HIGH STRENGTH STEEL SHEET HAVING EXCELLENT FORMABILITY AND METHOD FOR MANUFACTURING THE SAME

- JFE STEEL CORPORATION

A thin high strength steel sheet having excellent formability has a composition which includes, by mass %, 0.08 to 0.15% of C, 0.5 to 1.5% of Si, 0.5 to 1.5% of Mn, 0.01 to 0.1% of Al, and 0.005% or less of N to form a hot rolled sheet is conducted.

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Description
RELATED APPLICATIONS

This is a §371 of International Application No. PCT/JP2011/065415, with an inter-national filing date of Jun. 29, 2011 (WO 2012/002566 A1, published Jan. 5, 2012), which is based on Japanese Patent Application No. 2010-147419, filed Jun. 29, 2010, the subject matter of which is incorporated herein by reference.

TECHNICAL FIELD

This disclosure relates to a high strength steel sheet which is required to have excellent formability (stretch flangeability) and is suitable for use as a strength member or the like of automobile parts, and a method of manufacturing the high strength steel sheet.

BACKGROUND

In recent years, improvement in fuel economy of vehicles has become an important issue from the point of view of preservation of the global environment. Thus, trends toward increasing the strength of the materials used, decreasing the thickness of the structural members used, and reducing the weight of automobile bodies have accelerated. As materials to be used, high strength steel sheets having a tensile strength of 540 MPa or higher are particularly required. However, since increasing the strength of steel sheets leads to a decrease in formability, high strength steel sheets having excellent formability are in demand. This demand is particularly high for steel sheets having a small thickness (thin steel sheets).

To meet such demand, various multi phase steel sheets have been proposed such as a dual phase steel sheet (DP steel sheet) having a dual phase microstructure composed of a ferrite phase and a martensite phase and a steel sheet having a multi phase microstructure including a ferrite phase, a martensite phase, and a bainite phase.

For example, Japanese Unexamined Patent Application Publication No. 63-293121 describes a method of manufacturing a high strength cold rolled steel sheet having excellent local ductility. According to that method, a cold rolled steel sheet having a composition including 0.08 to 0.30% of C, 0.1 to 2.5% of Si, 0.5 to 2.5% of Mn, and 0.01 to 0.15% of P is subjected to recrystallization annealing at a temperature equal to or higher than an Ac1 point, forcibly air-cooled to a temperature region ranging from an Ar1 point to 600° C., rapidly cooled at a cooling rate equal to or higher than 100° C./s to form a multi phase microstructure composed of a ferrite phase and a low-temperature transformed phase, and overaged at a temperature in the range of 350° C. to 600° C. so that a ratio Hv (L)/Hv (α) of the hardness Hv (L) of the low-temperature transformed phase to ferrite hardness Hv (α), which is obtained by a predetermined relational expression, is in the range of 1.5 to 3.5. According to the technique described in JP '121, the volume fraction of the low-temperature transformed phase is increased by increasing a quenching start temperature and then the overaging is performed at a temperature of 350° C. to 600° C. to precipitate C in the ferrite, soften the low-temperature transformed phase, and thereby reduce the ratio Hv (L)/Hv (α) and improve local elongation.

Japanese Unexamined Patent Application Publication No. 05-112832 describes a method of manufacturing a high tensile hot rolled steel sheet with a low yield ratio and excellent corrosion resistance, the method including hot-rolling a steel slab containing 0.02 to 0.25% of C, 2.0% or less of Si, 1.6 to 3.5% of Mn, 0.03 to 0.20% of P, 0.02% or less of S, 0.05 to 2.0% of Cu, 0.005 to 0.100% of sol. Al, and 0.008% or less of N to form a hot rolled coil, pickling the hot rolled coil, and annealing the hot rolled coil at a temperature of 720° C. to 950° C. by a continuous annealing line. According to the technique described in JP '832, it is possible to manufacture a high tensile hot rolled steel sheet which maintains a low yield ratio, high ductility, and excellent hole expandability, exhibits excellent corrosion resistance, and has a multi phase microstructure.

Japanese Unexamined Patent Application Publication No. 10-60593 describes a high strength cold rolled steel sheet with an excellent balance between the strength and the stretch flangeability. This high strength cold rolled steel sheet has a composition containing 0.03 to 0.17% of C, 1.0% or less of Si, 0.3 to 2.0% of Mn, 0.010% or less of P, 0.010% or less of S, and 0.005 to 0.06% of Al and satisfying C (%)>(3/40)×Mn, has a microstructure composed of a ferrite phase and a second phase including mainly bainite or pearlite, and satisfies (Vickers hardness of second phase)/(Vickers hardness of ferrite phase)<1.6. The high strength cold rolled steel sheet described in JP '593 is obtained by an annealing treatment followed by an overaging treatment at a temperature of 500 to 250° C. In that annealing treatment, steel (slab) having the above-described composition is hot-rolled, coiled at a temperature equal to or lower than 650° C., pickled, cold-rolled, soaked at a temperature equal to or higher than an A1 point and equal to or lower than (A3 point +50° C.), gradually cooled to a temperature T1 in the range of 750° C. to 650° C. at a rate of 20° C./s or lower, and cooled T1 to 500° C. at a rate of 20° C./s or higher.

However, the technique described in JP '121 has problems in that a continuous annealing facility which can perform rapid cooling (quenching) after recrystallization annealing is required, and addition of large amounts of alloy elements is required to suppress a rapid decrease in strength due to the overaging at a high temperature.

In the technique described in JP '832, it is essential to add large amounts of P and Cu in combination. However, when a large amount of Cu is contained, hot formability decreases, and when a large amount of P is contained, steel is embrittled. In addition, P shows a marked tendency to segregate in the steel, and this segregated P causes problems such as a decrease in the stretch flangeability of the steel sheet and embrittlement of a welded portion.

The high strength cold rolled steel sheet described in JP '593 has excellent stretch flangeability. However, at a strength as high as 540 MPa or higher, the elongation is less than 26% and a problem occurs in that the elongation sufficient for maintaining desired excellent formability cannot be ensured.

