U.S. patent application number 09/727580 was filed with the patent office on 2001-11-15 for alpha + beta type titanium alloy, process for producing titanium alloy, process for coil rolling, and process for producing cold-rolled coil of titanium alloy.
This patent application is currently assigned to KABUSHIKI KAISHA KOBE SEIKO SHO. Invention is credited to Fujii, Masamitsu, Furutani, Kazumi, Kida, Takayuki, Oyama, Hideto.
Application Number | 20010041148 09/727580 |
Document ID | / |
Family ID | 32096928 |
Filed Date | 2001-11-15 |
United States Patent
Application |
20010041148 |
Kind Code |
A1 |
Oyama, Hideto ; et
al. |
November 15, 2001 |
Alpha + beta type titanium alloy, process for producing titanium
alloy, process for coil rolling, and process for producing
cold-rolled coil of titanium alloy
Abstract
A high strength and ductility .alpha.+.beta. type titanium
alloy, comprising at least one is isomorphous .beta. stabilizing
element in a Mo equivalence of 2.0-4.5 mass %, at least one
eutectic .beta.stabilizing element in an Fe equivalence of 0.3-2.0
mass %, Si in an amount of 0.1-1.5 mass %, and C in an amount of
0.01-0.15% mass, and has a .beta. transformation temperature no
lower than 940.degree. C.
Inventors: |
Oyama, Hideto;
(Takasago-shi, JP) ; Kida, Takayuki; (Osaka,
JP) ; Furutani, Kazumi; (Takasago-shi, JP) ;
Fujii, Masamitsu; (Tokyo, JP) |
Correspondence
Address: |
OBLON SPIVAK MCCLELLAND MAIER & NEUSTADT PC
FOURTH FLOOR
1755 JEFFERSON DAVIS HIGHWAY
ARLINGTON
VA
22202
US
|
Assignee: |
KABUSHIKI KAISHA KOBE SEIKO
SHO
3-18 Wakinohama-cho 1-chome Kobe
Chuo-ku
JP
651-0072
|
Family ID: |
32096928 |
Appl. No.: |
09/727580 |
Filed: |
December 4, 2000 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
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09727580 |
Dec 4, 2000 |
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09317897 |
May 25, 1999 |
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6228189 |
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Current U.S.
Class: |
420/420 ;
148/669 |
Current CPC
Class: |
C22F 1/183 20130101;
C22C 14/00 20130101 |
Class at
Publication: |
420/420 ;
148/669 |
International
Class: |
C22C 014/00; C22F
001/18 |
Foreign Application Data
Date |
Code |
Application Number |
May 26, 1998 |
JP |
10-144558 |
Nov 12, 1998 |
JP |
10-322673 |
Claims
What is claimed is:
1. An .alpha.+.beta. titanium alloy comprising at least one
isomorphous .beta.-stabilizing element in a Mo equivalence of
2.0-4.5 mass %, at least one eutectic .beta.-stabilizing element in
an Fe equivalence of 0.3-2.0 mass %, Si in an amount of 0.1-1.5
mass %, and C in an amount of 0.01-0.15 mass %, and has a P
transformation temperature no lower than 940.degree. C.
2. The .alpha.+.beta. titanium alloy according claim 1, wherein an
Al equivalence is more than 3 mass % and less than 6.5 mass %.
3. The .alpha.+.beta. titanium alloy according claim 2, wherein
those elements of Al equivalence are entirely Al.
4. The .alpha.+.beta. titanium alloy according to claim 1, which
substantially contains Mo in an amount of 1.0-3.0 mass %, V in an
amount of 1.0-2.0 mass %, Fe in an amount of 0.3-1.0 mass %, Al in
an amount of 3.5-5.5 mass %, Si in an amount of 0.2-0.5 mass %, and
C in an amount of 0.02-0.15 mass %, with the remainder being Ti and
inevitable impurities.
5. The .alpha.+.beta. titanium alloy according to claim 1, which
contains O as an additional element such that the amount of
Mo-equivalenve, the amount of Fe-equivalence, and the content of 0
satisfy the following inequality [1]. 7.0 mass
%.ltoreq.(Mo-equivalence+2.5.times.Fe-equivalenc- e+40.times.O mass
%).ltoreq.19 mass % [1]
6. The .alpha.+.beta. titanium alloy according to claim 1, which
further contains a platinum group element in an amount of 0.03-0.2
mass %.
7. A process for hot-rolling the titanium alloy of any of claims 1
to 6, said process comprising heating the titanium alloy at a
temperature (T1) which satisfies the following inequatity [2] and
then rolling it. [.beta.-transus-20.degree. C.-(770.times.C mass
%).degree. C.].ltoreq.T1<.beta.-transus [2]
8. A process for rolling the titanium alloy of any of claims 1 to
6, said process comprising annealing the titanium alloy at a
temperature (T2) which satisfies the following inequality [3] and
then rolling it, thereby producing a coil of titanium alloy.
[.beta.-transus-270.degree.
C.].ltoreq.T2.ltoreq.(.beta.-transus-50.degree. C.) [3]
9. The process for rolling to produce a coil according to claim 8,
wherein rolling is carried out under a tension of 49-392 MPa such
that the draft is no lower than 20%.
10. The process for rolling to produce a coil according to claim 8,
wherein rolling is repeated more than once, with annealing in the
cc+p region intervening between consecutive rolling steps.
11. A process for annealing a cold-rolled coil of the titanium
alloy of any of claims 1 to 6, characterized in that the heating
temperature for annealing is higher than the temperature at which
work hardening due to cold-rolling is relieved and lower than the
.beta. transus but excludes the temperature range in which a alloy
of brittle hexagonal crystals emerges, thereby improving the
elongation in the transverse direction of the rolled strip of the
titanium alloy.
12. A process of annealing a coil cold-rolled strip of the titanium
alloy of any of claims 1 to 6, wherein annealing is carried out at
the temperature (T3) which satisfies the inequality [4] below so as
to give a coil rolled titanium alloy strip superior in bending
properties. (.beta.-transus-130.degree.
C.).ltoreq.T3.ltoreq.(.beta.-transus-15.degre- e. C.) [4]
13. A process of annealing a coil cold-rolled strip of the titanium
alloy of claim 4, wherein annealing is carried out at a temperature
no lower than 850.degree. C. and no higher than 963.degree. C. so
as to give a coil rolled titanium alloy strip superior in bending
properties.
Description
BACKGROUND OF THE INVENTION
[0001] 1. Field of the Invention
[0002] The present invention relates to a high strength titanium
alloy which has high strength, excellent weldability (i.e.,
ductility in heat affected zone (HAZ) after welding, the same
meaning hereinafter) and good ductility to make the production of
strips possible. The present invention relates to a titanium alloy
coil-rolling process and a process for producing a coil-rolled
titanium strip, in which the titanium is the above-mentioned
titanium alloy.
[0003] 2. Related Art
[0004] Titanium and its alloys are light, and excellent in
strength, toughness and corrosion-resistance. Recently, therefore,
they have widely been made practicable in the fields of the
aerospace industry, the chemical industry and the like. However,
titanium alloys are materials which are generally not so good in
workability, so that costs for forming and working are very high,
as compared with other materials. For example, Ti--6Al--4V, a
typical .alpha.+.beta. type alloy, is a material which is difficult
to work at room temperature. Thus, it is said that the alloy can
hardly be made into a coil by cold rolling.
[0005] For this reason, at the time of rolling the Ti--6Al--4V
alloy into a sheet form, a manner called pack-rolling is adopted.
That is, the pack-rolling is a manner of stacking Ti--6Al--4V alloy
sheets obtained by hot rolling in the form of layers, putting the
sheets into a box made of mild steel, and hot rolling the sheets
packed into the box under heat-retention for keeping its
temperature more than a given temperature to produce a thin plate.
In this process, however, a mild steel cover for making a pack and
pack welding are necessary. Moreover, in order to block bonding of
titanium alloy strips themselves, a releasing agent must be
applied. In such a manner, the pack-rolling process requires very
troublesome works and great cost, as compared with cold rolling.
Additionally, the temperature range suitable for hot rolling is
limited, to cause many restrictions in working.
[0006] On the contrary, Japanese Patent Application Laid-Open Nos.
3-274238 and 3-166350 discloses that the contents of Al, V and Mo
in the parent material of titanium are defined and at least one
alloying element selected from Fe, Ni, Co and Cr is comprised
therein in an appropriate amount, so that a titanium alloy can be
obtained which has a strength substantially equal to that of the
Ti--6Al--4V alloy and are superior to the Ti--6Al--4V alloy in
superplasticity and hot workability.
[0007] Japanese Patent Application Laid-Open Nos. 7-54081 and
7-54083 disclose a titanium alloy in which the Al content is
reduced up to a level of 1.0-4.5%, the V content is limited to
1.5-4.5%, the Mo content is limited to 0.1-2.5%, and optionally a
small amount of Fe or Ni is comprised thereinto, thereby keeping
high strength and raising cold workability and weldability (in
particular, HAZ after welding).
[0008] This titanium alloy has both cold workability and high
strength, and further has improved weldability, and thus is an
excellent alloy. However, in these inventions, flow-stress during
plastic deformation is suppressed because of the necessity of
ensuring excellent cold workability. Thus, its strength is
considerably low. If the strength is raised, its cold workability
drops. For this reason, production of cold strips are substantially
impossible. Incidentally, in recent years, customers I demands of
high strength and high ductility to titanium alloys have been
becoming more and more strict. Thus, titanium alloys are desired to
be improved still more.
SUMMARY OF THE INVENTION
[0009] Paying attention to the above-mentioned situation, the
inventors have made the present invention. The subject of the
present invention is an .alpha.+.beta. type titanium alloy, and an
object thereof is to provide an .alpha.+.beta. type titanium alloy
having excellent strength and cold workability, and further having
ductility making it possible to produce strips in coil. Another
object of the present invention is to establish a continuous
rolling technique based on coil-rolling by devising working
conditions, and provide a process for obtaining a titanium alloy
having excellent workability and strength by annealing after the
coil-rolling.
