Patent Application
20160060744
The present invention relates to a case-hardening steel and a case-hardened steel member, and particularly to a case-hardening steel, which has excellent cold forgeability and capable of obtaining excellent temper softening resistance after a carburizing treatment or a carbonitriding quenching and tempering treatment, and is suitable for components for automobiles, construction machines, and industrial machines, and a case-hardened steel member.
Priority is claimed on Japanese Patent Application No. 2013-087857, filed Apr. 18, 2013, the content of which is incorporated herein by reference.
Transmission gears used in automobiles, construction machines, etc. and reduction gears used in industrial machines, etc. are mainly configured by gears. These components can be obtained by using a medium carbon alloy steel such as JIS SCr 420, JIS SCM 420, etc. as a material, forming the material into the shape of a component by hot forging, cutting, cold forging, or a combination of these processes and then subjecting the material to a surface-hardening treatment such as carburized quenching and tempering, etc. Among the components, the component formed by cold forging is subjected to spheroidizing annealing before cold forging in order to improve die life by softening of the material. The problem when the cold forging is performed is to prevent cracks from initiating during the cold forging and improve die life. Accordingly, when both suppression of formation of inclusions serving as an origin of crack initiation and softening of material can be achieved, the cost for cold forging can be reduced.
On the other hand, there has been a strong demand for increasing the strength of gears to achieve high output performance and improvement of fuel efficiency of automobiles, etc. . . . In the related art, in order to increase the strength of these components, a technique of improving the tooth root bending fatigue strength of gears, which is a big problem in high-strengthening, has been developed. However, in recent years, along with the expansion of application of hard shot peening capable of rapidly increasing tooth root bending fatigue strength, the main point of the problem to achieve high-strengthening of gears is shifted from improvement of tooth root bending fatigue strength to improvement of pitching strength.
In order to enhance (improve) the pitching strength, it is effective to improve the temper softening resistance of a carburized layer in a gear. As a method of improving temper softening resistance, a technique of improving the components of steel is proposed. For example, in Patent Document 1, it is disclosed that when the amounts of Si, Cr, and Mo are defined and the total amount of these elements is more than a predetermined value, temper softening resistance increases. However, when the total amount of these elements increases, the hardness of the material before cold forging is increased and deformation resistance is increased. In addition, for example, in Patent Document 2, it is disclosed that when the amount of Si is more than 0.15%, deformation resistance during cold forging is increased. As described above, generally, when the amount of each component of steel (particularly, Si) is increased, the effect of improving temper softening resistance is obtained. However, the hardness of the material is increased. That is, improving temper softening resistance and ensuring cold forgeability are in a trade-off relationship. Therefore, it has been desired to develop steel achieving both high temper softening resistance and high cold forgeability. In Patent Document 3, there is disclosed a method of realizing both high temper softening resistance and high cold forgeability by increasing temper softening resistance by increasing the amounts of Si and Cr and then limiting the total amount of Si, Mn, Cr, and Mo to a value that is determined by a predetermined relational expression or less. However, the technique disclosed in Patent Document 3 does not consider prevention of crack initiation during cold forging. Therefore, there is a problem of that cracks initiate from inclusions when large inclusions are present in the portion in which the working ratio increases during cold forging and there is still much room for improvement. In Patent Documents 4 to 11, a steel for machine structure in which the size of the inclusions is limited is disclosed. However, there is no description of cold forging in any of the above Patent Documents. In Patent Document 12, there is disclosed a steel bar and a wire rod achieving both high cold forgeability and high machinability by limiting the size of sulfide-based inclusions, oxide-based inclusions, nitride-based inclusions, or composite inclusions thereof. However, there is no description of a technique of improving temper softening resistance. In Patent Document 13, a steel for vacuum carburizing or vacuum carbonitriding is disclosed. In the steel, the maximum equivalent circle diameter of oxides, composite inclusions mainly composed of oxides and nitrides, and composite inclusion mainly composed of nitrides, which is expressed by [(πLW/4)0.5] when a cumulative distribution function estimated by an extreme value statistical method is 99%, is 35 μm or less. However, applying vacuum carburizing or vacuum carbonitriding is premised in Patent Document 13.
Accordingly, there is need for a steel which both improves temper softening resistance and ensures cold forgeability (prevention of crack initiation and prevention of an increase in the material hardness).
[Patent Document 1] Japanese Unexamined Patent Application, First Publication No. 2003-231943
[Patent Document 2] Japanese Unexamined Patent Application, First Publication No. H6-299241
[Patent Document 3] Japanese Unexamined Patent Application, First Publication No. 2006-199993
[Patent Document 4] Japanese Unexamined Patent Application, First Publication No. 2001-234275
[Patent Document 5] Japanese Unexamined Patent Application, First Publication No. 2001-131685
[Patent Document 6] Japanese Unexamined Patent Application, First Publication No. 2001-131686
[Patent Document 7] Japanese Unexamined Patent Application, First Publication No. 2003-269460
[Patent Document 8] Japanese Unexamined Patent Application, First Publication No. 2006-63402
[Patent Document 9] Japanese Unexamined Patent Application, First Publication No. 2007-289979
[Patent Document 10] Japanese Unexamined Patent Application, First Publication No. 2004-143550
[Patent Document 11] Japanese Unexamined Patent Application, First Publication No. 2005-154886
[Patent Document 12] Japanese Unexamined Patent Application, First Publication No. 2007-63589
[Patent Document 13] Japanese Unexamined Patent Application, First Publication No. 2010-150566
The present invention has been made in consideration of the above circumstances and an object thereof is to provide a case-hardening steel having excellent cold forgeability and temper softening resistance and a case-hardened steel member made from the case-hardening steel.
In the present invention, excellent temper softening resistance means that the hardness of a carburized layer after tempering at 300° C. is higher than the hardness of JIS SCr 420 or JIS SCM 420.