It could therefore be helpful to provide a high strength steel sheet having a small sheet thickness of about 1.0 to 1.8 mm and excellent formability and a method for manufacturing the high strength steel sheet. The “high strength” means that the steel sheet has a tensile strength TS equal to or higher than 540 MPa and preferably equal to or higher than 590 MPa. The “excellent formability” means that the elongation E1 is equal to or greater than 30% (with a JIS No. 5 test piece) and a hole expanding ratio λ, in a hole expanding test based on the Japan Iron and Steel Federation Standard JFST 1001-1996 is equal to or higher than 80%.

SUMMARY

We thus provide:

    • (1) A high strength steel sheet having excellent formability which has a composition including, by mass %, 0.08 to 0.15% of C, 0.5 to 1.5% of Si, 0.5 to 1.5% of Mn, 0.1% or less of P, 0.01% or less of S, 0.01 to 0.1% of Al, 0.005% or less of N, and the balance Fe with inevitable impurities, the steel sheet having a microstructure composed of a ferrite phase which is a main phase and a second phase including at least pearlite, wherein an area fraction of the ferrite phase is in the range of 75 to 90% and an area fraction of the pearlite is in the range of 10 to 25% with respect to the entire microstructure, an average grain size of the pearlite is 5 μm or smaller, and an area fraction of the pearlite is 70% or greater with respect to the total area of the second phase.
    • (2) The high strength steel sheet according to (1), in which the composition further includes one or more selected from among 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.
    • (3) The high strength steel sheet according to (1) or (2), in which the composition further includes one or two selected from between 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb by mass %.
    • (4) The high strength steel sheet according to any one of (1) to (3), in which the composition further includes 0.0003 to 0.0050% of B by mass %.
    • (5) The high strength steel sheet according to any one of (1) to (4), in which the composition further contains one or two selected from between 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu by mass %.
    • (6) The high strength steel sheet according to any one of (1) to (5), in which the composition further contains one or two selected from between 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM by mass %.
    • (7) A method for manufacturing a high strength steel sheet having excellent formability, including: a hot rolling step of hot-rolling a steel having a composition which includes, by mass %, 0.08 to 0.15% of C, 0.5 to 1.5% of Si, 0.5 to 1.5% of Mn, 0.1% or less of P, 0.01% or less of S, 0.01 to 0.1% of Al, 0.005% or less of N, and the balance Fe with inevitable impurities to form a hot rolled sheet; and a continuous annealing step including an annealing treatment of pickling the hot rolled sheet, and holding the pickled hot rolled sheet in a first temperature region of an Ac1 transformation point to an Ac3 transformation point for 5 to 400 s, and a cooling treatment of cooling the sheet at an average cooling rate of 5° C./s or higher from the first temperature region to 700° C. after the annealing treatment and adjusting a residence time in a second temperature region of 700° C. to 400° C. in the range of 30 to 400 s by using a continuous annealing line.
    • (8) The method for manufacturing a high strength steel sheet according to (7), in which the hot rolling step includes heating the steel at a temperature in the range of 1100° C. to 1280° C., hot-rolling the heated steel with a finish hot rolling temperature of 870° C. to 950° C. to form a hot rolled sheet, and coiling the hot rolled sheet at a coiling temperature of 350° C. to 720° C. upon completion of the hot rolling.
    • (9) The method for manufacturing a high strength steel sheet according to (7) or (8), in which a cooling time in a temperature region of 700° C. to 550° C. in the second temperature region is 10 s or longer.
    • (10) The method for manufacturing a high strength steel sheet according to any one of (7) to (9), in which the composition further includes one or more selected from among 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.
    • (11) The method for manufacturing a high strength steel sheet according to any one of (7) to (10), in which the composition further includes one or two selected from between 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb by mass %.
    • (12) The method for manufacturing a high strength steel sheet according to any one of (7) to (11), in which the composition further includes 0.0003 to 0.0050% of B by mass %.
    • (13) The method for manufacturing a high strength steel sheet according to any one of (7) to (12), in which the composition further includes one or two selected from between 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu by mass %.
    • (14) The method for manufacturing a high strength steel sheet according to any one of (7) to (13), in which the composition further includes one or two selected from between 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM by mass %.

A high strength steel sheet with excellent formability having a high strength, i.e., a tensile strength TS of 540 MPa or higher, an elongation E1 of 30% or greater, and a stretch flangeability λ of 80% or higher can be easily manufactured at a low cost, and thus our steel sheets and methods have particularly significant industrial advantages. Our steel sheets and methods also have an effect of significantly contributing to reduction of manufacturing cost, improvements of productivity, etc., since cold rolling can be omitted. When the steel sheet is applied to parts of automobile bodies, it can significantly contribute to the weight-reduction of automobile bodies.

DETAILED DESCRIPTION

We discovered that a microstructure composed of a ferrite phase as a main phase and a second phase including mainly fine pearlite can be formed by subjecting a hot rolled sheet in which the amounts of alloy elements are adjusted within an appropriate range to an annealing treatment which includes heating to an appropriate dual phase temperature region and an appropriate cooling treatment without cold rolling, and thus a desired high strength can be ensured, formability is significantly improved, and a high strength steel sheet having excellent formability with desired elongation and a desired hole expanding ratio can be obtained.

The detailed mechanism with regard to a significant improvement in formability by directly performing an appropriate annealing treatment on a hot rolled sheet without cold rolling has not been clear until now. However, we believe the following:

    • When a hot rolled sheet is subjected to an annealing treatment of heating the sheet to a dual phase temperature region without cold rolling, the only transformation that occurs during the annealing heating is from α to γ and new recrystallization does not occur. In this case, transformation from α to γ only occurs preferentially in portions with a high C concentration, and a more uniform microstructure can be obtained. In addition, C that diffuses rapidly is redistributed into α and γ up to an equilibrium composition during the annealing treatment. Therefore, presumably, precipitation of film cementite at grain boundaries has been suppressed, which particularly contributes to improvements in stretch flangeability. In contrast, when the hot rolled sheet is cold-rolled and then subjected to an annealing treatment, recrystallization and α to γ transformation competitively occur during the annealing heating, and thus the microstructure tends to become nonuniform and a significant improvement in formability is not readily expected.

