U.S. patent application number 15/559950 was filed with the patent office on 2018-04-19 for steel for nitrocarburizing and nitrocarburized component, and methods of producing same.
This patent application is currently assigned to JFE STEEL CORPORATION. The applicant listed for this patent is JFE STEEL CORPORATION. Invention is credited to Takashi IWAMOTO, Masayuki KASAI, Kunikazu TOMITA.
Application Number | 20180105919 15/559950 |
Document ID | / |
Family ID | 56977170 |
Filed Date | 2018-04-19 |
United States Patent
Application |
20180105919 |
Kind Code |
A1 |
KASAI; Masayuki ; et
al. |
April 19, 2018 |
STEEL FOR NITROCARBURIZING AND NITROCARBURIZED COMPONENT, AND
METHODS OF PRODUCING SAME
Abstract
The steel for nitrocarburizing has a chemical composition that
contains C: 0.01% or more and less than 0.20%, Si: 1.0% or less,
Mn: 1.5% or more and 3.0% or less, P: 0.02% or less, S: 0.06% or
less, Cr: 0.30% or more and 3.0% or less, Mo: 0.005% or more and
0.40% or less, V: 0.02% or more and 0.5% or less, Nb: 0.003% or
more and 0.20% or less, Al: 0.010% or more and 2.0% or less, Ti:
more than 0.005% and less than 0.025%, N: 0.0200% or less, Sb:
0.0005% or more and 0.02% or less, and the balance consisting of Fe
and incidental impurities; and a steel microstructure that contains
bainite phase in an area ratio of more than 50%.
Inventors: |
KASAI; Masayuki;
(Chiyoda-ku, Tokyo, JP) ; IWAMOTO; Takashi;
(Chiyoda-ku, Tokyo, JP) ; TOMITA; Kunikazu;
(Chiyoda-ku, Tokyo, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
JFE STEEL CORPORATION |
Chiyoda-ku, Tokyo |
|
JP |
|
|
Assignee: |
JFE STEEL CORPORATION
Chiyoda-ku, Tokyo
JP
|
Family ID: |
56977170 |
Appl. No.: |
15/559950 |
Filed: |
March 24, 2016 |
PCT Filed: |
March 24, 2016 |
PCT NO: |
PCT/JP2016/001721 |
371 Date: |
September 20, 2017 |
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C21D 8/0226 20130101;
C22C 38/002 20130101; C22C 38/38 20130101; C22C 38/26 20130101;
C22C 38/48 20130101; C22C 38/58 20130101; C21D 7/13 20130101; C22C
38/20 20130101; C21D 2211/004 20130101; C23C 8/26 20130101; C22C
38/28 20130101; C22C 38/46 20130101; C21D 1/74 20130101; C22C 38/24
20130101; C22C 38/30 20130101; C22C 38/008 20130101; C23C 8/32
20130101; C22C 38/001 20130101; C22C 38/32 20130101; C21D 1/06
20130101; C22C 38/06 20130101; C22C 38/02 20130101; C22C 38/50
20130101; C21D 8/005 20130101; C22C 38/22 20130101; C22C 38/54
20130101; C21D 6/004 20130101; C21D 8/0263 20130101; C22C 38/60
20130101; C21D 2211/002 20130101; C22C 38/42 20130101; C21D 6/005
20130101; C21D 6/008 20130101; C22C 38/44 20130101; C21D 1/20
20130101 |
International
Class: |
C23C 8/32 20060101
C23C008/32; C22C 38/58 20060101 C22C038/58; C22C 38/50 20060101
C22C038/50; C22C 38/48 20060101 C22C038/48; C22C 38/46 20060101
C22C038/46; C22C 38/44 20060101 C22C038/44; C22C 38/06 20060101
C22C038/06; C22C 38/02 20060101 C22C038/02; C22C 38/42 20060101
C22C038/42; C22C 38/54 20060101 C22C038/54; C22C 38/00 20060101
C22C038/00; C21D 8/00 20060101 C21D008/00; C21D 6/00 20060101
C21D006/00 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 24, 2015 |
JP |
2015-061400 |
Claims
1. A steel for nitrocarburizing comprising: a chemical composition
that contains, in mass %, C: 0.01% or more and less than 0.20%, Si:
1.0% or less, Mn: 1.5% or more and 3.0% or less, P: 0.02% or less,
S: 0.06% or less, Cr: 0.30% or more and 3.0% or less, Mo: 0.005% or
more and 0.40% or less, V: 0.02% or more and 0.5% or less, Nb:
0.003% or more and 0.20% or less, Al: 0.010% or more and 2.0% or
less, Ti: more than 0.005% and less than 0.025%, N: 0.0200% or
less, Sb: 0.0005% or more and 0.02% or less, the balance consisting
of Fe and incidental impurities, with the chemical composition
satisfying either one of the following relations: in a case where
the C content is 0.01% or more and 0.10% or less,
(S/32)/(Ti/48)+(N/14)/(Ti/48).ltoreq.13.0, and in a case where the
C content is more than 0.10% and less than 0.20%,
2(S/32)/(Ti/48)+3(N/14)/(Ti/48).ltoreq.35.0; and a steel
microstructure that contains bainite phase in an area ratio of more
than 50%.
2. The steel for nitrocarburizing according to claim 1, wherein the
steel composition further contains, in mass %, one or more selected
from the group consisting of B: 0.0100% or less, Cu: 0.3% or less,
Ni: 0.3% or less, W: 0.3% or less, Co: 0.3% or less, Hf: 0.2% or
less, Zr: 0.2% or less, Pb: 0.2% or less, Bi: 0.2% or less, Zn:
0.2% or less, and Sn: 0.2% or less.
3-4. (canceled)
5. A component comprising: a core part comprising the chemical
composition and the steel microstructure as recited in claim 1; and
a surface layer part comprising a chemical composition with high
nitrogen and carbon contents relative to the chemical composition
of the core part.
6. The component according to claim 5, wherein precipitates
containing V and Nb are dispersed in the bainite phase.
7. A method of producing a steel for nitrocarburizing, comprising:
subjecting a steel to hot working with a heating temperature of
950.degree. C. or higher and a finishing temperature of 800.degree.
C. or higher, the steel comprising a chemical composition that
contains, in mass %, C: 0.01% or more and less than 0.20%, Si: 1.0%
or less, Mn: 1.5% or more and 3.0% or less, P: 0.02% or less, S:
0.06% or less, Cr: 0.30% or more and 3.0% or less, Mo: 0.005% or
more and 0.40% or less, V: 0.02% or more and 0.5% or less, Nb:
0.003% or more and 0.20% or less, Al: 0.010% or more and 2.0% or
less, Ti: more than 0.005% and less than 0.025%, N: 0.0200% or
less, Sb: 0.0005% or more and 0.02% or less, the balance consisting
of Fe and incidental impurities, with the chemical composition
satisfying either one of the following relations: in a case where
the C content is 0.01% or more and 0.10% or less,
(S/32)/(Ti/48)+(N/14)/(Ti/48).ltoreq.13.0, and in a case where the
C content is more than 0.10% and less than 0.20%,
2(S/32)/(Ti/48)+3(N/14)/(Ti/48).ltoreq.35.0; and then cooling the
steel at a cooling rate of higher than 0.4.degree. C./s at least in
a temperature range of 700.degree. C. to 550.degree. C.
8. The method of producing a steel for nitrocarburizing according
to claim 7, wherein the steel composition further contains, in mass
%, one or more selected from the group consisting of B: 0.0100% or
less, Cu: 0.3% or less, Ni: 0.3% or less, W: 0.3% or less, Co: 0.3%
or less, Hf: 0.2% or less, Zr: 0.2% or less, Pb: 0.2% or less, Bi:
0.2% or less, Zn: 0.2% or less, and Sn: 0.2% or less.
9-10. (canceled)
11. A method of producing a component, comprising: processing the
steel for nitrocarburizing obtainable by the method as recited in
claim 7 into a desired shape; and then subjecting the steel for
nitrocarburizing to nitrocarburizing treatment at 550.degree. C. to
700.degree. C. for 10 minutes or more.
12. A component comprising: a core part comprising the chemical
composition and the steel microstructure as recited in claim 2; and
a surface layer part comprising a chemical composition with high
nitrogen and carbon contents relative to the chemical composition
of the core part.
13. The component according to claim 12, wherein precipitates
containing V and Nb are dispersed in the bainite phase.
14. A method of producing a component, comprising: processing the
steel for nitrocarburizing obtainable by the method as recited in
claim 8 into a desired shape; and then subjecting the steel for
nitrocarburizing to nitrocarburizing treatment at 550.degree. C. to
700.degree. C. for 10 minutes or more.
Description
TECHNICAL FIELD
[0001] The present disclosure relates to steel for nitrocarburizing
and component obtained from the steel for nitrocarburizing, and
methods of producing these. The components according to the
disclosure exhibit hot forgeability and excellent fatigue
properties after nitrocarburizing treatment and are suitable for
use as components for automobiles and construction machinery.
BACKGROUND
[0002] Since excellent fatigue properties are desired for machine
structural components, such as automobile gears, surface hardening
is generally performed. Examples of well-known surface hardening
treatment include carburizing treatment, induction quench
hardening, and nitriding treatment.
[0003] Among these, in carburizing treatment, C is immersed and
diffused in high-temperature austenite region and a deep hardening
depth is obtained. Therefore, carburizing treatment is effective in
improving fatigue strength. However, since heat treatment
distortion occurs by carburizing treatment, it was difficult to
apply such treatment to components that require severe dimensional
accuracy from the perspective of noise or the like.
[0004] Further, in induction quench hardening, since quenching is
performed on a surface layer part by high frequency induction
heating, heat treatment distortion is generated, and therefore
results in poor dimensional accuracy as in the case with
carburizing treatment.
[0005] On the other hand, in nitriding treatment, surface hardness
is increased by immersing and diffusing nitrogen in a relatively
low temperature range at or below the Ac.sub.1 transformation
temperature, and therefore there is no possibility of heat
treatment distortion such as mentioned above. However, there were
problems that the treatment requires a long time of 50 hours to 100
hours, and then it is necessary to remove brittle compound layers
on the surface layer after performing the treatment.
