Evaluation of microstructural e•ects on corrosion ... 2000 - Corrosion Scie… · passivation compared with the ingot. In die-cast and ingot, hydrogen evolution took place preferentially

Evaluation of microstructural e�ects oncorrosion behaviour of AZ91D magnesium

alloy

Rajan Ambat*, Naing Naing Aung, W. Zhou

School of Mechanical and Production Engineering, Nanyang Technological University, Nanyang Avenue,

Singapore 639798, Singapore

Received 3 September 1999; accepted 13 November 1999

Abstract

The e�ect of microconstituents on the corrosion and electrochemical behaviour ofAZ91D alloy prepared by die-casting and ingot casting route has been investigated in 3.5%

NaCl solution at pH 7.25. The experimental techniques used include constant immersiontechnique, in-situ corrosion monitoring, and potentiodynamic polarisation experiments.Surface examination and analytical studies were carried out using optical and scanning

electron microscopy, EDX and XRD. The corrosion behaviour of microconstituents namelyprimary a, eutectic a and b phases was signi®cantly di�erent. Coring of aluminum showedin¯uence on corrosion behaviour more signi®cantly in ingot material. Areas with

aluminium concentration less than about 8% were found to be prone to corrosion attackcompared with either those with higher amount of aluminium or b phase. Die-cast materialwith smaller grain size and ®ne b phase o�ered marginally lower corrosion rate and betterpassivation compared with the ingot. In die-cast and ingot, hydrogen evolution took place

preferentially on b phase. XRD pattern of non-corroded and corroded surface revealed theremoval of b phase from alloy surface during corrosion. The corrosion products for ingotconsisted of Mg(OH)2 with small amounts b phase, magnesium-aluminium oxide and

MgH2 while for die-cast, the product showed a highly amorphous structure. 7 2000Elsevier Science Ltd. All rights reserved.

Keywords: Magnesium alloys; AZ91D; Ingot; Die-cast; Microstructure; Corrosion

0010-938X/00/$ - see front matter 7 2000 Elsevier Science Ltd. All rights reserved.

PII: S0010 -938X(99)00143 -2

Corrosion Science 42 (2000) 1433±1455

* Corresponding author. Tel.: +65-790-4963; fax: +65-791-1857.

E-mail address: [email protected] (R. Ambat).

1. Introduction

The need for weight reduction, particularly in portable microelectronics,telecommunication, aerospace and automobile sector has stimulated engineers tobe more adventurous in their choice of materials. Magnesium and its alloys,with one quarter of the density of steel and only two thirds that of aluminium,and a strength to weight ratio that far exceeds either, ful®ls the role admirablyas an `ultra light' alloy. In the microelectronics industry, magnesium alloys havebecome promising materials for many components in hand phone, CVD/DVDchassis, computer disk drives and camera casings. The ultra lightness of thesealloys is also attractive for transport industry where weight saving pays manydividends.

Although a wide variety of applications can be envisaged for Mg alloys, theuse at present is limited mainly due to inferior corrosion properties. Thestandard electrochemical potential of magnesium is ÿ2.4 V (NHE), eventhough in aqueous solutions it shows a potential of ÿ1.5 V due to theformation of Mg(OH)2 ®lm. Consequently, magnesium dissolves rapidly inaqueous solutions by evolving hydrogen below pH 11.0, the equilibrium pHvalue for Mg(OH)2. As a result, researchers all over the world has shown muchinterest to study the corrosion behaviour of Mg alloys and to develop protectivemeasures.

Magnesium alloys with Al and Zn have (mainly AZ91 series) found wide spreadapplication in the automobile sector [1,2]. The spectrum includes severalapplications like steering wheel, dashboard mounting bracket, gearbox housingand chassis [2]. Corrosion resistance is one of the prerequisites for all theseapplications.

