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Evolution of microstructure in AZ91 alloy processed
by high-pressure torsion
Ahmed S. J. Al-Zubaydi
1,2
•Alexander P. Zhilyaev
3,4
•Shun C. Wang
1
•
P. Kucita
1
•Philippa A. S. Reed
1
Received: 25 September 2015 / Accepted: 6 December 2015 / Published online: 21 December 2015
ÓThe Author(s) 2015. This article is published with open access at Springerlink.com
Abstract An investigation has been conducted on AZ91
magnesium alloy processed in high-pressure torsion (HPT)
at 296, 423 and 473 K for different numbers of turns. The
microstructure has altered significantly after processing at
all processing temperatures. Extensive grain refinement has
been observed in the alloy processed at 296 K with
apparent grain sizes reduced down to 35 nm. Segmentation
of coarse grains by twinning has been observed in the alloy
processed at 423 K and 473 K with average apparent grain
sizes of 180 nm and 250 nm. Substantial homogeneity in
microhardness has been observed in the alloy processed at
296 K compared to that found at 423 K and 473 K. The
ultrafine-grained AZ91 alloy exhibited a significant
dependence of the yield strength on grain size as shown by
the microhardness measurements, and it obeys the expected
Hall–Petch relationship. The alloying elements, fraction of
nano-sized particles of b-phase, and the dominance of basal
slip and pyramidal modes have additional effects on the
strengthening of the alloy processed at 296 K.
Introduction
Magnesium alloys are promising alternatives to replace
denser materials, such as steel and aluminium alloys, with
the objective of meeting requirements to save fuel by
manufacturing light weight/high strength parts [1]. The
mechanisms of deformation in magnesium alloys at room
temperature are basal slip and twinning, which result in a
limitation in their workability at room temperature [2]. The
limited ductility and workability of these alloys can be
improved at higher temperatures by the activation of
additional slip systems [1]. Thermo-mechanical processing
is used to improve the workability of these alloys, although
such processing is associated with grain growth and a
greater consumption of energy [3]. Several processing
routes have been introduced to achieve optimization of the
microstructure, and these routes include dynamic recrys-
tallization under high-temperatures in ECAP processing
[4], HPT processing [5,6], ECAP processing at relatively
low temperatures assisted by a back-pressure [7], or
through the use of a higher channel angle of pressing die in
ECAP processing [8]. The majority of the earlier work in
SPD processing of magnesium alloys, especially for AZ91
alloy, has been conducted using ECAP at elevated tem-
peratures (C473 K) [2,4,9] with resultant grain refinement
being achieved in the micrometre range. The AZ91 alloy
(Mg–9wt%Al–1wt%Zn–0.3wt%Mn) is a common alloy in
&Ahmed S. J. Al-Zubaydi
asaz1e11@soton.ac.uk
Alexander P. Zhilyaev
alexz@anrb.ru
Shun C. Wang
wangs@soton.ac.uk
P. Kucita
pk4v07@soton.ac.uk
Philippa A. S. Reed
p.a.reed@soton.ac.uk
1
Materials Research Group, Faculty of Engineering and the
Environment, University of Southampton,
Southampton SO17 1BJ, UK
2
Branch of Materials Science, Department of Applied
Sciences, University of Technology, Baghdad, Iraq
3
Institute for Problems of Metals Superplasticity, Russian
Academy of Sciences, Khalturina 39, Ufa, Russia 450001
4
Research Laboratory for Mechanics of New Nanomaterials,
St. Petersburg State Polytechnical University, St. Petersburg,
Russia 195251
123
J Mater Sci (2016) 51:3380–3389
DOI 10.1007/s10853-015-9652-2
Content courtesy of Springer Nature, terms of use apply. Rights reserved.
the Mg–Al–Zn family. This alloy has a good strength-to-
density ratio, good corrosion resistance and ease of pro-
duction and machining [3]. To date, only one investigation
has been conducted on Mg–9wt%Al alloy [6] using HPT at
room temperature. The development of microstructure and
microhardness across horizontal and vertical cross-sections
of AZ91 samples processed by HPT has not been reported
to date. This research describes the microstructural
homogeneity and development of microhardness in AZ91
alloy after processing by HPT at different processing
temperatures. The dislocation density, distribution of b-
phase and Hall–Petch relationship have also been
investigated.
