WO2018061101A1 - Steel - Google Patents

Abstract

For the steel according to an embodiment of the present invention, the chemical components contain, in mass% units, C: 0.15%–0.40%, Mn: 0.10%–1.50%, S: 0.002–0.020%, Ti: 0.005%–0.050%, B: 0.0005–0.0050%, Bi: 0.0010%–0.0100%, P: 0.020% or less, N: 0.0100% or less, Si: 0% to less than 0.30%, Cr: 0–1.50%, Al: 0–0.050%, Mo: 0–0.20%, Cu: 0–0.20%, Ni: 0–0.20% and Nb: 0–0.030%, the balance being made of Fe and impurities.

Description

steel

The present invention relates to steel.

Cold forging (including rolling) can improve the surface texture and dimensional accuracy of the product compared to hot forging, and also has a good yield, so it is a relatively small machine such as a bolt. It is widely applied as a manufacturing method for parts. When manufacturing machine parts by cold forging, carbon steel and alloys for medium-carbon mechanical structures specified in JIS G 4051, JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc. as materials. Steel is often used as a final product through manufacturing processes such as hot wire rolling, annealing (or spheroidizing annealing), wire drawing, cold forging, quenching, and tempering. The general manufacturing process is characterized by adding an annealing or spheroidizing annealing process before cold forging. The reason why annealing or spheroidizing annealing is added before cold forging is the reason why medium carbon carbon steel and alloy steel remain hot rolled (that is, when air cooled without heat treatment after hot rolling) ), The hardness of the rolled material is high, and the die is worn at the time of cold forging, resulting in an increase in manufacturing cost, and cracking occurs at the time of cold forging because the ductility of the material is insufficient with hot rolling. This is because there is a manufacturing problem such as a reduction in yield due to facilitation.

However, since annealing is very costly, it has been demanded to develop a steel material that can omit the annealing process in order to reduce the manufacturing cost of parts. In view of such a demand, so-called bolt boron steel in which a small amount of B is added to steel has been developed (for example, Patent Document 1 and Patent Document 3). The feature of boron steel is that it reduces the carbon content of steel and the addition amount of alloy elements such as Cr and Mo, thereby reducing the hardness of the wire as it is hot rolled and improving the ductility. There is no need to compensate for the decrease in hardenability due to the reduction in the addition amount of the alloy element by the effect of improving hardenability by adding a small amount of B without increasing the hardness of the rolled material.

In order to express the effect of improving hardenability by adding a trace amount of B, it is necessary that B is in a solid solution state in austenite. On the other hand, when solid solution nitrogen is present in the steel, BN is formed, and the amount of solid solution B (solid solution B in the steel) decreases, thereby improving the hardenability of B. Will be lost. For this reason, in boron steel, it is a common practice to suppress the formation of BN by adding Ti having a strong affinity for N to previously fix N in the steel as TiN. For example, Patent Document 4 describes that the precipitation of BN is suppressed by setting Ti / N (mass% ratio) to 4 or more. In principle, if Ti / N is set to 3.42 or more, precipitation of BN can be suppressed.

However, the general boron steel as described above tends to generate so-called coarse grains in which some austenite crystal grains are coarsened by abnormal grain growth during quenching heating as compared with conventional steels. In a part in which coarse grains are generated, deterioration in dimensional accuracy due to an increase in heat treatment strain generated during quenching, and deterioration in part characteristics such as impact value, fatigue strength, and delayed fracture characteristics occur. Therefore, especially in the case of high strength bolts having a tensile strength of 800 MPa or more, preventing the generation of coarse grains is a major practical issue. In order to suppress the generation of coarse grains due to such abnormal grain growth, a large number of pinning particles (precipitates, etc.) are dispersed in the structure in order to pin the grain boundaries of austenite crystal grains. It is effective to disperse a large amount of particles.

There are two main reasons why coarse grains are likely to occur in boron steel.

(1) When boron steel is used as a component material, the annealing process after cold forging of boron steel is omitted, so that boron steel is directly heated from the cold-worked structure to the austenite region. In this case, since the austenite crystal grains are excessively refined and the crystal grain size is partially nonuniform due to the influence of cold working, some crystal grains are likely to cause abnormal grain growth.

(2) In the above boron steel, N in the steel is fixed as TiN by the addition of Ti, so that AlN that effectively acts as pinning particles in carbon steel and alloy steel as conventional steel is not generated. In addition, since TiN is coarser than AlN, it cannot be finely dispersed, and it is difficult to secure the number of pinning particles necessary for preventing coarse particles.

The above factor (1) is unavoidable in order to omit the annealing process. Therefore, how to secure the number of pinning particles in boron steel to improve the factor (2) It has been a point of prevention.

Under such circumstances, techniques for preventing the occurrence of coarse grains of boron steel have been proposed. For example, Patent Document 5 and Patent Document 6 describe using TiC and Ti (CN), which are finer precipitates than TiN, instead of AlN and TiN as pinning particles. In these techniques, in order to ensure the number of pinning particles necessary for preventing coarse grains, TiC and Ti (CN) having a diameter of 0.2 μm or less in steel before quenching and after hot rolling are used. Is dispersed in a total number of 20/100 μm ² or more. By dispersing a large amount of such fine precipitates before quenching and heating, these precipitates function as pinning particles that pin the austenite grain boundaries during quenching and heating. Since this technology makes it possible to stably prevent the generation of coarse grains in boron steel, steel to which this technology is applied is currently widely used as an inexpensive bolt steel material that can omit the annealing process.

