JP2005076056A - NON-HEAT TREATED Cu PRECIPITATION TYPE HIGH TENSILE STRENGTH STEEL SHEET EXCELLENT IN DUCTILITY, AND ITS PRODUCTION METHOD - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a Cu precipitation type high tensile strength steel sheet obtainable as non-heat treated without requiring aging treatment at high temperatures for a long time, excellent in ductility and having a tensile strength (TS) of ≥570 MPa. <P>SOLUTION: The steel sheet satisfies a prescribed componential composition, also satisfies expression (1): -0.1≤([Mn]-[Si+P])/[Cu]≤1.4 äwherein, [Mn]=Mn content(mass%)/54.9, [Si+P]=(Si content(mass%)/28.0)+(P content(mass%)/31.0), and [Cu]=Cu content(mass%)/63.5}, and also has a metallic structure of a ferrite single phase or of ferrite in ≥90%, by area ratio, and pearlite and/or bainite. The steel sheet is obtainable by performing cooling in the temperature range of at least <580°C at a cooling rate of ≤0.2°C/s after hot rolling. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention relates to a non-tempered Cu precipitation-type high-tensile steel plate excellent in ductility and a method for producing the same, and a steel plate having a tensile strength (TS) of 570 MPa or more used as a pillar for buildings or bridges, and for automobiles. The present invention relates to a steel plate applicable to a thin steel plate (hot-rolled steel plate) having the same strength used for a body and the like and a manufacturing method thereof.

  Conventionally, the following methods have been adopted as techniques for increasing the ductility of high-strength steel sheets. In order to ensure high strength, the microstructure is mainly composed of bainite or martensite, and after heating to a two-phase region of γ (austenite) + α (ferrite) in the manufacturing process, quenching is performed to obtain a hard phase (bainite or martensite). ) Introducing a ferrite structure, which is a soft phase, followed by tempering to remove thermal stresses and dislocations introduced during quenching after two-phase heat treatment, thereby reducing ductility, toughness and yield ratio There is a method of reducing the above (for example, Non-Patent Document 1, Patent Document 1). However, in the above method, it is necessary to provide a heat treatment step in a two-phase region, resulting in a decrease in productivity and an increase in manufacturing cost.

  Another method to increase the ductility of high-strength steel sheets is to use the TRIP phenomenon during tensile deformation by forming isothermal bainite transformation in the temperature range of 350-450 ° C after rolling and forming residual γ in the steel. (For example, Patent Documents 2 to 4). However, even in this method, it is necessary to provide a heat treatment step for isothermal bainite transformation, which causes an increase in cost.

  In addition, there is a method of increasing the strength by precipitating and strengthening the precipitation of carbides in the ferrite main structure. Specifically, the strength is increased by carrying out a tempering treatment to precipitate precipitates, and the strength and ductility are both achieved by reducing the dislocation density in the ferrite structure. However, in the case of precipitation strengthening, it is known that the size and dispersion state of precipitates have a significant effect on strength, and particularly when carbide is used, it tends to precipitate on grain boundaries, and the distribution state is uneven. It is known that it is easy to become. Also, since the diffusion of carbon is very fast, it is difficult to control the growth of precipitates. Therefore, non-uniformity in the dispersion state of carbides and variations in the size of precipitates are likely to cause variations in mechanical properties in the same plane of the steel sheet, in order to suppress non-uniformity in the size and distribution state of the carbides, This is complicated because the heat treatment temperature and time must be strictly controlled.

  By the way, a Cu precipitation type high strength steel sheet whose base material structure is a ferrite structure is a high strength steel sheet mainly composed of a hard structure such as bainite and martensite as described above, precipitation by carbides represented by TiC, NbC, and the like. Compared with a high-strength steel plate using reinforcement, there is little variation in strength and excellent ductility. This is different from the case where the Cu particles are carbide, softer than iron as a base structure, and even when sheared by dislocation at the time of deformation, it deforms plastically without breaking, so dislocation loops hardly remain around the precipitate, As a result, it is said that work hardening hardly occurs.

However, in general, in order to obtain a high-strength steel sheet using Cu precipitation strengthening, it is necessary to perform an aging treatment for a long time of 1 to 5 hours in a temperature range of 500 to 600 ° C. in order to sufficiently precipitate Cu. There is a problem that productivity is bad.
JP-A-53-23817 JP-A 61-272321 JP 62-188729 A JP-A-4-333524 Kunitake Tatsuro and Otani Yasuo “Mixed Structure and Toughness in Steel”, Journal of the Japan Institute of Metals, 1975, Vol. 14, No. 9, p. 689-697

  The present invention has been made in view of such circumstances, and its purpose is to maintain a non-tempered condition without requiring a heat treatment such as an aging treatment requiring a long time, and a tensile strength (TS) of 570 MPa. An object of the present invention is to provide a method for obtaining a Cu precipitation type high strength steel sheet having excellent ductility and the Cu precipitation type high strength steel sheet.