First, reasons for limitations on the composition of a steel sheet will be described. Hereinafter, mass % will be simply expressed by % unless otherwise noted.

C: 0.08 to 0.15%

C is an element that contributes to an increase in the strength of a steel sheet and effectively acts on formation of a multi phase microstructure composed of a ferrite phase and a second phase other than the ferrite phase. 0.08% or more of C is required to ensure a desired high strength, i.e., a tensile strength of 540 MPa or higher. When more than 0.15% of C is present, spot weldability is lowered and formability such as ductility is lowered. Therefore, the C content is 0.08 to 0.15% and preferably 0.10 to 0.15%.

Si: 0.5 to 1.5%

Si is an element that dissolves in steel and effectively acts to strengthen the ferrite, and also contributes to an improvement in ductility. 0.5% or more of Si is required to ensure a desired high strength, i.e., a tensile strength of 540 MPa or higher. When an excessive amount more than 1.5% of Si is present, generation of red scale and the like is accelerated, the surface quality of a steel sheet is lowered, and chemical treatability is lowered. When an excessive amount of Si is contained, resistance weldability is deteriorated with an increase in electric resistance in resistance welding. Therefore, the Si content is 0.5 to 1.5% and preferably 0.7 to 1.2%.

Mn: 0.5 to 1.5%

Mn is an element that contributes to an increase in the strength of a steel sheet and effectively acts on formation of a multi phase microstructure. 0.5% or more of Mn is required to obtain such an effect. When more than 1.5% of Mn is present, a martensite phase is easily formed in the course of cooling in the annealing, and thus formability, particularly, stretch flangeability, is lowered. Therefore, the Mn content is 0.5 to 1.5% and preferably 0.7 to 1.5%.

P: 0.1% or less

P is an element that dissolves in steel and acts to increase the strength of a steel sheet. However, P shows a marked tendency to segregate to grain boundaries and lowers the bonding power of grain boundaries. This results in a decrease in formability and concentration of P in the surface of the steel sheet, thereby decreasing chemical treatability and corrosion resistance. Such an adverse effect of P is notably shown when more than 0.1% of P is present. Therefore, the P content is 0.1% or less. The P content is preferably decreased to 0.1% or less as much as possible to avoid such an adverse effect of P. However, the excessive decrease leads to a rise in manufacturing cost, and thus the P content is preferably about 0.001% or more.

S: 0.01% or less

S mainly forms sulfides (inclusions) such as MnS in steel and lowers formability of a steel sheet, particularly, local elongation. In addition, the presence of sulfides (inclusions) also lowers weldability. Such an adverse effect of S is notably shown when more than 0.01% of S is present. Therefore, the S content is 0.01% or less. The S content is preferably decreased to 0.01% or less as much as possible to avoid such an adverse effect of S. However, the excessive decrease leads to a rise in manufacturing cost, and thus the S content is preferably about 0.0001% or more.

Al: 0.01 to 0.1%

Al is an element that acts as a deoxidizing agent and is necessary to improve the cleanliness of a steel sheet. Furthermore, Al effectively acts to improve the yield of carbide-forming elements. 0.01% or more of Al is required to obtain such an effect. When less than 0.01% of Al is present, Si-based inclusions which serve as starting points for delayed fracture are not sufficiently removed, and thus the risk of occurrence of delayed fracture is increased. However, when more than 0.1% of Al is present, the above-described effect is saturated and thus the effect matching the content cannot be expected, resulting in economic disadvantages. In addition, formability is lowered and the tendency to generate surface defects is increased. Therefore, the Al content is 0.01 to 0.1% and preferably 0.01 to 0.05%.

N: 0.005% or less

The N content is preferably decreased as much as possible since N is an intrinsically harmful element, but up to 0.005% of N can be permitted. Therefore, the N content is 0.005% or less. Since excessively decreasing the N content leads to a rise in manufacturing cost, the N content is preferably about 0.0001% or more.

The above-described components are basic components. However, in addition to the basic components, one or more selected from among 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo, and/or one or two selected from between 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb, and/or 0.0003 to 0.0050% of B, and/or one or two selected from between 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu, and/or one or two selected from between 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM can be present in accordance with need.

One or More Selected from Among 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo

All of Cr, V, and Mo are elements that increase the strength of a steel sheet and contribute to formation of a multi phase microstructure, and one or more selected in accordance with need can be present. It is desired that 0.05% or more of Cr, 0.005% or more of V, and 0.005% or more of Mo be contained to obtain such an effect. When more than 0.5%, more than 0.2%, and more than 0.2% of Cr, V, and Mo, respectively, are present, it is difficult to form the desired amount of pearlite in the cooling treatment after the annealing treatment, and thus a desired multi phase microstructure cannot be ensured, thereby lowering stretch flangeability and formability. Therefore, when Cr, V, and/or Mo are present, the Cr content is preferably 0.05 to 0.5%, the V content is preferably 0.005 to 0.2%, and the Mo content is preferably 0.005 to 0.2%.

One or Two Selected from Between 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb

Both of Ti and Nb are elements that increase the strength of a steel sheet by precipitation strengthening, and one or two selected in accordance with need can be present. It is desired that 0.01% or more of Ti and 0.01% or more of Nb be contained, respectively, to obtain such an effect. When more than 0.1% of Ti and more than 0.1% of Nb are present, formability and shape fixability are lowered. Therefore, when Ti and/or Nb is present, the Ti content is preferably 0.01 to 0.1% and the Nb content is preferably 0.01 to 0.1%.

B: 0.0003 to 0.0050%

B is an element that segregates to austenite grain boundaries and acts to suppress formation and growth of ferrite from the grain boundaries. B can be present in accordance with need. It is desired that 0.0003% or more of B be present. However, when more than 0.0050% of B is present, formability is lowered. Therefore, when B is present, the B content is preferably 0.0003 to 0.0050%. In addition, to obtain the above-described effect of B, it is necessary to suppress formation of BN, and Ti is preferably present together with B.

