Note: Descriptions are shown in the official language in which they were submitted.
TiAl INTERNETALLIC ARTICLES AND
METHOD OF MAXING SAME
Field of the Invention
The present invention relates to a method of
maXing articles based on TiAl intermetallic materials
and, more particularly, to TiAl intermetallic base
articles having a net or near-net shape for an
intended service application and having improved
strength.
:
:
~ 15 Background of the Invention
: :,
For the past several years, extensive
research has been devoted to the development of
. ~.. intermetallio materials, suoh as titanium aluminides,
for u e in the manufacture of light weight structural
comp~nents capable of withstanding high
temperatures/stresses. Suchl components are
represented, for example, by blades, vanesj disks,
sha~ts, casings and other components of the turbine
~: 25 s~ction of modern gas turbine engines where higher gas
: and resultant~component temperatures are desired to
incr~ase engine thrust/efficiency and oth~r
applications reguiring light weight, high temperature
materials.
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: ~ :
Pt ~
P 310 Howmet 2
Intermetallic materials, such as gamma
titanium aluminide, exhibit improved high temperature
mechanical properties, including high strength-to-
weight ratios, and oxidation resistance relative to
conven~ional hi.gh temperature titanium alloys.
However, general exploitation of these intermetallic
materials has been limited by the lack o~ strength,
room temperature ductility, and toughness, as well as
the technical challenges associated with processing
and fabricating the material into the complex end-use
shapes that are exemplified, for example, by the
: aforementioned turbine componen~s.
The Kampe et al U.S. Patent 4,915,905 issued
April 10, 1990 describes in detail the development of
various metallurgical processing techniques for
~ improving the lou (room) temperaturP duc~ility and
toughness of intermetallic material~ and increasing
their high temperature strength. The Kampe et al '905
patent relates to the rapid solidification of metalIic
matrix composites. In particular, in this patent, an
intermetallic-second phase composite is formed; for
example, by reacting second phase-forming constituents
in the presence of a solvent metal, to form in~situ
: ~
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20~73~3
P 310 Howmet 3
precipitated second phase paxticles, such as boride
dispersoids, within an inter~etallic-containing
matrix, such as titanium aluminide. The
intermetallic-second phase composite is then subjected
to rapid solidification to produce a rapidly
solidified composite. Thus, for example, a composite
comprising in-situ precipitated TiB2 particles within a
: titanium aluminide matrix may be formed and then
rapidly solidlfled to producs a rapidly solidified
: 10 powder o the composite. The powder is then
consolidated by such consolidation techniques as hot
isostatic pressing, hot extrusion and superplastic
.. ~. ....... .. ~orging to provide near~inal (i.e., neax-net) shapes.
U.S. Patent 4,836~g82 to Brupbacher et al
also relates to the rapid solidification of metal
matrix composites wherein second phase-forming
: con~tituents are reacted in the presence of a solvent
metal to ~orm in-situ precipitated second phase
particles, such as TiB2 or TiC, within the solvent
me~al, such as aluminum. .
U.S. Patents 4,774,052 and 4,916,029 to
Nagle et al are specifically directed toward the
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P-310 Howmet 4
2 ~ ~ 7 3 !1 3
production of metal matrix second phase composites in
which the metallic matrix compris~s an intermetallic
material, such as titanium aluminide. In one
embodiment, a first composite is formed which
comprises a disparsion of second phase particles, such
as TiBz, within a metal or alloy matrix, such as Al~
This composite i9 then introduced into an additional
.
metal which is reactive with the matrix to form an
intermetallic matrix. For example, a first composite
comprising a dispersion of TiB2 particl~s within an Al
matrix may be introduced into molten titanium to form
a final composite compri~ing TiB2 dispersed within a
_ titanium aluminide matrix. U.S. Patent 4,9~5,903 to
Brupbacher et al describes a modification of the
method~ tau~ht in the aforemen~ioned Nagle et al
pat~nts.
'
An attempt to improve room temperature
: ductility by alloying intermetallic materials with one
or more metals in combination with certain plastic
forming techniques is disclosed in the Blackburn U.S.
Patent 4,294,615 wherein vanadium was added to a TiAl
composition to yield a modified composition of Ti-31
to 36% Al-0 to 4% V. The modified composition was
,~
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P-310 Howmet 5 2~7373
; melted and isothermally forged to shape in a heated
die at a slow deforma~ion rate necessitated by the
dependency of ductility of the intermetallic material
: on strain rate. The isothermal forging process is
carried out at above 1000C such that special die
material~ (e.g., a Mo alloy known as TZM) must be
used. Generally, it is extremely difficult to process
TiAl intermetallic materlals in this way as a result
:of their high strength, high temperature nature and
the dependence of their ductility on strain rate.
A series of U.S. patents comprising U.S.
. . Patents 4,836,983; 4,842,817; 4,842,819; 4,842,820;
4,857,268; 4,879,092; 4,897,127; 4,~902,474; and
4,916,028, have described a1:tempts to make gamma TiAl
iDtermetalli materials having both a modified
stoichiometri~c ratio of Ti/Al and one or more alloyant
additions to improve room temperature strength and
; ductility. In making cylindrical shapes from these
modified compositions, the alloy was typically first
made into an in~ot by electro-arc melting. The ingot
was melted and melt spun to form rapidly solidified ~:
ribbon. The ribbon was placed in a suitable container
and hot isostatically pressed (HIP'ped) to form a
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P-310 Howmet ~ 3~3
consolidated cylindrical plug. The plug was placed
axially into a central opening of a billet and sealed
therein. The billet was heated to 975C for 3 hours
and extruded through a die to provide a reduction of
about 7 to 1. Samples ~rom the extruded plug were
removed from the billet and heat treated and aged.
