Microstructure and mechanical behavior of AZ91 Mg alloy processed by equal channel angular pressing

Article abstract of DOI:10.1016/j.jallcom.2004.10.050

A fine-grained AZ91 alloy was prepared by equal channel angular pressing (ECAP). The evolution of the microstructure and the deformation behavior were investigated as a function of the number of ECAP passes. At room temperature the tensile strength of the ECA pressed specimens is higher than that of the initial state due to the grain refinement and the increase of the dislocation density during ECAP. Above 200 °C the ductility increases due to the increase of the relative fraction of c + a dislocations and also owing to the breakage of the rod-like Al₁₂Mg₁₇ precipitates during ECAP processing.

Full text of DOI:10.1016/j.jallcom.2004.10.050

Journal of Alloys and Compounds 394 (2005) 194–199
Microstructure and mechanical behavior of AZ91 Mg alloy
processed by equal channel angular pressing
a,b
c,∗
c
´
K. Mathis , J. Gubicza , N.H. Nam
a
Department of Metal Physics, Charles University, Ke Karlovu 5, 12116 Praha 2, Czech Republic
b
Department of General Physics, Eo¨tvo¨s University, P.O. Box 32, H-1518 Budapest, Hungary
c
Department of Solid State Physics, Eo¨tvo¨s University, P.O. Box 32, H-1518 Budapest, Hungary
Received 28 September 2004; accepted 20 October 2004
Available online 10 December 2004
Abstract
A fine-grained AZ91 alloy was prepared by equal channel angular pressing (ECAP). The evolution of the microstructure and the deformation
behavior were investigated as a function of the number of ECAP passes. At room temperature the tensile strength of the ECA pressed specimens
is higher than that of the initial state due to the grain refinement and the increase of the dislocation density during ECAP. Above 200 ◦C the
ductility increases due to the increase of the relative fraction of c + a dislocations and also owing to the breakage of the rod-like Al12Mg17
precipitates during ECAP processing.
© 2004 Elsevier B.V. All rights reserved.
Keywords: Metals; Dislocations and disclinations; X-ray diffraction
1. Introduction
gular pressing (ECAP) which results in bulk, homogeneous
submicron or nanocrystalline microstructure without reduc-
tion of the cross section of the billet [2]. During the repetitive
pressing very high strains are achieved. The microstructure
developed after ECAP depends on many factors, e.g. rotation
of specimens between consecutive passes, number of passes
through the die (the increase of total strain) and temperature
of deformation [3,4].
The aim of the present paper is to study the effect of ECAP
on the microstructure and the mechanical properties of AZ91
magnesium alloy in the temperature range from room tem-
perature to 300 ◦C.
Magnesiumalloys, beingamongthelighteststructuralma-
terials, are very attractive in many applications. Their excel-
lent strength to weight ratio predestines Mg-based alloys for
applications as structural components in automobile and air-
craft industries. The disadvantages of these materials are the
poor workability because of their hexagonal structure and
the degradation of mechanical properties at elevated temper-
atures [1].
It is well known that materials with small grain size have
excellent mechanical properties. Reduction of the mean grain
size is expected to increase the yield strength (and also the
ultimate tensile strength) at room temperature and to promote
superplastic deformation at higher strain rates and/or lower
temperatures than those conventionally used for large grain
size materials. Ultrafine grained microstructure in metals can
be produced by severe plastic deformation (SPD) techniques.
One of the most often used SPD method is equal channel an-
2. Experimental procedures
2.1. Sample preparation and experimental techniques
The main alloying elements in the AZ91 magnesium alloy
studied here are 9 wt.% Al, 1 wt.% Zn and 0.2 wt.% Mn. The
alloy was solution heat-treated for 18 h at 413 ◦C. The heat-
treated samples with dimesions of 10 mm × 10 mm × 60 mm
∗
Corresponding author. Tel.: +36 1 372 2876; fax: +36 1 372 2868.
E-mail address: gubicza@ludens.elte.hu (J. Gubicza).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved.
doi:10.1016/j.jallcom.2004.10.050

