JP5729829B2 - High-strength steel sheet for warm forming excellent in ductility and deep drawability in warm and its manufacturing method - Google Patents

Description

The present invention relates to a high-strength steel sheet suitable for automobile members, and in particular, to a high-strength steel sheet for warm forming that is excellent in press formability, particularly, warm ductility and deep drawability, and a method for producing the same.

  For example, steel sheets used for automobile parts and the like are required to have high strength for the purpose of collision safety and fuel efficiency reduction by reducing the weight of the vehicle body, and excellent formability is required for processing into parts with complex shapes. The For this reason, in high-strength steel sheets, further improvements in ductility and deep drawability are particularly desired.

  In response to such needs, many high-strength steel sheets with improved ductility have been proposed in terms of materials based on various structural control concepts. Among them, as a steel plate having both high strength and ductility, a TRIP type (Transformation Induced Plasticity) steel plate has attracted attention. A TRIP type steel sheet is a steel sheet in which an austenite structure remains in the steel, and excellent ductility can be obtained by increasing the work hardening rate by inducing transformation of retained austenite into martensite due to stress due to work deformation.

  For example, in Patent Document 1, C: 0.16% to 0.72%, Si: 3.0% or less, Mn: 0.5% to 3.0%, P: 0.1% or less, S: 0.07% or less, Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, and the area ratio of martensite to the entire steel sheet structure is 10% or more. 90% or less, the amount of retained austenite is 5% or more and 50% or less, the area ratio of bainitic ferrite in the upper bainite is 5% or more, and the area ratio of the martensite to the entire steel sheet structure is the retained austenite. The total amount of bainitic ferrite in the upper bainite and the area ratio with respect to the entire steel sheet structure is 60% or more, and the area ratio of polygonal ferrite with respect to the entire steel sheet structure is 10% or less ( Percent included) is satisfied, and TRIP-type bainitic ferrite steel sheet the mean C amount is 2.00% 0.70% or more of the retained austenite is disclosed.

  In Patent Document 2, C: 0.10 to 0.20%, Si: 0.8 to 2.5%, Mn: 1.5 to 2.5%, Al: 0.01 to 0.10%, P: containing less than 0.1% (excluding 0%), S: containing less than 0.002% (excluding 0%), balance: satisfying iron and inevitable impurities, and the structure is at least bay Nitritic ferrite and retained austenite are included, and the area ratio to the whole structure is bainitic ferrite: 70% or more, retained austenite: 2 to 20%, polygonal ferrite and / or quasi-polygonal ferrite: 15% or less TRIP type bainitic ferrite steel sheet satisfying (excluding 0%) and having a proportion of retained austenite having an average particle size of 5 μm or less in the retained austenite of 60% or more is disclosed. That.

  The high-tensile steel sheets disclosed in Patent Documents 1 and 2 above are TRIP-type bainitic ferrite steel sheets that have both high strength and ductility, but are related to formability at room temperature, and the warm formability is considered. It has not been.

  From the molding surface, a warm or hot molding method has been proposed, and Non-Patent Document 1 discloses a technique related to warm molding. Among these, it has been shown that when a TRIP type steel sheet having a tensile strength of 780 MPa is formed at 80 to 230 ° C., the ductility is improved, but in a steel sheet having a tensile strength of 980 MPa or more, warm ductility is improved. It is not clear what kind of organization it is desirable to do.

JP 2010-65272 A JP 2008-7854 A

Press Technology, Vol. 42, No. 12, November 2004, p. 34-38

Accordingly, an object of the present invention is to provide a high-strength steel sheet for warm forming that has improved ductility and deep drawability, particularly ductility and deep drawability in warm (100 to 500 ° C.) with respect to a high-strength TRIP type bainitic ferrite steel sheet, and It is in providing the manufacturing method.

The invention described in claim 1
% By mass (hereinafter the same for chemical components)
C: 0.10 to 0.30%,
Si: more than 1.0% to 3.0%,
Mn: 1.0 to 3.0%
P: 0.10% or less (including 0%),
S: 0.010% or less (including 0%),
N: 0.0020 to 0.0300% or less,
Al: 0.0010 to 0.1%
And the remainder has a component composition consisting of iron and inevitable impurities,
Microstructure, all tissues sum of bainitic ferrite and bainite at an area ratio with respect to 65% or more,
Residual austenite: 5% or more,
Total of martensite and retained austenite: 35% or less,
Polygonal ferrite: 10% or less (including 0%)
Other than the above, the rest consists of 5% or less (including 0%),
The carbon concentration in the residual austenite is 1.3% or less, and the packet interface distance bainitic ferrite and / or bainite Ri der than 1.4 [mu] m,
It is a high-strength steel sheet for warm forming that is excellent in ductility and deep drawability in warm, characterized by having a tensile strength at room temperature of 980 MPa or more .

The invention described in claim 2
After the steel slab having the component composition according to claim 1 is hot-rolled to form a cold-rolled steel sheet by cold rolling, the cold-rolled steel sheet is then annealed at austenite single-phase region temperature for 15 to 600 seconds, upon cooling to the cooling stop temperature specified by retaining a temperature range of from 420 to 490 ° C., and cooled by controlling the average cooling rate until at least 550 ° C. to 5 ° C. / sec or higher and held at the holding temperature range for 300 seconds or more according It is a manufacturing method of the high strength steel plate for warm forming which is excellent in the ductility and deep drawability in warm of the claim | item 1 .

