and average grain sizes of the above microstructure phases. Table 3 also lists the total area fractions (M+γ) of martensite and retained austenite.
A JIS No. 5 test piece for tensile test (JIS Z 2201) was taken from each of the hot-rolled steel sheets such that the tensile direction of the test piece was perpendicular to the rolling direction. A tensile test was conducted in conformity with JIS Z 2241 at a strain rate of 10-3/s in order to determine tensile strength. In accordance with aspects of the present invention, an evaluation of “Passed” was given when the tensile strength was 980 MPa or more. Table 3 lists the results.
A test specimen having the dimensions and shape illustrated in the FIGURE was taken from each of the hot-rolled steel sheets such that the longitudinal direction of the test specimen was perpendicular to the rolling direction. A plane bending fatigue test was conducted in conformity with JIS Z 2275. The stress loading mode was such that the stress ratio R was −1 and the frequency f was 25 Hz. The amplitude of loading stress was changed in six stages, and the number of stress cycles applied until rupture occurred was measured. An S-N curve was determined, and a fatigue strength (fatigue limit) at 2×106 cycles was calculated. In accordance with aspects of the present invention, an evaluation of “excellent fatigue resistance” was given when a fatigue limit ratio calculated by dividing the fatigue limit by the tensile strength determined in the tensile test was 0.50 or more. Table 3 lists the results.
8.3
0.35
18
7.5
0.38
69
10
92
0.34
100
53
47
0.41
0.47
65
27
0.43
9.0
0.33
15
85
96
0.36
100
7.6
0.38
69
28
3
31
0.38
a
b
0.42
c
0.39
d
73
24
0.48
e
100
0.36
f
0.47
g
48
52
0.33
h
70
27
0.47
i
j
0.30
k
66
34
0.46
Invention Examples were all high-strength steel sheets having a tensile strength of 980 MPa or more and excellent fatigue resistance. In contrast, Comparative Examples, which were outside the scope of the present invention, did not have a tensile strength of 980 MPa or more or excellent fatigue resistance.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2022-055658 | Mar 2022 | JP | national |
This is the U.S. National Phase application of PCT/JP2023/011909, filed Mar. 24, 2023 which claims priority to Japanese Patent Application No. 2022-055658, filed Mar. 30, 2022, the disclosures of these applications being incorporated herein by reference in their entireties for all purposes.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2023/011909 | 3/24/2023 | WO |
