U.S. patent application number 12/085304 was filed with the patent office on 2009-11-12 for steel for warm working, warm working method using the steel, and steel material and steel component obtainable therefrom.
Invention is credited to Tadanobu Inoue, Yuuji Kimura, Kotobu Nagai, Kaneaki Tsuzaki.
Application Number | 20090277539 12/085304 |
Document ID | / |
Family ID | 38048742 |
Filed Date | 2009-11-12 |
United States Patent
Application |
20090277539 |
Kind Code |
A1 |
Kimura; Yuuji ; et
al. |
November 12, 2009 |
Steel for Warm Working, Warm Working Method Using the Steel, and
Steel Material and Steel Component Obtainable Therefrom
Abstract
There are provided a steel for warm working, to be subjected to
warm working as various structures, components of cars, and the
like, a warm working method thereof, and a steel material and a
steel component obtainable from the warm working method. [Solving
Means] A steel is to have a particle dispersion type fiber
structure formed in the matrix by warm working. The steel is
characterized in that the total amount of the dispersed
second-phase particles at room temperature is 7.times.10.sup.-3 or
more in terms of volume fraction, and the Vickers hardness (HV) is
equal to or larger than the hardness H of the following equation
(2): H=(5.2-1.2.times.10.sup.-4.lamda.).times.10.sup.2 (2) when the
steel is subjected to any of annealing, tempering, and aging
treatments in the as-unworked state under conditions such that a
parameter .lamda. expressed by the following equation (1):
.lamda.=T(log t+20)(T; temperature(K), t; time(hr)) (1) is
1.4.times.10.sup.4 or more in a prescribed temperature range of
350.degree. C. or more and Ac1 point or less. This steel is taken
as the steel for warm working.
Inventors: |
Kimura; Yuuji; (Tsukuba-shi,
JP) ; Inoue; Tadanobu; (Tsukuba-shi, JP) ;
Tsuzaki; Kaneaki; (Tsukuba-shi, JP) ; Nagai;
Kotobu; (Tsukuba-shi, JP) |
Correspondence
Address: |
WENDEROTH, LIND & PONACK, L.L.P.
1030 15th Street, N.W.,, Suite 400 East
Washington
DC
20005-1503
US
|
Family ID: |
38048742 |
Appl. No.: |
12/085304 |
Filed: |
November 21, 2006 |
PCT Filed: |
November 21, 2006 |
PCT NO: |
PCT/JP2006/323248 |
371 Date: |
May 21, 2008 |
Current U.S.
Class: |
148/504 ;
148/320 |
Current CPC
Class: |
C21D 9/46 20130101; C21D
8/105 20130101; C21D 9/525 20130101; C21D 9/0075 20130101; C22C
38/04 20130101; C21D 8/0231 20130101; C22C 38/02 20130101; C21D
8/0431 20130101; C21D 8/0205 20130101; C22C 38/18 20130101; C21D
8/065 20130101; C21D 9/0093 20130101 |
Class at
Publication: |
148/504 ;
148/320 |
International
Class: |
C21D 11/00 20060101
C21D011/00; C21D 8/02 20060101 C21D008/02; C21D 8/06 20060101
C21D008/06; C22C 38/00 20060101 C22C038/00 |
Foreign Application Data
Date |
Code |
Application Number |
Nov 21, 2005 |
JP |
2005-336331 |
Claims
1-15. (canceled)
16. A steel material having a particle dispersion fibrous grain
structure, wherein average grain diameter of minor axes of fibrous
ferrite crystals constituting the matrix structure is 3 .mu.m or
less, second-phase particles are finely dispersed in the matrix
structure at a volume fraction of 7.times.10.sup.-3 or more and
12.times.10.sup.-2 or less, and Vickers hardness at room
temperature is HV 3.7.times.10.sup.2 or more.
17. The steel material according to claim 16, wherein the matrix
structure is composed of fibrous ferrite crystals of which average
grain diameter of minor axes is 1 .mu.m or less.
18. The steel material according to claim 16, wherein the matrix
structure is composed of fibrous ferrite crystals of which average
grain diameter of the minor axes is 0.5 .mu.m or less.
19. The steel material according to claim 16, wherein average
particle diameter of major axes of the dispersed second-phase
particles is 0.1 .mu.m or less.
20. A steel for warm-working, from which the steel material
according to claim 16 is made through warm-working, containing
alloy elements which disperse the second-phase particles by
performing one of the heat treatments selected from annealing,
tempering and aging in an as-unworked state under conditions such
that a parameter .lamda. expressed by the following equation (1):
.lamda.=T(log t+20) (T; temperature(K), t; time(hr)) (1) is
1.4.times.10.sup.4 or more in a prescribed temperature range of
350.degree. C. or more and Ac1 point or less, wherein the steel
disperses one or both of second-phase particles formed by the heat
treatment and second-phase particles originally existing before the
heat treatment in total amount by which the dispersed second-phase
particles at room temperature after the heat treatment is
7.times.10.sup.-3 or more and 12.times.10.sup.-2 or less in terms
of volume fraction, and the Vickers hardness (HV) after the heat
treatment is equal to or larger than the hardness H of the
following equation (2) at room temperature:
H=(5.2-1.2.times.10.sup.4.lamda.).times.10.sup.2 (2)
21. The steel according to claim 20, wherein 80% by volume or more
of a matrix structure consists of a single structure of martensite
or bainite, or a mixed structure thereof.
22. The steel according to claim 20, wherein the steel contains C:
0.70 wt % or less, Si: 0.05 wt % or more, Mn: 0.05 wt % or more,
Cr: 0.01 wt % or more, Al: 0.5 wt % or less, O: 0.3 wt % or less,
N: 0.3 wt % or less, and the rest consisting essentially of Fe and
unavoidable impurities.
23. The steel according to claim 22, wherein the steel contains in
place of Fe one or more of elements selected from the group
consisting of Mo: 5.0 wt % or less, W: 5.0 wt % or less, V: 5.0 wt
% or less, Ti: 3.0 wt % or less, Nb: 1.0 wt % or less, and Ta 1.0
wt % or less.
24. The steel according to claim 22, wherein the steel contains in
place of Fe one or both of elements selected from the group
consisting of Ni: 0.05 wt % or more, and Cu: 2.0 wt % or less.
25. A steel plate formed by warm-working the steel according to
claim 20, wherein a particle dispersion fibrous grain structure is
formed in at least a surface layer of the steel plate.
26. A steel wire rod formed by warm-working the steel according to
claim 20, wherein a particle dispersion fibrous grain structure is
formed in at least a surface layer of the steel wire rod.
27. A steel bolt formed by warm-working the steel according to
claim 20, wherein a particle dispersion fibrous grain structure is
formed in at least a surface layer of a screw portion.
28. A method of making the steel material having a particle
dispersion fibrous grain structure, wherein average grain diameter
of minor axes of fibrous ferrite crystals constituting the matrix
structure is 3 .mu.m or less, second-phase particles are finely
dispersed in the matrix structure at a volume fraction of
7.times.10.sup.-3 or more and 12.times.10.sup.-2 or less, and
Vickers hardness at room temperature is HV 3.7.times.10.sup.2 or
more, wherein the steel according to claim 20 is warm-worked in a
prescribed shape under conditions such that a parameter .lamda.
expressed by the following equation (1): .lamda.=T(log t+20) (T;
temperature(K), t; time(hr)) (1) is 1.4.times.10.sup.4 or more in a
prescribed temperature range of 350.degree. C. or more and Ac1
point or less.
29. A method of making a steel plate from a steel material having a
particle dispersion fibrous grain structure, wherein average grain
diameter of minor axes of fibrous ferrite crystals constituting the
matrix structure is 3 .mu.m or less, second-phase particles are
finely dispersed in the matrix structure at a volume fraction of
7.times.10.sup.-3 or more and 12.times.10.sup.-2 or less, and
Vickers hardness at room temperature is HV 3.7.times.10.sup.2 or
more, wherein the steel according to claim 20 is warm worked in a
prescribed shape under conditions such that a parameter .lamda.
expressed by the following equation (1): .lamda.=T(log t+20) (T;
temperature(K), t; time(hr)) (1) is 1.4.times.10.sup.4 or more in a
prescribed temperature range of 350.degree. C. or more and Ac1
point or less, and wherein a particle dispersion fibrous grain
structure is formed in at least a surface layer of the steel
plate.
30. A method of making a steel wire rod from a steel material
having a particle dispersion fibrous grain structure, wherein
average grain diameter of minor axes of fibrous ferrite crystals
constituting the matrix structure is 3 .mu.m or less, second-phase
particles are finely dispersed in the matrix structure at a volume
fraction of 7.times.10.sup.-3 or more and 12.times.10.sup.-2 or
less, and Vickers hardness at room temperature is HV
3.7.times.10.sup.2 or more, wherein the steel according to claim 20
is warm worked in a prescribed shape under conditions such that a
parameter .lamda. expressed by the following equation (1):
.lamda.=T(log t+20) (T; temperature(K), t; time(hr)) (1) is
1.4.times.10.sup.4 or more in a prescribed temperature range of
350.degree. C. or more and Ac1 point or less and wherein a particle
dispersion fibrous grain structure is formed in at least a surface
layer of the steel wire rod.
31. A method of making a steel bolt from the steel material having
a particle dispersion fibrous grain structure, wherein average
grain diameter of minor axes of fibrous ferrite crystals
constituting the matrix structure is 3 .mu.m or less, second-phase
particles are finely dispersed in the matrix structure at a volume
fraction of 7.times.10.sup.-3 or more and 12.times.10.sup.-2 or
less, and Vickers hardness at room temperature is HV
3.7.times.10.sup.2 or more, wherein the steel according to claim 20
is warm worked in a prescribed shape under conditions such that a
parameter .lamda. expressed by the following equation (1):
.lamda.=T(log t+20) (T; temperature(K), t; time(hr)) (1) is
1.4.times.10.sup.4 or more in a prescribed temperature range of
350.degree. C. or more and Ac1 point or less and wherein a particle
dispersion of fibrous grain structure is formed in at least a
surface layer of a screw portion.
32. A steel product produced by cut-machining the steel material
according to claim 16.
33. A steel product produced by cut-machining the steel plate
according to claim 25.
34. A steel product produced by cut-machining the steel wire rod
according to claim 26.
35. A steel product produced by cut-machining the steel bolt
according to claim 27.
Description
TECHNICAL FIELD
[0001] The present invention relates to a steel to be worked into
various structures, components of cars, and the like for use. More
particularly, it relates to a steel for warm working to be
subjected to warm working, and a warm working method thereof, and a
steel material and a steel component obtainable from the warm
working method.
BACKGROUND ART
[0002] In recent years, there has been a demand for a tougher and
higher performance high strength steels than ever with an increase
in size of structures, and a reduction in weight of car components
or the like. As the means for improving the toughness of the steel,
conventionally, there are generally known: (1) reduction of
impurity elements such as P and S causing embrittlement, (2)
refinement and reduction of inclusions, (3) addition of alloy
elements, (4) reduction of carbon, (5) crystal grain refinement,
(6) refinement of dispersed second-phase particles such as carbide
particles, and the like.
[0003] Out of these, crystal grain refinement receives attention
because of the following: it has both the effects of reduction of
concentration of stress to the grain boundary and dilution at the
grain boundary of impurity elements, and it can raise the
brittleness fracture stress simultaneously with an increase in
yield stress. For example, in recent years, an attempt has been
made to reduce the ferrite grain size as ultrafine as 1 .mu.m or
less in low carbon steel allowing for conservation of natural
resources or recyclability, and thereby to achieve strengthening
and the longer life of the steel.
[0004] However, the studies on the crystal grain refinement of a
low carbon ferrite steel up to this point concentrate on the
strength level of 1000 MPa or less (e.g., Non-Patent Document 1-2;
Patent Document 1-2). This is due to the following. In order to
obtain a high strength of 1000 MPa or more only by refinement of
ferrite grains, the crystal grains are required to be reduced in
size to as ultrafine as 0.5 .mu.m or less. Thus, the reduction in
grain size to as ultrafine as 0.5 .mu.m or less is very difficult
with a thermo-mechanical treatment of a steel intended for mass
production. Further, on a laboratory scale, ultrafine grains of 0.5
.mu.m or less can be obtained with a super heavy deformation method
such as MM (Non-Patent Document 3) or ARB (Non-Patent Document 4)
of powder metallurgy. However, such an ultrafine grain steel shows
almost no uniform elongation, and the elongation is mostly caused
by ununiform deformation due to necking, resulting in a large
reduction in ductility. The same early plasticity instability as
this is also observed in a pure iron wire dislocation strengthened
by wire drawing (Non-Patent Document 5).
