JP5262740B2 - Case-hardened steel with excellent coarse grain prevention and fatigue characteristics during carburizing and its manufacturing method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide case hardening steel in which the generation of coarse grains can be stably suppressed in a carburizing and hardening process, the occurrence of strain and bending after the carburizing and hardening can be prevented, further, the generation of coarse grains can be prevented even in high temperature carburizing, and sufficient strength properties such as rolling fatigue strength can be obtained, and to provide a production method thereof. <P>SOLUTION: The case hardening steel is characterized in that it contains 0.05 to 0.2% Ti and the other specified components in specified ranges, the content of Al is limited to &lt;0.005% and the content of N is limited to &lt;0.0051%, the precipitation amount of AlN after hot rolling is suppressed, or further, the structural fraction of bainite after the hot rolling is limited to &le;30%, or further, the ferrite crystal grain size number after the hot rolling is controlled to No. 8 to 11 prescribed in JIS G0552, or further, in the cross-section in the longitudinal direction in the matrix after the hot rolling, the maximum diameter of Ti based precipitates by the statics of extremes measured under the following conditions is controlled to &le;40 &mu;m. <P>COPYRIGHT: (C)2010,JPO&amp;INPIT

Description

  The present invention relates to a case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing, and a method for producing the same.

  Gears, bearing parts, rolling parts, shafts, constant velocity joint parts are usually made of medium carbon alloy steel for machine structure as defined in JIS G 4052, JIS G 4104, JIS G 4105, JIS G 4106, etc. It is manufactured in a process of carburizing and quenching after being processed into a predetermined shape by cold forging (including rolling) or hot forging-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 becoming stronger, and the application targets and uses of carburized parts manufactured in the cold forging-carburizing process have increased significantly in recent years. A major problem with carburized parts is the reduction of heat treatment distortion. This is because when carburized parts are applied to the shaft, the function as the shaft is impaired if the carburized parts are bent and deformed due to heat treatment distortion, and when carburized parts are applied to gears and constant velocity joint parts, This is because when heat treatment distortion increases, it causes noise and vibration. Here, the largest cause of the heat treatment distortion generated in the carburized parts is based on coarse particles generated during carburizing. In order to suppress the coarse grains generated during carburizing, annealing is conventionally performed after cold forging and before carburizing and quenching. However, especially in recent years, from the viewpoint of cost reduction, the tendency to omit annealing has become stronger. Therefore, even when annealing is omitted, a steel material that does not generate coarse grains in the carburized component is strongly demanded.

  On the other hand, high-depth carburization is performed on bearing parts and rolling parts to which high surface pressure is applied among gears, bearing parts, and rolling parts. Since deep carburization usually takes a long time of several tens of hours to several tens of hours, shortening the carburizing time is an important issue from the viewpoint of reducing energy consumption. Increasing the carburizing temperature is effective for shortening the carburizing time. The normal carburizing temperature is about 930 ° C. On the other hand, when so-called high-temperature carburizing is performed in the temperature range of 990 to 1090 ° C, coarse grains are generated, and necessary fatigue characteristics, rolling fatigue characteristics, etc. are obtained. The problem of not being able to occur. Therefore, there is a demand for case-hardened steel that does not generate coarse grains even at high temperature carburization, that is, suitable for high temperature carburization. Gears, bearing parts, and rolling parts that are loaded with such high surface pressures are often large parts, and are usually "bar steel-hot forging-heat treatment such as normalization if necessary-cutting-carburizing and quenching-polishing if necessary". Manufactured in a process. In order to suppress the generation of coarse grains during carburizing, it is necessary to build in an appropriate material in order to suppress coarse grains in the state after hot forging, that is, in the state of a hot forged member. However, for that purpose, it is necessary to build in an appropriate material in order to suppress coarse grains in the state of the material of the steel bar wire.

  As a technique for stably suppressing the coarse grains of the case-hardened steel in the past, the coarse grains that contain a predetermined amount of Al and N and that optimize the state of the ferrite band of the cross-sectional structure parallel to the hot rolling direction A case-hardened steel excellent in grain prevention characteristics is disclosed (for example, see Patent Document 1). However, in the disclosed technique of Patent Document 1, the ability to suppress coarse grains may not be able to be stably exhibited for parts manufactured through the spheroidizing annealing-cold forging process, and also in high-temperature carburizing. The reality is that the generation of coarse particles may not be suppressed.

The disclosed technique of Patent Document 2 contains Ti: 0.10 to 0.30% by mass% and N: less than 0.01% in addition to a predetermined amount of C, Si, etc. A method for producing case-hardening steel is disclosed in which hot rolling is performed in a temperature range of 1250 to 1400 ° C. and heating is performed at product rolling heating Ac ₃ to 1050 ° C. The disclosed technique of Patent Document 3 discloses a technique for improving the rolling fatigue life and the rotating bending fatigue life by finely dispersing Ti carbide in the case-hardened steel composed of the components as in Patent Document 2. The case-hardened steel disclosed in Patent Documents 2 and 3 may not be able to stably exhibit the ability to suppress coarse grains. Further, as apparent from the above examples, the steel contains 0.0055% or more of N, and TiN coarsens due to product rolling and heating: Ac _{3 to} 1050 ° C. For this reason, this coarse TiN serves as a starting point for rolling fatigue and pitching fatigue, and thus has a drawback that sufficient fatigue characteristics cannot be obtained.

Further, in Patent Document 4, in addition to a predetermined amount of C, Si, or the like, Ti: more than 0.1 to 0.2%, N: 0.015% or less, and the prior austenite grain size is JIS in mass%. No. G0551 High-strength case-hardened steel having a martensite structure refined to 11 or more, or N: 0.020% or less in mass%, “Ti: 0.05 to 0.2%, V: 0.02 to 0.10 %,
Nb: 0.02 to 0.1% ”, and the former austenite grain size is JIS G0551 No. A high-strength case-hardened steel having a martensite structure refined to 11 or more is disclosed.