The corrosion behaviour of cast magnesium±aluminium alloys could dependconsiderably on microstructure (presence of b phase and coring) and theenvironment to which it is exposed. Microstructural features vary with the methodof processing that gives rise to di�erent corrosion behaviours. Even though fewstudies have been reported [3±5] earlier on the e�ect of microstructure on thecorrosion and electrochemical behaviour of AZ91 alloy, a considerable amount ofambiguity still exists. More importantly, coring during solidi®cation that leads tounequal distribution of aluminium and zinc in the microstructure on corrosionbehaviour is not understood properly. Further, in most of these studies, theinterpretation on the e�ect of several microconstituents is based more onelectrochemical data. The present investigation is concentrated on studying thee�ect of microconstituents on corrosion behaviour of AZ91D alloy prepared bydie-casting and ingot casting route by a combination of surface morphology andanalytical studies. This includes constant immersion experiments, in-situ corrosionanalysis, and surface morphology studies by optical and scanning electronmicroscopy and analytical studies by EDX and XRD. The electrochemicalparameters for the alloys were deduced by potentiodynamic polarizationexperiments.

R. Ambat et al. / Corrosion Science 42 (2000) 1433±14551434

2. Experimental method

2.1. Material

The AZ91D die-cast material (die-cast) was available in the cylindrical rod formwith a diameter of010 mm. Circular specimens of thickness 2 mm were sliced outfrom the rod. For ingot material (ingot), a specimen size of 1 cm � 1 cm � 2 mmwas used for all the experiments. The chemical composition of the alloy is given inTable 1.

2.2. Constant immersion testing

For constant immersion testing, the specimens were polished successively on®ner grades of emery papers up to 1000 level. All the specimens were initiallycleaned using the procedure of ASTM standard G-I-72 [6]. The polished andpreweighed specimens were exposed to the solution (150 ml, 3.5% NaCl) forvarious intervals of time. Final cleaning of the specimen at the end of theexperiment was done by dipping it in a solution of 15% CrO3 + 1% AgCrO4 in100 ml of water at boiling condition. An acetone washing followed this. Theweight loss was measured after each experiment and the corrosion rate wascalculated in millimeters per year. In each case duplicate experiments wereconducted and showed that the results fall within the expected error.

2.3. Electrochemical testing

Electrochemical polarisation experiments were carried out using an Auto LabCorrosion measurement system. Electrodes for this purpose were prepared byconnecting a wire to one side of the sample and covering with cold setting resin.The opposite surface of the specimen was exposed to the solution. The exposedarea was about 1 cm2. The specimens were given a metallographic polishing priorto each experiment, followed by washing with distilled water and acetone. Apolarisation test was carried out in a corrosion cell containing 150 ml of 3.5%NaCl using a standard three electrode con®guration: saturated colomel as areference with a platinum electrode as counter and the sample as the workingelectrode. Specimens were immersed in the test solution and a polarization scanwas carried out towards more noble values at a rate of 1 mV sÿ1, after allowing asteady state potential to develop.

Table 1

Chemical composition of the AZ91D alloy (in wt%)

Al Mn Ni Cu Zn Ca Si K Fe Mg

9.1 0.17 0.001 0.001 0.64 < 0.01 < 0.01 < 0.01 < 0.001 Bal

R. Ambat et al. / Corrosion Science 42 (2000) 1433±1455 1435

2.4. In-situ corrosion observation

Usually a humid atmosphere causes the formation of tiny water droplets on thesurface of materials. In marine atmospheres, these droplets contain a signi®cantamount of chloride ions. Several electrochemical cells nucleate and propagatewithin this tiny droplet. In the present investigation, a droplet of 3.5% NaCl wasput on the surface of the material and the corrosion phenomena within this tinydroplet was monitored continuously through a ZEISS model optical microscope.The features at the surface of the specimen were recorded as a function of time.The passivation behaviour was also studied by scratching technique.

2.5. Surface morphology Ð analytical work

The microstructure and surface morphology of the corroded surface werecharacterized using optical (ZEISS Model) and Scanning electron microscope(JEOL Model No. 5600LV). Concentration pro®ling of Al and Zn in the materialwas carried out using EDX facility (Oxford Model) attached to the scanningelectron microscope. X-ray di�raction technique (Philips, Model No. PW1830)was used for analysing corroded surface and corrosion products.

