U.S. patent application number 14/232198 was filed with the patent office on 2014-05-29 for method for manufacturing alloy containing transition metal carbide, tungsten alloy containing transition metal carbide, and alloy manufactured by said method.
This patent application is currently assigned to TOHOKU UNIVERSITY. The applicant listed for this patent is Hideo Arakawa, Hiroaki Kurishita, Satoru Matsuo. Invention is credited to Hideo Arakawa, Hiroaki Kurishita, Satoru Matsuo.
Application Number | 20140147327 14/232198 |
Document ID | / |
Family ID | 47629237 |
Filed Date | 2014-05-29 |
United States Patent
Application |
20140147327 |
Kind Code |
A1 |
Kurishita; Hiroaki ; et
al. |
May 29, 2014 |
METHOD FOR MANUFACTURING ALLOY CONTAINING TRANSITION METAL CARBIDE,
TUNGSTEN ALLOY CONTAINING TRANSITION METAL CARBIDE, AND ALLOY
MANUFACTURED BY SAID METHOD
Abstract
The present invention relates to the development of an alloy
material with significantly improved low-temperature brittleness,
recrystallization brittleness, and irradiation brittleness by the
introduction of a recrystallization microstructure into an alloy,
particularly a tungsten material, to significantly strengthen a
weak grain boundary of the recrystallization microstructure. The
present invention comprises the steps of: mechanically alloying at
least one species selected from a group-IVA, VA, or VIA transition
metal carbide and a metallic raw material; sintering base powders
obtained through the mechanically alloying step, by using a hot
isostatic press; and performing plastic deformation of at least 60%
on the alloy obtained through the sintering step, at a strain rate
between 10.sup.-5 s.sup.-1 and 10.sup.-2 S.sup.-1 and at a
temperature between 500.degree. C. and 2,000.degree. C. It is
therefore possible to obtain an alloy material with significantly
improved low-temperature brittleness, recrystallization
brittleness, and irradiation brittleness.
Inventors: |
Kurishita; Hiroaki; (Miyagi,
JP) ; Arakawa; Hideo; (Miyagi, JP) ; Matsuo;
Satoru; (Miyagi, JP) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
Kurishita; Hiroaki
Arakawa; Hideo
Matsuo; Satoru |
Miyagi
Miyagi
Miyagi |
|
JP
JP
JP |
|
|
Assignee: |
TOHOKU UNIVERSITY
Miyagi
JP
|
Family ID: |
47629237 |
Appl. No.: |
14/232198 |
Filed: |
July 27, 2012 |
PCT Filed: |
July 27, 2012 |
PCT NO: |
PCT/JP2012/069190 |
371 Date: |
January 10, 2014 |
Current U.S.
Class: |
419/18 ; 419/17;
75/236; 75/240 |
Current CPC
Class: |
C22C 1/05 20130101; C22F
1/18 20130101; B22F 2998/10 20130101; B22F 3/15 20130101; C22C
32/0052 20130101; B22F 3/15 20130101; B22F 2998/10 20130101; C22C
1/1084 20130101; C22C 27/04 20130101 |
Class at
Publication: |
419/18 ; 419/17;
75/236; 75/240 |
International
Class: |
C22C 1/05 20060101
C22C001/05; C22C 27/04 20060101 C22C027/04 |
Foreign Application Data
Date |
Code |
Application Number |
Jul 29, 2011 |
JP |
2011-166630 |
Claims
1. A method for manufacturing an alloy, characterized by having a
step for mechanical alloying of a metal raw material and at least
one selected from carbides of group IVA, VA, or VIA transition
metals, a step for sintering the raw material powder obtained in
said mechanical alloying step using hot isostatic pressing, and a
step for subjecting the alloy obtained in said sintering step to
grain boundary sliding based plastic deformation of 60% or greater
at 500 to 2000.degree. C. and a strain rate of 10.sup.-5 s.sup.-1
to 10.sup.-2 s.sup.-1.
2. The method for manufacturing an alloy according to claim 1,
characterized by having a step in which said transition metal
carbide and the metal raw material are degassed by heating prior to
said mechanical alloying step.
3. A tungsten alloy comprising 0.25 to 5 mass % of at least one
type selected from carbides of a group IVA, VA, or VIA transition
metals, the tungsten alloy characterized in that the oxygen content
is 950 ppm by mass or less, the nitrogen content is 60 ppm by mass
or less, 80% or more of an observed cross sectional area in the
tungsten phase is recrystallized equiaxed grains with grain
diameters of 0.05 to 10 .mu.m, the ductile-brittle transition
temperature determined by three-point flexure is 500K or less, and
plastic deformation due to grain boundary sliding is possible at or
above this temperature.
4. The tungsten alloy according to claim 3, characterized in that
90% or greater of the azimuths of the carbide present in the
tungsten alloy structure and the azimuths of the tungsten matrix
are in the (Kurdjumov-Sachs) azimuth relationship: {111} W//{110}
transition metal carbide <110> W//<111> transition
metal carbide.
5. The tungsten alloy according to claim 3, characterized in that
the full width at half maximum for reflection of the (220)
diffraction planes is 3.degree. or less as determined by X-ray
diffraction, or that there are 50 or fewer dislocations within
crystal grains as determined by transmission electron
microscopy.
6. The tungsten alloy according to claim 3, characterized in that
the maximum bend strength determined by three-point flexure is 1470
MPa or greater.
7. The alloy that is manufactured by the manufacturing method
according to claim 1.
8. The tungsten alloy according to claim 4, characterized in that
the full width at half maximum for reflection of the (220)
diffraction planes is 3.degree. or less as determined by X-ray
diffraction, or that there are 50 or fewer dislocations within
crystal grains as determined by transmission electron
microscopy.
9. The tungsten alloy according to claim 4, characterized in that
the maximum bend strength determined by three-point flexure is 1470
MPa or greater.
10. The tungsten alloy according to claim 5, characterized in that
the maximum bend strength determined by three-point flexure is 1470
MPa or greater.
11. The alloy that is manufactured by the manufacturing method
according to claim 2.
Description
TECHNICAL FIELD
[0001] The present invention relates to a method for manufacturing
an alloy containing transition metal carbide, a tungsten alloy
containing transition metal carbide, and an alloy manufactured by
said method. In particular, the present invention relates to a
method for manufacturing an alloy that manifests superplasticity
due to grain boundary sliding when the alloy is made to undergo
superplastic deformation, that exhibits high recrystallization
fracture strength, that has little decrease in strength or
ductility, even when heated to high temperatures due to its
recrystallized structure, and which has dramatically remedied
low-temperature embrittlement, recrystallization embrittlement, and
neutron irradiation embrittlement, as well as an alloy that has
been manufactured by this manufacturing method, in particular, a
tungsten alloy.
BACKGROUND ART
[0002] Tungsten and tungsten alloys have melting points of as high
as 3410.degree. C. which are the highest of any metal. These
materials thus provide a great many advantages that are
unparalleled by other metals. However, the materials have not been
used for structure due to an inability to resolve problems with
persisting embrittlement (low-temperature embrittlement,
recrystallization embrittlement, and irradiation embrittlement),
which has hampered practical use of these materials as
high-temperature structural materials in extreme environments.
[0003] These embrittlement phenomena all result from very weak
crystal grain boundaries and a tendency for fracture to originate
from the grain boundaries which are referred to as "grain boundary
embrittlement". The cause of grain boundary embrittlement is that
tungsten is a metal having an extremely high degree of covalent
bond character, and the grain boundaries are substantially weaker
(tend to fracture) due to their high energies. Additionally,
interstitial gas elements contained in air such as nitrogen and
oxygen have extremely low solubility in tungsten, and thus tend to
segregate and precipitate at the grain boundaries, which further
weakens the grain boundaries and promotes embrittlement.
