Corrosion Science 52 (2010) 1035–1041
Contents lists available at ScienceDirect
Corrosion Science
journal homepage: www.elsevier.com/locate/corsci
Effect of heat treatment on corrosion behaviour of magnesium alloy AZ91D
in simulated body fluid
Wei Zhou *, Tian Shen, Naing Naing Aung
School of Mechanical and Aerospace Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore
a r t i c l e
i n f o
Article history:
Received 12 June 2008
Accepted 18 November 2009
Available online 27 November 2009
Keywords:
A. Alloy
A. Magnesium
B. Weight loss
B. Polarisation
C. Intergranular corrosion
a b s t r a c t
The research explored ways of improving corrosion behaviour of AZ91D magnesium alloy through heat
treatment for degradable biocompatible implant application. Corrosion resistance of heat-treated samples is studied in simulated body fluid at 37 °C using immersion and electrochemical testing. Heat treatment significantly affected microgalvanic corrosion behaviour between cathodic b-Mg17Al12 phase and
anodic a-Mg matrix. In T4 microstructure, dissolution of the b-Mg17Al12 phase decreased the cathodeto-anode area ratio, leading to accelerated corrosion of a-Mg matrix. Fine b-Mg17Al12 precipitates in T6
microstructure facilitated intergranular corrosion and pitting, but the rate of corrosion was less than
those of as-cast and T4 microstructures.
Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction
Degradable biocompatible materials play an essential role in
load-bearing implants due to the risk of local inflammation of traditional implants such as stainless steels, titanium alloys, and cobalt–chromium-based alloys. Unfortunately, currently used
absorbable polymer materials have unsatisfactory mechanical
properties. In comparison, magnesium alloys are attractive candidates because they are ultralight alloys possessing mechanical
properties similar to those of natural bones and a natural ionic
presence with significant functional roles in biological systems
[1–4]. However, fast degradation via corrosion in the electrolytic
environment of the body constitutes the main shortcoming of
magnesium alloy implants [5–10]. Their degradation rates vary
over a range of 3 orders of magnitude [11]. Therefore, improving
corrosion resistance is an important issue for the application of
magnesium alloys as biodegradable load-bearing implants.
The corrosion resistance of magnesium alloys, particularly
AZ91D alloy, can be improved by different surface coating techniques such as electro plating [12], electroless plating [13], anodizing [14], laser surface cladding [15] and laser surface melting [16].
It is also possible to control the degradation process of magnesium
alloy through Zn and Mn alloying, purification and anodization
[17,18]. A Zn-containing magnesium alloy with a small amount
of manganese can be a potential biodegradable alloy [17]. If the alloy starts biodegradation too early, an anodized coating can be ap-
plied to delay it. If the alloy biodegrades too fast, it can be purified
to bring down the rate [17]. However, some surface coating techniques could have a negative impact on the implants due to their
possible release of toxic metallic ions.
In AZ91D alloy, the a-Mg matrix corrodes due to its very negative free corrosion potential and there is the tendency for the corrosion rate of the a-Mg phase to be accelerated by microgalvanic
coupling between anodic a-Mg phase and cathodic b-Mg17Al12
phase [19–25]. However, the b-Mg17Al12 phase may act as a barrier
against corrosion propagation if it is in the form of a continuous
network [20,22]. The distribution, configuration and size of the bMg17Al12-phase can be changed, which may result in different corrosion rates.
The corrosion resistance of the alloy may be enhanced by heat
treatment [22–25]. In this case, heat treatment does not normally
increase the corrosion barrier effect of the b-Mg17Al12 phase. Solution heat treatment dissolves the b-Mg17Al12 phase and removes
its barrier effect. However, the b-Mg17Al12 precipitates produced
by ageing were effective to reduce microgalvanic corrosion of the
adjacent a-Mg phase [22]. The objective of the research was to
study how heat treatment could be explored to improve corrosion
resistance of AZ91D alloy in simulated body fluid.
2. Materials and methods
2.1. Materials
* Corresponding author. Tel.: +65 6790 4700; fax: +65 6791 1859.
E-mail addresses: [email protected], [email protected] (W. Zhou).
