JP2007154289A - METHOD FOR PRODUCING HIGH IMPACT RESISTANT STEEL PIPE EXCELLENT IN DELAYED FRACTURING CHARACTERISTIC OF 1,700 MPa OR MORE OF TENSILE STRENGTH - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing a high impact resistant steel pipe excellent in delayed fracturing characteristic of ≥1,700 MPa tensile strength. <P>SOLUTION: The method for producing the high impact resistant steel pipe excellent in the delayed fracturing characteristic of ≥1700MPa tensile strength TS, is performed as the followings, that is, this steel is composed, by mass, of 0.19-0.35% C, 0.1-0.3% Si, 1.0-1.6% Mn, ≤0.025% P, ≤0.02% S, 0.01-0.05% Al, 0.001-0.004% B, 0.01-0.1% V, 0.005-0.05% Ti, ≤0.01% N, 0.001-0.01% Mg, 0.0001-0.008% O contained as the essential components, and the balance Fe with inevitable impurities, and the cooling speed till 1,500-1,400°C at the casting time is always ≥5°C/min and an induction-heating is applied to 850-1,050°C after rolling and pipe-making, and the water-cooling quenching is performed from under the austenite structure, at ≥100°C/s cooling speed. <P>COPYRIGHT: (C)2007,JPO&INPIT

Description

  The present invention relates to a method for producing a high impact resistant steel pipe having excellent delayed fracture characteristics with a tensile strength of 1700 MPa or more.

  As shown in Patent Document 1, a method of manufacturing a high-strength, high-impact steel pipe is formed from a steel plate made of steel containing a large amount of C, and then guided to 900 ° C. or higher through a work coil for high frequency heating. It is heated and manufactured by water-cooling and quenching from the austenite state to obtain a high-strength steel pipe having a tensile strength TS mainly composed of a martensite structure of 1700 MPa or more.

  However, it is known that the delayed fracture characteristics deteriorate as the strength of the steel increases. In particular, the deterioration is remarkable when the tensile strength TS is 1500 MPa or more.

  As a technique for improving delayed fracture characteristics, there is a wide range of methods for reducing the amount of diffusible hydrogen in steel by performing heat treatment such as tempering after quenching or annealing to precipitate matched precipitates and trapping hydrogen in the precipitates. Are known. For example, as shown in Patent Document 2, an element that precipitates carbonitrides such as V, Cr, Mo, and W and Mg are added and kept at 200 ° C. to 700 ° C. after annealing to form a hydrogen trap site. There are techniques for depositing alloy carbides or nitrides to improve delayed fracture characteristics. At this time, Mg is added in order to deposit carbonitride which becomes hydrogen trap sites uniformly and finely. If the precipitates are quenched, they are not compatible with the base metal and cannot be hydrogen trap sites. However, when tempering, the precipitates become hydrogen trap sites that are consistent with the base material. Precipitate. Patent Document 2 applies this technique.

  Although it is considered that delayed fracture characteristics can be improved by applying this technique even to a high impact steel pipe, if heat treatment is performed again after quenching, the strength is reduced and the heat treatment cost is increased.

JP 2003-129170 A Japanese Patent Laid-Open No. 2003-166035

  As described above, the high-impact-resistant steel pipe mainly composed of a martensite structure exceeding 1700 MPa has poor delayed fracture characteristics, and there is a concern of delayed fracture cracking before collision when installed as a member. A technique for improving delayed fracture characteristics without heat treatment has not yet been established. In addition, when a conventional technique for improving delayed fracture characteristics such as tempering heat treatment is applied, the strength is lowered and the heat treatment cost is increased.

  The present invention has been made in view of this point, and an object thereof is to improve the delayed fracture characteristics of a high impact-resistant steel pipe having a tensile strength of 1700 MPa or more without lowering the strength and increasing the heat treatment cost.