[0005] In general, strengthening of a steel unfavorably largely
reduces various characteristics such as ductility, toughness,
delayed fracture resistance characteristics, fatigue
characteristics, formability. Particularly, when a very versatile
low alloy martensitic steel is strengthened to 1.2 GPa or more, it
is remarkably reduced in toughness, delayed fracture resistance
characteristics, and the like. For this reason, the high strength
steel is largely prevented from being put into practical use. Under
such circumstances, there is a strong demand for simultaneously
achieving higher strength and higher toughness, and the improvement
of the delayed fracture resistance characteristics of a low alloy
steel. However, according to conventional findings, as the means
for improving the fracture characteristics of the low alloy
martensitic steel, the following are conceivable: (a) high
temperature tempering avoiding the tempering embrittlement
temperature range in the vicinity of 500.degree. C., (b)
prior-austenite grain refinement, (c) ausforming, and (d) formation
of fibrous structure, or combinations thereof. However, the
application of these means faces the following problems.
[0006] (a) High Temperature Tempering
[0007] High temperature tempering is carried out at about
550.degree. C. or more, and A1 point or less. According to this,
there are advantages as follows: (1) the internal stress introduced
upon quenching can be largely reduced accompanied by recovery of
dislocation; (2) the coherent precipitate (e.g., film-like
cementite) reducing the fracture toughness can be made incoherent
(be spheroidized), and other advantages. For this reason, for a
steel for a mechanical structure particularly requiring toughness,
tempering is generally carried out in the vicinity of 650.degree.
C. However, within such a temperature range, the dispersed
second-phase particles also grow with ease during tempering, and
hence the reduction in strength of the steel is inevitable.
Further, in the related art, there is adopted a method in which
large quantity of carbon is added to increase the amount of carbide
precipitated, thereby to increase the strength. However, the
toughness is reduced. Therefore, strengthening only by high
temperature tempering has its limitation. The ones of which
strengthening can be achieved even by high temperature tempering
are limited to the steels in which a large amount of special alloy
elements have been added such as maraging steels (Non-Patent
Documents 6-10).
[0008] (b) Crystal Grain Refinement
[0009] For strengthening of the steel, it is indispensable to
refine prior-austenite grains so as to ensure enough toughness. As
the method of austenite grain refinement, there are (1) a method by
recrystallization of processed austenite, and (2) a method using
phase transformation. Out of these, the thermo-mechanical treatment
for performing an austenitizing treatment after processing a
martensite structure within the cold or warm range, classified into
the latter, was thought to be capable of refining austenite grains
most effectively (Non-Patent Documents 11 and 12). For example, it
is known (Patent Document 3) that refinement of austenite grains of
several micrometers or less improves the toughness of the tempered
martensitic steel (Non-Patent Document 13), and improves the
delayed fracture characteristics (Patent Document 4). With
refinement of austenite, as the crystal grains become finer, the
grain growth rate also increases. For this reason, it is a
particularly important point how the grain growth is controlled in
austenite. Under such circumstances, in the related art, the
following are generally applied: dispersion of pinning grains
effective for suppressing the growth of austenite, reduction of the
austenitization temperature, austenitization for a rapid short time
using high-frequency heating, and the like. However, it is very
difficult to suppress the growth of ultrafine austenite grains. In
actuality, grain refinement reaches a limit at about several
micrometers. Whereas, excessive refinement of crystal grains
promotes the diffusion type phase transformation at the grain
boundary, which unfavorably makes hardening difficult, and also
causes other problems. Thus, the process window for austenite grain
refinement is relatively narrow.
[0010] (c) Ausforming
[0011] Ausforming is a treatment in which an austenitized steel is
quenched to a metastable austenite range, and processed at the
temperature, followed by quench hardening, thereby to cause
martensite or bainite transformation, and then, tempering is
carried out. It has a feature of being capable of strengthening the
steel without much impairing the toughness. With the ausforming, it
is considered as follows. The effects such as (1) refinement of
packets or blocks regarded as effective crystal grains, (2)
succession of dislocation from worked austenite to martensite, and
(3) pinning of dislocation by carbon atoms or carbides occur in an
overlapping manner, thereby to strengthen the steel. In recent
years, improved ausforming in which working is carried out within
the high-temperature metastable austenite range has been applied to
a medium carbon low alloy steel. Thus, improvements of the fatigue
and delayed fracture characteristics have been reported. Further,
the main factors of the improvements of characteristics by improved
ausforming are considered to be refinement of the matrix structure,
suppression of formation of the coarse grain boundary cementite due
to introduction of the grain boundary unevenness (Non-Patent
Document 14) or formation of the texture (Non-Patent Document 15).
However, ausforming is working of the austenite structure.
Therefore, it requires that the alloy components and the
thermo-mechanical treatment conditions should be strictly adjusted
so as to prevent the metastable austenite phase from undergoing
proeutectoid ferrite transformation or pearlite transformation
during working. Further, there is also another problem that
quenching crack is caused during cooling after working.
Accordingly, the applicable members are also limited to those in
simple shapes such as a plate and a rod.
[0012] (d) Fibrous Structure Formation
[0013] For enhancing the toughness of a steel, it is also effective
to effect the formation of a fiber structure in the inside by cold
or warm working. This has already been proposed for the steel
subjected to an ausforming treatment (Non-Patent Documents 16 and
17), the heavy cold drawing high-strength low carbon wire rod
(Patent Document 5), piano wire, pure iron wire (Non-Patent
Document 5), and the like.
[0014] For working of steel materials, cold working is the main
process of member formation today because it can mass-produce
members in complicated shapes such as bolts with high dimensional
accuracy. However, for such steel materials as those having a
tensile strength of more than 1.2 GPa, cold forging thereof is very
difficult because of the strength. For this reason, the members in
which a fiber structure is formed by the cold molding process as
described above are limited to wire rods and the like.
[0015] On the other hand, many attempts have also been made until
now on warm working in a two phase range of a ferrite phase and a
carbide at Ac1 point or less. For example, there are known a method
and a forming method in which a material for a high strength member
and a high strength steel material is prepared; the material is
warm worked so as to form a member in a desired geometrical shape
in such a state as to substantially hold or enhance the strength
characteristics of the material; as a result, a fiber structure is
formed, thereby to make a high strength steel structure member with
a tensile strength of at least 1 GPa (Patent Document 6).
[0016] Further, there is also known a manufacturing method of a
formed product characterized by the following: a material having an
ultrafine structure is warm worked or cold worked, and a steel
material including drawn ferrite grains with a minor axis of 3
.mu.m or less is used as a material; it is not subjected to a
refining treatment, and only forming is carried out, and a refining
treatment is not carried out (Patent Document 3).
[0017] Whereas, for a steel having a duplex phase structure such as
a tempered martensite structure, warm working is applied for the
purpose of obtaining a worked structure before a quench hardening
treatment for refining a reverse-transformed austenite (Non-Patent
Documents 11 and 12). The strength of the steel is achieved by a
refining treatment following warm working. For this reason, an
attempt has not been made to use the material in the as-warm worked
state of the tempered martensite structure.
[0018] Further, for warm straightening working of a high carbon
steel having a carbon content of 0.7 wt % or more, an over 1.8 GPa
class wire rod can be obtained. However, the elongation of the wire
rod is as low as around 6% (Non-Patent Document 18).
[0019] Non-Patent Document 1: TETSU TO HAGANE, 85 (1999), P.
620
[0020] Non-Patent Document 2: ISIJ International, 44 (2004), P.
1063
[0021] Non-Patent Document 3: SOSEI TO KAKOU (Journal of the Japan
Society for technology of plasticity), 41 (2000), P. 13
[0022] Non-Patent Document 4: TETSU TO HAGANE, 88 (2002), P.
359
[0023] Non-Patent Document 5: ASM, 62 (1969), P. 623
[0024] Non-Patent Document 6: Trans. ASM, 61 (1968), P. 798
[0025] Non-Patent Document 7: Metal. Trans., 1 (1970), P. 2011
[0026] Non-Patent Document 8: Mat. Sci. Tech., 19 (2003), P.
117
[0027] Non-Patent Document 9: Mat. Sci. Tech., 7 (1991), P.
1082
[0028] Non-Patent Document 10: Mat. Sci. Eng., A398 (2005), P.
367
[0029] Non-Patent Document 11: TEKKOU NO KESSYOURYU CYOUBISAIKA
BUKAI Report (The Iron and Steel Institute of Japan), (1991), P.
64
[0030] Non-Patent Document 12: Proc. First International Conference
on Advanced Structural Steels, (2002), P. 65
[0031] Non-Patent Document 13: Ultrafine-Grain Metals, Proc. the
16th sagamore Army Materials Conference, (1969), P. 138
[0032] Non-Patent Document 14: CAMP-ISIJ, 12 (1999), P. 565
[0033] Non-Patent Document 15: CAMP-ISIJ, 12 (1999), P.
1045-1048
[0034] Non-Patent Document 16: ASM, 55 (1962), P 654
[0035] Non-Patent Document 17: Met. Trans., 1 (1970), P 3037
[0036] Non-Patent Document 18: J. Japan Inst. Metals, 32 (1968), P.
289
[0037] Patent Document 1: JP-A-2004-285437
[0038] Patent Document 2: JP-A-2005-194547
[0039] Patent Document 3: JP-A-2004-60046
[0040] Patent Document 4: JP-A-11-80903
[0041] Patent Document 5: JP-B-6-53915
[0042] Patent Document 6: U.S. Pat. No. 5,236,520
DISCLOSURE OF THE INVENTION
Problems that the Invention is to Solve
[0043] As described above, prior-austenite grain refinement and
ausforming are important toughness-enhancing technology of a steel,
and studies and inventions thereon add up to massive amounts.
However, in these processes, quench hardening and tempering are
basic, so that strengthening receives restrictions by the problems
of hardenability and quenching crack, and the problem of temper
brittleness. Further, with an increase in strength, the amount of
the dispersed second-phase particles of carbides or the like
necessary for strengthening also increases, which makes softening
by spheroidizing or the like difficult. Whereas, when particularly
carbides become coarse by annealing, unfavorably, cracks occur in
the inside of the material in the process of forming the material
into a component by cold forging or the like, or other problems
occur. For these reasons, as far as the conventional strengthening
process by quench hardening and tempering, it is considered
impossible to largely improve the characteristics of such an over
1.5 GPa class high strength steel as to result in a tensile
strength of 1.2 GPa or more, and further, as to make softening
difficult, and to put the steel into practical use.
[0044] Further, the studies and inventions regarding warm working
up to this point mainly aim at forming into a member and
manufacturing of the prior structure. For this reason, in most
cases, a relatively soft matrix structure such as a ferrite or
pearlite structure with a low deformation resistance, or a
martensite structure subjected to tempering at high temperatures is
used as a starting material. Thus, under such conditions as to
result in reduction of the deformation resistance, warm working is
carried out. Further, the fine duplex structure is not formed in
consideration of the dispersion state of the dispersed second-phase
particles and the thermal stability. Thus, there has not yet been
implemented a high strength member with a tensile strength of 1.2
GPa or more after warm working, and excellent in ductility,
toughness, delayed fracture characteristics, and the like.
Particularly for the duplex phase structure steel such as a
tempered martensitic steel having a tensile strength of 1.2 GPa or
more at room temperature, there is a possibility that warm working
cannot be carried out because of its strength. For this reason,
application of warm working thereto has been regarded almost
impossible conventionally.
[0045] Under such circumstances, it is an object of the present
invention to provide a steel for warm working which resolves the
problems as described above, and which can form a particle
dispersion type fibrous structure for obtaining a high strength
steel having a tensile strength of 1.2 GPa or more, being excellent
in ductility and delayed fracture resistance characteristics, and
has been tremendously improved in toughness by warm working, and a
warm working method using the same. Further, it is another object
to provide steel materials such as a steel plate and a rod steel,
and steel components such as a bolt and cut machined products
having the foregoing characteristics obtained therefrom.
Means for Solving the Problems
[0046] The present inventors conducted a close study in order to
resolve the foregoing problems. As a result, they made the
following inventions.
[0047] First: a steel for warm working, being to have a particle
dispersion type fiber structure formed in the matrix by warm
working, the steel characterized by including an alloy element
or/and dispersed second-phase particles such that the total amount
of the dispersed second-phase particles at room temperature is
7.times.10.sup.-3 or more in terms of volume fraction, and the
steel characterized by having a Vickers hardness (HV) of equal to
or larger than the hardness H of the following equation (2):
H=(5.2-1.2.times.10.sup.-4.lamda.).times.10.sup.2 (2)
when the steel is subjected to any heat treatment of annealing,
tempering, and aging treatments in the as-unworked state under
conditions such that a parameter .lamda. expressed by the following
equation (1):
.lamda.=T(log t+20) (T; temperature(K), t; time(hr)) (1)
is 1.4.times.10.sup.4 or more in a prescribed temperature range of
350.degree. C. or more and Ac1 point or less.