Experimental materials and procedures
AZ91 alloy (Mg–9 %Al–1 %Zn) in the form of an extru-
ded rod was used in this work, the alloy was supplied by
Magnesium Elektron Co. (Manchester, UK). Thin discs
were made of the extruded rod with thicknesses of 1.5 mm
and final thicknesses of 0.85 mm. The HPT processing was
conducted at 296, 423 and 473 K using a HPT facility that
has been previously discussed in detail elsewhere [10]. The
HPT processing was conducted under a quasi-constrained
condition at a speed of 1 rpm using an applied pressure of
3.0 GPa for differing numbers of turns: N=1/2, 1, 5 and
10 turns. The as-received and processed microstructures
were observed using optical microscopy (OM, OLYMPUS-
BX51, Japan) and scanning electron microscopy (SEM,
JEOL JSM-6500F, Japan). Subsequently, a transmission
electron microscope (TEM, JEOL JEM-3010) was used for
microstructural observation of the alloy after HPT pro-
cessing. The chemical compositions of the as-received and
processed alloy were analysed using energy-dispersive
spectroscopy (EDS). The area fraction and average size of
the b-phase particles in the as-received alloy and processed
alloy were determined by ImageJ software using a point
count technique [11]. X-ray diffraction was used to deter-
mine the crystallite size and dislocation density in the
processed alloy using an XRD facility (D2 Phaser, Ger-
many). The diffraction data were analysed using Rietveld
refinement based software program (MAUD). Microstruc-
tural observations and microhardness testing were con-
ducted over the horizontal and vertical cross-sections that
are illustrated schematically in Fig. 1a, b. The microhard-
ness measurements of the processed disc were conducted
using a Vickers microhardness tester (FM-300, Japan) and
using an applied load of 100 gf and a dwell time of 15 s.
The microhardness data were recorded at separation dis-
tances of 0.3 and 0.1 mm throughout the entire horizontal
and vertical cross-sections, as reported earlier [5,12].
Experimental results
The microstructure of the AZ91 magnesium alloy prior to
and after HPT processing is shown in Fig. 2. The as-re-
ceived AZ91 alloy has an average grain size of 30 lm and
an average value of Vickers microhardness of 70 ±5. The
initial and processed microstructures consist of two main
phases: a-Mg matrix, b-phase and Al
8
Mn
5
particles as
shown in Fig. 2a, b. The chemical analysis obtained by
EDS of alloying elements in the alloy processed at 296 K
for N=5 turns is shown in Fig. 3. The alloy constituents
were identical before and after HPT as shown earlier [13].