U.S. Patent 4,916,028 (included in the
series of patents listed above) also refers to
. 10 processing the TiAl basa alloys as modi~ied to include
: C, Cr and Nb additions by ingot metallurgy to achieve
~: desirable combinations of ductility, ætrength and
. other properties at a lower processing cost than the
aforementioned rapid solidification approach. In
particular, the ingot metallurgy approach described in
the '02~ paten involves melting the modified alloy
and~solidifying it into a hockey puck-shaped ingot of
simple geometry and small size (e.g., 2 inches in
diameter and .5 inch thick), homogenizing the ingot at
20 1250C for 2 hours, enclosing the ingot in a steel
annulus, and then hot forging the annulus/ring
assembly to provide a 50% reduction in ingot
thickness. Tensile specimens cut from the ingot were
annealed at various temperatures above 1225C prior to
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P-310 Howmet 7 2~r~`7373
t~nsile testing. Tensile specimens prepared by this
ingot metallurgy approach exhibited lower yield
strengths but greater ductility than specimens
prepared by the rapid solidification approach.
Despite the improvements described hereabove
~ in the ductility and strength of intermetallic
: materials, there is a continuing desire and need in
the high performance material-using industries,
~: 10 especially in the ga~ turbine engine industry, for
intermetallic materials with improved properties or
combinations of properties and also ~or manufacturing
-- technology that will allow t:he fabrication of such
intermetallic materials into usable, complex
engineered end-use shapes on a relatively high volume
basis at much lower cost. It is an object o~ the
~ ~ present invention to satisfy these desires and needs.
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Summary of the Invention
The present invention involves a method of
making titanium aluminide base intermetallic articles
having a net or near-net shape for intended service
application and having improved strength. The method
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P-310 Howmet 8 ~7373
of the present invention involves forming a
titanium-aluminum melt comprising (in atomic %) Ti in
an amount of about 40% to about 52%, Al in an amount
of about 44% to about 52~, and one or more of Cr, C,
Ga, Mo, Mn, Nb, Ni~ Si, Ta, V, and W each in an amount
o~ about 0.05% to about 8~. Boride dispersoids are
provided in the melt in an amount of at least about .5
volume % of the melt. Preferably, a low volume % of
boride dispersoids in the range of abou~ . 5 to about
2.0 volume % is provided in the malt.
The dispersoid-containing melt i5 cast and
; , . solidified-in a mold cavity of a ceramic investment
mold wherein the mold cavity is configured in the net
or near-net shape o~ the article to ~e cast. The melt
is solidified in a manner to yield a cra~k-free, net
or near net shape cast article compriRinq a titanium
aluminide-containing matrix (e.g., gamma TiAl) havlng :
a grain size of about 50 to about 250 microns as a
result of grain ref;nement from the boride dispersoids
being-distributed throughout the melt during
solidification. The melt is solidified in the mold at
a cooling rate sufficiently fast to avoid migration of
the boride dispersoids to the grain boundaries during
P-310 Howmet 9 2~7~7~
solidification and yet sufficiently slow to avoid
cracking of th~ article. A cooling rate in the range
of about 102 tG about 103F/second is preferred to
this end. Following solidification, the net or
near-net shape, investment cast article may be
subjected to a consolidation operation to olose any
porosity in the as-cast condition. The consolidated
article may then be heat treated to provide at least a
partially equiaxed grain morphology.
In one embodiment of the invention, the
; boride dispersoids are provided in the melt by
- -introducin~ a preformed ~oride master material to the
melt. In another embodiment of the invention, the
boride dispersoids are provided in the melt by
introducing an effective amount of elemental boron in
the melt to form the desired volume % of borides
in situ therein. Regardless of how the boride
dispersoids are~provided in the melt, the melt is
maintained at a selected superheat temperature for a
given melt hold time prior to casting to avoid
deleterious coarsening (growth) of the boride
particles ~dispersoids) present in the melt.
7 3 ~ ~
P-310 Howmet 10
The present invention also invulves a
titanium aluminide base article having a net or
near-net investment cast shape for intended service
application and a titanium aluminide-containing matrix
(e.g., gamma TiAl) consisting essentially of (in
atomic %) about 40% to about 52% Ti, about 44% to
about 52% Al and one or more of Cr, C, Ga, Mo, Mn, Nb,
5i, Ta, V and W each included in an amount of about
~ 0.05% to about 8%. The matrix includes at least about
: 10 .5 volume % boride dispersoids distributed uniformly
throughout and a fine, equiaxed, grain structure have
: a grain size of about 10 to about 250 microns.
~- Preferably, the article, as consolidated and heat
treated to provide:the partially equiaxed grain
structure, exhibits a yield strength at room
~: temperature ~70F) of at least about 55 ksi and a
tensile ductility at room temperatura of at least
~.
~ bout 0.5% (measured by the ASTM E8M test procedure).
', :
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20: Thus, the present invention has as a
: particular purpose to provide net or near-net shape
:
articles o~ a TiAl base intermetallic material
modified by the addition of selected
alloyant(s)/dispersoids and~formed to shape by
,. .
P-310 Howmet 11 2~7373
investment casting in a crack-free condition treatable
by consolidation/heat treatment to exhibit improved
strength and ductility at room temperature. The
method of the invention provides an alternative to
much more costly techniques heretofore used to
fabricate TiAl base intermetallics.
