Journal of Alloys and Compounds 475 (2009) 730–734
Contents lists available at ScienceDirect
Journal of Alloys and Compounds
journal homepage: www.elsevier.com/locate/jallcom
Microstructure and mechanical properties of rapidly solidiﬁed NiAl–Cr(Mo)
eutectic alloy doped with trace Dy
L.Y. Sheng a,b , J.T. Guo a,∗ , Y.X. Tian a , L.Z. Zhou a , H.Q. Ye b
a
b
Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
a r t i c l e
i n f o
Article history:
Received 8 May 2008
Received in revised form 28 July 2008
Accepted 31 July 2008
Available online 11 September 2008
Keywords:
Intermetallics
Rapid-solidiﬁcation
Microstructure
Mechanical properties
a b s t r a c t
The Ni–33Al–28Cr–6Mo eutectic alloy doped with trace Dy was prepared by conventionally casting and
injection casting techniques. Microstructure examination revealed that with the addition of Dy, the Ni5 Dy
phase precipitated along the NiAl and Cr(Mo) phases interface in the intercellular region. Compared
with the conventional-cast alloy, the microstructure of injection-cast alloy got well optimized, which
was characterized by the thin interlamellar spacing, high proportion of eutectic cell area and ﬁne homogeneous distributed Ni5 Dy phase. Furthermore, the mechanical tests showed that the room temperature
mechanical properties of the injection casting alloy improve signiﬁcantly.
© 2008 Elsevier B.V. All rights reserved.
1. Introduction
As a kind of potential high temperature structural materials,
NiAl owns many advantages such as high melting point, low density
and excellent capacity of heat transmission. However, limited room
temperature (RT) ductility and toughness as well as poor elevated
temperature strength seriously hinder its commercial application
[1–3]. The previous researches show that the addition of refractory metals like V, Mo, Cr and Re can improve its RT toughness and
elevated temperature strength by directional solidiﬁcation of its
quasi-binary or ternary eutectic systems [3–5]. Among all the NiAlbased eutectic alloys, Ni–33Al–28Cr–6Mo (NiAl–Cr(Mo) for short)
eutectic alloy had been regarded as one of the most reasonable
choice, because of a relatively good combination between elevated
temperature strength, melting point and RT toughness [6]. A well
directionally solidiﬁed version of this kind of alloy yields a RT fracture toughness value of 21.5 MPa m1/2 [7]. However, little research
has been performed to improve its RT ductility, which is a critical
parameter for practical application. It is well known that rare earth
elements (REEs) are beneﬁt to improve the strength of grain boundaries, and the recent studies have successfully added REEs Dy into
the NiAl–Cr(Mo) alloy to improve its mechanical properties and oxidation resistance [8,9]. However, the segregation of Dy in eutectic
cell boundary result in the formation of massive hard phase which
∗ Corresponding author. Tel.: +86 24 23971917; fax: +86 24 83978045.
E-mail address: [email protected] (J.T. Guo).
0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved.
doi:10.1016/j.jallcom.2008.07.109
is harmful to the mechanical properties of the NiAl–Cr(Mo) alloy.
Therefore, for the Dy-doped NiAl–Cr(Mo) alloy, a feasible way to
reduce the segregation of Dy elements is to increase the cooling
rate.
As a kind of rapid-solidiﬁcation method, injection casting can
produce high cooling rate about 102 K/s, which is much higher than
that of the conventionally casting. The advantages of the injection
casting can be applied to reﬁne the microstructure and reduce the
segregation of alloys and then improve their ductility and toughness. So in the present work, the injection casting technique is used
to fabricate trace Dy-doped Ni–33Al–28Cr–6Mo eutectic alloy, in
order to improve its RT mechanical properties.
2. Experimental procedure
The master alloy of Ni–33Al–28Cr–6Mo (at.%) containing 0.15% Dy
(NiAl–Cr(Mo)–Dy for short) were prepared by induction melting with starting materials of 99.99% Ni, 99.9% Al, 99.9% Cr, 99.9% Mo and 99.9% Dy, respectively.