[0004] As shown in FIG. 1(a), with common metals, almost the entire
temperature range is the ductile temperature region, because of the
fact that plastic deformation (permanent set) occurs prior to
break. On the other hand, as shown in FIG. 1(b), because tungsten
has covalent bonding in which the directionality of the interatomic
bonds is extremely high, the grain boundaries are substantially
weaker, ductile-brittle transition occurs, and the ductile-brittle
transition temperature ("DBTT" below) is also high. This phenomenon
thus becomes extreme with lower temperatures at which there is a
precipitous increase in the Peierls stress (yield strength)
required for screw dislocations to glide in tungsten
(low-temperature embrittlement), and the phenomenon is even more
pronounced with recrystallized structure in which extremely weak
grain boundaries are formed (recrystallization embrittlement).
Moreover, when lattice defects are introduced by high-energy
particle irradiation using neutrons or the like, such irradiation
induced defects accumulate inside the crystal grains or at the
grain boundaries and impede dislocation slip, resulting in the
promotion of grain boundary embrittlement (irradiation
embrittlement).
[0005] Consequently, in order to simultaneously remedy
low-temperature embrittlement, recrystallization embrittlement, and
irradiation embrittlement, it is necessary to introduce
recrystallized microstructures containing high densities of sink
sites (sinks; crystal grain boundary or dispersed particles) that
can permit the material to tolerate irradiation induced defects and
to convert the weak grain boundaries in the recrystallized
microstructure to strong grain boundaries that resist fracture.
[0006] The inventors of the present invention, in order to resolve
problems with low temperature embrittlement and embrittlement of
tungsten due to neutron irradiation and recrystallization, carried
out manufacture of W-TiC having ultrafine crystal grains by a hot
isostatic pressing (HIP) method and by mechanical alloying (MA) in
Ar and H.sub.2 atmospheres. Effects such as an appreciable
improvement in room temperature toughness were found to occur, and
these results were published (non-patent documents 1, 2). However,
the tungsten materials manufactured by the above methods still were
not adequate for practical use.
[0007] On the other hand, a known method for improving toughness
and the like of high-melting (refractory) metals has been to
increase creep resistance by the introduction of 0.005 to 10 mass %
of one or more types of compounds or mixtures selected from the
group consisting of oxides, nitrides, carbides, borides, silicates,
or aluminates with particle diameters of .ltoreq.1.5 .mu.m and
having melting points of 1500.degree. C. or greater into
high-melting metals such as Mo, W, Nb, Ta, V, and Cr (refer to
patent document 1). However, patent document 1 discloses an
improvement in the heat resistance and creep resistance of
high-melting metals at high temperatures, not a remedy for
low-temperature embrittlement, recrystallization embrittlement, or
irradiation embrittlement.
[0008] In addition, the inventors of the present invention also
applied for a patent (refer to patent document 2) based on the
discovery that dispersion of 0.05 to 5 mol % of ultrafine particles
of group IVa transition metal carbides with particle diameters of
10 nm or less in molybdenum alloy and restricting the crystal grain
diameter to 1 .mu.m or less enables the strength of the molybdenum
alloy to increase, less loss of strength, even when heated at
high-temperature, and an alleviation of low-temperature
embrittlement, recrystallization embrittlement, and neutron
irradiation embrittlement. However, the molybdenum described in
patent document 2 is a material that exhibits ductility at room
temperature, even as a pure metal, and has completely different
properties and manufacture conditions in comparison to tungsten,
which is an extremely brittle material having a high melting point
that is 800.degree. C. higher than that of molybdenum.
[0009] In addition, with molybdenum, it is necessary to introduce
work-deformed structure by plastic working (hammering (forging),
rolling, or the like) in order to improve ductility in patent
document 2, resulting in a decrease in recrystallization
temperature and anisotropy. With tungsten, on the other hand, the
issue is ductility improvement in a recrystallized equiaxed
structure that is in a recrystallized state with absolutely no
work-deformed structure and therefore no anisotropy. The two cases
are thus substantially different.
PRIOR ART DOCUMENTS
[0010] Patent Document [0011] Patent document 1: Japanese Laid-open
Patent Publication No. 1-502680 [0012] Patent document 2: Japanese
Laid-open Patent Publication No. 8-85840
[0013] Non-Patent Documents [0014] Non-patent document 1: Collected
abstracts of the Japan Institute of Metals and Materials Vol. 148,
p. 235 [0015] Non-patent document 2: Collected abstracts of the
Japan Institute of Metals and Materials Vol. 143, p. 322
DISCLOSURE OF THE INVENTION
[0016] Problems to be Solved by the Invention
[0017] The inventors of the present invention carried out
painstaking investigations and discovered that, when an alloy
containing transition metal carbide that has been produced by
mechanical alloying (MA) and hot isostatic pressing (HIP) is
additionally subjected to a strengthening treatment for the
recrystallized random grain boundaries employing grain boundary
sliding by superplastic deformation in order to reinforce the
recrystallized grain boundaries in the alloy, the weak grain
boundaries in the recrystallized microstructure can be dramatically
strengthened. As a result, there is a dramatic resolution of
low-temperature embrittlement, recrystallization embrittlement, and
irradiation embrittlement. In addition, although the strengthening
treatment for random recrystallized grain boundaries employing
grain boundary sliding by superplastic deformation can be applied
to any alloy that exhibits superplastic deformation due to grain
boundary sliding, the novel discovery was made that this method is
effective for reducing the brittleness of tungsten, which is the
most brittle material among the metals. The present invention was
realized based on this new knowledge.
[0018] Specifically, an aim of the present invention is to provide
a method for manufacturing alloys containing transition metal
carbide having dramatically resolved low-temperature embrittlement,
recrystallization embrittlement, and irradiation embrittlement. An
additional aim of the present invention is to provide an alloy that
is manufactured by this manufacturing method. An additional aim of
the present invention is to provide a tungsten alloy containing a
transition metal carbide having dramatically resolved
low-temperature embrittlement, recrystallization embrittlement, and
irradiation embrittlement.
Means for Resolving the Problems
[0019] The present invention is described below and relates to a
method for manufacturing an alloy containing a transition metal
carbide, a tungsten alloy containing a transition metal carbide,
and an alloy that is manufactured by this manufacturing method.
[0020] (1) A method for manufacturing an alloy, characterized by
having a step for mechanically alloying a metal raw material and at
least one selected from carbides of group IVA, VA, or VIA
transition metals, a step for sintering the raw material powder
obtained in the mechanical alloying step (i.e., mechanically
alloyed powder) using hot isostatic pressing (HIP), and a step for
subjecting the alloy obtained in the sintering step to superplastic
deformation due to grain boundary sliding of 60% or greater at 500
to 2000.degree. C. and at a strain rate of 10.sup.-5 to 10.sup.-2
s.sup.-1.
[0021] (2) The method for manufacturing an alloy according to (1),
characterized by having a step in which the transition metal
carbide and the metal raw material are degassed by heating prior to
the mechanical alloying step.
[0022] (3) A tungsten alloy comprising 0.25 to 5 mass % of at least
one type selected from carbides of a group IVA, VA, or VIA
transition metals, the tungsten alloy characterized in that the
oxygen content is 950 ppm by mass or less, the nitrogen content is
60 ppm by mass or less, 80% or more of the tungsten phase observed
in a sectioned surface area is recrystallized to equiaxed grains
with grain diameters of 0.05 to 10 .mu.m, the ductile-brittle
transition temperature determined by three-point flexure is 500K or
less, and plastic deformation is possible at or above this
temperature.
[0023] (4) The tungsten alloy according to (3), characterized in
that 90% or greater of the azimuths of the carbide present in the
tungsten alloy structure and the azimuths of the tungsten matrix
are in the (Kurdjumov-Sachs) azimuth relationship: {111} W//{110}
transition metal carbide <110> W//<111> transition
metal carbide.
[0024] (5) The tungsten alloy according to (3) or (4),
characterized in that the full width at half maximum for reflection
of the (220) diffraction planes is 3.degree. or less as determined
by X-ray diffraction, or that there are 50 or fewer dislocations
within crystal grains as determined by transmission electron
microscopy.
[0025] (6) The tungsten alloy according to any of (3) to (5),
characterized in that the maximum bend strength determined by
three-point flexure is 1470 MPa or greater.