[0006] Therefore, nitrocarburizing treatment in which treatment is
performed at a treatment temperature almost equal to nitriding
treatment temperature and in a shorter treatment time was
developed, and in recent years, such treatment has been widely used
for machine structural components and the like. During this
nitrocarburizing treatment, N and C are simultaneously infiltrate
and diffused in a temperature range of 500.degree. C. to
600.degree. C. to harden the surface, and the treatment time can be
made half of what is required for conventional nitriding
treatment.
[0007] However, whereas the above-mentioned carburizing treatment
enables to increase the core hardness by quench hardening,
nitrocarburizing treatment does not increase core hardness since
the treatment is performed at a temperature at or below the
transformation point of steel. Therefore, fatigue strength of the
nitrocarburized material is inferior compared to the carburized
material.
[0008] In order to improve the fatigue strength of the
nitrocarburized material, quenching and tempering are usually
performed before nitrocarburizing to increase the core hardness.
However, the resulting fatigue properties cannot be considered
sufficient. Furthermore, this approach increases production costs
and reduces mechanical workability.
[0009] To address these issues, JPH559488A (PTL 1) proposes a steel
for nitrocarburizing which can exhibit high bending fatigue
strength after subjection to nitrocarburizing treatment by adding
Ni, Al, Cr, Ti, and the like to the steel.
[0010] Specifically, this steel is subjected to nitrocarburizing
treatment, whereby the core part is age hardened with Ni--Al- or
Ni--Ti-based intermetallic compounds or Cu compounds, while in the
surface layer part, for example, nitrides or carbides of Cr, Al,
Ti, and the like are caused to precipitate and harden in the
nitride layer to improve bending fatigue strength.
[0011] JP200269572A (PTL 2) proposes a steel for nitrocarburizing
which provides excellent bending fatigue properties after
subjection to nitrocarburizing treatment by subjecting a steel
containing 0.5% to 2% of Cu to extend forging by hot forging, and
then air cooling the steel so as to have a microstructure mainly
composed of ferrite with solute Cu dissolved therein, and then
causing precipitation hardening of Cu during nitrocarburizing
treatment at 580.degree. C. for 120 minutes and precipitation
hardening of carbonitrides of Ti, V and Nb.
[0012] JP2010163671A (PTL 3) proposes a steel for nitrocarburizing
obtained by dispersing Ti--Mo carbide, and further dispersing
carbides containing one or more of Nb, V, and W.
[0013] JP2013166997A (PTL 4) proposes a steel material for
nitrocarburizing that exhibits excellent fatigue strength by
providing a steel containing V and Nb with a microstructure in
which bainite is dominantly present prior to nitriding to suppress
precipitation of V and Nb carbonitrides, and these carbonitrides
are caused to precipitate upon nitriding to increase core
hardness.
CITATION LIST
Patent Literature
[0014] PTL 1: JPH559488A
[0015] PTL 2: JP200269572A
[0016] PTL 3: JP2010163671A
[0017] PTL 4: JP2013166997A
[0018] PTL 5: JP5567747B
[0019] PTL 6: JP201132537A
SUMMARY
Technical Problem
[0020] However, in the nitrocarburized steel described in PTL 1,
although bending fatigue strength is improved by precipitation
hardening of Ni--Al- or Ni--Ti-based intermetallic compounds, Cu,
and the like, workability cannot be said to be sufficiently
secured, and a high Ni content leads to the problem of increased
production costs.
[0021] The steel for nitrocarburizing in PTL 2 requires Cu, Ti, V,
and Nb to be added to the steel in relatively large amounts, and
has the problem of high production costs.
[0022] Further, the steel for nitrocarburizing in PTL 3 contains Ti
and Mo in relatively large amounts, and also has the problem of
high cost.
[0023] In the case of the steel materials for nitriding in PTLs 4
and 5, to ensure machinability by cutting, the increase of bainite
hardness is suppressed by reducing C content. Hardenability
decreases as the C content decreases, which makes it difficult to
form a bainitic microstructure. To compensate for this, Mn, Cr, and
Mo, which are effective for improving hardenability, are added to
the steel to promote the formation of a bainitic microstructure. In
the case of a rolled material being produced by continuous casting,
however, surface defects, called "continuous cast cracks", easily
form on the surface of the cast steel, leading to the problem of
reduced manufacturability.
[0024] In addition, the steel for nitriding in JP201132537A (PTL 6)
has a problem that surface cracks are liable to occur during
continuous casting, resulting in poor manufacturability.
[0025] It could thus be helpful to provide a steel for
nitrocarburizing whose mechanical workability before
nitrocarburizing treatment is guaranteed by ensuring fatigue
resistance without causing the steel to harden before subjection to
nitrocarburizing treatment, and a method of producing the same. It
could also be helpful to provide a nitrocarburized component whose
fatigue properties can be improved by increasing the surface
hardness through nitrocarburizing treatment after machining.
Solution to Problem
[0026] In order to solve the above problems, we intensely
investigated the influence of the chemical composition and
microstructure of steel.
[0027] As a result, we discovered that by arranging a steel to have
a chemical composition properly that contains V and Nb in
appropriate amounts, Sb in small amounts, and a steel
microstructure that contains bainite phase in an area ratio of more
than 50%, the resulting steel may have excellent mechanical
workability, and that after the steel being subjected to
nitrocarburizing treatment, fine precipitates containing V and Nb
at their cores are caused to dispersedly precipitate to increase
core hardness, and excellent fatigue properties can be
obtained.
[0028] The present disclosure was completed through additional
examination based on the above discoveries.
[0029] Specifically, the primary features of this disclosure are as
described below.
[0030] 1. A steel for nitrocarburizing comprising: a chemical
composition that contains (consists of), in mass %, C: 0.01% or
more and less than 0.20%, Si: 1.0% or less, Mn: 1.5% or more and
3.0% or less, P: 0.02% or less, S: 0.06% or less, Cr: 0.30% or more
and 3.0% or less, Mo: 0.005% or more and 0.40% or less, V: 0.02% or
more and 0.5% or less, Nb: 0.003% or more and 0.20% or less, Al:
0.010% or more and 2.0% or less, Ti: more than 0.005% and less than
0.025%, N: 0.0200% or less, Sb: 0.0005% or more and 0.02% or less,
the balance consisting of Fe and incidental impurities, with the
chemical composition satisfying either one of the following
relations:
[0031] in a case where the C content is 0.01% or more and 0.10% or
less,
(S/32)/(Ti/48)+(N/14)/(Ti/48).ltoreq.13.0, and
in a case where the C content is more than 0.10% and less than
0.20%,
2(S/32)/(Ti/48)+3(N/14)/(Ti/48).ltoreq.35.0; and
a steel microstructure that contains bainite phase in an area ratio
of more than 50%.
[0032] 2. The steel for nitrocarburizing according to 1., wherein
the steel composition further contains, in mass %, one or more
selected from the group consisting of B: 0.0100% or less, Cu: 0.3%
or less, and Ni: 0.3% or less.
[0033] 3. The steel for nitrocarburizing according to 1. or 2.,
wherein the steel composition further contains, in mass %, one or
more selected from the group consisting of W: 0.3% or less, Co:
0.3% or less, Hf: 0.2% or less, and Zr: 0.2% or less.
[0034] 4. The steel for nitrocarburizing according to 1., 2., or
3., wherein the steel composition further contains, in mass %, one
or more selected from the group consisting of Pb: 0.2% or less, Bi:
0.2% or less, Zn: 0.2% or less, and Sn: 0.2% or less.
[0035] 5. A component comprising: a core part comprising the
chemical composition and the steel microstructure as recited in any
one of 1. to 4.; and a surface layer part comprising a chemical
composition with high nitrogen and carbon contents relative to the
chemical composition of the core part.
[0036] 6. The component according to 5., wherein precipitates
containing V and Nb are dispersed in the bainite phase.
[0037] 7. A method of producing a steel for nitrocarburizing,
comprising: subjecting a steel to hot working with a heating
temperature of 950.degree. C. or higher and a finishing temperature
of 800.degree. C. or higher, the steel comprising a chemical
composition that contains (consists of), in mass %, C: 0.01% or
more and less than 0.20%, Si: 1.0% or less, Mn: 1.5% or more and
3.0% or less, P: 0.02% or less, S: 0.06% or less, Cr: 0.30% or more
and 3.0% or less, Mo: 0.005% or more and 0.40% or less, V: 0.02% or
more and 0.5% or less, Nb: 0.003% or more and 0.20% or less, Al:
0.010% or more and 2.0% or less, Ti: more than 0.005% and less than
0.025%, N: 0.0200% or less, Sb: 0.0005% or more and 0.02% or less,
and the balance consisting of Fe and incidental impurities, with
the chemical composition satisfying either one of the following
relations:
[0038] in a case where the C content is 0.01% or more and 0.10% or
less,
(S/32)/(Ti/48)+(N/14)/(Ti/48).ltoreq.13.0, and
in a case where the C content is more than 0.10% and less than
0.20%,
2(S/32)/(Ti/48)+3(N/14)/(Ti/48).ltoreq.35.0; and
then cooling the steel at a cooling rate of higher than 0.4.degree.
C./s at least in a temperature range of 700.degree. C. to
550.degree. C.
[0039] 8. The method of producing a steel for nitrocarburizing
according to 7., wherein the steel composition further contains, in
mass %, one or more selected from the group consisting of B:
0.0100% or less, Cu: 0.3% or less, and Ni: 0.3% or less.
[0040] 9. The method of producing a steel for nitrocarburizing
according to 7. or 8., wherein the steel composition further
contains, in mass %, one or more selected from the group consisting
of W: 0.3% or less, Co: 0.3% or less, Hf: 0.2% or less, and Zr:
0.2% or less.
[0041] 10. The method of producing a steel for nitrocarburizing
according to 7., 8., or 9., wherein the steel composition further
contains, in mass %, one or more selected from the group consisting
of Pb: 0.2% or less, Bi: 0.2% or less, Zn: 0.2% or less, and Sn:
0.2% or less.