0010-938X/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved.
doi:10.1016/j.corsci.2009.11.030
The material studied was an as-cast ingot of AZ91D alloy with the
following chemical composition (in wt.%): Al–9.1, Mn–0.17,
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W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
Zn–0.64, Si–60.01, Fe-60.001, Cu–60.01, Ni-0.001, other impurities
60.02, Mg-the rest. Solution treatment (T4) of the alloy was carried
out at 445 °C for 24 h in argon atmosphere followed by water quench
at 25 °C. Ageing treatment (T6) of the solution-treated samples was
performed at 200 °C for 8 h, 16 h and 24 h to produce three different
aged microstructures. The as-cast and heat-treated samples were
cut into samples of 10 mm 10 mm 2 mm in dimension for
immersion test and potentiodynamic polarisation study.
2.2. SBF preparation
The bioactivity of bone implant materials is usually tested in vitro
using simulated body fluid (SBF). The composition of common SBF
differs from that of blood plasma in that it has a higher Cl and a lower HCO3 concentration, which could affect the composition of
in vitro formed bone-like apatite. In this research, the high stability
SBF solution was prepared by mixing stable concentrated solutions,
which increase the reproducibility of in vitro tests due to negligible
changes of pH during preparation.
The SBF solution was prepared by pipetting calculated amounts
of concentrated solutions of KCl (59.64 g/l), NaCl (116.88 g/l), NaHCO3 (45.37 g/l), MgSO47H2O (49.30 g/l), TRIS (tris-hydroxymethyl
aminomethane (121.16 g/l), NaN3 (100 g/l) and KH2PO4 (27.22 g/l)
into double distilled water to prevent precipitation of homogeneously nucleated calcium phosphates or other phases and to minimize changes in pH during preparation [26]. The SBF had the
following composition (in ml/l): KCl-5, NaCl-50, NaHCO3-50,
MgSO47H2O-5, CaCl2-25, TRIS (tris-hydroxymethyl aminomethane) + HCl-50, NaN3-10 and KH2PO4-5. Concentration (in mmol/l)
of various ions in the SBF was: 142 Na+, 5 K+, 2.5 Ca2+, 1 Mg2+, 109
Cl, K+, 27 HCO3, 1 SO42, 1 HPO42. The pH of human blood plasma
ranges from 7.3 to 7.4 at 37 °C [26], so the pH of SBF was adjusted to
7.6–7.7 at 25 °C which is equal to a pH of 7.3–7.4 at 37 °C, by adding
concentrated HCl.
2.3. Immersion test
ðW b W a þ BÞ 1000
At
Corrosion rateðmm y1 Þ ¼
1
R
0:274 q
ð2Þ
q = density of alloy.
2.4. Electrochemical test
The electrochemical corrosion behaviour of the as-cast and
heat-treated samples was studied using a potentiostat/galvanostat
corrosion measurement system (EG&G model 263A). The samples
were encapsulated into epoxy resin so that only a surface with
the dimension of 10 mm 10 mm was exposed to the 300 ml of
the solution. Prior to the experiment the samples were ground
up to 4000 grit emery paper, followed by washing with distilled
water and acetone. A potentiodynamic polarisation test was carried out in a standard electrochemical cell using three electrodes
configuration. The sample was the working electrode. A saturated
calomel electrode (SCE) was used as a reference electrode and a
platinum electrode was used as a counter electrode. The test cell
was placed into the water bath having environmental temperature
to 37 °C.
Potentiodynamic polarisation curves were measured immediately after immersion of the sample in the solution. The cathodic
polarisation scan was started from 2000 mV to the steady state
corrosion potential and then polarizing in an anodic direction at
a scan rate of 1 mV s1. The icorr in mA cm2 can be related to the
icorr in mm y1 using Eq. (3).
Corrosion rateðmm y1 Þ ¼
3:28M
icorr
nq
ð3Þ
n = number of electrons freed by the corrosion reaction
M = atomic mass.
2.5. Analysis of corroded surface
The corrosion rate was determined by weight loss rate from the
immersion test according to ASTM-G31-72 [27]. The as-cast and
different heat-treated samples were immersed in 300 ml solution
of SBF for various periods of 8 h, 24 h, 72 h and 168 h. Separate
sample was used to determine the weight loss for each exposure.