[0026] (7) The alloy that is manufactured by the manufacturing
method according to (1) or (2).
Effect of the Invention
[0027] In accordance with the present invention, transition metal
carbide and alloy powder are treated by mechanical alloying (MA)
method and hot isostatic pressing (HIP) method, and superplastic
deformation that can maximally utilize grain boundary sliding is
used in order to foster and optimize carbide grain boundary
precipitation and grain boundary segregation in recrystallized
micrograin structures. As a result, (1) the grain boundary strength
of the alloy in the recrystallized structure is significantly
improved, particularly the grain boundary strength (grain boundary
bond strength) of tungsten, allowing high strength and high
toughness to be manifested, (2) there is little chance for
recrystallization embrittlement because the material undergoes
little structural change when heated at high temperatures due to
the original recrystallized state, resulting in extremely little
loss of strength or ductility, (3) irradiation embrittlement can be
greatly resolved, and (4) effects can be obtained such as a
suitable decrease in yield strength, because the tungsten alloy
crystal grain diameters grow to about 0.05 to 10 .mu.m when
tungsten is used as the metal matrix, and a tungsten alloy thus can
be produced which can undergo plastic deformation, even near room
temperature.
BRIEF DESCRIPTION OF THE DRAWINGS
[0028] FIG. 1 shows the relationship between strength and
temperature for a normal metal and tungsten;
[0029] FIG. 2 shows the principle of superplastic deformation;
[0030] FIG. 3 schematically shows plastic working with the
objective of introducing work-deformed structures with dislocations
as carriers, resulting in a decrease in recrystallization
temperature and anisotropy;
[0031] FIG. 4 schematically shows the GSMM step;
[0032] FIG. 5 shows the three-point bending behavior at a
temperature of 400 K in Embodiment 4 (DBTT: 310 K) and Embodiment 6
(DBTT: 420 K);
[0033] FIG. 6 shows the three-point bending behavior at 300 K in
Embodiment 4;
[0034] FIG. 7 shows the X-ray diffraction pattern of Embodiment 2
(GSMM treated) and Comparative Example 1 (not GSMM treated);
[0035] FIG. 8 is a photograph which shows the transmission electron
micrographs of Comparative Example 1 and Embodiment 2;
[0036] FIG. 9 shows the X-ray diffraction patterns of Embodiment 5
(GSMM treated) and the as-HIP prior to the GSMM treatment in
Embodiment 5;
[0037] FIG. 10 is a photograph showing the transmission electron
micrograph of the tungsten alloy of Embodiment 2.
MODE FOR CARRYING OUT THE INVENTION
[0038] The present invention is characterized in that an alloy is
manufactured by a step for degassing a raw material by heating as
necessary, a step for subjecting the raw material that is obtained
in the degassing step to mechanical alloying (MA; also referred to
as "MA step" below), a step for sintering the raw material powder
obtained in the mechanical alloying step using hot isostatic
pressing (HIP; also referred to below as "HIP step"), and a step
for subjecting the alloy obtained in the sintering step to a
recrystallization random grain boundary strengthening treatment
carried out by superplastic deformation that can maximally utilize
grain boundary sliding (also referred to below as "GSMM step,"
where "GSMM" is an abbreviation for grain boundary sliding-based
microstructure modification). In addition, the present invention is
more specifically characterized in that the alloy that has been
manufactured by this method is a tungsten alloy. The present
invention shall be described in greater detail below.
[0039] The raw materials that are used in the present invention
will first be described. The transition metal carbide that is used
in the present invention refers to a carbide of a transition metal
that is selected from group IVA, VA, or VIA. Titanium carbide,
zirconium carbide, niobium carbide, tantalum carbide, and the like
are particularly preferred, because these transition metal elements
rapidly diffuse and react with carbon and tend to form carbides
before the formation of brittle W.sub.2C, and because the carbides
that are formed are thermally stable. These transition metal
carbides of group IVA, VA, and VIA (referred to below simply as
"transition metal carbides") may be used individually, or multiple
carbides may be used in combinations.
[0040] The added amount of transition metal carbide with respect to
the alloy is preferably 0.25 to 5 mass %. If the added amount of
transition metal carbide is less than 0.25 mass %, then the grain
boundary strengthening effects or the migration inhibitory effects
of the grain boundaries at high temperatures will be poor, and the
effect of an increase in recrystallization temperature or the
effect of inhibiting the production of coarse crystal grains
subsequent to recrystallization will be poor. There will also be
insufficient remedy in low-temperature embrittlement,
recrystallization embrittlement, and neutron irradiation
embrittlement, as well as insufficient increase in high-temperature
strength. On the other hand, the alloy will tend to have an
undesirable increase in brittleness if the added amount of
transition metal carbide exceeds 5 mass %.
[0041] Examples of alloy raw materials other than the transition
metal carbides include one or more selected from tungsten,
molybdenum, vanadium, yttrium, chrome, niobium, tantalum, titanium,
zirconium, hafnium, and the like, or stainless steel, steel, and
the like. However, the manufacturing method of the present
invention is particularly useful for group VIA transition metals
such as tungsten. The alloy raw material powder preferably has a
Fischer particle diameter of 2 .mu.m or greater. Although described
in detail in the manufacturing methods presented below, the reason
is that high concentrations of oxygen or nitrogen in the alloy that
has been manufactured result in 1) impeding of transition metal
carbide grain boundary precipitation/segregation which is necessary
for dramatically resolving low-temperature embrittlement,
recrystallization embrittlement, and irradiation embrittlement, 2)
promotion of the formation of W.sub.2C, which is itself brittle and
acts as an origin for fracture, and 3) formation of pores by oxygen
and nitrogen which act as origins for fracture. For this reason, in
order to strengthen the weak grain boundaries in recrystallized
microstructures within the alloy, it is essential to reduce the
oxygen and nitrogen contents of the alloy raw material powder. It
is preferable to carry out the degassing step described below and
to provide the raw material with the particle diameter described
above. However, if atmospheric control is carried out strictly so
as to suppress the admixture of impurities, then this requirement
of 2 .mu.m or greater may not be needed, and 1 .mu.m or less may be
sufficient.
[0042] The respective steps of the manufacturing method of the
present invention are described below. The degassing step in which
the raw material powder is heated is carried out in order to
decrease the final content of oxygen and nitrogen impurities in the
alloy. In the raw material powder preparation stage, this is done
in order to thoroughly eliminate the air (in particular moisture)
contained in the raw material powder. The degree of damage caused
by the nitrogen or oxygen is different depending on the metal
material, and so the degassing conditions of the degassing step may
be appropriately adjusted in accordance with the metal material.
With vanadium, for example, the oxygen or nitrogen will be absorbed
to produce a solid solution, even when heated in an ultra-high
vacuum. This results in embrittlement (environmental
embrittlement), and so the degassing step is carried out at a
fairly low temperature or not at all. In addition, with SUS316L,
there is no need to strictly carry out the step. With tungsten, on
the other hand, oxygen or nitrogen that remains in the alloy
precipitates or segregates at the weak recrystallized grain
boundaries as described above, promoting grain boundary
embrittlement (recrystallization embrittlement). Along therewith,
pores are formed which act as origins of fracture. Thus, for
example, when commercial tungsten powder is used as a raw material,
the raw material powder is preferably placed in a container at the
time of preparation of the raw material powder (a boat made of Mo
or the like that is used as a powder carrier), and the raw material
powder is subjected to a degassing treatment at 800 to 1500.degree.
C. by evacuation to 10.sup.-4 Pa or less. However, the degassing
step can be omitted, for example, by using an ultra-high purity W
powder manufactured by PLANSEE Japan Ltd. or other tungsten raw
material that already has sufficiently low oxygen and nitrogen
concentrations, the raw material is sealed in an inert gas or
reducing gas atmosphere (gas that has been purified to a level at
which there is negligible water content or the like that contains
the impurities) and an MA step or the like is carried out, thereby
removing the admixed oxygen and nitrogen.
