JP3764627B2 - Case-hardened boron steel for cold forging that does not generate abnormal structure during carburizing and its manufacturing method - Google Patents

Description

【００７０】
熱間加工後の棒鋼から、ＡｌＮの析出量を化学分析により求めた。また、圧延後の棒鋼の組織観察を行い、フェライトの結晶粒度番号、ベイナイトの組織分率を求めた。さらに、圧延後の棒鋼のビッカース硬さを測定した。また、一部の試験片について、圧延方向に平行な断面のフェライトバンドの評点を求めた。
さらに、圧延ままの棒鋼から、据え込み試験片を作成し、冷間加工性の指標として、冷間変形抵抗と限界据え込み率を求めた。冷間変形抵抗は相当歪み１．０における変形抵抗で代表させた。
【００７１】
次に、直径３０ｍｍの棒鋼を削りだし、９４０×４時間の条件で浸炭処理を行い、硬さ分布と組織調査を行った。深さ０．４５ｍｍ位置での硬さと不完全焼入れ組織の有無を求めた。不完全焼入れ組織の分率が５％以下の場合に「不完全焼入れ組織：無し」と判定した。
【００７２】
また、圧延ままの棒鋼から、据え込み試験片を作成し、圧下率５０％の据え込みを行った後、浸炭シミュレーションを行った。浸炭シミュレーションの条件は、９１０℃〜１０１０℃に５時間加熱−水冷である。その後、切断面に研磨−腐食を行い、旧オーステナイト粒径を観察して粗粒発生温度（結晶粒粗大化温度）を求めた。浸炭処理は通常９３０〜９５０℃の温度域で行われるため、粗粒発生温度が９５０℃以下のものは結晶粒粗大化防止特性に劣ると判定した。なお、旧オーステナイト粒度の測定はＪＩＳ Ｇ ０５５１に準じて行い、４００倍で１０視野程度観察し、粒度番号５番以下の粗粒が１つでも存在すれば粗粒発生と判定した。
【００７３】
さらに、直径３０ｍｍの棒鋼を削りだし、直径２２ｍｍへ引き抜きを行った後、９４０℃×４時間の条件で浸炭焼入れを行い、γ粒度を測定した。
【００７４】
これらの調査結果を熱間加工条件とあわせて表２に示す。
【００７５】
【表２】

【００７６】
比較例１９、２０はＪＩＳのＳＣｒ４２０およびＳＣＭ４２０の特性であるが、本発明例の冷間変形抵抗は、比較例１９、２０に比較して顕著に小さく、また限界据え込み率も優れている。また、浸炭後の深さ０．４５ｍｍ位置において、本発明例では、不完全焼入れ組織が実質的になく、同位置の硬さもＨＶ６５０以上と良好である。さらに、本発明例の結晶粒粗大化温度は９７０℃以上であり、通常の上限の浸炭条件である９５０℃では、粗大粒の発生を防止できることが明らかである。
【００７７】
次に、表２において、比較例１０はＳｉの含有量が本願規定の範囲を上回った場合であり、本発明例に比較して、冷間加工性は劣る。
【００７８】
比較例１１はＭｏの含有量が本願規定の範囲を下回った場合であり、浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低い。
【００７９】
比較例１２はＴｉの含有量が本願規定の範囲を下回った場合であり、粗大粒防止特性は劣り、また、焼入れ性が低下して、浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低い。
比較例１３はＢの含有量が本願規定の範囲を下回った場合であり、浸炭材の焼入れ性が不足し、深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低い。
【００８０】
比較例１４はＳの含有量が本願規定の範囲を上回った場合であり、比較例１５はＶの含有量が本願規定の範囲を上回った場合であり、比較例１６はＭｇの含有量が本願規定の範囲を上回った場合であり、いずれも浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低い。
【００８１】
比較例１７はＮの含有量が本願規定の範囲を上回った場合であり、浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低く、また、Ｔｉの析出物が粗大になり、粗大粒防止特性も劣る。
【００８２】
比較例１８はＮｂの含有量が本願規定の範囲を上回った場合であり、粗大粒防止特性は劣るとともに、熱間加工後の硬さが高くなり、冷間加工性が本発明例に比較して劣る。
【００８３】
次に、比較例２１は、熱間圧延加熱温度が本願規定の範囲を下回り、圧延材のＡｌＮの析出量が本願規定の範囲を上回り、熱間加工後の硬さが本願規定の範囲を下回った場合であり、粗大粒発生温度は低い。また、比較例２２は熱間圧延時の仕上げ温度が本願規定の範囲を下回り、熱間加工後の硬さが本願規定の範囲を下回り、熱間圧延後のフェライトの結晶粒度番号が本願規定の範囲を上回った場合であり、浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低く、また粗大粒発生温度は９１０℃と実用上問題のあるレベルである。比較例２３、２４は熱間圧延に引き続く冷却速度が本願規定の範囲を上回り、ベイナイトの組織分率が本願規定の範囲を上回った場合であり、冷間加工性および粗大粒防止特性ともに顕著に劣り、浸炭後の深さ０．４５ｍｍ位置において、不完全焼入れ組織が発生し、同位置の硬さも低い。
【００８４】
（実施例２）
実施例１で製造した鋼水準Ａ〜ＨおよびＴ、Ｕの熱間圧延棒鋼について、球状化焼鈍を行った後、実施例１と同様の方法で冷間加工性調査、９４０℃×４時間浸炭における深さ０．４５ｍｍでの不完全焼入れ層の有無と硬さの調査および結晶粒粗大化特性の調査を行った。これらの調査結果をまとめて表３に示す。
【００８５】
【表３】

【００８６】
比較例３８、３９はＪＩＳのＳＣｒ４２０およびＳＣＭ４２０の球状化焼鈍材の特性であるが、本発明例の冷間加工性は、球状化焼鈍後もＳＣｒ４２０およびＳＣＭ４２０に比較して優れている。また、浸炭後の深さ０．４５ｍｍ位置において、本発明例では、不完全焼入れ組織が実質的になく、同位置の硬さもＨＶ６５０以上と良好である。さらに、本発明例の結晶粒粗大化温度は９９０℃以上であり、本発明鋼は、球状化焼鈍後も、通常の上限の浸炭条件である９５０℃において粗大粒の発生を防止できることが明らかである。
【００８７】
なお、本発明鋼は、冷間鍛造の前にその他の焼鈍工程を経由する場合においても、優れた冷間加工性と不完全焼入れ組織防止特性、粗大粒防止特性を有する。
【００８８】
【発明の効果】
本発明の冷間鍛造用肌焼きボロン鋼とその製造方法を用いれば、冷間鍛造時には冷間加工性に優れ、同時に冷間鍛造工程で製造しても、浸炭時に粗大粒の発生と表面から深さ０．２〜０．７ｍｍに生成する不完全焼入れ組織の生成を防止することができ、これにより、必要な強度特性や寸法・形状の精度を確保することができる。そのため、これまで、粗大粒や不完全焼入れ組織の問題から冷鍛化が困難であった部品の冷鍛化が可能になり、さらに冷鍛後の焼鈍を省略することも可能になり、本発明による産業上の効果は極めて顕著なるものがある。
【図面の簡単な説明】
【図１】縞状組織の程度を数量的に表示する金属組織の写真である。
【図２】添加Ｂ量と９４０℃浸炭材の深さ０．４５ｍｍでの不完全焼入れ組織分率の関係について解析した一例を示す図である。
【図３】Ｓ量と９４０℃浸炭材の深さ０．４５ｍｍでの不完全焼入れ組織分率の関係について解析した一例を示す図である。
【図４】Ｖ量と９４０℃浸炭材の深さ０．４５ｍｍでの不完全焼入れ組織分率の関係について解析した一例を示す図である。
【図５】Ｍｇ量と９４０℃浸炭材の深さ０．４５ｍｍでの不完全焼入れ組織分率の関係について解析した一例を示す図である。
【図６】Ｎ量と９４０℃浸炭材の深さ０．４５ｍｍでの不完全焼入れ組織分率の関係について解析した一例を示す図である。
【図７】熱間加工後の硬さと結晶粒粗大化温度の関係について解析した一例を示す図である。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a case-hardened boron steel for cold forging that does not generate an abnormal structure during carburizing and a method for producing the same.