  From the background as described above, the present inventors have studied the component elements including V and Mg, and specified the casting conditions and quenching conditions, so that delayed fracture characteristics can be obtained without lowering the strength and increasing the heat treatment cost. I came up with a way to improve it. That is, by increasing the casting speed, finely dispersing the Mg compound or composite compound, and precipitating TiN with the Mg compound as a nucleus during casting, the coarse TiN that becomes the starting point of delayed fracture cracks at the grain boundaries is finely divided. It is possible to disperse, to reduce the cracking susceptibility of the grain boundary, and to disperse VC in the grains using Mg compound as a nucleus, and to lower the quenching temperature, to refine the γ grain size, and to reduce the grain boundary area It has been found that the delayed fracture characteristics of a high-impact steel pipe can be improved sufficiently because the cracking susceptibility at the grain boundaries can be reduced by increasing the number. The gist of the present invention is as follows.

  [1]. By mass%, C: 0.19 to 0.35%, Si: 0.1 to 0.3%, Mn: 1.0 to 1.6%, P: 0.025% or less, S: 0.02 %: Al: 0.01-0.05%, B: 0.001% -0.004%, V: 0.01% -0.1%, Ti: 0.005-0.05%, N : 0.01% or less, Mg: 0.001% to 0.01%, O: 0.0001% to 0.008% as essential components, with the balance being Fe and inevitable impurities, The cooling rate from 1500 ° C. to 1400 ° C. is always 5 ° C./min or more, and after pipe forming after rolling, induction heating is performed to 850 ° C. to 1050 ° C., and the water cooling quenching is performed at a cooling rate of 100 ° C./s or more from the austenite state. Of high impact resistant steel pipe with excellent delayed fracture characteristics with a tensile strength TS of 1700 MPa or more. Method.

  [2]. Further, by mass%, Nb: 0.005 to 0.05%, Cu: 0.005 to 0.5%, Cr: 0.005 to 0.5%, Mo: 0.1 to 0.5%, Ni : 0.1 to 0.5%, Ca: 0.01% or less, rare earth element (REM): one or more selected from the group of 0.1% or less 1] The manufacturing method of the high impact-resistant steel pipe excellent in the delayed fracture characteristic of the tensile strength of 1700 MPa or more.

  [3]. The steel pipe satisfies martensite volume fraction ≧ 95%, average γ (former austenite) particle size ≦ 5 μm, and TiN precipitate average particle size ≦ 1.8 μm, according to [1] and [2] A method for producing a high impact steel pipe excellent in delayed fracture characteristics with a tensile strength of 1700 MPa or more.

  According to the present invention, the delayed fracture characteristics of a high impact steel pipe can be sufficiently improved without lowering the strength and increasing the heat treatment cost. When applied to a door impact beam or bumper beam, it is attached with some processing. However, there is no fear of delayed fracture. Also, it can be applied to a part to which some new stress is applied.

Hereinafter, the present invention will be described in detail.
First, the first invention and the second invention will be described. The reasons for limiting the chemical components of steel in the present invention will be described below. C is an essential component for strengthening martensite itself and improving the hardness, and at least 0.19% is necessary to obtain a tensile strength TS of 1700 MPa or more. However, if C is excessive, the martensite structure becomes brittle and causes cracking that breaks during quenching, so the content is made 0.35% or less.

  Si and Mn are components that promote martensitic transformation from austenite at the time of quenching, and are less than the respective ranges of Si: 0.1 to 0.3% and Mn: 1.0 to 1.6%. And hardenability falls and sufficient intensity is not obtained. Conversely, exceeding the above range is not preferable because it causes burning cracks and segregation.

  Ti is an important element that has the effect of refining crystal grains and suppressing grain growth by forming fine carbides, and further improving the hardenability by fixing N. In order to obtain such an effect, at least 0.005% is necessary. However, if added excessively, the carbide coarsens and causes toughness deterioration, and the TiN particle size also coarsens and delayed fracture characteristics deteriorate, so the content is made 0.05% or less. N not only lowers the hardenability but also coarsens the TiN grain size and causes the deterioration of delayed fracture characteristics, so the content is made 0.01% or less.

  B is a component that suppresses the precipitation of ferrite. However, when it is combined with N contained as a gas component in steel and becomes BN, its effect is lost, so the content is made 0.001% or more. However, if it exceeds 0.004%, segregated inclusions are formed. P and S become segregated inclusions and make the martensite structure brittle, so P: 0.025% or less and S: 0.02% or less are required. Al is a deoxidizing material. If it is less than 0.01%, the deoxidation effect is insufficient, and if it exceeds 0.05%, the oxide becomes an intercrystalline inclusion, which is not preferable.

