JP4941252B2 - Case-hardened steel for power transmission parts - Google Patents

Description

  The present invention relates to a case hardening steel for power transmission parts, and more particularly, to a case hardening steel for power transmission parts having excellent surface fatigue strength. More specifically, the present invention relates to a case hardening steel for power transmission parts having excellent pitching strength, spalling strength and low cycle bending fatigue strength.

  Conventionally, power transmission parts such as automobile gears, shafts, and constant velocity joints are formed by machining alloy steel materials for machine structures specified in JIS G 4053 (2003) into a predetermined shape by processing such as forging or cutting. It is manufactured by subjecting it to carburizing and quenching and carbonitriding and quenching.

  Conventionally, in the above parts, it has been emphasized to improve the strength against pitching, which is a kind of surface pressure fatigue.

  However, since it is said that the temperature of the contact surface at the time of use of the component rises to about 300 ° C., recently there has been a demand for suppressing temper softening on the surface layer at the time of use. Due to the increase, the stress applied to the power transmission component has also increased, so that in addition to the pitching strength, there has been a great demand for providing a spalling strength and a low cycle bending fatigue strength superior to those of the prior art.

  Therefore, in order to meet the above-described demand, for example, Patent Documents 1 to 3 propose steels to which Si is added as a material for the above parts.

  Specifically, Patent Document 1 has a specific chemical composition such as Si of 0.5 to 2.0% by mass%, [Q = 34140-605Si (%) + 183Mn (%) + 136Cr (%) + 122Mo (%)] A technique related to “skin-hardened steel” in which the value of activation energy Q of carbon diffusion in the steel defined by the formula is 34000 (kca1) or less and which can shorten the carburizing time as compared with SCr420 steel Has been.

  Patent Document 2 has a specific chemical composition such as Si of 0.35 to 1.3% by mass%, the austenite grain size of the carburized layer is 7 or more, and the surface carbon content is 0.9 to 0.9. A technique relating to “a carburized material excellent in rolling fatigue characteristics” of 1.5% and having a surface retained austenite amount of 25 to 40% is disclosed.

  Patent Document 3 has a specific chemical composition such as Si of 0.01 to 2.3% by mass%, and the precipitation amount of Nb (CN) after hot rolling is 0.005% or more, The amount of precipitation of AlN is limited to 0.015% or less, the structure fraction of bainite is 30% or less, the ferrite band score of the structure of the cross section parallel to the hot rolling direction is 1 to 5, and the hardness A technique relating to “high-temperature carburizing steel” in which HV is in a range defined by component parameters is disclosed.

JP 2005-36269 A JP 2000-54069 A JP 2001-279383 A

  In the case-hardened steel disclosed in Patent Document 1, B must be contained in order to obtain sufficient core hardness, and a sufficient hardening depth can be obtained when gas carburized and quenched. I can't. For this reason, in particular, when applied to power transmission parts used by gas carburizing and quenching, sufficient low cycle bending fatigue strength may not be obtained.

  The carburized material disclosed in Patent Document 2 pays attention to the surface layer of the bearing component in particular, and does not consider the core hardness, so that the core hardness is low when applied to a power transmission component. For this reason, sufficient low cycle bending fatigue strength may not be obtained.

  The steel for high-temperature carburizing disclosed in Patent Document 3 is a technique proposed under the technical idea that Ni is effective for imparting strength and hardenability to steel, like Cr and Mo. However, Ni is an element that suppresses the penetration of carbon during carburizing compared to Cr and Mo, and is disadvantageous for increasing the carburizing depth. Therefore, Ni is proposed under the technical idea as described above. Even if steel is used for power transmission parts, sufficient spalling strength cannot be obtained.

  As described above, the techniques proposed so far are only for the purpose of improving the structure after carburizing and part of strength characteristics such as increasing the carburizing depth. For this reason, when using conventional case-hardened steel as a raw material, it is possible to satisfy the demand for recent power transmission parts that want to have good characteristics in all of pitching strength, spalling strength and low cycle bending fatigue strength. could not.

  Therefore, the object of the present invention is to achieve the above three good by carburizing and quenching steel for power transmission parts capable of ensuring good characteristics in all of the pitching strength, spalling strength and low cycle bending fatigue strength, especially gas carburizing and quenching. It is to provide a case hardening steel for power transmission parts that can secure the characteristics.

  In order to improve the pitching strength of the power transmission component, generally, the Si content of the case-hardened steel as the material is increased for the purpose of increasing the softening resistance when the contact surface temperature is raised to about 300 ° C. during use. Is said to be effective.

  However, increasing the Si content has adverse effects as shown in the following (a) to (c).

  (A) Since the activity of C (carbon) in the steel increases, the carburization depth decreases.

(B) for A ₃ point of the steel is increased, the ferrite is generated in the core portion of the component during carburizing, microstructure after quenching becomes two phase structure of ferrite and martensite.

  (C) Oxides are generated at the grain boundaries on the surface during carburizing.

  As is clear from the above (a) to (c), in case hardening steel for power transmission parts, simply increasing the Si content is good in all of pitching strength, spalling strength and low cycle bending fatigue strength. It is difficult to satisfy the demand for recent power transmission components that want to have unique characteristics.

  That is, with respect to (a), when the carburization depth decreases, the hardness at the position of 0.3 mm depth from the surface corresponding to the starting position of the spalling failure decreases, so that hardening occurs from the inside before pitting occurs. Spalling that peels off the layer is likely to occur.

  Regarding (b), when soft ferrite is generated in the core, it is easily plastically deformed, so that the low cycle bending fatigue strength is significantly reduced.

  Regarding (c), since the grain boundary oxide layer becomes the starting point of the low cycle bending fatigue strength, the low cycle bending fatigue strength decreases.

  Therefore, for case hardening steel for power transmission parts, it is important not only to increase the Si content and improve the pitching strength, but also to design components that can increase the spalling strength and low cycle bending fatigue strength. It is.

  Therefore, the present inventors studied various chemical compositions for melting the various steels to increase the spalling strength and the low cycle bending fatigue strength.

