Patent Application
20130186522
The present invention relates to a carburizing steel having excellent cold forgeability, which can be suitably used in, for example, automobiles and other various industrial equipment, and a method for manufacturing the same.
As vehicle weight is reduced more and more for better energy saving, there is an increasing demand for gears of smaller sizes for use in automobiles or the like, and the gears have to bear increasingly higher load exerted thereon. Furthermore, such gears have to bear increasingly higher load exerted thereon as the engine output is made higher and higher. Durability of a gear is primarily determined by degrees of gear tooth root fracture caused by bending fatigue and pitting fatigue fracture of gear tooth surface.
Conventionally, gears have been manufactured by preparing gear materials with case hardening steel, which is specified in JIS G 4053 (2003) as SCM420H, SCM822H and so on, and subjecting the gear materials to surface treatment, such as carburization. However, since such gears are not endurable in use under high stress conditions, attempts have been made to improve strength against bending fatigue of a gear tooth root and pitching resistance by changing steel materials, changing heat treatment methods, and furthermore, using work hardening processes on surfaces, and so on.
For example, JP 07-122118 B (PTL 1) discloses a method including: reducing Si content in steel and controlling Mn, Cr, Mo and Ni contents to suppress generation of grain boundary oxide layers on surfaces that would otherwise be generated after carburizing heat treatment and thereby suppress generation of cracks; suppressing generation of incompletely-quenched layers to suppress a reduction in surface hardness and thereby enhance fatigue strength; and further adding Ca to control stretching of MnS that would contribute to generation and propagation of cracks.
In addition, JP 2945714 B (PTL 2) discloses a method of enhancing resistance to temper softening by using a steel material to which Si is added in an amount of 0.25% to 1.50% as a material.
In addition, a material of parts to be produced by cold forming of bar materials, such as automobile parts, is required to have high cold forgeability. In view of this, it has been practiced to subject steel to spheroidizing heat treatment to spheroidize carbide in the steel to improve cold forgeability thereof.
For example, JP 4392324 B (PTL 3) discloses a method of obtaining a steel material that has low, and uniform hardness after spheroidizing annealing by controlling as-rolled microstructures, performing wire drawing and drawing processes at an area reduction rate of 28% or more, and performing subsequent spheroidizing annealing.
PTL 1: JP 07-122118 B
PTL 2: JP 2945714 B
PTL 3: JP 4392324 B
However, any of these techniques disclosed in the patent literatures, PTL 1, PTL 2 and PTL 3, have problems as mentioned below.
That is, according to PTL 1, it is indeed possible to suppress generation of bending fatigue cracks at tooth root because grain boundary oxide layers and incompletely-quenched layers are reduced with decreasing Si content.
However, merely reducing Si content alone leads to a reduction in resistance to temper softening. As a result, it is no longer possible to suppress softening of surfaces due to tempering caused by frictional heat at tooth root, and therefore, the gear becomes susceptible to pitching and fracture would propagate from tooth root side to tooth surface side.
In PTL 2, Si content is increased for the purpose of enhancing resistance to temper softening, in which case, however, deformation resistance increases during cold working, which is considered unsuitable for cold forging applications.
In addition, in PTL 3, an extra step is required to carry out a wire drawing process prior to spheroidizing annealing, which leads to increased cost.
Further, the form of as-rolled microstructures affects the microstructures and hardness after spheroidizing heat treatment. Particularly, in the case of relatively coarse ferrite and pearlite microstructures, there has been another problem in that appropriate spheroidized microstructures can be obtained within a limited control range, which makes it difficult to obtain stable microstructures.
The present invention has been developed in view of the above-described current situation, and an object of the present invention is to propose such a carburizing steel that exhibits high bending fatigue strength at tooth root of the resulting gears as compared to that of conventional gears, that also has excellent contact fatigue property, that is suitably used as a material for high strength gears and so on, that can establish spheroidizing annealed microstructures at low cost in a relatively easy manner, and furthermore, that has excellent cold forgeability and can be manufactured by mass production, as well as an advantageous method for manufacturing the same.
