Intermetallics 17 (2009) 400–403
Contents lists available at ScienceDirect
Intermetallics
journal homepage: www.elsevier.com/locate/intermet
Modiﬁcation of NiAl–Cr(Mo)–0.15Hf alloy by Sc addition
Y. Xie a, b, J.T. Guo a, *, Y.C. Liang a, L.Z. Zhou a, H.Q. Ye b
a
b
Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, PR China
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, PR China
a r t i c l e i n f o
a b s t r a c t
Article history:
Received 22 June 2008
Received in revised form
12 September 2008
Accepted 25 November 2008
Available online 24 December 2008
The effect of Sc addition on the microstructure and room temperature compressive properties of NiAl–
28Cr–5.85Mo–0.15 (at pct) Hfx (wt pct) Sc (x ¼ 0,0.05, 0.1, 0.2, 0.3) alloys was investigated. The results
show that appropriate Sc addition (no more than 0.10 wt pct) leads to the reﬁnement of interlamellar and
intercellular spacings of the eutectic NiAl/Cr(Mo) cell, and the improvement of the compressive ductility
and ultimate compressive strength at room temperature. When the addition of Sc is more than
0.10 wt pct, the typical NiAl/Cr(Mo) cell structure becomes broken. With the fragment of Cr(Mo) rods
embedded in NiAl matrix instead of the alternating NiAl and Cr(Mo) plates, which damages the
compressive properties. In addition, when the Sc addition content increases to 0.20 wt pct, Sc-containing
phase is found and tentatively identiﬁed as ScO.
Ó 2008 Elsevier Ltd. All rights reserved.
Keywords:
A. Nickel aluminides, based on NiAl
B. Microalloying
B. Mechanical properties at ambient
temperature
D. Microstructure
1. Introduction
The ordered intermetallic compound NiAl is regarded as
a potential candidate for high temperature structural utilizations
because of its high melting point (Tm ¼ 1921 K), substantially lower
density (5.9 g/cm3) than commercial Ni-based superalloys (about
8 g/cm3), high thermal conductivity (above 6 W/m K), and excellent
oxidation resistance at temperature above 1273 K [1–3]. However,
the industrial applications of NiAl alloy are limited by two major
drawbacks. The ﬁrst one is poor strength and creep resistance at
high temperature; the second one is the serious scarcity in fracture
toughness and ductility at room temperature [1,2].
As for the ﬁrst drawback, Hf is an effective element for solid
solution strengthening and formation of strengthening Heusler
phase (Ni2AlHf) in NiAl alloy at elevated temperature [4,5]. DS
NiAl–Cr(Mo)–0.1Hf shows higher stress rupture strength than that
of superalloy Rene 80 at elevated temperature [6]. Whereas the
strength improves at the expense of the loss of room temperature
ductility.
With respect to the second drawback, there does not exist any
marked improvement except the perfect DS NiAl–Cr eutectic alloy
which possesses room temperature fracture toughness more than
20 MPam1/2. In order to obtain the perfect DS microstructure, the
growth rate must be lower than 25 mm/h [7,8] which is generally
* Correspondence to: J.T. Guo, Institute of Metal Research, Chinese Academy of
Sciences, 72 Wenhua Road, Shenyang 110016, PR China. Tel.: þ86 24 23971917;
fax: þ86 24 83978045.
E-mail address: [email protected] (J.T. Guo).
0966-9795/$ – see front matter Ó 2008 Elsevier Ltd. All rights reserved.
doi:10.1016/j.intermet.2008.11.018
too much slow for economic commercial production and industrial
application. Aoki and Izumi [9] found that B microalloying signiﬁcantly improves the room temperature ductility of polycrystalline
Ni3Al intermetallic. It indicates that microalloy may be a promising
approach to ameliorate the brittle NiAl alloy. Scandium belongs to
the rare element group which possesses special physical and
chemical features, such as surface activity. Guo et al. have investigated the effects of several kinds of rare earth elements (like Y, La,
Ce, Nd and Dy) on the microstructure and mechanical properties of
NiAl alloys [10–13]. Compared to other rare earth elements, Sc is
the ﬁrst element with lowest density in rare earth group. And it is
forecasted as ‘‘Boron-like’’ element in Mendeleev elements’ period
table. Therefore, the present work focuses on microstructure and
mechanical properties of NiAl–Cr(Mo)–0.15Hf alloy modiﬁed by
different Sc contents.