The non-tempered Cu precipitation type high-tensile steel sheet having a tensile strength of 570 MPa or more according to the present invention is in mass% (hereinafter the same), C: 0.06% or less (0% not included), Si: 0.1 0.6%, Mn: 0.7 to 1.6%, P: 0.02% or less (including 0%), S: 0.01% or less (including 0%), Al: 0.06% Or less (not including 0%), Cu: 0.7-2%, Ni: 0.3-2%, N: 0.008% or less (not including 0%), and −0.1 ≦ ([ Mn] − [Si + P]) / [Cu] ≦ 1.4 (1)
{In Formula (1),
[Mn] = Mn content (% by mass) /54.9
[Si + P] = [Si content (mass%) / 28.0] + [P content (mass%) / 31.0]
[Cu] = Cu content (% by mass) /63.5
Show}
In which the metal structure is a single phase of ferrite or a ferrite and pearlite and / or bainite having an area ratio of 90% or more, and a temperature range lower than at least 580 ° C. after hot rolling is 0.2 ° C. / It is characterized by being obtained by cooling at a cooling rate of s or less.

  It is preferable that the ferrite crystal grains have an equivalent circle average grain size of 7 μm or less because ductility and low temperature toughness can be further improved.

  The steel plate of the present invention may further contain Cr: 0.5% or less (not including 0%) and / or Mo: 0.5% or less (not including 0%) as other elements. Moreover, in order to improve hardenability etc., Nb: 0.03% or less (not including 0%), Ti: 0.05% or less (not including 0%), V: 0.05% or less (not including 0%) ), Or B: 0.003% or less (0% not included).

  Furthermore, Ca: 0.0005 to 0.005%, Zr: 0.0002 to 0.005%, or Mg: 0.0005 to 0.005% may be included.

  The present invention also defines a method for producing such a non-tempered Cu precipitation type high-tensile steel sheet, and a temperature range below at least 580 ° C. after hot rolling is set at a cooling rate of 0.2 ° C./s or less. It is characterized by cooling at

  In order to make the ferrite crystal grains finer, in the hot rolling, after heating to a temperature of Ac3 point or higher, cooling at a cooling rate of 0.2 ° C / s or higher in an austenite recrystallization temperature range of 900 ° C or higher. However, it is preferable to perform rolling with a cumulative rolling reduction of 40% or more.

  According to the present invention, Cu precipitation can be almost completed in the cooling process after hot rolling without providing a heat treatment step separately, and the base material structure is mainly composed of a ferrite structure, which is as high as 570 MPa or more. A steel sheet having high strength and excellent ductility can be obtained. Therefore, it efficiently manufactures steel plates used as building or bridge pillars with excellent fracture resistance such as earthquake resistance, and high-strength steel plates with excellent workability such as press working used in automobile bodies. be able to.

  The present inventors have focused on Cu precipitation-type high-tensile steel sheets as steel sheets that can simultaneously increase strength and ductility, and as described above, Cu precipitation-type high tensions can be obtained without performing heat treatment such as aging treatment for Cu precipitation. In order to obtain a steel sheet, the inventors conducted intensive research on the behavior of Cu in steel up to Cu precipitation and coexisting elements that affect the behavior.

  First, various investigations were conducted on the case where an element (coexisting element) considered to be unrelated to Cu precipitation was present in pure iron together with Cu, and the influence of the coexisting element on Cu precipitation was investigated. As a result, it was understood that the time until Cu deposition shifted to the short time side as the atomic ratio (ratio of the number of atoms) of the alloy elements excluding Si and P and C increased.

  One of the experimental results confirming this is shown in FIG. FIG. 1 is obtained by treating the three steel types shown in Table 1 as shown in the heat treatment history chart of FIG. 2 under the conditions shown in (i) to (iii) below, and measuring the Vickers hardness of the obtained steel. It is a thing.

  (I) A steel plate subjected to homogenization at 1100 ° C. is rolled to a thickness of 7 mm under heating and rolling conditions of heating temperature: 1050 ° C. and rolling finishing temperature: 950 ° C.

  (Ii) The sample was subjected to a solution treatment of Cu at 850 ° C. and then water-cooled.

  (Iii) Aging treatment at 500 to 600 ° C. with various holding times.

  From FIG. 1, it can be seen that Cu precipitation is promoted by coexisting Mn rather than Cu alone in pure iron, and Cu precipitation is further promoted by coexistence of Mn and Si.

  Such a phenomenon can be considered as follows. That is, when Cu atoms are present alone, since the Cu atoms are larger in size than the iron atoms, lattice strain is generated so as to form a compressive stress field around the Cu atoms (in addition, from the iron atoms). There are many elements with large atomic sizes, but Cu is the closest to iron atoms among them).