[0043] Degassing is preferably carried out using a degassing time
of 120 min at 800.degree. C. or greater, or a degassing time of 90
min or greater at 950.degree. C. or greater, as a guideline. If the
degassing temperature is less than 800.degree. C., then desorption
of gas will be insufficient, whereas reactions with the degassing
container (boat made of Mo or the like) will tend to occur if the
temperature exceeds 1500.degree. C., resulting in initiation of
aggregation of the raw material powder and additional undesirable
effects in subsequent steps.
[0044] With tungsten alloy, the oxygen content in the manufactured
alloy is 950 ppm or less, preferably 850 ppm or less, more
preferably 300 ppm or less, and the nitrogen content is 60 ppm or
less, preferably 50 ppm or less. If the oxygen and nitrogen
contents of the alloy are below these values, then the production
of a dense alloy will be possible. The oxygen or nitrogen content
of the tungsten alloy at the raw material powder stage is about
three times that of the final alloy that has been manufactured.
Consequently, in regard to process management, it is preferable for
process management to be carried out so that oxygen is contained at
about 3000 ppm or less and nitrogen is contained at about 180 ppm
or less in the raw material powder upon completion of the degassing
step.
[0045] An MA step is carried out after the degassing step. With
this MA step, operations extending from ball-milling of the raw
material powder through sealing of the ball-milled powder in a
capsule to produce a HIPed compact are preferably carried out in an
inert gas or reducing gas atmosphere in order to prevent the
admixture of oxygen or nitrogen. Ar, helium, neon, and the like are
examples of inert gases, and hydrogen and the like are examples of
reducing gasses.
[0046] The MA step is a step in which high mechanical energy is
imparted to the raw material alloy powder and the transition metal
carbide, thereby decomposing the transition metal carbide and
bringing about solid dissolution thereof in a uniform solid
solution in the matrix phase alloy structure. Simultaneously, the
ultrafine powder of the mother phase (alloy) is produced. This step
is carried out, for example, using a device such as a triaxial
vibrating ball mill, a planetary ball mill, or an attriter. The MA
treatment normally involves introducing balls and raw material
powder and/or pre-alloy powder into a pot (container) and rotating
or vibrating this pot on a ball mill support stand, thereby
imparting high mechanical energy to the raw material powder and/or
pre-alloy powder. As a result, the different element species that
have been added can be forcibly solid-dissolved, even in systems in
which solid dissolution will not occur under the equilibrium
conditions. In addition, extremely fine crystal grains (10 to 30
nm) can be produced at around room temperature. In order to remove
impurities from the ball surfaces and the inner walls of the pot
which are used in the MA step, the pot and balls alone may be
heated under vacuum for 3 to 10 h at 150 to 200.degree. C. prior to
introducing the raw material powder into the pot.
[0047] The specific treatment conditions for the MA step, such as
the treatment time, the rotation rate, the ball material and
diameter, the ratio of the total ball mass to total raw material
powder mass, and the ratio of total internal container volume to
total ball volume, may be determined appropriately so that the
transition metal carbide is uniformly mixed, decomposed and
solid-dissolved in the alloy, so that ultrafine matrix phase metal
crystals are produced, and so that effects of admixture of
container and ball material into the raw material powder during the
MA step is inhibited (suppression of admixed amounts to negligible
levels, or use of materials that will not affect subsequent
material characteristics, even if they become admixed).
[0048] The HIP step involves isostatic pressing of the MA powder
produced in the MA step using Ar gas and carrying out sintering at
a comparatively low temperature at which grain growth of the
ultrafine alloy powder produced in the MA step does not easily
occur, while preventing exposure to atmosphere constituted by gas
impurities that are harmful to the alloy. As a result, the
transition metal carbide that is forcibly solid-dissolved during
the MA step segregates and precipitates, thereby preventing grain
growth of the ultrafine particles due to a pinning effect, while
also producing equiaxed ultrafine grains of alloy matrix phase in
which transition metal carbide has segregated/precipitated at the
grain boundaries without strain resulting from recrystallization.
Specifically, the MA powder is sealed in the inert gas or reducing
gas atmosphere described above in a metal container made of soft
(mild) steel, SUS, Ti, Nb, Ta, or the like. After removing the
sealed gas by producing a high vacuum (typical evacuation level of
10.sup.-4 to 10.sup.-6 Pa), the material is sintered for 1 to 5 h
at 1350 to 1400.degree. C. and 100 to 1000 MPa, thereby producing
an alloy having the structure described above. In order to remove
impurities and the like on the inner walls of the metal container
that is used in the HIP step, the metal container alone may be
heated under vacuum for 1 to 3 h at 500 to 1000.degree. C. prior to
introduction of the MA powder into the metal container.
[0049] The GSMM step is a step in which the weak grain boundaries
in the recrystallized microstructure are replaced with strong
transition metal carbide heterophase interfaces or strong grain
boundaries in which the constitutive elements of the transition
metal carbide have precipitated or segregated. Consequently, when
the transition metal carbide precipitates or segregates at the
grain boundaries, the grain boundary bond strength is increased,
having the beneficial effect of increasing the fracture strength
and remedying embrittlement. In addition, the GSMM step has the
effect of increasing the crystal grain diameter to an appropriate
size, decreasing yield strength (flow strength), and manifesting
ductility (decreasing (relaxing) the grain boundary load); removing
residual gas pores that tend to act as origins for fracture (1 to
3% remaining after HIP), as well as strengthening the interfaces
(boundary interfaces) with different precipitates in dispersion
strengthened alloys containing precipitates (e.g., vanadium or
stainless steel). In the present invention, the effect of grain
boundary sliding at high temperatures is used in order to promote
and optimize grain boundary precipitation and segregation of the
transition metal carbide. FIG. 2 is a diagram showing the principle
of superplastic deformation by grain boundary sliding. Grain
boundary sliding refers to crystal grain displacement/movement in a
state in which the equiaxed condition is maintained, as indicated
in FIG. 2 (2).fwdarw.(3).fwdarw.(4) when shear stress .tau. is
applied to the crystal structure in FIG. 2 (1). As a result of an
extremely large number of repetitions of this type of grain
boundary sliding, the transition metal carbide precipitates or
segregates at the grain boundaries, having the effect of increasing
the fracture strength at the weak recrystallized grain boundaries
until it surpasses the yield strength (flow strength), allowing
apparent alloy deformation.
[0050] However, grain boundary sliding is non-uniform deformation
that conversely promotes embrittlement along with grain boundary
displacement due to crack formation at grain boundary triple points
(an example being high-temperature embrittlement typically seen
with copper alloys and the like). It is thus extremely important,
in the present invention, for the deformation amount at break to be
extremely large, and to employ superplastic deformation which can
maximally utilize grain boundary sliding. As stated above, grain
boundary sliding is non-uniform deformation that typically promotes
embrittlement due to crack formation at grain boundary triple
points in conjunction with grain boundary displacement. However, by
carrying out superplastic deformation of the present invention
using constant conditions described below for temperature and
strain rate (a quantity obtained by dividing the speed at which a
sample piece is deformed by the size of the sample piece to convert
to strain), a relaxation (accommodation) mechanism operating with
grain boundary sliding prevents the formation of cracks, and
elongation of several hundreds of percent is produced. In order to
promote and optimize transition metal carbide grain boundary
precipitation and segregation, it is more effective to carry out
GSMM for a longer period of time than for a shorter period of
time.
[0051] As stated above, superplastic deformation is a deformation
mode in which elongation of several hundreds of percent occurs due
to grain boundary sliding, and equiaxed crystal grains are
essentially maintained even after deformation. Thus, it is possible
to "promote and optimize transition metal carbide grain boundary
precipitation and segregation though relative motion or rotation of
crystal grains by active grain boundary sliding over a long period
of time, and to maintain an isotropic recrystallized structure with
little anisotropy." This "recrystallized random grain boundary
strengthening treatment carried out by superplastic deformation
that can maximally utilize grain boundary sliding" is designated,
in the present invention, by GSMM (grain boundary sliding-based
microstructural modification).
