Patent Application
20240384379
The present invention relates to a steel sheet that is suitably press formed into complicated shapes through a press-forming process for use in, for example, automobiles and home appliances and has excellent chemical convertibility; to a member obtained using the steel sheet; and to methods for manufacturing them.
Against the backdrop of heightened regulations on CO2 emissions worldwide, there is a growing demand for weight reduction of automobile bodies by increasing the strength of steel sheets for automobiles. Thus, 980 MPa or higher grade high strength steel sheets are increasingly applied to bodies and seat parts. In general, increasing the strength of steel sheets is accompanied by a decrease in press formability, such as ductility and stretch flangeability, thereby increasing the occurrence frequency of cracking at the time of press forming and lowering the degree of freedom in shape. Thus, the application is limited to parts having simple shapes. In order to apply high strength steel sheets to parts having complicated shapes, it is important to increase the strength of steel sheets while maintaining or enhancing formability.
Against the above background, techniques for enhancing the ductility of steel sheets have led to the development of TRIP steel in which retained austenite (retained γ) is dispersed in the microstructure of the steel sheets.
For example, Patent Literature 1 discloses a manufacturing method involving austempering treatment (a treatment in which the steel is cooled from a single-phase annealing temperature or a two-phase annealing temperature to a bainite transformation temperature and is isothermally held to form retained γ while utilizing bainite transformation during the isothermal holding or the cooling). Specifically, a steel sheet including C: 0.10 to 0.45%, Si: 0.5 to 1.8%, and Mn: 0.5 to 3.0% is annealed and is thereafter subjected to aging treatment at a temperature in the range of 350 to 500° C. for 1 to 30 minutes to form retained γ. According to the disclosure, a high ductility steel sheet having TS: 80 kgf/mm2 or more and TS×EL: 2500 kgf/mm2-% or more can be obtained.
Patent Literature 2 discloses that a steel sheet containing C: 0.10 to 0.25%, Si: 1.0 to 2.0%, and Mn: 1.5 to 3.0% is annealed, cooled to 450 to 300° C. at a rate of 10° C./s or more, and held for 180 to 600 seconds. In this manner, the microstructure is controlled so that the volume fraction of retained γ will be 5% or more and the area fractions of bainitic ferrite and polygonal ferrite will be 60% or more and 20% or less, respectively. According to the disclosure, a steel sheet excellent in both ductility: EL and stretch flangeability: A can be obtained.
Patent Literature 3 describes a manufacturing method involving Q&P treatment (a treatment in which the steel is cooled from a single-phase annealing temperature or a two-phase annealing temperature to a temperature ranging from a martensite start temperature: Ms to a martensite finish temperature: Mf, thereby forming a martensite microstructure, and the steel is then reheated to partition carbon from the martensite microstructure to non-transformed γ, thereby forming retained γ). Specifically, a steel sheet having a specific chemical composition is annealed, cooled to a range of temperatures of 150 to 350° C., and subsequently reheated to and held at 350 to 600° C. to form a microstructure including ferrite, tempered martensite, and retained γ. According to the disclosure, the steel sheet thus obtained has excellent ductility and stretch flangeability.
The austempering treatment described in Patent Literatures 1 and 2, and the Q&P treatment described in Patent Literature 3 are both heat treatments for producing TRIP steel sheets. The Q&P treatment is suitable for the manufacturing of steel sheets having higher strength because the treatment forms tempered martensite that contributes to increasing the strength.
Patent Literature 4 discloses a manufacturing method that improves the Q&P treatment described above. Specifically, steel in the course of post-annealing cooling is held at a temperature of 470 to 405° C. for 14 to 200 seconds to concentrate carbon into non-transformed γ while utilizing upper bainite transformation, and the steel is thereafter cooled to a temperature of Ms—90 to Ms—180 (° C.) to induce martensite transformation and is reheated to partition carbon from the martensite microstructure to non-transformed γ, thereby forming retained γ. The steel sheet thus obtained has both high ductility and excellent stretch flangeability.
In the above techniques, the steel sheet contains a large amount of silicon in order to promote efficient concentration of carbon into non-transformed γ. On the other hand, steel sheets used for press-formed members are incorporated into structures, such as automobiles, after being painted. Chemical conversion is performed on the steel sheets in order to impart good paintability. In chemical conversion, the presence of an oxide film as a superficial portion of the steel sheet destroys the uniformity of crystal grains that are attached by the chemical conversion and serves as a factor that deteriorates paintability. To address this, a continuous annealing furnace used in the manufacturing of steel sheets usually performs pickling treatment as a pretreatment to improve the chemical convertibility. Depending on the components in the steel sheets, particularly when the steel sheets have a high Si content, the chemical convertibility is significantly deteriorated by a Si-containing surface oxide layer that remains after pickling.
In order to solve the above problem, for example, Patent Literature 5 discloses that excellent chemical convertibility can be imparted even to a high-Si content steel sheet by steps in which the steel sheet is pickled by being continuously immersed in a mixed acid solution including an oxidative first acid and a non-oxidative second acid, and is repickled by being continuously immersed in an acid solution including a non-oxidative third acid.
While the conventional TRIP steel described in Patent Literature 1 has excellent El, its stretch flange formability is very low.
In the technique described in Patent Literature 2, the microstructure is mainly bainitic ferrite and includes a small amount of ferrite. Because of this composition, the steel sheet is excellent in stretch flange formability but is not necessarily high in ductility.
The technique described in Patent Literature 3 realizes relatively high ductility and high stretch flange formability compared to the conventional TRIP steels and steels making use of bainitic ferrite. However, difficulties are encountered in the formation of hard-to-form parts, such as center pillars, and further enhancements in ductility are required. In view of the application to hard-to-form parts, further improvements are demanded in ductility, in particular, uniform elongation and local elongation at the same time. The uniform elongation is a ductility indicator El that indicates the amount of elongation to the onset of necking and is written as U. El. The local elongation is the amount of elongation obtained by subtracting the uniform elongation from the total elongation: T. El, and is written as L. El. The L. El needs to be increased while maintaining the U. El.
The technique described in Patent Literature 4 can provide a steel sheet having high ductility and excellent stretch flangeability by holding the steel sheet in the course of post-annealing cooling in such a manner that upper bainite transformation is utilized, and by subsequently performing the Q&P treatment and reheating followed by bainite transformation. However, the improvement in local elongation is still insufficient to satisfy both bend formability and bulging formability that are required simultaneously in the formation of hard-to-form parts. Because the steel sheet contains a large amount of silicon to promote the partitioning of carbon from martensite to non-transformed γ in the Q&P treatment, a special pickling technique, such as one described in Patent Literature 5, is required.
The pickling described in Patent Literature 5 is a technique that imparts excellent chemical convertibility to steel sheets. However, the technique involves high running costs and causes a cost problem when applied to pickling of various kinds of steel sheets in a single continuous annealing furnace. Thus, the establishment of other techniques has been desired.
As discussed above, the conventional techniques are still insufficient and are incapable of imparting excellent chemical convertibility to steel sheets while ensuring high ductility and excellent stretch flange formability at the same time.
Aspects of the present invention have been made to solve the problems discussed above. It is therefore an object according to aspects of the present invention to provide a steel sheet that has 980 MPa or higher tensile strength and achieves high ductility, excellent stretch flange formability, and excellent chemical convertibility; a related member; and methods for manufacturing them.
Here, 980 MPa or higher tensile strength means that a JIS No. 5 test piece for tensile test has a tensile strength of 980 MPa or more when tested by a tensile test in accordance with the provisions of JIS Z2241 (2011) in the tensile direction perpendicular to the rolling direction at a crosshead speed of 10 mm/min.
Furthermore, high ductility means that a JIS No. 5 test piece for tensile test satisfies tensile strength (TS)×total elongation (T. El)≥18000 MPa-% or more when tested by a tensile test in accordance with the provisions of JIS Z2241 (2011) in the tensile direction perpendicular to the rolling direction at a crosshead speed of 10 mm/min.
Furthermore, excellent stretch flange formability means that the hole expansion ratio λ is 45% or more when tested by a hole expansion test in accordance with JFST (The Japan Iron and Steel Federation Standard) 1001.
Furthermore, good chemical convertibility means that a steel sheet is covered with a chemical conversion coating microstructure on all the faces when the steel sheet is subjected to sulfuric acid electrolytic pickling for 2 seconds at a current density of 20 to 35 A/dm2, degreasing (treatment temperature: 40° C., treatment time: 120 seconds, spray degreasing), surface conditioning (pH: 9.5, treatment temperature: room temperature, treatment time: 20 seconds), and chemical conversion using a zinc phosphate chemical conversion solution (temperature of the chemical conversion solution: 35° C., treatment time: 120 seconds).
The present inventors carried out extensive studies on approaches to imparting high ductility and excellent stretch flange formability and have obtained the following conclusions.