[0042] 11. A method of producing a component, comprising:
processing the steel for nitrocarburizing obtainable by the method
as recited in any one of 7. to 10. into a desired shape; and then
subjecting the steel for nitrocarburizing to nitrocarburizing
treatment at 550.degree. C. to 700.degree. C. for 10 minutes or
more.
Advantageous Effect
[0043] The present disclosure enables producing a steel for
nitrocarburizing that is excellent in mechanical workability with
an inexpensive chemical composition. By subjecting the steel for
nitrocarburizing to nitrocarburizing treatment, it is possible to
obtain a component having fatigue properties comparable to or
better than, for example, JIS SCr420 steel subjected to carburizing
treatment. Therefore, the component disclosed herein is very useful
when applied to mechanical structural components such as automotive
parts.
BRIEF DESCRIPTION OF THE DRAWING
[0044] FIG. 1 schematically illustrates the steps carried out to
produce a nitrocarburized component.
DETAILED DESCRIPTION
[0045] The following describes the present disclosure in
detail.
[0046] Firstly, reasons for limiting the chemical composition to
the aforementioned ranges in the disclosure will be described. The
% representations below indicating the chemical composition are in
mass % unless stated otherwise.
[0047] C: 0.01% or More and Less than 0.20%
[0048] C is added for the purpose of bainite phase formation and
for securing strength. However, if the C content is less than
0.01%, it is not possible to obtain a sufficient amount of bainite
phase and the amount of V and Nb precipitates formed after
nitrocarburizing treatment is insufficient, making it difficult to
guarantee sufficient strength. Therefore, the C content is set to
0.01% or more. On the other hand, if the C content is 0.20% or
more, the formed bainite phase increases in hardness, thereby
causing mechanical workability and fatigue properties to
deteriorate. Therefore, the C content is set to less than 0.20%.
More preferably, the C content is 0.04% or more and 0.18% or
less.
[0049] Si: 1.0% or Less
[0050] Si is added for its usefulness for deoxidation and bainite
phase formation purposes. If the Si content is more than 1.0%,
machinability by cutting and cold workability deteriorate due to
solid solution hardening of the ferrite and bainite phases.
Therefore, the Si content is set to 1.0% or less. The Si content is
preferably 0.8% or less, and more preferably 0.7% or less. For Si
to effectively contribute to deoxidation, it is preferable to set
the Si content to 0.01% or more.
[0051] Mn: 1.5% or More and 3.0% or Less
[0052] Mn is added for its usefulness for bainite phase formation
and strength enhancement purposes. However, if the Mn content is
less than 1.5%, less bainite phase forms, and V and Nb precipitates
are caused to form before nitrocarburizing treatment, resulting in
increased hardness before nitrocarburizing. Additionally, such a
low Mn content decreases the absolute amount of V and Nb
precipitates remaining after nitrocarburizing treatment, and ends
up lowering the hardness after nitrocarburizing, making it
difficult to guarantee sufficient strength. Therefore, the Mn
content is set to 1.5% or more. If it exceeds 3.0%, however,
continuous casting cracks are more likely to occur, causing
machinability by cutting and cold workability to deteriorate.
Therefore, the Mn content is set to 3.0% or less. The Mn content is
preferably in a range of 1.5% to 2.5%.
[0053] P: 0.02% or Less
[0054] P segregates at austenite grain boundaries, and lowers grain
boundary strength, thereby making continuous casting cracks more
likely to occur. This also lowers strength and toughness.
Therefore, the P content is desirably kept as small as possible,
yet a content of up to 0.02% is tolerable. As setting the content
of P to be less than 0.001% requires a high cost, it suffices in
industrial terms to reduce the P content to 0.001%.
[0055] S: 0.06% or Less
[0056] S is a useful element that forms MnS in the steel to improve
machinability by cutting. S content exceeding 0.06%, however,
causes deterioration of toughness. Therefore, the S content is set
to 0.06% or less. Further, S content exceeding 0.06% makes
continuous casting cracks more likely to occur. Therefore, the S
content is set to 0.04% or less.
[0057] For S to achieve an effect of improving machinability by
cutting, the S content is preferably set to 0.002% or more.
[0058] Cr: 0.30% or More and 3.0% or Less
[0059] Cr is added for its usefulness for the purpose of bainite
phase formation. Cr also has an effect of forming nitrides through
nitrocarburizing and improving surface hardness. However, if the Cr
content is less than 0.30%, less bainite phase forms, and V and Nb
precipitates are caused to form before nitrocarburizing treatment,
resulting in increased hardness before nitrocarburizing. Such a low
Cr content also decreases the absolute amount of V and Nb
precipitates remaining after nitrocarburizing treatment, and ends
up lowering the hardness after nitrocarburizing, making it
difficult to guarantee sufficient strength. Therefore, the Cr
content is set to 0.30% or more. On the other hand, Cr content
exceeding 3.0% lowers hot ductility, and causes hardening to
deteriorate machinability by cutting. Therefore, the Cr content is
set to 3.0% or less. The Cr content is preferably 0.5% or more and
2.0% or less, and more preferably 0.5% or more and 1.5% or
less.
[0060] Mo: 0.005% or More and 0.40% or Less
[0061] Mo increases hardenability and facilitates bainite phase
formation. Consequently, Mo has an effect of causing formation of
fine V and Nb precipitates and increasing the strength of the
nitrocarburized material. Therefore, Mo is one of the important
elements for the present disclosure. Mo is also effective in
bainite phase formation. To obtain the strength increasing effect,
the Mo content is set to 0.005% or more. On the other hand, Mo
content exceeding 0.40% lowers hot ductility and makes the cast
steel more prone to continuous casting cracks, and results in a
rise in component cost as Mo is an expensive element. Therefore,
the Mo content is set in a range of 0.005% to 0.40%. The Mo content
is preferably in a range of 0.015% to 0.3%, and more preferably in
a range of 0.04% to less than 0.2%.
[0062] V: 0.02% or More and 0.5% or Less
[0063] V is an important element that forms fine precipitates with
Nb due to the temperature rise during nitrocarburizing to thereby
increase core hardness and improve strength. To obtain this effect,
the V content is 0.02% or more. On the other hand, if the V content
exceeds 0.5%, precipitates become coarser, the strength increasing
effect is insufficient, and cracking is promoted during continuous
casting. Therefore, the V content is 0.5% or less. The V content is
preferably in a range of 0.03% to 0.3%, and more preferably in a
range of 0.03% to 0.25%.
[0064] Nb: 0.003% or More and 0.20% or Less
[0065] Nb forms fine precipitates with V due to the temperature
rise during nitrocarburizing and increases core hardness, and is
thus very effective in increasing fatigue strength. To obtain this
effect, the Nb content is set to 0.003% or more. On the other hand,
if the Nb content exceeds 0.20%, precipitates become coarser, the
strength increasing effect is insufficient, and cracking is
promoted during continuous casting. Therefore, the Nb content is
set to 0.20% or less. The Nb content is preferably in a range of
0.02% to 0.18%.
[0066] Al: 0.010% or More and 2.0% or Less
[0067] Al is a useful element for improving surface hardness and
effective hardened case depth after nitrocarburizing treatment, and
thus is intentionally added. Al is also a useful element for
inhibiting the growth of austenite grains during hot forging to
yield a finer microstructure and increased toughness. From this
perspective, the Al content is set to 0.010% or more. However,
adding Al beyond 2.0% does not increase this effect, but instead
promotes cracking during continuous casting and results in a rise
in component cost, which is disadvantageous. Therefore, the Al
content is set to 2.0% or less. Preferably, the Al content is more
than 0.020% and no more than 1.5%. More preferably, the Al content
is more than 0.020% and no more than 1.2%.
[0068] Ti: More than 0.005% and Less than 0.025%
[0069] Ti is a useful element for preventing the occurrence of
cooling cracks during continuous casting and surface cracks during
bending/bend restoration when using a bending continuous casting
machine, and is intentionally added in a range exceeding 0.005%. If
the Ti content is 0.025% or more, however, coarse TiN is generated
and fatigue strength decreases. Therefore, the Ti content is set to
less than 0.025%. The Ti content is preferably more than 0.012% and
no more than 0.023%, and more preferably in a range of 0.015% to
0.022%.
[0070] N: 0.0200% or Less
[0071] N is a useful element for forming carbonitrides in the steel
and improving the strength of the nitrocarburized material, and is
preferably added in an amount of 0.0020% or more. If the N content
exceeds 0.0200%, however, the resulting carbonitrides coarsen and
the toughness of the steel material decreases. In addition, the
cast steel suffers surface cracks, resulting in degradation of cast
slab quality. Therefore, the N content is set to 0.0200% or less.
The N content is preferably 0.0180% or less.
[0072] Sb: 0.0005% or More and 0.02% or Less
[0073] Sb has an effect of suppressing grain boundary oxidation and
surface cracking during casting, hot rolling, and hot forging, and
improving the surface quality of the product. This effect is
inadequate when the Sb content is below 0.0005%. On the other hand,
adding Sb beyond 0.02% does not increase this effect, but instead
results in a rise in component cost and causes Sb to segregate at
grain boundaries or otherwise, causing degradation in the toughness
of the base steel. Therefore, when added, the Sb content is set to
0.0005% or more and 0.02% or less. The Sb content is preferably
0.0010% or more and 0.01% or less.
[0074] Further, in the present disclosure, it is necessary to
satisfy the following formula in accordance with the C content:
in a case where the C content is 0.01% or more and 0.10% or
less,
(S/32)/(Ti/48)+(N/14)/(Ti/48).ltoreq.13.0, or
in a case where the C content is more than 0.10% and less than
0.20%,
2(S/32)/(Ti/48)+3(N/14)/(Ti/48).ltoreq.35.0.
[0075] We investigated the cause of cracking in the steel during
continuous casting, and found that precipitation of coarse MnS to
ferrite formed at grain boundaries during continuous casting is
responsible for causing cracking. We therefore studied how to
suppress the precipitation of MnS to ferrite at grain boundaries,
and revealed that the precipitation of MnS is closely related to
the contents of C, Ti, S, and N in the steel and that cracking
during continuous casting can be suppressed by adjusting the
contents of these elements to suppress the precipitation of MnS to
ferrite at grain boundaries. In other words, for C, Ti, S, and N,
by setting the parameters within the above ranges, it is possible
to cause S to precipitate as Ti carbosulfides to suppress
precipitation of coarse MnS to ferrite formed at grain boundaries
during continuous casting, and cast cracking can be reduced.