[0002]
[Prior art]
Gears, shafts, and CVJ parts are usually made of medium carbon alloy steel for machine structural use as defined in JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc., and cold forging (also rolling) -It is manufactured in a process of carburizing and quenching after being processed into a predetermined shape by cutting. Cold forging has good surface texture and dimensional accuracy of the product, has a lower manufacturing cost than hot forging, and has a good yield, so parts that were conventionally manufactured by hot forging are now cold forged. The tendency to switch is increasing, and the number of carburized parts manufactured in the cold forging-carburizing process has increased significantly in recent years. Here, when switching from hot forging to cold forging, it is important to reduce the cold deformation resistance of steel and to improve the critical compression ratio. This is because the former is for ensuring the life of the forging tool, and the latter is for preventing cracking of the steel during cold forging. There is boron steel as a steel material suitable for such cold forging. However, since boron steel generates two types of abnormal structures as described below at the time of carburizing, it is rarely used for carburizing. There are two types of abnormal structures generated during carburizing in boron steel: (1) coarse grains and (2) incompletely quenched structures formed to a depth of about 0.2 to 0.7 mm from the surface. In parts where coarse particles are generated, heat treatment distortion occurs after carburizing and quenching. For example, in gears and shaft parts, if this carburizing distortion is large, noise and vibration are caused. Moreover, although the incompletely hardened structure | tissue produced | generated to the depth of 0.2-0.7 mm normally exhibits a black structure, when this incompletely hardened structure | tissue produces | generates, it will cause a strength fall. There has never been proposed a technique for preventing the latter incompletely hardened structure with respect to case-hardened boron steel. Several techniques for preventing the generation of coarse grains during carburizing heating of case-hardened boron steel have been proposed.
[0003]
For example, Japanese Patent Application Laid-Open No. Sho 61-217553 aims to produce TiC by pinning the grain boundaries by setting the amounts of Ti and N to 0.02 <Ti-3.42N. However, the ability of the steel to suppress coarse grains is unstable, and the reality is that the generation of coarse grains during carburization cannot be suppressed depending on the steel manufacturing process. Moreover, since the steel adds a large amount of Ti with respect to the N amount, a large amount of TiC is generated, and therefore, cracks and scratches are likely to occur during the manufacture of the steel material, and it is hard and cold worked in the state of the material. It has disadvantages such as poor performance. Further, there is no mention of a technique for preventing the incompletely quenched structure generated at a depth of 0.2 to 0.7 mm.
[0004]
Japanese Patent Laid-Open No. 63-103052 discloses a case-hardening steel for cold forging with a reduced amount of Si and Mn and containing N amount: 0.008% or less and Nb: 0.01-0.20. Has been. However, the steel also has an unstable ability to suppress coarse grains, and depending on the manufacturing process of the steel material, the generation of coarse grains may or may not be possible. The reality is that it cannot be suppressed. Further, as apparent from the examples, the steel has N content in the range of 0.005 to 0.008 except for one steel type. Even at this level of N content, as described later, the grain coarsening characteristics It has an adverse effect. In addition, one steel type of the embodiment of the present invention has a low N content of 0.002%, but Nb is added in a large amount of 0.05%, so that a large amount of NbC is produced, and therefore the state of the material It is hard and cold workability is not good. Further, there is no mention of a technique for preventing the incompletely quenched structure generated at a depth of 0.2 to 0.7 mm.
[0005]
Japanese Patent Laid-Open No. 9-78184 discloses that in steel containing a specific amount of Al, Nb, N, Nb precipitates present in the material after hot rolling or hot forging, or a composite composition of Nb and Al. The number of precipitates is 5 / 10μm²A case-hardened steel excellent in cold workability and coarsening characteristics of crystal grains as described above is shown. However, the steel also has an unstable ability to suppress coarse grains, and the reality is that the generation of coarse grains during carburization cannot be suppressed depending on the steel manufacturing process. As apparent from the examples, the steel has a high N content in the range of 0.006 to 0.010. As described later, a large amount of N is added to the grain coarsening characteristics. This is considered to have an adverse effect. In addition, it is said that precipitates composed of a composite composition of Nb and Al exist in the material after hot rolling or hot forging. This is because the precipitates of Al are not dissolved during heating in hot rolling or hot forging. This is considered to be one of the reasons why the ability to suppress coarse grains is unstable. Further, there is no mention of a technique for preventing the incompletely quenched structure generated at a depth of 0.2 to 0.7 mm.
[0006]
[Problems to be solved by the invention]
In the disclosed method as described above, the generation of coarse particles cannot be stably suppressed in the carburizing and quenching process, and the cold workability is partially insufficient. Furthermore, the technique for preventing the incompletely hardened structure | tissue produced to the depth of 0.2-0.7 mm from the surface is not solved at all. The present invention solves such problems and is excellent in cold workability, and can prevent generation of coarse grains during carburization and generation of an incompletely quenched structure formed to a depth of 0.2 to 0.7 mm from the surface. The present invention provides a case-hardened steel and a manufacturing method thereof.
