Patent Application
20190233926
The present invention relates to a steel sheet and a plated steel sheet.
Recently, the reduction in weight of various members aiming at the improvement of fuel efficiency of automobiles has been demanded. In response to this demand, thinning achieved by an increase in strength of a steel sheet to be used for various members and application of light metal such as an Al alloy to various members have been in progress. The light metal such as an Al alloy is high in specific strength as compared to heavy metal such as steel. However, the light metal is significantly expensive as compared to the heavy metal. Therefore, the application of light metal such as an Al alloy is limited to special uses. Thus, the thinning achieved by an increase in strength of a steel sheet has been demanded in order to apply the reduction in weight of various members to a more inexpensive and broader range.
The steel sheet to be used for various members of automobiles is required to have not only strength but also material properties such as ductility, stretch-flanging workability, burring workability, fatigue endurance, impact resistance, and corrosion resistance according to the use of a member. However, when the steel sheet is increased in strength, material properties such as formability (workability) deteriorate generally. Therefore, in the development of a high-strength steel sheet, it is important to achieve both these material properties and the strength.
Concretely, when the steel sheet is used to manufacture a part having a complex shape, for example, the following workings are performed. The steel sheet is subjected to shearing or punching, and is subjected to blanking or hole making, and then is subjected to press forming based on stretch-flanging and burring mainly or bulging. The steel sheet to be subjected to such workings is required to have good stretch flangeability and ductility.
Further, in order to prevent deformation caused when collision of an automotive part occurs, it is necessary to use a steel sheet having a high yield stress as a material of the part. However, as the steel sheet has a higher yield stress, the steel sheet tends to be poor in ductility. Accordingly, the steel sheet to be used for various members of automobiles is also required to have both the yield stress and the ductility.
In Patent Reference 1, there is described a high-strength hot-rolled steel sheet excellent in ductility, stretch flangeability, and material uniformity that has a steel microstructure having 95% or more of a ferrite phase by area ratio and in which an average particle diameter of Ti carbides precipitated in steel is 10 nm or less. However, in the case where a strength of 480 MPa or more is secured in the steel sheet disclosed in Patent Reference 1, which has 95% or more of a soft ferrite phase, it is impossible to obtain sufficient ductility.
Patent Reference 2 discloses a high-strength hot-rolled steel sheet excellent in stretch flangeability and fatigue property that contains Ce oxides, La oxides, Ti oxides, and Al2O3 inclusions. Further, Patent Reference 2 describes a high-strength hot-rolled steel sheet in which an area ratio of a bainitic⋅ferrite phase is 80 to 100%. Patent Reference 3 discloses a high-strength hot-rolled steel sheet having reduced strength variation and having excellent ductility and hole expandability in which the total area ratio of a ferrite phase and a bainite phase and the absolute value of a difference in Vickers hardness between a ferrite phase and a second phase are defined.
Further, there is a compound structure steel sheet in which a hard phase such as bainite or martensite and a soft phase such as ferrite excellent in ductility are combined conventionally. Such a steel sheet is called a dual phase (Dual Phase) steel sheet. The dual phase steel sheet is good in uniform elongation in response to strength and is excellent in the strength-ductility-balance. For example, Patent Reference 4 describes a high-strength hot-rolled steel sheet having good stretch flangeability and impact property that has a structure composed of polygonal ferrite+upper bainite. Further, Patent Reference 5 describes a high-strength steel sheet that has a structure composed of three phases of polygonal ferrite, bainite, and martensite, is low in yield ratio, and is excellent in the strength-elongation-balance and stretch flangeability.
When a conventional high-strength steel sheet is formed by pressing in cold working, cracking sometimes occurs from an edge of a portion to be subjected to stretch flange forming during forming. This is conceivable because work hardening advances only in the edge portion due to the strain introduced into a punched end face at the time of blanking.
As an evaluation method of a stretch flangeability test of the steel sheet, a hole expansion test has been used. However, in the hole expansion test, a test piece leads to a fracture in a state where a strain distribution in a circumferential direction little exists. In contrast to this, when the steel sheet is worked into a part shape actually, a strain distribution exists. The strain distribution affects a fracture limit of the part. Thereby, it is estimated that even in a high-strength steel sheet that exhibits sufficient stretch flangeability in the hole expansion test, performing cold pressing sometimes causes cracking.
Patent References 1 to 5 disclose a technique to improve material properties by defining structures. However, it is unclear whether sufficient stretch flangeability can be secured even in the case where the strain distribution is considered in the steel sheets described in Patent References 1 to 5.
Patent Reference 1: International Publication Pamphlet No. WO2013/161090
Patent Reference 2: Japanese Laid-open Patent Publication No. 2005-256115
Patent Reference 3: Japanese Laid-open Patent Publication No. 2011-140671
Patent Reference 4: Japanese Laid-open Patent Publication No. 58-42726
Patent Reference 5: Japanese Laid-open Patent Publication No. 57-70257
An object of the present invention is to provide a steel sheet and a plated steel sheet that are high in strength, have good ductility and stretch flangeability, and have a high yield stress.
According to the conventional findings, the improvement of the stretch flangeability (hole expansibility) in the high-strength steel sheet has been performed by inclusion control, homogenization of structure, unification of structure, and/or reduction in hardness difference between structures, as described in Patent References 1 to 3. In other words, conventionally, the improvement in the stretch flangeability has been achieved by controlling the structure to be observed by an optical microscope.
However, it is difficult to improve the stretch flangeability under the presence of the strain distribution even when only the structure to be observed by an optical microscope is controlled. Thus, the present inventors made an intensive study by focusing on an intragranular misorientation of each crystal grain. As a result, they found out that it is possible to greatly improve the stretch flangeability by controlling the proportion of crystal grains each having a misorientation in a crystal grain of 5 to 14° to all crystal grains to 20 to 100%.
Further, the present inventors found out that the structure of the steel sheet is composed to contain two types of crystal grains that are different in precipitation state (number density and size) of precipitates in a crystal grain, thereby making it possible to fabricate a steel sheet excellent in the strength-ductility-balance. This effect is estimated to be due to the fact that the structure of the steel sheet is composed so as to contain crystal grains with relatively small hardness and crystal grains with large hardness, to thereby obtain such a function as a Dual Phase practically without existence of martensite.
The present invention was completed as a result that the present inventors conducted intensive studies repeatedly based on the new findings relating to the above-described proportion of the crystal grains each having a misorientation in a crystal grain of 5 to 14° to all the crystal grains and the new findings obtained by the structure of the steel sheet being composed to contain two types of crystal grains that are different in number density and size of precipitates in a crystal grain.
The gist of the present invention is as follows.
(1)
A steel sheet, includes:
a chemical composition represented by, in massa,
C: 0.008 to 0.150%,
Si: 0.01 to 1.70%,
Mn: 0.60 to 2.50%,
Al: 0.010 to 0.60%,
Ti: 0 to 0.200%,
Nb: 0 to 0.200%,
Ti+Nb: 0.015 to 0.200%,
Cr: 0 to 1.0%,
B: 0 to 0.10%,
Mo: 0 to 1.0%,
Cu: 0 to 2.0%,
Ni: 0 to 2.0%,
Mg: 0 to 0.05%,
REM: 0 to 0.05%,
Ca: 0 to 0.05%,
Zr: 0 to 0.05%,
P: 0.05% or less,
S: 0.0200% or less,
N: 0.0060% or less, and
balance: Fe and impurities; and
a structure represented by, by area ratio,
ferrite: 5 to 95%, and
bainite: 5 to 95%, in which
when a region that is surrounded by a grain boundary having a misorientation of 15° or more and has a circle-equivalent diameter of 0.3 μm or more is defined as a crystal grain, the proportion of crystal grains each having an intragranular misorientation of 5 to 14° to all crystal grains is 20 to 100% by area ratio, and
hard crystal grains A in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1016 to 1×1019 pieces/cm3 and soft crystal grains B in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1015 pieces/cm3 or less are contained, and the volume % of the hard crystal grains A/(the volume % of the hard crystal grains A+the volume % of the soft crystal grains B) is 0.1 to 0.9.