In order to solve the above-described problem, the present inventors have conducted an intensive investigation on adjustment of the chemical composition to be suitable for improving temper softening resistance and control of size of inclusions, which is required to prevent crack initiation during cold forging. As a result, it has been found that cracks can be prevented from initiating during cold forging due to the fact that (i) Si and Cr have a significant action of increasing the temper softening resistance of a carburized layer, (ii) the hardness of the steel after spheroidizing annealing depends on the total amount of Si, Mn, Cr and Mo and the contribution rate of each element is different, and (iii) the size of non-metallic inclusions present in the steel, particularly, the size of sulfide-based inclusions is appropriately limited, etc., and the present invention has been completed.
The gist of the present invention is as follows.
(1) According to an aspect of the invention, there is provided a case-hardening steel includes, as a chemical composition, by mass %:
C: 0.05% to 0.30%;
Si: 0.40% to 1.5%;
Mn: 0.2% to 1.0%;
S: 0.001% to 0.050%;
Cr: 1.0% to 2.0%;
Mo: 0.02% to 0.8%;
Al: 0.001% to 0.20%;
N: 0.003% to 0.03%;
Nb: 0% to 0.10%;
Cu: 0% to 0.2%;
Ni: 0% to 1.5%;
V: 0% to 0.20%;
Ca: 0% to 0.0050%;
Mg: 0% to 0.0050%;
Sb: 0% to 0.050%;
P: limited to 0.030% or less;
O: limited to 0.0020% or less;
Ti: limited to 0.005% or less; and
a balance consisting of Fe and impurities,
wherein the following Equations (α) and (β) are satisfied,
in inclusion evaluation using an extreme value statistical method, when an estimated area S is 30,000 mm2, an estimated value of the maximum size (√area)S of sulfide-based inclusions present in the estimated area S is 49 μm or less and an estimated value of the maximum size (√area)Ox of oxide-based inclusions present in the estimated area S is 80 μm or less, and
the number of sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm per 1 mm2 is limited to 200 or less,
12×Si(%)+25×Mn(%)+Cr(%)+2×Mo(%)≦25 (α),
31×Si(%)+15×Mn(%)+23×Cr(%)≧50 (β),
here, Si(%), Mn(%), Cr(%) and Mo(%) represent the amount of each element by mass % in Equations (α) and (β).
(2) The case-hardening steel according to (1), may include, as the chemical composition, by mass %:
Nb: 0.015% to 0.10%.
(3) The case-hardening steel according to (1), may include, as the chemical composition, by mass %:
Si: 0.55% to 1.5%.
(4) The case-hardening steel according to any one of (1) to (3), may include, as the chemical composition, by mass %, either or both of,
Cu: 0.001% to 0.2%; and
Ni: 0.001% to 1.5%.
(5) The case-hardening steel according to any one of (1) to (4), may include, as the chemical composition, by mass %:
V: 0.01% to 0.20%.
(6) The case-hardening steel according to any one of (1) to (5), may include, as the chemical composition, by mass %, either or both of,
Ca: 0.0001% to 0.0050%; and
Mg: 0.0001% to 0.0050%.
(7) The case-hardening steel according to any one of (1) to (6), may include, as the chemical composition, by mass %:
Sb: 0.0001% to 0.050%.
(8) The case-hardening steel according to any one of (1) to (7), in which the microstructure has a spheroidized carbide structure.
(9) According to another aspect of the invention, there is provided a case-hardened steel member made from the case-hardening steel according to any one of (1) to (8); in which a surface hardened layer that is formed by a carburizing quenching and tempering treatment or a carburizing nitriding quenching and tempering treatment.
According to the embodiments of the present invention, it is possible to provide a case-hardening steel having higher hardness of a carburized layer after tempering at 300° C. than the hardness of JIS SCr 420 or JIS SCM 420 and excellent cold forgeability and a case-hardened steel member. That is, it is possible to provide a case-hardening steel having excellent temper softening resistance and cold forgeability and providing a case-hardened steel member. In addition, the use of the case-hardening steel and the case-hardened steel member enables reduction of the production cost of gears and contribution to high output performance and improvement of fuel efficiency, etc. for automobiles, construction machines, and industrial machines.
As a result of the research, the present inventors have found the following (a) to (d).
(a) The limit of cold forgeability (limit of hardness before cold forging) can be determined based on indices of the amounts of each of Si, Mn, Cr and Mo in consideration of the hardness increasing action of each element.
The present inventors have conducted spheroidizing annealing (SA) on plural types of steel obtained by adding various alloy elements to 0.2% C (steel including C in an amount of 0.2%) to quantitatively evaluate the effect of each alloy element on hardness after spheroidizing annealing. When spheroidizing annealing is performed, carbides in steel constituting pearlite, etc. are spheroidized and the microstructure has a spheroidized carbide structure. When the carbides are spheroidized, the spacing between carbides which becomes an obstacle to dislocation motion is increased and the hardness is decreased. Thus, it is desirable to spheroidize carbides.
As a result of the investigation, the present inventors have found that the hardness of steel after spheroidizing annealing can be expressed by the form of the left side of the following Equation (1). The reason that the coefficients of Si and Mn are relatively high is that these alloy elements are solid-soluted into ferrite and the hardness of a spheroidizing-annealed material is increased by solute strengthening. On the other hand, the reason that the coefficients of Cr and Mo are relatively low is that these alloy elements, which are concentrated in cementite during spheroidizing annealing or precipitated in the form of alloy carbides, are not so much contribute to solute strengthening and these carbides relatively contribute little to precipitation strengthening because the size of the carbide is large.
The present inventors have found that when the value of the left side of the following Equation (1) is 25 or less, the hardness of the steel after spheroidizing annealing is not excessively increased, and when the value of the left side of the following Equation (1) is more than 25, the hardness of the steel after spheroidizing annealing is excessively increased and thus cold forgeability is deteriorated.
12×Si(%)+25×Mn(%)+Cr(%)+2×Mo(%)≦25 (1)
Here, Si(%), Mn(%), Cr(%) and Mo(%) represent the amounts (mass %) of each component in steel in Equation (1).
