Intermetallics 19 (2011) 137e142
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Intermetallics
journal homepage: www.elsevier.com/locate/intermet
Microstructure and mechanical properties of Ni3Al and Ni3Ale1B alloys fabricated
by SHS/HE
J.T. Guo a, *, L.Y. Sheng a, b, c, **, Y. Xie a, Z.X. Zhang c, V.E. Ovcharenko d, H.Q. Ye a
a
Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
Peking University, Beijing 100871, China
c
PKU-HKUST Shenzhen-HongKong Institution, Shenzhen 518057, China
d
Institute of Strength Physics and Materials Science, Russian Academy of Sciences, Siberian Branch, Tomsk, Russia
b
a r t i c l e i n f o
a b s t r a c t
Article history:
Received 12 June 2008
Accepted 31 July 2008
The well-densified Ni3Al alloys without and with boron addition were fabricated by self-propagation
high-temperature synthesis and hot extrusion (SHS/HE) technology. Microstructure investigation
showed that Ni3Al and Ni3Al-1B alloys contained fine grain structure. Analysis of X-ray spectra as well as
transmission electron microscopy studies revealed that three phases present in all alloys: g-Ni, Ni3Al and
dispersoids of a- Al2O3 and g- Al2O3. However, b-NiAl, Ni3B phase and twinned Ni3Al crystal are observed
in the Ni3Al-1B alloy. In addition, dislocations with high density exist in all alloys. The mechanical test
showed that the B addition leads to obvious improvement in yield strength and compressive ductility,
and compared with the ones synthesized by combustion, SHS/HE synthesized Ni3Al and Ni3Al-1B alloys
exhibit more excellent mechanical properties.
Ó 2010 Elsevier Ltd. All rights reserved.
Keywords:
A. Nickel aluminides, based on Ni3Al
B. Mechanical properties at ambient
temperature
C. Reaction synthesis
D. Microstructure
1. Introduction
Nickel aluminide intermetallic alloys (NiAl and Ni3Al) have
received considerable attention for high-temperature structural
and coating applications, for examples, as heat shields for
combustion chambers and as first-row vanes in industrial gas
turbines [1,2]. In addition, numerous alloys based upon Ni3Al have
been developed with broad utilizations ranging from furnace rolls
and radiant burner tubes for steel production to heat treating
fixtures, forging dies, and corrosion-resistant parts for chemical
industries [3,4]. This is because strong bonding between aluminum
and nickel, which persists at elevated temperatures, yields excellent properties competitive with those of superalloys and ceramics,
such as high melting point, low densities, high strength, as well as
good corrosion and oxidation resistance [1e4]. In spite of these
attractive properties, however, low ductility, brittle fracture, and
processing problems were the major disadvantages of nickel aluminides [3e5]. The brittleness of Ni3Al stems from an environmental effect, such as the hydrogen generated through the
* Corresponding author. Tel.: þ86 24 23971917; fax: þ86 24 83978045.
** Corresponding author. Tel.: þ86 013751855451; fax: þ86 24 83978045.
E-mail addresses: [email protected] (J.T. Guo), [email protected] (L.Y. Sheng).
0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved.
doi:10.1016/j.intermet.2010.08.027
reduction of moisture in air by aluminum in the aluminides [4,6]. A
major breakthrough to resolve this issue was the discovery of the
dramatic effects of boron addition on ductility improvement for
Ni3Al at ambient and high temperatures. Liu et al. [7] found that
40e50% tensile ductility can be achieved in the Ni3Al alloy with the
addition of small amounts of boron up to 0.4 wt.%. Interestingly, as
the Al content of B-doped Ni3Al is decreased to below 25 at.%, the
ductility increases significantly and the fracture mode changes
from brittle intergranular to ductile transgranular. However, the
ductility drops drastically to about 5% when the Al content is
increased to 25 at.% or higher [7].
Recently, Morsi [1] reviewed a number of novel processes applied
to the reaction synthesis of NieAl intermetallics. Among them,
combustion synthesis with the advantages of time and energy savings
has been recognized as a promising alternative to the conventional
methods of producing advanced materials, including carbides,
borides, nitrides, hydrides, and intermetallics, etc. [8e11]. Combustion synthesis of the Ni3Al intermetallic can be conducted in either of
two modes, the self-propagating high-temperature synthesis (SHS)
[12,13] and the thermal explosion [14,15]. By using combustion
synthesis in the SHS mode with powder compacts, Lebrat and Varma
[12] found that higher green density and preheating temperature led
to fully reacted product with a well-developed microstructure.