The temperature of the solution was kept at 37 °C by water bath.
Before each test, the sample was ground on progressively finer
grades of emery papers up to 4000 grit and then weighed. All
immersion tests were carried out without agitation or circulation
and without disturbing corrosion system.
In order to remove the corrosion products with minimal dissolution of base alloy, chemical cleaning of the corroded sample was
carried out in boiling 15% CrO3 + 1% AgCrO4 solution for 1 min. An
acetone washing followed this. The weight loss was measured after
each experiment and the corrosion rate (R) was calculated in mg
cm2 day1 using Eq. (1). The values of corrosion rate were converted from mg cm2 day1 to mm y1 using Eq. (2). The tests were
repeated three times to obtain the reproducible results.
R¼
t = exposure time, day.
ð1Þ
where:
R = corrosion rate, mg cm2day1
Wb = weight of test sample before test, g
Wa = weight of test sample after test, g
B = weight loss of blank, g (the average weight loss from three
unused and clean sample was used as the blank correction)
A = surface area of sample, cm2
Microstructure and morphology of the corroded surfaces were
analysed using Carl Zeiss Axioskop 2 optical microscope and scanning electron microscopy (SEM, Jeol model 5600 LV) coupled with
energy dispersed spectroscopy (EDS) system.
3. Results and discussion
3.1. Microstructures
The as-cast AZ91D microstructure has typically a primary a-Mg
matrix and a divorced eutectic phase distributed along the grain
boundaries (Fig. 1(a)). The close-up view clearly shows that the eutectic consists of large b-Mg17Al12 phase particles and the eutectic
a-Mg phase (Fig. 1(b)). The eutectic a-Mg phase is supersaturated
with Al and it can transform to form a fine lamellar structure.
Microstructure of the alloy was changed during the process of
heat treatment. The T4 heat treatment dissolved the b-Mg17Al12
phase and produced a microstructure consisting the supersaturated
a-Mg phase. However, some residual small b-Mg17Al12 phases were
still observed (Fig. 1(c)). During T6 treatment, b-Mg17Al12 phase precipitated along the grain boundary and within grains of the supersaturated a-Mg phase. Ageing made Al atoms diffuse towards grain
boundaries to form precipitates of b-Mg17Al12 phase, and this process reduced the aluminium concentration in the a-Mg matrix. In
this case, the degree of homogeneity of the b-Mg17Al12 precipitates
distribution and the aluminium content in the a-Mg matrix was different between the T6 microstructures (Fig. 1(d–f)). The aluminium
content in the a-Mg matrix of T6-16 h microstructure was found to
W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
1037
Fig. 1. Optical micrographs showing the microstructures of AZ91D alloy before and after heat treatment: (a) as-cast, (b) closed-up view of (a), (c) T4, (d) T6-8 h, (e) T6-16 h
and (f) T6-24 h.
decrease from 9% to 3%. This result confirmed with other researcher’s
result [28].
3.2. Effect of heat treatment on corrosion rates
Variation of corrosion rate with immersion time for the ascast and heat-treated samples is presented in Fig. 2. For short
immersion period, 8 h, the corrosion rate of T4 sample was
the lowest compared with other samples (Fig. 2(a)). However,
its corrosion rate shifted to the highest value after long exposure, 168 h (Fig. 2(b)). The corrosion rates of T6 samples were
higher than those of the as-cast and T4 microstructure at the
initial exposure of 8 h (Fig. 2(a)), but the rates slowed down
to the lower values after 168 h (Fig. 2(b)). Among T6 samples,
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W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
1.0
2.0
-1
Corrosion rate (mg cm day )
8 h immersion
0.8
0.2
0.4
T4
T6-8 h
T6-16 h
T6-24 h
-1
10
20
100 µm
15
5
10
100 µm
5
-1
As-cast
25
-2
0.4
30
Corrosion rate (mm y )
1.2
-1
0.6
100 µm
15
Corrosion rate (mg cm day )
-2
1.6
Corrosion rate (mm y )
0.8
(a)
0
8
15
6
10
4
T6-8 h
T6-16 h
T6-24 h
(b)
the corrosion rate of T6-16 h sample was the lowest (Fig. 2(a)
and (b)).
