U.S. patent application number 15/576177 was filed with the patent office on 2018-08-16 for steel plate and method of production of same.
This patent application is currently assigned to NIPPON STEEL & SUMITOMO METAL CORPORATION. The applicant listed for this patent is NIPPON STEEL & SUMITOMO METAL CORPORATION. Invention is credited to Takashi ARAMAKI, Motonori HASHIMOTO, Kazuo HIKIDA, Ken TAKATA, Kengo TAKEDA, Toshimasa TOMOKIYO, Yasushi TSUKANO.
Application Number | 20180230582 15/576177 |
Document ID | / |
Family ID | 57393441 |
Filed Date | 2018-08-16 |
United States Patent
Application |
20180230582 |
Kind Code |
A1 |
TAKEDA; Kengo ; et
al. |
August 16, 2018 |
STEEL PLATE AND METHOD OF PRODUCTION OF SAME
Abstract
Low carbon steel plate excellent impact resistance
characteristics after carburizing and quenching and after
tempering, characterized by having a predetermined chemical
composition, an average grain size of carbides of 0.4 .mu.m to 2.0
.mu.m, an area ratio of pearlite of 6% or less, a ratio of a number
of carbides at the ferrite grain boundaries to the number of
carbides inside the ferrite grains of over 1, and a Vickers
hardness of 100HV to 180HV.
Inventors: |
TAKEDA; Kengo; (Tokyo,
JP) ; HIKIDA; Kazuo; (Tokyo, JP) ; TAKATA;
Ken; (Tokyo, JP) ; HASHIMOTO; Motonori;
(Tokyo, JP) ; TOMOKIYO; Toshimasa; (Tokyo, JP)
; TSUKANO; Yasushi; (Tokyo, JP) ; ARAMAKI;
Takashi; (Tokyo, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
NIPPON STEEL & SUMITOMO METAL CORPORATION |
Tokyo |
|
JP |
|
|
Assignee: |
NIPPON STEEL & SUMITOMO METAL
CORPORATION
Tokyo
JP
|
Family ID: |
57393441 |
Appl. No.: |
15/576177 |
Filed: |
May 25, 2016 |
PCT Filed: |
May 25, 2016 |
PCT NO: |
PCT/JP2016/065509 |
371 Date: |
November 21, 2017 |
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C22C 38/46 20130101;
C22C 38/002 20130101; C22C 38/005 20130101; C22C 38/001 20130101;
C21D 2211/005 20130101; C22C 38/42 20130101; C21D 8/0205 20130101;
C22C 38/008 20130101; C22C 38/02 20130101; C22C 38/44 20130101;
C22C 38/06 20130101; C21D 9/46 20130101; C22C 38/04 20130101; C22C
38/60 20130101; C22C 38/54 20130101; C22C 38/50 20130101; C22C
38/48 20130101; C21D 8/0463 20130101 |
International
Class: |
C22C 38/60 20060101
C22C038/60; C21D 8/04 20060101 C21D008/04; C22C 38/42 20060101
C22C038/42; C22C 38/44 20060101 C22C038/44; C22C 38/46 20060101
C22C038/46; C22C 38/48 20060101 C22C038/48; C22C 38/50 20060101
C22C038/50; C22C 38/54 20060101 C22C038/54; C22C 38/06 20060101
C22C038/06; C22C 38/04 20060101 C22C038/04; C22C 38/02 20060101
C22C038/02; C22C 38/00 20060101 C22C038/00 |
Foreign Application Data
Date |
Code |
Application Number |
May 26, 2015 |
JP |
2015-106745 |
Claims
1. A steel plate being a low carbon steel plate having a chemical
composition consisting of, by mass %, C: 0.10 to 0.40%, Si: 0.01 to
0.30%, Mn: 0.30 to 1.00%, Al: 0.001 to 0.10%, Cr: 0.50 to 2.00%,
Mo: 0.001 to 1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020%
or less, O: 0.020% or less, Ti: 0.010% or less, B: 0.0005% or less,
Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or less, Nb:
0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or
less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0.050% or less, Ca:
0.050% or less, Y: 0.050% or less, Zr: 0.050% or less, La: 0.050%
or less, Ce: 0.050% or less, and a balance of Fe and impurities,
wherein a metal structure of the steel plate has: a carbide grain
size of 0.4 to 2.0 .mu.m; a pearlite area ratio of 6% or less; and
a ratio of the number of carbides at ferrite grain boundaries to
the number of carbides inside ferrite grains of more than 1, and
the steel plate has a Vickers hardness of 100 HV to 180 HV.
2. A method of production of the steel plate according to claim 1,
the method of production comprising: hot-rolling a steel slab with
a chemical composition according to claim 1 to obtain a hot rolled
steel plate, the hot-rolling in which finish hot-rolling is
completed in a 650.degree. C. to 950.degree. C. temperature region;
coiling the hot rolled steel plate at 400.degree. C. to 600.degree.
C.; pickling the coiled hot rolled steel plate, and subjecting a
first stage annealing to the pickled hot rolled steel plate, the
first stage annealing in which the pickled hot rolled steel plate
is heated, at a heating rate of 30.degree. C./hour to 150.degree.
C./hour, up to an annealing temperature of 650.degree. C. to
720.degree. C. and the steel plate is held for 3 hours to 60 hours;
then subjecting a second stage annealing to the hot rolled steel
plate, the second stage annealing in which the hot rolled steel
plate is heated, at a heating rate of 1.degree. C./hour to
80.degree. C./hour, up to an annealing temperature of 725.degree.
C. to 790.degree. C. and the steel plate is held for 3 hours to 50
hours as second stage annealing; and cooling the annealed hot
rolled steel plate to 650.degree. C. at a cooling rate of 1.degree.
C./hour to 100.degree. C./hour.
Description
TECHNICAL FIELD
[0001] The present invention relates to steel plate and a method of
production of the same.
BACKGROUND ART
[0002] Steel plate containing, by mass %, carbon in an amount of
0.1 to 0.4% is being used as a material for gears, clutches, and
other drive system parts of automobiles by being used press-formed,
enlarging holes, bent, drawn, thickened, and thinned and cold
forged by combinations of the same from a blank. Compared with
conventional hot forging etc., with cold forging, there is the
problem that the amount of strain accumulated in the material
becomes higher, cracks of the material and buckling at the time of
shaping are invited, and deterioration of the part characteristics
is caused.
[0003] In particular, to obtain wear resistance, after the shaped
material is carburized, quenched, and tempered, residual stress is
caused by the heat treatment, so formation and growth of fracture
from the cracked parts and buckled parts are invited. To use such a
part for the drive system, an impact resistance characteristic is
sought for preventing fracture due to brittleness in the face of
the large instantaneous load applied to the start of engagement of
the gears at the time of startup etc., so excellent cold
forgeability and impact resistance characteristic after
carburizing, quenching, and tempering are being demanded from the
above steel plate.
[0004] Up to now, various proposals have been made regarding arts
for improving the cold forgeability of steel plate and the impact
resistance characteristic after carburization (for example, see
PLTs 1 to 5).
[0005] For example, PLT 1 discloses, as steel for machine
structural use improving toughness by suppressing coarsening of
crystal grains in carburization heat treatment, steel for machine
structural use containing, by mass %, C: 0.10 to 0.30%, Si: 0.05 to
2.0%, Mn: 0.10 to 0.50%, P: 0.030% or less, S: 0.030% or less, Cr:
1.80 to 3.00%, Al: 0.005 to 0.050%, Nb: 0.02 to 0.10%, and N:
0.0300% or less and having a balance of Fe and unavoidable
impurities, having a structure before cold working comprised of
ferrite and pearlite structures, and having an average grain size
of ferrite grains of 15 .mu.m or more.
[0006] PLT 2 discloses, as steel excellent in cold workability and
carburizing and quenching ability, steel containing C: 0.15 to
0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol. Al: 0.02% or
less, N: 0.006% or less, and B: 0.005 to 0.050%, having a balance
of Fe and unavoidable impurities, and having a structure mainly
comprised of ferrite phases and graphite phases.
[0007] PLT 3 discloses a steel material for carburized bevel gear
use excellent in impact strength, a high toughness carburized bevel
gear, and a method of production of the same.
[0008] PLT 4 discloses steel for carburized part use having
excellent workability while suppressing coarsening of crystal
grains even with subsequent carburization and having an excellent
impact resistance characteristic and impact fatigue resistance
characteristic in a part produced by spheroidal annealing, then a
cold forging and a carburizing, quenching, and tempering
process.
[0009] PLT 5 discloses as cold tool steel for plasma carburization
use a steel containing C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn:
0.05 to 1.50%, and V: 1.8 to 6.0%, further containing one or more
of Ni: 0.10 to 2.50%, Cr: 0.1 to 2.0%, and Mo: 3.0% or less, and
having a balance of Fe and unavoidable impurities.
CITATION LIST
Patent Literature
[0010] PLT 1: Japanese Patent Publication No. 2013-040376A
[0011] PLT 2: Japanese Patent Publication No. 06-116679A
[0012] PLT 3: Japanese Patent Publication No. 09-201644A
[0013] PLT 4: Japanese Patent Publication No. 2006-213951A
[0014] PLT 5: Japanese Patent Publication No. 10-158780A
SUMMARY OF INVENTION
Technical Problem
[0015] The structure of the steel for machine structural use of PLT
1 is a structure of ferrite+pearlite. This structure, compared with
a ferrite+cementite structure, has a large hardness, so wear of the
die in cold forging cannot be suppressed and the steel cannot
necessarily be said to be steel for machine structural use
excellent in cold forgeability.
[0016] In the steel of PLT 2, the graphitization treatment of the
cementite requires annealing at a high temperature. A drop in the
yield and an increase in the manufacturing costs cannot be
suppressed.
[0017] The method of production of PLT 3 requires further hot
forging after cold forging and carburizing. Since hot forging is
essential, this is not a method of production leading to
fundamentally lower costs.
[0018] It is unclear if the steel for carburized part use of PLT 4
can exhibits similar effects in cold forging given a large strain.
Furthermore, the specific form of the structure and method of
control of the structure are also unclear, so this cannot be said
to be steel exhibiting excellent workability even in the plate
forging growing in use in recent years and other shaping by forging
cold while giving a large strain.
[0019] PLT 5 does not disclose at all the findings and art relating
to the optimum components and form of structure for improving the
formability of steel, in particular cold forgeability.
[0020] The present invention, in consideration of the above prior
art, has as its problem the provision of steel plate excellent in
cold forgeability and impact resistance characteristic after
carburizing, quenching, and tempering, in particular suitable for
obtaining a high cycle gear or other part by forming a plate and of
a method of production of the same.
Solution to Problem
[0021] To solve the above problem and obtain steel plate suitable
for a material such as a drive system part, it is understood that
in a steel plate containing the C required for raising the
hardenability, enlargement of the ferrite in grain size,
spheroidization of the carbides (mainly cementite) to a suitable
grain size, and reduction of the pearlite structures are
preferable. This is due to the following reasons.
[0022] A ferrite phase is low in hardness and high in ductility.
Therefore, in a structure mainly comprised of ferrite, it becomes
possible to increase the grain size so as to raise the formability
of the material.
[0023] Carbides, by being made to suitably disperse in the metal
structure, can maintain the formability of the material while
imparting an excellent wear resistance and rolling fatigue
characteristic, so provides a structure essential for drive system
parts. Further, the carbides in the steel plate are strong
particles obstructing slip.
[0024] By forming carbides at the ferrite grain boundaries, it is
possible to prevent propagation of slip exceeding the crystal grain
boundaries and suppress the formation of shear zones. Thus the cold
forgeability is improved and, simultaneously, the formability of
steel plate is also improved.
[0025] However, cementite is a hard, brittle structure. If a
laminar structure with ferrite present, that is, in the state of
pearlite, the steel becomes hard and brittle, so it has to be
present in a spheroidal form. If considering the cold forgeability
and the occurrence of fractures at the time of forging, its grain
size has to be a suitable range.
[0026] However, no method of production for realizing the above
structure has been disclosed up to now. Therefore, the inventors
intensively researched a method of production for realizing the
above structure.
[0027] As a result, they discovered the following: To make the
metal structure of the steel plate after coiling after hot rolling
a bainite structure of fine pearlite or fine ferrite with small
lamellar spacing in which cementite is dispersed, the steel plate
is coiled at a relatively low temperature (400.degree. C. to
550.degree. C.). By coiling at a relatively low temperature, the
cementite dispersed in the ferrite also easily becomes spheroidal.
Next, the cementite is partially made spheroidal by annealing at a
temperature just under the Ac1 point as first stage annealing.
Next, as second stage annealing, part of the ferrite grains is left
while part is transformed to austenite by annealing at a
temperature between the Ac1 point and Ac3 point (so-called dual
phase region of ferrite and austenite). By then making the
remaining ferrite grains grow while slowly cooling the steel while
using these as nuclei to transform the austenite to ferrite, it is
possible to obtain large ferrite phases and make cementite
precipitate at the grain boundaries to realize the above
structure.
[0028] That is, the method of production of steel plate
simultaneously satisfying hardenability and formability is
difficult to realize even if designing the hot rolling conditions,
annealing conditions, etc. as single processes. It was discovered
that this can be realized by optimization by a so-called integral
process of hot rolling, an annealing process, etc.
[0029] Further, improvement of the drawability at the time of cold
forging requires the reduction of plastic anisotropy. It was
discovered that for such improvement, adjustment of the hot rolling
conditions is important.
[0030] The present invention was made based on these discoveries
and has as its gist the following:
[0031] (1) A steel plate being low carbon steel plate having a
chemical composition containing, by mass %, C: 0.10 to 0.40%, Si:
0.01 to 0.30%, Mn: 0.30 to 1.00%, Al: 0.001 to 0.10%, Cr: 0.50 to
2.00%, Mo: 0.001 to 1.00%, P: 0.020% or less, S: 0.010% or less, N:
0.020% or less, O: 0.020% or less, Ti: 0.010% or less, B: 0.0005%
or less, Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or
less, Nb: 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W:
0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0.050% or
less, Ca: 0.050% or less, Y: 0.050% or less, Zr: 0.050% or less,
La: 0.050% or less, and Ce: 0.050% and having a balance of Fe and
impurities, the metal structure of the steel plate having a carbide
grain size of 0.4 to 2.0 .mu.m, a pearlite area ratio of 6% or
less, and a ratio of a number of carbides at the ferrite grain
boundaries to the number of carbides inside the ferrite grains of
over 1, the steel plate having a Vickers hardness of 100 HV to 180
HV.
[0032] (2) A method of production of the steel plate according to
(1), the method of production comprising the steps of: hot rolling
a steel slab of a chemical composition according to claim 1,
completing finish hot rolling in a 650.degree. C. to 950.degree. C.
temperature region to obtain a hot rolled steel plate; coiling the
hot rolled steel plate at 400.degree. C. to 600.degree. C.;
pickling the coiled hot rolled steel plate and heating the pickled
hot rolled steel plate by a 30.degree. C./hour to 150.degree.