In the austempering treatment, bainite transformation at around 400° C. causes partitioning of carbon into non-transformed austenite until T0 composition is reached in which the free energies of fcc phases and bcc phases are equal to each other, and thereafter the bainite transformation stops. The coarse and thermally unstable non-transformed austenite forms a hard martensite microstructure or mechanically unstable retained γ at the time of final cooling, with the result that stretch flangeability is deteriorated. Thus, it is difficult to satisfy ductility and stretch flangeability at the same time by the austempering treatment.
On the other hand, in the Q&P treatment (for example, Q&P treatment in which the steel is held during post-annealing cooling), thermally unstable non-transformed γ undergoes martensite transformation at a cooling stop temperature of Ms to Mf and the martensite is tempered during subsequent reheating. Thus, the difference in hardness between hard phases and soft phases is reduced, and the steel exhibits excellent stretch flangeability and also attains enhanced ductility at the same time. This shows an advantage of the Q&P treatment (for example, Q&P treatment in which the steel is held during post-annealing cooling).
However, the Q&P treatment (for example, Q&P treatment in which the steel is held during post-annealing cooling) too requires that the steel contain large amounts of carbon and silicon in order to ensure that retained γ will be formed. That is, high-cost pickling treatment is necessary in order to impart chemical convertibility. An alloy design that involves, in particular, a reduced amount of silicon is necessary. However, reducing the amount of silicon lowers ductility. Thus, a process is desired that can offer high ductility even when the steel sheet has a small amount of silicon. Incidentally, the chemical convertibility described here is defined as a characteristic with which the steel sheet after a general pickling process can be treated to attain a coating weight and uniformity offering satisfactory paintability. The general pickling process is, for example, sulfuric acid pickling, but the pickling process is not limited.
As a result of extensive studies on the heating process before annealing, the present inventors have found that a soft ferrite microstructure is formed and acicular γ is formed adjacent thereto by heating a steel containing specific components at a specific heating rate. The present inventors have found that this acicular γ contributes to the partitioning of carbon and the formation of retained γ in microstructure formation during the cooling process, and allows even a low-Si steel sheet to exhibit excellent ductility. The findings are based on the outlines below. Here, the term low-Si means, although not particularly limited to, that the Si content is less than 1.60 mass %.
As described above, excellent uniform elongation and local elongation can be obtained at the same time by making use of acicular austenite formed in the heating process and by utilizing upper bainite transformation before martensite transformation. As a result, high ductility and excellent stretch flange formability can be imparted to even a steel sheet having a reduced amount of silicon. Furthermore, the chemical convertibility of the steel sheet can also be improved.
Aspects of the present invention have been made based on the above knowledge. Specifically, aspects of the present invention include the following:
Si/Mn<0.50 Formula (1)
According to aspects of the present invention, a steel sheet is provided that has 980 MPa or higher tensile strength and achieves high ductility, excellent stretch flange formability, and good chemical convertibility. A related member, and methods for manufacturing them are also provided. The steel sheet according to aspects of the present invention is suitably used in press forming of complicated shapes produced in the press forming process for use in, for example, automobiles and home appliances.
Hereinafter, embodiments of the present invention will be described in detail. The present invention is not limited to the embodiments discussed below.
A steel sheet according to aspects of the present invention has a chemical composition including, in mass %, C: 0.10 to 0.24%, Si: 0.4% or more and less than 1.60%, Mn: 2.0 to 3.6%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, and N: less than 0.015%, the chemical composition satisfying formula (1) below, the balance being Fe and incidental impurities. The steel sheet includes a microstructure in which the area fraction of polygonal ferrite is 5% or more and 25% or less, the area fraction of upper bainite is 5% or more and 50% or less, the volume fraction of retained austenite is 3% or more and 20% or less, the area fraction of fresh martensite is 12% or less (including 0%), the total of the area fractions of tempered martensite and lower bainite is 10% or more and 50% or less, and the area fraction of a remaining microstructure is 5% or less. The microstructure is such that the ratio of the total number of fresh martensite grains and retained austenite grains having an equivalent circular diameter of less than 0.8 μm is 50% or more relative to the number of all fresh martensite grains and all retained austenite grains, and the ratio of fresh martensite grains and retained austenite grains having an aspect ratio of 2.0 or more and an equivalent circular diameter of 0.8 μm or more is 30% or more relative to the number of fresh martensite grains and retained austenite grains having an equivalent circular diameter of 0.8 μm or more.
Si/Mn<0.50 Formula (1)
In formula (1), Si and Mn indicate the Si content (mass %) and the Mn content (mass %), respectively.
The steel sheet according to aspects of the present invention will be described below in the order of its chemical composition and its steel microstructure.
The steel sheet according to aspects of the present invention includes the components described below. In the following description, the unit “%” for the contents of components means “mass %”.
Carbon is added in order to control the hardenability of the steel sheet, the strength of martensite, and the volume fraction of retained γ to desired ranges. If the C content is less than 0.10%, the strength of the steel sheet and the ductility of the steel sheet cannot be sufficiently ensured.
Thus, the C content is limited to 0.10% or more. The C content is preferably 0.12% or more, more preferably 0.14% or more, and still more preferably 0.16% or more. If the C content exceeds 0.24%, the toughness of welds is deteriorated. Thus, the C content is limited to 0.24% or less. In order to enhance ductility and the toughness of spot welds, the C content is preferably 0.22% or less. In order to further improve the toughness of spot welds, the C content is more preferably 0.20% or less.
Si: 0.4% or More and Less than 1.60%
Silicon is added in order to effectively enhance ferrite strength, to suppress the formation of carbides in martensite and bainite, and to stabilize retained γ and thereby enhance ductility. From these points of view, the Si content is limited+ to 0.4% or more. In order to enhance ductility, the Si content is preferably 0.5% or more. The Si content is more preferably 0.6% or more. If the Si content is 1.60% or more, chemical convertibility is significantly deteriorated. Thus, the Si content is limited to less than 1.60%. The Si content is preferably 1.30% or less, and more preferably 1.20% or less. The Si content is still more preferably less than 1.0%.
Manganese ensures predetermined hardenability, suppresses ferrite transformation, and ensures the desired area fraction of tempered martensite and/or bainite to ensure strength. Furthermore, manganese is concentrated into γ during ferrite/γ two-phase annealing and lowers the Ms temperature of retained γ to stabilize the retained γ and thereby to improve ductility. Furthermore, similarly to silicon, manganese suppresses the formation of carbides in bainite and enhances ductility. Furthermore, manganese increases the volume fraction of retained γ to enhance ductility. From these points of view, manganese is an important element in accordance with aspects of the present invention. In order to obtain these effects, the Mn content is limited to 2.0% or more. In order to enhance hardenability, the Mn content is preferably 2.1% or more. The Mn content is more preferably 2.2% or more. If, on the other hand, the Mn content exceeds 3.6%, bainite transformation is significantly retarded to make it difficult to ensure high ductility. Furthermore, more than 3.6% manganese makes it difficult to suppress the formation of massive coarse γ and massive coarse martensite, and deteriorates stretch flange formability. Thus, the Mn content is limited to 3.6% or less. In order to ensure high ductility by promoting bainite transformation, the Mn content is preferably 2.8% or less.
Si/Mn<0.50 Formula (1)
Silicon oxide is a surface oxide on the steel sheet that significantly deteriorates chemical convertibility. Thus, Si/Mn is limited to less than 0.50 in order to form Mn-containing oxide that is readily soluble in an acid solution. That is, in accordance with aspects of the present invention, formula (1) is limited to Si/Mn<0.50. In formula (1), Si and Mn indicate the Si content (mass %) and the Mn content (mass %), respectively. When the formula is satisfied, chemical convertibility can be imparted in the dew point range of −50° C. or above and −30° C. or below. Si/Mn is preferably 0.40 or less, and more preferably 0.35 or less.
Phosphorus is an element that strengthens steel, but much phosphorus deteriorates spot weldability. Thus, the P content is limited to 0.02% or less. In order to improve spot weldability, the P content is preferably 0.01% or less. The P content may be nil. From the point of view of manufacturing cost, the P content is preferably 0.001% or more.
Sulfur is an element that is effective in improving scale exfoliation in hot rolling and effective in suppressing nitridation during annealing, but sulfur deteriorates spot weldability and local elongation. To eliminate or reduce the deterioration, the S content is limited to 0.01% or less. In accordance with aspects of the present invention, the contents of C, Si, and Mn are high and spot weldability tends to be deteriorated. In order to improve spot weldability, the S content is preferably 0.0020% or less, and more preferably less than 0.0010%. The S content may be nil. From the point of view of manufacturing cost, the S content is preferably 0.0001% or more.