As a result of a keen study to solve the aforementioned problems, the inventers of the present invention have reached findings below.
a) By optimizing Si, Mn and Cr contents in a steel material to enhance resistance to temper softening, and, through this optimization, suppressing softening due to heat generation at gear contact surfaces, it is possible to suppress generation of cracks that would otherwise occur at tooth surfaces at the time of driving gears.
b) Regarding grain boundary oxide layers that may cause bending fatigue and fatigue cracks, by adding Si, Mn and Cr in an amount above a certain level, the direction in which the grain boundary oxide layers grow changes from the depth direction to the direction in which surface density increases. Accordingly, there are no oxide layers growing in the depth direction that would otherwise cause bending fatigue and fatigue cracks, and hence the grain boundary oxide layers less likely cause bending fatigue and fatigue cracks.
c) As mentioned in a) and b) above, while Si, Mn and Cr are effective for improving resistance to temper softening and for controlling grain boundary oxide layers, the contents of Si, Mn and Cr should be tightly controlled to achieve both effects.
d) To facilitate spheroidizing of carbides and improve cold forgeability, the contents of C, Si, Mn and Cr should be tightly controlled. Among others, it is particularly effective to add Cr in a large amount.
e) For stable spheroidizing of carbides, it is important that as-rolled microstructures are fine ferrite-pearlite microstructures. Therefore, the inventors of the present invention applied the spheroidizing heat treatment conditions as shown in
It should be noted that each steel subjected to the experiments contains basic elements that meet the following requirements and preferred conditions.
f) Further, cold forgeability is affected by microstructures, which in turn are strongly affected by those microstructures before annealing in addition to the above-described spheroidizing annealing conditions. That is, investigations were made on such microstructures before annealing for the proportion of ferrite-pearlite microstructures and the ferrite grain size.
As shown in
As used herein, in the experiment shown in
The present invention is based on the aforementioned findings.
That is, the arrangement of the present invention is summarized as follows:
[1] A carburizing steel having a composition containing, in mass %,
3.1≧{([% Si]/2)+[% Mn]+[% Cr]}≧2.2 (1)
[% C]−([% Si]/2)+([% Mn]/5)+2[% Cr]≧3.0 (2)
2.5≧[% Al]/[% N]≧1.7 (3)
where [% M] represents content (in mass %) of element M.
Besides, the above-described carburizing steel is subjected to cold forging, where it is processed to the shape of various parts after carburizing process. While it is preferable that the steel is subjected to spheroidizing annealing prior to this cold forging, it may be directly subjected to cold forging without undergoing spheroidizing annealing, depending on the degree of working required, and so on.
[2] The carburizing steel according to [1] above, wherein the composition further contains, in mass %, at least one element selected from
[3] A method of manufacturing a carburizing steel, the method comprising:
subjecting a steel material to hot working, where the steel material is heated to 1160° C. or higher and lower than 1220° C., the steel material having a composition containing, in mass %,
After suspending the hot working within a temperature range of Ar3 point or higher, cooling the steel material to 450° C. or lower;
then reheating the steel material to a temperature higher than 900° C., but not higher than 970° C., to resume the hot working;
terminating the hot working under a condition of a total reduction rate being 70% or more after the reheating; and
then cooling the steel material at a rate of 0.1° C./s to 1.0° C./s from 800° C. to 500° C.:
3.1≧{([% Si]/2)+[% Mn]+[% Cr]}≧2.2 (1)
[% C]−([% Si]/2)+([% Mn]/5)+2[% Cr]≧3.0 (2)
2.5≧[% Al]/[% N]≧1.7 (3)
where [% M] represents content (in mass %) of element M.
[4] The method of manufacturing a carburizing steel according to [3] above, wherein the composition further contains, in mass %, at least one element selected from
According to the present invention, a carburizing steel that has not only excellent bending fatigue properties at tooth root, but also excellent contact fatigue properties of tooth surfaces, when worked to gears, for example, may be obtained in a step associated with cold forging under mass production conditions.