  According to the present invention, it is possible to improve ductility and deep drawability, in particular, it is possible to remarkably improve warm ductility and deep drawability, and a high-strength TRIP-type bainitic ferrite steel sheet having more excellent formability. Now available. The steel sheet of the present invention is excellent in ductility even in the cold and has a low yield ratio, and is therefore suitable for cold forming.

It is a crystal orientation map measured by EBSP. It is the elements on larger scale of the crystal orientation map of FIG. It is another partial enlarged view of the crystal orientation map of FIG. It is a figure which shows the relationship between a packet interface space | interval, and intensity | strength and elongation (TSxEL). It is a figure which shows the relationship between austemper temperature and a packet interface space | interval. It is a figure which shows the relationship between a packet interface space | interval and EL. It is a figure which shows the relationship between tensile test temperature and TS. It is a figure which shows the relationship between test temperature and forming speed, and the maximum deep drawing depth in a deep drawing test. It is a figure for demonstrating typically the balance of force in deep drawing. It is a figure which shows the relationship between a tensile speed and TS.

  The inventors of the present invention have a TRIP type steel sheet that has both high strength and ductility, in particular, a high strength TRIP type bainitic ferrite steel sheet mainly composed of bainitic ferrite and / or bainite. In order to improve the deep drawability, we have intensively studied the relationship between the structure and ductility and deep drawability. As a result, the fraction of retained austenite and its solute carbon concentration are controlled within a specific range (residual austenite is 5% or more and carbon concentration in the retained austenite is 1.3% or less), and the bay in the structure By controlling the packet interface spacing of nitrite ferrite and / or bainite (hereinafter, “packet interface spacing of bainitic ferrite and / or bainite” is also simply referred to as “packet interface spacing”) to 1.4 μm or more. In addition, the TRIP effect of retained austenite can be further promoted, and ductility, in particular, ductility during warming can be remarkably improved. In addition, the dependence of the material strength during warming on the processing speed is increased, and the deep drawability during warming. As a result, it was found that the present invention can be improved, and the present invention has been completed based on the findings.

  Hereinafter, the structure characterizing the steel sheet of the present invention will be described first.

<Bay total of bainitic ferrite and bainite: 65% on more than>
Since the bainitic ferrite and / or bainite structure is uniform and fine and rich in ductility, and has a high dislocation density and high strength, the strength-formability balance can be increased by using the matrix. In order to obtain a strength of 980 MPa or more, 65% or more of bainitic ferrite and / or bainite is required in the structure. In addition, in order to obtain excellent ductility in the warm, 5% or more of retained austenite is necessary, which generates bainitic ferrite and / or bainite in the austemper and solidifies in the untransformed austenite phase. What is necessary is just to concentrate molten carbon and to lower Ms temperature to below room temperature. When the temperature and time of the austemper is insufficient and the transformation from austenite to bainitic ferrite and / or bainite is less than 65% in area ratio, the above solute carbon concentration is insufficient and 5% or more of retained austenite Cannot be generated. Therefore, the sum of bainitic ferrite and bainite is 65% or more. Preferably it is 70% or more, More preferably, it is 75% or more.

<Residual austenite: 5% or more>
Excellent ductility in the warm is achieved by increasing the work hardening rate due to the transformation of the retained austenite into a hard phase (work-induced transformation) when the material is processed by including the retained austenite in the metal structure. Is done. When the retained austenite is 5% or less, the amount of increase in work hardening rate becomes insufficient, so that excellent ductility in the warm cannot be obtained. Therefore, the retained austenite is 5% or more. Preferably, it is 7% or more, more preferably 9% or more.

<Total of martensite and retained austenite: 35% or less>
If the total of martensite (including those in which retained austenite is processing-induced transformation) and retained austenite is more than 35% in the structure, the ductility decreases due to the increase in strength, and excellent ductility in the warm is obtained. Disappear. Therefore, the total of martensite and retained austenite is 35% or less. Preferably, it is 30% or less, more preferably 25% or less.

<Polygonal ferrite: 10% or less (including 0%)>
Polygonal ferrite does not contain dislocations and is relatively soft, which is advantageous for improving the ductility in the warmth, but greatly reduces the strength. If there is more than 10% polygonal ferrite, sufficient strength cannot be achieved. Therefore, the polygonal ferrite is 10% or less (including 0%). Preferably, it is 7% or less, more preferably 4% or less.

<Organization other than the above as the balance: 5% or less (including 0%)>
Examples of the remaining structure other than the above include pearlite and pseudo-pearlite in which retained austenite is decomposed, but each is relatively soft, which is advantageous for improving the ductility in the warmth, but greatly reduces the strength. If pearlite or pseudo pearlite exceeding 5% is present, sufficient strength cannot be obtained. Therefore, the remaining structure is 5% or less (including 0%). Preferably, it is 3% or less.