'
The advantag4s of the present invention will
become more readily understood by consideration of the
following detailed description and examples.
BrieP Description of the Drawinqs
:
Figure 1 is a flow sheet illustrating one
embodiment of the method of the invention.
Figures 2A through 2F are photomicrographs
of in~estment castings of Alloys A through E,
respectively, illustrating the effect of increasing
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~ 20 boron in the melt on grain refinement.
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Figures 3A-3B are photomicrographs of the
microstructures of investment castings illustrating
the affect of heat treatment under different
P-310 Howmet 12 ~7~73
conditions on grain morphology.
Figures 4A-4C are photomicrographs
illustrating ths boride dispersoids present in a
particular Alloy D investment casting.
Figures 5A-SC are photomicrographs
illustrating the boride dispersoids present in a
particular Alloy E investment casting.
Figures 6A-6C are photomicrographs
illustratinq the boride dispersoids present in a
- - particular Alloy F investment casting.
' -
lS Figures 7A-7F are photomicrographs of
investment castings illustrating the effect of
increasiny borides (added by ma~t~r borid~ material)
in the melt on grain refinement.
~':
: ~ 20 Figure 8A 8F are photomicrographs of the
as-cast microstructures of the investment castings of
Figs. 6A-6F.
Figure 9A-9F are photomicrographs of the hot
.
2~ 7 3
P-310 Howmet 13
isostatically pressed microstructures of the
investment castings of Figs. 7A-7F.
Figures lOA-lOB are photomicrographs o~ the
microstructures of investment castings illustrating
the effect of heat treatment under different
conditions on grain morphology.
: Figures llA-llC are photomicrographs
illustrating the boride dispersoids present in a
particular Alloy lXD invsstment casting (as-cast).
': '
Figures 12A-12C are photomicrographs
illustrating the boride dispersoids present in a
particular Alloy 2XD investment casting (as-cast).
Figures 13A-13C are photomicrographs
illustrating the boride dispersoids present in a
particular Alloy 3XD investment casting (as-cast).
Figures 14A-14C are photomicrographs
illustrating the boride ~ispersoids present in a
particular Alloy 5Xd investment casting (as-cast).
~73~3
P-310 Howmet 1~
Figures 15A-15C are photomicrographs
illustrating boride particle~ extracted from the Alloy
2XD investment casting (as-cast~.
Figures 16A-16C are photomicrographs
illustrating boride particles extracted from the Alloy
3XD investment casting (as-cast).
:`:
: ~ Figure 17 is a schematic illustration of
boride particles o~ various morphology that occur in
the investment castings.
~ ~ - Detailed Description of the Invention
.: .
The present invention relates to net or
near-net shape articles comprised of a titanium
aluminide base intermetallic material modified by the
:~ addition of selected alloyant(s)/dispersoids and ~
~ormed to shape by investment casting in a crack-free, . .
fine grained condition treatable by consolidationtheat
treatment to exhibit improved strength at room
temperature. Titanium-aluminum base alloys employed
in practicing the present invention consist
essentially of, by atomic %, about 40% to about 52%
P~310 Howmet 15 2 ~ ~ 7 3 7 3
Ti, about 4~% to about 52% Al and one or more of the
alloyants Cr, C, ~a, Mo, Mn, Nb, Ni, Si, Ta, V,and W
each in an amount of about 0.05% to about 8%. The
listed alloyants are provided in the base composition
as a result of their beneficial effect on ductility
when present in certain combinations and/or
concen~rations.
preferred base alloy ~or use in practicing
the present invention consists essentially of, by
.
atomic %, about 44% to about 50% Ti, about 46% to
about 49% Al and at least one of Cr, C, &a, Mo, Mn,
~ Nb, Ni, si, Ta, V,and W wherein Cr, Ga, Mo, Mn, Nb,
;~ : Ta, V and W, when present, are each included in an
amount of about 1% to about 5% and wherein C, Ni and
Si, when present, are each 1ncluded in an amount of
about O 05% to about 1.0%. Two or more of the
;~ ; alloyants~Cr, C, Ga, Mo, Mn, Nb, Si, Ta, V and W are
: : present in an even more preferred embodiment within
the concentration ranges given. Although the present
: invention is~not limited to a particular base
composition within the ranges set forth hereabove,
certain specific preferred base compositions are
described in the Examples set forth hereinbelow.
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P-310 Howmet 16 ~37373
Referring to Figure 1, the various steps
involved in practicing one embodiment of the method of
the invention are illus rated. In this embodiment, a
melt of the TiAl base alloy is formed in a suitable
container, such as a crucible, by a variety of melting
techniques including, but not limited to, vacuum arc
melting (VAR), vacuum induction melting (VIM~,
induction skull melting (ISR~, electron beam melting
~: (EB?, and plasma arc melting (PAM). In the vacuum arc
: 1~ melting technique, an electrode .is fabricated of the
base alloy composition and is melted by direct
electrical arc heating (i.e., an arc astablished
between the electrode and the crucible) into an
undexlying non-reactive crucible. An actively cooled
copper crucible is useful in this regardO Vacuum
induction melting involves heating and melting a
charge of the base alloy in a non-reactive,~ refractory
::crucibla by induction heating th charge usin~ a
surrounding electrically energized induction coil.
Induction skull melting involves inductively heating
and melting a charge o~ the base alloy in a
water-cooled, segmented, non contaminating copper
crucible surrounded by a suitable induction coil.
Electron beam melting and plasma melting involve
;
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P-310 Howmet 17 ~73~3
melting using a configuration of electron beam(s) or a
plasma plume directed on a charge in an actively
: cooled copper crucible. These melting techniques are
known genarally in the art of melting o~ metals and
alloys.