(2)
The steel sheet according to (1), in which
a tensile strength is 480 MPa or more,
the product of the tensile strength and a limit form height in a saddle-type stretch-flange test is 19500 mm·MPa or more, and
the product of a yield stress and ductility is 10000 MPa·% or more.
(3)
The steel sheet according to (1) or (2), in which
the chemical composition contains, in massa, one type or more selected from the group consisting of
Cr: 0.05 to 1.0%, and
B: 0.0005 to 0.10%.
(4)
The steel sheet according to any one of (1) to (3), in which
the chemical composition contains, in mass %, one type or more selected from the group consisting of
Mo: 0.01 to 1.0%,
Cu: 0.01 to 2.0%, and
Ni: 0.01% to 2.0%.
(5)
The steel sheet according to any one of (1) to (4), in which
the chemical composition contains, in mass %, one type or more selected from the group consisting of
Ca: 0.0001 to 0.05%,
Mg: 0.0001 to 0.05%,
Zr: 0.0001 to 0.05%, and
REM: 0.0001 to 0.05%.
(6)
A plated steel sheet, in which
a plating layer is formed on a surface of the steel sheet according to any one of (1) to (5).
(7)
The plated steel sheet according to (6), in which
the plating layer is a hot-dip galvanizing layer.
(8)
The plated steel sheet according to (6), in which
the plating layer is an alloyed hot-dip galvanizing layer.
According to the present invention, it is possible to provide a steel sheet that is high in strength, has good ductility and stretch flangeability, and has a high yield stress. The steel sheet of the present invention is applicable to a member required to have strict ductility and stretch flangeability while having high strength.
Hereinafter, there will be explained embodiments of the present invention.
[Chemical Composition]
First, there will be explained a chemical composition of a steel sheet according to the embodiment of the present invention. In the following explanation, “%” that is a unit of the content of each element contained in the steel sheet means “mass %” unless otherwise stated. The steel sheet according to this embodiment has a chemical composition represented by C: 0.008 to 0.150%, Si: 0.01 to 1.70%, Mn: 0.60 to 2.50%, Al: 0.010 to 0.60%, Ti: 0 to 0.200%, Nb: 0 to 0.200%, Ti +Nb: 0.015 to 0.200%, Cr: 0 to 1.0%, B: 0 to 0.10%, Mo: 0 to 1.0%, Cu: 0 to 2.0%, Ni: 0 to 2.0%, Mg: 0 to 0.05%, rare earth metal (REM): 0 to 0.05%, Ca: 0 to 0.05%, Zr: 0 to 0.05%, P: 0.05% or less, S: 0.0200% or less, N: 0.0060% or less, and balance: Fe and impurities. Examples of the impurities include one contained in raw materials such as ore and scrap, and one contained during a manufacturing process.
“C: 0.008 to 0.150%”
C bonds to Nb, Ti, and so on to form precipitates in the steel sheet and contributes to an improvement in strength of steel by precipitation strengthening. When the C content is less than 0.008%, it is impossible to sufficiently obtain this effect. Therefore, the C content is set to 0.008% or more.
The C content is preferably set to 0.010% or more and more preferably set to 0.018% or more. On the other hand, when the C content is greater than 0.150%, an orientation spread in bainite is likely to increase and the proportion of crystal grains each having an intragranular misorientation of 5 to 14° becomes short. Further, when the C content is greater than 0.150%, cementite harmful to the stretch flangeability increases and the stretch flangeability deteriorates. Therefore, the C content is set to 0.150% or less. The C content is preferably set to 0.100% or less and more preferably set to 0.090% or less.
“Si: 0.01 to 1.70%”
Si functions as a deoxidizer for molten steel. When the Si content is less than 0.01%, it is impossible to sufficiently obtain this effect. Therefore, the Si content is set to 0.01% or more. The Si content is preferably set to 0.02% or more and more preferably set to 0.03% or more. On the other hand, when the Si content is greater than 1.70%, the stretch flangeability deteriorates or surface flaws occur. Further, when the Si content is greater than 1.70%, the transformation point rises too much, to then require an increase in rolling temperature. In this case, recrystallization during hot rolling is promoted significantly and the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes short. Further, when the Si content is greater than 1.70%, surface flaws are likely to occur when a plating layer is formed on the surface of the steel sheet. Therefore, the Si content is set to 1.70% or less. The Si content is preferably set to 1.60% or less, more preferably set to 1.50% or less, and further preferably set to 1.40% or less.
“Mn: 0.60 to 2.50%”
Mn contributes to the strength improvement of the steel by solid-solution strengthening or improving hardenability of the steel. When the Mn content is less than 0.60%, it is impossible to sufficiently obtain this effect. Therefore, the Mn content is set to 0.60% or more. The Mn content is preferably set to 0.70% or more and more preferably set to 0.80% or more. On the other hand, when the Mn content is greater than 2.50%, the hardenability becomes excessive and the degree of orientation spread in bainite increases. As a result, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes short and the stretch flangeability deteriorates. Therefore, the Mn content is set to 2.50% or less. The Mn content is preferably set to 2.30% or less and more preferably set to 2.10% or less.
“Al: 0.010 to 0.60%”
Al is effective as a deoxidizer for molten steel. When the Al content is less than 0.010%, it is impossible to sufficiently obtain this effect. Therefore, the Al content is set to 0.010% or more. The Al content is preferably set to 0.020% or more and more preferably set to 0.030% or more. On the other hand, when the Al content is greater than 0.60%, weldability, toughness, and so on deteriorate. Therefore, the Al content is set to 0.60% or less. The Al content is preferably set to 0.50% or less and more preferably set to 0.40% or less.
“Ti: 0 to 0.200%, Nb: 0 to 0.200%, Ti +Nb: 0.015 to 0.200%”
Ti and Nb finely precipitate in the steel as carbides (TiC, NbC) and improve the strength of the steel by precipitation strengthening. Further, Ti and Nb form carbides to thereby fix C, resulting in that generation of cementite harmful to the stretch flangeability is suppressed. Further, Ti and Nb can significantly improve the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° and improve the stretch flangeability while improving the strength of the steel. When the total content of Ti and Nb is less than 0.015%, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes short and the stretch flangeability deteriorates. Therefore, the total content of Ti and Nb is set to 0.015% or more. The total content of Ti and Nb is preferably set to 0.018% or more. Further, the Ti content is preferably set to 0.015% or more, more preferably set to 0.020% or more, and further preferably set to 0.025% or more. Further, the Nb content is preferably set to 0.015% or more, more preferably set to 0.020% or more, and further preferably set to 0.025% or more. On the other hand, when the total content of Ti and Nb is greater than 0.200%, the ductility and the workability deteriorate and the frequency of cracking during rolling increases. Therefore, the total content of Ti and Nb is set to 0.200% or less. The total content of Ti and Nb is preferably set to 0.150% or less. Further, when the Ti content is greater than 0.200%, the ductility deteriorates. Therefore, the Ti content is set to 0.200% or less. The Ti content is preferably set to 0.180% or less and more preferably set to 0.160% or less. Further, when the Nb content is greater than 0.200%, the ductility deteriorates. Therefore, the Nb content is set to 0.200% or less. The Nb content is preferably set to 0.180% or less and more preferably set to 0.160% or less.
“P: 0.05% or less”
P is an impurity. P deteriorates toughness, ductility, weldability, and so on, and thus a lower P content is more preferable. When the P content is greater than 0.05%, the deterioration in stretch flangeability is prominent. Therefore, the P content is set to 0.05% or less. The P content is preferably set to 0.03% or less and more preferably set to 0.02% or less. The lower limit of the P content is not determined in particular, but its excessive reduction is not desirable from the viewpoint of manufacturing cost. Therefore, the P content may be set to 0.005% or more.