K. Ma´this et al. / Journal of Alloys and Compounds 394 (2005) 194–199
195
were pressed through an ECAP die. The angle between the
intersecting channels was 90◦. The ECAP was carried out at
constant temperature of 270 ◦C. The number of ECAP passes
was two, four, and eight. The samples coated with molybde-
num disulfide (MoS2) were pressed at a rate of 5 mm/min
using route C (i.e. the samples were rotated around their lon-
gitudinal axis by 180◦ after each pass) [4]. The total strain,
εn, in the specimen after N passes can be estimated as [4,5]:
in the specimen. The q1 and q2 values for the eleven possible
ˇ
slip systems in Mg have been calculated according to Kuzel
and Klimanek [8] and listed in Table 2 in reference [9]. The
11 dislocation slip systems can be classified into three groups
¯
based on their Burgers vectors: b1 = 1/3ꢀ2 1 1 0ꢁ (ꢀaꢁ-type),
¯
b2 = ꢀ0 0 0 1ꢁ (ꢀcꢁ-type) and b3 = 1/3ꢀ2 1 1 3ꢁ (ꢀc + aꢁ-type).
There are 4, 2 and 5 slip systems in the ꢀaꢁ, ꢀcꢁ and ꢀc + aꢁ
Burgersvectorgroups, respectively. Acomputerprogramwas
elaborated to determine the Burgers vector population from
φ
(m)
(m)
the measured values of q1 and q2 [10]. First the pro-
εn = 1.15N cot ,
(1)
2
gram selects some slip systems from one of the three groups.
where φ is 90◦. The values of the total strain were 2.3, 4.6
and 9.2 for two, four, and eight passes, respectively.
¯
¯
In the second step the calculated Ch k 0q1 and Ch k 0q2 values
of the selected slip systems are averaged with equal weights,
¯
Tensile specimens were machined from the initial sample
and also from the rods produced by ECAP. These specimens
had rectangular cross section of 1 mm × 2 mm and gauge
length of 9 mm. The tensile tests were carried out by an MTS
machine in the temperature range from 20 to 300 ◦C at con-
where Ch k 0 is the average contrast factor corresponding to
the h k 0-type of reflections. This procedure is carried out for
each Burgers vector group. The relative fractions of the three
groups, hi (i = 1, 2, 3) are calculated by solving the following
three equations:
stant strain rate of 5 × 10−4 s−1
.
3
ꢀ
1
The phase composition of the samples was studied by X-
ray diffraction using a Philips Xpert θ–2θ powder diffrac-
tometer with Cu K␣ radiation. The microstructure of the
specimens was investigated by X-ray peak profile analy-
sis, optical microscopy (Olympus) and scanning electron-
microscopy (SEM, Tesla B525). The X-ray diffraction peak
profiles were measured by a high-resolution diffractometer
(Nonius FR591) with rotating Cu anode and a Ge monochro-
mator (Cu K␣1 radiation, λ = 0.15406 nm) [6]. The profiles
were recorded by a linear position sensitive gas-flow detec-
q1(m)
=
=
hibi (Ch k 0q1)i,
(2)
(3)
2
¯
P
i=1
3
ꢀ
1
q2(m)
hibi (Ch k 0q2)i,
2
¯
P
i=1
and
3
ꢀ
hi = 1
(4)
i=1
¨
tor (OED 50 Braun, Munich). The instrumental broadening
¯
¯
was negligible compared to the physical broadening of the
profiles therefore instrumental correction was not applied.
where (Ch k 0q1)i and (Ch k 0q2)i are obtained by averag-
ing for the ith Burgers vector group and P is given as
3
ꢁ
2
¯
P =
hibi (Ch k 0)i. If the three hi weights have positive
2.2. Evaluation of X-ray diffraction data
i=1
values the program stores them as one of the possible solu-
tions. After examining all the possible combinations of the
slip systems, ranges of the three weights are obtained as the
final solution. The detailed description of the procedure is
given in reference [9].
The measured physical X-ray peak profiles were evaluated
by means of the multiple whole profile (MWP) fitting proce-
dure, described in detail in reference [7]. In this method, the
Fourier coefficients of the experimental profiles are fitted by
the theoretical Fourier transforms calculated on the basis of
a model of the microstructure. In this model, the crystallites
have spherical shape and log–normal size distribution. The
lattice strains are assumed to be caused by dislocations. The
procedure has six fitting parameters for hexagonal crystals:
(i) the median and the variance, m and σ, of the log–normal
size distribution function; (ii) the density and the arrange-
ment parameter of dislocations, ρ and M; and (iii) the q1
and q2 parameters in the contrast factors of dislocations. The
volume-weighted mean crystallite size, ꢀxꢁvol, was calculated
from m and σ using the formula given in reference [7]. The
magnitude of M gives the strength of the dipole character of
dislocations: a higher M value corresponds to a weaker dipole
character and a weaker screening of the displacement fields of
dislocations. The q1 and q2 parameters of the contrast factors
depend on the character of dislocations and therefore enable
the determination of the prevailing dislocation slip systems
3. Results and discussion
3.1. Influence of ECAP processing on the microstructure
The ECAP results in a considerable decrease of the grain
size of AZ91 alloy. The process of grain refinement can be
seen in the optical micrographs of Fig. 1. Fig. 1a shows the
initial microstructure after the solution heat treatment. The
decomposition of the grain-structure owing to ECAP begins
in the grain boundary area, as it can be seen in Fig. 1b, where
the microstructure of the specimen deformed by two ECAP
passes is shown. After four passes, further grain refinement
can be observed (Fig. 1c). Finally, the sample subjected to
eight passes has a homogeneous, fine grained microstruc-
ture, as it can be seen in Fig. 1d. Mussi [11] has made trans-
mission electron microscopy (TEM) investigations on these