. ' :
Although the present invention is not
limited to any particular melting process, certain
;~ specific melting processes are described in the
: 10 Examples set forth hereinbelow.
~: Referring again to Figure l, the melt of the
TiAl base alloy in the container (crucible) is
provided with boride disper~;oids in an amount o~ at
least about 0.5 volume % prior to casting of the melt
in an investment mold to be described in detail
herebelow. Typically, the boride dispersoids comprise
simple titanium borides (TiB2) and/or complex borides
such as (Ti,M)XBy where M is Nb, W, Ta or other
al}oyant. Although varying amounts of the boride
disparsoids may be used depending upon the end-use
properties desired ~or the cast article, relatively
low boride dispersoids levels of about 0.5 to about
20.0 volume % are useful in practicing the invention
P-310 ~owmet 18 ~ 3 7
to achieve the desired grain refinement effects in the
casting as well as s~rength and ductility improvements
upon further treatment of the casting. Boride
dispersoid levels above the upper limit set ~orth tend
; 5 to reduce ductility and thus are not preferxed. In
accordance with the invention, optimum strength and
ductility are achieved when the boride dispersoid
level is preferably about 0.5 to about 2.0 volume % of
the melt or cast article.
}O
The TiAl base alloy melt described hereabove
can be provided with the desired level of boride
dispersoids in a variety of ways including the
addition of a boride mastar material to the melt in
accordance with U.S. Patents 4,751,048 and 4,916,030,
the teachings of which are incorporated herein by
reference. In particular, a porous sponge having a
relatively high concentration of boride particles
(e.g., ~iB~) is introduced and incorporated in the TiAl
base melt to provide a lower concentration of boride
particles therein~ Of course, the concentration of
boride particles in the sponge is chosen to yleld a
selectad lower concentration of the particles in the
melt; for example, at least about .5 volume ~ boride
.
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P-310 ~owmet 19 20~7373
dispersoids in the melt. Boride mastPr materials
(i.e. sponges) useful in practicing the present
invention are available from Martin Marietta
Corporation, Bethasda, Md and its licensees.
The TiAl base alloy melt also can be
provided with the desired level of boride dispersoids
by providing an effective amount o~ elemental boron in
the melt to form and precipitate the aforementioned
simple and/or complex titanium boride particles
in-situ therein. When using the VAR melting process
to form the TiAl base melt, elemental boron can be
: -provided-in the melt by dispersing elemental boron in
the VAR electrode with the other alloyants as
described in the Examples hierebelow. When the
electrode is melted into thle underlying crucible, the
TiAl base composition and the boron are brought
together in the melt so that the boron can react with
metals in the melt to precipitatP simple borides
(e.g., TiB2) and/or complex borid s (e.g., Ti,Nb)XBy
in the melt. When using the vacuum induction,
induction skull, electron beam and plasma mel~ing
processes referred to hereabove, the elemental boron
an be provided in the melt hy blending with the
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P-310 Howmet 20 ~7373
~ initial alloyants of the charge to be melted or by
: addition to the already melted alloy charge.
~
,
Other methods of providing the desired level
- 5 of boride dispersoids in the melt are described in U.
~: S. Patent ~,915,052 and 4,916,029, although the
present invention is not limited to any particular
technique in this regard.
Importantly, the dispersoid-containing TiAl
base alloy melt is maintained at a selected superheat
temperature (~or a given melt hold time prior to :~ -
- casting) to avoid growth of th~ boride particles
present in the melt to a harmful size. Namely, the
~ 15 superheat of the melt is mai.ntained sufficienkly low
; so as to avoid formation of deleterious TiB needles
(whiskers) having a length ~reater than about 50
microns~ These TiB needles form from the existing TiB2
particles in the melt by particle growth pxocesses and
are quité harmful to the properties,:especially the
ductility, of the casting. In general, the superheat
temperature of the melt is maintained at the melting
temperature of the TiAl base composition plus about 25
to 200F thereabove to this end. Temperature
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P-310 Howmet 21 ~37373
maintenance in this manner ~osters the presence of
blocky (e.g., equiaxed), lacey and/or small needles
(less than about 50 microns length) of TiB2 in the
melt~ Such boride particles are illustrated
schematically in Fi~. 17.
.
Pre~erably, the dispersoid-containing TiAl
base alloy melt is stirred in the crucible prior to
casting. When tha aforementioned VAR, VIM, ISR and
other melting techniques are used, the melt is stirred
in the crucible by the action o~ an induction heating
coil on the melt. Melt stirring in this manner
aintains a homogenous ~elt with the boride
d~spersoids distributed uniformly throughout.
Melting and casting o~ the TiAl base alloy
containing the boride dispersoids is conducted under
relative vacuum (e.g., 1 micron vacuum) or under inert
: atmosphere (e.g., .5 atmosphere Ar) to minimize ~:
70 contamination of the melt.
The dispersoid-containing TiAl base alloy
melt is cast into a non-reactive, ceramic investment
mold having one or more mold cavities configured in
.
P-310 Howmet 22
2~7373
the net or near-net shape of the article to be cast.
