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Microstructure Investigations of Ball Milled Materials
Laboratory of Atomic Imaging of Solids, Institute of Metal Research, Chinese Academy of Sciences,
Shenyang 110015, People’s Republic of China
mechanical alloying; high-resolution electron microscopy; nanocrystalline; deformation twinning; solid solution
HREM and FEG TEM were emphasized and extensively used to follow the most
subtle changes in the structure and composition of ball-milled Cu, Fe-Cu, and thermally decomposed
Fe60Cu40. Some significant results are obtained and summarized as follows: HREM shows that the
deformation of ball-milled copper proceeds mainly by twinning and shear bands (SBs) formation.
The nano-grains formed during ball milling (BM) contain a high density of dislocations. The grain
boundaries (GBs) of nanocrystalline (NC) Cu prepared by BM are ordered, curved, and strained, but
disordering, lattice distortion, and nanovoids in local regions were frequently observed. Nanoscale
composition analysis on mechanically alloyed Fe16Cu84 shows that the average Fe content in both
the interior of grains and the GBs is close to the designed composition, which proves that a
supersaturated solid solution has really formed. However, the Fe content is rather inhomogeneous
between the larger and smaller grains, which infers the inhomogeneous mixing of Fe and Cu during
mechanical alloying (MA). NC structure and the mechanical force-enhanced fast diffusion are the
reasons of the formation of supersaturated solid solutions in immiscible systems with positive
enthalpy of mixing. HREM observations carried out with the thermally decomposed Fe60Cu40 solid
solution show that the Nishiyama (N-W) or Kurdyumov-Sachs (K-S) orientation relationships exist
between a-Fe and Cu. Energy dispersive X-ray spectra (EDXS) results show that the Cu content in
these a-Fe grains reaches as high as 9.5 at.% even after heating to 1,400°C, which is even higher
than the maximum solubility of Cu in g-Fe at 1,094°C. Microsc. Res. Tech. 40:101–121, 1998.
r 1998 Wiley-Liss, Inc.
Mechanical alloying (MA)/ball milling (BM) (often
refers to the milling of single component powders) has
attracted considerable interest from the materials science community in recent years. This is mainly due to
the wide range of materials exhibiting intriguing properties and structures that can be fabricated using this
method. Many materials such as oxide dispersion
strengthened (ODS) alloys (Benjamin, 1970), intermetallics (Morris and Morris, 1996; Nash et al., 1995),
amorphous and nanocrystalline materials (Fecht, 1995;
Koch et al., 1983; Schwarz and Koch, 1986; Shingu et
al., 1988), and a number of nonequilibrium structures
(Gente et al., 1993; Huang et al., 1994a; Ma and
Atzmon, 1991; Schwarz, 1996; Shingu and Ishihara,
1995; Shingu et al., 1990) have been synthesized by
MA. Although the underlying fundamentals of the MA
process are increasingly well understood (Koch, 1995),
many details during the mechanical treatment of solids
remain unclear. For example, it is well known that
repeated mechanical deformation and cold-welding often lead to the formation of nanocrystals, but the
formation mechanisms and the structural characterization of the nanocrystals are poorly understood. This is
mainly due to the difficulty in preparing specimens
from mechanically alloyed powders for direct transmission electron microscopy (TEM) observations. The simplest method involves suspending the powder in a
volatile liquid and pipetting the suspension on to a
carbon support film. Usually, TEM images can only be
obtained from the edges of mechanically alloyed powders. However, the microstructure of the edges may
differ from that in the interior of the particles, and for
powders with a large size distribution the small fraction may not be representative of the whole. A complex
method is that the powders are consolidated by hot
compaction or by extrusion, but the original structure
during such a process may be destroyed. Most commonly, powder and nickel are plated over a substrate
and the resultant composite is electrochemically polished (Field and Fraser, 1978; Kang and Benn, 1987;
Yang et al., 1987). Recently, two novel techniques
[powder-fixing method (Huang et al., 1994b) and powder-dispersing method (Huang et al., 1994c)] for specimen preparation of mechanically alloyed powders were
developed. Using the two new techniques some original
microstructural information of ball milled materials
was obtained in our laboratory in the past 6 years and is
reviewed in this paper. This type of work has been less
commonly reported due to difficulties in specimen preparation but is, in fact, very important since many results
obtained by MA or BM are rather puzzling and in many
cases direct observations seem to be very necessary. It
is expected that TEM or high-resolution electron micros-
Contract grant sponsor: National Natural Science Foundation of China.