[0010] The high strength and ductility .alpha.+.beta. type titanium
alloy of the present invention for overcoming the above-mentioned
problems comprises at least one isomorphous .beta. stabilizing
element in a Mo equivalence of 2.0-4.5 mass %, at least one
eutectic .beta. stabilizing element in an Fe equivalence of 0.3-2.0
mass %, and Si in an amount of 0.1-1.5 mass %. (Hereinafter, %
means % mass unless specified otherwise.) In the titanium alloy, a
preferred Al equivalence, including Al as an a stabilizing element,
is more than 3% and less than 6.5%. If C is further comprised
thereinto in an amount of 0.01-0.15%, the strength property of the
alloy becomes more excellent. In addition, incorporation with a
platinum group element improves corrosion resistance. It is
important in view of rollability that the .beta. transus (T.beta.)
should be no lower than 940.degree. C.
[0011] The process for producing titanium alloy according to the
present invention is characterized in that a hot-rolling method
suitable for said titanium alloy is specified. The process consists
of heating the titanium alloy at a temperature (T1) satisfying the
following inequality [2] and then performing rolling.
[.beta.-transus-20.degree. C.-(770.times.C mass %).degree.
C.].ltoreq.T1<.beta.-transus [2]
[0012] The rolling method according to the present invention is
applicable to the continuous production of coil strip from the
above-mentioned titanium alloy. It consists of annealing a titanium
alloy plate or sheet at a temperature (T2) which satisfies the
following equation [3] and then performing rolling to produce
coiled strip.
[.beta.-transus-270.degree.
C.].ltoreq.T2.ltoreq.(.beta.-transus-50.degree- . C.) [3]
[0013] At the time of the coil-rolling, preferably the tension for
the coil-rolling ranges from 49 to 392 MPa and the rolling ratio
for the coil-rolling is 20% or more. If the coil-rolling is
performed plural times in a manner that an annealing step in the
.alpha.+.beta. temperature range intervenes therebetween, the total
rolling reduction can be raised as the occasion demands. Thus, even
a thin plate can easily be obtained.
[0014] Furthermore, the process for producing a titanium alloy
strip according to the present invention is a process of specifying
annealing suitable for cold-rolled strips after the cold-rolling of
the above-mentioned .alpha.+.beta. type titanium alloy. The process
is characterized by improving transverse elongation of a
cold-rolled titanium strip by selecting a heating temperature at
the time of annealing from temperatures which are not less than
temperature for relieving work-hardening at the time of
cold-rolling and are temperatures, in the range of temperatures not
more than .beta. transus (T.beta.), for promptly avoiding
temperature ranges causing brittleness resulting from the formation
of brittle hexagonal crystal a, so as to perform the annealing.
[0015] The above-mentioned titanium alloy is used to perform the
annealing, so as to easily obtain a titanium alloy strip having a
tensile strength after the annealing of 900 MPa or more, an
elongation of 4% or more, and [longitudinal (coil-rolling
direction)]/[transverse (direction perpendicular to the
coil-rolling direction) elongation]of 0.4-1.0.
BRIEF DESCRIPTION OF THE DRAWINGS
[0016] FIG. 1 is a graph showing the relationship between 0.2%
proof strength and elongation, after annealing in the .beta.
temperature range (corresponding to the properties in HAZ after
welding).
[0017] FIG. 2 is a phase diagram of a titanium alloy.
[0018] FIG. 3 is a view for explaining the coil-rolling process of
the present invention, referring to .alpha. phase diagram.
[0019] FIG. 4 is a graph showing the relationship between annealing
temperature, and strength and elongation obtained in Experiment
Examples.
[0020] FIG. 5 is a graph showing the relationship between annealing
temperature, and strength and elongation obtained in other
Experiment Examples.
[0021] FIG. 6 is a view conceptually showing the relationship
between annealing temperature and elongation that the inventors
have ascertained.
[0022] FIG. 7 is a view showing the relationship of ductility of
the transformed .beta. phase (i.e., the a phase) in the titanium
alloy, in the light of phase diagram in an .alpha.+.beta. type
titanium alloy.
[0023] FIG. 8 is a graph showing the relationship between 0.2%
proof strength and elongation after annealing in the .alpha.+.beta.
temperature range.
[0024] FIG. 9 is a graph showing the relation between tensile
strength and the value of
[Mo-equivalence+2.5.times.Fe-equivalence+40.times.0%].
[0025] FIG. 10 is a graphical representation of the results of the
experiment example (large scale) showing the relation between the
annealing temperature and the tensile strength and elongation in
the transverse direction.
[0026] FIG. 11 is a graphical representation of the results of the
experiment example (large scale) showing the relation between the
annealing temperature and the tensile strength and elongation in
the longitudinal direction.
[0027] FIG. 12 is a graphical representation of the results of the
experiment example (large scale) showing the relation between the
annealing temperature and the minimum bending radius.
[0028] FIG. 13A is a graph showing difference in ability to keep
passive state between pure titanium and the titanium alloy of the
present invention.
[0029] FIG. 13B is a graph showing difference in corrosion speed
between pure titanium and the titanium alloy of the present
invention.
[0030] FIG. 13C is a graph showing difference in resistance to
crevice corrosion between pure titanium and the titanium alloy of
the present invention.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0031] The .alpha.+.beta. type titanium alloy of the present
invention has a basic composition wherein the contents of
isomorphous .beta. stabilizing element and eutectic .beta.
stabilizing element are defined, and preferably Al equivalence
including Al, which is an .alpha. stabilizing element, is defined.
The .alpha.+.beta. type titanium alloy is an alloy wherein an
appropriate amount of Si is comprised into the basic composition
and preferably an appropriate amount of C is comprised as another
element thereinto, so as to give excellent strength property and
cold workability, thereby having high strength and simultaneously
making the production of coils possible. The following will
describe reasons of defining the contained percentages of the
above-mentioned respective elements.
[0032] At Least One Isomorphous .beta. Stabilizing Element: Mo
Equivalence of 2.0-4.5%:
[0033] The isomorphous .beta. stabilizing elements such as Mo cause
an increase in the volume fraction of the .beta. phase, and is
solved into the .beta. phase to contribute to a rise in strength.
Moreover, the isomorphous .beta. stabilizing elements have a nature
that they are solved into the parent material of titanium to
produce fine equiaxial microstructure easily. They are useful
elements from the standpoint of enhancing strength-ductility
balance. In order to exhibit such effects of the isomorphous .beta.
stabilizing elements effectively, they should be comprised in an
amount of 2.0% or more, and preferably 2.5% or more. However, if
the amount is too large, ductility after .beta. annealing decreases
and further corrosion of the titanium alloy increases. Thus, it
becomes difficult to remove TiO.sub.2 scales produced in the
annealing after cold rolling and an oxygen-solved ground metal,
called an .alpha.-case, so that the workability falls.
Additionally, the density of the whole of the titanium alloy is
heightened to damage the property of a high specific strength which
the titanium alloy originally has. Therefore, the above-mentioned
amount should be 4.5% or less, and preferably 3.5% or less.
[0034] The most typical element among all isomorphous .beta.
stabilizing elements is Mo. However, V, Ta, Nb and the like have
substantially the same effect as that of Mo. In the case wherein
these elements are contained, the Mo equivalence [Mo+{fraction
(1/1.5)}.times.V+1/5.times.Ta- +{fraction (1/3.6)}.times.Nb] ,
including these elements, should be adjusted into the range of
2.0-4.5%. However, Nb and Ta are less effective in
.beta.-stabilization per unit amount added. Therefore, they should
be added in a large amount to attain the same degree of
stabilization; moreover, they are expensive. It is recommended that
they are substituted with Mo and V. V is less expensive than Mo to
achieve the same degree of .beta.-stabilization. However, V added
alone decreases the T.beta. excessively. Consequently, the
desirable amount is 1.0-3.0% for Mo and 1.0-2.0% for V.
[0035] At Least One Eutectic .beta. Stabilizing Element: Fe
Equivalence of 0.3-2.0%:
[0036] The eutectic .beta. stabilizing elements such as Fe cause
improvement in strength by addition of a small amount thereof.
Moreover, they have the effect of improving hot workability.
Furthermore, cold workability is enhanced, particularly when Mo and
Fe coexist, but this reason is unclear at present. In order to
exhibit such effects effectively, Fe should be contained in an
amount of 0.3% or more, and preferably 0.4% or more. However, if
the amount is too large, ductility after .beta. annealing is
greatly lowered and further segregation becomes remarkable at the
time of ingot-making to reduce the stability of quality. The amount
should be 2.0% or less and preferably 1.5% or less.
[0037] Cr, Ni, Co and the like have substantially the same effect
as that of Fe. Thus, in the case that Cr and the like are
contained, the Fe equivalence
[Fe+1/2.times.Cr+1/2.times.Ni+{fraction (1/1.5)}.times.Co+{fraction
(1/1.5)}.times.Mn], including these elements, should be adjusted
into the range of 0.3-2.0%. However, it is recommended to replace
all of them by Fe, because Fe is cheapest and Cr slightly decreases
tensile strength. The minimum amount of Fe should preferably be
0.3% in view of the effect of improving hot-rollability and
strengthening. The maximum amount of Fe should preferably be 1.0%,
because Fe in an excessive amount causes remarkable segregation in
the Vacuum arc remelting (VAR) process.
[0038] Al Equivalence: More than 3%, and Less than 6.5%
[0039] Al is an element which contributes, as an
.alpha.-stabilizing element, to the improvement in strength. If the
Al content is 3% or less, the strength of the titanium alloy is
insufficient. However, if the Al content is 6.5% or more, the limit
cold-reduction is lowered so that it becomes difficult to make the
alloy into a coil. Additionally, the cold workability as a coil
product is also lowered so as to increase the number of cold
working steps and annealing steps until the alloy is rolled up to a
predetermined thickness. Thus, a rise in cost is caused.
Considering the strength-cold workability balance, preferably the
lower limit and the upper limit of the Al equivalence are 3.5% and
5.5%, respectively.
[0040] In the present invention, Sn and Zr also exhibit the effect
as an a stabilizing element in the same way as Al. Therefore, in
the case that these elements are contained, the Al equivalence
[Al+1/3.times.Sn+1/6.tim- es.Zr], including these elements, should
be desirably adjusted into the range of more than 3% and less than
6.5%. However, in the case where Sn and Zr are contained as the
a-stabilizing elements of Al equivalence, it is recommended to
replace all of them by Al because they have an adverse effect on
cold-rollability.
[0041] Typical examples of preferable .alpha.+.beta. type titanium
alloys satisfying the requirement of the above-mentioned
composition used as a base titanium alloy in the present invention
includes Ti--(4-5%)Al--(1.5-3%)Mo--(1-2%)V--(0.3-2.0%)Fe, in
particular Ti--4.5%Al--2%Mo--1-6%V--0.5%Fe.