- Net shape castings require no machining to achieve
final print dimensions/tolerances~ Near-net shape
castings may require only a minor machining operation
of the casting, or portion thereof, to provide final
print dimensions/tolerances. Investment molds used in
~ practicing the invention are made in accordance with
;~ conventional mold forming processes wherein a ~ugative
pattern (e.g., a wax pattern) having the near-net
shape to be cast is repeatedly dipped in a ceramic
~ slurry, stuccoed with ceramic particulate and then
; dried to build up a suitable shell mold about the
pattern. After the desired thickness of the shell
mold is formed, the pattern is removed ~rom the mold,
leaving one or more mold cavities therein. When wax
patterns axe used, the patterns can be removed by
known dewaxing techniques, such as steam autoclave
dewaxing, ~lash dewaxing in a furnace and the like.
After pattern removal, the shall mold is treated at
elevated temperatures to remove absorbed water and
gases there~rom. Although the invention 1s not
limited to any particular mold formation process,
certain specific mold formation processes are set
forth in the Examples herebelow.
,
P-310 Howmet 23 2~737~ :
The investment mold is made from ceramic
materials which will be substantially nonreactive with
the TiAl base alloy melt so as not react with and
contaminate the melt. In particular, the mold
facecoat that contacts the melt typically comprises a
ceramic material selected Prom zirconia, yttria and
the like to this end~ The mold coats subsequently
applied to the facecoat (i.e., the backup coats) may
be selected from a variety of ceramic materials
depending upon the particular casting application
involved. The investment mold may be made in various
configurations as needed for a particular casting
application.
Referring to Figure 1, the dispersoid-
containing TiAl base alloy mslt at the appropriate
superheat temperature is cast (e.g., poured) from the
melting crucible into a preheated investment mold and
; solidified therein to form a net or near-net shape,
cast article whose microstructure will be describad in
detail herebelow~ The melt may be gravity or
countergravity cast into an investment mold that is
stationary or that is rotat~d as, ~or example, in
centri~ugal casting processes. Regardless of the
,
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P-310 Howmet 24
2~3~3
~- casting me~hod employed, the cooling (freezing) rate
of the melt and cooling rate of the casting are
controlled so as to be fast enough to prevent
migration and segregation of the boride dispersoids to
the grain boundaries and yet slow enough to avoid
cracking of the solidified casting. The cooling rate
employed will depend upon the melt superheat, the
section size of the casting to be produced, the
configuration of the casting to be produced, the
particular TiAl base alloy composition, the loading
level o~ dispersoids in the melt as well as other
factors. In general, cooling rates of about 102 to
- about 1030F per second are employed to this end. Such
cooling rates are typically achieved by placing the
melt-filled investment mold in a bed of refractory
material (e.g., Al~03) and allowing the melt to
~solidify to ambient temperature. Once the casting has
cooled to ambient temperature ~or other demold
temperature~, the casting and the investment mold are
separated in usual manner, such as by vibration.
: ,
Referring again to Figure 1, following
separation o~ the mold and~the casting, the casting
may be subjected to a consolidation operation to close
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P-310 Howmet 25 ~7373
any porosity in the casting. Preferably, the casting
is hot isostatically pressed at, for example,
2100-2400F and a pressure o~ 10-45 ksi for 1-10 hours
depending on the size of the casting, to close any
: 5 porosity present in casting. Thereafter, the HIP'ped
casting i~ heat treated to provide at least a
partially equiaxed grain structure in lieu of the
lamellar grain structure present in the as-cast
microstructure. Heat treat parameters of 1600-2500F
for 1-75 hours may be used. Of course, other
consolidation processes/parameters and heat treatment
procasses/parameters can be employed in practicing the
invention.
The titanlum alum:inide base casting produced
in accordance with the present invention is
characterized as having a net or near-net shape for
the intended servic~ application and a predominantly
gamma TiAl intermetallic matrix corresponding in
composition to that of the base composition. The
matrix exhibits a fine, as-cast grain structure of
lamellar morphology and a grain siz~ within the range
of about 10 to about 250 microns, preferably about 50
to about 150 microns. The matrix may include other
.:
P-310 Howmet 26 2~7373
titanium aluminide phases (e.g., Ti3A1 or TiAl3) in
minor amounts such as up to about 15.0 volume %. The
as-ca~t lamellar grain structure is changed to a
partially equiaxed grain structure by the subsequent
heat treatment operation.
~s will become apparent from the Examples
~ set forth herebelow, a certain minimu= level of boride
; dispersoids, such as at least about 0.5 volume %
dispersoids, must be uniformly distributed throughout
the melt during solidification in order to achieve a
grain refinement effect that yields as-cast and heat
treated grain sizes in the aforementioned ranges for
strength enhancement purposes. Dispersoid levels
below the minimum level are .ineffective to produce the
fine as-cast grain sizes required for improved
strength. The dispersoids are distributed generally
uniformly throughout the as-cast matrix (as shown in
~: Figs. 5, 6, 13 and 14) and are not segregated at the
grain boundaries.
As will also becoma apparent from the
Examples set for~h herebelow, the boride dispersoids
are present in the matrix in various morphologies
including a) ribbon shapes generally 0.1-2.0 microns
,, . ;
~, .
P-310 Howmet 27 20~7373
thick, 0.2-5.0 microns wide and 5.0-1000 microns long,
b) blacky (equiaxed) shapes generally of 0.1-50O0
microns average size (major particle dimension), c)
needle shapes generally 0.1-5.0 microns wide and
5.0-50.0 microns long, and d) acicular shapes
generally 1.0-10.0 microns wide and 5.0-30.0 microns
long. These various dispersoids particle shapes are
illustrated schematically in Fig. 17. As mentioned
hereabove, large TiB needles having a length greater
than about 50 microns are to be avoided in the matrix
so as not to adversely affect the ductility of the
casting.