One or Two Selected from Between 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu

Both of Ni and Cu are elements that act to increase the strength of a steel sheet and also act to promote internal oxidation to thereby improve adhesion of the coating. Ni and Cu can be selected and present in accordance with need. It is desired that 0.05% or more of Ni and 0.05% or more of Cu be present, respectively, to obtain such an effect. However, when more than 0.5% of Ni and more than 0.5% of Cu are present, it is difficult to form a desired amount of pearlite in the cooling treatment after the annealing treatment, and thus a desired multi phase microstructure cannot be ensured and stretch flangeability and formability are lowered. Therefore, when Ni and/or Cu is present, the Ni content is preferably 0.05 to 0.5% and the Cu content is preferably 0.05 to 0.5%.

One or Two Selected from Between 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM

Both of Ca and REM are elements that contribute to controlling the form of sulfides. They act to spheroidize the form of sulfides and suppress the adverse effects of sulfides on the formability, particularly, stretch flangeability. It is desired that 0.001% or more of Ca and 0.001% or more of REM be present. However, when more than 0.005% of Ca and more than 0.005% of REM are present, the amount of inclusions increases and surface defects and internal defects occur frequently. Therefore, when Ca and/or REM is present, the Ca content is preferably 0.001 to 0.005%, and the REM content is preferably 0.001 to 0.005%.

The balance other than the above-described components includes Fe and inevitable impurities.

The steel sheet has the above-described composition and has a microstructure composed of a ferrite phase as a main phase and a second phase including at least pearlite.

In the steel sheet, the area fraction of the ferrite phase, i.e., the main phase, with respect to the entire microstructure is 75 to 90%. When the area fraction of the ferrite phase is lower than 75%, desired elongation and a desired hole expanding ratio cannot be obtained and formability is lowered. On the other hand, when the area fraction of the ferrite phase exceeds 90%, the area fraction of the second phase is lowered and a desired high strength cannot be obtained. Therefore, the area fraction of the ferrite phase which is the main phase is limited to the range of 75 to 90% and preferably 80 to 90%.

In the steel sheet, at least pearlite is included in the second phase. The area fraction of the pearlite is 10 to 25% with respect to the entire microstructure. When the area fraction of the pearlite is lower than 10%, a desired hole expanding ratio cannot be obtained and stretch flangeability and formability are lowered. On the other hand, when the area fraction of the pearlite exceeds 25%, the number of interfaces between the ferrite phase and the pearlite increases and voids are easily formed during the forming. Accordingly, stretch flangeability is lowered and formability is lowered.

The pearlite is fine grains having an average grain size of 5 μm or less. When the average grain size of the pearlite is large, that is, exceeding 5 μm, stress concentration occurs at the pearlite grains (interfaces) in forming the steel sheet and microvoids are formed. Accordingly, stretch flangeability is lowered and formability is lowered. Therefore, the average grain size of the pearlite is limited to 5 μm or less and preferably 4.0 μm or less.

The second phase of the microstructure of the steel sheet is a phase that includes at least pearlite and mainly composed of pearlite, area fraction of which is 70% or more of the total area of the second phase. When the area fraction of pearlite is less than 70% with respect to the total area of the second phase, the amount of a hard martensite or bainite phase, or retained γ becomes too large, and thus formability is easily lowered. Therefore, the area fraction of pearlite is limited to 70% or greater and preferably 75 to 100% with respect to the total area of the second phase.

The second phase may include bainite, martensite, retained austenite (retained γ) and the like, in addition to pearlite. However, particularly, since bainite and martensite are hard phases and retained γ is transformed into martensite during the forming, bainite, martensite, and retained austenite lower formability. Therefore, it is desired that the amounts of the bainite, martensite and retained austenite are as small as possible, and the area fraction of these with respect to the entire microstructure is preferably 5% or less in total and more preferably 3% or less in total.

Next, a preferred method for manufacturing the steel sheet will be described.

A steel having the above-described composition is used as a starting material. It is not necessary to particularly limit the method for manufacturing the steel. However, from the point of view of productivity, molten steel having the above-described composition is preferably refined through a general refining method using a steel converter, an electric furnace or the like, and formed into a steel such as a slab through a common casting method such as a continuous casting method. An ingot making-slabbing method, a thin-slab casting method, and the like can also be applied.

A steel having the above-described composition is hot-rolled into a hot rolled sheet. The hot rolling step preferably includes heating the steel at a temperature in the range of 1100° C. to 1280° C., hot rolling the heated steel with a finish hot rolling temperature of 870° C. to 950° C. to form a hot rolled sheet, and, upon completion of the hot rolling, coiling the hot rolled sheet at a coiling temperature of 350° C. to 720° C.

When the heating temperature of the steel is lower than 1100° C., deformation resistance becomes too high, and thus a rolling load becomes excessive and it becomes difficult to perform the hot rolling in some cases. On the other hand, when the heating temperature is higher than 1280° C., the crystal grains become too coarse, and thus a desired fine steel sheet microstructure cannot be easily obtained even when hot rolling is performed. Therefore, the heating temperature for hot rolling is preferably in the range of 1100° C. to 1280° C. and more preferably lower than 1280° C.

When the finish hot rolling temperature is lower than 870° C., ferrite (α) and austenite (γ) are formed during the rolling, and a banded microstructure is easily formed in the steel sheet. This banded microstructure remains even after annealing, and sometimes causes generation of anisotropy in the obtained steel sheet characteristics and lowers the formability. On the other hand, when the finish hot rolling temperature is higher than 950° C., the microstructure of the hot rolled sheet becomes coarse, and thus a desired microstructure cannot be obtained even after annealing in some cases. Therefore, the finish hot rolling temperature is preferably in the range of 870° C. to 950° C.