That is, steels A to U having the chemical composition shown in Table 1 were melted in a 50 kg vacuum melting furnace to produce an ingot. In Table 1, the element symbol in the formula is expressed as the content in mass% of the element, and [Fn1 = Cr-2 × Si], [Fn2 = 959− ( 203 × C ^0.5 + 30 × Mn ) ] described later and The values of Fn1 to Fn3 represented by [Fn3 = 7 × Mn−10 × (Si + Cr−1.8)] are also shown.

  Steel A in Table 1 is a steel corresponding to SCM420 defined in JIS G 4053 (2003). Steel B and Steel C are based on steel A and have a carbon content of mass%. , 0.25% and 0.29%. Further, the steels D to U are obtained by changing the contents of C, Si, Mn, P, S, Cu, Ni, Cr, Mo, and N in various ways.

  The above ingot was heated to 1250 ° C. and subjected to hot forging to obtain a round bar having a diameter of 35 mm.

  Next, the round bar having a diameter of 35 mm was heated to 925 ° C., held for 60 minutes, and then subjected to a heat treatment for air cooling to room temperature.

  A four-point bending test piece having the shape shown in FIG. 1 is cut out from the center of the round bar having a diameter of 35 mm obtained in this way, and gas carburized and quenched under the conditions shown in FIG. % Was subjected to gas carburizing treatment at 930 ° C. and then quenched in oil at an oil temperature of 60 ° C. to investigate the influence of chemical composition on spalling strength and low cycle bending fatigue strength.

  In addition, the dimension of the test piece in FIG. 1 is a thing in mm unit, and the processing time of the gas carburizing in FIG. 2 was 180 min.

  In order not to cause spalling, the hardness at the position of 0.3 mm from the surface is desirably more than 700 in terms of Vickers hardness (hereinafter referred to as “Hv hardness”). The four-point bending test piece was crossed at a position of 20 mm from the end, and the Hv hardness (hereinafter also referred to as “0.3 mm position hardness”) at a depth of 0.3 mm from the surface was measured. In addition, a micro Vickers hardness meter was used for the measurement, the test force was set to 4.9N, three points were measured, and the arithmetic average was performed.

  Table 2 shows the measurement results of the 0.3 mm position hardness.

  Comparing the 0.3 mm position hardness of steels A to C in Table 2, it can be seen that the hardness increases as the C content of the base material increases. That is, as the C content of the base material increases, the C activity in the steel decreases, so the amount of C entering from the surface increases, and as a result, the 0.3 mm position hardness increases.

  Next, in order to examine the influence of alloy components other than C, the effect of lowering the activity of C in the steel with the same level of C content and the same level of 0.3 mm position hardness. Cr was marked and, conversely, Si that markedly increased the activity of C was arranged.

  Specifically, a steel having a C content of 0.25% and a 0.3 mm position hardness of more than 700 and not more than 720 in terms of Hv hardness, that is, steel B, steel E, steel M, steel N 3 and Steel P are arranged in FIG. 3 with the Cr content on the vertical axis and the Si content on the horizontal axis. In addition, “%” on the vertical axis and the horizontal axis in FIG. 3 represents “mass%”.

  As a result, it was found from FIG. 3 that the steel having a 0.3 mm position hardness of more than 700 and 720 or less in Hv hardness is plotted on a gradient of [Cr / Si = 2].

  That is, it was found that the 0.3 mm position hardness can be defined by the value of [Fn1 = Cr-2 × Si] if the C content is the same. Note that Cr and Si in the above formula of Fn1 are the amounts of Cr and Si in mass% contained in the base material, respectively.

  Then, next, regarding steel whose C content of the base material in mass% is 0.25%, that is, steel B, steel D to G, steel J to Q, and steel U, the 0.3 mm position hardness is plotted on the vertical axis. Further, the above [Cr-2 × Si] values are arranged in FIG. In FIG. 4, “%” on the horizontal axis represents “mass%”.

  As a result, from FIG. 4, the 0.3 mm position hardness and the value of [Cr−2 × Si] are substantially proportional, and the 0.3 mm position hardness, that is, the Hv hardness at the depth of 0.3 mm from the surface. In order to exceed 700, it was found that the value of [Fn1 = Cr-2 × Si] should exceed 0.4.

  Regarding the low cycle bending fatigue strength, the following investigation was conducted.

That is, using the four-point bending test piece subjected to gas carburizing and quenching, a constant bending load was repeatedly applied under the conditions of a stress ratio of 0.1 and a frequency of 5 Hz so that a tensile stress was always applied to the notch. The number of repetitions when the test piece broke was investigated. Then, an SN curve was obtained by performing a test in which the value of the bending load was changed, and the bending load breaking at 10 ⁴ times was evaluated as the low cycle bending fatigue strength.

  Further, a four-point bending test piece subjected to gas carburizing and quenching was traversed at a position 20 mm from the end, and the core hardness was measured and the microstructure was observed. The core hardness was measured by using a micro Vickers hardness meter, measuring three points with a test force of 4.9 N, and calculating the arithmetic average.

  Table 2 shows the core hardness and low cycle bending fatigue strength measured as described above.

  FIG. 5 shows the relationship between the core hardness and the low cycle bending fatigue strength for each steel type shown in Table 1.

  From the comparison of steels A to C in Table 2, as with the 0.3 mm position hardness, the core hardness becomes harder as the C content of the base material increases, and the low cycle bending fatigue strength increases accordingly. I understand that

  Further, from FIG. 5, it can be seen that the low cycle bending fatigue strength generally has a linear relationship with the core hardness except for the three points “white triangle” and “black square”. That is, in order to improve the low cycle bending fatigue strength, it is necessary to increase the hardness of the core.

  However, as described above, there are three points in FIG. 5 that deviate from the linear relationship. Therefore, the test piece whose core hardness was measured was mirror-polished and then corroded with nital to observe the microstructure of the core. did. Table 2 also shows the results of observation of the microstructure of the core.