Though the synthesis of Ni3Al has been extensively studied, the
porosity is still its main problem. The hot extrusion could solve the
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J.T. Guo et al. / Intermetallics 19 (2011) 137e142
trouble, but it resulted in the crystal coarsening. The self-propagating high-temperature synthesis and hot extrusion (SHS/HE)
technology could well densify the alloy and also avoid the crystal
coarsening. So in the present study, the investigation of microstructure and mechanical properties of boron and chromium doped
Ni3Al alloys fabricated by SHS/HE technique was carried out.
2. Experimental
Fig. 1. Schematic diagram of the SHS/HE synthesis system.
Fig. 2. X-ray diffraction patterns of Ni3Al and Ni3Al-1B alloys prepared by SHS/HE.
Two kinds of powders, including of powders of nickel (with an
average particle size of 0.98 mm) and aluminum (1.45 mm), were
used as the initial materials. Besides the main constituents, boron
powders (3 mm) and chromium powders (0.92 mm) were used as
the additive. Powder mixtures with nominal composition (Ni:
Al ¼ 3:1) were dry mixed in a boll mill for 10 h. The mixed powders
were put into the SHS/HE synthesis system, as shown in Fig. 1. The
mixed powder was pressed into compacts before heating, and then
the induction coil heated the reaction puncheon rapidly to 480 C to
start the reaction synthesis. A thermal pair was put in the SHS/HE
synthesis system to detect the temperature of synthesis system.
When the temperature increases dramatically, it indicates the
reaction synthesis begins. Then 2 s later, a force of 400 MPa was
loaded on the reaction puncheon aiming to extrude the synthesized
alloy out of the reaction floor through a hole with diameter of
6 mm.
The samples for microstructure observation and compression
test were cut from the center of the extruded bars. The resultant
phases in the different alloys were characterized by X-ray diffraction (XRD) with a Cu Ka radiation at 40 kV and 40 mA. Microstructural characterization of all alloys were carried out on
OLYMPUS GX41 Optical microscope (OM). Samples for OM observations were prepared by conventional methods of mechanical
polishing and chemical etching with an acidic mixture (CH3COOH/
HNO3/HCl ¼ 8:4:1). The foils for transmission electron microscope
(TEM) observation were prepared by the conventional twin jet
polishing technique using an electrolyte of 10% perchloric acid in
methnol at 20 C after mechanical polishing to 50 mm and cutting
into disc with a diameter of 3.0 mm. The TEM observation was
performed by a JEM-2010 transmission electron microscope operated at 200 kV.
Microhardness measurement was carried out on a Vickers
microhardness tester (MHV-2000)using a load of 150 g and a dwell
time of 15 s. Seven measurements were performed to evaluate an
average value. The compressive specimens with size of 4 4 6 mm3
were cut from the different alloys and all surfaces were mechanically
ground with 600-grit SiC abrasive prior to compression test. The
Fig. 3. Optical micrograph of the SHS/HE synthesized alloy; (a) the Ni3Al alloy, (b) the Ni3Al-1B alloy.
J.T. Guo et al. / Intermetallics 19 (2011) 137e142
compression was conducted in Gleeble-1500 test machine at room
temperature (RT), with an initial strain rates of 1 103 s-1.
3. Results and discussion
3.1. Microstructure characteristics
X-ray diffraction patterns show that the elemental powders
have been transformed to the Ni3Al phase after SHS/HE processing,
as shown in Fig. 2. The B addition has not resulted in obvious
change, except the appearance of (221) peak in the Ni3Al-1B alloy.
The microstructure of all materials after SHS/HE synthesis is
shown in Fig. 3. It is clear all samples were well densified and few
porosities exist. All alloys are composed of coarse and fine grain
with size from 200 nm to 10 mm. Small irregular Al2O3 particles
distribute on the grain boundaries and within the grains. And in
few zone Al2O3 particles agglomerate. The origin of Al2O3 dispersoids can be explained by either the fracture of the oxide layer
covering the original powder particles or powder oxidation during
the milling process. The formation of Al2O3 dispersoids and their
effect on the formation of a microstructure with coarse and fine
grains have been explained elsewhere [12,13]. With the addition of
B the Ni3Al grain has a trend to becoming coarse.
