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4 Surface Modification of Magnesium and its Alloys for Biomedical Applications
Researchers interested in Mg for biodegradable applications are urged to consult the
prior literature, particularly the recent reviews (Atrens, Liu, Zainal Abidin, 2011;
Atrens, Liu, Zainal Abidin, & Song, 2011). It is not the intention to repeat (or even
summarise) herein all the information contained therein. Researchers should also consult our other prior reviews on Mg corrosion (Atrens, Winzer, Dietzel, Srinivasan, &
Song, 2011; Atrens, Dietzel, Srinivasan, Winzer, & Kannan, 2011, Atrens, Shi, &
Song, 2011; Song & Atrens, 1999; Song & Atrens, 2000; Song & Atrens, 2003;
Winzer et al., 2005), papers on measurement of Mg corrosion (Liu, Qiu, Zhao,
Song, & Atrens, 2008; Liu et al., 2009; Shi & Atrens, 2011; Shi, Liu & Atrens,
2010, Shi, Prasad, & Atrens, 2012a; Song, Atrens, & StJohn, 2001), the Mg corrosion
mechanism (Atrens & Dietzel, 2007; Atrens, 2013; Cao et al., 2013a; Cao, Shi, Song,
Liu, & Atrens, 2013b; Liu et al., 2008, Liu, Schmutz, Uggowitzer, Song, & Atrens
et al., 2010; Qiao, Shi, Hort, Zainal Abidin, & Atrens, 2012; Shi & Atrens, 2013a;
Shi et al., 2012; Shi, Jia, & Atrens, 2012b; Shi, Cao, Song, & Liu, 2013b; Song,
Atrens, St. John, Nairn, & Lang, 1997; Song, Atrens, StJohn, Wu, Nairn, 1997;
Song, Atrens, Wu, & Zhang, 1998; Song, Atrens, & Dargusch, 1999; Zhao, Liu,
Song, & Atrens et al., 2008a; Zhao, Liu, Song, & Atrens, 2008b), galvanic corrosion
(Atrens et al., 2011; Jia et al., 2004; Jia, Atrens, Song, & Muster 2005; Jia, Song, &
Atrens, 2005; Jia, Song, & Atrens 2006; Jia, Song, & Atrens, 2007; Shi et al.,
2012c, Winzer et al., 2005), stress corrosion cracking (SCC) (Atrens et al., 2011;
Atrens et al., 2011; Bobby Kannan, Dietzel, Blawert, Atrens, & Lyon, 2008; Shi
et al., 2012b; Song, Blawert, Dietzel, & Atrens., 2005; Winzer, Atrens, Dietzel,
Song, & Kainer, 2007a,b,c; Winzer, Atrens, Dietzel, Song, & Kainer, 2008a,b,c;
Winzer, et al., 2008), surface films (Liu et al., 2009; Liu et al., 2010; Seyeux et al.,
2009), and flammability (Liu et al., 2012; Prasad, Shi, & Atrens, 2012a,b). There is
also much recent literature (Alvarez-Lopez et al., 2010; Bakhsheshi-Rad, Abdul-
Kadir, Idris, & Farahany et al., 2012;Capek & Vojtech, 2013; Choudhary & Singh
Raman, 2013; Han et al., 2012; Kalb et al., 2012; Liu et al., 2010; Minarik et al., in
press; Pu et al., 2012; Sun, Li, & Fang et al., 2011; Yang et al., 2012; Zhou, Shen,
& Aung et al., 2010) on these topics, and there are also much valuable data in the
ASM Handbook (Shaw & Wolfe, 2005).
This chapter focuses on recent developments, under the following subhead ings:
(1) corrosion of Mg alloys, (2) ultra-pure Mg, (3) measurement of Mg corrosion, (4)
in vitro and in vivo comparison, (5) Mg corrosion mecha nism, and (6) galvanic
corrosion. The important points are summarised in the final section.
