CN109790602B

CN109790602B - steel

Info

Publication number: CN109790602B
Application number: CN201680089493.8A
Authority: CN
Inventors: 久保田学; 志贺聪; 长谷川一
Original assignee: Nippon Steel and Sumitomo Metal Corp
Current assignee: Nippon Steel Corp
Priority date: 2016-09-28
Filing date: 2016-09-28
Publication date: 2021-03-02
Anticipated expiration: 2036-09-28
Also published as: EP3521469A1; JP6798557B2; EP3521469A4; JPWO2018061101A1; KR20190041502A; US20190256957A1; WO2018061101A1; CN109790602A

Abstract

本发明的一技术方案涉及的钢，化学成分以质量％为单位含有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％以下、N：0.0100％以下、Si：0％以上且小于0.30％、Cr：0～1.50％、Al：0～0.050％、Mo：0～0.20％、Cu：0～0.20％、Ni：0～0.20％、和Nb：0～0.030％，余量包含Fe和杂质。The steel according to one aspect of the present invention contains, in mass %, 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% or more and less than 0.30%, Cr: 0-1.50%, Al: 0-0.050%, Mo: 0- 0.20%, Cu: 0 to 0.20%, Ni: 0 to 0.20%, and Nb: 0 to 0.030%, and the balance contains Fe and impurities.

Description

Steel

Technical Field

The present invention relates to steel.

Background

Cold forging (including roll forming) is widely used as a method for manufacturing relatively small machine parts such as bolts because it can improve the surface skin, dimensional accuracy, and the like of products and can improve the yield as compared with hot forging. When machine parts are manufactured by cold forging, for example, a carbon steel or alloy steel for machine structure having a medium carbon as defined in JIS G4051, JIS G4052, JIS G4104, JIS G4105, JIS G4106 and the like is used as a raw material, and a final product is often manufactured through manufacturing steps such as wire hot rolling-annealing (or spheroidizing annealing) -wire drawing-cold forging-quenching and tempering. The above-described conventional manufacturing process is characterized by adding an annealing or spheroidizing annealing step before cold forging. The reason why the additional annealing or spheroidizing annealing is performed before the cold forging is because there are the following manufacturing problems: carbon steel and alloy steel of medium carbon have high hardness of rolled material in a hot rolled state (that is, in a case where the steel is air-cooled without heat treatment after hot rolling), and the loss of a die at the time of cold forging is significant, so that the manufacturing cost is high; in the hot rolled state, the ductility of the material is insufficient, and cracks are likely to occur during cold forging, so that the yield is reduced; and so on.

However, since annealing requires a very large cost, it is necessary to develop a steel material that can omit the annealing step in order to reduce the manufacturing cost of the component. In response to such a demand, boron steel for so-called bolts has been developed in which a small amount of B is added to a steel material (for example, patent documents 1 and 3). The boron steel is characterized in that: by reducing the carbon content of the steel and the amount of the alloying elements such as Cr and Mo, the hardness of the hot-rolled wire rod is reduced and the ductility is improved, so that annealing is not necessary and the effect of improving the hardenability by the addition of a trace amount of B which does not increase the hardness of the rolled material is utilized to compensate for the reduction in hardenability caused by the reduction in the amount of the alloying elements added.

In order to exhibit the effect of improving hardenability by adding a small amount of B, it is necessary that B be in a solid solution state in austenite. On the other hand, when nitrogen in a solid solution state is present in the steel, BN is generated, and the amount of solid solution B (B solid dissolved in the steel) decreases, whereby the effect of B on improving hardenability is lost. Therefore, the following work is generally performed: in boron steel, N in steel is fixed as TiN in advance by adding Ti having a strong affinity with N, and the generation of BN is suppressed. For example, patent document 4 describes that precipitation of BN is suppressed by setting Ti/N (mass% ratio) to 4 or more. In principle, if Ti/N is 3.42 or more, precipitation of BN can be suppressed.

However, in the conventional boron steel, in comparison with conventional steels, some of the austenite grains tend to grow abnormally large and coarsen, that is, so-called coarse grains tend to be generated during quenching and heating. The component having coarse grains is deteriorated in dimensional accuracy due to an increase in heat treatment strain generated during quenching, and is deteriorated in component characteristics such as an impact value, fatigue strength, and delayed fracture characteristics. Therefore, in particular, in a high-strength bolt having a tensile strength of 800MPa or more, it is a large practical problem to prevent the generation of coarse crystal grains. In order to suppress the generation of such coarse grains due to abnormal grain growth, it is effective to disperse a large number of pinning particles (precipitates and the like) in the structure, that is, to disperse a large number of fine particles in order to pin the grain boundaries of austenite grains.

The following two reasons why coarse grains are likely to be generated in boron steel are mainly described below.

(1) In the case of using boron steel as a component material, since the annealing step after cold forging of boron steel is omitted, boron steel is directly heated from the cold worked structure to the austenite region. In this case, because the austenite grains are excessively refined and the grain diameter is partially uneven due to the influence of cold working, abnormal grain growth is likely to occur in some of the grains.