2. Experimental
The alloys used for this work were arc melted by a nonconsumable tungsten electrode under an argon atmosphere in
a water-cooled copper crucible from starting materials of Ni
(99.9 at pct), Al (99.9 at pct), Mo (99.9 at pct), Cr (99.5 at pct), Hf
(99.5 at pct) and Al–Sc alloy (Sc, 1.92 wt pct). Each alloy button was
turned over and remelted at least three times to ensure a homogeneous specimen. Since weight losses were generally less than
0.5%, the compositions of the alloys were considered to be equal to
their nominal compositions, shown in Table 1. The specimens were
encapsulated in quartz tube and then heat treated at 1523 K for
24 h.
Y. Xie et al. / Intermetallics 17 (2009) 400–403
3. Results
Table 1
Nominal compositions of the alloys.
Alloy No.
0
1
2
3
4
401
Ni/Al
(1:1)
Atomic fraction, %
Cr
Mo
Hf
Mass fraction, %
SC
Bal.
Bal.
Bal.
Bal.
Bal.
28.00
28.00
28.00
28.00
28.00
5.85
5.85
5.85
5.85
5.85
0.15
0.15
0.15
0.15
0.15
0.00
0.05
0.10
0.20
0.30
Microstructural characterization was analyzed by scanning
electron microscope (SEM) and transmission electron microscope
(TEM). SEM, coupled with energy dispersive X-ray spectroscopy
(EDX), was performed with 20 kV accelerating voltage using a JEOL
JSM-6301F ﬁeld emission gun (FEG) scanning electron microscope.
TEM analysis was conducted by a JEOL 2000 FXII TEM operating at
200 kV. Samples for TEM were thinned mechanically to a thickness
of about 50 mm and then ion milled to the ﬁnal thickness.
Compression specimens with 4 mm 4 mm 6 mm in size
were taken by electron discharge machine (EDM). All the EDM’ed
surfaces were ground to 1000 grit by abrasive papers. The
compressive testing was conducted in air with a Gleeble 1500
testing machine at room temperature under the initial strain rate of
1.94 103 s1. The autographically recorded load–time curves
were converted to true stress–strain curves by taking constant
volume into account.
The SEM micrographs of NiAl–28Cr–5.85Mo–0.15Hf alloys with
the trace addition of 0, 0.05, 0.1, 0.2 and 0.3 wt pct Sc after
homogenization treatment are presented in Fig. 1. No. 0–2 alloys
exhibit typical eutectic cell consisting of alternating black NiAl
phase and gray Cr(Mo) phase. The lamellar Cr(Mo) plates in each
cell emanate radially from the cell interior to the cell boundaries
and become thicker at the periphery of the cell. With the increasing
Sc addition from 0 to 0.1 wt pct, the microstructure including the
interlamellar spacing and the intercellular spacing reﬁnes. The
interlamellar spacing in the cell and the intercellular spacing of the
No. 0 alloy are w0.8 mm and 8–25 mm, respectively. While in the No.
2 alloy, these two kinds of spacing reduce to w0.7 mm and 3–13 mm
respectively. For the alloys with the addition of Sc more than
0.1 wt pct, the typical cell structure is breached. Especially in the
No. 4 alloy, the cell structure has completely disappeared. The
microstructure can be characterized by the block Cr(Mo) phase
which is embedded in the NiAl matrix. Compared to the Cr(Mo)
plate in the eutectic cell of alloy No. 2 (Fig. 1(c)), the Cr(Mo) rods in
alloy No. 4 are coarsen (Fig. 1(e)). In addition, when the Sc content
increases to 0.2 wt pct, a white Sc-containing phase forms
(Fig. 1(d)–(e)), and is further identiﬁed as ScO compound through
TEM (Fig. 2). Its lattice parameters are: a ¼ b ¼ c ¼ 0.448 nm, with
Fm3m space group [14,15].