C./hour heating rate to a 650.degree. C. to 720.degree. C.
annealing temperature and holding it there for 3 hours to 60 hours
as first stage annealing; then heating the hot rolled steel plate
to an annealing temperature of 725.degree. C. to 790.degree. C. by
a heating rate of 1.degree. C./hour to 80.degree. C./hour and
holding the steel plate for 3 hours to 50 hours as second stage
annealing; and cooling the annealed hot rolled steel plate to
650.degree. C. by a cooling rate of 1.degree. C./hour to
100.degree. C./hour.
[0033] According to the present invention, it is possible to
provide a steel plate excellent in cold forgeability and impact
resistance characteristic after carburizing, quenching, and
tempering, in particular one suitable for obtaining a high cycle
gear or other part by forming a plate.
BRIEF DESCRIPTION OF DRAWINGS
[0034] FIGS. 1A to 1C are views schematically showing a summary of
the cold forging test and the form of a crack introduced by cold
forging. FIG. 1A shows a disk-shaped test material cut out from a
hot rolled steel plate, while FIG. 1B shows the shape of a test
material after cold forging, while FIG. 1C shows the
cross-sectional shape of the test material after cold forging.
[0035] FIG. 2 is a view schematically showing a summary of a drop
weight test evaluating the impact resistance characteristic of a
sample performing carburizing, quenching, and tempering.
[0036] FIG. 3 is a view showing a relationship among a ratio of a
number of carbides at the grain boundaries to the number of
carbides in the grains, the crack length of the cold forging test
piece, and the impact resistance characteristic after carburizing,
quenching, and tempering.
[0037] FIG. 4 is a view showing another relationship among a ratio
of a number of carbides at the grain boundaries to the number of
carbides in the grains, the crack length of the cold forging test
piece, and the impact resistance characteristic after carburizing,
quenching, and tempering.
DESCRIPTION OF EMBODIMENTS
[0038] Below, the present invention will be explained in detail.
First, the reasons for limitation of the chemical composition of
the steel plate of the present invention will be explained. Here,
the "%" according to the chemical composition means "mass %".
[0039] C: 0.10 to 0.40%
[0040] C is an element forming carbides in steel and effective for
strengthening the steel and refining the ferrite grains. To
suppress the formation of a matte surface in cold working and
secure surface beauty of a cold forged part, suppression of
coarsening of the ferrite grain size is essential, but if less than
0.10%, the carbides become insufficient in volume fraction and
coarsening of the carbides during annealing can no longer be
suppressed, so C is made 0.10% or more. Preferably it is 0.11% or
more.
[0041] On the other hand, if exceeding 0.40%, the carbides increase
in volume fraction, a large amount of cracks are formed acting as
starting points of breakage at the time of an instantaneous load
and a drop in the impact resistance characteristic is invited, so C
is made 0.40% or less. Preferably it is 0.38% or less.
[0042] Si: 0.01 to 0.30%
[0043] Si is an element which acts as a deoxidizing agent and
further has an effect on the form of the carbides. To reduce the
number of carbides in the ferrite grains giving the deoxidizing
effect and increase the number of carbides at the ferrite grain
boundaries, it is necessary to use two-stage step type annealing to
produce austenite phases during annealing, make the carbides
dissolve once, then gradually cool the structure to promote the
formation of carbides at the ferrite grain boundaries.
[0044] If Si exceeds 0.30%, the ferrite falls in ductility,
fractures are easily formed at the time of cold forging, and the
cold forgeability and impact resistance characteristic after
carburizing, quenching, and tempering deteriorate, so Si is made
0.30% or less. Preferably it is 0.28% or less.
[0045] Si is preferably as low as possible, but reduction to less
than 0.01% invites a large increase in refining costs, so Si is
made 0.01% or more. Preferably it is 0.02% or more.
[0046] Mn: 0.30 to 1.00%
[0047] Mn is an element controlling the form of carbides in
two-stage step type annealing. If less than 0.30%, in the gradual
cooling after second stage annealing, it becomes difficult to form
carbides at the ferrite grain boundaries, so Mn is made 0.30% or
more. Preferably it is 0.33% or more.
[0048] On the other hand, if exceeding 1.00%, the toughness after
carburizing, quenching, and tempering falls, so Mn is made 1.00% or
less. Preferably it is 0.96% or less.
[0049] Al: 0.001 to 0.10%
[0050] Al is an element acting as a deoxidizing agent of steel and
stabilizing ferrite. If less than 0.001%, the effect of addition is
not sufficiently obtained, so Al is made 0.001% or more. Preferably
it is 0.004% or more.
[0051] On the other hand, if exceeding 0.10%, the number ratio of
carbides at the grain boundaries is lowered and an increase in
crack length at the time of cold forging is invited, so Al is made
0.10% or less. Preferably it is 0.09% or less.
[0052] Cr: 0.50 to 2.00%
[0053] Cr and Mo are elements which improve the toughness. Cr is an
element effective for stabilization of carbides at the time of heat
treatment. If less than 0.50%, it becomes difficult to cause
carbides to remain at the time of carburization, coarsening of the
austenite grain size at the surface layer is invited, and a drop in
the impact resistance characteristic is caused, so Cr is made 0.50%
or more. Preferably it is 0.52% or more.
[0054] On the other hand, if exceeding 2.00%, the amount of Cr
concentrating at the carbides increases and a large amount of fine
carbides remain in the austenite phases produced by the two-stage
step type annealing, carbides remain in the grains after gradual
cooling, the hardness increases and number ratio of carbides at the
grain boundaries fall and the cold forgeability falls, so Cr is
made 2.00% or less. Preferably it is 1.94% or less.
[0055] Mo: 0.001 to 1.00%
[0056] Mo is an element effective for control of the form of
carbides. If less than 0.001%, the effect of addition is not
sufficiently obtained, so Mo is made 0.001% or more. Preferably it
is 0.017% or more.
[0057] On the other hand, if exceeding 1.00%, Mo concentrates in
the carbides and stable carbides increase in the austenite phase as
well, so after gradual cooling, carbides are present in the grains
as well, an increase in hardness and drop in number ratio of
carbides at the grain boundaries are invited, and the cold
forgeability falls, so Mo is made 1.00% or less. Preferably it is
0.94% or less.
[0058] The following elements are impurities and have to be
controlled to certain amounts or less.
[0059] P: 0.020% or Less
[0060] P is an element segregating at the ferrite grain boundaries
and suppressing the formation of carbides at the grain boundaries.
The smaller amount is preferable. The content of P may also be 0,
but a long time is required for refining in order to make the
purity a high one of less than 0.0001% in a refining process and a
large increase in the manufacturing cost is invited, so the de
facto lower limit is 0.0001 to 0.0013%.
[0061] On the other hand, if exceeding 0.020%, the number ratio of
carbides at the grain boundaries falls and the cold forgeability
falls, so P is made 0.020% or less. Preferably it is 0.018% or
less.
[0062] S: 0.010% or Less
[0063] S is an impurity element forming MnS and other nonmetallic
inclusions. The nonmetallic inclusions form starting points of
formation of fractures at the time of cold forging, so the smaller
the S, the better. The content of S may also be 0, but to lower S
to less than 0.0001%, the refining costs greatly increase, so the
de facto lower limit is 0.0001 to 0.0012%.
[0064] On the other hand, if exceeding 0.010%, an increase is
invited in the crack length at the time of cold forging, so S is
made 0.010% or less. Preferably it is 0.009% or less.
[0065] N: 0.020% or Less
[0066] N is an element segregating at the ferrite grain boundaries
and suppressing the formation of carbides at the grain boundaries.
The smaller amount is preferable. The content of N may also be 0,
but if reducing it to less than 0.0001%, the refining costs greatly
increase, so the de facto lower limit is 0.0001 to 0.0006%.
[0067] On the other hand, if exceeding 0.020%, even if performing
dual phase region annealing and gradual cooling, the ratio of the
number of carbides at the ferrite grain boundaries with respect to
the number of carbides in the ferrite grains becomes less than 1
and the cold forgeability fall, so N is made 0.020% or less.
Preferably it is 0.017% or less.
[0068] O: 0.0001 to 0.020%
[0069] O is an element forming oxides in the steel. The oxides
present in the ferrite grains become sites for production of
carbides, so the smaller the amount, the better. The content of 0
may also be 0, but if reducing 0 to less than 0.0001%, the refining
costs greatly increase, so the de facto lower limit is 0.0001 to
0.0006%.
[0070] On the other hand, if exceeding 0.020%, the ratio of the
number of carbides at the ferrite grain boundaries with respect to
the number of carbides in the ferrite grains becomes less than 1
and the cold forgeability falls, so 0 is made 0.020% or less.
Preferably it is 0.017% or less.
[0071] Ti: 0.010% or Less
[0072] Ti is an element important for control of the form of the
carbides. It is an element by which, by inclusion in a large
amount, formation of carbides in the ferrite grains is promoted.
The smaller amount is preferable. The content of Ti may also be 0,
but if reducing it to less than 0.0001%, the refining costs greatly
increase, so the de facto lower limit is 0.0001 to 0.0006%.
[0073] On the other hand, if over 0.010%, the ratio of the number
of carbides at the ferrite grain boundaries to the number of
carbides inside the ferrite grains becomes less than 1 and the cold
forgeability falls, so Ti is made 0.010% or less. Preferably it is
0.007% or less.
[0074] B: 0.0005% or Less
[0075] B is an element effective for control of slip of
dislocations at the time of cold forging. By inclusion of a large
amount, activity of the slip system is limited, so the smaller the
amount of B, the better. The content of B may also be 0. Fine care
is required for detection of less than 0.0001% of B. Depending on
the analysis device, it is below the lower limit of detection.
[0076] On the other hand, if exceeding 0.0005%, cross slip of
dislocations at the shear zone formed by the cold forging is
suppressed. Strain concentrates locally and fractures are formed,
so B is made 0.0005% or less. Preferably it is 0.0005% or less.
[0077] Sn: 0.050% or Less
[0078] Sn is an element entering from the steel starting materials
(scraps). The smaller amount is preferable. The content of Sn may
also be 0, but if reducing it to less than 0.001%, the refining
costs greatly increase, so the de facto lower limit is 0.001 to
0.002%.
[0079] On the other hand, if exceeding 0.050%, the ferrite becomes
brittle and the cold forgeability falls, so Sn is made 0.050% or
less. Preferably, it is 0.048% or less.
[0080] Sb: 0.050% or Less
[0081] Sb, like Sn, is an element entering from the steel starting
materials (scraps). Sb segregates at the grain boundaries and
lowers the number ratio of carbides at the grain boundaries, so the
smaller the amount, the better. The content of Sb may also be 0,
but if reducing it to less than 0.001%, the refining costs greatly
increase, so the de facto lower limit is 0.001 to 0.002%.
[0082] On the other hand, if exceeding 0.050%, the cold
forgeability falls, so Sb is made 0.050% or less. Preferably, it is
0.048% or less.
[0083] As: 0.050% or Less
[0084] As is an element which enters from the steel starting
materials (scraps) like Sn and Sb. As segregates at the grain
boundaries and lowers the number ratio of carbides at the grain
boundaries, so the content is preferably small. The content of As
may also be 0, but if reducing it to less than 0.001%, the refining
cost greatly increases, so the de facto lower limit is 0.001 to
0.002%.
[0085] On the other hand, if over 0.050%, the number ratio of
carbides at the grain boundaries falls and the cold forgeability
falls, so As is made 0.050% or less. Preferably it is 0.045% or
less.
[0086] The steel plate of the present invention has the above
elements as basic elements, but may further contain the following
elements for the purpose of improving the cold forgeability and
other characteristics. The following elements are not essential for
obtaining the effects of the present invention, so the contents may
also be 0.
[0087] Nb: 0.10% or Less
[0088] Nb is an element effective for control of the form of the
carbides. Further, it is an element refining the structure and
contributing to improvement of the toughness. If less than 0.001%,
the effect of addition is not sufficiently obtained, so Nb is
preferably made 0.001% or more. More preferably, it is 0.002% or
more.
[0089] On the other hand, if over 0.10%, a large number of fine Nb
carbides precipitate, the strength excessively rises, and, further,
the number ratio of carbides at the grain boundaries falls and the
cold forgeability falls, so Nb is made 0.10% or less. Preferably it
is 0.09% or less.
[0090] V: 0.10% or Less
[0091] V, like Nb, is an element effective for control of the form
of the carbides. Further, it is an element refining the structure
and contributing to improvement of the toughness. If less than
0.001%, the effect of addition is not sufficiently obtained, so V
is preferably made 0.001% or more. More preferably, it is 0.004% or
more.
[0092] On the other hand, if over 0.10%, a large number of fine V
carbides precipitate, the strength excessively rises, and, further,
the number ratio of carbides at the grain boundaries falls and the
cold forgeability falls, so V is made 0.10% or less. Preferably, it
is 0.09% or less.
[0093] Cu: 0.10% or Less
[0094] Cu is an element forming fine precipitates and contributing
to improvement of the strength. If less than 0.001%, the effect of
improvement of the strength is not sufficiently obtained, so Cu is
preferably made 0.001% or more. More preferably, it is 0.008% or
more.
[0095] On the other hand, if over 0.10%, red hot embrittlement
occurs during hot rolling and the productivity falls, so Cu is made
0.10% or less. Preferably, it is 0.09% or less.
[0096] W: 0.10% or Less
[0097] W, like Nb and V, is an element effective for control of the
form of the carbides. If less than 0.001%, the effect of addition
is not sufficiently obtained, so W is preferably made 0.001% or
more. More preferably, it is 0.003% or more.
[0098] On the other hand, if over 0.10%, a large number of fine W
carbides precipitate, the strength excessively rises, and, further,
the number ratio of carbides at the grain boundaries falls and the
cold forgeability falls, so W is made 0.10% or less. Preferably, it
is 0.08% or less.
[0099] Ta: 0.10% or Less
[0100] Ta, like Nb, V, and W, is an element effective for control
of the form of the carbides. If less than 0.001%, the effect of
addition is not sufficiently obtained, so Ta is preferably made
0.001% or more. More preferably, it is 0.007% or more.
[0101] On the other hand, if over 0.10%, a large number of fine Ta
carbides precipitate, the strength excessively rises, and, further,
the number ratio of carbides at the grain boundaries falls and the
cold forgeability falls, so Ta is made 0.10% or less. Preferably,
it is 0.09% or less.
[0102] Ni: 0.10% or Less
[0103] Ni is an element effective for improvement of the impact
resistance characteristic of parts. If less than 0.001%, the effect
of addition is not sufficiently obtained, so Ni preferably is made
0.001% or more. More preferably it is 0.002% or more.
[0104] On the other hand, if over 0.10%, the number ratio of
carbides at the grain boundaries falls and the cold forgeability
falls, so Ni is made 0.10% or less. Preferably, it is 0.09% or
less.