[0042] Si: 0.1-1.5%
[0043] The .alpha.+.beta. type titanium alloy having the basic
composition that satisfies the content requirements of the
isomorphous p stabilizing element, the eutectic .beta. stabilizing
element, and the Al equivalence has an excellent cold workability
exhibiting a limit cold-reduction of about 40% or more. Thus, the
alloy can be made into a coil. However, its strength property and
weldability are not necessarily sufficient. The alloy cannot meet
the recent demand of enhancing strength.
[0044] However, it has been ascertained that if Si is contained in
an amount of 0.1-1.5% into the .alpha.+.beta. type alloy of the
above-mentioned basic composition, it is possible to heighten
remarkably the strength property and the property (strength and
ductility) in HAZ after welding, as a titanium alloy, without
lowering ductility necessary for making the alloy into a coil.
[0045] In other words, Si has an effect of raising the strength
property in the state that Si hardly has a bad influence on
cold-reduction of the .alpha.+.beta. type titanium alloy.
Furthermore, Si exhibits an effect of raising the strength and
ductility in HAZ after welding. By such addition of an appropriate
amount of Si, it is possible to obtain an alloy wherein the
strength and ductility of the titanium alloy parent material are
raised still more and further the HAZ after welding have strength
and ductility of a high level.
[0046] In order to exhibit such effects of Si more effectively, it
is necessary that Si is contained in an amount within a very
restrictive range of 0.1-1.5%. If the Si content is insufficient,
the strength tends to be short. Moreover, the effect of the
improvement in the strength-ductility balance of the welded zone
also becomes insufficient. On the other hand, if the Si content is
more than 1.5%, the cold-reduction becomes poor so that a coil
cannot easily be produced. Considering the above-mentioned
advantages and disadvantages of Si, preferably the lower limit and
the upper limit of the Si content are 0.2% and 1.0%, respectively.
The more preferable upper limit of Si is 0.5%, because Si in excess
of 0.5% suffers from poor cold-rollability.
[0047] Si in an Amount up to 0.5% Greatly Improves
Cold-Rollability.
[0048] C: 0.01-0.15%
[0049] Carbon (C) has an effect of enhancing the strength property
of the .alpha.+.beta. type titanium alloy still more while keeping
excellent ductility thereof, and an effect of enhancing the
strength in HAZ after welding remarkably with a little drop in the
ductility thereof. Such effects of the addition of C make the
strength and the ductility of the titanium alloy parent material
far higher, and also makes the strength and the ductility of the
HAZ even higher. Also, C is an essential element to raise the
.beta.-transus above 940.degree. C. so that the hot-rolling
temperature is set up as high as possible.
[0050] In order to exhibit such effects of C more effectively, it
is necessary that C is contained in an amount within a very
restrictive range of 0.01-0.15%. If the C content is insufficient,
the strength is insufficient, the increase of .beta.-transus is
also insufficient. On the other hand, if the C content is over
0.15%, cold-reduction is damaged by remarkable
precipitation-hardening of carbides such as TiC to block
coil-rolling. Considering such advantages and disadvantages of C,
preferably the lower limit and the upper limit of the C content are
0.02% and 0-12%, respectively.
[0051] In the present invention, if a small amount of 0 (oxygen) is
comprised thereto, as well as Si and C, the strength can be raised
still more in the state that the oxygen hardly has a bad influence
on coil-formation of the titanium alloy and its ductility. Thus, it
is preferable for oxygen to be comprised. Such an effect of oxygen
is exhibited by its very small amount. In order to exhibit the
effect more surely, oxygen is comprised in an amount of preferably
about 0.07% or more, and more preferably about 0.1% or more.
However, if the oxygen content is too large, the cold workability
drops. Besides, the ductility also drops by an excessive rise in
the strength. The oxygen content should be 0.25% or less and
preferably 0.18% or less. Incidentally, a strip produced by
unidirectional rolling has a decreased strength in the longitudinal
direction due to anisotropy. For the strip to have a strength
higher than 900 MPa in the longitudinal direction, it is necessary
to take the effect of oxygen quantity into account. It is essential
that the total amount of Mo-equivalence +2.5 x Fe-equivalence
+40.times.0% should be higher than 7.0%. If the total amount
exceeds 19%, the titanium alloy is so poor in ductility that it is
incapable of rolling. As the total amount exceeds 16.2%, the
titanium alloy begins to decrease in cold-rollability. Therefore,
the upper limit is 19%, and the preferred upper limit is 16.2%.
[0052] Reasons why such effects and advantages as above are
exhibited in the present invention by comprising an appropriate
amount of Si, C plus such an amount of Si, or further an
appropriate amount of oxygen into the .alpha.+.beta. type titanium
alloy as a base are not necessarily made clear, but the following
reasons can be considered.
[0053] That is, the reason why the strength property can be
improved without damaging the cold-reduction can be considered as
follows. Although Si is solved into the .beta. phase to contribute
to the strength, Si is not a factor for reducing the ductility very
much. Even if Si is comprised over its solubility limit, silicide
is formed so that the concentration of Si in the P phase is kept
not more than a given level. Therefore, if the Si content is
controlled into the range that the ductility is not reduced by the
excessive formation of silicide, the alloy keeps a high ductility
and simultaneously has an improved strength property.
[0054] If Si is comprised in an appropriate amount, silicide formed
in the .beta. phase as described above causes the suppression of a
phenomenon that the grain in the HAZ after welding is made coarse.
Additionally, Ti is trapped by the precipitation of silicide so
that the P phase is stabilized, or the retained .beta. phase
increases by the transformation-suppressing effect of solved Si. It
appears that these effects are cooperated to improve
weldability.
[0055] Carbon is solved into the a phase to contribute to the
improvement in the strength, but does not become a factor for
reducing the ductility of the a phase very much. In addition, if C
is comprised over its solubility limit, a carbide is formed so that
the concentration of C in the .alpha. phase is kept not more than a
certain level. Therefore, it appears that if the C content is
controlled into the range that the ductility is not reduced by the
excessive of carbide, the alloy keeps a high ductility and
simultaneously has an improved strength property. Incidentally, Si
and C produce the effect of enhancing the heat resistance of the
titanium alloy in addition to the above-mentioned effects.
[0056] Furthermore, O is solved into both of the .alpha. phase and
the .beta. phase (the solved amount is larger in the .alpha.
phase), to exhibit solution-hardening effect. However, if the
solved amount becomes large in either phase, the ductility is
reduced. Thus, the oxygen content should be controlled into a very
small amount as described above.
[0057] .beta.-transus Higher than 940.degree. C.
[0058] In hot-rolling at a temperature (for .alpha.+.beta. region)
lower than the .beta.-transus, which is essential for the equiaxial
structure, the titanium alloy is remarkably subject to edge
cracking due to temperature drop that occurs as hot-rolling
proceeds if the heating temperature is lower than 900.degree. C.
Edge cracking extremely lowers yields. On the other hand, the
temperature of the heating furnace inevitably deviates about
.+-.20.degree. C. from the aimed value on account of limited
control precision. Therefore, it is necessary that the lowest
.beta.-transus should be 940.degree. C.
[0059] Small amounts of other elements than the above may be
comprised as inevitable impurity elements into the titanium alloy
of the present invention. However, so far as they do not hinder the
property of the alloy of the present invention, these elements are
allowable to be comprised. The titanium alloy may be incorporated
with other elements than mentioned above so that it has additional
characteristic properties without altering its original ones
ascribed to the present invention. Examples of such elements
include platinum group elements (such as Pb, Ru, Ir, and In, about
0.03-0.2%) which improve corrosion resistance, P (less than about
0.05%) which improves heat resistance, and N (less than about
0.03%) which improves strength.
[0060] Platinum Group Element: 0.03-0.2%
[0061] It is generally known that titanium improves in corrosion
resistance by incorporation with a platinum group element. This
also applies to the titanium alloy of the present invention. The
titanium alloy incorporated with more than 0.05% of Ru (which is
the cheapest among platinum group elements) is comparable to or
better than pure titanium in corrosion resistance, without adverse
effect by Ru on its hot-workability, cold-workability, and
strength. This effect levels off when the amount of Ru exceeds
0.2%. The upper limit of the amount of Ru should preferably be
0.2%, more preferably less than 0.1%, because Ru is more expensive
than common elements. Pt and Ir in a smaller amount are as
effective as Ru in improving corrosion resistance.
[0062] The .alpha.+.beta. type titanium alloy of the present
invention wherein the constituent elements are specified as above
has a basic composition wherein the contents of the isomorphous
.beta. stabilizing element and the eutectic .beta. stabilizing
element are defined, and preferably Al equivalence is defined. The
.alpha.+.beta. type titanium alloy is an alloy wherein an
appropriate amount of Si is comprised into this basic composition
or optionally an appropriate amount of C or O is comprised
thereinto so as to have a high level strength property and
simultaneously an excellent ductility making the production of
coils possible, and further have an excellent weldability.
Specifically, the alloy has a 0.2% proof strength after annealing
in the .alpha.+.beta. temperature range of 813 MPa or more, a
tensile strength of about 882 MPa or more, and a limit
cold-reduction of 40% or more.
[0063] Even in the case of .alpha.+.beta. type titanium alloys, if
the alloys have a limit cold-reduction of less than 40%, at the
time of producing the alloys continuously into coils the number of
repeated cold rolling-annealing steps becomes large so that costs
become unsuitable for the actual situation. In addition,
recrystallized microstructure cannot easily be obtained, resulting
in a problem that the transverse and longitudinal anisotropy as a
strip material becomes larger. However, the alloy having a limit
cold-reduction of 40% or more can be made into coils without any
difficulty by a continuous method. Costs can be greatly reduced by
the improvement in productivity.
[0064] The limit cold-reduction herein means a reduced ratio of a
strip thickness in such a limit state that, after the step wherein
a small crack is produced but the propagation of the crack stops at
a certain level (for example, about 5 mm), the crack starts to
propagate up to the surface of the strip, from an industrial
standpoint.
[0065] Hot-rolling to produce coiled strips from the .alpha.+.beta.
titanium alloy of the present invention should be carried out under
the following conditions.
[0066] Prior to hot-rolling, the titanium alloy should be heated at
a temperature (T1) which satisfies the following inequality [2] so
that coiled strips with a minimum of edge cracking are produced in
high yields.