196
K. Ma´this et al. / Journal of Alloys and Compounds 394 (2005) 194–199
Fig. 1. Optical micrographs of the microstructure for (a) the initial state and the specimens deformed by ECAP for (b) two, (c) four, and (d) eight passes.
specimens. It was found that the initial grain size of 40 ␮m
decreased to 1.2 ␮m after eight ECAP passes. It should be
noted that TEM experiments also revealed that in the latter
sample the grains were divided into smaller subgrains with
the size of 0.1 ␮m.
dislocation cells (or subgrains), which are separated from
each other by small misorientation, typically under 1–2◦. At
the same time it was found previously [7,12] that these dis-
location cells could only be observed separately by electron
microscopy if highly magnified TEM images were studied
very carefully. The usual TEM investigation of SPD metals
gives the grain size, which is higher than the dislocation cell
or subgrain size obtained by X-ray diffraction peak profile
analysis.
The dislocation density increases by a factor of five as a
result of eight ECAP passes (see Table 1). It was found, that
in the initial state the relative fraction of the ꢀaꢁ-type Burg-
ers vectors is the highest. The abundance of this dislocation
type beside the ꢀcꢁ- and ꢀc + aꢁ-type dislocations can be ex-
plained by its smallest formation energy. After high temper-
ature ECAP the relative fraction of ꢀc + aꢁ-type dislocations
increases, which is in agreement with theoretical calculations
and former experimental observations [13].
The results of the X-ray peak profile analysis confirm the
microscopic observation of the refinement of the microstruc-
ture due to ECAP. Table 1 shows the volume-weighted mean
crystallite size, the dislocation density and the relative frac-
tions of the three Burgers vector types for the initial state
and the sample deformed by eight passes. The crystallite size
decreases due to ECAP, which is consistent with the micro-
scopic observations. It is worth to note that after eight passes
the crystallite size (97 nm) determined by X-ray peak profile
analysis is in good agreement with the subgrain size obtained
by TEM (∼100 nm) [11]. Generally, the mean crystallite size
determined by X-rays is lower than the grain size observed
in the TEM images which has been already reported for SPD
materials previously [7]. This can be explained by the fact
that the grains in SPD materials are divided into subgrains
and/or dislocation cells, which are separated from each other
by low angle grain boundaries. The crystallite size in SPD
metals obtained by X-ray diffraction is equivalent to the mean
size of the domains, which scatter X-rays coherently. Con-
sequently, X-ray diffraction makes a difference between the
The X-ray diffractogram of the initial specimen after so-
lution heat-treatment is shown in Fig. 2a. Only the reflections
of a Mg-based solid solution appear in the diffractogram. The
position of the diffraction lines are shifted to higher angles
compared with pure Mg due to the reduction of the lattice pa-
rameters owing to the alloying Al atoms. At the same time,
the high temperature ECA pressing result in the formation of
Table 1
The microstructural parameters of AZ91 alloy for the initial state and after eight ECAP passes
ꢀxꢁvol (nm)
ρ (×1014 m−2
)
ꢀaꢁ-type dislocations (%)
ꢀcꢁ-type dislocations (%)
ꢀc + aꢁ-type dislocations (%)
Initial state
563
0.4
68–86
0–15
0–14
Eight ECAP
97
2.0
56–58
0–2
38–42

