Corrosion Science 52 (2010) 1035–1041

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Corrosion Science

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Effect of heat treatment on corrosion behaviour of magnesium alloy AZ91Din simulated body fluid

Wei Zhou *, Tian Shen, Naing Naing AungSchool 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 a b s t r a c t

Article history:Received 12 June 2008Accepted 18 November 2009Available online 27 November 2009

Keywords:A. AlloyA. MagnesiumB. Weight lossB. PolarisationC. Intergranular corrosion

0010-938X/$ - see front matter � 2009 Elsevier Ltd. Adoi:10.1016/j.corsci.2009.11.030

* Corresponding author. Tel.: +65 6790 4700; fax: +E-mail addresses: WZhou@Cantab.Net, mwzhou@n

The research explored ways of improving corrosion behaviour of AZ91D magnesium alloy through heattreatment for degradable biocompatible implant application. Corrosion resistance of heat-treated sam-ples is studied in simulated body fluid at 37 �C using immersion and electrochemical testing. Heat treat-ment significantly affected microgalvanic corrosion behaviour between cathodic b-Mg17Al12 phase andanodic a-Mg matrix. In T4 microstructure, dissolution of the b-Mg17Al12 phase decreased the cathode-to-anode area ratio, leading to accelerated corrosion of a-Mg matrix. Fine b-Mg17Al12 precipitates in T6microstructure facilitated intergranular corrosion and pitting, but the rate of corrosion was less thanthose of as-cast and T4 microstructures.

� 2009 Elsevier Ltd. All rights reserved.

1. Introduction

Degradable biocompatible materials play an essential role inload-bearing implants due to the risk of local inflammation of tra-ditional implants such as stainless steels, titanium alloys, and co-balt–chromium-based alloys. Unfortunately, currently usedabsorbable polymer materials have unsatisfactory mechanicalproperties. In comparison, magnesium alloys are attractive candi-dates because they are ultralight alloys possessing mechanicalproperties similar to those of natural bones and a natural ionicpresence with significant functional roles in biological systems[1–4]. However, fast degradation via corrosion in the electrolyticenvironment of the body constitutes the main shortcoming ofmagnesium alloy implants [5–10]. Their degradation rates varyover a range of 3 orders of magnitude [11]. Therefore, improvingcorrosion resistance is an important issue for the application ofmagnesium alloys as biodegradable load-bearing implants.

The corrosion resistance of magnesium alloys, particularlyAZ91D alloy, can be improved by different surface coating tech-niques such as electro plating [12], electroless plating [13], anodiz-ing [14], laser surface cladding [15] and laser surface melting [16].It is also possible to control the degradation process of magnesiumalloy through Zn and Mn alloying, purification and anodization[17,18]. A Zn-containing magnesium alloy with a small amountof manganese can be a potential biodegradable alloy [17]. If the al-loy starts biodegradation too early, an anodized coating can be ap-

ll rights reserved.

65 6791 1859.tu.edu.sg (W. Zhou).

plied to delay it. If the alloy biodegrades too fast, it can be purifiedto bring down the rate [17]. However, some surface coating tech-niques could have a negative impact on the implants due to theirpossible release of toxic metallic ions.

In AZ91D alloy, the a-Mg matrix corrodes due to its very nega-tive free corrosion potential and there is the tendency for the cor-rosion rate of the a-Mg phase to be accelerated by microgalvaniccoupling between anodic a-Mg phase and cathodic b-Mg17Al12

phase [19–25]. However, the b-Mg17Al12 phase may act as a barrieragainst corrosion propagation if it is in the form of a continuousnetwork [20,22]. The distribution, configuration and size of the b-Mg17Al12-phase can be changed, which may result in different cor-rosion rates.

The corrosion resistance of the alloy may be enhanced by heattreatment [22–25]. In this case, heat treatment does not normallyincrease the corrosion barrier effect of the b-Mg17Al12 phase. Solu-tion heat treatment dissolves the b-Mg17Al12 phase and removesits barrier effect. However, the b-Mg17Al12 precipitates producedby ageing were effective to reduce microgalvanic corrosion of theadjacent a-Mg phase [22]. The objective of the research was tostudy how heat treatment could be explored to improve corrosionresistance of AZ91D alloy in simulated body fluid.