  When the surface layer was further observed, precipitates were observed along the grain boundaries in Steel U. For this reason, the said site | part was analyzed using EPMA (electron beam microanalyzer). As a result, it was confirmed that C and Cr were concentrated in the precipitate, and it was found to be Cr carbide.

  As a result of these micro observations, the following matters became clear.

  First, in FIG. 5, what is indicated by “black squares” is steel U, which was broken at the same time as cracks were generated, and Cr carbide was generated at the grain boundaries as described above.

  That is, when Cr carbide is generated at the grain boundary, the low cycle bending fatigue strength is significantly reduced. Therefore, it is necessary to suppress the formation of Cr carbide.

  In addition, the production | generation of Cr carbide is not dependent only on Cr content in steel, but is influenced also by Si content. That is, since Si has the effect of reducing the surface C concentration during carburizing by increasing the activity of C in the steel, Si has the effect of suppressing the formation of Cr carbides. Such Cr carbide is recognized in the steel U having the value of [Cr-2 × Si] of 2.00 and the steel O having a value of [Cr-2 × Si] of 1.50. It was found that the value of [Fn1 = Cr-2 × Si] needs to be 1.8 or less.

  Secondly, what is indicated by “white triangles” in FIG. 5 is steel H and steel T, and the core structure is a two-phase structure of ferrite and martensite. That is, when soft ferrite is generated in the core, the low cycle bending fatigue strength is significantly reduced. Therefore, it is necessary to suppress the formation of ferrite in the core.

In addition, the ferrite of a core part arises from the austenite single phase area | region during a carburizing process, becoming a two phase area | region of a ferrite and austenite. Therefore, in order to suppress the formation of ferrite, it is sufficient to lower the _three point A of the steel.

Si, which is an important element of the present invention, is an element that raises the A ₃ point. Therefore, even when the Si content is high, it is necessary to A ₃ point to component design preform so as not to exceed the carburizing temperature.

Therefore, Andrews Iron Steel Inst. 203 (1965), p. Based on equation A ₃ points are reported in 721, captures the effects of Si on the constant term, for the Table 1 Steel, calculate the value of ^{[Fn2 = 959- (203 × C 0.5} + 30 × Mn) ] did. Note that C and Mn in the above formula of Fn2 are respectively the amounts of C and Mn in mass% contained in the base material.

As a result, the ferrite of the core part was recognized in the steel H of 856 and the steel T of 850 having the value of [959− ( 203 × C ^0.5 + 30 × Mn ) ], and [959− ( 203 × C ^0.5 + 30 ×]. Since the value of Mn ) ] was not observed in steel A of 844, it was found that the value of Fn2 should be less than 850 in order to suppress the formation of ferrite in the core.

  Thirdly, as shown in FIG. 5, even when the low cycle bending fatigue strength and the core hardness show a linear relationship, it is clear that they are roughly divided into two groups above and below the least square line as a boundary. It was.

  Accordingly, when the steel types having a linear relationship between the low cycle bending fatigue strength and the core hardness were investigated in more detail, it was found that the grain boundary oxide depth of the surface layer affects the low cycle bending fatigue strength.

  That is, in FIG. 5, the “white circle” mark is a steel having a grain boundary oxide layer depth of 20 μm or less, while the “black circle” mark is a steel having a grain boundary oxide layer depth of more than 20 μm. From FIG. 5, the steel having a grain boundary oxide layer depth of 20 μm or less and the steel having a grain boundary oxide layer depth of more than 20 μm are roughly divided above and below the straight line with a least square line as a boundary. It has been clarified that the low cycle bending fatigue strength tends to decrease as the depth increases, and that a higher low cycle bending fatigue strength can be obtained when the grain boundary oxide layer depth is 20 μm or less.

  Therefore, next, 16 steel types were melted and the same treatment as described above was performed, and various relationships between the grain boundary oxide layer depth and the contents of Si, Cr and Mn were investigated.

  As a result, it turned out that it can arrange | position like FIG. 6 and FIG. In FIG. 6 and FIG. 7, “%” on the horizontal axis represents “mass%”.

That is, when the value of [Si + Cr], which is the sum of the contents of Si and Cr in mass%, is in the range of 0.5 to 3.5, as shown in FIG. If it becomes shallower and the grain boundary oxide layer depth is y,
y = −10 × (Si + Cr ) +19
In addition, when the content of Mn in mass% is in the range of 0.3 to 2.2%, as shown in FIG. 7, the grain boundary oxide layer depth is deep with the content. If the grain boundary oxide layer depth is y,
y = 7 × Mn−1
It was found that it can be expressed by the following formula.

  Therefore, when the above results are linearly combined and the relationship with the grain boundary oxide layer depth is arranged with the parameter [Fn3 = 7 × Mn−10 × (Si + Cr−1.8)], Si is 0.25% or more. In this range, it became as shown in FIG.

  In FIG. 8, the “white circle” mark indicates that the grain boundary oxide layer depth is 20 μm or less, and the “black circle” mark indicates that the grain boundary oxide layer depth exceeds 20 μm. From FIG. 8, it was found that by setting the value of Fn3 to 8 or less, the grain boundary oxide layer depth can be set to 20 μm or less, and the low cycle bending fatigue strength can be improved.

    This invention is completed based on said knowledge, The summary exists in the case hardening steel for power transmission components shown to following (1)-(3).

(1) By mass%, C: 0.15 to 0.35%, Si: 0.40 to 1.10%, Mn: 0.50 to 1.50%, S: 0.01 to 0.05% , Cr: 1.20 to 2.60%, Al: 0.010 to 0.050% and N: 0.010 to 0.025%, and the contents of C, Si, Mn and Cr are The values of Fn1 to Fn3 represented by the following formulas (1) to (3) satisfy 0.4 <Fn1 ≦ 1.8, Fn2 <850 and Fn3 ≦ 8, respectively, and the balance is made of Fe and impurities. P, Cu, Ni and V in the impurities are respectively P: 0.05% or less, Cu: 0.10% or less, Ni: 0.10% or less and V: 0.005% or less. Hardened steel for power transmission parts.
Fn1 = Cr-2 × Si (1),
Fn2 = 959− ( 203 × C ^0.5 + 30 × Mn ) (2),
Fn3 = 7 * Mn-10 * (Si + Cr-1.8) ... (3).
Here, the element symbol in the formulas (1) to (3) represents the content in mass% of the element.