(b) The temper softening resistance (hardness after tempering at 300° C.) of steel (particularly, carburized layer) can be expressed by indices of the amounts of each of Si, Mn and Cr in consideration of the temper softening resistance increasing action of each element.
Si, Mn and Cr have a significant action of increasing the temper softening resistance of the carburized layer. This is because when Si, Mn and Cr are contained in the steel, iron carbides precipitated during tempering are prevented from being coarsened. In order to quantitatively evaluate the effect of each alloy element, the present inventors have simulated a carburized layer and conducted tempering on plural types of steel obtained by adding various alloy elements to 0.8% C at 300° C. to investigate the effect of various alloy elements on hardness after tempering (hardness after tempering at 300° C.). As a result, the present inventors have found that the action of increasing the hardness of the carburized layer after tempering at 300° C. by each alloy element can be expressed by the form of the left side of the following Equation (2). In addition, it has been found that when the value of the left side is 50 or more, the hardness after tempering at 300° C. is apparently increased compared to the hardness of a general carburized component, and excellent pitching strength can be obtained, and when the value is less than 50, the pitching strength is not sufficiently improved.
31×Si(%)+15×Mn(%)+23×Cr(%)≧50 (2)
Here, Si(%), Mn(%) and Cr(%) represent the amounts (mass %) of each component in steel in Equation (2).
Accordingly, both an increase in temper softening resistance and a decrease in material hardness (ensuring cold forgeability) can be achieved by containing Si, Mn, Cr and Mo in a range satisfying the above Equations (1) and (2) at the same time.
c) It is possible to prevent cracks from initiating during cold forging by limiting the size of non-metallic inclusions (sulfide-based inclusions, oxide-based inclusions and nitride-based inclusions), particularly, sulfide-based inclusions, present in the steel.
The large inclusions present in the steel serve as an origin of crack initiation. Thus, it is necessary to evaluate the distribution condition of the material of the component in a wide range for stable mass production on an industrial scale. The presence of the large inclusions serving as an origin of crack initiation can be estimated by an “extreme value statistical method”. The extreme value statistical method refers to a method of estimating the maximum particle size of inclusions (√area) present in the population or an arbitrary area (or volume) by selecting plural test pieces from a given population, measuring the maximum size of inclusions present in the individual test pieces by microscopy, and plotting the square root of the area to extreme value probability paper. As a specific method of applying the extreme value statistical method to the evaluation of non-metallic inclusions in the steel, for example, the extreme value statistical method can be performed according to the method described in non-patent documents, Metal fatigue: effects of small defects and nonmetallic inclusions, written by Murakami Yukitaka, etc. In the embodiment, the following method is used. (i) Optical microscope observation is performed in 30 view fields respectively for each test material by setting the area of one view field (inspection reference area: S0) to, for example, 10 mm×10 mm, so as to avoid overlapping of the area S0. (ii) The maximum particle size of inclusions present in the respective 30 view fields is measured and the square root of the area (√area) is plotted to extreme value probability paper. (iii) The maximum particle size of inclusions (√area) is estimated by setting an estimated area S to 30,000 mm2.
For the measurement of the inclusions, it is necessary to measure the size of respective oxide inclusions and sulfide inclusions. This is because the particle distribution of oxides and the particle distribution of sulfides are different from each other and the evaluation has to be performed separately. The extreme value statistical method is relatively simple and has a high reliability.
(d) The presence frequency of sulfide-based inclusions is high. Therefore, in order to prevent cracks from initiating during cold forging, it is necessary to limit the number of sulfide-based inclusions per unit area of a given size or larger (number density) in addition to the maximum size estimated by the extreme value statistical method.
Hereinafter, a case-hardening steel according to one embodiment of the present invention (referred to as a case-hardening steel according to an embodiment in some cases) and a case-hardened steel member according to one embodiment of the present invention (referred to as a case-hardened steel member according to an embodiment in some cases) will be described in detail. First, the reason for limiting the components of the case-hardening steel according to the embodiment will be described. The components represent the components of a core which is not affected by an increase in the amount of carbon due to carburizing of the surface portion. The symbol % of amounts of components means mass %.
(C: 0.05% to 0.30%)
C is an essential element to obtain core strength of a carburizing-quenched component. In addition, the amount of C determines the hardness of the core and also affects the depth of an effective hardened layer of a carburized layer. Here, in the embodiment, the lower limit of the amount of C is set to 0.05%. However, when the amount of C is excessive, toughness is deteriorated. Therefore, the upper limit of the amount of C is set to 0.30%. The amount of C is more desirably 0.10% to 0.25%.
(Si: 0.40% to 1.5%)
Si is an effective element for improving the temper softening resistance of the carburized layer. Therefore, the lower limit of the amount of Si is set to 0.40%. However, when the amount of Si is excessive, the hardness after spheroidizing annealing is increased and cold forgeability is deteriorated. Therefore, the upper limit of the amount of Si is set to 1.5%. The amount of Si is desirably 0.45% to 1.0%. When improving the temper softening resistance while limiting a cost increase, it is more desirable to set the lower limit of the amount of Si to 0.55%.
(Mn: 0.2% to 1.0%)
Mn is an effective element for improving the hardenability of the steel. In addition, Mn improves hot ductility by fixing S in the steel in the form of MnS and prevents scratches from being formed in the process of producing steel (continuous casting and hot rolling). Further, MnS improves machinability. In order to obtain these effects, the lower limit of the amount of Mn is set to 0.2%. However, when the amount of Mn is excessive, the hardness after spheroidizing annealing is increased and cold forgeability is deteriorated. Therefore, the upper limit of the amount of Mn is set to 1.0%. The amount of Mn is desirably 0.4% to 0.7%.