  Here, when a Si atom having a size smaller than that of an iron atom is adjacent to a Cu atom, since the Si atom is an atom smaller than an iron atom, a lattice strain is formed so as to form a tensile stress field around the Si atom. . Therefore, it is considered that the influence of the lattice distortion is canceled and the influence on the surroundings (the influence caused by the occurrence of the lattice distortion) is reduced (hereinafter, such a state is referred to as “pairing”).

  When the pairing occurs, Cu diffusion is suppressed and Cu is unlikely to precipitate. However, as described above, when only Cu is present in pure iron, the pairing is not formed, so that precipitation is promoted. Conceivable. However, it cannot be said from FIG. 1 that Cu precipitation is promoted. This is because, when an element larger than an iron atom such as Cu is dissolved, the stress field formed around the element is a compressive stress field and repels each other, and Cu precipitation is suppressed. it is conceivable that. In addition, when only Cu is present alone, there is repulsion between Cu, but the permissible region of Cu is wide compared to the case where other elements are present, and as a result, the diffusion distance is increased and precipitation is difficult. It is also considered as a cause.

  On the other hand, in the case where only Mn, which is an element larger than an iron atom, coexists with Cu as in the Cu-Mn steel type shown in Table 1, precipitation is observed from FIG. You can see that it is promoted. This is presumably because precipitation is promoted because the allowable region of Cu is narrowed by the presence of Mn atoms and the diffusion distance of Cu is shortened.

  However, as shown in FIG. 1, Cu precipitation is further accelerated by simultaneously adding Si and Mn. This is because even when Si and Mn form a pairing, the stress field is not completely cancelled, and a stress field specific to each atom remains on the opposite side (outside) of the pairing formation. It is presumed that the compressive stress field remaining around Cu is attracted to the tensile stress field remaining around, thereby forming a state of promoting the diffusion of Cu atoms.

  In other words, the state where atoms smaller in size than iron atoms exist in a free state is not preferable for Cu precipitation, but there exists a pairing of atoms larger than iron atoms other than Cu, such as Mn, and atoms smaller than iron atoms. This is very effective in the present invention.

  From this point of view, the present inventors divided into a group of elements such as Si and P having a smaller atomic size than the iron atoms constituting the base material and a group of elements having a larger atomic size than the iron atoms. Study was carried out.

  Examples of elements having a size smaller than that of iron atoms include C and S. However, these elements do not contribute to the above-described action because they are interstitial and solid solution elements in iron. Therefore, in the present invention, Si and P are controlled as substitutional solid solution elements smaller in size than iron atoms.

  On the other hand, Cr and Ni are examples of elements that are larger in atomic size than iron atoms and are dissolved in substitutional form in iron. However, these elements have almost the same atomic size as iron atoms. Almost no pairing with Si or P is formed. Therefore, it hardly shows solid solution strengthening, and even when added, it does not contribute to improving the strength, so that it is not necessary to consider pairing with Si or P. In addition, Al and Ti are elements that form nitrides, and Mo, Nb, V, and the like are elements that form stable carbides. These elements have already formed precipitates in the Cu precipitation temperature range and are solid. Since it is not in a molten state, it is considered that it does not contribute to the above action. Therefore, as a substitutional solid solution element having a size larger than that of iron atoms, Mn that does not form a precipitate in the Cu precipitation temperature range was set as a control target.

  By utilizing the effect of Si, Mn, etc. on Cu as described above, Cu particles are precipitated by the self-holding heat of the steel plate during cooling after rolling or during winding processing without performing aging treatment separately. I found that it was almost complete.

The present inventors further studied the appropriate ratio of Si, Mn, etc. in order to efficiently express the above-described effects, and as a parameter,
([Mn]-[Si + P]) / [Cu]
{[Mn] = [Mn content (mass%)] / 54.9,
[Si + P] = [Si content (mass%)] / 28.0+ [P content (mass%)] / 31.0,
[Cu] = [indicates Cu content (% by mass)] / 63.5}
I found out that I should control.

  FIG. 3 shows the relationship between the above ([Mn] − [Si + P]) / [Cu] (hereinafter sometimes referred to as “parameter value”) and tensile strength, and FIG. 4 shows the parameter value and elongation. It is the graph which showed the relationship. These are results obtained by using the steel type A of the examples described later and under the conditions shown in FIGS.

  When the value of [[Mn]-[Si + P]) / [Cu] is extremely smaller than 0 (that is, when there are a lot of elements smaller in size than iron atoms), Cu is paired with a smaller element. Since it is easy to form a ring and Cu diffusion is suppressed and Cu precipitation hardly occurs, the strength cannot be increased efficiently. From FIG. 3, in order to obtain a tensile strength of 570 MPa or more, the parameter value should be −0.1 or more, preferably 0 or more, and 0.5 or more in order to obtain a higher strength. I understand.