[0007]
[Means for Solving the Problems]
In the case-hardened boron steel, the present inventors have earnestly investigated the technology for preventing the generation of coarse grains during carburizing and preventing the formation of an incompletely quenched structure formed to a depth of 0.2 to 0.7 mm from the surface. The point was clarified.
[0008]
(1) It was clarified that the incompletely quenched structure generated at a depth of 0.2 to 0.7 mm from the surface is a kind of lower bainite-like structure, which is generated by the following mechanism. During carburizing, nitrogen penetrates from the surface simultaneously with carbon, and austenite is refined in the hardened layer. Therefore, since BN is generated, the solid solution B is reduced, and at the same time, the austenite grains are refined, so that the contribution to the hardenability of B is reduced in the hardened layer. The amount of C is about 0.8% at the outermost surface, while the amount of C is about 0.4 to 0.6% at a depth of 0.2 to 0.7 mm from the surface. Since the amount is low, the hardenability is lowered at a depth of 0.2 to 0.7 mm from the surface, and the cooling rate is slow. Therefore, an incompletely hardened structure which is a kind of lower bainite-like structure is generated at a depth of 0.2 to 0.7 mm from the surface. The transformation nucleus of the incompletely quenched structure transformation at this time is various oxides or a composite precipitate of MnS and V (CN). As the oxide, Mg-based oxides are particularly harmful.
[0009]
(2) The present inventors have newly found that the following measures are effective in order to prevent the formation of an incompletely quenched structure generated at a depth of 0.2 to 0.7 mm.
[0010]
(1) Oxygen, Mg, S, N, and V are reduced, and various oxides that are transformation nuclei of incompletely quenched structure transformation, composite precipitates of MnS and V (CN) are reduced.
[0011]
(2) The appropriate addition amount of B for improving hardenability is usually around 0.001% in general boron steel, but in the steel of the present invention, it compensates for the decrease in solid solution boron accompanying BN formation in the hardened layer. Therefore, a large amount is added to 0.002% or more. Further, a large amount of Ti is added to fix the invading N.
[0012]
(3) At a depth of 0.2 to 0.7 mm from the surface, in order to compensate for the decrease in hardenability due to the decrease in solid solution boron, it is difficult to generate oxides (becomes transformation nuclei of the above incompletely quenched structure). Difficult) Add Cr and Mo as essential elements. Mo has the effect of substituting the grain boundary segregation of Mo for the reduction of the grain boundary segregation of boron due to the decrease in solid solution boron, so the addition of a small amount of Mo is particularly effective.
[0013]
(4) The finer the austenite grains, the more pronounced the effect of the decrease in the amount of dissolved boron on the hardenability. Therefore, it is important to prevent the austenite grains from becoming too fine. It is necessary that the previous structure does not substantially contain a bainite structure and that the ferrite structure has a grain size number of 11 or less.
[0014]
(3) Next, in order to prevent coarse grains, which is another problem in preventing abnormal carburizing structure, Ti carbonitrides mainly composed of TiC instead of TiN, or Nb carbon mainly composed of NbC It is effective to finely precipitate nitride during carburizing.
[0015]
(4) The following method was newly discovered as a method for finely depositing Ti carbonitride mainly composed of TiC or Nb carbonitride mainly composed of NbC during carburizing.
[0016]
(1) In order to use Ti carbonitrides mainly composed of TiC or Nb carbonitrides mainly composed of NbC as pinning particles, it is necessary to finely disperse these precipitates in a large amount during carburizing and quenching. There is. For this purpose, TiC and NbC precipitates need to be once solutionized at the time of rolling and heating when hot working a steel bar or wire. If the amount of N is high and a large amount of TiN remains during rolling and heating, NbC and TiC form a composite precipitate with TiN, making it difficult to form a solution. In the cooling process after hot rolling, NbC and TiC are deposited on coarse TiN, and fine dispersion of NbC and TiC is hindered. Therefore, it is necessary to reduce the N amount as much as possible. In addition, coarse AlN and Al during rolling and heating₂O_ThreeThe presence of an oxide such as the same has the same adverse effect as TiN described above. Therefore, it is necessary to make AlN into solution during rolling and heating. Here, if AlN is in solution during rolling and heating, precipitation of AlN hardly occurs in the hot rolling-cooling process of steel bars and wires. Therefore, by regulating the precipitation amount of AlN after hot working, it is possible to confirm the solution state of AlN during rolling and heating.
[0017]
{Circle around (2)} If heating is performed under the condition that AlN can be dissolved during rolling and heating, it is possible to temporarily precipitate NbC and TiC precipitates. Therefore, by regulating the precipitation amount of AlN after hot working, it is possible to confirm that the precipitates of NbC and TiC were once formed into a solution during rolling and heating.
[0018]
(3) Further, in order to stably exhibit the pinning effect of Ti carbonitride mainly composed of TiC or Nb carbonitride mainly composed of NbC, it is constant in the matrix after hot working. It is necessary to finely precipitate more TiC or more NbC than the amount. For this purpose, it is necessary to precipitate the phase interface during the diffusion transformation from austenite during the cooling process during hot working. If bainite is generated in the structure as it is hot-worked, it becomes difficult to precipitate the phase interface of TiC and NbC, so it is essential to make the structure substantially free of bainite. Therefore, it is an essential requirement to obtain a predetermined hardness according to the addition amount of Ti or Nb while being hot-worked by precipitation hardening.
[0019]
  4). Furthermore, the generation characteristics of coarse grains depend on the degree of the striped structure called a ferrite band that is recognized in a cross section parallel to the hot rolling direction of the steel material after hot working. Here, the grade of the ferrite band is graded in seven stages of 1 to 7 in “Metal Society of Japan Journal Vol. 34, No. 9, page 961” published by the Japan Institute of Metals in 1970 (FIG. 1). . That is, the above-mentioned Journal of the Japan Institute of Metals, Vol. 34, No. 9, pages 957-962, as the title says, “About the effect of austenite grain size and forging ratio on ferrite stripe structure” is described, In the left column, lines 7 to 8 on page 961, “Photo. 4 standard photo was created in order to quantitatively display the degree of the striped structure” is described, and “Photo. 4 Classifications of ferrite bands (× 50 × 2/3 × 5/6) ”includes 1 to 7 reference photographs. The score indicates that the smaller the score number is, the lighter the ferrite band is, and the higher the score number is, the more prominent the ferrite band is. In order to suppress coarse grains, the score of the ferrite band defined in the above-mentioned Journal of the Japan Institute of Metals, Vol. 34, page 961, of the cross-sectional structure parallel to the hot rolling direction should be 1-5.preferable. This is because when the ferrite band is prominent, such as a ferrite band score of 6 or more, the pearlite structure is continuously connected. This is because coarse particles are generated.