(2) In the boron steel described above, since N in the steel is fixed as TiN by the addition of Ti, AlN that effectively functions as pinning particles in conventional carbon steel and alloy steel is not generated, and since TiN is coarser than AlN, it is not possible to finely disperse TiN, and it is difficult to secure the number of pinning particles necessary to prevent coarse crystal grains.

Since the annealing step is omitted, the factor of (1) above cannot be avoided, and therefore, how to secure the number of pinning particles in boron steel becomes the main factor of improving (2) to prevent the generation of coarse crystal grains.

Under such circumstances, techniques for preventing the generation of coarse crystal grains in boron steel have been proposed. For example, patent documents 5 and 6 describe that instead of AlN and TiN, TiC and ti (cn) which are precipitates finer than TiN are used as pinning particles. In these techniques, in order to secure the number of pinning particles necessary for preventing coarse crystal grains, there are defined: the total number of the dispersed particles in the steel before quenching and heating and after hot rolling was 20 particles/100. mu.m²TiC and Ti (CN) having a diameter of 0.2 μm or less. By dispersing such fine precipitates in a large amount in advance before quenching and heating, these precipitates function as pinning particles for pinning austenite grain boundaries during quenching and heating. According to this technique, coarse grains can be stably prevented from being generated in boron steel, and steel to which this technique is applied is now widely used as an inexpensive steel material for bolts, which can omit an annealing step.

However, the above-described techniques have disadvantages. That is, when a large amount of fine TiC and Ti (CN) are dispersed in the structure after hot rolling, there is a side effect that the hardness of ferrite is increased due to precipitation strengthening by fine precipitate particles, and thus there is a problem that the softening effect of the hot rolled material by boron tempering is reduced. That is, when the amount of fine TiC and Ti (CN) is increased, the generation of coarse crystal grains is suppressed, but the hardness of the rolled material is increased by precipitation strengthening, and the life of the cold forging die is reduced. Conversely, if the amount of TiC and Ti (CN) is reduced, the hardness of the rolled material can be reduced, but coarse crystal grains are formed. That is, when fine TiC or Ti (CN) is used, there is a contradictory relationship between suppression of generation of coarse crystal grains and suppression of hardness of the rolled material before cold forging. Therefore, it is difficult to achieve both complete softening of the rolled material and stable suppression of coarse grains by using only the above-described technique.

Patent document 7 also describes the same technical idea as the above-described technique for preventing the generation of coarse crystal grains in boron steel. That is, the relationship among the contents of Ti, Nb, Al, and N is set within a certain range, whereby carbonitrides of these elements are dispersed in the steel to prevent the coarsening of crystal grains. Patent document 7 also describes that the machinability is improved by adding 0.01% or more of Bi. However, patent document 7 discloses only an effect of improving the machinability as an effect of Bi. The relationship between Bi and the coarsening characteristics of crystal grains is not described at all. Since Bi is added for the purpose of improving the machinability, patent document 7 only studies on the case where a large amount of Bi is added. In this case, as described in patent document 7, there is a concern that the hot workability may be reduced by the Bi addition.

Patent document 8 discloses a case hardening steel that exhibits superior grain coarsening resistance even when carburized at high temperatures, and also exhibits superior cold workability even without soft annealing, as compared to conventional examples. However, patent document 8 also proposes only: as means for securing the property of resisting coarsening of crystal grains, fine Ti carbides, Ti-containing composite carbides, and the like are used. In patent document 8, since the hot rolling temperature is extremely low in order to ensure cold workability, the productivity of the steel for case hardening is impaired.

Prior art documents

Patent document

Patent document 1 Japanese patent application laid-open No. Hei 5-339676

Patent document 2 Japanese examined patent publication No. 5-63524

Patent document 3 Japanese patent application laid-open No. Sho 61-253347

Patent document 4 Japanese patent application laid-open No. Hei 3-47918

Patent document 5 Japanese patent No. 3443285

Patent document 6 Japanese patent No. 3490293

Patent document 7 Japanese laid-open patent publication No. 2000-328189

Patent document 8 Japanese laid-open patent publication No. 2006-265704

Disclosure of Invention

One of the problems of the cold forging steel is: in order to improve cold forgeability of steel and productivity of steel, the steel is kept soft without annealing after hot rolling and before cold forging and without using production conditions that impair productivity. Another problem of the cold forging steel is: in order to impart high strength to a mechanical part, high hardenability is required after cold forging. Another problem of the cold forging steel is: in order to prevent deterioration of the mechanical parts such as dimensional accuracy, impact value, fatigue strength, and delayed fracture characteristics, generation of coarse crystal grains is suppressed during quenching after cold forging. As described above, the conventional techniques cannot solve all of these problems at the same time. As a means for suppressing the generation of coarse crystal grains, TiC and ti (cn) proposed in the prior art are used, because steel after hot rolling and before cold forging is hardened by precipitation strengthening, cold forgeability and productivity of the steel are impaired.