2. Materials and methods

2.1. Materials

The material studied was an as-cast ingot of AZ91D alloy with thefollowing chemical composition (in wt.%): Al–9.1, Mn–0.17,

1036 W. Zhou et al. / Corrosion Scien

Zn–0.64, Si–60.01, Fe-60.001, Cu–60.01, Ni-0.001, other impurities60.02, Mg-the rest. Solution treatment (T4) of the alloy was carriedout at 445 �C for 24 h in argon atmosphere followed by water quenchat 25 �C. Ageing treatment (T6) of the solution-treated samples wasperformed at 200 �C for 8 h, 16 h and 24 h to produce three differentaged microstructures. The as-cast and heat-treated samples werecut into samples of 10 mm � 10 mm � 2 mm in dimension forimmersion test and potentiodynamic polarisation study.

2.2. SBF preparation

The bioactivity of bone implant materials is usually tested in vitrousing simulated body fluid (SBF). The composition of common SBFdiffers from that of blood plasma in that it has a higher Cl� and a low-er HCO3

� concentration, which could affect the composition ofin vitro formed bone-like apatite. In this research, the high stabilitySBF solution was prepared by mixing stable concentrated solutions,which increase the reproducibility of in vitro tests due to negligiblechanges of pH during preparation.

The SBF solution was prepared by pipetting calculated amountsof concentrated solutions of KCl (59.64 g/l), NaCl (116.88 g/l), NaH-CO3 (45.37 g/l), MgSO4�7H2O (49.30 g/l), TRIS (tris-hydroxymethylaminomethane (121.16 g/l), NaN3 (100 g/l) and KH2PO4 (27.22 g/l)into double distilled water to prevent precipitation of homoge-neously nucleated calcium phosphates or other phases and to mini-mize changes in pH during preparation [26]. The SBF had thefollowing composition (in ml/l): KCl-5, NaCl-50, NaHCO3-50,MgSO4�7H2O-5, CaCl2-25, TRIS (tris-hydroxymethyl aminometh-ane) + 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+, 109Cl�, K+, 27 HCO3

�, 1 SO42�, 1 HPO4

2�. The pH of human blood plasmaranges from 7.3 to 7.4 at 37 �C [26], so the pH of SBF was adjusted to7.6–7.7 at 25 �C which is equal to a pH of 7.3–7.4 at 37 �C, by addingconcentrated HCl.

2.3. Immersion test

The corrosion rate was determined by weight loss rate from theimmersion test according to ASTM-G31-72 [27]. The as-cast anddifferent heat-treated samples were immersed in 300 ml solutionof SBF for various periods of 8 h, 24 h, 72 h and 168 h. Separatesample 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 finergrades of emery papers up to 4000 grit and then weighed. Allimmersion tests were carried out without agitation or circulationand without disturbing corrosion system.

In order to remove the corrosion products with minimal disso-lution of base alloy, chemical cleaning of the corroded sample wascarried out in boiling 15% CrO3 + 1% AgCrO4 solution for 1 min. Anacetone washing followed this. The weight loss was measured aftereach experiment and the corrosion rate (R) was calculated in mgcm�2 day�1 using Eq. (1). The values of corrosion rate were con-verted from mg cm�2 day�1 to mm y�1 using Eq. (2). The tests wererepeated three times to obtain the reproducible results.

R ¼ ðWb �Wa þ BÞ � 1000A� t

ð1Þ

where:

R = corrosion rate, mg cm�2day�1

Wb = weight of test sample before test, gWa = weight of test sample after test, gB = weight loss of blank, g (the average weight loss from threeunused and clean sample was used as the blank correction)A = surface area of sample, cm2

t = exposure time, day.

Corrosion rateðmm y�1Þ ¼ 10:274� q

R ð2Þ

q = density of alloy.