Sol. Al: Less than 1.0%
Aluminum is added for the purpose of deoxidization or for the purpose of stabilizing retained γ as a substitute for silicon. The lower limit of the sol. Al content is not particularly limited. For stable deoxidization, the sol. Al content is preferably 0.01% or more. On the other hand, 1.0% or more sol. Al significantly lowers the strength of the base material and also deteriorates chemical convertibility. Thus, the sol. Al content is limited to less than 1.0%. In order to obtain high strength, the sol. Al content is preferably less than 0.20%, and more preferably 0.10% or less.
N: Less than 0.015%
Nitrogen is an element that forms nitrides, such as BN, AlN, and TiN, in steel. This element lowers the hot ductility of steel and lowers the surface quality. Furthermore, in B-containing steel, nitrogen has a harmful effect in eliminating the effect of boron through the formation of BN. The surface quality is significantly deteriorated if the N content is 0.015% or more. Thus, the N content is limited to less than 0.015%. The N content may be nil. From the point of view of manufacturing cost, the N content is preferably 0.0001% or more.
The balance after the above components is Fe and incidental impurities. The steel sheet according to aspects of the present invention preferably has a chemical composition that contains the basic components described above, with the balance consisting of Fe and incidental impurities.
In addition to the above components, the chemical composition of the steel sheet according to aspects of the present invention may appropriately include, in place of part of Fe and incidental impurities, one or two of optional elements selected from the following (A) and (B):
Niobium is preferably added in order to reduce the size of the microstructure and enhance the defect resisting characteristics of spot welds. Furthermore, niobium may be added to produce an effect of reducing the size of the steel microstructure and increasing the strength, an effect of promoting bainite transformation through the grain size reduction, an effect of improving bendability, and an effect of enhancing delayed fracture resistance. In order to obtain these effects, the Nb content is preferably 0.002% or more, but the lower limit is not particularly limited. The Nb content is more preferably 0.004% or more, and still more preferably 0.010% or more. However, adding much niobium results in excessive precipitation strengthening and low ductility. Furthermore, the rolling load is increased and castability is deteriorated. Thus, when niobium is added, the Nb content is limited to 0.2% or less. The Nb content is preferably 0.1% or less, more preferably 0.05% or less, and still more preferably 0.03% or less.
Titanium is preferably added in order to reduce the size of the microstructure and enhance the defect resisting characteristics of spot welds. Furthermore, titanium fixes nitrogen in steel as TiN to produce an effect of enhancing hot ductility and an effect of allowing boron to produce its effect of enhancing hardenability. In order to obtain these effects, the Ti content is preferably 0.002% or more, but the lower limit is not particularly limited. In order to fix nitrogen sufficiently, the Ti content is more preferably 0.008% or more. The Ti content is still more preferably 0.010% or more. On the other hand, more than 0.2% titanium causes an increase in rolling load and a decrease in ductility by an increased amount of precipitation strengthening. Thus, when titanium is added, the Ti content is limited to 0.2% or less. The Ti content is preferably 0.1% or less, and more preferably 0.05% or less. In order to ensure high ductility, the Ti content is still more preferably 0.03% or less.
Vanadium may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in martensite and upper/lower bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide. In order to obtain these effects, the V content is preferably 0.003% or more, but the lower limit is not particularly limited. The V content is more preferably 0.005% or more, and still more preferably 0.010% or more. On the other hand, much vanadium significantly deteriorates castability. Thus, when vanadium is added, the V content is limited to 0.2% or less. The V content is preferably 0.1% or less. The V content is more preferably 0.05% or less.
Boron advantageously facilitates the formation of a predetermined area fraction of tempered martensite and/or bainite. Furthermore, residual solute boron enhances delayed fracture resistance. In order to obtain these effects of boron, the B content is preferably 0.0002% or more. The B content is more preferably 0.0005% or more. The B content is still more preferably 0.0010% or more. If, on the other hand, the B content exceeds 0.01%, the effects are saturated, and further hot ductility is significantly lowered to invite surface defects. Thus, when boron is added, the B content is limited to 0.01% or less. The B content is preferably 0.0050% or less. The B content is more preferably 0.0030% or less.
Copper enhances the corrosion resistance in automobile use environments. Furthermore, corrosion products of copper cover the surface of the steel sheet and effectively suppress penetration of hydrogen into the steel sheet. Copper is an element that is mixed when scraps are used as raw materials. By accepting copper contamination, recycled materials can be used as raw materials and thereby manufacturing costs can be reduced. From these points of view and further from the point of view of enhancing delayed fracture resistance, the Cu content is preferably 0.05% or more, but the lower limit is not particularly limited. The Cu content is more preferably 0.10% or more. On the other hand, too much copper invites surface defects. Thus, when copper is added, the Cu content is limited to 0.2% or less.
Similarly to copper, nickel is an element that acts to enhance corrosion resistance. Furthermore, nickel also acts to eliminate or reduce the occurrence of surface defects that tend to occur when the steel contains copper. In order to obtain these effects, the Ni content is preferably 0.01% or more, but the lower limit is not particularly limited. The Ni content is more preferably 0.04% or more, and still more preferably 0.06% or more. On the other hand, adding too much nickel can instead cause surface defects because scales are formed nonuniformly in a heating furnace, and also increases the cost. Thus, when nickel is added, the Ni content is limited to 0.2% or less.
Chromium may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. In order to obtain these effects, the Cr content is preferably 0.01% or more, but the lower limit is not particularly limited. The Cr content is more preferably 0.03% or more, and still more preferably 0.06% or more. On the other hand, too much chromium deteriorates pitting corrosion resistance. Thus, when chromium is added, the Cr content is limited to 0.4% or less.
Molybdenum may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. In order to obtain these effects, the Mo content is preferably 0.01% or more. The Mo content is more preferably 0.03% or more, and still more preferably 0.06% or more. On the other hand, molybdenum significantly deteriorates the chemical convertibility of the cold rolled steel sheet. Thus, when molybdenum is added, the Mo content is limited to 0.15% or less.
Magnesium fixes oxygen as MgO and contributes to improvement in delayed fracture resistance. Thus, the Mg content is preferably 0.0002% or more. The Mg content is more preferably 0.0004% or more, and still more preferably 0.0006% or more. On the other hand, much magnesium deteriorates surface quality and bendability. Thus, when magnesium is added, the Mg content is limited to 0.0050% or less. The Mg content is preferably 0.0025% or less, and more preferably 0.0010% or less.
Calcium fixes sulfur as CaS and contributes to improvements in bendability and delayed fracture resistance. Thus, the Ca content is preferably 0.0002% or more. The Ca content is more preferably 0.0005% or more, and still more preferably 0.0010% or more. On the other hand, much calcium deteriorates surface quality and bendability. Thus, when calcium is added, the Ca content is limited to 0.0050% or less. The Ca content is preferably 0.0035% or less, and more preferably 0.0020% or less.
Tin suppresses oxidation and nitridation of a superficial portion of the steel sheet and thereby eliminates or reduces the loss of the C and B contents in the superficial portion. Furthermore, the elimination or reduction of the loss of the C and B contents leads to suppressed formation of ferrite in the superficial portion of the steel sheet, thus increasing strength and improving fatigue resistance. From these points of view, the Sn content is preferably 0.002% or more. The Sn content is more preferably 0.004% or more, and still more preferably 0.006% or more.
The Sn content is further preferably 0.008% or more.
If, on the other hand, the Sn content exceeds 0.10%, castability is deteriorated. Furthermore, tin is segregated at prior γ grain boundaries to deteriorate delayed fracture resistance. Thus, when tin is added, the Sn content is limited to 0.10% or less. The Sn content is preferably 0.04% or less, and more preferably 0.03% or less.
Antimony suppresses oxidation and nitridation of a superficial portion of the steel sheet and thereby eliminates or reduces the loss of the C and B contents in the superficial portion. Furthermore, the elimination or reduction of the loss of the C and B contents leads to suppressed formation of ferrite in the superficial portion of the steel sheet, thus increasing strength and improving fatigue resistance. From these points of view, the Sb content is preferably 0.002% or more. The Sb content is more preferably 0.004% or more, and still more preferably 0.006% or more. If, on the other hand, the Sb content exceeds 0.10%, castability is deteriorated and segregation occurs at prior γ grain boundaries to deteriorate delayed fracture resistance. Thus, when antimony is added, the Sb content is limited to 0.10% or less. The Sb content is preferably 0.04% or less, and more preferably 0.03% or less.
Rare earth metals are elements that spheroidize the shape of sulfides and thereby eliminate or reduce adverse effects of sulfides on stretch flange formability, thus improving stretch flange formability. In order to obtain this effect, the REM content is preferably 0.0005% or more. The REM content is more preferably 0.0010% or more, and still more preferably 0.0020% or more.
If, on the other hand, the REM content exceeds 0.0050%, the effect of improving stretch flange formability is saturated. Thus, when rare earth metals are added, the REM content is limited to 0.0050% or less.