[0105] Mg: 0.050% or Less
[0106] Mg is an element which can control the form of sulfides by
addition in a trace amount. If less than 0.0001%, the effect of
addition is not sufficiently obtained, so Mg preferably is made
0.0001% or more. More preferably it is 0.0008% or more.
[0107] On the other hand, if over 0.050%, the ferrite becomes
brittle and the cold forgeability falls, so Mg is made 0.050% or
less. Preferably it is 0.049% or less.
[0108] Ca: 0.050% or Less
[0109] Ca, like Mg, is an element which can control the form of
sulfides by addition in a trace amount. If less than 0.001%, the
effect of addition is not sufficiently obtained, so Ca preferably
is made 0.001% or more. More preferably it is 0.003% or more.
[0110] On the other hand, if over 0.050%, coarse Ca oxides are
formed and become starting points of fracture at the time of cold
forging, so Ca is made 0.050% or less. Preferably it is 0.04% or
less.
[0111] Y: 0.050% or Less
[0112] Y, like Mg and Ca, is an element which can control the form
of sulfides by addition in a trace amount. If less than 0.001%, the
effect of addition is not sufficiently obtained, so Y preferably is
made 0.001% or more. More preferably it is 0.003% or more.
[0113] On the other hand, if over 0.050%, coarse Y oxides are
formed and become starting points of fracture at the time of cold
forging, so Y is made 0.050% or less. Preferably it is 0.031% or
less.
[0114] Zr: 0.050% or Less
[0115] Zr, like Mg, Ca, and Y, is an element which can control the
form of sulfides by addition in a trace amount. If less than
0.001%, the effect of addition is not sufficiently obtained, so Zr
preferably is made 0.001% or more. More preferably it is 0.004% or
more.
[0116] On the other hand, if over 0.050%, coarse Zr oxides are
formed and become starting points of fracture at the time of cold
forging, so Zr is made 0.050% or less. Preferably it is 0.045% or
less.
[0117] La: 0.050% or Less
[0118] La is an element effective for control of the form of
sulfides by addition in a trace amount. Further, it is an element
which segregates at the grain boundaries and lowers the number
ratio of carbides at the grain boundaries. If less than 0.001%, the
effect of control of the form is not sufficiently obtained, so La
is preferably made 0.001% or more. More preferably, it is 0.003% or
more.
[0119] On the other hand, if over 0.050%, the number ratio of
carbides at the grain boundaries falls and the cold forgeability
falls, so La is made 0.050% or less. Preferably it is 0.047% or
less.
[0120] Ce: 0.050% or Less
[0121] Ce, like La, is an element able to control the form of
sulfides by addition in a trace amount. Further, it is an element
which segregates at the grain boundaries and lowers the number
ratio of carbides at the grain boundaries. If less than 0.001%, the
effect of control of the form is not sufficiently obtained, so Ce
is preferably made 0.001% or more. More preferably, it is 0.003% or
more.
[0122] On the other hand, if exceeding 0.050%, the number ratio of
carbides at the grain boundaries falls and the cold forgeability
falls, so Ce is made 0.050% or less. Preferably it is 0.046% or
less.
[0123] Note that, the remainder of the chemical composition of the
steel plate of the present invention is comprised of Fe and
unavoidable impurities.
[0124] Next, the structure of the steel plate of the present
invention will be explained.
[0125] The structure of the steel plate of the present invention is
substantially a structure comprised of ferrites and carbides. The
carbides include cementite (Fe.sub.3C) which is a compound of iron
and carbon, a compound obtained by substituting Mn, Cr, etc. for
the Fe atoms in the cementite, and alloy carbides (M.sub.23C.sub.6,
M.sub.6C, MC, etc., where M is Fe and other metal elements).
[0126] When forming steel plate into a predetermined part shape, a
shear zone is formed at the macrostructure of the steel plate and
slip deformation occurs concentrated near the shear zone. In slip
deformation, along with proliferation of dislocations, a region of
a high dislocation density is formed near the shear zone. Along
with the increase in the amount of strain imparted to the steel
plate, slip deformation is promoted and the dislocation density
increases.
[0127] In cold forging, strong working is performed with an
equivalent strain exceeding 1. For this reason, in conventional
steel plate, it was not possible to prevent the formation of voids
and/or cracks along with the increase in dislocation density and
was difficult to improve the cold forgeability.
[0128] To solve this difficult problem, it is effective to suppress
the formation of a shear zone at the time of forming. From the
viewpoint of the microstructure, formation of a shear zone can be
understood as the phenomenon of slip occurring at a certain one
grain crossing the crystal grain boundary and being continuously
propagated to the adjoining grain. Accordingly, to suppress the
formation of a shear zone, it is necessary to prevent propagation
of slip crossing crystal grain boundaries.
[0129] The carbides in steel plate are strong particles inhibiting
slip. By forming carbides at the ferrite grain boundaries, it
becomes possible to suppress the formation of a shear zone and
improve the cold forgeability.
[0130] To obtain such an effect, carbides have to be made to
disperse in the metal structure in suitable sizes. Therefore, the
average particle size of carbides is made 0.4 .mu.m to 2.0 .mu.m.
If the particle size of the carbides is less than 0.4 .mu.m, the
steel plate remarkably increases in hardness and the cold
forgeability falls. More preferably it is 0.6 .mu.m or more.
[0131] On the other hand, if the average particle size of the
carbides exceeds 2.0 .mu.m, at the time of cold forming, the
carbides form starting points of fractures. More preferably, it is
1.95 .mu.m or less.
[0132] Further, cementite, a carbide of iron, has a hard and
brittle structure. If present in the form of pearlite, which is a
layered structure with ferrite, the steel becomes hard and brittle.
Therefore, pearlite has to be reduced as much as possible. In the
steel plate of the present invention, the area ratio is made 6% or
less.
[0133] Pearlite has a unique lamellar structure, so can be
discerned by observation by an SEM or optical microscope. By
calculating the region of the lamellar structure at any
cross-section, the area ratio of the pearlite can be found.
[0134] Based on theory and principle, cold forgeability is
considered to be strongly affected by the rate of coverage of the
ferrite grain boundaries by carbides. High precision measurement is
sought, but measurement of the rate of coverage of ferrite grain
boundaries by carbides in a three-dimensional space requires serial
sectioning SEM observation using an FIB to repeatedly cut a sample
for observation in a scanning electron microscope or 3D EBSP
observation. A massive measurement time is required and technical
knowhow has to be built up.
[0135] The inventors clarified this and searched for a simpler,
higher precision indicator for evaluation and as a result
discovered that it is possible to evaluate the cold forgeability by
using the ratio of the number of carbides at the ferrite grain
boundaries to the number of carbides in the ferrite grains as an
indicator and that if the ratio of the number of carbides at the
ferrite grain boundaries to the number of carbides in the ferrite
grains exceeds 1, the cold forgeability remarkably rises.
[0136] Note that, buckling, folding, and twisting of the steel
plate occurring at the time of cold forging occur due to
localization of strain accompanying the formation of a shear zone,
so similarly by forming carbides at the ferrite grain boundaries to
reduce the formation of a shear zone and localization of strain, it
is possible to suppress buckling, folding, and twisting.
[0137] The carbides are observed by a scanning electron microscope.
Before observation, the sample for observation of the structure is
polished by wet polishing by Emery paper and a diamond abrasive
having an average particle size of 1 .mu.m, the observed surface is
polished to a mirror finish, then a saturated picric acid-alcohol
solution is used to etch the structure.
[0138] The magnification of the observation was made 3000.times.
and images of eight fields of 30 .mu.m.times.40 .mu.m at a plate
thickness 1/4 layer were captured at random. The obtained
structural images were analyzed by image analyzing software such as
one made by Mitani Shoji (Win ROOF) to measure in detail the areas
of the carbides contained in those regions. The circle equivalent
diameters (=2.times.(area/3.14)) were found from the areas of the
carbides and the average value was made the particle size of the
carbides.
[0139] Note that, to keep down the effect of measurement error due
to noise, carbides with an area of 0.01 .mu.m.sup.2 or less are
excluded from the coverage of the evaluation.
[0140] The number of carbides which present at the ferrite grain
boundaries are counted, the number of carbides at the grain
boundaries are subtracted from the total number of carbides, and
the number of carbides in the ferrite grains are found. Based on
the measured number, the number ratio of carbides at the grain
boundaries with respect to the carbides inside the ferrite grains
is calculated.
[0141] By making as the structure after annealing a structure with
ferrite grains of a size of 3.0 .mu.m to 50.0 .mu.m, it is possible
to improve the cold forgeability. If the size of the ferrite grains
is less than 3 .mu.m, the hardness increases and fractures and
cracks easily form at the time of cold forging, so the ferrite
grain size is preferably 3.0 .mu.m or more. More preferably it is
7.5 .mu.m or more.
[0142] On the other hand, if the ferrite grain size is over 50.0
.mu.m, the number of carbides on the crystal grain boundaries
suppressing the propagation of slip is decreased and the cold
forgeability falls, so the ferrite grain size is preferably 50.0
.mu.m or less. More preferably it is 37.9 .mu.m or less.
[0143] The ferrite grain size is measured by using the
above-mentioned procedure to polish the observed surface of the
sample for observation of structure to a mirror finish, then
observing the structure of the observed surface etched by a 3%
nitric acid-alcohol solution by an optical microscope or scanning
electron microscope and applying the line segment method to the
captured image.
[0144] By making the Vickers hardness of the steel plate 100 HV to
180 HV, it is possible to improve the cold forgeability and impact
resistance characteristic after carburizing, quenching, and
tempering. If the Vickers hardness is less than 100 HV, buckling
easily occurs during cold forging, folding and twisting occur at
the buckled part, and the impact resistance characteristic falls,
so the Vickers hardness is made 100 HV or more. Preferably it is
110 HV or more.
[0145] On the other hand, if the Vickers hardness exceeds 180 HV,
the ductility falls, internal cracking easily occurs at the time of
cold forging, and the impact resistance characteristic
deteriorates, so the Vickers hardness is made 180 HV or less.
Preferably it is 170 HV or less.
[0146] Next, the method of evaluation of the cold forgeability will
be explained.
[0147] FIGS. 1A to 1C schematically show an outline of the cold
forging test and form of a crack introduced by cold forging. FIG.
1A shows a disk-shaped test material cut out from a hot rolled
steel plate, FIG. 1B shows the shape of a test material after cold
forging, and FIG. 1C shows the cross-sectional shape of a cold
forged test material.
[0148] As shown in FIGS. 1A to 1C, from a plate thickness 5.2 mm
hot rolled steel plate, a diameter 70 mm disk-shaped test material
1 was cut out (see FIG. 1A) and deep drawn so as to prepare a
cup-shaped test material with a bottom surface of a diameter of 30
mm (not shown). Next, a one shot forming press made by Mori
Ironworks was used to thicken the vertical wall parts of a
cup-shaped test material by a thickening ratio of 1.54 (=8 mm/5.2
mm) (cold forging) to prepare a cup-shaped test material 2 with a
diameter of 30 mm, a height of 30 mm, and a vertical wall thickness
of 8 mm (see FIG. 1B).
[0149] The thickened cup-shaped test material 2 is cut by a wire
cut electrical discharge machine made by FANUC so that the
cross-section of the diameter part appeared (see FIG. 1C. The cut
surface is polished to a mirror finish and the presence of a
fracture 3 at the cut surface was confirmed. The ratio of the
maximum length of fracture L present at the vertical wall parts to
the thickness of the vertical wall part after thickening (=maximum
length of fracture L/thickness of vertical wall part after
thickening 8 mm) is measured. This measured value is used to
evaluate the cold forgeability.
[0150] Note that, even if the initial plate thickness is other than
5.2 mm, if adjusting the diameter of the cut out disk-shaped test
material so that the height of the vertical wall after thickening
becomes 30 mm and forming the material by a thickening ratio of the
same 1.54, it is possible to reproduce the results of evaluation
without regard as to the initial plate thickness, so the hot rolled
steel plate covered by the present invention is not limited to a
plate thickness 5.2 mm hot rolled steel plate. The present
invention can improve the cold forgeability and the impact
resistance characteristic after carburizing, quenching, and
tempering even in a general plate thickness (2 to 15 mm) hot rolled
steel plate.
[0151] Next, the method of production of the present invention will
be explained. The technical idea of the method of production of the
present invention is to integrally manage the rolling conditions
and annealing conditions when producing steel plate from a steel
slab of the above-mentioned chemical composition so as to improve
the cold forgeability and the impact resistance characteristic
after carburizing, quenching, and tempering.
[0152] The features of the method of production of the present
invention will be explained next.
[0153] Features of Hot Rolling
[0154] Molten steel having the required chemical composition is
continuously cast into a slab. The slab is used for hot rolling as
is in accordance with an ordinary method or is cooled once, then
heated and used for hot rolling. The finish hot rolling is ended in
the 650.degree. C. to 950.degree. C. temperature region. The hot
rolled steel plate after finish rolling is cooled on the ROT and
coiled by a coiling temperature of 400.degree. C. to 600.degree.
C.
[0155] Features of Annealing
[0156] The hot rolled steel plate is pickled, then held at two
temperature regions as two-stage step type annealing, but at that
time, in the first stage annealing, the hot rolled steel plate is
heated until the annealing temperature by a 30.degree. C./hour to
150.degree. C./hour heating rate and held at a 650.degree. C. to
720.degree. C. temperature region for 3 hours to 60 hours for
annealing.
[0157] At the next second stage annealing, the hot rolled steel
plate is heated until the annealing temperature by a 1.degree.
C./hour to 80.degree. C./hour heating rate and held at a
725.degree. C. to 790.degree. C. temperature region for 3 hours to
50 hours for annealing.
[0158] Next, the annealed hot rolled steel plate is cooled down to
650.degree. C. by a cooling rate of 1.degree. C./hour to
100.degree. C./hour, then is cooled down to room temperature.
[0159] Due to the linkage between these hot rolling conditions and
annealing conditions, low carbon steel plate excellent in cold
forgeability and impact resistance characteristic after
carburizing, quenching, and tempering can be obtained.
[0160] Below, the conditions of steps of the method of production
of the present invention will be specifically explained.
[0161] Hot Rolling
[0162] Finish hot rolling temperature: 650.degree. C. to
950.degree. C.
[0163] Coiling temperature: 400.degree. C. to 600.degree. C.
[0164] Molten steel having the required chemical composition is
continuously cast into a slab. The slab is used for hot rolling as
is or cooled once, then heated. The finish hot rolling is ended in
the 650.degree. C. to 950.degree. C. temperature region. The hot
rolled steel plate is coiled at 400.degree. C. to 600.degree.
C.
[0165] The slab heating temperature is preferably 1300.degree. C.
or less, while the heating time where the slab is held at a
temperature of the slab surface layer of 1000.degree. C. or more is
preferably 7 hours or less.