  (2) The case-hardened steel for power transmission parts as described in (1) above, which contains Mo: 0.40% or less in mass% instead of part of Fe.

  (3) The above (1), characterized in that, instead of a part of Fe, by mass%, one or two of Ti: 0.10% or less and Nb: 0.10% or less are contained. Or the case hardening steel for power transmission components as described in (2).

  Hereinafter, the inventions related to the case-hardening steel for power transmission parts (1) to (3) are referred to as “present invention (1)” to “present invention (3)”, respectively. Also, it may be collectively referred to as “the present invention”.

  By using the case-hardened steel for power transmission parts of the present invention, a power transmission part having good characteristics in all of pitching strength, spalling strength and low cycle bending fatigue strength can be obtained.

  Hereinafter, each requirement of the present invention will be described in detail. In addition, “%” of the content of the chemical component means “mass%”.

C: 0.15-0.35%
C is an important element for improving the spalling strength and the low cycle bending fatigue strength by increasing the carburizing depth and further increasing the hardness at the 0.3 mm position and the core hardness during carburizing and quenching. However, if the C content is less than 0.15%, the effect of increasing the carburization depth and improving the low cycle bending fatigue strength is small. On the other hand, if the content exceeds 0.35%, the toughness of the core portion decreases, so that fracture occurs at the same time as cracking, and low cycle bending fatigue strength deteriorates. Therefore, the content of C is set to 0.15 to 0.35%. A desirable range for the C content is 0.20 to 0.30%.

  In addition, in the content of C in the above range, the value of Fn2 represented by the above formula (2) needs to satisfy Fn2 <850.

Si: 0.40 to 1.10%
Si is an important element for increasing the temper softening resistance at about 300 ° C. When the Si content is less than 0.40%, the above effect is small and the surface pressure strength (pitching strength) is lowered. However, since Si is also an element that increases the activity of C in the austenite during carburizing, if it exceeds 1.10%, the carburizing property is remarkably deteriorated. Therefore, the Si content is set to 0.40 to 1.10%. The desirable range of Si content is 0.45 to 0.90%.

  The Si content is within the above range, and the values of Fn1 and Fn3 represented by the above formulas (1) and (3) satisfy 0.4 <Fn1 ≦ 1.8 and Fn3 ≦ 8, respectively. It is necessary to satisfy.

Mn: 0.50 to 1.50%
Mn is an element that lowers the activity of C in the austenite during carburizing and promotes carburizing, and at the same time forms MnS and increases machinability. However, when the Mn content is less than 0.50%, the above effect is small. On the other hand, when Mn exceeds 1.50%, not only the above effects are saturated, but also the hardness after normalization increases, and as a result, the machinability is significantly deteriorated. Therefore, the Mn content is set to 0.50 to 1.50%. A desirable range of the Mn content is 0.60 to 1.20%.

  In addition, the content of Mn needs to satisfy Fn2 <850 and Fn3 ≦ 8, respectively, in the values within the above range, the values of Fn2 and Fn3 represented by the above formulas (2) and (3).

S: 0.01 to 0.05%
S is an essential element for forming MnS together with Mn. MnS formed in a state where a certain amount of Mn is contained is necessary for ensuring machinability. In order to obtain this effect, the S content needs to be 0.01% or more. However, when the S content is excessive, the hot ductility is lowered, and particularly when the content exceeds 0.05%, the hot ductility is significantly lowered. Therefore, the content of S is set to 0.01 to 0.05%. In addition, it is preferable that content of S shall be 0.01-0.03%.

Cr: 1.20 to 2.60%
Cr is an element that lowers the activity of C in the austenite during carburizing and promotes carburizing. However, when the Cr content is less than 1.20%, the above effect is small. On the other hand, when Cr is contained in excess of 2.60%, it becomes easy to produce Cr carbide at the grain boundary during carburizing, and lowers the low cycle bending fatigue strength. Therefore, the content of Cr is set to 1.20 to 2.60%. A desirable range of the Cr content is 1.30 to 2.30%, and a more desirable range is 1.50 to 2.00%.

  In addition, Cr content needs to satisfy | fill 0.4 <Fn1 <= 1.8 also in Fn1 value represented by said Formula (1) in said range.

Al: 0.010 to 0.050%
Al combines with N in steel to form AlN, and has the effect of improving the toughness of the carburized layer by refining crystal grains during carburization. In order to obtain this effect, the Al content needs to be at least 0.010%. However, if the Al content exceeds 0.050%, the machinability deteriorates due to the formation of hard Al ₂ O ₃ . Therefore, the content of Al is set to 0.010 to 0.050%. In addition, it is preferable that content of Al shall be 0.010 to 0.040%.

N: 0.010 to 0.025%
N combines with Al to form AlN and refines crystal grains during carburizing, thereby improving the toughness of the carburized layer. In order to obtain this effect, N must have a content of at least 0.010%. However, even if the content exceeds 0.025%, the above effect is saturated, and a so-called “nest” is easily formed in the interior during melting. Therefore, the N content is set to 0.010 to 0.025%. A desirable range of the N content is 0.012 to 0.020%.

Fn1 value: more than 0.4 and 1.8 or less As described above, Si is an important element for increasing the temper softening resistance, and Cr is an important element for promoting carburization. However, when the value of Fn1 represented by the above formula (1) is 0.4 or less, the carburizing property is remarkably deteriorated. Therefore, the 0.3 mm position hardness is set to 700 or more in terms of Hv hardness. It is not possible to reduce the spalling strength. On the other hand, if the value of Fn1 exceeds 1.8, the amount of Cr carbide generated at the grain boundaries increases, leading to a decrease in low cycle bending fatigue strength. Therefore, the value of Fn1 represented by the above equation (1), that is, [Fn1 = Cr−2 × Si] is set to satisfy 0.4 <Fn1 ≦ 1.8. Note that the value of Fn1 is desirably 0.5 ≦ Fn1 ≦ 1.5, and a more desirable range is 0.6 ≦ Fn1 ≦ 1.2.

