Materials Science and Engineering A250 (1998) 152 – 157
Microstructure of a rapidly solidified Al–4V–2Fe ultrahigh
strength aluminum alloy
K. Hono a,*, Y. Zhang b, T. Sakurai b, A. Inoue b
a
Materials Physics Di6ision, National Research Institute for Metals, 1 -2 -1 Sengen, Tsukuba 305 -0047, Japan
b
Institute for Materials Research, Tohoku Uni6ersity, Sendai 980 -8577, Japan
Abstract
The microstructure of a rapidly solidified Al–4V – 2Fe alloy has been examined in detail by means of transmission electron
microscopy and atom probe field ion microscopy. It is composed of a mixture of a-Al and amorphous phases. a-Al grains are
relatively large (400 nm), but nanoscale domains of amorphous phase are trapped in the fcc grains. The interfaces between the
a-Al and amorphous phases are very irregular and solute elements are partitioned into the amorphous phase. The evolution of
such a microstructure is explained on the basis of an unstable interface due to constitutional supercooling during the solidification
process. Enrichment of Fe is observed at some interfaces. Solutes rejected during crystallization of a-Al are believed to be enriched
at the advancing a/liquid interface and this solute-enriched liquid solidifies as amorphous phase, resulting in nanoscale amorphous
regions embedded in fcc grains. © 1998 Elsevier Science S.A. All rights reserved.
Keywords: Microstructure; Al–4V–2Fe; Transmission electron microscopy; Atom probe field ion microscopy
1. Introduction
The recent discovery of nanocrystalline Al – TM–LN
(TM, late transition metals; LN, lanthanide elements)
alloys demonstrated that the tensile strength of aluminum alloys can reach as high as 1500 MPa [1,2],
which is more than two times higher than the strongest
age hardenable wrought aluminum alloy. One drawback of the series of Al – TM – LN alloy is their use of
lanthanide elements, which makes them too expensive
for possible commercial applications. Recently, Inoue
et al. [3,4] reported that a nanoscale mixture of fcc and
amorphous phase can be produced by rapidly solidifying Al–V–TM (TM= Fe, Co, Ni) alloys, which exhibit
ultimate tensile strengths of 1300 MPa. Depending
on the composition of the alloys, the constituent phases
in the microstructure vary, but they are composed of a
nanoscale mixture of a-Al and amorphous phase or a
mixture of a-Al, amorphous phase and icosahedral
phase. As these alloys contain only transition metal
elements as solutes, their potential as commercial materials appears to be higher than the lanthanide containing alloys.
* Corresponding author. E-mail: [email protected]
0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved.
PII S0921-5093(98)00552-8
The highest tensile strength of the rapidly solidified
Al–4V–2Fe alloy was reported to be 1250 MPa with a
microstructure composed of a nanoscale mixture of
a-Al and amorphous phases. According to the original
paper by Inoue et al. [4], the nanoscale amorphous
particles are embedded in a-Al particles. As main features of the microstructure, they reported the absence
of distinct a/amorphous interfaces and no partitioning
of solute between these two phases. However, the absence of partitioning between the a-Al and the amorphous phases is difficult to explain based on the fact
that this alloy has two distinct phases. Thus, this paper
seeks to clarify the microstructural features of a rapidly
solidified Al–4V–2Fe alloy and to elucidate the mechanism of the microstructural evolution during the solidification process.