[0048] Second: the steel for warm working, characterized in that
80% by volume or more of the matrix structure is any single
structure of martensite and bainite, or a mixed structure
thereof.
[0049] Third: the steel for warm working, characterized by
including, in chemical composition, C, 0.70 wt % or less, Si: 0.05
wt % or more, Mn: 0.05 wt % or more, Cr: 0.01 wt % or more, Al: 0.5
wt % or less, O: 0.3 wt % or less, and N: 0.3 wt % or less, and the
balance being substantially Fe and inevitable impurities, in the
claim.
[0050] Fourth: the steel for warm working, characterized by further
including, one or two or more selected from a group consisting of
Mo: 5.0 wt % or less, W: 5.0 wt % or less, V: 5.0 wt % or less, Ti:
3.0 wt % or less, Nb: 1.0 wt % or less, and Ta: 1.0 wt % or
less.
[0051] Fifth: the steel for warm working, characterized by further
including one or two of Ni: 0.05 wt % or more and Cu: 2.0 wt % or
less.
[0052] Sixth: a warm working method characterized by performing
warm working for imparting a stain of 0.7 or more in a temperature
range of 350.degree. C. or more and Ac1 point -20.degree. C. or
less on any of the foregoing steels for warm working.
[0053] Seventh: the warm working method, characterized in that
after performing warm working, an aging treatment is performed in a
temperature range of 350.degree. C. or more and Ac1 point or
less.
[0054] Eighth: a steel material which is a steel having a particle
dispersion type fiber structure obtained by warm working any of the
foregoing steels for warm working, characterized in that the
average grain diameter of the minor axes of a fibrous ferrite
crystal forming the matrix structure is 3 .mu.m or less, the
second-phase particles are finely dispersed in the matrix structure
at a volume fraction of 7.times.10.sup.-3 or more, and the Vickers
hardness at room temperature is HV 3.7.times.10.sup.2 or more.
[0055] Ninth: the steel material, characterized by having a matrix
structure including a fibrous crystal of which the average grain
diameter of the minor axes is 1 .mu.m or less.
[0056] Tenth: the steel material being any of the foregoing steel
materials, characterized by having a matrix structure including a
fibrous crystal of which the average grain diameter of the minor
axes is 0.5 .mu.m or less.
[0057] Eleventh: the steel material being any of the foregoing
steel materials, characterized in that the average particle
diameter of the major axes of the dispersed second-phase particles
is 0.1 .mu.m or less.
[0058] Twelfth: a steel plate obtained by warm working any of the
foregoing steels for warm working into a plate form, characterized
by including a fiber structure formed in at least the surface layer
part thereof.
[0059] Thirteenth: a wire rod steel obtained by warm working any of
the foregoing steels for warm working into a bar form or a wire
form, characterized by including a fiber structure formed in at
least the surface layer part thereof.
[0060] Fourteenth: a bolt obtained by warm working any of the
foregoing steels for warm working, characterized by including a
fiber structure formed in at least the surface layer part of the
screw part.
[0061] Fifteenth: a steel component characterized by being the one
obtained by working the steel material according to any of the
foregoing items into a component by cutting.
ADVANTAGE OF THE INVENTION
[0062] As for the steel for warm working of the first invention,
the softening resistance when the steel is heated, namely, the
thermal stability and the total amount of the matrix structure and
the dispersed second-phase particles are controlled. As a result, a
particle dispersion type fibrous structure can be formed when the
steel is subjected to warm working, and the Vickers hardness after
warm working can be set at 3.7.times.10.sup.2 or more. As a result,
there is provided a steel for warm working capable of being
tremendously improved in toughness while keeping the tensile
strength of 1.2 GPa or more at ordinary temperatures.
[0063] In accordance with the second invention, the structure of
the steel for warm working as the prior working structure is
transformed into an ultrafine duplex structure including dispersed
second-phase particles such as carbide particles finely dispersed
therein by using martensite transformation or bainite
transformation. As a result, it becomes possible to effect the
formation of a fiber structure even in the inside with efficiency
when the steel is subjected to warm working. In addition to this,
it becomes possible to largely improve the delayed fracture
resistance characteristics.
[0064] In accordance with the third invention, the alloy
composition excellent in cost efficiency and recyclability can
achieve strengthening of the steel obtained when the steel is
subjected to warm working.
[0065] In accordance with the fourth invention, it is possible to
disperse dispersed second-phase particles which are finer and
excellent in hydrogen trapping property. Further, it is possible to
strengthen the steel material obtained when the steel is subjected
to warm working, and to largely enhance the toughness in a low
temperature range, and the delayed fracture resistance
characteristics.
[0066] In accordance with the fifth invention, it is possible to
improve the toughness further to a lower temperature range.
[0067] In accordance with the sixth invention, while working the
steel for warm working into a desired shape, a fiber structure can
be formed to obtain a high toughness. Incidentally, as the
equipment, warm working equipment which has been conventionally put
into practical use can be utilized. Therefore, the invention has a
very high practical utility.
[0068] In accordance with the seventh invention, by performing an
aging treatment with the fiber structure finely held, it is
possible to manufacture a steel showing less variations in
characteristics than the sixth invention.
[0069] In accordance with the eighth invention, there is
implemented a steel material which not only has high toughness, but
also has also been improved in secondary workability by formation
of the fine fiber structure.
[0070] In accordance with the ninth invention, a dense fiber
structure with an average spacing of the minor axes of 1 .mu.m or
less is developed, and in accordance with the tenth invention, a
dense fiber structure with an average spacing of the minor axes of
0.5 .mu.m or less is developed. Thus, there are implemented steel
materials which have been much more enhanced in strength,
toughness, and workability than before warm working.
[0071] In accordance with the eleventh invention, by controlling
the average particle diameter of the major axes of the dispersed
second-phase particles to 0.1 .mu.m or less, it is possible to
implement much more strengthening and enhancement of the toughness
with dispersion of a small amount of the dispersed second-phase
particles.
[0072] In accordance with the twelfth and thirteenth inventions,
the steel material not only has a high toughness and tensile
strength, but also has a secondary workability. For this reason,
there are implemented a steel plate and a steel rod wire which have
been tremendously enhanced in practical utility, usable for
manufacturing various components and products.
[0073] In accordance with the fourteenth invention, there is
implemented a bolt excellent in impact strength and delayed
fracture resistance, in which a fiber structure is formed in the
root of the thread of the screw part to which a stress particularly
concentrates.
[0074] In accordance with the fifteenth invention, even a high
strength component in a complicated shape is provided as the one
excellent in impact strength and delayed fracture resistance.
[0075] As described above, there is provided a high strength steel
multiphased by fine dispersion of a small amount of dispersed
second-phase particles. Particularly, even an ultrahigh strength
steel which is hard to soften and is hard to form is applied with
prescribed deformation in a temperature range in which the
deformation resistance is reduced and no cracks occur in the
material, to be formed into a prescribed shape (thin plate, thick
plate, wire rod, or component). As a result, conventional
spheroidizing and quench hardening and tempering treatments after
component forming are omitted. At the same time, an ultrafine
duplex phase structure is developed into a fibrous form. Thus,
there is provided a high strength steel largely improved in the
ductility, particularly the toughness, and the delayed fracture
resistance characteristics in the relation of trade-off balance
with high strength, and a member thereof.
[0076] The foregoing effects are due to the following
mechanisms:
[0077] (a) CUltra Grain Refinement of Crystal and Formation of
Fibrous Matrix Structure by Warm Working
[0078] The following finding has been reached. A material
satisfying given specific conditions can form a particle dispersion
type fiber structure far more excellent in toughness and delayed
fracture resistance characteristics even than conventional
ausformed steels in the member. Namely, the pinning effect due to
fine dispersion or precipitation of the second-phase particles is
effectively used. Thus, in a temperature range in which recovery of
the dislocation introduced by deformation appropriately occurs, but
primary recrystallization or remarkable grain growth does not
occur, the material is deformed, and applied with prescribed
strain, thereby to refine crystal grains. As a result, it is
possible to form an ultrafine grain duplex phase structure which is
low in internal stress and has no starting point for occurrence of
cracks. Particularly, in such ultrafine grains, a fiber structure
having a further narrower crystal grain boundary spacing is
developed. As a result, it is possible to suppress not only the
occurrence of cracks but also the propagation of cracks, and
thereby to largely enhance the fracture toughness.
[0079] (b) Refinement of Coarse Second-Phase Particles
[0080] Even such coarse dispersed second-phase particles as to
cause occurrence of cracks with cold working can be deformed
relatively easily without occurrence of cracks with warm working.
Thus, by utilizing decomposition and reprecipitation of the
dispersed second-phase particles formed particularly during
working, coarse film-like precipitates considered as the cause of
grain boundary cracking can be not only spheroidized but also
finely dispersed to be utilized for strengthening.
[0081] (c) Dispersion of Ultrafine Alloy Carbides, Intermetallic
Compounds, and the Like
[0082] Alloy elements high in carbide forming ability such as Mo,
V, W, Ta, Ti, and Nb form nano-size alloy carbides such as Mo2C,
V4C3, W2C, TaC, NbC, and TiC in a temperature range in the vicinity
of 500.degree. C. to 600.degree. C. independently from already
existing cementite. For this reason, addition of these alloy
elements is effective for strengthening of the steel. The maximum
value of precipitation strengthening due to these nano-size alloy
carbides is obtained in the transition range of the strengthening
mechanism of from Cutting to Orowan mechanism. However, at such an
aging stage, much coherent strain occurs around the precipitates,
so that the toughness of the steel is reduced. For this reason, in
general, the steel is tempered to the sufficiently overaged state
of these carbides even somewhat sacrificing the strength of the
steel. On the other hand, when the dynamic precipitation of these
alloy carbides due to warm working is utilized, it is also possible
to effect incoherent precipitation of the carbides without causing
much growth of the carbides even in the precipitation transition
temperature range. Namely, it is also possible to make the maximum
use of precipitation strengthening of the alloy carbides due to the
Orowan mechanism. Further, the same effects can also be expected
for precipitation of intermetallic compounds including the alloy
elements, and Ni, Al, or the like, nitrides, oxides, Cu particles,
and the like.
BRIEF DESCRIPTION OF DRAWINGS
[0083] FIG. 1 is a view showing one example of a thermo-mechanical
treatment pattern.
[0084] FIG. 2 is a view showing one example of the
thermo-mechanical treatment pattern.
[0085] FIG. 3 is a view showing one example of the
thermo-mechanical treatment pattern.
[0086] FIG. 4 is a view showing the relationship of the temper
hardness and T(log t+20)=.lamda., where T denotes the temper
temperature (K), and t denotes the temper time(hr).
[0087] FIG. 5 is a view showing a 500.degree. C. warm worked
structure (ultrafine fiber structure).
[0088] FIG. 6 is a view showing the relationship between the
tensile strength and the impact value (U notch).
[0089] FIG. 7 is a view showing the relationship between the
tensile strength and the absorption energy (V notch).
[0090] FIG. 8 is a view showing the relationship between the
absorption energy and the test temperature.
[0091] FIG. 9 is a photographic view showing one example of the
fracture form of a B steel subjected to a Charpy impact test (U
notch).
[0092] FIG. 10 is a view showing the relationship between the
hardness of the warm worked material and the aging temperature.
[0093] FIG. 11 is a view showing an ultrafine fiber structure
formed in the center part of a plate material.
[0094] FIG. 12 is a view showing an ultrafine fiber structure
formed in the surface layer part of a rod material.
BEST MODE FOR CARRYING OUT THE INVENTION
[0095] The present invention has the features as described above.
However, below, the requirements of the invention and the like will
be described in details.
[0096] A steel for warm working of the present invention is a steel
which is to have a particle dispersion type fiber structure formed
in the matrix by warm working, the steel characterized by including
an alloy element or/and dispersed second-phase particles such that
the total amount of the dispersed second-phase particles at room
temperature is 7.times.10.sup.-3 or more in terms of volume
fraction, and the steel characterized by having a Vickers hardness
(HV) of equal to or larger than the hardness H of the following
equation (2):
H=(5.2-1.2.times.10.sup.4.lamda.).times.10.sup.2 (2)
when the steel is subjected to any heat treatment of annealing,
tempering, and aging treatments in the as-unworked state under
conditions such that a parameter .lamda. expressed by the following
equation (1):
.lamda.=T(log t+20) (T; temperature(K), t; time(hr)) (1)
is 1.4.times.10.sup.4 or more, and preferably 1.5.times.10.sup.4 or
more in a prescribed temperature range of 350.degree. C. or more
and Ac1 point or less. Thus, the steel for warm working of the
invention changes in the dispersion state of the dispersed
second-phase particles and the matrix structure during warm working
to be performed thereon. Therefore, it is configured such that the
lower limit of the equation (2) is set with respect to the hardness
(structure) of the non-worked material obtained from the heat
treatment simulating the heat history of warm working. Namely, as
described below, the structure state is expressed by the
hardness.