*Correspondence to: H.Q. Ye, Laboratory of Atomic Imaging of Solids, Institute
of Metal Research, Chinese Academy of Sciences, Shenyang 110015, China.
E-mail: [email protected]
Received 20 February 1996; accepted in revised form 14 March 1996
copy (HREM) could shed some new light on the existing
An irregular copper powder (99.0% 1 pure) with a
particle size smaller than 70 µm was milled in a
vibratory ball mill in an argon atmosphere. Some
Fig. 1. SEM micrograph of copper powders after 20 hours of
milling, showing the formation of spheres.
powders after 20 hours of ball milling have agglomerated to spheres or flat fragments with diameters of
about 2–2.5 mm due to cold-welding, as shown in
Figure 1. Some of them were mechanically ground,
dimpled, and ion thinned for HREM observations.
During MA of Fe-Cu system, two different methods
were used to prepare the TEM specimens. For mechanically alloyed Fe16Cu84, the powder-dispersing procedure was used, which involved mechanical grinding
and subsequent ion beam thinning of the small agglomerates formed during MA (Huang et al., 1994c). The
specimens for Fe60Cu40 powders after 1, 3, and 60 hours
of MA were prepared by the powder-fixing method
(Huang et al., 1994b), which involved mixing the powders into a small amount of epoxy resin and then a piece
of metal net was embedded into the mixture. The
as-prepared samples were flattened and solidified for a
few hours, and then were mechanically ground, dimpled,
and finally thinned by ion beam milling. The MA
experiment was performed in a planetary ball mill
under the protection of an argon atmosphere.
The mechanically alloyed Fe60Cu40 powders were
heated in a differential thermal analyzer (DTA). After
DTA heating to 1,400°C and then cooling to room
temperature the powders have agglomerated to a small
Fig. 2. TEM micrographs of copper powders after 20 hours of
milling. a: Bright field image. b: Dark field image. c: Corresponding
Fig. 3. HREM images showing complex configurations of higher order twins. 1–5 indicate the five twin
variants. Arrowheads in b indicate the grain boundary.
Fig. 4. a: HREM image showing the lenticular morphology and the thickening of the multiple twins;
the steps at one of the TB are denoted by arrowheads. b: Local magnification of shear bands in a, showing
the dislocations piling up in it. c: Local magnification of a, showing a low-angle grain boundary produced
in the shear bands. Notice the 51116 planes of region A deviating by about 5°.
cylindric block with length 5 mm and radius 1.5 mm.
The agglomerate was cut into thin slices, polished, and
ion beam thinned for TEM observations. HREM observations were carried out with a JEM-2000EXII electron
microscopy operated at 200 kV. The energy dispersive
X-ray spectroscopy (EDXS) was performed in a Hitachi
HF-2000 FEG TEM with a probe diameter of about 1
nm in the microanalysis mode.
Ball Mill of Pure Copper
(Huang et al., 1995, 1996a)
The grain size obtained from X-ray diffraction (XRD)
using the Scherrer formula for the unmilled copper is
340 nm. From TEM observations, after 20 hours of
milling the grain size varies from region to region, but
generally it falls in the range of 10–100 nm. Typical
bright and dark field images with their corresponding
electron diffraction pattern (EDP) are shown in Figure
2a–c, respectively. The different grain size obtained in
one specimen infers that the microstructure is inhomogeneous through the specimen. HREM revealed that a
number of deformation twins were generated in these
Cu grains. Figure 3 shows higher order twins formed in
the Cu grains. In Figure 3a the five segments do not
arrange in the manner of fivefold symmetry. In order to
fill the whole space, many internal distortions were
produced, as evidenced by mismatch dislocations at the
interface between segments 1 and 2, and the multiple
twins in segment 2 as well as the edge dislocation in
segment 5, etc. Figure 3b shows other high-order twins.
The strained regions in segments 4 and 5 cross continuously through the twin boundary (TB). In some cases,
twinning takes place on only one set of twinning planes.
