JP2006233294A - Low yield-ratio high-strength steel sheet having excellent baking hardening property and its production method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a low yield-ratio high-strength steel sheet having a high tensile strength of 440 to 640 MPa and also a low yield ratio satisfying ≤55% YR, and, further, while maintaining excellent baking hardenability, having formability and surface quality applicable to automobile inside and outside plates, and to provide its production method. <P>SOLUTION: The low yield-ratio high-strength steel sheet comprises, by mass, 0.01 to <0.040% C, ≤1.0% Si, 0.3 to 1.6% Mn, <0.07% P, ≤0.03% S, 0.01 to 0.10% Al and ≤0.01% N, satisfying 1.3%≤Mn+1.29Cr+3.29Mo≤2.1% (wherein, Cr:≤0.5% and Mo:≤0.5%), and the balance Fe with inevitable impurities, and has a steel sheet structure comprising a ferritic phase of ≥70% and a martensitic phase of 1 to 15% by a volume fraction. <P>COPYRIGHT: (C)2006,JPO&NCIPI

Description

  TECHNICAL FIELD The present invention relates to a low yield ratio high strength steel sheet having excellent bake hardenability and a method for producing the same, which can be applied to, for example, automotive outer plate parts mainly formed by overhanging such as doors, hoods, and roofs.

  2. Description of the Related Art In recent years, steel sheets for automobiles have been made thinner and thinner for the purpose of improving fuel efficiency by reducing the weight of the vehicle body and increasing the strength for improving safety. Generally, increasing the strength of a steel sheet has the problem of causing deterioration of formability, but in order to overcome this problem, it is soft and easy to form at the time of press forming, and an increase in strength is obtained in the baking coating process after press forming. Steel plates having so-called bake hardening (BH) properties have been developed. In particular, in the case of an automotive outer plate part, the BH property is regarded as important because an effect of increasing the resistance against local residual dent of the final product can be obtained by utilizing this characteristic. This BH steel sheet is a technology that hardens due to the strain aging phenomenon using solid solution C and N. On the other hand, the aging of the remelted C and N progresses even at room temperature, resulting in an increase in yield strength and yield point elongation. Therefore, when it is held for a long time, stretcher strain is generated during molding, which may impair the surface appearance. In order to achieve both aging resistance and BH characteristics which are such contradictory characteristics, for example, Patent Document 1 discloses that C is obtained by optimizing the addition amount of at least two kinds of Si, Mn, and P in low carbon steel. A method for suppressing segregation to dislocation and improving aging resistance is disclosed.

  Patent Document 2 discloses a method of imparting a high BH property with non-aging at room temperature by utilizing a low-temperature transformation phase capable of introducing a large amount of mobile dislocations during transformation in ultra-low carbon steel. Yes. Furthermore, Patent Document 3 discloses a method for obtaining a steel sheet having an ultra-high strength and excellent BH property by increasing the volume ratio of the martensite phase.

However, each of the above conventional techniques has the following problems.
The method of Patent Document 1 defines the upper limit of solid solution strengthening elements such as Si, Mn, and P, and uses a low-carbon steel plate to which no alloying element is added. It is difficult to ensure a strength of 440 MPa or more. In addition, since the yield ratio is high, for example, the load during pressing is large and not only the pressing itself is difficult, but also so-called shape defects such as surface distortion after molding and an increase in the amount of springback occur. It is difficult to say that it has press formability that can be applied to automotive outer plate parts that are subject to the above.

  In addition, since the method of Patent Document 2 is based on extremely low carbon steel, in order to achieve the required high strength of 440 MPa or more, a large amount of a solid solution strengthening element such as Si or P is added. In order to increase the hardness of the low temperature transformation phase, it is necessary to add a large amount of Mn, Cr or the like. Further, from the viewpoint of stabilizing the austenite phase and obtaining a single phase structure of a low temperature transformation phase, it is necessary to add a large amount of Mn, Cr and the like.

  However, when Si and P are added in a large amount, red scale is generated, non-plating, chemical conversion treatment failure, and the like occur, and the obtained steel sheet is hardly in a state of having good surface properties. Furthermore, since the steel sheet for an outer panel intended by the present invention is required to have extremely excellent surface quality, addition of a large amount of Mn and Cr is disadvantageous from the viewpoint of surface properties and causes a significant increase in manufacturing cost.