(a) Structure of Steel for Warm Working
[0097] In order to simultaneously implement strengthening and
enhancement of the toughness of a duplex phase structure steel by
warm working, it is important that strengthening by dispersion of
the dispersed second-phase particles in as small amount as possible
and as fine as possible, and refinement of the matrix structure and
the formation of the fiber structure can be simultaneously carried
out. Then, in order to implement the formation of the ultrafine
duplex phase structure, fine dispersion or fine dispersive power of
the dispersed second-phase particles in the steel for warm working
which is a material is important.
[0098] In the present invention, for the fine dispersion or fine
dispersive power of the second-phase particles, the following three
patterns can be considered:
(i) In the steel for warm working, the second-phase particles have
already been dispersed; (ii) In the steel for warm working, the
second-phase particles are not dispersed, but during warm working,
one type or two or more types of second-phase particles
precipitate, and after a working treatment, a particle dispersion
type fiber structure is formed; and (iii) In the steel for warm
working, the second-phase particles have already been dispersed,
but during warm working, other particles than these
precipitate.
[0099] Then, dispersion (precipitation) strengthening by the
dispersed second-phase particles depends upon the dispersion
conditions such as the volume fraction of the secondary phase
duplex particles, and the size, hardness, and shape of the
particles. When dispersion strengthening is caused by the Orowan
mechanism, from the following equation (A) (TEKKOU NO
SEKISYUTSUSEIGYO METARAJII SAIZENNSENN, (the Iron and Steel
Institute of Japan) (2001) P. 69), the amount of dispersion
strengthening increases with a decrease in particle diameter (d)
and with an increase in volume fraction (f). Namely, the dispersion
conditions (and the dispersive power) of the dispersed second-phase
particles have close relation with the hardness.
.DELTA..sigma.=(3.2 Gb)/[(0.9 f.sup.-1/2-0.8)d] (A)
[0100] where G denotes the shear modulus of the steel, 80 GPa, and
b denotes the Burgers vector, 0.25 nm.
[0101] However, when the particles is excessively reduced in
diameter than a critical particle diameter, the dislocation ceases
to be pinned by the particles. Thus, the particles come to be
sheared by dislocation, so that the Orowan mechanism ceases to
hold. In the so-called Cutting mechanism, in which particles are
sheared by dislocation, the amount of dispersion strengthening
increases with an increase in particle diameter. Namely, the
minimum particle diameter with which the Orowan mechanism holds can
provide the maximum amount of dispersion strengthening. The minimum
particle diameter capable of achieving the maximum dispersion
strengthening depends upon the hardness of the particles, and
decreases in inverse proportion to the hardness of the particles
(TEKKOU NO SEKISYUTSUSEIGYO METARAJII SAIZENNSENN, (the Iron and
Steel Institute of Japan) (2001) P. 69). Therefore, when comparison
is made in terms of the same volume fraction, the minimum particle
diameter with which the Orowan mechanism holds decreases with an
increase in hardness of the particles. Accordingly, the maximum
amount of particle dispersion strengthening also increases.
[0102] For example, it is known that TiC is capable of carrying out
effective dispersion particle strengthening because of its higher
hardness and smaller density among alloy carbides. Now, assuming
that TiC can provide a minimum particle diameter to which the
Orowan mechanism is applicable of 7 nm, a particle dispersion
strengthening amount of about 0.9 GPa (TS (GPa) is nearly equal to
0.0032 HV, HV 2.8.times.10.sup.2) is expectable in dispersion at a
volume fraction of 7.times.10.sup.-3. Incidentally, with a density
of TiC of 4.94 Mg/m.sup.3, an atomic weight of Ti of 47.9, and an
atomic weight of C of 12, Ti and C necessary for precipitating TiC
at a volume fraction of 7.times.10.sup.-3 are in an amount of 0.35
wt % and 0.087 wt %, respectively. In addition, the strength of the
matrix of the practical ferrite steel is about 0.3 GPa (about
0.9.times.10.sup.2 in HV). Therefore, the room temperature strength
of the steel including the TiC dispersed in the ferrite matrix is
expected to be 1.2 GPa or more (HV 3.7.times.10.sup.2 or more).
Accordingly, considering the ideal dispersion conditions for TiC,
with the dispersed particles to which the Orowan mechanism is
applicable, a size of 7 nm or more can sufficiently satisfy an HV
of 3.7.times.10.sup.2 only by dispersion strengthening at a volume
fraction as small as 7.times.10.sup.-3. The same effects as with
this are also expectable for the dispersed second-phase particles
including a carbonitride, an intermetallic compound, an oxide, Cu
particles, and the like. Then, as such dispersed second-phase
particles, for example, specifically, there can be considered
carbides such as Mo2C, V4C3, W2C, TaC, NbC, and TiC, oxides such as
Fe3O4, Fe2O3, Al2O3, Cr2O3, SiO2, and Ti2O3, nitrides such as AlN,
CrN, and TiN, intermetallic compounds such as Ni3Ti, NiAl, TiB,
Fe2Mo, Ni3Nb, and Ni3Mo, metal particles such as Cu particles, and
the like.
[0103] Incidentally, it is known that metal carbide particles of
Mo, Ti, or the like generally have a size of around 10 nm, and can
effectively effect strengthening even by dispersion in an amount as
small as less than 10.times.10.sup.-3 in volume fraction. However,
the size of the dispersed second-phase particles and the
distribution in the matrix structure also varies according to
segregation of alloy elements or the like, or other factors.
Therefore, in the invention, considering such that a fine crystal
structure can be obtained with stability by warm working even when
there are variations in distribution of the dispersed second-phase
particles, the volume fraction at room temperature of the dispersed
second-phase particles is specified at 7.times.10.sup.-3 or more.
Incidentally, for a low alloy martensitic steel or bainite steel,
in view of the fact that the average grain diameter of a general
cementite (Fe3C) prior to warm working is several tens nanometers
or more, the volume fraction of the dispersed second-phase
particles is preferably set at 20.times.10.sup.-3 or more.
[0104] Whereas, the upper limit of the volume fraction of the
dispersed second-phase particles has no particular restriction in
effecting strengthening. However, in view of the toughness, it is
preferably set at 12.times.10.sup.-2 or less. Further, particle
dispersion strengthening by the Orowan mechanism is expected to
become remarkable in the region of several tens nanometers or less
from the equation (A). Thus, with the dispersion conditions of the
dispersed second-phase particles having an average particle
diameter of more than 0.5 .mu.m, it is difficult to obtain a
strength of 1.2 GPa or more. Therefore, the average particle
diameter of the dispersed second-phase particles is desirably 0.5
.mu.m or less, and more preferably 0.1 .mu.m or less as the steel
for warm working.
[0105] However, the foregoing conditions are predicated upon the
fact that the dispersed second-phase particles do not grow even in
the temperature range equal to or higher than that of the third or
more stage of tempering of 350.degree. C. In other words, in order
for the steel to have a strength of 1.2 GPa or more even after warm
working, it becomes a necessary condition that, during heating and
working, and after working, in addition to the matrix structure,
particularly, the dispersed second-phase particles do not undergo
Ostwald ripening, resulting in a reduction of the strength.
Therefore, when the thermal stability of the structure is evaluated
using .lamda. expressed by the following equation (1) commonly
known as a temper parameter as an index, conceivably, it is a
necessary and sufficient condition for the prior working structure,
i.e., the steel for warm working of the invention to show such a
softening resistance that the Vickers hardness (HV) at room
temperature when the steel is subjected to any heat treatment of
annealing, tempering, and aging in the as-unworked state under the
condition .lamda..gtoreq.1.4.times.10.sup.4 in a prescribed
temperature range of 350.degree. C. or more and Ac1 point or less
is equal to or higher than the hardness H given by the following
equation (2).
.lamda.=T(log t+20) (1)
where T denotes the temperature (K), and t denotes the time
(h).
H=(5.2-1.2.times.10.sup.4.lamda.).times.10.sup.2 (2)
[0106] Incidentally, the wording "in a prescribed temperature
range" denotes that the foregoing conditions may be satisfied at
any temperature of from 350.degree. C. to the Ac1 point, and means
that the foregoing conditions are not required to be satisfied over
the whole temperature range. In other words, also in the case where
when an aging or tempering treatment is carried out, the material
undergoes remarkable aging hardening or secondary hardening to have
a hardness of H or higher only in a given temperature range within
the foregoing range, it can serve as the steel for warm working of
the invention.
[0107] Herein, for example, for the TiC, the particle growth
suppressing effect will be considered. The stable crystal grain
diameter D resulting from the normal particle growth of the ferrite
structure including TiC dispersed therein (d=7.times.10.sup.-3
.mu.m, volume fraction f=7.times.10.sup.-3, B=4/9 to 4/3) is
estimated from the commonly well known Zener's relational
expression (D=B.times.d/f). Then, it is found to be about 0.4 to
1.3 .mu.m. In other words, such stable crystal grains are
established for the normal grain growth of recrystallized grains.
Therefore, the following can be sufficiently expected: with warm
working in a lower temperature range than the recrystallization
temperature, a fibrous structure with an average grain size of 3
.mu.m or less is obtained by the two effects of refinement of the
matrix structure by impartment of a predetermined strain and grain
boundary pinning by TiC.
[0108] In order to obtain an ultrafine duplex phase structure
having a tensile strength of 1.2 GPa or more by warm working based
on precipitation strengthening due to the ideal dispersion
conditions of a TiC carbide in this manner, the necessary and
sufficient condition of the prior working structure is as follows:
the lower limit value of the volume fraction of the dispersed
second-phase particles is set at 7.times.10.sup.-3, and the steel
after any heat treatment of annealing, tempering, and aging under
the condition T(log t+20).gtoreq.1.4.times.10.sup.4 has a hardness
of HV.gtoreq.(5.2-1.2.times.10.sup.4.lamda.).times.10.sup.2.
Namely, the invention has the following features as a steel for
warm working: fine dispersion or precipitation of the dispersed
second-phase particles in the matrix structure as particle
dispersion strengthening particles, and the structure control to
enhance the thermal stability of the dispersed second-phase
particles.
[0109] As for the structure of the steel for warm working of the
present invention as described up to this point, during the
treatment of warm working, the dispersion conditions of the
dispersed second-phase particles and the matrix structure variously
change. Therefore, although not limited by the room-temperature
structure form, all the steels with a strength of 1.2 GPa or more
except for the steels having a pearlite structure as the main
structure can be considered as the steel for warm working. As such
ones, for example, specifically, for martensitic steels (tempered
martensite structures), there are JIS-G4053 low alloy steels,
JIS-G-4801 spring steels, secondary hardened steels with a hardness
equal to or higher than this, maraging steels, TRIP steels, and
ausformed steels.
[0110] Then, the second steel for warm working of the invention is
configured such that 80 percent by volume or more of the matrix
structure is any single structure of martensite and bainite or a
mixed structure thereof. This is due to the following fact. It is
elucidated from a recent study that the width of the block regarded
as the effective crystal grain of martensite is 1 .mu.m or less in
a medium carbon low alloy steel (Scripta Mater., 49 (2003), P.
1157). Thus, by subjecting the tempered martensite structure
including a carbide or the like finely dispersed therein to warm
working, it is possible to form a fiber structure with efficiency.
In addition, the bainite structure also has a needle-like or
plate-like structure form including a carbide finely dispersed
therein. Thus, also when this is taken as the prior working
structure, it is possible to obtain a fibrous structure similarly.
In the steel for warm working of the invention, it is a preferred
form that a single structure of any of such martensite and bainite
or a mixed structure thereof accounts for 90 percent by volume or
more of the matrix structure.
[0111] Particularly, in order to keep the strength of 1.2 GPa or
more after warm working with stability, it is desirable that a
martensite or bainite structure having a temper softening
resistance equal to or higher than that of a tempered martensitic
steel of JIS-SCM430 steel is included in a volume of 80% or more.
Incidentally, 20% by volume or less structure other than martensite
or bainite and a mixed structure thereof may be accounted for by
any structure such as a ferrite, pearlite, or austenite structure.
This is for the following reason. Such a ferrite, pearlite, or
austenite structure, or the like decomposes or disappears, or
changes into a microstructure during a warm thermo-mechanical
treatment. Therefore, when it is present in an amount of 20% by
volume or less, it is judged as no problem.
(b) Chemical Composition
[0112] The third to fifth steels for warm working of the invention
are alloy designed based on the foregoing findings. The gist
resides in a steel for warm working characterized by containing, as
the chemical composition, C, 0.70 wt % or less, Si: 0.05 wt % or
more, Mn: 0.05 wt % or more, Cr: 0.01 wt % or more, Al: 0.5 wt % or
less, O: 0.3 wt or less, and N: 0.3 wt % or less, and the balance
substantially being Fe and inevitable impurities. Whereas, the
following and the like can be considered. The steel for warm
working further contains one or two or more selected from a group
consisting of Mo: 5.0 wt % or less, W: 5.0 wt % or less, V: 5.0 wt
% or less, Ti: 3.0 wt % or less, Nb: 1.0 wt % or less, and Ta: 1.0
wt % or less, or contains one or two or more of Ni: 0.05 wt % or
more and Cu: 2.0 wt % or less. Below, the reasons why the component
structure of the steel in the invention is restricted will be
described.