The melted alloy was casted into rods with 30 mm in diameter. These rods fabricated
by the conventionally casting technique were cut into slices. Some of them were
investigated at as-cast state, and the remaining ones were crushed for injection
casting. The injection casting experiment was conducted with water-cooled copper
mold method, which was usually utilized to prepare bulk metallic glasses, having
signiﬁcant undercooling capacity.
Microstructural characterization and fracture surface investigation of alloys fabricated by conventionally casting and injection casting were carried out by S-3400
scanning electron microscope (SEM) with energy dispersive spectrometer (EDS) and
the compositions of constitute phases were detected by EPMA-1610 electronic probe
microanalysis (EPMA). 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 methanol at −20 ◦ C after mechanical polishing
L.Y. Sheng et al. / Journal of Alloys and Compounds 475 (2009) 730–734
731
Fig. 1. (a) SEM back-scattered electron image of the CC alloy, (b) eutectic cell of the CC alloy, (c) EDS of Dy rich phase, (d) TEM micrograph of NiAl and Cr(Mo) precipitated
particles (inset micrograph shows the SADP of NiAl–Cr(Mo)), (e) SEM back-scattered electron image of the IC alloy and (f) eutectic cell of the IC alloy.
to 50 ␮m 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.
The microhardness measurement was carried out 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 mm × 4 mm × 6 mm were cut from the conventional-cast and injection-cast alloys
by electro-discharge machining (EDM) and all surfaces were mechanically ground
with 600-grit SiC abrasive prior to compression test. The compression tests were
conducted in Gleeble-1500 test machine at room temperature , with an initial strain
rates of 1 × 10−3 s−1 .
3. Results and discussion
3.1. Microstructure characteristics
The typical microstructures of the NiAl–Cr(Mo)–Dy alloy prepared by conventionally casting and injection casting techniques
are shown in Fig. 1. The microstructure of the conventional-cast
(CC) alloy is mainly composed of eutectic cell and intercellular zone
732
L.Y. Sheng et al. / Journal of Alloys and Compounds 475 (2009) 730–734
with a small amount of white phases distributing along cell boundaries, as shown in Fig. 1(a) and (b), respectively. The EDS tests reveal
that the white phases are rich of Dy. In eutectic cells, black NiAl
and gray Cr(Mo) plates exhibit the radial emanating pattern from
cell interior to cell boundaries, while in intercellular zone, coarser
NiAl dendrites and Cr(Mo) phases exhibit irregular shape. Many primary NiAl phases are shown in Fig. 1(a), indicating that the trace Dy
addition may has inﬂuence on the eutectic point of the alloy. TEM
observation shows that a lot of Cr particles have precipitated in NiAl
phase, and in Cr(Mo) phase many small NiAl particles precipitate as
well, as shown in Fig. 1(d). The SEM observation on the injectioncast (IC) alloy exhibit that the alloy possesses a ﬁner microstructure,
compared with the CC alloy, as shown in Fig. 1(e). Moreover, the Dy
rich phase becomes ﬁner and distributes more homogeneously. The
intercellular zone is narrower than that of the CC alloy; while in primary NiAl phase, Cr precipitates are seldom observed, as shown in
Fig. 1(f).
The different microstructures between the CC and IC alloys are
mainly resulted by different cooling rate. The cooling rate of injection casting is between 102 and 103 K/s, which is much higher than
that of conventionally casting, so it will inevitably results in a deviation of the eutectic point of the NiAl–Cr(Mo)–Dy alloy. And the
decrease of the amount of primary NiAl phases in the IC alloy compared with the CC alloy conﬁrms the assuming above. Raj et al.
investigated the directionally solidiﬁed NiAl–Cr(Mo) eutectic alloy
and found that the eutectic cell size and lamellar spacing decreased
with increasing growth rate from 12.7 to 508 mm/h and the average width of the intercellular region was essentially independent
of growth rate and varied between 20 and 25 ␮m [10]. While in the
present study, with increasing cooling rate the thickness of NiAl and
Cr(Mo) plates and the intercellular spacing decrease greatly. The
reason can be attributed to the more embryos getting the chance
to grow with high cooling rate. Furthermore, the great undercooling in front of the liquid/solid (L/S) interface would make the crystal
growth velocity vertical to the L/S interface higher than the one parallel to L/S interface [11]. As a result, the experiment alloy exhibits
a much ﬁne lamellar microstructure. Furthermore the high cooling
rate suppresses the diffusion of Dy elements, so the Dy distribution
becomes more homogeneous.
