HIGH-STRENGTH STEEL PLATE AND MANUFACTURING METHOD THEREOF

Disclosed is a high-strength sheet including: C: 0.15% to 0.35% by mass, total of Si and Al: 0.5% to 3.0% by mass, Mn: 1.0% to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, with the balance being Fe and inevitable impurities, wherein the steel structure satisfies that: a ferrite fraction is 5% or less, the total fraction of tempered martensite and tempered bainite is 60% or more, the amount of retained austenite is 10% or more, MA has an average size of 1.0 μm or less, retained austenite has an average size of 1.0 μm or less, retained austenite having a size of 1.5 μm or more accounts for 2% or more of the total amount of retained austenite, and a scattering intensity at the q value of 1 nm−1 in X-ray small angle scattering is 1.0 cm1 or less.

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Description
TECHNICAL FIELD

The present disclosure relates to a high-strength sheet that can be used in various applications including automobile parts.

BACKGROUND ART

In order to realize both weight reduction and collision safety, steel sheets used for automobile parts and the like are required to achieve both improvement in strength and improvement in impact resistance properties.

For example, Patent Document 1 discloses a high-strength steel sheet in which an attempt is made to improve impact resistance properties by heating a slab to 1,210° C. or higher and controlling the hot-rolling conditions to form fine TiN particles having a size of 0.5 μm or less, thereby suppressing the formation of AlN particles having a particle size of 1 μm or more that act as a starting point of low temperature fracture.

Patent Document 2 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by forming a network structure in which 50% or more of a ferrite grain size is in contact with a hard phase while defining the C amount to more than 0.45% and 0.77% or less, the Mn amount to 0.1% or more and 0.5% or less and the Si amount to 0.5% or less, and defining each addition amount of Cr, Al, N and O.

Patent Document 3 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by adding 3.5 to 10% of Mn, thereby adjusting the amount of retained austenite to 10% or more and an average interval of retained austenite to 1.5 μm or less.

Patent Document 4 discloses a high-strength sheet that has a tensile strength of 980 to 1,180 MPa and also exhibits satisfactory deep drawability.

PRIOR ART DOCUMENT

Patent Document

  • Patent Document 1: JP 5240421 B1
  • Patent Document 2: JP 2015-105384 A
  • Patent Document 3: JP 2012-251239 A
  • Patent Document 4: JP 2009-203548 A

DISCLOSURE OF THE INVENTION Means or Solving the Problems

In order to realize further weight reduction, steel sheets used for automobile parts are required to have sufficient strength and impact resistance properties while being made thinner. Thus, steel sheets having higher tensile strength and excellent impact properties are required.

In various applications including automobile parts, steel sheets are required to have not only high tensile strength and impact properties, but also excellent strength-ductility balance, high yield ratio, excellent deep drawability and excellent hole expansion ratio.

Specifically, the followings are required for each of the tensile strength, the strength-ductility balance, the yield ratio, the deep drawing properties and the hole expansion ratio.

The tensile strength is required to be 980 MPa or higher. In order to increase stress that can be applied during use, there is a need to have high yield strength (YS), in addition to high tensile strength (TS). From the viewpoint of ensuring collision safety or the like, there is a need to increase the yield strength of the steel sheet and to attain properties capable of suppressing fracture during deformation in order to stably exhibit strength properties upon collision. Therefore, specifically, there is required, together with a yield ratio (YR=YS/TS) of 0.75 or more, an increase in thickness reduction ratio of the fracture portion during a tensile test as an evaluation index substituting fracture properties. A joint strength of the spot welded portion is also required as basic performances of the steel sheet for automobiles. Specifically, a cross tensile strength of the spot welded portion is required to be 6 kN or higher.

Regarding the strength-ductility balance, the product (TS×EL) of TS and total elongation (EL) is required to be 20,000 MPa % or higher. In order to ensure the formability during parts forming, it is also required that LDR showing deep drawability is 2.05 or more and the hole expansion ratio λ showing expansion properties is 20% or more.

However, it is difficult for the high-strength sheets disclosed in Patent Documents 1 to 4 to satisfy all of these requirements, and there has been required a high-strength steel sheet that can satisfy all of these requirements.

The embodiment of the present invention has been made to respond to these requirements, and it is an object thereof to provide a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), LDR, hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test, and cross tensile strength (SW cross tension) of the spot welded portion are at a high level, and a manufacturing method thereof.

Means for Solving the Problems

Aspect 1 of the present invention provides a high-strength sheet including:

C: 0.15% by mass to 0.35% by mass,

total of Si and Al: 0.5% by mass to 3.0% by mass,

Mn: 1.0% by mass to 4.0% by mass,

P: 0.05% by mass or less, and

S: 0.01% by mass or less, with the balance being Fe and inevitable impurities,

in which the steel structure satisfies that:

a ferrite fraction is 5% or less,

the total fraction of tempered martensite and tempered bainite is 60% or more,

the amount of retained austenite is 10% or more,

MA has an average size of 1.0 μm or less,

retained austenite has an average size of 1.0 μm or less,

retained austenite having a size of 1.5 μm or more accounts for 2% or more of the total amount of retained austenite, and

a scattering intensity at the q value of 1 nm−1 in X-ray small angle scattering is 1.0 cm−1 or less.

Aspect 2 of the present invention provides the high-strength sheet according to aspect 1, in which the C amount is 0.30% by mass or less.

Aspect 3 of the present invention provides the high-strength sheet according to aspect 1 or 2, in which the Al amount is less than 0.10% by mass.

Aspect 4 of the present invention provides a method for manufacturing a high-strength sheet, which includes:

preparing a rolled material including: C: 0.15% by mass to 0.35% by mass, total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, with the balance being Fe and inevitable impurities;

heating the rolled material to a temperature of an Ac3 point or higher, thereby austenitizing the rolled material;

after the austenitizing, cooling the material between 650° C. and 500° C. at an average cooling rate of 15° C./sec or more and less than 200° C./sec, followed by retention at a temperature in a range of 300° C. to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more and less than 300 seconds;

after the retention, cooling the material from a temperature of 300° C. or higher to a cooling stopping temperature between 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more;

heating the material from the cooling stopping temperature to a reheating temperature in a range of 300° C. to 500° C. at an average heating rate of 30° C./sec or more;

holding at the reheating temperature so as to satisfy a tempering parameter P of 10,000 to 14,500 defined in the equation (1) and a holding time of 1 to 300 seconds; and

after the holding, cooling from the reheating temperature to 200° C. at an average cooling rate of 10° C./sec or more.


P=T×(20+log(t/3600))  (1)

where T: reheating temperature (K) and t: holding time (seconds).

Aspect 5 of the present invention provides the manufacturing method according to aspect 4, in which the retention includes holding at a constant temperature in a range of 300° C. to 500° C.

Aspect 6 of the present invention provides the manufacturing method according to aspect 4 or 5, in which the tempering parameter is 11,000 to 14,000 and the holding time is 1 to 150 seconds.

Effects of the Invention

According to the embodiment of the present invention, it is possible to provide a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), LDR, hole expansion ratio (A), thickness reduction ratio (RA) of the fracture portion during a tensile test (impact resistance properties), and cross tensile strength (SW cross tension) of the spot welded portion are at a high level, and a manufacturing method thereof.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.