However, the above techniques have drawbacks. That is, when fine TiC and Ti (CN) are dispersed in a large amount in the structure after hot rolling, there is a side effect that the hardness of ferrite increases due to precipitation strengthening by fine precipitate particles. The problem is that the softening effect of the hot-rolled material due to boron steel is reduced. That is, when the amount of fine TiC or Ti (CN) is increased, the generation of coarse grains can be suppressed, but the life of the cold forging die is reduced by increasing the hardness of the rolled material by precipitation strengthening. . Conversely, when the amount of fine TiC or Ti (CN) is suppressed, the hardness of the rolled material can be suppressed, but coarse grains are generated. That is, when fine TiC or Ti (CN) is used, the suppression of the generation of coarse grains and the suppression of the hardness of the rolled material before cold forging are in a trade-off relationship. Therefore, it is difficult to achieve both the softening of the rolled material and the suppression of stable coarse grains completely only by the above technique.

Patent Document 7 also describes a technical idea similar to the technique for preventing the generation of coarse grains of the boron steel. That is, it is a technique for preventing the coarsening of crystal grains by dispersing the carbonitrides of these elements in steel by making the relationship among the contents of Ti, Nb, Al, and N within a certain range. Patent Document 7 further describes the effect of improving the machinability by adding 0.01% or more of Bi. However, in Patent Document 7, only the effect of improving the machinability is disclosed as the effect of Bi. There is no description of the relationship between Bi and the coarsening characteristics of crystal grains. Since Bi is added for the purpose of improving the machinability, Patent Document 7 only considers adding a relatively large amount of Bi. In this case, as described in Patent Document 7, there is a concern about a decrease in hot workability due to Bi addition.

Patent Document 8 discloses a case-hardening steel that exhibits excellent grain coarsening characteristics even when carburized at a higher temperature than conventional examples, and exhibits excellent cold workability without soft annealing. Steel for case hardening intended to provide is disclosed. However, Patent Document 8 proposes only the use of fine Ti carbides, Ti-containing composite carbides, and the like as means for ensuring the crystal grain coarsening resistance. In Patent Document 8, the hot rolling temperature is extremely low in order to ensure cold workability, and thus the productivity of case hardening steel is impaired.

Japanese Laid-Open Patent Publication No. 5-339676 Japanese Patent Publication No. 5-63524 Japanese Unexamined Patent Publication No. 61-253347 Japanese Unexamined Patent Publication No. 3-47918 Japanese Patent No. 3443285 Japanese Patent No. 3490293 Japanese Unexamined Patent Publication No. 2000-328189 Japanese Unexamined Patent Publication No. 2006-265704

One of the challenges of steel for cold forging is to improve productivity without annealing after hot rolling and before cold forging in order to improve the cold forgeability of steel and the productivity of steel. It is to keep the steel soft without using damaging manufacturing conditions. Another problem with cold forging steel is to exhibit high hardenability after cold forging in order to impart high strength to machine parts. A further problem with steel for cold forging is the generation of coarse grains during quenching after cold forging in order to prevent deterioration of dimensional accuracy, impact value, fatigue strength, delayed fracture characteristics, etc. of machine parts. It is to suppress. As mentioned above, the prior art cannot solve all of these simultaneously. The use of TiC and Ti (CN) proposed in the prior art as a means for suppressing the generation of coarse grains hardens the steel after hot rolling and before cold forging by precipitation strengthening. Impair productivity.

The present invention has been made in view of the above problems. That is, the present invention suppresses the generation of coarse grains during quenching without using Ti carbides and Ti carbonitrides such as TiC and Ti (CN), thereby improving productivity, cold forgeability, and quenching. It is an object to provide a steel excellent in all mechanical properties.

The gist of the present invention is as follows.

(1) In the steel according to one embodiment of the present invention, the chemical composition is unit mass%, C: 0.15% to 0.40%, Mn: 0.10% to 1.50%, S: 0.00. 002 to 0.020%, Ti: 0.005% to 0.050%, B: 0.0005 to 0.0050%, Bi: 0.0010% to 0.0100%, P: 0.020% or less, N: 0.0100% or less, Si: 0% or more and less than 0.30%, Cr: 0 to 1.50%, Al: 0 to 0.050%, Mo: 0 to 0.20%, Cu: 0 to It contains 0.20%, Ni: 0 to 0.20%, and Nb: 0 to 0.030%, with the balance being Fe and impurities.
(2) In the steel according to (1), the chemical component is unit mass%, Si: 0.01% or more and less than 0.30%, Cr: 0.01 to 1.50%, and Al: One or more selected from the group consisting of 0.001 to 0.050% may be contained.
(3) In the steel according to (1) or (2), the chemical component is unit mass%, Mo: 0.02 to 0.20%, Cu: 0.02 to 0.20%, Ni One or two or more selected from the group consisting of: 0.02 to 0.20% and Nb: 0.002 to 0.030% may be contained.
(4) The steel according to any one of (1) to (3) may have an N fixed index I _FN defined by the following formula 1 of 0 or more.
I _FN = [Ti] −3.5 × [N] (Formula 1)
Here, [Ti] is the Ti content in unit mass%, and [N] is the N content in unit mass%.
(5) above (1) to (4) Steel according to any one of, Ti-Nb-based precipitates generated index _{I P,} which is defined by Equation 2 below may also be 0.0100 or less .
I _P = 0.3 × [Ti] + 0.15 × [Nb] − [N] (Formula 2)
Here, [Ti] is the Ti content in unit mass%, [Nb] is the Nb content in unit mass%, and [N] is the N content in unit mass%.