Moreover, although the method of patent document 3 makes it possible to give high BH property by producing the martensite phase of 30 to 95% by volume ratio, the increase in a martensite phase fraction is simultaneously high intensity | strength. And formability deteriorates.
Japanese Patent Laid-Open No. 5-105985 JP-A-6-116650 JP-A-3-87320

  The present invention has been made in view of such circumstances, and has a tensile strength as high as 440 MPa or higher and 640 MPa or lower, a low yield ratio of YR of 55% or lower, and further excellent bake hardenability, while being an automobile. An object of the present invention is to provide a low-yield-ratio high-strength steel sheet having formability and surface quality applicable to inner and outer plate applications and a method for producing the same.

In order to solve the above-described problems, the present invention provides, in mass%, C: 0.01% or more and less than 0.040%, Si: 1.0% or less, Mn: 0.3 to 1.6%, P: 0 0.07% or less, S: 0.03% or less, Al: 0.01 to 0.10%, N: 0.01% or less, further satisfying the following formula (1), the balance being Fe A low yield ratio high-strength steel sheet with excellent bake-hardening characteristics, characterized by comprising a ferrous phase and a martensite phase content of 70% or more and a martensite phase as a steel sheet structure. To do.
1.3% ≦ Mn + 1.29Cr + 3.29Mo ≦ 2.1% (1)
However, in the formula (1), Mn, Cr, and Mo represent the content of each element, Cr: 0.5% or less, Mo: 0.5% or less.

  Moreover, this invention is the said steel plate in the mass%, and also V: 0.001-0.1%, B: 0.0001-0.01%, Ti: 0.001-0.1%, Nb : Provided is a low yield ratio high-strength steel sheet excellent in bake hardening characteristics, characterized by containing one or more of 0.001 to 0.1%.

Furthermore, the present invention is a method of cold rolling a steel hot-rolled sheet having any one of the above-mentioned chemical composition and annealing in a temperature range of Ac ₁ point or more and Ac ₃ point or less, and thereafter 3-20 ° C / sec. Low yield ratio high strength steel plate with excellent bake hardening characteristics characterized by primary cooling to a temperature range of 550 to 750 ° C. at a cooling rate, and secondary cooling to 200 ° C. or less at a cooling rate of 100 ° C./sec or more. A manufacturing method is provided.

  The present inventors have intensively studied the component design for obtaining a low-yield ratio high-strength steel sheet having good bake-hardening characteristics, a tensile strength of 440 to 640 MPa, and a YR of 55% or less.

  As a result, when the amount of added C increases, it becomes easy to increase the amount of dissolved C, but at the same time, the second phase such as martensite phase and carbide tends to increase, and the workability deteriorates. From the viewpoint of avoiding local strain concentration at the time of machining by C, it is effective that the C amount is the minimum necessary amount for a predetermined strength, and the martensite phase fraction is 1 to 15% by volume ratio. I found it.

  Moreover, in order to obtain a martensite phase with the minimum necessary amount of C, it is extremely important to control the hardenability improving elements Mn, Cr, and Mo. Mn, Cr, and Mo can stabilize the γ phase to obtain an effect of improving hardenability, but the degree of the effect is different. Therefore, paying attention to the weighted total amount reflecting this degree, as a result of repeated studies, in order to obtain a martensite structure and satisfy the desired bake hardenability and formability, Mn, Cr, Mo It was newly found out that there is an optimum range for the weighted total amount Mn + 1.29Cr + 3.29Mo.

  Furthermore, as a manufacturing condition for maximizing the bake hardening characteristics, it is effective to increase the secondary cooling rate in the continuous annealing process, and at the same time, the amount of addition of the hardenability improving element can be reduced. Therefore, it is possible to have excellent processability and surface properties while having excellent BH characteristics.

  From the above examination results, the conditions for obtaining a low yield ratio high strength steel sheet having good bake-hardening characteristics, tensile strength of 440 to 640 MPa, and YR of 55% or less became clear.