Fn2 values: the content of 850 less than C, Si and Mn affects _three points A steel. Since the A ₃ point is a ferrite phase is generated in the core during the carburizing to rise, after quenching is also the core portion is a two-phase structure of ferrite and martensite, low cycle bending fatigue strength deteriorates. The present invention, since the active use of Si having an effect of increasing temper softening resistance which increases the _three points A, for the control of A _3-point, it is necessary to control the content of C and Mn . And when the value of Fn2 represented by said Formula (2) is 850 or more, a ferrite produces | generates in a core part and low cycle bending fatigue strength falls. For this reason, the value of Fn2 represented by the above formula (2), that is, [Fn2 = 959− ( 203 × C ^0.5 + 30 × Mn ) ] is determined to satisfy Fn2 <850. The value of Fn2 is preferably Fn2 ≦ 840.

Fn3 value: 8 or less Mn, Si and Cr all affect the grain boundary oxide layer depth during gas carburization. If the value of Fn3 represented by the above formula (3) exceeds 8, the grain boundary oxide layer depth becomes deep due to the effect of Mn, and the low cycle bending fatigue strength deteriorates. Therefore, the value of Fn3 represented by the above equation (3), that is, [Fn3 = 7 × Mn−10 × (Si + Cr−1.8)] satisfies Fn3 ≦ 8. The value of Fn3 is preferably Fn3 ≦ 5.0.

  In the present invention, it is necessary to limit the contents of P, Cu, Ni, and V in the impurity to a specific value or less. This will be described below.

P: 0.05% or less In the present invention, the content of P as an impurity must be suppressed to 0.05% or less. That is, P causes a decrease in hot ductility, and particularly when its content exceeds 0.05%, the decrease in hot ductility becomes significant. Therefore, in the present invention, the content of P in the impurities is set to 0.05% or less. In addition, it is preferable that content of P in an impurity shall be 0.03% or less.

Cu: 0.10% or less In the present invention, the content of Cu as an impurity must be suppressed to 0.10% or less. Cu raises the activity of C in the austenite at the time of carburizing, prevents an increase in 0.3 mm position hardness, and lowers the spalling strength. In particular, when the Cu content exceeds 0.10%, the influence is remarkable. Therefore, in the present invention, the content of Cu in the impurities is set to 0.10% or less. In addition, it is preferable that content of Cu in an impurity shall be 0.05% or less.

Ni: 0.10% or less In the present invention, Ni must be contained as an impurity in the same manner as Cu, and its content must be suppressed to 0.10% or less. Ni raises the activity of C in the austenite at the time of carburizing, prevents an increase in 0.3 mm position hardness, and lowers the spalling strength. In particular, when the content exceeds 0.10%, the influence is remarkable. Therefore, in the present invention, the content of Ni in the impurities is set to 0.10% or less. In addition, it is preferable that content of Ni in an impurity shall be 0.05% or less.

V: 0.005% or less In the present invention, the content of V as an impurity must be suppressed to 0.005% or less. That is, V combines with C and N to form VC and VCN, and these carbides and carbonitrides are effective in refining austenite grains at the beginning of the temperature raising process of the carburizing process, but on the contrary The crystal grains become too small and the driving force for crystal grain growth becomes too large, causing abnormal grain growth. In particular, when the V content exceeds 0.005%, abnormal grain growth becomes remarkable, and the low cycle bending fatigue strength decreases. Therefore, in the present invention, the content of V in the impurities is set to 0.005% or less.

  For the above reasons, the case-hardening steel for power transmission parts according to the present invention (1) has C: 0.15-0.35%, Si: 0.40-1.10%, Mn: 0.50-1 .50%, S: 0.01 to 0.05%, Cr: 1.20 to 2.60%, Al: 0.010 to 0.050% and N: 0.010 to 0.025% In addition, the contents of C, Si, Mn, and Cr are 0.4 <Fn1 ≦ 1.8, Fn2 <850, and Fn1 to Fn3 represented by the above formulas (1) to (3), respectively. Fn3 ≦ 8 is satisfied, the balance is made of Fe and impurities, and P, Cu, Ni and V in the impurities are respectively P: 0.05% or less, Cu: 0.10% or less, Ni: 0.10% or less And V: specified to be 0.005% or less.

In addition, the case-hardened steel for power transmission parts according to the present invention (1) is replaced with a part of the Fe, if necessary,
First group: Mo: 0.40% or less,
2nd group: One or two of Ti: 0.10% or less and Nb: 0.10% or less,
One or more elements of each group can be selectively contained.

  That is, one or more of the elements of either the first group or the second group may be included as an arbitrary element.

  Hereinafter, the above optional elements will be described.

First group: Mo: 0.40% or less Mo, which is an element of the first group, has an effect of improving the hardenability of the carburized layer and improving the surface pressure strength. In order to obtain this effect, the Mo content is preferably 0.05% or more. However, when the Mo content exceeds 0.40%, the reduction of the machining becomes remarkable. Therefore, the Mo content is set to 0.40% or less. When Mo is contained, the amount of Mo is preferably 0.05 to 0.40%, and more preferably 0.05 to 0.30%.

2nd group: Ti: 0.10% or less and Nb: 0.10% or less, and 2nd group elements Ti and Nb have the effect of increasing the toughness of the carburized layer. In order to acquire this effect, you may contain said element. Hereinafter, the second group of elements will be described in detail.