K. Ma´this et al. / Journal of Alloys and Compounds 394 (2005) 194–199
197
Fig. 2. X-ray diffractograms for (a) the initial state and (b) after two ECAP
passes.
Mg17Al12 precipitates from the supersaturated Mg(Al, Zn,
Mn) solid solution, as it can be seen for two passes in Fig. 2b.
It is worth to note that a common precipitation heat treat-
ment lasts at least for 8 h [1] while the duration of two ECAP
passes is approximately 50 min. The precipitation could oc-
cur within such a short time because it was facilitated by the
severe plastic deformation. Zhang et al. [14] observed simi-
lar rapid precipitation during high temperature creep test of
Mg–Al alloys. Disk shape continuous and discontinuous pre-
cipitates with bcc structure and rod-like shape form during
ECAP [15,16] and they play an important role during de-
formation. Fracture is mostly initiated along them [17] and
the fine dispersion of precipitates could improve ductility
at high temperature [18]. The shape of precipitates changes
during the ECAP processing. Fig. 3a and b shows the SEM
micrographs characterizing the microstructure in the weakly
and the severely deformed parts of ECAP processed speci-
mens, respectively. The first image is taken from the front
part of the specimen after four passes. The second picture is
obtained from the internal part of the sample deformed by
eight ECAP passes. Comparing the two micrographs it can
be established that the discontinuous Mg17Al12 precipitates
initially have a long rod-like shape which are broken into
smaller parts during ECAP and their dispersion is relatively
homogeneous.
Fig. 3. SEM images on (a) the weakly deformed front part of the sample
after four passes and (b) the severely strained internal part of the specimen
after eight passes.
3.2. Deformation behaviour of ECAP processed samples
Fig. 4. The true stress—true strain curves for the initial state and for the
specimens subjected to ECAP obtained by tension at (a) room temperature
and (b) 300 ◦C.
Fig. 4a and b shows the true stress—true strain curves
obtained by tension for the initial state and the ECAP pro-