TEM has been employed to identify the Dy rich phase. The bright
ﬁeld image and corresponding selected area diffraction pattern
(SADP) of [0 1 2 1] zone axis are shown in Fig. 2. The result reveals
that the Dy rich phase can be determined as Ni5 Dy, which has a
hexagonal crystal structure with a = 0.4856 nm, c = 0.3969 nm and
the space group of P6/mmm.
TEM observation on the IC alloy ﬁnds that there are abundant
interface dislocation networks along the NiAl and Cr(Mo) phase
boundaries, as shown in Fig. 3. Such high-density interface dislocation networks well demonstrate the extension of solid solubility.
As shown in Table 1, in the IC alloy the amount of Ni and Al solid
soluted in Cr(Mo) phase is higher than that of Cr and Mo in NiAl
phase. It is no doubt that this will increase the difference of crystal lattice parameters between NiAl and Cr(Mo) phases, which can
result in more interface dislocations along boundaries between
NiAl and Cr(Mo) phases. According to the report of Probst-hein et
al. [12], such high-density interface dislocations are beneﬁcial to
improve the strength of the IC alloy. In addition, a great amount of
ﬁne NiAl precipitates with an average size of 20 nm are observed in
the Cr(Mo) phase, which also demonstrates that the high cooling
rate inhibits element diffusions.
The compositions of constituent phases in the CC and IC alloys
have been detected by EMPA, as shown in Table 1. The results reveal
that the Cr content of primary NiAl in the IC alloy is more than
four times of that in the CC alloy. And the contents of Ni and Al
in Cr(Mo) phase in the IC alloy are much higher than that in the
Fig. 2. (a) TEM bright ﬁeld micrograph of Ni5 Dy phase in the CC alloy and (b) SADP
of Ni5 Dy phase with beam direction B = [0 1 2 1].
Fig. 3. Interface dislocation networks along the NiAl and Cr(Mo) phase boundaries
in the IC alloy.
L.Y. Sheng et al. / Journal of Alloys and Compounds 475 (2009) 730–734
733
Table 1
Compositions of the constituent phases in the CC and IC alloys (at.%)
Alloys
Phase
Ni
Al
Cr
Mo
Dy
CC
Primary NiAl
NiAl
Cr(Mo)
Ni5 Dy
46.70
48.90
4.43
57.80
49.77
48.24
7.82
20.63
3.30
2.86
77.75
2.61
–
–
10.00
–
0.23
–
–
18.97
IC
Primary NiAl
NiAl
Cr(Mo)
Ni5 Dy
46.05
47.60
16.69
70.92
44.23
45.72
15.82
10.08
9.10
6.68
62.60
4.50
–
–
4.89
–
0.62
–
–
14.50
CC alloy as well. According with the research carried out by Kim et
al. [13], when the cooling rate is high, it is difﬁcult for elements to
diffuse from the L/S interface, which leads to more elements solid
soluting in phases. In the present investigated eutectic alloy, the
reﬁned lamellar microstructure results in more NiAl and Cr(Mo)
phases interfaces, which effectively repress the elements diffusion
across the NiAl–Cr(Mo) interfaces, combining the characteristics of
the eutectic solidiﬁcation. As a result, more solid soluted elements
stayed in the phases.
3.2. Microhardness
The results of microhardness tests reveal that the hardness of
primary NiAl in the IC alloy is 20% higher than that of the CC alloy,
as shown in Fig. 4. Moreover, the hardness of the NiAl–Cr(Mo)
eutectic of the IC alloy is higher than that of the CC alloy as well.