In accordance with aspects of the present invention, the rare earth metals indicate scandium (Sc) with atomic number 21, yttrium (Y) with atomic number 39, and lanthanide elements from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71. The REM concentration in accordance with aspects of the present invention is the total content of one, or two or more elements selected from the above rare earth metals.
When the content of any of the above optional components is below the lower limit, the optional element present below the lower limit does not impair the advantageous effects according to aspects of the present invention. Thus, such an optional element below the lower limit content is regarded as an incidental impurity.
Next, the steel microstructure of the steel sheet according to aspects of the present invention will be described.
In order to ensure high ductility, the area fraction of polygonal ferrite is limited to 5% or more. The polygonal ferrite is preferably 8% or more, and more preferably 11% or more. On the other hand, the area fraction of polygonal ferrite is limited to 25% or less in order to obtain predetermined strength. The polygonal ferrite is more preferably 23% or less.
Upper bainite is bainite involving a low level of carbide precipitation. Upper bainite partitions carbon into surrounding non-transformed γ and thus can be used to form retained γ with high working stability. In addition, upper bainite has hardness intermediate between those of ferrite and martensite, and the presence of such a microstructure having intermediate hardness enhances local elongation. Furthermore, upper bainite transformation from acicular γ formed by annealing promotes the formation of retained γ having a high aspect ratio. Thus, 5% or more upper bainite is required at a strength level where the tensile strength (TS) is 980 MPa or more. The area fraction of upper bainite is therefore limited to 5% or more. The area fraction of upper bainite is preferably 6.0% or more, and more preferably 7.0% or more. On the other hand, excessive formation of upper bainite lowers strength. Thus, the area fraction is limited to 50% or less. The area fraction of upper bainite is preferably 45% or less, and more preferably 40% or less.
In order to ensure high ductility, the volume fraction of retained γ is limited to 3% or more relative to the whole of the steel microstructure. The volume fraction of retained γ (the amount of retained γ) is preferably 3.0% or more, more preferably 5% or more, and still more preferably 7% or more. This amount of retained γ includes the amounts of retained γ generated adjacent to upper bainite and retained γ generated adjacent to martensite and lower bainite. If the amount of retained γ is excessively large, strength is lowered and stretch flange formability is significantly lowered. Thus, the volume fraction of retained γ is limited to 20% or less. The volume fraction of retained γ is preferably 15% or less, and more preferably 13% or less. Incidentally, the “volume fraction” may be regarded as the “area fraction”.
Fresh martensite is a microstructure that lowers local elongation, but can offer enhanced strength when formed within a range not detrimental to bendability and flangeability. From this point of view, the area fraction of fresh martensite is limited to range from 0% up to 12%. The area fraction of fresh martensite may be 12.0% or less.
In accordance with aspects of the present invention, lower bainite is formed when the steel sheet is over-aged and held at 500° C. or below and 350° C. or above. Furthermore, tempered martensite is formed when the martensite microstructure formed by cooling to a second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above is later tempered by over-aging and holding in a range of temperatures of 350 to 550° C. for 20 to 3000 seconds.
While upper bainite involves a low level of carbide precipitation, tempered martensite and lower bainite have carbides precipitated in their microstructures. Thus, the amount of carbon partitioned to non-transformed γ is reduced. Tempered martensite and lower bainite, however, broaden the To composition at low temperatures to bring about enrichment of carbon to non-transformed γ or further reduce the amount of fresh martensite occurring during the final cooling. It is therefore necessary to control these microstructures to obtain retained γ having high working stability.
If the total of the area fractions of tempered martensite and lower bainite exceeds 50%, the precipitation of carbides is promoted and the required amount of retained γ cannot be obtained to make it impossible to obtain desired ductility. Thus, in accordance with aspects of the present invention, the total of the area fractions of tempered martensite and lower bainite is limited to 50% or less.
The total of these area fractions is preferably 45% or less, and more preferably 40% or less.
If, on the other hand, the total of the area fractions of tempered martensite and lower bainite is less than 10%, strength becomes insufficient, and an increased amount of fresh martensite is formed during the final cooling to cause deterioration in flangeability. Thus, in accordance with aspects of the present invention, the total of the area fractions of tempered martensite and lower bainite is limited to 10% or more. The total of these area fractions is preferably 13% or more, and more preferably 16% or more.
The remaining microstructure is a microstructure other than polygonal ferrite, upper bainite, retained austenite, fresh martensite, tempered martensite, and lower bainite, and includes, for example, pearlite. A pearlite microstructure inhibits efficient carbon partitioning and suppresses the formation of retained γ, thus causing a decrease in ductility. In accordance with aspects of the present invention, influences on the material quality can be ignored as long as the area fraction of the remaining microstructure is 5% or less. Thus, the upper limit of the area fraction of the remaining microstructure is limited to 5%. The area fraction of the remaining microstructure may be 0%.
Ratio of the Total Number of Fresh Martensite Grains and Retained γ Grains Having an Equivalent Circular Diameter of Less than 0.8 μm Relative to the Number of all Fresh Martensite Grains and all Retained Austenite Grains: 50% or More
Fresh martensite grains and retained austenite grains having an equivalent circular diameter of less than 0.8 μm are unlikely to serve as stress concentration sites during local deformation and do not contribute to void formation. Thus, they are microstructures that do not deteriorate local ductility and flangeability.
Excellent local elongation and flangeability can be obtained in accordance with aspects of the present invention when the total number of fresh martensite grains and retained austenite grains having an equivalent circular diameter of less than 0.8 μm is 50% or more of the number of all fresh martensite grains and all retained austenite grains.
Thus, in accordance with aspects of the present invention, the ratio of the total number of fresh martensite grains and retained γ grains having an equivalent circular diameter of less than 0.8 μm is limited to 50% or more relative to the number of all fresh martensite grains and all retained austenite grains. That is, formula (A) below is satisfied.
100×(total number of fresh martensite grains and retained γ grains having an equivalent circular diameter of less than 0.8 μm)/(number of all fresh martensite grains and all retained γ grains)≥50(%) Formula (A)
The ratio of the left side specified by formula (A) is preferably 55% or more.
In order to obtain the above microstructure, one or both of tempered martensite and lower bainite may be formed in the microstructure. Tempered martensite may be obtained sufficiently by cooling the steel sheet to a second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above. Lower bainite may be obtained sufficiently by over-aging and holding the steel sheet in a range of temperatures of 350 to 550° C. for 20 to 3000 seconds.
Ratio of the Total Number of Fresh Martensite Grains and Retained γ Grains Having an Aspect Ratio of 2.0 or More and an Equivalent Circular Diameter of 0.8 μm or More Relative to the Number of Fresh Martensite Grains and Retained γ Grains Having an Equivalent Circular Diameter of 0.8 μm or More: 30% or More
Stress concentration during local deformation can be reduced and the formation of voids can be suppressed to attain enhancements in local ductility and flangeability when fresh martensite grains and/or retained austenite grains having an equivalent circular diameter of 0.8 μm or more have a high aspect ratio of the fresh martensite grains and/or the retained austenite grains.
The area fraction of such fresh martensite grains and/or retained austenite grains can be increased by ensuring that acicular austenite formed in the heating process and surrounded by a soft ferrite microstructure is transformed into bainite in the subsequent cooling process. In accordance with aspects of the present invention, desired formability can be obtained when the grains having an equivalent circular diameter of 0.8 μm or more and an aspect ratio of 2.0 or more represent 30% or more of the total number of fresh martensite grains and retained austenite grains having an equivalent circular diameter of 0.8 μm or more.
Thus, in accordance with aspects of the present invention, the ratio of the total number of fresh martensite grains and retained austenite grains having an aspect ratio of 2.0 or more and an equivalent circular diameter of 0.8 μm or more is limited to 30% or more relative to the number of fresh martensite grains and retained austenite grains having an equivalent circular diameter of 0.8 μm or more. That is, formula (B) below is further satisfied in addition to formula (A) described hereinabove. The ratio of the left side specified by formula (B) is preferably 35% or more.
100×(total number of fresh martensite grains and retained γ grains having an aspect ratio of 2.0 or more and an equivalent circular diameter of 0.8 μm or more)/(number of fresh martensite grains and retained γ grains having an equivalent circular diameter of 0.8 μm or more)≥30(%) Formula (B)
The microstructure of the steel sheet obtained is measured in the following manner.
The steel sheet is cut to give an observation specimen so that a cross section perpendicular to the steel sheet surface and parallel to the rolling direction will be observed. The through-thickness cross section is etched with 1 vol % Nital. Microstructure images of 3000 μm2 or larger regions are photographed at thickness t/4 locations with a scanning electron microscope (SEM) at a magnification of 2000 times. The images are analyzed to determine items (i) to (iv) below. Incidentally, the letter t indicates the sheet thickness and the letter w indicates the sheet width.