When the coiling temperature after the hot rolling is lower than 350° C., bainitic ferrite, bainite, martensite and the like are formed and the hot rolled microstructure tends to become hard and nonuniform in grain size. In the subsequent annealing step, the microstructure tends to be nonuniform in grain size due to this hot rolled microstructure, and desired formability cannot be obtained in some cases. On the other hand, when the coiling temperature is high, that is, higher than 720° C., it becomes difficult to ensure uniform mechanical characteristics over the entire steel sheet in the longitudinal direction and in the width direction of the steel sheet. Therefore, the coiling temperature is preferably 350° C. to 720° C. and more preferably 500° C. to 680° C.

The hot rolled sheet obtained through the hot rolling step is pickled according to a common method to remove scales on surfaces of the steel sheet, and then directly subjected to a continuous annealing step that includes an annealing treatment and a subsequent cooling treatment in a continuous annealing line without cold-rolling the hot rolled sheet.

The annealing treatment is a process in which the sheet is held in a first temperature region of an Ac1 transformation point to an Ac3 transformation point for 5 to 400 s.

When the temperature (heating temperature) in the first temperature region of the annealing treatment is lower than the Ac1 transformation point or when the holding time (annealing time) in the first temperature region is shorter than 5 s, carbides in the hot rolled sheet are not sufficiently dissolved and/or a sufficient α-to-γ transformation may not occur or the α-to-γ transformation does not occur at all. Accordingly, a desired multi phase microstructure cannot be obtained by the subsequent cooling treatment, and thus a steel sheet having ductility and stretch flangeability that satisfy desired elongation and a desired hole expanding ratio cannot be obtained. On the other hand, when the heating temperature of the annealing treatment is high, that is, higher than the Ac3 transformation point, coarsening of austenite grains is notably shown, the microstructure formed by the subsequent cooling treatment is coarsened, and the formability is thereby decreased in some cases. In addition, when the holding time (annealing time) in the first temperature region is longer than 400 s, the amount of time for the treatment is increased, the amount of consumed energy is increased, and the manufacturing cost is increased. Therefore, the annealing treatment is limited to a process in which holding is performed for 5 to 400 s in the first temperature region of the Ac1 transformation point to the Ac1 transformation point.

A value calculated using Expression (1) below is used as the Ac1 transformation point of each steel sheet and a value calculated using Expression (2) below is used as the Ac3 transformation point. When a steel sheet does not contain all elements set forth in the expressions, the contents of elements that are not contained are assumed to be zero in performing the calculations.


Ac1 Transformation Point (° C.)=723+29.1Si−10.7Mn−16.9Ni+16.9Cr+6.38W+290As  (1)


Ac3 Transformation Point (° C.)=910−203√C+44.7Si−30Mn+700P+400Al−15.2Ni−11Cr−20Cu+31.5Mo+104V+400Ti+13.1W+120As  (2)

where C, Si, Mn, Ni, Cr, W, As, C, P, Al, Cu, Mo, V, and Ti represents contents (mass %) of the respective elements.

The cooling treatment after the annealing treatment is a process of cooling the sheet at an average cooling rate of 5° C./s or higher from the above-described first temperature region to 700° C. and adjusting the residence time in a second temperature region of 700° C. to 400° C. over 30 to 400 s.

When the average cooling rate from the first temperature region to 700° C. is lower than 5° C./s, the amount of formed ferrite becomes too large. As a result, a desired multi phase microstructure is not obtained, the formability is lowered, and a desired tensile strength (540 MPa or higher) cannot be ensured in some cases. Therefore, the average cooling rate from the first temperature region to 700° C. is limited to 5° C./s or higher, preferably 20° C./s or lower, and more preferably 5 to 15° C./s.

The residence time in the second temperature region of 700° C. to 400° C. is an important factor for the formation of pearlite included in the second phase. The “residence time” means the length of time the sheet remains in the above-described second temperature region. This covers the case where the sheet is held at a specific temperature in the second temperature region, a case where the sheet is cooled in the second temperature region at a specific cooling rate, and a case where the sheet is cooled by the combination of the two cases. When the residence time in the second temperature region is shorter than 30 s, pearlite transformation does not occur or the amount of formed pearlite is insufficient, and thus a desired multi phase microstructure cannot be obtained. On the other hand, when the residence time in the second temperature region is long, that is, longer than 400 s, productivity is lowered. Therefore, the residence time in the second temperature region is 30 to 400 s and preferably 150 s or shorter. To secure a desired amount of pearlite, the cooling time in a temperature region of 700° C. to 550° C. in the second temperature region is preferably 10 s or longer, that is, the average cooling rate in the temperature region of 700° C. to 550° C. is preferably 15° C./s or lower. When the cooling time in the temperature region of 700° C. to 550° C. is shorter than 10 s, pearlite is not sufficiently formed, a desired multi phase microstructure is not obtained, and desired formability cannot be obtained in some cases.

Hereinafter, our steel sheets and methods will be more specifically described on the basis of the Examples. The disclosure is, however, not limited to these Examples.

EXAMPLES

Molten steels each having a composition shown in Table 1 were refined and formed into steels by a common method. The steels were hot-rolled at the heating temperatures and the finish hot rolling temperatures shown in Table 2 to form 1.6 mm-thick hot rolled sheets. Upon completion of the hot rolling, the hot rolled sheets were coiled at the coiling temperatures shown in Table 2. Thereafter, pickling was performed. Some of the hot rolled sheets (sheet thickness: 3.2 mm) were subjected to pickling and then to cold rolling with a rolling reduction of 50% to form 1.6 mm-thick cold rolled sheets, which were used as the Comparative Examples.

Under the conditions shown in Table 2, the obtained hot rolled sheets or cold rolled sheets were further subjected to a continuous annealing step that includes an annealing treatment of heating the sheets to a temperature in the first temperature region and holding the sheets thereat, and a cooling treatment of cooling the sheets at average cooling rates shown in Table 2 from the temperature in the first temperature region to 700° C., cooling the sheets at average cooling rates (cooling time) shown in Table 2 from 700° C. to 550° C. in a second temperature region, and adjusting the residence time in the second temperature region of 700° C. to 400° C. to the time shown in Table 2 so as to form annealed sheets. The transformation points of the respective steel sheets shown in Table 2 are values calculated using the above-described Expressions (1) and (2).