[0020]
(5) As a method for realizing the above, the following points were clarified.
[0021]
(1) In order to limit the precipitation amount of AlN as much as possible in the state of the steel material after hot working, it is necessary to increase the rolling heating temperature.
[0022]
(2) It is necessary to optimize the finishing temperature and cooling conditions after rolling in order to reduce the restriction of the amount of bainite structure of the steel material after hot working and the degree of ferrite band.
[0023]
(3) In order to finely precipitate Ti carbonitride mainly composed of a certain amount or more of TiC or Nb carbonitride mainly composed of NbC in the hot-worked steel material in advance, the rolling heating temperature Can be dispersed in a large amount by heat-treating these precipitates once to form a solution, and gradually cooling the precipitation temperature range of these precipitates after hot rolling.
[0024]
The present invention has been completed based on the above new findings.
[0025]
  That is, the claims of the present invention1, 2In the invention of the present invention,
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.8%
Cr:0.83~ 1.5%,
Mo: 0.005 to 0.3%,
B:0.0031~ 0.006%,
Al: 0.015 to 0.05%,
Ti: 0.01 to 0.1%
Containing
If necessary,
Nb: 0.002 to 0.05%
Containing
P: 0.025% or less (including 0%),
S:0.013% Or less (including 0%),
V: 0.01% or less (including 0%),
Mg: 0.03% or less (including 0%),
N: 0.005% or less (including 0%),
O: 0.002% or less (including 0%)
Limit each to
The balance consists of iron and inevitable impurities,
When the precipitation amount of AlN is limited to 0.01% or less, the ferrite grain size number is 8 to 11, the structure fraction of bainite is 10% or less, and the hardness index H is defined by the following formula: Hardness is HV, H-20 to H + 30It is characterized byIt is to use case-hardened boron steel for cold forging that does not generate an abnormal structure during carburizing.
      H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr%
          + 136.7Mo% + 708Ti% + 599Nb%
[0026]
  Claims of the invention3When manufacturing the above steel, the heating temperature is 1150 ° C. or higher, the hot rolling finishing temperature is 840 to 1000 ° C., and the temperature range of 800 to 500 ° C. is 1 ° C./second or less following the hot rolling. By using a method for producing case-hardened boron steel for cold forging that does not produce an abnormal structure during carburizing, characterized by hot working on wire or steel bar under conditions of slow cooling at a cooling rate ofis there.
[0027]
By using the steel and the manufacturing method of the present invention, boron steel is excellent in cold workability during cold forging, and generates coarse grains and a depth of 0.2 to 0.7 mm from the surface during carburizing. Generation of a completely quenched structure can be prevented.
[0028]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be described in detail.
[0029]
First, the reasons for limiting the components will be described.
[0030]
C is an effective element for imparting the necessary strength to the steel, but if it is less than 0.1%, the necessary tensile strength cannot be secured, and if it exceeds 0.3%, it becomes hard and cold work is performed. The core portion toughness after carburization is deteriorated as well as the property is deteriorated, so it is necessary to be within the range of 0.1 to 0.3%.
[0031]
Si is an element effective for deoxidation of steel and is an element effective for imparting necessary strength and hardenability to steel and improving temper softening resistance. However, if it is less than 0.01%, the effect is ineffective. It is enough. On the other hand, if it exceeds 0.15%, the hardness is increased and the cold forgeability is deteriorated. For the above reasons, the content needs to be in the range of 0.01 to 0.15%. The preferred range is 0.02 to 0.1%.
[0032]
Mn is an element effective for deoxidation of steel and is an element effective for imparting necessary strength and hardenability to the steel, but if less than 0.2%, the effect is insufficient, 0.8% If it exceeds 1, the increase in hardness will be caused and the cold forgeability will deteriorate, so it is necessary to be within the range of 0.2% to 0.8%. The preferred range is 0.3-0.6%.
[0033]
  Cr is an effective element for imparting strength and hardenability to steel. Since oxygen penetrates from the surface during carburizing, the elements added to ensure hardenability are usually oxidized and the contribution to hardenability is diminished, but Cr is less likely to be oxidized compared to Mn, It is effective in preventing the formation of an incompletely quenched structure that is generated at a depth of the hardened layer of 0.2 to 0.7 mm. The effect of improving hardenability is0.83Addition of less than% is insufficient, and addition over 1.5% leads to an increase in hardness and deteriorates cold forgeability. For the above reasons, the content is0.83It must be in the range of ~ 1.5%. The preferred range is0.83~ 1.3%.
[0034]
Mo is also an effective element for imparting strength and hardenability to steel. Mo, like Cr, is an element that is not easily oxidized by oxygen that enters during carburization, and is effective in preventing the formation of an incompletely quenched structure formed to a depth of 0.2 to 0.7 mm from the surface. In particular, Mo has the effect of increasing the hardenability by grain boundary segregation, similar to B, and Mo has the effect of substituting the grain boundary segregation of boron for reducing the grain boundary segregation of boron by reducing the solid solution boron. Even a very small amount of addition is effective in preventing the formation of an incompletely quenched structure at a depth of 0.2 to 0.7 mm. For such effects, addition of less than 0.005% is insufficient, and addition over 0.3% leads to an increase in hardness and deteriorates cold forgeability. For the above reasons, the content needs to be in the range of 0.005 to 0.3%. The preferred range is 0.005 to 0.2%.
[0035]
B is added for the following three points. (1) In steel bar / wire rolling, boron iron carbide is generated in the cooling process after rolling, thereby increasing the growth rate of ferrite and promoting softening during rolling. (2) When carburizing and quenching, impart hardenability to the steel. (3) Improve the fatigue strength and impact strength of carburized parts by improving the grain boundary strength of the carburized material. Here, the effects of (2) and (3) must be present in the form of solute boron in the steel, but the austenite grains become fine because nitrogen penetrates from the surface during carburizing, and BN The solid solution boron is reduced. Therefore, the contribution of boron to the hardenability is reduced, and an incompletely hardened structure is easily generated at a depth of 0.2 to 0.7 mm.
[0036]
  FIG. 2 shows the effect of the amount of added B on the fraction of the incompletely quenched structure at 940 ° C. × 4 hours at a depth of 0.45 mm. Incompletely quenched structure0.0031With more than B additionReduced to nearly 0%. in this way,0.0031If the addition is less than%, the above effect is insufficient. On the other hand, even if a large amount of B is added, the above effect is saturated, and on the contrary, adverse effects such as deterioration of ductility appear. In particular, when it exceeds 0.006%, such an adverse effect becomes remarkable. For the above reasons, the content is0.0031It is necessary to be within the range of ˜0.006%. The preferred range is0.0031-0.004%.