“S: 0.0200% or less”
S is an impurity. S causes cracking at the time of hot rolling, and further forms A-based inclusions that deteriorate the stretch flangeability. Thus, a lower S content is more preferable. When the S content is greater than 0.0200%, the deterioration in stretch flangeability is prominent. Therefore, the S content is set to 0.0200% or less. The S content is preferably set to 0.0150% or less and more preferably set to 0.0060% or less. The lower limit of the S content is not determined in particular, but its excessive reduction is not desirable from the viewpoint of manufacturing cost. Therefore, the S content may be set to 0.0010% or more.
“N: 0.0060% or less”
N is an impurity. N forms precipitates with Ti and Nb preferentially over C and reduces Ti and Nb effective for fixation of C. Thus, a lower N content is more preferable. When the N content is greater than 0.0060%, the deterioration in stretch flangeability is prominent. Therefore, the N content is set to 0.0060% or less. The N content is preferably set to 0.0050% or less. The lower limit of the N content is not determined in particular, but its excessive reduction is not desirable from the viewpoint of manufacturing cost. Therefore, the N content may be set to 0.0010% or more.
Cr, B, Mo, Cu, Ni, Mg, REM, Ca, and Zr are not essential elements, but are arbitrary elements that may be contained as needed in the steel sheet up to predetermined amounts.
“Cr: 0 to 1.0%”
Cr contributes to the strength improvement of the steel. Desired purposes are achieved without Cr being contained, but in order to sufficiently obtain this effect, the Cr content is preferably set to 0.05% or more. On the other hand, when the Cr content is greater than 1.0%, the above-described effect is saturated and economic efficiency decreases. Therefore, the Cr content is set to 1.0% or less.
“B: 0 to 0.10%”
B increases the hardenability and increases a structural fraction of a low-temperature transformation generating phase being a hard phase. Desired purposes are achieved without B being contained, but in order to sufficiently obtain this effect, the B content is preferably set to 0.0005% or more. On the other hand, when the B content is greater than 0.10%, the above-described effect is saturated and economic efficiency decreases. Therefore, the B content is set to 0.10% or less.
“Mo: 0 to 1.0%”
Mo improves the hardenability, and at the same time, has an effect of increasing the strength by forming carbides. Desired purposes are achieved without Mo being contained, but in order to sufficiently obtain this effect, the Mo content is preferably set to 0.01% or more. On the other hand, when the Mo content is greater than 1.0%, the ductility and the weldability sometimes decrease. Therefore, the Mo content is set to 1.0% or less.
“Cu: 0 to 2.0%”
Cu increases the strength of the steel sheet, and at the same time, improves corrosion resistance and removability of scales. Desired purposes are achieved without Cu being contained, but in order to sufficiently obtain this effect, the Cu content is preferably set to 0.01% or more and more preferably set to 0.04% or more. On the other hand, when the Cu content is greater than 2.0%, surface flaws sometimes occur. Therefore, the Cu content is set to 2.0% or less and preferably set to 1.0% or less.
“Ni: 0 to 2.0%”
Ni increases the strength of the steel sheet, and at the same time, improves the toughness. Desired purposes are achieved without Ni being contained, but in order to sufficiently obtain this effect, the Ni content is preferably set to 0.01% or more. On the other hand, when the Ni content is greater than 2.0%, the ductility decreases. Therefore, the Ni content is set to 2.0% or less.
“Mg: 0 to 0.05%, REM: 0 to 0.05%, Ca: 0 to 0.05%, Zr: 0 to 0.05%”
Ca, Mg, Zr, and REM all improve toughness by controlling shapes of sulfides and oxides. Desired purposes are achieved without Ca, Mg, Zr, and REM being contained, but in order to sufficiently obtain this effect, the content of one type or more selected from the group consisting of Ca, Mg, Zr, and REM is preferably set to 0.0001% or more and more preferably set to 0.0005% or more. On the other hand, when the content of Ca, Mg, Zr, or REM is greater than 0.05%, the stretch flangeability deteriorates. Therefore, the content of each of Ca, Mg, Zr, and REM is set to 0.05% or less.
“Metal Microstructure”
Next, there will be explained a structure (metal microstructure) of the steel sheet according to the embodiment of the present invention. In the following explanation, “%” that is a unit of the proportion (area ratio) of each structure means “area %” unless otherwise stated. The steel sheet according to this embodiment has a structure represented by ferrite: 5 to 95% and bainite: 5 to 95%.
“Ferrite: 5 to 95%”
When the area ratio of the ferrite is less than 5%, the ductility deteriorates to make it difficult to secure properties required for automotive members and so on generally. Therefore, the area ratio of the ferrite is set to 5% or more. On the other hand, when the area ratio of the ferrite is greater than 95%, the stretch flangeability deteriorates or it becomes difficult to obtain sufficient strength. Therefore, the area ratio of the ferrite is set to 95% or less.
“Bainite: 5 to 95%”
When the area ratio of the bainite is less than 5%, the stretch flangeability deteriorates. Therefore, the area ratio of the bainite is set to 5% or more. On the other hand, when the area ratio of the bainite is greater than 95%, the ductility deteriorates. Therefore, the area ratio of the bainite is set to 95% or less.
The structure of the steel sheet may contain martensite, retained austenite, pearlite, and so on, for example. When the area ratio of structures other than the ferrite and the bainite is greater than 10% in total, the deterioration in stretch flangeability is concerned. Therefore, the area ratio of the structures other than the ferrite and the bainite is preferably set to 10% or less in total. In other words, the area ratio of the ferrite and the bainite is preferably set to 90% or more and more preferably set to 100% in total.
The proportion (area ratio) of each structure can be obtained by the following method. First, a sample collected from the steel sheet is etched by nital. After the etching, a structure photograph obtained at a ¼ depth position of the sheet thickness in a visual field of 300 μm×300 μm is subjected to an image analysis by using an optical microscope. By this image analysis, the area ratio of ferrite, the area ratio of pearlite, and the total area ratio of bainite and martensite are obtained. Then, a sample etched by LePera is used, and a structure photograph obtained at a ¼ depth position of the sheet thickness in a visual field of 300 μm×300 μm is subjected to an image analysis by using an optical microscope. By this image analysis, the total area ratio of retained austenite and martensite is obtained. Further, a sample obtained by grinding the surface to a depth of ¼ of the sheet thickness from a direction normal to a rolled surface is used, and the volume fraction of retained austenite is obtained through an X-ray diffraction measurement. The volume fraction of the retained austenite is equivalent to the area ratio, and thus is set as the area ratio of the retained austenite. Then, the area ratio of martensite is obtained by subtracting the area ratio of the retained austenite from the total area ratio of the retained austenite and the martensite, and the area ratio of bainite is obtained by subtracting the area ratio of the martensite from the total area ratio of the bainite and the martensite. In this manner, it is possible to obtain the area ratio of each of ferrite, bainite, martensite, retained austenite, and pearlite.
In the steel sheet according to this embodiment, in the case where a region surrounded by a grain boundary having a misorientation of 15° or more and having a circle-equivalent diameter of 0.3 μm or more is defined as a crystal grain, the proportion of crystal grains each having an intragranular misorientation of 5 to 14° to all crystal grains is 20 to 100% by area ratio. The intragranular misorientation is obtained by using an electron back scattering diffraction (EBSD) method that is often used for a crystal orientation analysis. The intragranular misorientation is a value in the case where a boundary having a misorientation of 15° or more is set as a grain boundary in a structure and a region surrounded by this grain boundary is defined as a crystal grain.
The crystal grains each having an intragranular misorientation of 5 to 14° are effective for obtaining a steel sheet excellent in the balance between strength and workability. The proportion of the crystal grains each having an intragranular misorientation of 5 to 14° is increased, thereby making it possible to improve the stretch flangeability while maintaining desired strength of the steel sheet. When the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° to all the crystal grains is 20% or more by area ratio, desired strength and stretch flangeability of the steel sheet can be obtained. It does not matter that the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° is high, and thus its upper limit is 100%.