Test pieces were taken from the obtained annealed sheets, and observation of microstructure, a tensile test, and a hole expanding test were performed thereon. The test methods were as follows.

(1) Observation of Microstructure

A test piece for observation of microstructure was taken from an obtained annealed sheet. A cross-section (L cross-section) parallel to the rolling direction was polished and corroded with a nital solution, and the microstructure was observed in three or more fields by using a scanning electron microscope (magnification: 3000) and photographed to determine the type of the microstructure and an area fraction of each phase with respect to the entire microstructure. Moreover, an area fraction of the total area of the second phase with respect to the entire microstructure was calculated. The average crystal grain size of the pearlite included in the second phase was also calculated. Regarding the average crystal grain size of the pearlite, an area of each pearlite grain was measured, an equivalent circle diameter was calculated from the area, the obtained equivalent circle diameters of the grains were arithmetically averaged, and the arithmetically averaged value was used as the average crystal grain size of the pearlite grains. The number of the measured pearlite grains was equal to or more than 20. An area fraction of the pearlite with respect to the total area of the second phase was also calculated.

(2) Tensile Test

A JIS No. 5 test piece was taken from an obtained annealed sheet so that the tensile direction is coincident with a direction perpendicular to the rolling direction. The tensile test was performed on the basis of the provisions of JIS Z 2241 and tensile characteristics (yield point YP, tensile strength TS, and elongation E1) were determined.

(3) Hole Expanding Test

A 100 mm-square test piece for a hole expanding test was taken from an obtained annealed sheet. The hole expanding test was performed on the basis of the Japan Iron and Steel Federation Standard JFST 1001-1996, and a hole expanding ratio λ (%) was determined.

The obtained results are shown in Table 3.

TABLE 1 Chemical Composition (mass %) Steel No. C Si Mn P S Al N Cr, V, Mo Ti, Nb B Ni, Cu Ca, REM Remarks A 0.10 0.72 1.22 0.014 0.002 0.043 0.0038 — — — — — Example B 0.09 0.72 1.21 0.023 0.002 0.035 0.0028 Cr: 0.35 — — — — Example C 0.09 1.30 0.80 0.013 0.002 0.036 0.0029 — — — — — Example D 0.14 1.02 1.32 0.015 0.001 0.035 0.0025 — — — — — Example E 0.09 1.55 1.24 0.013 0.001 0.043 0.0034 — — — — — Comparative Example F 0.15 0.23 1.40 0.015 0.002 0.042 0.0041 — — — — — Comparative Example G 0.14 1.02 1.62 0.012 0.001 0.039 0.0034 — — — — — Comparative Example H 0.14 0.71 2.00 0.012 0.001 0.041 0.0036 — — — — — Comparative Example I 0.15 0.80 0.75 0.010 0.001 0.038 0.0029 Mo: 0.1 — — — — Example J 0.10 1.02 1.23 0.015 0.001 0.037 0.0039 V: 0.1 — — — — Example K 0.13 1.00 0.82 0.015 0.002 0.037 0.0040 — Nb: 0.031 — — — Example L 0.14 1.41 0.88 0.013 0.001 0.040 0.0028 — Ti: 0.022 B: 0.0012 — — Example M 0.14 1.02 0.84 0.015 0.001 0.037 0.0039 — — — Ni: 0.2 — Example N 0.14 0.81 0.82 0.015 0.002 0.042 0.0034 — — — — — Example O 0.12 1.43 0.91 0.011 0.002 0.042 0.0041 — — — — REM: 0.002 Example P 0.07 1.00 1.20 0.014 0.003 0.038 0.0035 — — — — — Comparative Example Q 0.16 1.01 1.24 0.013 0.002 0.043 0.0035 — — — — — Comparative Example R 0.13 0.40 1.32 0.012 0.001 0.036 0.0038 — — — — — Comparative Example

TABLE 2 Continuous annealing step Hot rolling step Annealing Finish treatment Cooling treatment hot Coil- First temperature Average Heating rolling ing Cold region cooling Trans- tem- tem- tem- roll- Sheet Tem- rate up Second temperature region formation Steel pera- pera- pera- ing Thick- pera- to Cooling Cooling Residence point Sheet Steel ture ture ture Yes/ ness ture Holding 700° c.* time** rate** time*** Ac1 Ac3 no. no. (° c.) (° c.) (° c.) No (mm) (° C.) time (s) (° c./s) (s) (° C./s) (s) (° C.) (° C.) Remarks 1 A 1200 900 600 No 1.6 800 100 10 15 10  35 731 868 Example 2 B 1200 900 600 No 1.6 820 100 15 15 10 150 737 873 Example 3 B 1200 900 600 No 1.6 850 100 15 15 10 300 737 873 Example 4 B 1200 900 600 Yes 1.6 820 100 15 15 10 150 737 873 Comparative Example 5 C 1200 920 560 No 1.6 820 100 10 15 10  35 752 907 Example 6 C 1200 920 560 No 1.6 800 100 10 15 10  20 752 907 Comparative Example 7 D 1200 900 560 No 1.6 820 100 10 15 10 150 739 865 Example 8 D 1200 900 560 Yes 1.6 820 100 10 15 10 150 739 865 Comparative Example 9 E 1200 920 560 No 1.6 850 100  3 15 10 150 755 908 Comparative Example 10 F 1200 900 600 No 1.6 800 100 10 15 10  35 715 827 Comparative Example 11 F 1200 900 600 Yes 1.6 800 100 10 15 10 150 715 827 Comparative Example 12 G 1200 900 600 No 1.6 800 100 10 20  8  70 735 855 Comparative Example 13 H 1200 900 600 No 1.6 800 100 10 20  8  70 722 831 Comparative Example 14 I 1200 900 640 No 1.6 820 100 10 15 10 150 738 871 Example 15 I 1200 900 640 No 1.6 820 100 10   0.2 750     0.4 738 871 Comparative Example 16 J 1250 900 600 No 1.6 820 100 15 15 10 150 740 890 Example 17 K 1250 900 560 No 1.6 820 150 10 30  5  35 743 882 Example 18 K 1250 900 560 No 1.6 820 100 10  5 30  20 743 882 Comparative Example 19 K 1250 900 560 No 1.6 715 100 10 30  5 150 743 882 Comparative Example 20 L 1250 920 600 No 1.6 820 150 10 15 10  35 755 905 Example 21 L 1250 920 600 No 1.6 820 150 10 15 10 150 755 905 Example 22 L 1250 920 600 No 1.6 820 150 10 15 10 300 755 905 Example 23 L 1250 920 600 No 1.6 920 150 10 15 10 150 755 905 Comparative Example 24 M 1200 920 600 No 1.6 840 100 15 20  8  70 740 877 Example 25 M 1200 900 600 No 1.6 840 100 15 15 10 300 740 877 Example 26 N 1200 900 640 No 1.6 820 100 15 15 10 150 738 873 Example 27 O 1200 920 600 No 1.6 820 100 15 15 10 150 755 901 Example 28 P 1200 920 600 No 1.7 800 100 10 15 10  35 739 890 Comparative Example 29 P 1200 920 600 No 1.6 800 100 10  5 30  10 739 890 Comparative Example 30 Q 1200 900 640 No 1.6 820 100 10 15 10  70 739 863 Comparative Example 31 R 1200 900 640 No 1.6 820 100 10 15 10  70 721 838 Comparative Example *Cooling rate (average) from temperature in first temperature region to 700° C. **Cooling time and average cooling rate (° C./s) between 700° C. and 550° C. ***Residence time between 700° C. and 400° C.