[0052] FIG. 3 is a diagram that schematically shows the concept of
"plastic working performed to introduce work-deformed structures
using dislocations as a carrier, resulting in anisotropy and a
decrease in recrystallization temperature," which has been widely
used for increasing toughness in metals including tungsten. The
term "dislocation" refers to linear lattice defects, and a
characteristics of plastic working include (1) that specific
crystallographic planes can perform sliding motion under small
stress in specific crystallographic directions, (2) that
dislocations can multiply anew through the slip process, (3) that
extremely strong interactions occur with areas having an elastic
strain field (e.g., the dislocations are entirely enclosed by an
elastic strain field due to elastic strain that arises at around
the periphery of different atomic species with different sizes).
For this reason, when deformation progresses as a result of tensile
stress applied to a material as shown in FIG. 3(1), 3(2), slip
arises in the material as shown in FIG. 3(3), dislocations in the
material multiply (specifically, the dislocation density
increases), and, as a result, the stress required for the
dislocations to undergo additional slipping, in other words, the
stress required for plastic deformation of the alloy, increases,
and reaches the fracture strength, resulting in break as shown in
FIG. 3(4). "Structures in which the dislocation density has
increased due to plastic deformation" are referred to as
work-deformed structures or worked structures, but the increase in
dislocation density amounts to an increase in the internal strain
field of the material (crystal) and produces a condition of high
internal energy. Thus, materials having this high internal energy
tend to release this internal energy, and so when heat is applied
(the temperature is increased), the internal energy is released
with just a slight amount of energy (slight increase in
temperature); one process for this release being recrystallization.
In most cases, the upper limit of elongation by "plastic working
performed to introduce deformation-processed structures using
dislocations as a carrier and to produce anisotropy and a decrease
in the recrystallization temperature" is approximately several tens
of percentage points, and is considerably less than 100%,
particularly with elongation.
[0053] On the other hand, with superplastic deformation occurring
by grain boundary sliding, recrystallized structure having little
strain is maintained even after deformation, and the internal
energy is not substantially increased. Thus, with the GSMM
treatment of the present invention, although the crystal grains
grow (increase by roughly a factor of ten) because the treatment is
carried out at a temperature that is higher than the HIP
temperature, grain growth leads to decrease in the total area of
crystal grain boundaries and hence internal energy because crystal
grain boundaries are a high energy region.
[0054] As stated above, GSMM in the present invention is a new
structure control technique whereby increased durability is
manifested by strengthening the weak recrystallized grain
boundaries which are a cause of grain boundary embrittlement. The
plastic work referred to above is substantially different in
principle, and the post-treatment alloy fracture strength and
elongation at break are also completely different.
[0055] The GSMM step, as shown in FIG. 4 involves sandwiching the
alloy that is produced in the HIP step between heat resistant, high
strength ceramics or ceramics composite plates (typically BN-SiC
composite material plates) and applying a pressure at a strain rate
of 10.sup.-5 s.sup.-1 to 10.sup.-2 s.sup.-1 at a high temperature
of 500 to 2000.degree. C. (40 to 50% or more higher than the
melting points of the respective alloys measured in absolute
temperature) to carry out plastic deformation due to grain boundary
sliding at 60% or greater. The temperature is preferably suitably
adjusted in accordance with the melting points of the respective
alloys, as stated above. For example, 1200 to 2000.degree. C. is
preferred for tungsten and molybdenum, but 1400 to 2000.degree. C.
is additionally preferred for tungsten. In addition, a temperature
of 800 to 1500.degree. C. is preferred for vanadium and SUS316.
With tungsten, there are cases where the alloy will break during
compression deformation if the temperature is less than
1400.degree. C., whereas exceeding 2000.degree. C. is undesirable
because the equipment used for industrial manufacture will increase
in size. In addition, if the strain rate is slower than 10.sup.-5
s.sup.-1, effects will be obtained, but an excessively long
processing time will be required, which is industrially
disadvantageous. It is undesirable for the strain rate to be
greater than 10.sup.-2 s.sup.-1 due to the danger of alloy
fracture. Carrying out plastic deformation of 60% or greater means
that the elongation (deformation) of the test piece due to plastic
deformation is 60% or greater. Elongation is expressed as the
elongation length of a test piece (.DELTA.L) divided by the initial
length (L), multiplied by 100 in order to obtain a percentage. The
material may instead be subjected to tensile deformation, shear
deformation, or the like may be used, provided that the
aforementioned temperature, strain rate, and plastic deformation
can be provided.
[0056] The transition metal carbide is necessary in order for the
crystal grains of the alloy matrix phase to be maintained in fine
grain sizes and in order to manifest superplastic deformation. In
addition, the heterophase interface between the transition metal
carbide and the alloy mother phase (matrix) satisfies the
Kurdjumov-Sachs azimuth (orientation) relationship, and thus
high-strength hetero-phase interfaces are formed. When tungsten is
used as the alloy raw material, 90% or greater of the azimuths of
the transition metal carbide present in the tungsten alloy
structure and the azimuths of the tungsten matrix are in the
(Kurdjumov-Sachs) azimuth relationship: {111} W//{110} transition
metal carbide <110> W//<111> transition metal carbide.
If 10% or more of the transition metal carbide particles do not
satisfy the Kurdjumov-Sachs azimuth relationship, then it will not
be possible to obtain sufficient maximum flexural strength at room
temperature (approximately 1470 MPa).
[0057] In addition, with the alloy that has been manufactured by
the manufacturing method of the present invention, growth occurs
until the crystal particle diameter of the tungsten alloy is about
0.05 to 10 .mu.m. As a result, an effect is produced whereby the
yield strength is decreased to an optimal level, and a tungsten
alloy can be produced that can undergo plastic deformation near
room temperature. For tungsten alloys, when the alloy is
manufactured by the manufacturing method of the present invention,
the three-point bending ductile-brittle transition temperature (nil
ductility temperature; DBTT) can be decreased to about 500K, and
thus plastic deformation is possible at or above the
ductile-brittle transition temperature.
[0058] The crystal grain diameter can be determined as the average
grain diameter by using commercial image processing software (e.g.,
Image Pro) to carry out image processing on photographs that are
typically taken by a transmission electron microscope from the
center part of a sample cross section. The average grain diameter
can be determined only for the tungsten matrix phase. Because
averaged data could be obtained by counting the tungsten crystal
grains over a surface area ratio of 80% or greater, statistical
determinations were carried out.
[0059] For less than 20% of the surface area, counting the tungsten
crystal grains was difficult (the grains were fine and too
numerous, and it was difficult to determine where the crystal grain
boundaries were, because it was difficult to see the borders of the
crystal grains constituting the crystal grain boundaries). However,
if the average grain diameter for tungsten is calculated over a
region constituting 80% or more of the surface area, then the
characteristics of the various materials can be elucidated. The
crystal grain diameter can be calculated as a stable average grain
diameter by counting 300 or more tungsten crystal grains and
calculating the surface area. As necessary, the crystal grain
diameter can be measured over a broad region of 80% or more of the
total field of numerous photographs taken with a transmission
electron microscope. As a result, 80% or more of the crystal grains
that can be measured should be in the grain diameter range of 0.05
to 10 .mu.m.
[0060] If the average grain diameter is less than 0.05 .mu.m,
plastic deformation will be extremely difficult, because the yield
strength will become extremely high, resulting in decreased work
and manufacture yields, which is industrially disadvantageous. On
the other hand, if the average grain diameter exceeds 10 .mu.m,
superplastic deformation will not readily occur. In order to allow
plastic deformation in the vicinity of room temperature, it is
necessary to carry out suitable adjustments so that a suitable work
ratio is produced during plastic deformation (in other words,
during GSMM treatment) that is carried out in order to increase
toughness. The temperature during the GSMM treatment may be
decreased in order to produce a smaller average grain diameter, or
the temperature during GSMM treatment may be increased in order to
produce a larger average grain diameter.