198
K. Ma´this et al. / Journal of Alloys and Compounds 394 (2005) 194–199
Fig. 6. The volume-weighted mean crystallite size and the dislocation den-
sity obtained after tensile test at different temperatures for eight ECAP spec-
imen.
Chuvildeev et al. [20], who has found a maximum for both
the elongation and the internal friction at 300 ◦C.
The crystallite size and the characteristic parameters of
the dislocation structure were determined for the eight ECAP
specimen after tension at different temperatures. The volume-
weighted mean crystallite size and the dislocation density as
a function of the testing temperature are plotted in Fig. 6.
Comparing Figs. 5a and 6, one can deduce that the tensile
strength behaves similarly than the dislocation density as a
function of the temperature of deformation. Plotting the ten-
sile strength versus the dislocation density, a monotonic re-
lationship can be observed between the two quantities (see
Fig. 7). This means that the changes of the tensile strength be-
tween room temperature and 300 ◦C can be attributed mainly
to the changes of the dislocation density. The high dislocation
densities after room temperature and 100 ◦C deformation are
owing to the dislocation multiplication processes. Disloca-
tion – forest dislocations, dislocation-precipitate and dislo-
cation – grain boundary interactions are responsible for the
hardening effects [21,22]. Above 100 ◦C the dislocation den-
sity strongly decreases due to the recovery processes. The dy-
namic recovery causes strain softening during deformation.
Beside the recrystallization, annihilation of dislocations due
to the cross-slip of screw dislocations and/or climb of edge
dislocations can be taken into account as the mechanisms
responsible for softening [23,24].
Fig. 5. The tensile strength (a) and the elongation to failure (b) as a function
of the deformation temperature.
cessed AZ91 specimens at room temperature and 300 ◦C, re-
spectively. The tensile strength and the elongation to failure
obtained from the stress-strain curves are plotted as a function
of the deformation temperature in Fig. 5a and b, respectively.
The tensile strength was determined as the maximum stress
on the true stress–true strain curves and it is denoted by σmax
.
The elongation to failure is given as the engineering strain in
percent. It can be seen, that the ECAP procedure results in
the increase of the tensile strength at room temperature as
a consequence of the reduction of the grain size as well as
the increase of the dislocation density. At the same time at
high temperature (T ≥ 200 ◦C), where the recovery processes
become important, the tensile strength decreases and the duc-
tility increases owing to ECAP. At 300 ◦C the elongation of
the ECAP processed materials achieves the superplastic re-
gion (i.e. the engineering strain is higher than 100%). The
superplastic behavior is also proved by the value of 0.45 for
the strain rate sensitivity parameter (m = ∂ ln σ/∂ ln ε˙) ob-
tained by means of strain rate change experiments. From the
microstructural analysis and the mechanical testing it can be
deduced that at high temperatures the large elongation and
the low strength of the ECAP processed specimens compared
to the initial state are most probably caused by two phenom-
ena: (i) the breakage of Al12Mg17 precipitates into smaller
parts which facilitates the dislocation motion (e.g. cross-slip)
resulting in the dynamic recovery of the microstructure; (ii)
the relatively small grain size which enables the material to
deform by the mechanism of grain boundary sliding. Grain
boundary sliding plays an important role in superplasticity.
The high resolution electron microscopy (HREM) investi-
gation of Mabuchi et al. [19] shows, that the high tempera-
ture deformation of AZ91 alloy takes place partly by grain
boundary sliding. This phenomenon was also confirmed by
The activity of non-basal slip systems may also plays an
important role in strain softening [25]. Screw dislocations
of ꢀc + aꢁ-type can move to the next slip planes by double
Fig. 7. The monotonic relationship between the dislocation density (ρ) and
the tensile strength (σmax) for eight ECAP specimen deformed at different
temperatures. The arrow indicates the direction corresponding to the increase
of the deformation temperature.

K. Ma´this et al. / Journal of Alloys and Compounds 394 (2005) 194–199
199
Acknowledgements
This work was supported by the Grant Agency of
Academy of Sciences of the Czech Republic under Grant
A2041203. The authors are grateful for the financial support
of the Hungarian Scientific Research Fund (OTKA) under
Contract Numbers F-047057, T-042714 and T-043247.
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