  Mg is the most important component element among the additive elements of the present invention. Mg is present as Mg oxide and Mg sulfide after solidification, and other precipitates are finely precipitated with the Mg compound as a nucleus. In the present invention, coarse TiN that is the starting point of cracks at grain boundaries is finely dispersed to reduce cracking susceptibility at grain boundaries, and V carbides and nitrides are dispersed and precipitated to suppress the growth of γ grains during quenching. It is important to add at least 0.001% or more of Mg to obtain the effect. However, if Mg is excessive, the Mg compound is coarsened, and the precipitate that precipitates using the Mg compound is also coarsened, so the upper limit is made 0.01% or less.

  V is an important component element among the additive elements of the present invention. Addition of 0.01% or more is necessary to obtain the effect of dispersing and precipitating V carbide and V nitride in the grains with Mg compound as the nucleus and pinning and suppressing the growth of γ grains during quenching. is there. However, if added excessively, V carbides and nitrides become coarse and may become the starting point of delayed fracture cracking, so the upper limit is made 0.1% or less.

  O is an essential element for generating Mg oxide. The amount of oxygen finally remaining in the steel is less than 0.0001%, so the number of oxides is not sufficient, so 0.0001% is set as the lower limit. On the other hand, when it exceeds 0.008%, coarse oxides increase, resulting in a decrease in the base material and HAZ toughness. Therefore, the upper limit is made 0.008% or less.

  Nb is a precipitation strengthening component that improves the strength by generating precipitates in the martensite structure and preventing the passage of dislocations. Cu, Cr, Mo, and Ni are solid solution strengthening components that improve the strength by forming a solid solution in the martensite structure and preventing the passage of dislocations. Cr and Mo also act as precipitation strengthening components. These components contribute to an increase in strength, but increase the cost and excessive addition becomes segregation inclusions. Therefore, Nb: 0.005 to 0.05%, Cu: 0.005 to 0.5%, Cr: 0.005 to 0.5%, Mo: 0.1 to 0.5%, and Ni: 0.1 to 0.5% are preferable.

  Ca and rare earth elements (REM) are components that contribute to the control of the morphology of inclusions, but excessive addition leads to harmful segregation leading to the destruction of the martensite structure, so Ca: 0.01% or less, REM: 0 .1% or less is appropriate. These Nb, Cu, Cr, Mo, Ni, Ca, and rare earth elements (REM) are not essential components, but are optional components that are added as necessary. For example, Y, La, Ce, or Sm can be used as the rare earth element (REM).

  Next, the reason for limiting the cooling rate during casting will be described. The freezing point of steel containing 0.19 to 0.35% of C is about 1520 ° C., and after the solidification, the formation of Mg compound and precipitation of TiN or the like with the core occur around 1500 ° C. to 1400 ° C. Therefore, since the form of the precipitate changes at a cooling rate in the vicinity of 1500 ° C. to 1400 ° C., it is important to limit the cooling rate in this temperature range. In order to refine the precipitate sufficiently to improve the delayed fracture characteristics, it is necessary to always set the cooling rate from 1500 ° C. to 1400 ° C. during casting to 5 ° C./min or more. When the cooling rate is slower than this rate, the Mg compound grows coarsely, and as a result, precipitates such as TiN that precipitates as a nucleus also coarsen. Therefore, the cooling rate was limited to the above conditions.

  After the casting, the steel plate that has been hot-rolled and, in some cases, cold-rolled is produced in the same manner as a normal steel plate manufacturing method. In the production method of the present invention, when the final quenching is performed, the austenite region is heated. Therefore, the temperature condition during the hot rolling does not greatly affect the characteristics of the product. However, if the milling temperature after hot rolling becomes too low, low temperature transformation structures such as bainite are mixed, and the strength becomes high and it becomes difficult to form a pipe. Therefore, the milling temperature is preferably about 650 ° C. In addition, the pipe can be formed by any method such as UO, spiral, and forge welding. However, economically, it is most desirable to form a pipe by roll forming and to make a steel pipe by electro-welding.