2.4. Electrochemical test

The electrochemical corrosion behaviour of the as-cast andheat-treated samples was studied using a potentiostat/galvanostatcorrosion measurement system (EG&G model 263A). The sampleswere encapsulated into epoxy resin so that only a surface withthe dimension of 10 mm � 10 mm was exposed to the 300 ml ofthe solution. Prior to the experiment the samples were groundup to 4000 grit emery paper, followed by washing with distilledwater and acetone. A potentiodynamic polarisation test was car-ried out in a standard electrochemical cell using three electrodesconfiguration. The sample was the working electrode. A saturatedcalomel electrode (SCE) was used as a reference electrode and aplatinum electrode was used as a counter electrode. The test cellwas placed into the water bath having environmental temperatureto 37 �C.

Potentiodynamic polarisation curves were measured immedi-ately after immersion of the sample in the solution. The cathodicpolarisation scan was started from �2000 mV to the steady statecorrosion potential and then polarizing in an anodic direction ata scan rate of 1 mV s�1. The icorr in mA cm�2 can be related to theicorr in mm y�1 using Eq. (3).

Corrosion rateðmm y�1Þ ¼ 3:28Mnq

icorr ð3Þ

n = number of electrons freed by the corrosion reactionM = atomic mass.

ce 52 (2010) 1035–1041

2.5. Analysis of corroded surface

Microstructure and morphology of the corroded surfaces wereanalysed using Carl Zeiss Axioskop 2 optical microscope and scan-ning electron microscopy (SEM, Jeol model 5600 LV) coupled withenergy dispersed spectroscopy (EDS) system.

3. Results and discussion

3.1. Microstructures

The as-cast AZ91D microstructure has typically a primary a-Mgmatrix and a divorced eutectic phase distributed along the grainboundaries (Fig. 1(a)). The close-up view clearly shows that the eu-tectic consists of large b-Mg17Al12 phase particles and the eutectica-Mg phase (Fig. 1(b)). The eutectic a-Mg phase is supersaturatedwith Al and it can transform to form a fine lamellar structure.

Microstructure of the alloy was changed during the process ofheat treatment. The T4 heat treatment dissolved the b-Mg17Al12

phase and produced a microstructure consisting the supersaturateda-Mg phase. However, some residual small b-Mg17Al12 phases werestill observed (Fig. 1(c)). During T6 treatment, b-Mg17Al12 phase pre-cipitated along the grain boundary and within grains of the supersat-urated a-Mg phase. Ageing made Al atoms diffuse towards grainboundaries to form precipitates of b-Mg17Al12 phase, and this pro-cess reduced the aluminium concentration in the a-Mg matrix. Inthis case, the degree of homogeneity of the b-Mg17Al12 precipitatesdistribution and the aluminium content in the a-Mg matrix was dif-ferent between the T6 microstructures (Fig. 1(d–f)). The aluminiumcontent in the a-Mg matrix of T6-16 h microstructure was found to

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 hand (f) T6-24 h.

W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041 1037

decrease from 9% to 3%. This result confirmed with other researcher’sresult [28].

3.2. Effect of heat treatment on corrosion rates

Variation of corrosion rate with immersion time for the as-cast and heat-treated samples is presented in Fig. 2. For short

immersion period, 8 h, the corrosion rate of T4 sample wasthe lowest compared with other samples (Fig. 2(a)). However,its corrosion rate shifted to the highest value after long expo-sure, 168 h (Fig. 2(b)). The corrosion rates of T6 samples werehigher than those of the as-cast and T4 microstructure at theinitial exposure of 8 h (Fig. 2(a)), but the rates slowed downto the lower values after 168 h (Fig. 2(b)). Among T6 samples,

(b)

(a)

As-cast T4 T6-8 h T6-16 h T6-24 h

2

4

6

8

10

12

14

5

10

15

20

25

30

Corrosion rate (m

m y

-1)

Cor

rosi

on ra

te (m

g cm

-2 d

ay-1)

168 h immersion

As-cast T4 T6-8 h T6-16 h T6-24 h

0.2

0.4

0.6

0.8

1.0

0.4

0.8

1.2

1.6

2.0

Corrosion rate (m

m y

-1)

Cor

rosi

on ra

te (m

g cm

-2 d

ay-1) 8 h immersion

Fig. 2. Corrosion rates for as-cast and heat-treated samples in SBF: (a) after 8 h and(b) after 168 h.