The true stress–strain curves for Sc-doped and Sc-free alloys
generated by room temperature compression test are presented in
Fig. 1. SEM BSE (back scattered electron) images of NiAl–Cr(Mo)–Hf alloys containing various Sc content (wt pct): (a) 0, (b) 0.05, (c) 0.1, (d) 0.2, (e) 0.3.
402
Y. Xie et al. / Intermetallics 17 (2009) 400–403
Fig. 2. (a) Bright ﬁeld TEM images of ScO phase in NiAl–28Cr–5.85Mo–0.15Hf–0.3Sc alloy, and the corresponding diffraction patterns (b) along zone axis [001], (c) along zone axis
[012].
Fig. 3. The specimens were compressed to fracture. As shown in
Fig. 3, the onset of the crack growth is indicated by arrow. Fig. 4
shows the compressive properties of various Sc-doped alloys at
room temperature. For all Sc-doped alloys, the compressive
ductility and strength at room temperature increase with Sc addition when the Sc content is not more than 0.1 wt pct, while above
this critical content the mechanical properties of the Sc-doped
alloys deteriorate. It is clearly that the 0.10 wt pct Sc-doped alloy
attained the best ductility of about 35% and the highest compressive strength of 1600 MPa at room temperature. Therefore,
0.10 wt pct Sc addition is speculated as the best content to improve
the room temperature compressive properties of NiAl–Cr(Mo)–
0.15Hf alloy.
4. Discussions
The rare earth elements are surface active, so the additional Sc
can decrease the S–L surface tension (sLS). On the assumption of
globular crystal nucleus, the critical radius (r*) of crystal nucleus
and the critical nucleation energy (DG*) are given in Ref. [16].
r* ¼
2sLS
DGm
(1)
Fig. 3. The true stress–true strain curves of the NiAl–Cr(Mo)–Hf alloys containing
various Sc content tested at room temperature. (‘‘/’’ indicates the onset of crack
growth).
s3LS
16
DG ¼ p
3
DG2m
*
!
(2)
where DGm is the Gibbs free energy difference between solid and
liquid per unit volume. Due to the decrease of S–L surface tension
(sLS), both the critical radius (r*) of crystal nucleus and the critical
nucleation power (DG*) diminish. It means that the nucleation
becomes easy and the number of crystallization nuclei will
increase. Therefore, appropriate content Sc can give rise to the
microstructural reﬁnement. As shown in Fig. 1(a)–(c), the interlamellar spacing and the intercellular spacing become ﬁne.
When the addition of Sc element is up to 0.3 wt pct, the typical
NiAl/Cr(Mo) eutectic cells are transformed into the Cr(Mo) rods
embedded in NiAl matrix. This marked microstructural change is
caused by adding a certain amount of Sc which can bring out
a disturbance of the solidiﬁcation process. In NiAl–Cr(Mo) eutectic
system, the solid solubility of Sc element in NiAl phase outclasses
that in Cr(Mo) phase [17]. It indicates that the distribution coefﬁcient of Sc in NiAl/Cr(Mo) is far more than 1. The high distribution coefﬁcient of Sc leads to the enrichment of Sc at the S/L
interface of Cr(Mo) phase during the solidiﬁcation, which hinders
the continuous growth of Cr(Mo) phase. However, the NiAl phase
grows faster than Cr(Mo) phase because of the less Sc element at
the S/L interface of NiAl phase. Therefore, the laggard Cr(Mo)
Fig. 4. Compressive strain and ultimate compressive stress as a function of Sc content
at room temperature.
Y. Xie et al. / Intermetallics 17 (2009) 400–403
phase is separated into the block rods by the fast-growing NiAl
phase.