The present invention will be further described below with reference to the accompanying drawings, wherein:
The present invention will be described in detail below.
First, the reasons why the chemical compositions of the steel materials are restricted to the aforementioned ranges in the present invention will be described below. As used herein, “%” used for elements of a steel sheet represents “mass %,” unless otherwise stated.
0.1%≦C≦0.35%
Carbon (C) content in steel needs to be 0.1 mass % or more in order to enhance hardness at the center portion thereof by quenching after carburizing treatment. However, C content in steel exceeding 0.35 mass % decreases toughness of the core portion of the steel. Therefore, the C content is limited in the range of 0.1% to 0.35%, preferably 0.1% to 0.3%.
0.01%≦Si≦0.22%
Silicon (Si) is an element that enhances resistance to softening within a temperature range of 200° C. to 300° C., which is expected to be reached by gears or the like while rotating. To obtain this effect, it is essential to add Si in an amount of at least 0.01%, preferably 0.03% or more. On the other hand, however, Si is a ferrite-stabilizing element that raises the Ac3 transformation point if added excessively, and facilitates generation of ferrite at the core portion having low carbon content within a normal quenching temperature range, which leads to a reduction in deterioration in strength. In addition, excessive addition of Si results in hardening of a steel material before being carburized and deteriorates cold forgeability, which is considered disadvantageous. To this extent, Si content of 0.22% or less does not have such an adverse effect as described above. Therefore, the Si content is limited to a range of 0.01% to 0.22%, preferably 0.03% to 0.22%.
0.3%≦Mn≦1.5%
Manganese (Mn) is an element that effectively improves quench hardenability and needs to be added to steel in an amount of at least 0.3%. However, Mn facilitates formation of abnormally carburized layers and excessive addition of Mn leads to excessive retained austenite and lower hardness. Therefore, the upper limit of the Mn content is to be 1.5%, preferably 0.4% to 1.2%, more preferably 0.6% to 1.2%.
1.35%≦Cr≦3.0%
Chromium (Cr) is an element that effectively improves not only quench hardenability but also resistance to temper softening. However, Cr content in steel below 1.35% is less effective in terms of achieving this effect. On the other hand, if Cr content in steel exceeds 3.0%, the effect of enhancing resistance to softening that is attained by adding Cr reaches a saturation point, and facilitates, rather suppresses, formation of abnormally carburized layers. Therefore, the Cr content is limited in the range of 1.35% to 3.0%, preferably 1.35% to 2.6%.
P≦0.018%
Phosphorus (P) is an element that exists in a segregated manner at crystal grain boundaries and deteriorates toughness of steel. Accordingly, lower content of P in steel is more desirable, although the presence of P in steel is acceptable up to 0.018%, preferably not more than 0.016%. The P content may be 0%, where possible, although it is generally difficult to do so.
S≦0.02%
Sulfur (S) is an element that exists as a sulfide inclusion in steel and effectively improves machinability of the steel. However, excessive addition of S in steel leads to a deterioration in fatigue strength of the steel. Therefore, the upper limit of S content is to be 0.02 mass %. From the viewpoint of machinability, S may also be contained in an amount of 0.004% or more.
0.015%≦Al≦0.05%
Aluminum (Al) is an element that is bonded to nitrogen (N) to form AlN and contributes to refinement of austenite crystal grains. To obtain this effect, Al needs to be added in an amount of 0.015% or more, preferably 0.018% or more. However, Al content in steel exceeding 0.05% facilitates generation of Al2O3 inclusions, which harmfully affects fatigue strength of the steel. Therefore, the Al content in steel is limited in the range of 0.015% to 0.05%, preferably 0.015% to 0.037%.