[0113] C: C forms a carbide particle, and is the most effective
component for strengthening. However, when the content exceeds 0.70
wt %, the toughness degradation is caused. For this reason, the
content is set at 0.70 wt % or less. In order that strengthening is
sufficiently expectable, the content is preferably 0.08 wt % or
more, and more preferably 0.15 wt % or more.
[0114] Si: Si is an effective element for enhancing the strength of
the steel by deoxidation and solid solution in ferrite, and finely
dispersing cementite. Therefore, the content is set at 0.05 wt % or
more inclusive of the one added as a deoxidizer, and to remain in
the steel. The upper limit is not particularly restricted for
strengthening. However, in view of the workability of the steel
material, the content is preferably set at 2.5 wt % or less.
[0115] Mn: Mn is an effective element for reducing the
austenitization temperature, and refining austenite. In addition,
it is an effective element for the hardenability, and being
dissolved in a solid solution form in cementite and suppressing
coarsening of cementite. When the content is less than 0.05 wt %,
desired effects cannot be obtained. Therefore, the content is set
at 0.05 wt % or more. More preferably, the content is 0.2 wt % or
more. The upper limit is not particularly restricted for
strengthening. However, in view of the toughness of the resulting
steel material, the content is preferably set at 3.0 wt % or
less.
[0116] Cr: Cr is an effective element for improving the
hardenability, and it is an element having a strong action of being
dissolved in a solid solution form in cementite and delaying the
growth of cementite. Further, it is also one of the important
elements in the present invention for forming a high Cr carbide
which is thermally more stable than cementite, or improving the
corrosion resistance by being added in a relatively larger amount.
Therefore, Cr is required to be contained at least in an amount of
0.01 wt % or more. It is contained in an amount of preferably 0.1
wt % or more, and more preferably 0.8 wt % or more.
[0117] Al: Al is an effective element for deoxidization and forming
an intermetallic compound with an element such as Ni and enhancing
the strength of the steel. However, excessive addition reduces the
toughness. Therefore, the content is set at 0.5 wt % or less.
Incidentally, when the intermetallic compound of Al and other
elements, nitride or oxide of Al, or the like is not used as the
dispersion strengthening particles, the content is preferably set
at 0.02 wt % or less, and further restrictively 0.01 wt % or
less.
[0118] O: O (oxygen) effectively acts not as an inclusion but as
grain growth preventing or dispersion strengthening particles when
it can be finely and uniformly dispersed as an oxide. However, when
oxygen is contained excessively, the toughness is reduced.
Therefore, the content is set at 0.3 wt % or less. When the oxide
is not used as the dispersion strengthening particles, the content
is preferably set at 0.01 wt % or less.
[0119] N: N (nitrogen) effectively acts as grain growth preventing
or dispersion strengthening particles when it can be finely and
uniformly dispersed as a nitride. However, when nitrogen is
contained excessively, the toughness is reduced. Therefore, the
content is set at 0.3 wt % or less. When the nitride is not used as
the dispersion strengthening particles, the content is preferably
set at 0.01 wt % or less.
[0120] Mo: Mo is an effective element for strengthening the steel
in the invention. It not only improves the hardenability of the
steel, but also is dissolved in a small amount in a solid solution
form also in cementite to make cementite thermally stable.
Particularly, completely separately from cementite, it newly causes
separate nucleation of an alloy carbide on dislocation in the
matrix phase, thereby to effect secondary hardening, resulting in
strengthening of the steel. Further, the formed alloy carbide is
effective for grain refinement, and is also effective for
replacement of hydrogen. Therefore, Mo is contained in an amount of
preferably 0.1 wt % or more, and more preferably 0.5 wt % or more.
However, it is an expensive element. In addition, excessive
addition forms a coarse undissolved carbide or intermetallic
compound to degrade the toughness. Therefore, the upper limit of
the content is set at 5 wt %. From the viewpoint of cost
efficiency, the content is preferably set at 2 wt % or less.
[0121] Incidentally, also for W, V, Ti, Nb, and Ta, the same
effects as with Mo are exerted. The respective upper limits of the
content are set. Further, the composite addition of these elements
is effective in finely dispersing dispersion strengthening
particles.
[0122] Ni: Ni is an element effective for improving the
hardenability, and effective for reducing the austenitization
temperature, and refining austenite, improving the toughness, and
improving the corrosion resistance. Further, it is also an
effective element for forming an intermetallic compound with Ti or
Al, and precipitation strengthening the steel when it is contained
in a proper amount. When the content is less than 0.01 wt %,
desired effects cannot be obtained. Therefore, the content is set
at 0.01 wt % or more. Ni is more preferably contained in an amount
of 0.2 wt % or more. The upper limit is not particularly
restricted. However, Ni is an expensive element, and hence it is
preferably contained in an amount of 9 wt % or less.
[0123] Cu: Cu is a detrimental element causing hot brittleness. But
on the other hand, when it is added in a proper amount, it causes
precipitation of fine Cu particles at 500.degree. C. to 600.degree.
C., thereby to strengthen the steel. When it is added in a large
amount, it causes hot brittleness. Therefore, the content is set at
2 wt % or less which is roughly the maximum amount of solid
solution into ferrite.
[0124] Incidentally, when strengthening due to precipitation of a
fine intermetallic compound is intended, it is also effective that
Co is contained in an amount of 15 wt % or less.
[0125] P (phosphorus) and S (sulfur) are not particularly
specified. However, P or S reduces the grain boundary strength, and
hence it is an element which is desired to be removed as much as
possible. Each content is preferably set at 0.03 wt % or less.
[0126] Incidentally, for other elements than the foregoing ones,
various elements are allowed to be contained in such an amount as
not to reduce the effects of the invention.
(c) Preparation of Steel for Warm Working
[0127] Incidentally, as the methods for manufacturing the steel for
warm working as described above, for example, various ones can be
considered according to the methods for manufacturing martensite
structures or bainite structures of JIS standard, and the like.
These are not limited to dissolution and forging methods. For
example, other manufacturing methods such as powder metallurgy can
also be used. Specifically, for example, the following procedure or
the like is also possible. By using a technique such as a ball
milling method, most undissolved compounds such as oxides in the
steel are steel powder dispersed with a size of nanometer size, and
then, (ISIJ International, 39 (1999), p 176), such a mechanical
milled powder is consolidated and formed in a proper temperature
range to obtain an objective bulk body.
(d) Warm Working
[0128] The warm working method of the invention is characterized by
subjecting any steel for warm working described above to warm
working for applying 0.7 or more strain in the temperature range of
350.degree. C. or more and Ac point -20.degree. C. or less. It can
also be considered that after performing warm working, an aging
treatment is performed in the temperature range of 350.degree. C.
or more and Ac1 point or less. According to such warm working, the
following advantages can be obtained:
[0129] (1) Recovery of dislocation moderately occurs, and crystal
grain refinement can be achieved, and the internal stress can be
reduced;
[0130] (2) Diffusion of alloy elements becomes relatively easy, and
decomposition and re-precipitation of the dispersed second-phase
particles of carbides or the like remarkably occur, which enables
refinement of the structure; and
[0131] (3) The deformation resistance (high temperature hardness)
of the steel is remarkably reduced, so that the steel can be formed
without occurrence of cracks and the like.
[0132] As for such a working temperature, more specifically, for
example, in the case of a medium carbon low alloy steel for use as
a steel for general mechanical structures, including a martensite
structure as the matrix, it can be set at 350.degree. C. or more
roughly corresponding to the third stage of tempering in which
cementite precipitates. Particularly, in order to effectively use
an alloy carbide, an intermetallic compound, Cu, or the like as the
dispersed second-phase particles, it is desirable that working is
carried out in the temperature range of 500.degree. C. to
650.degree. C. which is the precipitation temperature of the
second-phase particles.
[0133] On the other hand, in the portion which has undergone
austenite transformation during working, phase transformation such
as pearlite transformation or martensite transformation is effected
during the cooling process. As a result, there is a high
possibility that such an inhomogeneous structure as to cause the
occurrence of cracks is formed. Further, also in view of an
increase in temperature due to working heat generation, the upper
limit temperature of working is set at Ac1 point -20.degree. C.
However, as the combination of the working temperature and time of
the material, when the hardness is arranged by a tempering
parameter .lamda., the combination such that the Vickers hardness
at room temperature when the material is subjected to any of
annealing, tempering, and aging treatments in the as-unworked state
does not become 3.7.times.10.sup.2 or less in HV is preferable in
order to obtain a strength of 1.2 GPa or more after warm working.
Particularly, for working in a high temperature range, the time
required for working is required to be shorten in view of the
softening resistance and the heating time of the material.
[0134] The degree to which the structure is developed depends upon
the prior working structure, the working temperature, and the
strain amount. In other words, the necessary strain amount varies
according to the prior working structure or the working
temperature. Therefore, although the strain amount cannot be
strictly specified, it is preferable that a strain of 0.7 or more,
and more preferably 1 or more is imposed when a fibrous structure
is desired to be formed in the inside of the material. As for the
steel for warm working having a martensite or bainite structure in
which prior-austenitegrains have been elongated in fine fibers by
previously applying working in the unrecrystailization temperature
of austenite, it is possible to homogeneously form a fine fiber
structure by application with a strain amount of less than 1.
However, in most cases, it is desirable that the strain amount is
preferably 1 or more, and more preferably 1.5 or more.
[0135] At this step, the strain to be imparted may be introduced
not only in a single working pass, but in a plurality of divided
passes. Further, the direction of working is not constantly limited
to the same direction. Still further, the inter-pass time is also
not particularly restricted. Furthermore, the process also includes
imposing a prescribed strain not over the entire region of the
material to be worked, but on a specific region (e.g., the surface
layer requiring strengthening, or R part of a component). However,
the actual strain amount can be understood only after considering
the material characteristics of the material to be worked, the
friction conditions (e.g., the type or the presence or absence of a
lubricant) of the roll (mold for forging) and the material to be
worked, the deformation of the roll (mold for forging), rolling
(forging) rate, the rolling (forging) temperature, and the like.
Particularly, when component forming is carried out by forging, it
is necessary that ununiform strain has been introduced.
Accordingly, it is desirable to estimate the amount of strain with
high precision numerical analysis technology. However, in general,
when the cumulative rolling reduction is 45% or more for plate
rolling intended for the plane strain state, or when the cumulative
reduction of area is 45% or more for wire rod rolling, it can be
considered that a strain of 0.7 or more has been introduced into
the entire region of the material to be worked. Incidentally, when
the cumulative rolling reduction or the cumulative reduction of
area is 58% or more, it can be considered that a strain of 1 or
more has been introduced into the entire region of the material to
be worked. However, for example, even when the rolling reduction
(reduction of area) is less than 45%, a strain of 0.7 or more may
be introduced into the entire region or into a specific region of
the material to be worked under the influence of friction or the
like. Therefore, in that case, it is necessary to quantitatively
study the amount of introduced strain by numerical analysis.
(e) Steel Material
[0136] The steel material of the invention is a steel obtained by
warm working a steel for warm working as described above. It is
characterized in the following respects: it has a matrix structure
including a fibrous crystal having an average grains size of the
minor axis of 3 .mu.m or less; the second-phase particles are
finely dispersed in the matrix structure at a volume fraction of
7.times.10.sup.-3 or more at room temperature; and the Vickers
hardness at room temperature is 3.7.times.10.sup.2 or more in HV.
Incidentally, it can be understood that the matrix structure in the
steel material of the invention includes a fibrous ferrite crystal
with an expansion degree (aspect ratio) of more than 2, and
typically with an aspect ratio of 5 or more, in which the
second-phase particles are finely dispersed.
[0137] It is known that the effect of the crystal grain refinement
exerted on the mechanical characteristics of the steel becomes
remarkable in the crystal grain region of several micrometers or
less. In the invention, the upper limit of the average spacing
(i.e., minor axis average grain diameter) of the matrix structure
including the fibrous crystal is set at 3 .mu.m. Incidentally,
herein, the word "crystal grains" denotes the crystal grains
surrounded by the grain boundaries with a crystal orientation
difference of 15.degree. or more. On the other hand, when the
average particle diameter of the major axis of the dispersed
second-phase particles is larger than 0.3 .mu.m, particle
dispersion strengthening can be hardly expected. Further, for the
steel of 1.2 GPa or more, there is a high possibility that the
toughness is remarkably degraded. Accordingly, the average particle
diameter of the major axes is desirably 0.3 .mu.m or less.