A cumulative strain at the final three stages of finish rolling is controlled as will be described later, and thereby crystal misorientation occurs in grains of ferrite and bainite. The reason for this is considered as follows. By controlling the cumulative strain, dislocation in austenite increases, dislocation walls are made in an austenite grain at a high density, and some cell blocks are formed. These cell blocks have different crystal orientations. It is conceivable that austenite that has a high dislocation density and contains the cell blocks having different crystal orientations is transformed, and thereby, ferrite and bainite also include crystal misorientations even in the same grain and the dislocation density also increases. Thus, the intragranular crystal misorientation is conceived to correlate with the dislocation density contained in the crystal grain. Generally, the increase in the dislocation density in a grain brings about an improvement in strength, but lowers the workability. However, the crystal grains each having an intragranular misorientation controlled to 5 to 14° make it possible to improve the strength without lowering the workability. Therefore, in the steel sheet according to this embodiment, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° is set to 20% or more. The crystal grains each having an intragranular misorientation of less than 5° are excellent in workability, but have difficulty in increasing the strength. The crystal grains each having an intragranular misorientation of greater than 14° do not contribute to the improvement in stretch flangeability because they are different in deformability among the crystal grains.
The proportion of the crystal grains each having an intragranular misorientation of 5 to 14° can be measured by the following method. First, at a ¼ depth position of a sheet thickness t from the surface of the steel sheet (¼ t portion) in a cross section vertical to a rolling direction, a region of 200 μm in the rolling direction and 100 μm in a direction normal to the rolled surface is subjected to an EBSD analysis at a measurement pitch of 0.2 μm to obtain crystal orientation information. Here, the EBSD analysis is performed by using an apparatus that is composed of a thermal field emission scanning electron microscope (JSM-7001F manufactured by JEOL Ltd.) and an EBSD detector (HIKARI detector manufactured by TSL Co., Ltd.), at an analysis speed of 200 to 300 points/second. Then, with respect to the obtained crystal orientation information, a region having a misorientation of 15° or more and a circle-equivalent diameter of 0.3 μm or more is defined as a crystal grain, the average intragranular misorientation of crystal grains is calculated, and the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° is obtained. The crystal grain defined as described above and the average intragranular misorientation can be calculated by using software “OIM Analysis (registered trademark)” attached to an EBSD analyzer.
The “intragranular misorientation” in this embodiment means “Grain Orientation Spread (GOS)” that is an orientation spread in a crystal grain. The value of the intragranular misorientation is obtained as an average value of misorientations between the reference crystal orientation and all measurement points in the same crystal grain as described in “Misorientation Analysis of Plastic Deformation of Stainless Steel by EBSD and X-ray Diffraction Methods,” KIMURA Hidehiko, et al., Transactions of the Japan Society of Mechanical Engineers (series A), Vol. 71, No. 712, 2005, p. 1722-1728. In this embodiment, the reference crystal orientation is an orientation obtained by averaging all the measurement points in the same crystal grain. The value of GOS can be calculated by using software “OIM Analysis (registered trademark) Version 7.0.1” attached to the EBSD analyzer.
In the steel sheet according to this embodiment, the area ratios of the respective structures observed by an optical microscope such as ferrite and bainite and the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° have no direct relation. In other words, for example, even if there are steel sheets having the same area ratio of ferrite and the same area ratio of bainite, they are not necessarily the same in the proportion of the crystal grains each having an intraqranular misorientation of 5 to 14°. Accordingly, it is impossible to obtain properties equivalent to those of the steel sheet according to this embodiment only by controlling the area ratio of ferrite and the area ratio of bainite.
The steel sheet according to this embodiment contains hard crystal grains A in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1016 to 1×1019 pieces/cm3 and soft crystal grains B in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1015 pieces/cm3 or less, and the volume % of the hard crystal grains A/(the volume % of the hard crystal grains A+the volume % of the soft crystal grains B) is 0.1 to 0.9. The total of the volume % of the hard crystal grains A and the volume % of the soft crystal grains B is preferably set to 70% or more and more preferably set to 80% or more. In other words, when the volume % of crystal grains dispersed with a number density of greater than 1×1015 pieces/cm3 and less than 1×1016 pieces/cm3 is greater than 30%, it is sometimes difficult to obtain properties equivalent to those of the steel sheet according to this embodiment. Thus, the volume % of the crystal grains dispersed with a number density of greater than 1×1015 pieces/cm3 and less than 1×1016 pieces/cm3 is preferably set to 30% or less and more preferably set to 20% or less.
The size of the “precipitates or clusters” in the hard crystal grains A and the soft crystal grains B is a value obtained by measuring the maximum diameter of each of plural precipitates by a later-described measurement method and obtaining the average value of measured values. The maximum diameter of the precipitates is defined as a diameter in the case where the precipitate or cluster has a spherical shape, and is defined as a diagonal length in the case where it has a plate shape.
The precipitates or clusters in the crystal grain contribute to improvement of strengthening of the steel sheet. However, when the maximum diameter of the precipitates exceeds 8 nm, strain concentrates in precipitates in a ferrite structure at the time of working of the steel sheet to be a generation source of voids and thereby the possibility of deterioration in ductility increases, and thus it is not preferred. The lower limit of the maximum diameter of the precipitates does not need to be limited in particular, but it is preferably set to 0.2 nm or more in order to stably sufficiently exhibit the effect of improving the strength of the steel sheet obtained by a pinning force of dislocations in the crystal grain.
The precipitates or clusters in this embodiment are preferably formed of carbides, nitrides, or carbonitrides of one type or more of precipitate-forming elements selected from the group consisting of Ti, Nb, Mo, and V. Here, the carbonitride means a precipitate combined with carbide into which nitrogen is mixed and carbide. Further, in this embodiment, precipitates other than the carbides, nitrides, or carbonitrides of the above-described precipitate-forming element/precipitate-forming elements are allowed to be contained in a range not impairing the properties equivalent to those of the steel sheet according to this embodiment.
In the steel sheet according to this embodiment, the number densities of the precipitates or clusters in the crystal grains of the hard crystal grains A and the soft crystal grains B are limited based on the following mechanism in order to increase both a tensile strength and ductility of the target steel sheet.
As the number density of the precipitates in the crystal grains increases in both the hard crystal grains A and the soft crystal grains B, the hardness of each crystal grain is conceived to increase. On the coriLrary, as the number density of precipitated carbides in the crystal grains decreases in both the hard crystal grains A and the soft crystal grains B, the hardness of each crystal grain is conceived to decrease. In this case, elongation (total elongation, uniform elongation) of each crystal grain increases, but the contribution to strength decreases.
When the hard crystal grains A and the soft crystal grains B are substantially the same in the number density of the precipitates in the crystal grains, the elongation in response to the tensile strength decreases, failing to obtain a sufficient strength-ductility-balance (YP X El). On the other hand, in the case where the difference in number density of the precipitates in the crystal grains between the hard crystal grains A and the soft crystal grains B is large, the elongation in response to the tensile strength increases to be able to obtain a good strength-ductility-balance. The hard crystal grain A plays a role in increasing the strength mainly. The soft crystal grain B plays a role in increasing the ductility mainly. The present inventors experimentally found out that in order to obtain a steel sheet having a good strength-ductility-balance (YP×El), it is necessary to set the number density of the precipitates in the hard crystal grains A to 1×1016 to 1×1019 pieces/cm3 and set the number density of the precipitates in the soft crystal grains B to 1×1015 pieces/cm3 or less.
When the number density of the precipitates in the hard crystal grains A is less than 1×1016 pieces/cm3, the strength of the steel sheet becomes insufficient, failing to obtain the strength-ductility-balance sufficiently. Further, when the number density of the precipitates in the hard crystal grains A exceeds 1×1019 pieces/cm3, the effect of improving the strength of the steel sheet obtained by the hard crystal grains A is saturated to become the cause of an increase in cost due to an added amount of the precipitate-forming element/precipitate-forming elements, or toughness of ferrite or bainite deteriorates and the stretch flangeability deteriorates in some cases.
When the number density of the precipitates in the soft crystal grains B exceeds 1×1015 pieces/cm3, the ductility of the steel sheet becomes insufficient, failing to obtain the strength-ductility-balance sufficiently. For the above reasons, in this embodiment, the number density of the precipitates in the hard crystal grains A is set to 1×1016 to 1×1019 pieces/cm3 and the number density of the precipitates in the soft crystal grains B is set to 1×1015 pieces/cm3 or less.