2.6. Corrosive media

A solution of 3.5% NaCl with a pH of 7.25 was used for all the experiment. Allthe experiments were conducted at room temperature. The solution was preparedusing A.R. grade NaCl in distilled water.

3. Experimental results

3.1. Microstructure

Fig. 1 shows the microstructure of the die-cast and ingot. The microstructureconsisted of primary a, eutectic a and b phase (marked in the ®gure). Die-castshowed much ®ner grains and b phases. A considerable amount of porosity couldbe observed on die-cast. On the other hand, ingot showed coarse grain structurewith bigger b particles.

3.2. Constant immersion studies

The variation of corrosion rate as a function of exposure time for AZ91D die-cast and ingot is given in Figs. 2 and 3. The ingot showed an initial decrease incorrosion rate followed by an increase upto 6 days of exposure. However, in thecase of die-cast, the corrosion rate showed a steady decrease with exposure timeexcept for 24 h exposure where it showed a slight upward trend. The corrosionrate for ingot was slightly higher than die-cast at all exposure times.


Figs. 2 and 3 also show the typical surface features of the corroded surface afterexposure. Corrosion started initially at localised sites, preferentially in both caseson primary a phase (more details given later). For ingot material this is clearlyvisible on the surface while these features are not so clear for die-cast probablydue to ®ne grain size. The localised attack invaded the entire surface withcontinued exposure to give a general corrosion pattern (after 6 days of exposure)with several b particles remained on the surface una�ected while the entire matrixwas dissolved. The degree of attack revealed the aggressiveness of the corrosivemedium.

Fig. 1. Microstructures of AZ91D alloy (a) ingot and (b) die-cast.


In Fig. 4 a comparison is made between the corrosion of ingot and die-castmaterials at three di�erent pH conditions. At 7.25 pH condition, the plot alsogives the di�erence in corrosion rate in distilled water and in 3.5% NaCl solution.The importance of these three pH conditions is as follows. At 2.0 pH, as it is

Fig. 2. Corrosion rate for AZ91D ingot as a function of exposure time.

Fig. 3. Corrosion rate for AZ91D die-cast as a function of exposure time.


evident from the Eh±pH diagram for magnesium and aluminium [7], theseelements are not stable. At 7.25 pH, magnesium is unstable (so that it cancorrode) but aluminium is passive and at 12.0 pH, aluminium is active whilemagnesium is passive. The corrosion rates for the materials in these threesolutions showed that the acidic solutions are very aggressive while the highlyalkaline solutions o�er better corrosion protection. An increase in pH from 2.0 to7.25 decreased the corrosion rate by 10 times and the rates were almost similar fordie-cast and ingot. The addition of 3.5% NaCl had increased the corrosion rateby about four times both for ingot and die-cast.

To understand the e�ect of microconstituents in greater detail, severalspecimens were immersed in 3.5% NaCl solution for a very short interval of time(starting from few seconds to 15 min) and observed under SEM. Selectedmicrographs are given in Fig. 5. For ingot, in general, the corrosion was initiatedon primary a phase while the eutectic a and b phases were completely protected,although few areas were observed with severe corrosion at the b interface leadingto undermining of the same. Die-cast showed (Fig. 5(d)) a cellular type of networkafter corrosion. It can be clearly seen that the corrosion was preferentially startedat the grain interior surrounded by b phase. Corrosion was found to be verysevere at these locations leading to deep pits in contrast to ingot where thecorrosion attack was more spread out.

Fig. 4. Corrosion rate for ingot and die-cast at three di�erent pH conditions.