<Carbon concentration in the retained austenite: 1.3% or less>
The retained austenite becomes more stable as the carbon concentration is higher as the temperature is higher (the process-induced transformation is less likely to occur). Therefore, the retained austenite is more stabilized in warm compared to room temperature. At this time, if the amount of carbon in the retained austenite exceeds 1.3%, the stability of the retained austenite becomes too high, and no work-induced transformation occurs in the warm state, so that excellent ductility in the warm state cannot be obtained. Therefore, the carbon concentration in the retained austenite is set to 1.3% or less. Preferably, it is 1.15% or less, more preferably 1.0% or less.

<Bainitic ferrite and / or bainite packet interface spacing> 1.4 μm>
As described above, by promoting the TRIP effect of retained austenite, the amount of increase in work hardening rate is increased and ductility is improved. This is achieved by increasing the stress distributed from the parent phase to the retained austenite and making it more prone to work-induced transformation. Residual austenite exists mainly at the packet interface of bainitic ferrite and / or bainite, and the number of retained austenite that can exist per unit interface length is approximately constant. Therefore, widening the packet interface spacing reduces the total interface length, reduces the number of austenite and increases the individual size. As a result, the stress of each retained austenite during processing increases, and the processing-induced transformation of the retained austenite is promoted, so that excellent ductility in the warm can be obtained. This is achieved by setting the austemper at a relatively high temperature for a long time and setting the packet interface interval to 1.4 μm or more. When the packet interface interval is less than 1.4 μm, the number of retained austenite is increased and finely distributed, so that the stress distributed to the retained austenite is reduced. As a result, processing-induced transformation during processing does not occur sufficiently, and excellent ductility during warming cannot be obtained. Therefore, the packet interface spacing of bainitic ferrite and / or bainite (the definition will be described later in the section <Measurement of packet interface spacing of bainitic ferrite and / or bainite>) is set to 1.4 μm or more. Preferably, it is 1.5 μm or more, more preferably 1.6 μm or more.

  Next, the component composition which comprises this invention steel plate is demonstrated. Hereinafter, all the units of chemical components are mass%.

[Component composition of the steel sheet of the present invention]
<C: 0.10 to 0.30%>
C is an essential element for securing high strength and securing retained austenite. Specifically, it is an important element for dissolving sufficient C 2 in the austenite phase and leaving the desired austenite phase even at room temperature. If C is less than 0.10%, it is difficult to obtain a high strength of 980 MPa or more, and C is 0.10% or more. Preferably it is 0.13% or more. However, if C is excessive, weldability deteriorates, so the C content is limited to 0.30% or less. Preferably it is 0.25% or less, More preferably, it is 0.19% or less.

<Si: more than 1.0% to 3.0% or less>
In addition to being useful as a solid solution strengthening element, Si is an element that effectively suppresses the formation of carbides by decomposition of residual austenite. From such a viewpoint, the Si amount is set to exceed 1.0% in the present invention. Preferably it is 1.3% or more, More preferably, it is 1.5% or more. However, if Si exceeds 3.0%, the surface properties are deteriorated, so it is suppressed to 3.0% or less. Preferably it is 2.0% or less, More preferably, it is 1.8% or less.

<Mn: 1.0 to 3.0%>
Like Si, Mn is useful as a solid solution strengthening element and is an element necessary for stabilizing austenite and obtaining desired retained austenite. In order to exhibit such an action effectively, it is necessary to contain 1.0% or more. Preferably it is 1.5% or more, More preferably, it is 2.0% or more. On the other hand, if the amount of Mn is excessive, it may cause cracks in the slab, so the content is made 3.0% or less. Preferably it is 2.8% or less, more preferably 2.6% or less.

<P: 0.10% or less (including 0%)>
P is unavoidably present as an impurity element, segregates at the prior austenite grain boundaries, and embrittles the grain boundaries. Preferably it is 0.05% or less, More preferably, it is 0.02% or less.

<S: 0.010% or less (including 0%)>
S is also unavoidably present as an impurity element, forms MnS inclusions, and causes toughness deterioration and weld cracking, so 0.010% or less. Preferably it is 0.007% or less, More preferably, it is 0.005% or less.

<N: 0.0020 to 0.0300%>
N is an unavoidable impurity in the high-strength steel sheet according to the present embodiment, and if it is excessively contained, coarse nitrides are precipitated and workability is deteriorated. For this reason, it is desirable to reduce the N content as much as possible. However, if it is 0.0300% or less, even a high-strength material as intended in the present invention does not adversely affect workability. For this reason, N content shall be 0.0300% or less. Preferably it is 0.0200% or less, More preferably, it is 0.0100% or less. Note that, if N is less than 0.0020%, a large increase in manufacturing cost is caused, so that the lower limit is about 0.0020% from the viewpoint of manufacturing cost.

<Al: 0.0010 to 0.1%>
Al is an element added for deoxidation in steel, and 0.0010% or more is necessary for deoxidation with Al. Preferably it is 0.01% or more, More preferably, it is 0.03% or more. On the other hand, Al combines with N to form AlN, but if it exceeds 0.1%, too much AlN will deteriorate the ductility, so 0.1% is made the upper limit. Preferably it is 0.07% or less, More preferably, it is 0.05% or less.

  Next, the preferable manufacturing method for obtaining this invention steel plate is demonstrated below.