0 50 100 1500

5

10

15

0

5

10

15

20

25

30

Corrosion rate (m

m y

-1)Cor

rosi

on ra

te ( m

g cm

-2 d

ay-1)

Immersion time (h)

100 µm

100 µm

100 µm

Fig. 3. Extent of corrosion for T4 samples in SBF as a function of exposure time.

100 µm

100 µm

100 µm

0 50 100 150

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

1.0

1.5

2.0

2.5

3.0

3.5

4.0

Corrosion rate (m

m y

-1)

Cor

rosi

on r

ate

(mg

cm-2da

y-1)

Immersion time (h)

Fig. 4. Extent of corrosion for T6-16 h samples in SBF as a function of exposuretime.

1038 W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041

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 func-tion of exposure time are shown in Figs. 3 and 4. The corrosion rateand the corresponding area corroding of the samples increasedwith time of exposure to the corroding solution. Corrosion startedaround localised sites at grain boundaries in the T4 sample and theattack invaded the entire surface with continued exposure. After168 h immersion, T4 sample was highly damaged, as shown inFig. 3. For T6-16 h sample, corrosion accelerated with increasingexposure time but the extent of corrosion was lower in T6-16 hsample (Fig. 4) compared with T4 sample (Fig. 3). The results sug-gest that the heat treatment which gave the best resistance to cor-rosion in SBF was T6 rather than T4.

3.3. Effect of microstructure on corrosion

In order to understand the effect of heat treatment on corrosionmechanism in greater detail, the corroded surfaces of samples for8 h and 72 h exposure were carefully analysed under SEM. SelectedSEM micrographs are shown in Figs. 5 and 6.

Corrosion of the as-cast microstructure initiated at the primarya-Mg matrix in the eutectic region as indicated by arrows inFig. 5(a). In the case of T4 microstructure, localised attack was ob-served around residual b-Mg17Al12 phase at grain boundaries(Fig. 5(b)). For the T6 microstructures, corrosion occurred preferen-

tially along the grain boundary and some pits were found withingrains as indicated by arrows in Fig. 5(c–e).

Corroded morphologies of the samples after longer exposureclearly showed that the changes in distribution, configurationand size of the b-Mg17Al12-phase due to heat treatment resultedin different corrosion behaviours (Fig. 6). In the as-cast microstruc-ture, the b-Mg17Al12 phase is highly cathodic to the a-Mg phaseand can thus act as an effective cathode to cause microgalvaniccorrosion [19–25]. The b-Mg17Al12 phase contains much more alu-minium compared with the a-Mg phase. According to our previousstudy [20], the variation of the concentration of aluminium is inthe range of about 35% in the b-Mg17Al12 phase to about 6% inthe primary a-Mg phase. The region with less than 8% aluminiumcould be corroded preferentially [20]. Therefore, the lower alumin-ium content of the primary a-Mg matrix was the initiation site ofcorrosion (Fig. 6(a)).

In the T4 microstructure, there was a meta-stable, partially pro-tective film on high aluminium content of the a-Mg matrix. Thisfilm prevented corrosion and the result therefore showed the

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 .

W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041 1039

lowest corrosion rate in the initial exposure (Figs. 2(a) and 3).However, the dissolution rate of localised corrosion could be fasteronce the protective film was broken down due to weak localisedsites of residual b-Mg17Al12 phase (Fig. 6(b)). Besides residual b-Mg17Al12 phase, Al–Mn and Al–Mn–Fe intermetallics could alsogreatly accelerate the corrosion [29]. As a result, the corrosion ratesignificantly increased after longer exposure (Figs. 2(b), 3 and6(b)).

For T6 microstructure, the corrosion mechanism appeared to beinfluenced by the precipitation of b-Mg17Al12 in relation to theamount of aluminium content in the a-Mg matrix. The corrosionpits initiated at the anodic a-Mg matrix adjacent to the cathodicb-Mg17Al12 precipitates. The microgalvanic action due to the b-Mg17Al12 precipitates at grain boundary led to intergranular corro-sion. The T6 microstructural features gave the tendency to occurintergranular corrosion as well as pitting corrosion (Fig. 6(c)).