0.008%≦N≦0.015%
Nitrogen (N) is an element that is bonded to Al to form AlN and contributes to refinement of austenite crystal grains. Accordingly, it is necessary to add N in an amount of 0.008% or more to obtain this effect, although an appropriate amount of N to add is determined by the quantitative balance with Al. However, excessive addition of N causes air bubbles in steel ingots during solidification and/or leads to a deterioration in forgeability. Therefore, the upper limit of N content is to be 0.015%, preferably in the range of 0.010% to 0.015%.
O≦0.0015%
Oxygen (O) is an element that exists as an oxide inclusion in steel and impairs fatigue strength. Therefore, lower content of O in steel is more desirable, although the presence of O in steel is acceptable up to 0.0015%. The O content may be 0%, where possible, although it is generally difficult to do so.
While appropriate composition ranges have been described with respect to the basic elements of the present invention, to implement the present invention, it is inadequate that each element only satisfies the above-described range. Rather, regarding the contents of C, Si, Mn, Cr, Al and N, it is also necessary to satisfy the following formulas (1), (2) and (3):
3.1≧{([% Si]/2)+[% Mn]+[% Cr]}≧2.2 (1)
[% C]−([% Si]/2)+([% Mn]/5)+2[% Cr]≧3.0 (2)
2.5≧[% Al]/[% N]≧1.7 (3)
where [% M] represents content (in mass %) of element M.
The formula (1) is a factor that affects quench hardenability and temper softening resistancy. If the formula (1) has a value below 2.2, it is not possible to produce a sufficient effect of improving quench hardenability and temper softening resistancy, which results in insufficient fatigue strength.
On the other hand, if the formula (1) has a value above 3.1, the above-described improvement effect reaches a saturation point, even leading to a deterioration in cold-rolling workability.
In addition, the formula (2) is a factor that affects the ease to spheroidize carbides, where spheroidizing of carbides becomes easier to perform when the formula (2) is 3.0 or more. Combining this composition with the above findings e) and f) may offer outstanding cold forgeability after spheroidizing annealing.
Further, the formula (3) is a factor that affects refinement of austenite crystal grains. If the formula (3) has a value below 1.7, there is a poor refining effect, leading to insufficient fatigue strength. On the other hand, if the formula (3) has a value above 2.5, crystal grains coarsen easily and fatigue strength becomes insufficient, and furthermore, workability deteriorates due to solute Al and solute N.
In addition to the aforementioned basic components of the present invention, the composition of the steel of the present invention may optionally contain other elements as described below appropriately.
Cu≦1.0%
Cupper (Cu) is an element that effectively improves the strength of a base material. However, Cu content in steel exceeding 1.0% causes hot shortness and impairs the surface texture of a steel material. Therefore, the Cu content is to be 1.0% or less. Cu is preferably added in an amount of 0.01% or more.
Ni≦0.5%
Nickel (Ni) is an element that effectively improves the strength and toughness of a base material, but is expensive. Therefore, Ni content in steel is to be 0.5% or less. Ni is preferably added in an amount of 0.01% or more.
Mo≦0.5%
Molybdenum (Mo) is an element, as is the case with nickel (Ni), that effectively improves the strength and toughness of a base material, but is expensive. Therefore, Mo content in steel is to be 0.5% or less. The Mo content may be 0.2% or less. The Mo content is preferably 0.05% or more.
V≦0.5%
Vanadium (V), like silicon (Si), is an element that is useful for enhancing resistance to temper softening. However, if V content exceeds 0.5%, the effect attained by adding V reaches a saturation point. Therefore, the V content in steel is to be 0.5% or less. V is preferably added in an amount of 0.01% or more.
Nb≦0.06%
Niobium (Nb) is an element, as is the case with V and Si, that is useful for enhancing resistance to temper softening. However, if Nb content in steel exceeds 0.06%, the effect attained by adding Nb reaches a saturation point. Therefore, the Nb content is to be 0.06% or less. Nb is preferably added in an amount of 0.007% or more.