[Preferred production method of the steel sheet of the present invention]
In order to manufacture the steel plate as described above, first, steel having the above component composition is melted and hot rolled after being formed into a slab by ingot forming or continuous casting. As preferable hot rolling conditions, the slab heating temperature (SRT) is 1000 to 1300 ° C., the rolling end temperature (FDT) is 870 to 950 ° C., and the winding temperature (CT) is 350 to 720 ° C. . After hot rolling is completed, pickling is performed and then cold rolling is performed. The cold rolling rate (CR) is preferably 40 to 90%.

  And after the said cold rolling, it anneals continuously.

[Annealing conditions]
As the annealing conditions, the cold-rolled steel sheet was annealed at the temperature of the austenite single-phase region for 15 to 600 seconds, and then cooled to the cooling stop temperature determined in the holding temperature region of 420 to 490 ° C., and the average cooling to at least 550 ° C. Cooling is performed by controlling the speed to 5 ° C./s or higher, and the temperature is maintained for 300 seconds or longer in the holding temperature range.

<Annealing at austenite single phase region temperature for 15 to 600 seconds>
Annealing at the temperature of the austenite single phase region is to suppress the formation of polygonal ferrite, to have bainitic ferrite and / or bainite as a parent phase structure, and to obtain a sufficient amount of retained austenite with an appropriate carbon concentration, It takes 15 seconds or more to promote reverse transformation from ferrite to austenite. However, if it exceeds 600 seconds, the structure becomes coarse, and sufficient bainitic ferrite and / or bainite amount and residual austenite amount cannot be obtained during the subsequent holding at 420 to 490 ° C., and excellent warm ductility is achieved. Therefore, 600 seconds or less is set. The annealing is preferably performed for 50 to 400 seconds, more preferably 100 to 300 seconds. The temperature of the austenite single-phase region refers to the A3 point or higher. Here, the A3 point can be calculated from the following calculation formula described on page 273 of “Leslie Steel Material Science” (Maruzen Co., Ltd., issued May 31, 1985). However, those that do not correspond to the components defined in the present invention are excluded from the formula.
^{Ac3 = 910-203 × [C] 0.5} -15.2 × [Ni] + 44.7 × [Si] + 104 × [V] + 31.5 × [Mo] + 13.1 × [W] -30 × [ Mn] -11 × [Cr] −20 × [Cu] + 700 × [P] + 400 × [Al] + 400 × [Ti]

<Cooling at an average cooling rate of 5 ° C./second or more until at least 550 ° C.>
If the average cooling rate to at least 550 ° C. is less than 5 ° C./second, more than 10% of polygonal ferrite is formed. Therefore, in order to make polygonal ferrite 10% or less, it is necessary to cool by controlling the average cooling rate to 5 ° C./second or more to at least 550 ° C. Preferably it is 10 degreeC / second or more, More preferably, it is 15 degreeC / second or more.

<Hold for 300 seconds or more in the holding temperature range of 420-490 ° C>
When the holding temperature range is less than 420 ° C., the nucleation frequency of bainitic ferrite and / or bainite increases, and the microstructure becomes finer, so that the packet interface interval becomes less than 1.4 μm, and the processing-induced transformation of residual austenite is suppressed. , Excellent ductility in the warm cannot be obtained. On the other hand, when the holding temperature range exceeds 490 ° C., the retained austenite is decomposed and 5% or more of retained austenite cannot be obtained, and excellent ductility in the warm state cannot be obtained. The holding temperature is not necessarily constant, and may be varied as long as it is within a temperature range of 420 to 490 ° C. In addition, when the holding time is less than 300 seconds, the time for growth of the generated bainitic ferrite and / or bainite is insufficient, and a packet interface interval of 1.4 μm or more cannot be obtained, which is excellent in warm conditions. The ductility cannot be obtained. Therefore, it is necessary to hold for 300 seconds or more in the holding temperature range of 420 to 490 ° C. The holding temperature is preferably 430 to 480 ° C, more preferably 440 to 470 ° C. The holding time is preferably 350 seconds or more, more preferably 400 seconds or more. In addition, if it is in the said holding temperature range, 65% or more of the total amount of bainitic ferrite and bainite and 35% or less of the total amount of martensite and residual γ will be achieved.

  Hereinafter, the present invention will be described more specifically with reference to experimental examples.However, the present invention is not limited by the following experimental examples, but may be appropriately modified within a range that can meet the purpose described above and below. It is also possible to implement, and they are all included in the technical scope of the present invention.

[Experimental Example 1]
After melting the steel of the components shown in Table 1 in a converter and forming a slab by continuous casting, a hot rolled steel sheet having a thickness of 3.2 mm was obtained by hot rolling, and then the surface scale was removed by pickling. Cold-rolled to 1.2 mm, and then annealed as follows to prepare various samples (sample Nos. 1 to 5).
<Hot rolling process>
Heating temperature: Hold at 1200 ° C for 3 hours Finishing temperature: 890 ° C
Winding temperature: 550 ° C
<Cold rolling process>
Cold rolling ratio: 62.5% (3.2 mm → 1.2 mm)
<Continuous annealing process>
Table 2 shows the annealing conditions.

  About each sample after annealing, the metal structure and the tensile property at room temperature and 300 degreeC were investigated with the following method.