[0037]
Al is added as a deoxidizer. If it is less than 0.015%, the effect is insufficient. On the other hand, if it exceeds 0.05%, AlN remains without solution during rolling and heating, and becomes a precipitation site for precipitates of Ti and Nb, which inhibits fine dispersion of these precipitates, Promotes coarsening. For the above reasons, the content needs to be in the range of 0.015 to 0.05%. The preferred range is 0.025 to 0.04%.
[0038]
Ti combines with N in the steel to produce TiN, and is added for the purpose of preventing precipitation of BN by fixing solid solution N, that is, securing solid solution B. Further, the remaining Ti that has produced TiN is added to produce fine TiC in the steel and thereby refine γ grains during carburization. Ti is added for these two purposes. However, if less than 0.01%, the effect is insufficient. On the other hand, when Ti is added in excess of 0.1%, precipitation hardening due to TiC becomes remarkable, and cold workability is remarkably deteriorated. For the above reasons, the content needs to be in the range of 0.01 to 0.1%. The preferred range is 0.02 to 0.05%.
[0039]
Next, in the present invention, Nb is contained as necessary. Nb forms NbC-based Nb (CN), and is an element effective for suppressing coarsening of crystal grains during carburizing heating. If it is less than 0.002%, the effect is insufficient. On the other hand, if it exceeds 0.05%, the hardness of the material becomes hard and the cold forgeability deteriorates, and it becomes difficult to form a solution during heating of the steel bar and wire rod. For the above reasons, the content needs to be in the range of 0.002 to 0.05%. The preferred range is 0.005 to 0.03%.
[0040]
Since P is an element that increases deformation resistance during cold forging and deteriorates toughness, cold forgeability deteriorates. Further, since the fatigue strength is deteriorated by embrittlement of the grain boundaries of the parts after quenching and tempering, it is desirable to reduce them as much as possible. Therefore, it is necessary to limit the content to 0.025% or less. The preferred range is 0.015% or less.
[0041]
S forms MnS in the steel and becomes a transformation nucleus of an incompletely quenched structure generated to a depth of 0.2 to 0.7 mm during carburizing, and promotes the formation of this abnormal structure.
[0042]
  FIG. 3 shows the influence of the amount of S on the fraction of the incompletely quenched structure at 940 ° C. × 4 hours at a depth of 0.45 mm. Incompletely hardened structure reduces the amount of S0.013% Or lessReduced to nearly 0%. In other words, the adverse effect of S as described above is that the amount of S is0.013%, It becomes particularly prominent.0.013% Or less is required.
[0043]
V forms V (CN) in the steel and, like MnS, becomes a transformation nucleus of an incompletely quenched structure generated at a depth of 0.2 to 0.7 mm during carburizing, and promotes the formation of this abnormal structure.
[0044]
FIG. 4 shows the influence of the amount of V on the incompletely quenched structure fraction at a depth of 0.45 mm of the carburized material at 940 ° C. for 4 hours. The incompletely hardened structure becomes 10% or less when the V amount is 0.01% or less. From this, the adverse effect of V becomes particularly prominent when the V content exceeds 0.01%, so the content needs to be 0.01% or less.
[0045]
Mg forms an oxide in steel and, like MnS, becomes a transformation nucleus of an incompletely quenched structure that is formed at a depth of 0.2 to 0.7 mm during carburizing, and promotes the formation of this abnormal structure. In particular, Mg-based oxides also serve as BN precipitation sites, so Mg promotes the formation of an incompletely quenched structure even in a small amount.
[0046]
FIG. 5 shows the influence of the amount of Mg on the fraction of the incompletely quenched structure at 940 ° C. × 4 hours at a depth of 0.45 mm. The incompletely hardened structure becomes 10% or less when the Mg content is 0.03% or less. From this, the adverse effect of Mg becomes particularly significant when the Mg content exceeds 0.03%, so the content needs to be 0.03% or less. The preferred range is 0.01% or less.
[0047]
It is desirable to limit N as much as possible for the following three reasons. (1) B is added for the purpose of improving hardenability and strengthening grain boundaries as described above. However, since the effect of these B is manifested only in the state of solid solution B in steel, the amount of N is reduced. Therefore, it is essential to suppress the generation of BN. BN also has an adverse effect of promoting the formation of an incompletely quenched structure at a depth of 0.2 to 0.7 mm. (2) When V is present, V (CN) is formed in the steel, and as described above, the formation of an incompletely quenched structure that is generated at a depth of 0.2 to 0.7 mm during carburization is promoted. Since this phenomenon becomes more conspicuous as the amount is higher, it is essential to reduce the amount of N in order to prevent this abnormal tissue. (3) N, when combined with Ti in the steel, produces coarse TiN that hardly contributes to grain control, which becomes the precipitation site of Ti precipitates mainly composed of TiC, and the carburization of these Ti carbonitrides. Inhibits fine precipitation at the time and promotes the formation of coarse particles.
[0048]
FIG. 6 shows the influence of the amount of N on the fraction of the incompletely quenched structure at 940 ° C. × 4 hours at a depth of 0.45 mm. The incompletely hardened structure becomes 10% or less by making the N amount 0.005% or less. As described above, the adverse effect of N as described above becomes particularly remarkable when the amount of N exceeds 0.005%. For the above reasons, the content needs to be 0.005% or less.
[0049]
O forms Al-based or Mg-based oxide inclusions in the steel. When a large amount of oxide inclusions are present in the steel, it becomes a transformation nucleus of an incompletely quenched structure generated to a depth of 0.2 to 0.7 mm during carburizing, and promotes the formation of this abnormal structure. Moreover, it becomes a TiC precipitation site, and TiC precipitates coarsely during hot working, and the coarsening of crystal grains cannot be suppressed during carburizing. Such an adverse effect is particularly noticeable when the amount of O exceeds 0.002%, so the amount of O needs to be 0.002% or less. The preferred range is 0.0015% or less.
