JP5719545B2 - High strength thin steel sheet with excellent elongation and press forming stability - Google Patents

Description

  The present invention relates to a high-strength thin steel sheet excellent in elongation and press-forming stability, which is suitable for structural materials such as automobiles, which are mainly used after being pressed, and a method for producing the same.

  High press workability and strength are required for steel plates used in automobile body structures. In particular, elongation is the main cause of deterioration of formability of high tensile strength because it decreases as the strength increases, although it is the most important property in press molding.

  In order to solve such a problem, in Patent Documents 1 and 2, the retained austenite phase is left in the steel sheet and plastic deformation-induced transformation (TRIP effect) is used, and the elongation is very high despite high strength. Is disclosed. These increase the amount of C and Si, stabilize the retained austenite phase, and stably remain even at room temperature.

  As a technique for further effectively utilizing such a TRIP effect, Patent Document 3 discloses that the austenite phase residual rate at the maximum stress point is hydroformed at a temperature of 60% to 90%, thereby expanding the pipe. A hydroforming technique is disclosed in which the rate is 150% higher than at room temperature.

  Patent Document 4 discloses a technique for improving deep drawability in TRIP steel by heating a mold.

  However, in the former, the processing is limited to pipes, and in the latter, the mold heating is costly in order to obtain a sufficient effect, so the application range is limited.

  On the other hand, in order to effectively exhibit the TRIP effect at a lower temperature from the steel sheet side, there is further addition of C, but the added C is not only concentrated in the austenite phase but also precipitated as coarse carbides. Resulting in. Then, conversely, problems such as material deterioration due to a decrease in the amount of retained austenite and generation of cracks at the time of hole expansion starting from carbides occur.

Japanese Patent Laid-Open No. 61-217529 Japanese Patent Laid-Open No. 5-59429 JP 2004-330230 A JP 2007-1111765 A

  The present invention has been made to solve the conventional problems, and is a technique for stably securing excellent press formability in retained austenitic steel, and is excellent in elongation and press stability. Another object is to provide a high-strength thin steel sheet and a method for producing the same.

  Residual austenitic steel controls the ferrite transformation and bainite transformation during annealing, and increases the C concentration in the austenitic phase, leaving the austenitic phase in the steel structure of the product. By the TRIP effect of this residual austenitic phase, high elongation is achieved. It is a high strength steel plate with

  However, since this TRIP effect is temperature-dependent, changes in the mold temperature during press forming deteriorate the stability of the press formability of the steel sheet, resulting in unstable cracks in the early or late stages of the press. was there.

  As a result of intensive studies to make the TRIP effect work independently of the press molding temperature, the present inventors have found that the problem can be overcome by uniformly dispersing the retained austenite phases having different forms. In addition, by increasing the amount of martensite in the retained austenitic steel as compared with a normal steel sheet, a technique for uniformly dispersing an austenitic phase having a different stability than ever has been found, and the present invention has been completed.

(Reference 1) M. Takahashi: IS3-2007, (2007), 47-50.
That is, the high-strength thin steel sheet excellent in stretch molding stability of the present invention is
(1) In mass%,
C: 0.05% or more, 0.35% or less,
Si: 0.05% or more, 2.0% or less,
Mn: 0.8% or more, 3.0% or less,
P: 0.0010% or more, 0.1% or less,
S: 0.0005% or more, 0.05% or less,
N: 0.0010% or more and 0.010% or less Al: 0.01% or more and 2.0% or less, and has a steel composition composed of the balance iron and inevitable impurities, the microstructure is area ratio, ferrite The total austenite phase is 10% or more and 90 % or less, the residual austenite phase is 5% or more and 30% or less in the area ratio, the martensite phase is 5% or more and 20% or less in the area ratio, and the residual austenite phase is lath. In the elongation and press forming stability, the area ratio γi of the island-like residual austenite phase and the area ratio γ of the total residual austenite phase satisfy the following formula (1). Excellent high strength steel plate.