[0166] If the heating temperature exceeds 1300.degree. C. or the
heating time exceeds 7 hours, the decarburization of the slab
surface layer becomes remarkable. At the time of heating before
hardening, the austenite grains of the surface layer abnormally
grow and the impact resistance characteristic falls, so the heating
temperature is preferably 1300.degree. C. or less and the heating
time is preferably 7 hours or less. More preferably, the heating
temperature is 1280.degree. C. or less, while the heating time is 6
hours or less.
[0167] The finish hot rolling is ended at 650.degree. C. to
950.degree. C. in temperature. If the finish hot rolling
temperature is less than 650.degree. C., the rolling load
remarkably rises due to the increase of the deformation resistance
of the steel material. Furthermore, the amount of roll wear
increases and the productivity falls, so the finish hot rolling
temperature is made 650.degree. C. or more. Preferably it is
680.degree. C. or more.
[0168] On the other hand, if the finish hot rolling temperature
exceeds 950.degree. C., bulky scale is formed during passage
through the ROT (Run Out Table). Due to the scale, flaws form at
the surface of the steel plate. At the time of cold forging and/or
at the time when an impact load is applied after carburizing,
quenching, and tempering, cracks form starting from the flaws and
the impact resistance characteristic falls, so the finish hot
rolling temperature is made 950.degree. C. or less. Preferably it
is 920.degree. C. or less.
[0169] The cooling rate when cooling the hot rolled steel plate on
the ROT is preferably 10.degree. C./sec to 100.degree. C./sec. If
the cooling rate is less than 10.degree. C./sec, in the middle of
the cooling, it is not possible to suppress the formation of bulky
scale and the formation of flaws due to the same and the impact
resistance characteristic falls, so the cooling rate is preferably
10.degree. C./sec or more. More preferably it is 20.degree. C./sec
or more.
[0170] On the other hand, if cooling the hot rolled steel plate
from the surface layer to the inside part of the steel plate by a
cooling rate of over 100.degree. C./sec, the outermost layer part
is excessively cooled and bainite, martensite, and other low
temperature transformed structures are formed at the outermost
layer part.
[0171] After coiling, when the 100.degree. C. to room temperature
hot rolled steel plate is paid out, fine cracks form at low
temperature transformed structures. It is difficult to remove the
cracks in the following pickling process and cold rolling process.
At the time of cold forging and/or at the time when an impact load
is applied after carburizing, quenching, and tempering, cracks grow
starting from those cracks and invite a drop in the impact
resistance characteristic, so the cooling rate is preferably
100.degree. C./sec or less. More preferably it is 80.degree. C./sec
or less.
[0172] Note that, the cooling rate indicates the cooling ability
received from the cooling facilities in a water spray section at
the time when being cooled on the ROT down to the target
temperature of coiling from the time when the hot rolled steel
plate after finish hot rolling is water cooled at a water spray
section after passing through a non-water spray section. It does
not show the average cooling rate from the starting point of water
spray to the temperature at which the steel plate is coiled up by
the coiler.
[0173] The coiling temperature is made 400.degree. C. to
600.degree. C. This is a temperature lower than the general coiling
temperature. By coiling the hot rolled steel plate produced in the
above-mentioned condition in this temperature range, the structure
of the steel plate can be made a bainite structure in which
carbides are dispersed in fine ferrite.
[0174] If the coiling temperature is less than 400.degree. C., the
austenite, which was not transformed before coiling, transforms to
hard martensite. At the time of discharging the coiled hot rolled
steel plate, cracks form at the surface layer and the impact
resistance characteristics fall, so the coiling temperature is made
400.degree. C. or more. Preferably, it is 430.degree. C. or
more.
[0175] On the other hand, if the coiling temperature exceeds
600.degree. C., pearlite with a large lamellar spacing is formed,
high thermal stability bulky needle shaped carbides are formed, and
needle shaped carbides remain even after two-stage step type
annealing. At the time of cold forging, cracks occur and grow
starting from these needle shaped carbides, so the coiling
temperature is made 600.degree. C. or less. Preferably it is
570.degree. C. or less.
[0176] The hot rolled steel plate produced under the above
conditions was pickled, then held in two temperature regions for
two-stage step type annealing. By treating the hot rolled steel
plate by two-stage step type annealing, the carbides are controlled
in stability and the formation of carbides at the ferrite grain
boundaries is promoted.
[0177] First, the technical idea of two-stage step type annealing
will be explained.
[0178] By performing the first stage annealing in a temperature
region of the Ac1 point or less, the carbides are made to coarsen
and added metal elements are made to concentrate to raise the
thermal stability of the carbides. After that, the steel is raised
to the Ac1 or more in temperature to form austenite in the
structure, the fine carbides in the ferrite grains are made to
dissolve into the austenite, and coarse carbides are left in the
austenite.
[0179] By the subsequent gradual cooling, the austenite is
transformed to ferrite and raises the concentration of carbon in
the austenite. Along with gradual cooling, carbon atoms are
adsorbed at the carbides remaining in the austenite, the carbides
and austenite cover the grain boundaries of the ferrite, and,
finally, it becomes possible to form a structure with a large
amount of carbides present at the ferrite grain boundaries. For
this reason, it is clear that the structure of the present
invention cannot be formed by just simple annealing.
[0180] Below, the specific annealing conditions will be
explained.
[0181] First Stage Annealing
[0182] Heating rate up to annealing temperature: 30.degree. C./hour
to 150.degree. C./hour
[0183] Annealing temperature: 650.degree. C. to 720.degree. C.
[0184] Holding time at annealing temperature: 3 hours to 60
hours
[0185] The heating rate up to the first stage annealing temperature
is made 30.degree. C./hour to 150.degree. C./hour. If the heating
rate is less than 30.degree. C./hour, time is required for raising
the temperature and the productivity falls, so the heating rate is
made 30.degree. C./hour or more. Preferably, it is 40.degree.
C./hour or more.
[0186] On the other hand, if the heating rate is over 150.degree.
C./hour, the temperature difference between the outer
circumferential part and the inside part of the coil increases,
scratches and seizing occur due to the difference in heat
expansion, and relief shapes are formed at the steel plate surface.
At the time of cold forging, cracks occur starting from the relief
shapes and invite a drop in cold forgeability and a drop in impact
resistance characteristic after carburizing, quenching, and
tempering, so the heating rate is made 150.degree. C./hour or less.
Preferably, it is made 120.degree. C./hour or less.
[0187] The annealing temperature in the first stage annealing
(first stage annealing temperature) is made 650.degree. C. to
720.degree. C. If the first stage annealing temperature is less
than 650.degree. C., the carbides becomes insufficient in stability
and it becomes difficult to form carbides remaining in the
austenite in the second stage annealing, so the first stage
annealing temperature is made 650.degree. C. or more. Preferably it
is 670.degree. C. or more.
[0188] On the other hand, if the annealing temperature exceeds
720.degree. C., before the carbides rise in stability, austenite is
formed and it becomes impossible to control the above-mentioned
changes in structure, so the annealing temperature is made
720.degree. C. or less. Preferably it is 700.degree. C. or
less.
[0189] The holding time in the first stage annealing (first stage
holding time) is made 3 hours to 60 hours. If the first stage
holding time is less than 3 hours, the carbides become insufficient
in stability and it becomes difficult to form carbides remaining in
the second stage annealing, so the first stage holding time is made
3 hours or more. Preferably it is 10 hours or more.
[0190] On the other hand, if the first stage holding time exceeds
60 hours, no further improvement in stability of the carbides can
be expected. Furthermore, a drop in productivity is invited, so the
first stage holding time is made 60 hours or less. Preferably it is
50 hours or less.
[0191] Second Stage Annealing
[0192] Heating rate up to annealing temperature: 1.degree. C./hour
to 80.degree. C./hour
[0193] Annealing temperature: 725.degree. C. to 790.degree. C.
[0194] Holding time at annealing temperature: 3 hours to 50
hours
[0195] After finishing being held at the first stage annealing, the
hot rolled steel plate is heated up to the annealing temperature by
a heating rate of 1.degree. C./hour to 80.degree. C./hour. If
cooling without performing this second stage annealing, the ferrite
grain size does not become larger and the ideal structure cannot be
obtained.
[0196] In the second stage annealing, austenite is produced and
grows from the ferrite grain boundaries. By slowing the heating
rate, it is possible to suppress formation of nuclei of austenite.
In the structure obtained after gradual cooling, it becomes
possible to raise the rate of coverage of the grain boundaries of
the carbides. For this reason, the heating rate at the second stage
annealing is preferably small.
[0197] If the heating rate is less than 1.degree. C./hour, time is
required for raising the temperature and the productivity falls, so
the heating rate is made 1.degree. C./hour or more. Preferably it
is 10.degree. C./hour or more.
[0198] On the other hand, if the heating rate exceeds 80.degree.
C./hour, the temperature difference between the outer
circumferential part and inside part of the coil increases. Due to
the large difference in heat expansion due to deformation,
scratches and seizing occur and relief shapes are formed at the
surface of the steel plate. At the time of cold forging, cracks
form starting from the relief shapes leading to a drop in cold
forgeability and a drop in impact resistance characteristic after
carburizing, quenching, and tempering, so the heating rate is made
80.degree. C./hour or less.
[0199] The annealing temperature in the second stage annealing
(second stage annealing temperature) is made 725.degree. C. to
790.degree. C. If the second stage annealing temperature is less
than 725.degree. C., the amount of production of austenite becomes
smaller. After cooling after the second stage annealing, the number
ratio of carbides at the ferrite grain boundaries falls and,
further, the ferrite grain size becomes smaller. For this reason,
the second stage annealing temperature is made 725.degree. C. or
more. Preferably it is 735.degree. C. or more.
[0200] On the other hand, if the second stage annealing temperature
exceeds 790.degree. C., it becomes difficult to form carbides
remaining in the austenite and control to the above-mentioned
change of structure becomes difficult, so the second stage
annealing temperature is made 790.degree. C. or less. Preferably it
is 780.degree. C. or less.
[0201] The holding time in the second stage annealing (second stage
holding time) is made 1 hour to 50 hours. If the second stage
holding time is less than 1 hour, the amount of austenite which is
produced is small, the carbides in the ferrite grains are not
sufficiently dissolved, it becomes difficult to increase the number
ratio of carbides at the ferrite grain boundaries, and, further,
the ferrite grains become smaller in size, so the second stage
holding time is made 1 hour or more. Preferably, it is 5 hours or
more.
[0202] On the other hand, if the second stage holding time exceeds
50 hours, it is difficult to make carbides remain in the austenite,
so the second stage holding time is made 50 hours or less.
Preferably it is 45 hours.
[0203] Cooling After Annealing
[0204] Cooling stop temperature: 650.degree. C.
[0205] Cooling rate: 1.degree. C./hour to 100.degree. C./hour
[0206] After finishing being held at the second stage annealing,
the annealed hot rolled steel plate is gradually cooled down to
650.degree. C. by a cooling rate of 1.degree. C./hour to
100.degree. C./hour. If the stop temperature of gradual cooling
exceeds 650.degree. C., due to the cooling rate subsequently
exceeding 100.degree. C./hour down to room temperature,
nontransformed austenite transforms to pearlite or bainite, the
hardness increases, and the cold forgeability falls, so the cooling
stop temperature is made 650.degree. C.
[0207] To cool the austenite formed in the second stage annealing
and transform it to ferrite and to make carbon be adsorbed at the
carbides remaining in the austenite, a slower cooling rate is
preferable. If the cooling rate is less than 1.degree. C./hour, the
time required for cooling increases and the productivity falls, so
the cooling rate is made 1.degree. C./hour or more. Preferably it
is 10.degree. C./hour or more.
[0208] On the other hand, if the cooling rate exceeds 100.degree.
C./hour, austenite transforms to pearlite and the steel plate
increases in hardness so a drop in cold forgeability and a drop in
impact resistance characteristics after carburizing, quenching, and
tempering are invited, so the cooling rate is made 100.degree.
C./hour or less. Preferably it is 90.degree. C./hour.
[0209] Here, the cooling stop temperature is the temperature where
the cooling rate should be used for control. If cooling down to
650.degree. C. by a cooling rate of 1.degree. C./hour to
100.degree. C./hour, the cooling down to 650.degree. C. or less is
not particularly limited.
[0210] Note that, the atmosphere of the annealing is not limited to
any specific atmosphere. For example, it may be any of an
atmosphere of 95% or more of nitrogen, an atmosphere of 95% or more
of hydrogen, and an air atmosphere.
[0211] As explained above, according to the method of the present
invention of integrally managing the hot rolling conditions and
annealing conditions and controlling the structure of the steel
plate, it is possible to produce low carbon steel plate exhibiting
excellent cold forgeability in cold forging combining drawing and
thickening and, furthermore, excellent in impact resistance
characteristics after carburizing, quenching, and tempering.
EXAMPLES
[0212] Next, examples will be explained, but the levels in the
examples are illustrations of conditions employed for confirming
the workability and effects of the present invention. The present
invention is not limited to these illustrations of conditions. The
present invention can employ various conditions so long as not
deviating from the gist of the present invention and achieving the
object of the present invention.
[0213] A continuously cast slab (steel ingot) having a chemical
composition shown in Table 1 was heated at 1240.degree. C. for 1.8
hours, then was used for hot rolling. The finish hot rolling was
ended at 890.degree. C., the steel was cooled on a ROT by a
45.degree. C./sec cooling rate down to 520.degree. C. and was
coiled up at 510.degree. C. to produce a hot rolled coil with a
plate thickness of 5.2 mm.