  Further, in the steel sheet according to the present invention, the required strength is mainly obtained by strengthening the structure by the martensite phase, so that a large amount of Si and P, which are solid solution strengthening elements, is not required, and superior in terms of surface properties. It is.

  The present invention having the above-described configuration has been completed based on the above knowledge, and in a semi-very low carbon steel sheet, in order to improve bake hardening characteristics, a hardenability improving element is optimized and formability is improved. From this point of view, the martensite phase fraction is controlled within a specific range, but such a concept does not exist in the prior art.

  According to the present invention, it has formability suitable for automobile outer plate parts mainly composed of overhang molding such as doors, hoods, roofs, and the like, and can ensure equivalent dent resistance when promoting weight reduction. Provided is a low-yield ratio high-strength steel sheet having excellent bake hardenability and capable of being industrially manufactured with existing equipment and having a tensile strength of 440 to 640 MPa and a YR of 55% or less. be able to.

Hereinafter, the present invention will be described in detail.
First, the chemical component composition of the steel sheet will be described. In the following description,% indicates mass%.

C: 0.01% or more and less than 0.040% C is one of the extremely important elements in the present invention, and is very effective in generating a low-temperature transformation phase and increasing the strength. However, when the C content is 0.040% or more, the workability is significantly lowered, and the weldability is further deteriorated. Therefore, the C content is less than 0.040%. Preferably it is 0.03% or less. On the other hand, in order to form a low temperature transformation phase having a constant volume ratio and introduce sufficient mobile dislocations or to secure a certain amount of solid solution C, it is necessary to contain a certain amount of C. 01% or more.

Si: 1.0% or less Si is an effective element for stably obtaining a composite structure. However, when the Si content exceeds 1.0%, the surface properties and the chemical conversion treatment properties are remarkably lowered. Therefore, the Si content is 1.0% or less. Preferably it is 0.8% or less, More preferably, it is 0.5% or less.

Mn: 0.3 to 1.6%
Mn is an extremely important element for the formation of a low-temperature transformation phase. The effect of improving hardenability and the fixing of S in the steel as MnS cause maturation that occurs due to the grain boundary embrittlement of S. It is an essential element having the action of preventing slab cracking during hot rolling, and it is necessary to contain 0.3% or more in order to effectively exhibit the above action. It is preferably 0.6% or more, and more preferably 1.0% or more from the viewpoint of obtaining a low yield ratio by obtaining a hard low-temperature transformation phase. However, if the Mn content exceeds 1.6%, the slab cost increases significantly, and the yield strength that increases with the appearance defect and the increase in strength causes the press formability to deteriorate and the shape defect to occur. Therefore, the Mn content is set to 1.6% or less.

P: Less than 0.07% P is an effective element for increasing the strength and stabilizing the low-temperature transformation phase. However, when the P content is 0.07% or more, the alloying rate of the galvanized layer is reduced, which causes plating defects and non-plating, and segregates at the grain boundaries of the steel sheet, resulting in resistance to secondary work brittleness. Deteriorate. Therefore, the P content is less than 0.07%.

S: 0.03% or less S is segregated at the grain boundary during hot rolling and causes slab cracking, so the surface flaw generation rate is increased. Therefore, by adding Mn, S is fixed as MnS. However, excessive MnS serves as a starting point of voids during processing, which causes a decrease in workability. Accordingly, it is desirable that the S content is small. If the S content exceeds 0.03%, the workability is remarkably deteriorated, so the S content is set to 0.03% or less.

Al: 0.01 to 0.1%
Al has a function of reducing inclusions in the steel as a deoxidizing element. However, when the Al content is less than 0.01%, the above-described action cannot be stably obtained. On the other hand, if the amount of Al exceeds 0.1%, cluster-like alumina inclusions increase and workability deteriorates. Therefore, the Al content is set within a range of 0.01 to 0.1%.

N: 0.01% or less N is better from the viewpoint of workability and aging properties. When the N content exceeds 0.01%, ductility and toughness deteriorate due to the formation of excessive nitride. Therefore, the N content is 0.01% or less.