TABLE 3 Microstructure P Tensile characteristics Fraction of Microstructure (area fraction) average Yield Tensile Stretch- Steel Main Second Phase grain strength strength Elongation flange- sheet Steel phase P size YS TS El ability no. no. Type F* P* M* Ret. γ* Total fraction** (μm) (MPa) (MPa) (%) λ (%) Remarks 1 A F + P + M 87.7 12.1 0.2 — 12.3 98 3.1 436 545 33.4 127 Example 2 B F + P + M 84.5 15.3 0.2 — 15.5 99 3.0 442 552 32.1 116 Example 3 B F + P + M 84.1 15.8 0.1 — 15.9 99 3.2 450 558 31.6 121 Example 4 B F + P + M 87.9  3.2 8.9 — 12.1 26 2.8 458 599 31.8 67 Comparative Example 5 C F + P + M + Ret. γ 83.6 15.2 0.1 1.1 16.4 93 3.2 453 553 34.5 132 Example 6 C F + M 84.8  0.0 15.2 — 15.2  0 — 488 724 28.2 65 Comparative Example 7 D F + P + M 82.6 13.2 4.2 — 17.4 76 3.4 452 624 31.7 108 Example 8 D F + P + M 82.2 10.6 7.2 — 17.8 60 3.2 448 625 32.4 69 Comparative Example 9 E F + P + M 86.7  1.2 12.1 — 13.3  9 2.1 387 624 33.6 58 Comparative Example 10 F F + P + M 88.8 10.6 0.6 — 11.2 95 2.9 425 518 34.2 85 Comparative Example 11 F F + P + M 88.2 10.9 0.9 — 11.8 92 3.2 429 526 33.7 83 Comparative Example 12 G F + M 85.9  0.0 14.1 — 14.1  0 — 425 638 30.2 56 Comparative Example 13 H F + M 83.6  0.0 16.4 — 16.4  0 — 463 672 28.9 51 Comparative Example 14 I F + P + M 85.0 14.8 0.2 — 15.0 99 2.9 449 564 33.1 114 Example 15 I F + M 54.8  0.0 45.2 — 45.2  0 — 653 1082 13.8 32 Comparative Example 16 J F + P + M 83.6 16.2 0.2 — 16.4 99 3.7 462 592 32.4 118 Example 17 K F + P + M 83.4 15.1 1.5 — 16.6 91 3.4 472 611 31.8 111 Example 18 K F + M 84.8  0.0 15.2 — 15.2  0 — 502 752 24.7 54 Comparative Example 19 K F + C 90.4  0.0 — — 0.0  0 — 402 535 31.1 79 Comparative Example 20 L F + P + M + 85.9 11.5 1.4 1.2 14.1 82 3.8 453 603 32.5 105 Example Retained γ 21 L F + P + M + Ret. γ 85.3 12.4 1.2 1.1 12.3 84 3.9 468 601 32.6 111 Example 22 L F + P + M 86.5 12.2 1.3 — 13.5 90 3.9 462 597 32.4 104 Example 23 L F + P + M 83.0 15.8 1.2 — 17.0 93 6.5 488 625 31.1 75 Comparative Example 24 M F + P + M 87.3 12.5 0.2 — 12.7 98 3.2 482 614 31.5 117 Example 25 M F + P 86.1 13.9 — — 13.9 100  3.4 469 586 32.0 120 Example 26 N F + P + M 86.8 13.1 0.1 — 13.2 99 3.0 442 564 32.9 118 Example 27 O F + P + M + Ret. γ 86.0 12.4 0.2 1.4 14.0 89 3.2 467 608 31.8 109 Example 28 P F + P + M 91.7  8.2 0.1 — 8.3 99 3.1 385 482 35.1 105 Comparative Example 29 P F + P + M 87.8  5.8 6.4 — 12.2 48 2.2 425 572 32.4 62 Comparative Example 30 Q F + P + M 90.9  4.9 4.2 — 9.1 54 2.2 524 694 26.8 69 Comparative Example 31 R F + P 89.7 10.3 — — 10.3 100  3.1 409 528 32.9 110 Comparative Example *F: ferrite, P: pearlite, M: martensite, Retained γ (Ret. γ): retained austenite, C: cementite **P fraction: P area/second phase total area ratio

In all of our Examples, high strength steel sheets having excellent formability with a high strength, i.e., a tensile strength TS of 540 MPa or higher, high ductility, i.e., elongation E1 of 30% or greater, and excellent stretch flangeability, i.e., hole expanding ratio λ of 80% or higher are obtained. In contrast, in the Comparative Examples outside our range, a desired high strength is not obtained, desired elongation is not obtained, or a desired hole expanding ratio λ is not obtained, and thus formability is lowered.