  Moreover, in patent document 5, Ti: 0.05-0.2% contains a specific range other specific component in the mass%, N: It restrict | limits to less than 0.0051% by mass%, or further, Nb by mass%. : Containing less than 0.04%, limiting the precipitation amount of AlN after hot rolling to 0.01% or less, or further limiting the structure fraction of bainite after hot rolling to 30% or less, or Further, the ferrite grain size number after hot rolling is set to Nos. 8 to 11 specified in JIS G0552, or, further, extreme values measured under the following conditions in the longitudinal section in the matrix of the steel after hot rolling There is disclosed a case-hardened steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing, in which the maximum diameter of Ti-based precipitates is 40 μm or less according to statistics.

  However, the Ti-added coarse-grained steel as shown in Patent Documents 1 to 5 described above has a problem that machinability deteriorates due to the addition of a large amount of Ti. There are two main causes of machinability deterioration: 1) an increase in material hardness, and 2) a decrease in MnS, which has a high machinability improvement effect due to the formation of Ti-based carbonitrides.

JP-A-11-106866 JP-A-11-92863 JP-A-11-92824 JP 2003-34843 A JP-A-2005-240175

  In the disclosed technologies of Patent Documents 1 to 5 as described above, the generation of coarse grains cannot be stably suppressed in the carburizing and quenching process, and the generation of distortion and bending cannot be stably prevented, or It was difficult to apply because of concern about deterioration of machinability. This invention solves such a problem, and provides the case hardening steel excellent in the coarse grain prevention characteristic at the time of carburizing at the time of carburizing with small heat processing distortion, and the fatigue characteristic, and its manufacturing method.

  In order to solve the above-mentioned problems, the present inventors diligently investigated the controlling factors for the coarsening of crystal grains and the method for improving the machinability deterioration due to the addition of a large amount of Ti for suppressing the coarsening. Was revealed.

(1) To prevent coarsening of crystal grains during carburizing, it is effective to finely precipitate Ti-based precipitates mainly composed of TiC and TiCS during carburizing rather than using AlN or NbN as pinning particles. is there. Or, in combination with the above, the coarse grain prevention characteristics are further improved by finely precipitating Nb carbonitrides mainly composed of NbC during carburizing.

(2) The following method was newly found as a method for finely depositing the Ti-based precipitate or the Nb carbonitride described above during carburizing.

[1] In order to utilize the Ti-based precipitates or the Nb carbonitrides as pinning particles during carburizing, it is necessary to finely disperse these precipitates in a large amount during carburizing and quenching. For this purpose, the Ti-based precipitate or the Nb carbonitride needs to be once solutionized at the time of rolling and heating when hot rolling a steel bar or wire. If the amount of N is high and a large amount of TiN remains during rolling and heating, the Ti-based precipitate becomes a composite precipitate mainly composed of TiN, which makes it difficult to form a solution. Further, in the cooling process after hot rolling, TiC, TiCS or further NbC is deposited on coarse TiN, and the fine dispersion of these precipitates is hindered. Therefore, it is necessary to reduce the N amount as much as possible.

[2] If coarse AlN is present during rolling and heating, it 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 formed into a solution at the time of rolling and heating, precipitation of AlN hardly occurs in the hot rolling-cooling process of the steel bars and wires. Therefore, by limiting the amount of Al and regulating the amount of AlN deposited after hot rolling, it is possible to confirm the solution state of AlN during the heating by rolling.

[3] If heating is performed under the condition that AlN can be dissolved during rolling and heating, it is possible to temporarily form Ti-based precipitates or further NbC precipitates. Therefore, by regulating the amount of AlN deposited after hot rolling, it is possible to confirm that Ti-based precipitates or further NbC precipitates were once formed into a solution during rolling and heating.

[4] Further, in order to stably exhibit the pinning effect of Ti-based precipitates or further NbC precipitates, it is necessary to finely precipitate these precipitates in the matrix after hot rolling. is there. For this purpose, it is necessary to cause phase interface precipitation during the diffusion transformation from austenite during the cooling process during hot rolling. If bainite is generated in a structure that has been hot-rolled, it becomes difficult to precipitate the phase interface of the above-mentioned precipitates. Therefore, it is essential to make the structure substantially free of bainite.

(3) In order to limit the precipitation amount of AlN as much as possible in the state of the steel material after hot rolling, that is, in order to solutionize Ti-based precipitates or further NbC precipitates during hot rolling heating, rolling It is necessary to increase the heating temperature. Furthermore, the amount of Al was limited to less than 0.005%, and the amount of precipitated AlN was suppressed.

(4) In order to finely precipitate Ti-based precipitates or further NbC precipitates in the steel material after hot rolling, the rolling heating temperature and the cooling conditions after rolling may be optimized. That is, by increasing the rolling heating temperature, Ti-based precipitates or further NbC precipitates are once dissolved in the matrix, and after hot rolling, the precipitation temperature range of Ti-based precipitates or further NbC precipitates is gradually increased. By cooling, a large amount of these carbonitrides can be finely dispersed.

(5) Furthermore, if the ferrite grains of the steel material after hot rolling are excessively fine, coarse grains are likely to be generated during carburizing heating, so that optimization of the rolling finishing temperature is also a point.

(6) In Ti-added steels, Ti precipitates are the starting point for fatigue failure, so fatigue characteristics, particularly rolling fatigue characteristics, are likely to deteriorate. However, Ti precipitates are reduced by reducing N and increasing the hot rolling temperature. By reducing the maximum size, fatigue characteristics can be improved, and both coarse grain prevention characteristics and fatigue characteristics can be achieved.