[0138] Particularly, the effects of crystal grain refinement
becomes especially remarkable in the region in which the average
crystal grain diameter is 1 .mu.m or less; and particle dispersion
strengthening due to the Orowan mechanism, in the region in which
the average particles diameter is 0.1 .mu.m or less. Accordingly,
in order to effectively use strengthening by crystal fiber
formation and particle dispersion strengthening in a superposed
manner, it is effective that the minor axis average grain diameter
of the fibrous crystal is set at 1 .mu.m or less, and further 0.5
.mu.m or less. Then, it is more preferable that the average
particle diameter of the major axes of the dispersed second-phase
particles is also set at 0.1 .mu.m or less, and further 0.05 .mu.m
or less according to the refinement of the matrix structure.
[0139] To such a steel material for warm working, other than the
foregoing strengthening mechanisms, strengthening mechanisms such
as solid solution strengthening and dislocation strengthening can
also be applied. The effects of superposition of these
strengthening mechanisms lead to provide such a high performance
material as unpredictable with simple addition of the strengthening
mechanisms.
[0140] Such fine fiber structures can be formed by warm forming of
rod wire materials, screw members of bolts, and the like, including
plate materials. Particularly even when the cumulative strain
amount is small, a fiber structure can be formed in the surface
layer part which has locally undergone intense deformation or the
like. This can largely improve the characteristics of various
components and the desirable parts.
[0141] Below, examples will be shown by reference to the
accompanying drawings, and embodiments of the present invention
will be described in more details. Of course, this invention is not
limited to the following examples, and it is needless to say that
various forms are possible for the details.
EXAMPLES
[0142] Table 1 shows the steel components (A to K, M, N, and O)
within the scope of the invention and the steel component (L)
outside the scope. Incidentally, in examples, carbides are used as
the dispersed second-phase particles. Table 2 shows the volume
fractions of alloy carbides dispersible as the dispersed
second-phase particles and cementites for the steels of the
compositions of Table 1. The steels of the examples cover
martensitic steels of from SCM435 to 2 GPa class secondary hardened
steels except for Co-added maraging steels.
TABLE-US-00001 TABLE 1 Chemical composition (wt %) Steel type C Si
Cr Mn Ni Mo Cu Nb P S T-Al T-N O A 0.20 1.95 1.01 0.21 -- 1.02 --
-- 0.001 0.0010 0.010 0.0022 0.0007 B 0.39 1.98 1.04 0.21 -- 1.05
-- -- <0.001 <0.001 0.041 0.0020 <0.005 C 0.59 1.99 0.98
0.20 -- 1.01 -- -- <0.002 0.0007 0.004 0.0016 0.0005 D 0.21 0.09
2.01 0.21 2.02 1.01 -- -- 0.010 0.0020 0.011 <0.01 <0.004 E
0.22 0.09 2.00 0.20 2.00 1.00 -- 0.034 0.009 0.0020 0.007 <0.01
<0.004 F 0.40 0.08 2.00 0.20 2.01 1.03 -- -- 0.010 0.0020 0.009
<0.01 <0.004 G 0.57 1.96 1.02 0.16 -- 0.002 -- -- <0.001
<0.001 0.041 0.0018 <0.0005 H 0.35 0.20 1.10 0.70 0.25 0.20
0.053 -- 0.006 0.0020 0.009 0.0070 0.0040 I 0.40 0.20 1.03 0.77
0.07 0.17 0.040 -- 0.023 0.0080 -- 0.0050 0.0010 J 0.38 0.97 2.00
0.46 3.01 0.95 0.010 -- 0.018 0.0040 0.013 0.0010 0.0040 K 0.38
0.24 2.00 0.45 1.50 0.95 <0.01 -- 0.017 0.0040 0.012 0.0010
0.0040 L 0.22 0.01 <0.001 <0.005 0.01 <0.002 <0.001 --
<0.002 0.0020 0.003 0.0014 0.0018 M 0.62 1.97 0.01 0.20 2.00
0.98 -- -- <0.002 0.0007 0.003 0.0036 0.0006 N 0.40 0.10 2.00
0.20 1.00 0.70 -- -- 0.015 0.0030 0.020 -- -- O 0.20 0.25 0.15 0.50
3.00 3.00 -- 0.015 0.0030 0.020 -- --
TABLE-US-00002 TABLE 2 AMOUNT OF ALLOY CARBIDE AND CEMENTITE
DISPERSED (calculated value) Total volume Steel Type fMo.sub.2C
.times. 10.sup.3 fNbC .times. 10.sup.3 fFe.sub.3C .times. 10.sup.3
fraction .times. 10.sup.3 A 9.6 -- 21.0 31 B 9.9 -- 49.5 59 C 9.5
-- 79.5 89 D 9.5 -- 22.5 32 E 9.4 0.41 24.0 34 F 9.7 -- 51.0 61 G
-- -- 85.5 86 H 1.9 -- 51.0 53 I 1.6 -- 58.5 60 J 8.9 -- 48.0 57 K
8.9 -- 48.0 57 L -- -- 33.0 33 M 9.2 -- 84.0 93 N 6.6 -- 51.0 58 O
28.2 -- 3.0 31 Total volume fraction (f) of secondary phase
particles in each steel > 7 .times. 10.sup.-3
Total volume fraction (f) of secondary phase particles in each
steel >7.times.10.sup.-3
[0143] Carbides of various stoichiometric compositions are present
in actual steels according to the components of the steel and the
heat treatment conditions. For this reason, it is difficult and not
practical to strictly measure the volume fraction of the dispersed
second-phase particles by chemical analysis or structure
observation. Under such circumstances, the present inventors
determined the volume fraction of each carbide from the known
theoretical density determined by the structural analysis or the
like of the carbide (KAGAKU DAIJITENN, TOKYO KAGAKU DOJIN Co. Ltd.,
(1989), P. 1361 to 1363). The approximate expressions and the like
are as shown in Table 3.
TABLE-US-00003 TABLE 3 Approximate expression for determining the
volume fraction (f) of MxCy type carbide fMxCy =
.rho.Fe/.rho.MxCy/{XM*/(XM* + 12Y)} (wt % M)/100 M*: atomic weight
For example, fFe.sub.3C = .rho.Fe/.rho.Fe.sub.3C/{12/(3 * 55.85 +
12)} (wt % C)/100 = 0.15 (wt % C) fMo.sub.2C = 0.0094 (wt % Mo)
fNbC = 0.012(wt % Nb) fTiC = 0.020 (wt % Ti) where Density of
ferrite iron: .rho.Fe = 7.86 Mg/m.sup.3 Density of Fe.sub.3C;
.rho.Fe.sub.3C = 7.72 Mg/m.sup.3 Density of Mo.sub.2C;
.rho.Mo.sub.2C = 8.90 Mg/m.sup.3 Density of NbC; .rho.NbC = 7.78
Mg/m.sup.3 Density of TiC; .rho.TiC = 4.94 Mg/m.sup.3 The density
of each carbide is cited from the following document: (KAGAKU
DAIJITENN, TOKYO KAGAKU DOJIN Co. Ltd., (1989))
[0144] For calculation, it is assumed that the alloy elements
respectively combine with carbons in order of decreasing carbide
forming ability (Nb>Mo>Cr>Fe, and the like) to form
carbides. As for Nb or Mo, it is well known that it is an element
which tends to form its specific carbide in the steel and is less
likely to be dissolved in cementite. Thus, precipitation of NbC or
Mo2C is assumed. However, for a G steel or a L steel, Mo in an
amount of 0.002 wt % can be sufficiently dissolved in a solid
solution form in cementite, and hence it is excluded from the
estimation of the volume fraction of the Mo carbide. As for Cr,
when Cr is added in a large amount, it forms a carbide such as
M23C6 or M7C3 with a high Cr concentration. However, when Cr is
added in the amount of this example, there is a low possibility
that Cr is dissolved in a solid solution form in cementite to form
the alloy carbides. Therefore, the volume fraction of the alloy
carbides of Cr is excluded from the estimation.
[0145] The most important thing herein is the following: The amount
of carbide serving as the dispersion strengthening particles to be
dispersed depends upon the carbon content in a medium carbon low
alloy steel. Particularly, when there is no possibility that a
alloy carbide with a sufficiently large density with respect to
cementite is formed, or when the amount of elements for forming
alloy carbides to be added is small, the amount of the second-phase
particles to be dispersed is roughly determined by the amount of
cementite. Namely, as shown in Table 2, in the steels having a C
content of 0.2 wt % or more used in examples of Table 1, the total
amount of the volume fractions of the second phase sufficiently
exceeds 7.times.10.sup.-3.
[0146] In FIGS. 1, 2, and 3, steps of the thermo-mechanical
treatment applied in examples are shown. This process basically
includes (1) a solution treatment and working for reducing coarse
undissolved carbides; (2) a quench hardening treatment and
tempering for obtaining a tempered martensite or bainite structure
as a structure of the steel for warm working of the invention; and
(3) warm working also serving as shape forming into a component.
Incidentally, in the thermo-mechanical treatment pattern 1 of FIG.
1, refinement of reverse transformed austenite grains due to
austenitizing at low temperatures following the solution treatment
is taken into consideration. In the pattern 2 of FIG. 2, quench
hardening from recrystallized austenite resulting from hot working
following the solution treatment or the unrecrystallized austenite
(elongated austenite) structure resulting from warm working is
taken into consideration. FIG. 3 shows a quench hardening process
from the worked austenite (elongated austenite) structure by an
ausforming treatment in a metastable austenite range. With these
thermo-mechanical treatment processes, it is possible to obtain a
microstructure by warm working with a smaller cumulative strain
amount for finer crystal grains. Particularly, as the prior working
structure for developing a fiber structure with efficiency, it is
most effective to employ the martensite resulting from fine
unrecrystallized austenite (elongated austenite) as the prior
structure.
[0147] First, a square bar of about 40 mm square.times.120 mm in
length cut from a hot rolled steel sheet or a forged material was
subjected to the steps up to the quench hardening treatment in the
thermo-mechanical treatment patterns 1, 2, and 3, thereby to obtain
a martensite single structure close to nearly 100% by volume. This
corresponds to one example of the steel for warm working of the
invention. Then, the square bar was heated to a prescribed
temperature for 0.5 hour, and tempered. Then, it was subjected to
warm rolling working to a prescribed reduction of area by the use
of a groove roll, and applied with a strain, and air cooled.
[0148] For the structure of the resulting steel material, the cross
section in parallel with the rolling direction (RD) was polished
and observed by the use of an optical microscope, a transmission
electron microscope (TEM), and a FE-SEM with EBSP analyzer. The
polished surface was corroded with picric acid alcohol, and the
prior-austenite grain boundary was revealed. Thus, the
prior-austenite grain diameter was determined according to the
comparison method or the cutting method specified in JIS G 0552.
The average particle diameter of the dispersed second-phase
particles was determined in the following manner. By the use of TEM
or SEM, 3 or more visual fields were observed at a magnification of
10000 times to 100000 times to measure the length of each major
axis of a total of 250 or more particles. Incidentally, when some
particles combine with each other and agglomerate, these are
regarded as one particle. The maximum particle diameter is allowed
to correspond to the length of the major axis of the largest
carbide of the carbides measured. As for the average grain
diameters of the minor axes and the major axes of the elongated
grains in the fiber structure, according to EBSP analysis, the
average section lengths of the minor axes and the major axes of the
elongated crystal grains having a crystal orientation difference of
15.degree. or more were measured with a cutting method (see FIG.
5).
[0149] The hardness of the resulting steel material was measured
under a load of 20 kg and for a holding time of 15 s by means of a
Vickers hardness tester according to the testing method specified
in JIS Z 2244.
[0150] The tensile test was performed at ordinary temperatures by
means of an Instron type tensile test machine according to the
testing method specified in JIS Z 2241, for 1) a JIS No. 14
proportional test piece with a parallel part diameter of 3.5 mm, a
length of 24.5 mm, and a mark-to-mark distance of 17.5 mm, or with
a size of 6 mm, a length of 42 mm, and a mark-to-mark distance of
30 mm, or for 2) a JIS No. 4 sub-size test piece with a parallel
part diameter of 10 mm, a length of 45 mm, and a mark-to-mark
distance of 35 mm. The cross head speeds were 0.5 mm/min and 10
mm/min for 1) JIS No. 14A and 2) JIS No. 4, respectively. The
elongation was measured until rupture by mounting an extensometer
on each test piece.
[0151] The impact test was performed according to the testing
method specified in JIS Z 2242 for a U notch or V notch test piece
of 55 mm in length and 10 mm in height and width manufactured by
cutting machining from a steel material of 1.8 cm.sup.2 or more in
cross section.