TABLE-US-00001 TABLE 1 Chemical composition (mass %) No. C Si Mn P
S Al Cr Mo N O Ti B Remarks A 0.12 0.07 0.85 0.0154 0.0084 0.031
0.527 0.636 0.0068 0.0003 0.0095 0.0004 Invention steel B 0.13 0.03
0.76 0.0069 0.0046 0.011 1.483 0.017 0.0057 0.0122 0.0039 0.0001
Invention steel C 0.17 0.18 0.34 0.0027 0.0019 0.004 0.996 0.204
0.0002 0.0030 0.0004 0.0000 Invention steel D 0.21 0.17 0.81 0.0133
0.0073 0.032 0.563 0.173 0.0086 0.0166 0.0030 0.0002 Invention
steel E 0.23 0.24 0.64 0.0169 0.0095 0.046 1.934 0.731 0.0094
0.0063 0.0043 0.0003 Invention Steel F 0.25 0.06 0.84 0.0189 0.0099
0.062 1.043 0.195 0.0048 0.0196 0.0033 0.0002 Invention Steel G
0.26 0.19 0.45 0.0112 0.0075 0.017 1.024 0.003 0.0053 0.0176 0.0088
0.0002 Invention Steel H 0.30 0.05 0.95 0.0100 0.0012 0.091 0.586
0.591 0.0034 0.0170 0.0058 0.0004 Invention Steel I 0.34 0.28 0.92
0.0043 0.0017 0.078 1.905 0.860 0.0086 0.0188 0.0013 0.0000
Invention Steel J 0.36 0.01 0.51 0.0162 0.0006 0.069 0.705 0.943
0.0016 0.0125 0.0093 0.0002 Invention Steel K 0.39 0.10 0.85 0.0013
0.0074 0.033 1.635 0.736 0.0131 0.0143 0.0049 0.0002 Invention
Steel L 0.07 0.21 0.76 0.0166 0.0005 0.013 0.609 0.842 0.0050
0.0163 0.0086 0.0004 Comparative Steel M 0.11 0.19 0.78 0.0211
0.0069 0.002 1.379 0.941 0.0183 0.0186 0.0087 0.0003 Comparative
Steel N 0.14 0.24 0.79 0.0169 0.0020 0.040 1.328 1.071 0.0174
0.0155 0.0063 0.0002 Comparative Steel O 0.15 0.13 0.58 0.0018
0.0098 0.031 0.449 0.291 0.0181 0.0171 0.0024 0.0003 Comparative
Steel P 0.20 0.19 0.31 0.0025 0.0090 0.108 0.525 0.762 0.0195
0.0172 0.0084 0.0003 Comparative Steel Q 0.24 0.19 1.09 0.0196
0.0081 0.057 0.774 0.066 0.0066 0.0017 0.0039 0.0003 Comparative
Steel R 0.25 0.25 0.90 0.0099 0.0109 0.100 0.849 0.652 0.0178
0.0067 0.0017 0.0003 Comparative Steel S 0.27 0.31 0.36 0.0050
0.0063 0.094 0.822 0.006 0.0019 0.0099 0.0019 0.0003 Comparative
Steel T 0.29 0.26 0.63 0.0120 0.0049 0.003 2.236 0.011 0.0153
0.0092 0.0073 0.0001 Comparative Steel U 0.45 0.19 0.64 0.0009
0.0025 0.072 1.150 0.008 0.0109 0.0154 0.0063 0.0004 Comparative
Steel V 0.18 0.16 0.28 0.0129 0.0100 0.048 0.917 0.961 0.0085
0.0001 0.0005 0.0002 Comparative Steel W 0.15 0.21 0.89 0.0076
0.0009 0.091 1.293 0.005 0.0058 0.0209 0.0042 0.0000 Comparative
Steel X 0.25 0.26 0.46 0.0070 0.0052 0.029 1.089 0.610 0.0162
0.0161 0.0105 0.0001 Comparative Steel Y 0.28 0.20 0.76 0.0112
0.0026 0.083 1.193 0.713 0.0171 0.0083 0.0030 0.0006 Comparative
Steel Z 0.25 0.14 0.60 0.0033 0.0004 0.047 0.744 0.176 0.0208
0.0035 0.0054 0.0004 Comparative Steel
[0214] The hot rolled coil was pickled, the coil was loaded into a
box-type annealing furnace, the atmosphere was controlled to 95%
hydrogen-5% nitrogen, then the coil was heated from room
temperature up to 705.degree. C. by a heating rate of 100.degree.
C./hour, was held at 705.degree. C. for 36 hours to make the
temperature distribution inside the coil uniform, then was heated
by a 5.degree. C./hour heating rate up to 760.degree. C., and,
furthermore, was held at 760.degree. C. for 10 hours, then was
cooled down to 650.degree. C. by a 10.degree. C./hour cooling rate,
then was furnace cooled down to room temperature to prepare a
sample for evaluation of the characteristics.
[0215] The structure of the sample was observed by the
above-mentioned method. The crack length at the sample after cold
forging was measured by the above-mentioned method.
[0216] The carburization of the thickened sample was performed by
gas carburization. To make the carbon disperse from the inside of
the furnace atmosphere gas through the surface layer of the sample
to the inside of the steel, the sample was treated by holding it at
940.degree. C. for 120 minutes inside a furnace controlled to a
carbon potential of 0.5 mass % C, then was furnace cooled down to
room temperature.
[0217] Next, the sample was heated from room temperature to
840.degree. C., then was held for 20 minutes and quenched in
60.degree. C. oil. The hardened sample was held at 170.degree. C.
for 60 minutes, then air cooled for tempering to prepare a
carburized, quenched, and tempered sample.
[0218] The carburized, quenched, and tempered sample was evaluated
for impact resistance by a drop weight test. FIG. 2 schematically
shows an outline of the drop weight test for evaluating the impact
resistance characteristic of a carburized, quenched, and tempered
sample. The bottom of the cup of a carburized, quenched, and
tempered cup-shaped sample 4 was fastened by a fixture. On a side
surface of the cup, a weight 2 kg dropping weight (top side width:
50 mm, bottom side width: 10 mm, height: 80 mm, and length: 110 mm)
was allowed to freely drop from 4 m above to give an approximately
80J impact on the vertical wall part of the sample 4. The sample
was examined for the presence of any cracking and was evaluated for
the impact resistance characteristic.
[0219] A sample with no fracture or breakage observed as a result
of free dropping was evaluated as excellent in impact resistance
characteristics, that is, "OK", while a sample with a fracture or
breakage observed was evaluated as inferior in impact resistance,
that is, "NG".
[0220] Table 2 shows the results of measurement and results of
evaluation of the carbide size, pearlite area ratio, ferrite grain
size, Vickers hardness, ratio of the number of carbides at the
ferrite grain boundaries to number of carbides in the ferrite
grains, ratio of maximum crack length to plate thickness at the
vertical wall parts, and impact resistance characteristic in the
prepared samples.
TABLE-US-00002 TABLE 2 No. of Ratio carbides of Ferrite Pearlite at
grain maximum Carbide grain area Vickers boundaries/No. crack
Impact size size rate hardness of carbides length resistance
(.mu.m) (.mu.m) (%) (HV) inside grains (%) characteristic Remarks
A-1 1.11 23 1.2 106.0 8.67 1.5 OK Invention Steel B-1 0.92 18.5 0.9
106.1 5.72 2.0 OK Invention Steel C-1 0.87 23.5 1.3 107.4 4.87 2.3
OK Invention Steel D-1 1.07 19.4 1.3 117.4 7.65 2.3 OK Invention
Steel E-1 0.78 15 1.3 126.7 1.08 3.8 OK Invention Steel F-1 0.99
17.2 0.7 115.7 6.75 2.4 OK Invention Steel G-1 0.88 19.5 0.3 117.5
4.90 2.9 OK Invention Steel H-1 1.09 17.7 0.5 118.0 8.44 2.2 OK
Invention Steel I-1 0.84 13.1 0.5 140.2 1.76 3.4 OK Invention Steel
J-1 0.98 18.3 1.1 114.7 6.68 2.3 OK Invention Steel K-1 0.9 14.9
0.7 127.3 5.63 3.3 OK Invention Steel L-1 1.08 26.2 0.6 92.2 9.29
15.5 NG Comparative Steel M-1 0.96 19.7 1.1 114.4 0.12 15.4 NG
Comparative Steel N-1 0.96 18.1 8.8 191.7 0.18 22.9 NG Comparative
Steel O-1 1.06 22.7 0.1 107.4 7.16 1.8 NG Comparative Steel P-1
0.97 25.2 0.6 110.0 0.61 13.7 NG Comparative Steel Q-1 1.08 18.1
0.9 123.9 8.00 2.6 NG Comparative Steel R-1 1.03 17 1.2 128.8 7.06
18.8 NG Comparative Steel S-1 0.85 21.4 1.0 124.3 4.23 13.6 NG
Comparative Steel T-1 0.69 14.1 13.1 232.5 0.24 26.2 NG Comparative
Steel U-1 0.89 15.4 1.5 135.0 21.15 3.2 NG Comparative Steel V-1
0.9 25.8 0.3 105.3 0.67 11.5 NG Comparative Steel W-1 0.94 17.1 0.7
122 0.82 14.3 NG Comparative Steel X-1 0.89 19.3 0.1 121.8 0.51
15.3 NG Comparative Steel Y-1 0.95 16.1 0.4 126.8 6.06 15.4 NG
Comparative Steel Z-1 1 19.1 0.8 116.1 0.68 14.2 NG Comparative
Steel
[0221] As shown in Table 2, in Invention Steels A-1, B-1, C-1, D-1,
E-1, F-1, G-1, H-1, I-1, J-1, and K-1, in each case, the ratio of
the number of carbides at the ferrite grain boundaries to the
number of carbides in the ferrite grains is over 1, the Vickers
hardness is 100 HV to 180 HV, and the cold forgeability and impact
resistance characteristics after carburizing, quenching, and
tempering are excellent.
[0222] As opposed to this, in Comparative Steel L-1, the amount of
C is low and the hardness before cold forging is less than 100 HV,
so the cold forgeability is low. In Comparative Steels M-1, P-1,
and Z-1, P, Al, and N are excessively contained and, at the second
stage annealing, the amount of segregation at the y/a interfaces is
large, so formation of carbides at the grain boundaries is
suppressed.
[0223] In Comparative Steel S-1, Si is excessively contained and
ductility of the ferrite is low, so the cold forgeability is low.
In Comparative Steels N-1 and T-1, Mo and Cr are excessively
contained, so carbides finely disperse inside the ferrite grains
and the hardness exceeds 180 HV. In Comparative Steel Q-1, Mn is
excessively contained, so the impact resistance characteristic
after carburizing, quenching, and tempering is remarkably low.
[0224] In Comparative Steel 0-1, the amount of Cr is small and the
austenite grains at the surface layer abnormally coarsen at the
time of carburization, so the impact resistance is low. In
Comparative Steel R-1, S is excessively contained, so coarse MnS is
formed in the steel and the cold forgeability is low. In
Comparative Steel U-1, C is excessively contained, so coarse
carbides form inside the thickened layer of the steel and coarse
carbides remain even after the carburizing and quenching, so the
impact resistance characteristic is low.
[0225] In Comparative Steel V-1, the amount of Mn is small and the
carbides are hard to raise in stability, so the cold forgeability
and impact resistance characteristic after carburizing, quenching,
and tempering are low. In Comparative Steels W-1 and X-1, 0 and Ti
are excessively contained, so the oxides and TiC present in the
ferrite grains become site for formation of carbides in gradual
cooling after dual phase region annealing, the formation of
carbides at the grain boundaries is suppressed, and the cold
forgeability is low. In Comparative Steel Y-1, B is excessively
contained, so the cold forgeability is low.
[0226] Next, to investigate the effects of the manufacturing
conditions, slabs having the chemical compositions A, B, C, D, E,
F, G, H, I, J, and K shown in Table 1 were hot rolled and annealed
under the conditions shown in Table 3 to prepare annealed samples
of hot rolled plates of a thickness of 5.2 mm.
[0227] Table 4 shows the results of measurement and results of
evaluation of the carbide size, pearlite area ratio, ferrite grain
size, Vickers hardness, ratio of the number of carbides at the
ferrite grain boundaries to number of carbides in the ferrite
grains, ratio of maximum crack length to plate thickness at the
vertical wall parts, and impact resistance characteristics in the
prepared samples.
TABLE-US-00003 TABLE 3 Hot rolling conditions Annealing conditions
Finish ROT 1st stage 2nd stage Cooling hot cooling Heating Heating
rate Heating Soaking rolling rate Coiling rate Holding Holding rate
Holding Holding down to temp. time temp. (.degree. C./ temp.
(.degree. C./ temp. time (.degree. C./ temp. time 650.degree. C.
(.degree. C.) (hours) (.degree. C.) sec) (.degree. C.) hour)
(.degree. C.) (hours) hour) (.degree. C.) (hours) (.degree.
C./hour) Remarks A-2 -- -- 837.1 63 449.2 143.7 653.1 44.1 48.1
745.2 49.7 26.8 Inv. steel B-2 1162 0.5 905 35 409.3 101.3 706.5
32.9 2.2 773.5 13.6 84.6 Inv. steel C-2 1196 2.5 882.3 64 388 67.9
695.6 36.5 19.5 735.7 20.7 33.8 Comp. steel D-2 1060 5.6 826 21
496.8 55.9 683.2 49.2 31.4 789 41.4 99.5 Inv. steel E-2 1241 3.4
639 33 526.7 84.5 687.5 43 40.5 775.1 4.6 72.4 Comp. steel F-2 1271
0.8 714.7 91 412.1 71.4 709.2 63.1 23.8 769.7 14.4 10.8 Comp. steel
G-2 1265 5.6 844.4 52 627 133.6 671.7 44.2 28.9 750 39.7 92.1 Comp.
steel H-2 1287 5.9 752.9 30 567.5 139.4 709 22 16.4 731.9 28.9 57.7
Inv. steel I-2 1030 5.6 850.3 100 541.1 44.5 695.4 51.9 29.9 788.7
30.5 76.7 Inv. steel J-2 1258 0.6 741.9 76 543.9 35.9 720 33.7 78.8
773 1.3 1.4 Inv. steel K-2 1152 5.9 703 41 516.2 44.8 681.1 46.5
2.8 721 35.3 72 Comp. steel A-3 1138 4.5 824.7 53 472.1 85.9 687.2
35.8 88 773 29.2 92.8 Comp. steel B-3 1254 1.7 717.5 70 526.9 42.6
715.1 15.6 48 740.3 1.8 6.7 Comp. steel C-3 1292 2.5 870.5 35 535.8
36.2 669 47.7 37.9 784.7 15.9 95 Inv. steel D-3 1010 1.4 812.3 69
403.5 124.9 709.2 11 72.3 734.5 8.8 0.4 Comp. steel E-3 1007 0.8
689.5 93 466.5 158 700 54.6 55.8 775.8 45.6 16.7 Comp. steel F-3
1062 3.6 656.1 25 567.7 126.5 699.9 42.7 24.9 731.2 29.5 13.1 Inv.
steel G-3 1137 2.8 851.7 48 446.5 94.4 682.2 14.8 77.1 757.8 42.6
16.4 Inv. steel H-3 1089 5.6 767.6 10 468.4 146.7 685.1 49.1 8
768.9 45.6 34.8 Inv. steel I-3 1288 1.7 682.3 55 443 109.2 683.1
39.2 63.7 769.9 29.9 116 Comp. steel J-3 1037 5.6 879 53 519.9 59.1
719.2 59.8 77.5 727 57 47.3 Comp. steel K-3 -- -- 664.6 85 407.9
75.7 665.4 32 40.2 777 21.4 32 Inv. steel A-4 1164 2.5 710.6 80
503.2 135.8 669.3 47.1 71.3 731.3 25.8 69.3 Inv. steel B-4 1136 4.8
695.9 28 445 90.1 701.4 21 0.2 776.1 42.9 87.4 Comp. steel C-4 1046
1.1 667.3 80 405.1 54.9 710.4 47.8 56.4 804 17.1 97.5 Comp. steel
D-4 1253 4.5 708.1 29 493.9 74 737 57.3 38.4 755.7 33.6 55.4 Comp.