0.7 ≧ γi / γ ≧ 0.3 Formula (1)
(2) In the steel sheet described in (1),
A high strength thin steel sheet excellent in elongation and press forming stability, characterized in that an average C concentration in an austenite phase is 0.7% or more and 1.5% or less in mass%.
(3) In the steel sheet described in (1) or (2), the total of the ferrite phase and the bainite phase is 50% or more in terms of volume fraction with respect to the entire structure, and press forming High strength thin steel sheet with excellent stability.
(4) In the steel sheet described in any one of (1) to (3),
% By mass
Mo: A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing 0.02% or more and 0.5% or less.
(5) In the steel sheet described in any one of (1) to (4),
% By mass
Nb: 0.01% or more, 0.10% or less,
Ti: 0.01% or more, 0.20% or less,
V: 0.005% or more, 0.10% or less,
Cr: 0.1% or more, 5.0% or less,
W: A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more selected from 0.01% or more and 5.0% or less.
(6) In the steel sheet described in any one of (1) to (5),
% By mass
A high-strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more selected from Ca, Mg, Zr, and REM in an amount of 0.0005% to 0.05%.
(7) In the steel sheet described in any one of (1) to (6),
% By mass
Cu: 0.04% or more, 2.0% or less,
Ni: 0.02% or more, 1.0% or less,
B: A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more selected from 0.0003% or more and 0.007% or less.
(8) A method for producing a steel sheet according to any one of (1) to (7),
For casting slabs, after finishing casting, or once cooled to 1100 ° C or lower, when re-heating to 1100 ° C or higher and performing hot rolling, the finishing temperature ends at 850 ° C or higher and 970 ° C or lower. Then, after cooling at a speed V satisfying the formula (2), winding is performed in a temperature range of 550 ° C. or less, pickling, and then cold rolling of 30% or more, and the highest temperature during annealing is Ac1. As described above, after annealing at Ac3 or lower, the sample is cooled to a temperature range of 200 ° C. or higher and 480 ° C. or lower at an average cooling rate of 1 ° C./second or higher and 200 ° C./second or lower. A method for producing a high-strength thin steel sheet excellent in elongation and press forming stability, characterized in that holding is performed below.

Vc ₉₀ -10>V> Vc ₉₀ -30 Formula (2)
50> Vc ₉₀ > 30 Formula (3)
Vc ₉₀ = 10 ^{(3.69-0.75β)} [° ^C./s ] Formula (4)
β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (5)
(9) A method for producing a steel sheet according to any one of (1) to (7),
When the cast slab is cast as it is or after being cooled to 1100 ° C or lower and then re-heated to 1100 ° C or higher to perform hot rolling, the finishing temperature ends at 850 ° C or higher and 970 ° C or lower. Then, after cooling at a speed V satisfying the formula (2), winding is performed in a temperature range of 550 ° C. or less, pickling, and then cold rolling of 30% or more, and the highest temperature during annealing is Ac1. As described above, after annealing at Ac3 or lower, the sample is cooled to a temperature range of 200 ° C. or higher and 480 ° C. or lower at an average cooling rate of 1 ° C./second or higher and 200 ° C./second or lower. A method for producing a high-strength thin steel sheet excellent in elongation and press forming stability, characterized in that the following holding is performed, followed by immersion in a hot dip galvanizing bath.

Vc ₉₀ -10>V> Vc ₉₀ -30 Formula (2)
50> Vc ₉₀ > 30 Formula (3)
Vc ₉₀ = 10 ^{(3.69-0.75β)} [° ^C./s ] Formula (4)
β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (5)
(10) In the method for producing a steel sheet described in (9), elongation and press forming stability are characterized in that alloying treatment is performed in a range of 465 ° C. or higher and 580 ° C. or lower after immersion in a hot dip galvanizing bath. For producing high-strength thin steel sheets with excellent properties.

  The high-strength thin steel sheet according to the present invention ensures extremely high elongation despite the high strength due to the TRIP effect of retained austenite, and uniformly disperses the retained austenite phase having various forms. It is possible to stably ensure high formability at the time.

  In addition, the basis of the present invention, to ensure a stable and extremely high elongation and high formability during pressing is obtained by controlling the form of austenite phase and the martensite phase fraction, Since the effect can be continued within the range of conditions, it can also be applied to an electroplated steel sheet.

  The present invention is not affected by casting conditions. For example, a casting method (continuous casting or ingot casting) and a slab thickness are not significantly affected, and a special casting-hot rolling method such as a thin slab may be used.

It is a figure showing that the stability of ductility at the time of warm forming improves by optimizing the amount of island-like austenite. It is a figure showing that the stability of the ductility at the time of warm forming improves by optimizing the amount of martensite.