(Metal structure)
<Martensite and retained austenite>
1/4 part of the plate thickness in the rolling direction cross section of the sample taken from the vicinity of 1/2 part of the coil longitudinal direction and 1/2 part of the width direction is subjected to Nital corrosion, and any three places are subjected to a scanning electron microscope (SEM). And observed at 3000 times.
Martensite, retained austenite, and mixtures thereof are all gray in SEM and can be distinguished. Using a photograph taken by SEM observation (3000 times), an area ratio was calculated by a point calculation method with a measurement region of 50 × 50 μm and a measurement interval of 0.1 μm. In the point calculation method, a total of 20 grid lines are drawn in the SEM photograph 50 μm × 50 μm area, 10 grid lines at equal intervals in each of the vertical and horizontal directions, and exist at the points where these grid lines intersect (lattice points). The martensite, retained austenite, and the mixture thereof were discriminated according to the above criteria, and the counting result was divided by 100, which is the total number of lattice points. And the average value of the area ratio computed from the SEM photograph of three places selected arbitrarily was calculated | required.

<Polygonal ferrite>
One-quarter of the plate thickness in the cross section in the rolling direction of the sample taken from about 1/2 part of the coil longitudinal direction and 1/2 part of the width direction is subjected to Nital corrosion, and any three places are observed with an optical microscope at 1000 times. did.
Polygonal ferrite is a massive ferrite with no dislocation or very little substructure, and can be distinguished because it corrodes white. Using a photograph taken by observation with an optical microscope (1000 times), an area ratio was calculated by a point calculation method with a measurement region of 50 × 50 μm and a measurement interval of 0.1 μm. In the point calculation method, a total of 20 grid lines are drawn, 10 at regular intervals in each of the vertical and horizontal directions in the optical micrograph 50 μm × 50 μm area, and the points (lattice points) where these grid lines intersect. The existing polygonal ferrite was discriminated according to the above criteria, and the counting result was divided by 100 which is the total number of lattice points. And it measured similarly in the 3 visual fields selected arbitrarily, and calculated | required the average value.

<Bainitic ferrite and / or bainite>
1/4 part of the plate thickness in the rolling direction cross section of the sample taken from the vicinity of 1/2 part of the coil longitudinal direction and 1/2 part of the width direction is subjected to Nital corrosion, and any three places are subjected to a scanning electron microscope (SEM). And observed at 3000 times.
The bainitic ferrite and / or bainite structure referred to in the present invention means a structure having a substructure with a high dislocation density, and both are plate-like and dark gray or black in SEM photographs. The area ratio was calculated in the same manner as described for martensite and retained austenite.
At this time, in the SEM photograph, the polygonal ferrite looks the same (plate-like dark gray or black) and is often difficult to discriminate. Therefore, from the area ratio of the dark gray or black portion determined from the SEM photograph, The area ratio of nail ferrite was subtracted to obtain the area ratio of bainitic ferrite and / or bainite structure.

<Residual austenite>
The measurement object is an arbitrary measurement region (about 20 × 20 μm) on a plane parallel to the rolling surface of ¼ part of the plate thickness in a sample taken from ½ part in the coil longitudinal direction and ½ part in the width direction. And However, when polishing up to the measurement surface, electrolytic polishing is preferably performed in order to prevent transformation of retained austenite due to mechanical polishing. Next, X-rays are irradiated using an X-ray diffractometer. In the present invention, the target was MoKα, and the conditions were an acceleration voltage of 50 kV and an acceleration current of 250 mA. Integration of bcc (ferrite martensite) (200) plane, (211) plane and fcc (austenite) (200) plane, (220) plane, (311) plane by X-ray analysis using Mo Kα ray The reflection intensity was measured, and the area ratio of residual austenite (fcc) was measured using the diffraction peak intensity ratio of bcc and fcc as the area fraction of the phase.

<Other organizations>
The total area (100%) was obtained by subtracting the area ratio occupied by the structure (martensite and retained austenite, bainitic ferrite and / or bainite, polygonal ferrite).

<Carbon concentration in retained austenite>
The carbon concentration in the retained austenite is determined by integral reflection of ferrite (200) plane, (211) plane and austenite (200) plane, (220) plane, (311) plane by X-ray analysis using Mo Kα ray. The strength was measured and calculated by the method described in “Journal of The Iron and Steel Institute, 1968, 206, p. 60”.