The present invention has been made in view of the above problems. That is, an object of the present invention is to provide a steel excellent in all of manufacturability, cold forgeability, and mechanical properties after quenching by suppressing generation of coarse crystal grains at the time of quenching without using Ti carbide such as TiC and Ti (cn) and Ti carbonitride.

The gist of the present invention is as follows.

(1) The steel according to one embodiment of the present invention contains, in terms of mass%, C: 0.15-0.40%, Mn: 0.10 to 1.50%, S: 0.002-0.020%, Ti: 0.005-0.050%, B: 0.0005 to 0.0050%, Bi: 0.0010-0.0100%, P: 0.020% or less, N: 0.0100% or less, Si: 0% or more and 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 to 0.030%, the balance comprising Fe and impurities.

(2) The steel according to the item (1), wherein the chemical component may contain, in mass%, a chemical component selected from the group consisting of Si: 0.01% or more and less than 0.30%, Cr: 0.01 to 1.50%, and Al: 0.001 to 0.050% of at least one of the above two compounds.

(3) The steel according to the above (1) or (2), wherein the chemical component may contain, in mass%, a chemical component selected from the group consisting of Mo: 0.02 to 0.20%, Cu: 0.02 to 0.20%, Ni: 0.02 to 0.20%, and Nb: 0.002-0.030% of at least one of the above-mentioned two compounds.

(4) The steel according to any one of the above (1) to (3), wherein the N fixation index I defined by the following formula 1 is adopted_FNMay be 0 or more.

I_FN＝[Ti]－3.5×[N]… (formula 1)

Wherein [ Ti ] is a Ti content in terms of mass%, and [ N ] is an N content in terms of mass%.

(5) The steel according to any one of the above (1) to (4), wherein a Ti-Nb-based precipitate formation index I defined by the following formula 2 is used_PMay be 0.0100 or less.

I_P＝0.3×[Ti]+0.15×[Nb]－[N]… (formula 2)

Wherein [ Ti ] is a Ti content in terms of mass%, [ Nb ] is a Nb content in terms of mass%, and [ N ] is an N content in terms of mass%.

According to the present invention, it is possible to provide a steel which can be softened before cold forging and can suppress the generation of coarse crystal grains during quenching after cold forging. The steel according to the present invention is excellent in manufacturability because it does not cause cracking during casting, rolling, and the like, and can be produced under conditions within a range that does not put a burden on production facilities. By applying the steel according to the present invention to a cold-forged part, the wear of the die during cold forging can be suppressed, and the life of the die can be increased. Further, since the application of the steel according to the present invention to a cold-forged part can reduce the cost of a high-priced die, it can contribute to the reduction of the manufacturing cost of a high-strength bolt having a tensile strength of 800MPa or more in particular. Furthermore, the steel according to the present invention is also excellent in machinability. Therefore, the present invention contributes greatly to the industry.

Detailed Description

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

(a) The steel according to the present embodiment contains, in terms of mass%, C: 0.15-0.40%, Mn: 0.10 to 1.50%, S: 0.002-0.020%, Ti: 0.005-0.050%, B: 0.0005 to 0.0050%, Bi: 0.0010-0.0100%, P: 0.020% or less, N: 0.0100% or less, Si: 0% or more and 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 to 0.030%, the balance comprising Fe and impurities.

(b) The steel according to the item (a), wherein the chemical component may contain, in mass%, a chemical component selected from the group consisting of Si: 0.01% or more and less than 0.30%, Cr: 0.01 to 1.50%, and Al: 0.001 to 0.050% of at least one of the above two compounds.

(c) The steel according to the above (a) or (b), wherein the chemical component may contain, in mass%, a chemical component selected from the group consisting of Mo: 0.02 to 0.20%, Cu: 0.02 to 0.20%, Ni: 0.02 to 0.20%, and Nb: 0.002-0.030% of at least one of the above-mentioned two compounds.

(d) The steel according to any one of the above (a) to (c), wherein the N fixation index I defined by the following formula 1 is adopted_FNMay be 0 or more.

I_FN＝[Ti]－3.5×[N]… (formula 1)

(e) The steel according to any one of the above (a) to (d), wherein a Ti-Nb precipitate formation index I defined by the following formula 2 is used_PMay be 0.0100 or less.

I_P＝0.3×[Ti]+0.15×[Nb]－[N]… (formula 2)

Further, by performing bolt working, quenching, and tempering by a known method on the steel according to the present embodiment, a bolt without coarse crystal grains can be obtained with excellent productivity.

The present inventors have studied about a technique for suppressing the generation of coarse crystal grains, which is different from a conventional technique for finely dispersing particles such as TiC and ti (cn), which are particles that cause a significant increase in ferrite hardness due to precipitation strengthening, thereby causing an increase in steel hardness and impairing cold workability of steel. The above-described characteristics are based on the following findings that the present inventors have conducted intensive studies on a technique for suppressing abnormal grain growth of austenite grains during quenching and heating of steel.