The processed microstructure at 296 K showed extensive
grain refinement, and the original decoration of the grain
boundaries by b-phase disappeared with increasing number
of turns as shown in Fig. 2b, c. The b-phase fragmented
into nano-sized particles as observed in Fig. 2b–d and
appears aligned along the direction of torsional straining. A
strong degree of grain refinement after processing at 296 K
was observed with an apparent grain size down to 500 and
50 nm observed after N=1/2 and 1 turn, respectively, as
shown in Fig. 2e, f. A reduction in the crystallite size from
60 to 35 nm was found with increasing number of turns up
to N=10 turns. The processed microstructures at 296 K
across the vertical cross-sections are shown in Fig. 4. The
microstructure seems slightly deformed with the presence
of twinning as shown in Fig. 4a. Shear bands decorated by
the b-phase were observed aligned parallel to the radial
direction across the vertical cross-section as observed in
Fig. 4b. Recorded peaks by XRD as shown in Fig. 5are
Fig. 1 An illustration of a HPT disc shows athe horizontal cross-
section, and bThe vertical cross-section. These cross-sections were
used in the microstructural and microhardness observations. The
arrow from the centre to the edge refers to the longitudinal (radial)
direction, whereas the arrow from the upper surface to lower surface
refers to the through-thickness (vertical) direction
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prismatic planes 10
10ðÞ,11
20ðÞ,ð20
20Þ, basal plane
ð0002Þand pyramidal planes ð10
11Þ,ð10
12Þ,ð10
13Þ. The
microstructures of the alloy processed at 423 and 473 K are
shown in Fig. 6. The samples showed twinning, and the
distribution of twinning increased and spread gradually
with increasing number of turns. The microstructures were
effectively refined by the segmentation of the coarse grains
by twinning as observed in Fig. 6a, b. However, grain
growth has been observed at 473 K with increasing number
of turns up to N=5 turns as shown in Fig. 6b. The
apparent area fraction has increased (which may reflect a
sampling effect once the second phase is more homoge-
neously distributed), and the average size of the b-phase
particles has been refined down to 200 nm in the pro-
cessed alloy compared to the as-received alloy as shown
in Fig. 7. A gradual development in the microhardness
over the horizontal and vertical cross-sections has been
achieved with increasing number of turns up to N=10
turns as shown in Figs. 8and 9. The distributions of
microhardness were relatively lower for the alloy pro-
cessed at 423 and 473 K than at 296 K. A significant
increase in the microhardness has been observed as shown
Fig. 2 Microstructural observations using SEM for athe as-received
alloy, bthe alloy processed for N=1 turn (296 K), cthe alloy
processed for N=10 turns (296 K) and dthe nano-sized particles of
b-phase in the alloy processed for N=10 turns (296 K), and TEM
observation of the alloy processed for eN=1/2 turn (296 K) and
fN=1 turn (296 K). The corresponding numbers (1,2,3) in the
micrograph arepresent the lamellar, agglomerate forms of the b-
phase (Mg
17
Al
12
) and Al
8
Mn
5
particle, respectively
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in Fig. 10, with increasing equivalent strain imposed
during HPT for the alloy processed at 296 K. A signifi-
cant dependency of the microhardness on the crystallite
size of the AZ91 alloy processed at 296 K is shown in
Fig. 11. The lower processing temperature leads to finer
crystallite size, higher microhardness and dislocation
density, and at elevated temperatures, these outcomes
decreased significantly as the number of turns increased
as shown in Fig. 12.
Discussion
Feasibility of HPT processing of AZ91 magnesium
alloy
The TEM and XRD revealed the occurrence of extensive
grain refinement in the AZ91 alloy due to the imposition of a
very high plastic strain by HPT at 296 K. However, for the
sample processed for N=1/2 turn, it is noteworthy that the
Fig. 3 The chemical analysis with weight fractions of the alloy processed at 296 K for N=5 turns showing aa-Mg matrix, bb-phase
(Mg
17
Al
12
), cAl
8
Mn
5
particle
Fig. 4 The microstructures of the alloy processed at 296 K as observed along the vertical cross-sections for aN=1 turn and bN=5 turns. The
black and white arrows refer to the twinning and shear bands decorated by the b-phase, respectively
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value of the crystallite size obtained by XRD was signifi-
cantly lower than the apparent grain size measured by TEM.
This difference between the measurements via XRD and
TEM is expected in SPD-processed materials, because the
grains in these materials are made of subgrains and/or dis-
location cells. Thus, coherent scattering of the X-ray from
these substructures represents the (smaller) mean crystallite
size rather than grains which can be more easily observed in
TEM [14]. The feasibility of HPT processing at 296 K for
the AZ91 magnesium alloy can be attributed to the presence
of hydrostatic pressure, which prevents propagation of
fracture during processing [6–8]. Furthermore, the geometry
of the processing zone constrains the alloy within a specific
volume as illustrated earlier and thus activation of twinning
[8,15,16]. The XRD observations indicate the orientation of
the processed microstructure towards twinning and basal
deformation modes under HPT conditions that facilitate
processing at room temperature [16]. The unidirectional
nature of straining during HPT processing may have con-
tributed to re-orientation of the microstructure towards easy
slip [17]. The twinning activity has persisted in the pro-
cessed alloy at 296 K with increasing number of turns,
which confirms its accommodation for the higher imposed
strain produced by HPT [18].