(S: 0.001% to 0.050%)
S has an effect of improving machinability by forming MnS in the steel. In order to obtain this effect, the lower limit of the amount of S is set to 0.001%. However, when the amount of S is excessive, the amount of MnS, etc., so-called sulfide-based inclusions, is increased and the size thereof is coarsened. As described later, when the number of coarsened sulfide-based inclusions is large, the coarsened sulfide-based inclusions serve as an origin of crack initiation during cold forging. Therefore, the upper limit of the amount of S is set to 0.050%. The amount of S is desirably 0.005% to 0.020%.
(Cr: 1.0% to 2.0%)
Cr is an effective element not only for improving hardenability but also for improving temper softening resistance. Additionally, even when the amount of Cr is relatively large, the hardness after spheroidizing annealing is less affected by Cr. Therefore, the lower limit of the amount of Cr is set to 1.0%. However, when the amount of Cr is more than 2.0%, the effect of improving temper softening resistance is saturated and thus the upper limit of the amount of Cr is set to 2.0%. The amount of Cr is desirably 1.3% to 1.6%.
(Mo: 0.02% to 0.8%)
Mo is an effective element for improving hardenability. Si, Mn and Cr cause the hardenability of the surface portion to be deteriorated by being selectively oxidized in the surface portion of the steel during carburizing heating in some cases. In such cases, a layer which is not completely hardened during quenching is formed and causes deterioration in bending fatigue strength and pitching strength. On the other hand, since Mo is less likely to be oxidized compared to the above elements, Mo is effective for reducing the incompletely hardened layer of the surface portion. In order to obtain the effect, the lower limit of the amount of Mo is set to 0.02%. However, when the amount of Mo is excessive, the hardness after spheroidizing annealing is increased and cold forgeability is deteriorated. Thus, the upper limit of the amount of Mo is set to 0.8%. The amount of Mo is desirably 0.05% to 0.5%.
(Al: 0.001% to 0.20%)
Al has an effect of refining austenite grains by forming fine nitrides in the steel. In order to obtain the effect, the lower limit of the amount of Al is set to 0.001%. However, when the amount of Al is more than 0.20%, the effect becomes saturated. Therefore, the upper limit of the amount of Al is set to 0.20%. The amount of Al is desirably 0.015% to 0.050%.
(N: 0.003% to 0.03%)
N has an effect of refining austenite grains by forming nitrides with Al, Nb, or V in the steel. In order to obtain the effect, the lower limit of the amount of N is set to 0.003%. However, when the amount of N is excessive, the hot ductility of the steel is deteriorated and scratches are remarkably formed in the process of producing steel (continuous casting and hot rolling). Therefore, the upper limit of the amount of N is set to 0.03%. The amount of N is desirably 0.007% to 0.02%.
(P: 0.030% or less)
P is an impurity element and is an element which deteriorates the toughness of the steel. Therefore, the amount of P is limited to 0.030% or less. The amount of P is desirably limited to 0.020% or less.
(O: 0.0020% or less)
O is an impurity element and forms oxides with Al, Si, etc. When the amount of O is increased, the amount of so-called oxide-based inclusions is increased and the size thereof is also coarsened. As described later, when coarsened oxide-based inclusions are present, the coarsened oxide -based inclusions serve as an origin of crack initiation during cold forging. Therefore, the amount of O is limited to 0.0020% or less. The amount of O is desirably limited to 0.0015% or less and more desirably to 0.0005% or less.
(Ti: 0.005% or less)
Ti is an element which enters unavoidably and forms nitrides such as TiN in the embodiment. When the amount of Ti is increased, the amount of so-called nitride-based inclusions is increased and the size thereof is coarsened. When coarsened nitride-based inclusions are present, the coarsened nitride-based inclusions serve as an origin of crack initiation during cold forging. Therefore, the amount of Ti is limited to 0.005% or less. The amount of Ti is desirably limited to 0.003% or less.
The case-hardening steel according to the embodiment basically contains the above-described chemical composition and may further contain the following components. The following elements are not necessarily contained in the steel. Therefore, it is not necessary to particularly limit the lower limit of the amount of each component and the lower limit thereof is 0%.
(Cu: 0.2% or less)
Cu is an effective element for improving hardenability similar to Mo. In addition, Cu is an element which is less likely to be oxidized and an effective element for reducing the incompletely hardened layer of the surface portion. In order to obtain these effects, the lower limit of the amount of Cu is desirably set to 0.001%. However, when the amount of Cu is excessive, the hot ductility of the steel is deteriorated and defects are remarkably formed in the process of producing steel (continuous casting and hot rolling). Therefore, the upper limit of the amount of Cu is set to 0.2%. When Cu is contained in the steel and Ni whose amount is about a half of the amount of Cu needs to be contained at the same time, the deterioration of hot ductility is reduced. The amount of Cu is desirably 0.05% to 0.15%.
(Ni: 1.5% or less)
Ni is an effective element for improving hardenability similar to Mo and Cu. In addition, Ni is an element which is less likely to be oxidized and is an effective element for reducing the incompletely hardened layer of the surface portion. In order to obtain these effects, it is desirable to set the lower limit of the amount of Ni to 0.001%. However, since Ni is an element which significantly affects the cost, the upper limit of the amount of Ni is set to 1.5%. The amount of Ni is more desirably 0.05% to 1.0%.
(Nb: 0.10% or less)
Nb has the effect of refining austenite grains by forming fine carbides and nitrides in the steel. In order to obtain the effect, the lower limit of the amount of Nb is desirably set to 0.001%. Particularly, since austenite grains are likely to be coarsened when normalizing or annealing is not performed after cold forging, when the carburizing temperature is a high temperature which is 930° C. or higher, etc., it is effective to increase the amount of Nb carbonitrides to prevent coarsening. Therefore, it is more desirable to set the lower limit of the amount of Nb to 0.015%. However, when the amount of Nb is more than 0.10%, the effect becomes saturated. Therefore, the upper limit of the amount of Nb is set to 0.10%. The upper limit of the amount of Nb is desirably 0.050%.