[0061] One point that should be noted in regard to the description
of the present invention is that superior characteristics (e.g.,
fracture strength and ductility) can be obtained in a structure in
which thorough (full) recrystallization has occurred, because
equiaxed crystal grains that are not anisotropic are produced in
the metal structure. The term "equiaxed crystal grains" in the
present invention means that the aspect ratio (ratio of the
longitudinal and transverse crystal grain lengths) is 2 or less
regardless of the cross-section when the metal structure is viewed
two-dimensionally.
Embodiments
[0062] With the methods for manufacturing the alloys and the
manufactured alloys of the embodiments described below, as shown in
FIG. 4, simple compressive deformation was carried out along one
axis, but deformation is not restricted to simple compression,
provided that superplastic deformation allowing maximal utilization
of grain boundary sliding can be realized. Depending on the shape
of the alloy product that is desired, for example, reduction by
rolling can be employed for sheet-form materials, for example.
[0063] Characterization of transition metal carbide amount required
for manifesting superplasticity
Experiment 1
[0064] TiC powder with an average particle diameter of 0.7 .mu.m
(manufactured by Soekawa Chemical Co., Ltd.) was added to tungsten
powder with an average particle diameter of 4 .mu.m (Manufactured
by A.L.M.T. Corp.) using the Fischer method. The material was
introduced into a molybdenum boat in a hydrogen atmosphere and was
then subjected to a degassing treatment by heating for 1.5 h at
950.degree. C. under high vacuum (<1.times.10.sup.-4 Pa). Next,
the material was subjected to a mechanical alloying (MA) treatment
by mixing for 70 h at a vibration frequency of 360 cycles/min (6
Hz) in a TZM (titanium, zirconium-containing molybdenum alloy)
container (pot) using a tri-axial vibrating ball mill (TKMAC 1200,
manufactured by Topology Systems). In order to characterizes the
appropriate TiC powder addition range, eight MA treatment sample
runs were carried out with TiC powder contents of 0 to 6.0 mass
%.
[0065] Next, the MA-treated powder was introduced into a molybdenum
boat and was heated for 1.5 h at 950.degree. C. under high vacuum
in order to degas the hydrogen that had admixed in the TiC powder
and the tungsten during the MA treatment. This degassed powder was
then sealed in an HIP capsule (mild steel), and the container
(capsule) was vacuum-sealed before carrying out a HIP treatment for
3 h at 1350.degree. C. and 196 MPa in argon gas to obtain a
sintered body. The resulting sintered body is referred to as
"as-HIPed compact".
[0066] Pieces having dimensions of 0.4 mm.times.4 mm.times.16 mm
were then wire-cut from the as-HIPed compact (parallel part length
5 mm; I-shaped flat sheet-form tensile test piece similar to the
test piece shown in FIG. 1 of T. Kuwabara, H. Kurishita, M.
Hasegawa, Development of an Ultra-Fine Grained V-1.7 mass % Y alloy
Dispersed with Yttrium Compounds Having Superior Ductility and High
Strength, Mater. Sci. Eng. A 417 (2006) 16-23). The entire surface
was mechanically polished with waterproof paper (to #1500), the
four edges were then chamfered, and the piece was mounted on a
tensile test fixture and subjected to high-temperature tensile
testing. The tensile fixture was a test-piece shoulder-bearing" (R
part) type whereby alignment is ensured by a system in which the
compressive load on the fixture is converted to tensile load on the
test piece, allowing one-touch mounting of the test piece on the
fixture. Heating of the test piece was carried out by
high-frequency induction heating using a graphite susceptor. The
surface temperature of the test piece was continually observed and
recorded using a two-color radiation thermometer (Chino, model
1R-AQ). The tensile test was carried out using an Instron model
81362 Electrically Actuated Tester at temperatures of 1500.degree.
C., 1600.degree. C., and 1700.degree. C., an initial strain rate of
5.times.10.sup.-4/s (cross-head speed: 0.0025 mm/s), and an
evacuation level of 5.times.10.sup.-4 Pa. Load and elongation (%)
were measured during the tensile test. The oxygen and nitrogen
concentrations in the sample were measured by infrared absorption
on a LECO-TC600 device using a thermal conductivity method. With
all of the samples, the oxygen concentration was 850 ppm or less,
and the nitrogen concentration was 50 ppm or less. The results are
shown in Table 1. In the table, >160 denotes that no break
occurred, even at a deformation of 160%.
TABLE-US-00001 TABLE 1 1500.degree. C. 1600.degree. C. 1700.degree.
C. TiC content Elongation Elongation Elongation Sample No. Mass %
(%) (%) (%) 1 0 2 5 5 2 0.15 3 10 30 3 0.25 70 105 >160 4 0.5
>160 >160 >160 5 1.1 >160 >160 >160 6 1.5 >160
>160 >160 7 5 >70 >100 >160 8 6 10 30 60
Experiment 2
[0067] The same test as in Experiment 1 was carried out, with the
exception that the hydrogen in Experiment 1 was changed to argon,
and nine samples were used in which the TiC content was varied. The
results are shown in Table 2
TABLE-US-00002 TABLE 2 1500.degree. C. 1600.degree. C. 1700.degree.
C. TiC content Elongation Elongation Elongation Sample No. Mass %
(%) (%) (%) 9 0 2 5 5 10 0.25 3 7 7 11 0.5 30 40 60 12 0.7 50 70
>160 13 0.8 70 110 >160 14 1.1 120 >160 >160 15 1.5 70
>160 >160 16 5 50 70 >160 17 6 10 30 60
[0068] Experiment 1 and Experiment 2 above show that the TiC amount
required for manifestation of superplasticity at 1600 to
1700.degree. C. (elongation at break:100% or greater) is 0.25 to 5
mass % with the as-HIPed compacts produced from powder that was
MA-treated in hydrogen atmosphere, and 0.7 to 5 mass % with those
produced using an argon atmosphere. If the TiC amount is below
these ranges, then weak grain boundaries occur in great numbers
among the grain boundaries of the tungsten phase, and there are few
grains of a second phase that inhibit grain boundary movement. For
this reason, grain growth is rapid in the tungsten phase, resulting
in the production of large-size crystal grains. The TiC phase is
essential for rotation and movement of crystal grains during grain
boundary sliding and for maintaining fine equiaxed crystal grains
that are required for superplastic deformation. For this reason, if
the amount of TiC phase is small, then when grain boundary sliding
non-uniform deformation arises in high-temperature tensile testing,
grain boundary cracks will form and grow, and the elongation at
break will be low.
[0069] Conversely, if the amount of TiC exceeds these ranges, then
the contact frequency between TiC phases will increase, and the
proportion of TiC/TiC interfaces will increase. In comparison to
the tungsten matrix phase, the TiC phase has low plastic
deformation capability, and it is thought that TiC/TiC interfaces
also do not readily slide. Consequently, overloading of the harmony
of the tungsten phase occurs in relation to continuous grain
boundary sliding of the tungsten grains, and grain boundary
(interfacial) cracking arises, thereby also decreasing the
elongation at break.
Experiment 3
[0070] An experiment was carried out in the same manner as in
Experiment 1, with the exception that the material was changed to
titanium carbide of Experiment 1, and zirconium carbide, niobium
carbide, tungsten carbide, or mixtures thereof were added while
varying the content thereof. In addition, high-temperature tensile
testing was carried out only at 1600.degree. C. The results are
shown in Table 3.
TABLE-US-00003 TABLE 3 1600.degree. C. Carbide content in tungsten
Elongation (%) ZrC 0.3 mass % 120 ZrC 4.7 mass % >160 NbC 0.32
mass % 130 NbC 4.5 mass % >160 TaC 0.28 mass % 130 TaC 3.3 mass
% >160 TaC 5.0 mass % >160 ZrC 0.3 mass % + TaC 2 mass %
>160 NbC 0.3 mass % + TiC 2 mass % >160 TiC 1 mass % + TaC 2
mass % >160 TaC 0.1 mass % 2 TiC 0.08 mass % + TaC 0.03 5 mass %
ZrC 6 mass % Could not be performed
[0071] As is clear from Table 3, even when transition metal
carbides other than TiC were added to the alloy at about 0.25 to 5
mass %, it was clear that there was an improvement in ductile
characteristics.