<Measurement of packet interface spacing of bainitic ferrite and / or bainite>
The measurement of the packet interval of bainitic ferrite and / or bainite can be performed at an arbitrary value of 1/4 part of the plate thickness in the rolling direction cross section of the sample taken from about 1/2 part of the coil longitudinal direction and about 1/2 part of the width direction. 3 locations, each of which has a packet interface map (packet interface definition: crystal orientation difference of 2 °) in a region of 50 μm × 50 μm (1000 times), and using an EBSD (Electron back-scattered diffraction) measuring device manufactured by TSL It was measured. More specifically, as illustrated in FIG. 1, a packet interface map (50 μm × 50 μm, measurement pitch 1.5 μm) measured by EBSD is divided into C1 to C4, L1 at equal intervals of 10 μm in each of vertical and horizontal directions. A total of eight lattice lines L to L4 were drawn, and the distances (for example, d1, d2, d3,..., Dn) between the packet interfaces intersecting these lattice lines were measured. 2 and 3 are enlarged views on the left side of the L1 lattice line and on the right side of the L4 lattice line on the C2 lattice line of FIG. 1, respectively.
In image analysis, a region where Confidence Index (CI value), which is a value indicating the reliability of the crystal orientation measurement result by EBSD, is considered to be martensite because it is considered that there are many dislocations and the like. Therefore, the point where the lattice line collides with the region where the CI value is 0.1 or less inside the bucket is also regarded as the bucket interface, and the distance between the packet interfaces is measured as shown by d2 and d3 in FIG.
When a single grid line passes through the same packet a plurality of times, as shown in FIG. 3, the distance between the packet interfaces is measured for all of them, and the distance between all the packet interfaces is measured. The data was added to the total number as it was, and the average value was obtained. For example, the average value of the distance between the packet interfaces on the lattice line C2 was determined as dC2 = (d1 + d2 + d3 +... + Dn) / n. In this way, the average value (dC1, dC2, dC3, dC4, dL1, distance between the packet interfaces on all eight grid lines (C1, C2, C3, C4, L1, L2, L3, L4) dL2, dL3, dL4) were determined, and the average value d (d = (dC1 + dC2 + dC3 + dC4 + dL1 + dL2 + dL3 + dL4) / 8) was determined. Then, the measurement results (average value d) at each of the three locations were further averaged, and the average value was defined as “packet interface distance of bainitic ferrite and / or bainite” of each sample.
In this example, a schottky field emission scanning electron microscope (JSM-6500F manufactured by JEOL Ltd.) was used for FE-SEM, and an HIKARI (EBSD detection camera, manufactured by EDAX / TSL) was used for EBSD. OIM Data Collection (ver. 5.2) was used for the measurement, and OIM Analysis (ver. 5.2) was used for the analysis.

(Tensile strength, elongation)
<Tensile test at room temperature>
The tensile test at room temperature was conducted in accordance with JIS Z2241, using a No. 13 B tensile test piece (gauge length 50 mm, width 12.5 mm) described in JIS Z2201, under the conditions of a crosshead speed of 10 mm / min. The temperature was 20 ° C. This was carried out using an autograph (Shimadzu Corporation).

<Warm tensile test>
The warm tensile test was performed using a JIS 13B tensile test piece (gauge length 50 mm, width 12.5 mm) at a crosshead speed of 10 mm / min and a test temperature of 300 ° C. The temperature is measured by bringing a thermocouple into contact with three parallel parts of the test piece (the center of the test piece and 7 mm inward from both ends of the parallel part), and the test piece is brought to a predetermined temperature using an infrared heating furnace (ULVAC Riko). Heated. In order to eliminate the temperature distribution in the test piece, a tensile test was performed after holding at a predetermined temperature for 15 minutes or more.

  Tables 3 and 4 show the results of investigating the metal structure and the tensile properties at room temperature and 300 ° C.

From Table 3, samples 2 to 4, 8 to 10, 12, and 13 have component compositions within an appropriate range and appropriate annealing conditions, and target structures are obtained. However, although the component composition is within the appropriate range, the sample No. In 1, 5, 6, and 7, the packet interface interval is small and is less than 1.4 μm. Sample No. In No. 1, the retention temperature is low, the nucleation frequency of bainitic ferrite is high, many parts of the structure become bainitic ferrite, and the retained austenite area ratio is 4.9%, which is below the target. This is because bainitic ferrite is generated from the prior austenite grain boundaries, grows while discharging carbon, and the untransformed austenite region remains in the structure as retained austenite. Therefore, the retained austenite decreases as bainitic ferrite increases. This is probably because of this. Sample No. In No. 5, the carbon concentration in the retained austenite is high. As in the case of No. 1, since the holding temperature is low, the nucleation frequency of bainitic ferrite is high, and the amount of bainitic ferrite and the retained austenite are no. In addition to being almost the same amount as No. 1, no. Compared to the case of 1 (380 ° C.), the holding temperature range is higher (400 ° C.), so the carbon discharge from bainitic ferrite to austenite is faster and the concentration is promoted. Sample No. In No. 6, since the holding temperature is lower than the preferred range, the carbon concentration in the retained austenite is high, and the packet interface interval is narrow. Both of these tissue changes increase the stability of residual γ, so the warm TRIP effect is suppressed and the warm TS × EL is low. Sample No. 7, since the holding temperature is lower than the preferred range, the packet interface interval is narrowed, thereby increasing the stability of residual γ and suppressing the TRIP effect in the warm, so that TS × EL in the warm Is low. In the temperature range of 420 ° C. or higher, the nucleation frequency of bainitic ferrite decreases as the temperature rises, so that the suppression of the formation becomes remarkable, the increase in the volume fraction of retained austenite is large, and the relative solidification. The dissolved carbon concentration tends to decrease.
On the other hand, although the annealing conditions are appropriate, Sample No. In No. 11, since the amount of residual γ produced is insufficient (less than 5%), the increase in work hardening rate due to the TRIP effect is small, and the TS × EL in the warm state is low.