According to the present invention, it is possible to provide steel that can achieve both softening before cold forging and suppression of generation of coarse grains during quenching after cold forging. Moreover, the steel according to the present invention is excellent in manufacturability because it is not cracked during casting or rolling, and can be manufactured under conditions that do not place a load on the manufacturing equipment. By applying the steel according to the present invention to a cold forged part, wear of the mold during cold forging can be suppressed, and the life of the mold can be improved. In addition, by applying the steel according to the present invention to cold forged parts, the cost of expensive dies can be reduced, which can contribute to the reduction of the manufacturing cost of high-strength bolts having a tensile strength of 800 MPa or more. it can. Furthermore, the steel according to the present invention is excellent in machinability. Therefore, the present invention has a great industrial contribution.

The steel according to one embodiment of the present invention will be described. The steel according to this embodiment has the following characteristics.

(A) In the steel according to the present embodiment, the chemical composition is unit mass%, C: 0.15% to 0.40%, Mn: 0.10% to 1.50%, S: 0.002 to 0.020%, Ti: 0.005% to 0.050%, B: 0.0005 to 0.0050%, Bi: 0.0010% to 0.0100%, P: 0.020% or less, N: 0.0100% or less, Si: 0% or more and less than 0.30%, Cr: 0 to 1.50%, Al: 0 to 0.050%, Mo: 0 to 0.20%, Cu: 0 to 0.0. 20%, Ni: 0 to 0.20%, and Nb: 0 to 0.030%, with the balance being Fe and impurities.
(B) In the steel described in (a) above, the chemical component is unit mass%, Si: 0.01% or more and less than 0.30%, Cr: 0.01 to 1.50%, and Al: One or more selected from the group consisting of 0.001 to 0.050% may be contained.
(C) In the steel described in (a) or (b), the chemical component is unit mass%, Mo: 0.02 to 0.20%, Cu: 0.02 to 0.20%, Ni One or two or more selected from the group consisting of: 0.02 to 0.20% and Nb: 0.002 to 0.030% may be contained.
(D) The steel according to any one of the above (a) to (c) may have an N fixed index I _FN defined by the following formula 1 of 0 or more.
I _FN = [Ti] −3.5 × [N] (Formula 1)
Here, [Ti] is the Ti content in unit mass%, and [N] is the N content in unit mass%.
(E) above (a) ~ according to any one of (d) steel, Ti-Nb-based precipitates generated index _{I P,} which is defined by Equation 2 below may also be 0.0100 or less .
I _P = 0.3 × [Ti] + 0.15 × [Nb] − [N] (Formula 2)
Here, [Ti] is the Ti content in unit mass%, [Nb] is the Nb content in unit mass%, and [N] is the N content in unit mass%.
Moreover, the bolt which does not generate | occur | produce a coarse grain with excellent productivity is obtained by performing bolt processing, quenching, and tempering with respect to the steel which concerns on this embodiment by a well-known method.

The inventors have noticed a significant increase in ferrite hardness due to precipitation strengthening, and hence TiC and Ti (CN), etc., which are particles that cause an increase in steel hardness and impair the cold workability of steel. We studied a technology for suppressing the generation of coarse grains, which is different from the conventional technology for finely dispersing the particles. The above features are based on the following findings obtained by the present inventors by earnestly studying the technology for suppressing abnormal grain growth of austenite grains during quenching and heating of steel.

(1) With a very small amount of Bi of 0.0100% or less, abnormal grain growth of austenite crystal grains during quenching heating can be suppressed, and a cold-worked part having excellent dimensional accuracy and mechanical properties can be obtained.

(2) Austenite without depending on precipitates (TiC, Ti (CN), NbC) conventionally used as pinning particles (ie, without impairing the cold workability of steel) due to the effect of Bi described above. Abnormal grain growth of crystal grains can be suppressed. Thereby, the hardness of the rolling material after hot rolling can be suppressed and the cold workability of steel can be improved.

(3) On the other hand, if the Bi content exceeds 0.0100%, the hot ductility of the steel decreases, so that cracks and flaws are likely to occur in the steel manufacturing process (casting, rolling process, etc.). It was found that the yield decreased. Furthermore, it was also found that when the Bi content exceeds 0.0100%, grain boundary embrittlement occurs in the steel after quenching, and the mechanical properties of the steel are impaired. Therefore, it was also found that the Bi content is essential in the steel according to the present embodiment, but the content needs to be suppressed to an extremely low level.

Hereinafter, the steel according to the present embodiment will be described in detail.
First, chemical components of the steel of the present invention will be described. Hereinafter, the unit “%” regarding the chemical component indicates “% by mass”.