  The high-strength thin steel sheet according to the present invention focuses on dispersing the retained austenite phase having different stability in the retained austenitic steel, and controls the form of the retained austenite phase and the fraction of the martensite phase. And achieved a high level of both elongation and press stability.

  First, the reason for defining the organization of the present invention will be described.

In order to concentrate carbon in the austenite phase, it is necessary to mainly have a ferrite phase and a bainite phase. In order to obtain the effect, it is necessary that the ferrite phase and the bainite phase are present in total 10%. On the other hand, as described later, the residual austenite phase and the martensite phase must be present in a total of 10 % or more, so the upper limit was made 90 %.

  In addition, the present invention is an invention related to a steel sheet characterized by high ductility, and in order to increase ductility, it is necessary to sufficiently secure the α phase, and the total of the ferrite phase and the bainite phase should be 50%. The effect is enhanced by.

  It is necessary to contain 5% or more of the retained austenite phase. However, in order to further improve the stability, it is desirable to contain 10% or more of retained austenite. On the other hand, when the residual austenite phase is present in excess of 30%, secondary processing cracks become a problem. Therefore, the upper limit was made 30%.

  The C concentration in the residual austenite phase is preferably 0.7% or more. This is because if the C concentration is less than 0.7%, the retained austenite phase becomes unstable and transforms into a martensite phase at the initial stage of molding, and cannot contribute to improvement of ductility. On the other hand, if the C concentration is more than 1.5%, the amount is too stable, the amount of transformation of the austenite phase to the martensite phase during molding is small, and the improvement in ductility is small. Therefore, the upper limit is made 1.5% or less.

  Moreover, if it is in said range, even if it contains 5% or less of a pearlite phase, a material will not deteriorate significantly, Therefore It is desirable for a pearlite phase to be 5% or less.

  In the high-strength steel sheet and the method for producing the same according to the present invention, the fraction of martensite phase, the type of retained austenite phase, and their ratio are important. In order to ensure press stability, it is necessary to disperse a retained austenite phase that can always exhibit the TRIP effect with respect to changes in mold temperature during continuous pressing. The highly stable austenite phase contributes to high press formability when the mold temperature at the initial stage of continuous pressing is relatively low, and the low stability austenite phase is high press forming in the latter half of the continuous press where the mold temperature increases. Contributes to sex. For this reason, it is necessary to disperse residual austenite phases having different stability.

As a result of intensive studies, the present inventors have found that the martensite phase has an area ratio of 5% or more and 20% or less, the residual austenite phase has a lath shape and an island shape, and the area ratio γi of _the island-like residual austenite phase _, And a high-strength steel sheet excellent in elongation and press forming stability, characterized in that the area ratio γ of the total retained austenite phase satisfies the following formula.

0.7 ≧ γi / γ ≧ 0.3 Formula (1)
When the island-like retained austenite phase has low stability and the mold temperature is relatively high, it contributes to high press formability, and the lath-like retained austenite phase has high stability and the mold temperature is relatively low Furthermore, it contributes to high press formability. As will be described later, excellent stability is exhibited when the formula (1) is satisfied. This means that the same amount of lath-like and island-like retained austenite phase needs to be present.

  On the other hand, in the high-strength thin steel sheet according to the present invention, the martensite phase is larger than that of ordinary retained austenitic steel. The cause of this is not clear, but it is thought as follows. In the case of warm forming, a supersaturated carbon amount is expelled from the martensite phase into the surrounding austenite phase to stabilize the austenite phase. In the case of heating for about 10 minutes, the effect is considered to occur at 150 ° C. or higher where carbon easily diffuses, and the ductility peak due to the island-like retained austenite phase shifts to the low temperature side.

  On the other hand, since the lath-like residual austenite phase has a peak at 50 to 100 ° C., the effect of concentration of C from the martensite phase is small, and the shift of the ductile peak due to the lath-like residual austenite phase to the low temperature side is small.

  Therefore, in a normal TRIP steel with few martensite phases, a valley due to two ductile peaks existed at 150 ° C., whereas when a martensite phase exists, the two peaks approach and 150 ° C. It is considered that the ductility can be increased and the ductility can be increased at any temperature of 100 to 200 ° C.