As for the structure in this embodiment, the ratio of the volume % of the hard crystal grains A to the entire volume of the structure of the steel sheet {the volume % of the hard crystal grains A/(the volume % of the hard. crystal grains A +the volume % of the soft crystal grains B)} is in a range of 0.1 to 0.9. The volume % of the hard crystal grains A to the entire volume of the structure of the steel sheet is set to 0.1 to 0.9, thereby obtaining the strength-ductility-balance of the target steel sheet stably. When the ratio of the volume % of the hard crystal grains A to the entire volume of the structure of the steel sheet is less than 0.1, the strength of the steel sheet decreases, resulting in a difficulty in securing strength, which is a tensile strength of 480 MPa or more. When the ratio of the volume % of the hard crystal grains A exceeds 0.9, the ductility of the steel sheet becomes short.
Incidentally, in the steel sheet according to this embodiment, the fact that the structure is the hard crystal grains A or the soft crystal grains B and the fact that the structure is bainite or ferrite do not always correspond to each other. In the case where the steel sheet according to this embodiment is a hot-rolled steel sheet, for example, the hard crystal grains A are likely to be bainite mainly and the soft crystal grains B are likely to be ferrite mainly. However, ferrite in large amounts may be contained in the hard crystal grains A of the hot-rolled steel sheet, or bainite in large amounts may be contained in the soft crystal grains B. The area ratio of bainite or ferrite in the structure and the proportion of the hard crystal grains A and the soft crystal grains B can be adjusted by annealing or the like.
In the structure of the steel sheet according to this embodiment, the maximum diameter of the precipitates or clusters in the crystal grains and the number density of the precipitates or clusters with a maximum diameter of 8 nm or less can be measured by using the following method.
It is difficult to, though depending on a defect density in the structure, measure the amount of the precipitates with a maximum diameter of 8 nm or less in the crystal grains by an observation by means of a transmission electron microscope (TEM) generally. Therefore, it is preferred to measure the maximum diameter and the number density of the precipitates in the crystal grains by using a three-dimensional atom probe (3D-AP) method suitable for observing the precipitates with a maximum diameter of 8 nm or less. Further, the observation method by means of the 3D-AP is preferred in order to accurately measure the maximum diameter and the number density of the clusters smaller in size out of the precipitates.
The maximum diameter and the number density of the precipitates or clusters in the crystal grains can be measured as follows, for example, by using the observation method by means of the 3D-AP. First, a bar-shaped sample of 0.3 mm×0.3 mm×10 mm is cut out from the steel sheet to be measured and is worked into a needle shape by electropolishing to be set as a sample. By using this sample, half a million atoms or more are measured by the 3D-AP in an arbitrary direction in a crystal grain and are visualized by a three-dimensional map to be quantitatively analyzed. Such a measurement in an arbitrary direction is performed on 10 or more different crystal grains and the maximum diameter of precipitates contained in each of the crystal grains and the number density of precipitates with a maximum diameter of 8 nm or less (the number of precipitates per volume of an observation region) are obtained as average values. As the maximum diameter of the precipitates in the crystal grain, out of precipitates each having an apparent shape, a bar length of bar-shaped one, a diagonal length of plate-shaped one, and a diameter of spherical-shaped one are set. Out of the precipitates, clusters smaller in size in particular are not apparent in terms of their shapes in many cases, and thus the maximum diameters of the precipitates and the clusters are preferably determined by a precise size measurement method utilizing field evaporation of a field-ion microscope (FIM) or the like.
The arbitrary crystal grains and the measurement results in arbitrary directions as above make it possible to find a precipitation state of the precipitates in each crystal grain and distinguish crystal grains with different precipitation states of precipitates from one another, and find a volume ratio of these.
Further, in addition to the above-described measurement method, it is also possible to use a field-ion microscope (FIM) method, which enables a broader visual field, in combination. The FIM is a method of two-dimensionally projecting a surface electric field distribution by applying a high voltage to a needle-shaped sample and introducing an inert gas. Generally, precipitates in a steel material provide lighter or darker contrast than a ferrite matrix. Field evaporation of a specific atomic plane is performed one atomic plane by one atomic plane to observe occurrence and disappearance of contrast of precipitates, thereby making it possible to accurately estimate the size of the precipitate in a depth direction.
In this embodiment, the stretch flangeability is evaluated by a saddle-type stretch-flange test method using a saddle-type formed product.
In a conventional hole expansion test used as a test method coping with the stretch flangeability, the sheet leads to a fracture with little or no strain distributed in a circumferential direction. Therefore, the strain and the stress gradient around a fractured portion differ from those at an actual stretch flange forming time. Further, in the hole expansion test, evaluation is made at the point in time when a fracture occurs penetrating the sheet thickness, or the like, resulting in that the evaluation reflecting the original stretch flange forming is not made. On the other hand, in the saddle-type stretch-flange test used in this embodiment, the stretch flangeability considering the strain distribution can be evaluated, and thus the evaluation reflecting the original stretch flange forming can be made.
According to the steel sheet according to this embodiment, a tensile strength of 480 MPa or more can be obtained. That is, an excellent tensile strength can be obtained. The upper limit of the tensile strength is not limited in particular. However, in a component range in this embodiment, the upper limit of the practical tensile strength is about 1180 MPa. The tensile strength can be measured by fabricating a No. 5 test piece described in JIS-Z2201 and performing a tensile test according to a test method described in JIS-Z2241.
According to the steel sheet according to this embodiment, the product of the tensile strength and the limit form height in the saddle-type stretch-flange test, which is 19500 min·Pa or more, can be obtained. That is, excellent stretch flangeability can be obtained. The upper limit of this product is not limited in particular. However, in a component range in this embodiment, the upper limit of this practical product is about 25000 mm·MPa.
According to the steel sheet according to this embodiment, the product of a yield stress and ductility, which is 10000 MPa·% or more, can be obtained. That is, an excellent strength-ductility-balance can be obtained.
Next, there will be explained a method of manufacturing the steel sheet according to the embodiment of the present invention. In this method, hot rolling, first cooling, and second cooling are performed in this order.
“Hot Rolling”
The hot rolling includes rough rolling and finish rolling. In the hot rolling, a slab (steel billet) having the above-described chemical composition is heated to be subjected to rough rolling. A slab heating temperature is set to SRTmin° C. expressed by Expression (1) below or more and 1260° C. or less.
SRTmin=[7000/{2.75−log([Ti]×[C])}×273)+10000/{4.29−log([Nb]×[C])}−273)]/2 (1)
Here, [Ti], [Nb], and [C] in Expression (1) represent the contents of Ti, Nb, and C in mass %.
When the slab heating temperature is less than SRTmin° C., Ti and/or Nb are/is not sufficiently brought into solution. When Ti and/or Nb are/is not brought into solution at the time of slab heating, it becomes difficult to make Ti and/or Nb finely precipitate as carbides (TiC, NbC) and improve the strength of the steel by precipitation strengthening. Further, when the slab heating temperature is less than SRTmin° C., it becomes difficult to fix C by formation of the carbides (TiC, NbC) to suppress generation of cementite harmful to a burring property. Further, when the slab heating temperature is less than SRTmin° C., the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° is likely to be short. Therefore, the slab heating temperature is set to SRTmin° C. or more. On the other hand, when the slab heating temperature is greater than 1260° C., the yield decreases due to scale-off. Therefore, the slab heating temperature is set to 1260° C. or less.
By the rough rolling, a rough bar is obtained. Thereafter, by finish rolling, a hot-rolled steel sheet is obtained. The cumulative strain at the final three stages (final three passes) in the finish rolling is set to 0.5 to 0.6 in order to set the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° to 20% or more, and then later-described cooling is performed. This is due to the following reason. The crystal grains each having an intragranular misorientation of 5 to 14° are generated by being transformed in a paraequilibrium state at relatively low temperature. Therefore, the dislocation density of austenite before transformation is limited to a certain range in the hot rolling, and at the same time, the subsequent cooling rate is limited to a certain range, thereby making it possible to control generation of the crystal grains each having an intragranular misorientation of 5 to 14°.