Polygonal ferrite (recrystallized F) and upper bainite (UB) are both gray in SEM images but can be distinguished by their shapes. An exemplary SEM image is illustrated in
Fresh martensite and retained γ are both white in SEM images and cannot be distinguished. Thus, retained γ was measured separately by a method described later. The total area fraction of fresh martensite and retained γ is measured from the SEM image by a point count method in accordance with ASTM E562-11 (2014), and the area fraction of retained γ measured by the method described later is subtracted from the total area fraction to determine the area fraction of fresh martensite. The total area fraction of fresh martensite and retained γ is measured by the point count method with respect to 5 locations, the measurement results being averaged, and the volume fraction of retained γ measured by the method described later is subtracted from the average value to give the area fraction of fresh martensite.
(iii) Tempered Martensite and/or Lower Bainite
Tempered martensite and lower bainite are carbide-containing microstructures that are seen as white fine microstructures in SEM images. These two microstructures can be distinguished by more microscopic observation but are difficult to distinguish by SEM images. Thus, in accordance with aspects of the present invention, tempered martensite and lower bainite are defined as the same microstructure, and the total area fraction of tempered martensite and lower bainite is measured by a point count method in accordance with ASTM E562-11 (2014). The results measured at 5 locations are averaged to give the total area fraction of tempered martensite and lower bainite.
The area fractions of polygonal ferrite, upper bainite, fresh martensite, retained γ, tempered martensite, and lower bainite measured by the above methods are subtracted from 100%. The difference is defined as the area fraction of the remaining microstructure.
The steel sheet is polished by ¼ sheet thickness and is further polished by 0.1 mm by chemical polishing. The exposed face is analyzed with an X-ray diffractometer using MoKα radiation to measure the integrated reflection intensities of (200) plane, (220) plane, and (311) plane of FCC iron (y), and of (200) plane, (211) plane, and (220) plane of BCC iron (ferrite). The volume fraction of retained γ is determined from the intensity ratio of the integrated reflection intensity of the planes of FCC iron (y) to the integrated reflection intensity of the planes of BCC iron (ferrite). In accordance with aspects of the present invention, the volume fraction of retained γ can be regarded as the area fraction of retained γ.
Equivalent Circular Diameter and Aspect Ratio of Fresh Martensite Grains and/or Retained γ Grains
The steel sheet is cut to give an observation specimen so that a cross section parallel to the rolling direction will be observed. The microstructure on the through-thickness cross section is exposed by etching with LePera etchant. Microstructure images of 10000 μm2 or larger regions are photographed at thickness t/4 locations with a laser microscope (LM) at a magnification of 1000 times. Lepera etching is color etching. Fresh martensite grains and/or retained γ grains are extracted by showing fresh martensite and/or retained γ in white contrast, and image analysis is performed to measure the equivalent circular diameter and the aspect ratio of the fresh martensite grains and/or the retained γ grains.
Of all the grains obtained, the number of grains having an equivalent circular diameter of less than 0.8 μm is determined, and the ratio of those grains to the number of all the grains is calculated.
Of all the grains obtained, the number of grains having an equivalent circular diameter of 0.8 μm or more is measured. Of those grains, the number of grains having an aspect ratio of 2.0 or more is determined. The ratio is calculated of the grains having an aspect ratio of 2.0 or more and an equivalent circular diameter of 0.8 μm or more to all the grains having an equivalent circular diameter of 0.8 μm or more.
Next, an embodiment of a method for manufacturing a steel sheet according to aspects of the present invention will be described in detail. Unless otherwise specified, the temperatures of heating or cooling of steel, for example, a steel slab (a steel material) or a steel sheet, described below means the surface temperature of the steel, for example, the steel slab (the steel material) or the steel sheet.
A method for manufacturing a steel sheet according to aspects of the present invention includes, after hot rolling and pickling are performed on a steel slab having the chemical composition described hereinabove, a cold rolling step of performing a cold rolling treatment on the hot rolled steel sheet to produce a cold rolled steel sheet, and an annealing step of performing an annealing treatment on the cold rolled steel sheet to produce a steel sheet. The cold rolling step is such that the cold rolled steel sheet is obtained by performing the cold rolling treatment in such a manner that the cumulative cold rolling reduction ratio is 30 to 85%, and the rolling reduction ratio in a first pass is 5% or more and less than 25%, thereby controlling the area fraction of the total of microstructures having {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100}<011> orientation to 35% or more and 75% or less relative to all bcc phase microstructures. The annealing step is such that the annealing treatment includes: heating the cold rolled steel sheet at an average heating rate of 0.5 to 15° C./sec or less in a range of temperatures of 500° C. or above and Ac or below, to an annealing temperature T being 840° C. or below and satisfying 0.6≤(T−Ac1)/(Ac3−Ac1)<1.0; after the heating, soaking and holding the steel sheet at the annealing temperature T in a furnace atmosphere having a dew point Td of −50° C. or above and −30° C. or below, thereby producing a steel sheet having a number density of acicular austenite microstructures of 5 microstructures/1000 μm2 or more; subsequently performing first cooling of cooling the steel sheet at an average cooling rate of 6.0° C./sec or more in a range of temperatures of 750 to 550° C., to a first cooling stop temperature Tc1 of 550° C. or below and 400° C. or above; after the first cooling, subjecting the steel sheet to first holding at the first cooling stop temperature Tc1 for 25 seconds or more; after the first holding, performing second cooling of cooling the steel sheet at an average cooling rate of 3.0 to 80° C./s in a range of temperatures of 350° C. or below and 200° C. or above, to a second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above; subjecting the steel sheet to second holding at the second cooling stop temperature Tc2 for 2 to 20 seconds; after the second holding, subjecting the steel sheet to over-aging and holding in a range of temperatures of 350 to 500° C. for 20 to 3000 seconds; and after the over-aging and holding, performing third cooling of cooling the steel sheet.
The steps will be described below.
In accordance with aspects of the present invention, for example, hot rolling in the hot rolling step may be performed in such a manner that the steel slab having the chemical composition described hereinabove is reheated and then rolled, that the steel slab from continuous casting is subjected to hot direct rolling without heating, or that the steel slab from continuous casting is heat treated for a short time and then rolled. The hot rolling may be performed in accordance with a conventional procedure. For example, the slab heating temperature may be 1100° C. or above and 1300° C. or below; the soaking time may be 20 to 30 minutes; the finish rolling temperature may be Ar3 transformation temperature (° C.) or above and Ar3 transformation temperature (° C.)+200° C. or below; and the coiling temperature may be 400 to 720° C. In order to eliminate or reduce thickness variations and to ensure high strength stably, the coiling temperature is preferably 430 to 530° C.
The steel slab (the steel material) may be produced by any smelting method without limitation. A known smelting technique, such as a converter or an electric arc furnace, may be used. Secondary refining may be performed in a vacuum degassing furnace.
In the pickling treatment step, the hot rolled steel sheet from the hot rolling step is subjected to a pickling treatment. The pickling treatment conditions are not particularly limited, and pickling treatment conditions in known production methods may be adopted.
Cumulative cold rolling reduction ratio: 30 to 85% If the rolling reduction ratio (the cumulative cold rolling reduction ratio) in the cold rolling treatment is less than 30%, recrystallization is not sufficiently promoted and acicular γ discussed in accordance with aspects of the present invention is not formed sufficiently. Furthermore, the desired cold rolled texture does not develop, and the total of microstructures having {111} <0-11> orientation, {111} <11-2>orientation, {211} <0-11> orientation, and {100} <011> orientation described later does not reach 35% or more relative to all the bcc phase microstructures. Thus, the rolling reduction ratio in cold rolling is limited to 30% or more. The rolling reduction ratio (the cumulative cold rolling reduction ratio) is preferably 40% or more, and more preferably 50% or more. On the other hand, the rolling reduction ratio (the cumulative cold rolling reduction ratio) is 85% or less from the point of view of cold rolling load or further from the point of view of material quality.
In the cold rolling step, the number of passes is not particularly limited, and may be, for example, 5 passes. The cumulative cold rolling reduction ratio (the thickness reduction ratio) indicates (1−(sheet thickness after cold rolling (after final pass)/sheet thickness before cold rolling)×100.
Rolling Reduction Ratio in the First Pass: 5% or More and Less than 25%
From the point of view of operability, the rolling reduction ratio in the first pass is limited to 5% or more. Because the sheet temperature at the time of the first pass of cold rolling is low, 25% or more rolling reduction ratio in the first pass applies shear strain components to the material being cold rolled, and the desired texture does not develop and acicular γ is not formed. Thus, the rolling reduction ratio in the first pass is limited to 5% or more and less than 25%.
Incidentally, the rolling reduction ratio (the thickness reduction ratio) in the first pass indicates (1−(sheet thickness after first pass of cold rolling)/(sheet thickness before cold rolling))×100.