Figure 4a shows such a grain that was fully deformed
by multiple twinning. Most of the twin lamellas, which
Fig. 4.
are thickened by passage of some steps, are curved and
their tendency to be lenticular has become obvious. The
TBs are not one atomic plane thick but extend over a
few layers, and in many regions they have developed to
bands with a width of 2–5 nm. Figure 4b is a higher
magnification of a TB in Figure 4a, and it shows clearly
that a number of dislocations are piling up in it. It is
hereby suggested that the described structure is in fact
the ‘‘shear bands’’ (SBs). Figure 5 shows SBs formation
in another grain. There are four SBs in the interior of
this grain with a width of 5–10 nm and an average
spacing of 10 nm running parallel to the slip planes. In
SB 1, microtwins can also be found. The lattice image in
most of the SBs has heavily distorted, indicating that
severe plastic deformation has occurred in these regions. It should be mentioned that the lower right part
of this grain was not deformed, and a misorientation of
about 10° has been produced between the deformed
region and the undeformed one. Thus, a low-angle grain
boundary (LAGB) was formed.
With the development of the SBs, subgrains tend to
form in these bands, and this has been accepted as a
typical mechanism for grain size reduction during
plastic deformation (Fecht et al., 1990; Hansen, 1982)
(hereafter referred to as the first way). We indeed
observed some LAGBs in some of the SBs: a typical
example is shown in Figure 4c, in which the lattice
image of region A has rotated an angle of about 5° with
respect to the twinning planes, and there is a strong
tendency to form subgrains. In addition to the first way,
some LAGBs were found to form at the tip of the SBs
(Fig. 5) or the TBs (hereafter referred to as the second
way). Figure 6a shows a HREM image of four grains.
Deformation is rather inhomogeneous through these
four grains. Grain 4 was deformed by multiple twinning, while grain 1, 2, and 3 were not deformed. Grain 1
has rotated an angle of about 4° with respect to grain 4.
From a local magnification of the grain boundary (GB)
as shown in Figure 6b, it can be seen that a number of
dislocations are piling up in the GB region. It is
Fig. 5. HREM image showing the shear bands (denoted by 1–4), and the low-angle grain boundary
formed at the tip of the shear bands (indicated by the black arrowheads). Disordered and lattice rotation
regions are marked by black dots and a white arrow, respectively, the short and long black lines indicating
the microtwins and the misorientation between the two neighboring grains, respectively.
interesting that most of these dislocations are in pairs,
and generally they belong to two types (denoted by ‘‘A’’
and ‘‘B’’). The Burgers vector bA 5 (1/2)[101] or
(1/2)[011], while bB 5 (1/2)[011] or (1/2)[101], and they
both have an angle of 60° with respect to the dislocation
With continuous refinement of the grain size, the
number of GBs increased rapidly. Figure 7 shows a
HREM image of a GB. Most of this GB is ordered, but
lattice distortion in local regions can also be detected
(see regions ‘‘P’’ and ‘‘O’’). Figure 8a is a typical HREM
image containing a number of grains and GBs. All the
GBs in Figure 8a are ordered, curved, and show a
strained appearance. Figure 8b is a local magnification
of grain 1 in Figure 8a; a number of dislocations were
found in the interior of this grain. All the observed
dislocations are also in pairs. Detailed analyses found
that these dislocations belong to two different types
(marked by ‘‘A’’ and ‘‘B’’) and the Burgers vector bA 5
(1/2)[011] or (1/2)[101], and both have an angle of QA 5
60° with dislocation lines. So dislocation A is a mobile
one. However, bB 5 (1/2)[110], QB 5 90°, and it is
It was believed for many years that a face-centercubic (f.c.c.) metal does not deform by twinning, since
they have enough slip systems to accommodate deformation. In fact, mechanical twins were never observed in
copper when deformed at room temperature by cyclic
fatigue (Liu et al., 1994; Winter et al., 1981) or tensile
stress (Hansen and Ralph, 1982). However, Blewitt et
al. (1957) showed that twins were formed during the
tensile testing of certain copper crystals at low temperatures. Since this initial work, several workers have
reported additional observations of this mode of deformation.