That is, the cumulative strain at the final three stages in the finish rolling and the subsequent cooling are controlled, thereby making it possible to control the nucleation frequency of the crystal grains each having an intragranular misorientation of 5 to 14° and the subsequent growth rate. As a result, it is possible to control the area ratio of the crystal grains each having an intragranular misorientation of 5 to 14° in a steel sheet to be obtained after cooling. More concretely, the dislocation density of the austenite introduced by the finish rolling is mainly related to the nucleation frequency and the cooling rate after the rolling is mainly related to the growth rate.
When the cumulative strain at the final three stages in the finish rolling is less than 0.5, the dislocation density of the austenite to be introduced is not sufficient and the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes less than 20%. Therefore, the cumulative strain at the final three stages is set to 0.5 or more. On the other hand, when the cumulative strain at the final three stages in the finish rolling exceeds 0.6, recrystallization of the austenite occurs during the hot rolling and the accumulated dislocation density at a transformation time decreases. As a result, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes less than 20%. Therefore, the cumulative strain at the final three stages is set to 0.6 or less.
The cumulative strain at the final three stages in the finish rolling (εeff.) is obtained by
Expression (2) below.
εeff.=Σεi(t,T) (2)
Here,
εi(t,T)=εi0/exp{(t/τR)2/3},
τR=τ0·exp(Q/RT),
τ0=8.46×10−9,
Q=183200J,
R=8.314J/K·mol,
εi0 represents a logarithmic strain at a reduction time, t represents a cumulative time period till immediately before the cooling in the pass, and T represents a rolling temperature in the pass.
When a finishing temperature of the rolling is set to less than Ar3° C., the dislocation density of the austenite before transformation increases excessively, to thus make it difficult to set the crystal grains each having an intragranular misorientation of 5 to 14° to 20% or more. Therefore, the finishing temperature of the finish rolling is set to Ar3° C. or more.
The finish rolling is preferably performed by using a tandem rolling mill in which a plurality of rolling mills are linearly arranged and that performs rolling continuously in one direction to obtain a desired thickness. Further, in the case where the finish rolling is performed using the tandem rolling mill, cooling (inter-stand cooling) is performed between the rolling mills to control the steel sheet temperature during the finish rolling to fall within a range of Ar3° C. or more to Ar3+150° C. or less. When the maximum temperature of the steel sheet during the finish rolling exceeds Ar3+150° C., the grain size becomes too large, and thus deterioration in toughness is concerned.
The hot rolling is performed under such conditions as above, thereby making it possible to limit the dislocation density range of the austenite before transformation and obtain a desired proportion of the crystal grains each having an intragranular misorientation of 5 to 14°.
Ar3 is calculated by Expression (3) below considering the effect on the transformation point by reduction based on the chemical composition of the steel sheet.
Ar3=970-325×[C]+33×[Si]+287×[P]+40×[Al]−92×([Mn]+[Mo]+[Cu])−46×([Cr]+[Ni]) (3)
Here, [C], [Si], [P], [Al], [Mn], [Mo], [Cu], [Cr], and [Ni] represent the contents of C, Si, P, Al, Mn, Mo, Cu, Cr, and Ni in mass % respectively. The elements that are not contained are calculated as 0%.
“First Cooling, Second Cooling”
After the hot rolling, the first cooling and the second cooling of the hot-rolled steel sheet are performed in this order. In the first cooling, the hot-rolled steel sheet is cooled down to a first temperature zone of 600 to 750° C. at a cooling rate of 10° C./s or more. In the second cooling, the hot-rolled steel sheet is cooled down to a second temperature zone of 450 to 650° C. at a cooling rate of 30° C./s or more. Between the first cooling and the second cooling, the hot-rolled steel sheet is retained in the first temperature zone for 1 to 10 seconds. After the second cooling, the hot-rolled steel sheet is preferably air-cooled.
When the cooling rate of the first cooling is less than 10° C./s, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short. Further, when a cooling stop temperature of the first cooling is less than 600° C., it becomes difficult to obtain 5% or more of ferrite by area ratio, and at the same time, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short. Further, when the cooling stop temperature of the first cooling is greater than 750° C., it becomes difficult to obtain 5% or more of bainite by area ratio, and at the same time, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short.
When the retention time at 600 to 750° C. exceeds 10 seconds, cementite harmful to the burring property is likely to be generated. Further, when the retention time at 600 to 750° C. exceeds 10 seconds, it is often difficult to obtain 5% or more of bainite by area ratio, and further, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short. When the retention time at 600 to 750° C. is less than 1 second, it becomes difficult to obtain 5% or more of ferrite by area ratio, and at the same time, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short.
When the cooling rate of the second cooling is less than 30° C./s, cementite harmful to the burring property is likely to be generated, and at the same time, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short. When a cooling stop temperature of the second cooling is less than 450° C. or greater than 650° C., the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° becomes short.
The upper limit of the cooling rate in each of the first cooling and the second cooling is not limited, in particular, but may be set to 200/s or less in consideration of the facility capacity of a cooling facility.
It is effective to set a temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling to 30 to 250° C. When the temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling is less than 30° C., the volume % of the hard crystal grains A to the entire volume of the structure of the steel sheet (the volume % of the hard crystal grains A/(the volume % of the hard crystal grains A +the volume % of the soft crystal grains B)} becomes less than 0.1. Therefore, the temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling is set to 30° C. or more, preferably set to 40° C. or more, and more preferably set to 50° C. or more. When the temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling exceeds 250° C., the volume % of the hard crystal grains A to the entire volume of the structure of the steel sheet becomes greater than 0.9. Therefore, the temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling is set to 250° C. or less, preferably set to 230° C. or less, and more preferably set to 220° C. or less.
Further, the temperature difference between the cooling stop temperature of the first cooling and the cooling stop temperature of the second cooling is set to 30 to 250° C., and thereby the structure contains the hard crystal grains A in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1016 to 1×1019 pieces/cm3 and the soft crystal grains B in which precipitates or clusters with a maximum diameter of 8 nm or less are dispersed in the crystal grains with a number density of 1×1015 pieces/cm3 or less.
In this manner, it is possible to obtain the steel sheet according to this embodiment.
In the above-described manufacturing method, the hot rolling conditions are controlled, to thereby introduce work dislocations into the austenite. Then, it is important to make the introduced work dislocations remain moderately by controlling the cooling conditions. That is, even when the hot rolling conditions or the cooling conditions are controlled independently, it is impossible to obtain the steel sheet according to this embodiment, resulting in that it is important to appropriately control both of the hot rolling conditions and the cooling conditions. The conditions other than the above are not limited in particular because well-known methods such as coiling by a well-known method after the second cooling, for example, only need to be used. Further, temperature zones for precipitation are separated, thereby making it possible to disperse the above-described hard crystal grains A and soft crystal grains B.
Pickling may be performed in order to remove scales on the surface. As long as the hot rolling and cooling conditions are as above, it is possible to obtain the similar effects even when cold rolling, a heat treatment (annealing), plating, and so on are performed thereafter.
In the cold rolling, a reduction ratio is preferably set to 90% or less. When the reduction ratio in the cold rolling exceeds 90%, the ductility sometimes decreases. This is conceivably because the hard crystal grains A and the soft crystal grains B are greatly crushed by the cold rolling, and recrystallized grains at an annealing time after the cold rolling encroach on both portions that were the hard crystal grains A and the soft crystal grains B after the hot rolling and are no longer the crystal grains having two types hardnesses. The cold rolling does not have to be performed and the lower limit of the reduction ratio in the cold rolling is 0%. As above, an intact hot-rolled original sheet has excellent formability. On the other hand, on dislocations introduced by the cold rolling, solid-dissolved Ti, Nb, Mo, and so on collect to precipitate, thereby making it possible to improve a yield point (YP) and a tensile strength (TS). Thus, the cold rolling can be used for adjusting the strength. A cold-rolled steel sheet is obtained by the cold rolling.