Mn + 1.29Cr + 3.29Mo: 1.3-2.1%
(However, Cr: 0.5% or less, Mo: 0.5% or less)
Mn, Cr, and Mo are elements that improve the hardenability, and it is extremely important to optimally control in order to generate a low-temperature transformation phase. Mn + 1.29Cr + 3.29Mo is a weighted total amount weighting the hardenability contribution of each of these elements, and when this weighted total amount is less than 1.3%, DP ( Dual Phase) It becomes difficult to obtain a structure, and the aging resistance and BH characteristics are impaired due to the lack of movable dislocations introduced during martensitic transformation. Furthermore, the yield ratio is high, and not only the press working itself becomes difficult, but also shape defects are likely to occur. From the viewpoint of stably obtaining excellent BH characteristics, 1.6% or more is preferable. As the weighted total amount is increased, the number of movable dislocations introduced at the time of transformation increases, and a good BH characteristic can be obtained at the time of baking coating, but the weighted total amount exceeds 2.1%. Not only saturates the effect, but also causes an increase in production cost due to the addition of a large amount of alloying elements, and the yield strength increases as the strength increases, and the press formability is significantly lowered. In the present invention, Mn is an essential addition, and Cr and Mo are added as necessary. However, when the content of Cr and Mo exceeds 0.5%, not only the effect is saturated but also the production cost is increased. Therefore, when these are added, the upper limit is 0.5%.

  In the present invention, in addition to the above, one or more of V, B, Ti, and Nb may be contained in the following range as a selection component.

V: 0.001 to 0.1%
V is an element for improving hardenability, and can be added as necessary to stably generate a low-temperature transformation phase, and the effect is effectively exhibited at 0.001% or more. However, even if the content exceeds 0.1%, the effect is not only saturated but also disadvantageous in terms of cost. Therefore, when adding V, the content is made 0.001 to 0.1%.

B: 0.0001 to 0.01%
B is an element effective for improving hardenability, and can be added as necessary to stably obtain a low-temperature transformation phase, and the effect is effectively exhibited at 0.0001% or more. However, even if the content exceeds 0.1%, an effect commensurate with the cost cannot be obtained. Therefore, when adding B, the content is made 0.0001 to 0.01%.

Ti, Nb: 0.001 to 0.1% respectively
Ti and Nb are effective elements for forming carbonitrides and improving deep drawability, and the effect is effectively exhibited at 0.001% or more. However, even if the content exceeds 0.1%, the effect is saturated, the recrystallization temperature during annealing is increased, and the productivity is lowered. Therefore, when adding Ti and Nb, the content is made 0.001 to 0.1%, respectively.

  In addition to the above elements and the remaining Fe, various impurity elements and trace addition elements essential in the manufacturing process are inevitably mixed in the manufacturing process. Such inevitable impurities are particularly effective for the effects of the present invention. It does not affect and is allowed.

Next, the structure of the steel plate will be described.
In the present invention, in addition to satisfying the above chemical component composition, the steel sheet structure needs to have a volume fraction of 70% or more of ferrite phase and 1 to 15% of martensite phase.

  In the present invention, in order to have high formability, it is important to obtain a structure mainly composed of a soft ferrite phase and having a martensite phase fraction suppressed as compared with conventional DP steel. However, the martensite phase requires a certain amount of martensite phase because the mobile dislocations introduced at the time of transformation are fixed by solid solution C and N, so that a high BH property can be obtained. Need to control rate. From such a viewpoint, the volume fraction of the ferrite phase is set to 70% or more, and the volume fraction of the martensite phase is set to 1 to 15%. The volume fraction of the ferrite phase is preferably 80% or more. Further, the volume fraction of the martensite phase is preferably 10% or less, more preferably 8% or less from the viewpoint of moldability. Other structures may include residual γ phase, bainite phase, and carbide.

Next, manufacturing conditions will be described.
The steel sheet of the present invention is obtained by cold rolling a steel hot-rolled sheet having the above-mentioned chemical composition and annealing in a temperature range of Ac ₁ point or more and Ac ₃ point or less, and then at a cooling rate of 3 to 20 ° C./sec. It can be obtained by performing primary cooling to a temperature range of 550 to 750 ° C., and further performing secondary cooling to 200 ° C. or less at a cooling rate of 100 ° C./sec or more.