The balance of the composition of each steel material is composed of Fe and incidental impurities. For example, although not particularly added herein, the steel may also contain B as an impurity, provided that the B content is on the order of less than 0.0003%.
Further, in addition to the above-described adjustment of the chemical compositions, it is also necessary to control microstructures of the steel before spheroidizing annealing of a material.
Total Microstructure Proportion of Ferrite and Pearlite≧85%
Deformation resistance increases with increasing a proportion of of bainite in microstructures before spheroidizing annealing, which impairs cold forgeability. Therefore, it is necessary to control the total proportion of ferrite and pearlite microstructures to be 85% or more to reduce the proportion of bainite. The upper limit of the total microstructure proportion of ferrite and pearlite may be 100%.
Since the present invention uses steels that have high hardenability and satisfy the above formulas, such as the formula (1) and so on, it is difficult to ensure a sufficient amount of ferrite and pearlite as described above when using a normal manufacturing method. However, the condition of (ferrite+pearlite=85% or more) may be satisfied by adjusting the heating temperature during rolling, total reduction rate, and cooling rate.
Average Ferrite Grain Size≦25 μm
Microstructures before spheroidizing annealing have a large impact on the properties after spheroidizing annealing. That is, the microstructures before spheroidizing annealing having a ferrite grain size of more than 25 μm leads to a deterioration in cold forgeability after spheroidizing process.
Particularly, in view of its large impact on the limit upset ratio, the average ferrite grain size is to be 25 μm or less. Practically, but not necessarily in technical sense, a lower limit is about 5 μm.
Next, conditions under which the present invention is manufactured will be described.
In the present invention, the following operations should be performed:
heating a steel material having the above-described preferred chemical composition to at least 1160° C. and lower than 1220° C.; After suspending rolling within a temperature range of Ar3 point or higher, cooling the steel material to 450° C. or lower; then reheating the steel material to a temperature higher than 900° C., but not higher than 970° C.; terminating the hot rolling under a condition of a total reduction rate being 70% or more after the reheating; and cooling the steel material at a rate of 0.1° C./s to 1.0° C./s from 800° C. to 500° C.
In the following, reasons for the above-described limitations on the operational conditions will be described.
Steel Material Heating Temperature (First Stage): 1160° C. or Higher and Lower than 1220° C.
In the present invention, a steel material is heated to a temperature of 1160° C. or higher due to the need to dissolve AlN to a sufficient extent from an as-solidified state. However, an excessively high heating temperature worsens scale loss, deteriorates surface texture, increases fuel cost, and so on. Therefore, the first stage heating temperature is to be lower than 1220° C.
Cooling to 450° C. or Lower after the Completion of Hot Working within a Temperature Range of Ar3 Point or Higher
In this hot working step, preferably in the hot rolling step, working is suspended at Ar3 point or higher, cooling the steel material to 450° C. or lower so that the as-cast microstructures are altered to provide ferrite-pearlite microstructures. It is also advantageous from the viewpoint of obtaining ferrite-pearlite microstructures that the hot working is performed at a reduction rate of 50% or more. Regarding the finish cooling temperature, any practical values may be selected by taking into account reheating cost, and so on, without having to set a particular lower limit. Also, regarding the reduction rate of the hot working, any practical values may be selected considering facility load, and so on, without having to set a particular upper limit.
Steel Material Reheating Temperature (Second Stage): Higher than 900° C., but Not Higher than 970° C.
The steel material is reheated to a temperature of 970° C. or lower, because the as-rolled microstructures need to be fine ferrite-pearlite microstructures in order to obtain spheroidizing annealed microstructures and low hardness. Above 970° C., AlN precipitates coarsely; whereas at 970° C. or lower, AlN precipitates finely, which is also effective for suppressing grain coarsening during carburizing. However, AlN does not precipitate to a sufficient extent when the steel material is heated to 900° C. or lower. Therefore, the second stage heating temperature is to be higher than 900° C., preferably 920° C. or higher.