[0050]
Next, in the present invention, the precipitation amount of AlN after hot working is limited to 0.01% or less. The reason for this limitation will be described below. When coarse AlN is present during rolling and heating, it becomes a precipitation site for Ti precipitates mainly composed of TiC, and the Ti precipitates are coarsely deposited after hot working, making it impossible to suppress the coarsening of crystal grains during carburizing. Therefore, it is necessary to solutionize AlN during rolling and heating. Here, if AlN is in solution at the time of rolling and heating, precipitation of AlN hardly occurs in the hot rolling-cooling process of steel bars and wires. Therefore, it is possible to confirm that AlN has been sufficiently formed into a solution during rolling and heating by regulating the precipitation amount of AlN after hot working. In order to utilize Ti precipitates as pinning particles, TiC precipitates need to be once in solution during rolling and heating. If heating is performed under conditions that allow AlN to form a solution during rolling and heating, the TiC precipitate can be once formed into a solution. Therefore, by regulating the precipitation amount of AlN after hot working, it is possible to confirm that the TiC precipitates were once in solution during rolling and heating. Similarly to TiC, it is possible to confirm that NbC precipitates were once formed into solution during rolling and heating by regulating the precipitation amount of AlN after hot working. When the precipitation amount of AlN exceeds 0.01%, the above effect is insufficient, and there is a concern about the generation of coarse particles practically. For the above reasons, the precipitation amount of AlN after hot working is limited to 0.01% or less. The preferred range is 0.005% or less.
[0051]
Next, in the present invention, the ferrite crystal grain size number after hot working is set to 8 to 11, and the reason for this limitation will be described below. If the ferrite grains after hot working are excessively fine, the austenite grains are excessively refined during carburizing. The finer the austenite grain, the more pronounced the effect of the decrease in solid solution boron due to BN formation on the hardenability, and the incompletely hardened structure is more likely to be formed at a depth of 0.2 to 0.7 mm, and the hardened layer is quenched Low. If the ferrite crystal grain size exceeds 11, generation of an incompletely hardened structure at a depth of 0.2 to 0.7 mm becomes remarkable, and there is a concern about deterioration of strength as a carburized material. In addition, when the ferrite crystal grain size number after hot working is coarser than less than 8, the ductility of the hot worked material is deteriorated and the cold forgeability is deteriorated. For the above reasons, the ferrite crystal grain size number after hot working needs to be in the range of 8-11.
[0052]
Next, in the present invention, the structure fraction of bainite after hot working is limited to 10% or less. The reason for this limitation will be described below. When a bainite structure is mixed in the steel material after hot working, austenite grains are excessively refined during carburizing. As described above, as the austenite grains are finer, the effect of the decrease in solid solution boron due to the generation of BN on the hardenability becomes more prominent, and an incompletely hardened structure is easily generated at a depth of 0.2 to 0.7 mm. Furthermore, when a bainite structure is mixed in the steel material after hot working, it causes coarse grains during carburizing heating. In addition, suppression of bainite contamination is also desirable from the viewpoint of improving cold workability. These adverse effects become particularly significant when the bainite structural fraction exceeds 10%. For the above reasons, it is necessary to limit the structural fraction of bainite after hot working to 10% or less. The preferred range is 5% or less.
[0053]
Next, in the present invention, when the hardness index H is defined by the following formula, the hardness of the hot-worked material is limited to the range of H-20 to H + 30 in HV. The reason for this limitation will be described below.
[0054]
H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr% + 136.7Mo% + 708Ti% + 599Nb%
(However, when Nb is not contained, Nb steel is 0.)
[0055]
The present invention is characterized in that TiC-based Ti carbonitride or further NbC-based Nb carbonitride is finely dispersed during carburizing in order to prevent coarse grains during carburizing. In order to utilize Ti or Nb for this purpose, it is necessary to precipitate these precipitates at the phase interface during the ferrite transformation from austenite in the cooling process after hot working. In order to deposit these precipitates at the phase interface, it is essential to limit the bainite transformation in the cooling process after hot working as described above. In the state where bainite is not formed, if Ti and Nb precipitates are precipitated at the phase interface, the hardness increases due to precipitation hardening, but for the above reasons, the lower limit value of the hardness of the steel material is limited according to the amount of Ti and Nb. This makes it possible to finely disperse Ti and Nb precipitates at the time of carburizing and to prevent coarse grains. From the above technical idea, the hardness index determined by the component system was introduced to define the lower limit of the hardness of the hot-worked material. The hardness index H is an index that formulates the influence of the alloy component on the hardness of the hot-worked material, and its unit is HV. As preconditions for defining the hardness index H, it is assumed that the hot-worked material does not substantially contain a bainite structure, and that Ti and Nb are all added to contribute to precipitation strengthening.
[0056]
FIG. 7 shows the relationship between the hardness of the hot-worked material manufactured under various manufacturing conditions and the coarse grain generation temperature. The hardness index H of this steel material is 154. The coarse grain generation temperature was determined by carrying out a carburization simulation by holding at a reduction rate of 50% and holding at each temperature for 5 hours. When the hardness of the hot-worked material is HV and less than H-20, the crystal grain coarsening temperature is lowered. On the other hand, when the hardness of the hot-worked material is increased, the cold workability is deteriorated, but the influence becomes particularly remarkable when the hardness exceeds H + 30. For the above reasons, the hardness of the hot-worked material is limited to the range of H-20 to H + 30 by HV. A preferable range is a range of H-20 to H + 20.
[0057]
In addition, the hardness (HV) prescribed | regulated by this invention is the hardness of the outermost layer except the surface decarburization layer of a hot work material.
[0060]
Next, hot working conditions will be described.
[0061]
The steel composed of the above-described components of the present invention is melted by a usual method such as a converter, an electric furnace, etc., and the components are adjusted. A rolling material to be rolled is used.
[0062]
Next, claim 4 of the present invention is such that the heating temperature is 1150 ° C. or higher, the finishing temperature of hot rolling is 840 to 1000 ° C., and the temperature range of 800 to 500 ° C. is 1 ° C./second or less following hot rolling. Hot working on wire or bar under conditions of slow cooling at the cooling rate.
[0063]
First, the heating temperature is set to 1150 ° C. or higher for the following reason. When the heating temperature is less than 1150 ° C., AlN and TiC or further NbC cannot be once dissolved in the matrix during heating, and the amount of precipitates contributing to fine TiC or further NbC grain control after hot working And the generation of coarse grains during carburization cannot be suppressed. Therefore, it is necessary to heat at a temperature of 1150 ° C. or higher during hot rolling.
[0064]
Next, the finishing temperature of hot rolling is set to 840 to 1000 ° C. for the following reason. When the finishing temperature is less than 840 ° C., the ferrite grain size becomes as fine as 11 or more, and the ferrite band becomes more prominent as the rating exceeds 5. The coarse grains and the depth of 0.2 to 0.7 mm are formed during the subsequent carburizing. Incompletely hardened structure tends to occur. On the other hand, if the finishing temperature exceeds 1000 ° C., the ferrite crystal grains become coarse and the cold forgeability deteriorates. For these reasons, the hot rolling finishing temperature is set to 840 to 1000 ° C. The preferred range is 850-960 ° C.