[0076] In addition to the basic components described above, the
chemical composition in the present disclosure may optionally
further contain: one or more selected from the group consisting of
B: 0.0100% or less, Cu: 0.3% or less, and Ni: 0.3% or less; one or
more selected from the group consisting of W: 0.3% or more, Co:
0.3% or less, Hf: 0.2% or less, and Zr: 0.2% or less; or one or
more selected from the group consisting of Pb: 0.2% or less, Bi:
0.2% or less, Zn: 0.2% or less, and Sn: 0.2% or less. The reasons
for the addition of each element will be described below.
[0077] B: 0.0100% or Less
[0078] B has an effect of improving hardenability and promoting the
formation of bainite microstructure. Thus, B is preferably added in
an amount of 0.0003% or more. If the B content is exceeds 0.0100%,
however, B precipitates as BN, the hardenability improving effect
is saturated, and the component cost rises. Therefore, when added,
the B content is set to 0.0100% or less. The B content is
preferably 0.0005% or more and 0.0080% or less.
[0079] Cu: 0.3% or Less
[0080] Cu is a useful element for forming an intermetallic compound
with Fe, Ni, or the like during nitrocarburizing treatment and
increasing the strength of the nitrocarburized material by
precipitation hardening, and is also effective for formation of
bainite phase. When the Cu content exceeds 0.3%, hot workability
decreases. Therefore, the Cu content is set to 0.3% or less. The Cu
content is preferably in a range of 0.05% to 0.25%.
[0081] Ni: 0.3% or Less
[0082] Ni has an effect of increasing hardenability and suppressing
low-temperature brittleness. A Ni content exceeding 0.3% not only
cause a rise in hardness and adversely affect machinability by
cutting, but also is disadvantageous in terms of cost. Therefore,
the Ni content is set to 0.3% or less. The Ni content is preferably
in a range of 0.05% to 0.25%.
[0083] W: 0.3% or Less, Co: 0.3% or Less, Hf: 0.2% or Less, Zr:
0.2% or Less
[0084] W, Co, Hf, and Zr are effective elements for improving the
strength of the steel, and are each preferably added in an amount
of 0.01% or more. However, adding W and Co beyond 0.3% and Hf and
Zr beyond 0.2% decreases the toughness. Therefore, the upper limit
is 0.3% for W and Co and 0.2% for Hf and Zr. Preferably, the
content is W: 0.01% to 0.25%, Co: 0.01% to 0.25%, Hf: 0.01% to
0.15%, and Zr: 0.01% to 0.15%.
[0085] Pb: 0.2% or Less, Bi: 0.2% or Less, Zn: 0.2% or Less, Sn:
0.2% or Less
[0086] Pb, Bi, Zn, and Sn are effective elements for improving the
machinability by cutting of the steel, and each can preferably be
added in an amount of 0.02% or more. However, addition beyond 0.2%
decreases strength and toughness. Therefore, the upper limit for
each added element is 0.2%.
[0087] It suffices for the chemical composition of the steel to
contain the above-described elements and the balance of Fe and
incidental impurities, yet the chemical composition preferably
consists of the above-described elements and the balance of Fe and
incidental impurities.
[0088] Next, the steel microstructure of the steel for
nitrocarburizing according to the disclosure will be described.
[0089] [Bainite Phase: More than 50% in Area Ratio]
[0090] In the present disclosure, it is vital that the steel
microstructure contains bainite phase in an area ratio of more than
50% with respect to a whole volume of the steel microstructure.
[0091] The present disclosure intends to improve the fatigue
strength after nitrocarburizing treatment by dispersing and
precipitating V and Nb during nitrocarburizing treatment to
increase the hardness of the nitride layer and the core part. In
other words, if V and Nb precipitates are present in large amounts
prior to nitrocarburizing treatment, this is disadvantageous from
the viewpoint of machinability by cutting at the time of cutting
work that is normally performed before nitrocarburizing. Further,
in the bainite transformation process, V and Nb precipitates are
less easily formed in the matrix phase as compared to the
ferrite-pearlite transformation process. Therefore, the steel
microstructure of the steel for nitrocarburizing according to the
disclosure, i.e., the steel microstructure before nitrocarburizing
treatment is mainly composed of bainite phase. Specifically, the
area ratio of bainite phase is set to more than 50%, preferably
more than 60%, and more preferably more than 80%, and may be 100%,
with respect to the whole volume of the steel microstructure.
[0092] Possible microstructures other than the bainite phase
include ferrite phase and pearlite phase, yet it is understood that
such microstructures are preferably as less as possible.
[0093] Here, the phase area ratio is determined by polishing, and
then etching with nital, the cross sections parallel to the rolling
direction (L-sections) of test pieces sampled from the obtained
steels for nitrocarburizing, and then observing the microstructures
of the cross sections under an optical microscope or a scanning
electron microscope (SEM) (microstructure observation under an
optical microscope at 200 times magnification) to identify the
phase type.
[0094] [Precipitates Containing V and Nb Dispersed in the Bainite
Phase]
[0095] In the nitrocarburized component according to the
disclosure, the steel for nitrocarburizing disclosed herein is
preferably subjected to nitrocarburizing treatment so that
precipitates containing V and Nb are dispersed in the bainite
phase. The reason is that by causing precipitates containing V and
Nb to be dispersed in the microstructure at the core part other
than the nitrocarburized portion at the surface layer part,
hardness increases and the fatigue strength after nitrocarburizing
treatment is significantly improved.
[0096] The term "core part" used herein refers to a region
excluding the surface compound layer and the hardened layer formed
as a result of nitrocarburizing. However, it is preferable to cause
precipitates containing V and Nb to disperse throughout the bainite
phase, rather than only in the core part.
[0097] Further, precipitates containing V and Nb in the bainite
phase preferably have a mean particle size of less than 10 nm, and
the number of such precipitates to be dispersed is preferably at
least 500 per unit area (1 .mu.m.sup.2) in order for the
precipitates to contribute to strengthening by precipitation after
nitrocarburizing treatment. The measurement limit for the diameter
of precipitates is around 1 nm.
[0098] It is noted here that a component obtained by
nitrocarburizing treatment has a nitrocarburized layer on the
surface layer. In such component, a surface layer part (a part
other than the core part) has a chemical composition that has
higher carbon and nitrogen contents than those in the core
part.
[0099] Next, methods of producing the steel for nitrocarburizing
and the nitrocarburized component according to the disclosure will
be described.
[0100] FIG. 1 illustrates a typical process for producing a
nitrocarburized component using a steel bar as the steel for
nitrocarburizing disclosed herein. In the figure, S1 is steel bar
production step, where a steel bar is used as the material, S2 is
steel bar transportation step, and S3 is product (nitrocarburized
component) finish step.
[0101] Firstly, in the steel bar production step (S1), a cast steel
is hot rolled into a semi-finished product and hot rolled into a
steel bar. The steel bar then goes through quality inspection
before it is shipped.
[0102] Then, after being transported (S2), in the product
(nitrocarburized component) finish step (S3), the steel bar is cut
into a predetermined dimension, subjected to hot forging or cold
forging, formed into a desired shape (such as the shape of a gear
or a shaft component) by cutting work such as drill boring or lathe
turning as necessary, and then subjected to nitrocarburizing
treatment to obtain a product.
[0103] Alternatively, the hot rolled material may be directly
subjected to cutting work such as lathe turning or drill boring to
form a desired shape before subjection to nitrocarburizing
treatment to obtain a product. In the case of hot forging, hot
forging may be followed by cold straightening. In addition, the
final product may be subjected to coating treatment such as
painting or plating.
[0104] According to the method of producing the steel for
nitrocarburizing disclosed herein, at the time of hot working right
before nitrocarburizing treatment, it is possible to obtain a
microstructure composed mainly of bainite phase as mentioned above
and to suppress the formation of V and Nb precipitates by setting a
specific heating temperature and a specific working temperature for
hot working.
[0105] The phrase "hot working right before nitrocarburizing
treatment" refers to either hot rolling or hot forging. However,
hot forging may be performed after hot rolling. Of course, hot
rolling may be followed by cold forging.
[0106] In the case of the hot working right before nitrocarburizing
being hot rolling, in other words, if hot forging is not performed
after hot rolling, the hot rolling needs to satisfy a set of
conditions given below.
[0107] [Rolling Heating Temperature: 950.degree. C. or Higher]
[0108] In the hot rolling, to prevent coarse carbonitrides from
forming on the material being rolled and lowering fatigue strength,
carbides remaining undissolved after dissolution are caused to
dissolve and form a solute. If the rolling heating temperature is
below 950.degree. C., it is difficult for the carbides remaining
undissolved after dissolution to dissolve and form a solute.
Therefore, the rolling heating temperature is set to 950.degree. C.
or higher, and preferably 960.degree. C. to 1250.degree. C.
[0109] [Rolling Finishing Temperature: 800.degree. C. or
Higher]
[0110] When the rolling finishing temperature is below 800.degree.
C., a ferrite phase forms, which is disadvantageous in obtaining a
microstructure that contains bainite phase in an area ratio of more
than 50% with respect to the whole volume of the microstructure
before nitrocarburizing treatment. The rolling load also increases.
Therefore, the rolling finishing temperature is set to 800.degree.
C. or higher. Regarding the upper limit, when the rolling finishing
temperature exceeds 1100.degree. C., crystal grains coarsen,
causing degradation in surface characteristics at the time of
cutting work after the hot rolling, cold forgeability, and the
like. Therefore, the rolling finishing temperature is preferably up
to 1100.degree. C.
[0111] [Cooling Rate after Rolling at Least in a Temperature Range
of 700.degree. C. to 550.degree. C.: Higher than 0.4.degree.