Annealing temperature: Ac ₁ point or more Ac ₃ point or less In order to obtain the microstructure of the ferrite phase and the low temperature transformation phase, the annealing temperature needs to be ripened to an appropriate temperature. When the annealing temperature is less than Ac ₁ point, an austenite phase is not generated and a low temperature transformation phase cannot be obtained. On the other hand, when the annealing temperature exceeds the Ac ₃ point, the ferrite phase is entirely austenitized, so that the properties such as formability obtained by recrystallization deteriorate. Therefore, annealing temperature shall be Ac ₁ point or more and Ac ₃ point or less. Preferably, the temperature is Ac ₁ point or more and Ac ₁ point + 100 ° C. or less, and more preferably Ac ₁ point + 80 ° C. or less.

Primary cooling: cooling rate 3-20 ° C / sec, cooling stop temperature 550-750 ° C
The cooling rate and the cooling stop temperature of the primary cooling must be appropriately controlled in order to suppress pearlite precipitation and to secure the volume ratio of austenite. When the cooling rate exceeds 20 ° C./sec, the two-phase separation of the ferrite phase and the austenite phase may not sufficiently proceed. In that case, a hard low temperature transformation phase cannot be obtained, and desired characteristics cannot be obtained. On the other hand, when the primary cooling rate is less than 3 ° C./sec, pearlite transformation occurs and formability deteriorates. Therefore, the cooling rate of the primary cooling is 3 to 20 ° C./sec. On the other hand, when the cooling stop temperature is higher than 750 ° C., secondary cooling is performed in a state where the volume ratio of austenite is high, so that the volume ratio of the low-temperature transformation phase becomes large and the workability is deteriorated. On the other hand, when the cooling stop temperature is lower than 550 ° C., bainite transformation is likely to occur during cooling, and it becomes difficult to obtain a desired martensite volume ratio. Therefore, the cooling stop temperature of the primary cooling is set in the range of 550 to 750 ° C.

Secondary cooling: Cooling rate of 100 ° C./sec or more and cooling stop temperature of 200 ° C. or less Martensite when the cooling rate of secondary cooling is less than 100 ° C./sec or the cooling stop temperature exceeds 200 ° C. The driving force for the transformation becomes insufficient, and the desired martensite phase fraction cannot be obtained. In addition, when the cooling rate is insufficient, the amount of solid solution C in the ferrite phase decreases, and desired BH characteristics cannot be obtained. In order to obtain the BH characteristics required by the present invention, the above-described cooling rate is required. However, this secondary cooling condition is also a subcooling condition necessary and sufficient for obtaining the driving force for martensitic transformation. It is possible to reduce the amount of alloying elements added to DP steel, and at the same time, good workability can be obtained while having excellent BH characteristics. Therefore, the secondary cooling rate is set to 100 ° C./sec or more, and the cooling stop temperature is set to 200 ° C. or less. Preferably, the cooling rate is 200 ° C./sec or more and the cooling stop temperature is 100 ° C. or less. At this time, the cooling method is not limited, but it is possible to satisfy such secondary cooling conditions by, for example, quenching in jet water. In addition, when it hold | maintains at primary cooling stop temperature, a pearlite transformation will start and a desired martensite fraction may not be obtained, Therefore It is preferable to perform secondary cooling following primary cooling. It is desirable to carry out within 10 seconds from the end of the next cooling.

  Moreover, it is preferable to perform a tempering process at 400 ° C. or lower in order to recover toughness after the secondary cooling. When the tempering treatment temperature exceeds 400 ° C., the strength is remarkably lowered, and further, the solid solution C is precipitated, so that the BH characteristics are lowered. Desirably, it is 300 degrees C or less.

  The cold-rolled steel sheet obtained as described above may be further subjected to temper rolling. However, the addition of large pressure control causes a decrease in ductility and an increase in yield stress, so a pressure control rate of 1.5% or less is desirable.