 MODE FOR CARRYING OUT THE INVENTION

The inventors of the present application have intensively studied and found that it is possible to obtain a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), LDR, hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test (impact resistance properties), and cross tensile strength (SW cross tension) of the spot welded portion are at a high level by allowing the steel structure (metal structure) to satisfy that: a ferrite fraction is 5% or less, the total fraction of tempered martensite and tempered bainite is 60% or more, the amount of retained austenite (γ) is 10% or more, MA has an average size of 1.0 μm or less, retained austenite has an average size of 1.0 μm or less, retained austenite having a size of 1.5 μm or more accounts for 2% or more of the total amount of retained austenite, and a scattering intensity at the q value of 1 nm−1 in X-ray small angle scattering is 1.0 cm−1 or less, in a steel including predetermined components.

1. Steel Structure

The steel structure of the high-strength sheet according to the embodiment of the present invention will be described in detail below.

In the following description of the steel structure, there are cases where mechanisms capable of improving various properties by having such the structure are described. It should be noted that these mechanisms are those envisaged by the inventors of the present application based on the findings currently obtained, but do not limit the technical scope of the present invention.

(1) Ferrite Fraction: 5% or Less

Ferrite generally has excellent workability but has a problem such as low strength. As a result, a large amount of ferrite leads to a decrease in yield ratio. Therefore, a ferrite fraction was set at 5% or less (5 volume % or less).

The ferrite fraction is preferably 3% or less, and more preferably 1% or less.

The ferrite fraction can be determined by observing with an optical microscope and measuring the white region by the point counting method. By such a method, it is possible to determine the ferrite fraction by an area ratio (area %). The value obtained by the area ratio may be directly used as the value of the volume ratio (volume %).

(2) Total Fraction of Tempered Martensite and Tempered Bainite: 60% or More

By setting the total fraction of tempered martensite and tempered bainite at 60% or more (60 volume % or more), it is possible to achieve both high strength and high hole expansion properties. The total fraction of tempered martensite and tempered bainite is preferably 70% or more.

It is possible to determine the amounts of tempered martensite and tempered bainite (total fraction) by performing SEM observation of a Nital-etched cross-section, measuring a fraction of MA (i.e., the total of retained austenite and martensite as quenched) and subtracting the above-mentioned ferrite fraction and MA fraction from the entire steel structure.

(3) Amount of Retained Austenite: 10% or More

The retained austenite causes the TRIP phenomenon of being transformed into martensite due to strain induced transformation during working such as press working, thus making it possible to obtain large elongation. Martensite thus formed has high hardness. Therefore, excellent strength-ductility balance can be obtained. By setting the amount of retained austenite at 10% or more (10 volume % or more), it is possible to realize TS×EL of 20,000 MPa % or more and excellent strength-ductility balance.

The amount of retained austenite is preferably 15% or more.

In the high-strength sheet according to the embodiment of the present invention, most of retained austenite exists in the form of MA. MA is abbreviation of a martensite-austenite constituent and is a composite (complex structure) of martensite and austenite.

It is possible to determine the amount of retained austenite by calculating a diffraction intensity ratio of ferrite (including tempered martensite and untempered martensite in X-ray diffraction) and austenite by X-ray diffraction, followed by calculation. As an X-ray source, Co-Kα ray can be used.

(4) Average Size of MA: 1.0 μm or Less

MA is a hard phase and the vicinity of matrix/hard phase interface acts as a void forming site during deformation. The larger the size of MA, the more strain concentration occurs at the matrix/hard phase interface, thus easily causing fracture from voids formed in the vicinity of the matrix/hard phase interface as a starting point.

Therefore, it is possible to improve the hole expansion ratio λ by decreasing the size of MA, especially the average size of MA to 1.0 μm or less, thereby suppressing fracture.

The average size of MA is preferably 0.8 μm or less.

It is possible to determine the average size of MA by observing a Nital-etched cross-section in three or more fields of view at a magnification of 3,000 times with SEM, drawing a straight line of 200 μm or more in arbitrary position in the micrograph, measuring the length of intercept where the straight line crosses MA, and calculating the average of the intercept lengths.

(5) Average Size of Retained Austenite: 1.0 μm or Less, and Retained Austenite having Size of 1.5 μm or More: Accounting for 2% or More of Total Amount of Retained Austenite

It has been found that excellent deep drawability can be obtained by setting an average size of retained austenite at 1.0 μm and setting a ratio (volume ratio) of retained austenite having a size of 1.5 μm or more to the entire retained austenite at 2% or more.

If the incoming stress of the flange portion is smaller than the tensile stress of the vertical wall portion formed during deep drawing, drawing is easily advanced, and thus satisfactory deep drawability can be obtained. Regarding the deformation behavior of the flange portion, since compressive stress is applied from the board surface direction and circumference, formation occurs in a state where isotropic compressive stress is applied. Meanwhile, martensitic transformation is accompanied by volume expansion, so that martensitic transformation hardly occurs under isotropic compressive stress. Therefore, strain induced martensitic transformation of retained austenite at the flange portion is suppressed to reduce work hardening.

As a result, deep drawability is improved. As the size of retained austenite increases, the greater effect of suppressing martensitic transformation is exhibited.

In order to increase the tensile stress of the vertical wall portion formed by deep drawing, there is a need to maintain a high work hardening rate during deformation. Unstable retained austenite that easily undergoes strain induced transformation under relatively low stress and stable retained austenite that does not undergo strain induced transformation unless under high stress are allowed to coexist to cause strain induced transformation over a wide stress range, thus making it possible to maintain a high work hardening rate during deformation. Therefore, a study was made to obtain a steel structure containing a predetermined amounts of each of unstable coarse retained austenite and stable fine retained austenite. Thus, the inventors of the present invention have found that a high work hardening rate is maintained during deformation by setting the average size of retained austenite at 1.0 μm and setting the ratio (volume ratio) of the amount of retained austenite having a size of 1.5 μm or more to the total amount of retained austenite at 2% or more, thus making it possible to obtain excellent deep drawability (LDR).

As mentioned above, when retained austenite undergoes strain induced transformation, the TRIP phenomenon occurs and high elongation can be obtained. Meanwhile, the martensitic structure formed by strain induced transformation is hard and acts as a starting point of fracture. Larger martensite structure easily acts as the starting point of fracture. It is also possible to obtain the effect of suppressing fracture by setting the average size of retained austenite at 1.0 μm or less to reduce the size of martensite formed by strain induced transformation.

It is possible to determine the average size of retained austenite and the ratio of the amount of retained austenite having a size of 1.5 μm or more to the total amount of retained austenite by creating a Phase map using the electron back scatter diffraction patterns (EBSD) method that is a crystal analysis method using SEM. An area of each austenite phase (retained austenite) is obtained from the obtained Phase map and a circle equivalent diameter (diameter) of each austenite phase is obtained from the area, and then an average of the obtained diameter is taken as the average size of retained austenite. It is possible to obtain the ratio of retained austenite having a size of 1.5 μm or more to the entire austenite by integrating the area of the austenite phase having an equivalent circle diameter of 1.5 μm or more to determine the ratio of austenite phase to the total area of the austenite phase. The thus obtained ratio of the retained austenite having a size of 1.5 μm or more to the entire austenite is the area ratio and is equivalent to the volume ratio.

(6) Scattering Intensity at q Value of 1 nm−1 in X-ray Small Angle Scattering: 1.0 cm−1 or Less

X-ray small angle scattering means that the size distribution of fine particles (e.g., cementite particles dispersed in a steel sheet) contained in the steel sheet can be obtained by irradiating the steel sheet with X-rays and measuring scattering of X-rays transmitted through the steel sheet. In the steel sheet according to the embodiment of the present invention, it is possible to determine the size distribution of cementite particles that are fine particles dispersed in tempered martensite by X-ray small angle scattering. Specifically, in X-ray small angle scattering, it is possible to analyze the size and the fraction of cementite particles using the q value and the scattering intensity.