[.beta.-transus-20.degree. C.-(770.times.Cmass %).degree.
C.]<T.sub.1<.beta.-transus [2]
[0067] If the heating temperature is lower than
[.beta.-transus-20.degree. C.-(770.times.C mass %).degree. C.], the
titanium alloy suffers edge cracking remarkably due to temperature
fall during hot-rolling. In actual tandem rolling with one heating
stage from a slab (thicker than 100 mm, for example) into a 4-mm
thick coiled sheet, serration-like edge cracking occurs in the
lateral direction (longer than about 60 mm). Such edge cracks have
to be trimmed away together with the uncracked portion (more than
20 mm wide); otherwise, the sheet is very likely to break in the
cold-rolling step. By contrast, edge cracks will be smaller than 30
mm at the most if the heating temperature is higher than
[.beta.-transus-20.degree. C.-(770.times.C mass %).degree. C.]. In
this case, trimming up to 10 mm beyond edge cracks is enough to
greatly reduce the possibility of breaking in the course of cold
rolling. The higher is the heating temperature, the more decreases
the depth of edge cracks. However, heating at a temperature above
the .beta.-transus brings about rapid oxidation and transfer from
the equiaxial structure into the acicular structure, thereby making
the sheet liable to surface cracking and internal cracking in the
course of cold rolling. Therefore, the heating temperature should
be lower than the .beta.-transus. Moreover, the heating temperature
should preferably be lower than .beta.-transus minus 10.degree. C.
in consideration of the fact that the .beta.-transus varies from
one place to another due to macroscopic segregation. In this way it
is possible to produce in very high yields the desired titanium
alloy sheet without edge cracking.
[0068] Incidentally, in the present invention, a high level
strength property can be kept and simultaneously an excellent
cold-reduction making the production of coils possible can be
ensured by specifying the basic composition of the .alpha.+.beta.
type titanium alloy and simultaneously specifying the Si content,
or further the C or O content as described above. From further
investigations on requirements for surer assurance of the strength
property in HAZ after welding of such titanium alloys, it has been
ascertained that the alloy wherein the relationship between the
0.2% proof strength (YS) and the elongation (EL) satisfies the
following inequality [5] is good in the strength-elongation balance
in the HAZ after welding and stably exhibits a high weldability.
This matter will be in detailed described, referring to FIG. 1, in
Examples described later.
6.9.times.(YS-835)+245.times.(El-8.2).gtoreq.0 [5]
[0069] The following will describe a coil-rolling process for
producing the .alpha.+.beta. type titanium alloy of the present
invention efficiently and continuously.
[0070] At the time of coil-rolling the above-mentioned titanium
alloy, a strip of the titanium alloy is annealed at the temperature
[T2] satisfying the inequality [3] below, and then coil-rolled to
produce coils efficiently and continuously. Furthermore, at the
time of the coil-rolling, it is preferred to adjust the tension
into the range of 49-392 MPa and set a rolling ratio to 20% or
more. If the coil-rolling is performed plural times in a manner
that an annealing step in the .alpha.+.beta. Temperature range
intervenes therebetween, the total rolling reduction can be
heightened as the occasion demands. Even a thin plate can easily be
obtained.
(.beta. transus-270.degree. C.).ltoreq.T2.ltoreq.(.beta.
transus-50.degree. C.) [3]
[0071] The heat treatment conditions are very important
requirements for performing the coil-rolling easily.
[0072] That is, the criterion of the microstructure which controls
mechanical properties of titanium alloys is phase diagram as shown
in FIG. 2. (Its vertical axis represents temperature, and its
horizontal axis represents the amount of .beta.-stabilizing
elements.) As the contained percentage of the .beta. stabilizing
elements in the titanium alloy increases, the .beta. transus drops
in the form of a parabola. Therefore, at the time of heat-treating
titanium alloys, their microstructure varies remarkably dependently
on whether the heat temperature is set up to a higher temperature
than the .beta. transus of the respective alloys, or a lower
temperature than it.
[0073] The inventors paid attention to the .beta. transus of
titanium alloys and the change in their microstructure by heat
treatment temperature, and considered that, concerning the
.alpha.+.beta. type alloy of the present invention, a
microstructure suitable for cold rolling would be obtained by
setting appropriate heat treatment conditions. Thus, the inventors
have been researching from various standpoints. As a result
thereof, it has been found that if the titanium alloy strip having
the composition according to the present invention is subjected to
annealing at a temperature (T2) satisfying the following inequality
[3], its microstructure can be made up to a microstructure
comprising .alpha. phase+metastable .beta. phase or orthorhombic
martensite (acc) and having a very high ductility so that
coil-rolling can easily be performed.
(.beta. transus-270.degree. C.).ltoreq.T2.ltoreq.(.beta.
transus-50.degree. C.) [3]
[0074] As described in, for example, "METALLURGICAL TRANSACTIONS A,
VOLUME 10A, JANUARY 1979, P. 132-134", the .beta. transus of Ti
alloys which are objects of coil-rolling can be obtained from, for
example, the following equation [6], which is well known as a
calculating equation of the .beta. transus obtained from the
amounts of alloying elements contained in the titanium alloys:
the .beta. transus=872+23.4.times.Al %-7.7.times.Mo %-12.4.times.v
%-14.3.times.Cr %-8.4.times.Fe % [6]
[0075] Referring to a phase diagram of FIG. 3, reasons for setting
the annealing temperature conditions for which the .beta. transus
is an index will be made clear in the following.
[0076] In connection with FIG. 3, the inventors ascertained the
following in the case of annealing .alpha.+.beta. type titanium
alloy A. When annealing temperature (T2) is set within the range
"(.beta. transus-270.degree. C.)-(.beta. transus-50.degree. C.)",
the obtained microstructure becomes a structure comprising primary
.alpha. phase+metcastable .beta. phase or orthorhmbic martensite
(.alpha.") and having a very high ductility so as to have an
excellent workability making satisfactory cold rolling possible.
On. the other hand, in the low temperature range wherein the
annealing temperature (T2) does not reach (.beta.
transus-270.degree. C.), the microstructure of the alloy becomes an
age-hardened microstructure wherein the a phase is finely
precipitated in the .beta. matrix. Thus, its ductility becomes poor
so that its workability deteriorates extremely. On the contrary, in
the temperature range wherein the annealing temperature (T2) is
from (the .beta. transus-50.degree. C.) to the .beta. transus,
martensite (.alpha.') having a low ductility is produced in the
metallic microstructure after annealing and cooling so that good
workability cannot be obtained as well. When annealing is performed
at a higher temperature than the .beta. transus, .beta. grains get
coarse so that cold workability unfavorably decreases.
[0077] Based on the above-mentioned finding, a first characteristic
of the coil-rolling process of the present invention is that the
.alpha.+.beta. type alloy of the present invention is made up to
have a high ductility microstructure comprising primary a phase
+metastable p phase or orthorhombic martensite (.alpha.") by
annealing the alloy within the temperature range of (.beta.
transus-270.degree. C.)-(.beta. transus-50.degree. C.)", so that
the coil-rolling of the alloy is made easy. The time necessary for
annealing within the temperature range is not especially limited.
However, in order to make the whole of any treated titanium alloy
strip into the microstructure, the time is preferably 3 minutes or
more, and more preferably about 1 hour or more.
[0078] Conditions of coil-rolling performed after suitable
annealing as described above are not especially limited. Concerning
especially preferred conditions, however, tension is 49-392MPa, and
rolling reduction is 20% or more.
[0079] Namely, in coil-rolling, tension is applied to a material to
be rolled in its rolling directions in order to heighten rolling
efficiency, and it is effective at the time of coil-rolling the
above-mentioned .alpha.+.beta. type titanium alloy that the rolling
tension is controlled into a suitable range. The rolling tensile
strength herein means a value obtained by dividing the tension at
the time of the rolling by the sectional area of the titanium alloy
strip, and is generated by a winding reel for coils arranged before
and after a rolling roll. That is, if the rolling tension is
changed, the tension for winding coils during the rolling and after
the rolling can also be changed accordingly.
[0080] The .alpha.+.beta. type titanium alloy of the present
invention has a higher strength and lower Young s modulus than pure
titanium so that spring-back is liable to arise. Thus, if the
rolling tensile strength is low, winding of coils easily gets loose
so that production efficiency is reduced and further scratches are
easily generated between layers of the strip by the loose winding.
Thus, the yield of products tends to be reduced. For such a reason,
the rolling tension is set to 49 MPa or more, and preferably 98MPa
or more.
[0081] Incidentally, in the above-mentioned .alpha.+.beta. type
titanium alloy having a higher strength than pure titanium and
equiaxial microstructure, in particular fracture resistance is low
so that crack propagation arises easily. Thus, it is feared that
coil failure arises from a small edge crack produced in the
rolling, as a starting point. Therefore, in order not to promote
the outbreak of edge cracks and the propagation thereof, the
rolling tension is set up to 392 MPa or less, and preferably 343
MPa or less.
[0082] The rolling reduction is set up to about 20% or more and
preferably about 30% or more. This is because a rolling reduction
of less than 20% is disadvantageous for the improvement in
productivity and makes it impossible to give plastic strain
necessary and sufficient for making the alloy up to equiaxial
microstructure in the annealing step after the rolling. If the
alloy is not made up to the equiaxial microstructure, the
strength-ductility balance falls. Thus, such a case is unfavorable
for the material property of the alloy. The upper limit of the
rolling reduction varies in accordance with difference in the
property of particular alloys. The upper limit is set up to about
80% or less, and preferably about 70% or less in order to prevent
the increase in flow stress by work-hardening and the propagation
of edge cracks.
[0083] In the above-mentioned coil-rolling, in the case of some
rolling reduction, the alloy may be rolled up to a target thickness
by only one coil rolling step after annealing. If the rolling
reduction for one rolling step is excessively raised, there arises
problems, for example, the increase in flow stress by work-
hardening, and the propagation of edge cracks. Generally,
therefore, in the rolling process, coil-rolling is stepwise
performed in such a manner that plural annealing steps intervene in
the rolling process. In order to raise the strength-ductility
balance, it is effective that the .alpha.+.beta. titanium alloy is
made up to fine equiaxial microstructure. In order to realize the
equiaxial microstructure effectively, it is preferred that the
rolling step under the above-mentioned suitable conditions is
performed plural times in such a manner that an annealing step in
the .alpha.+.beta. temperature range intervenes therebetween than
rolling is performed one time at a large rolling reduction and then
annealing is performed.