(1) By using Bi in such an extremely small amount of 0.0100% or less, abnormal grain growth of austenite grains during quenching and heating is suppressed, and a cold-worked component excellent in dimensional accuracy, mechanical properties, and the like can be obtained.

(2) The above-described effect of Bi can suppress abnormal grain growth of austenite grains without relying on precipitates (TiC, Ti (CN), NbC) conventionally used as pinning particles (i.e., without impairing cold workability of steel). This can suppress the hardness of the rolled material after hot rolling and improve the cold workability of the steel.

(3) On the other hand, it was clarified that: when 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 production process (casting, rolling, etc.), and the yield of the steel decreases. But also clarified: 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. Thus, it was also clarified that: in the steel according to the present embodiment, although Bi is required to be contained, the content thereof needs to be suppressed to an extremely low level.

Hereinafter, the steel according to the present embodiment will be described in detail.

First, the chemical composition of the steel of the present invention will be described. Hereinafter, the unit "%" of the chemical component represents "% by mass".

[C：0.15～0.40％]

C is an element required to improve the strength of steel having a tempered martensite structure. In order to achieve a tensile strength of 800MPa or more after quenching, 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, and the life of the cold forging die is significantly reduced. Therefore, the upper limit of the C content is set to 0.40%. The preferred upper limit of the C content is 0.35%, 0.34%, 0.33%, or 0.30%.

[Mn：0.10～1.50％]

Mn is an element effective for improving hardenability of steel. In order to ensure hardenability required for obtaining martensite by quenching, the Mn content needs to be 0.10% or more. The lower limit of the preferred 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, and therefore the life of the die for cold forging is significantly reduced. Therefore, the upper limit of the Mn content is set to 1.50%. The preferred upper limit of the Mn content is 1.30%, 1.00%, or 0.80%.

[S：0.002～0.020％]

S has the following effects: as MnS, TiS, and Ti₂C₂S is present in the steel and acts as pinning particles during quenching and heating, thereby suppressing abnormal grain growth of austenite grains. Therefore, the S content needs to be 0.002% or more. The lower limit of the preferable S content is 0.003%.

However, in the steel according to the present embodiment, since Bi is used to suppress abnormal grain growth, the S content is sufficient even if it is smaller than that in the prior art. When the S content exceeds 0.020%, S embrittles the prior austenite grain boundary of the steel after quenching to make the steel resistant to delayed fractureCracking characteristics (hydrogen embrittlement resistance) are reduced. And, due to the above-mentioned Ti₂C₂Since S is a particle that impairs the machinability of steel, if the S content exceeds 0.020%, the machinability of steel may deteriorate. Therefore, the S content needs to be limited to 0.020% or less. The upper limit of the S content is preferably 0.015%, 0.010%, or 0.005%. [ Ti: 0.005-0.050%]

Ti has the following effects: form a compound with C, N, S in steel as TiN, Ti (CN), TiC, TiS, Ti₂C₂Ti-based inclusions such as S are present in the steel and act as pinning particles during quenching and heating, thereby suppressing abnormal grain growth of austenite grains. Since Ti has a strong affinity for the solid solution N in the steel, Ti is an element that is extremely effective for suppressing the generation of BN by fixing the solid solution N in the steel as TiN. In boron steel, in order to ensure the content of solid-solution B effective for improving hardenability, it is necessary to suppress the generation of BN. Therefore, the Ti content needs to be 0.005% or more. The lower limit of the Ti content is preferably 0.010%, 0.015%, or 0.020%.

However, in the steel according to the present embodiment, since abnormal grain growth is suppressed by Bi, the Ti content is sufficient even if it is smaller than that in the prior art. When the Ti content exceeds 0.050%, precipitation strengthening occurs in Ti-based inclusion particles, and the hardness of the rolled material after hot rolling becomes too high, so that the life of the die for cold forging is significantly reduced. In order to increase the content of Ti-based inclusion particles and to suppress the hardness of the rolled material after hot rolling, it is necessary to lower the hot rolling temperature, but this is not preferable from the viewpoint of productivity, equipment life, and the like. When the content of Ti is increased, a large amount of Ti particles damaging the machinability of the steel are generated₂C₂S causes deterioration in machinability, and therefore it is difficult to apply cutting work to the steel according to the present embodiment. Therefore, the upper limit of the Ti content is set to 0.050%. The Ti content is preferably 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 improvement of hardenability of steel when contained in a trace amount, and can achieve the effect of improving hardenability without increasing the hardness of the rolled material after hot rolling and before cold forging, and can increase the hardness after cold forging and quenching. B is an element necessary for boron steel for bolts in particular. In addition, B has the effect of strengthening the prior austenite grain boundaries by segregating at the prior austenite grain boundaries, thereby inhibiting grain boundary destruction. In order to obtain the above-described effects, the B content needs to be 0.0005% or more. The lower limit of the B content is preferably 0.0010%, 0.0012%, or 0.0015%.

On the other hand, when the B content exceeds 0.0050%, the effect thereof is saturated. Therefore, the B content is set to 0.0050% or less. The upper limit of the B content is preferably 0.0030%, 0.0025%, 0.0020%, or 0.0018%.