The q value is an index of the size of particles (e.g., cementite particles) in the steel sheet. The “q value of 1 nm−1” corresponds to cementite particles having a particle size of about 1 nm. The scattering intensity is an index of the volume fraction of particles (e.g., cementite particles) in the steel sheet. The larger the scattering intensity, the larger the volume fraction of cementite becomes.

The scattering intensity at a certain q value semi-quantitatively indicates the volume fraction of cementite particles of the size corresponding to the q value. For example, the scattering intensity at the q value of 1 nm−1 semi-quantitatively indicates the volume fraction of fine cementite particles having a size of about 1 nm.

In other words, large scattering intensity at the q value of 1 nm−1 indicates large volume fraction of fine cementite particles having a size of about 1 nm. In the steel sheet in which “the scattering intensity at the q value of 1 nm−1 is 1.0 cm−1 or less”, it means that the volume fraction of fine cementite particles having a size of about 1 nm existing in the steel sheet is a predetermined value (the value corresponding to the scattering intensity of 1.0 cm−1) or less. As described later, it is considered that the steel sheet in which “the scattering intensity at the q value of 1 nm−1 is 1.0 cm−1 or less” is excellent in collision resistance properties since the volume fraction of cementite having a size of about 1 nm is suppressed to a low value.

In high-ductility steel containing retained γ, it is preferable that no cementite ideally exists in a state where carbon is concentrated in retained austenite. Fine cementite having a grain size of about 1 nm dispersed in the steel material hinders dislocation migration, thus enabling degradation of the deformability of the steel material. Therefore, in the steel material having a large volume fraction of cementite having a grain size of about 1 nm, fracture during deformation is promoted, thus enabling degradation of collision resistance properties.

In the steel sheet according to the embodiment of the present invention, by suppressing the volume fraction of fine cementite to a low value, more specifically, by setting the scattering intensity at the q value of 1 nm−1 at 1 cm−1 or less, fine carbide formed in laths of tempered martensite is reduced to enhance the deformability in martensite. Thus, fracture of the steel sheet upon collision is suppressed to improve collision resistance properties of the steel sheet.

X-ray small angle scattering was measured using a Nano-viewer, Mo tube manufactured by Rigaku Corporation. A 3 mmφ disk-shaped sample was cut out from the steel sheet and samples having a thickness of 20 μm were cut out from the vicinity of the thickness of ¼ and then used. Data at the q value of 0.1 to 10 nm−1 were collected. Among them, absolute intensity was determined for the q value of 1 nm−1.

(7) Other Steel Structure:

In the present description, steel structures other than the above-mentioned ferrite, tempered martensite, tempered bainite retained austenite and cementite are not specifically defined. However, pearlite, untempered bainite, untempered martensite and the like may exist, in addition to the steel structures such as ferrite. As long as the steel structure such as ferrite satisfies the above-mentioned structure conditions, the effects of the present invention are exhibited even if pearlite or the like exists in the steel.

2. Composition

The composition of the high-strength sheet according to the embodiment of the present invention will be described below. Main elements C, Si, Al, Mn, P and S will be described. Note that all percentages as unit with respect to the composition are by mass.

(1) C: 0.15 to 0.35%

Carbon (C) is an element indispensable for ensuring properties such as high strength-ductility balance (TS×EL balance) by increasing the amount of desired structure, especially retained γ. In order to effectively exhibit such effect, there is a need to add C in the amount of 0.15% or more. However, the amount of more than 0.35% is not suitable for welding. The amount is preferably 0.18% or more, and more preferably 0.20% or more. The amount is preferably 0.30% or less. If the C amount is 0.25% or less, welding can be easily performed.

(2) Total of Si and Al: 0.5 to 3.0%

Si and Al each have the effect of suppressing the precipitation of cementite, thus remaining retained austenite. In order to effectively exhibit such effect, there is a need to add Si and Al in the total amount of 0.5% or more. If the total amount of Si and Al exceeds 3.0%, the deformability of the steel is degraded, thus degrading TS×EL. The total amount is preferably 0.7% or more, and more preferably 1.0% or more. The total amount is preferably 2.5% or less.

Note that Al may be added in the amount enough to function as an deoxidizing element, i.e., less than 0.10% by mass. For the purpose of suppressing the formation of cementite to increase the amount of retained austenite, Al may be added in a larger amount of 0.7% by mass or more.

(3) Mn: 1.0 to 4.0%

Mn suppresses the formation of ferrite. In order to effectively exhibit such effect, there is a need to add Mn in the amount of 1.0% or more. If the amount exceeds 4.0%, MA becomes coarse, thus degrading hole expansion properties. The amount is preferably 1.5% or more, and more preferably 2.0% or more. The amount is preferably 3.5% or less.

(4) P: 0.05% or Less

P inevitably exists as an impurity element. If more than 0.05% of P exists, EL and A are degraded. Therefore, the content of P is set at 0.05% or less (including 0%). Preferably, the content is 0.03% or less (including 0%).

(5) S: 0.01% or Less

S inevitably exists as an impurity element. If more than 0.01% of S exists, sulfide-based inclusions such as MnS are formed and act as a starting point of cracking, thus degrading λ. Therefore, the content of S is set at 0.01% or less (including 0%). The content is preferably 0.005% or less (including 0%).

(6) Balance

In a preferred embodiment, the balance is composed of iron and inevitable impurities. It is permitted to mix, as inevitable impurities, trace elements (e.g., As, Sb, Sn, etc.) incorporated according to the conditions of raw materials, materials, manufacturing facilities and the like. There are elements whose content is preferably as small as possible, like P and S, that are therefore inevitable impurities in which the composition range is separately defined as mentioned above. Therefore, “inevitable impurities” constituting the balance as used herein means the concept excluding elements whose composition range is separately defined.

However, it is not limited to this embodiment. As long as properties of the high-strength steel sheet according to the embodiment of the present invention can be maintained, any other element may be further included.

3. Properties

As mentioned above, regarding the high-strength sheet according to the embodiment of the present invention, all of TS, YR, TS×EL, LDR, λ, collision resistance properties and SW cross tension are at a high level. These properties of the high-strength sheet according to the embodiment of the present invention will be described in detail below.

(1) Tensile Strength (TS)

The high-strength sheet has TS of 980 MPa or higher. Preferably, TS is 1,180 MPa or higher. If TS is lower than 980 MPa, excellent fracture properties can be more surely obtained, but it is not preferable since withstand load upon collision decreases.

(2) Yield Ratio (YR)

The high-strength sheet has an yield ratio of 0.75 or more. This makes it possible to realize a high yield strength combined with the above-mentioned high tensile strength and to use the final product obtained by working such as deep drawing under high stress. Preferably, the high-strength sheet has a yield ratio of 0.80 or more.

(3) The Product (TS×EL) of TS and Total Elongation (EL)

TS×EL is 20,000 MPa % or more. By having TS×EL of 20,000 MPa % or more, it is possible to obtain high-level strength-ductility balance that has both high strength and high ductility. Preferably, TS×EL is 23,000 MPa % or more.