[C: 0.15-0.40%]
C is an element necessary for increasing the strength of steel having a tempered martensite structure. In order to set the tensile strength after quenching to 800 MPa or more, the C content needs to be 0.15% or more. The lower limit of the preferred C content is 0.17%, 0.19%, or 0.23%.
On the other hand, if the C content exceeds 0.40%, the hardness of the rolled material after hot rolling becomes too high, so the life of the cold forging die is significantly reduced. Therefore, the upper limit of C content is 0.40%. The upper limit of the preferable C content is 0.35%, 0.34%, 0.33%, or 0.30%.

[Mn: 0.10 to 1.50%]
Mn is an element effective for improving the hardenability of steel. In order to ensure the hardenability necessary for obtaining martensite by quenching, the Mn content needs to be 0.10% or more. The lower limit of the preferable Mn content is 0.20%, 0.35%, or 0.40%.
On the other hand, if the Mn content exceeds 1.50%, the hardness of the rolled material after hot rolling and before cold forging becomes too high, so that the life of the cold forging die is significantly reduced. Therefore, the upper limit of the Mn content is 1.50%. The upper limit of the preferable Mn content is 1.30%, 1.00%, or 0.80%.

[S: 0.002 to 0.020%]
S exists in steel as MnS, TiS, and Ti ₂ C ₂ S, and has the effect of suppressing abnormal grain growth of austenite crystal grains by acting as pinning particles during quenching heating. For this reason, it is necessary to make S content 0.002% or more. The lower limit of the preferable S content is 0.003%.
However, since the steel according to the present embodiment uses Bi to suppress abnormal grain growth, the S content may be smaller than that of the prior art. Furthermore, if the S content exceeds 0.020%, S causes embrittlement of the prior austenite grain boundaries of the steel after quenching, and deteriorates delayed fracture resistance (hydrogen embrittlement resistance). In addition, since Ti ₂ C ₂ S described above is a particle that impairs the machinability of the steel, if the S content exceeds 0.020%, the machinability of the steel may be deteriorated. Therefore, it is necessary to limit the S content to 0.020% or less. Preferably, the upper limit of the S content is 0.015%, 0.010%, or 0.005%.

[Ti: 0.005% to 0.050%]
Ti forms a compound with C, N, and S in steel and exists in steel as Ti-based inclusions such as TiN, Ti (CN), TiC, TiS, and Ti ₂ C ₂ S. By acting as a stop particle, it has the effect of suppressing abnormal grain growth of austenite grains. Further, Ti has a strong affinity for solute N in steel, and is therefore an extremely effective element for preliminarily fixing solute N in steel as TiN and suppressing the formation of BN. In boron steel, it is necessary to suppress the formation of BN in order to ensure the content of solute B that is effective in improving hardenability. Therefore, the Ti content needs to be 0.005% or more. The lower limit of the preferable Ti content is 0.010%, 0.015%, or 0.020%.
However, since the steel according to this embodiment uses Bi to suppress abnormal grain growth, the Ti content may be smaller than that of the prior art. Furthermore, if the Ti content exceeds 0.050%, Ti-based inclusion particles cause precipitation strengthening, and the hardness of the rolled material after hot rolling becomes too high. The service life is significantly reduced. In order to suppress the hardness of the rolled material after hot rolling while increasing the content of Ti-based inclusion particles, it is necessary to lower the hot rolling temperature, which means productivity, equipment life, etc. This is not preferable. Furthermore, when the Ti content is increased, a large amount of Ti ₂ C ₂ S, which is a particle that impairs the machinability of the steel, is produced, and the machinability is deteriorated. Becomes difficult. Therefore, the upper limit of the Ti content is 0.050%. A preferable Ti content is 0.040% or less, 0.030% or less, less than 0.030%, or 0.025% or less.

[B: 0.0005 to 0.0050%]
B is an element that contributes to improving the hardenability of steel when contained in a trace amount, and improves hardenability without increasing the hardness of the rolled material after hot rolling and before cold forging. An effect can be acquired and the hardness after cold forging and hardening can be increased. B is an essential element particularly for boron steel for bolts. Further, B has an effect of suppressing grain boundary fracture by segregating at the prior austenite grain boundaries and strengthening the prior austenite grain boundaries. In order to obtain the above effect, the B content needs to be 0.0005% or more. Preferably, the lower limit of the B content is 0.0010%, 0.0012%, or 0.0015%.
On the other hand, when the B content exceeds 0.0050%, the effect is saturated. Therefore, the B content is 0.0050% or less. Preferably, the upper limit of the B content is 0.0030%, 0.0025%, 0.0020%, or 0.0018%.

[Bi: 0.0010% to 0.0100%]
There has been no detailed study of the influence of a minute amount of Bi of about 0.0010% to 0.0100% on the structure during steel quenching. The present inventors have found that a small amount of Bi has an effect of preventing the generation of coarse grains by suppressing abnormal grain growth of austenite crystal grains during quenching heating. Moreover, since Bi content required in order to suppress abnormal grain growth is very small, the above-mentioned effect of Bi that suppresses the generation of coarse grains during quenching heating reduces the hardness of the rolled material after hot rolling. The inventors have also found that it can be obtained without increasing. In order to obtain the above effect, the Bi content needs to be 0.0010% or more. The lower limit of Bi content is preferably 0.0020%, 0.0025%, or 0.0030%.
On the other hand, when the Bi content exceeds 0.0100%, not only the effect is saturated, but also the hot ductility of the steel decreases, so cracks and flaws occur in the steel manufacturing process (casting, rolling process, etc.). It becomes easier and the yield decreases. Furthermore, if the Bi content exceeds 0.0100%, grain boundary embrittlement occurs in the steel after quenching, and the mechanical properties of the steel are impaired. Therefore, the Bi content is 0.0100% or less. The Bi content is preferably less than 0.0100%, 0.0080% or less, or 0.0060% or less.