Smith (1958) found that copper after shocks of above
about 200 kbars shows numerous mechanical twins,
which, however, make no significant contribution to
hardening. Johari and Thomas (1964) showed that
after a critical pressure the mode of deformation
changed from slip to microtwinning in explosively
deformed Cu. The critical pressure depends on the
stacking fault energy. Matthews (1970) pointed out that
the deformation of copper thin film under tensile
stresses was accompanied by the generation of lenticu-
Fig. 6. a: HREM image showing a low-angle grain boundary formed at the tip of multiple twins. Black
dots indicate the grain boundaries, black lines show a misorientation of about 5° between grain 1 and 4. b:
Local magnification of the grain boundary between grains 1 and 4 showing a number of dislocations piling
up on the grain boundary. A and B indicate the two types of dislocations.
Fig. 7. HREM image showing the ordered state of a grain boundary. Arrows indicate the grain
boundary. P and O show the lattice distortion regions.
lar twins whose growth seemed to take place by the
migration of twinning dislocations.
It appears from the present state of research on
mechanical twinning that the latter is favored by
deformation at low temperatures or very high strain
rates as well as by lowering stacking fault energy (by
alloying) (Thornton and Mitchell, 1962).
The following factors may account for the generation
of mechanical twins in ball milled copper (Huang et al.,
1995, 1996a). Firstly, the pressure induced by BM was
larger than the critical shear stress for twinning.
Secondly, the grain size reaches a critical value below
which twinning is the preferred mode of deformation.
Finally, the strain rate induced by BM is rather high.
The resultant materials from MA are usually
nanocrystals. MA produces the nanostructure not by
cluster assembly but by the structural decomposition of
coarser-grained structures as a result of severe plastic
deformation. While the structural characteristic of the
nanostructured materials (NC-materials) produced by
other methods, such as gas-condensation/vacuum compaction (Gleiter, 1990), have been studied extensively,
that produced by MA was rarely reported. We got some
structural information of NC-copper produced by MA at
an atomic level. Most of the investigated nanograins
are heavily strained and contain a high density of
dislocations. This is in contrast to the NC-Cu prepared
by sliding wear (Ganapathi and Rigney, 1990), in which
no structural defects were detected inside grains. This
is possibly caused by the different grain size of the
resultant materials. In our case the grain size of the
resultant NC-Cu is 10–100 nm, while in the latter one,
a grain size of 4–5 nm was reported. Generally, if the
grain size is too small, the possibility of activating a
dislocation is rare according to the Hall-Petch relationship.
The GB structure in NC-materials is still controversial in terms of whether a ‘‘gas-like’’ structure can exist
in the GBs of NC-materials. Most of the research
accomplished to date shows that the GB structure of
NC-materials is ordered and is similar to those in
coarse-grained materials (Ganapathi and Rigney, 1990;
Li et al., 1993; Siegel, 1994; Suryanarayana, 1995). In
ball milled copper, the conclusions are similar to those
drawn above. All the investigated GBs are ordered,
curved, and strained. The lattice distortion and nanovoids were frequently encountered in the matrix nearby
GBs, and GBs contain a high density of dislocations in
several cases. It can be concluded that there is no
characteristic difference between the nanograined and
coarse-grained GBs in ball milled copper.
Mechanical Alloying in an Immiscible
Fe-Cu System
A number of studies have been focused on the
FeXCu(100-X) system since it was shown that the miscibility of Fe in Cu can be greatly enhanced, 0 , X , 60,
through MA (Corespo et al., 1993; Eckert et al., 1993;
Jiang et al., 1993; Ma and Atzmon, 1995; Ma et al.,
1993; Marci et al., 1993; Uensh et al., 1992; Yavari et
al., 1992). However, direct observation on the mechanically alloyed Fe-Cu alloys has not been reported. In the
present paper, the mechanically alloyed Fe16Cu84 and
Fe60Cu40 were directly observed by HREM. Both atomic
Fig. 8. a: HREM image showing the nanocrystals formed in ball milled copper. Arrowhead indicates
the strain and the lattice distortion. b: Local magnification of grain 1 in a showing two types of
dislocations (60° and 90° dislocations marked by A and B, respectively). The Burgers circuits of each were
Fig. 9. TEM micrographs of MA Fe16Cu84. a: Bright field image. b:
Dark field image. c: Corresponding EDP.
Fig. 10. TEM micrographs of MA Fe60Cu40. a: Bright field image. b:
Dark field image. c: Corresponding EDP.
level structure and nano-scale composition information
were obtained.
Figures 9 and 10 are the TEM micrographs corresponding to specimen Fe16Cu84 and Fe60Cu40, respectively. Both the EDPs show a set of fcc phases, which
agrees well with the XRD results (Huang et al., 1996c).