(4) Deep Drawability (LDR)

LDR is an index used for evaluation of the deep drawability. D/d is referred to as a limiting drawing ratio (LDR), where d denotes a diameter of a cylinder obtained in cylindrical drawing and D denotes a maximum diameter of a disk-shaped steel sheet (blank) capable of obtaining a cylinder without causing fracture by one deep drawing process. More specifically, a disk-shaped sample having a thickness of 1.4 mm and various diameters is subjected to cylindrical deep drawing using a die having a punch diameter of 50 mm, a punch angle radius of 6 mm, a die diameter of 55.2 mm and a die angle radius of 8 mm. It is possible to determine LDR by finding a maximum sample diameter (maximum diameter D) among the sample diameters of the disc-shaped sample that was completely deep-drawn without causing fracture.

The high-strength sheet according to the embodiment of the present invention has LDR of 2.05 or more, and preferably 2.10 or more, and has excellent deep drawability.

(5) Hole Expansion Ratio (λ)

A hole expansion ratio λ is determined in accordance with Japan Iron and Steel Federation Standard JFS T1001. A punched hole having a diameter d0 (d0=10 mm) was formed in a test piece and a punch having a tip angle of 60° was pushed into this punched hole, and a diameter d of the punched hole at the time when the generated cracking penetrated the thickness of the test piece was measured, and then the hole expansion ratio is calculated by the following equation.


λ(%)={(d−d0)/d0}×100

The high-strength sheet according to the embodiment of the present invention has a hole expansion ratio λ of 20% or more, and preferably 30% or more. This makes it possible to obtain excellent workability such as press formability.

(6) Thickness Reduction Ratio in Tensile Test (R5 Tensile Thickness Reduction Ratio)

Using a test piece provided with an arcuate notch having a radius of 5 mm on a No. 5 test piece, a tensile test was performed at a deformation rate of 10 mm/min and the test piece was fractured. Then, the fracture surface was observed and the value (t1/t0) obtained by dividing a thickness t1 in a thickness direction of the fracture surface by an original thickness t0 was taken as a thickness reduction ratio.

The thickness reduction rate in this test is 50% or more, preferably 52% or more, and more preferably 55% or more. This makes it possible to obtain a steel sheet having excellent impact resistance properties since the steel sheet is hardly fractured even if it deforms greatly upon collision.

(7) Cross Tensile Strength of Spot Welding

Cross tensile strength of spot welding was evaluated in accordance with JIS Z 3137. Two 1.4 mm-thick steel sheets laid one upon another were used as the conditions of spot welding. Using a dome radius type electrode, spot welding was performed under a welding pressure of 4 kN by increasing a current by 0.5 kA in a range from 6 kA to 12 kA, and the current value (minimum current value) at which dust was generated during welding was examined. A cross joint spot-welded at a current that is 0.5 kA lower than the minimum current value was used as a test piece for measurement of a cross tensile strength. Samples having a cross tensile strength of 6 kN or higher were rated “Good”. The cross tensile strength is preferably 8 kN or higher, and more preferably 10 kN or higher.

When the cross tensile strength is 6 kN or higher, it is possible to obtain parts having high bonding strength during welding when automobile parts and the like are manufactured from the steel sheet.

4. Manufacturing Method

The method for manufacturing a high-strength sheet according to the embodiment of the present invention will be described below.

The inventors of the present application have found that the above-mentioned desired steel structure is attained by subjecting a rolled material with predetermined composition to a heat treatment (multi-step austempering treatment) mentioned later, thus obtaining a high-strength steel sheet having the above-mentioned desired properties.

Details will be described below.

FIG. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.

The rolled material to be subjected to the heat treatment is usually produced by cold-rolling after subjecting to hot-rolling. However, the process is not limited thereto, and the rolled material may be produced by any one of hot-rolling and cold-rolling. The conditions of hot-rolling and cold-rolling are not particularly limited.

(1) Austenitizing Treatment

As shown in [1] and [2] of FIG. 1, a rolled material is heated to a temperature of an Ac3 point or higher and heated for a predetermined heating time, thereby austenitizing the rolled material. The heating time at this heating temperature is, for example, 1 to 1,800 seconds. The upper limit of the heating temperature is preferably the Ac3 point or higher and the Ac3 point+100° C. or lower. This is because grain coarsening can be suppressed by setting at the temperature of the Ac3 point+100° C. or lower. The heating temperature is more preferably the Ac3 point+10° C. or higher and the Ac3 point+90° C. or lower, and still more preferably the Ac3 point+20° C. or higher and the Ac3 point+80° C. or lower. This is because the formation of ferrite can be more completely suppressed and grain coarsening can be more surely suppressed by more complete austenitizing.

Heating during austenitizing shown in [1] of FIG. 1 may be performed at an arbitrary heating rate, and the average heating rate is preferably 1° C./sec or more, and more preferably 20° C./sec.

(2) Cooling and Retention at Temperature in Range of 300° C. to 500° C.

After the austenitizing, cooling is performed, followed by retention at a temperature in a range of 300° C. to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more and less than 300 seconds, as shown in [5] of FIG. 1.

Cooling is performed at an average cooling rate of 15° C./sec or more and less than 200° C./sec between at least 650° C. and 500° C. This is because the formation of ferrite during cooling is suppressed by setting the average cooling rate at 15° C./sec or more. It is also possible to prevent the occurrence of excessive thermal strain due to rapid cooling by setting the cooling rate at less than 200° C./sec. Preferred example of such cooling includes cooling to a rapid cooling starting temperature of 650° C. or higher at relatively low average cooling rate of 0.1° C./sec or more and 10° C./sec or less, as shown in [3] of FIG. 1, followed by cooling from the rapid cooling starting temperature to a retention starting temperature of 500° C. or lower at an average cooling rate of 20° C./sec or more and less than 200° C./sec, as shown in [4] of FIG. 1.

Retention is performed at a temperature in a range of 300° C. to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more. In other words, the material is left to stand at a temperature in a range of 300° C. to 500° C. in a state where the cooling rate is 10° C./sec or less for 10 seconds or more. The state where the cooling rate is 10° C./sec or less also includes the case of holding at a substantially constant temperature (i.e., cooling rate is 0° C./sec), as shown in [5] of FIG. 1.

This retention enables partial formation of bainite. Since bainite has solid solubility limit of carbon that is lower than that of austenite, carbon exceeding the solid solubility limit is discharged from bainite, thus forming a region of austenite in which carbon is concentrated around bainite.

After cooling and reheating mentioned later, this region becomes somewhat coarse retained austenite. By forming this somewhat coarse retained austenite, it is possible to enhance the deep drawability as mentioned above.

If the retention temperature is higher than 500° C., since the carbon-concentrated region excessively increases, not only retained austenite but also MA becomes coarse, thus decreasing the hole spreading ratio. Meanwhile, if the retention temperature is lower than 300° C., the carbon-concentrated region decreases and the amount of coarse retained austenite becomes insufficient, thus degrading the deep drawability.

If the retention time is less than 10 seconds, the area of the carbon-concentrated region decreases and the amount of coarse retained austenite becomes insufficient, thus degrading the deep drawability. Meanwhile, if the retention time is 300 seconds or more, since the carbon-concentrated region excessively increases, not only retained austenite but also MA becomes coarse, thus decreasing the hole expansion ratio.

If the cooling rate during retention is more than 10° C./sec, since sufficient bainite transformation does not occur, sufficient carbon-concentrated region is not formed, leading to insufficient amount of coarse retained austenite.

Therefore, retention is performed at a temperature in a range of 300° C. to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more. Retention is preferably performed at a temperature in a range of 320 to 480° C. at a cooling rate of 8° C./sec or less for 10 seconds or more and, during the retention, holding is preferably performed at a constant temperature for 3 to 80 seconds.