Ti: 0.10% or less Ti combines with C and N in steel to form carbonitrides, which, like AlN, is effective in refining crystal grains during carburizing and improves the toughness of the carburized layer There is an action to make. In order to obtain such an effect, the Ti content is preferably 0.005% or more. However, when the Ti content exceeds 0.10%, hard TiN is generated, and the machinability is significantly lowered. Therefore, the Ti content is set to 0.10% or less. When Ti is contained, the amount of Ti is preferably 0.005 to 0.10%.

Nb: 0.10% or less Nb also combines with C and N in steel to form carbonitrides, like Ti, and has the effect of refining crystal grains during carburization, similar to AlN. It has the effect of improving the toughness of the. In order to obtain such effects, the Nb content is preferably 0.005% or more. However, when the Nb content exceeds 0.10%, coarse precipitates are generated, and abnormal grain growth tends to occur. Therefore, the Nb content is 0.10% or less. The content of Nb when contained is preferably 0.005 to 0.10%.

  In addition, said Ti and Nb can be contained only in any 1 type of them, or 2 types of composites.

  For the above reasons, the case hardening steel for power transmission parts according to the present invention (2) is replaced with a part of Fe of the case hardening steel for power transmission parts according to the present invention (1). It was specified that it contained Mo.

  Further, the case hardening steel for power transmission parts according to the present invention (3) is replaced with a part of Fe of the case hardening steel for power transmission parts according to the present invention (1) or the present invention (2). It was defined to contain one or two of Ti and Nb which are group elements.

  Hereinafter, the present invention will be described in more detail with reference to examples.

  Steels 1 to 22 having chemical compositions shown in Table 3 were melted in a 50 kg vacuum melting furnace to produce ingots.

  In addition, the steels 1-8 in Table 3 are steels having a chemical composition within the range defined by the present invention. On the other hand, steels 9 to 22 are steels of comparative examples whose chemical compositions deviate from the conditions specified in the present invention. Among the steels of this comparative example, steel 9 is SCM420 defined by JIS G 4053 (2003).

  Each ingot was subjected to a treatment for 30 minutes at 1250 ° C., and then hot forged to obtain a round bar having a diameter of 35 mm. In addition, the finishing temperature of hot forging was made not to drop below 1000 ° C., and cooling after hot forging was allowed to cool in the air.

  Next, the above-mentioned round bar having a diameter of 35 mm is heated to 1250 ° C., held for 60 minutes, and then subjected to a heat treatment that is allowed to cool to room temperature in the atmosphere, thereby suppressing microsegregation, and further heated to 925 ° C. After maintaining for 120 minutes, it was cooled to room temperature at a cooling rate of 10 to 30 ° C./min, and the microstructure was changed to a mixed structure of ferrite and pearlite.

  From the central part of the round bar having a diameter of 35 mm obtained as described above, a small roller test piece used in the two-cylinder rolling fatigue test having the shape shown in FIG. 9, a four-point bending test piece having the shape shown in FIG. The indicated round bar specimen was cut out. In addition, the dimension of the test piece in FIG. 9 and FIG. 10 is also a unit of mm.

  Each of the above test pieces was subjected to gas carburizing and quenching under the heat treatment conditions shown in FIG. 11 after grinding the test surface. In FIG. 11, “1.0% C” and “0.8% C” indicate that the carbon potential is 1.0% and 0.8%, respectively, and “80 ° C. OQ” indicates the oil temperature. It represents quenching in 80 ° C. oil.

  The small roller test piece produced as described above was subjected to a two-cylinder rolling fatigue test under the conditions shown in Table 4, and the surface pressure fatigue strength was investigated.

  In addition, as a large roller test piece used for the two-cylinder rolling fatigue test, SCM822 defined in JIS G 4053 (2003) was machined, then gas carburized and quenched, and the surface layer was ground by 50 microns.

  Specifically, the material was hot forged to a diameter of 150 mm, heated to 1250 ° C., held for 60 min, and then subjected to a heat treatment that was allowed to cool to room temperature in the air, thereby suppressing microsegregation. Thereafter, the mixture was further heated to 925 ° C. and held for 120 minutes, and then cooled to room temperature at a cooling rate of 10 ° C./min, so that the structure was a mixed structure of ferrite and pearlite.

  Next, the material subjected to the above treatment was machined into a roller having a diameter of 130 mm and a width of 20 mm with a crowning with a radius of 150 mm. The roller was subjected to gas carburizing treatment at 930 ° C. with a target of the total carburizing depth of 1.5 mm, and then quenched into oil at an oil temperature of 60 ° C. Thereafter, the crowning surface was polished by 50 μm to finish a large roller test piece.

  In addition, as shown in Table 4 above, the surface pressure fatigue strength is determined under the condition that the load during the test is constant under the condition that the rotation speed of the small roller is 1000 rpm and the slip ratio of the large roller is 80%. A cylindrical rolling fatigue test was performed. At this time, a commercially available ATF (automatic transmission oil) was discharged from the reverse direction of rotation of the test piece to the contact portion under the conditions of an oil temperature of 40 ° C. and 2 liters / min.

In the test, the surface pressure was set to 2800 MPa, and the two-cylinder rolling fatigue test was continued until fatigue peeling occurred, or until fatigue peeling did not occur until 2.0 × 10 ⁷ times. Neither spalling nor spalling fatigue strength is assumed. Pitching and spalling were distinguished from the cross-sectional shape of the fracture surface. That is, the surface where the peeling surface is slanted in the depth direction from the surface is called pitching, and the surface where the peeling surface is almost parallel to the original surface is called spalling.

In addition, using a four-point bending test piece, a four-point bending fatigue test was conducted under conditions of a stress ratio of 0.1 and a repetition rate of 5 Hz so that a tensile stress was always applied to the notch, and low cycle bending fatigue characteristics were investigated. Then, 10 ⁴ times bending fatigue strength was determined from the SN diagram.

The target of the bending fatigue strength The 10 ⁴ times as above 1000MPa, 10 ⁴ times bending fatigue strength in the case of more than 1000MPa to the target, was excellent in low cycle bending fatigue strength.