Generally speaking, the increase of the hardness of the IC alloy
should be mostly attributed to the solid solubility extension. As
shown above, in the IC alloy, the 9% Cr solid solution in primary
NiAl is more than its solid solubility in equilibrium solidiﬁcation
state [14]. As is well known, in the NiAl alloy Cr is one of mainly
solid solution strengthening elements and prefers to occupy the
site of Al [15,16]. The investigation of Frommeyer et al. exhibits
that the substituting of Cr for Al decreases the lattice parameter of NiAl slightly [17], which results in lattice distortion, and
then strengthens the alloy. Similar with that of the primary NiAl,
the increase of hardness of NiAl–Cr(Mo) eutectic of the IC alloy
should be ascribe to the solid solution strengthening effect. Furthermore the ﬁne lamellar microstructure and abundant interface
dislocations contribute much to improve the microhardness of
NiAl–Cr(Mo) eutectic, which can impede the deformation of the
eutectic.
Fig. 5. True stress–true strain curves of the CC and IC alloys at RT with an initial
strain rate of 1.0 × 10−3 s−1 .
3.3. Compressive properties
The true stress–true strain compressive curves at RT of the CC
and IC alloys and their mechanical tests data are shown in Fig. 5 and
Table 2, respectively. Obviously, the true stress–true strain curves
of CC and IC alloys almost exhibit the similar trend with continuous
hardening after yield. But the slope of the compressive curve during
elastic deformation of the IC alloy is higher than that of the CC
alloy, which implies that the IC alloy has a relative higher Young’s
modulus. As listed in Table 2, the RT compressive strain and yield
strength of the IC alloy are 18% and 1362 MPa, which increase by
63.6% and 37.9%, compared with the CC alloy.
In order to interpret the signiﬁcant mechanical properties
improvement of the IC alloy, the micro-mechanisms of deformation and fracture are discussed as follows. Fracture surfaces after
compression testing were observed by SEM, as displayed in Fig. 6(a)
and (b). The micrographs exhibit that the CC and IC alloys possess
almost different fracture morphologies. The fracture surface of the
CC alloy are mainly debonding along the boundaries of NiAl and
Cr(Mo) phases, with few cleavage. Moreover, it is observed cracks
extend along the eutectic cell boundary, which indicate that the
adhesion of eutectic cell boundaries is still weak. However, to the
IC alloy, the fracture morphology exhibits smooth cleavage in the
primary NiAl phase, and cleavage of NiAl plates in NiAl–Cr(Mo)
Table 2
Results of compressive tests at RT with an initial rate of 1.0 × 10−3 s−1
Fig. 4. Microhardness of the CC and IC alloys.
Alloys
Yield strength
(MPa)
Compressive
strength (MPa)
Compressive
strain (%)
CC
IC
987
1362
1935
2274
11
18
734
L.Y. Sheng et al. / Journal of Alloys and Compounds 475 (2009) 730–734
Fig. 7. TEM bright ﬁeld micrograph of Dy2 O3 oxide (inset micrograph shows its
SADP).
4. Conclusions
Fig. 6. Typical fracture surfaces of RT compressive samples: (a) the CC alloy, (b) the
IC alloy (the A in (b) indicates the cleavage of primary NiAl and the B refer to the
deformed Cr(Mo)).
(1) With trace Dy added, the Ni5 Dy phase with hexagonal crystal structure forms along NiAl–Cr(Mo) phase boundary in the
intercellular region.
(2) The injection casting reﬁnes the eutectic cell and intercellular
region signiﬁcantly and the extension of solid solute content
occurs in the IC alloy, due to the high cooling rate. Furthermore,
by the injection casting technique, Ni5 Dy particles become ﬁne
and distribute homogeneously.
(3) Compared with the CC alloy, the RT compressive ductility and
yield strength of the IC alloy both get signiﬁcant improvement,
which can be attribute to the reﬁned microstructure and homogeneous distributed Ni5 Dy phases.
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