[P: 0.020% or less]
P is an impurity and is an element that embrittles the old γ grain boundary and lowers the delayed fracture resistance (hydrogen embrittlement resistance) of the steel. Therefore, it is necessary to limit the P content to 0.020% or less. Preferably, the upper limit of the P content is 0.015%, 0.013%, or 0.010%.
Since P is not required to solve the problem of the steel according to the present embodiment, the lower limit value of the P content is 0%. However, in order to suppress the cost of the refining process for reducing the P content, the lower limit value of the P content may be 0.001%.

[N: 0.0100% or less]
When N forms a compound with B and is present in the steel as BN, the amount of solid solution B is reduced, and the effect of improving the hardenability by B is impaired. Since N is harmful in the steel according to this embodiment, the lower limit of the N content is 0%. However, in order to suppress the cost of the refining process for reducing the N content, the lower limit value of the N content may be 0.0001%, 0.0005%, or 0.0010%.
When the N content is large, the Ti content necessary for fixing N in the steel as TiN increases, so it is desirable to reduce the N content as much as possible. Therefore, it is necessary to limit the N content to 0.0100% or less. Preferably, the upper limit of the N content is 0.0070%, 0.0050%, or 0.0040%.

The spring steel according to the present embodiment may further contain one or more selected from the group consisting of Si, Cr, and Al as necessary, within a range described below. However, since Si, Cr, and Al are not essential, the lower limit of the content of each of Si, Cr, and Al is 0%.

[Si: 0% or more and less than 0.30%]
As described above, in the steel according to the present embodiment, the lower limit value of the Si content is 0%. However, Si is an element effective in improving the hardenability of steel and improving the temper softening resistance of martensite. In order to obtain the above effect, the Si content is preferably more than 0% or 0.01% or more. The lower limit value of the Si content may be 0.05% or 0.15%.
However, when the Si content is 0.30% or more, the amount of increase in the hardness of the steel (rolled material) after hot rolling and before cold forging increases, so the life of the die for cold forging decreases. To do. Therefore, the Si content is less than 0.30%. The upper limit of the preferable Si content is 0.27%, 0.25%, or 0.20%.

[Cr: 0 to 1.50%]
As described above, in the steel according to the present embodiment, the lower limit value of the Cr content is 0%. However, Cr is an effective element for improving the hardenability of steel and improving the temper softening resistance of martensite. In order to obtain the above effect, the Cr content is preferably more than 0% or 0.01% or more. The lower limit of the Cr content may be 0.10%, 0.20%, or 0.30%.
On the other hand, if the Cr content exceeds 1.50%, the hardness of the rolled material after hot rolling and before cold forging becomes too high, so the life of the cold forging die is significantly reduced. Therefore, the upper limit of the Cr content is 1.50%. The upper limit of the preferable Cr content is 1.20%, 1.00%, or 0.80%.

[Al: 0 to 0.050%]
Al is an element effective for deoxidation of steel. However, when deoxidation is performed with other elements (Si, Ti, etc.), it is not necessarily included. Therefore, the lower limit of the Al content is 0%. However, in order to obtain the deoxidation effect by Al, it is preferable to contain 0.001% or more, 0.005% or more, or 0.010% or more.
On the other hand, when the Al content exceeds 0.050%, problems such as generation of coarse inclusions and deterioration of the toughness of steel become significant. Therefore, even when Al is contained, the upper limit of the Al content is 0.050%. The upper limit of the Al content is preferably 0.040%, 0.030%, or 0.025%.

The spring steel according to the present embodiment may further contain one or more selected from the group consisting of Mo, Cu, Ni, and Nb as necessary, within a range described below. However, since Mo, Cu, Ni, and Nb are not essential, the lower limit of the contents of Mo, Cu, Ni, and Nb is 0%.

[Mo: 0 to 0.20%]
As described above, in the steel according to the present embodiment, the lower limit value of the Mo content is 0%. However, Mo is an element that contributes to improving the hardenability of steel even if its content is small. In order to obtain the above effect, the Mo content is preferably set to 0.02% or more. More preferably, the lower limit of the Mo content is 0.03%, 0.04%, or 0.05%.
On the other hand, since Mo is an expensive alloy element, if the Mo content exceeds 0.20%, it is disadvantageous in terms of manufacturing cost. Therefore, even when Mo is contained, the Mo content is set to 0.20% or less. Preferably, the upper limit of the Mo content is 0.16%, 0.13%, or 0.10%.

[Cu: 0 to 0.20%]
As described above, in the steel according to the present embodiment, the lower limit value of the Cu content is 0%. However, Cu is an element that improves the corrosion resistance of steel. In order to obtain the above effect, the Cu content is preferably set to 0.02% or more. More preferably, the lower limit of the Cu content is 0.05%.
On the other hand, when the Cu content exceeds 0.20%, the hot ductility of the steel is lowered, and problems such as impaired productivity during continuous casting become remarkable. Therefore, even when Cu is contained, the Cu content is set to 0.20% or less. Preferably, the upper limit of Cu content is 0.15%, 0.10%, or 0.08%.