  From Table 4, it can be seen from the mechanical properties at warm (300 ° C.) that the sample interface with a packet interface spacing of about 1.5 to 1.7 μm. 2 to 4, 8 to 10, 12, and 13 have TS of about 1000 to 1100 MPa and EL of 23% to 31%. This is the same as sample No. The EL at 300 ° C. is improved by up to about 16% (about 15% → about 31%), while the room temperature EL of 2 to 4, 8 to 10, 12, and 13 is about 15 to 18%. On the other hand, the sample No. with a packet interface interval of 1.2 μm. Sample No. 1 having a TS of about 1160 MPa, an EL of 18%, and a packet interface interval of 1.36 μm. Sample No. 5 having a TS of about 1046 MPa, an EL of 21%, and a packet interface spacing of 1.26 μm. Sample No. 6 having a TS of about 1170 MPa, an EL of 17%, and a packet interface spacing of 1.38 μm. The TS of 7 is about 1080 MPa and the EL is 20%, which is an improvement of 4 to 7% with respect to the EL at room temperature of 11%, 15%, 12% and 16%, respectively. Thus, when the packet interface interval is 1.4 μm or more, the EL improvement margin is twice or more (16% vs. 4 to 7%) as compared with the case of less than 1.4 μm. When the packet interface interval is 1.4 μm or more, the elongation is remarkably improved because the distribution of retained austenite in the structure becomes coarse, the stress distributed to each retained austenite by processing increases, It is thought that transformation-induced plasticity (TRIP) is more likely to occur than at room temperature. On the other hand, if the packet interface interval is less than 1.4 μm, the distribution of retained austenite in the structure becomes fine, and the stress distributed to each retained austenite by processing decreases, so transformation induced plasticity (TRIP) even in warm conditions. ) Is less likely to occur.

  FIG. 4 shows the relationship between the packet interface spacing and the strength / ductility (TS × EL), FIG. 5 shows the relationship between the austempering temperature and the packet interface spacing, and FIG. 6 shows the relationship between the ductility (EL) and the packet interface spacing.

  As shown in FIG. 4, the strength / ductility (TS × EL) is about 21000 to 22000 MPa% when the packet interface interval is less than 1.4 μm. As a result, the packet interface interval is 1.7 μm, and TS × EL is about 30000 MPa%.

  As shown in FIG. 5, by heating to a single phase region, cooling at a predetermined cooling rate, and controlling the austemper temperature to 420 ° C. or higher, the packet interface interval can be made 1.4 μm or higher.

As shown in FIG. 6, the EL at 300 ° C. is about 20% when the packet interface interval is less than 1.4 μm, but it is remarkably improved when the packet interface interval is 1.4 μm or more, and the packet interface interval is 1.7 μm. EL is over 30%. On the other hand, the EL at room temperature is 11% when the packet interface interval is 1.2 μm, but the EL is improved to almost 15% when the packet interface interval is 1.3 to 1.7 μm.
When the packet interface interval is less than 1.4 μm, the EL difference between 300 ° C. and room temperature is 6 to 7%, and this difference is considered to be an improvement in EL due to higher temperatures. However, when the packet interface interval is 1.4 μm or more, EL improvement (12 to 16%) far exceeding the EL improvement allowance (6 to 7%) due to high temperature is recognized.

[Experimental example 2]
Next, in order to examine warm deep drawability, the sample Nos. Shown in Tables 3 and 4 above were used. 3 (invented steel plate) and Sample No. The following experiment was performed using both samples 5 (comparative steel plates).

<Investigation of effects of processing temperature and processing speed on warm characteristics (TS, EL)>
First, in order to grasp the influence of the processing temperature and the processing speed on the warm characteristics (TS, EL), the test temperature is set to 20 at each of two tensile speeds (crosshead speeds) of 10 mm / min and 1000 mm / min. The tensile test was carried out by sequentially changing between ˜300 ° C. The tensile test conditions are the same as in [Experimental Example 1].

  The test results are shown in Table 5 and FIG.

From Table 5 and FIG. 7, the inventive steel plate (Sample No. 3) has a marked decrease in tensile strength (TS) in a specific temperature range (around 200 ° C.) compared to the comparative steel plate (Sample No. 5), It can be seen that the higher the tensile speed, the more the degree of decrease in the tensile strength (TS), that is, the higher the tensile speed dependency of the tensile strength (TS).
It is confirmed that the material strength after the heating test is not lower than the material strength before the test heating at any test temperature.

<Investigation of warm drawability>
Therefore, in order to investigate warm deep drawability, punch diameter 50 mm (punch shoulder R5 mm), die diameter 54 mm (die shoulder R7 mm), blank diameter 103 mm, wrinkle holding force 1224 kN, processing speed 0.1 mm / s or 10 mm A deep drawing test was performed at a test temperature of 22 ° C. or 200 ° C. under the conditions of / s.

  The test results are shown in Table 6 and FIG.

  From Table 6 and FIG. 8, the invention steel plate (sample No. 3) has only the same level of deep drawability at room temperature as compared with the comparative steel plate (sample No. 5), while warm, It is recognized that the deep drawability is greatly improved in a specific temperature range (around 200 ° C.) in which the tensile strength is remarkably reduced, and the improvement effect becomes more remarkable as the molding speed is increased.

<Assumed mechanism for improving deep drawability in warm conditions>
As described above, the mechanism by which the deep drawability during warming in the inventive steel sheet is remarkably improved was studied.