  On the other hand, when the value of ([Mn] − [Si + P]) / [Cu] is large (that is, when many Mn atoms larger than Cu are present), small elements (Si, P) and Cu are Since the lattice distortion is more relaxed when forming a pairing with Mn having a larger size than when forming a pairing, the pairing with Mn is formed preferentially. As a result, Cu becomes single and easily diffuses. In addition, Cu is attracted to the compressive stress field remaining around the Si and P forming the pairing and easily diffuses, so it is considered that the time until the completion of the precipitation shifts to the short time side.

  This is clear from FIG. 1 because the time until the Cu deposition is completed is shorter as the value of ([Mn] − [Si + P]) / [Cu] is larger.

  However, if the value of ([Mn]-[Si + P]) / [Cu] becomes too large, the elongation decreases as shown in FIG. This is thought to be due to the fact that the solid solution strengthening remarkably progresses due to the excess of Mn present alone, resulting in a decrease in elongation. From FIG. 4, it can be seen that the above parameters should be made 1.4 or less in order to achieve an elongation of 27.5% or more. In order to further increase the elongation and achieve excellent ductility, the parameter may be set to 1.0 or less, more preferably 0.8 or less, particularly 0.5 or less.

  From these results, the present inventors control the components of the steel material so that the above parameters satisfy the range of the following formula (1) in order to obtain a highly ductile Cu precipitation-type high-tensile steel sheet by non-tempering. I found that it was good.

−0.1 ≦ ([Mn] − [Si + P]) / [Cu] ≦ 1.4 (1)
{In Formula (1),
[Mn] = [Mn content (% by mass)] / 54.9
[Si + P] = [Si content (mass%)] / 28.0+ [P content (mass%)] / 31.0
[Cu] = [indicates Cu content (% by mass)] / 63.5}
In addition, [] in the said Formula (1) calculated | required each atomic number obtained by dividing the mass% of an additional element by each atomic weight, and has specifically shown the atomic ratio of the additional element.

  In order to effectively exhibit such effects and to have other characteristics, it is very effective to satisfy the following component composition and to manufacture a steel material having the component composition by the following method. It is.

<About component composition>
C: 0.06% or less (excluding 0%)
Although it is an element necessary for ensuring the strength of the steel material, if it is excessively contained, work hardening due to carbide (including cementite) occurs, and the amount of formation of the second phase structure increases, leading to a decrease in ductility. Therefore, it is necessary to suppress it to 0.06% or less. From the viewpoint of securing a good strength-ductility balance, it is preferably contained within a range of 0.02% to 0.05%.

Si: 0.1-0.6%
Si is an element that contributes to an increase in the strength of the base metal, and also has a role as a deoxidizer for molten steel. In order to exhibit these effects, it is good to contain 0.1% or more. However, as described above, Si has a function of forming a pairing with Cu and suppressing the precipitation of Cu. Therefore, if excessively contained, Cu precipitation is inhibited. Moreover, since it also causes deterioration of weldability and base material toughness, the upper limit of Si content is set to 0.6%. Preferably it is 0.4% or less, More preferably, it is 0.25% or less.

Mn: 0.7 to 1.6%
Mn is an element utilized next to C because it has a role of increasing the strength of the base material and is inexpensive. Further, as described above, since the Mn atom is a substitutional solid solution element having a size larger than that of the iron atom, it has a function of forming a pairing with the Si atom or the P atom in the steel and promoting the precipitation of Cu. Therefore, it should be contained so that the number of atoms is 0.7% or more (preferably 1.0% or more) and equal to or more than Si and P in the steel. That is, it is contained so that ([Mn] − [Si + P]) ≧ 0 in the above formula (1).

  However, if excessively contained, Cu precipitation is promoted, but ductility decreases due to segregation of Mn and increase in solid solution strengthening. Moreover, since weldability and base material toughness are also deteriorated, it is limited to 1.6% or less (preferably 1.4% or less).

P: 0.02% or less (including 0%)
P (phosphorus) is an element contained as an inevitable impurity in steel, and is not an element added intentionally. If it is contained in a large amount in steel, weldability deteriorates, so it is necessary to reduce it as much as possible. Further, similarly to Si, P atoms are substitution-type solid solution strengthening elements smaller than iron atoms, and therefore when Cu is contained in a large amount, precipitation of Cu is inhibited. Therefore, P must be suppressed to 0.02% or less, preferably 0.015% or less, and more preferably 0.01% or less.

S: 0.01% or less (including 0%)
S is an element contained as an impurity like P, and is not an element added intentionally. When contained in a large amount in steel, inclusions are formed to deteriorate the base metal toughness and weldability. Therefore, S must be suppressed to 0.01% or less, preferably 0.006% or less, more preferably 0.003% or less.