  However, if a martensite phase in an amount exceeding 20% is present, it is difficult to cause the residual austenite phase to be present in an amount of 5% or more with the C content in the range specified by the present invention, and as a result, press formability deteriorates. Therefore, the upper limit was made 20%.

  Below, the reason for limitation of the chemical component of the high-strength thin steel sheet concerning this invention is demonstrated.

  C is an essential element from the viewpoint of securing strength and as a basic element for stabilizing the austenite phase. If C is less than 0.05%, the strength is not satisfied, and a residual austenite phase is not formed. On the other hand, if C exceeds 0.35%, the strength is excessively increased and the ductility is insufficient, so that it cannot be used as an industrial material. Moreover, spot weldability is remarkably deteriorated.

When high elongation is required, C is desirably 0.2% or more. On the other hand, when weldability is required, C is preferably 0.25% or less.
Si is an element that delays the formation of cementite phase in addition to ensuring strength, and is an element effective for residual austenite phase formation. Therefore, it is an element that is usually added to ensure ductility. It is. However, even if added over 2.0%, the effect is saturated, and embrittlement tends to occur. When hot dip galvanizing property and ease of chemical conversion treatment are required, 1.5% or less is desirable. On the other hand, if the addition is less than 0.05%, the effect of suppressing the cementite phase cannot be obtained. Therefore, 0.05% is set as the lower limit. When the amount of Al added to obtain the same effect as Si is 0.1% or less, addition of 1% or more is desirable.
If Mn is less than 0.8%, the strength is not satisfied, and the formation of retained austenite becomes insufficient, resulting in deterioration of ductility. Further, when Mn exceeds 3.0%, the hardenability is improved, so that the martensite phase is generated instead of the retained austenite phase, and the strength is easily increased, thereby increasing the variation of the product, Due to lack of ductility, it cannot be used as an industrial material.

  Therefore, the range of Mn in the present invention is 0.8% or more and 3.0% or less. In terms of material, it is preferably 1.0 to 2.4%.

  P is added according to the strength level required as an element for increasing the strength of the steel sheet. However, if the amount added is large, it segregates to the grain boundary and deteriorates the local ductility. In addition, the weldability is deteriorated.

  Therefore, the P upper limit is set to 0.1%. On the other hand, if it is less than 0.0010%, the deterioration effect of P can be ignored.

  S is an element that degrades local ductility and weldability by generating MnS, and is preferably an element that does not exist in steel. Therefore, the upper limit is made 0.05%. On the other hand, if it is less than 0.0005%, the cost increases.

  Al, like Si, has an effect of promoting the formation of a ferrite phase and is one of important elements that can also suppress a cementite phase. That is, it has the effect of stabilizing the retained austenite phase. This effect cannot be expected with Al addition of less than 0.01%. On the other hand, even if Al is added excessively, the above effect is saturated and the steel is embrittled, so 2.0% was made the upper limit.

  In consideration of hot dip galvanizing properties, Al degrades this, so the upper limit is desirably 1.8%.

  N is an element that is inevitably included, but if it is contained in a large amount, not only the aging property is deteriorated but also the amount of precipitated AlN is increased and the effect of Al addition is reduced. % Or less is preferable. Further, unnecessarily reducing N increases the cost in the steelmaking process, so it is usually preferable to control the N to 0.0010% or more.

  Mo is an element that suppresses the formation of a pearlite phase in steel, and is an especially important element when the cooling rate during annealing is slow, or when reheating is performed by alloying treatment of plating or the like. In order to obtain this effect, the minimum addition amount of Mo was set to 0.02%. If it is less than 0.02%, the formation of pearlite is not suppressed, and the retained austenite ratio is reduced. On the other hand, excessive addition of Mo may deteriorate ductility and chemical conversion properties, so the upper limit was made 0.5%. In order to obtain a higher strength-ductility balance, the content is preferably 0.3% or less.

  Nb, Ti, V, Cr, and W are elements that generate fine carbides, nitrides, or carbonitrides, and are effective in securing strength. Therefore, one or more of them may be added as necessary. Is possible. In order to achieve this, it is necessary to add 0.01% for Nb, Ti, and W, 0.005% for V, and 0.1% for Cr.

  On the other hand, excessive addition increases the strength too much and lowers the ductility, so Nb is 0.10% or less, Ti is 0.20% or less, V is 0.10% or less, and Cr is 5.0%. Hereinafter, W needs to be 5.0% or less.