Next, the reason for limiting the quenching conditions will be described. The A ₃ transformation point of steel containing 0.19 to 0.35% of C is about 800 ° C. to 850 ° C., and in order to obtain sufficient strength, the entire steel pipe is quenched from a temperature range above the A ₃ transformation point. It is important that the lower limit of the induction heating temperature is 850 ° C. On the other hand, if the induction heating temperature is too high, the austenite crystal grains become coarse, and not only a sufficient strength cannot be obtained after quenching, but if the crystal grains become coarse, the grain interfacial area decreases and the delayed fracture characteristics deteriorate. Therefore, the upper limit of the induction heating temperature is 1050 ° C. Further, when the cooling rate from the austenite region is low, it does not become a high-strength martensite structure but transforms into an intermediate structure such as bainite or remains as retained austenite with inferior delayed fracture characteristics without being transformed. Therefore, the cooling rate must be a sufficiently fast cooling rate that transforms the entire steel pipe into martensite, and the lower limit is set to 100 ° C./s.

  Next, the third invention will be described. First, the reason for limiting the martensite volume ratio will be described. The portion that did not become martensite by quenching becomes retained austenite. If this retained austenite is present, not only is the strength lowered, but the delayed fracture characteristics are also degraded. Therefore, it is desirable that the retained austenite is as small as possible. However, since it is difficult to completely eliminate the retained austenite when manufacturing with an actual machine, the martensite volume ratio is set to 95%. As a method for that, it is relatively easy to add this by adding an alloy element having high hardenability such as Si or Mn as described above, or by sufficiently increasing the cooling condition during quenching as described above. Can be achieved.

  Next, the reason for limiting the average γ particle size will be described. For the same strength, the smaller the average γ particle size, the smaller the martensite particle size after transformation, so the particle size area increases and the amount of allowable intrusion hydrogen until cracking occurs increases. Tend to get better. This trend is known but has not yet been quantitatively evaluated. The present inventors conducted a delayed fracture test on samples with the average γ grain size changed by changing the quenching temperature of the same steel type, and quantitatively evaluated the relationship between the average γ grain size and delayed fracture characteristics. did. The results are shown in FIG. 1, and it was found that the average γ grain size suddenly delayed from about 5 μm and the fracture characteristics changed. When the grain size is halved, the interfacial area of the grain is about doubled, so that the delayed fracture characteristics tend to be about doubled. That is, the average γ grain size and delayed fracture characteristics are in an inversely proportional relationship, and the delayed fracture characteristics suddenly change from around 5 μm, so the average γ (former austenite) grain size was set to 5 μm or less. Moreover, the strength increases as the average γ particle size decreases. Further, since the toughness is improved when the grain size is reduced, crack propagation after the occurrence of delayed fracture cracking is also suppressed.

  Next, the reason for limiting the average particle size of the TiN precipitate will be described. In general, delayed fracture breaks starting from the grain boundary, but if a precipitate such as coarse TiN is present at the grain boundary, the stress intensity factor increases and stress concentration occurs, and hydrogen intrusion into that part is also significant. As a result, delayed fracture characteristics deteriorate. As a result of observing the crack fracture surface after delayed fracture by SEM, the present inventors have found that many steels with TiN precipitate average particle size> 1.8 μm are the origin of cracking of TiN. In many cases, TiN precipitates having an average grain size ≦ 1.8 μm often cracked from grain boundaries, particularly where dislocations are concentrated. That is, when the average particle size of TiN precipitates is> 1.8 μm, TiN induces delayed fracture cracking. However, in order to have sufficient delayed fracture characteristics, the average particle size of TiN precipitates was specified to be 1.8 μm or less.

  By the production method of the present invention, a high impact-resistant steel pipe having a TS of 1700 MPa or more and sufficient delayed fracture characteristics can be obtained without heat treatment after quenching, and installed in automobile doors, bumpers, roofs, pillars, lockers, etc. There is no worry about breaking and cracking. Moreover, it becomes possible to install in the site | part to which some new stress is loaded. This technology can be applied not only to steel pipes but also to steel plates.