3.3. In-situ corrosion observation

The anodic and cathodic activity under a droplet of 3.5% NaCl on the surfaceof the die-cast and ingot was recorded using an optical microscope. Theseexperiments were done on etched (so that the microstructural features can beclearly seen) and unetched surfaces. Figs. 6 and 7 show some of the sequentialpictures taken during in-situ observation of corrosion on etched surfaces. In allthe cases, as can be seen from the pictures, hydrogen evolution startedimmediately after the solution was dropped (see arrows in Figs. 6(b) and 7(b)).However, these nucleation sites were concentrated to a few places, rather than onentire b surface. As shown in Fig. 6, for ingot, at a few locations hydrogenevolution was so severe that the bubbles were coalesced and there was vigorousreaction underneath (see arrow in Fig. 6(c)). Similar sites were found on die-castthough the pictures shown (Fig. 7) did not have these. The activity was sustainedat these sites for a few minutes with the expulsion of corrosion product that was

Fig. 5. Scanning electron micrographs showing the e�ect of microconstituents on corrosion behaviour

of AZ91D ingot and die-cast material in 3.5% NaCl solution. (a), (b) and (c) ingot: (a) 5 min exposure,

(b) 15 min exposure, (c) 15 min exposure, back scattered image showing corrosion free areas near to bphase and (d) die-cast after 15 min, exposure; back scattered image showing b network and corrosion

with in the grain.


visible in the neighbouring areas (see arrows in Fig. 6(d) and (e)). On die-cast, theactivity at these sites was subsided much faster than that on ingot. It is alsoimportant to note that this type of vigorous hydrogen evolution did not happen atall hydrogen evolution sites. Larger pits were observed in these areas aftercleaning the surface (Fig. 6(f)). Fig. 7(f) shows corrosion at the water±air interface

Fig. 6. In-situ corrosion observation of AZ91D ingot: micro-corrosion features as a function of time.

(a) 0 min, (b) 3 min, (c) 10 min, (d) 22 min, (e) 30 min and (f) surface after cleaning.


for die-cast. These areas have undergone severe corrosion compared to the surfaceinside the drop. A similar behaviour was observed for ingot.

In general, the unetched surface was found to be less active compared to theetched one. For the unetched surface, the surface of the specimens tarnished withtime due to corrosion. For ingot, corrosion was initiated at primary a phase.

Fig. 7. In-situ corrosion observation of AZ91D die-cast: micro-corrosion features as a function of time.

(a) 0 min, (b) 3 min, (c) 10 min, (d) 30 min, (e) surface after cleaning and (f) corrosion at the

boundary.


In another experiment, after a speci®c interval of time, the surface of thespecimen under the droplet was scratched with a thin glass rod. This broke theoxide ®lm that was covering the surface of the alloy. For ingot, the repassivationof these areas took a long time while on die-cast; these scratches were repassivatedquickly.

3.4. Aluminium and zinc concentration pro®ling using EDX

It was suspected that the di�erence in corrosion behaviour with in the a phaseis due to the variation in aluminium content. In order to understand this in detailconcentration pro®ling of Al and Zn was carried out at several locations starting

Fig. 8. Concentration pro®les for Al and Zn along the line shown in the micrographs for AZ91D ingot.


from b phase to interior of the microstructure using EDX. Figs. 8 and 9 gives theconcentration pro®le of the elements along the lines shown in micrographs. It hasto be remembered that the EDX analysis gives the concentration of the elementnot exactly on the surface, but with in the interacting volume of the electronbeam. It is possible that, in such cases, interference from features below thesurface. To take care of this fact, a number of measurements were carried out atdi�erent locations (Fig. 8) and the interpretation is based on majority of the data.In general, for ingot, it can be seen that the b phase and the neighbouring areas(probably eutectic a phase) consisted of higher percentage of aluminium and zinc,and the concentration decreases as we move away from the b phase. The width ofthis aluminium rich region (>8%) adjacent to the b phase varied at di�erent sites.The concentration of aluminium typically varied between 035% at the b phase to08±6% near or with in the primary a phase. In highly corroded region of themicrostructure (probably primary a� found to have an aluminium concentrationR8±6% with negligible amount of zinc. In Fig. 8, curve (C ) represents theconcentration pro®le of aluminium between two b phases in a specimen exposedto corrosive environment for 5 min. Comparing the pro®le with the micrograph, itis clear that the corrosion was nucleated at areas where the concentration ofaluminium was less than 8%. Analysis of the EDX results led to the conclusionthat the composition of the b phase is 1 Mg17 Al11 Zn0.5. The concentrationpro®les for a non-corroded and corroded sample was similar.