Fig. 5 XRD diffraction patterns for athe as-received alloy and bthe alloy processed at 296 K for N=10 turns
Fig. 6 The microstructures of the alloy as observed across the horizontal cross-sections after HPT processing at a423 K (N=5 turns) and
b473 K (N=5 turns)
Fig. 7 The area fraction and average size of the b-phase particles in
the as-received alloy and processed alloy at 296 K for different
number of turns
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Fig. 8 The colour-coded maps of the microhardness over the
horizontal cross-sections of the AZ91 discs processed forf aN=1/
2 turn (296 K), bN=1 turn (296 K), cN=5 turns (296 K),
dN=10 turns (296 K), eN=10 turns (423 K) and fN=10 turns
(473 K). The small inset in the figure shows the scale of the
microhardness with regard to each colour (Color figure online)
Fig. 9 The colour-coded maps of the microhardness distributions
over the vertical cross-sections of the AZ91 discs processed for
N=10 turns at 296 K (upper), 423 K (centre) and 473 K (lower).
The small inset in the figure shows the scale of the microhardness
with regard to each colour (Color figure online)
Fig. 10 Correlation of the measured microhardness with the equiv-
alent strain imposed by HPT processing for the alloy processed at
296 K for different number of turns
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Grain refinement in AZ91 alloy
The relatively high content of aluminium in the AZ91
magnesium alloy leads to a significant reduction in the
stacking fault energy through the solute–dislocation inter-
action and results in smaller grain sizes under SPD pro-
cessing [19]. The effect of dynamic recovery was absent as
the alloy has been processed at room temperature. It is
anticipated that the homogeneity developed gradually with
further straining at room temperature as mentioned by
several investigators [20–23]. The grain refinement in the
processed alloy at 423 K has developed efficiently by
twinning intersections and the grain subdivision mecha-
nism. At this temperature, dynamic recrystallization was
absent or had a minor effect on the refinement process
compared to the twinning activity. It is likely that dynamic
recrystallization may have contributed to grain refinement
in the processed alloy at 473 K. However, the formation
and fragmentation of twinning appears to be the dominant
mechanism for refinement at 473 K. The HPT-processed
alloy at 423 and 473 K has significantly refined apparent
grain sizes of 180 and 250 nm, respectively, which are
finer than in the previously reported ECAP [20,24–26],
FSP [27] and ARB-processed alloys [28]. In the afore-
mentioned SPD techniques, grain refinement occurs mainly
by dynamic recrystallization with resultant microstructures
of micrometre size grains. The severe levels of deformation
in the alloy and the deformation incompatibility between a-
Mg matrix and b-phase have resulted in fragmentation of
the b-phase [29]. The significant dispersion of nano-sized
particles of the b-phase during processing had a pinning
effect on grain growth at a higher number of turns and
elevated temperatures [23]. The alloy processed at 296 K
showed microstructural homogeneity at the initial stage of
HPT processing rather than the heterogeneity observed in
the alloy processed at 423 and 473 K, which required
further processing turns and/or higher processing temper-
ature to achieve a reasonable homogeneity [10]. The tem-
perature rise expected during HPT processing at room
temperature does not exceed 293 K for samples processed
Fig. 11 The Hall–Petch relationship for the ultrafine-grained AZ91
alloy in the current work and for AZ31 and AZ61 alloys processed by
HPT and ECAP
Fig. 12 The overall variation in the average acrystallite size,
bdislocation density and cmicrohardness for the AZ91 alloy after
HPT at different processing temperatures
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at 296 K for N=10 turns. This value of temperature rise
has been calculated using the equation stated in [30], and is
similar to the experimental value (290 K) measured
directly from the thermocouple located in the upper anvil.