In deep drawing, as schematically shown in FIG. 9, the force acting on the flange portion of the blank (material) is
(1) Circumferential compressive force component accompanying shrinkage deformation of the flange portion (2) Friction force component from the die and blank holder (3) Bending and bending back force component from the die shoulder Is the molding load (P).

  And while the fracture resistance of the punch shoulder is larger than the molding load (P), the blank (material) is continued without breaking and deep drawing can be performed.

  On the other hand, the deformation speed in each part of the blank (material) at the time of deep drawing is greatly different. In the punch shoulder, after the initial stage of processing until the blank (material) conforms to the punch shoulder, the deformation speed is It becomes relatively smaller than the flange part.

  Here, FIG. 10 shows the data at the test temperature of 200 ° C. in Table 5 rearranged in the relationship between the tensile speed and the tensile strength. Like the invention steel plate in the figure, the material whose strength of processing is more dependent on the processing speed (that is, the material whose degree of decrease in the material strength is greater with respect to the increase in the processing speed), the forming load (P), that is, the deformation speed. The stress required to continue the deformation of the relatively large flange portion becomes low, while the breaking strength (TS) of the punch shoulder portion where the deformation speed is relatively small is maintained high. ) Is large, the processing can be continued without breaking, and the deep drawability becomes good.

  Moreover, although it is not necessarily clear why the tensile strength decreases and the tensile speed dependency increases in a specific temperature range (around 200 ° C.) in the invention steel plate, it is considered as follows.

That is, as described above, the invented steel sheet increases the packet interface interval by increasing the austemper temperature in order to ensure warm ductility. From this, the steel sheet according to the present invention increases the size of retained austenite, and as described above, the stress distribution due to the processing to retained austenite is increased, thereby promoting the transformation to martensite. On the other hand, bainitic ferrite, which is the parent phase, softens as the temperature rises. Therefore, when the tensile temperature rises from room temperature to around 200 ° C, the effect of softening the parent phase becomes stronger, and stress distribution to residual austenite is reduced. Reduce temporarily. In particular, when the packet interface interval is increased as in the case of the invention steel plate, the degree of softening of the base material in the vicinity of 200 ° C. increases, and the transformation to martensite is greatly suppressed. As a result, it is considered that the tensile strength near 200 ° C. is lower than the tensile strength at room temperature.
Further, in the steel sheet of the invention, when the tensile speed is increased in this temperature range, the transformation speed cannot follow the processing speed, the tensile strength is further reduced, and the tensile speed dependency is increased.

  However, if the tensile temperature is further increased from around 200 ° C, the work hardening by dynamic strain aging exceeds that of the softening of the parent phase due to the temperature rise, so the parent phase strength increases again despite the high temperature. As the stress distribution to the retained austenite is strengthened, the processing-induced transformation is accelerated. For this reason, in the temperature range from around 250 ° C. to around 300 ° C., the elongation is improved and the tensile strength is recovered.

<Other effects of invention steel plate>
As described above, the inventive steel sheet is excellent in formability having both ductility and deep drawability, but the strength of such an inventive steel sheet is relatively lowered with respect to its tensile strength at room temperature. By performing warm forming in a temperature range (for example, 100 to 250 ° C .; see FIG. 7), the processing-induced transformation of retained austenite is suppressed. Therefore, when press-molding automobile parts using the invention steel plate, the amount of retained austenite in the parts can be made larger than parts manufactured by ordinary cold pressing by processing in such a temperature range. It becomes possible. As a result, even if the part is deformed at the time of collision, an effect of increasing the crack generation limit due to the deformation can be obtained due to the large amount of retained austenite.
Moreover, since the tensile strength is reduced by warm working in the above temperature range, it naturally has an effect of reducing the load applied to the press.

Claims (2)

% By mass (hereinafter the same for chemical components)
C: 0.10 to 0.30%,
Si: more than 1.0% to 3.0% or less Mn: 1.0 to 3.0%
P: 0.10% or less (including 0%),
S: 0.010% or less (including 0%),
N: 0.0020 to 0.0300% or less,
Al: 0.0010 to 0.1%
And the remainder has a component composition consisting of iron and inevitable impurities,
Microstructure, all tissues sum of bainitic ferrite and bainite at an area ratio with respect to 65% or more,
Residual austenite: 5% or more,
Total of martensite and retained austenite: 35% or less,
Polygonal ferrite: 10% or less (including 0%)
Other than the above, the rest consists of 5% or less (including 0%),
The carbon concentration in the residual austenite is 1.3% or less, and the packet interface distance bainitic ferrite and / or bainite Ri der than 1.4 [mu] m,
A high-strength steel sheet for warm forming that has excellent ductility and deep drawability in warm, characterized by having a tensile strength at room temperature of 980 MPa or more . After the steel slab having the component composition according to claim 1 is hot-rolled to form a cold-rolled steel sheet by cold rolling, the cold-rolled steel sheet is then annealed at austenite single-phase region temperature for 15 to 600 seconds, upon cooling to the cooling stop temperature specified by retaining a temperature range of from 420 to 490 ° C., and cooled by controlling the average cooling rate until at least 550 ° C. to 5 ° C. / sec or higher and held at the holding temperature range for 300 seconds or more according Item 2. A method for producing a high-strength steel sheet for warm forming , which is excellent in warm ductility and deep drawability according to Item 1 .