  Steel pipes having the minimum cooling rate during casting, induction heating temperature, and quenching cooling rate conditions shown in Table 2 were prepared from steels having the compositions shown in Table 1. Table 2 shows TS, YS, EL, martensite volume fraction, average γ particle diameter, and TiN average particle diameter obtained as a result of the tensile test and micro observation. The results of the delayed fracture test are also listed in Table 2. The delayed fracture test was performed by cutting out a part of a steel pipe and charging it with hydrogen while applying a constant strain in a sulfuric acid solution, and confirming whether or not cracking occurred in 120 minutes. As for the strain, 0.4% strain was applied, no crack was generated, and ○ was generated, and x was generated. In Tables 1 and 2, parts that depart from the scope of the present invention are underlined.

  As described above, the high impact steel pipe of the present invention is characterized by having a TS of 1700 MPa or more and exhibiting sufficient delayed fracture characteristics. Martensite volume fraction is a factor that affects both TS and delayed fracture characteristics. The average γ particle size is also a factor that affects both TS and delayed fracture characteristics. The TiN average particle size is a factor that affects delayed fracture characteristics.

  As shown in Table 2, all of the steel pipes of the present invention examples have a TS of 1700 MPa or more and sufficient delayed fracture characteristics. Compared to the inventive example, the steel pipe 9 of the comparative example is out of the scope of the present invention for both V and Mg. Since the addition of V is not sufficient, the average γ particle size is increased, and since the addition of Mg is not sufficient, the TiN average particle size is also increased. Therefore, it is considered that sufficient delayed fracture characteristics could not be obtained. In the steel pipe 10 of the comparative example, C deviates from the scope of the present invention. Although it has sufficient delayed fracture characteristics, since C is not sufficient, the strength of the martensite structure is low, and sufficient TS cannot be obtained. The steel pipe 11 of the comparative example has a minimum cooling rate at the time of casting deviating from the scope of the present invention. Therefore, the TiN average particle size has become larger than the range of the present invention. As a result, sufficient delayed fracture characteristics cannot be obtained. This is probably because the Mg compound grows coarsely because the cooling rate at the time of casting is slow, and TiN that precipitates using it as a nucleus also coarsens. The steel pipe 12 of the comparative example has a quenching cooling rate that deviates from the scope of the present invention. Since the cooling rate is slow, the martensite volume fraction becomes small. Therefore, sufficient TS cannot be obtained. Further, since there is residual austenite that cannot be martensitic transformed, sufficient delayed fracture characteristics cannot be obtained.

It is the figure which showed the relationship between an average (gamma) particle size, and the time required for crack generation in the delayed fracture test which charges hydrogen in a sulfuric acid solution while loading 0.4% strain.

Claims (3)

  By mass%, C: 0.19 to 0.35%, Si: 0.1 to 0.3%, Mn: 1.0 to 1.6%, P: 0.025% or less, S: 0.02 %: Al: 0.01-0.05%, B: 0.001% -0.004%, V: 0.01% -0.1%, Ti: 0.005-0.05%, N : 0.01% or less, Mg: 0.001% to 0.01%, O: 0.0001% to 0.008% as essential components, with the balance being Fe and inevitable impurities, The cooling rate from 1500 ° C. to 1400 ° C. is always 5 ° C./min or more, and after pipe forming after rolling, induction heating is performed to 850 ° C. to 1050 ° C., and the water cooling quenching is performed at a cooling rate of 100 ° C./s or more from the austenite state. Of high impact resistant steel pipe with excellent delayed fracture characteristics with a tensile strength TS of 1700 MPa or more. Method.   Further, by mass%, Nb: 0.005 to 0.05%, Cu: 0.005 to 0.5%, Cr: 0.005 to 0.5%, Mo: 0.1 to 0.5%, Ni : 0.1 to 0.5%, Ca: 0.01% or less, rare earth element (REM): 1 type or 2 or more types selected from the group of 0.1% or less A method for producing a high impact resistant steel pipe having excellent delayed fracture characteristics with a tensile strength TS of 1700 MPa or more.   3. The tensile strength according to claim 1, wherein the steel pipe satisfies a martensite volume fraction ≧ 95%, an average γ (former austenite) particle size ≦ 5 μm, and a TiN precipitate average particle size ≦ 1.8 μm. A method for producing a high impact-resistant steel pipe having an excellent delayed fracture property with a TS of 1700 MPa or more.

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