  Further, the round bar test piece was crossed at a position of 20 mm from the end, and the 0.3 mm position hardness and the core hardness were measured. In addition, a micro Vickers hardness meter was used for the measurement, the test force was set to 4.9 N, and three points each were measured and arithmetically averaged.

  Further, after re-polishing the test surface after the hardness measurement, the corrosion time in nital was changed, and the grain boundary oxide layer depth was measured and the microstructure of the core was observed.

  Table 5 shows the results of the above investigations.

  From Table 5, it is clear that in the case of Test Nos. 1 to 8 that satisfy the conditions of the present invention, pitching and spalling do not occur, the surface fatigue strength is excellent, and the low cycle bending fatigue strength is also excellent.

  On the other hand, in the case of test numbers 9 to 22 of comparative examples that deviate from the conditions defined in the present invention, the target of the present invention has not been reached.

In the case of test number 9, the value of Fn1 of steel 9, that is, the value of [Cr-2 × Si] is less than the range specified in the present invention, so that the carburizing property is deteriorated and the 0.3 mm position hardness is reduced to 690. Thus, the contact pressure fatigue form becomes spalling, and the two-cylinder rolling fatigue life is as short as 1.6 × 10 ⁷ and does not reach 2.0 × 10 ⁷ times. Further, since the value of Fn3, that is, the value of [7 × Mn−10 × (Si + Cr−1.8)] exceeded the range defined in the present invention, the grain boundary oxide layer depth became as deep as 24 μm, so that the low cycle The bending fatigue strength deteriorates and is as low as 930 MPa, and the target of 1000 MPa or more is not reached.

In the case of test number 10, the C content of steel 10 is below the range defined in the present invention, and the value of Fn2, that is, the value of [959− ( 203 × C ^0.5 + 30 × Mn ) ] is defined in the present invention. From this, ferrite is generated in the core, and in addition, the value of Fn3, that is, the value of [7 × Mn−10 × (Si + Cr−1.8)] exceeds the range specified in the present invention, and grain boundary oxidation. Since the layer depth is as deep as 22 μm, the low cycle bending fatigue strength deteriorates, is as low as 880 MPa, and does not reach the target of 1000 MPa or more. Furthermore, since the value of Fn1, that is, the value of [Cr-2 × Si] is also below the range specified in the present invention, the carburizing property is deteriorated, the 0.3 mm position hardness is reduced to 680, and the surface pressure fatigue mode is reduced. As a result of spalling, the two-cylinder rolling fatigue life was as short as 1.1 × 10 ⁶ and did not reach 2.0 × 10 ⁷ times.

  In the case of test number 11, since the C content of steel 11 exceeds the range specified in the present invention, the fracture occurred at the same time as crack generation, the low cycle bending fatigue strength deteriorated, the low 950 MPa, and the target of 1000 MPa or more was reached. Not.

In the case of test number 12, since the value of Fn1 of steel 12, that is, the value of [Cr-2 × Si] is below the range specified in the present invention, the 0.3 mm position hardness is as low as 650, and the surface pressure fatigue mode The two-cylinder rolling fatigue life is as short as 9.0 × 10 ⁶ and does not reach 2.0 × 10 ⁷ times.

In the case of test number 13, since the Si content of steel 13 is below the range specified in the present invention, the pitching strength is deteriorated, the two-cylinder rolling fatigue life is as short as 9.8 × 10 ^6, and 2.0 × 10 Has not reached ⁷ times. Moreover, since the value of Fn3 of steel 13, that is, the value of [7 × Mn−10 × (Si + Cr−1.8)] exceeds the range defined in the present invention, the grain boundary oxide layer depth becomes as deep as 25 μm. It was. For this reason, the low cycle bending fatigue strength is also deteriorated, being as low as 900 MPa and not reaching the target of 1000 MPa or more.

In the case of test number 14, in addition to the Mn content of steel 14 being below the range specified in the present invention, the value of Fn2, that is, the value of [959− ( 203 × C ^0.5 + 30 × Mn ) ] is specified in the present invention. Therefore, ferrite is generated in the core portion, so that the low cycle bending fatigue strength is deteriorated, is as low as 900 MPa, and does not reach the target of 1000 MPa or more. Moreover, since the value of Fn1 of steel 14, that is, the value of [Cr-2 × Si] is below the range defined in the present invention, the 0.3 mm position hardness is reduced to 690, and the surface pressure fatigue mode becomes spalling. The two-cylinder rolling fatigue life is as short as 9.5 × 10 ^6, and has not reached 2.0 × 10 ⁷ times.

  In the case of test number 15, since the value of Fn1 of steel 15, that is, the value of [Cr-2 × Si] exceeds the range specified in the present invention, Cr carbide is generated at the grain boundary, and fracture occurs simultaneously with crack generation. However, the low cycle bending fatigue strength deteriorates and is as low as 980 MPa and does not reach the target of 1000 MPa or more.

In the case of test number 16, since the value of Fn1 of steel 16, that is, the value of [Cr-2 × Si] is below the range defined in the present invention, the 0.3 mm position hardness is reduced to 673, and the surface pressure The fatigue form is spalling, and the two-cylinder rolling fatigue life is as short as 9.0 × 10 ⁶ and does not reach 2.0 × 10 ⁷ times. Further, since the value of Fn3 of steel 16, that is, the value of [7 × Mn-10 × (Si + Cr−1.8)] exceeds the range defined in the present invention, the grain boundary oxide layer depth becomes as deep as 30 μm. Thus, the low cycle bending fatigue strength deteriorates, is as low as 950 MPa, and does not reach the target of 1000 MPa or more.

In the case of test number 17, in addition to the Cu content of steel 17 exceeding the range specified by the present invention, the value of Fn1, that is, the value of [Cr-2 × Si] is below the range specified by the present invention. The carburizability deteriorates and the 0.3 mm position hardness is as low as 680. As a result, the contact pressure fatigue form becomes spalling, the two-cylinder rolling fatigue life is as short as 1.1 × 10 ^7, and has not reached 2.0 × 10 ⁷ times.