From the bright and dark field images it can be seen
that the grain size varies from 10 to 100 nm. Figure 11
shows a HREM image of 3 grains in MA Fe16Cu84: in the
large grain, i.e., grain 2, small twin lamellas can be
found. However, the structure of the smaller grains 1
and 3 is almost perfect. These GBs are curved and
slightly strained; lattice distortion or nanovoids can
also be detected in local regions, as marked by circles in
Figure 11. In many cases, the size of the Fe grains can
be as small as 2–5 nm, and these small Fe grains tend to
insert into the interior of the Cu grains, as shown in
Figure 12. Figure 13 is a HREM image containing a
Fig. 11. HREM image of GB structures in MA Fe16Cu84. 1–3 denote the number of grains. Arrowheads
indicate the GBs. Circles mark nanovoids or local disordering in the GBs.
Fig. 12.
HREM image of MA Fe16Cu84 showing an Fe grain inserting into the center of a Cu grain.
Fig. 13. HREM image of GB structures in MA Fe16Cu8. Circle indicates a small Fe grain. 1–5 denote
the number of grains.
number of grains and GBs in an Fe16Cu84 specimen.
These grains are flattened and some nano-scale layered
structures have developed. Even though the diffraction
peak from the bcc phase has completely disappeared in
the XRD of this specimen, small Fe domains can still be
detected in the HREM image shown in Figure 13, as
indicated by a circle.
Figure 14 shows the HREM image of Fe60Cu40 after 1
hour MA. The grain size has reached nanometer scale.
The Cu grains are coarse and their size is about 5–30
nm; while the Fe grains are very small (less than 5 nm),
and some domains with only a few atomic planes were
also detected, as indicated by circles in Figure 14. After
3 hours MA, some large Cu grains can still be found, as
shown in Figure 15. It can be seen that this large Cu
grain is heavily strained and the deformation occurred
in a rather inhomogeneous way. The center of Figure 15
is a highly deformed region of about 10 nm in width
that extends throughout the grain. The HREM image of
this region reveals microstructures consisting of individual grains with a diameter of about 10 nm, which are
slightly rotated with respect to each other at a rotation
angle of less than 20°. After 60 hours MA, most of the
grains are very small and their size is less than 5 nm.
Although no bcc diffraction is observed in both the EDP
and XRD, some ultra-fine Fe domains can still be
detected, as marked by the circles in Figure 16.
The composition of both the grain interior and the
GBs of specimen Fe16Cu84 was analyzed by EDXS.
Typical EDXS spectra from the grain interior and the
Fig. 14.
HREM image of Fe60Cu40 after 1 hour of MA. Circles show some small Fe domains.
GBs are shown in Figure 17. The result showed that the
average Fe content (about 16%) in both the grain
interior and the GBs is very close to the designed
composition (16%), but the Fe content in both cases is
rather inhomogeneous; generally, the smaller the grain
size, the higher the Fe content. However, in most cases
the Fe content falls in the range of about 5 to 20%,
which infers that most Fe atoms have dissolved in the
Cu lattice and true alloying has occurred.
A standard method for identifying the phase that is
formed under MA is XRD. In many investigations, the
lattice parameters of phases such as solid solutions are
determined from the XRD patterns and plotted as a
function of composition. Continuous variation of the
lattice parameters with composition, such as a Vegard
law dependence, is commonly taken as an indication
that the material is single phase. However, the question
arises as to the homogeneity of such solid solution
phases. While the structure of a solid solution is readily
identified by XRD, data about its homogeneity are not
easily obtained. This has been demonstrated by X-ray
investigations on Co/Cu multi-layers and on decomposed Co-Cu solid solutions (Michaelsen, 1995). Although these materials consist of separate fcc Cu and
fcc Co regions, they exhibit a diffraction pattern that
can hardly be distinguished from the XRD of a solid
solution having the same overall composition, if the Cu
and Co regions are coherent and their size is smaller
than about 5–8 nm. Michaelsen (1995) therefore concluded that it is not possible to determine the supersaturation of a solid solution using a conventional XRD.