Embodiment 1
[0072] An as-HIPed compact material was produced in the same manner
as with sample no. 5 of Experiment 1, with the exception that the
degassing conditions involved heating for 1.5 h at 1050.degree. C.
under high vacuum (1.times.10.sup.-4 Pa). Next, the as-HIPed
material that had been produced was wire-cut to produce a sintered
body with a diameter of about 9 to 10 mm and a height of about 20
mm. In order to strengthen the weak random grain boundaries by
superplastic deformation that can maximally utilize grain boundary
sliding, a disk shape material was produced by compression
deformation to a thickness of about 3.5 mm (diameter of about 21 to
23 mm) at a temperature of 1650.degree. C. and a strain rate of 0.5
to 2.times.10.sup.-4 s.sup.-1 (superplastic behavior tends to more
readily occur with slower strain rates, and so a rate was selected
at which experiments were most easily carried out, while
considering response (increase in flow stress) exhibited by the
material as a result of incremental increases in strain rates).
Heating of the sintered body was carried out by high-frequency
induction heating under evacuation using a graphite susceptor. An
Instron model R1362 Electrically Actuated Tester was used for
high-temperature compressive deformation. A piece with dimensions
of 1 mm.times.1 mm.times.20 mm was cut perpendicularly in the
compressive direction from this disk shape material, and the
surfaces and edges were polished with water-proof emery paper out
to #1500, thereby producing a bending test piece. The oxygen
concentration was 40 ppm and the nitrogen concentration was 30 ppm
in the test piece, as measured by infrared absorption on a
LECO-TC600 device using a thermal conductivity method. Next, the
test piece was subjected to three-point bend testing at a
temperature range of room temperature to 600.degree. C. with a
cross-head speed of 0.001 mm/s in an atmosphere produced by a flow
of high-purity Ar containing 4% H.sub.2. The three-point bend test
was carried out using a Servopulser EHF-2 model fatigue testing
machine, manufactured by Shimadzu Corporation (load capacity of 5
ton), connecting a span .+-.2.5 mm LVDT (Linear Variable
Differential Transformer) to the actuator head and attaching a
shear-type load sensor with a load capacity of 5 kN directly below
a load cell with a capacity of 5 ton. Control of testing was
carried out using a static test application program. An infrared
heating furnace (ULVAC-RIKO, Inc.) was used for heating the test
pieces, and measurement of the test piece temperature and
atmosphere (at a location separated by a few millimeters from the
test piece) was carried out in advance on a dummy test piece having
a spot-welded thermocouple. In actual testing, the temperature of
the atmosphere was controlled and measured. Flexural strength was
measured at room temperature, and the minimum value of the measured
averages of five bending test pieces was taken as the minimum
flexural strength, whereas the maximum value was taken as the
maximum flexural strength. The DBTT was determined by recording the
variation of measurements on the plastic strain at respective
temperatures while increasing the testing temperature by roughly 50
increments starting from room temperature. The temperature found by
extrapolating to a plastic strain of zero using a linear
approximation was taken as the DBTT. In determining a single DBTT,
it is necessary to measure the plastic strain while varying the
testing temperature, and measurements were carried out by preparing
three to five bar shape test pieces having the same impurity
concentrations and structures.
Embodiments 2 to 7
[0073] The degassing conditions involved heating for 1.5 h at
950.degree. C. in Embodiment 2, heating for 1 h at 950.degree. C.
in Embodiment 3, heating for 1 h at 900.degree. C. in Embodiment 4,
heating for 1.5 h at 850.degree. C. in Embodiment 5, heating for 1
h at 850.degree. C. in Embodiment 6, and heating for 1 h at
800.degree. C. in Embodiment 7. With the exception that the oxygen
amount and nitrogen amount in the tungsten alloy were changed, test
pieces were produced out using the same procedure as in Embodiment
1. The oxygen amount, nitrogen amount, minimum flexural strength
and maximum flexural strength at room temperature, and DBTT were
measured.
Comparative Example 1
[0074] Test pieces were prepared from as-HIPed compacts without
carrying out a compressive deformation treatment for GSMM, and
measurements were performed using the same procedure as in
Embodiment 2.
Comparative Example 2
[0075] With the exception that the TiC content was changed to 1.1
mass % and the degassing treatment was not carried out, test pieces
were prepared by the same procedure as in Embodiment 1, and
measurements were performed.
[0076] The results of measurements in Embodiments 1 to 7 and
Comparative Examples 1 and 2 are shown in Table 4.
TABLE-US-00004 TABLE 4 Minimum Maximum Oxygen Nitrogen flexure
flexure amount amount TiC strength strength DBTT (ppm) (ppm) (%)
(MPa) (MPa) (K) Embodiments 1 40 30 1.1 2800 3200 210 2 160 30 1.1
2700 2800 230 3 230 40 1.1 2690 2940 240 4 610 40 1.1 1840 2380 310
5 850 50 1.1 1450 1620 330 6 870 140 1.1 1240 1500 420 7 950 60 1.1
1340 1470 500 Comparative 1 160 30 1.1 1610 2160 850* examples 2
2120 180 1.1 1000 1260 .gtoreq.630
[0077] As is clear from Table 4, the tungsten alloy flexural
strength increased as the concentrations of oxygen and nitrogen
decreased when the material was subjected to a GSMM treatment. In
addition, it was clear that the DBTT decreased dramatically when
the as-HIPed material was subjected to the GSMM treatment, and
ductility was obtained even at low temperatures.
[0078] Three-Point Bending Testing
[0079] FIG. 5 shows the three-point bending behavior at a
temperature of 400K for Experiment 4 (DBTT: 310 K) and Experiment 6
(DBTT: 420K). FIG. 6 shows the three-point bending behavior for
Experiment 4 at 300K. As is clear from FIGS. 5 and 6, break
(fracture) occurred without manifestation of ductility at lower
temperatures than the DBTT temperatures of the resulting alloys,
and thus it was clear that the amounts of oxygen and nitrogen must
be decreased in addition to carrying out the GSMM treatment (by
compression).
[0080] X-Ray Diffraction Pattern Measurement
[0081] FIG. 7 compares the X-ray diffraction patterns for
Experiment 2 (GSMM treatment performed) and Comparative Example 1
(GSMM treatment not performed). From a comparison of the two, a
large difference in intensity was seen with the TiC peak,
indicating that TiC precipitation progressed during the GSMM
treatment. This was confirmed by transmission electron
microscopy.
[0082] Transmission Electron Micrographs
[0083] FIG. 8(1) is a transmission electron micrograph of
Comparative Example 1, and FIG. 8(2) is a transmission electron
micrograph of Embodiment 2. FIG. 8(3) is an enlarged view of the
portion indicated by ".rarw." in FIG. 8(2). As is clear from the
micrographs, the tungsten alloy that had been subjected to the GSMM
treatment was confirmed to have experienced TiC grain boundary
precipitation in the alloy. It was simultaneously confirmed that
the TiC constituent elements had undergone solid dissolution and
segregated at the grain boundaries.
[0084] X-Ray Diffraction Analyses
[0085] FIG. 9 compares the X-ray diffraction patterns of Embodiment
5 (GSMM treated) and the as-HIPed material prior to GSMM treatment
in Embodiment 5. It is clear that TiC precipitation similarly
progressed as a result of GSMM treatment, but that the
ductility-impeding (i.e., easily fractured) carbide W.sub.2C, had
formed, in contrast to FIG. 7. It is thought that this material is
produced as a result of an increase in oxygen level, which provides
oxygen distribution between the TiC and tungsten and results in a
reaction between some of the carbon that has dissociated from the
TiC with the surrounding tungsten.