1.2 Corrosion of magnesium alloys
Mg is the most active of the engineering materials. Mg has a high driving force for
corrosion, weakly checked by the poorly protective corrosion product films that
typically form on the surface of Mg in most corrosion situations (Song & Atrens,
1999, 2000, 2003).
Mg alloys corrode faster than high-purity (HP) Mg because of the microgalvanic
acceleration caused by second phases. In this context, HP Mg is defined as Mg with
Revolutionising biodegradable biomaterials e significance of magnesium and its alloys 5
an impurity level lower than the tolerance limit (Liu et al., 2009). Thus, purification
may produce an alloy with a corrosion rate lower than that of an impure alloy. Nevertheless, the HP Mg alloy has a corrosion rate greater than that of HP Mg.
Particularly problematic are second phases associated with the impurity elements,
particularly iron (Fe). Fe is easily picked up during melt processing in Fe containers
or hot or cold working with Fe tools. These Fe-rich phases, and particles, can increase
the corrosion rate by orders of magnitude, even for concentrations of Fe in the Mg
alloy in the parts per million (ppm) range (Liu et al., 2008, 2009). This is clear
from Figure 1.1 (Atrens et al., 2011; Liu et al., 2009), in which the corrosion of
low-purity Mg (280 ppm Fe) and HP Mg (45 ppm) is characterised using hydrogen
evolution during an immersion test. The corrosion rate is high if the Fe content is
above the impurity limit, whereas the corrosion rate is low for HP Mg.
One molecule of hydrogen is produced, durin g Mg corrosion in an immersion test,
by each corroded Mg atom, as is apparent from the overall corrosion reaction:
Mg þ H
þ
þ H2/MgþþþOHþ H
2
(1.1)
The measurement of the evolved hydrogen is an easy and convenient means to characterise Mg corrosion (Song et al., 2001), and measurement is possible of (1) the
instantaneous corrosion rate, (2) how the instantaneous corrosion rate changes with
time, and (3) the average corrosion rate over the whole exposure period. Measurement
of the corrosion rate using hydrogen evolution is much easier than measurement of the
corrosion rate (1) by weight loss or (2) using electrochemical measurements. Weight
loss measurements require considerable care and experimental skill to ensure (1) that
all the corrosion products have been removed after the corrosion exposure without the
110
100
90
)
2
80
70
60
50
40
evolution (mL/cm
2
30
H
20
10
0
0 10203040506070
Figure 1.1 The corrosion of low purity (280 ppm Fe) and HP (45 ppm) Mg was characterised
using hydrogen evolution. One molecule of hydrogen is produced by each corroded Mg, as is
apparent from the overall corrosion reaction. The corrosion rate is high if the Fe content is above
the impurity limit, whereas the corrosion rate is low for HP Mg. Liu et al. (2009).
High-purity Mg
Low-purity Mg
Immersion time (h)
6 Surface Modification of Magnesium and its Alloys for Biomedical Applications
removal of any Mg metal, and (2) that the amount of moisture on the specimen is
similar when the specimen is weighed before and after corrosion exposure. Electrochemical measurements require a different skill set, and they also require the
appropriate electrochemical apparatus. Electrochemical methods include Tafel extrapolation of polarisation curves, electrochemical impedance spectroscopy (EIS), and
polarisation resistance. Furthermore, anyone using electrochemical methods of measurement for Mg corrosion should be aware of the known limitations of these techniques, which are unique and specifictoMg(Atrens et al., 2011; Cao et al.,
2013a,b; Shi & Atrens, 2011; Shi et al., 2010, 2012a; Song & Atrens, 2003). Hydrogen
evolution measurements also need care and may lead to corrosion rates lower than the
actual corrosion rate, particularly for (1) HP Mg and (2) Mg alloy corrosion in synthetic body fluids (Zainal Abidin et al., 2011; Zainal Abidin et al., 2013). In these
cases, not all the evolved hydrogen is collected because some hydrogen dissolves in
the Mg metal.