Complementary measurements have to be performed
in order to prove the homogeneity of a supersaturated
alloy. In the present experiment, EDXS result shows
that the average Fe contents in the interior of most fcc
grains in the FCB specimen are close to the designed
composition, which is certainly direct evidence for the
formation of supersaturated solid solutions. However,
the mixing is not homogeneous since the data of the Fe
contents in both the interior of grains and the GBs are
rather scattering.
The mechanism by which a solid solution with positive heat of mixing is formed upon milling is still a
subject of controversy. Yavari et al. (1992) addressed the
interfacial energy as a driving force for homogenization. They found that the chemical component of the
interfacial energy is insufficient to dissolve nanograins
of average size as determined by XRD, and only the tail
of the grain size distribution with r , 1 nm will dissolve
initially. They then suggested that codeformation re-
Fig. 15. HREM image of Fe60Cu40 after 3 hours of MA, showing the generation of subgrains in the
shear band (denoted by 1–3).
Fig. 16.
HREM image of Fe60Cu40 after 60 hours of MA. Circles indicate some small Fe domains.
Fig. 18. The chemical contribution schem of the Fe/Cu interfaces to
the enthalpy of a Fe60Cu40 composite as a function of the size d of the
Fe and Cu regions.
Fig. 17. EDXS spectra of MA Fe16Cu84 in the interior grains (a) and
in the GBs (b).
sults in the formation of thin Fe layers, which, by
undergoing necking to striction, generate smaller fragments. Their calculations show that when the fragments reach diameters as small as 2 nm, the Fe-Cu
interfacial energy, which has a large chemical component, results in their dissolution. Further deformation
during MA then generates more such fragments that
also dissolve until a complete solution is obtained. This
kind of tip radii or fragments were indeed frequently
observed in our HREM images (Figs. 12–14 and 16),
which are surely direct evidence to support Yavari et
al.’s model. According to this model, in going from a
NC-mixture of bcc-Fe and fcc-Cu crystals to a fcc solid
solution of equal grain size, the interfacial energy is
reduced by DE 5 6Vmschem/d, where Vm is the molar
volume and schem is the chemical contribution to the
interfacial energy. DHmix 5 DE 1 TDSmix, where DHmix
and DSmix are the enthalpy and entropy of mixing,
respectively. The refinement of the microstructure enhances the chemical contribution of the interface enthalpy as shown in Figure 18 for different sizes of the Fe
and Cu regions. It can be seen that the grain size down
to 1 nm raises the enthalpy of a Fe60Cu40 composite
above 17 KJ/mol, which is sufficient to allow for a phase
transformation of the composite into a supersaturated
solid solution; there is still additional enthalpy from the
heavy mechanical deformation, which is not considered. Thus, the enthalpy due to the contribution of
alloying at the interface and mechanical deformation
may be large enough to overcome the calculated positive enthalpy of mixing of 12 KJ/mol for Fe60Cu40. It is
useful to compare the results of MA with that of plastic
deformation under pressure in Fe-Cu. Teplov et al.
(1995) recently reported that NC structure and nonequilibrium solid solutions of Fe80Cu20 and Fe20Cu80
were obtained after severe plastic deformation by shear
under pressure. Based on grain rolling, a mechanism of
superplastic deformation has been proposed. The main
idea of such a mechanism is that plastic deformation of
a fine-grain polycrystal is fully accommodated by the
GBs. In the meantime, very fast volume diffusion may
also exist (Gryaznov et al., 1992).
Very fast GB diffusion and volume diffusion occur
during the cold forced superplastic deformation by
shear under pressure. Obviously, diminishing the grain
sizes to minimal value, fast diffusion, and high internal
stress in the fine-grain materials also take place during
MA. During the early stage of MA, the grain size
decreases rapidly to a steady value after only a few
Fig. 19.
Susceptibility of Fe60Cu40 solid solution during heating and subsequent cooling at 10°C/min.
sliding. This can increase the diffusion coefficient tremendously. The volume diffusion coefficient Dv may be
calculated by the expression (Gryaznov et al., 1992):
4Dv · t 5 d2, where t is the time of plastic deformation, d
is the grain size. For example, the typical milling time
is about 1 hour from a NC-Cu and -Fe composite to a fcc
solid solution, but at any instant of time during MA
only a small fraction of the powder is undergoing
impact or shear due to ball motion. Thus the deformation time t is also only a small fraction of the total
milling time. Supposing t 5 600 seconds, and the grain
size before alloying is about 5 nm; substituting t and d
to the above equation, we obtain Dv 5 10216cm2/sec,
which is several orders of magnitude higher than the
volume coefficient of the diffusion of Cu in Fe at 300 K
(Dv 5 10242cm2/sec) (Anand and Agarwala, 1966; Teplov
et al., 1995). So NC structure and mechanical driven
fast diffusion are the reasons for the formation of solid
solution in immiscible systems with positive enthalpy
of mixing.