;:
Consolidated and heat treated TiAl
intermetallic base investment castings in accordance
with the invention typically exhibit a yield strength
at room temperature (70F) of at lea~t about 55 ksi
and a ductility~at room temperature o~ at least 0.5%
as measured by the ASTM E8M test procedure.
Consolida ed and heat treated TiAl intermetallic base
investment castings of the invention having the
aforementioned even more pre~erred composition
typically exhibit a yield strength at room temperature
(70F) of at least about 60 ksi and a ductility at
` ' ,. ..
:
.
P-310 Howmet 28
2~5~73
room temperature of at least about 1.0% as measured by
the same ~STM test procedure. These room temperature
properties represent a substantial improvement over
the room temperature properties demonstrated
heretofore by investment cast TiAl intermetallic
materials which have not been modified by addition of
borides or boron.
- ~ The following Examples are offered to
illustrate the invention in further de~ail without
limiting the scope thereof.
- EXA~PLE 1
This example illustrates practice of one
embodiment of the invention wherein elemental boron is
provided in the TiAl base alloy melt in order to form
boride di~persoids in-situ therein. Various amounts
of elemental boron were provided in the TiAl base melt
to determine the dependence of grain refinement on the
amount of boride dispersoids present in the melt. The
following melt compositions were prepared by the VAR
melting process referred to hereabove:
~ . . .
P-310 Howmet 29 2~7373
Alloy A--- Ti 47.1% Al-2.1% Nb-1.6~ Mn-0.047% B
(0.04 v/o borides3
Alloy B -- Ti-47.8% Al-2.1% Nb-2.4% Mn-0.11% B
(0~07 v/o borides)
Alloy C--- Ti-46.9~ Al-2.0% Nb-1.7~ Mn 0.17% B
(0.13 v/o borides)
10 Alloy D--- Ti-47.2% Al-2.0% Nb-1.5% Mn-0.3~ B
(0.27 v/o borides or 0.30 atomic % B)
- Alloy E--- Ti-48.4~ Al-2.0~ Nb-1.5% Mn-1.0% B
(0.70 v/o borides or 1.0 atomic ~ B~
Alloy F~-- Ti-45 . 3~ Al-l . 9% Nb~ % Mn-2.49% B
(1.94 v/o borides or 2.5 atomic % Bj
: ~ :
~ : A cylindrical electrode of each of these TiAl base
; 20 alloy compositions was prepared by cold pressing Ti
sponge, Al pellets~ A1/Nb master alloy chunks, Al/Mn
master alloy chunks and elemental boron powder in the
appropriate amounts in a Ti tube. The cold pressed
body was subjected to a ~irst meltiny operation to
' ' :' ' ;
~ "~ ., '
~ P-~10 Howmet 30 2~57373
produce an ingot. The ingot was grit blasted and then
remelted again to produce the electrode~ Each
electrode was then VAR melted into a copper crucible
to form a TiAl base alloy melt in which elemental
boron was present.
Each TiAl alloy melt was maintained at a
superheat temperatuxe of about 25F above the melting
point by VAR melting prior to casting. Agitation
during VAR melting also acted to stir the melt prior
to casting. Each melt was poured from the crucible
into a preheated (600F) cexamic investment mold
compri~ing a Zr203 mold facecoat for contacting the
melt and nine backup coats of Al203. Each mold
included five mold cavities in the shape of cylinders
having the following dimensions: 0.625 inch diameter x
8;inches long. Each melt was melted and cast into the:
mold under 7 microns vacuum. Each melt-filled molds
was placad in~a bed o~ Al203 (to a depth of about 8
inches) and allowed to cool to ambient temperature
:
over a period of about Z hours. Each mold and the
cylindrical-shaped casting were then separated.
Figures 2A-2F illustrate the effect of boron
:: . :' ~, :
P-310 Howmet 31
2~737'~
concentration (expressed in atomic %) of the base
alloy composition and of volume % boride dispersoids
in the castings on the as-cast grain structure. It is
evident that little or no grain refinement was
observed in Figs. 2A through Fig. 2D for the Alloy A,
B, C and D castings. on the other hand, dramatic
grain refinement was present in the Alloy E and F
castings as shown in Fig. 2E and Fig. 2F. The
transition from no observed grain re~inement to
dramatic grain refinement occurred between Alloy D
(0.3 atomic % B) and Alloy E (1.0 atomic % B). The
grain size of Alloy E casting and Alloy F casting were
about 50 t~ about 150 microns, respectively.
.
Alloy E castings were hot isostatically
pressed at 2300F and 25 ksi. for 4 hour~ and then
subjected to different heat treatments to determine
: response of the as-cast lamellar grain struc~ure to
dif~erent temperatures. Figures 3A and 3B illustrate
20~ the change in grain structure from lamellar to
partially equiaxed after heat treatments at 2100F and
1850F with the same time-at-temperature and gradual
~urnace cool (GFC). The change from lamellar to
partially equiaxed grain structure is evident in both
" . ' ,.
,~ . ~ - :.':
,
~7~73
P-310 Howmet 32
Figs. 3A,3B.
Figures 4A-4C, 5A-5C, and 6A-6C illustrate
the effects of boron concentration on the appearance
of boride dispersoids in Alloys D, E and F,
respectively, as consolidated/heat treated. Three
different known electron microprobe techniques were
used to vi w the dispersoids; namely, the secondary
technique, the back sca~ter technique and the boron
: 10 dot map. Based upon these Figures, the solubility of
borcn in the Ti-Al-Nb-Mn compositions set forth above
appears to be less than 0.05 atomic % B.