Nevertheless, it is best practice to use a number on independent measurements of
the Mg corrosion rate and to convert each independent measurement to the same units
so that they can be easily compared. The independent measurements should all have
the same quantitative value if they are measuring the same quantity.
The Fe tolerance limit can be understood by reference to the Mg-Fe phase diagram,
as illustrated in Figure 1.2 (Atrens et al., 2011; Liu et al., 2009). The Mg-Fe phase
diagram is a eutectic diagram. The eutectic composition corresponds to 180 ppm
(0.018 wt%) Fe. This eutectic composition corresponds to the tolerance limit for
as-cast Mg. The microstructure contains the body centred cubic (BCC) Fe rich phase
for the solidification of a casting with a Fe content greater than the eutectic Fe content,
and consequently the corrosion rate is high. The solidification of a casting containing
less than 180 ppm is expected, under normal solidification, to result in alpha-Mg containing Fe in solid solution, and consequently the corrosion rate is low under these conditions, as is evident from Figure 1.1. However, Figure 1.2 indicates that such a casting
is in a metastable state with respect to the Fe in solid solution if the Fe content is greater
than a few ppm. Heat treatment at elevated temperatures leads to higher corrosion rates
due to precipitation of Fe rich phases, as shown in Figure 1.3. This implies that corrosion rates are expected to be high for Mg-X alloys solution heat treated to dissolve the
alloying elements (X), because the solution heat treatment causes the precipitation of
the Fe rich phases.
Furthermore, Figure 1.2 indicates that a casting with a Fe content less than the Fe
tolerance limit for as-cast Mg can also be of low purity and can have high corrosion
rate, if the Fe rich phase forms during solidification, despite the low Fe content.
This is clear from the data in Figure 1.4 (Qiao et al., 2012). Figure 1.4 presents data
for HP (40 ppm) Mg characterised using hydrogen evolution, specimens 26 and 27
compared with corrosion of specimens 19 and 28, which came from a Helmholtz Zentrum Geesthacht (HZG) ingot (30 ppm). Specimens 26 and 27 were HP Mg. In
contrast, specimens 19 and 28 did not behave consistent with the behaviour of HP
Mg, despite their Fe content being below the Fe tolerance limit for Mg castings. Moreover, the corrosion morphology was also not consistent with the corrosion of HP Mg as
is evident from Figure 1.5.
Revolutionising biodegradable biomaterials e significance of magnesium and its alloys 7
700
675
650
625
T[C]
600
575
Mg–Fe
550
0 0.01
Mg
HCP
Liquid
0.02
w%(Fe)
Liquid + BCC
HCP + BCC
0.03 0.04
Figure 1.2 The Mg-Fe phase diagram is a eutectic phase diagram with the eutectic composition
corresponding to 180 ppm Fe. This eutectic composition corresponds to the tolerance limit for
as-cast Mg. The microstructure contains the BCC Fe rich phase for the solidification of a Mg
casting with a Fe content greater than the eutectic Fe content, and consequently the corrosion rate
is high. The solidification of a casting containing less than 180 ppm is expected under normal
solidification to result in alpha-Mg containing the Fe in solid solution, and consequently the
corrosion rate is low under these conditions, as is evident from Figure 1.1. However, Figure 1.2
indicates that such a casting is in a metastable state with respect to the Fe in solid solution. Heat
treatment at elevated temperatures leads to higher corrosion rates due to precipitation of Fe rich
phases, as shown in Figure 1.3. This implies that corrosion rates are expected to be high for
Mg-X alloys solution heat treated to dissolve the alloying elements, because the solution heat
treatment causes the precipitation of the Fe rich phases. Liu et al. (2009).
45
40
)
2
35
30
25
20
15
10
evolution volume (mL/cm
2
H
5
0
0 50 75 100 125 150 175 200 22525
Mg
Mg-24h at 550 ºC
Mg-48h at 550 ºC
Immersion time (h)
Figure 1.3 The corrosion of high-purity (45 ppm) Mg characterised using hydrogen evolution.