2. Experimental
An Al–4V–2Fe (at.%) alloy ribbon specimen was
prepared by the single roller melt spinning process in an
argon atmosphere. The specimen was 1 mm wide and
15 mm thick. The circumferential velocity of the
melt-spinning roll was 40 ms − 1. In this study, only the
K. Hono et al. / Materials Science and Engineering A250 (1998) 152–157
as-quenched microstructure was examined, since the
highest strength was obtained in the as-quenched specimen [4]. TEM specimens were prepared by ion thinning
using Fischione Model 3000 ion mill, using 4 kV ion
energy and 4 mA ion current. For transmission electron
microscopy (TEM), Philips CM200 and JEOL JEM4000EX were employed. For local concentration characterization, an energy compensated time-of-flight atom
probe (APFIM) was employed. Needle shaped specimens for APFIM analysis were prepared by electropolishing a square rod of 15 mm ×15 mm × 10 mm which
were made by mechanical grinding the ribbon. The
atom probe field ion microscope (APFIM) used in the
present study was an energy compensated time-of-flight
atom probe. The field ion microscopy (FIM) images
were observed using Ne as an imaging gas at 25–40 K.
The atom probe analysis was performed at 35 K in
1× 10 − 8 Pa vacuum with a pulse fraction (Vp/Vdc)
of 0.15 and a pulse repetition rate of 100 Hz. Separation of Al and Fe ions cannot be accurately made by
mass analysis, because some of the ions overlap. Most
of Al ions are detected as 27Al2 + and 27Al + having
mass-to-charge ratio, m/n, at 13.5 and 27, respectively.
Fe atoms are detected as doubly charged ions, Fe2 + ,
having m/n at 27, 28 and 28.5. Thus, 27Fe2 + and 27Al +
overlap each other, and it is impossible to separate
153
these. Fortunately, the isotope abundance of 54Fe is
5.8%, thus only 5.8% of Fe atoms will be incorrectly
assigned as Al. Al atoms tend to form hydride ions
even under an ultrahigh vacuum condition, and a small
portion of ions with m/n= 28 is expected to be AlH +
which overlap with 56Fe2 + . Thus, the atom probe
analysis results presented in this paper is qualitative,
rather than quantitative.
3. Results
Fig. 1 shows a bright field electron micrograph and
the corresponding selected area diffraction pattern
(SADP). A crystal grain approximately 30 nm across is
observed with dark contrast. The SADP taken from
this region suggests that the selected area is composed
of almost a single orientation, but with a slightly rotated misorientation with respect to the [001] zone axis.
The weak halo ring indicates that an amorphous phase
is present within the selected area. Fig. 2 illustrates a
bright field image and microdiffraction patterns using a
nominal 5 nm converged electron beam taken from the
indicated regions. Most of the regions have the same
orientation. Some microdiffraction patterns taken from
within the grain indicate that amorphous phase remains
Fig. 1. Bright field electron micrograph of a rapidly solidified Al – 4V – 2Fe alloy and the selected area diffraction pattern taken from the darkly
imaging grain.
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K. Hono et al. / Materials Science and Engineering A250 (1998) 152–157
Fig. 2. Microdiffraction patterns taken from the darkly imaging grain. Nominal size of 5 nm electron beam was used.
within the grain. One microdiffraction pattern shows
that the orientation is different from [001], and this
indicates that the darkly imaged grain is not a perfect
single crystal. The microdiffraction pattern taken from
the bright mottled region outside the grain indicates it
is largely amorphous, but some crystal reflection with
totally different orientation is observed. This means
that a crystal which does not satisfy the Bragg condition is mixed with the amorphous phase in this region,
and this causes a mottled contrast instead of the featureless contrast typical of amorphous phases.
Fig. 3 shows a high resolution image of this specimen
taken from the [001] zone of a grain. From the crystallized region, a fringe contrast corresponding to the
[001] zone can be observed. The lack of the fringe
contrast and the typical isotropic maze-like pattern
expected from the amorphous structure indicates the
presence of amorphous phase. Some amorphous regions appear to be surrounded by the interconnected
a-Al phase. Such an irregular interface suggests that
interfacial instability occurs during the growth of the
fcc interconnected crystal. Fig. 4 shows an enlarged
HREM micrograph obtained from a relatively clear
a/amorphous interface. It is interesting to note that the
K. Hono et al. / Materials Science and Engineering A250 (1998) 152–157
Fig. 3. HREM image of a rapidly solidified Al–4V–2Fe alloy.