C./s]
[0112] When the cooling rate after rolling at least in a
temperature range of 700.degree. C. to 550.degree. C. is
0.4.degree. C./s or lower, fine precipitates are formed and
hardened before molding of components, resulting in increased
cutting resistance during cutting work, and the tool life
decreases. Therefore, at least in a temperature range of
700.degree. C. to 550.degree. C., which is the temperature range in
which fine precipitates form, the cooling rate after rolling is set
above the critical cooling rate of 0.4.degree. C./s at which fine
precipitates are obtained. Regarding the upper limit, if it exceeds
200.degree. C./s, a hard martensite phase forms and machinability
is greatly reduced. Therefore, the cooling rate after rolling in
this temperature range is preferably up to 200.degree. C./s.
[0113] In addition, in the case of the hot working right before
nitrocarburizing treatment being hot forging, in other words, if
hot forging is performed either alone or after hot rolling, the hot
forging needs to satisfy a set of conditions given below. When hot
rolling is performed before the hot forging, the hot rolling does
not necessarily have to satisfy the above-described conditions as
long as the below-described conditions are satisfied by the hot
forming.
[0114] [Forging Heating Temperature: 950.degree. C. or Higher]
[0115] In the hot forging, in order to form bainite phase in an
area ratio of more than 50% with respect to the whole volume of the
microstructure, and to suppress the formation of fine precipitates
from the perspective of cold straightening and machinability by
cutting after the hot forging, the heating temperature during the
hot forging is set to 950.degree. C. or higher. The heating
temperature is preferably from 960.degree. C. to 1250.degree.
C.
[0116] [Forging Finishing Temperature: 800.degree. C. or
Higher]
[0117] When the forging finishing temperature is below 800.degree.
C., a ferrite phase forms, which is disadvantageous in obtaining a
microstructure that contains bainite phase in an area ratio of more
than 50% with respect to the whole volume of the microstructure
before nitrocarburizing treatment. The forging load also increases.
Therefore, the forging finishing temperature is set to 800.degree.
C. or higher. Regarding the upper limit, when the forging finishing
temperature exceeds 1100.degree. C., crystal grains coarsen,
causing degradation in surface characteristics at the time of
cutting work after the hot forging. Therefore, the forging
finishing temperature is preferably up to 1100.degree. C.
[0118] [Cooling Rate after Forging at Least in a Temperature Range
of 700.degree. C. to 550.degree. C.: Higher than 0.4.degree.
C./s]
[0119] When the cooling rate at least in a temperature range of
700.degree. C. to 550.degree. C. after forging is 0.4.degree. C./s
or lower, fine precipitates are formed and hardened before molding
of components, resulting in increased cutting resistance during
cutting work, and the tool life decreases. Therefore, at least in a
temperature range of 700.degree. C. to 550.degree. C., which is the
temperature range in which fine precipitates form, the cooling rate
after forging is set above the critical cooling rate of 0.4.degree.
C./s at which fine precipitates are obtained. With respect to the
upper limit, if it exceeds 200.degree. C./s, a hard martensite
phase forms and machinability is greatly reduced. Therefore, the
cooling rate after forging in this temperature range is preferably
up to 200.degree. C./s.
[0120] Then, the materials thus rolled or forged may be subjected
to cutting work and the like to have the shape of a component, and
subsequently to nitrocarburizing treatment under a set of
conditions below.
[0121] [Nitrocarburizing Treatment Conditions]
[0122] To form fine precipitates, nitrocarburizing treatment is
preferably performed at a nitrocarburizing temperature in a range
of 550.degree. C. to 700.degree. C. for a duration of 10 minutes or
more. The reason why the nitrocarburizing temperature is set from
550.degree. C. to 700.degree. C. is that if the nitrocarburizing
temperature is below 550.degree. C., a sufficient amount of
precipitates cannot be obtained, while if the nitrocarburizing
temperature is above 700.degree. C., it reaches the austenite
region and makes and nitrocarburizing difficult to perform. The
nitrocarburizing temperature is more preferably in a range of
550.degree. C. to 630.degree. C.
[0123] Since N and C are introduced and diffused at the same time
in nitrocarburizing treatment, nitrocarburizing treatment may be
performed in a mixed atmosphere of nitriding gas such as NH.sub.3
or N.sub.2 and carburizing gas such as CO.sub.2 or CO, for example
in an atmosphere of NH.sub.3: N.sub.2: CO.sub.2=50:45:5.
EXAMPLES
[0124] Examples of the present disclosure will be specifically
described below.
[0125] Steels (ID 1 to ID 51) having the compositions presented in
Tables 1 and 2 were made into cast steels, each being 8000 mm long
and having a cross section of 300 mm.times.400 mm, using a
continuous casting machine. At that time, each steel was checked
for cracks on the surface. Specifically, surface observation was
performed in the longitudinal direction of each cast steel, and the
presence or absence of cracks having a length of 10 mm or more was
assessed. The number of cracks formed on the surface of the cast
steel was counted per 1 m.sup.2 of each cast steel, and based on
the assessment criteria, A: no crack, B: 1-4 cracks/m.sup.2, and C:
5 or more cracks/m.sup.2, cases A and B were scored as passed.
TABLE-US-00001 TABLE 1 Steel ID C Si Mn P S Cr Mo V Nb Al Ti N Sb
Others Formula 1 *) Category 1 0.043 0.05 1.75 0.014 0.015 1.15
0.070 0.08 0.063 0.035 0.013 0.0090 0.0005 -- 4.10 Example 2 0.131
0.11 1.45 0.012 0.019 1.25 0.006 0.14 0.074 0.031 0.015 0.0044
0.0006 -- 6.82 Example 3 0.179 0.25 2.42 0.012 0.017 1.32 0.105
0.12 0.051 0.028 0.009 0.0053 0.0007 -- 11.72 Example 4 0.065 0.35
1.65 0.015 0.015 0.35 0.123 0.09 0.248 0.123 0.016 0.0048 0.0007 --
2.43 Example 5 0.045 0.65 1.75 0.010 0.016 1.26 0.051 0.12 0.176
0.350 0.009 0.0055 0.0012 -- 4.76 Example 6 0.069 0.06 1.76 0.012
0.015 2.35 0.250 0.20 0.083 0.250 0.024 0.0053 0.0018 -- 1.69
Example 7 0.053 0.07 1.55 0.008 0.016 1.35 0.256 0.18 0.034 0.265
0.021 0.0045 0.0009 -- 1.88 Example 8 0.078 0.06 1.76 0.012 0.015
0.85 0.124 0.20 0.084 0.250 0.006 0.0123 0.0052 -- 10.78 Example 9
0.054 0.07 1.64 0.008 0.016 1.65 0.090 0.34 0.199 0.265 0.013
0.0045 0.0185 -- 3.03 Example 10 0.092 0.86 1.82 0.010 0.023 2.78
0.183 0.35 0.203 0.025 0.015 0.0056 0.0056 -- 3.58 Example 11 0.034
0.05 1.77 0.012 0.017 1.74 0.084 0.04 0.124 0.010 0.011 0.0046
0.0098 -- 3.75 Example 12 0.045 0.06 1.69 0.008 0.015 1.26 0.126
0.18 0.004 0.350 0.018 0.0036 0.0008 B: 0.0005 1.94 Example 13
0.063 0.12 2.26 0.015 0.016 1.34 0.254 0.13 0.148 0.550 0.022
0.0192 0.0010 Cu: 0.1 4.08 Example 14 0.180 0.31 1.69 0.018 0.033
1.13 0.180 0.15 0.039 1.420 0.006 0.0054 0.0006 Cu: 0.1, 25.76
Example Ni: 0.15 15 0.271 0.09 1.65 0.014 0.024 1.25 0.142 0.14
0.069 0.045 0.013 0.0065 0.0012 -- 10.68 Comparative Example 16
0.084 1.26 1.65 0.012 0.024 1.16 0.111 0.16 0.078 0.035 0.011
0.0073 0.0053 -- 5.55 Comparative Example 17 0.149 0.25 1.02 0.013
0.019 0.98 0.074 0.15 0.101 0.980 0.012 0.0047 0.0008 -- 8.78
Comparative Example 18 0.129 0.09 3.24 0.018 0.025 1.46 0.068 0.21
0.049 0.135 0.018 0.0052 0.0009 -- 7.14 Comparative Example 19
0.119 0.13 1.61 0.028 0.018 1.65 0.079 0.13 0.064 0.088 0.015
0.0063 0.0012 -- 7.92 Comparative Example 20 0.051 0.13 2.13 0.018
0.075 1.34 0.231 0.25 0.063 0.153 0.011 0.0124 0.0013 -- 14.09
Comparative Example 21 0.162 0.13 1.61 0.018 0.022 0.25 0.094 0.10
0.057 0.325 0.015 0.0032 0.0006 -- 6.59 Comparative Example 22
0.090 0.13 1.61 0.018 0.022 3.51 0.118 0.11 0.054 0.325 0.015
0.0032 0.0012 -- 2.93 Comparative Example 23 0.152 0.05 1.64 0.016
0.015 1.46 0.004 0.12 0.075 0.053 0.012 0.0056 0.0010 -- 8.55
Comparative Example 24 0.112 0.22 1.73 0.014 0.034 2.01 0.096 0.01
0.052 0.024 0.006 0.0152 0.0009 -- 43.06 Comparative Example 25
0.179 0.31 1.48 0.011 0.008 2.25 0.132 0.55 0.069 0.024 0.019
0.0065 0.0007 -- 4.78 Comparative Example 26 0.169 0.06 1.66 0.015
0.026 1.48 0.062 0.14 0.002 0.824 0.015 0.0088 0.0014 -- 11.23
Comparative Example 27 0.132 0.06 1.66 0.008 0.015 1.96 0.059 0.06
0.281 0.065 0.013 0.0051 0.0011 -- 7.50 Comparative Example 28
0.151 0.06 1.66 0.012 0.026 1.34 0.057 0.13 0.054 0.004 0.011
0.0088 0.0009 -- 15.32 Comparative Example 29 0.088 0.06 1.66 0.015
0.026 1.22 0.063 0.09 0.046 2.340 0.014 0.0103 0.0015 -- 5.31
Comparative Example 30 0.118 0.13 1.92 0.016 0.028 1.16 0.164 0.14
0.088 0.030 0.005 0.0059 0.0008 -- 28.94 Comparative Example 31
0.142 0.26 1.95 0.023 0.022 1.39 0.212 0.15 0.076 0.025 0.032