Table 1 sets forth strength and ductility
properties of tho Alloy A, B, D, and E castin~s after
HIP'ing at 2300F and 25 ksi for 4 hours followed by
heat treatment at 1850F for 50 hours in an inert
~atmosphere. Included for comparison purposes in Table
i9 a base alloy (T1-48~Al-2%Nb-2%Mn-0%~) HIP'ed
~: 20 using the same parameters and heat treated to a
qimilar microstructure. Tensile t~sts were conducted
at room (70F) temperatuxe in accordance with ASTM E8M
test procedure and at 1500F in accordance with ~STM
E21 test procedure.
::
37~3
P-310 Howmet 33
Table 1
TEST TEMP. I~S Y~ ELONG.
(F) (KSI) (KSI) (%)
Base Alloy 70 58.0 40.01.7
1500 50.0 37.030.0
Alloy A 70 62.2 52 8l 0
1500 54.4 4~ 644 7
Alloy ~ 70 52.2 46.10.6
1500 62.~ 45.26.8
Alloy D 70 54.3 50 00 5
1500 61.3 39 717 1
Alloy E 70 69.4 59.20.7
15~0 66.1 45.220.7
This combination strength and ductility properties
represent significa~t improvements over those
obtainable heretofore in the casting of gamma titanium
aluminide (TiAl).
:
:
EXAMPLE 2
: : This example illustrates practice of another
:: : embodiment of the invention wherein preformed boride
dispersoids (TiB2) are provided in the TiAl base alloy
melt by adding a master boride material thereto. The
master boride material comprised a porous sponge
having 70 weight % of borides (TiB2) in an Al matrix
.
'
3 ~ 3
P-310 Howmet 34
metal. Various amounts o~ the sponge material were
added to tha TiAl base alloy melt so as to determine
the dependence of grain refinement on the amount
(volume %) of boride dispersoids present in the melt.
The following melt compositions were prepared by the
VAR melting process referred to hereabove:
Alloy OXD--~ Ti-45.4~ Al-1.9% Nb-1.4% Mn-O vol.%
TiB2 ( at.% ~)
; Alloy lXD--- Ti-45.4% A1-1,9% Nb-1.4~ Mn- 0.1 vol.%
TiB2 (.17 at.% B or 0.1 volume % borides)
.
Alloy 2XD--- Ti-46.1% Al-1.8% Nb-1,6% Mn-0.4 vol.% TiB2
(.50 at.% B or 0.4 volume % borides)
.
~: Alloy 3XD--- Ti-47.7% A1-2.0% Nb-2.0% Mn-l.O vol.% TiB2
(1.40 at. % B or 1.0 volume % borides)
20 Alloy 4XD--- Ti-44.2~ Al-2.0~ N~-1.4% Mn-2.0 vol.% TiB2
(2.59 at. % B)
:
Alloy 5XD--- Ti-45.4% Al-1.9% Nb 1.6% Mn-~.6 vol.% TiB2
(5.97 at.% B or 4.6 volume % borides)
,
',
~7373
P-310 Howmet 35
Interstitial concentrations in these alloys are set
forth below:
INTERSTITIALS (ppm wt%)
O N H
Alloy OXD--- 716 42 6
:Alloy lXD--- 632 58 9
Alloy 2XD -~ 684 68 14
Alloy 3XD--- 538 47 10
Alloy 4XD -- 795 90 10
Alloy 5XD--- 654 48 13
Each of these TiAl base alloy compositions
was fabricated into a cylindrical alectrode by the
procedure described hereina]bove for Example 1. After
double melti~g as described above, each electrode was
subjected to a surface treatment operation using a SiC
grinding tool, grit blasting (or alternatively
chemical milling operatlon using 10~ HF aqueous
solution as an etchant) to remove surface oxidation
therefrom. About a .020 inch depth was removed from
the electrode. Each electrode was then VAR melted by
direct electric arc heating into a copper crucible to
form a TiAl base alloy melt to which the preformed
~ .
2 ~ 7 3
P-310 Howmet 36
ma~ter sponge was added.
:
Each TiAl alloy melt was maintained at a
superheat temperature of about 25 F above the alloy
melting point by electric arc melting prior to
castinq. Each melt was poured from the crucible into
a preheated (600F) ceramic investment mold comprising
a Zr2o3 mold facecoat for contacting the melt and nine
~backup coats of Al2030 Each mold included five mold
cavities in the shape of cylinder~ haviny the
following dimensions: 0.625 inch diameter x 8 inches
long. Each melt was melted and cast into the mold
- under a 7-micron vacuum. Each melt-filled mold was
placed in a bed of Al203 (to a depth of about 8 inches)
and allowed to cool to ambient temperature over a
period of about 2 hours. Each mold and the
cylindrical-shaped castings were then separated.
: Figures 7A-7F illustrate the effect of
2Q :boride loading (volume %) on ~he as-cast grain
structure of ~lloys lXD through 5XD, respectively. It
is evident from Figs. 7A through 7C, that little or no
grain refinement was observed for the Alloy OXD, lXD
and 2XD castings. on the other hand, dramatic grain
:
,: ;; :
2~737~
P-310 Howmet 37
refinement was present in the Alloy 3XD, 4XD and 5XD
castings as shown in Figs. 7D through 7F. The
transition from no observed grain refinement to
dramatic grain refinement oscurred between Alloy 2XD
(0.4 vol. % TiB2) and Alloy 3XD ~1.0 vol.% TiB2). The
grain size of Al.loy 3XD, 4XD and 5XD castings was
about 50 to about 150 microns.