After heat treatment the corrosion rate has increased significantly. Liu et al. (2009).
8 Surface Modification of Magnesium and its Alloys for Biomedical Applications
10
9
8
7
–1
6
5
4
3
2
Corrosion rate, mm∙y
1
0
–200 20406080
100 120 140 160 180 200
Time, h
No.26 (7d)
No.27 (7d)
No.19 (7d)
No.28 (7d)
Figure 1.4 The corrosion of HP (40 ppm) Mg characterised using hydrogen evolution,
specimens 26 and 27 compared with corrosion of specimens 19 and 28, which came from an
HZG ingot (30 ppm). Specimens 19 and 28 did not behave consistent with the behaviour of HP
Mg, despite their Fe content being below the normal Fe tolerance limit for Mg castings.
Moreover, the corrosion morphology was also not consistent with the corrosion of HP Mg, as is
evident from Figure 1.5. Qiao et al. (2012).
Figure 1.5 Typical corrosion morphology for specimens from the HZG Mg ingot (30 ppm) was
not consistent with the behaviour of HP Mg, despite their Fe content being below the normal Fe
tolerance limit for Mg castings. Qiao et al. (2012).
Revolutionising biodegradable biomaterials e significance of magnesium and its alloys 9
Figure 1.1 illustrates a further important point (Atrens et al., 2011; Liu et al., 2009).
The data for low-purity Mg clearly shows that the corrosion rate is not linear with time.
This is common for Mg alloys. The corrosion rate is smal l initially, then the corrosion
rate accelerates to a steady-state corrosion rate. This is a common observation, both in
vitro in laboratory studies and in vivo (Remennik et al., 2011). Figure 1.4 shows a
similar behaviour for specimens 19 and 28 (Qiao et al., 2012). The corrosion rate
was initial ly small, then accelerated and achieved a steady-state corrosion rate that
was an order of magnitude greater than the initial corrosion rate.
1.3 Ultra-pure magnesium
The corrosion of Mg alloys is currently too fast for applications as a biodegradable
implant (Atrens et al., 2011; Atrens et al., 2011; Song & Atrens, 2003). Furthermor e,
the corrosion rate of all exist ing Mg alloys is faster than that of ultra-HP Mg (Atrens
et al., 2011;u Atrens et al., 2011; Song & Atrens, 2003). There are two possible future
approaches: (1) development of Mg alloys with corrosion rates lower than that of
HP Mg or (2) surface modification of existing Mg alloys to reduce the corrosion
rate. This second approach is the theme of many of the contributions to this book.
This section is related to the development of Mg alloys with corrosion rates lower
than that of HP Mg.
Our recent work discovered how to produce ultra-HP Mg alloys, when ultra-pure
Mg alloys are defined as having a Fe content less than 2 ppm. Figure 1.6 shows
how molten Mg can be purified with zirconium (Zr) (Prasad et al., 2012c) to produce
ultra-HP (UP) Mg-X alloys. Zr is added to the molten Mg alloy containing Fe. Zr reacts with Fe, and forms FeZr
as also shown by Qian et al. (2001), because their density is significantly higher than
molten Mg. The melt composition attains the equilibrium composition corresponding
to the maximum solubility of Fe and Zr in the molten Mg at the reaction temperature.
Figure 1.6(a) shows that 2 ppm Fe (0.0002% Fe) is in equilibrium with 0.2% Zr and
FeZr
in molten Mg at 660C. Thermodynamics indicate that Zr similarly removes
2
aluminium (Al), manganese (Mn), silicon (Si), nickel (Ni), cobalt (Co), tin (Sn), and
antimony (Sb) from molten Mg (Prasad et al., 2012). Our preliminary research (Prasad
et al., 2012c) showed that (1) the melt treatment with Zr works in purifying molten Mg
and (2) the resultant melt composition is in good agreement with theoretical expectations as shown in Figure 1.6(b). The isothermal sections in Figure 1.6(a) and (b) were
calculated using the Pandat thermodynamic package computherm.