Retention is more preferably performed at a temperature in a range of 340 to 460° C. at a cooling rate of 3° C./sec or less for 10 seconds or more and, during the retention, holding is performed a constant temperature for 5 to 60 seconds.

(3) Cooling to Cooling Stopping Temperature between 100° C. or Higher and Lower than 300° C.

After the above-mentioned retention, as shown in [6] of FIG. 1, cooling is performed from a second cooling starting temperature of 300° C. or higher to a cooling stopping temperature of 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more. In one of preferred embodiments, as shown in [6] of FIG. 1, the above-mentioned retention end temperature (e.g., holding temperature shown in [5] of FIG. 1) is taken as the second cooling starting temperature.

This cooling causes martensitic transformation while leaving the above-mentioned carbon-concentrated region as austenite. By controlling the cooling stopping temperature at a temperature in a range of 100° C. or higher and lower than 300° C., final amount of retained austenite is controlled by adjusting the amount of austenite remaining without being transformed into martensite.

If the cooling rate is less than 10° C./sec, the carbon-concentrated region expands more than necessarily during cooling and MA becomes coarse, thus decreasing the hole spreading ratio. If the cooling stopping temperature is lower than 100° C., the amount of retained austenite becomes insufficient. As a result, TS increases but EL decreases, leading to insufficient TS×EL balance.

If the cooling stopping temperature is 300° C. or higher, coarse unmodified austenite increases and remains even after the subsequent cooling. Finally, the size of MA becomes coarse, thus decreasing the hole expansion ratio λ.

The cooling rate is preferably 15° C./° C. or higher, and the cooling stopping temperature is preferably 120° C. or higher and 280° C. or lower. The cooling rate is more preferably. 20° C./sec or more, and the cooling stopping temperature more preferably 140° C. or higher and 260° C. or lower.

As shown in [7] of FIG. 1, holding may be performed at the cooling stopping temperature. In the case of holding, the holding time is preferably 1 to 600 seconds. Even if the holding time increases, there is almost no influence on properties. However, the holding time of more than 600 seconds degrades the productivity.

(4) Reheating to Temperature in Range of 300° C. to 500° C.

As shown in [8] of FIG. 1, heating is performed from the above cooling stopping temperature to a reheating temperature in a range of 300° C. to 500° C. at a reheating rate of 30° C./sec or more. Rapid heating enables a decrease in retention time in a temperature range where precipitation and growth of carbide are promoted, thus making it possible to suppress the formation of fine carbide. The reheating rate is preferably 60° C./sec or more, and more preferably 70° C./sec.

Such rapid heating can be achieved by a method such as high-frequency heating or electric heating.

After reaching the reheating temperature, as shown in [9] of FIG. 1, holding is performed at the same temperature. At that time, it is preferred that a tempering parameter P represented by the following equation (1) is set at 10,000 or more and 14,500 or less and the holding time is set at 1 to 150 seconds. The tempering parameter P of the steel sheet of the present embodiment is represented by the following equation (1):


P=T(K)×(20+log(t/3600)  (1)

where T is a tempering temperature (K) and t is a holding time (seconds).

During reheating, redistribution of carbon that is supersaturatedly sold-soluted in martensite occurs. Specifically, two phenomena, i.e. carbon diffusion from martensite to austenite and precipitation of carbide (cementite) in martensite laths. Among two phenomena, the precipitation of carbide easily occurs when holding is performed at low temperature for a long time. Even in the case of holding at high temperature, carbide is precipitated when the heating rate is low or the holding time is too long. Meanwhile, since carbon diffusion from martensite to austenite strongly depends on the diffusion rate, carbon diffusion can be sufficiently performed by a heat treatment at high temperature in a short time.

Particles of cementite existing in martensite easily act as a starting point of collision fracture and can degrade collision resistance properties. Therefore, in the case of reheating, it is desired that a reheating treatment is performed to promote carbon diffusion from martensite to austenite while suppressing the precipitation of carbide (cementite) in martensite laths. Thus, it is effective to perform rapid heating and a heat treatment at high temperature in a short time.

In order to obtain desired tensile strength by causing sufficient carbon diffusion, there is a need to control the tempering parameter P as a factor of a combination of temperature and time within a given range.

When the tempering parameter P is less than 10,000, carbon diffusion from martensite to austenite does not sufficiently occur and austenite becomes unstable, thus failing to ensure the amount of retained austenite, leading to insufficient TS×EL balance. If the tempering parameter P is more than 14,500, the formation of carbide cannot be prevented even by a short-time treatment, thus failing to ensure the amount of retained austenite, leading to degradation of TS×EL balance. Even if the tempering parameter is appropriate, carbide is formed in martensitic laths if the heating rate is too low and heating time is too long, so that crack propagation easily occurs during collision deformation, thus degrading collision resistance properties. The amount of carbide in martensite laths can be determined from the scattering intensity of X-ray small angle scattering.

If the reheating temperature is lower than 300° C., diffusion of carbon becomes insufficient, thus failing to obtain sufficient amount of retained austenite, leading to degradation of TS×EL. If the reheating temperature is higher than 500° C., retained austenite is decomposed into cementite and ferrite, thus failing to ensure properties because of insufficient retained austenite.

If holding is not performed or the holding time is less than 1 second, carbon diffusion may be insufficient, similarly. Therefore, it is preferred to hold at a reheating temperature for 1 second or more. If the holding time is more than 150 seconds, carbon may precipitate as cementite, similarly. Therefore, the holding time is preferably 150 seconds or less.

The reheating temperature is preferably 320 to 480° C., and more preferably 340 to 460° C.

The tempering parameter P is preferably 10,500 to 14,500, and the holding time at this time is preferably 1 to 150 seconds. The tempering parameter P is more preferably 11,000 to 14,000, and the holding time at this time is preferably 1 to 100 seconds, and more preferably 1 to 60 seconds.

After reheating, as shown in [10] of FIG. 1, cooling may be performed to the temperature of 200° C. or lower, for example, room temperature. The average cooling rate to 200° C. or lower is preferably 10° C./sec.

The high-strength sheet according to the embodiment of the present invention can be obtained by the above-mentioned heat treatment.

There is a possibility for person with ordinary skill in the art who came into contact with the method of manufacturing a high-strength steel sheet according to the embodiment of the present invention described above to obtain the high strength steel sheet according to the embodiment of the present invention by trial and error, using a manufacturing method different from the above-mentioned method.

Examples 1. FABRICATION OF SAMPLES

After producing a cast material with the chemical composition shown in Table 1 was produced by vacuum melting, this cast material was hot-forged to form a steel sheet having a thickness of 30 mm and then hot-rolled. In Table 1, an Ac3 point calculated from the composition was also shown.

Although the conditions of hot-rolling do not have a substantial influence on the final structure and properties of the present patent, a steel sheet having a thickness of 2.5 mm was produced by multistage rolling after heating to 1,200° C. At this time, the end temperature of hot-rolling was set at 880° C. After cooling to 600° C. at 30° C./sec, cooling was stopped and the steel sheet was inserted into a furnace heated to 600° C., held for 30 minutes and then furnace-cooled to obtain a hot-rolled steel sheet.