steel E-4 1018 5.0 870.2 54 560.4 141.7 710.9 9.2 19.7 748 16.2
96.5 Inv. steel F-4 1229 6.2 962 71 475.1 99.3 672.1 24 28.4 736.3
24.1 7.9 Comp. steel G-4 1139 6.4 749.8 19 535 144.8 639 39.3 18
783.9 32.5 91.3 Comp. steel H-4 1119 4.8 731.7 66 456.1 13 668.3
46.6 55.8 744.7 19 43.8 Comp. steel I-4 1101 4.2 686 33 569.4 73.6
656.1 31.1 20.3 740.4 31.5 24.2 Inv. steel J-4 1042 0.6 750.8 69
545.5 40 719.4 1.4 71.9 729.1 38.3 11.9 Comp. steel K-4 1268 4.5
912.2 44 521.5 92.5 652.2 54.7 39.2 745.5 37.4 48.1 Inv. steel
TABLE-US-00004 TABLE 4 No. of Ratio carbides of Ferrite Pearlite at
grain maximum Carbide grain area Vickers boundaries/No. crack
Impact size size rate hardness of carbides length resistance
(.mu.m) (.mu.m) (%) (HV) inside grains (%) characteristic Remarks
A-2 0.71 19.6 1.2 103.5 9.19 1.3 OK Invention Steel B-2 0.73 22.8
0.7 101.6 1.90 2.9 OK Invention Steel C-2 1 18.4 1.0 111.4 6.83 2.1
NG Comparative Steel D-2 0.57 25.0 1.4 106.5 2.16 3.3 OK Invention
Steel E-2 0.54 13.7 0.9 126.6 2.57 4.3 OK Comparative Steel F-2
2.23 38.7 0.8 98.4 4.24 12.9 NG Comparative Steel G-2 0.28 16.2 1.1
120.0 1.56 14.9 NG Comparative Steel H-2 0.55 10.2 0.4 127.3 1.69
5.2 OK Invention Steel I-2 0.57 18.4 0.2 125.5 1.89 4.8 OK
Invention Steel J-2 1.95 29.5 1.1 104.4 18.06 0.7 OK Invention
Steel K-2 0.76 8.8 0.7 136.3 0.08 17.9 NG Comparative Steel A-3
0.59 23.5 0.1 97.9 2.45 22.6 NG Comparative Steel B-3 0.7 8.3 0.1
129.3 0.16 16.9 NG Comparative Steel C-3 0.48 25.6 1.1 104.3 2.71
2.8 OK Invention Steel D-3 2.34 16.9 0.1 128.6 21.84 21.5 NG
Comparative Steel E-3 0.86 28.4 0.6 107.9 3.32 5.0 NG Comparative
Steel F-3 0.58 10.3 0.2 124.9 1.71 5.0 OK Invention Steel G-3 0.72
25.5 1.4 107.8 7.56 1.8 OK Invention Steel H-3 0.78 22.1 1.4 105.2
3.76 2.4 OK Invention Steel I-3 0.49 16.6 9.5 186.2 2.05 20.5 NG
Comparative Steel J-3 1.18 17.9 9.6 189.7 1.52 14.3 NG Comparative
Steel K-3 0.74 19.1 0.5 113.9 9.88 1.8 OK Invention Steel A-4 0.47
11.1 0.3 119.0 2.75 3.7 OK Invention Steel B-4 0.57 26.2 0.8 93.1
2.08 2.5 OK Comparative Steel C-4 0.84 40.2 10.6 195.0 1.59 19.2 NG
Comparative Steel D-4 1.26 26.4 8.9 196.5 1.23 14.5 NG Comparative
Steel E-4 0.37 11.2 0.7 131.7 1.29 6.2 OK Invention Steel F-4 0.7
11.6 1.4 121.7 13.91 6.2 NG Comparative Steel G-4 0.48 24.6 12.4
211.2 3.03 18.4 NG Comparative Steel H-4 0.62 12.6 0.3 120.9 5.65
2.9 OK Comparative Steel I-4 0.25 9.7 0.0 143.3 2.06 5.7 OK
Invention Steel J-4 0.34 13.1 14.2 218.2 3.13 18.7 NG Comparative
Steel K-4 0.38 12.4 0.1 125.8 3.44 3.8 OK Invention Steel
[0228] In Comparative Steel E-3, the finish hot rolling temperature
is low, the rolling load increases, and the productivity falls. In
Comparative Steel D-2, the finish hot rolling temperature is high
and scale flaws form at the surface of the steel plate, so when
subjected to a wear resistance test after quenching and tempering,
fractures and peeling occur starting from the scale flaws and the
wear resistance characteristic falls. In Comparative Steel F-2, the
cooling rate at the ROT (Run Out Table) is slow and a drop of
productivity and formation of scale flaws are invited.
[0229] In Comparative Steel C-4, the cooling rate at the ROT is
100.degree. C./sec and the outermost layer part of the steel plate
is excessively cooled, so fine cracks formed at the outermost layer
part. In Comparative Steel C-2, the coiling temperature is low,
large amounts of bainite, martensite, and other low temperature
transformed structures are formed causing embrittlement, fractures
frequently form at the time of pay out from the hot rolled coil,
and the productivity falls. Furthermore, in a sample taken from a
cracked piece, the wear resistance characteristic is low.
[0230] In Comparative Steel G-2, the coiling temperature is high,
bulky pearlite of lamellar spacing is formed in the hot rolled
structure, the needle shaped coarse carbides become high in thermal
stability, and the above carbides remain in the steel plate even
after two-stage step type annealing, so the machinability is low.
In Comparative Steel H-4, the heating rate in the first stage
annealing of the two-stage step type annealing is slow, so the
productivity is low.
[0231] In Comparative Steel E-3, the heating rate in the first
stage annealing is fast, so the temperature difference between the
inside part and inside and outside circumferential parts of the
coil becomes larger, scratches and seizing occur due to the
difference in thermal expansion, and, when used for evaluating and
testing the wear resistance characteristic after quenching and
tempering, cracks and peeling occur from the flaw parts and the
wear resistance characteristic falls.
[0232] In Comparative Steel G-4, the holding temperature in the
first stage annealing (annealing temperature) is low, the
coarsening treatment of carbides at the Ac1 point or less is
insufficient, and the carbides are insufficient in thermal
stability, so carbides remaining at the second stage annealing are
reduced and pearlite transformation in the structure after gradual
cooling cannot be suppressed, so the machinability is low.
[0233] In Comparative Steel D-4, the holding temperature in the
first stage annealing (annealing temperature) is high, austenite is
formed during the annealing, and the carbides cannot be raised in
stability, so pearlite is formed after annealing, the Vickers
hardness exceeds 180 HV, and the machinability is low. In
Comparative Steel J-4, the holding time in the first stage
annealing is short and the stability of carbides cannot be raised,
so the machinability is low.
[0234] In Comparative Steel F-2, the holding time in the first
stage annealing is long, the productivity is low, and further
seizing flaws occur and the wear resistance characteristic is low.
In Comparative Steel B-4, the heating rate in the second stage
annealing in the two-stage step type annealing is slow, so the
productivity is low. In Comparative Steel A-3, the heating rate at
the second stage annealing is fast, so the temperature difference
between the inside part and outer circumferential part of the coil
become greater, scratches and seizing occur due to the large
difference in heat expansion due to deformation, and the wear
resistance characteristic after quenching and tempering is low.
[0235] In Comparative Steel K-2, the holding temperature in the
second stage annealing (annealing temperature) is low, the amount
of production of austenite is small, and the ratio of number of
carbides at the ferrite grain boundaries cannot be increased, so
the machinability is low. In Comparative Steel C-4, the holding
temperature at the second stage annealing (annealing temperature)
is high and dissolution of the carbides during the annealing is
promoted, so it becomes difficult to form carbides at the grain
boundaries after the gradual cooling and further pearlite is
produced, the Vickers hardness exceeds 180 HV, and the
machinability is low.
[0236] In Comparative Steel J-3, the holding time at the second
stage annealing is long and dissolution of the carbides is
promoted, so the machinability is low. In Comparative Steel D-3,
the cooling rate from second stage annealing to 650.degree. C. is
slow, the productivity is low, coarse carbides are formed in the
structure after gradual cooling, cracks are formed starting from
the coarse carbides at the time of cold forging, and the cold
forgeability falls. In Comparative Steel 1-3, the cooling rate from
second stage annealing to 650.degree. C. is fast, the pearlite
transformation occurs at the time of cooling, and the hardness
increases, so the cold forgeability is low.
[0237] Next, to investigate the allowable contents of the other
elements, continuously cast slabs (steel ingots) having the
chemical compositions shown in Table 5 and Table 6 (continuation of
Table 5) were heated at 1240.degree. C. for 1.8 hours, then were
used for hot rolling. The finish hot rolling was ended at
890.degree. C., the steels were cooled on a ROT by a 45.degree.
C./sec cooling rate down to 520.degree. C. and were coiled up at
510.degree. C. to produce hot rolled coils with a plate thickness
of 5.2 mm.
TABLE-US-00005 TABLE 5 Chemical composition (mass %) C Si Mn P S Al
N O Ti Cr Mo B Nb V AA 0.13 0.01 0.96 0.0076 0.0063 0.011 0.0139
0.0112 0.0043 0.509 0.869 0.0001 0.028 AB 0.16 0.25 0.70 0.0063
0.0087 0.083 0.0162 0.0119 0.0028 0.618 0.680 0.0001 0.029 AC 0.18
0.11 0.97 0.0007 0.0043 0.073 0.0067 0.0076 0.0095 1.199 0.402
0.0004 AD 0.22 0.29 0.69 0.0145 0.0020 0.007 0.0077 0.0005 0.0015
1.400 0.807 0.0003 0.012 AE 0.22 0.21 0.61 0.0067 0.0072 0.002
0.0008 0.0093 0.0072 1.130 0.422 0.0002 0.004 AF 0.27 0.16 0.60
0.0098 0.0032 0.079 0.0011 0.0011 0.0042 1.197 0.010 0.0003 0.070
AG 0.28 0.04 0.79 0.0075 0.0035 0.049 0.0116 0.0137 0.0028 0.862
0.802 0.0001 0.040 AH 0.28 0.07 0.43 0.0064 0.0058 0.006 0.0089
0.0025 0.0036 1.346 0.510 0.0000 0.046 0.065 AI 0.30 0.10 0.80
0.0047 0.0045 0.061 0.0196 0.0027 0.0057 0.961 0.002 0.0002 0.022
0.094 AJ 0.31 0.27 0.60 0.0121 0.0046 0.077 0.0105 0.0153 0.0029
1.649 0.013 0.0001 0.002 0.084 AK 0.36 0.07 0.60 0.0093 0.0040
0.028 0.0016 0.0022 0.0043 1.722 0.009 0.0004 AL 0.36 0.29 0.89
0.0187 0.0055 0.048 0.0190 0.0114 0.0031 0.881 0.893 0.0004 0.092
AM 0.37 0.06 0.84 0.0002 0.0002 0.072 0.0102 0.0092 0.0020 0.581
0.345 0.0001 0.015 0.037 AN 0.37 0.19 0.45 0.0139 0.0049 0.033
0.0050 0.0147 0.0055 0.540 0.077 0.0001 0.039 0.012 AO 0.39 0.29
1.00 0.0176 0.0009 0.068 0.0142 0.0176 0.0070 1.951 0.330 0.0002 AP
0.39 0.27 0.82 0.0075 0.0025 0.025 0.0116 0.0076 0.0077 0.981 0.387
0.0002 0.062 AQ 0.39 0.17 0.43 0.0194 0.0034 0.014 0.0037 0.0036
0.0014 1.480 0.626 0.0002 0.018 Chemical composition (mass %) Cu W
Ta Ni Sn Sb As Mg Ca Y Zr La Ce Remarks AA 0.095 0.021 0.028 0.034
0.029 0.001 0.047 0.020 Invention steel AB 0.082 0.086 0.053 0.086
0.016 0.024 0.043 0.005 0.028 0.034 0.046 Invention steel AC 0.008
0.031 0.029 0.008 0.002 0.045 0.011 0.039 0.006 Invention steel AD
0.019 0.019 0.053 0.013 0.029 0.034 0.026 Invention steel AE 0.054
0.003 0.085 0.027 0.002 0.001 0.003 0.026 0.031 Invention steel AF
0.038 0.014 0.048 0.006 0.042 0.039 0.009 0.013 Invention steel AG
0.076 0.073 0.038 0.009 0.049 0.017 0.045 0.026 Invention steel AH
0.076 0.007 0.002 0.024 0.019 0.012 0.029 0.011 Invention steel AI
0.092 0.012 0.041 0.038 0.008 0.010 0.014 Invention steel AJ 0.086
0.080 0.048 0.027 0.041 0.021 0.030 0.001 Invention steel AK 0.003
Invention steel AL 0.095 0.042 0.078 0.005 0.006 0.046 0.032
Invention steel AM 0.058 0.062 0.048 0.011 0.046 0.006 0.002 0.021
0.042 Invention steel AN 0.067 0.093 0.021 0.006 0.044 0.019 0.016
0.001 0.040 Invention steel AO 0.005 Invention steel AP Invention
steel AQ 0.068 0.002 0.005 0.029 0.033 0.031 0.004 0.023 Invention
steel
TABLE-US-00006 TABLE 6 (Continuation of Table 5) Chemical
composition (mass %) C Si Mn P S Al N O Ti Cr Mo B Nb V AR 0.12
0.13 0.95 0.0195 0.0083 0.06 0.0132 0.0033 0.0064 0.554 0.014
0.0005 0.045 0.006 AS 0.15 0.11 0.52 0.0122 0.0050 0.008 0.0177
0.0011 0.0008 0.773 0.632 0.0001 0.002 0.038 AT 0.15 0.06 0.70
0.0115 0.0060 0.036 0.0139 0.0100 0.0083 0.663 1.046 0.0005 0.001
AU 0.15 0.09 0.57 0.0165 0.0029 0.048 0.0154 0.0040 0.0074 0.811
0.314 0.0002 AV 0.16 0.18 0.54 0.0058 0.0064 0.086 0.0051 0.0079
0.0070 1.406 0.161 0.0002 0.116 AW 0.19 0.27 0.67 0.0199 0.0009
0.083 0.0004 0.0009 0.0054 0.513 0.728 0.0002 0.032 AX 0.19 0.24
1.00 0.0050 0.0010 0.084 0.0016 0.0162 0.0095 0.904 0.841 0.0004
0.080 0.077 AY 0.24 0.24 0.53 0.0094 0.0005 0.018 0.0060 0.0032
0.0084 1.688 0.811 0.0003 0.044 AZ 0.24 0.03 0.42 0.0050 0.0018
0.088 0.0076 0.0008 0.0002 1.216 0.844 0.0004 0.067 0.025 BA 0.28
0.08 0.58 0.0005 0.0021 0.045 0.0094 0.0047 0.0032 2.037 0.076
0.0000 0.050 BB 0.30 0.03 0.60 0.0044 0.0094 0.064 0.0015 0.0021
0.0075 1.285 0.513 0.0002 0.065 BC 0.31 0.02 0.57 0.0130 0.0074
0.041 0.0050 0.0157 0.0001 1.379 0.004 0.0001 0.032 0.051 BD 0.33
0.30 0.84 0.0140 0.0095 0.057 0.0085 0.0169 0.0030 1.648 0.182
0.0000 BE 0.34 0.29 0.58 0.0004 0.0087 0.072 0.0011 0.0092 0.0019
1.094 0.007 0.0003 0.091 0.079 BF 0.34 0.29 0.94 0.0149 0.0024
0.022 0.0167 0.0036 0.0050 1.561 0.856 0.0003 0.097 BG 0.36 0.36
0.60 0.0157 0.0088 0.086 0.0198 0.0064 0.0012 0.934 0.268 0.0001 BH
0.36 0.28 0.63 0.0099 0.0091 0.032 0.0098 0.0051 0.0029 1.624 0.011
0.0000 0.047 0.044 BI 0.37 0.14 1.17 0.0014 0.0048 0.094 0.0151
0.0113 0.0003 1.210 0.003 0.0000 0.085 BJ 0.39 0.19 0.95 0.0071
0.0012 0.094 0.0014 0.0186 0.0082 1.317 0.849 0.0002 0.051 0.118 BK
0.45 0.17 0.61 0.0115 0.0055 0.054 0.0024 0.0167 0.0086 0.688 0.118
0.0000 0.043 Chemical composition (mass %) Cu W Ta Ni Sn Sb As Mg
Ca Y Zr La Ce Remarks AR 0.075 0.084 0.024 0.027 0.049 0.052
Comparative steel AS 0.025 0.055 0.041 0.003 0.058 0.001 0.040
0.044 0.013 0.026 Comparative steel AT 0.012 0.048 0.011 0.011
0.031 0.026 0.029 Comparative steel AU 0.098 0.095 0.047 0.008
0.037 0.053 0.011 Comparative steel AV 0.078 0.005 0.048 0.027
0.035 0.040 0.006 Comparative steel AW 0.131 0.021 0.036 0.012
0.025 0.004 0.036 0.040 0.015 0.006 Comparative steel AX 0.087
0.045 0.011 0.040 0.014 0.043 0.058 0.020 0.041 0.046 Comparative
steel AY 0.014 0.095 0.025 0.005 0.052 0.027 0.040 0.027 0.027
Comparative steel AZ 0.097 0.073 0.106 0.009 0.044 0.017 0.020
0.022 0.025 0.025 0.014 Comparative steel BA 0.031 0.036 0.020
0.008 0.027 0.012 0.050 0.026 0.017 Comparative steel BB 0.051
0.091 0.072 0.005 0.055 0.019 0.012 0.002 Comparative steel BC
0.034 0.003 0.105 0.064 0.005 0.005 0.029 0.028 0.049 Comparative
steel BD 0.059 Comparative steel BE 0.099 0.019 0.049 0.038 0.043
0.062 0.028 0.022 0.032 Comparative steel BF 0.023 0.049 0.079
0.049 0.001 0.001 0.002 0.051 Comparative steel BG 0.070 0.005
0.040 0.047 0.016 0.023 0.032 0.002 Comparative steel BH 0.107
0.014 0.068 0.045 0.003 0.014 0.018 Comparative steel BI 0.008
0.028 0.017 0.029 0.006 0.009 0.027 0.037 0.038 Comparative steel
BJ 0.077 0.096 0.076 0.012 0.046 0.036 0.047 0.011 0.022 0.043
0.011 Comparative steel BK 0.083 0.048 0.016 0.025 0.039 0.004
0.046 Comparative steel
[0238] The hot rolled coils were pickled, the hot rolled coils were
loaded into a box-type annealing furnace, the atmosphere was
controlled to 95% hydrogen-5% nitrogen, then the coils were heated
from room temperature up to 705.degree. C. by a heating rate of
100.degree. C./hour, were held at 705.degree. C. for 36 hours to
make the temperature distribution inside the coils uniform, then
were heated by a 5.degree. C./hour heating rate up to 760.degree.