Besides the modiﬁcation of microstructure, the additional Sc
also exerts an inﬂuence on the room temperature compressive
properties of the present alloys. Sc belongs to the active element
which will react with the impurities from raw materials, such as S,
O, etc. On one hand, the alloy is puriﬁed for the removal of impurities. The deformation of substrate becomes easy. On the other
hand, the products formed by the reaction between Sc and impurities can be regarded as the non-spontaneous nuclei of crystallization which will increase the number of the crystallization nuclei.
In addition, the previous investigations found that the rare earth
elements tend to segregate to the grain boundary or phase
boundary, remove the impurities at the grain boundary and reduce
or eliminate the harmful effect of impurities on the cohesive
strength [10]. The above factors lead to the improvement of
compressive properties.
However, excess Sc results in the weakening of the compressive
properties. It can be ascribed to the changed microstructure. The
typical NiAl/Cr(Mo) eutectic cell structure breaks into block Cr(Mo)
phase embedded in the NiAl matrix when the additional Sc is up to
0.2 wt pct. Moreover, the formation of the ScO phase is also harmful
to the compressive ductility, although the inﬂuence is small owing
to the slim volume of scandium oxide.
5. Conclusions
(1) When the additional Sc content is not more than 0.1 wt pct, the
interlamellar spacing and the intercellular spacing decrease
with the increasing Sc content. The microstructural reﬁnement
leads to the improvement of room temperature compressive
properties.
403
(2) Excess Sc addition breaks the typical NiAl/Cr(Mo) eutectic cell
structure. The disappearance of alternating NiAl and Cr(Mo)
plates plays a decisive role on the deterioration of compressive
properties. Furthermore, too much Sc addition gives rise to the
formation of ScO phase which also does harm to compressive
properties.
(3) The present alloy doped with 0.10 wt pct Sc possesses the best
ductility of about 35% and the highest compressive strength of
1600 MPa at room temperature. Therefore, 0.10 wt pct Sc
addition is speculated as the appropriate amount to improve
the room temperature ductility of NiAl–Cr(Mo)–0.15Hf alloy.
References
[1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
[15]
[16]
[17]
Miracle DB. Acta Metall Mater 1993;41:649–84.
Noebe RD, Bowman RR, Nathal MV. Int Mater Rev 1993;38:193–232.
Barrett CL. Oxid Met 1988;30:361–90.
Whittenberger JD, Locci IE, Darolia R, Bowman RR. Mater Sci Eng A
1999;268(1–2):165–83.
Takeyama M, Liu CT. J Mater Res 1990;5(6):1189–96.
Guo JT, Cui CY, Qi YH, Ye HQ. J Alloy Compd 2002;343(1–2):142–50.
Raj SV, Locci IE. Intermetallics 2001;9(3):217–27.
Whittenberger JD, Raj SV, Locci IE, Salem JA. Intermetallics 1999;7(10):1159–
68.
Aoki K, Izumi O. J Jpn Inst Met 1979;43(12):1190–6.
Guo JT, Huai KW, Gao Q, Ren WL, Li GS. Intermetallics 2007;15(5-6):
727–33.
Li HT, Guo JT, Huai KW, Ye HQ. J Cryst Growth 2006;290(1):258–65.
Ren WL, Guo JT, Li GS, Zhou JY. Mater Lett 2003;57(8):1374–9.
Ren WL, Guo JT, Li GS, Zhou JY. J Mater Sci Technol 2004;20(2):163–6.
Massalski TB, Okamoto H, Subramanian PR, Kalprzak L. Binary alloy phase
diagrams. Ohio: ASM International; 1990.
Power Diffraction File-PDF #780719 Software PCPDWIN, version 2.02; 1999.
Christian JW. The theory of transformation in metals and alloys. 2nd ed. New
York: Pergamon Press; 1975.
Xie Y, Guo JT, Zhou LZ, Liang YC, Ye HQ. Acta Meta Sini 2008;44(5):
529–34.