The extent of corrosion for T4 and T6 samples in SBF as a function of exposure time are shown in Figs. 3 and 4. The corrosion rate
and the corresponding area corroding of the samples increased
with time of exposure to the corroding solution. Corrosion started
around localised sites at grain boundaries in the T4 sample and the
attack invaded the entire surface with continued exposure. After
168 h immersion, T4 sample was highly damaged, as shown in
Fig. 3. For T6-16 h sample, corrosion accelerated with increasing
exposure time but the extent of corrosion was lower in T6-16 h
sample (Fig. 4) compared with T4 sample (Fig. 3). The results suggest that the heat treatment which gave the best resistance to corrosion in SBF was T6 rather than T4.
3.3. Effect of microstructure on corrosion
In order to understand the effect of heat treatment on corrosion
mechanism in greater detail, the corroded surfaces of samples for
8 h and 72 h exposure were carefully analysed under SEM. Selected
SEM micrographs are shown in Figs. 5 and 6.
Corrosion of the as-cast microstructure initiated at the primary
a-Mg matrix in the eutectic region as indicated by arrows in
Fig. 5(a). In the case of T4 microstructure, localised attack was observed around residual b-Mg17Al12 phase at grain boundaries
(Fig. 5(b)). For the T6 microstructures, corrosion occurred preferen-
1.8
1.6
3.5
3.0
100 µm
1.4
2.5
1.2
100 µm
1.0
2.0
0.8
-1
Fig. 2. Corrosion rates for as-cast and heat-treated samples in SBF: (a) after 8 h and
(b) after 168 h.
4.0
100 µm
Corrosion rate (mm y )
5
T4
150
2.0
-1
20
-2
-1
-2
10
As-cast
100
Fig. 3. Extent of corrosion for T4 samples in SBF as a function of exposure time.
-1
Corrosion rate (mg cm day )
25
Corrosion rate (mm y )
12
2
50
Immersion time (h)
Corrosion rate (mg cm day )
168 h immersion
14
0
0
30
1.5
0.6
1.0
0
50
100
150
Immersion time (h)
Fig. 4. Extent of corrosion for T6-16 h samples in SBF as a function of exposure
time.
tially along the grain boundary and some pits were found within
grains as indicated by arrows in Fig. 5(c–e).
Corroded morphologies of the samples after longer exposure
clearly showed that the changes in distribution, configuration
and size of the b-Mg17Al12-phase due to heat treatment resulted
in different corrosion behaviours (Fig. 6). In the as-cast microstructure, the b-Mg17Al12 phase is highly cathodic to the a-Mg phase
and can thus act as an effective cathode to cause microgalvanic
corrosion [19–25]. The b-Mg17Al12 phase contains much more aluminium compared with the a-Mg phase. According to our previous
study [20], the variation of the concentration of aluminium is in
the range of about 35% in the b-Mg17Al12 phase to about 6% in
the primary a-Mg phase. The region with less than 8% aluminium
could be corroded preferentially [20]. Therefore, the lower aluminium content of the primary a-Mg matrix was the initiation site of
corrosion (Fig. 6(a)).
In the T4 microstructure, there was a meta-stable, partially protective film on high aluminium content of the a-Mg matrix. This
film prevented corrosion and the result therefore showed the
W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
1039
Fig. 5. SEM micrographs showing the corroded morphologies of the samples in SBF after 8 h exposure: (a) as-cast, (b) T4, (c) T6-8 h, (d) T6-16 h and (e) T6-24 h .
lowest corrosion rate in the initial exposure (Figs. 2(a) and 3).
However, the dissolution rate of localised corrosion could be faster
once the protective film was broken down due to weak localised
sites of residual b-Mg17Al12 phase (Fig. 6(b)). Besides residual bMg17Al12 phase, Al–Mn and Al–Mn–Fe intermetallics could also
greatly accelerate the corrosion [29]. As a result, the corrosion rate
significantly increased after longer exposure (Figs. 2(b), 3 and
6(b)).