Ultra-pure Mg-X alloys (where X may be elements such as dysprosium [Dy] and
gandolinium [Gd]) may be stainless after solution heat treatment (SHT) to dissolve
the alloying element into solid solution. Dy and Gd have high solubilities in Mg
and may produce surface films (Atrens et al., 2011; Atrens et al., 2011; Hort et al.,
2010; Song & Atrens, 2003; Yang et al., 2012) on these Mg-X alloys that are substan-
tially more protective than those that form naturally on UP Mg. This is the mechanism
that produces stainless steels. A highly protective surface passive film forms on stainless steels for a chromium (Cr) content greater than 10.5% (Bruesch et al., 1984, 1985;
particles. These particles settle to the bottom of the melt,
x
10 Surface Modification of Magnesium and its Alloys for Biomedical Applications
0.005
(a)
0.004
0.003
w%(Fe)
0.002
0.001
0
Mg
0
Liquid
Liquid + Fe
0.05
0.1
2
Zr
0.15
0.2
w%(Zr)
Liquid + FeZr
0.25
0.3
T = 660 °C
2
0.35
0.4
(b)
Figure 1.6 (a) Isothermal section through the Mg-Fe-Zr phase diagram at 660C showing the
maximum solubility of Fe and Zr in liquid Mg at 660
measurements with the expectation from the phase diagram at 675
0.007
0.006
0.005
0.004
0.003
w%(Fe)
0.002
0.001
0
Mg
T = 675 °C
2
3
4
0
0.004
0.008
w%(Zr)
Liquid + Fe2Zr
Liquid
0.012
C, (b) comparison of experimental
0.016
C There was good agreement
0.02
between experiment and the theoretical thermodynamic predictions. The isothermal sections in
(a) and (b) were calculated using the Pandat thermodynamic package. Prasad et al. (2012).
Heine & Kirchheim, 1990; Jin & Atrens, 1987, 1988; Kirchheim et al., 1989; Lim &
Atrens, 1991, 1992a,b,c; Olefjord & Elfstrom, 1982; Schneider et al., 1990). In
comparison, red rust (iron oxide) is not protective for a lower Cr content. On this basis,
we suggested MgCr
(Song & Atrens, 2003) in 2003. However, extensive efforts using
y
rapid solidification by Uhlenhaut (2008) were not successful, because molten Cr and
Mg formed immiscible liquids (like oil and water). Now our recent work (Prasad et al.,
2012c) opened up the possibility to produce stainless Mg using conventional
metallurgy.
Revolutionising biodegradable biomaterials e significance of magnesium and its alloys 11
Ultra-HP Mg-X alloys (X ¼ Zn, Sn, Pb) are expe cted to have low corrosion rates,
significantly lower than those of ultra-pure Mg, because the electrochemical properties of the Mg-X alloys may be cha nged compared with the ultra-HP Mg sufficiently
by making the cathodic hydrogen evolution reaction much more difficult than
on ultra-HP Mg (Frankel et al., 2013; Pourbaix, 1966). This is an entirely new
approach to the development of Mg alloys with corrosion rates lower than those of
ultra-HP Mg.
Previously, the intrinsic corrosion behavi ou r of thes e Mg-X a lloys co ul d not be
assessed. Assessment of their intrinsic corrosion behaviour requires Mg-X alloy production (not possible previously in ultra-HP form ) and solution heat treatm ent so that
all the alloying elements are in solid solution. But the intrinsic corrosion behaviour is
masked by a Fe impurity concentration above the Fe tolerance limit (Liu et al., 2008,
2009; Song & Atrens, 2003). Above the Fe tolerance limit , t he corrosion rate is high,
because the corrosion behaviour is dominated by the presence of Fe-rich particles.