0.0123 0.0011 -- 6.02 Comparative Example 32 0.083 0.02 0.95 0.011
0.025 1.45 0.149 0.012 0.012 0.018 0.008 0.0252 0.0010 -- 15.49
Comparative Example 33 0.050 0.03 1.46 0.010 0.016 1.34 0.101 0.15
0.163 0.350 0.011 0.0123 0.0002 -- 6.02 Comparative Example 34
0.206 0.33 0.81 0.014 0.021 1.15 0.001 0.005 0.001 0.027 0.001
0.0130 -- -- -- Conventional Example *) Formula 1 [C: 0.01% or more
and 0.10% or less] (S/32)/(T1/48) + (N/14)/(Ti/48) = 0 to 13.0 [C:
more than 0.10% and less than 0.20%] 2(S/32)/(Ti/48) +
3(N/14)/(Ti/48) = 0 to 35.0
TABLE-US-00002 TABLE 2 (mass %) Steel ID C Si Mn P S Cr Mo V Nb Al
Ti N Sb Others Formula 1 *) Category 35 0.055 0.03 3.00 0.008 0.06
1.28 0.072 0.12 0.065 0.023 0.012 0.0096 0.0056 -- 10.2 Example 36
0.043 0.02 1.85 0.012 0.018 1.25 0.050 0.10 0.049 0.025 0.006
0.0102 0.0040 W: 0.3 10.3 Example 37 0.145 0.05 1.51 0.010 0.020
2.51 0.103 0.15 0.190 0.031 0.012 0.0088 0.0060 Co: 0.3 12.5
Example 38 0.102 0.20 1.98 0.015 0.017 1.85 0.123 0.18 0.156 0.028
0.013 0.0065 0.0008 Hf: 0.2, 9.1 Example Zr: 0.2 39 0.088 0.03 1.86
0.013 0.015 0.68 0.089 0.12 0.210 1.230 0.023 0.0152 0.0070 Pb: 0.1
3.2 Example 40 0.124 0.75 1.75 0.010 0.016 1.26 0.051 0.15 0.176
0.350 0.009 0.0055 0.0012 Bi: 0.2 11.6 Example 41 0.045 0.02 1.54
0.012 0.013 2.21 0.062 0.45 0.123 0.045 0.015 0.0122 0.0040 Zn: 0.2
4.1 Example 42 0.011 0.05 2.98 0.009 0.045 1.89 0.080 0.17 0.034
0.265 0.021 0.0045 0.0009 Sn: 0.2 3.9 Example 43 0.088 0.02 1.82
0.012 0.024 1.11 0.200 0.03 0.036 0.050 0.006 0.0123 0.0012 -- 13.0
Example 44 0.124 0.55 1.75 0.011 0.025 1.24 0.051 0.18 0.106 0.032
0.008 0.0199 0.0045 -- 35.0 Example 45 0.180 0.05 1.55 0.023 0.018
1.35 0.410 0.15 0.086 0.032 0.015 0.0065 0.0085 -- 8.1 Comparative
Example 46 0.045 0.05 1.56 0.008 0.005 1.08 0.050 0.16 0.049 0.026
0.005 0.0030 0.0052 -- 3.6 Comparative Example 47 0.155 0.03 1.78
0.010 0.006 1.62 0.075 0.12 0.150 0.032 0.005 0.0081 0.0034 -- 20.3
Comparative Example 48 0.092 0.04 1.63 0.010 0.015 1.55 0.096 0.05
0.078 0.025 0.026 0.0200 0.0065 -- 3.5 Comparative Example 49 0.045
0.03 1.98 0.018 0.031 1.63 0.060 0.03 0.045 0.050 0.008 0.0171
0.0005 -- 13.1 Comparative Example 50 0.165 0.23 2.31 0.018 0.017
1.35 0.063 0.12 0.120 0.025 0.006 0.0155 0.0126 -- 35.1 Comparative
Example 51 0.123 0.12 1.54 0.012 0.013 2.35 0.062 0.36 0.111 0.046
0.015 0.0186 0.0004 -- 15.4 Comparative Example *) Formula 1 [C:
0.01% or more and 0.10% or less] (S/32)/(Ti/48) + (N/14)/(Ti/48) =
0 to 13.0 [C: more than 0.10% and less than 0.20%] 2(S/32)/(Ti/48)
+ 3(N/14)/(Ti/48) = 0 to 35.0
[0126] Each cast steel was subjected to soaking at 1200.degree. C.
for 30 minutes and hot rolled into a semi-finished product having a
rectangular cross section with sides of 150 mm. Then, each cast
steel was hot rolled under the conditions including heating
temperature and rolling finishing temperature, as presented in
Tables 3 and 4, to obtain a steel bar of 60 mm .phi.. Then, each
cast steel was cooled to room temperature with the cooling rate in
the temperature range of 700.degree. C. to 550.degree. C. being
adjusted as presented in Tables 3 and 4, and used as the material
as hot rolled. It is noted here that Steel ID 34 is steel
equivalent to JIS SCr 420.
[0127] Each material as hot rolled was further subjected to hot
forging under the conditions presented in Tables 3 and 4 to obtain
a steel bar of 30 mm .phi., which in turn was cooled to room
temperature with the cooling rate in the temperature range of
700.degree. C. to 550.degree. C. being adjusted as presented in
Tables 3 and 4.
[0128] For the hot forged materials thus obtained, some of which
were as hot rolled, the machinability was evaluated by an outer
periphery turning test. As test pieces, either the hot forged
materials or the materials as hot rolled in a situation in which
hot forging was not performed were cut to a length of 200 mm. As
the cutting tool, CSBNR 2020 was used as the folder and SNGN 120408
UTi20 high-speed tool steel was used for the tip (CSBNR 2020 and
SNGN 120408 UTi20 are both manufactured by Mitsubishi Materials
Corporation). The conditions of the outer circumferential turning
test were as follows: cut depth 1.0 mm, feed rate 0.25 mm/rev,
cutting speed 200 m/min, and no lubricant. For an evaluation item,
the tool life was defined as the time until the tool wear (flank
wear) reached 0.2 mm.
[0129] In addition, microstructure observation and hardness
measurement were performed on the hot forged materials or the
materials as hot rolled in a situation in which hot forging was not
performed. In the microstructure observation, the type of phases
was identified and the area ratio of each identified phase was
determined with the above-described method.
[0130] In the hardness measurement, hardness HV was determined by
averaging the results of measuring hardness at five locations, each
being one-fourth the diameter from the surface of the test piece
(which is hereinafter considered as the core part) with a test load
of 2.94 N (300 gf) using a Vickers hardness meter in accordance
with JIS Z 2244.
[0131] Regarding Steel Nos. 1 to 33, after subjection to the
above-described hot forging, the test pieces were further subjected
to nitrocarburizing treatment. Steel ID 1 includes cases where hot
forging was not performed, in which case nitrocarburizing treatment
was performed after hot rolling. On the other hand, regarding the
hot forged materials with Steel ID 34, carburizing treatment was
performed for comparison.
[0132] Nitrocarburizing treatment was performed by heating the
steel samples to a temperature range of 525.degree. C. to
620.degree. C. in an atmosphere of NH.sub.3: N.sub.2:
CO.sub.2=50:45:5 and retaining them for 3.5 hours.
[0133] On the other hand, carburizing treatment was performed by
carburizing the test pieces at 930.degree. C. for 3 hours, holding
them at 850.degree. C. for 40 minutes, oil quenching them, and
further tempering them at 170.degree. C. for 1 hour.
[0134] The materials thus obtained by being subjected to
nitrocarburizing treatment and carburizing heat treatment were
further subjected to microstructure observation, hardness
measurement, and fatigue property evaluation.
[0135] In the microstructure observation, as it was before
nitrocarburizing treatment, the type of phases was identified and
the area ratio of each identified phase was determined with the
above-described method.
[0136] In the hardness measurement, measurement was made of the
surface hardness of each of the above-described heat-treated
materials at a depth of 0.05 mm from the surface, and of the core
hardness at the core part. In the surface hardness measurement and
core hardness measurement, surface hardness HV and core hardness HV
were determined by respectively averaging the results of measuring
the hardness at the core part at six locations with a test load of
2.94 N (300 gf) using a Vickers hardness meter in accordance with
JIS Z 2244. Measurement was further made of the depth of the
hardened layer, which was defined as the depth from the surface at
which HV of 520 is obtained.
[0137] Further, from the core parts of the nitrocarburized
materials and the carburized materials, test pieces were prepared
by twin-jet electropolishing for transmission electron microscope
observation, and precipitates on the test pieces were observed
under a transmission electron microscope with acceleration voltage
of 200 V. Further, the compositions of the observed precipitates
were determined with an energy-dispersive X-ray spectrometer
(EDX).
[0138] For fatigue property evaluation, a roller pitching test was
conducted, and fatigue strength after 10.sup.7 cycles was
determined. Fatigue test pieces were sampled from the materials as
hot rolled or the hot forged materials as described above in
parallel with their longitudinal direction. Each test piece had a
parallel portion of 26 mm .phi..times.28 mm long and a grip portion
of 24 mm.phi.. Each test piece was then subjected to
nitrocarburizing treatment. For those test pieces that were rated B
or C regarding the presence or absence of cracks on the surface of
the cast steel, test pieces were sampled from locations other than
where cracks occurred. In each roller pitching test piece, 26 mm
rolling contact surface was left as nitrocarburized (without
polishing). In the roller pitching test, the slip rate was -40%,
automatic transmission oil (Mitsubishi ATF SP-III) was used as the
lubricating oil, and the oil temperature was 80.degree. C. As large
rollers, carburized quenched products of SCM 420H with crowning R
of 150 mm were used.
[0139] Tables 3 and 4 present the results of the above tests. Nos.
1-19 and 50-59 are our examples, Nos. 20-48 and 60-66 are
comparative examples, and No. 49 is a conventional example in which
a steel equivalent to JIS SCr420 was subjected to carburizing
treatment.