  In the case of test number 18, since the V content of steel 18 exceeds the range specified in the present invention, abnormal grain growth occurs during carburizing and quenching, the low cycle bending fatigue strength deteriorates, and the target is as low as 950 MPa and 1000 MPa or more. Not reached.

In the case of test number 19, since the value of Fn1 of steel 19, that is, the value of [Cr-2 × Si] is below the range defined in the present invention, the carburization depth is shallow, and the surface pressure fatigue mode becomes spalling, The two-cylinder rolling fatigue life is as short as 6.3 × 10 ⁶ and does not reach 2.0 × 10 ⁷ times.

In the case of the test number 20, the value of Fn2 of the steel 20, that is, the value of [959− ( 203 × C ^0.5 + 30 × Mn ) ] exceeds the range specified in the present invention, so that ferrite is generated in the core portion. The low cycle bending fatigue strength deteriorates, is as low as 890 MPa, and does not reach the target of 1000 MPa or more. Moreover, since the value of Fn1 of steel 20, that is, the value of [Cr-2 × Si] is below the range defined by the present invention, the 0.3 mm position hardness is reduced to 690, and the surface pressure fatigue mode is spalling. Thus, the two-cylinder rolling fatigue life is as short as 9.4 × 10 ⁶ and does not reach 2.0 × 10 ⁷ times.

In the case of the test number 21, the value of Fn3 of the steel 21, that is, the value of [7 × Mn−10 × (Si + Cr−1.8)] exceeds the range specified in the present invention, so that the grain boundary oxide layer is 26 μm. For this reason, the low cycle bending fatigue strength deteriorates and is as low as 880 MPa and does not reach the target of 1000 MPa or more. Moreover, since the value of Fn1 of steel 21, that is, the value of [Cr-2 × Si] is below the range defined in the present invention, the 0.3 mm position hardness is lowered to 685, and the surface pressure fatigue mode is spalling. Thus, the two-cylinder rolling fatigue life is as short as 9.0 × 10 ⁶ and does not reach 2.0 × 10 ⁷ times.

In the case of the test number 22, since the Ni content of the steel 22 exceeds the range specified in the present invention, the carburization property is deteriorated and the 0.3 mm position hardness is as low as 690. As a result, the contact pressure fatigue form becomes spalling, the two-cylinder rolling fatigue life is as short as 1.0 × 10 ^7, and has not reached 2.0 × 10 ⁷ times.

  By using the case-hardened steel for power transmission parts of the present invention, a power transmission part having good characteristics in all of pitching strength, spalling strength and low cycle bending fatigue strength can be obtained.

It is a figure which shows the shape of a four-point bending test piece. The unit of the test piece dimension is “mm”. It is a figure which shows the conditions of the gas carburizing quenching given to the four-point bending test piece of steel AU. In the figure, “0.8%” indicates that the carbon potential is 0.8%, and “60 ° C. OQ” indicates that the carbon is quenched in oil at an oil temperature of 60 ° C. For steel B, steel E, steel M, steel N and steel P with a C content of the base metal of 0.25% and a 0.3 mm position hardness of more than 700 and 720 or less in Hv hardness, its Cr content It is the figure which arranged the quantity on the vertical axis and arranged Si content on the horizontal axis. “%” On the vertical axis and horizontal axis in the figure represents “mass%”. For steel B, steel D to G, steel J to Q, and steel U with a C content of the base metal of 0.25%, the hardness of 0.3 mm is plotted on the vertical axis, and the value of [Cr-2Si] is plotted horizontally It is the figure arranged for the axis. “%” On the horizontal axis in the figure represents “mass%”. It is a figure which arranges and shows the relation between core part hardness and low cycle bending fatigue strength. It is a figure which arranges and shows the relation between the value of [Si + Cr] and the grain boundary oxide layer depth. “%” On the horizontal axis in the figure represents “mass%”. It is a figure which rearranges and shows the relationship between content of Mn and the grain boundary oxide layer depth. “%” On the horizontal axis in the figure represents “mass%”. It is a figure which rearranges and shows the relationship between the value of [Fn3 = 7 * Mn-10 * (Si + Cr-1.8)] and the grain boundary oxide layer depth. It is a figure which shows the shape of the small roller test piece used for the two-cylinder rolling fatigue test of an Example. The unit of the test piece dimension is “mm”. It is a figure which shows the shape of the round bar test piece cut out in the Example. The unit of the test piece dimension is “mm”. It is a figure which shows the conditions of the gas carburizing quenching implemented in the Example. In the figure, “1.0% C” and “0.8% C” indicate that the carbon potential is “1.0%” and “0.8%”, respectively, and “80 ° C. OQ” indicates oil. It represents quenching in oil at a temperature of 80 ° C.

Claims (3)

In mass%, C: 0.15 to 0.35%, Si: 0.40 to 1.10%, Mn: 0.50 to 1.50%, S: 0.01 to 0.05%, Cr: 1.20 to 2.60%, Al: 0.010 to 0.050% and N: 0.010 to 0.025%, and the contents of C, Si, Mn and Cr are the following ( The values of Fn1 to Fn3 represented by the formulas 1) to (3) satisfy 0.4 <Fn1 ≦ 1.8, Fn2 <850 and Fn3 ≦ 8, respectively, and the balance is composed of Fe and impurities. P, Cu, Ni, and V are P: 0.05% or less, Cu: 0.10% or less, Ni: 0.10% or less, and V: 0.005% or less, respectively. Case-hardened steel for parts.
Fn1 = Cr-2 × Si (1)
Fn2 = 959− ( 203 × C ^0.5 + 30 × Mn ) (2)
Fn3 = 7 × Mn−10 × (Si + Cr−1.8) (3)
Here, the element symbol in the formulas (1) to (3) represents the content in mass% of the element.   The case-hardening steel for power transmission parts according to claim 1, characterized by containing Mo: 0.40% or less in mass% instead of part of Fe.   3. Instead of a part of Fe, it contains one or two of Ti: 0.10% or less and Nb: 0.10% or less in mass%. Case-hardened steel for power transmission parts.

Images