(V: 0.20% or less)
V has the effect of refining austenite grains by forming fine carbides and nitrides in the steel. In order to obtain the effect, the lower limit of the amount of V is desirably set to 0.01%. However, when the amount of V is more than 0.20%, the effect becomes saturated. Therefore, the upper limit of the amount of V is set to 0.20%. The amount of V is more desirably 0.05% to 0.15%.
(Ca: 0.0050% or less)
Ca has the effect of preventing so-called sulfide-based inclusions from serving as an origin of crack initiation during cold forging by refining so-called sulfide-based inclusions. In order to obtain the effect, it is desirable to set the lower limit of the amount of Ca to 0.0001%. However, when the amount of Ca is more than 0.0050%, the effect becomes saturated. Therefore, the upper limit of the amount of Ca is set to 0.0050%. The amount of Ca is more desirably 0.0005% to 0.0015%.
(Mg: 0.0050% or less)
Mg has the effect of preventing so-called sulfide-based inclusions from serving as an origin of crack initiation during cold forging by refining the sulfide-based inclusions. In order to obtain the effect, it is desirable to set the lower limit of the amount of Mg to 0.0001%. However, when the amount of Mg is more than 0.0050%, the effect becomes saturated. Therefore, the upper limit of the amount of Mg is set to 0.0050%. The amount of Mg is more desirably 0.0005% to 0.0015%.
(Sb: 0.050% or less)
Sb has the effect of suppressing decarburization during hot rolling and spheroidizing annealing. In order to obtain the effect, it is desirable to set the lower limit of the amount of Sb to 0.0001%. However, when the amount of Sb is more than 0.050%, the effect becomes saturated. Therefore, the upper limit of the amount of Sb is set to 0.050%. The amount of Sb is more desirably 0.001% to 0.010%.
Next, in the case-hardening steel according to the embodiment, the amounts of Si, Mn, Cr and Mo will be described from the viewpoint of cold forgeability and temper softening resistance.
In the case-hardening steel according to the embodiment, from the viewpoint of cold forgeability, it is necessary to control the amounts of Si, Mn, Cr and Mo to satisfy the following Equation (1), that is, to make the value of the left side of the following Equation (1) be 25 or less. This is because the limit of the cold forgeability (hardness before cold forging) of a spheroidizing-annealed material has to be determined in consideration of the effect of each of Si, Mn, Cr and Mo on the hardness of the spheroidizing-annealed material. In the left side of the following Equation (1), the reason that the coefficients of each element of Si, Mn, Cr and Mo are different is that the level of the contribution of the elements to cold forgeability (hardness before cold forging) is different.
A desirable range of the left side of the following Equation (1) is 24.5 or lower and a more desirable range thereof is 23 or less.
12×Si(%)+25×Mn(%)+Cr(%)+2×Mo(%)≦25 (1)
In addition, in the case-hardening steel according to the embodiment, from the viewpoint of temper softening resistance, it is necessary to control the amounts of Si, Mn and Cr to make the value of the left side of the following Equation (2) is 50 or more. In powertrain components such as gears and CVT, heat is locally generated at a position in which the powertrain components are brought into contact with other components by the contact while being used and are subjected to tempering to be softened. This softening is a controlling factor of the deterioration of pitching fatigue properties. Accordingly, in order to improve pitching fatigue properties, it is effective to improve hardness after tempering at 300° C. which is the temper softening resistance of the carburized layer. When the value of the left side of the Equation (2) is 50 or more, pitching fatigue properties are improved. The value of the left side is desirably 53 or more and more desirably 55 or more.
31×Si(%)+15×Mn(%)+23×Cr(%)≧50 (2)
Next, in the case-hardening steel according to the embodiment, the size and the number of the sulfide-based inclusions will be described.
The sulfide-based inclusions in the embodiment are inclusions containing S and refer to for example, MnS, CaS, MgS, (Mn,Ca,Mg)S, TiS, Ti(C,S), FeS, etc.
In the case-hardening steel according to the embodiment, when the inclusion evaluation is performed using an extreme value statistical method, it is required that the estimated value of (√area)S which is the maximum size of sulfide-based inclusions present in the estimated area S=30,000 mm2 is 49 μm or less, and the number of the sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm per 1 mm2 is 200 or less.
When the steel is subjected to deformation processing by cold forging with a high working degree and large sulfide-based inclusions are present in the steel, the interface between the inclusion and the matrix serves as an origin of crack initiation and finally, cracks grow to large cold forging cracks in some cases. However, as long as the size of the sulfide-based inclusions is 49 μm or less, the inclusions do not serve as the origin of crack initiation and are harmless. On the other hand, the sulfide-based inclusions having a size of more than 49 μm serve as an origin of crack initiation. Therefore, the upper limit of (√area)S is set to 49 μm.
Since the amount of the sulfide-based inclusions is large compared to the amount of oxide-based inclusions or nitride-based inclusions, the presence frequency of the sulfide-based inclusions is high. In addition, since the sulfide-based inclusions are stretched into a long and narrow shape by hot working, the sulfide-based inclusions significantly affect the cold forging cracks. For example, (√area)S of the sulfide-based inclusions having a length of 20 μm and a thickness of 2 μm is 6.3 μm and is smaller than the maximum size of sulfide-based inclusions (49 μm) limited in the above description. However, when the number of the sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm per 1 mm2 is more than 200, it is similar to the state in which large inclusions are present and crack initiation occurs frequently during cold working. Accordingly, regarding the sulfide-based inclusions, it is necessary to define not only the size of inclusions but also the number of inclusions whose size is equal to or larger than a predetermined size. That is, it is necessary to define the number of the sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm per 1 mm2 to be 200 or less. When the length and the thickness of the sulfide-based inclusions or the number of the sulfide-based inclusions are out of the above ranges, cracks are likely to initiate. When the size of the sulfide-based inclusions is measured, the major axis is the length and the minor axis is the thickness.