  The steel can further contain one or more of Ca, Mg, Zr, and REM (rare earth elements) alone or in total from 0.0005% to 0.05%. Ca, Mg, Zr, and REM control the shape of sulfides and oxides to improve local ductility and hole expandability. For this purpose, it is necessary to add one or more of these elements alone or in total of 0.0005% or more. However, excessive addition deteriorates workability, so the upper limit was made 0.05%.

  The steel is further selected from the group consisting of Cu: 0.04% or more, 2.0% or less, Ni: 0.02% or more, 1.0% or less, B: 0.0003% or more, 0.007% or less. 1 type or 2 types or more can be contained. These elements can delay the transformation and increase the strength of the steel. However, when Cu is less than 0.04%, Ni is less than 0.02%, and B is less than 0.0003%, the hardenability is weak, and ferrite at high temperatures. In order to promote phase formation, the necessary strength cannot be obtained. On the other hand, if the addition exceeds this range, the hardenability becomes too strong, and the ferrite phase and bainite phase transformation slows, so that the C concentration to the residual austenite phase is delayed.

  In addition to the above elements, steel contains elements inevitably mixed such as Sn and As, and is made of the remaining iron.

  Below, the manufacturing method of the high intensity | strength thin steel plate which concerns on this invention is demonstrated.

Hot rolling can be performed by either a normal hot rolling process or a continuous hot rolling process in which slabs are joined and rolled in finish rolling. The steel having the component ranges described above is cast, and the obtained slab is hot-rolled after being heated or reheated. The reheating temperature is preferably 1100 ° C. or higher in order to dissolve the carbide. Moreover, it is desirable to set it as 1300 degrees C or less from the surface of productivity. The rolling end temperature in the subsequent hot rolling is desirably 850 ° C. or higher and 970 ° C. or lower from the viewpoint of productivity, sheet thickness accuracy, or anisotropy improvement.

  The subsequent cooling rate V (° C./s) is a rate satisfying the following formula, and the winding temperature is 550 ° C. or less.

Vc ₉₀ -10>V> Vc ₉₀ -30 Formula (2)
50>Vc90> 30 Formula (3)
Vc ₉₀ = 10 ^{(3.69-0.75β)} [° ^C./s ] Formula (4)
β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (5)
The reason for this will be described below.

  Since the structure of the hot-rolled sheet is inherited by the structure after cold-rolling-annealing, the structure control of the hot-rolled sheet is important when the form of the structure after annealing is controlled as in the formula (1). .

  By making the structure of the hot-rolled sheet a composite structure of a ferrite phase, a pearlite phase, and a (bainite phase and a martensite phase), the structure after annealing can satisfy the formula (1). This is to take over the structure of the hot-rolled sheet even in the structure after annealing.

  The structure of the ferrite phase-pearlite phase becomes a coarse austenite phase during annealing, and tends to become an island-like austenite phase during subsequent overaging. On the other hand, the portion that was a bainite phase at the time of hot-rolled sheet becomes a relatively small austenite phase at the time of annealing, so in the subsequent overaging, a coarse island-like austenite phase remains when the bainite transformation proceeds. Hard to do.

Therefore, it is necessary to form a structure in which the pearlite phase and the bainite phase are mixed in hot rolling. In order to produce the hot-rolled sheet, cooling is performed at a speed satisfying the formula (2), and then winding is performed at a temperature of less than 550 ° C. Vc ₉₀ in the formula (2) is an index indicating the hardenability determined by the components of the steel material, and is a cooling rate at which the martensite phase becomes 90% when continuously cooled.

  By satisfy | filling Formula (2), a ferrite phase-pearlite phase can be made 30 to 70%, and when it wound up after that, in a temperature range of 550 degreeC, a bainite phase and a martensite phase are 70- As a result, in the structure after annealing, a form of a retained austenite phase satisfying the formula (1) is obtained.

  Cold rolling may be performed to obtain a necessary plate thickness. However, when the cold rolling rate is less than 30%, the entire structure after annealing or annealing plating becomes coarse grains, and ductility and bendability deteriorate. Therefore, the cold rolling rate is preferably 30% or more.