[Bi：0.0010～0.0100％]

The effect of a small amount of Bi of about 0.0010 to 0.0100% on the structure of steel during quenching has not been studied in detail so far. The inventor finds that: the trace amount of Bi suppresses abnormal grain growth of austenite grains during quenching and heating, thereby having the effect of preventing the generation of coarse grains. Further, the present inventors have found that the above-described effect of Bi for suppressing the generation of coarse grains during quenching and heating can be obtained without increasing the hardness of the rolled material after hot rolling because the content of Bi necessary for suppressing abnormal grain growth is small. In order to obtain the above-described effects, the Bi content needs to be 0.0010% or more. The lower limit of the Bi content is preferably 0.0020%, 0.0025%, or 0.0030%.

On the other hand, when the Bi content exceeds 0.0100%, the effect is saturated and the hot ductility of the steel is lowered, so that cracks and flaws are likely to occur in the steel production process (casting, rolling, etc.), and the yield is lowered. 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, the Bi content is set to 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 original γ grain boundary and reduces delayed fracture resistance (hydrogen embrittlement resistance) of the steel. Therefore, the P content needs to be limited to 0.020% or less. The upper limit of the P content is preferably 0.015%, 0.013%, or 0.010%.

Since P is not necessary to solve the problems of the steel according to the present embodiment, the lower limit of the P content is 0%. However, in order to reduce the cost of the refining step for reducing the P content, the lower limit of the P content may be set to 0.001%.

[ N: 0.0100% or less ]

When N forms a compound with B and is present as BN in steel, the amount of B in solid solution is reduced, and the effect of B on improving hardenability is impaired. Since N is harmful in the steel according to the present embodiment, the lower limit of the N content is 0%. However, in order to suppress the cost of the refining step for reducing the N content, the lower limit of the N content may be set to 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, and therefore it is desirable to reduce the N content as much as possible. Therefore, the N content needs to be limited to 0.0100% or less. The upper limit of the N content is preferably 0.0070%, 0.0050%, or 0.0040%.

The spring steel according to the present embodiment may further contain 1 or two or more selected from Si, Cr, and Al in the ranges described below as necessary. 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 of the Si content is 0%. However, Si is an element effective for improving the hardenability of steel and improving the temper softening resistance of martensite. In order to obtain the above-described effects, the Si content is preferably set to be greater than 0%, or 0.01% or more. The lower limit of the Si content may be set to 0.05% or 0.15%.

However, when the Si content becomes 0.30% or more, the increase in hardness of the steel (rolled material) after hot rolling and before cold forging becomes large, and therefore the life of the die for cold forging is reduced. Therefore, the Si content is set to less than 0.30%. The preferred upper limit of the Si content is 0.27%, 0.25%, or 0.20%.

[Cr：0～1.50％]

As described above, in the steel according to the present embodiment, the lower limit of the Cr content is 0%. However, Cr is an element effective for improving the hardenability of steel and improving the temper softening resistance of martensite. In order to obtain the above-described effects, the Cr content is preferably set to more than 0%, or 0.01% or more. The lower limit of the Cr content may be set to 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, and therefore the life of the die for cold forging is significantly reduced. Therefore, the upper limit of the Cr content is set to 1.50%. The upper limit of the preferable Cr content is 1.20%, 1.00%, or 0.80%.

[Al：0～0.050％]

Al is an element effective for deoxidizing steel, but may not necessarily be contained when deoxidizing with other elements (Si, Ti, etc.). 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, if the Al content exceeds 0.050%, coarse inclusions are formed, and problems such as a decrease in toughness of the steel become significant. Therefore, even when Al is contained, the upper limit of the Al content is set to 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 1 or more selected from Mo, Cu, Ni, and Nb in a range described later as necessary. However, since Mo, Cu, Ni, and Nb are not essential, the lower limit of the content of each of Mo, Cu, Ni, and Nb is 0%.

[Mo：0～0.20％]

As described above, in the steel according to the present embodiment, the lower limit of the Mo content is 0%. However, Mo is an element that contributes to improvement in hardenability of steel even if the content thereof is small. In order to obtain the above-described effects, the Mo content is preferably 0.02% or more. The lower limit of the Mo content is more preferably 0.03%, 0.04%, or 0.05%.

On the other hand, since Mo is a high-priced alloy element, when 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. The upper limit of the Mo content is preferably 0.16%, 0.13%, or 0.10%.

[Cu：0～0.20％]

As described above, in the steel according to the present embodiment, the lower limit of the Cu content is 0%. However, Cu is an element that improves the corrosion resistance of steel. In order to obtain the above-described effects, the Cu content is preferably 0.02% or more. The lower limit of the Cu content is more preferably 0.05%.

On the other hand, if the Cu content exceeds 0.20%, the hot ductility of the steel decreases, and the problem of deterioration of the manufacturability during continuous casting becomes significant. Therefore, even when Cu is contained, the Cu content is set to 0.20% or less. The upper limit of the Cu content is preferably 0.15%, 0.10%, or 0.08%.