Fig. 20. XRD patterns of the as-milled Fe60Cu40 and after annealing at different temperatures.
hours (generally less than 5 hours) (Eckert et al., 1993;
Huang et al., 1994a) due to the mobility of dislocations.
Further deformation can be accommodated only by GB
Thermal Decomposition of Mechanically Alloyed
NC-Fe60Cu40 (Huang et al., 1996b)
Figure 19 shows the susceptibility of Fe60Cu40 during
decomposition heating to 915°C and subsequently cooling. The susceptibility tends to vanish around the Curie
temperature Tc 5 250°C of the supersaturated fcc-FeCu
solid solution before increasing as decomposition sets in
above 300°C. The decomposition finishes at about 450°C.
Abrupt changes of the susceptibility are observed at
about 640 = 760°C on heating and from 800 to 640°C on
Figure 20 shows the XRD patterns of Fe60Cu40 after
milling and subsequently annealing at different temperatures. It is seen that after heating to 300°C, the bcc
phase appears. The Fe60Cu40 specimen after DTA heating to 1,400°C and cooling to room temperature was
characterized by TEM. Figure 21 shows the general
TEM morphology of some Fe grains. The four Fe grains
shown in Figure 21a with a size of about 200–400 nm
are nearly spherical in shape; in some cases they
contain many dislocations (indicated by arrowhead).
Those shown in Figure 21b and c with a size of about
500–600 nm take an approximately cubic form, and
some structural defects can also be detected in the
center of the Fe grain shown in Figure 21b. Figure 22a
shows a typical EDP between the a- and g-phases. One
can see that the N-W orientation relationship, i.e.,
(110)a//(111)g; [00 1]a//[01 1] g, exists between them,
which can also be clearly identified from the HREM
image shown in Figure 22b. In several cases, another
orientation relationship, the K-S orientation relationship, was also encountered, and a typical example is
shown in Figure 23. In this case, the K-S orientation
relationship holds with: (110)a//(111)g; [1 1 1]a//[01 1]g.
Even after DTA heating to 1,400°C and subsequent
cooling to room temperature, some small Fe grains
were still detected. Figure 24 shows a low magnification
HREM image of an Fe grain with approximately spherical morphology and a size of about 50 nm. A number of
edge dislocations are realized in the interior of this
grain. The strained region at the interface extends to
about 2 nm. The corresponding EDP in Figure 24b
shows that there also exists the K-S orientation relationship between Fe and the matrix.
The composition of both the a-Fe and fcc Cu phases
was analyzed by EDXS in a HF-2000 FEG-TEM with a
probe diameter of 1 nm. The result showed that the Fe
content in the matrix is about 1 at.%, whereas the Cu
content within the Fe grains reaches as high as 9.5%,
which even exceeds the highest solubility of Cu in g-Fe
(8.5% at 1,094°C). A typical EDXS spectrum of an Fe
grain is shown in Figure 25.
From the thermomagnetic measurement and heat
treatment results, the decomposition of the fcc Fe60Cu40
supersaturated solid solution could take place in the
following way, on heating: fcc Fe60Cu40 (I) =300=460°C aFe(Cu) 1 g-Fe(Cu) 1 Cu(Fe) (II) =640=760°C g-Fe(Cu) 1
Cu(Fe) (III) =1,081°C g-Fe(Cu) 1 Cu(melted) (IV) and upon
the subsequent cooling: g-Fe(Cu) 1 Cu(Fe)(melted)
(IV)=1,126°C g-Fe(Cu) 1 Cu(Fe) (V) =800=640°C a-Fe(Cu) 1
Cu(Fe) (VI).