  The annealing conditions are specified as follows. The temperature rising rate is desirably 2 ° C./s or more. In the case of 2 ° C./s or less, the particle size becomes coarse and ductility and bendability are greatly deteriorated. The maximum heating temperature needs to be Ac1 or more and Ac3 or less. If it is less than Ac1, there is no residual austenite phase, and residual austenitic steel cannot be built. On the other hand, when it becomes Ac3 or more, the grains become coarse, and as described above, the form of the retained austenite phase cannot be controlled to satisfy the formula (1).

  Therefore, the upper limit was set to Ac3. Thereafter, slow cooling may be performed in order to adjust the ferrite phase fraction. Subsequently, after heating or slow cooling, rapid cooling is performed at 1 ° C./s or more. In the case of rapid cooling at less than 1 ° C./s, many pearlite phases appear, and a retained austenite phase and martensite phase cannot be secured, so that the strength and ductility are greatly deteriorated.

  More preferably, in order to suppress pearlite transformation, 20 ° C./s or higher is preferable. On the other hand, even if it exceeds 200 ° C./s, the effect is saturated, and the temperature controllability of the cooling end point temperature, which is most important for the formation of the retained austenite phase, is remarkably deteriorated. For this reason, the cooling rate after annealing is set to 1 ° C./s or more and 200 ° C./s or less on average.

  The cooling end point temperature and subsequent holding is an important technique for controlling the bainite phase formation and determining the C concentration of the residual austenite phase. When the cooling end point temperature is less than 200 ° C., a large amount of martensite phase appears, the steel strength becomes excessively high, and it becomes difficult to leave the austenite phase, resulting in extremely large elongation deterioration. On the other hand, when the temperature exceeds 480 ° C., the bainite transformation is delayed, and in addition, the cementite phase is generated during the holding, and the concentration of C in the retained austenite phase is lowered.

  Therefore, the average cooling stop temperature is 1 ° C./s to 200 ° C./s, and the holding temperature is 200 ° C. to 480 ° C.

  Thereafter, overaging treatment is performed. This temperature need not match the cooling stop temperature. By raising the temperature above the cooling stop temperature, the bainite transformation rate is improved, productivity is increased, or after a part of the martensite phase appears, the tempered martensite phase may be present by raising the temperature. it can. The tempered martensite phase has a lower strength than the martensite phase, but has a higher strength structure than the bainite phase, so it can reduce the hardness difference between the structures while having high strength, This technology leads to improved hole expansion.

  If the retention time of the overaging treatment is less than 1 second, the bainite transformation does not occur sufficiently, C concentration becomes insufficient, and a retained austenite phase in the range specified in the present invention cannot be secured, while the martensite phase is Appearing in large amounts, it greatly exceeds the fraction of martensite phase within the range defined in the present invention.

  On the other hand, if the time exceeds 1000 seconds, a cementite phase is generated in the austenite phase, so that the C concentration tends to decrease. Cementite phase precipitation reduces the fraction of martensite and austenite phases, which in turn degrades strength and ductility. Degradation of spreadability is large.

  Therefore, the holding time is 1 second or more and 1000 seconds or less. It is important to set the fractions of the retained austenite phase and martensite phase to values within the range defined in the present invention.

  The present technology can also be applied to a hot dip galvanized steel sheet, in which case the steel is immersed in a hot dip galvanizing tank after being held at 350 ° C. to 480 ° C. Moreover, this technique can also perform an alloying process after immersion.

  Moreover, the steel plate according to the present invention may be plated. The plating conditions may be a regular method that can ensure corrosion resistance. In other words, 5 to 12% is suitable for the Si concentration in the bath for aluminum plating, and 0.1 to 50% is suitable for the Al concentration in the bath for zinc plating. At this time, the alloying process of plating is performed in the range of 465 ° C. or higher and 580 ° C. or lower. When the temperature is lower than 465 ° C., alloying is insufficient, and when it exceeds 580 ° C., it becomes an overalloy and the corrosion resistance is remarkably deteriorated.

  In addition, when the amount of Si in the steel is large, if the plating is generated on the surface of the steel, the plating may not be successfully applied. Therefore, by setting the dew point to −20 ° C. or higher, the oxide is not inside but on the surface. It is possible to ensure plating properties by precipitating.