[Ni：0～0.20％]

As described above, in the steel according to the present embodiment, the lower limit of the Ni content is 0%. However, Ni is an element that improves the corrosion resistance of steel and is also an element effective for improving the toughness of steel. In order to obtain the above-described effects, the Ni content is preferably 0.02% or more. The lower limit of the Ni content is more preferably 0.03%, 0.04%, or 0.05%.

On the other hand, since Ni is an alloy element at a high price, 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. The upper limit of the Ni content is preferably 0.15%, 0.12%, 0.10%, or 0.08%.

[Nb：0～0.030％]

As described above, in the steel according to the present embodiment, the lower limit of the Nb content is 0%. However, Nb has the following effects: the compound is present in the steel as Nb-based inclusions such as NbC or TiNb (CN) with C in the steel, and acts as pinning particles to suppress abnormal grain growth of austenite grains during quenching and heating. In order to obtain the above-described effects, the Nb content is preferably 0.002% or more. The lower limit of the Nb content is more preferably 0.003%, 0.005%, or 0.006%.

On the other hand, if the Nb content exceeds 0.030%, not only the effect is saturated, but also the Nb-based inclusions are precipitation-strengthened, thereby impairing the manufacturability at the time of continuous casting. Alternatively, in this case, since the Nb-based inclusions cause precipitation strengthening, the hardness of the rolled material after hot rolling is too high. Therefore, if the Nb content exceeds 0.030%, problems such as a reduction in manufacturability and a significant reduction in the life of the die for cold forging become significant. Therefore, even when Nb is contained, the Nb content is set to 0.030% or less. The upper limit of the Nb content is preferably 0.015%, 0.013%, or 0.010%.

The steel according to the present embodiment contains the above alloy components, and the balance (the remainder) of the chemical components thereof includes Fe and impurities. In the present embodiment, the impurities are components mixed by raw materials such as ores and scraps or other factors in the industrial production of steel materials, and are components in such an amount that does not impair the operational effects of the steel according to the present embodiment.

[ N fixed index I_FN: preferably 0 or more]

In order to obtain the effect of the above-described B content, it is necessary to suppress the generation of BN by reducing N (solid solution N) dissolved in the steel. Therefore, it is desired to reduce the amount of solid solution N by stably fixing N in the form of TiN by adding Ti to steel while reducing the content of N in the steel. In order to fix N by Ti to obtain the above-mentioned effects, it is preferable to use an N fixation index I defined by the following formula 1_FNIs set to 0 or more. N may also be fixed to index I_FNThe lower limit of (b) is set to 0.0005, 0.0010, 0.0014, or 0.0050. However, the N-fixation index I is not particularly limited_FNAs long as the Ti content and the N content are controlled as described aboveIn the above range, the steel according to the present embodiment is also softened before cold forging, and the generation of coarse crystal grains during quenching can be suppressed.

I_FN＝[Ti]－3.5×[N]… (formula 1)

In formula 1, Ti and N represent the Ti content and N content in the steel in terms of mass%, and these elements are not included and may be represented as 0%.

[ Ti-Nb-based precipitate formation index I_P: preferably 0.0100 or less]

As described above, it is preferable that: the amount of dissolved N is reduced by fixing N as TiN using Ti. However, it is not preferable to contain Ti in an amount exceeding the amount required for fixing TiN. As described above, Ti also bonds with C, S, and the like 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 also has the same action as Ti.

Specifically, fine TiC, Ti (CN), NbC, TiNb (CN), and Ti as precipitates present in steel₂C₂Ti-Nb precipitates such as S have the following effects: the pinning particles suppress abnormal grain growth of austenite grains during quenching heating, thereby suppressing generation of coarse grains. However, when these Ti — Nb precipitate particles are dispersed in a large amount in the structure after hot rolling, there are the following side effects: the hardness of ferrite increases due to precipitation strengthening caused by fine precipitate particles. Therefore, when these Ti — Nb-based precipitate particles are excessively dispersed in a large amount in the steel, the hardness of the rolled material after hot rolling becomes excessively high, and thus the problem of a significant reduction in the life of the die for cold forging and the like becomes significant. Further, as described above, Ti₂C₂S 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 hot rolling, it is desirable that the Ti-Nb precipitate formation index I calculated by the following formula 2_PIs set to 0.0100 or less. The Ti-Nb precipitate formation index I_PIs set to 0.0075 or less, less than 0.0050, 0.0045 or less, 0.0040 or less, or 0.0035 or less. However, the Ti-Nb precipitate formation index I is not particularly limited_PAs long as the Ti content, the Nb content, and the N content are controlled within the above ranges, the steel according to the present embodiment is also softened before cold forging, and the generation of coarse crystal grains during quenching can be suppressed.

I_P＝0.3×[Ti]+0.15×[Nb]－[N]… (formula 2)

In formula 2, [ Ti ], [ N ], and [ Nb ] represent the Ti content, N content, and Nb content in the steel in mass%, and these elements are not included and may be represented as 0%.