C., and, furthermore, were held at 760.degree. C. for 10 hours,
then were cooled down to 650.degree. C. by a 10.degree. C./hour
cooling rate, then were furnace cooled down to room temperature to
prepare samples for evaluation of the characteristics.
[0239] Note that, the structures of the samples were observed by
the above-mentioned method while the crack lengths present in the
samples after cold forging were measured by the above-mentioned
method.
[0240] Table 7 shows the results of measurement and results of
evaluation of the carbide size, pearlite area ratio, ferrite grain
size, Vickers hardness, ratio of the number of carbides at the
ferrite grain boundaries to number of carbides in the ferrite
grains, ratio of maximum crack length to plate thickness at the
vertical wall parts, and impact resistance characteristics in the
prepared samples.
TABLE-US-00007 TABLE 7 No. of carbides at t grain Ratio boundaries/
of Ferrite Pearlite No. of maximum Carbide grain area Vickers
carbides crack Impact size size rate hardness inside length
resistance (.mu.m) (.mu.m) (%) (HV) grains (%) characteristic
Remarks AA-1 1.14 23.5 1.2 103.1 9.53 0.8 OK Invention steel AB-1
1.04 20 0.7 119.7 7.19 1.5 OK Invention steel AC-1 0.99 17.5 0.7
117.2 7.03 1.4 OK Invention steel AD-1 0.89 16.4 0.3 127.0 5.25 2.0
OK Invention steel AE-1 0.94 17.8 0.7 119.5 5.81 1.7 OK Invention
steel AF-1 0.88 16.8 1.4 121.6 4.71 1.9 OK Invention steel AG-1
1.02 17.6 0.0 114.8 7.35 1.3 OK Invention steel AH-1 0.86 19.1 0.7
111.6 4.56 1.6 OK Invention steel AI-1 1 17.4 0.0 119.7 6.75 1.6 OK
Invention steel AJ-1 0.79 15.3 0.8 132.5 3.60 2.5 OK Invention
steel AK-1 0.79 15.2 1.4 122.3 3.66 2.1 OK Invention steel AL-1
1.03 16.1 0.8 136.0 7.59 2.0 OK Invention steel AM-1 1.06 17.6 0.3
120.6 7.63 1.5 OK Invention steel AN-1 0.97 19.1 0.1 124.2 5.46 1.9
OK Invention steel AO-1 0.83 13.4 1.0 142.6 4.61 2.6 OK Invention
steel AP-1 0.98 16 0.2 135.4 6.32 2.2 OK Invention steel AQ-1 0.81
17.3 1.1 126.3 3.88 2.2 OK Invention steel AR-1 1.1 22.8 0.7 111.3
0.15 14.9 NG Comparative steel AS-1 0.99 22 1.0 105.6 0.28 14.0 NG
Comparative steel AT-1 1.07 21.3 7.8 188.0 0.19 22.6 NG Comparative
steel AU-1 0.98 21.2 1.3 106.4 6.89 21.9 NG Comparative steel AV-1
0.83 18.1 13.1 243.1 0.44 26.4 NG Comparative steel AW-1 0.94 19.2
0.6 122.5 0.04 16.5 NG Comparative steel AX-1 1.06 17.5 0.4 126.0
10.11 21.4 NG Comparative steel AY-1 1.06 16.7 1.5 123.4 4.01 14.8
NG Comparative steel AZ-1 0.81 19.8 0.4 107.8 0.89 12.4 NG
Comparative steel BA-1 0.87 15.2 11.4 234.4 0.25 26.3 NG
Comparative steel BB-1 0.73 16.7 1.0 115.2 0.11 15.5 NG Comparative
steel BC-1 0.9 16.8 9.8 212.6 0.94 21.2 NG Comparative steel BD-1
0.88 14.1 1.2 138.2 4.53 13.4 NG Comparative steel BE-1 0.9 16.4
0.6 133.4 4.83 11.2 NG Comparative steel BF-1 0.85 14.6 1.2 137.2
0.73 16.2 NG Comparative steel BG-1 0.9 16.5 0.7 138.8 4.97 13.8 NG
Comparative steel BH-1 0.93 14.9 10.6 228.7 0.85 23.4 NG
Comparative steel BI-1 0.91 16.6 0.3 128.4 7.24 3.0 NG Comparative
steel BJ-1 0.82 14.5 11.4 240.6 0.24 26.8 NG Comparative steel BK-1
1.01 16.6 0.8 131.5 22.08 1.1 NG Comparative steel
[0241] As shown in Table 7, in each of Invention Steels AA-1, AB-1,
AC-1, AD-1, AE-1, AF-1, AG-1, AH-1, AI-1, AJ-1, AK-1, AL-1, AM-1,
AN-1, AO-1, AP-1, and AQ-1, the ratio of the number of carbides at
the ferrite grain boundaries to the number of carbides in the
ferrite grains exceeds 1, the Vickers hardness is 100 HV to 180 HV,
and the cold forgeability and impact resistance characteristic
after carburizing, quenching, and tempering are excellent.
[0242] As opposed to this, in each of Comparative Steels AR-1,
AS-1, AW-1, AZ-1, BB-1, and BF-1, La, As, Cu, Ni, Sb, and Ce are
excessively contained and the amount of segregation at the y/a
interface becomes greater at the time of second stage annealing, so
formation of carbides at the grain boundaries is suppressed. In
Comparative Steel BG-1, Si is excessively contained and the
ductility of the ferrite is low, so the cold forgeability is
low.
[0243] In each of Comparative Steels AT-1, AV-1, BA-1, BC-1, BH-1,
and BJ-1, Mo, Nb, Cr, Ta, W, and V are respectively excessively
contained, so carbides finely disperse inside the ferrite grains
and the hardness exceeds 180 HV. In Comparative Steel BF-1, Mn is
excessively contained, so the impact resistance characteristic
after carburizing, quenching, and tempering is remarkably low.
[0244] In each of Comparative Steels AU-1, AX-1, AY-1, and BE-1,
Zr, Ca, Mg, and Y are respectively excessively contained, coarse
oxides or nonmetallic inclusions are formed in the steel, cracks
form starting from the coarse oxides or coarse nonmetallic
inclusions at the time of cold forging, and the cold forgeability
falls. In Comparative Steel BD-1, Sn is excessively contained, the
ferrite becomes brittle, and the cold forgeability is low. In
Comparative Steel BK-1, C is excessively contained, so coarse
carbides form at the inside of the increased thickness part of the
steel, coarse carbides remain even after carburizing and quenching,
and the impact resistance characteristic also falls.
[0245] Next, to investigate the effects of the manufacturing
conditions, slabs having the chemical compositions of AA, AB, AC,
AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP, and AQ shown in
Table 5 were hot rolled and annealed under the conditions shown in
Table 8 to fabricate annealed samples of hot rolled plates of
thicknesses of 5.2 mm.
TABLE-US-00008 TABLE 8 Hot rolling conditions Annealing conditions
Finish ROT 1st stage 2nd stage hot cooling Heating Heating Cooling
rate Heating Soaking rolling rate Coiling rate Holding Holding rate
Holding Holding down to temp. time temp. (.degree. C./ temp.
(.degree. C./ temp. time (.degree. C./ temp. time 650.degree. C.
(.degree. C.) (hours) (.degree. C.) sec) (.degree. C.) hour)
(.degree. C.) (hours) hour) (.degree. C.) (hours) (.degree.
C./hour) Remarks AA-2 1143 0.4 836.3 11 403.5 124 644 31.9 36.2 774
14.2 47.5 Comp. steel AB-2 1199 4.5 939.4 17 435.8 123.4 692.3 38.9
52.5 766.7 34.7 35.3 Inv. steel AC-2 1043 4.2 641 61 415.4 75.8
703.9 50.4 24.2 741.7 15.1 88.1 Comp. steel AD-2 1272 0.8 812.8 65
467.4 96.7 692.7 59.9 39.6 796 7.3 60.5 Comp. steel AE-2 1164 0.5
902.3 53 554 131.8 670.2 15.7 78.3 734.5 49.8 66.4 Inv. steel AF-2
1238 4.5 832 42 435.3 40.5 671.4 2.7 29.3 757.4 11.8 6.1 Comp.
steel AG-2 1226 3.4 830.5 14 503.9 171 688.5 28.1 14.6 747 10.8 51
Comp. steel AH-2 1050 1.4 833 48 558.2 102 706.7 22.2 26.9 744.5
49.7 80.7 Inv. steel AI-2 1067 2.8 690.7 26 452 45.7 662.2 31.5
53.6 737.7 10.2 29.6 Inv. steel AJ-2 1008 1.4 850.8 53 564.3 70.8
693.9 29.9 4.9 787.8 41.6 83.3 Inv. steel AK-2 1166 4.8 729.7 23
510 43.9 667.2 13.6 32 753 7.5 40.4 Inv. steel AL-2 1150 1.7 802.3
60 516.2 95.6 660.4 47.1 21.8 742.7 30 36.1 Inv. steel AM-2 1083
1.4 779.6 66 534.9 86.2 685.7 25.5 22.8 754 28.5 15.4 Inv. steel
AN-2 1084 1.1 799.6 96 552.8 74.1 712 52.5 53.4 739 43.4 17.4 Inv.
steel AO-2 1194 2.2 710.1 16 592.8 42.5 689.2 50.7 67.4 748 27.1
86.6 Inv. steel AP-2 1101 2.5 800.4 61 480.4 119 702 12.1 71.1
781.4 27.7 123 Comp. steel AQ-2 -- -- 842.1 24 403.5 43.2 693.2
29.7 23.7 773.9 31.1 48.3 Inv. steel AA-3 1173 3.4 758.8 45 521.1
81.2 686.8 18.4 8.6 751 21.3 27 Inv. steel AB-3 1023 2.8 834.7 10
500.4 114 677.3 27.6 59.8 745.4 21.5 64.7 Inv. steel AC-3 1296 0.6
725.4 87 416.2 142.7 682.2 58.8 40.3 740.1 10.7 4.9 Inv. steel AD-3
1143 3.9 946.1 91 588.6 36.6 677.3 53.1 8.8 737.3 6.9 72 Inv. steel
AE-3 1110 0.6 874.7 44 454.5 106.6 678.1 31.5 57.6 770 18.1 42.7
Inv. steel AF-3 1054 5.9 684.6 97 526.2 49.1 693.9 32.5 44.5 777.1
28.3 40.4 Inv. steel AG-3 1163 4.2 773 63 501.3 52.9 701.8 9.8 18.2
751.3 44.4 43.7 Inv. steel AH-3 1122 0.5 884.9 36 613 84.2 686.6
58.5 2.9 731.3 33.9 39 Comp. steel AI-3 1245 7.0 797.5 44 433.1
136.2 698.7 54.9 91 778 48.4 60.6 Comp. steel AJ-3 1281 4.8 895.7
100 562.9 127 673.4 18.9 13.2 748 28.7 78.6 Inv. steel AK-3 1164
5.0 910.4 86 486.1 123.9 653.8 56.4 4.4 758.9 2.1 41.5 Comp. steel
AL-3 1212 0.8 691 87 486.5 61.8 693.4 30.6 15.3 709 23.5 85.9 Comp.