Fig. 6. SEM micrographs showing different corrosion mechanisms on the corroded surfaces: (a) as-cast, (b) T4 and (c) T6-16 h.

-1.8 -1.6 -1.4 -1.2 -1.0

0.01

0.1

T4As-cast

T6-24 h

T6-16 h

T6-8 h

I (Ac

m-2)

E (V/SCE)

Fig. 7. Potentiodynamic polarisation curves for as-cast and heat-treated samples inSBF at 37 �C.

1040 W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041

However, their extent of corrosion was lower compared with theas-cast and T4 microstructure (Figs. 2(b), 4 and 6(c)). The isolatedfine b-Mg17Al12 precipitates do not lead to an obvious loss of cor-rosion because they act as a small cathode connected to the largeanode a-Mg matrix [22]. In this case, the corrosion resistance ofthe T6-16 h microstructure was slightly better than the T6-8 hand T6-24 h microstructure. The reason could be due to its morehomogeneous distribution of the b-Mg17Al12 precipitates as wellas the aluminium content in the a-Mg matrix.

3.4. Electrochemical behaviour of heat-treated microstructures

The electrochemical corrosion behaviour of the samples in SBFis shown in Fig. 7. The values of corrosion potential (Ecorr), Tafelslope (bc) and the corrosion current density (icorr) for each polarisa-tion curve are summarized in Table 1. The Ecorr values of heat-trea-ted samples were shifted to less negative values showing morecathodic behaviour compared with the as-cast one. It should beespecially noted that the T6-16 h sample showed less negative va-lue of Ecorr. The bc values were similar for all samples, indicatingthat the same electrochemical reactions occurred. The cathodiccurrents from the polarisation curves were much higher for allheat-treated samples at all potentials. On the other hand, the T6-16 h microstructure was the least active anodically with the lowesticorr value (Table 1).

The polarisation curves were used to explore the relationshipbetween electrochemical measurements of the corrosion rate,

based on the corrosion current at the free corrosion potential (Ta-ble 1), and direct measurements using weight loss (Fig. 2). Therewas a good correlation between the corrosion rates determinedfrom icorr and those from weight loss data only for short immersiontime. The corrosion rate from the Tafel extrapolation may relate tothe onset of corrosion, whereas the corrosion rate from the weight

Table 1Ecorr, bc and icorr values for as-cast and heat-treated samples in SBF.

Sample Ecorr (V/SCE) bc (V dec�1) icorr (mA cm�2) Corrosion rate (mm y�1)

As-cast �1.31 0.186 0.039 0.85T4 �1.24 0.176 0.028 0.61T6-8 h �1.29 0.189 0.066 1.44T6-16 h �1.19 0.195 0.027 0.59T6-24 h �1.29 0.192 0.074 1.62

W. Zhou et al. / Corrosion Science 52 (2010) 1035–1041 1041

loss/hydrogen evolution measurement relates to corrosion aver-aged over a considerable time period and includes corrosion someconsiderable time after corrosion onset, when the corrosion is wellestablished [30].

4. Conclusions

(1) Heat treatment significantly changed the corrosion resis-tance of AZ91D alloy in SBF. Compared with the as-castcondition, the T6 treatment reduced the corrosion rate by30–60%, and the T4 treatment increased the rate consider-ably over long immersion time of 168 h though it providedthe lowest corrosion rate at short immersion time of 8 h.

(2) Dissolution of the b-Mg17Al12 phase in T4 microstructuredecreased the cathode-to-anode area ratio leading to highlylocalised corrosion in the a-Mg matrix. Intergranular corro-sion and pitting were the predominant corrosion mechanismsin the T6 microstructure. The homogeneous distribution ofthe b-Mg17Al12 precipitates as well as the aluminium contentin the a-Mg matrix affected the rate of corrosion in the T6microstructure.

(3) The length of the exposure tests may be too short to com-pare with the expected time of the alloy’s application inhuman body. It would be desirable to conduct furtherresearches in the future on how T6 heat treatment influ-ences the corrosion process over much longer periods.

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