[0152] The hydrogen embrittlement characteristics were evaluated at
room temperature at a cross head speed of 0.005 mm/min using a slow
strain rate tensile test machine for each notch test piece of 10 mm
in diameter, 6 mm in notch bottom diameter, and 4.9 in stress
concentration coefficient. For the hydrogen embrittlement test, the
test was carried out after setting the following conditions. The
average hydrogen amount in the test piece was changed by 72-hour
cathode charge with varying charged solution and current density,
to perform Cd plating. This prevents hydrogen in the test piece
from dissipating. The analysis of hydrogen was carried out by a
thermal desorption analysis using a quadrupole mass spectrometer
for a sample from which Cd-plating had been removed. Thus, the
hydrogen to be released up to 300.degree. C. was defined as
diffusive hydrogen, and determined.
[0153] Table 4 summarizes the manufacturing conditions and the
structure form of each steel for warm working, and the quench
hardening and tempering conditions of the unworked material and the
hardness thereof, and the results of evaluation of the suitability
as the steel for warm working of the invention.
[0154] FIG. 4 shows the relationship between T(log t+20)=.lamda.
and the hardness of the tempered martensitic steel in the
as-unworked state.
TABLE-US-00004 TABLE 4 Manufacturing conditions and structure form
of steel for warm working, and temper hardness of unworked material
Manufacturing conditions and structure form of steel for warm
working Temper hardness Quench of unworked hardening (1) Quench
Average material Solution Reduction or .gamma. Reduction hardening
prior .gamma. Quench T(logt + treatment of area transformation of
area temperature grain hardening 20) = Tempered Steel temperature
(1): e1 temperature (2) (2) diameter hardness .lamda. .times.
hardness type 1) (.degree. C.) .times. 10.sup.-2 (%) (.degree. C.)
.times. 10.sup.-2 e2 (%) (.degree. C.) .times. 10.sup.-2 (.mu.m)
.times. 10.sup.-3 (HV) .times. 10.sup.-2 10.sup.-4 (HV) .times.
10.sup.-4 A 2 12.0 40 9.0 -- -- 0.06 5.0 1.55 4.4 For working 1 B 1
12.0 40 9.2 -- -- 0.03 6.7 1.55 5.5 For working 2 2 12.0 40 9.0 --
-- 0.06 6.8 1.55 6.2 For working 3 1.75 5.0 C 2 12.0 40 9.0 -- --
0.06 8.7 1.55 6.2 For working 4 1.75 5.5 D 1 12.0 40 8.0 -- -- 0.01
4.8 1.55 3.7 For working 5 E 1 12.0 40 8.0 -- -- 0.009 4.6 1.55 4.1
For working 6 F 1 12.0 40 8.0 -- -- 0.007 6.7 1.55 4.4 For working
7 G 1 12.0 40 12.0 -- -- 0.5 8.0 1.55 5.0 For working 8 1.75 4.1 H
1 12.0 40 9.2 -- -- 0.04 6.0 1.55 3.8 For working 9 I 1 -- -- 11.0
-- -- 0.1 6.5 1.56 4.0 For working 10 J 2 12.0 40 9.0 -- -- 0.06
6.7 1.55 5.3 For working 11 K 2 12.0 40 9.0 -- -- 0.06 6.7 1.55 5.0
For working 12 L 1 -- -- 10.5 -- -- 0.07 4.8 1.36 2.7 Comp. Example
1 M 1 12.0 0 12.0 -- -- 0.5 8.5 1.55 5.4 For working 13 1.75 4.4 O
2 12.0 40 9.0 -- -- 0.06 4.8 1.55 4.2 For working 14 3 12.0 40 10.2
0 5.8 0.06 4.6 1.55 4.2 For working 15 33 5.8 0.05* 5.5 1.55 4.7
For working 16 55 5.7 0.04* 5.6 1.55 4.9 For working 17 70 5.8
0.03* 5.7 1.55 5.0 For working 18 1) Thermo-mechanical treatment
pattern *For ausformed (elongated grain) structure, the grain
diameter of the minor axis is measured.
[0155] As for the L steel of the comparative material, the volume
fraction of cementite is 33.times.10.sup.-3. However, the alloy
elements specified in the invention are not properly contained.
Therefore, cementite is not thermally stable, and grows with ease
by heating. Accordingly, with a tempering treatment at
.lamda.=1.4.times.10.sup.4 or more, the hardness of the L steel is
less than H=(5.2-1.2.times.10.sup.-4.lamda.) indicated with a
broken line in the diagram. Thus, with warm working at 350.degree.
C. or more, HV 3.7.times.10.sup.2 cannot be achieved for the L
steel.
[0156] FIG. 5 shows an example of analysis of the structure of the
material obtained by subjecting an I steel to .gamma.
transformation at 11.0.times.10.sup.2.degree. C., followed by water
cooling, and subjecting it to a tempering treatment at
5.0.times.10.sup.2.degree. C. for 1.5 hours, followed by warm
groove rolling. Incidentally, the cumulative strain amount imparted
at this step is 2.4, and the hardness is HV 3.7.times.10.sup.2. As
indicated from the EBSP analysis diagram (a) and the TEM photograph
(b) of the Bcc phase, there is obtained an ultrafine fiber
structure in which spheroidal carbide is dispersed in a fibrously
elongated ferrite phase matrix. By the EBSP analysis, the average
grain diameter of the minor axes of crystal grains having a crystal
orientation difference of 15.degree. or more was measured with a
cutting method. As a result (c), the average grain diameter of the
minor axes of the elongated crystal grains was found to be 0.3
.mu.m. However, in this steel, the fiber structure has been
developed in a complicated form, so that it was not possible to
measure the average grain diameter of the major axes. On the other
hand, the 287 carbide particle diameters (major axis lengths) were
measured by TEM. As a result, the average particle diameter of the
carbide was found to be 0.06 .mu.m, and the maximum diameter
thereof was found to be 0.2 .mu.m (d).
[0157] The inverse pole figure with respect to the rolling
direction (RD) indicates that there is formed a fiber structure in
which <011>//RD texture has developed. Incidentally, also for
other developed steels, the same textures were formed. The cleavage
plane of the Bcc iron is {100}. Therefore, the formation of such a
<011> fiber structure is considered to be very effective for
fracture due to tensile deformation in the direction of fiber axis,
flexural deformation receiving flexural moment along the direction
of fiber, or the like.
[0158] Table 5 shows the relationship between the warm working
conditions and the structure and hardness of the resulting warm
worked material. Incidentally, T and t in the table denote the
working temperature and the working treatment time shown in FIGS. 1
to 3, respectively.
TABLE-US-00005 TABLE 5 Warm working conditions and structure form
of worked material Structure form of warm worked material Warm
working conditions Average Maximum Working Cumulative carbide
carbide temperature equivalent particle particle (.degree. C.)
.times. T(logt + 20) = Reduction strain diameter diameter Steel
type 1) 10.sup.-2 .lamda. .times. 10.sup.-4 area e (%) amount:
.epsilon. *1 (.mu.m) *2 (.mu.m) *2 For working 1 A 2 5.0 1.55 76
1.7 <0.1 0.07 For working 2 B 1 5.0 1.55 76 1.7 <0.1 -- For
working 3 2 5.0 1.55 76 1.7 <0.1 0.1 5.0 1.55 12 0.2 <0.1 --
5.0 1.55 27 0.4 <0.1 -- 5.0 1.55 51 0.8 <0.1 -- 6.0 1.75 76
1.7 <0.1 0.3 7.0 1.96 76 1.7 0.2 0.85 For working 4 C 2 5.0 1.55
76 1.7 <0.1 0.15 6.0 1.75 76 1.7 <0.1 -- 7.0 1.96 76 1.7 0.2
1.2 For working 5 D 1 5.0 1.55 76 1.7 <0.1 0.4 For working 6 E 1
5.0 1.55 76 1.7 <0.1 0.3 For working 7 F 1 5.0 1.55 76 1.7
<0.1 0.3 For working 8 G 1 5.0 1.55 76 1.7 <0.1 0.4 6.0 1.75
76 1.7 0.2 0.6 7.0 1.95 76 1.7 0.3 0.8 For working 9 H 1 5.0 1.55
76 1.7 <0.1 0.25 5.0 1.55 57 1.0 <0.1 0.3 For working I 1 5.0
1.56 87 2.4 0.06 0.2 10 6.0 1.76 87 2.4 0.08 0.25 7.0 1.96 87 2.4
0.11 0.48 For working J 2 5.0 1.55 76 1.7 <0.1 0.2 11 For
working K 2 5.0 1.55 76 1.7 <0.1 0.25 12 For working M 1 5.0
1.56 76 1.7 <0.1 -- 13 6.0 1.76 76 1.7 <0.1 -- For working O
2 5.0 1.55 76 1.7 <0.1 <0.1 14 For working 3 5.0 1.55 79 1.8
<0.1 <0.1 15 For working 5.0 1.55 68 1.3 <0.1 <0.1 16
Structure form of warm worked material Minor axis Major axis
average grain average grain Worked diameter of diameter of
Elongation material elongated elongated rate hardness Steel type
grains L1 (.mu.m) grains L2 (.mu.m) (L2/L1) (HV) .times. 10.sup.-2
For working 1 A 0.4 Immeasurable -- 5.2 Example 1 For working 2 B
0.3 Immeasurable -- 4.9 Example 2 For working 3 0.3 Immeasurable --
5.6 Example 3 XX Immeasurable -- 5.3 Comp. Example 2 --
Immeasurable -- 5.4 Example 4 -- Immeasurable -- 5.2 Example 5 0.3
Immeasurable -- 4.6 Example 6 0.5 -- -- 3.3 Comp. Example 3 For
working 4 C 0.3 Immeasurable -- 6.0 Example 7 0.4 Immeasurable --
4.8 Example 8 0.9 5.4 6 3.6 Comp. Example 4 For working 5 D 0.3
Immeasurable -- 4.6 Example 9 For working 6 E 0.3 Immeasurable --
4.6 Example 10 For working 7 F 0.3 Immeasurable -- 4.9 Example 11
For working 8 G 0.4 Immeasurable -- 4.4 Example 12 0.5 Immeasurable
-- 3.8 Example 13 0.7 1.6 2 3.2 Comp. Example 5 For working 9 H 0.5
Immeasurable -- 3.8 Example 14 1.4 Immeasurable -- 3.8 Example 15
For working I 0.3 Immeasurable -- 3.7 Example 16 10 0.4
Immeasurable -- 3.1 Comp. Example 6 0.7 2.8 4 2.4 Comp. Example 7
For working J 0.3 Immeasurable -- 5.3 Example 17 11 For working K
0.3 Immeasurable -- 4.9 Example 18 12 For working M 0.3
Immeasurable -- 5.7 Example 19 13 0.7 Immeasurable -- 4.3 Example
20 For working O 0.9 Immeasurable -- 5.6 Example 21 14 For working
0.8 Immeasurable -- 5.7 Example 22 15 For working 1.0 Immeasurable
-- 5.9 Example 23 16 1) Thermo-mechanical treatment pattern *1
.epsilon. = 2/ 3(In(1/(1 - e/100)) *2 TEM observation or SEM
observation
[0159] The hardness of the worked material largely depends on the
temper softening resistance. When the comparison is made in the
same .lamda.=T(log t+20), a steel having a larger temper softening
resistance provides a worked material with a higher hardness.
Particularly, with a worked material of HV 4.0.times.10.sup.2 or
more, the matrix structure has been refined to as ultrafine as 0.5
.mu.m or less in average width. In a worked material of HV
4.0.times.10.sup.2 or more, very fine particles are densely
dispersed. For this reason, it was not possible to strictly
determine the average carbide particle diameter. However, when
comparison was made with the one including relatively large
particles such as the I steel of FIG. 5, it was possible to judge
the diameter as less than 0.1 .mu.m.
[0160] However, even in the steel for warm working high in temper
softening resistance, for example, with working in a high
temperature range of 700.degree. C., carbide particles and the like
grow with ease during working. Therefore, it becomes difficult to
obtain a warm worked material of HV 3.7.times.10.sup.2 or more
(Comparative Examples 3, 4, 5, and 7). Therefore, in working within
such a high temperature range, it is desirable that short-time
heating and working are combined and carried out so as to prevent
carbides and the like from growing by the use of, for example, high
frequency heating. Further, with working within a high temperature
range in the vicinity of 700.degree. C., grain growth becomes more
likely to occur, and hence the proportion of the crystal grains
with a small aspect and a relatively large grain diameter
increases. As a result, the expansion degree is reduced. For
example, in Comparative Examples 4, 5, and 7, the expansion degrees
were measured to be 6, 2, and 4, respectively. As for Examples, it
was not possible to measure the expansion degree. However, in
comparison in the structure with Comparative Example 7, it was
possible to judge the expansion degree as 6 or more.
[0161] Tables 6 and 7 summarize Examples and Comparative Examples
for the mechanical properties. Incidentally, in the tables, UE and
VE denote the absorption energies of the U notch and V notch test
pieces, respectively.