(7) Further, by adjusting the oxide composition, 1) the oxide has a low melting point and is softened in the cutting environment to improve machinability, and 2) the oxide is a sulfide inclusion. By using the crystal precipitation nuclei, it is possible to finely disperse sulfide inclusions that reduce the amount of limit compressive strain that generates cracks under cold forging, thereby improving cold forgeability.

  The present invention has been made based on the above novel findings, and the gist of the present invention is as follows.

That is, in the invention according to claim 1, the chemical component is mass%, C: 0.1 to 0.6%, Si: 0.02 to 1.5%, Mn: 0.3 to 1.8% , S: 0.001 to 0.15%, Al: less than 0.005%, Ti: 0.05 to 0.2%, N: less than 0.0051%, P: 0.025% or less, O: 0 .0025% or less, Cr: 0.4-2.0%, Mo: 0.02-1.5%, Ni: 0.1-3.5%, V: 0.02-0.5% limiting contain one or more 0.0002 to 0.005 percent, the remainder Ri is Do iron and unavoidable impurities, after hot rolling, the structural fraction of bainite below 30%:, B In addition, the ferrite grain size number is 8 to 11 defined in JIS G0552, and the maximum diameter of the Ti-based precipitate is 4 in the longitudinal section in the steel matrix. And wherein the Der Rukoto following 0μm.

The invention according to claim 2 is the invention according to claim 1, wherein the chemical component is mass%, and Zr is 0.0003 to 0.0050% and Mg is 0.0003 to 0.0050%. Or it contains 2 types, It is characterized by the above-mentioned.

  The invention described in claim 3 is characterized in that, in the invention described in claim 1 or 2, the chemical component is contained in mass% and contains Nb: less than 0.04%.

The invention according to claim 4 is the invention according to any one of claims 1 to 3, wherein the heat holding time is 10 minutes or more at a temperature of 1150 ° C. or higher and hot-rolled to a wire rod or steel bar. The finishing temperature of rolling is set to 840 to 1000 ° C., the steel has a ferrite grain size number of 8 to 11 specified in JIS G0552, and the temperature range of 800 to 500 ° C. is 1 ° C./second after hot rolling. Slow cooling is performed at the following cooling rate so that the steel has a bainite structure fraction of 30% or less after hot rolling.

The invention according to claim 5 is a carburized part that is processed into a part shape using the case-hardened steel according to any one of claims 1 to 3 .

  By using the case-hardened steel excellent in the coarse grain prevention characteristics and fatigue characteristics during carburizing and the manufacturing method thereof according to the present invention, even when parts are manufactured in the cold forging process, the coarsening of crystal grains during carburizing is suppressed. Therefore, the fatigue strength characteristics can be improved, and the deterioration of dimensional accuracy due to quenching distortion can be extremely reduced as compared with the conventional case. For this reason, it becomes possible to cold forge parts that have been difficult to cold forge due to the problem of coarse grains, and it is also possible to omit annealing after cold forging. Moreover, even if this steel material is applied to a part manufactured in the hot forging process, generation of coarse grains can be prevented even in high-temperature carburization, and sufficient strength characteristics such as rolling fatigue characteristics can be obtained. Moreover, according to the case-hardened steel to which the present invention is applied, the machinability, which has been one of the problems of the coarse grain-preventing steel, can be exhibited excellent machinability. Cutting workability can be obtained. As described above, the industrial effects of the present invention are extremely remarkable.

It is a figure which shows a Charpy impact test piece.

  Hereinafter, as a form for carrying out the present invention, case hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing will be described in detail.

  First, the reasons for limiting chemical components in case-hardened steel to which the present invention is applied will be described. Hereinafter, the mass% in the composition is simply described as%.

C: 0.1 to 0.6%
C is an element effective for giving steel the necessary strength, but if it is less than 0.1%, the required tensile strength cannot be secured, and if it exceeds 0.6%, it becomes hard and cold work is performed. In addition to the deterioration of the properties, the core toughness after carburizing deteriorates, so the C content needs to be in the range of 0.1 to 0.6%.

Si: 0.02 to 1.5%
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 0.02%, the effect is ineffective. It is enough. On the other hand, if it exceeds 1.5%, the hardness is increased and the cold forgeability is deteriorated. For the above reasons, the content needs to be in the range of 0.02 to 1.5%. The suitable range of steel materials that undergo cold working is 0.02 to 0.3%. In particular, when emphasizing cold forgeability, it is desirable to make the range 0.02 to 0.15%. On the other hand, Si is an element effective for increasing the grain boundary strength. Further, in bearing parts and rolling parts, it is an element effective for extending the life by suppressing structural changes and material deterioration during the rolling fatigue process. Therefore, when aiming at high intensity | strength, the range of 0.2 to 1.5% is suitable. In particular, in order to obtain a high level of rolling fatigue strength, it is desirable to be in the range of 0.4 to 1.5%. In addition, the effect of suppressing the structural change and material deterioration in rolling fatigue process of bearing parts and rolling parts due to the addition of Si is 30 to 40% of the retained austenite amount (common name, residual γ amount) in the structure after carburizing. Especially at times. In order to control the residual γ amount within this range, it is effective to perform a so-called carburizing and nitriding treatment. The carburizing and nitriding process is a process of performing nitriding in the process of diffusion after carburizing. Conditions under which the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. In this case, it is desirable that the carbon potential at the time of carburizing is in the range of 0.9 to 1.3%.