The temperature of the heat treatment (annealing) after the cold rolling is preferably set to 40° C. or less. At the time of annealing, complicated phenomena such as strengthening by precipitation of Ti and Nb that did not precipitate sufficiently at the hot rolling stage, dislocation recovery, and softening by coarsening of precipitates occur. When the annealing temperature exceeds 840° C., the effect of coarsening of precipitates is large, the precipitates with a maximum diameter of 8 nm or less decrease, and at the same time, the proportion of the crystal grains each having an intragranular crystal misorientation of 5 to 14° becomes short. The annealing temperature is more preferably set to 820° C. or less and further preferably set to 800° C. or less. The lower limit of the annealing temperature is not set in particular. As described above, this is because the intact hot-rolled original sheet that is not subjected to annealing has excellent formability.
On the surface of the steel sheet in this embodiment, a plating layer may be formed. That is, a plated steel sheet can be cited as another embodiment of the present invention. The plating layer is, for example, an electroplating layer, a hot-dip plating layer, or an alloyed hot-dip plating layer. As the hot-dip plating layer and the alloyed hot-dip plating layer, a layer made of at least one of zinc and aluminum, for example, can be cited. Concretely, there can be cited a hot-dip galvanizing layer, an alloyed hot-dip galvanizing layer, a hot-dip aluminum plating layer, an alloyed hot-dip aluminum plating layer, a hot-dip Zn—Al plating layer, an alloyed hot-dip Zn—Al plating layer, and so on. From the viewpoints of platability and corrosion resistance, in particular, the hot-dip galvanizing layer and the alloyed hot-dip galvanizing layer are preferable.
A hot-dip plated steel sheet and an alloyed hot-dip plated steel sheet are manufactured by performing hot dipping or alloying hot dipping on the aforementioned steel sheet according to this embodiment. Here, the alloying hot dipping means that hot dipping is performed to form a hot-dip plating layer on a surface, and then an alloying treatment is performed thereon to form the hot-dip plating layer into an alloyed hot-dip plating layer. The steel sheet that is subjected to plating may be the hot-rolled steel sheet, or a steel sheet obtained after the cold rolling and the annealing are performed on the hot-rolled steel sheet. The hot-dip plated steel sheet and the alloyed hot-dip plated steel sheet include the steel sheet according to this embodiment and have the hot-dip plating layer and the alloyed hot-dip plating layer provided thereon respectively, and thereby, it is possible to achieve an excellent rust prevention property together with the functional effects of the steel sheet according to this embodiment. Before performing plating, Ni or the like may be applied to the surface as pre-plating.
When the heat treatment (annealing) is performed on the steel sheet, the steel sheet may be immersed in a hot-dip galvanizing bath directly after being subjected to the heat treatment to form the hot-dip galvanizing layer on the surface thereof. In this case, the original sheet for the heat treatment may be the hot-rolled steel sheet or the cold-rolled steel sheet. After the hot-dip galvanizing layer is formed, the alloyed hot-dip galvanizing layer may be formed by reheating the steel sheet and performing the alloying treatment to alloy the galvanizing layer and the base iron.
The plated steel sheet according to the embodiment of the present invention has an excellent rust prevention property because the plating layer is formed on the surface of the steel sheet. Thus, when an automotive member is reduced in thickness by using the plated steel sheet in this embodiment, for example, it is possible to prevent shortening of the usable life of an automobile that is caused by corrosion of the member.
Note that the above-described embodiments merely illustrate concrete examples of implementing the present invention, and the technical scope of the present invention is not to be construed in a restrictive manner by these embodiments. That is, the present invention may be implemented in various forms without departing from the technical spirit or main features thereof.
Next, examples of the present invention will be explained. Conditions in the examples are examples of conditions employed to verify feasibility and effects of the present invention, and the present invention is not limited to the examples of conditions. The present invention can employ various conditions without departing from the spirit of the present invention to the extent to achieve the objects of the present invention.
Steels having chemical compositions illustrated in Table 1 and Table 2 were smelted to manufacture steel billets, the obtained steel billets were heated to heating temperatures illustrated in Table 3 and Table 4 to be subjected to rough rolling in hot working and then subjected to finish rolling under conditions illustrated in Table 3 and Table 4. Sheet thicknesses of hot-rolled steel sheets after the finish rolling were 2.2 to 3.4 mm. Each blank column in Table 1 and Table 2 indicates that an analysis value was less than a detection limit. Each underline in Table 1 and Table 2 indicates that a numerical value thereof is out of the range of the present invention, and each underline in Table 4 indicates that a numerical value thereof is out of the range suitable for the manufacture of the steel sheet of the present invention.
0.162
2.73
3.21
0.250
0.007
0.282
0.231
1210
985
1270
1125
846
0.44
0.71
1055
0.61
Ar3 (° C.) was obtained from the components illustrated in Table 1 and Table 2 by using Expression (3).
Ar3=970−325×[C]+33×[Si]+287×[P]+×[Al]−92×([Mn]+[Mo]+[Cu])−46×([Cr]+[Ni]) (3)
The cumulative strain at the final three stages was obtained by Expression (2)
εeff.=Σεi(t,T) (2)
Here,
εi(t,T)=εi0/exp{(t/τR)2/3},
τR=τ0·exp(Q/RT),
τ0=8.46×10−9,
Q=183200J,
R=8.314J/K·mol,
εi0 represents a logarithmic strain at a reduction time, t represents a cumulative time period till immediately before the cooling in the pass, and T represents a rolling temperature in the pass.
Next, under conditions illustrated in Table 5 and Table 6, first cooling, retention in a first temperature zone, and second cooling were performed, and hot-rolled steel sheets of Test No. 1 to 44 were obtained.
The hot-rolled steel sheet of Test No. 21 was subjected to cold rolling at a reduction ratio illustrated in Table 5 and subjected to a heat treatment at a heat treatment temperature illustrated in Table 5, and then had a hot-dip galvanizing layer formed thereon, and further an alloying treatment was performed to thereby form an alloyed hot-dip galvanizing layer (GA) on a surface. The hot-rolled steel sheets of Test No. 18 to 20, and 44 were subjected to a heat treatment at heat treatment temperatures illustrated in Table 5 and Table 6. The hot-rolled steel sheets of Test No. 18 to 20 were subjected to a heat treatment, and then had hot-dip galvanizing layers (GI) each formed thereon. Each underline in Table 6 indicates that a numerical value thereof is out of the range suitable for the manufacture of the steel sheet of the present invention.
540
790
0
15
350
680
260
860
Then, of each of the steel sheets (the hot-rolled steel sheets of Test No. 1 to 17 and 22 to 43, the heat-treated hot-rolled steel sheets of Test No. 18 to 20, and 44, and a heat-treated cold-rolled steel sheet of Test No. 21), structural fractions (area ratios) of ferrite, bainite, martensite, and pearlite and a proportion of crystal grains each having an intragranular misorientation of 5 to 14° were obtained by the following methods. Results thereof are illustrated in Table 7 and Table 8. The case where martensite and/or pearlite are/is contained was described in the column of “BAINITE AREA RATIO” in the table in parentheses. Each underline in Table 8 indicates that a numerical value thereof is out of the range of the present invention.
“Structural Fractions (Area Ratios) of Ferrite, Bainite, Martensite, and Pearlite”
First, a sample collected from the steel sheet was etched by nital. After the etching, a structure photograph obtained at a ¼ depth position of the sheet thickness in a visual field of 300 μm×300 μm was subjected to an image analysis by using an optical microscope. By this image analysis, the area ratio of ferrite, the area ratio of pearlite, and the total area ratio of bainite and martensite were obtained. Next, a sample etched by LePera was used, and a structure photograph obtained at a ¼ depth position of the sheet thickness in a visual field of 300 μm×300 μm was subjected to an image analysis by using an optical microscope. By this image analysis, the total area ratio of retained austenite and martensite was obtained. Further, a sample obtained by grinding the surface to a depth of ¼ of the sheet thickness from a direction normal to a rolled surface was used, and the volume fraction of the retained austenite was obtained through an X-ray diffraction measurement. The volume fraction of the retained austenite was equivalent to the area ratio, and thus was set as the area ratio of the retained austenite. Then, the area ratio of martensite was obtained by subtracting the area ratio of the retained austenite from the total area ratio of the retained austenite and the martensite, and the area ratio of bainite was obtained by subtracting the area ratio of the martensite from the total area ratio of the bainite and the martensite. In this manner, the area ratio of each of ferrite, bainite, martensite, retained austenite, and pearlite was obtained.