Our recent research (Liu et al., 2009) showed that the Fe tolerance limit is
w2 ppm after heat treatment of the type needed to bring all the alloying elements
into solid solution. That means that the Fe impurity element dominates the corrosion
behaviour, and the corr osio n rate is high, if the Fe concentration is greater than
w2 ppm for a Mg-X alloy, heat-treated to bring all the alloying element into solid
solution. And such low Fe concentrations had not been achieved for these alloys
previously. These alloys in the as-cast condition were available in HP form, because
Fe concentrations between 40 ppm and 180 ppm for cast alloys (Liu et a l., 2009)are
readily achieved in industrial practice. However, their corrosion rate is high in the
cast condition because of galvanic acceleration by the second phase in the twophase cast microstructure. And after heat treatment of these alloys with Fe contents
greater than w2 ppm, the corrosion behaviour is dominated by the Fe impurity
element.
There is no alternative method to produce ultra-high purity (UP) Mg-X alloys, with
X ¼ Dy and Gd. The alloying elements (particularly Dy and Gd) are not available in
sufficient purity to allow production of UP Mg-X alloys by conventional melting
with UP Mg produced by vacuum distillation as produced by Uggowitzer (2012) or
Hort (2013). The Fe content is typically over 1000 ppm for master alloys contai ning
Dy and Gd. Furthermore, vacuum distillation cannot produce Mg-X alloys because
of the high melting point of the alloying elements.
Mg-X alloys (X ¼ Dy, Gd) have been shown to be biocompatible (Feyerabend
et al., 2010; Hort et al., 2010; Yang et al., 2012).
The logical and feasible next step is to produce ultra-pure Mg-X alloys and to study
their corrosion behaviour with the expectation that these alloys will produce Mg alloys
that have corrosion rates lower than or equal to that of HP Mg.
1.4 Measurement of magnesium corrosion
Weight-loss experiments, properly carried out, are routinely used to measure long-term
corrosion rates . Moreover, weight loss can be used to measure both in vitro corrosion
p
12 Surface Modification of Magnesium and its Alloys for Biomedical Applications
rates and in vivo corrosion rates. The best and most reliable comparison would be if
these weight-loss measurements, both in vivo and in vitro, were carried out by the
same experimenter or by the same research group. Then, it would be expected that
there would be similar errors in both the in vitro and the in vivo measurements, for
measurements carried out by the same experimenter. It is more difficult to carry out
comparisons between groups.
In contrast, Tafel extrapolation of polarisation curves (typically measured soon after
specimen immersion in the solution) (Atrens et al., 2011; Shi et al., 2010; Song &
Atrens, 2003) has not yielded values of corrosion rate for Mg that have agreed with
other independent measurements of corrosion rate, like weight loss or hydrogen evolution. This is despite the fact that Tafel extrapolation is routinely used for the measurement of the corrosion rate for most corroding metals (Jones, 1992). There are
some well-known reasons why Tafel extrapolation has not yielded good measurements
for Mg corrosion (Shi & Atrens, 2011; Shi et al., 2012b). In the present circumstances,
it must be considered the responsibility of any researcher who wishes to use Tafel
extrapolation, to be aware of the known limitation and to provide supporting data
from at least one of the other methods for the measurement of the corrosion rate of
Mg alloys. Furthermore, any research that provides Mg corrosion data based only
on Tafel extrapolation must be viewed critically.
One reason for the lack of agreement between Tafel extrapolation and other
measurements of corrosion for Mg alloys is that Tafel extrapolation is often
carried out for polarisation curves measured soon after specimen immersion in the
solution (Kirkland et al., 2012). The reason for the problem is obvious from the
data in Figure 1.7 (Zainal Abidin et al., 2011). The open symbols show the corrosion
)
–2
Hydrogen evolution volume (mL cm
60
50
40
30
20
10
H
AZ91 1
2
H
AZ91 2
2
AZ91 3
H
2
pH AZ91 1
pH AZ91 2
pH AZ91 3
0
50
0
100 150 200 250 300 350
Immersion time, t, (hour)
9
8
H
7
6
5
Figure 1.7 Corrosion behaviour as characterised by hydrogen evolution (open symbols), as well
as the solution pH (corresponding closed symbols) during the immersion test for AZ91 in Nor’s
solution (CO
e bicarbonate buffered Hank’s solution, with the CO2partial pressure equal to
2
0.009 atm). Zainal Abidin et al. (2011).