[Ni: 0 to 0.20%]
As described above, in the steel according to the present embodiment, the lower limit value of the Ni content is 0%. However, Ni is an element that improves the corrosion resistance of steel and is also an effective element for improving the toughness of steel. In order to obtain the above effects, the Ni content is preferably 0.02% or more. More preferably, the lower limit of the Ni content is 0.03%, 0.04%, or 0.05%.
On the other hand, since Ni is an expensive alloy element, if the Ni content exceeds 0.20%, it is disadvantageous in terms of manufacturing cost. Therefore, even when Ni is contained, the Ni content is 0.20% or less. Preferably, the upper limit of the Ni content is 0.15%, 0.12%, 0.10%, or 0.08%.

[Nb: 0 to 0.030%]
As described above, in the steel according to the present embodiment, the lower limit value of the Nb content is 0%. However, Nb forms a compound with C in steel and exists in steel as Nb-based inclusions such as NbC or TiNb (CN), and suppresses abnormal growth of austenite grains as pinning particles during quenching heating. Has the effect of In order to obtain the above effect, the Nb content is preferably set to 0.002% or more. More preferably, the lower limit of Nb content is 0.003%, 0.005%, or 0.006%.
On the other hand, when the Nb content exceeds 0.030%, not only the effect is saturated, but also the Nb-based inclusions cause precipitation strengthening, so that the manufacturability during continuous casting is impaired. Alternatively, in this case, Nb-based inclusions cause precipitation strengthening, so that the hardness of the rolled material after hot rolling becomes too high. Therefore, when the Nb content exceeds 0.030%, problems such as a decrease in manufacturability and a significant decrease in the life of a cold forging die become prominent. Therefore, even when Nb is contained, the Nb content is set to 0.030% or less. Preferably, the upper limit of Nb content is 0.015%, 0.013%, or 0.010%.

The steel according to this embodiment contains the above alloy components, and the balance of the chemical components contains Fe and impurities. In the present embodiment, impurities are components mixed due to raw materials such as ores and scraps and other factors when industrially producing steel materials, and do not impair the effects of the steel according to the present embodiment. Means what is the amount.

[N fixed index I _FN : preferably 0 or more]
In order to obtain the above-described effect of containing B, it is necessary to suppress the formation of BN by reducing N (solid solution N) dissolved in the steel. Therefore, it is desirable to reduce the content of N in the steel and to stably fix N in the form of TiN by containing Ti in the steel, thereby reducing the amount of solute N. In order to obtain the above effect by fixing N with Ti, it is preferable to set the N fixed index _IFN defined by the following formula 1 to 0 or more. The lower limit of the N fixed index _{I FN} 0.0005,0.0010,0.0014, or may be 0.0050. However, the steel according to the present embodiment is softened before cold forging as long as the Ti content and the N content are controlled within the above-described range even if the N fixed index _IFN is not particularly limited. The generation of coarse grains during quenching can be suppressed.
I _FN = [Ti] −3.5 × [N] (Formula 1)
In addition, [Ti] and [N] in the above formula 1 indicate the Ti content and N content in the steel in unit mass%, and 0% when these elements are not contained.

[Ti—Nb-based precipitate formation index I _P : preferably 0.0100 or less]
As described above, it is preferable to reduce the amount of solid solution N by fixing Ti as TiN using Ti. However, it is not preferable to contain an amount of Ti that exceeds the amount necessary to fix TiN. As described above, Ti combines with C and S to form fine precipitates, and these fine precipitates may adversely affect the properties of the steel according to the present embodiment. The present inventors have also found that Nb has the same function as Ti.

Specifically, fine TiC, Ti (CN), NbC, TiNb (CN), and Ti—Nb-based precipitates such as Ti ₂ C ₂ S, which are precipitates existing in steel, are pinned during quenching heating. By suppressing abnormal grain growth of austenite crystal grains as stopping grains, it has the effect of suppressing the generation of coarse grains. However, when these Ti—Nb-based precipitate particles are dispersed in a large amount in the structure after hot rolling, there is a side effect that the hardness of the ferrite increases due to precipitation strengthening by fine precipitate particles. . Therefore, when these Ti—Nb-based precipitate particles are dispersed in an excessively large amount in the steel, the hardness of the rolled material after hot rolling becomes too high. Problems such as a significant decrease in the service life of the battery become noticeable. Furthermore, as described above, Ti ₂ C ₂ S causes deterioration of machinability. Therefore, in the steel according to the present embodiment, it is preferable to limit the amount of these Ti—Nb-based precipitate particles.

In order to suppress the hardness after rolling after the hot rolling, it is preferable that the Ti-Nb-based precipitates generated index I _P, which is calculated by the following equation 2 and 0.0100 or less. The Ti-Nb-based precipitates generated index _{I P} 0.0075 or less, less than 0.0050, 0.0045 or less, 0.0040 or less, or 0.0035 may be less. However, even without particularly limiting the Ti-Nb-based precipitates generated index I _P, Ti content within the range described above, Nb content, and N as long as the content is controlled, according to the present embodiment Steel is softened before cold forging and can suppress the generation of coarse grains during quenching.
I _P = 0.3 × [Ti] + 0.15 × [Nb] − [N] (Formula 2)
[Ti], [N] and [Nb] in the above formula 2 indicate the Ti content, the N content, and the Nb content in the steel in unit mass%, and when these elements are not contained. 0%.