Next, a suitable method for producing the steel of the present embodiment will be described.

In order to produce the steel of the present embodiment, the steel having the above chemical components is melted in a converter, and is subjected to a secondary refining step as necessary to produce a cast slab by continuous casting. The cast slab is reheated and cogging-rolled to produce a raw material (billet) for wire rod rolling having a cross section of, for example, 162mm square (length 162mm × width 162 mm). Then, the billet is heated at a temperature of about 1000 to 1280 ℃, and then wire-rolled to form a wire rod shape having a diameter of 6 to 20 mm. Thereafter, the film was wound in a roll shape by a winding device in a hot state, and then cooled to room temperature. Thus, the steel of the present embodiment is obtained.

Further, in the steel according to the present embodiment, since the amount of Ti-based precipitated particles that cause precipitation strengthening is suppressed, in the method for producing steel according to the present embodiment, it is not necessary to lower the hot rolling temperature to increase the load on the hot rolling equipment in order to suppress the hardness of the steel, and defects such as cracks and flaws due to an increase in hardness are less likely to occur in the steel. The steel according to the present embodiment can be suppressed in hardness without annealing after hot rolling. Therefore, the steel according to the present embodiment is excellent in terms of high productivity.

According to the steel of the present embodiment, softening before cold forging and suppression of generation of coarse crystal grains during quenching can be achieved. The steel of the present embodiment is free from cracks during casting and rolling, and is excellent in manufacturability.

The hardness of the steel according to the present embodiment is not particularly limited, since 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 the present embodiment is preferably Hv180 or less, more preferably Hv170 or less, or Hv160 or less. The lower limit of the hardness of the steel according to the present embodiment is not particularly limited, but may be considered to be substantially about Hv130 or about Hv140 in terms of the chemical composition thereof. The steel according to the present embodiment can have a hardness within the above-described appropriate range without annealing after hot rolling. The steel according to the present embodiment is also excellent in machinability.

In addition, in the case where the steel according to the present embodiment is subjected to tempering treatment in which the steel is heated to a temperature of, for example, 840 to 1100 ℃ and held for 30 minutes, then quenched with water or oil, and further heated and held at a temperature of 150 to 450 ℃, the tensile strength can be increased to 800MPa or more. Therefore, the steel according to the present embodiment is suitable as a material for a member requiring high strength. However, when the steel according to the present embodiment is used as a steel for quenching, the heat treatment conditions are not particularly limited and may be appropriately selected depending on the application.

The use of the steel according to the present embodiment is not particularly limited, but the steel is preferably applied to high-strength machine parts, particularly high-strength bolts, which are manufactured by cold forging and quenching. When the steel according to the present embodiment having high cold forgeability is used as a material for a high-strength machine component, the wear of the die during cold forging can be suppressed, and the life of the die can be increased. In addition, since the cost of the expensive die can be reduced, it is possible to contribute to a reduction in the manufacturing cost of a high-strength bolt having a tensile strength of 800MPa or more in particular.

Examples

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

Firstly, smelting a furnace with a furnace table1-1 and tables 1-2, and further continuously casting the steel to obtain a cast slab. In tables 1-1 and 1-2, the contents of the elements whose contents are not more than the impurity level are indicated by blanks, and the N fixation index I is calculated_FNAnd an index I of formation of Ti-Nb-based precipitates_PWhen it is "0 mass%". In tables 1-1 and 1-2, values outside the range specified in the present invention are underlined. It was confirmed whether or not surface cracking occurred in the thus-obtained cast slab. When the crack on the surface of the cast slab was confirmed, the surface of the cast slab was observed after removing the scale on the surface of the cast slab by notch inspection (check scarf), and the crack depth was examined. When a crack having a depth of 1mm or more is detected on the surface of the cast slab, it is judged that there is a crack on the surface of the cast slab during continuous casting, and it is judged as "defective" with respect to the productivity. The results of the manufacturability evaluation are shown in tables 2-1 and 2-2.

The cast slab was subjected to soaking diffusion treatment and cogging rolling as required to obtain a material (billet) for wire rod rolling having a cross section of 162mm square (length 162mm × width 162 mm). Then, the billet is heated at a temperature of about 1000 to 1280 ℃, and then wire rod rolling is performed, thereby producing a wire rod (steel for springs) having a diameter of 10 mm.

Samples for vickers hardness measurement were cut from the rolled wire rods. Specifically, a sample having a cross section including the central axis of the wire rod is cut in a direction parallel to the rolling direction. After the cut cross section was polished, the vickers hardness of 1/4 depth parts (1/4 parts) of the wire diameter from the surface of the wire was measured. The test load was 10kgf, and the average value obtained by measuring 4 points was shown in tables 2-1 and 2-2 as "hardness after rolling", and this was used as an index for predicting the life of the cold forging die. The samples having hardness of the rolled material exceeding HV180 did not exhibit sufficient effect of improving the life of the cold forging die, and therefore the cold forgeability was judged as "failed". The results of the cold forgeability evaluation are shown in tables 2-1 and 2-2.