[0086] Confirmation of Equiaxed Recrystallized Grains
[0087] With the tungsten alloy of Embodiment 2, a thin film having
a small perforation was formed at the center by electrolytic
polishing (TenuPol) at a thickness of about 50 .mu.m and a diameter
of 3 mm, whereupon the material was observed with a transmission
electron microscope (JEOL 2000) at an acceleration voltage of 200
kV. FIG. 10(1) shows the observation direction of the transmission
electron microscope, and FIG. 10(2) shows a micrograph of the
sample as observed from above (specifically, from a direction
parallel to the compression direction). FIG. 10(3) is a micrograph
of the sample as observed from the side (from a direction
perpendicular to the compression direction). In all cases, clear
images were obtained with an observational magnification of roughly
10,000.times.. As is clear from FIG. 10, the crystal grains were
equiaxed grains, with crystal grain aspect ratios in the range of 1
to 2.
[0088] In addition, upon observing the thin film under diffraction
conditions and at additionally high resolution, almost no
dislocations were observed in the crystal grains, and the number of
dislocations with the observed crystal grains was extremely low, at
about 1 to 3, in many cases. From the results of observations
described above, it is clear that the structure of the articles of
the present invention is a recrystallized structure. Dislocations
are present at 1000 or greater in the crystal grains of tungsten
that has not been recrystallized, including work-deformed
structure. In contrast, it was found that the there are 50 or less
dislocations in the crystal grains of recrystallized tungsten. It
became clear that, if this condition is satisfied, then the
characteristics of unstrained tungsten crystal grains are
exhibited.
[0089] In addition, an unstrained condition was confirmed based on
XRD measurements using a Rigaku RAD II-B device. From the results
of XRD measurements, although the effect of fine crystal grains was
obtained, the diffraction width of the diffraction peaks increased
with increasing strain in the non-recrystallized state. For
example, by investigating XRD results obtained by measurements on
stress-relieved commercial tungsten material and tungsten alloy of
Embodiment 2 under conditions of 40 kV and 30 mA using a Cu target
and a 1.degree. slit, it was clear that if the full width at half
maximum exceeds 3.degree. in (220) diffraction of tungsten with a
lattice constant of 0.11188 nm, strain remains and the material
does not have a recrystallized structure (commercially available
pure tungsten subjected to stress relief treatment), whereas the
material has a recrystallized structure without strain if the full
width at half maximum is 3.degree. or less.
[0090] Confirmation of Grain Diameter
[0091] The tungsten alloys produced in Embodiments 1 to 7 were
used, thin films having a small perforation was formed at the
center by electrolytic polishing (TenuPol) at a thickness of about
50 .mu.m and a diameter of 3 mm, whereupon the materials were
observed with a transmission electron microscope (JEOL 2000) at an
acceleration voltage of 200 kV. In the transmission electron
micrographs of all of the embodiments, the crystal grain diameters
could be measured for 80% or more of the entire field of the
micrographs, and it was confirmed that 80% or more of the crystal
grains that were measured were in the grain diameter range of 0.05
to 10 .mu.m.
[0092] Confirmation of Carbide Orientation and Tungsten Matrix
Orientation in Tungsten Alloy Structure
[0093] In contrast that the case described in "Confirmation of
grain diameter" above, bright-field images, dark-field images, and
selected area diffraction patterns at magnifications of several
hundred thousand were taken in numerous fields including a single
carbide grain within a tungsten matrix phase. The orientation
relationship between the carbide and the tungsten matrix phase were
thus analyzed. As a result, in the transmission electron
micrographs from all of the embodiments, it was confirmed that the
carbide present in the tungsten alloy structure and the tungsten
matrix satisfied the Kurdjumov-Sachs orientation relationship:
{111}W//{110} transition metal carbide <110> W//<111>
transition metal carbide.
Embodiment 8
[0094] After weighing and blending raw material powders at a
vanadium:yttrium:tungsten:TiC mass ratio of 89.8:1.4:8.0:0.8, the
material was placed in a Mo boat, and a degassing treatment was
carried out for 1 h at 200.degree. C. Next, a container and balls
used for MA treatment (material: TZM (Mo-0.5 Ti-0.1 Zr)) was baked
for 10 h at 150 to 200.degree. C. under high vacuum, whereupon the
blended raw material powder was introduced into the container along
with the balls, and an MA treatment was carried out using a
tri-axial vibrating ball mill for 70 h in a purified hydrogen
atmosphere. In order to eliminate the hydrogen that had been
admixed from the atmosphere during MA, a dehydrogenation treatment
was carried out for 1 h at 600.degree. C. under a vacuum of
1.times.10.sup.-4 Pa or less. Subsequently, the MA-treated vanadium
alloy powder was sealed in a hydrogen atmosphere in an HIP capsule
(soft (mild) steel) that had been degassed by heating under vacuum
at 900.degree. C., and, while degassing under vacuum at room
temperature, the HIP capsule was vacuum-sealed under a high vacuum
(2.times.10.sup.-5 Pa). Consequently, the HIP capsule interior was
in a highly evacuated tightly-sealed condition. This material was
then subjected to an HIP treatment for 3 h at 196 MPa and
1000.degree. C. in argon gas to produce a sintered body with a
relative density of 99.5% or more. Tensile test pieces were then
cut out from the sintered body in the same manner as in Embodiment
1, and superplastic deformation allowing maximal utilization of
grain boundary sliding was employed. In order to strengthen weak
portions, such as interface boundaries between precipitates and
vanadium matrix phase that generally work as crack initiation
sites, a GSMM treatment was carried out at temperature of
1300.degree. C. and a strain rate of 0.5 to 2.times.10.sup.-4
s.sup.-1. The resulting test piece was subjected to tensile testing
under conditions of room temperature and an initial strain rate of
1.times.10.sup.-3/s using a Servopulser EHF-2 model device
manufactured by Shimadzu Corporation. The yield strength (=0.2%
proof strength), tensile strength, uniform elongation and
elongation at break (total elongation) were measured.
Comparative Example 3
[0095] Production of test pieces and measurements were carried out
using the same procedure as in Embodiment 8, with the exception at
a GSMM treatment was not carried out.
Embodiment 9
[0096] SUS316L (316L stainless steel alloy powder supplied by
Hoganas AB) and TiC were used at a SUS316L:TiC mass ratio of 98:2
as alloy raw material powder. A degassing treatment was carried out
for 1.5 h at 450.degree. C. After an MA treatment, a
dehydrogenation treatment was carried out for 1.5 h at 450.degree.
C. The heating temperature while sealed under vacuum in the HIP
capsule was 750.degree. C., and a HIP treatment was carried out for
3 h at 850 to 900.degree. C. In addition, a GSM treatment was
carried out at 950.degree. C. With these exceptions test pieces
were produced and measurements were carried out in the same manner
as in Embodiment 8.
Comparative Example 4
[0097] Production of test pieces and measurements were carried out
using the same procedure as in Embodiment 9, with the exception
that a GSMM treatment was not used.
[0098] The results of measurement in Embodiments 8 to 9 and in
Comparative Examples 3 and 4 are shown in Table 5.
TABLE-US-00005 TABLE 5 0.2% Proof Tensile Uniform Elongation
strength strength elongation at break (GPa) (GPa) (%) (%)
Embodiment 8 0.66 0.7 14 20 9 0.84 1.13 20 23 Comparative 3 0.71
0.72 5 9 Example 4 0.68 0.95 10 10
[0099] As is clear from Table 5, with vanadium alloys and stainless
steel alloys that were subjected to the GSMM treatment, uniform
elongation and elongation at break were at least doubled, and it
was clear that the GSMM treatment can improve the ductility
characteristics of various types of metals or alloys, including
tungsten.
INDUSTRIAL APPLICABILITY
[0100] Because low-temperature embrittlement, recrystallization
embrittlement, and irradiation embrittlement of alloys,
particularly tungsten, can be dramatically remedied by subjecting
the alloy to a GSMM treatment, new avenues for the utilization of
alloys, particularly tungsten, are expected in regard to use in
extreme environments involving exposure to severe thermal loads,
such as with high-temperature structural materials, molybdenum
substitute materials, plasma-facing materials for international
thermonuclear experimental reactor (ITER), high-temperature test
fixtures, and solid rotating targets for spallation neutron
sources.
* * * * *