[0084] The following will describe a process for producing a
cold-rolled strip, suitable for the .alpha.+.beta. type alloy of
the present invention.
[0085] The inventors have succeeded in improving elongation of in
particular the transverse direction (direction perpendicular to the
cold coil-rolling direction) along which ductility is extremely
reduced in the cold coil-rolling step, and heightening
deformability while keeping a high strength by selecting such an
annealing condition. The structural feature of the present
invention and its effect and advantage will be described
hereinafter, following details of experiments.
[0086] The inventors eagerly researched the .alpha.+.beta. type
titanium alloy making cold coil-rolling possible, according to the
present invention, in order to make clear the influence on the
ductility and the strength in the longitudinal direction (identical
to the coil-rolling direction) and the transverse direction by
annealing conditions after cold coil-rolling.
[0087] As a result, it was ascertained that as shown in attached
FIGS. 4 and 5 (both in the case of small scale), proof strength and
tensile strength are not affected very much by annealing
temperature, but concerning in particular transverse elongation
(along the transverse direction, a drop in ductility by cold
coil-rolling becomes the most serious problem), specific tendency
is exhibited in accordance with the annealing temperature. In
short, in the above-mentioned alloy system, the transverse
elongation shows a minimum value by some annealing temperature
(about 850.degree. C. in FIG. 4, and about 800.degree. C. in FIG.
5). The transverse elongation tends to rise in all annealing
temperature ranges above and below the above-mentioned
temperature.
[0088] The inventors further pursued a reason why the
above-mentioned specific tendency is exhibited, so as to make the
following fact clear.
[0089] In general, annealing after cold coil-rolling is carried out
to relieve work-hardening generated by the cold coil-rolling by
recrystallization based on heating and recover the transverse
ductility lowered mainly by the cold rolling. It is considered that
such ductility-improving effect by recrystallization is improved
still more as the annealing temperature is higher.
[0090] The alternate long and short dash line in FIG. 6
conceptually shows the relationship between annealing temperature
and ductility that is generally recognized. In the low temperature
range wherein the annealing temperature after cold rolling is about
600.degree. C. or less, the effect of improving the transverse
ductility is hardly recognized. When the annealing temperature is
raised up to about 700.degree. C. or more, the ductility is
recovered to some extent. As the annealing temperature is raised
thereafter, the recovery of the ductility advances. When the
annealing temperature is raised to not less than the .beta. transus
(T.beta.), complete recrystallization arises so that anisotropy is
cancelled. Thus, it appears that the ductility is remarkably
improved.
[0091] Concerning the .alpha.+.beta. type titanium alloy of the
present invention, however, the inventors examined the relationship
between annealing temperature and elongation after cold
coil-rolling by experimentally producing an ingot of small scale
and using a cold-rolled sample. As a result, the following were
ascertained. As shown by solid lines A and B in FIG. 6, in the
range of the annealing temperature of about 800.degree. C. or less,
both of the longitudinal elongation (solid line A) and the
transverse elongation (solid line B) are improved by the evolution
of recovery of dislocation as the temperature rises. This fact is
the same as the recognition in the prior art. When the annealing
temperature is raised to more than about 800.degree. C., the
elongations drop abruptly. When the annealing temperature is
further raised thereafter, the elongations again rise abruptly.
Such a specific tendency is exhibited. It was ascertained that such
a specific tendency is remarkably exhibited in the case of the
.alpha.+.beta. type titanium alloy of the present invention.
[0092] This tendency can be explained on the basis of phase diagram
of the .alpha.+.beta. type titanium alloy as shown in FIG. 7 and
change in the microstructure of the titanium alloy. That is, FIG. 7
is a diagram (result from small scale) showing the relationship of
the ductility of the transformed .beta. phase (i.e., the a phase)
in the titanium alloy, in the light of the phase diagram of the
.alpha.+.beta. type titanium alloy. The a phase wherein the amount
of the stabilizing elements is relatively small has a hexagonal
structure which is relatively excellent in ductility. On the other
hand, as the amount of .beta. stabilizing elements increases,
brittle hexagonal crystal is produced at some amount as a
borderline so that the ductility drops abruptly. When the amount of
.beta. stabilizing elements increases still more thereafter, an
orthorhombic crystal having a relatively high ductility is formed.
As a result, its yield stress and tensile strength drop but its
ductility tends to rise again. In summary, the ductility of the
.alpha.+.beta. type titanium alloy varies considerably, dependently
on the difference in the crystal structure resulting from the
change in the amount of .beta. stabilizing elements. It is
important to prevent the emergence of the brittle hexagonal crystal
which is generated just before the emergence of the orthorhombic
crystal by controlling the alloy composition.
[0093] As is evident from the tendency shown in FIGS. 6 and 7, the
ductility of the .alpha.+.beta. type titanium alloy after cold
coil-rolling is not simply decided by the annealing temperature for
recrystallization for relieving work-hardening. The ductility is
remarkably affected by the crystal structure of the titanium alloy
as well. As a result from a synergetic effect of these, the
following is considered. Even in the case that the annealing
temperature for recrystallization is raised as shown in FIG. 6,
when the transformed .beta. phase turns mainly into brittle
hexagonal crystal, its ductility drops abruptly. After the time
when the brittle hexagonal crystal structure turns into an ductile
orthorhombic structure having a high ductility, the ductility of
the alloy is abruptly recovered again by the evolution of
recrystallization based on annealing.
[0094] As described above, the present invention is based on the
verification of the fact that the ductility of the .alpha.+.beta.
type titanium alloy after cold coil-rolling is not simply decided
by the annealing temperature for recrystallization for relieving
work-hardening and the ductility is remarkably affected by the
crystal structure of the titanium alloy as well. In short, the
characteristic of the present invention is in that when
work-hardening is relieved by annealing the cold coil-rolled
.alpha.+.beta. type titanium alloy to raise the ductility, the
annealing temperature is controlled to avoid temperature range
causing the brittle phase production based on the emergence of the
brittle hexagonal crystal as much as possible, thereby heightening
the elongation surely to obtain excellent deformability.
[0095] At this time, as shown in region X in FIG. 7, even in the
region wherein the alloy composition of the .beta. phase causes the
emergency of the brittle hexagonal crystal at the time of heating
for annealing, if under the temperature not causing the emergency
of the brittle hexagonal crystal the material is slowly cooled (for
example, cooling in the furnace), the change in the microstructure
of the titanium alloy changes along the .beta. transus (T.beta.) to
suppress the emergency of the brittle hexagonal crystal. If its
temperature range is avoided and usual cooling (for example, air
cooling) is carried out, an annealed material having a high
performance can be obtained.
[0096] Thus, the .alpha.+.beta. type titanium alloy of the present
invention obtained by avoiding the brittle range and being annealed
as described above has a tensile strength of 900 MPa or more, and
further has an elongation of 4% or more, and exhibits an
anisotropy, that is, (longitudinal elongation)/(transverse
elongation) of about 0.4-1.0 by great recovery of the transverse
elongation. This makes it possible to obtain an annealed material
having excellent deformability in the longitudinal and transverse
directions.
[0097] Incidentally, FIG. 7 shows the relationship between
annealing temperature and elongation at the time of annealing a
cold-rolled strip comprising, for example, an .alpha.+.beta. type
titaniumalloy of Ti--4.5%Al--2%Mo--1.6%V--0.5%Fe. As shown in FIG.
7, brittle hexagonal crystal makes its appearance at about
850.degree. C. Therefore, when the cold coil-rolled titanium alloy
having this composition is annealed, it is necessary that the
annealing temperature is controlled out of the temperature which
causes the brittle hexagonal crystal, preferably withinthe
temperature range of 760-825.degree. C. or 875-T.beta..degree.
C.
[0098] Even in the same .alpha.+.beta. type titanium alloys of the
present invention, their brittle hexagonal crystal production
temperature range varies according to conditions such as
composition, production scale of coil cold-rolled strip, and
cooling rate.
[0099] For example, with attention given to the fact that coil
cold-rolled strips produced from ingots of large quantities vary in
ductility and strength in the longitudinal and transverse
directions depending on how they are annealed after coil
cold-rolling, researches were made into the effect of annealing
conditions. The results are shown in FIGS. 10 to 13. It is noted
that the annealing temperature (about 925.degree. C. in FIG. 10)
detrimental to elongation in the transverse direction tends to be
higher than that in small-scale operation, and scarcely recognized.
Although the result of strength and transverse elongation is slight
different between large scale and small scale, the phenomenon of
the change in the microstructure of titanium alloy during the
annealing process is thought to be the same. The fact that the
results in large scale differ from small scale is due to a
difference in working conditions which arises from the production
of ingots in large quantities, a difference in cooling rate of
cold-rolled strips, and so on. As the result, it was found that
annealing for the titanium alloy industrially produced in large
quantities should be carried out at temperatures in the range of
(.beta.-transus-130.degree.
C.).ltoreq.T3.ltoreq.(.beta.-transus-15.degree. C.) so that the
resulting products have good bending properties.
[0100] Therefore, at the time of carrying out the present
invention, it is preferred to make sure of this temperature range
beforehand according to the conditions such as the scale of
production of coil cold-rolled strips and then control annealing
temperature to be out of this temperature range. And the titanium
alloy industrially produced in large quantities should be carried
out at temperatures in the range of (.beta.-transus-130.degree.
C.).ltoreq.T3.ltoreq.(.beta.-transus-15.degre- e. C.). In this way,
an annealed material having a high strength and an improved
transverse elongation can be surely obtained.
[0101] At this time, the annealing must be performed at the
above-mentioned high rolling reduction for some kind of cold rolled
product. In this case, however, softening annealing is performed
one or plural times on the way of the rolling. Thus, while
work-hardening is relieved, the titanium alloy is cold rolled into
any thickness. In all case, the titanium alloy of the present
invention has a higher elongation than conventional .alpha.+.beta.
titanium alloys, so that it can be coil-rolled without the
above-mentioned pack-rolling. The alloy keeps a high strength and
simultaneously exhibits an excellent deformability by subsequent
annealing.
[0102] The thus obtained .alpha.+.beta. type titanium alloy of the
present invention can be made into coils for its excellent cold
workability, and further can easily be manufactured into any form
such as a wire, a rod or a tube regardless of the cold workability.