Total Reduction Rate During Hot Working: 70% or More
If a total reduction rate during hot working after the reheating, i.e., a total of reduction rates in a working step after the reheating is low, crystal grains coarsen to reduce the proportion of ferrite after cooling, coarse grains are produced more easily during carburizing, and furthermore, the hardness of the worked material increases. Therefore, the total reduction rate is to be 70% or more. Regarding the reduction rate, any practical values may be selected considering facility load, and so on, without having to set a particular upper limit.
As used herein, the term “reduction rate” means a thickness reduction rate where a steel material to be obtained by hot working is a steel sheet, while meaning an area reduction rate where a steel material to be obtained by hot working is a steel bar or wire rod.
Cooling Rate from 500° C. to 800° C.: 0.1° C./s to 1.0° C./s
If the cooling rate is lower than 0.1° C./s from 800° C. to 500° C. during a cooling process after the hot working, the ferrite grain size increases, which results in coarse ferrite-pearlite microstructures. On the other hand, if the cooling rate is higher than 1.0° C./s, the proportion of ferrite after cooling decreases, which provides bainite and ferrite-pearlite composite microstructures. Therefore, the cooling rate within this temperature range is limited in the range of 0.1° C./s to 1.0° C./s.
Each carburizing steel obtained by the above-described manufacturing method is desirably subjected to spheroidizing annealing and subsequent cold forging. The spheroidizing annealing is preferably performed at 760° C. to 820° C. for 2 to 15 hours or so. However, the present invention may still offer excellent cold forgeability even if spheroidizing annealing is performed at a relatively low temperature, in particular, about 740° C. to 760° C. Besides, the microstructures after the spheroidizing annealing results from separating and spheroidizing the plate-like cementite in the layered pearlite of the previous microstructures. While the base microstructures are ferrite, the previous microstructures are generally succeeded by the subsequent microstructures, because the previous ones are retained at the austenite-ferrite dual phase region during the heating stages. Each steel that has been cold forged to the shape of predetermined parts is subjected to carburizing heat treatment in a conventional manner. The resulting members after the carburizing heat treatment have surfaces that are mainly composed of martensite microstructures (or tempered martensite microstructures when members have been subjected to tempering processes).
Steel samples having different chemical compositions shown in Table 1 were prepared by steelmaking in a 100 kg vacuum melting furnace to produce cast steel products, which in turn were subjected to rolling under the hot working and cooling conditions as shown in Table 2 to be finished to steel bars. That is, the steel samples were subjected to the first stage hot working where they were heated at respective heating temperatures as shown in Table 2, then cooled to 450° C. or lower, and thereafter subjected to the second stage hot working where they were heated, rolled and cooled under the conditions as shown in Table 2 relating to the heating temperature, total reduction rate and cooling rate, to thereby obtain steel bars. The resulting steel bars were evaluated for their microstructure proportion and average ferrite grain size, cold workability, spheroidizing heat treatability, properties of carburized portions, and fatigue properties. The evaluation was made under the following conditions.
(1) Microstructure Proportion and Average Ferrite Grain Size
A position of a diameter×¼ from a steel surface (¼ D position) in a cross-section along longitudinal direction (L direction) of each steel bar was first mirror polished at and then etched by nital. Thereafter, a magnified (400×) image of the cross-section was observed by image analysis to determine the microstructure proportion (area ratio) of ferrite and pearlite microstructures and the average ferrite grain size.
(2) Evaluation of Cold Workability (Cold Forgeability)
The cold workability of each steel bar was evaluated in terms of deformation resistance and limit upset ratio.
Specifically, the deformation resistance was determined by: collecting a test piece (diameter: 10 mm, height: 15 mm) from a region ranging from a steel surface to a ¼ D position of each as-rolled steel bar; and then measuring compression load at 70% upset forging by using a 300 t (3000 kN) press machine according to the deformation resistance measuring method recommended by The Japan Society for Technology of Plasticity, based on end face confined compression.