Mn: 0.3 to 1.8%
Mn is an element effective for deoxidation of steel and is an element effective for imparting the necessary strength and hardenability to the steel, but if less than 0.3%, the effect is insufficient, 1.8% If it exceeds, the effect is not only saturated, but also the hardness is increased and the cold forgeability is deteriorated, so it is necessary to be within the range of 0.3% to 1.8%. The preferred range is 0.5-1.2%. In addition, when importance is attached to cold forgeability, it is desirable to set it as 0.5 to 0.75% of range.

S: 0.001 to 0.15%
S forms MnS in the steel and is added for the purpose of improving the machinability. However, if it is less than 0.001%, its effect is insufficient. On the other hand, if it exceeds 0.15%, the effect is saturated, and rather, grain boundary segregation occurs, leading to grain boundary embrittlement. For these reasons, the S content needs to be in the range of 0.001 to 0.15%. In addition, in bearing parts and rolling parts, since MnS deteriorates the rolling fatigue life, it is necessary to reduce S as much as possible, and it is desirable to make it in the range of 0.001 to 0.01%.

Al may be added as a deoxidizer, but the remaining amount of Al is less than 0.005%. If over 0.005%, the Al-based oxide present in the steel hinders machinability. Unless the Al content is less than 0.005%, hard inclusion alumina (so-called corundum) is produced and remains in the steel. This hard inclusion promotes abrasive wear on the cutting tool, lowers the machinability, and is extremely disadvantageous for controlling MnS in steel. In general, it is known that inclusions represented by MnS are likely to become nuclei at the time of crystallization of MnS, but alumina is less effective and often exists alone. On the other hand, produces a very stable alumina composite oxide or the like and combined addition of 0.005% or more be added other oxide formation elemental Z r in terms of softening oxides remain rigid is there.

  This greatly affects manufacturability on steelmaking. That is, it is said that the alumina-based oxide is likely to adhere to or clog the nozzle used for continuous casting. Therefore, it was not suitable for mass production by continuous casting and was not suitable for making stable materials.

  Thus, the limitation of the amount of Al is important from the viewpoint of high quality and quality stability in terms of machinability, steelmaking manufacturability, and sulfide control.

  Al is preferably suppressed to 0.003% or less, more preferably 0.001% or less, whereby the oxides in steel are more difficult Si-Mn oxides, Ca, Zr, Mg, etc. Complex oxides with other oxide-generating elements are generated, and it becomes easier to generate softer or more suitable oxides for controlling MnS.

Ti: 0.05 to 0.2%
Ti is added in order to produce fine TiC and TiCS in the steel and thereby to refine the γ grains during carburization. However, if it is less than 0.05%, the effect is insufficient. On the other hand, when Ti is added in excess of 0.2%, precipitation hardening due to TiC becomes remarkable, cold workability is remarkably deteriorated, and precipitates mainly composed of TiN become remarkable, and rolling fatigue characteristics deteriorate. For the above reasons, the content needs to be in the range of 0.05 to 0.2%. The preferred range is 0.05 to less than 0.1%. In the steel and hot forged member of the present invention, carbon and nitrogen that enter during carburizing heating react with solute Ti, and a large amount of fine Ti (CN) precipitates in the carburized layer. Therefore, in bearing parts and rolling parts, these Ti (CN) contributes to the improvement of rolling fatigue life. Therefore, in bearing parts and rolling parts, in the case of directing a particularly high level of rolling fatigue life, the carbon potential at the time of carburizing should be set higher in the range of 0.9 to 1.3%, or It is effective to perform so-called carburizing and nitriding treatment. The carburizing and nitriding treatment is a treatment in which nitriding is performed in the course of the diffusion treatment after carburizing as described above, but the conditions that the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. is there.

  The present inventor has found that when Ti is in the range of 0.05 to 0.2%, MnS is fine and reduced through the formation of TiCS, thereby improving the impact value.

N: Less than 0.0051% When N is combined with Ti in steel, it produces coarse TiN that hardly contributes to grain control. This is TiC, TiCS-based Ti-based precipitates, NbC, NbC-based Nb (CN). It becomes a precipitation site, inhibits fine precipitation of these Ti-based precipitates and Nb carbonitrides, and promotes the formation of coarse particles. The above adverse effect is particularly remarkable when the N amount is 0.0051% or more. For the above reasons, the content needs to be less than 0.0051%.

P: 0.025% or less P is an element that increases deformation resistance during cold forging and deteriorates toughness, so that 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.

O: 0.0025% or less In the high Ti steel as in the present invention, O forms Ti-based oxide inclusions in the steel. When a large amount of oxide inclusions are present in the steel, TiC precipitates, and TiC precipitates coarsely during hot rolling, making it impossible to suppress grain coarsening during carburizing. Therefore, it is desirable to reduce the amount of O as much as possible. For the above reasons, the content needs to be limited to 0.0025% or less. The preferred range is 0.0020% or less. In bearing parts and rolling parts, oxide inclusions are the starting point of rolling fatigue failure, so the rolling life is improved as the O content is lower. Therefore, it is desirable to limit the O content to 0.0012% or less in bearing parts and rolling parts.

  Moreover, this invention contains 1 type (s) or 2 or more types of Cr, Mo, Ni, V, and B prescribed | regulated by the following component range.

Cr: 0.4-2.0%
Cr is an element effective for imparting strength and hardenability to steel. Further, in bearing parts and rolling parts, the amount of residual γ after carburizing is increased, and the structure changes and material deterioration occurs during rolling fatigue. It is an element that is effective in extending the life by suppressing the above. If it is less than 0.4%, the effect is insufficient, and if it exceeds 2.0%, the hardness is increased and the cold forgeability deteriorates. For the above reasons, the content needs to be in the range of 0.4 to 2.0%. A preferable range is 0.7 to 1.6%. The effect of suppressing the structural change and material deterioration in the rolling fatigue process of bearing parts and rolling parts due to the addition of Cr is particularly great when the amount of residual γ in the structure after carburizing is 30 to 40%. In order to control the amount of residual γ within this range, it is effective to perform a so-called carburizing and nitriding treatment so that the surface nitrogen concentration is in the range of 0.2 to 0.6%.