Revolutionising biodegradable biomaterials e significance of magnesium and its alloys 13
of AZ91, as characterised by hydrogen evolution, of AZ91 in Nor’s solution (CO
bicarbonate buffered Hank’s solution, with the CO
partial pressure equal to
2
2
0.009 atm). This shows that the corrosion rate of AZ91 was initial ly low during an incubation period, whereas afterwards the corrosion rate accelerated to a steady state
value of w18 mm y
1
. Corrosion rates varied by a factor of two for identical specimens tested under identical conditions. This variability is common in the study of
Mg corrosion (Atrens et al., 2011; Shi & Atrens, 2011; Zhao et al., 2008a). In cases
such as that illustrated in Figure 1.7, there is little relationship between the initial
corrosion rate and the steady-state corrosion rate.
Similarly, it has been reported that the corrosion rate increases in vivo. Witte et al.
(2005) found the in vivo corrosion rate for AZ91, implanted into the femora (bone) of
guinea pigs, was P
E¼ 0.43 mm y1(over 6 weeks) and P
r6w
¼ 1.3 mm y1(over
r18w
18 weeks; where the subscript “r” indicates the corrosion rate was evaluated from the
decrease in size of the implant). The corrosion rate increased with increasing exposure
time. Similarly, Xu et al. (2007), for as-cast Mg-1.2Mn-1Zn implanted into rat femora,
reported corrosion rates of P
¼ 0.3 mm y1(over 9 weeks) and P
r9w
¼ 0.7 mm y
r18w
1
(over 18 weeks), consistent with those of Witte et al. (2005), which also increased with
increasing exposure time.
Similarly, the corrosion rate accelerated for the data for low-purity Mg in Figure 1.1,
the HP Mg after heat treatment in Figure 1.3, specimens 19 and 28 in Figure 1.4. In all
these cases, the measurement of corrosion rate soon after specimen immersion does not
provide a good estimate of the long-term corrosion, which is what is needed if the
service life is to be predicted.
It is for this reason that the advice of Kirkland et al. (2012) is not appropriate
that the corrosion rate (by Tafel extrapolation of polarisation curves) should be
assessed from polarisation curves measured soon after specimen immersion in the
solution.
Fishing-line specimens and plug-in specimens were developed (Sh i & Atrens,
2011) in response to the known issues (Atrens et al., 2011; Shi et al., 2010; Song &
Atrens, 2003) with the measurement of Mg corrosion using electrochemical methods.
Fishing-line specimens were designed as the most minimalist possible specimen
mount and identified the issues of crevice corrosion for Mg during immersion tests
(Shi & Atrens, 2011).
Plug-in specimens were subsequently designed, Figure 1.8 (Shi & Atrens, 2011), to
have no crevice and to allow reliable polarisation curves to be measured. Three reasons
were identified (Shi & Atrens, 2011) why Tafel extrapolation had previously not
yielded corrosion rate measurements in agreement with independent measurements:
(1) crevice corrosion can occur in the specimen mount when the specimen is mounted
in a metallurgical mount or similar as widely used in Mg corrosion studies, Figure 1.9;
(2) Tafel extrapolation has often been used to measure corrosion behaviour soon after
specimen immersion. This initial corrosion behaviour may not correlate with steady
state corrosion behaviour, either in vivo or in vitro, see previous section also;
(3) Figure 1.10 shows that at times longer than 14 days, the corrosion rate increased
significantly, but this increase in corrosion rate was not measured by the Tafel
e
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