[0065]
Next, following the hot rolling, the temperature range of 800 to 500 ° C. is gradually cooled at a cooling rate of 1 ° C./second or less for the following reason. When the cooling rate exceeds 1 ° C./s, the structure fraction of bainite increases, the precipitation amount of fine TiC and other precipitates after hot working becomes insufficient, and coarse grains tend to be generated during carburizing. Furthermore, when the structure fraction of bainite is increased, the hardness of the rolled material is significantly increased and the cold forgeability is deteriorated. Therefore, the cooling rate is limited to 1 ° C./second or less. The preferred range is 0.7 ° C./s or less. In addition, as a method of reducing the cooling rate, a method of installing a heat insulating cover or a heat insulating cover with a heat source behind the rolling line and thereby performing slow cooling can be mentioned.
[0066]
In the present invention, the size of the slab, the cooling rate during solidification, and the ingot rolling conditions are not particularly limited, and any conditions may be used as long as the requirements of the present invention are satisfied. In addition, the steel of the present invention is not only used in the process of forming as-rolled steel bars into parts by cold forging, but also during the cold forging process when passing through an annealing process or warm / hot forging before cold forging. In the case of including an annealing process, it can also be applied to a case where the part is formed in a warm / hot forging process and a part is formed in a cutting process.
[0067]
【Example】
Hereinafter, the effects of the present invention will be described more specifically by way of examples.
[0068]
Example 1
Converter molten steel having the composition shown in Table 1 was continuously cast, and a rolling raw material of 162 mm square was obtained through a batch rolling process as necessary. Subsequently, a steel bar having a diameter of 34 mm was manufactured by hot working. The comparative steels T and U are JIS SCr420 and SCM420.
[0069]
[Table 1]

[0070]
From the steel bar after hot working, the precipitation amount of AlN was determined by chemical analysis. Moreover, the structure of the steel bar after rolling was observed, and the grain size number of ferrite and the structure fraction of bainite were obtained. Furthermore, the Vickers hardness of the steel bar after rolling was measured. Moreover, the score of the ferrite band of the cross section parallel to a rolling direction was calculated | required about some test pieces.
Furthermore, an upsetting test piece was prepared from the rolled steel bar, and the cold deformation resistance and the limit upsetting rate were obtained as indicators of cold workability. The cold deformation resistance was represented by the deformation resistance at an equivalent strain of 1.0.
[0071]
Next, a steel bar having a diameter of 30 mm was cut out, carburized under conditions of 940 × 4 hours, and the hardness distribution and the structure were examined. The hardness at a depth of 0.45 mm and the presence or absence of an incompletely quenched structure were determined. When the fraction of the incompletely quenched structure was 5% or less, it was determined that “incompletely quenched structure: none”.
[0072]
In addition, an upsetting test piece was prepared from the rolled steel bar, and after upsetting at a reduction rate of 50%, carburization simulation was performed. The conditions for the carburizing simulation are heating to 910 ° C. to 1010 ° C. for 5 hours and water cooling. Thereafter, the cut surface was polished and corroded, and the prior austenite grain size was observed to determine the coarse grain generation temperature (crystal grain coarsening temperature). Since the carburizing process is normally performed in a temperature range of 930 to 950 ° C., it was determined that a coarse grain generation temperature of 950 ° C. or lower is inferior in crystal grain coarsening prevention characteristics. The prior austenite particle size was measured in accordance with JIS G 0551, observed at 400 times for about 10 fields of view, and if there was at least one coarse particle having a particle size number of 5 or less, it was determined that coarse particles were generated.
[0073]
Further, a steel bar having a diameter of 30 mm was cut out and drawn to a diameter of 22 mm, followed by carburizing and quenching under conditions of 940 ° C. × 4 hours, and the γ particle size was measured.
[0074]
These investigation results are shown in Table 2 together with hot working conditions.
[0075]
[Table 2]

[0076]
Comparative Examples 19 and 20 are the characteristics of JIS SCr420 and SCM420, but the cold deformation resistance of the inventive example is significantly smaller than that of Comparative Examples 19 and 20, and the limit upsetting rate is also excellent. Moreover, in the depth 0.45 mm position after carburizing, in the example of the present invention, there is substantially no incompletely quenched structure, and the hardness at the same position is also good at HV650 or more. Furthermore, the crystal grain coarsening temperature of the present invention example is 970 ° C. or higher, and it is clear that the generation of coarse grains can be prevented at 950 ° C., which is a normal upper limit carburizing condition.
[0077]
Next, in Table 2, Comparative Example 10 is a case where the Si content exceeds the range specified in the present application, and the cold workability is inferior compared to the inventive example.
[0078]
In Comparative Example 11, the Mo content falls below the range specified in the present application, and an incompletely hardened structure is generated at a depth of 0.45 mm after carburization, and the hardness at the same position is also low.
[0079]
Comparative Example 12 is a case where the Ti content falls below the range specified in the present application, the coarse grain prevention characteristics are inferior, and the hardenability is lowered, and incomplete quenching at a depth of 0.45 mm after carburizing. A tissue is generated and the hardness at the same position is low.
Comparative Example 13 is a case where the content of B is below the range specified in the present application, the hardenability of the carburized material is insufficient, an incompletely hardened structure is generated at a depth of 0.45 mm, and the hardness at the same position is also Low.
[0080]
Comparative Example 14 is the case where the S content exceeds the range specified in the present application, Comparative Example 15 is the case where the V content exceeds the range specified in the present application, and Comparative Example 16 has the Mg content of the present application. It is a case where it exceeds the specified range, and in any case, an incompletely quenched structure is generated at a depth of 0.45 mm after carburization, and the hardness at the same position is low.
[0081]
Comparative Example 17 is a case where the N content exceeds the range specified in the present application. At a depth of 0.45 mm after carburization, an incompletely quenched structure is generated, the hardness at the same position is low, and Ti Precipitates become coarse and coarse grain prevention properties are also poor.
[0082]
Comparative Example 18 is a case where the Nb content exceeds the range specified in the present application, and the coarse grain prevention characteristics are inferior, the hardness after hot working is increased, and the cold workability is compared with the inventive example. Inferior.