Mo: 0.02 to 1.5%
Mo has an effect of imparting hardenability by adding, and 0.02% or more is necessary to obtain the effect. Further, in bearing parts and rolling parts, it is an element effective in increasing the residual γ amount after carburizing and extending the life by suppressing the structural change and material deterioration in the rolling fatigue process. However, if added over 1.5%, the hardness increases, and the machinability and cold forgeability deteriorate. For the above reasons, the content needs to be in the range of 1.5% or less. The preferred range is 0.5% or less. As with Cr, the effect of suppressing the structure change and material deterioration during rolling fatigue of bearing parts and rolling parts due to the addition of Mo is the so-called carburizing and nitriding treatment, and the amount of residual γ in the structure after carburizing. Is particularly large at 30-40%.

Ni: 0.1 to 3.5%
Ni has an effect of imparting hardenability by adding, and 0.1% or more is necessary to obtain the effect. If added over 3.5%, the hardness is increased, and the machinability and cold forgeability deteriorate. For the above reasons, the content needs to be in the range of 3.5% or less. The preferred range is 2.0% or less.

V: 0.02-0.5%
V has the effect of imparting hardenability when added, and 0.02% or more is necessary to obtain the effect. If added over 0.5%, the hardness is increased, and the machinability and cold forgeability deteriorate. For the above reasons, the content needs to be in the range of 0.5% or less. The preferred range is 0.2% or less.

B: 0.0002 to 0.005%
B has the effect of imparting hardenability when added, and 0.0002% or more is necessary to obtain the effect. Furthermore, in B, 1) the effect of increasing the growth rate of ferrite by forming boron iron carbide in the cooling process after rolling in steel bar / wire rolling, and promoting softening while being rolled, 2) carburizing By improving the grain boundary strength of the material, it also has the effect of improving fatigue strength and impact strength as carburized parts. However, when B is added over 0.005%, the effect is saturated, and adverse effects such as deterioration of impact strength are concerned. Therefore, the content needs to be within the range of 0.005% or less. The preferred range is 0.003% or less.

Further in the present invention may contain a Z r, Mg, Nb that will be defined by the components within the following range.

Zr: 0.0003 to 0.0050%
Zr is a deoxidizing element and generates an oxide, but also has an interrelationship with MnS by generating sulfides. Zr-based oxides tend to become nuclei for crystallization / precipitation of MnS. Therefore, it is effective for distributed control of MnS. Further, in an environment where the amount of Al is as low as 0.0005%, this effect is emphasized by the increase of the Zr-based oxide, and the adverse effect of MnS that degrades fatigue strength can be suppressed.

The amount of Zr added is preferably more than 0.003% in order to aim at spheroidization of MnS. However, in order to finely disperse, addition of 0.0003 to 0.0050% is preferable. As the product, the latter is practically preferable from the viewpoint of production in terms of quality stability (component yield and the like), that is, 0.0003 to 0.0050% in which MnS is finely dispersed. If it is 0.0002% or less, the effect of adding Zr is hardly observed.

Mg: 0.0003 to 0.0050%
Mg is a deoxidizing element and generates an oxide, but also has an interrelationship with MnS by generating sulfides. Mg-based oxides tend to become nuclei for crystallization / precipitation of MnS. Moreover, since the sulfide becomes a composite sulfide of Mn and Mg, the deformation is suppressed and spheroidized. Therefore, it is effective for dispersion control of MnS. This effect is preferably in an environment where the amount of Al is as low as 0.0005%, and this effect is emphasized by increasing the amount of Mg-based oxide, and the adverse effect of MnS that degrades fatigue strength can be suppressed. In order to express these effects, it is necessary to add at least 0.0003% of Mg. However, when the amount of Mg exceeds 0.0050%, such effects are saturated.

Nb: less than 0.04% Next, claim 3 of the present invention contains Nb: less than 0.04% in addition to claims 1 and 2, and the reason for this limitation will be described below. Nb combines with C and N in steel during carburizing heating to form Nb (CN), and is an element effective for suppressing coarsening of crystal grains. By adding Nb, the effect of “preventing coarse grains due to Ti-based precipitates” becomes more effective. This is because Nb dissolves in the Ti-based precipitate and suppresses the coarsening of the Ti-based precipitate. Therefore, within the range of the addition amount of the present invention, the effect increases depending on the addition amount of Nb, but in a small amount addition such as less than 0.03%, less than 0.02%, and even less than 0.01%. However, compared with the case where Nb is not added, the coarse grain prevention characteristic is remarkably improved. However, Nb addition causes deterioration of machinability, cold forgeability, and carburization characteristics. In particular, when the amount of Nb added is Nb: 0.04% or more, the hardness of the material becomes hard and the machinability and cold forgeability deteriorate, and it is difficult to form a solution during heating of the steel bar and wire rod. Become. For the above reasons, the content needs to be within a range of less than 0.04%. The preferred range when workability such as machinability and cold forgeability is important is less than 0.03%. In addition to workability, the preferred range when carburization is important is less than 0.02%. Furthermore, the preferable range in the case where carburizability is particularly emphasized is less than 0.01%. In order to achieve both the coarse grain prevention characteristics and the workability, it is recommended that the amount of Nb added be adjusted according to the amount of Ti added. For example, the preferable range of Ti + Nb is 0.07 to less than 0.17%. Particularly in high-temperature carburizing and cold forged parts, the desirable range is more than 0.091% and less than 0.17%.