In order to simulate the influence of wire drawing and cold forging (cold working) when a wire rod is processed into a bolt shape, the wire rod is subjected to cold drawing with a reduction of area of 70%, then heated at 840 to 1100 ℃ for 30 minutes, and water-cooled quenched to freeze an austenite structure as a prior austenite grain boundary of a martensite structure. Then, the quenched sample was tempered in a temperature range of a point a1 or less as necessary, and a sample having a cross section including the center of the drawn material was cut in a direction parallel to the rolling and drawing directions. The cut sample was polished in cross section, and the prior austenite grain boundaries were developed by corrosion, and the prior austenite grain sizes after quenching and tempering were measured by observation with an optical microscope. The prior austenite grain size was measured in accordance with JIS G0551. The field of view of measurement was 10 fields or more at a magnification of 400, and it was determined that coarse grains were generated even when 1 large grain having an austenite grain size of no more than 5 existed in the sample. The critical (lowest) heating temperature at which coarse grains are generated, which is clarified by observation and measurement of the prior austenite grain size of the sample heated to various temperatures, is defined as the grain coarsening temperature of the sample and is used as an index of the grain coarsening resistance. The sample having a crystal grain coarsening temperature of 900 ℃ or lower was judged as "failed" because of poor crystal grain coarsening resistance. The results of the measurement of the crystal grain coarsening temperature are shown in Table 2-1 and Table 2-2.

As is clear from tables 2-1 and 2-2, a1 to a32 which are examples of the present invention are excellent in cold forgeability because the hardness of the wire rod after rolling is low and the life of the die for cold forging can be expected to be improved, and also excellent in manufacturability because coarse crystal grains are not generated even when heated at a temperature exceeding 900 ℃ in quenching and heating after cold working and because surface cracks of the cast slab are not generated in continuous casting. Furthermore, all of the inventive examples A1 to A32 after heat treatment for measuring the prior austenite grain size described above have a tensile strength of 800MPa or more.

In contrast, in the case of comparative example, any of the cold forgeability, coarse grain prevention property, and manufacturability is inferior. That is, since the amount of Bi added is too large, B1 to B4 have low hot ductility and poor manufacturability. B is5 to B7 had poor coarse grain prevention properties because Bi was not added or the amount of addition was too small. B8 and B9 have a Ti content too large or a small N content relative to the Ti content, and the Ti-Nb precipitate formation index I_PThe excess amount is large, and therefore the hardness of the rolled wire rod is high, and the cold forgeability is poor.

TABLE 2-1

Tables 2 to 2

Industrial applicability

According to the present invention, it is possible to provide steel that can be softened during cold forging and can suppress the generation of coarse grains during quenching after cold forging. The steel according to the present invention is excellent in manufacturability because it does not cause cracking during casting or rolling and can be produced in a range that does not place a burden on production facilities. By applying the steel according to the present invention to a cold-forged part, the wear of the die during cold forging can be suppressed, and the life of the die can be increased. Further, since the application of the steel according to the present invention to a cold-forged part can reduce the cost of a high-priced die, it can contribute to the reduction of the manufacturing cost of a high-strength bolt having a tensile strength of 800MPa or more in particular. Furthermore, the steel according to the present invention is also excellent in machinability. Therefore, the present invention contributes greatly to the industry.

Claims

1. A steel characterized by comprising, in mass%, as chemical components

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: more than 0.0010% and less than 0.0100%,

P: less than 0.020%,

N: less than 0.0100%,

Si: more than 0% and less than 0.30%,

Cr：0～1.50％、

Al：0～0.050％、

Mo：0～0.20％、

Cu：0～0.20％、

Ni：0～0.20％、

Nb: 0 to 0.030%, and

the balance Fe and impurities, and

n fixed index I defined by the following formula 1_FNIs a compound of a formula (I) or more than 0,

a Ti-Nb precipitate formation index I defined by the following formula 2_PIs less than 0.0100, I_FN＝[Ti]－3.5×[N]… formula 1

I_P＝0.3×[Ti]+0.15×[Nb]－[N]… formula 2

Wherein [ Ti ] is a Ti content in terms of mass%, [ N ] is an N content in terms of mass%, and [ Nb ] is an Nb content in terms of mass%.

2. The steel according to claim 1, wherein the chemical component contains in mass% a chemical component selected from the group consisting of

Si: more than 0.01 percent and less than 0.30 percent,

Cr: 0.01 to 1.50%, and

Al：0.001～0.050％

1 or more than two of them.

3. The steel according to claim 1, wherein the chemical component contains in mass% a chemical component selected from the group consisting of

Mo：0.02～0.20％、

Cu：0.02～0.20％、

Ni: 0.02 to 0.20%, and

Nb：0.002～0.030％

1 or more than two of them.

4. The steel according to claim 2, wherein the chemical component comprises in mass% a chemical component selected from the group consisting of

Mo：0.02～0.20％、

Cu：0.02～0.20％、

Ni: 0.02 to 0.20%, and

Nb：0.002～0.030％

1 or more than two of them.