In the present invention, when the slab is hot-rolled, it may be rolled after being reheated in a heating furnace, or may be directly fed without being heated. Moreover, it is good to implement hot rolling finish rolling temperature at ₃ or more points of Ar. The cold pressure rate may be 50 to 85% within the normal operating range.

  In the present invention, the cold-rolled steel sheet as described above may be subjected to electrogalvanizing or hot-dip galvanizing, and in the case of hot-dip galvanized steel, an alloying treatment may be performed. In addition, such a plated steel sheet may be further subjected to an organic film treatment after plating.

Examples of the present invention will be described below.
Example 1
No. shown in Table 1. After melting 1 to 15 steels, slabs were produced by continuous casting. Of these, No. 1 to 9 are within the scope of the present invention. As for 10-15, any of C amount, Mn amount, Cr amount, and Mn + 1.29Cr + 3.29Mo is outside the scope of the present invention.

After heating these slabs to 1200 ° C, a finishing press is performed at a temperature of ₃ points or more at Ar, a hot-rolled sheet having a thickness of 3.2 mm at a winding temperature of 550 ° C, pickled, and then cooled to a thickness of 0.75 mm. Rolled for a while. Subsequently, in the continuous annealing step, annealing was performed at 770 ° C., primary cooling was performed at a cooling rate of 8 ° C./sec and a cooling stop temperature of 600 ° C., and then secondary cooling was performed to room temperature by water cooling at a cooling rate of 1000 ° C./sec. . Thereafter, a tempering treatment was performed for 10 minutes at 150 ° C. 1-15 annealed plates were obtained.

  About the obtained annealing board, evaluation of mechanical characteristics, evaluation of BH property, structure observation, and evaluation of formability were performed. Mechanical properties were evaluated by collecting JIS No. 5 tensile test pieces from the annealed plates. The BH property was evaluated based on the amount of increase in YP when a JIS No. 5 tensile test piece was sampled and subjected to heat treatment at 170 ° C. for 20 minutes after adding 2% pre-strain and then subjected to a tensile test again. In addition, the specimen was subjected to Nital corrosion, and the central portion of the plate thickness was continuously observed at a magnification of 2000 times by SEM observation of a 100 μm × 200 μm field of view, and the ferrite phase fraction and martensite phase fraction were measured.

  Further, the formability was evaluated by using a punch with an 800 mmR cylindrical surface after collecting a sample having a blank angle of 160 mm, and ◯ when no cracks or wrinkles occurred, and x when occurring.

  Table 2 shows the mechanical properties, BH properties, ferrite phase fraction, martensite phase fraction, and moldability evaluated as described above. No. which is a comparative example. No. 10 had a C content lower than the range of the present invention, so a martensitic phase structure was not obtained and the BH characteristics were inferior. Moreover, No. which is a comparative example. No. 11 had a C content exceeding the range of the present invention, had a high martensite phase fraction, and was inferior in moldability.

  No. C amount is outside the specified range. Except for 10 and 11, the influence of the total weight Mn + 1.29Cr + 3.29Mo of Mn, Cr and Mo on the BH amount was investigated. The result is shown in FIG. As shown in FIG. 1, Mn + 1.29Cr + 3.29Mo was 1.3% or more and the BH amount was 60 MPa or more. The amount of BH of 60 MPa or more is the amount of BH necessary for securing equivalent dent resistance properties when reducing the weight by applying high-tensile steel to steel plates currently applied to the outer panel of automobiles. is there. Moreover, when Mn + 1.29Cr + 3.29Mo is 1.6% or more, the amount of BH is 80 MPa or more, which is more preferable. On the other hand, No. 1 with Mn + 1.29Cr + 3.29Mo of less than 1.3%. In No. 14, a martensite phase structure was not obtained, the amount of BH was reduced, and press formability was inferior due to high YR. In addition, No. in which Mn + 1.29Cr + 3.29Mo exceeds 2.1%. No. 13 shows that the effect of improving the BH characteristics is saturated, and the formability is inferior due to the addition of a large amount of alloy elements.

  In addition, each of No. 1 containing a large amount of Mn and Cr. 12, no. No. 15 had good BH characteristics but poor moldability.