[0140] As is clear from Tables 3 and 4, Examples 1-19 and 50-59 are
all superior in fatigue strength as compared to Conventional
Example 49 subjected to carburizing treatment. Examples 1-19 and
50-59 also exhibit better machinability by cutting before
nitrocarburizing treatment than Conventional Example No. 49.
[0141] Furthermore, as a result of observing precipitates with a
transmission electron microscope and investigating compositions of
the precipitates with an energy dispersive X-ray spectroscope
(EDX), it was confirmed that in the nitrocarburized materials of
Examples 1-19 and 50-59, at least 500 per 1 .mu.m.sup.2 fine
precipitates containing V and Nb and having a particle size of less
than 10 nm were formed and dispersed in the bainite phase. From
this result, it is considered that the nitrocarburizing materials
according to the disclosure exhibited high fatigue strength due to
the fine precipitates.
[0142] By contrast, for Comparative Example Nos. 20-48 in which the
chemical composition or the obtained steel microstructure was
outside the range of this disclosure, many cracks occurred during
continuous casting or fatigue strength or machinability was
inferior.
[0143] Specifically, for No. 20, since the heating temperature
during hot rolling was low, precipitates were not dissolved
sufficiently and the fatigue properties were inferior. Besides, due
to a high proportion of F+P microstructure, the machinability by
cutting after hot rolling was also low.
[0144] For No. 21, since the finishing temperature of hot rolling
was too low, the bainite fraction of the microstructure was low and
the machinability by cutting was inferior. In addition, since the
proportion of F+P microstructure was high, fine precipitates were
not formed after nitrocarburizing, and the fatigue properties were
thus inferior.
[0145] For Nos. 22 and 23, since the cooling rate after hot forging
was low, an appropriate amount of bainite phase was not obtained,
and only a small amount of fine precipitates was formed through
nitrocarburizing treatment, resulting in insufficient strengthening
by precipitation and lower fatigue strength compared to our
examples. The machinability by cutting was also low.
[0146] For No. 24, since the heating temperature of hot forging was
low, precipitates were not dissolved sufficiently and the fatigue
properties were inferior. Besides, due to a high proportion of F+P
microstructure, the machinability by cutting after hot rolling was
also low.
[0147] For No. 25, since the finishing temperature of hot forging
is too low, the bainite fraction of the microstructure is low and
the machinability by cutting is inferior. In addition, since the
proportion of F+P microstructure was high, fine precipitates were
not formed after nitrocarburizing, and the fatigue properties were
inferior.
[0148] For Nos. 26 and 27, since the cooling rate after hot forging
was low, an appropriate amount of bainite phase was not obtained,
and only a small amount of fine precipitates was formed through
nitrocarburizing treatment, resulting in insufficient strengthening
by precipitation and lower fatigue strength compared to our
examples. The machinability by cutting was also low.
[0149] For No. 28, since the nitrocarburizing temperature was low,
the depth of the hardened layer was small and the fatigue strength
was inferior.
[0150] For No. 29, since the nitrocarburizing treatment temperature
was high, nitrocarburizing was not sufficient, nor was
precipitation of fine precipitates adequate. Thus, the fatigue
strength was low.
[0151] For No. 30, since the C content exceeded the appropriate
range, the hot forged material increased in hardness before
subjection to nitrocarburizing treatment, and decreased in
machinability by cutting.
[0152] For No. 31, since the Si content exceeded the appropriate
range, the hot forged material increased in hardness before
subjection to nitrocarburizing treatment, and decreased in
machinability by cutting.
[0153] Regarding example No. 32, since the Mn content was below the
appropriate range, ferrite and pearlite phases were dominant in the
steel microstructure of the hot forged material before subjection
to nitrocarburizing treatment. Thus, V and Nb precipitates were
formed in the microstructure, the hardness before nitrocarburizing
treatment increased, and the machinability by cutting
decreased.
[0154] For No. 33, since the Mn content exceeded the appropriate
range, many cracks occurred during continuous casting. In addition,
a martensite phase was formed before nitrocarburizing treatment,
and the machinability by cutting was low.
[0155] For No. 34, since the P content exceeded the appropriate
range, many cracks occurred during continuous casting. The fatigue
strength was also low.
[0156] For No. 35, since the S content exceeded the appropriate
range and the value on the left side of the above Formula (1) was
outside the range of the present disclosure, many cracks occurred
during continuous casting.
[0157] For No. 36, since the Cr content was below the appropriate
range, ferrite and pearlite phases were dominant in the steel
microstructure of the hot forged material before subjection to
nitrocarburizing treatment. Accordingly, coarse V and Nb
precipitates were formed in the microstructure, the hardness before
nitrocarburizing treatment increased, and the fatigue strength
decreased.
[0158] For No. 37, since the Cr content exceeded the appropriate
range, many cracks occurred during continuous casting. In addition,
since the hardness after hot forging was high, the machinability by
cutting was inferior.
[0159] For No. 38, since the Mo content was below the appropriate
range, the hardenability decreased and the formation of the bainite
phase is insufficient. This resulted in a small amount of fine
precipitates formed after nitrocarburizing treatment and
insufficient core hardness. Accordingly, the fatigue strength was
low as compared with Conventional Example No. 49.
[0160] For No. 39, since the V content was below the appropriate
range, only a small amount of fine precipitates was formed through
nitrocarburizing treatment, and sufficient core hardness was not
obtained. Accordingly, the fatigue strength was low as compared
with Conventional Example No. 49.
[0161] For No. 40, since the V content exceeded the appropriate
range, many cracks occurred during continuous casting.
[0162] Regarding example No. 41, since the Nb content was below the
appropriate range, only a small amount of fine precipitates was
formed through nitrocarburizing treatment, and sufficient core
hardness was not obtained. Accordingly, the fatigue strength was
low as compared with Conventional Example No. 49.
[0163] For No. 42, the Nb content exceeded the appropriate range,
and many cracks occurred during continuous casting.
[0164] For No. 43, since the Al content was below the appropriate
range, neither sufficient surface strength nor an effective
hardened case depth were obtained after nitrocarburizing treatment,
and the fatigue strength was lower than that of Conventional
Example No. 49.
[0165] For No. 44, since the Al content exceeded the appropriate
range, many cracks occurred during continuous casting.
[0166] For No. 45, the Ti content did not satisfy the appropriate
range, many cracks occurred during continuous casting.
[0167] For No. 46, since the Ti content exceeded the appropriate
range, the fatigue strength was low.
[0168] For No. 47, since the N content exceeded the appropriate
range, many cracks occurred during continuous casting.
[0169] For No. 48, since the Sb content exceeded the appropriate
range, many cracks occurred during continuous casting.
[0170] For No. 60, since the Mo content exceeded the appropriate
range, many cracks occurred during continuous casting.
[0171] For Nos. 61 and 62, since the Ti content was below the
appropriate range, many cracks occurred during continuous
casting.
[0172] For No. 63, since the Ti content exceeded the appropriate
range, the fatigue strength was low.
[0173] For No. 64, since the value on the left side of the above
Formula (1) exceeded 13.0, many cracks occurred during continuous
casting.
[0174] For No. 65, since the value on the left side of the above
Formula (1) exceeded 35.0, many cracks occurred during continuous
casting.
[0175] In No. 66, since the Sb content was below the appropriate
range, many cracks occurred during continuous casting.
TABLE-US-00003 TABLE 3 Cracks Hot rolling conditions Hot forging
conditions on the Cooling Cooling Steel properties (before
nitrocarburizing treatment) surface Heating rate Heating Hot rate
Area of the temp. Hot rolling after temp. forging after ratio of
semi- for hot finishing hot for hot finishing hot Core bainite
Steel finished rolling temp. rolling forging temp. forging hardness
Steel phase No. ID product*4 (.degree. C.) (.degree. C.) (.degree.
C./s) (.degree. C.) (.degree. C.) (.degree. C./s) HV microstructure
(%) 1 1 A 1150 970 0.8 1200 1100 0.8 253 B dominant 92 2 2 A 1150
970 0.7 1200 1100 0.8 301 B dominant 98 3 3 A 1150 970 0.8 1200
1100 0.8 317 B dominant 97 4 4 A 1150 970 0.7 1200 1100 0.8 282 B
dominant 95 5 5 A 1150 970 0.8 1200 1100 0.7 280 B dominant 96 6 6
B 1150 970 0.8 1200 1100 0.8 285 B dominant 98 7 7 B 1150 970 0.6
1200 1100 0.8 279 B dominant 94 8 8 A 1150 970 0.8 1200 1100 0.8
284 B dominant 96 9 9 A 1150 970 0.8 1200 1100 0.5 252 B dominant
55 10 10 A 1150 970 0.5 1200 1100 0.8 291 B dominant 93 11 11 A
1050 900 0.8 1200 1100 0.6 281 B dominant 60 12 12 A 1150 970 0.8
1200 1100 0.8 277 B dominant 97 13 13 B 1050 910 0.8 1200 1100 0.5
290 B dominant 96 14 14 A 1080 920 0.8 1200 1100 0.8 275 B dominant
80 15 1 A 1150 970 0.8 960 840 0.8 249 B dominant 88 16 1 A 1150
970 0.8 1250 1050 0.8 256 B dominant 96 17 1 A 1150 970 0.8 -- --
-- 252 B dominant 96 18 1 A 960 810 0.8 -- -- -- 239 B dominant 87
19 1 A 1250 1050 0.8 -- -- -- 233 B dominant 75 20 1 A 930 815 0.8
-- -- -- 189 F + P + B 42 21 1 A 1150 750 0.8 -- -- -- 222 F + P +
B 24 22 1 A 1150 970 0.3 -- -- -- 210 F + P + B 38 23 1 A 1150 970
0.4 -- -- -- 215 F + P + B 45 24 1 A 1150 970 0.8 900 1100 0.8 195
F + P + B 45 25 1 A 1150 970 0.8 1200 750 0.8 256 F + P + B 24 26 1
A 1150 970 0.8 1200 1100 0.3 220 F + P + B 41 27 1 A 1150 970 0.8
1200 1100 0.4 231 F + P + B 48 28 1 A 1150 970 0.8 1200 1100 0.8
253 B dominant 92 29 1 A 1150 970 0.8 1200 1100 0.8 253 B dominant
92 30 15 B 1150 970 0.8 1200 1100 0.8 389 B dominant 95 31 16 B
1150 970 0.8 1200 1100 0.8 356 B dominant 85 32 17 A 1150 970 0.8
1200 1100 0.8 321 F + P 32 33 18 C 1150 970 0.8 1200 1100 0.8 363 B
dominant 96 34 19 C 1150 970 0.8 1200 1100 0.8 298 B dominant 96 35
20 C 1150 970 0.8 1200 1100 0.8 281 B dominant 95 36 21 A 1150 970
0.8 1200 1100 0.8 285 F + P + B 23 37 22 B 1150 970 0.8 1200 1100
0.8 401 M + B 56 38 23 A 1150 970 0.8 1200 1100 0.8 202 F + P + B
18 39 24 A 1150 970 0.8 1200 1100 0.8 191 B dominant 96 40 25 C
1150 970 0.8 1200 1100 0.8 320 B dominant 97 41 26 A 1150 970 0.8
1200 1100 0.8 223 B dominant 95 42 27 C 1150 970 0.8 1200 1100 0.8
280 B dominant 96 43 28 A 1150 970 0.8 1200 1100 0.8 307 B dominant
94 44 29 C 1150 970 0.8 1200 1100 0.8 292 B dominant 87 45 30 C
1150 970 0.8 1200 1100 0.8 299 B dominant 91 46 31 A 1150 970 0.8
1200 1100 0.8 306 B dominant 94 47 32 C 1150 970 0.8 1200 1100 0.8
285 B dominant 93 48 33 C 1150 970 0.8 1200 1100 0.8 278 B dominant
90 49 34 B 1150 970 0.8 1200 1100 0.8 220 F + P 75 Steel properties
(before nitro- Steel properties after nitrocarburizing treatment)
carburizing Nitro- Area treatment) carburizing Depth of ratio of
Tool treatment Surface hardened Core bainite Fatigue life temp.