The low value of temperature rise can be attributed to (1)
the heat loss from relatively small samples in contact with
the much larger HPT anvils and (2) due to the low strain
rates of deformation in HPT processing [31]. A further
factor, making the heat generated low, is the lower friction
expected between the relatively lower strength magnesium
alloy and the high strength (high speed tool steel) anvils
[15,31]. As a result, any temperature rise due to processing
is considered negligible and unlikely to produce any
occurrence of recrystallization or grain growth during
processing at room temperature [32].
Development of microhardness
The initial heterogeneity of microstructures leads to an
initial heterogeneity in the distribution of microhardness
[10,33,34]. The difference in grain sizes at the centre and
edge regions was diminished by further straining, where a
gradual evolution towards homogeneity was found in the
observed microstructure and microhardness at both centre
and edge regions at a higher number of turns [10,35]. The
existence of misalignment between the anvils at a high
number of turns causes an additional deformation at the
centre region of the processed disc, which appears as an
increase in the measured microhardness [36,37]. The
development of microhardness after HPT processing
depends on the stacking fault energy of the alloy [10,19].
The AZ91 alloy with a low stacking fault energy [38]
shows a slow rate of dynamic recovery during processing
at room temperature, thus strain hardening occurs at a fast
rate during processing [20,32]. The AZ91 alloy processed
in HPT showed an earlier saturation in the microhardness
distribution than for the AZ31 alloy [21] processed in HPT
at room temperature. The stacking fault energy is lower,
and the fraction of particles of b-phase is higher in the
AZ91 alloy than for the AZ31 alloy [38,39]. Therefore, the
evolution of grain refinement and strain hardening occurred
at faster rates in the AZ91 alloy than for the AZ31 alloy.
The overall microhardness values for the alloy processed at
473 K were significantly lower than for their counterparts
processed at 296 and 423 K, due to the variation in dislo-
cation density with processing temperature [14]. However,
the level and homogeneity of strengthening are still higher
when processing by HPT at elevated temperatures than
observed in ECAP [26], and FSP [27], where strengthening
has been lowered by dynamic recrystallization, over-ageing
and precipitate coarsening [26,27]. The microhardness
distributions in the AZ91 alloy are heterogeneous along the
through-thickness directions in the initial stage of defor-
mation. This is supported by the differences in
microstructural observations along this direction. A suffi-
cient high number of turns may reduce heterogeneity, by
filling the alloy in-between the anvils and achieving a
significant sticking condition which then increases the
deformation and microstructural homogeneity [5,12,15,
40,41]. The distribution of microhardness along the ver-
tical and horizontal cross-sections showed considerable
consistency for the current alloy processed at each specific
processing temperature. This indicates the development of
microstructural homogeneity with increasing imposed
strain at each condition [8]. This consistency in the AZ91
alloy has not been observed in the AZ31 alloy or AZ91
alloy processed by ECAP [26] and FSP [27]. This is
attributed to the difference in the aluminium content and
stacking fault energy in both alloys, which control the
extent of grain refinement, dislocation density, achieved
homogeneity and resultant mechanical properties [14,21,
38]. The behaviour of strain hardening and homogeneity of
microhardness in the AZ91 alloy follows a standard model
of hardness evolution with increasing equivalent strain
reported in earlier work [20].