The hot-rolled steel sheet was subjected to pickling to remove the scale on the surface, and then cold-rolled to reduce the thickness to 1.4 mm. This cold rolled sheet was subjected to a heat treatment to obtain samples. The heat treatment conditions are shown in Table 2. The number in parentheses, for example, [2] in Table 2 corresponds to the process of the same number in parentheses in FIG. 1. In Table 2, samples Nos. 1, 4, 7 and 26 are samples that were not retained at a temperature in a range of 300 to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more in the step corresponding to [5] of FIG. 1. Especially, samples Nos. 1 and 26 are samples (samples in which the steps corresponding to [5] and [6] in FIG. 1 were skipped) that were immediately cooled to 200° C. after starting rapid cooling at 700° C. Sample No. 9 is sample (sample in which the steps corresponding to [6] to [8] in FIG. 1 were skipped) that was cooled to a reheating temperature instead of cooling to a cooling stopping temperature between 100° C. or higher and lower than 300° C., followed by holding at the same temperature.

Reheating corresponding to [8] was performed by an electric heating method.

In Table 1 to Table 4, the numerical value with an asterisk (*) indicates that it deviates from the range of the embodiment of the present invention.

TABLE 1 Composition C Si Mn P S Al Si + Al Ac3 Steel % by % by % by % by % by % by % by point No. mass mass mass mass mass mass mass ° C. a 0.27 1.28 2.01 0.010 0.001 0.04 1.32 818 b 0.22 2.10 1.80 0.010 0.002 0.03 2.13 868 c *0.10 1.42 2.51 0.007 0.003 0.03 1.45 847 d 0.19 1.28 *5.20 0.012 0.003 0.02 1.30 732 e 0.21 1.54 *0.63 0.012 0.003 0.02 1.56 876 f 0.27 0.20 2.19 0.013 0.002 0.03 *0.23 761 g *0.48 1.51 1.67 0.007 0.002 0.03 1.54 800 h 0.29 3.20 1.90 0.006 0.001 0.04 *3.24 904 i 0.24 1.32 2.49 0.012 0.002 0.04 1.36 813 j 0.23 1.24 1.69 0.010 0.002 0.04 1.28 835 k 0.21 1.31 2.41 0.006 0.002 0.03 1.34 816 l 0.25 1.39 1.99 0.006 0.002 0.04 1.43 829 m 0.25 1.03 1.75 0.008 0.003 0.02 1.05 812 n 0.25 1.61 2.46 0.005 0.002 0.03 1.64 819 o 0.27 0.84 2.46 0.008 0.002 0.25 1.09 870 p 0.28 1.38 2.09 0.008 0.003 0.03 1.41 814 q 0.28 0.97 2.37 0.013 0.002 0.03 1.00 788

TABLE 2 Heat treatment conditions [3] [4] [6] [1] [1] [2] Slow Rapid cooling [4] [5] [5] [6] Cooling [7] [8] [8] [9] [10] Heating Heating Heating cooling starting Cooling Holding Holding Cooling stopping Holding Reheating Reheating Holding Cooling rate temperature time rate temperature rate temperature time rate temperature time rate temperature time rate No. Steel No. ° C./sec ° C. Sec ° C./scc ° C. ° C./sec ° C. Sec ° C./sec ° C. Sec ° C./sec ° C. Sec ° C./sec Parameter 1 a 10 850 120 10 700 28 *— *— *— 200 50 100 400 *300 10 12,734 2 a 10 850 120 10 700 28 400 *300 30 200 50 100 440 20 10 12,652 3 a 10 850 120 10 700 28 400 50 *1 200 50 100 400 *300 10 12,734 4 a 10 850 120 10 700 28 400 *3 30 200 50 100 440 20 10 12,652 5 a 10 850 120 10 700 28 *550 50 30 275 50 100 440 20 10 12,652 6 a 10 850 120 10 700 28 *250 50 30 125 50 100 440 20 10 12,652 7 a 10 850 120 10 700 28 400 *0 30 *350 50 100 440 20 10 12,652 8 a 10 *780 120 10 700 28 400 50 30 150 50 100 440 20 10 12,652 9 a 10 850 120 10 700 28 400 50 30 *400 — *— 400 *300 10 12,734 10 a 10 850 120 10 700 28 400 50 30 *20 50 100 440 20 10 12,652 11 a 10 850 120 — 850 28 400 50 30 200 50 *30 440 20 10 12,652 12 a 10 850 120 10 *580 28 400 50 30 200 50 90 400 50 10 12,210 13 a 10 850 120 10 700 28 400 50 30 200 50 100 350 50 10 11,303 14 a 10 850 120 10 700 *8 400 50 30 200 50 100 350 20 10 11,055 15 a 10 850 120 10 700 28 400 50 30 200 50 100 440 20 10 12,652 16 a 10 830 120 10 700 28 400 50 30 200 50 100 400 *300 10 12,734 17 a 10 850 120 10 700 28 400 50 30 200 50 *15 400 20 10 11,942 18 a 10 870 120 10 700 28 400 50 30 200 50 100 440 20 10 12,652 19 a 10 850 120 10 700 28 400 50 30 200 50 100 *550 20 10 *14,604 20 a 10 850 120 10 700 28 400 50 30 200 50 100 *250 20 10 *9,280 21 b 10 900 120 10 700 28 400 50 30 200 50 100 400 20 10 11,942 22 c 10 900 120 10 700 28 400 50 30 200 50 100 400 20 10 11,942 23 d 10 850 120 10 700 28 400 50 30 200 50 100 400 20 10 11,942 24 e 10 900 120 10 700 28 400 50 30 200 50 100 400 20 10 11,942 25 f 10 850 120 10 700 28 400 50 30 200 50 100 400 20 10 11,942 26 g 10 850 120 10 700 28 *— *— *— 200 50 100 400 20 10 11,942 27 h 10 940 120 10 700 28 400 50 30 200 50 100 400 20 10 11942 28 i 10 850 120 10 700 28 400 50 30 200 50 100 400 20 10 11942 29 j 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 30 k 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 31 l 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 32 m 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 33 n 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 34 o 10 920 120 — 920 28 400 50 30 200 50 100 400 20 10 11942 35 p 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942 36 q 10 850 120 — 850 28 400 50 30 200 50 100 400 20 10 11942

2. STEEL STRUCTURE

With respect to each sample, the ferrite fraction, the total fraction of tempered martensite and tempered bainite (mentioned as “tempered M/B” in Table 3), the amount of retained (amount of retained γ), the average size of MA, the average size of retained austenite (average size of retained γ), the ratio of retained austenite having a size of 1.5 μm or more to the entire austenite (mentioned as “ratio of retained γ having a size of 1.5 μm or more” in Table 3), and the scattering intensity at the q value of 1 nm−1 in X-ray small angle scattering were determined by the above-mentioned methods. In the measurement of the amount of retained austenite, a two-dimensional microfocused X-ray diffraction apparatus (RINT-RAPID II) manufactured by Rigaku Corporation was used. The results are shown in Table 3.

In this Example, the steel structure (balance structure) other than the steel structure shown in Table 3 is untempered martensite in samples excluding sample No. 9, or untempered bainite in sample No. 9.