  On the other hand, in addition to Mn + 1.29Cr + 3.29Mo, the other requirements are No. 1 to 9 have excellent BH characteristics with a BH amount of 60 MPa or more as described above, a tensile strength within the range of 440 to 640 MPa, and a YR of 55% or less and a low yield ratio. And it was confirmed that it has good moldability.

(Example 2)
No. shown in Table 1. After melting the steels 1 to 3, a slab was formed by continuous casting. The slab was hot-rolled and cold-rolled under the same conditions as in Example 1 and then subjected to continuous annealing under the conditions shown in Table 3. Tempering treatment for 10 minutes at 150 ° C. is performed. An annealed plate of 16 to 31 was obtained. About each of the obtained annealing board, evaluation of a mechanical characteristic, evaluation of BH property, structure | tissue observation, and evaluation of a moldability were performed like Example 1. FIG. Table 3 shows the evaluation results of mechanical properties, BH properties, ferrite phase fraction, martensite phase fraction, and moldability.

  As shown in Table 3, No. 1 as an example of the present invention. 16, 17, 21, 22, 23, 24, 27, and 28 all have excellent BH characteristics with a tensile strength in the range of 440 to 640 MPa and a BH amount of 60 MPa or more. Has a low yield ratio of 55% or less. However, no. Nos. 23 and 24 are No. Compared with 21 and 22, the yield ratio was high and the BH content was low. This is no. No. 23 secondary cooling stop temperature is as high as 180 ° C. This is considered to be because the secondary cooling rate of 24 was 150 ° C./sec, which was slower than other examples of the present invention, and the secondary cooling rate was 200 ° C./sec or more and the secondary cooling stop temperature was preferably 100 ° C. or less. .

  On the other hand, No. which is a comparative example. 18-20, 25, 26, 29-31 were all inferior in properties because the annealing conditions were outside the scope of the present invention, although the component composition was within the scope of the present invention. For example, no. Nos. 19, 25, and 30 show a martensite phase structure, which shows deterioration in BH characteristics and an increase in yield ratio. No. No. 29 had a high martensite phase fraction and poor moldability. No. In Nos. 18 and 20, the austenite phase was not stabilized and a large number of carbides were precipitated, so the yield ratio was high and the moldability was poor. No. No. 26 had a decreased amount of BH due to insufficient cooling. No. Since No. 31 was not cooled to 200 ° C. or lower during the secondary cooling, a hard martensite phase was not obtained, the yield ratio (YR) was high, and the moldability was poor.

The figure which shows the influence which Mn + 1.29Cr + 3.29Mo which is a weighted total amount of Mn, Cr, and Mo has on the amount of BH.

Claims (3)

In mass%, C: 0.01% or more and less than 0.040%, Si: 1.0% or less, Mn: 0.3 to 1.6%, P: less than 0.07%, S: 0.03% Hereinafter, Al: 0.01 to 0.10%, N: 0.01% or less, further satisfying the following formula (1), the balance consisting of Fe and inevitable impurities, A low-yield-ratio high-strength steel sheet excellent in bake hardening characteristics, comprising a volume fraction of 70% or more of a ferrite phase and 1 to 15% of a martensite phase.
1.3% ≦ Mn + 1.29Cr + 3.29Mo ≦ 2.1% (1)
However, in the formula (1), Mn, Cr, and Mo represent the content of each element, Cr: 0.5% or less, Mo: 0.5% or less.   Further, V: 0.001 to 0.1%, B: 0.0001 to 0.01%, Ti: 0.001 to 0.1%, Nb: 0.001 to 0.1% The low yield ratio high-strength steel sheet having excellent bake hardening characteristics according to claim 1, comprising one or more of them. A steel hot-rolled sheet having the chemical composition according to claim 1 or claim 2 is cold-rolled and annealed in a temperature range of Ac ₁ point or more and Ac ₃ point or less, and thereafter 3-20 ° C / sec. Low yield ratio high strength steel plate with excellent bake hardening characteristics characterized by primary cooling to a temperature range of 550 to 750 ° C. at a cooling rate, and secondary cooling to 200 ° C. or less at a cooling rate of 100 ° C./sec or more. Manufacturing method.

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