The decomposition of the fcc solid solution from stage
I to II has been investigated by many researchers
(Corespo et al., 1993; Drbohlav and Yavari, 1994; Jiang
et al., 1993; Marci et al., 1993). At the initial stage of
decomposition, precipitation begins with the formation
of the g-Fe phase. With prolonged annealing or with the
rise of temperature, g-Fe precipitates grow, losing their
fully coherency. When the annealing temperature is
further raised, stable a-Fe precipitates nucleate from
large g-Fe precipitates or directly from the dislocation
tangles of the fcc Cu phase, as well as at the grain
boundaries. In the meantime, due to excellent coherency between the g-phase and the matrix fcc phase,
some small g-precipitates will not transform to the
a-phase. The sluggish transformation of g- to a-Fe in a
Cu matrix, which only reaches completion after cold
working, was already reported by Easterling and Miekk
Fig. 21.
TEM morphology of Fe60Cu40 solid solution after DTA run.
With the rising of temperature, there is a a- = g-Fe
transformation at 640 = 760°C (II = III), as proved by
the abrupt decrease of the susceptibility shown in
Figure 19. The reverse transformation is martensitic
and occurs from 800 to 640°C. Obviously, the a t g
transformations occur in a wide temperature range.
The different grain size may be responsible for this
behavior. After DTA heating to 1,400°C, a different
grain size from 50 (Fig. 24) to 600 nm (Fig. 21) was
observed. It has been reported (Cech and Turnbull, 1956)
Fig. 22. TEM micrographs showing the N-W orientation relationship between the a- and g-phases. a:
EDP image. b: HREM image. Arrowheads indicate the interface between the a- and g-phases.
Fig. 23. TEM micrographs showing the K-S orientation relationship between the a- and g-phases. a:
EDP image. b: HREM image. Arrowheads indicate the interface between the a- and g-phases.
that the martensitic transformation temperature can be
decreased remarkably with decreasing grain size.
The orientation relationship between the austensite
and martensite in as-cast Fe-Cu alloys was established
by Easterling and Weatherly (1969), and they found
that the K-S orientation relationship exists between
the austenite and martensite. However, in thermally
treated MA Fe60Cu40, both the K-S and N-W orientation
relationships were found between a-Fe and Cu phases.
Due to excellent coherency between the g and Cu
phases, this orientation relationship should also represent the one between g- and a-Fe. Also, after DTA
heating to 1,400°C, the solubility of Cu in a-Fe exceeds
the highest solubility of Cu in g-Fe (8.5 at.% above
Fig. 24. TEM micrographs showing a nanometer-scale a-Fe(Cu) remained after being heated to
1,400°C (a), and the K-S orientation relationship between the a- and g-phases (b). The indexing of EDP is
the same as that of Figure 23.
1,094°C). This may be caused by the small grain size,
since the grain size in Fe60Cu40 after DTA run is still not
very large (50–600 nm).
The martensitic transformation in bulk Fe-Cu alloys
has been investigated by many researchers (Easterling
and Miekk, 1967), but due to limited solubility in
equilibrium between Cu and Fe, all the investigated
compositions are Fe-rich ones with the Cu content less
than 2 at.%. From these investigations, it is suggested
that diffusionless decomposition of austenite occurs
both by massive and martensitic mechanisms (Easterling and Miekk, 1967). The transformation tempera-
quite different from that in as-cast alloys, which is
possibly due to the different fabrication process and the
non-equilibrium microstructure (especially the high
content of Cu in Fe) of MA. Because of the lack of
substantial experimental data in as-cast alloys, we
could not satisfactorily compare our results with that of
the bulk materials at the present time.
This work was supported by the National Natural
Science Foundation of China, which is gratefully acknowledged.
Fig. 25.
A typical EDXS spectrum of the a-Fe phase.
Fig. 26. The massive and martensite transformation temperatures
in Fe-Cu alloy after Easterling and Miekk (1967), the allotropic
transformation start temperature after Hansen (1958), and the martensite transformation temperature obtained in the present work.
tures for both these reactions have been determined in
the bulk materials (Easterling and Miekk, 1967), and
are given in Figure 26. It is seen that there is no
experimental data corresponding to the Cu composition
in our alloy, but by extrapolating the Mass or Ms curves
to our composition (9.5% Cu), the transformation temperature should be less than 500°C, which is lower than
700 = 622°C in the Fe60Cu40 alloy. Even by extrapolating the a t g allotropic transformation line to 9.5%, a
temperature of 500°C might be obtained, which is still
lower than that in Fe60Cu40. It seems that the martensitic transformation in thermally treated MA Fe-Cu is
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