The results of TEM observations on the alloys are presented in
Fig. 4. The results show that the SHS/HE synthesized alloys are
mainly composed of Ni3Al phase. More observations find that
except the fine grains exhibited in the OM micrograph, finer Ni3Al
grains exist in the alloy, as shown in Fig. 4(a). Such refined microstructure should be attributed to the recrystallization resulted by
the hot extrusion. In addition, many fine Ni3Al particles are surrounded by dispersoids of Al2O3, as shown in Fig. 4(b). According to
139
the former research [12,16], this morphology of Al2O3 can handicap
the growth of Ni3Al phase, and contribute to the microstructural
refinement. Further observations on the Al2O3 finds that g-Ni
particles exist inside and around the Al2O3 particle, as shown in
Fig. 4(c) and (d). Except the gathering of Ni resulted in the formation of g-Ni, perhaps there is a reaction that remnant oxygen after
compact reacts with aluminum and forms Al2O3 oxide during the
synthesis process. This reaction leads to the deficient of Al and the
formation of g-Ni phase.
Further observations on the Al2O3 reveal that the Al2O3 particles
existing in the Ni3Al and Ni3Al-1B alloys have two different kind of
structure. Fig. 5 exhibits the morphology of the Al2O3 particles and
their selected area electron diffraction (SAED) patterns. One has the
hexagonal crystal structure with R3c space group and the cell
parameters are: a ¼ b ¼ 0.4758 nm, c ¼ 1.299 nm, as shown in Fig. 5
(a)e(c). This Al2O3 can be determined as the a- Al2O3, which is the
most stable one in all Al2O3. The other one has the Face-centered
crystal structure with Fd3m space group and the cell parameters
are: a ¼ b ¼ c ¼ 0.7948 nm, as shown in Fig. 5(d) and (e). This Al2O3
is determined as the g- Al2O3, which is a metastable phase and will
transform into a- Al2O3 when heat-treated above 1200 K. The
existence of the two kinds of Al2O3 should be attributed to the
synthesis process. The high pressure and high temperature in short
time lead to the formation of Al2O3 with different structure, and the
relative high cooling speed contributes to the reservation of the
Al2O3.
Except Ni3Al phase and Al2O3 particles, fine b-NiAl and Ni3B
particles are found in the Ni3Al-1B alloy with addition of B, as
shown in Fig. 6(a) and (b). The phenomenon has been reported in
the previous researches [7,17]; the boride will be formed in the
Ni3Al alloy, when the B addition is exceeding 1% (at.%). As is well
Fig. 4. TEM bright-field micrograph of Ni3Al alloy (a), Morphology of Al2O3 particles around Ni3Al (b), Formation of Ni rich particles inside and around of Al2O3 (c), Precipitated
Al2O3 particle around Ni rich particle (d) (Inset pictures show the SAED patterns of the corresponding phases).
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J.T. Guo et al. / Intermetallics 19 (2011) 137e142
known, the distribution of powders in powders metallurgy is not
uniform. So, if Al is gathering in some place, it is reasonable for bNiAl particles to precipitate. Further observation on the Ni3Al-1B
alloy finds the presence of twinned Ni3Al, as shown in Fig. 6(c).
Previous research on Ni3Al reveals that the annealing on highdeformed Ni3Al polycrystals leads to the formation of twinned
Ni3Al [18]. In the present investigation, the characteristics of SHS/
HE synthesis technique are similar with the former. The high
cooling speed after hot extrusion results in the formation of twinned Ni3Al.
The difference between the SHS/HE and the common combustion synthesis is the hot extrusion after the combustion synthesis.
Fig. 7 shows the morphology of the part of Ni3Al extruded out of the
mold. It is obvious that the Ni3Al matrix has deformed drastically.
The Ni3Al grains are elongated along the extrusion direction.
Compared with the top part of the SHS/HE synthesized alloys, the
size of the Ni3Al grains in the extrusion part is refined significantly.