hardness layer hardness Steel phase strength No. (s) (.degree. C.)
HV (mm) HV microstructure (%) (MPa) Remarks 1 630 575 816 0.17 263
B dominant 93 2550 Example 2 510 580 813 0.16 312 B dominant 98
2600 Example 3 471 600 826 0.20 326 B dominant 97 2650 Example 4
559 590 831 0.21 302 B dominant 92 2600 Example 5 563 595 823 0.20
298 B dominant 96 2650 Example 6 550 580 815 0.18 301 B dominant 97
2600 Example 7 564 575 819 0.19 304 B dominant 95 2650 Example 8
552 570 817 0.16 305 B dominant 94 2600 Example 9 632 570 820 0.15
256 B dominant 53 2550 Example 10 534 570 823 0.18 310 B dominant
95 2900 Example 11 561 570 843 0.20 299 B dominant 57 2550 Example
12 570 570 826 0.19 288 B dominant 99 2550 Example 13 538 570 821
0.17 306 B dominant 98 2600 Example 14 575 570 812 0.18 295 B
dominant 78 2850 Example 15 640 570 823 0.23 260 B dominant 84 2750
Example 16 623 570 819 0.22 263 B dominant 93 2650 Example 17 632
560 825 0.23 269 B dominant 95 2500 Example 18 665 570 823 0.21 251
B dominant 84 2550 Example 19 680 570 805 0.18 249 B dominant 74
2500 Example 20 405 570 786 0.19 223 F + P + B 40 1950 Comparative
Example 21 309 570 789 0.18 205 F + P + B 22 2150 Comparative
Example 22 186 570 796 0.18 208 F + P + B 35 2050 Comparative
Example 23 213 570 803 0.19 216 F + P + B 42 2100 Comparative
Example 24 299 570 650 0.14 278 F + P + B 40 2050 Comparative
Example 25 126 570 812 0.26 401 F + P + B 22 2000 Comparative
Example 26 133 570 788 0.22 202 F + P + B 36 2100 Comparative
Example 27 136 570 801 0.19 215 F + P + B 42 1950 Comparative
Example 28 630 400 602 0.05 222 B dominant 91 2050 Comparative
Example 29 630 710 599 0.08 213 B dominant 89 2050 Comparative
Example 30 126 570 812 0.26 401 B dominant 94 2400 Comparative
Example 31 374 570 825 0.23 345 B dominant 84 2650 Comparative
Example 32 133 570 788 0.18 371 F + P 32 2650 Comparative Example
33 224 570 803 0.25 370 B dominant 94 2400 Comparative Example 34
517 570 805 0.19 316 B dominant 94 1900 Comparative Example 35 561
570 814 0.20 301 B dominant 96 2550 Comparative Example 36 169 570
794 0.10 262 F + P + B 19 2150 Comparative Example 37 122 570 825
0.29 398 tempered M + B 39 2650 Comparative Example 38 162 570 792
0.09 211 F + P + B 18 2000 Comparative Example 39 510 570 816 0.08
178 B dominant 94 2100 Comparative Example 40 312 570 810 0.18 326
B dominant 95 2050 Comparative Example 41 505 570 808 0.17 222 B
dominant 97 2100 Comparative Example 42 493 570 810 0.18 295 B
dominant 96 2050 Comparative Example 43 495 570 702 0.07 321 B
dominant 95 1950 Comparative Example 44 533 570 845 0.29 324 B
dominant 94 2650 Comparative Example 45 515 570 815 0.18 319 B
dominant 93 2200 Comparative Example 46 498 570 804 0.14 275 B
dominant 95 2050 Comparative Example 47 551 270 810 0.15 270 B
dominant 94 2150 Comparative Example 48 568 570 816 0.17 272 B
dominant 94 2300 Comparative Example 49 265 -*3 730 1.02 344
tempered M + B 48 2400 Comparative Example *1 Underlined values are
outside of the range of the present disclosure. *2 Microstructural
symbols are the follow* phases: F: ferrite, P: pearlite, B:
bainite, M: martensite *3Carburizing treatment was performed.
*4Criteria for assessing cracks on the surface of the cast steel:
A: no crack, B: 1-4 cracks/m.sup.2, and C: 5 or more
cracks/m.sup.2.
TABLE-US-00004 TABLE 4 Steel properties Cracks (before
nitrocarburizing treatment) on the Hot rolling conditions Hot
forging conditions surface Heating Hot Cooling Heating Cooling Area
of the temp. rolling rate temp. Hot forging rate ratio of semi- for
hot finishing after hot for hot finighing after hot Core bainite
Steel finished rolling temp. rolling forging temp. forging hardness
Steel phase No. ID product*4 (.degree. C.) (.degree. C.) (.degree.
C./s) (.degree. C.) (.degree. C.) (.degree. C./s) HV microstructure
(%) 50 35 B 1150 970 0.6 1200 1100 0.7 262 B dominant 92 51 36 A
1150 970 0.8 1200 1100 0.8 296 B dominant 95 52 37 A 1150 970 0.7
1200 1100 0.8 363 B dominant 100 53 38 A 1150 970 0.5 1200 1100 0.5
335 B dominant 92 54 39 B 1150 970 0.8 1200 1100 0.8 280 B dominant
96 55 40 B 1150 970 0.8 1200 1100 0.8 301 B dominant 98 56 41 B
1150 970 0.8 1200 1100 0.6 256 B dominant 97 57 42 B 1150 970 0.8
1200 1100 0.8 237 B dominant 95 58 43 B 1150 970 0.8 1200 1100 0.8
280 B dominant 95 59 44 B 1150 970 0.7 1200 1100 0.7 301 B dominant
94 60 45 C 1150 970 0.8 1200 1100 0.8 333 B dominant 100 61 46 C
1150 970 0.8 1200 1100 0.8 256 B dominant 94 62 47 C 1150 970 0.8
1200 1100 0.8 319 B dominant 93 63 48 B 1150 970 0.8 1200 1100 0.8
283 B dominant 95 64 49 C 1150 970 0.8 1200 1100 0.8 256 B dominant
95 65 50 C 1150 970 0.8 1200 1100 0.8 324 B dominant 96 66 51 C
1150 970 0.8 1200 1100 0.8 300 B dominant 96 Steel properties
(before nitro- Steel properties (after nitrocaburizing treatment)
carburizing Nitro- Area treatment) carburizing Depth of ratio of
Tool treatment Surface hardened Core bainite Fatigue life temp.
hardness layer hardness Steel phase strength No. (s) (.degree. C.)
HV (mm) HV microstructure (%) (MPa) Remarks 50 1003 570 820 0.16
302 B dominant 91 2550 Example 51 625 590 834 0.15 356 B dominant
99 2950 Example 52 483 595 815 0.21 370 B dominant 96 3000 Example
53 543 585 812 0.20 299 B dominant 90 2850 Example 54 1205 570 800
0.19 350 B dominant 95 2500 Example 55 1360 565 811 0.17 265 B
dominant 98 2450 Example 56 1230 570 825 0.18 280 B dominant 96
2500 Example 57 750 570 821 0.20 257 B dominant 97 2550 Example 58
657 570 815 0.18 301 B dominant 95 2500 Example 59 614 600 810 0.19
230 B dominant 93 2550 Example 60 547 570 823 0.18 342 B dominant
99 2450 Comparative Example 61 709 570 805 0.17 272 B dominant 92
2400 Comparative Example 62 577 570 811 0.19 302 B dominant 91 2450
Comparative Example 63 652 570 822 0.16 274 B dominant 96 1950
Comparative Example 64 709 570 812 0.18 280 B dominant 95 2400
Comparative Example 65 565 570 816 0.16 341 B dominant 93 2450
Comparative Example 66 615 570 817 0.16 332 B dominant 94 2500
Comparative Example *1 Underlined values are outside of the range
of the present disclosure. *2 Microstructural symbols are the
following phases: F: ferrite, P: pearlite, B: bainite, M:
martensite *3 Carburizing treatment was performed. *4Criteria for
assessing cracks on the surface of the cast steel A: no crack B:
1-4 cracks/m.sup.2, and C: 5 or more cracks/m.sup.2.
* * * * *