The present alloy has both excellent strength property and
ductility, and further has good weldability as described above, and
its HAZ after welding has a high level ductility. For this reason,
the present alloy can widely be used as applications which are
subjected to welding until they are worked into final products, for
example, a plate for a heat-exchanger, Ti golf driver head
materials, welding tubes, various wires, rods, very fine wires.
EXAMPLES
[0103] The following will specifically describe the structural
features, and effects and advantages of the present invention.
However, the present invention is not limited by the following
Examples, and can be modified within the scope consistent with the
subject manner of the present invention described above and below.
All of them are included in the technical scope of the present
invention.
Example 1
[0104] Titanium alloy ingots (60.times.130.times.260 mm) having the
compositions shown in Table 1 were produced by button melting. The
ingots were then heated to the .beta. temperature range (about
110.degree. C.), and rolled to break down into sample plates of 40
mm thickness. Subsequently, the plates were kept in the .beta.
temperature range (about 1100.degree. C.) for 30 minutes and then
air-cooled. The plates were then heated in the a +0 temperature
range (900-920.degree. C.) below the .beta. transus and hot rolled
to produce hot rolled plates of 4.5 mm thickness. Thereafter, the
plates were again annealed in the .alpha.+.beta. temperature range
(about 760.degree. C.) for 30 minutes, and then their 0.2% proof
strength, tensile strength and elongation were measured. Their test
pieces were obtained by machining the surface of the sample plates
into pieces having a gage length of 50 mm and a parallel portion
width of 12.5 mm.
[0105] Next, test pieces for cold-rolling were subjected to
shot-blasting and picking to remove oxygen-rich layers on the
surfaces. These were used as cold rolling materials to continues to
be cold rolled by a rolling reduction amount of about 0.2 mm per
pass until cracks in the plate surfaces were introduced. Thus,
their cold-reduction was measured. In order to measure their
weldability, the respective sample plates were heated at
1000.degree. C., which was not less than the .beta. transus, for 5
minutes and then air-cooled, to examine tensile property in the
state of acicular microstructure.
[0106] The results are collectively shown in Table 2.
1TABLE 1 Mo Sym- equiv- Fe bol Alloy composition (the balance: Ti)
alence equivalence A 3.5Mo-0.8Cr-4.5Al-0.3Si 3.5 0.4 B
3.5Mo-O.5Fe-0.8Cr-4.5Al-0.- 3Si 3.5 0.9 C
2.5Mo-1.6V-0.6Fe-4.5Al-0.15Si-0.04C 3.6 0.6 D
2.5Mo-1.6V-0.6Fe-4.5Al-0.45Si-0.04C 3.6 0.6 E
2.5Mo-1.6V-0.6Fe-4.5Al-1.0Si-0.04C 3.6 0.6 F
2.5Mo-1.6V-0.6Fe-4.5Al-0.3Si-0.08C 3.6 0.6 G
4.5Mo-0.8Cr-4.5Al-0.3Si 4.5 0.4 H 2.5Mo-1.6V-0.6Fe-4.5Al-0.3Si-0.1-
2C 3.6 0.6 I 2.5Mo-1.6V-0.6Fe-4.0Al-0.3Si-0.04C 3.6 0.6 J
2.5Mo-1.6V-0.6Fe-5.0Al-0.3Si-0.04C 3.6 0.6 K
3.5Mo-0.5Fe-0.8Cr-4.5Al-0.3Si-0.05C 3.5 0.4 L
3.5Mo-0.5Fe-0.8Cr-4.5Al-0.3Si-0.1C 3.5 0.4 M
2Mo-1.6V-0.5Fe-4.5A1-0.3Si-0.03C 3.1 0.5 N 1Mo-1.6V-0.5Fe-4.5A1-0.-
3Si-0.03C 2.1 0.5 O 3.5Mo-0.8Cr-4.5A1 3.5 0.4 P
3.5Mo-0.5Fe-0.8Cr-4.5A1 3.5 0.5 Q 4.5Mo-0.8Cr-4.5A1 4.5 0.4 R
2.5Mo-1.6V-0.6Fe-4.5A1-0.04C 3.6 0.6 S 3.5Mo-0.5Fe-0.8Cr-3.0A1-0-
.3Si 3 0.9 T 2.5Mo-0.5Fe-0.8Cr-3.0Al-0.3Si 2.5 0.9 U
3.0Mo-0.5Fe-0.8Cr-3.0Al-0.3Si-0.05C 3.9 0.9 V
2.5Mo-1.6V-0.6Fe-4.5A1-1.5Si-0.04C 3.6 0.6 W
2.0Mo-1.6V-0.6Fe-6.5A1-0.3Si-0.04C 3.1 0.6 X
0.8Mo-1.6V-0.5Fe-4.5A1-0.3Si-0.03C 1.9 0.5 Y
3.5Mo-1.6V-0.5Fe-4.5A1-0.3Si-0.03C 4.6 0.5 Z
2Mo-1.6V-2.5Fe-4.5A1-0.3Si-0.03C 3.1 2.5
[0107]
2 TABLE 2 Tensile properties after .beta. annealing (Acicular,
corresponding to HAZ after welding) Tensile properties after
.alpha. + .beta. annealing 0.2% Proof Tension 6.9 x (YS- 0.2% Proof
Tension Cold reduction strength strength Elongation 835) + 245
.times. (El- strength strength Elongation Being made Symbol (MPa)
(MPa) (%) 8.2) (MPa) (MPa) (%) into a coil Note A 835 1010 8.2 0
882 937 15.5 .largecircle. (possible) B 936 1112 7.7 763 875 941
15.7 .largecircle. C 1069 1250 3.8 538 822 900 19.2 .largecircle. D
1121 1342 4.3 1019 885 963 17.8 .largecircle. E 1191 1356 1.2 739
933 1061 12.8 .largecircle. F 1087 1298 4.5 831 893 959 20.7
.largecircle. C 994 1156 5.8 507 891 946 15.0 .largecircle. H 992
1221 3.8 4 925 984 16.9 .largecircle. I 1032 1223 6.2 869 815 912
17.9 .largecircle. J 1164 1365 2.9 973 932 999 19.4 .largecircle. K
1044 1215 3.6 313 940 992 19.0 .largecircle. L 1080 1298 1.3 0 1085
1131 18.4 .largecircle. M 827 907 8.5 19 857 916 192 .largecircle.
N 814 885 9.1 78 821 894 19.5 .largecircle. O 775 974 10.1 53 785
861 22.6 .largecircle. Insufficient strength P 880 1024 6.3 -155
795 874 15.6 .largecircle. Insufficient strength and bad
weldability Q 899 1039 4.9 -369 767 835 21.2 .largecircle.
Insufficient strength and bad weldability R 1036 1249 1.3 -305 810
889 17.7 .largecircle. Insufficient strength and bad weldability S
751 920 11.5 227 652 781 16.5 .largecircle. Insufficient strength T
734 899 13.2 528 703 810 16.7 .largecircle. Insufficient strength U
1018 1238 3 -10 767 856 16.3 .largecircle. Insufficient strength
and bad weldability V 1223 1373 0 5 791 983 1103 8.1 X (impossible)
Bad cold-rollability W 1219 1429 0.3 715 975 1115 9.2 X Bad
cold-rollability X 797 858 10.5 300 799 868 19.5 .largecircle.
Insufficient strength Y 1081 1229 0.5 -190 1147 1179 18.9
.largecircle. Bad weldability Z 1099 1278 0 -190 1127 1229 17.4
.largecircle. Bad weldability
[0108] FIG. 1 shows, as a graph, the relationship between the 0.2%
proof strength and the elongation after .beta. annealing, which
corresponds to the physical property in HAZ after welding, among
the experimental data shown in Table 1.
[0109] In this graph, solid line Y is a line connecting the
relationship points between 0.2% proof strength and elongation of
other than comparative samples wherein their cold reduction was
represented by "x" (limit cold reduction: less than 40%). Broken
line X represents a relationship formula represented by
6.9.times.(YS-835)+245.times.(El-8.2)- .
[0110] As is evident from this graph, the solid line Y and the
broken line X cross each other at a point of a 0.2% proof strength
of 813 MPa. The inclination of the solid line Y (comparative
samples) in the area having a higher proof strength than this proof
strength is steeper than that of the broken line X. This graph
proves that in the high proof strength area of the comparative
samples, this elongation drops abruptly as the proof strength
rises. On the other hand, in Examples of the present invention all
of the relationship points between the proof strength and the
elongation are positioned in the right and upper area relative to
the broken line X. The drop in the elongation with the rise in the
proof strength is relatively small. Thus, it can be ascertained
that the samples of Examples had high strength and ductility.
[0111] FIG. 8 is a graph showing an arranged relationship between
the 0.2% proof strength and the elongation after .alpha.+.beta.
annealing. It can be understood from this graph that all of the
comparative samples do not reach a proof strength of 813 MPa but
all of the samples of Examples exhibit a proof strength more than
this value, and the material of the present invention has a high
strength and an excellent ductility.
Example 2
[0112] Titanium alloys having the compositions shown in Table 3
were produced in a melting state by vacuum arc melting and made
into ingots (their diameter: 100 mm). The ingots were then heated
to the .beta. temperature range (about 1000-1050.degree. C.), and
rolled to break down into sample plates of 15 mm thickness.
Subsequently, the plates were kept in the 13 temperature range
(about 1000-1050.degree. C.) for 30 minutes and then air-cooled.
The plates were then heated in the a +13 temperature range
(850.degree. C.), which was not more than the .beta. transus, and
hot rolled to produce hot rolled plates of 5.7 mm thickness.
Thereafter, the plates were again annealed in the .alpha.+.beta.
temperature range (630-890.degree. C.) for 5 minutes. Next, they
were subjected to shot-blasting and pickling to remove oxygen-rich
layers on the surfaces. These were used as cold rolling materials.
In the cold coil-rolling, the rolling reduction amount was 0.2 mm
per pass. In the rolling, tension was applied along the rolling
direction to roll the plates up to a predetermined rolling
reduction. After the rolling, the depth of edge cracks in the
plates was measured. Thereafter, the plates were annealed in the
.alpha.+.beta. temperature range and then were subjected to optical
microstructure observation of their cross sections.
[0113] The results are shown in Table 4.
[0114] The difference in sectional microstructures was observed
between the plates which were rolled one time up to a predetermined
thickness and then annealed, and the plates which were rolled three
times up to a predetermined thickness in a manner that annealing
intervened therebetween on the way of the rolling process and then
annealed. The results are shown in Table 5.