The effect of the equivalent strain on the Hall–Petch
relation and dislocation density
The increase in the equivalent strain resulted in an evolu-
tion in microstructure and a gradual development in the
microhardness [10]. The strength of the alloy in terms of its
microhardness improved significantly with grain refine-
ment at room temperature. This proportionality has been
expressed by the Hall–Petch relationship for hardness
measurements: Hv ¼H0þkHd1=2[42]. The effect of
grain refinement on the strength of the ultrafine-grained
alloy AZ91 alloy showed a significant consistency with this
Hall–Petch relationship. The material constants are
H0=76 MPa and kH=233 MPa lm
1/2
, which is rela-
tively higher than those found for AZ31 and AZ61 alloys
(H0=647–697 MPa and kH=118–170 MPa lm
1/2
)[7,
22,43]. Thus, the ultrafine-grained AZ91 alloy shows a
relatively higher level of hardness than for AZ31 and AZ61
alloys processed by HPT and ECAP processing at room
temperature and elevated temperatures [21,22,44]. The
difference in kHcan be attributed to the difference in
alloying constituents in the mentioned alloys, where the
high content of alloying element in the AZ91 alloy resulted
in a lowering of its stacking fault and thus a finer
microstructure and a higher dislocation density in the AZ91
alloy after processing than in the AZ61 and AZ31 alloy
[19,21,22]. The evolution in dislocation density with
increase of imposed strain in HPT has a major effect on the
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achieved strengthening in the AZ91 alloy. The evolution of
dislocation density is affected by the fraction of nano-sized
particles of b-phase, value of applied pressure in HPT, and
value of stacking fault energy. The widely distributed b-
phase fine particles are reported as acting as barriers for
mobile dislocations during deformation [39]. The high
value of applied pressure has also been reported to enhance
the obstruction of defect migration in the processed mate-
rial and then promotes the suppression of dislocation
annihilation [45,46]. The low stacking fault energy in the
AZ91 alloy leads to a significant inhibition of dislocation
cross-slip, and formation of a high density of planar arrays
of dislocations has also been reported [10,38].
Conclusions
1. AZ91 magnesium alloy has been effectively processed
in HPT processing at room temperature with an
ultrafine-grained microstructure down to 35 nm. The
alloy processed at 423 and 473 K has been signifi-
cantly refined by twinning segmentation of the original
grains into fine grains with average apparent grain
sizes of 180 and 250 nm, respectively.
2. Fragmentation and alignment of the b-phase in the
direction of torsional strain have been observed during
processing. This phase has been refined down to
nanometre sizes with a higher fraction as the number
of turns increased, indicating the very high level of
plastic deformation that is imparted to the alloy during
HPT.
3. Existence of twins at all processing temperatures and
their distribution was proportional to processing tem-
perature and the number of turns. The occurrence of
twinning has been induced by the need for re-orientation
of the microstructure towards the slip direction and to
accommodate severe plastic deformation.
4. Lower processing temperature has resulted in homoge-
nous microstructure and significant development of
strength. Higher processing temperatures have resulted
in heterogeneous microstructures especially in the
initial stages of HPT and this heterogeneity decreased
gradually at higher numbers of turns.
5. A considerable dislocation density has developed with
increasing the number of turns at lower processing
temperature rather than at higher processing temper-
atures. The values of dislocation density after HPT
were higher than earlier reported data for the same
alloy.
6. The ultrafine-grained AZ91 alloy follows the Hall–
Petch relationship, and this emphasizes the significant
dependence of strength on grain size. The higher
alloying content, fraction of nano-sized particles of
b-phase and the dominance of basal slip and pyramidal
modes after processing also have a significant effect on
the strengthening of the alloy processed at 296 K.
Acknowledgements One of the authors (Ahmed S. J. Al-Zubaydi)
is grateful to The Higher Committee for Education Development
(HCED) of the Government of Iraq for the provision of Ph.D.
scholarship.
Funding This work was supported in part by the Russian Science
Foundation under Grant No. 14-29-00199 (APZ).
Compliance with ethical standards
Conflict of Interest The authors declare that they have no conflict
of interest.
Open Access This article is distributed under the terms of the Crea-
tive Commons Attribution 4.0 International License (http://creative
commons.org/licenses/by/4.0/), which permits unrestricted use,
distribution, and reproduction in any medium, provided you give
appropriate credit to the original author(s) and the source, provide a link
to the Creative Commons license, and indicate if changes were made.
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