steel AM-3 1098 4.5 798.7 44 543.5 130.7 726 4.9 46 770.2 43.9 19.4
Comp. steel AN-3 1022 7.0 879.7 48 391 146.2 657 13.7 10.9 765.1
32.7 89.9 Comp. steel AO-3 1094 3.9 725.5 67 488.4 137.2 671.4 42.5
71.1 778.9 19.5 40.3 Inv. steel AP-3 1083 1.1 919.3 99 573.9 42.3
715.3 25.6 25.8 730.3 14.4 97.7 Inv. steel AQ-3 1096 3.1 743.6 22
527.4 113.3 652.7 46.3 13.6 751.3 14.3 0.3 Comp. steel AA-4 -- --
834.5 36 533.2 83.3 698 47.6 29.4 764.4 22.3 86.5 Inv. steel AB-4
1300 5.3 719.6 27 491.8 32 659.5 32.9 69 784.1 6.1 50.4 Inv. steel
AC-4 1049 0.6 898.5 63 540.1 36.5 717.1 39.5 64.7 782.3 5.9 14.8
Inv. steel AD-4 1264 6.4 872.2 100 530.3 104.7 707.5 6.8 47.5 747.7
19.4 12.5 Inv. steel AE-4 1160 5.0 808.3 55 556.7 67.8 655.3 38.3
72.6 735 28.4 45.8 Inv. steel AF-4 1011 3.4 885.6 60 402.3 24 662.5
45.8 17.2 742.7 36.9 55.7 Comp. steel AG-4 1044 0.4 716.6 98 496.5
55.1 684.3 23.5 67.4 756.7 41.2 22.4 Inv. steel AH-4 1286 5.0 696.9
92 444.3 65.3 656.9 44.4 46.8 750.3 28.5 21.8 Inv. steel AI-4 1054
1.7 754.6 90 452.1 55.4 705.2 51.8 14.6 775.6 39.6 68 Inv. steel
AJ-4 1233 5.9 772.2 22 536.2 63.2 675.9 26.3 3 756.9 66 23.7 Comp.
steel AK-4 1010 5.6 940.5 47 554.4 122.4 689.8 30.8 27.1 759.8 27
59 Inv. steel AL-4 1199 6.2 846.4 92 557.1 99.3 706.2 25.6 64 753.3
20.6 99.2 Inv. steel AM-4 1239 3.6 750.5 38 500.7 150 688.9 31.9
44.5 780.5 49.5 34.3 Inv. steel AN-4 1256 5.6 956 74 498 149.2
690.8 8 74.8 744.3 38.8 76.6 Comp. steel AO-4 1241 6.7 750.8 56 454
118.1 657.9 61.5 47.4 747.4 43.6 20 Comp. steel AP-4 1043 4.8 800.3
78 417.4 31.9 707 7.8 0.5 766.1 8.1 2.7 Comp. steel AQ-4 1032 2.2
793.4 43 559.4 64.8 718 13 51.2 748.4 17.1 15.9 Inv. steel
[0246] Table 9 shows the results of measurement and results of
evaluation of the carbide size, pearlite area ratio, ferrite grain
size, Vickers hardness, ratio of the number of carbides at the
ferrite grain boundaries to number of carbides in the ferrite
grains, ratio of maximum crack length to plate thickness at the
vertical wall parts, and impact resistance characteristic in the
prepared samples.
TABLE-US-00009 TABLE 9 No. of carbides at grain Ratio boundaries/
of Ferrite Pearlite No. of maximum Carbide grain area Vickers
carbides crack Impact size size rate hardness inside length
resistance (.mu.m) (.mu.m) (%) (HV) grains (%) characteristic
Remarks AA-2 0.78 23.8 9.1 195.2 13.70 15.6 NG Comparative Steel
AB-2 0.78 26.2 1.2 108.9 3.56 1.6 OK Invention Steel AC-2 0.74 13.0
0.0 120.8 1.35 3.2 OK Comparative Steel AD-2 0.75 23.2 10.3 214.7
2.81 18.8 NG Comparative Steel AE-2 0.17 14.7 0.1 121.4 1.26 5.6 OK
Invention Steel AF-2 0.97 18.0 13.1 217.2 8.98 13.5 NG Comparative
Steel AG-2 0.56 11.5 0.4 122.1 3.35 2.2 NG Comparative Steel AH-2
0.46 21.7 1.4 105.3 1.20 4.7 OK Invention Steel AI-2 0.49 9.4 0.2
132.6 6.21 2.0 OK Invention Steel AJ-2 0.44 20.8 1.0 125.0 1.92 4.7
OK Invention Steel AK-2 0.37 10.7 1.2 127.5 5.48 2.0 OK Invention
Steel AL-2 0.52 12.1 0.9 137.5 5.69 2.3 OK Invention Steel AM-2
0.81 16.5 1.4 116.4 9.20 1.3 OK Invention Steel AN-2 0.75 20.9 0.9
117.0 1.35 5.1 OK Invention Steel AO-2 0.43 10.0 1.5 144.9 1.34 6.9
OK Invention Steel AP-2 0.48 16.4 14.2 228.1 2.16 19.9 NG
Comparative Steel AQ-2 0.65 26.7 0.9 113.1 2.74 2.0 OK Invention
Steel AA-3 0.73 17.0 0.1 106.1 9.74 0.8 OK Invention Steel AB-3
0.48 14.3 0.7 124.1 3.28 2.3 OK Invention Steel AC-3 1.02 12.3 1.2
125.9 11.41 1.4 OK Invention Steel AD-3 0.36 7.5 1.3 147.6 1.18 7.7
OK Invention Steel AE-3 0.65 21.5 0.8 111.5 5.36 1.4 OK Invention
Steel AF-3 0.63 23.3 1.5 110.3 2.88 1.9 OK Invention Steel AG-3
0.57 17.0 0.1 110.2 3.90 1.6 OK Invention Steel AH-3 0.79 14.5 1.5
115.4 4.28 13.1 NG Comparative Steel AI-3 0.75 25.1 0.6 105.8 1.93
2.1 NG Comparative Steel AJ-3 0.21 14.8 1.5 129.8 1.27 6.2 OK
Invention Steel AK-3 0.46 9.0 0.8 133.6 0.66 10.6 NG Comparative
Steel AL-3 0.65 7.4 0.5 153.7 0.14 11.7 NG Comparative Steel AM-3
0.83 21.1 8.2 189.8 5.20 12.7 NG Comparative Steel AN-3 0.51 25.0
1.3 116.2 3.59 1.9 NG Comparative Steel AO-3 0.61 15.6 0.4 131.8
5.45 2.1 OK Invention Steel AP-3 0.57 8.8 1.0 146.8 1.11 6.8 OK
Invention Steel AQ-3 2.17 20.1 1.4 119.5 23.93 21.2 NG Comparative
Steel AA-4 0.66 19.7 1.0 100.3 2.07 1.8 OK Invention Steel AB-4
0.68 20.7 1.1 115.2 10.50 1.1 OK Invention Steel AC-4 1.04 24.4 0.7
107.4 4.35 1.4 OK Invention Steel AD-4 0.6 13.3 0.9 129.2 7.44 1.8
OK Invention Steel AE-4 0.24 12.2 0.5 126.5 1.52 5.4 OK Invention
Steel AF-4 0.36 18.3 0.3 115.8 1.62 2.7 OK Comparative Steel AG-4
0.73 19.3 1.1 107.0 7.13 1.1 OK Invention Steel AH-4 0.52 21.9 1.0
105.2 5.55 1.2 OK Invention Steel AI-4 0.78 24.1 0.7 106.3 1.74 2.2
OK Invention Steel AJ-4 0.44 22.4 8.4 199.8 2.88 13.7 NG
Comparative Steel AK-4 0.45 16.5 0.3 115.9 1.92 4.2 OK Invention
Steel AL-4 0.56 12.6 0.4 176.3 1.74 5.7 OK Invention Steel AM-4
0.77 22.3 0.1 109.1 3.38 1.7 OK Invention Steel AN-4 0.34 20.2 1.3
119.5 2.05 2.6 NG Comparative Steel AO-4 2.18 18.0 1.2 132.1 4.03
14.0 NG Comparative Steel AP-4 1.65 24.7 1.0 123.2 7.68 1.6 OK
Comparative Steel AQ-4 0.52 16.4 1.1 122.2 1.68 4.9 OK Invention
Steel
[0247] In Comparative Steel AC-2, the finish hot rolling
temperature is low and the productivity is low. In Comparative
Steel AN-4, the finish hot rolling temperature is high, scale flaws
form at the surface of the steel plate and cracks form from the
flaw parts when impact load was given after cold forging and
carburizing, quenching, and tempering, and the impact resistance
characteristic falls.
[0248] In Invention Steel AB-3, the cooling rate at the ROT is
slow, so a drop in productivity and formation of scale flaws are
invited. In Invention Steels AJ-3 and AD-4, the cooling rate at the
ROT is 100.degree. C./sec, the outermost layer part of the steel
plate is excessively cooled, and fine cracks are formed at the
outermost layer part.
[0249] In Comparative Steel AN-3, the coiling temperature is low,
large amounts of bainite, martensite, and other low temperature
transformed structures are produced resulting in embrittlement,
fractures frequently occur at the time of pay out of the hot rolled
coil, and the productivity falls. Furthermore, at the sample taken
from the cracked slab, the cold forging and impact resistance
characteristic after carburizing, quenching, and tempering are
inferior.
[0250] In Comparative Steel AH-3, the coiling temperature is high,
bulky pearlite of the lamellar spacing is formed in the hot rolled
structure, needle-shaped coarse carbides are high in thermal
stability, and even after two-stage step type annealing, the above
carbides remain in the steel plate, so the cold forgeability is
low.
[0251] In Comparative Steel AF-4, the heating rate in the first
stage annealing of the two-stage step type annealing is slow, so
the productivity is low. In Comparative Steel AG-2, the heating
rate in the first stage annealing is fast, so the difference in
temperature between the inside part and outer circumferential part
of the coil becomes larger, scratches and seizing due to the
difference in heat expansion occur, and the cold forging and impact
resistance characteristic after carburizing, quenching, and
tempering fall.
[0252] In Comparative Steel AA-2, the holding temperature in the
first stage annealing (annealing temperature) is low, the
coarsening of the carbides at the Ac1 point or less is
insufficient, the thermal stability of the carbides becomes
insufficient, the carbides remaining at the time of the second
stage annealing decrease, the pearlite transformation cannot be
suppressed in the structure after gradual cooling, and the cold
forgeability falls.
[0253] In Comparative Steel AM-3, the first stage holding
temperature (annealing temperature) is high, austenite is produced
during the annealing, the stability of the carbides cannot be
raised, and the cold forgeability and impact resistance
characteristic after carburizing, quenching, and tempering fall. In
Comparative Steel AF-2, the holding time in the first stage
annealing is short, the stability of the carbides cannot be raised,
and the cold forgeability is low. In Comparative Steel AO-4, the
holding time in the first stage annealing is long and the
productivity is low.
[0254] In Comparative Steel AP-4, the heating rate at the second
stage annealing in the two-stage step type annealing is slow, so
the productivity is low. In Comparative Steel AI-3, the heating
rate at the second stage annealing is fast, so the temperature
difference between the inside part and the outer circumferential
part of the coil become greater, scratches and seizing occur due to
the large difference in heat expansion due to transformation, and,
when an impact load is given after carburizing, quenching, and
tempering, fractures occur from the flaw parts and the impact
resistance characteristics fall.
[0255] In Comparative Steel AL-3, the holding temperature in the
second stage annealing (annealing temperature) is low, the amount
of production of austenite is small, it is not possible to increase
the number ratio of carbides at the ferrite grain boundaries, and
the cold forgeability falls. In Comparative Steel AD-2, the holding
temperature in the second stage annealing (annealing temperature)
is high, the dissolution of carbides during annealing is promoted,
and therefore it becomes difficult to cause the production of
carbides at the grain boundaries after gradual cooling, and the
cold forgeability and impact resistance characteristics after
carburizing, quenching, and tempering fall.
[0256] In Comparative Steel AJ-4, the holding time in the second
stage annealing is long and dissolution of carbides is promoted, so
the cold forgeability is low. In Comparative Steel AQ-3, the
cooling rate from the second stage annealing to 650.degree. C. is
slow so the productivity is low and coarse carbides are formed in
the structure after gradual cooling so cracks formed starting from
the coarse carbides at the time of cold forging and the cold
forgeability dropped. In Comparative Steel AP-2, the cooling rate
from the second stage annealing to 650.degree. C. was slow,
pearlite transformation occurred at the time of cooling, and the
hardness increased, so the cold forgeability fell.
[0257] Here, FIG. 3 shows the relationship among the ratio of the
number of carbides at the grain boundaries to the number of
carbides in the grains, and the crack length and impact resistance
characteristics of cold forging test pieces after carburizing,
quenching, and tempering.
[0258] From FIG. 3, it will be understood that if the number ratio
(=number of carbides at the grain boundaries/number of carbides in
grains) exceeds 1, it is possible to keep down the ratio of the
length of cracks introduced by cold forging and possible to obtain
excellent impact resistance after carburizing, quenching, and
tempering.
[0259] Further, FIG. 4 shows another relationship between the ratio
of the number of carbides at the grain boundaries to the number of
carbides in the grains and the crack length of cold forging test
pieces and impact resistance characteristic after carburizing,
quenching, and tempering. FIG. 4 is a view showing that it is
possible to keep down crack length even in steel plate to which
additional elements are added.
[0260] From FIG. 4, it will be understood that even if adding a
suitable range of elements to steel plate, if the number ratio
(=number of carbides at the grain boundaries/number of carbides in
grains) exceeds 1, it is possible to keep down the ratio of the
length of cracks introduced by cold forging and possible to obtain
excellent impact resistance after carburizing, quenching, and
tempering.
INDUSTRIAL APPLICABILITY
[0261] As explained above, according to the present invention, it
is possible to provide low carbon steel plate excellent in cold
forgeability and impact resistance characteristic after
carburizing, quenching, and tempering and a method of production of
the same. The steel plate of the present invention is, for example,
suitable as a material when forming a part by cold forging such as
plate working to obtain a high cycle gear or other part, so the
present invention has high industrial applicability.
REFERENCE SIGNS LIST
[0262] 1. disk-shaped test material [0263] 2. cup-shaped test
material [0264] 3. crack [0265] 4. sample [0266] 5. dropping weight
[0267] L. maximum length of crack
* * * * *