From the results it can be concluded that the Ni3Al grains are
fragmentized and then recrystallization in the following process.
But due to the short time in high temperature, the recrystallization
is incomplete. So in the extrusion part there are massive dislocations in many Ni3Al grains and many substructures.
Fig. 8 shows the configuration of dislocations in the SHS/HE
synthesized Ni3Al alloys. Fig. 8(a) and (b) shows the configuration
of randomly distributed dislocations and dislocation arrays inside
of Ni3Al grain. Most of the dislocations are nucleated from the grain
boundary, but some nucleate inside the Ni3Al grain. Dislocations
extend inside the grain and tend to cross the grain. In addition,
Fig. 5. TEM micrographs of the a- Al2O3 particle (a) and their SAED patterns [120] (b), [121] (c); TEM micrographs of the g- Al2O3 particle (d) and their SAED patterns [110] (e), [111] (f).
Fig. 6. b-NiAl particle in the Ni3Al-1B alloy (a), Ni3B particles precipitating in the Ni3Al-1B alloy (b), Twinned Ni3Al in the Ni3Al-1B alloy (c) (Inset pictures show the SAED patterns of
the corresponding phases).
dislocations piling up from the grain boundary and dislocations
tangling with high density inside grain are observed, as shown in
Fig. 8(c) and (d). The high density of dislocations in the grains
demonstrates that dislocation glide behavior dominates the entire
hot extrusion process.
3.2. Mechanical properties
The microhardness and compressive properties at room
temperature (RT) of the SHS/HE synthesized Ni3Al and Ni3Al-1B
alloys are shown in Table 1. With the addition of B, the microhardness and compressive properties of Ni3Al alloys increase significantly. Except the compressive strength the microhardness, yield
strength and compressive ductility of the Ni3Al-1B alloy are improve
more or less than 50%, compared with the Ni3Al alloy.
Compared with the combustion synthesized Ni3Al alloy, the RT
mechanical properties of SHS/HE synthesized Ni3Al alloy improves
obviously. With the B addition, the Ni3Al-1B alloy possesses higher
RT compressive properties than the combustion synthesized Ni3Al0.5B alloy, except the compressive strength. According with the
reference [19], the average grain size of the combustion synthesized Ni3Al-0.5B alloy is about 50 mm, which is much coarser than
the Ni3Al-1B alloy. It is well known that the fine microstructure can
strengthen the alloy without deteriorating its ductility. In the
present investigation, the SHS/HE synthesis technique not only
refines the microstructure significantly but also produces highdensity dislocations in the alloys. The dislocations and increased
grain boundaries can handicap the extension of the new dislocations and then repress the deformation of the alloy, which increase
the microhardness and yield strength. So the SHS/HE synthesized
alloys gain significant improvement in mechanical properties.
Table 1
Mechanical properties at RT of the SHS/HE and combustion synthesized Ni3Al alloys.
Fig. 7. TEM images of the extrusion part, showing deformation of the Ni3Al particles.
Alloy
Michrohardness Yield strength Compressive
Compressive
(HV)
(MPa)
strength (MPa) strain (%)
Ni3Al
Ni3Al-1B
Ni3Al [19]
Ni3Al-0.5B [19]
290
568
220
360
400
850
230
750
1390
1560
680
1800
24
40
15
28
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J.T. Guo et al. / Intermetallics 19 (2011) 137e142
Fig. 8. Dislocation and sub-boundary configuration in SHS/HE synthesized Ni3Al and Ni3Al-1B alloys.
4. Conclusions
1) The Ni3Al alloy with fine microstructure is fabricated by the
SHS/HE synthesis technique, which mainly contains g-Ni, Ni3Al
phases and dispersoids of a- Al2O3 and g- Al2O3.
2) With addition of B, b-NiAl and Ni3B rich particles form in the
Ni3Al-1B alloy. Moreover twinned Ni3Al phase are observed in
the Ni3Al-1B.
3) Compared with the combustion synthesized Ni3Al alloys, the
SHS/HE synthesized Ni3Al alloys own more excellent RT
mechanical properties. And in SHS/HE synthesized Ni3Al alloys,
the B-doped alloy possesses better RT mechanical properties
than the Ni3Al alloy.
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