Regarding MnS having a length of 20 μm or less, when the thickness is small, the limit is not applied to the thickness. However, when considering a case in which the thickness is very large, for example, a case in which MnS having a length of more than 20 μm is present, the thickness is a length and the length is a thickness, and thus, the limit is applied to the thickness.
The smaller the size of the sulfide-based inclusions and oxide-based inclusions is, the more desirable. Thus, the lower limit of the particle size thereof is 0 μm. In addition, a smaller number of the sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm are desirable. Thus, the lower limit of the number density is 0 piece/mm2.
Next, in the case-hardening steel according to the embodiment, the size of the oxide-based inclusions will be described.
The oxide-based inclusions referred to in the present invention are inclusions containing O and refer to, for example, Al2O3, CaO, Cr2O3, MnO, NbO, SiO2, MgO, ZrO2, TixOy, Nb2O5, FeOx, composite compounds thereof, etc.
In the case-hardening steel according to the embodiment, it is preferable that the estimated value of the maximum size (√area)Ox of oxide-based inclusions present in the estimated area S=30,000 mm2 is 80 μm or less in the inclusion evaluation using the extreme value statistical method.
This is because when the steel is subjected to deformation processing by cold forging with a high working degree and large oxide-based inclusions are present in the steel the interface between the inclusion and the matrix serves as an origin of crack initiation and finally, cracks grow to large cold forging cracks in some cases. However, oxide-based inclusions whose (√area)Ox is 80 μm or less do not serve as the origin of crack initiation and are harmless. On the other hand, the oxide-based inclusions having a size of more than 80 μm serve as an origin of crack initiation. Accordingly, it is necessary to define the size of the oxide-based inclusions as described above.
The case-hardened steel member according to the embodiment can be obtained by subjecting the above-described case-hardening steel to carburizing quenching and tempering treatment or a carburizing nitriding quenching and tempering treatment. That is, the case-hardened steel member is made from case-hardening steel. Therefore, the case-hardened steel member according to this embodiment has substantially the same chemical composition and inclusions as the chemical composition and inclusions of the case-hardening steel according to the above-described embodiment. Accordingly, in order to control the chemical composition and inclusions of the case-hardened steel member, the case-hardening steel may be controlled to have a predetermined chemical composition and inclusions.
However, since the case-hardened steel member is subjected to a carburizing quenching and tempering or a carburizing nitriding quenching and tempering treatment, a surface hardened layer is provided and the case-hardened steel member is different from the case-hardening steel in that the surface hardened layer is provided.
Preferable conditions for producing the case-hardening steel and the case-hardened steel member according to the embodiment will be described.
In the embodiment, in secondary refining, RH vacuum degassing is performed under the conditions of a total treatment time of 30 minutes or longer including a treatment time of 15 minutes or longer at a reduced pressure of 1 Torr or lower (refining process). When refining is performed under the above-described conditions, the size and the number of the oxide-based inclusions can be controlled to predetermined ranges. In addition, in the refining process, the chemical composition is adjusted to be within the above-described preferable range.
Next, the molten steel with the chemical composition adjusted in the refining process is subjected to continuous casting to obtain a slab (casting process). When continuous casting is used to obtain a slab, it is desirable that the casting rate is set to 0.45 m/min or more. When the casting rate is set to 0.45 m/min or more, the size and the number of the sulfide-based inclusions can be controlled to the above-described ranges. When the casting rate is less than 0.45 m/min, coarse sulfide-based inclusions are crystallized and precipitated during solidification of the steel. A desirable casting rate is 0.50 m/min to 1.5 m/min.
Further, when the steel is subjected to casting, it is desirable that the slab is cooled such that the cooling rate from the liquidus temperature to the solidus temperature is 5° C./min to 200° C./min at a ¼ portion of the slab in the thickness direction. When the cooling rate is less than 5° C./min, coarse sulfide-based inclusions are precipitated and also the productivity of continuous casting is deteriorated. Thus, the cooling rate of less than 5° C./min is not desirable. In addition, when the cooling rate is more than 200° C./min, cracks are more likely to initiate in the slab during continuous casting and thus the cooling rate of more than 200° C./min is not desirable.
The cooling condition is related to a secondary dendrite arm spacing. Therefore, by measuring the secondary dendrite arm spacing, the above-described cooling rate can be calculated. Specifically, the cooling rate can be obtained by calculation using the following Equitation (3) using the spacing between the secondary dendrite arms in the solidification structure of the solidified slab in the thickness direction.
Rc=(λ2/770)(−1/0.41) (3)
Rc: cooling rate (° C./min), and λ2: spacing between secondary dendrite arms (μm).
A bloom obtained by the above-described casting process is subjected to blooming to obtain a billet (blooming process). The heating temperature during blooming is desirably 1240° C. or higher since unavoidably formed coarse sulfides are temporarily solid-soluted in the matrix. The heating temperature is more desirably 1260° C. or higher. The reduction of area of the blooming leads to a reduction in the thickness of the sulfide-based inclusions and thus it is necessary to set the reduction of area to 40% or more. The reduction of area is desirably 45% or more. In addition, when the cooling rate during blooming or after blooming is low, the solid-soluted MnS is precipitated as a coarse sulfide again. Thus, it is necessary to set the cooling rate to 1240° C. to 1000° C. in the cooling process during blooming or after blooming to 0.7° C./s or more. The cooling rate is more desirably 1.5° C./s or more. The cooling rate is a cooling rate obtained from the actually measured value of the surface temperature.
In order to form the billet into a case-hardening steel (steel bar or wire rod), steel bar rolling or wire rod rolling is performed. At the time of heating in the steel bar rolling or wire rod rolling, to prevent MnS from growing and being coarsened, the heating temperature is desirably set to 1200° C. or lower. The heating temperature is more desirably 1000° C. to 1150° C. In addition, the total reduction of area until the process of rolling the billet into a steel bar or a wire rod is completed (total reduction of area in blooming and steel bar rolling or wire rod rolling) is set to 65% or more. When the total reduction is less than 65%, the thickness reduction along with the stretching of the sulfide-based inclusions is not sufficient and thus the number of sulfide-based inclusions having a large thickness which are harmful to cold forging crack initiation cannot be reduced. A suitable range of the total reduction of area is 90% or more.