When coarse AlN is present during rolling and heating, it becomes a precipitation site for Ti-based precipitates and Nb precipitates, and after hot rolling, Ti-based precipitates and Nb precipitates are coarsely precipitated, and the grains become coarse during carburizing. Can not be suppressed. 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. In order to utilize Ti-based precipitates and Nb precipitates as pinning particles, these precipitates also need to be dissolved once during rolling and heating. If heating is performed under conditions that allow AlN to form a solution during rolling and heating, the above precipitate can be once formed into a solution. Therefore, by regulating the precipitation amount of AlN after hot rolling, it is possible to confirm that the Ti-based precipitates and Nb precipitates were once in solution during rolling and heating.

Next, in claims 1 and 4 of the present invention, the structure fraction of bainite after hot rolling is limited to 30% or less. The reason for this limitation will be described below. If a bainite structure is mixed in the steel after hot rolling, coarse grains are generated during carburizing heating. In addition, suppression of bainite contamination is also desirable from the viewpoint of improving cold workability. These adverse effects become particularly prominent when the bainite structural fraction exceeds 30%. For the above reasons, it is necessary to limit the structure fraction of bainite after hot rolling to 30% or less. The preferred range when the carburizing conditions are severe for preventing coarse grains due to high-temperature carburizing or the like is 20% or less. The preferred range is 10% or less when the carburizing conditions are severe for preventing coarse grains via cold forging.

Next, in claims 1 and 4 of the present invention, the ferrite grain size number after hot rolling is set to Nos. 8 to 11 defined in JIS G0552. The reason for such limitation will be described below. If the ferrite grains after hot rolling are excessively fine, the austenite grains are excessively refined during carburizing. If the austenite grains become excessively fine, coarse grains are likely to be formed, and the tendency becomes prominent particularly when the ferrite crystal grain size exceeds # 11. Further, when the austenite grains become excessively fine so as to exceed No. 11 defined in JIS G0551, similar to the steel material described in Patent Document 4, problems such as insufficient strength due to deterioration of hardenability occur. On the other hand, when the ferrite crystal grain size number after hot rolling is coarser than the number 8 specified in JIS G0552, the ductility of the hot rolled material deteriorates and the cold forgeability deteriorates. For the above reasons, it is necessary to set the ferrite crystal grain size number after hot rolling within the range of 8 to 11 defined by JIS G0552.

In the first aspect of the present invention, in longitudinal section in the matrix of the steel after hot rolling, the inspection reference area: 100 mm ^2, the number of inspections 16 viewing area make predictions: 30,000 mm measured pole in ^two conditions The maximum diameter of the Ti-based precipitate by value statistics is 40 μm or less. The reason for this limitation will be described below. One of the required characteristics of the carburized parts targeted by the present invention is contact fatigue strength such as rolling fatigue characteristics and surface fatigue strength. If coarse Ti-based precipitates are present in the steel, it becomes a starting point for contact fatigue failure, and the fatigue characteristics deteriorate. The extreme value statistics, inspection reference area: 100 mm ^2, the number of inspections 16 viewing area make predictions: the maximum diameter of Ti-based precipitates when measured under the conditions of 30,000 mm ² is more than 40 [mu] m, on the particular contact fatigue properties The adverse effect of Ti-based precipitates becomes significant. For the above reasons, the maximum diameter of Ti-based precipitates according to extreme value statistics needs to be 40 μm or less. The method for measuring and predicting the maximum diameter of precipitates by extreme value statistics is, for example, the method described in Takayoshi Murakami “Effects of Metal Fatigue Microdefects and Inclusions” Yokendo pp 233-239 (1993). The present invention uses a two-dimensional inspection method that estimates the maximum precipitate observed within a certain area (predicted area: 30000 mm ² ) by two-dimensional inspection. Detailed measurement procedures are described in detail in the Examples section below.

  Next, hot rolling conditions in the method for producing the case hardening steel to which the present invention is applied will be described.

  The steel of the present invention composed of the above-described chemical components is melted by a usual method such as a converter or an electric furnace, the components are adjusted, and the wire or bar steel is heated through a casting process and, if necessary, a block rolling process. The rolled material is manufactured by rolling.

Next, Claim 4 of this invention heats this manufactured rolling raw material at the temperature of 1150 degreeC or more and the heat retention time for 10 minutes or more. If the heating condition is less than 1150 ° C. or the heat retention time is less than 10 minutes even if the heating temperature is 1150 ° C. or higher, Ti-based precipitates, Nb precipitates and AlN cannot be once dissolved in the matrix. . Therefore, a certain amount or more of Ti-based precipitates and Nb precipitates cannot be finely precipitated in advance on the steel material after hot rolling, and coarse Ti-based precipitates and Nb precipitates exist after hot rolling. In addition, the generation of coarse grains cannot be suppressed during carburizing. Therefore, in hot rolling, it is necessary to heat at a temperature of 1150 ° C. or higher for a heat retention time of 10 minutes or longer. The preferred range is a hot rolling temperature of 1180 ° C. or higher and a heat retention time of 10 minutes or longer.

Next, Claim 4 of this invention anneals the temperature range of 800-500 degreeC with the cooling rate of 1 degrees C / sec or less after hot rolling. When the cooling condition exceeds 1 ° C./second, the precipitation temperature range of the Ti-based precipitate can be passed only for a short time, and the amount of fine TiC-based precipitate after hot rolling becomes insufficient, Moreover, the structure fraction of bainite increases. Therefore, it becomes impossible to suppress the generation of coarse particles during carburizing. Further, if the cooling rate is high, the hardness of the rolled material increases and the cold forgeability deteriorates. Therefore, it is desirable to make the cooling rate as small as possible. The preferred range is 0.7 ° C./second 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.