TABLE-US-00006 TABLE 6 Tensile deformation characteristics at room
temperature of warm worked steel and impact absorption energy (J)
at each test temperature Room Low Low Room Low Low Tensile temper-
temper- temper- temper- temper- temper- Proof strength Uniform
Total Reduction TS .times. total ature ature ature ature ature
ature stress TS elongation elongation of area elongation toughness
toughness toughness toughness toughness toughness Steel type (GPa)
(GPa) (%) (%) (%) balance UE20 (J) UE-20 (J) UE-60 (J) VE20 (J)
VE-20 (J) VE-60 (J) Ex. 1 A 1.66 1.66 3.8 19.2 56 32 -- -- -- 244
249 136 Ex. 2 B 1.73 1.73 6.7 16.4 45 28 160 228 174 -- 148 -- Ex.
3 1.86 1.86 6.5 16.9 51 31 214 -- 229 226 293 248 Comp. 1.70 1.78
2.7 12.3 45 22 -- -- -- 18 16 -- Ex. 2 Ex. 4 1.65 1.74 5.7 13.6 47
24 -- -- -- -- 21 -- Ex. 5 1.74 1.79 6.5 15.0 49 27 -- -- -- 45 26
-- Ex. 6 1.36 1.46 9.2 18.0 52 26 167 -- 254 -- -- -- Ex. 7 C 2.04
2.08 6.0 11.3 33 24 132 84 39 -- -- -- Ex. 8 1.44 1.53 7.7 12.3 38
19 183 150 87 -- -- -- Ex. 9 D 1.53 1.53 5.1 16.8 58 26 181 197 197
-- 154 -- Ex. 10 E 1.53 1.53 4.5 14.1 60 22 181 -- 2.3 -- -- -- Ex.
11 F 1.68 1.68 3.9 15.2 51 26 119 136 136 -- 201 -- Ex. 12 G 1.40
1.49 7.8 13.9 49 21 132 132 89 -- -- -- Ex. 13 1.09 1.22 9.6 18.0
50 22 101 138 121 -- -- -- Ex. 14 H 1.27 1.27 6.9 17.9 52 23 156
143 150 -- 125 -- Ex. 15 1.22 1.27 6.0 15.2 53 19 98 104 90 -- --
-- Ex. 17 J 1.78 1.80 6.2 18.5 45 33 -- -- -- 306 300 185 Ex. 18 K
1.75 1.75 -- 15.6 44 27 -- -- -- 196 269 158 Ex. 19 M 1.90 1.92 5.2
9.6 30 18 115 89 45 -- 125 -- Ex. 20 1.30 1.37 8.1 13.9 45 19 94 79
-- -- -- -- Ex. 21 O 2.04 2.04 -- 10.3 54 21 -- -- -- -- 38 -- Ex.
22 2.01 2.01 -- 10.0 52 20 -- -- -- -- 27 -- Ex. 23 2.03 2.03 --
9.2 51 19 -- -- -- -- 28 --
TABLE-US-00007 TABLE 7 Tensile deformation characteristics at room
temperature of unworked steel (QT material) and impact absorption
energy (J) at each test temperature Room Low Low Room .gamma.
Tempering Tensile Total temper- temper- temper- temper- transfor-
temper- Proof stress Uniform elon- Reduction TS .times. total ature
ature ature ature mation ature TS strength elongation gation of
elongation toughness toughness toughness toughness temperature
(.degree. C.) .times. Steel type (GPa) (GPa) (%) (%) area (%)
balance UE20 (J) UE-20 (J) UE-60 (J) VE-20 (J) (.degree. C.)
.times. 10.sup.-2 10.sup.-2 Comp. A 1.04 1.37 6.9 15.3 54 21 -- --
-- 15 9.9 5.0 Ex. 8 Comp. B 1.42 1.65 5.9 15.2 54 25 -- -- -- 24
9.5 6.0 Ex. 9 Comp. B 1.49 1.77 4.5 10.2 31 18 -- -- -- 12 9.5 5.0
Ex. 10 Comp. B 1.50 1.74 4.5 12.0 48 21 32 35 -- -- 9.2 5.0 Ex. 11
Comp. C 1.62 2.06 4.7 9.1 31 19 21 21 13 -- 9.2 4.5 Ex. 12 Comp. C
1.66 2.00 4.7 9.1 32 18 21 17 -- -- 9.2 5.0 Ex. 13 Comp. C 1.58
1.80 6.0 12.7 44 23 33 24 21 -- 9.2 5.7 Ex. 14 Comp. C 1.14 1.30
8.5 15.9 46 21 43 -- -- -- 9.2 6.5 Ex. 15 Comp. G 1.59 1.79 4.6
10.3 36 18 26 -- -- -- 9.2 4.5 Ex. 16 Comp. H 1.13 1.21 5.2 18.1 60
22 -- -- -- 69 8.5 5.0 Ex. 17 Comp. H 0.94 1.02 7.3 21.0 63 21 --
-- -- 106 8.5 6.0 Ex. 18 Comp. H 1.10 1.21 4.4 13.0 55 16 77 68 38
-- 9.2 5.0 Ex. 19
[0162] The steels of the compositions herein shown have been
subjected to proper alloy design and heat treatments so that the
second-phase particles are finely dispersed as the steels for warm
working. Even unworked steels of Comparative Examples show a
tensile strength.times.total elongation balance of 16 or more.
However, when comparison is made for the same composition, the
developed steel subjected to warm working provides a larger tensile
strength.times.total elongation balance than Comparative Examples.
Further, even addition of carbon in an amount of about 0.2 wt %
provides an over 1.5 GPa class almost the same as the quench
hardening hardness when alloy elements such as Mo are properly
added. Further, the ductility is excellent. These receive attention
(A, D, and E steels). Further, O steel provides a 2 GPa class
ultrahigh strength steel.
[0163] On the other hand, the impact absorption energy indicates
that each developed steel has a far more excellent toughness than
conventional high strength steels to a low temperature range.
[0164] FIG. 6 summarizes the relationship between the tensile
strength and the impact value (U notch test piece) at room
temperature. Incidentally, in the diagram, there is also shown data
of the steel for mechanical structure specified in JIS (SHINNNIHON
CYUUTANZOU KYOUKAI: GENNBAYOU KIKAIKOUZOUYOU HAGANEZAIRYOU DATA
SHEETS (1995)). With conventional steels, the impact value is
largely reduced in the strength range of 1.2 GPa or more, and 70
J/cm.sup.2 or less at a strength of 1.5 GPa or more. In contrast,
the steels of the invention show a very high impact value of 150
J/cm.sup.2 or more particularly even at a strength of 1.5 GPa or
more.
[0165] FIG. 7 summarizes the relationship between the tensile
strength and the absorption energy at room temperature (V notch
test piece). Incidentally, in the diagram, there is also shown data
of the steel for mechanical structure specified in JIS (Institute
of Metal Material Technology Fatigue Data sheet Reference 5). The
invention is also superior in toughness in a high strength region
to conventional ausformed steels, fine grain steels, maraging
steels, and the like.
[0166] FIG. 8 shows the relationship between the test temperature
and the absorption energy. For example, from Example 1 and
Comparative Example 8, and Examples 3 and 5 and Comparative
Examples 10, it can be confirmed that materials with high
absorption energy are obtained by a working treatment.
Particularly, the following receives attention. For some developed
steels, the absorption energy in the vicinity of room temperature
is not only higher than that of the comparative steels, but also
shows such a specific temperature dependency as to show the maximum
value in a low temperature range and decrease. For example, for the
A steel of Example 1, or the B steel of Example 3, a peak is
observed in the vicinity of -40.degree. C.; for the F steel of
Example 11, in the vicinity of -100.degree. C. There are also some
partially fractured ones in the peak temperature range. Then, such
developed steels are, as shown in FIG. 9, characterized in that the
fractured surface shows such a fibrous form as upon breaking
bamboo. The phenomenon similar to this has been observed when an
ausformed 0.2 wt % C-3 wt % Ni-3% Mo steel (tensile strength; 1.6
GPa) was tested in the vicinity of 200.degree. C. However, the
absorption energy in the vicinity of ordinary temperatures has been
reduced down to about 33 J (Non-Patent Document 16). Further, also
for the steel obtained by the improved ausforming treatment of a
0.5 wt % C-0.9 wt % Mn-0.8 wt % Cr steel (5150 steel), when an
impact test is performed, fibrous fracture occurs, and the
improvement of toughness is observed. However, the maximum
absorption energy at ordinary temperatures is about 90 J at a
strength level of 1.5 GPa (Non-patent Document 17). Therefore, the
following are the findings unprecedented and worthy of special
mention. At a tensile strength of 1.2 GPa or more, as is the case
with the developed steels of this invention, the absorption energy
in the vicinity of normal temperatures is far higher than that of
existing ausformed steels, and in addition, the absorption energy
shows the maximum value in a low temperature range of -40.degree.
C. or less.
[0167] The excellent mechanical characteristics, particularly the
high impact characteristics of the developed steels as described up
to this point are largely due to the ultrafinefiber structure with
<011>//RD texture which has been densely developed by warm
working of the particle dispersion type duplex structure.
[0168] FIG. 10 shows the relationship between the hardness of the
warm working material and the aging temperature. As for the warm
working material to which a secondary hardening element such as Mo
has been added, it is also possible to keep the hardness up to a
high temperature, or to more enhance the strength than in the
as-warm worked state by an aging treatment.
[0169] FIG. 11 shows an example of observation of the ultrafine
fiber structure formed in the central part of the plate material
with warm rolling at 650.degree. C. of N steel by SEM.
[0170] FIG. 12 shows an example of observation of the ultrafine
fiber structure formed in the surface layer part of the locally
intense strained rod material by SEM.
[0171] Table 8 shows the results of the hydrogen embrittlement
resistance characteristic test.
TABLE-US-00008 TABLE 8 Results of hydrogen embrittlement resistance
characteristics evaluation Diffusive Hydrogen Tensile 0.7 hydrogen
Notch Notch embrittlement strength TS content strength strength
resistance Steel type (GPa) *1 (GPa) (ppm) (GPa) *2 (GPa) *3
characteristics Ex. 1 A 1.66 1.16 0.36 2.41 1.76 AA Ex. 3 B 1.86
1.30 0.36 2.59 1.78 AA Ex. 7 C 2.08 1.46 0.32 2.17 1.63 AA Comp.
Ex. C 1.80 1.26 0.34 2.22 1.79 AA 14 Comp. Ex. C 2.06 1.44 0.25
1.91 0.73 CC 12 Comp. Ex. G 1.79 1.25 0.31 2.00 0.77 CC 16 Comp.
Ex. I 1.40 0.98 0.31 2.05 0.61 CC 20 *1 Tensile strength (TS) of
smooth tensile test with no hydrogen charging *2 Notch tensile
strength of notch test piece not charged with hydrogen (Kt = 4.9)
*3 Notch tensile strength of notch test piece charged with hydrogen
(Kt = 4.9) Hydrogen embrittlement resistance characteristics; the
case where the notch tensile strength is 0.7 time or more the TS of
the smooth test piece at a hydrogen content of about 0.3 mass ppm
is indicated with AA.
[0172] Herein, a notch tensile test of a steel charged with
hydrogen in an amount of about 0.3 mass ppm was carried out by a
slow strain rate tensile test. The hydrogen embrittlement
resistance characteristics were evaluated according to whether the
notch tensile strength at this step is 0.7 time or more the tensile
strength of a smooth tensile test piece not charged with
hydrogen.
[0173] The developed steels satisfy these conditions even when the
tensile strength is at a high strength level of 1.6 GPa or more,
and it can be judged as being excellent in hydrogen embrittlement
resistance characteristics. Incidentally, Comparative Example 14 is
a steel for a high strength mechanical structure excellent in
delayed fracture invented in Japanese Patent Application No.
2001-264399.
INDUSTRIAL APPLICABILITY
[0174] In accordance with the present invention, as described in
details up to this point, there is provided a high strength steel
multiphased by fine dispersion of a small amount of dispersed
second-phase particles. Particularly, even an ultrahigh strength
steel which is hard to soften and is hard to form is applied with
prescribed deformation in a temperature range in which the
deformation resistance is reduced and no cracks occur in the
material, to be formed into a prescribed shape (thin plate, thick
plate, wire rod, or component). As a result, conventional
spheroidizing and quench hardening and tempering treatments after
component forming are omitted. At the same time, an ultrafine
duplex phase structure is developed into a fibrous form. Thus,
there is provided a high strength steel largely improved in the
ductility, particularly the toughness, and the hydrogen
embrittlement resistance characteristics in the relation of
trade-off balance with high strength, and a member thereof.
[0175] These are useful as a steel to be worked into various
structures, car components and the like for use, or these are
useful as members.
* * * * *