[0083]
  Next, in Comparative Example 21, the hot rolling heating temperature falls below the range specified in the present application, the precipitation amount of AlN in the rolled material exceeds the range specified in the present application, and the hardness after hot working falls below the range specified in the present application. In this case, the coarse particle generation temperature is low. Further, in Comparative Example 22, the finishing temperature at the time of hot rolling is lower than the range specified in the present application, the hardness after hot working is lower than the range specified in the present application,The grain size number of ferrite after hot rolling isThis is a case that exceeds the range specified in this application. At a depth of 0.45 mm after carburization, an incompletely quenched structure is generated, the hardness at the same position is low, and the coarse grain generation temperature is 910 ° C., which is a practical problem. It is a certain level. Comparative Examples 23 and 24 are cases in which the cooling rate following hot rolling exceeds the range specified in the present application, and the bainite structural fraction exceeds the range specified in the present application, and both the cold workability and the coarse grain prevention characteristics are remarkable. Inferior, an incompletely quenched structure occurs at a depth of 0.45 mm after carburization, and the hardness at that position is also low.
[0084]
(Example 2)
About hot-rolled steel bars of steel levels A to H, T, and U manufactured in Example 1, after spheroidizing annealing, a cold workability investigation and 940 ° C. × 4 hours carburizing in the same manner as in Example 1 The presence or absence of an incompletely hardened layer at a depth of 0.45 mm and the hardness and the grain coarsening characteristics were investigated. These survey results are summarized in Table 3.
[0085]
[Table 3]

[0086]
Comparative Examples 38 and 39 are characteristics of spheroidized annealing materials of JIS SCr420 and SCM420, but the cold workability of the inventive example is superior to SCr420 and SCM420 even after spheroidizing annealing. Moreover, in the depth 0.45 mm position after carburizing, in the example of the present invention, there is substantially no incompletely quenched structure, and the hardness at the same position is also good at HV650 or more. Furthermore, the crystal grain coarsening temperature of the present invention example is 990 ° C. or higher, and it is clear that the steel of the present invention can prevent the generation of coarse grains at 950 ° C., which is the normal upper limit carburizing condition, even after spheroidizing annealing. is there.
[0087]
The steel of the present invention has excellent cold workability, incomplete quenching structure preventing properties, and coarse grain preventing properties even when passing through other annealing steps before cold forging.
[0088]
【The invention's effect】
If the case-hardened boron steel for cold forging of the present invention and its manufacturing method are used, it is excellent in cold workability at the time of cold forging. Generation of an incompletely hardened structure generated at a depth of 0.2 to 0.7 mm can be prevented, and thereby necessary strength characteristics and accuracy of dimensions and shapes can be ensured. Therefore, it becomes possible to cold forge parts that have been difficult to cold forge due to the problem of coarse grains and incompletely quenched structure, and further, it is possible to omit annealing after cold forging, the present invention The industrial effect of is extremely remarkable.
[Brief description of the drawings]
FIG. 1 is a photograph of a metal structure that quantitatively displays the extent of a striped structure.
FIG. 2 is a diagram showing an example of analyzing the relationship between the amount of added B and the incompletely quenched structure fraction at a depth of 940 ° C. carburized material of 0.45 mm.
FIG. 3 is a diagram showing an example of analyzing the relationship between the amount of S and the incompletely quenched structure fraction at a depth of 940 ° C. carburized material of 0.45 mm.
FIG. 4 is a diagram showing an example of analyzing the relationship between the amount of V and the incompletely quenched structure fraction at a depth of 940 ° C. carburized material of 0.45 mm.
FIG. 5 is a diagram showing an example of analyzing the relationship between the amount of Mg and the fraction of incompletely quenched structure at a depth of 940 ° C. carburized material of 0.45 mm.
FIG. 6 is a diagram showing an example of analyzing the relationship between the amount of N and the incompletely quenched structure fraction at a depth of 940 ° C. carburized material of 0.45 mm.
FIG. 7 is a diagram showing an example in which the relationship between the hardness after hot working and the crystal grain coarsening temperature is analyzed.

Claims (3)

% By mass
C: 0.1 to 0.3%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.8%
Cr: 0.83 to 1.5%,
Mo: 0.005 to 0.3%,
B: 0.0031 to 0.006%,
Al: 0.015 to 0.05%,
Ti: 0.01 to 0.1%
Containing
P: 0.025% or less (including 0%),
S: 0.013 % or less (including 0%),
V: 0.01% or less (including 0%),
Mg: 0.03% or less (including 0%),
N: 0.005% or less (including 0%),
O: 0.002% or less (including 0%)
Limit each to
The balance consists of iron and inevitable impurities,
The precipitation amount of AlN is limited to 0.01% or less, the ferrite grain size number is 8 to 11, the bainite structure fraction is 10% or less, and the hardness (HV) is expressed by the following formula (1). A case hardening boron steel for cold forging that does not produce an abnormal structure during carburizing, which is characterized by satisfaction.
H-20 ≦ HV ≦ H + 30 (1)
However, H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr%
+ 136.7Mo% + 708Ti% C: 0.1 to 0.3% by mass%
Si: 0.01 to 0.15%,
Mn: 0.2 to 0.8%
Cr: 0.83 to 1.5%,
Mo: 0.005 to 0.3%,
B: 0.0031 to 0.006%,
Al: 0.015 to 0.05%,
Ti: 0.01 to 0.1%
Nb: 0.002 to 0.05%
Containing
P: 0.025% or less (including 0%),
S: 0.013 % or less (including 0%),
V: 0.01% or less (including 0%),
Mg: 0.03% or less (including 0%),
N: 0.005% or less (including 0%),
O: 0.002% or less (including 0%)
Limit each to
The balance consists of iron and inevitable impurities,
The precipitation amount of AlN is limited to 0.01% or less, the ferrite grain size number is 8 to 11, the bainite structure fraction is 10% or less, and the hardness (HV) is expressed by the following formula (1). A case hardening boron steel for cold forging that does not produce an abnormal structure during carburizing, which is characterized by satisfaction.
H-20 ≦ HV ≦ H + 30 (1)
However, H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr%
+ 136.7Mo% + 708Ti% + 599Nb% The steel comprising the component according to claim 1 or 2 is heated at a temperature of 1150 ° C. or higher, a hot rolling finishing temperature of 840 to 1000 ° C., followed by hot rolling at a temperature range of 800 to 500 ° C. The wire or bar steel is hot worked under the condition of slow cooling at a cooling rate of ℃ / second or less, the precipitation amount of AlN after hot working is limited to 0.01% or less, and the ferrite crystal grain size number is 8 to 11 The bainite structure fraction is 10% or less, and the hardness (HV) satisfies the following formula (1): Case hardening for cold forging that does not generate an abnormal structure during carburizing Boron steel manufacturing method.
H-20 ≦ HV ≦ H + 30 (1)
However, H = 273.5C% + 39.1Si% + 54.7Mn% + 30.4Cr%
+ 136.7Mo% + 708Ti% + 599Nb%

Images