The above-described case-hardening steel is further subjected to carburizing quenching and tempering treatment or a carburizing nitriding quenching and tempering treatment to obtain a case-hardened steel member. The carburizing quenching and tempering or carburizing nitriding quenching and tempering may be performed by known methods.
Hereinafter, the present invention will be further described using Examples. Converter molten steels having the compositions (chemical composition) shown in Tables 1-1 and 1-2 were subjected to RH vacuum degassing under the conditions shown in Table 2 and subsequently subjected to continuous casting under the conditions shown in Table 3, and then subjected to soaking as required. Through blooming, rectangular rolled steels (billets), having a square-shaped cross section whose length of one side is 162 mm were obtained. The balance in Tables 1-1 and 1-2 includes iron and impurities and the blank indicates that the component is intentionally not added.
Next, working was performed by hot rolling under the conditions shown in Table 4 to form the billets into steel bars and then some of the steel bars were subjected to spheroidizing annealing (SA) under the conditions of
Then, columnar test pieces having a diameter of 16 mm and a length of 24 mm were taken from the materials by cutting. The columnar test pieces were subjected to upset cold working under the conditions of an upset rate of 50% and a strain rate of 1.0. Next, in order to simulate carburizing, the cold-worked columnar test pieces were heated and retained at 950° C. for 5 hours and then immediately water-cooled to freeze the austenite structure after the carburizing simulation as the prior austenite grain boundary of the martensite structure. Next, the prior austenite grain structure in the cross section of each of the test pieces which was subjected to the carburizing simulation in the rolling direction was observed and the JIS grain size number was measured. A coarse particle was defined as a particle with a grain size number of 5 or lower in JIS G 0551 and even when one coarse particle was formed in the view field in the entire cross section, it was determined that coarse particle was present.
The case-hardening steel and the case-hardened steel member of the present invention may be subjected to SA, however, SA is not necessarily performed. When cold working is not performed or cold working is possible without performing SA in actual production of a component, SA may not be performed. In this case, the steel can be used as high-strength steel.
First, the Vickers hardness (measurement load of 10 kgf) at a ¼ depth position of the diameter of the steel bar and the casting raw material was measured according to JIS Z 2244. The number of measurement points is 4 per material and the average value was obtained. For the steel having a hardness of HV 155 or higher, it steel was determined that cold forgeability was deteriorated since the deformation resistance of the steel during cold forging was increased and the die life was remarkably reduced.
At the position in the vicinity of ¼ of the diameter of the steel bar or at the position in the vicinity of ¼ of the diameter of the casting raw material, optical microscope observation was performed and inclusion measurement was performed. The estimated value of the maximum size (√area)S of sulfide-based inclusion and the maximum size (√area)Ox of oxide-based inclusions present in the estimated area S=30,000 mm2 was set to 10 mm×10 mm in one visual field area (investigation reference area: S0) and optical microscope observation was performed in 30 view fields so as to avoid overlapping of the area S0. The maximum size (√area) of inclusions present in the respective 30 visual fields was measured and plotted to extreme value probability paper. The maximum size (√area) of inclusions was determined by estimating the maximum size of inclusions while setting the estimated area S to 30,000 mm2. In the inclusion measurement, each of oxides (oxide-based inclusions) and sulfides (sulfide-based inclusions) were independently evaluated.
The number of sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm in each visual field was measured while above-described inclusion measurement was performed. The numbers of the entire 30 visual fields were added and the added value was divided by the total measurement area (3,000 mm2) to measure the presence number of sulfide-based inclusions having a length of more than 20 μm and a thickness of more than 2 μm per area of 1 mm2.
Next, as an index of crack initiation in the steel during cold forging, the limit compression ratio was measured. Test pieces for limit compression ratio measurement (φ6 mm×9 mm, notch configuration: 30°, depth: 0.8 mm, radius of curvature at tip end: 0.15 mm) were prepared by taking the test pieces from the steel bars and casting raw materials in a direction parallel to in the longitudinal direction. In the measurement of the limit compression ratio, cold compression was performed using a binding die at a speed of 10 mm/min, the compression was stopped when fine cracks having a length of 0.5 mm or more initiated in the vicinity of the notch, the compression ratio was calculated when the compression is stopped, and the obtained compression ratio was set to a compression ratio at crack initiation. The test was performed (n=10) at one level to obtain a compression ratio of a cumulative failure probability of 50% and the obtained compression ratio was set to the index of the limit working ratio. Since the limit compression ratio of a SA material of JIS-SCr 420 was about 65%, the steel having a high value of 68% or more, which was apparently higher than the above value, was determined to have an excellent limit working ratio and contrarily the steel having a value of less than 68% was determined to have a deteriorated working ratio.
Next, the hardness after tempering at 300° C. which is an index of the pitching resistance of the component after carburizing was measured. In order to measure the hardness after tempering at 300° C., first, test pieces for carburizing (φ20 mm×30 mm) was taken from the steel bars as a material (SA material and non-SA material). Then, gas carburizing was performed by a gas type converter. The gas carburizing was performed under the condition of an atmosphere temperature of 950° C. and a retaining time of 5 hours, an atmosphere temperature of 850° C. and a retaining time of 0.5 hours, oil quenching at 130° C., a tempering temperature of 150° C. and a retaining time of 90 minutes at a gas potential of 0.8% in order. Next, for the investigation of the surface structure, a section in the vicinity of the center portion of the test piece in the longitudinal direction was cut in a direction orthogonal to the longitudinal direction and a microscope sample having a cross section was prepared. The sample was subjected to corrosion with 2% nital for structure observation and the surface portion of the carburized layer was observed with a microscope. The depth of an incompletely hardened layer formed on the surface portion of the carburized layer (a layer in which a non-martensite structure mainly composed of pearlite and/or bainite is present) was measured. When the depth of the insufficient hardened layer was deep, the pitching properties are deteriorated and the depth of the incompletely hardened layer of JIS-SCr 420 is about 25 μm. Thus, when the depth of the incompletely hardened layer was deeper than 25 μm, it was determined that the pitching properties were insufficiently improved.