Hereinafter, the present invention will be described in detail based on examples, but the conditions in the examples are one example of conditions adopted for confirming the feasibility and effects of the present invention. It is not limited to one condition example. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.
Example 1
In the present invention, tensile tests were performed at 100, 150 and 200 ° C., and those having a small difference between the maximum value and the minimum value of the ductility were evaluated as having high press stability.

  Steel having the component composition shown in Table 1 was manufactured, cooled and solidified, reheated to 1200 ° C., and subjected to normal rough rolling and finish rolling. The subsequent production conditions are shown in Table 2.

  After annealing, 1% skin pass rolling was performed for the purpose of suppressing yield point elongation.

  Tensile properties were evaluated at temperatures of 100 ° C., 150 ° C., and 200 ° C. as described above in C-direction tension of JIS No. 5 tensile test pieces. Identification of ferrite phase, pearlite phase and bainite phase in structure, observation of existing position and average particle diameter (average equivalent circle diameter) The cross section at right angles was corroded and quantified by optical microscope observation at 500 to 1000 times and SEM.

  The method for measuring the residual austenite ratio is carried out on the surface chemically polished to 1/4 thickness from the surface layer of the test material plate, and the strengths of (200) and (211) areas of the ferrite phase by monochromated MoKα rays, The residual austenite phase was quantified from the (200), (220) and (311) area strengths of the austenite phase.

  The classification and fraction of the morphology of the retained austenite phase were calculated based on the obtained image by the SEM-EBSP method. The austenite phase fraction of island-like austenite having a major axis / minor axis of less than 3 was γi, and the others were γl.

  The identification of the martensite phase and the occupancy rate were corroded with a repeller reagent and quantified with an optical microscope of 500 to 1000 times.

  The test results are as shown in Tables 3 and 4.

  Among the test steels, a, b, c, d are out of the range defined by the present invention, and even if the hot rolling conditions, cold rolling conditions and annealing conditions within the ranges defined by the present invention are satisfied, The structure of the range defined by the present invention cannot be satisfied, and the moldability is poor.

In Tables 2 to 4, D1, D4, and D5 indicate that the cooling rate after hot rolling is out of the range defined in the present invention, and the final structure does not satisfy the formula (1) and warm press stability. Is bad. When the cooling rate V exceeds Vc ₉₀ −10 (° C./s ), γl is large and is out of the range defined by the present invention, while it is lower than Vc ₉₀ −40 (° C./s ). Is out of the range defined by the present invention because γi increases.

  FIG. 1 shows the effect of γi / γ on the difference between the maximum and minimum ductility values of tensile tests at 100, 150 and 200 ° C. When the value is within the range defined by the present invention, it can be seen that the difference in ductility is small and the stability of warm forming is excellent.

  FIG. 2 shows the influence of the martensite phase fraction on the difference between the maximum and minimum ductility values of the tensile test at 100, 150 and 200 ° C. When the value is within the range defined by the present invention, the difference in ductility is small and the stability of warm forming is excellent, and the martensite phase fraction is 5%, which is a value outside the range defined by the present invention. Below, stability is bad.

  Further, in the range where the martensite phase exceeds 20%, the stability is relatively high, but there is no residual austenite phase and the ductility is poor.

  In the treatment numbers D6 and D9 in Tables 2 to 4, the annealing temperature is outside the range defined in the present invention, and the material is bad. The treatment number D6 is less than Ac1, no residual austenite phase exists in the final structure, and the ductility is poor. Moreover, in the case of process number D9, it becomes Ac3 or more, does not satisfy Formula (1), and the stability of ductility at the time of warm forming is low.

  According to the present invention, due to the TRIP effect of retained austenite, a high-strength thin steel sheet capable of ensuring extremely high elongation and stably ensuring high formability during pressing despite high strength. Can also be applied to electroplated steel sheets, and thus contributes greatly to the machinery industry including the steel industry and the automobile industry.