Amorphous regions surrounded by interconnected a-Al phase are
arrowed.
interface is not atomically smooth, but there are atomic
protrusions and concavities. This suggests that there is
an instability at the interface.
A FIM image of the as-quenched Al – 4V – 2Fe alloy
is shown in Fig. 5. By selected area atom probe analyses, the brightly imaging regions were confirmed to be
the amorphous phase containing higher concentration
of solutes. This image clearly shows that the amorphous and fcc phases are entangled with each other.
Because the TEM results have shown that the fcc
particle with mottled contrast in Fig. 1 is in fact a single
crystal containing nanoscale amorphous domains, the
fcc phase mixed with the amorphous phase is believed
to be interconnected. The darkly imaging regions in the
Fig. 4. HREM image of a/amorphous interface in a rapidly solidified
Al– 4V – 2Fe alloy. The image is Fourier filtered for clarity.
155
Fig. 5. Ne field ion image of a rapidly-solidified Al – 4V–2Fe alloy.
Brightly imaging region corresponds to the solute-rich amorphous
phase and the interconnected dimly imaging region corresponds to
the a-Al phase.
FIM image (Fig. 5) also appear to be interconnected.
Solute partitioning in the rapidly quenched Al–4V–
2Fe specimen has been examined by APFIM analysis as
shown in Fig. 6. In the concentration depth profiles, the
regions which contain mostly Al correspond to the
crystallized a-Al. The other regions enriched with V
and Fe correspond to the amorphous phase. The concentration of Fe and V in the a-Al phase was determined to be 2 and 3 at.%, respectively. As seen in Fig.
6, it is evident that the solutes partition during the
solidification process. The partitioning factors of Fe
and V in the amorphous phase with respect to the a-Al
are approximately five. In addition to this uniform
Fig. 6. Atom probe concentration depth profiles of Al, V and Fe in
a rapidly solidified Al – 4V – 2Fe alloy.
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K. Hono et al. / Materials Science and Engineering A250 (1998) 152–157
Fig. 7. Atom probe concentration depth profiles of Al, V and Fe in
a rapidly solidified Al–4V–2Fe alloy.
partitioning, we have found another type of solute
partitioning as shown in Fig. 7. This concentration
profile was obtained from the same batch of material,
but from a different region. Again, both V and Fe are
rejected from the a-Al phase. In this case, however, the
Fe is observed to have separated from the V and is
enriched at the interface between the V rich phase and
the a-Al phase. The concentration of V in the amorphous phase is : 12 at.%, while that in the a-Al is 2
at.%; thus the partitioning factor of V in the amorphous phase with respect to that in the a-Al phase is
approximately six. On the other hand, Fe is enriched at
the outer side of the amorphous phase near the amorphous/a interfaces. At both sides of the a/amorphous/a
interfaces, Fe is strongly enriched in the amorphous
phase near the interfaces. Inside the amorphous, the
concentration of Fe is as low as that in the fcc-Al
phase. This peculiar concentration profile indicates that
there are two types of interfaces in this microstructure.
4. Discussion
An earlier paper by Inoue et al. [3] described the
features of the melt-spun Al – 4V – 2Fe alloy as follows:
(1) homogeneous dispersion of nanoscale amorphous
particles in an isolated state; (2) absence of solute
partitioning. They also suggested that the amorphous
particles precipitate from the supercooled liquid. This
study, however, has clearly shown that there is partitioning of solutes between the fcc grains and the dispersed amorphous phase. This suggests that a-Al
crystallizes first from the liquid during solidification,
and the rejected solute stabilizes the liquid phase result-
ing in the formation of amorphous phase in the asquenched microstructure. The presence of amorphous
phase in the as-quenched microstructure suggests that
there is a deep eutectic in the ternary phase diagram.