The limit upset ratio was determined by measuring an upset ratio at a point in time when any crack occurred on an end portion of the test piece after the test piece had been subjected to a compression process according to the deformation resistance measuring method as described above.
It is considered that a deformation resistance of 918 MPa or lower and a limit upset ratio of 76% or more offer good cold workability.
(3) Evaluation of Spheroidizing Heat Treatability
The cold workability of each steel bar was evaluated in terms of hardness after spheroidizing heat treatment, deformation resistance, and limit upset ratio.
Specifically, as is the case with the evaluation of cold workability as described in item (2) above, the deformation resistance and limit upset ratio were determined by: collecting a test piece (diameter: 10 mm, height: 15 mm) from a region ranging from a steel surface to a diameter×¼ position in the radial direction of each as-rolled steel bar having a diameter D; and then subjecting the test piece to spheroidizing heat treatment to determine the deformation resistance and limit upset ratio thereof. The spheroidizing heat treatment was performed under two conditions (A) and (B) shown in
In addition, if the deformation resistance after the spheroidizing heat treatment (condition (A)) is 890 MPa or less, and if the limit upset ratio is 80% or more, then the steel bar is assumed to have good cold workability.
(4) Evaluation of Properties of Carburized Portions
Carburized-portion properties were evaluated in terms of presence or absence of coarse grains at carburized portions and oxidation depth at grain boundaries after carburization was performed at 930° C. for 7 hours with a carbon potential of 0.8%.
Specifically, “circle” indicates a case where no coarse grains occurred at any carburized portions, while “cross” represents a case where coarse grains occurred at any carburized portions.
Oxidation behavior at grain boundaries was evaluated by observing, with an optical microscope, a surface of each test specimen after being subjected to the carburizing process and measuring the oxidation depth at grain boundaries thereof. Specifically, a surface of each test specimen was observed with an optical microscope at 400× magnification to determine the maximum grain boundary oxidation depth from each field of view. Then, an average value from 10 fields of view was regarded as the grain boundary oxidation depth. If there are no coarse grains at carburized portions, and if the grain boundary oxidation depth is 10 μm or less, then the steel bar is assumed to have excellent properties of carburized portions.
(5) Evaluation of Fatigue Properties
Fatigue properties were evaluated in terms of rotating bending fatigue test specimen and surface fatigue strength.
Specifically, rotating bending fatigue test specimens and roller pitting fatigue test specimens for evaluation of surface fatigue strength were taken from as-rolled steel bars and subjected to experiments. Then, these test specimens were subjected to carburization at 930° C. for 7 hours with a carbon potential of 0.8%, followed by a heating and tempering process at 180° C. for one hour. The rotating bending fatigue test was conducted at a revolution speed of 1800 rpm to evaluate the fatigue strength at finite life after 107 cycles.
The roller pitting fatigue test was conducted under the conditions of slip rate: 40% and oil temperature: 80° C. to evaluate the fatigue strength at finite life after 107 cycles.
If the rotating bending fatigue strength is 806 MPa or more and the surface fatigue strength is 3250 MPa or more, then the steel bar is assumed to have good fatigue strength.
The obtained results are shown in Table 3.
It can be seen from Table 3 that the inventive examples obtained according to the present invention all show excellent cold workability in an as-rolled state and after being subjected to spheroidizing heat treatment, have a small grain boundary oxidation depth, involve no coarse grains at carburized portions, and furthermore, exhibit better rotating bending fatigue strength and surface fatigue strength than those of the comparative examples.
The present invention may provide a carburizing steel that is excellent in cold workability, rotating bending fatigue strength and surface fatigue strength. Accordingly, for example, such a carburizing steel that has not only excellent bending fatigue properties at tooth root, but also excellent contact fatigue properties of tooth surfaces when worked to gears may be obtained in a step associated with cold forging under mass production conditions.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2010-267779 | Nov 2010 | JP | national |
| 2011-218340 | Sep 2011 | JP | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/JP2011/006655 | 11/29/2011 | WO | 00 | 4/3/2013 |