Next, claim 4 of the present invention sets the finishing temperature of hot rolling to 840 to 1000 ° C. If the finishing temperature is less than 840 ° C., the ferrite crystal grain size becomes excessively fine, and coarse grains are likely to be generated during the subsequent carburizing. On the other hand, when the finishing temperature exceeds 1000 ° C., the hardness of the rolled material becomes hard and the cold forgeability deteriorates. For these reasons, the hot rolling finishing temperature is set to 840 to 1000 ° C. When used for cold forging without annealing, a range of 840 to 920 ° C is desirable, and a range of 920 to 1000 ° C is desirable otherwise.

  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.

  The present invention can be applied to both parts manufactured by the cold forging process and parts manufactured by hot forging. As an example of the hot forging process, there is a process of “bar steel-hot forging-heat treatment such as normalization if necessary-cutting-carburizing and quenching-polishing if necessary". By using the steel material of the present invention, for example, hot forging at a heating temperature of 1150 ° C. or higher, and then performing a normalization treatment as necessary, such as severe carburization in a temperature range of 950 ° C. to 1090 ° C. Even in carburizing and quenching heat treatment under conditions, the generation of coarse grains can be suppressed, and excellent material properties can be obtained. For example, in the case of bearing parts and rolling parts, excellent rolling fatigue characteristics can be obtained even when high-temperature carburizing is performed.

  In the present invention, carburizing conditions are not particularly limited. In bearing parts and rolling parts, when aiming at a particularly high level of rolling fatigue life, as described above, the carbon potential during carburizing should be set higher within the range of 0.9 to 1.3%. Alternatively, it is effective to perform a so-called carburizing and nitriding treatment. The carburizing and nitriding treatment is a treatment in which nitriding is performed in the course of the diffusion treatment after carburizing, and conditions under which the surface nitrogen concentration is in the range of 0.2 to 0.6% are appropriate. By selecting these conditions, a large amount of fine Ti (CN) precipitates in the carburized layer, and the introduction of 30 to 40% of residual γ contributes to the improvement of the rolling life.

  In the present invention, it is needless to say that carburized parts formed using the case-hardened steel having the above-described configuration and processed into a part shape are also included.

  Examples of the present invention will be described below.

  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, steel bars having a diameter of 24 to 30 mm were manufactured by hot rolling.

About the steel bar after hot rolling, micro observation was performed and the ferrite grain size was measured according to the measurement of the bainite fraction and the provisions of JIS G 0552. Furthermore, the Vickers hardness was measured and used as an index of cold workability.

The steel bar manufactured in the above process was subjected to spheroidizing annealing, then an upsetting test piece was prepared, and upsetting was performed at a reduction rate of 50%, and then a carburizing 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 the coarse grain generation temperature was 950 ° C. or lower and the crystal grain coarsening characteristics were inferior. 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. Moreover, the maximum diameter of Ti-based precipitates was estimated by the following method using the extreme value statistical method. Whether or not the precipitate is Ti-based was determined from the difference in contrast in the optical microscope. The validity of the discrimination method based on the difference in contrast was confirmed in advance with a scanning electron microscope equipped with an energy dispersive X-ray spectrometer. A region having an inspection reference area of 100 mm ² (10 mm × 10 mm region) in the longitudinal section of the test piece was prepared in advance for 16 fields of view. Then, a Ti-based maximum precipitate in each inspection reference area 100 mm ² was detected, and this was photographed 1000 times with an optical microscope. This was repeated 16 times for 16 visual fields of each inspection reference area 100 mm ² (that is, 16 visual fields were inspected). The diameter of the largest deposit in each inspection reference area was measured from the obtained photograph. In the case of an ellipse, the geometric mean of the major axis and minor axis was determined and used as the diameter of the precipitate. 16 data of the maximum precipitate diameter obtained were plotted on the extreme probability sheet by the method described in Takayoshi Murakami “Effects of Metal Fatigue and Micro Defects and Inclusions” Yokendo pp 233-239 (1993). The maximum precipitation distribution line (maximum precipitate diameter and linear function of the extreme value statistical normalization variable) is obtained, and the maximum precipitation distribution line is extrapolated to obtain the prediction area: 30000 mm ² The diameter was predicted. further,
After the hot rolling, the steel bar having a diameter of 24 to 30 mm was subjected to a heat treatment of normalizing and adjusting cooling to form a total ferrite-pearlite structure, and then a cylindrical test piece having a diameter of 22 to 28 mm and a height of 21 mm was cut and milled. The thing was used as the test piece for drill cutting. A drill drilling test was performed on the test piece for drill cutting under the cutting conditions shown in Table 2, and the machinability of each steel material of the present invention example and the comparative example was evaluated. At that time, as an evaluation index, a maximum cutting speed VL1000 capable of cutting to a cumulative hole depth of 1000 mm was adopted in the drill drilling test.

（NACHI通常ドリルとは株式会社不二越製の型番SD3.0のドリルを示す。）

(NACHI normal drill means a drill with model number SD3.0 manufactured by Fujikoshi Co., Ltd.)

Next, a cylindrical rolling fatigue test piece having a diameter of 12.2 mm was prepared from each steel material that had been cold forged at a reduction rate of 50%, under the conditions of 950 ° C. × 5 hours and carbon potential of 0.8%. Carburized. The temperature of the quenching oil is 130 ° C., and the tempering is 180 ° C. × 2 hours. For these carburized and quenched materials, the γ grain size of the carburized layer was investigated. Furthermore, the rolling fatigue characteristics were evaluated using a point contact type rolling fatigue tester (Hertz maximum contact stress 5884 MPa). As a measure of fatigue life, L ₁₀ life defined as “the number of stress repetitions until fatigue failure at a cumulative failure probability of 10% obtained by plotting test results on Weibull probability paper” was used.