The primary phase is a-Al, and when the primary
particles are nucleated during cooling, rejection of solute occurs from the a-Al. This causes solute enrichment at the advancing interface of the primary
aluminum grains. Such solute enrichment at the growing solid/liquid interface can produce constitutional
supercooling ahead of the advancing interface, and
make growth of a smooth interface unstable provided
the temperature gradient in the liquid is not too high.
Such constitutional supercooling often causes dendritic
growth of crystals, but in the present case, no preferential crystallographic growth has been observed. The
inability of a crystallographic dendritic morphology
may be explained by the growth direction being primarily determined by the solute mobility rather than the
difference in growth rate due to crystal growth anisotropy and thus the growth occurs in more randomly
oriented bursts.
As the SADP in Fig. 1 indicates, the grains of the
primary particles are not perfect single crystals (see
mottled contrast in the bright field image in a grain and
slight rotation in the SADP), but they contain slight
misorientations. This may be explained in terms of
slight misorientations developing within different
growth fronts similar to misorientations in dendrites.
This contrast in the grain is also explained by the
existence of nanoscale amorphous regions within the
grains. These amorphous regions are believed to be
trapped by the growth of the fcc grain. As solute
elements are rejected from the a-Al grains during their
growth, solute enriched liquid will be entrained in the
relatively large a-Al grains. Solute enrichment is believed to stabilize the amorphous phase, and these
solute-rich regions will remain amorphous when the
specimen is cooled to the room temperature.
In addition to these trapped amorphous regions, the
matrix amorphous will also remain amorphous as a
grain boundary phase. Thus, two types of interfaces
may be expected. One is the interface between the
entrained amorphous phase and the primary a-Al
grains and the other is the interface between the a-Al
grains and the matrix amorphous. The concentration of
the latter interface will maintain features of the growing
interface during the solidification process and such
interface will be far from the metastable equilibrium
between the amorphous and the primary phase which
will be achieved in the former interface. This may
explain two types of solute enrichment observed in the
APFIM analyses. However, the observation of enrichment of Fe in the amorphous phase near the interface
(Fig. 7) is puzzling. From the volume diffusion data in
the a-Al, the diffusivity of Fe is expected to be faster
K. Hono et al. / Materials Science and Engineering A250 (1998) 152–157
than that of V. Furthermore, the diffusivities of these
two elements are expected to be similar. Thus, enrichment of Fe near the interface cannot be explained
kinetically. We do not have any good explanation for
this observation and further study into the reproducibility of Fe enrichment at the interface is in progress.
5. Conclusion
The microstructural feature of a rapidly solidified
Al–4V–2Fe alloy has been investigated by TEM and
APFIM. The microstructure is composed of relatively
large primary crystals of a-Al (40 nm). Nanoscale
amorphous particles enriched with Fe and V are present
within the primary Al phase by entrainment of the
liquid phase during solidification. The interface between
the amorphous and the a-Al is irregular, suggesting that
interfacial instability occur during solidification due to
constitutional supercooling. At some amorphous/a-Al
interfaces, enrichment of Fe atoms is observed.
.
157
Acknowledgements
This study is partly supported by the NEDO International Joint Research Grant on ‘The Nanocrystalline
and Supercooled Liquid States of Alloys’. The authors
are grateful to Professor W.T. Reynolds at Virginia
Tech. and Dr P.J. Warren at University of Oxford for
valuable discussions. We also acknowledge Mr
Sasamori for preparation of the specimens used in the
present study.
References
[1] Y.-H. Kim, A. Inoue, T. Masumoto, Mater. Trans. JIM 32
(1990) 747.
[2] H. Chen, Y. He, G.J. Shiflet, S.J. Poon, Scr. Metall. Mater. 25
(1991) 1421.
[3] A. Inoue, H. Kimura, K. Sasamori, T. Masumoto, Mater. Trans.
JIM 36 (1995) 1219.
[4] A. Inoue, H. Kimura, K. Sasamori, T. Masumoto, Mater. Trans.
JIM 37 (1996) 1287.
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