In addition, in order to obtain the hardness after tempering at 300° C., tempering was further performed at a tempering temperature of 300° C. for a retaining time of 90 minutes. Then, a portion in the vicinity of the center portion of the test piece in the longitudinal direction was cut in a direction vertical to the longitudinal direction and the Vickers hardness of the cross section was measured. The hardness measurement position was set to a position having a depth of 50 μm from the surface and the measurement load was set to 300 gf. Further, the measurement was performed at 5 points per test piece and the average value was obtained. Since the temper hardness of JIS-SCr 420 at 300° C. is HV 640, the steel having a value of HV 670 or more, which was apparently higher than the above value, was determined to have excellent pitching properties and the steel having a value of less than HV 670 was determined to have insufficient pitching properties.
In Table 2, the effect of RH conditions was summarized. In RH condition No. 1-3 of Table 2, both of the total treatment time of RH vacuum degassing and the treatment time in a reduced atmosphere of 1 Torr or less were out of the desirable ranges. In addition, in RH condition No. 1-4, the treatment time in a reduced atmosphere of 1 Torr or less was out of the desirable range. Further, in RH condition No. 1-B, the total treatment time of RH vacuum degassing was out of the desirable range. In Production condition Nos. 20, 23, 42, a, b, c, d, e and f adopting these conditions, floating oxides in the molten steel were not sufficiently removed and the size of oxide-based inclusions present in the steel bar was large. As a result, the limit compression ratio was deteriorated. In contrast, in Production Nos. 1, 9 and 2 adopting RH condition Nos. 1-1, 1-2 and 1-A in which the RH conditions were appropriate, the size of oxide-based inclusions was small and the limit compression ratio of the SA material was good.
In Table 3, the effect of casting conditions was summarized. In Casting condition No. 2-8 of Table 3, the casting rate was out of the desirable range. In addition, in Casting condition No. 2-9, a cooling rate from the liquidus temperature to the solidus temperature at ¼ portion of the slab in the thickness direction was low and thus the size of sulfide-based inclusions present in the steel bar was large. As a result, all Production Nos. 64, 65, 66 and 67 adopting Casting condition No. 2-8 or 2-9, has a low limit compression ratio. In contrast, in Production Nos. No.1, 2 and 53 to 58 adopting Casting condition Nos. 2-1 to 2-7 in which the continuous casting condition was appropriate, the size of sulfide-based inclusions were small and the limit compression ratio was good.
In Table 4, the effect of rolling conditions was summarized. The total reduction of area of Rolling condition Nos. 3-6 and 3-B of Table 4 was out of the desirable range. As a result, in Production Nos. 68 and 69 adopting these conditions, the thickness of MnS was not sufficiently reduced by rolling and thus there were a large number of sulfide-based inclusions having a large thickness. Accordingly, in Production Nos. 68 and 69, the limit compression ratio was deteriorated. In contrast, in Production Nos. 1 and 59 to 63 adopting the rolling condition No. 3-1 to 3-5 and 3-A in which the total reduction of area of hot rolling was appropriate, the thickness was large, the number of stretched sulfide-based inclusions was small and the limit compression ratio was also good.
In Tables 5-1, 5-2, 6 and 7, the inclusion measurement results and properties of steels obtained under each production condition are shown. In Tables 5-1, 5-2 and 6, the result of a material which has been subjected to SA and in Table 7, the results of a material which has not been subjected to SA are shown.
As seen from Tables 5-1, 5-2 and 6, in Production Nos. 1 to 15and 53 to 63 in which all conditions were within the ranges of the present invention, all of the hardness after SA, the limit compression ratio, the hardness of the carburized layer after tempering at 300° C., and the thickness of the incompletely hardened layer were excellent. In addition, in Production Nos. 1, 8, 9 and 11 including Nb, coarse particles were not observed.
In contrast, in Production Nos.16 to 52, 64 to 69 and a to f in which at least one of the chemical composition and the production condition was out of the desirable range, any of the hardness after SA, the limit compression ratio, the hardness of the carburized layer after tempering at 300° C., or the thickness of the incompletely hardened layer did not satisfy the desired value. Further, in Production Nos. 20, 23, 31, 34, 42 and 45, the amount of O was large and the maximum √area of oxide-based inclusions was out of the range of the present invention. In addition, in Production Nos. 22, 33 and 44, the amount of S was beyond the range of the present invention, and thus the maximum √area of sulfide-based inclusions was out of the range of the present invention.
As seen from Table 7, even with respect to the material to which SA was not subjected, in Production Nos. 101 to 115 in which all conditions were within the range of the present invention, the hardness of the carburized layer after tempering at 300° C. and the thickness of the incompletely hardened layer were excellent. On the hand, in Production Nos. 116 to 118, 124, 126, 128, 129, 135, 136, 138, 139, 146, 148, 150 and 151 in which at least one of the chemical composition and the production condition was out of the desirable range, the hardness after tempering at 300° C. or the thickness of the incompletely hardened layer was deteriorated. The same tendency was applied to the SA material.
According to the case-hardening steel and the case-hardened steel member of the present invention, it is possible to provide a case-hardened steel having excellent temper softening resistance and cold forgeability and a case-hardened steel member. In addition, the use of the case-hardening steel and the case-hardened steel member enables reduction of the production cost of gears and contribution to high output performance and improvement of fuel efficiency, etc. for automobiles, construction machines, and industrial machines.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2013-087857 | Apr 2013 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2014/060800 | 4/16/2014 | WO | 00 |