TABLE 3 Steel structure Scattering Tempered Amount of Average size of Average size of intensity at Ferrite M/B retained γ MA retained γ Amount of retained γ having q value of 1 nm−1 No. Steel No. % % % μm μm size of 1.5 μm or more cm−1 1 a 0 70 16.4 0.65 0.78 *0.74 *2.43 2 a 0 69 16.7 *1.30 0.87 3.88 0.71 3 a 0 72 17.8 *14.3 0.82 4.22 *2.42 4 a 0 71 16.3 0.53 0.83 *1.21 0.73 5 a 0 66 18.0 *1.48 0.93 4.11 0.70 6 a 0 76 17.8 0.68 0.76 *0.66 0.73 7 a 0 *55 16.9 *1.50 *1.42 3.30 0.74 8 a *26.8 *47 17.0 0.60 0.79 3.78 0.68 9 a 0 *0 16.6 *1.65 *1.69 5.37 0.87 10 a 0 79 *4.5 0.50 0.57 *0.90 0.75 11 a 0 70 16.4 0.50 0.77 3.71 *12.7 12 a *20 *51 17.9 0.73 0.79 4.66 0.95 13 a 0 70 17.5 0.65 0.81 2.98 0.89 14 a *19 *52 18.1 0.65 0.80 2.94 0.69 15 a 0 70 16.8 0.50 0.84 2.78 0.69 16 a 0 70 17.6 0.68 0.82 3.50 *2.39 17 a 0 70 16.3 0.63 0.85 3.67 *14.2 18 a 0 70 18.1 0.58 0.79 3.93 0.74 19 a 0 73 *7.0 0.68 0.81 *0.80 *2.01 20 a 0 67 *6.3 0.53 0.78 *0.72 0.70 21 b 0 73 13.6 0.50 0.81 3.11 0.73 22 c 0 78 *6.3 0.53 0.81 *0.66 0.67 23 d 0 60 *8.6 0.53 0.84 3.59 0.67 24 e *26 *50 14.6 0.53 0.79 3.66 0.67 25 f 0 *51 *6.7 *1.33 0.86 *0.74 0.74 26 g 0 61 30.8 0.63 0.84 3.98 0.69 27 h 0 70 18.5 0.53 0.82 3.03 0.70 28 i 0 73 15.7 0.55 0.81 4.36 0.71 29 j 0 73 13.8 0.50 0.78 3.21 0.75 30 k 0 72 14.1 0.70 0.79 2.81 0.70 31 l 0 71 16.5 0.53 0.83 2.80 0.67 32 m 0 72 14.9 0.53 0.79 3.51 0.70 33 n 0 72 16.6 0.60 0.77 2.91 0.69 34 o 0 70 163 0.70 0.85 3.14 0.68 35 p 0 69 17.5 0.55 0.77 4.42 0.73 36 q 0 69 17.1 0.55 0.84 3.78 0.75

3. MECHANICAL PROPERTIES

With respect to the resulting samples, YS, TS and EL were measured using a tensile tester, and YR and TS×EL were calculated. By the above-mentioned methods, the hole expansion ratio λ, the deep drawability LDR, the cross tensile strength (SW cross tension) of the spot welded portion, and the R5 tensile thickness reduction ratio were determined. The results are shown in Table 4.

TABLE 4 Properties SW cross R5 tensile thickness YS TS EL TS × EL λ tension reduction ratio No. Steel No. MPa MPa YR % MPa % % LDR kN % 1 a 1,009 1,226 0.82 18.5 22,676 39 *1.8 14.3 *31.8 2 a 1,036 1,257 0.82 18.3 22,996 *14 2.30 14.8 55.9 3 a 1,009 1,218 0.83 19.3 23,449 *13 2.33 14 *31.5 4 a 989 1,202 0.82 18.6 22,367 31 *1.95 15.2 55.6 5 a 1,081 1,325 0.82 17.9 23,757 *14 2.32 1.52 55 6 a 955 1,140 0.84 20.8 23,718 45 *1.82 15.4 55.3 7 a 1,131 1,432 0.79 16.1 23,105 *14 2.51 15 54.9 8 a 619 *961 *0.64 24.1 23,160 21 2.29 15.1 58 9 a 601 *926 *0.65 22.4 20,742 *10 2.75 14.4 56.1 10 a 911 1,114 0.82 12.7 *14,105 42 *1.89 13.7 54 11 a 1,028 1,249 0.82 18 22,474 26 2.28 13.7 *41.6 12 a 699 *955 *0.70 22 21,010 21 2.32 13.6 51.7 13 a 1,113 1,346 0.83 17.3 23,247 45 2.24 13.7 54.8 14 a 646 *974 *0.66 22 21,428 28 2.25 13.9 55.4 15 a 1,042 1,264 0.82 18.1 22,873 40 2.25 14 56.5 16 a 1,006 1,221 0.82 19.3 23,584 34 2.29 15.5 *28.7 17 a 1,068 1,299 0.82 17.4 22,632 21 2.28 14.5 *42 18 a 1,019 1,232 0.83 19.1 23,575 21 2.30 15.4 54.7 19 a 819 1,014 0.81 15.8 *16,039 34 *2.03 19 *33.8 20 a 1,270 1,593 0.80 9.8 *15,551 28 *2.03 12.8 53.8 21 b 1,040 1,267 0.82 16.2 20,515 31 2.26 16.8 58.1 22 c 967 1,184 0.82 13.2 *15,594 45 *2.03 23.4 57.3 23 d 1,173 1,481 0.79 11.6 *17,161 41 2.36 18.2 54.8 24 e 624 *964 *0.65 22 21,208 28 2.30 17.8 58.3 25 f 1,204 1,584 0.76 10 *15,881 *15 *2.03 15.4 56 26 g 1,205 1,446 0.83 22.6 32,731 30 2.30 *3.4 58.1 27 h 1,105 1,335 0.83 13.7 *18,346 34 2.26 13 55.8 28 i 1,067 1,290 0.83 17 21,951 44 2.33 17.2 57.4 29 j 1,030 1,248 0.83 16.7 20,804 28 2.25 16.5 53.9 30 k 1,058 1,284 0.82 16.3 20,980 43 2.24 17.1 58.5 31 l 1,044 1,268 0.82 17.9 22,658 38 2.24 16 54.4 32 m 1,035 1,256 0.82 17.1 21,478 21 2.27 14.8 56.9 33 n 1,051 1,273 0.83 18 22,889 41 2.25 15.7 54.1 34 o 1,093 1,327 0.82 16.8 22,337 36 2.26 14 57.6 35 p 1,094 1,328 0.82 17.7 23,527 25 2.33 14.4 57.4 36 q 1,067 1,295 0.82 17.9 23,192 31 2.28 14.7 55.9

The results of Table 4 will be considered. Samples Nos. 13, 15, 18, 21 and 28 to 36 are Examples that satisfy all requirements (composition, manufacturing conditions and steel structure) defined in the embodiment of the present invention. All of these samples achieve the tensile strength (TS) of 980 MPa or higher, the yield ratio (YR) of 0.75 or more, TS×EL of 20,000 MPa % or higher, LDR of 2.05 or more, the hole expansion ratio (A) of 20% or more, the SW cross tension of 6 kN or higher, and the R5 tensile thickness reduction ratio (RA) of 50% or more.

To the contrary, sample No. 1 exhibited insufficient amount of retained austenite having a size of 1.5 μm or more, thus failing to obtain sufficient deep drawability since retention was not performed at a temperature in a range of 300° C. to 500° C. after austenitizing. Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 2 exhibited excessive average size of MA, thus failing to obtain sufficient hole expansion ratio since the holding temperature [5] was as long as 300 seconds.

Sample No. 3 exhibited excessive average size of MA, thus failing to obtain sufficient hole expansion ratio since the cooling rate [6] was as low as 1° C./sec. Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 4 exhibited insufficient amount of retained austenite having a size of 1.5 μm or more, thus failing to obtain sufficient deep drawability since the holding time [5] was as short as 3 seconds.

Sample No. 5 exhibited excessive average size of MA, thus failing to obtain sufficient hole expansion ratio and sufficient deep drawability since the holding temperature [5] was as high as 550° C.