Claims (10)

% By mass
C: 0.05% or more, 0.35% or less,
Si: 0.05% or more, 2.0% or less,
Mn: 0.8% or more, 3.0% or less,
P: 0.0010% or more, 0.1% or less,
S: 0.0005% or more, 0.05% or less,
N: 0.0010% or more and 0.010% or less Al: 0.01% or more and 2.0% or less, and has a steel composition composed of the balance iron and inevitable impurities, the microstructure is area ratio, ferrite The total austenite phase is 10% or more and 90 % or less, the residual austenite phase is 5% or more and 30% or less in the area ratio, the martensite phase is 5% or more and 20% or less in the area ratio, and the residual austenite phase is lath. In the elongation and press forming stability, the area ratio γi of the island-like residual austenite phase and the area ratio γ of the total residual austenite phase satisfy the following formula (1). Excellent high strength steel plate.
0.7 ≧ γi / γ ≧ 0.3 Formula (1) The steel sheet according to claim 1, further comprising:
A high strength thin steel sheet excellent in elongation and press forming stability, characterized in that an average C concentration in an austenite phase is 0.7% or more and 1.5% or less in mass%.   In the steel sheet according to claim 1 or claim 2, the total of the ferrite phase and the bainite phase is 50% or more in volume fraction with respect to the entire structure. Excellent high-strength thin steel sheet. In the steel sheet according to any one of claims 1 to 3, further in mass%,
Mo: A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing 0.02% or more and 0.5% or less. In the steel sheet according to any one of claims 1 to 4, further,
% By mass
Nb: 0.01% or more, 0.10% or less,
Ti: 0.01% or more, 0.20% or less,
V: 0.005% or more, 0.10% or less,
Cr: 0.1% or more, 5.0% or less,
W: A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more selected from 0.01% or more and 5.0% or less. In the steel plate described in any one of claims 1 to 5,
% By mass
A high-strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more selected from Ca, Mg, Zr, and REM in an amount of 0.0005% to 0.05%. In the steel sheet according to any one of claims 1 to 6,
% By mass
Cu: 0.04% or more, 2.0% or less,
Ni: 0.02% or more, 1.0% or less,
B: 0.0003% or more, 0.007% or less A high strength thin steel sheet excellent in elongation and press forming stability, characterized by containing one or more kinds selected from shells. It is a manufacturing method of the steel plate given in any 1 paragraph of Claims 1-7,
For casting slabs, after finishing casting, or once cooled to 1100 ° C or lower, when re-heating to 1100 ° C or higher and performing hot rolling, the finishing temperature ends at 850 ° C or higher and 970 ° C or lower. Then, after cooling at a speed V satisfying the formula (2), winding is performed in a temperature range of 550 ° C. or less, pickling, and then cold rolling of 30% or more, and the highest temperature during annealing is Ac1. As described above, after annealing at Ac3 or lower, the sample is cooled to a temperature range of 200 ° C. or higher and 480 ° C. or lower at an average cooling rate of 1 ° C./second or higher and 200 ° C./second or lower. A method for producing a high-strength thin steel sheet excellent in elongation and press forming stability, characterized in that holding is performed below.
Vc ₉₀ -10>V> Vc ₉₀ -30 Formula (2)
50> Vc ₉₀ > 30 Formula (3)
Vc ₉₀ = 10 ^{(3.69-0.75β)} [° ^C./s ] Formula (4)
β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (5) A method for producing a steel sheet according to any one of claims 1 to 7,
For casting slabs, after finishing casting, or once cooled to 1100 ° C or lower, when re-heating to 1100 ° C or higher and performing hot rolling, the finishing temperature ends at 850 ° C or higher and 970 ° C or lower. Then, after cooling at a speed V satisfying the formula (2), winding is performed in a temperature range of 550 ° C. or less, pickling, and then cold rolling of 30% or more, and the highest temperature during annealing is Ac1. As described above, after annealing at Ac3 or lower, the sample is cooled to a temperature range of 200 ° C. or higher and 480 ° C. or lower at an average cooling rate of 1 ° C./second or higher and 200 ° C./second or lower. A method for producing a high-strength thin steel sheet excellent in elongation and press forming stability, characterized in that the following holding is performed, followed by immersion in a hot dip galvanizing bath.
Vc ₉₀ -10>V> Vc ₉₀ -30 Formula (2)
50> Vc ₉₀ > 30 Formula (3)
Vc ₉₀ = 10 ^{(3.69-0.75β)} [° ^C./s ] Formula (4)
β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (5)   The method for producing a steel sheet according to claim 9, further comprising an alloying treatment in a range of 465 ° C or higher and 580 ° C or lower after immersion in a hot dip galvanizing bath, and excellent in elongation and press forming stability. A method for producing high strength thin steel sheets.

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