3TABLE 3 .beta. Al Mo V Fe Si O Ti transus 4.5 2.0 1.5 0.5 0.3 0.16
Balance 963.degree. C. (mass %)
[0115]
4TABLE 4 Rolling conditions Results Annealing Edge cracks Experi-
Rolling Rolling temperature .circleincircle.: less than 5 mm
Structure Total judgement ment tension reduction before
.largecircle.: 5 mm-10 mm after .largecircle.: Suitable No. (MPa)
(%) rolling X: 10 mm or more annealing X: Unsuitable 1 147 50 760
.circleincircle. Equiaxial .largecircle. 2 294 50 760
.circleincircle. Equiaxial .largecircle. 3 98 50 760
.circleincircle. Equiaxial .largecircle. 4 343 50 760
.circleincircle. Equinxinl .largecircle. 5 294 30 760
.circleincircle. Equiaxial .largecircle. 6 294 70 760
.circleincircle. Equiaxial .largecircle. 7 294 50 820
.circleincircle. Equiaxial .largecircle. 8 294 50 700
.circleincircle. Equiaxial .largecircle. 9 294 40* 630 X Equiaxial
X 10 294 30* 890 X Equiaxial X 11 441 50 760 X Equiaxial X 12 294
10 760 .circleincircle. Non- X equiaxial 13 294 85 760 X Equiaxial
X *Rolling load exceeded for a 50% rolling reduction of a target.
Thus, the rolling was stopped on the way.
[0116]
5 TABLE 5 Steps Total Structure Experiment Cold .alpha. + .beta.
Cold .alpha. + .beta. Cold .alpha. + .beta. rolling after the final
No. rolling 1 annealing Rolling 2 annealing Rolling 3 annealing
ratio annealing 14 40% Performed 40% Performed 40% Performed 78.5%
Fine equiaxial microstructure 15 80% Performed -- -- -- -- 80%
Partial equiaxial microstructure
[0117] The following can be understood from Tables 3-5.
[0118] Experiments Nos. 1-8: Examples satisfying all of the
requirements defined in the present invention. The microstructure
of the annealing was uniformly equiaxial and had a few edge cracks,
so as to be sufficiently suitable for practical use of
coil-rolling.
[0119] Experiments Nos. 9 and 10: Comparative Examples wherein the
temperature of the annealing before the rolling was out of the
defined range. Edge cracks were generated before the arrival to a
50% rolling reduction which was a rolling target. Thus, the rolling
was stopped when the rolling reduction was 40% or 30%. However,
considerably large edge cracks were observed. It is difficult that
the Comparative Examples were made practicable.
[0120] Experiment No. 11: Reference Example wherein a tension at
the time of the rolling was raised up to 45%. The tension was too
high, so that edge cracks were liable to be generated.
[0121] Experiment No. 12: Reference Example wherein the rolling
ratio at the time of the rolling was set to a low value. The
coil-rolling was able to be performed without any generation of
large edge cracks. However, a part of the microstructure after the
annealing became non-equiaxial. The strength-elongation balance was
bad.
[0122] Experiment No. 13: Reference Example wherein the rolling
reduction at the time of the rolling was raised up to 85%. Because
the rolling reduction was excessively high, large edge cracks were
observed.
[0123] Experiment No. 14: Example which was coil-rolled 3 times,
the rolling reduction per rolling being 40%, in a manner that
annealing intervened therebetween 2 times on the way. The
microstructure after the final annealing was fine equiaxial, and a
good coil which had no edge cracks and a good strength-elongation
balance was obtained.
[0124] Experiment No. 15: Example in which substantially the same
rolling as in Experiment No. 14 was performed by a single rolling
step without any annealing on the way. A part of the microstructure
after the annealing became non-equiaxial. The strength-elongation
balance was slightly bad.
[0125] Experiment 3-1
[0126] A Ti alloy ingot (80 mm.sup.T.times.200 mm.sup.W.times.300
mm.sup.L) of Ti--2%Mo--1.6%V--0.5%Fe--4.5%Al--0.3%Si--0.03% C was
produced by induction-skull melting, heated in the .beta.
temperature range (about 1100.degree. C.) and then rolled to break
down into sample plates of 40 mm thickness. Subsequently, the
plates were kept in the .beta. temperature range (about
1100.degree. C.) for 30 minutes and then air-cooled. The plates
were then hot rolled in the .alpha.+.beta. temperature range
(900-920.degree. C.), which was lower than the .beta. transus to
produce hot rolled plates of 4.5 mm thickness.
[0127] Next, the plates were annealed at 760.degree. C. for 30
minutes, and then they were subjected to shot-blasting and pickling
to prepare cold rolling materials. These were subjected to the
treatment of [40% cold rolling+annealing at 760.degree. C. for 5
minutes] two times to perform cold rolling up to a rolling
reduction of 40%. Thereafter, annealing was performed under
conditions shown in Table 6. The respective annealed products were
pickled to remove oxygen rich layers on their surfaces. Their
transverse and longitudinal 0.2% proof strength, tensile strength,
and elongations were measured. The result are shown in Table 6 and
FIG. 4.
6TABLE 6 Ti-2Mo-1.6V-0.5Fe-4.5Al-0.3Si-0.03C Meas- Annealing ured
0.2% Proof Tensile tempera- direc- strength strength ture (.degree.
C.) tion (MPa) (MPa) Elongation (%) Example 760 L 982 1096 10.4
Com- 850 L 991 1202 7.8 parative Example Example 900 L 1028 1239
7.2 Example 760 T 1073 1144 4.6 Example 800 T 1082 1128 4.6 Example
825 T 1014 1087 5.6 Com- 850 T 1082 1198 2 parative Example Example
900 T 1085 1164 5.8 Example 925 T 1095 1182 7.8 Example 950 T 1027
1143 10.6
[0128] As is clear from Table 6 and FIG. 4, it was ascertained that
in the .alpha.+.beta. type titanium alloy of the component systems
used in the present invention the transverse elongation (the
elongation in the direction perpendicular to the rolling direction)
decreased remarkably by the production of brittle hexagonal crystal
in the annealing temperature range of about 850.degree. C. Thus, it
can be understood that if the alloy was annealed in the temperature
range of 750-830.degree. C. or 900-950.degree. C., out of the
above-mentioned annealing temperature range, an annealed product
was obtained which kept high tensile strength and 0.2% proof
strength, and had an excellent elongation.
[0129] Experiment 3-2
[0130] A Ti alloy ingot (80 MM.sup.T.times.200 mm.sup.W.times.300
mm.sup.L) of Ti--3.5%Mo--0.5%Fe--4.5%Al--0.3%Si was produced by
induction-skull melting, and was heated in the .beta. temperature
range (about 1100 .degree. C. ) for 30 minutes and then rolled to
break down into sample plates of 40 mm thickness. Subsequently, the
plates were kept in the p temperature range (about 1100.degree. C)
and then air-cooled. The plates were then hot rolled in the
.alpha.+.beta. temperature range (900-920.degree. C.), which was
lower than the .beta. transus to produce hot rolled plates of 4.5
mm thickness.
[0131] Next, the plates were annealed at 760.degree. C for 30
minutes, and then they were subjected to shot-blasting and pickling
to prepare cold rolling materials. These were subjected to the
treatment of [40% cold rolling +annealing at 760.degree. C. for 5
minutes] two times to perform cold rolling up to a rolling
reduction of 40%. Thereafter, annealing was performed under
conditions shown in Table 1. The respective annealed products were
pickled to remove oxygen rich layers on their surfaces. Their
transverse and longitudinal 0.2% proof strength, tensile strength,
and elongations were measured. The result are shown in Table 7 and
FIG. 5.
7TABLE 7 Ti-3.5Mo-0.5Fe-4.5Al-0.3Si Annealing 0.2% Proof Tensile
temperature Measured strength strength Elongation (.degree. C.)
direction (MPa) (MPa) (%) Example 760 L 982 1096 10.4 Example 850 L
906 1125 7.8 Example 900 L 1051 1244 7.2 Example 760 T 1092 1142
5.2 Com- 800 T 1007 1059 2.4 parative Example Example 825 T 986
1077 5.6 Example 850 T 985 1103 6.4 Example 900 T 1058 1249 6
[0132] As is clear from Table 7 and FIG. 5, it was ascertained that
in the .alpha.+.beta. type titanium alloy of the component systems
used in the present invention the transverse elongation (the
elongation in the direction perpendicular to the rolling direction)
decreased remarkably by the production of brittle hexagonal crystal
in the annealing temperature range of about 800.degree. C. Thus, it
can be understood that if the alloy was annealed in the temperature
range of 760.degree. C. or lower, or 820-950.degree. C., out of the
above-mentioned annealing temperature range, an annealed product
was obtained which kept high tensile strength and 0.2% proof
strength, and had an excellent elongation.
Example 4
[0133] A 5-ton ingot of titanium alloy having an aimed composition
of Ti--2Mo--1.6V--0.5Fe--4.5Al--0.3Si--0.03C was prepared by the
VAR process. The ingot was made into a 140-mm thick slab by forging
and rolling in the beta phase. The slab was heated at
930.+-.20.degree. C. and then rolled down to a 4-mm thick sheet.
The sheet underwent annealing and cold rolling repeatedly to give a
1.2-mm thick cold-rolled strip. The ingot has the chemical
composition (in its top and bottom) as shown in Table 8, and the
sheet has the tensile properties in the rolling direction (in the
longitudinal direction) as shown in Table 8. Heat numbers in Table
8 having the same number of first four figures show that the
titanium alloy is charged and made into the ingot at the same
time.