For T6 microstructure, the corrosion mechanism appeared to be
influenced by the precipitation of b-Mg17Al12 in relation to the
amount of aluminium content in the a-Mg matrix. The corrosion
pits initiated at the anodic a-Mg matrix adjacent to the cathodic
b-Mg17Al12 precipitates. The microgalvanic action due to the bMg17Al12 precipitates at grain boundary led to intergranular corrosion. The T6 microstructural features gave the tendency to occur
intergranular corrosion as well as pitting corrosion (Fig. 6(c)).
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W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
Fig. 6. SEM micrographs showing different corrosion mechanisms on the corroded surfaces: (a) as-cast, (b) T4 and (c) T6-16 h.
3.4. Electrochemical behaviour of heat-treated microstructures
The electrochemical corrosion behaviour of the samples in SBF
is shown in Fig. 7. The values of corrosion potential (Ecorr), Tafel
slope (bc) and the corrosion current density (icorr) for each polarisation curve are summarized in Table 1. The Ecorr values of heat-treated samples were shifted to less negative values showing more
cathodic behaviour compared with the as-cast one. It should be
especially noted that the T6-16 h sample showed less negative value of Ecorr. The bc values were similar for all samples, indicating
that the same electrochemical reactions occurred. The cathodic
currents from the polarisation curves were much higher for all
heat-treated samples at all potentials. On the other hand, the T616 h microstructure was the least active anodically with the lowest
icorr value (Table 1).
The polarisation curves were used to explore the relationship
between electrochemical measurements of the corrosion rate,
0.1
T6-8 h
-2
I (Acm )
However, their extent of corrosion was lower compared with the
as-cast and T4 microstructure (Figs. 2(b), 4 and 6(c)). The isolated
fine b-Mg17Al12 precipitates do not lead to an obvious loss of corrosion because they act as a small cathode connected to the large
anode a-Mg matrix [22]. In this case, the corrosion resistance of
the T6-16 h microstructure was slightly better than the T6-8 h
and T6-24 h microstructure. The reason could be due to its more
homogeneous distribution of the b-Mg17Al12 precipitates as well
as the aluminium content in the a-Mg matrix.
0.01
T6-24 h
T6-16 h
As-cast
T4
-1.8
-1.6
-1.4
-1.2
-1.0
E (V/SCE)
Fig. 7. Potentiodynamic polarisation curves for as-cast and heat-treated samples in
SBF at 37 °C.
based on the corrosion current at the free corrosion potential (Table 1), and direct measurements using weight loss (Fig. 2). There
was a good correlation between the corrosion rates determined
from icorr and those from weight loss data only for short immersion
time. The corrosion rate from the Tafel extrapolation may relate to
the onset of corrosion, whereas the corrosion rate from the weight
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W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041
Table 1
Ecorr, bc and icorr values for as-cast and heat-treated samples in SBF.
Sample
Ecorr (V/SCE)
bc (V dec1)
icorr (mA cm2)
Corrosion rate (mm y1)
As-cast
T4
T6-8 h
T6-16 h
T6-24 h
1.31
1.24
1.29
1.19
1.29
0.186
0.176
0.189
0.195
0.192
0.039
0.028
0.066
0.027
0.074
0.85
0.61
1.44
0.59
1.62
loss/hydrogen evolution measurement relates to corrosion averaged over a considerable time period and includes corrosion some
considerable time after corrosion onset, when the corrosion is well
established [30].
4. Conclusions
(1) Heat treatment significantly changed the corrosion resistance of AZ91D alloy in SBF. Compared with the as-cast
condition, the T6 treatment reduced the corrosion rate by
30–60%, and the T4 treatment increased the rate considerably over long immersion time of 168 h though it provided
the lowest corrosion rate at short immersion time of 8 h.
(2) Dissolution of the b-Mg17Al12 phase in T4 microstructure
decreased the cathode-to-anode area ratio leading to highly
localised corrosion in the a-Mg matrix. Intergranular corrosion and pitting were the predominant corrosion mechanisms
in the T6 microstructure. The homogeneous distribution of
the b-Mg17Al12 precipitates as well as the aluminium content
in the a-Mg matrix affected the rate of corrosion in the T6
microstructure.
(3) The length of the exposure tests may be too short to compare with the expected time of the alloy’s application in
human body. It would be desirable to conduct further
researches in the future on how T6 heat treatment influences the corrosion process over much longer periods.
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