Next, the suitable manufacturing method of the steel of this embodiment is demonstrated.
In order to manufacture the steel of the present embodiment, the above-described chemical component steel is melted in a converter, and a slab is obtained by continuous casting through a secondary refining process as necessary. The slab is reheated and subjected to ingot rolling to obtain a wire rolling material (steel slab) having a cross section of, for example, 162 mm square (vertical 162 mm × lateral 162 mm). Next, the steel slab is heated at a temperature of about 1000 to 1280 ° C. and subsequently subjected to wire rod rolling to form a wire rod having a diameter of 6 to 20 mm. Then, after being hotly wound into a coil shape by a winding device, it is cooled to room temperature. In this way, the steel of this embodiment is obtained.

In the steel according to the present embodiment, the amount of Ti-based precipitated particles that cause precipitation strengthening is suppressed. Therefore, in the steel manufacturing method according to the present embodiment, hot rolling is performed in order to suppress the hardness of the steel. It is not necessary to apply a load to the hot rolling equipment by lowering the temperature, and defects such as cracks and wrinkles due to an increase in hardness are less likely to occur in the steel. Furthermore, the hardness of the steel according to the present embodiment is suppressed without performing annealing after hot rolling. Therefore, the steel according to this embodiment is excellent in terms of high productivity.

According to the steel of this embodiment, both softening before cold forging and suppression of the generation of coarse grains during quenching can be achieved. Moreover, the steel of this embodiment is excellent in manufacturability without cracking during casting or rolling.

The hardness of the steel according to the present embodiment is not particularly limited because it can be appropriately adjusted according to the application. However, when it is necessary to ensure cold forgeability, the hardness of the steel according to this embodiment is preferably Hv 180 or less, and more preferably Hv 170 or less, or Hv 160 or less. . The lower limit value of the hardness of the steel according to the present embodiment is not particularly limited, but is considered to be substantially about Hv130 or about Hv140 in view of its chemical composition. Even if it does not anneal after hot rolling, the steel which concerns on this embodiment can make the hardness into the above-mentioned suitable range. Moreover, the steel according to the present embodiment is excellent in machinability.

Further, the steel according to the present embodiment is heated to a temperature of, for example, 840 ° C. to 1100 ° C., held for 30 minutes, and then quenched under water cooling or oil cooling, and further in a temperature range of 150 ° C. to 450 ° C. When the tempering process is performed by heating and holding, the tensile strength can be 800 MPa or more. Therefore, the steel according to the present embodiment is suitable as a material for parts that require high strength. However, when using the steel which concerns on this embodiment as steel for hardening, heat processing conditions are not specifically limited, It can select suitably according to a use.

Although the use of the steel according to the present embodiment is not particularly limited, it is preferable that the steel is applied to high-strength mechanical parts manufactured by cold forging and quenching, particularly high-strength bolts. When the steel according to the present embodiment having high cold forgeability is used as a material for high-strength mechanical parts, wear of the mold during cold forging can be suppressed and the life of the mold can be improved. Moreover, since the cost of an expensive metal mold | die can be reduced, it can contribute to the reduction of the manufacturing cost of the high intensity | strength volt | bolt whose tensile strength is 800 MPa or more especially.

Next, the present invention will be described using examples, but the present invention is not limited to the following examples.

First, steels having the chemical components shown in Table 1-1 and Table 1-2 were melted by a converter and further cast into slabs by continuous casting. In Table 1-1 and Table 1-2, for elements whose content is less than or equal to the impurity level, the content display is left blank, and the N fixed index I _FN and Ti—Nb-based precipitate formation index I _P In the calculation of “0% by mass”, it was considered. In Table 1-1 and Table 1-2, values outside the specified range of the present invention are underlined. It was confirmed whether or not a slab surface crack occurred in the slab obtained in this way. In the confirmation of the slab surface crack, the scale of the slab surface was removed with a check scarf, and then the slab surface was observed to investigate the crack depth. When a crack having a depth of 1 mm or more was detected on the surface of the slab, the slab surface crack during continuous casting was determined as “present”, and the manufacturability was determined as “failed”. The manufacturability evaluation results are shown in Tables 2-1 and 2-2.

The cast slab was subjected to soaking diffusion treatment and partial rolling as necessary to obtain a wire rolling material (steel piece) having a cross section of 162 mm square (vertical 162 mm × lateral 162 mm). Next, the steel slab was heated at a temperature of about 1000 to 1280 ° C. and subsequently subjected to wire rod rolling to obtain a wire rod (spring steel) having a diameter of 10 mm.

A test piece for Vickers hardness measurement was cut out from the rolled wire. Specifically, a test piece having a cross section including the central axis of the wire in a direction parallel to the rolling direction was cut out. After the cut cross section was polished, the Vickers hardness of a portion (1/4 part) having a depth of 1/4 of the diameter of the wire from the surface of the wire was measured. The test load is 10 kgf, and the average value obtained by measuring four points is described as “hardness after rolling” in Tables 2-1 and 2-2. This is an index for predicting the life of a die for cold forging. did. About what the hardness of a rolling material exceeds HV180, since the sufficient improvement effect of the metal mold | die for cold forging was not acquired, it was determined that "cold forgeability" was "failed". The evaluation results of cold forgeability are shown in Tables 2-1 and 2-2.