The rolling temperature (the sheet temperature) in the first pass is preferably 20° C. or above and 40° C. or below. The rolling temperature in the first pass is determined by measuring the temperature of a portion of the steel sheet surface free from the lubricating oil after the first pass with a radiation thermometer. If the rolling temperature in the first pass is below 20° C. or if the rolling temperature in the first pass is above 40° C., the desired texture described hereinabove may not develop and acicular γ may not be formed. Thus, the rolling temperature in the first pass is preferably 20° C. or above and 40° C. or below.
Microstructure of the cold rolled steel sheet after cold rolling: The area fraction of the total of microstructures having {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100} <011> orientation is 35% or more and 75% or less relative to all the bcc phase microstructures.
Acicular γ has a specific crystallographic orientation relationship (Near Kurdjumov-Sachs relationship) with ferrite surrounding its nucleation sites.
By ensuring that the cold rolled steel sheet after cold rolling has a certain or higher area fraction of the total of microstructures having specific orientations, specifically, {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100} <011> orientation, relative to all the bcc phase microstructures, reverse transformed γ having the above specific orientations is formed easily between surrounding ferrite grains, and, as a result, a large amount of acicular γ is formed. In order to form a desired amount of acicular γ, it is necessary that the area fraction of the total of microstructures having {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100}<011> orientation be 35% or more relative to all the bcc phase microstructures. The fraction is preferably 40% or more. If, on the other hand, the area fraction of the total of microstructures having {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100} <011> orientation is more than 75% relative to all the bcc phase microstructures, anisotropy occurs in the material quality of the steel sheet. Thus, the area fraction of the total of microstructures having the above specific orientations is limited to 75% or less relative to all the bcc phase microstructures. The fraction is preferably 68% or less, and more preferably 65% or less.
In accordance with aspects of the present invention, the ratio of the area fraction of the total of microstructures having the above specific orientations to the area fraction of all the bcc phase microstructures can be brought to the desired range by subjecting the hot rolled steel sheet having the chemical composition described hereinabove to the cold rolling treatment with a cold rolling reduction ratio of 30 to 85% while controlling the rolling reduction ratio in the first pass to 5% or more and less than 25%.
The cold rolled steel sheet from the cold rolling step is cut to give a measurement specimen so that a cross section parallel to the rolling direction will be a measurement face. The measurement face is mechanically polished or electrolytically polished, and 80000 μm2 or larger regions are analyzed by SEM-EBSD (measurement conditions: WD: 20 mm, acceleration voltage: 20 kV). The area fraction is quantified of bcc phase microstructures in which {ND plane} <RD direction> rolling orientations are {111} <0-11> orientation, {111} <11-2> orientation, {211} <0-11> orientation, and {100}<011> orientation. The area fraction is expressed as a ratio to the area fraction of the bcc phases of all the orientations to evaluate the texture of the cold rolled steel sheet.
In the annealing step in accordance with aspects of the present invention, the cold rolled steel sheet from the cold rolling step is heated at an average heating rate (HR1) of 0.5 to 15° C./sec in a range of temperatures of 500° C. or above and Ac1 or below, to an annealing temperature T that is 840° C. or below and satisfies 0.6≤(T−Ac1)/(Ac3−Ac1)<1.0. After the heating, the steel sheet is soaked and held at the temperature T in a furnace atmosphere having a dew point Td of −50° C. or above and −30° C. or below, thereby giving a steel sheet having a number density of acicular γ microstructures of 5 microstructures/1000 μm2 or more. Subsequently, first cooling is performed in which the steel sheet is cooled at an average cooling rate of 6.0° C./sec or more in a range of temperatures of 750 to 550° C., to a first cooling stop temperature Tc1 of 550° C. or below and 400° C. or above. After the first cooling, the steel sheet is subjected to first holding at the first cooling stop temperature Tc1 for 25 seconds or more. After the first holding, second cooling is performed in which the steel sheet is cooled at an average cooling rate of 3.0 to 80° C./s in a range of temperatures of 350° C. or below and 200° C. or above, to a second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above. The steel sheet is then subjected to second holding at the second cooling stop temperature Tc2 for 2 to 20 seconds. After the second holding, the steel sheet is over-aged and held in a range of temperatures of 350 to 500° C. for 20 to 3000 seconds. After the over-aging and holding, third cooling is performed in which the steel sheet is cooled.
In accordance with aspects of the present invention, the cold rolled sheet that has the microstructure after the cold rolling step described above is heated at an appropriate heating rate to sufficiently promote recrystallization and thereafter acicular austenite is formed by heating of the steel sheet to the temperature T or by holding of the steel sheet at the temperature T. For this purpose, the average heating rate in a range of temperatures of 500° C. or above and Ac1 or below where austenite transformation does not occur is limited to 15° C./sec or less. The average heating rate is preferably 10° C./sec or less.
From the point of view of operation, the lower limit of the average heating rate is limited to 0.5° C./sec or more. The average heating rate is preferably 1.0° C./sec or more, and more preferably 1.5° C./sec or more.
Here, the average heating rate (° C./s) is calculated from ((Ac1 (° C.)−500° C.)/(heating time (sec) from 500° C. to Ac1 (° C.)).
Heating to the Annealing Temperature T that is 840° C. or Below and Satisfies 0.6≤(T−Ac1)/(Ac3−Ac1)<1.0
In accordance with aspects of the present invention, acicular austenite described later can be formed by heating the steel sheet to the temperature T (annealing temperature T) described later or by further holding the steel sheet at the annealing temperature T. If the steel sheet is heated to an austenite single-phase region of Ac3 (° C.) or above, acicular austenite coalesces with adjacent austenite, and the austenite morphology becomes equiaxed. Thus, in accordance with aspects of the present invention, the annealing needs to be two-phase annealing.
If the annealing temperature T is such that (T−Ac1)/(Ac3−Ac1) is less than 0.6, reverse transformation to austenite does not occur sufficiently and acicular austenite is not formed, and equiaxed austenite is exclusively formed along recrystallized ferrite grain boundaries. Furthermore, the amount of ferrite microstructures is so large that 980 MPa or higher strength may not be obtained.
Thus, the annealing temperature T is limited to satisfy 0.6≤(T−Ac1)/(Ac3−Ac1)<1.0. If the temperature T is above 840° C., good chemical convertibility cannot be obtained. Thus, the temperature T is limited to 840° C. or below.
If the dew point Td is below −50° C., good chemical convertibility cannot be obtained. If the dew point Td is above −30° C., good chemical convertibility cannot be obtained. Thus, the dew point Td is limited to −50° C. or above and −30° C. or below. The dew point Td is preferably −48° C. or above, and more preferably −45° C. or above. The dew point Td is preferably −32° C. or below, and more preferably −34° C. or below.
The soaking time at the annealing temperature T is not particularly limited but is preferably 25 to 350 seconds, and more preferably 50 to 300 seconds from the point of view of element partitioning during the two-phase annealing.
Incidentally, Ac1 (° C.) may be calculated from the formula below based on empirical rules.
Ac3 (° C.) may be calculated from the formula below based on empirical rules.
Incidentally, [X %] in the above formulas is the content (mass %) of component element X in the steel sheet and is “0” when the content is nil.
In accordance with aspects of the present invention, acicular γ is utilized to impart desired formability. Plenty of acicular austenite (acicular γ) promotes the formation of a large amount of retained γ having a high aspect ratio. In order to ensure that aspects of the present invention realize the desired formability, the number density of acicular γ microstructures formed by the heating to and the soaking and holding at the annealing temperature T needs to be 5 microstructures/1000 μm2 or more. By the nature of acicular γ, the upper limit is not limited and a larger number of acicular γ grains is more preferable.
In accordance with aspects of the present invention, the number density of acicular γ microstructures can be brought to the desired range by heating the cold rolled steel sheet having the chemical composition and the microstructure described hereinabove to the annealing temperature T in such a manner that the average heating rate in the range of temperatures of 500° C. or above and Ac or below is 0.5 to 15° C./sec or less, and by soaking and holding the steel sheet at the annealing temperature T in a furnace atmosphere satisfying the dew point Td.
When a microstructure formed at a high temperature is to be evaluated, a common practice is to freeze the microstructure by water-cooling and evaluate the microstructure that has been formed. In accordance with aspects of the present invention, it is important that acicular γ formed by treatments in the annealing step up to the soaking and holding at the annealing temperature T contribute, in the subsequent cooling process, to the formation of retained γ having a high aspect ratio and high working stability. Thus, the number density of the acicular γ microstructures is measured. The steel sheet is cut to give an observation specimen so that a cross section parallel to the rolling direction will be observed. The through-thickness cross section is etched with 1 vol % Nital. Microstructure images of 3000 μm2 or larger regions are photographed at thickness t/4 locations with a scanning electron microscope (SEM) at a magnification of 2000 times. The SEM image illustrated in
In the first cooling, ferrite transformation occurs predominantly in a range of temperatures of 750 to 550° C. In excessive ferrite transformation, acicular γ is transformed into ferrite. Thus, ferrite transformation is suppressed by controlling the average cooling rate in the range of temperatures of 750 to 550° C. to 6.0° C./sec or more. The average cooling rate is preferably 8.0° C./sec or more, and more preferably 10.0° C./sec or more.