8TABLE 8 Heat No. Position Direction Mo V Fe Al Si C O N AT4790-1 T
L 1.93 1.82 0.52 4.62 0.27 0.038 0.163 0.0044 AT4790-1 B L 2.09
1.33 0.48 4.07 0.27 0.03 0.156 0.0053 AT4790-2 T L 1.93 1.82 0.52
4.62 0.27 0.038 0.163 0.0044 AT4790-2 B L 2.09 1.33 0.48 4.07 0.27
0.03 0.156 0.0053 AT4962-3 T L 2.12 1.79 0.6 4.61 0.29 0.032 0.11
0.0048 AT4962-3 B L 2.25 1.41 0.52 4.29 0.3 0.028 0.13 0.0047
AT4962-4 T L 2.12 1.79 0.6 4.61 0.29 0.032 0.11 0.0048 AT4962-4 B L
2.25 1.41 0.52 4.29 0.3 0.028 0.13 0.0047 AT5038-3 T L 1.87 1.74
0.5 4.55 0.26 0.033 0.137 0.0046 AT5038-3 B L 2.12 1.62 0.51 4.63
0.27 0.034 0.141 0.0056 AT5038-4 T L 1.87 1.74 0.5 4.55 0.26 0.033
0.137 0.0046 AT5038-4 B L 2.12 1.62 0.51 4.63 0.27 0.034 0.141
0.0056 AT5066-1 T L 1.9 1.77 0.49 4.51 0.26 0.03 0.141 0.0041
AT5066-1 B L 2.12 1.52 0.48 4.38 0.26 0.03 0.143 0.0035 AT5066-2 T
L 1.9 1.77 0.49 4.51 0.26 0.03 0.141 0.0041 AT5066-2 B L 2.12 1.52
0.48 4.38 0.26 0.03 0.143 0.0035 AT5199 T L 1.87 1.64 0.51 4.35
0.27 0.032 0.098 0.0048 AT5199 B L 2.11 1.51 0.46 4.47 0.23 0.032
0.119 0.0066 0.2% Proof Tensile Mo eq + Strength/ Strength/
Elongation/ Mo eq + 2.5 .times. Fe eq + Heat No. H MPa MPa % 2.5
.times. Fe eq 0 .times. 40 AT4790-1 0.0053 916 994 11 4.44 10.98
AT4790-1 0.0056 955 1030 10 4.18 10.42 AT4790-2 0.0053 916 994 11
4.44 10.96 AT4790-2 0.0056 955 1030 10 4.18 10.42 AT4962-3 0.0062
918 1000 11 4.81 9.21 AT4962-3 0.006 897 981 9 4.49 9.69 AT4962-4
0.006 923 1011 10 4.81 9.21 AT4962-4 0.006 923 1010 10 4.49 9.69
AT5038-3 0.0073 896 973 10 4.28 9.76 AT5038-3 0.0076 900 977 10
4.48 10.12 AT5038-4 0.0074 884 966 10 4.28 9.76 AT5038-4 0.0086 898
977 11 4.48 10.12 AT5066-1 0.0079 907 1001 10 4.31 9.95 AT5066-1
0.0079 915 998 11 4.33 10.05 AT5066-2 0.0079 932 1007 9 4.31 9.95
AT5066-2 0.0079 916 998 9 4.33 10.05 AT5199 0.0051 837 929 10 4.24
8.16 AT5199 0.0048 868 955 10 4.27 9.03
[0134] In FIG. 9, the tensile strength of the cold-rolled strip in
this example is plotted against the amount (%) of
[Mo-equivalence+2.5.times.Fe- -equivalence+40.times.O %]. A good
correlation between them is noticed. It is apparent that the
tensile strength exceeds 900 MPa when the amount of
[Mo-equivalence+2.5.times.Fe-equivalence+40.times.O %] exceeds
7.0%.
[0135] Also, the sample of heat No. AT4790 in this example was
examined to see how ductility and strength in the longitudinal and
transverse directions are affected differently depending the
conditions under which annealing is carried out after coil
cold-rolling. To this end, a 1.2-mm thick cold-rolled strip was
produced by the method mentioned above, and then it was tested for
elongation in the longitudinal and transverse directions, proof
stress (at 0.2% permanent set), and tensile strength. The results
are shown in Tables 10 and 11. It was found that those ingots
produced in large quantities as in this example yield strips which
are slightly low in elongation in the transverse direction
(perpendicular to the rolling direction) if annealing is carried
out at about 925.degree. C. after coil cold-rolling. In other
words, it is apparent that in the case of annealing coil
cold-rolled strips in large scale, the annealing temperature
leading to a decrease in elongation in the transverse direction is
somewhat higher than in small scale and the annealing at an
increased temperature decreases elongation only a little in the
transverse direction.
[0136] The fact that the results in this example differ from those
in Example 3 is due to a slight difference in composition which
arises from the production of ingots in large quantities in this
example, a difference in scale of the production of coil
cold-rolled strips, and a difference in thickness (or cooling rate)
of cold-rolled strips.
[0137] Subsequently, the sample of heat No. AT4790, which is a
1.2-mm thick cold-rolled coil, was subjected to annealing at
different temperatures and ensuring bending test. The sample in
bending test was evaluated in terms of the minimum value of R/t,
where R is the bending radius and t is the thickness, (which is
called the minimum radius). The results are shown in FIG. 12. It is
noted that the sample (L in FIG. 12) which was bent such that the
bending axis is parallel to the rolling direction of the sample
subjected to annealing at temperatures in the range of 850.degree.
C. to 950.degree. C. has a small minimum radius (which implies good
bending properties). It is also noted that the ingot of heat No.
AT4790 has the .beta.-transus of 973.degree. C. at its top and
978.degree. C. at its bottom and hence it exhibits good bending
properties if it undergoes final annealing at a temperature between
(.beta.-transus-130.degree. C.) and (.beta.-transus-15.degree. C.).
The preferred annealing temperature for the alloy in this example
is 850-963.degree. C.
Example 5
[0138] A 20-mm thick slab was prepared by button arc melting from a
titanium alloy having a base composition of
Ti-2Mo-1.6V-0.5Fe-4.5Al-0.3Si- -0.03C and additionally containing
Ru in an amount of 0.05% or 0.08%. The ingot was heated at
1000.degree. C. for 30 minutes and then hot-rolled to give a 10-mm
thick plate. The plate was heated at 930.degree. C. and then
hot-rolled to give a 4-mm thick sheet. After annealing and
descaling, the sheet was cold-rolled until its thickness was
halved. It was found that cold-rolling was accomplished
successfully as in the case of titanium alloy containing no Ru.
After cold-rolling, the strip was annealed at 800.degree. C. for 10
minutes. The thus obtained samples were tested for tensile strength
and elongation in the longitudinal and transverse directions (twice
each). The results are shown in Table 9.
9TABLE 9 0.2% proof Ru Tensile strength, Tensile strength, content,
% direction MPa MPa Elongation, % 0.05 L 908 989 16 0.05 L 920 987
18 0.05 T -- 1042 13 0.05 T -- 1045 11 0.08 L 917 979 16 0.08 L 913
991 17 0.08 T -- 1017 12 0.08 T -- 1008 11
[0139] The samples shown in Table 9 (containing 0.05% Ru,
containing 0.08% Ru, and not containing Ru) were tested for
corrosion resistance, with pure titanium being a control.
[0140] First, each sample was immersed in an HCl solution to test
for ability to keep the passive state. Evaluation was made in terms
of the concentration of HCl solution at which the sample loses its
passive state. The results are shown in FIG. 13a. It is noted that
the sample containing 0.05% Ru or 0.08% Ru keeps the passive state
in the same way as pure titanium.
[0141] Then, the samples were tested for corrosion speed by
immersion in an aqueous solution containing 1 mol/L of NaCl and 1
mol/L of HCl. The corrosion speed in this aqueous solution was
compared with that in boiling water. The results are shown in FIG.
13b. It is noted that the corrosion speed of the Ru-containing
sample is about one half of that of pure titanium.
[0142] The samples were also tested for crevice corrosion (by the
multi-crevice method) in order to find the rate of corrosion
occurrence. After polishing with emery (#400) in wet process and
degreasing, specimens were immersed in a boiling aqueous solution
containing 20% of NaCl for 1 week. The number of incidences of
crevice corrosion that occurred was counted, and the ratio of that
number to the number of crevices was calculated. The results are
shown in FIG. 13c. It is noted that the Ru-containing samples are
superior to pure titanium in resistance to crevice corrosion.
Example 6
[0143] A 20-kg ingot was prepared from a titanium alloy having the
composition of Ti--2Mo--1.6V--0.5Fe--4.5Al--0.3Si--0.03C (with an
actually measured .beta.-transus of 963.degree. C.) or
Ti--2Mo--1.6V--0.5Fe--4.5Al--0.3Si--0.06C (with an
actuallymeasured.beta.-transus of 987.degree. C.). For comparison,
an ingot was prepared in the same way as above from a C-free
titanium alloy having the composition of
Ti--2Mo--1.6V--0.5Fe--4.5Al--0.3Si (with an actually measured
.beta.-transus of 940.degree. C.). The ingot was made into 36-mm
thick slab. The slab was made into a 4-mm thick sheet by
hot-rolling with single heating at a different temperature of
910.degree. C., 930.degree. C., or 950.degree. C. The rolled sheet
was examined for edge cracking. It was found that there is no
significant difference among the samples in occurrence of edge
cracking despite the fact that they differ in .beta.-transus by
23.degree. C. (equivalent to about 770.degree. C./%C) because they
differ in C content by 0.03%. In the case of heating at 910.degree.
C., edge cracking (about 3-5 mm deep) occurred; however, in the
case of heating at 930.degree. C. or 950.degree. C., no edge
cracking occurred. A probable reason for this is that C is an
interstitial element and hence it does not contribute so much to
solid-solution strengthening at high temperatures, with the result
that the a-phase keeps its high ductility even though it is heated
to a high temperature. Namely, it is noted that the C-free alloy is
liable to edge cracking at the temperature (910.degree. C.), which
is lower than (.beta.-transus-20.degree. C.=920.degree. C.),
whereas it is immune to edge cracking at the high temperature
(930.degree. C. or 950.degree. C.). It is concluded that the
C-containing sample contains more a-phase than the C-free sample at
the same temperature which is higher than
(.beta.-transus-20.degree. C.), but the increased a-phase does not
greatly affect the occurrence of edge cracking.
[0144] As described above, the present invention has a basic
composition wherein the contained percentages of the isomorphous
.alpha. stabilizing element and the eutectic .beta. stabilizing
element are defined, and a specified amount of Si, or additionally
a small amount of C or 0 is incorporated into the basic
composition. Thus, the present invention has a strength property
which is not inferior to Ti--6Al--4V alloys which have been most
widely used, and has remarkably raised cold workability, which is
insufficient in the conventional alloys, to make coil-rolling
possible. Moreover, the present invention can provide a titanium
alloy having all of remarkably improved strength and ductility in
HAZ after welding, and high workability, strength and
weldability.
[0145] Therefore, the titanium alloy of the present invention can
be used in various applications for its characteristics. The
present invention can be very useful used as, for example plates
for heat-exchangers by using, in particular, excellent
corrosion-resistance, lightness, heat conductivity and
cold-formability.
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