Sample No. 6 exhibited insufficient amount of retained austenite having a size of 1.5 μm or more, thus failing to obtain sufficient deep drawability since the holding temperature [5] was as low as 250° C.

Sample No. 7 exhibited insufficient total amount of tempered martensite and tempered bainite, and excessive average size of MA and excessive average size of retained austenite since the cooling stopping temperature [6] was as high as 350° C. As a result, sufficient hole expansion ratio and deep drawability could not be obtained.

Sample No. 8 exhibited excessive amount of ferrite and insufficient total amount of tempered martensite and tempered bainite, thus failing to obtain sufficient tensile strength and yield ratio since the heating temperature [1] was as low as 780° C.

Sample No. 9 did not form martensite and bainite, and exhibited excessive average size of MA and excessive average size of retained austenite since the cooling stopping temperature [6] was as high as 400° C. As a result, sufficient tensile strength and yield ratio could not be obtained. Furthermore, a small amount of carbide was formed because of holding at the same temperature for 300 seconds (holding time [9]). As a result, λ decreased.

Sample No. 10 exhibited reduced amount of retained γ and insufficient amount of retained austenite having a size of 1.5 μm or more since the cooling stopping temperature [5] was as low as 20° C. As a result, sufficient value of TS×EL and sufficient deep drawability could not be obtained.

Sample No. 11 precipitated carbide (cementite) since the reheating rate [8] was as low as 30° C./sec. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 12 exhibited excessive amount of ferrite and insufficient total amount of tempered martensite and tempered bainite, thus failing to obtain sufficient tensile strength and yield ratio since the rapid cooling starting temperature [4] was as low as 580° C.

Sample No. 14 exhibited excessive amount of ferrite, insufficient total amount of tempered martensite and tempered bainite, and excessive average size of MA since the cooling rate [4] was as low as 8° C./sec. As a result, sufficient tensile strength and yield ratio could not be obtained.

Sample No. 16 precipitated carbide (cementite) since the holding time [9] was as long as 300 seconds. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 17 precipitated carbide (cementite) since the reheating rate [8] was as low as 15° C./sec. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 19 exhibited the parameter increased to 14,604 since the reheating temperature [7] was as high as 550° C. Therefore, the amount of retained γ decreased and the amount of retained austenite having a size of 1.5 μm or more was insufficient. As a result, TS×EL and deep drawability were degraded. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 20 exhibited the parameter decreased to 9,280 since the reheating temperature [8] was as low as 250° C. Therefore, the sample exhibited insufficient carbon diffusion, decreased amount of retained γ, and insufficient amount of retained austenite having a size of 1.5 μm or more. As a result, TS×EL and deep drawability were degraded.

Sample No. 22 exhibited small C amount, insufficient amount of retained austenite and insufficient amount of retained austenite having a size of 1.5 μm or more, thus failing to obtain sufficient TS×EL and deep drawability.

Sample No. 23 exhibited large Mn amount and insufficient amount of retained austenite, thus failing to obtain sufficient TS×EL.

Sample No. 24 exhibited small Mn amount, excessive amount of ferrite, and insufficient total amount of tempered martensite and tempered bainite. As a result, sufficient tensile strength and yield ratio could not be obtained.

Sample No. 25 exhibited small amount of Si+Al, insufficient total amount of tempered martensite and tempered bainite, small amount of retained austenite, excessive average size of MA, and excessive average size of retained austenite. As a result, sufficient TS×EL, hole expansion ratio and deep drawability could not be obtained.

Sample No. 26 exhibited excessive C amount, thus failing to obtain sufficient SW cross tensile strength since retention was not performed at a temperature in a range of 300° C. to 500° C. after austenitizing. Sample No. 27 exhibited excessive amount of Si+Al, thus failing to obtain sufficient TS×EL.

4. CONCLUSION

In this way, it could be confirmed that, regarding the steel sheet that satisfies the composition and the steel structure defined in the embodiment of the present invention, all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), LDR, hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test, and cross tensile strength of the spot welded portion are at a high level.

It could be also confirmed that the manufacturing method according to the embodiment of the present invention enables the production of the steel sheet that satisfies the composition and the steel structure defined in the embodiment of the present invention.

This application claims priority based on Japanese Patent Application 2016-153107 filed on Aug. 3, 2016, the disclosure of which is incorporated by reference herein.

Claims

1. A high-strength sheet, comprising:

Fe,
C: 0.15% by mass to 0.35% by mass,
a total of Si and Al: 0.5% by mass to 3.0% by mass,
Mn: 1.0% by mass to 4.0% by mass,
P: 0.05% by mass or less, and
S: 0.01% by mass or less,
wherein the high-strength sheet comprises a steel structure wherein:
a ferrite fraction is 5% or less,
a total fraction of tempered martensite and tempered bainite is 60% or more,
an amount of retained austenite is 10% or more,
a martensite-austenite constituent has an average size of 1.0 μm or less,
the retained austenite has an average size of 1.0 μm or less,
retained austenite having a size of 1.5 μm or more accounts for 2% or more of a total amount of the retained austenite, and
a scattering intensity at a q value of 1 nm−1 in X-ray small angle scattering is 1.0 cm−1 or less.

2. The high-strength sheet according to claim 1, comprising:

0.30% by mass or less of C.

3. The high-strength sheet according to claim 1, comprising:

less than 0.10% by mass of Al.

4. A method for manufacturing a high-strength sheet, the method comprising:

preparing a rolled material comprising: Fe, C: 0.15% by mass to 0.35% by mass, a total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less,
heating the rolled material which has an AC3 point, to a temperature of the Ac3 point or higher, thereby austenitizing the rolled material;
after the austenitizing, cooling the rolled material between 650° C. and 500° C. at an average cooling rate of 15° C./sec or more and less than 200° C./sec, followed by retention at a temperature in a range of 300° C. to 500° C. at a cooling rate of 10° C./sec or less for 10 seconds or more and less than 300 seconds;
after the retention, cooling the rolled material from a temperature of 300° C. or higher to a cooling stopping temperature between 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more;
heating the material from the cooling stopping temperature to a reheating temperature in a range of 300° C. to 500° C. at an average heating rate of 30° C./sec or more;
holding at the reheating temperature T for a holding time of 1 to 150 seconds so as to satisfy a tempering parameter P of 10,000 to 14,500; and
after the holding at the reheating temperature, cooling from the reheating temperature T to 200° C. at an average cooling rate of 10° C./sec or more, wherein P=T×(20÷log(t/3600))  (1)
where T: reheating temperature (K) and t: holding time (seconds).

5. The method according to claim 4, wherein the retention comprises holding at a constant temperature in a range of 300° C. to 500° C.

6. The method according to claim 4, wherein the tempering parameter P is from 11,000 to 14,000 and the holding time t is from 1 to 150 seconds.

Patent History
Publication number: 20190218640
Type: Application
Filed: Jul 21, 2017
Publication Date: Jul 18, 2019
Applicant: Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) (Kobe-shi)
Inventors: Elijah KAKIUCHI (Kobe-shi), Toshio MURAKAMI (Kobe-shi), Shigeo OTANI (Kobe-shi), Yuichi FUTAMURA (Kakogawa-shi), Tadao MURATA (Kakogawa-shi)
Application Number: 16/321,544
Classifications
International Classification: C21D 9/46 (20060101); C21D 6/00 (20060101); C21D 1/34 (20060101); C22C 38/06 (20060101); C22C 38/04 (20060101); C22C 38/00 (20060101); C22C 38/02 (20060101);