Here, the average cooling rate (° C./sec) is calculated from (750° C. (cooling start temperature)−550° C. (finish cooling temperature))/(cooling time (sec) from cooling start temperature to cooling stop temperature).
The first cooling stop temperature Tc1 in the first cooling is a temperature for allowing upper bainite transformation to occur. If the first cooling stop temperature Tc1 is above 550° C., non-transformed austenite is transformed into ferrite and/or pearlite and the formation of retained austenite is suppressed, with the result that desired ductility cannot be ensured. If, on the other hand, the first cooling stop temperature Tc1 is below 400° C., non-transformed austenite is transformed into martensite and carbon (C) cannot be efficiently partitioned to non-transformed austenite, with the result that ductility is lowered. Thus, the first cooling stop temperature Tc1 is limited to 550° C. or below and 400° C. or above. The first cooling stop temperature Tc1 is preferably 500° C. or below. The first cooling stop temperature Tc1 is preferably 420° C. or above.
In the first holding after the first cooling, the holding time at the first cooling stop temperature Tc1 is limited to 25 seconds or more in order to allow bainite transformation to occur sufficiently. The holding time in the first holding is preferably 30 seconds or more, and more preferably 35 seconds or more. The holding time in the first holding is preferably 60 seconds or less, and more preferably 55 seconds or less. In the first holding after the first cooling, the first cooling stop temperature Tc1 may be modulated as long as the temperature is in the range of 550° C. or below and 400° C. or above.
Second cooling: The steel sheet is cooled at an average cooling rate of 3.0 to 80° C./s in a range of temperatures of 350° C. or below and 200° C. or above, to a second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above.
After the first holding, second cooling is performed. First, the steel sheet is cooled at an average cooling rate of 3.0 to 80° C./s in a range of temperatures of 350° C. or below and 200° C. or above. If the average cooling rate in the range of temperatures of 350° C. or below and 200° C. or above exceeds 80° C./s, the cooling is so rapid that the sheet shape is deteriorated. If, on the other hand, the average cooling rate is less than 3.0° C./s, martensite transformation and carbon partitioning compete with each other and non-transformed γ is stabilized. As a result, a large amount of fresh martensite is formed after the final cooling. Thus, in accordance with aspects of the present invention, the average cooling rate in the range of temperatures of 350° C. or below and 200° C. or above is limited to 3.0 to 80° C./s. The average cooling rate is preferably 60° C./sec or less, and more preferably 50° C./sec or less. The average cooling rate is preferably 5.0° C./sec or more, and more preferably 10.0° C./sec or more.
Here, the average cooling rate (° C./sec) is calculated from (350° C. (cooling start temperature)−200° C. (cooling stop temperature))/(cooling time (sec) from cooling start temperature to cooling stop temperature).
When the steel sheet is cooled to the second cooling stop temperature Tc2 of 320° C. or below and 150° C. or above, martensite transformation occurs. If the second cooling stop temperature Tc2 is above 320° C., martensite transformation does not occur and coarse fresh martensite grains and/or retained austenite is formed during the final cooling to cause deterioration in local elongation and flangeability. As a result, desired formability cannot be ensured. Thus, the upper limit of the second cooling stop temperature Tc2 is limited to 320° C. The second cooling stop temperature Tc2 is preferably 300° C. or below, and more preferably 280° C. or below. If, on the other hand, the second cooling stop temperature Tc2 is below 150° C., most of non-transformed austenite is transformed into martensite and little retained austenite is formed, with the result that ductility is deteriorated. Thus, the lower limit of the second cooling stop temperature Tc2 is limited to 150° C. The second cooling stop temperature Tc2 is preferably 170° C. or above, and more preferably 190° C. or above.
Second holding: The steel sheet is held at the second cooling stop temperature Tc2 for 2 to 20 seconds or less.
In the second holding after the second cooling, the holding time at the second cooling stop temperature Tc2 is limited to 2 seconds or more. In this manner, martensite transformation occurs sufficiently during cooling to the second cooling stop temperature Tc2. The martensite microstructure that is obtained is uniform in the width direction and the thickness direction, and the variations in material quality can be reduced. On the other hand, the holding time in the second holding is limited to 20 seconds or less from the point of view of operation. The holding time in the second holding is preferably 4 seconds or more, and more preferably 6 seconds or more. The holding time in the second holding is preferably 17 seconds or less, and more preferably 14 seconds or less.
In the second holding after the second cooling, the second cooling stop temperature Tc2 may be modulated as long as the temperature is in the range of 320° C. or below and 150° C. or above.
Over-aging and holding: The steel sheet is held in a range of temperatures of 350 to 500° C. for 20 to 3000 seconds.
The steel sheet is over-aged and held in a range of temperatures of 350 to 500° C. in order to transform non-transformed austenite to bainite and further to temper a martensite microstructure formed by the cooling to the second cooling stop temperature Tc2 and thereby to promote partitioning of carbon to non-transformed austenite. If the over-aging and holding temperature is above 500° C., retained austenite is decomposed, cementite is precipitated, or further part of the microstructure undergoes pearlite transformation, with the result that ductility is lowered. If, on the other hand, the over-aging and holding temperature is below 350° C., non-transformed austenite is not transformed and, in addition, carbon is not partitioned from martensite formed at the second cooling stop temperature Tc2, with the result that retained austenite having low mechanical stability is formed at the time of the final cooling. Thus, the range of temperatures at which the steel sheet is over-aged and held is limited to 350 to 500° C. When the holding time in the over-aging and holding is 20 seconds or more, non-transformed austenite is transformed to bainite and carbon is partitioned from martensite formed at the second cooling stop temperature Tc2, with the result that desired formability can be ensured. On the other hand, the holding time in the over-aging and holding is limited to 3000 seconds or less in view of operability.
Third cooling: After the over-aging and holding, the steel sheet is cooled.
After the over-aging and holding, the steel sheet is cooled to room temperature (10 to 30° C.). The steel sheet according to aspects of the present invention is thus obtained.
After the annealing step, for example, temper rolling with an elongation ratio of 0.05 to 0.5% may be performed. However, the post treatment is not particularly limited thereto.
The steel sheet according to aspects of the present invention that is obtained by the steel sheet manufacturing method according to aspects of the present invention preferably has a thickness of 0.5 mm or more. The thickness is preferably 2.0 mm or less.
Next, a member and a method for manufacture thereof according to aspects of the present invention will be described.
The member according to aspects of the present invention is obtained by subjecting the steel sheet according to aspects of the present invention to at least one working of forming and joining. The method for manufacturing a member according to aspects of the present invention includes a step of subjecting the steel sheet according to aspects of the present invention to at least one working of forming and joining to produce a member.
The steel sheet according to aspects of the present invention has a tensile strength of 980 MPa or more and has high ductility, excellent stretch flange formability, and good chemical convertibility. Thus, the member that is obtained using the steel sheet according to aspects of the present invention also has high strength and has high ductility, excellent stretch flange formability, and good chemical convertibility compared to the conventional high-strength members. Furthermore, weight can be reduced by using the member according to aspects of the present invention. Thus, for example, the member according to aspects of the present invention may be suitably used in an automobile body frame part. The member according to aspects of the present invention also includes a welded joint.
The forming may be performed using any common working process, such as press working, without limitation.
Furthermore, the joining may be performed using common welding, such as spot welding or arc welding, or, for example, riveting or caulking without limitation.
EXAMPLES of the present invention will be described below.
Steel slabs having a thickness of 250 mm and a chemical composition described in Table 1 were hot rolled (slab heating temperature: 1250° C., soaking time: 30 minutes, finish rolling temperature: Ar3+50° C., coiling temperature: 550° C.) and pickled. The hot rolled steel sheets obtained were cold rolled under conditions described in Table 2. Cold rolled steel sheets were thus manufactured.
Next, the cold rolled steel sheets were annealed in a continuous annealing line under conditions described in Table 2 and were then temper rolled with an elongation ratio of 0.2 to 0.4%. Evaluation steel sheets were thus manufactured. After over-aging and holding (isothermal holding), third cooling was performed, and the steel sheets were cooled to room temperature (20° C.).
I
1.71
0.50
J
3.72
K
0.255
L
1.62
M
0.089
N
0.55
1.2
35
35
22.6
22
850
0.3
28
22.2
0.5
0.5
−25
30
28.3
−55
I
J
K
L
M
N
0.0
2.4
570
2.4
3.3
300
350
530
120
2.4
3.2
3.3
250
0.0
360
The steel sheets obtained were evaluated in the following manner.