Al: 0.06% or less (excluding 0%)
Al has an effect as a deoxidizing material and an effect of forming a nitride to refine the base material structure. In order to exert such an effect, the content is preferably 0.02% or more, and more preferably 0.03% or more. However, if excessively contained, the base material toughness deteriorates, so the content is suppressed to 0.06% or less. Preferably it is 0.05% or less.

Cu: 0.7-2%
Cu is an essential element for obtaining the precipitation-strengthened steel sheet of the present invention. Since the solid solution amount in the aging treatment temperature range of 500 to 600 ° C. is about 0.7%, precipitation strengthening (curing) cannot be expected unless it is contained in an amount of 0.7% or more. Preferably it is 0.9% or more, more preferably 1.1% or more.

  On the other hand, if excessively contained, the weldability deteriorates and hot brittleness (rolling crack) is caused during rolling. Therefore, the Cu content is 2% or less, preferably 1.5% or less, more preferably 1.3%. % Or less.

Ni: 0.3-2%
Ni is an element effective for improving hardenability and low temperature toughness. In order to exhibit these effects, it is preferable to contain 0.3% or more. In particular, the inclusion of 0.5 times or more of the Cu content (mass%) also has an effect of suppressing hot brittleness due to Cu. From this point of view, [Ni content (mass%)] / [Cu content (mass%)] should be contained so that ≧ 0.5, more preferably [Ni content (mass%)]. / [Cu content (mass%)] ≧ 1.0. However, since Ni is expensive, it is contained in a range of 2% or less in order to suppress an increase in cost.

N: 0.008% or less (excluding 0%)
Since nitrogen in the air is taken in during the molten steel treatment, it is an element inevitably mixed in the steel. N forms nitrides with Al, Ti, Nb, V, etc., and promotes refinement of the base material structure. In order to exhibit such an effect, the content is preferably 0.003% or more, more preferably 0.004% or more. However, if it is contained excessively, the solute N increases, and particularly toughness deterioration of the welded portion is caused. Preferably it is 0.006% or less, More preferably, it is 0.005% or less.

  The preferred contained elements of the present invention are as described above, and the remaining component is substantially Fe. However, it is a matter of course that a small amount of inevitable impurities are allowed to be contained in the steel sheet, and the action of the present invention. It is also possible to actively contain other elements as described below within a range that does not adversely affect the elements.

Cr: 0.5% or less, and / or Mo: 0.5% or less Both Cr and Mo are elements that precipitate carbonitride and contribute to an increase in strength. Accordingly, Cr may be contained in an amount of 0.1% or more and Mo may be contained in an amount of 0.05% or more. However, if these elements are contained excessively, weldability and base metal toughness are deteriorated. Therefore, it is preferable to suppress each to 0.5% or less. Preferably each is 0.4% or less, more preferably 0.3% or less.

One or more selected from the group consisting of Nb: 0.03% or less, Ti: 0.05% or less, V: 0.05% or less, and B: 0.003% or less Nb increases strength by improving hardenability It is an element that has an effect, and is effective in reducing the coarsening and recrystallization of austenite grains during rolling, refining the ferrite grains after rolling, and increasing the strength through the formation of carbonitrides. It is. In order to exert such an effect, the content is preferably 0.002% or more, more preferably 0.005% or more, and still more preferably 0.01% or more.

  However, when Nb is contained excessively, weldability deteriorates and NbC (carbide) precipitates excessively to deteriorate ductility and base metal toughness. Therefore, the Nb content is preferably 0.03% or less, more preferably 0.025% or less, and still more preferably 0.02% or less.

  Ti, like Nb, has an effect of increasing strength by improving hardenability, and is an element effective for increasing the strength of a steel sheet by refining ferrite crystal grains by forming nitrides. In order to exert such an effect, the content is preferably 0.004% or more, more preferably 0.007% or more, and still more preferably 0.012% or more.

  On the other hand, when the Ti content is excessive, TiC (carbide) is excessively precipitated and the ductility and the base metal toughness are lowered. More preferably, it is 0.03% or less, More preferably, it is 0.02% or less.

  V, like Nb and Ti, is an element that is effective in increasing the strength by improving the hardenability and finer ferrite grains through the formation of nitrides and increasing the strength. In order to acquire this effect, it is good to make it contain 0.002% or more, More preferably, it is 0.01% or more, More preferably, it is 0.02% or more. On the other hand, if V is contained excessively, weldability is lowered, so it is preferable to keep it at 0.05% or less. More preferably, it is 0.04% or less, More preferably, it is 0.03% or less.