In addition, in order to simulate the effect of wire drawing or cold forging (cold working) when processing the wire into a bolt shape, after performing cold drawing with a surface reduction ratio of 70% on the wire, Heating was performed at a temperature of 840 ° C. to 1100 ° C. for 30 minutes, and quenching by water cooling was performed to freeze the austenite structure as a prior austenite grain boundary of the martensite structure. Thereafter, tempering is performed on the quenched specimen as necessary, and a specimen having a cross section including the center of the drawn material is cut out in a direction parallel to the rolling / pulling direction. It was. After the cross section of the cut specimen was polished, the prior austenite grain boundaries appeared by corrosion and observed with an optical microscope to measure the prior austenite grain size after quenching and tempering. The prior austenite grain size was measured according to JISG0551. The measurement field of view was 400 times magnification and 10 fields or more, and a test piece in which even one large crystal grain having a prior austenite grain size of No. 5 or less was present was determined to have coarse grains. The limit (minimum) heating temperature at which coarse grains are generated, which is clarified by observing and measuring the prior austenite grain size for test pieces heated to various temperatures, is the crystal grain coarsening temperature of the test piece. Defined and used as an index of the resistance to grain coarsening. Those having a crystal grain coarsening temperature of 900 ° C. or lower were judged to be “failed” because they had poor resistance to crystal grain coarsening. The results of measuring the crystal grain coarsening temperature are shown in Tables 2-1 and 2-2.

From Tables 2-1 and 2-2, A1 to A32, which are examples of the present invention, have low hardness of the wire after rolling and can be expected to improve the life of the cold forging die. Slab debris rate because no coarse grains are generated even when heated above 900 ° C during quenching and heating after cold working, and surface cracks of the slab do not occur during continuous casting It is clear that it is low and therefore is excellent in manufacturability. The inventive examples A1 to A32 after the heat treatment for measuring the prior austenite grain size described above all had a tensile strength of 800 MPa or more.

On the other hand, in the case of the comparative example, any of the cold forgeability, coarse grain prevention characteristics, and manufacturability is inferior. That is, since B1 to B4 have too much Bi added, the hot ductility is lowered and the manufacturability is poor. In B5 to B7, Bi was not added, or the addition amount was too small, so the coarse grain preventing properties were inferior. In B8 and B9, the added amount of Ti is too large, or the N content is small relative to the added Ti amount and the Ti—Nb-based precipitate formation index _IP is exceeded. It was inferior to the forgeability.

According to the present invention, it is possible to provide steel that can achieve both softening during cold forging and suppression of generation of coarse grains during quenching after cold forging. In addition, the steel according to the present invention is excellent in manufacturability because it is not cracked during casting or rolling, and can be manufactured under conditions that do not impose a load on manufacturing equipment. By applying the steel according to the present invention to a cold forged part, wear of the mold during cold forging can be suppressed, and the life of the mold can be improved. In addition, by applying the steel according to the present invention to cold forged parts, the cost of expensive dies can be reduced, which can contribute to the reduction of the manufacturing cost of high-strength bolts having a tensile strength of 800 MPa or more. it can. Furthermore, the steel according to the present invention is excellent in machinability. Therefore, the present invention has a great industrial contribution.

Claims (5)

Chemical component is unit mass%,
C: 0.15% to 0.40%,
Mn: 0.10% to 1.50%,
S: 0.002 to 0.020%,
Ti: 0.005% to 0.050%,
B: 0.0005 to 0.0050%,
Bi: 0.0010% to 0.0100%,
P: 0.020% or less,
N: 0.0100% or less,
Si: 0% or more and less than 0.30%,
Cr: 0 to 1.50%,
Al: 0 to 0.050%,
Mo: 0 to 0.20%,
Cu: 0 to 0.20%,
Ni: 0 to 0.20%, and Nb: 0 to 0.030%,
A steel characterized in that the balance consists of Fe and impurities. The chemical component is unit mass%,
Si: 0.01% or more and less than 0.30%,
Cr: 0.01 to 1.50%, and Al: 0.001 to 0.050%
The steel according to claim 1, comprising one or more selected from the group consisting of: The chemical component is unit mass%,
Mo: 0.02 to 0.20%,
Cu: 0.02 to 0.20%,
Ni: 0.02 to 0.20%, and Nb: 0.002 to 0.030%
The steel according to claim 1 or 2, comprising one or more selected from the group consisting of: The steel according to any one of claims 1 to 3, wherein the N fixed index _IFN defined by the following formula 1 is 0 or more.
I _FN = [Ti] −3.5 × [N] (Formula 1)
Here, [Ti] is the Ti content in unit mass%, and [N] is the N content in unit mass%. Steel according to any one of claims 1 to 4 or less is defined by equation 2 Ti-Nb-based precipitates generated index I _P is equal to or is 0.0100 or less.
I _P = 0.3 × [Ti] + 0.15 × [Nb] − [N] (Formula 2)
Here, [Ti] is the Ti content in unit mass%, [Nb] is the Nb content in unit mass%, and [N] is the N content in unit mass%.