Claims

1. A high strength steel sheet having excellent formability which has a composition comprising, by mass %: and the balance Fe with inevitable impurities, the steel sheet having a microstructure composed of a ferrite phase which is a main phase and a second phase including at least pearlite, wherein an area fraction of the ferrite phase is 75 to 90% and an area fraction of the pearlite is 10 to 25% with respect to the entire microstructure, an average grain size of the pearlite is 5 μm or smaller, and an area fraction of the pearlite is 70% or greater with respect to the total area of the second phase.

0.08 to 0.15% of C, 0.5 to 1.5% of Si,
0.5 to 1.5% of Mn, 0.1% or less of P,
0.01% or less of S, 0.01 to 0.1% of Al,
0.005% or less of N,

2. The high strength steel sheet according to claim 1, wherein the composition further comprises one or more selected from the group consisting of 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.

3. The high strength steel sheet according to claim 1, wherein the composition further comprises one or two selected from the group consisting of 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb by mass %.

4. The high strength steel sheet according to claim 1, wherein the composition further comprises 0.0003 to 0.0050% of B by mass %.

5. The high strength steel sheet according to claim 1, wherein the composition further comprises one or both of 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu by mass %.

6. The high strength steel sheet according to claim 1, wherein the composition further comprises one or both of 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM by mass %.

7. A method for manufacturing a high strength steel sheet having excellent formability, comprising: and the balance Fe with inevitable impurities to form a hot rolled sheet;

a hot rolling step of hot-rolling a steel having a composition which comprises, by mass %,
0.08 to 0.15% of C, 0.5 to 1.5% of Si,
0.5 to 1.5% of Mn, 0.1% or less of P,
0.01% or less of S, 0.01 to 0.1% of Al,
0.005% or less of N,
a continuous annealing step including an annealing treatment of pickling the hot rolled sheet, and holding the pickled hot rolled sheet in a first temperature region of an Ac1 transformation point to an Ac3 transformation point for 5 to 400 s; and
a cooling treatment of cooling the sheet at an average cooling rate of 5° C./s or higher from a first temperature region to 700° C. after the annealing treatment and adjusting a residence time in a second temperature region of 700° C. to 400° C. in the range of 30 to 400 s by using a continuous annealing line.

8. The method according to claim 7, wherein the hot rolling step includes heating the steel at a temperature of 1100° C. to 1280° C., hot-rolling the heated steel with a finish hot rolling temperature of 870° C. to 950° C. to form a hot rolled sheet, and coiling the hot rolled sheet at a coiling temperature of 350° C. to 720° C. upon completion of the hot rolling.

9. The method according to claim 7, wherein a cooling time in a temperature region of 700° C. to 550° C. in the second temperature region is 10 s or longer.

10. The method according to claim 7, wherein the composition further comprises one or more selected from the group consisting of 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.

11. The method according to claim 7, wherein the composition further one or two selected from the group consisting of 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb by mass %.

12. The method according to claim 7, wherein the composition further comprises 0.0003 to 0.0050% of B by mass %.

13. The method according to claim 7, wherein the composition further comprises one or two selected from the group consisting of 0.05 to 0.5% of Ni and 0.05 to 0.5% of Cu by mass %.

14. The method according to claim 7, wherein the composition further comprises one or two selected from the group consisting of 0.001 to 0.005% of Ca and 0.001 to 0.005% of REM by mass %.

15. The high strength steel sheet according to claim 2, wherein the composition further comprises one or two selected from the group consisting of 0.01 to 0.1% of Ti and 0.01 to 0.1% of Nb by mass %.

16. The high strength steel sheet according to claim 2, wherein the composition further comprises 0.0003 to 0.0050% of B by mass %.

17. The high strength steel sheet according to claim 3, wherein the composition further comprises 0.0003 to 0.0050% of B by mass %.

18. The method for manufacturing a high strength steel sheet according to claim 8, wherein a cooling time in a temperature region of 700° C. to 550° C. in the second temperature region is 10 s or longer.

19. The method for manufacturing a high strength steel sheet according to claim 8, wherein the composition further comprises one or more selected from the group consisting of 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.

20. The method for manufacturing a high strength steel sheet according to claim 9, wherein the composition further comprises one or more selected from the group consisting of 0.05 to 0.5% of Cr, 0.005 to 0.2% of V, and 0.005 to 0.2% of Mo by mass %.

Patent History
Publication number: 20130233453
Type: Application
Filed: Jun 29, 2011
Publication Date: Sep 12, 2013
Applicant: JFE STEEL CORPORATION (Tokyo)
Inventors: Kenji Kawamura (Chiyoda-ku), Hidetaka Kawabe (Chiyoda-ku), Kazuhiro Seto (Chiyoda-ku), Noriyuki Katayama (Chiyoda-ku)
Application Number: 13/704,781
Classifications
Current U.S. Class: With Working (148/602); With Flattening, Straightening, Or Tensioning By External Force (148/645); Ferrous (i.e., Iron Base) (148/320); Chromium Containing, But Less Than 9 Percent (148/333); Three Percent Or More Manganese Containing Or Containing Other Transition Metal In Any Amount (148/337); Beryllium Or Boron Containing (148/330); Copper Containing (148/332); Nickel Containing (148/336); Rare Earth Meal Containing (148/331)
International Classification: C21D 8/02 (20060101); C22C 38/14 (20060101); C22C 38/12 (20060101); C22C 38/00 (20060101); C22C 38/06 (20060101); C22C 38/04 (20060101); C22C 38/02 (20060101); C22C 38/18 (20060101); C22C 38/08 (20060101);