  B is a very effective element for increasing the hardenability and increasing the strength by containing a small amount. In order to obtain such an effect, the content is preferably 0.0002% or more, more preferably 0.001% or more, and still more preferably 0.0015% or more. However, if it is contained excessively, the low temperature toughness of the base material is lowered, so it is suppressed to 0.003% or less. More preferably, it is 0.0025% or less, More preferably, it is 0.002% or less.

One or more selected from the group consisting of Ca: 0.0005-0.005%, Zr: 0.0002-0.005%, and Mg: 0.0005-0.005% Ca is an interposition in steel In order to have the effect of improving the toughness of the base material by spheroidizing the material, and exhibiting this effect, 0.0005% or more is preferable. More preferably, it is 0.001% or more, More preferably, it is 0.0015% or more. On the other hand, if the Ca content is excessive, the toughness of the base material deteriorates, so the upper limit is preferably made 0.005%. More preferably, it is 0.004% or less, More preferably, it is 0.003% or less.

  Zr has the effect of improving the toughness of the base metal by spheroidizing the inclusions in the steel, like Ca. In order to exert such an effect, the content is preferably 0.0002% or more, more preferably 0.001% or more, and still more preferably 0.002% or more. However, if Zr is excessively contained, the toughness of the base material is deteriorated on the contrary, so it is suppressed to 0.005% or less. More preferably, it is 0.004% or less, More preferably, it is 0.003% or less.

  Mg combines with oxygen in the steel to form an oxide. The oxide is very stable even in a high temperature state and suppresses the coarsening of crystal grains in the weld heat affected zone. In order to exhibit this effect, Mg may be contained in an amount of 0.0005% or more. More preferably, it is 0.001% or more, More preferably, it is 0.002% or more. However, when the Mg content is excessive, the oxides increase and the base material toughness deteriorates. Therefore, it is good to restrain to 0.005% or less, More preferably, it is 0.004% or less, More preferably, it is 0.003% or less.

<About manufacturing method>
In order to effectively exhibit the above-described effects, the present inventors have conducted intensive research. FIG. 5 is a graph showing the relationship between the cooling rate and Vickers hardness for each steel type when the cooling rate after rolling is changed in the step shown in FIG. 6 using the three steel types shown in Table 1. is there.

  From FIG. 5, the decrease in Vickers hardness due to the increase in the cooling rate after rolling is more gradual in Cu-Mn-Si steel than in Cu steel and Cu-Mn steel, but also for the Cu-Mn-Si steel. It can be seen that the hardness decreases when the cooling rate is increased. This is considered to be because when the cooling rate is increased, the residence time in the temperature range of about 500 to 600 ° C. optimal for Cu precipitation is shortened, and Cu is not sufficiently precipitated.

  To sufficiently precipitate Cu in the cooling process after rolling and achieve a high hardness of 180 or more in Vickers hardness, a steel containing Si and Mn together with Cu is added at 0.2 ° C./s after hot rolling. It turns out that it is good to cool at the speed of the following (preferably 0.1 degrees C / s or less).

  In order to realize cooling at such a cooling rate, cooling is performed after rolling, or when a thin steel plate is manufactured, winding is performed after rolling.

  In FIG. 5 above, the cooling rate from 850 ° C. after rolling was examined. Cu precipitation occurs most effectively at about 500 to 600 ° C. Therefore, from the rolling to the Cu precipitation temperature range, accelerated cooling with a higher cooling rate may be performed, and the present inventors have set the lower limit temperature of the temperature range in which accelerated cooling is possible (accelerated cooling cooling). The stop temperature was investigated.

  FIG. 7 is a graph showing a relationship between the cooling stop temperature of water cooling (accelerated cooling) after rolling (cooling to room temperature after stopping accelerated cooling) and tensile strength or elongation. This is performed under the conditions shown in FIG. 7 using a steel type A of an example described later. From FIG. 7, it can be seen that in order to achieve excellent elongation of 27.5% or more and high strength of 570 MPa or more at the same time, the accelerated cooling needs to be at least up to 580 ° C.

  In other words, Cu can be sufficiently precipitated by cooling at least a temperature range below 580 ° C. at a cooling rate of 0.2 ° C./s or less. Preferably, the temperature range of 500 to 580 ° C. is cooled at a cooling rate of 0.2 ° C./s or less. Specifically, a method of cooling to room temperature after rolling, a method of cooling to room temperature after cooling to 580 ° C or higher, a method of cooling to room temperature after rolling, a method of winding after rolling, and water cooling to a temperature of 580 ° C or higher after rolling After that, a method of winding is exemplified.

  Further, as will be described later, in order to refine the ferrite crystal grains in order to ensure better ductility and low temperature toughness, hot rolling is preferably performed by the following method.