The steel sheet was cut to give an observation specimen so that a cross section perpendicular to the steel sheet surface and parallel to the rolling direction would be observed. The through-thickness cross section was etched with 1 vol % Nital. Microstructure images of 3000 μm2 or larger regions were photographed at thickness t/4 locations with a scanning electron microscope (SEM) at a magnification of 2000 times. The images were analyzed to determine items (i) to (iv) below. The results are described in Table 3. Incidentally, the letter t indicates the sheet thickness and the letter w indicates the sheet width.
Polygonal ferrite (recrystallized F) and upper bainite (UB) are both gray in SEM images but can be distinguished by their shapes. An exemplary SEM image is illustrated in
Fresh martensite and retained γ are both white in SEM images and cannot be distinguished. Thus, retained γ was measured separately by a method described later. The total area fraction of fresh martensite and retained γ was measured from the SEM image by a point count method in accordance with ASTM E562-11 (2014), and the area fraction of retained γ measured by the method described later was subtracted from the total area fraction to determine the area fraction of fresh martensite. The total area fraction of fresh martensite and retained γ was measured by the point count method with respect to 5 locations, the measurement results being averaged, and the volume fraction of retained γ measured by the method described later was subtracted from the average value to give the area fraction of fresh martensite.
(iii) Tempered Martensite and/or Lower Bainite
Tempered martensite and lower bainite are carbide-containing microstructures that are seen as white fine microstructures in SEM images. These two can be distinguished by more microscopic observation but are difficult to distinguish by SEM images. Thus, in accordance with aspects of the present invention, tempered martensite and lower bainite were defined as the same microstructure, and the total area fraction of tempered martensite and lower bainite was measured by a point count method in accordance with ASTM E562-11 (2014). The results measured at 5 locations were averaged to give the total area fraction of tempered martensite and lower bainite.
The area fractions of polygonal ferrite, upper bainite, fresh martensite, retained γ, tempered martensite, and lower bainite measured by the above methods were subtracted from 100%. The difference was defined as the area fraction of the remaining microstructure.
The steel sheet was polished by ¼ sheet thickness and was further polished by 0.1 mm by chemical polishing. The exposed face was analyzed with an X-ray diffractometer using MoKα radiation to measure the integrated reflection intensities of (200) plane, (220) plane, and (311) plane of FCC iron (γ), and of (200) plane, (211) plane, and (220) plane of BCC iron (ferrite). The volume fraction of retained γ was determined from the intensity ratio of the integrated reflection intensity of the planes of FCC iron (γ) to the integrated reflection intensity of the planes of BCC iron (ferrite). In accordance with aspects of the present invention, the volume fraction of retained γ can be regarded as the area fraction of retained γ.
When a microstructure formed at a high temperature is to be evaluated, a common practice is to freeze the microstructure by water-cooling and evaluate the microstructure that has been formed. In accordance with aspects of the present invention, it is important that acicular γ formed by treatments in the annealing step up to the holding at the temperature T contribute, in the subsequent cooling process, to the formation of retained γ having a high aspect ratio and high working stability. Thus, the number density of the acicular γ microstructures was measured. The steel sheet was cut to give an observation specimen so that a cross section parallel to the rolling direction would be observed. The through-thickness cross section was etched with 1 vol % Nital. Microstructure images of 3000 μm2 or larger regions were photographed at thickness t/4 locations with a scanning electron microscope (SEM) at a magnification of 2000 times. The SEM image illustrated in
(4) Equivalent Circular Diameter and Aspect Ratio of Fresh Martensite Grains and/or Retained γ Grains
The steel sheet was cut to give an observation specimen so that a cross section parallel to the rolling direction would be observed. The microstructure on the through-thickness cross section was exposed by etching with LePera etchant. Microstructure images of 10000 μm2 or larger regions were photographed at thickness t/4 locations with a laser microscope (LM) at a magnification of 1000 times. Lepera etching is color etching. Fresh martensite grains and/or retained γ grains were extracted by showing fresh martensite and/or retained γ in white contrast, and image analysis was performed to measure the equivalent circular diameter and the aspect ratio of the fresh martensite grains and/or the retained γ grains.
Of all the grains obtained, the number of grains having an equivalent circular diameter of less than 0.8 μm was determined, and the ratio of those grains to the number of all the grains was calculated.
Of all the grains obtained, the number of grains having an equivalent circular diameter of 0.8 μm or more was measured. Of those grains, the number of grains having an aspect ratio of 2.0 or more was determined. The ratio was calculated of the grains having an aspect ratio of 2.0 or more and an equivalent circular diameter of 0.8 μm or more to all the grains having an equivalent circular diameter of 0.8 μm or more. The results are described in Table 3.
The cold rolled steel sheet from the cold rolling step was cut to give a measurement specimen so that a cross section parallel to the rolling direction would be a measurement face. The measurement face was mechanically polished or electrolytically polished, and 80000 μm2 or larger regions were analyzed by SEM-EBSD (measurement conditions: WD: 20 mm, acceleration voltage: 20 kV). The area fraction was quantified of bcc phase microstructures in which {ND plane} <RD direction> rolling orientations were {111} <0-11> orientation, {111} <11-2> orientation, {211}<0-11> orientation, and {100} <011> orientation. The area fraction was expressed as a ratio to the area fraction of the bcc phases of all the orientations to evaluate the texture of the cold rolled steel sheet.
JIS No. 5 test pieces for tensile test were fabricated from the steel sheet so that the tensile direction would be perpendicular to the rolling direction. The test pieces were each subjected to a tensile test in accordance with the provisions of JIS Z2241 (2011). The crosshead speed in the tensile test was 10 mm/min. The measurement was performed twice, and the measured values were averaged to give the tensile strength (TS) of the steel sheet.
100 mm×100 mm test specimens were sampled. Three test specimens from each sampling location were tested by a hole expansion test in accordance with JFST (The Japan Iron and Steel Federation Standard) 1001. The results of the three measurements were averaged (total of values of three measurements (%)/3) to give the hole expansion ratio z (%)
In accordance with aspects of the present invention, the steel sheets were evaluated as having high strength when the tensile strength (TS) was 980 MPa or more.
The ductility El was evaluated as excellent when tensile strength (TS)×total elongation (T. El)≥18000 MPa-% or more. The stretch flange formability λ was evaluated as excellent when the hole expansion ratio z (%) 45%.
The steel sheet after annealing was electrolytically pickled with sulfuric acid for 2 seconds at a current density of 20 to 35 A/dm2 and was degreased and surface conditioned. Subsequently, chemical conversion was performed using a zinc phosphate chemical conversion solution. In the chemical conversion, the degreasing step involved treatment temperature: 40° C., treatment time: 120 seconds, and spray degreasing; the surface conditioning step involved pH: 9.5, treatment temperature: room temperature, and treatment time: 20 seconds; and the chemical conversion step involved chemical conversion solution temperature: 35° C. and treatment time: 120 seconds. The degreasing step, the surface conditioning step, and the chemical conversion step involved the following treatment agents, respectively: degreasing agent: FC-E2011, surface conditioning agent: PL-X, and chemical conversion solution: PALBOND PB-L3065, each manufactured by Nihon Parkerizing Co., Ltd. The surface chemical conversion microstructure was observed with respect to 10000 μm2 or larger regions by SEM at a magnification of 2000 times. The chemical convertibility was evaluated as o when the chemical conversion coating microstructure was present on all the faces, and was evaluated as x when the chemical conversion coating microstructure was visually found to be absent no matter how small the size. The results are described in Table 3.
54.6
21.2
11.4
24.6
76.1
22.7
18.3
6.0
53.1
27.2
16.7
I
33.6
25.7
J
32.0
K
33.6
L
12.2
M
N
51.6
22
14815
41
x
24
18
32
14683
24
17450
40
17
34
15790
x
24
29
14774
12
24
24
17657
24
26
15655
33
21
x
22
16000
41
x
32
33
29
x
x
22
28
As described in Table 3, the steel sheets according to aspects of the present invention were shown to have 980 MPa or higher tensile strength, high ductility, and excellent stretch flange formability and also to be excellent in chemical convertibility.
The steel sheets of INVENTIVE EXAMPLES have high strength, high ductility, excellent stretch flange formability, and good chemical convertibility. This has shown that members obtained by forming of the steel sheets of INVENTIVE EXAMPLES, members obtained by joining of the steel sheets of INVENTIVE EXAMPLES, and members obtained by forming and joining of the steel sheets of INVENTIVE EXAMPLES will have high strength, high ductility, excellent stretch flange formability, and good chemical convertibility similarly to the steel sheets of INVENTIVE EXAMPLES.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2021-160629 | Sep 2021 | JP | national |
This is the U.S. National Phase application of PCT/JP2022/033946, filed Sep. 9, 2022, which claims priority to Japanese Patent Application No. 2021-160629 filed Sep. 30, 2021, the disclosures of these applications being incorporated by reference in their entireties for all purposes.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2022/033946 | 9/9/2022 | WO |