These survey results are summarized in Table 3. The rolling fatigue life indicates the relative value of the L ₁₀ life of each material when the L ₁₀ life of Comparative Example 12 is 1.

  It is clear that the crystal grain coarsening temperature of the example of the present invention is 990 ° C. or higher, the γ grains of the 950 ° C. carburized material are finely sized, and have excellent rolling fatigue characteristics. Also, the machinability evaluated by VL1000, which is an index of machinability, is all good at 36 m / min or more in the examples of the present invention, and it is clear that the machinability is excellent.

On the other hand, Comparative Example 19 has the characteristics of JIS SCr420, but the coarse grain generation temperature is low, and the γ grains of the 950 ° C. carburized material are coarse. Comparative Example 29 is a case where the N content exceeds the range specified in the present application, and the maximum diameter of the Ti-based precipitate exceeds the range specified in the present application. Fatigue properties are not good. In Comparative Example 21, the Al content exceeded the range specified in the present application, and the machinability was inferior. Comparative Example 22 is a case where the Ti content falls below the range specified in the present application, and since the Ti pinning effect is small, it does not represent an effect for suppressing coarse grains. In Comparative Example 23, the Ca content exceeded the range specified in the present application, and the machinability was inferior. In Comparative Example 24, the components are within the range specified in the present application, but the cooling rate after hot rolling exceeds the range of the present invention, and the bainite structure fraction after hot rolling exceeds the range specified in the present application. Coarse grains are also produced. In Comparative Example 25, the finishing temperature was lower than the range of the present invention, and the ferrite crystal grain size after rolling became finer than the range of the present invention, and the coarse grain prevention characteristics were inferior. Comparative Example 26 is a case where the rolling heating temperature is below the range of the present invention, and the precipitation amount of AlN after hot rolling exceeds the range specified in the present application, the coarse grain prevention characteristics are inferior, and the rolling fatigue characteristics are also good. There wasn't.

  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 70 mm was manufactured by hot rolling. Using this steel bar as a raw material, hot forging was performed to finish a hot forged member having a diameter of 40 mm. The heating temperature of hot forging is 1100 ° C to 1290 ° C.

About the hot forging member manufactured at the process mentioned above, the normalization process was performed on the conditions of 900 degreeC x 1 hour heating and air cooling. Thereafter, a carburization simulation was performed under the condition of a heating time of 5 hours, and the coarse grain generation temperature was determined in the same manner as in Example 1. Further, after normalizing the hot forged member, a cylindrical rolling fatigue test piece having a diameter of 12.2 mm and a Charpy impact test piece with a 10R notch shown in FIG. Carburizing and quenching was performed at a potential of 1.0%. The temperature of the quenching oil is 130 ° C., and the tempering is 180 ° C. × 2 hours. The rolling fatigue life test was performed under the same conditions as in Example 1. The Charpy impact test was conducted at room temperature and organized by absorbed energy. These survey results are summarized in Table 3. The rolling fatigue life indicates the relative value of the L ₁₀ life of each material when the L ₁₀ life of Comparative Example 12 is 1.

  Moreover, as shown in Table 4, in the example of the present invention, the crystal grain coarsening temperature is over 1070 ° C. Further, the γ grains of the 1050 ° C. carburized material are fine grains of No. 8 or more, and the rolling fatigue life is extremely good, that is, twice or more as compared with the comparative example.

On the other hand, the comparative example deviated from the range of the requirements of the present invention as in Example 1, and the coarse grain prevention characteristics were inferior, and the rolling fatigue characteristics were lower than those of the present invention example.

Claims (5)

Chemical composition is mass%,
C: 0.1-0.6%
Si: 0.02 to 1.5%,
Mn: 0.3 to 1.8%
S: 0.001 to 0.15%,
Al: less than 0.005%,
Ti: 0.05 to 0.2%,
N: less than 0.0051%,
P: 0.025% or less,
O: 0.0025% or less,
Furthermore, Cr: 0.4-2.0%, Mo: 0.02-1.5%, Ni: 0.1-3.5%, V: 0.02-0.5%, B: 0.0. contain one or two or more from 0002 to 0.005%, Ri Do the balance iron and unavoidable impurities,
After hot rolling, the structure fraction of bainite is limited to 30% or less, and the ferrite grain size number is No. 8-11 as defined in JIS G0552. In the longitudinal section in the steel matrix, Ti-based precipitation excellent hardening steel in preventing coarse grains characteristics and fatigue characteristics at carburizing the maximum diameter of the object is characterized in der Rukoto below 40 [mu] m. Furthermore, the chemical composition is mass% ,
Z r: 0.0003~0.0050%, Mg: 0.0003~0.0050% of one or preventing coarse grains characteristics and fatigue characteristics when carburizing according to claim 1, characterized in that it comprises two Excellent case hardening steel.   Furthermore, a chemical component is mass%, and contains Nb: less than 0.04%, The case hardening steel excellent in the coarse grain prevention characteristic and fatigue characteristic at the time of carburizing of Claim 1 or 2 characterized by the above-mentioned. Heating is carried out at a temperature of 1150 ° C. or higher for a heat retention time of 10 minutes or more and hot-rolled to a wire or a steel bar, the hot rolling finishing temperature is set to 840 to 1000 ° C., and the ferrite grain size number is specified in JIS G0552 The steel was No. 11, and after the hot rolling, the temperature range of 800 to 500 ° C. was gradually cooled at a cooling rate of 1 ° C./second or less, and the structure fraction of bainite after hot rolling was 30% or less. The method for producing a case hardening steel excellent in coarse grain prevention characteristics and fatigue characteristics during carburizing according to any one of claims 1 to 3, wherein the steel is made of steel. A carburized part obtained by processing the case-hardened steel according to any one of claims 1 to 3 into a part shape.

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