  That is, at the time of hot rolling, it is first heated to a temperature of Ac3 point or higher. This is because the structure can be brought into an austenite single-phase state by heating to this temperature, and the structure defined in the subsequent steps can be formed. Preferably, heating is performed in a temperature range of (Ac3 point + 100 ° C.) or more and (Ac3 point + 250 ° C.) or less.

  Next, it is preferable to carry out rolling with a cumulative rolling reduction of 40% or more while cooling at a cooling rate of 0.2 ° C./s or more in an austenite recrystallization temperature range of 900 ° C. or more.

  The present invention is based on the premise that the metal structure is mainly composed of ferrite. Specifically, the ferrite single phase or ferrite with an area ratio of 90% or more (preferably 93% or more) and the remaining first The case where it consists of pearlite and / or bainite as a two-phase structure is mentioned. Thus, it is possible to ensure excellent ductility by using ferrite as a main component. To further improve ductility and secure low temperature toughness, it is effective to refine the ferrite crystal grains.

  FIG. 5 is a graph showing the relationship between the ferrite grain size and elongation or low-temperature toughness, which was produced using the steel type A in the examples described later under the conditions shown in FIG. Steels with different ferrite grain sizes are manufactured by changing the cooling rate and cumulative rolling reduction during rolling in the recrystallization temperature range.

  From FIG. 8, it can be seen that in order to obtain a steel sheet having an elongation of 27.5% or more and vTrs of −45 ° C. or less and excellent ductility and low-temperature toughness, it is effective to make the ferrite grain size 7 μm or less. .

  FIG. 9 shows the relationship between the cumulative rolling reduction at 900 ° C. or more and the ferrite grain size according to the cooling rate at the time of rolling in the austenite recrystallization temperature range of 900 ° C. or more. This was manufactured under the conditions shown in FIG. From FIG. 9, in order to obtain a ferrite structure with a ferrite grain size of 7 μm or less, the cooling rate during rolling in the austenite recrystallization temperature range of 900 ° C. or higher is set to 0.2 ° C./s or higher, and 900 ° C. or higher. It can be seen that it is effective to set the cumulative rolling reduction of 40% or more. Preferably, rolling is performed so that the cumulative rolling reduction is 50% or more.

  However, if the cooling rate is too high, the residence time in the austenite recrystallization region is shortened and sufficient recrystallization region rolling cannot be performed, so 0.5 ° C./s or less is preferable. A larger cumulative rolling reduction is desirable because the structure becomes finer.

  The present invention is not limited to other production conditions in the production process, and it is only necessary to adopt conditions that are usually performed for melting and casting of steel materials.

  The steel plate of the present invention exhibits the above-described effects without depending on the plate thickness, and as described above, a steel plate having a plate thickness of about 40 to 80 mm used as a construction material or a plate used for automobiles. It can be applied to any thin steel plate having a thickness of about 10 mm or less.

  EXAMPLES Hereinafter, the present invention will be described more specifically with reference to examples. However, the present invention is not limited by the following 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.

  After the steel materials shown in Table 2 were melted, hot rolling was performed under the conditions shown in Table 3 or Table 4 (heating conditions, rolling conditions, and cooling conditions) to obtain steel plates having the thicknesses shown in Table 3 or Table 4. .

  The metal structure and mechanical properties of the steel sheet thus obtained were examined. After the metal structure was corroded with 3% nital solution, the area ratio of ferrite and the average crystal grain size (equivalent circle equivalent diameter) of ferrite were examined using an image analyzer.

  For mechanical properties, tensile tests were conducted to determine the values of strength and elongation at break. When the plate thickness was less than 20 mm, a JIS Z 2201 No. 5 test piece was sampled, and when the plate thickness was 20 mm or more, a JIS Z 2201 No. 4 test piece was sampled. Further, low temperature toughness (vTrs) was examined by collecting JIS Z 2201 No. 4 test pieces from a ¼ portion of the thickness of a steel plate having a thickness of 10 mm or more. These results are also shown in Table 3 or Table 4.

  The following can be considered from Tables 2-4. No. shown in Table 3 1-18 satisfy | fills the requirements prescribed | regulated by this invention, and it turns out that intensity | strength and elongation are high and it is excellent also in low-temperature toughness. In contrast to this, No. Nos. 19 to 36 were inferior in mechanical properties or inferior in low-temperature toughness because the component composition specified in the present invention was out of the range or not manufactured by the specified method.

  That is, no. No. 21 was inferior in strength because it was subjected to accelerated cooling to a temperature range below the regulation after hot rolling. No. Nos. 28 and 30 were wound after hot rolling, but after hot rolling, accelerated cooling was performed to a temperature range lower than specified. Similar to 21, the strength was inferior.

  No. In No. 23, accelerated cooling was performed to a temperature range significantly lower than the specified range, so that the metal structure was not mainly composed of ferrite, resulting in inferior elongation. No. 27, the cooling rate during rolling at 900 ° C. or higher was too slow, and further, accelerated cooling after rolling was performed to a specified temperature or lower, so that the metal structure was mainly composed of bainite and martensite but not to ferrite. The result was inferior.