“Proportion of Crystal Grains Each Having an Intragranular Misorientation of 5 to 14°”
At a ¼ depth position of a sheet thickness t from the surface of the steel sheet (¼ t portion) in a cross section vertical to a rolling direction, a region of 200 μm in the rolling direction and 100 μm in a direction normal to the rolled surface was subjected to an EBSD analysis at a measurement pitch of 0.2 μm to obtain crystal orientation information. Here, the EBSD analysis was performed by using an apparatus composed of a thermal field emission scanning electron microscope (JSM-7001F manufactured by JEOL Ltd.) and an EBSD detector (HIKARI detector manufactured by TSL Co., Ltd.), at an analysis speed of 200 to 300 points/second. Next, with respect to the obtained crystal orientation information, a region having a misorientation of 15° or more and a circle-equivalent diameter of 0.3 μm or more was defined as a crystal grain, the average intragranular misorientation of crystal grains was calculated, and the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° was obtained. The crystal grain defined as described above and the average intragranular misorientation were calculated by using software “OIM Analysis (registered trademark)” attached to an EBSD analyzer.
Of each of the steel sheets (the hot-rolled steel sheets of Test No. 1 to 17 and 22 to 43, the heat-treated hot-rolled steel sheets of Test No. 18 to 20, and 44, and the heat-treated cold-rolled steel sheet of Test No. 21), the maximum diameter of precipitates or clusters in crystal grains and the number density of precipitates or clusters with a maximum diameter of 8 nm or less were measured by the following method. Further, the volume % of hard crystal grains A and the volume % of soft crystal grains B were calculated by using obtained measured values, to obtain the volume % of the hard crystal grains A/(the volume % of the hard crystal grains A +the volume % of the soft crystal grains B) (a volume ratio A/(A+B)}. Results thereof are illustrated in Table 7 and Table 8.
“Measurement of the Maximum Diameter of Precipitates or Clusters yn Crystal Grains and the Number Density of Precipitates or Clusters with a Maximum Diameter of 8 nm or Less”
The maximum diameter and the number density of precipitates or clusters in the crystal grains were measured as follows by using an observation method by means of a 3D-AP. A bar-shaped sample of 0.3 mm×0.3 mm×10 mm was cut out from the steel sheet to be measured and was worked into a needle shape by electropolishing to be set as a sample. By using this sample, half a million atoms or more were measured by the 3D-AP in an arbitrary direction in a crystal grain and were visualized by a three-dimensional map to be quantitatively analyzed. Such a measurement in an arbitrary direction was performed on 10 or more different crystal grains and the maximum diameter of precipitates contained in each of the crystal grains and the number density of precipitates with a maximum diameter of 8 nm or less (the number of precipitates per volume of an observation region) were obtained as average values. As the maximum diameter of the precipitates in the crystal grain, out of precipitates each having an apparent shape, a bar length of bar-shaped one, a diagonal length of plate-shaped one, and a diameter of spherical-shaped one were set. Out of the precipitates, clusters smaller in size in particular are not apparent in terms of their shapes in many cases, and thus the maximum diameters of the precipitates and the clusters were determined by a precise size measurement method utilizing field evaporation of a field-ion microscope (FIM).
Further, in addition to the above-described measurement method, a field-ion microscope (FIM) method enabling a broader visual field was used in combination. The FIM is a method of two-dimensionally projecting a surface electric field distribution by applying a high voltage to a needle-shaped sample and introducing an inert gas. Ones having lighter or darker contrast than a ferrite matrix were set as precipitates. Field evaporation of a specific atomic plane was performed one atomic plane by one atomic plane to observe occurrence and disappearance of the contrast of the precipitates, to thereby estimate the size of the precipitate in a depth direction.
11
100
15
19
18
13
18
96
10
0.04
17
98
18
0.02
13
0.97
11
12
0.05
10
0.96
0.96
0.97
10
Of each of the hot-rolled steel sheets of Test No. 1 to 17 and 22 to 43, the heat-treated hot-rolled steel sheets of Test No. 18 to 20, and 44, and the heat-treated cold-rolled steel sheet of Test No. 21, in a tensile test, a yield strength and a tensile strength were obtained, and by a saddle-type stretch-flange test, a limit form height of a flange was obtained. Then, the product of the tensile strength (MPa) and the limit form height (mm) was set as an index of the stretch flangeability, and the case of the product being 19500 mm·MPa or more was judged to be excellent in stretch flangeability. Further, the case of the tensile strength (TS) being 480 MPa or more was judged to be high in strength. Further, the case where the product of a yield stress (YP) and ductility (EL) is 10000 MPa·% or more was judged to be good in the strength-ductility-balance. Results thereof are illustrated in Table 9 and Table 10. Each underline in Table 10 indicates that a numerical value thereof is out of a desirable range.
As for the tensile test, a JIS No. 5 tensile test piece was collected from a direction right angle to the rolling direction, and this test piece was used to perform the test according to JISZ2241.
The saddle-type stretch-flange test was performed by using a saddle-type formed product in which a radius of curvature of a corner is set to R60 mm and an opening angle θ is set to 120° and setting a clearance at the time of punching the corner portion to 11%. The limit form height was set to a limit form height with no existence of cracks by visually observing whether or not a crack having a length of ⅓ or more of the sheet thickness exists after forming.
17430
18231
10091
460
10898
17414
17332
18581
18284
18417
18127
16371
17609
16365
16949
19051
17606
14277
14632
12994
In the present invention examples (Test No. 1 to 21), the tensile strength of 480 MPa or more, the product of the tensile strength and the limit form height in the saddle-type stretch-flange test of 19500 mm·MPa or more, and the product of the yield stress and the ductility of 10000 MPa·% or more were obtained.
Test No. 22 to 28 each are a comparative example in which the chemical composition is out of the range of the present invention. In Test No. 22 to 24 and Test No. 28, the index of the stretch flangeability did not satisfy the target value, In Test No. 25, the total content of Ti and Nb was small, and thus the stretch flangeability and the product of the yield stress (YP) and the ductility (EL) did not satisfy the target values. In Test No. 26, the total content of Ti and Nb was large, and thus the workability deteriorated and cracks occurred during rolling.
Test No. 28 to 44 each are a comparative example in which the manufacturing conditions were out of a desirable range, and thus one or more of the structures observed by an optical microscope, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14°, the number density of the precipitates in the hard crystal grains A, the number density of the precipitates in the soft crystal grains B, and the volume ratio {the volume % of the hard crystal grains A/(the volume % of the hard crystal grains A +the volume % of the soft crystal grains B) did not satisfy the range of the present invention. In Test No. 29 to 41 and Test No. 44, the proportion of the crystal grains each having an intragranular misorientation of 5 to 14° was small, and thus the product of the yield stress (YP) and the ductility (EL) and/or the index of the stretch flangeability did not satisfy the target values/target value. In Test No. 42 to 43, the volume ratio {A/(A+B)} was large, and thus the product of the yield stress (YP) and the ductility (EL) and the index of the stretch flangeability did not satisfy the target values.
According to the present invention, it is possible to provide a steel sheet that is high in strength, has good ductility and stretch flangeability, and has a high yield stress. The steel sheet of the present invention is applicable to a member required to have strict stretch flangeability while having high strength. The steel sheet of the present invention is a material suitable for the weight reduction achieved by thinning of automotive members and contributes to improvement of fuel efficiency and so on of automobiles, and thus has high industrial applicability.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2016-155090 | Aug 2016 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2017/028478 | 8/4/2017 | WO | 00 |