: Figures 8A-8F illustrate the as-cast
microstructures of the castings OXD-5XD, respectively.
Figures 9A-9F illustrate the as-HIP'ped
~: - microstructures of the castings OXD-5XD, respectively.
,
Alloy 3XD castings were hot isostatically
pressed at 2300F and 25 ksi for 4 hours and then
: subjected to different heat treatments to determine
response of the as-cast lamellar grain structure to
different temperatures. Figures lOA and lOB
illustrat- the change in grain structure from lamellar
to partially equiaxed after heat treatments at 2100F
and 1850F with the same time-at-temperature and
gradual furnace cool. The change from lamellar to
equiaxed grain structure is evident in both Figs.
P-310 Howmet 38 2~7~7~
lOA, 1OB.
: ,
Figures llA-llC, 12A-12C, 13A-13C and
14A-14C illustrate the effects of boron concentration
on the appearance of boride dispersoids in Alloys lXD,
2XD, 3XD, and 5XD, respectively, as-cast. Three
different known electron microprobe techniques were
used to view the dispersoids; namely, the secondary
techniqu~, the back scatter technique and the boron
~: 10 dot map.
: `
Figures 15A-15C and 16A-16C illustrate
various TiB2 particle shape~; extracted from Alloy 2XD
and 3XD, respectively.
Table 2 sets forth strength and ductility
~: : : propertie~ of the Alloy 2XD and 3XD castin~s after
~ :
HIP'ing:at 2300F and 25 ksi ~or 4 hours followed by
~ ~ ~ : heat treatment at 1850~F for 50 hours in an inert (Ar~
:~ :atmosphere. Tensile tests were conducted at room
- ~
. (70F) temperature and at 1500F in accordance with
: ASTM E8M and E21 test procedures, respectlvely.
.; .
P-310 Howmet 39 ~737`
Table 2
TEST TEMP. UTS YS ELONG. AVERAGE
tF~(KSI) (KSI) (%) GRAIN SIZE
Alloy 2XD 7062.2 51.0 1.0 1000 um
150065.8 45.2 1~.0
Alloy 3XD 7084.4 78.2 0.7 75 um
150060.6 48.~ 8.9
This combination of strength and ductility
properties represent significant improvements over
: those heretofore obtainable in the prior art cast
gamma (TiAl) titanium a1uminide alloys.
EX~MPLE 3
~ ..
This example illustrates practice of still
another embodiment o~ the invention wherein a charge
of Ti sponge, Al:pellets, AllMn master alloy chunks,
Al/Nb master alloy chunks and elemental boron powder
are melted using:the induction skull melting
procedure. In particular, the charge was melted in a
~egmented, water-cooled copper crucible such that a
~solidified metal skull ~ormed on the crucible surfaces
shortly after melting of the meltlng of the charge.
The charge was melted~by energization of an induction
coil positioned about the crucible (see U.S. Patent
,
2~373
P-310 Howmet 40
4,923,508) and was maintained at a superheat
temperature of about 50F above the alloy melting
point by induction heatiny. The melt was stirred as a
result of the induction heating.
The melt was poured from the crucible into a
preheated (600F) ceramic investment mold comprising a
Zr203 mold facecoat for contacting the melt and nine
back up coats of~Al2O3. Each mold included 5 mold
cavities in the shape of cylinders having the
following dimensions: 0.652 inch diameter x 8 inches
long. Each melt was melted under .5 atmosphere Ar and
- ~ast into ~he mold under 200 microns vacuum. Each
melt-~illed mold was placed in a bed of Al203 (to a
depth of about 8 inches) and allowed ~o cool to
ambient temperature over a period of about 2 hours.
Each moId and the cylindrical-shaped castings were
then separated.
The following melt compositions (in atomic
%) were ISR melted and investment cast as described
above:
. ~ .
' .
P-310 Howmet 41 2~737~
Alloy 1--- Ti-45.6~ Al-1.9% Nb-2.3% Mn-1.10% B
Alloy 2--- Ti-45.1% Al 1.9~ Nb-2.2% Mn-2.4% B
Fox comparison purposes, two alloys (XD0 and
XD7) were prepared in accordance with EXample 2 to
include 0 volume % and 7 volume % titanium borides.
Table 3 sets forth room temperature strength
: 10 and cluctility properties o~ Alloys 1-2 after HIP'ing
at 2300F and 25 ksi for 4 hours followed by heat
treatment at 1650F for 24 houxs in lnert (Ar)
at~o~phere. Alloys XD0 and XD7 (Ti-48%Al-2%Nb-2~Mn
: with 0 volume % and 7 volume~ % borides, respectively)
were HIP' ed using the same parameters and heat treated
to a similar microstructure. The room temperature
tensile tests were conducted pursuant to ASTM ~8M test
procedure.
' ~ -
. .
,
P-310 Howmet 42 2~7373
: :.
Table 3
ROOM TEMPERATURE TENSILE RESULTS
BORIDE/BORON YIELD ULTIMATE PLASTIC
AMOUNT STRENGTH STRENGTH ELONGATION
XD0 0 40.0 58.0 1.7 :
XD7 7 Vol.% 65.0 79.0 0.5 :
~: Alloy 1 1.10 At%B 74.0 89.0 1.3
: Alloy 2 2.40 At%B 75.0 86~0 0.9
::
While the invention has been described in
terms of specif ic em~odiments thereof, it is not
: intended to be :limited thereto but rather only to the
- extent set forth in the following claims.
.
:
~::
: : :
:: : :
,
. ; ' ~