  In order to ensure the excellent low temperature toughness, the ferrite crystal grain size should be within the range specified by the present invention. From 19, 22, 25, 26 and 31, it can be seen that cooling during rolling at 900 ° C. or higher is preferably performed at a prescribed cooling rate. No. From 20, 24, and 29, it can be seen that rolling is preferably performed at 900 ° C. or higher so that the cumulative rolling reduction during rolling is within a specified range.

  No. 32 to 36 deviate from the component composition defined in the present invention. No. Since No. 32 had too little Cu, sufficient precipitation hardening could not be achieved, resulting in poor mechanical properties.

  No. No. 33 is an example in which the amount of C is too large, and since the fraction of the second phase other than ferrite occupying the metal structure is large, the result is high strength but slightly inferior in elongation.

  No. Since 34 was outside the lower limit of the formula (1) defined in the present invention, the strength was insufficient. No. No. 35 could not secure sufficient strength because the Si content was too small. No. No. 36, which was wound after hot rolling, was outside the lower limit of the formula (1) defined in the present invention, so that the effect of the present invention was not fully exhibited and the strength was insufficient.

It is the graph which showed the relationship between an aging treatment and Vickers hardness according to steel classification. FIG. 2 is a history diagram of heat treatment performed in the experiment of FIG. 1. It is the graph which showed the relationship between a parameter value and tensile strength. It is the graph which showed the relationship between a parameter value and elongation. It is the graph which showed the relationship between the cooling rate after rolling, and Vickers hardness for each steel type. FIG. 6 is a history diagram of heat treatment performed in the experiment of FIG. 5. It is the graph which showed the relationship between the cooling stop temperature of water cooling after rolling, and tensile strength or elongation. 4 is a graph showing the relationship between ferrite grain size and elongation or low temperature toughness. It is the graph which showed the relationship between the cumulative reduction rate at the time of rolling at 900 degreeC or more, and the ferrite particle size according to the cooling rate at the time of the rolling.

Claims (7)

% By mass (the same applies below)
C: 0.06% or less (excluding 0%),
Si: 0.1 to 0.6%,
Mn: 0.7 to 1.6%,
P: 0.02% or less (including 0%),
S: 0.01% or less (including 0%),
Al: 0.06% or less (excluding 0%),
Cu: 0.7-2%,
Ni: 0.3-2%,
N: 0.008% or less (excluding 0%)
And −0.1 ≦ ([Mn] − [Si + P]) / [Cu] ≦ 1.4 (1)
{In Formula (1),
[Mn] = Mn content (% by mass) /54.9
[Si + P] = [Si content (mass%) / 28.0] + [P content (mass%) / 31.0]
[Cu] = Cu content (% by mass) /63.5
Show}
In which the metal structure is a ferrite single phase or a ferrite with a ratio of area of 90% or more and pearlite and / or bainite, and a temperature range below at least 580 ° C. after hot rolling is 0.2 ° C./s. A non-tempered Cu precipitation type high-tensile steel sheet having a tensile strength of 570 MPa or more, obtained by cooling at the following cooling rate.   2. The non-tempered Cu precipitation type high-tensile steel sheet according to claim 1, wherein the ferrite crystal grains have an equivalent circle average grain size of 7 μm or less. As other elements,
Cr: 0.5% or less (not including 0%) and / or Mo: 0.5% or less (not including 0%)
The non-tempered Cu precipitation type high-tensile steel sheet according to claim 1 or 2. As other elements,
Nb: 0.03% or less (excluding 0%),
Ti: 0.05% or less (excluding 0%),
V: 0.05% or less (not including 0%) or B: 0.003% or less (not including 0%)
The non-tempered Cu precipitation type high-tensile steel sheet according to any one of claims 1 to 3. As other elements,
Ca: 0.0005 to 0.005%,
Zr: 0.0002 to 0.005% or Mg: 0.0005 to 0.005%
The non-tempered Cu precipitation type high-tensile steel sheet according to any one of claims 1 to 4.   A method for producing the steel sheet according to any one of claims 1 to 5, wherein a temperature range lower than at least 580 ° C after hot rolling is cooled at a cooling rate of 0.2 ° C / s or less. A method for producing a non-tempered Cu precipitation-type high-tensile steel sheet.   In the hot rolling, after heating to a temperature of Ac3 point or higher, rolling at a cumulative reduction rate of 40% or higher while cooling at a cooling rate of 0.2 ° C / s or higher in an austenite recrystallization temperature range of 900 ° C or higher. The manufacturing method of the non-tempered Cu precipitation type high strength steel plate of Claim 6 to implement.

Images