In the case of die-cast (Fig. 9), the pattern was slightly di�erent due to the

Fig. 9. Concentration pro®les for Al and Zn along the lines shown in the micrographs for AZ91D die-

cast.


®neness of the microstructure. The EDX measurement was di�cult on thismaterial due to interference from bottom as the b phase was very thin. Adjacentto b phase, only a very thin layer was found with higher amount of aluminium.

3.5. X-ray di�raction studies

XRD patterns (Fig. 10 (a) and (b)) were taken at same area of the specimenbefore and after corrosion to get an idea about the amount of b phase removedfrom the surface due to corrosion. The corrosion products were removed beforesubjecting the specimen to XRD.

The surface of the unexposed alloys showed mainly peaks corresponding to Mgand b phase (Fig. 10(a)(i) and (b)(i)). After exposing to the corrosive media theintensity of the peaks corresponding b phase was lowered showing theundermining of many b phases (Fig. 10(a)(ii) and (b)(ii)). For ingot, the peaks forb phase was more intense compared to that of die-cast. One to one comparison ofintensities of all the peaks corresponding to b phase before and after corrosionrevealed that for ingot, the intensity has come down by an average of 8% whilefor die-cast the value was only about 2%. Keeping in mind that it was alloccurring at the surface and di�erence in the size of b phases, the amounts arequite high for both materials.

X-ray di�ractogram of the corrosion product (Fig. 11(a)) for ingot, extractedfrom the corrosive medium showed the presence of Mg(OH)2, b phase, MgH2 anda mixed oxide of magnesium and aluminium. The JCPDS card showed that theMgO:Al2O3 ratio of this mixed oxide is 1:2.5 [8].

However, it is interesting to note that the corrosion product analysis of die-castin the same corrosive media revealed an amorphous structure for corrosionproduct (Fig. 11(b)). Two broad peaks can be seen in the di�ractogram probablydue to Mg(OH)2 and the mixed oxide mentioned above.

3.6. Potentiodynamic polarisation

The potentiodynamic polarisation curves for AZ91D die-cast and ingot in 3.5%NaCl solution is shown in Fig. 12. Both materials did not show any passivity inthis solution. Table 1 shows the Ecorr, tafel slope and Icorr values for the materials.The die-cast is found to be slightly cathodic to ingot cast. Tafel slope is same forboth alloys whereas the Icorr is higher for die-cast compared to that of ingot. Thecathodic part of the curve reveals that the cathodic currents were much higher fordie-cast at all potentials. On the other hand, ingot was more active anodically.The surface of the specimen after polarisation studies showed surface featuressimilar to that in Fig. 5, except for ingot; some attack was found on eutectic aphase.


Fig. 10. X-ray di�ractogram of alloy surface: (a) ingot: (i) before corrosion and (ii) same area after

exposure to 3.5% NaCl for 6 days, (b) die-cast: (i) before corrosion and (ii) same area after exposure to

3.5% NaCl for 6 days.


Fig. 10 (continued)


Fig. 11. X-ray di�ractogram of corrosion product in 3.5% NaCl after exposure for 6 days: (i) ingot

and (ii) die-cast.


4. Discussion

The phase diagram for Mg±Al system is characterised by three microstructuralconstituents in the cast microstructure [9]. These are: primary a, eutectic a and bphase. These three microconstituents can be clearly seen in the microstructuregiven in Fig. 1(a). Macroscopically the a phase looks homogeneous, howevermicroscopically it is heterogeneous due to coring during solidi®cation. As a result,primary a contain much less aluminium compared to that of eutectic a and bphase. The width of this zone depends on solidi®cation rate [9]. The b phasereported to have more cathodic potential than that of a matrix [4]. That way inAZ91 alloys, the size and spatial distribution of b phase together with the coringmakes the alloy surface electrochemically heterogeneous. The corrosion behaviourof the material thus depends on how these microconstituents interact when it isexposed to aqueous environment. The above said parameters changes withprocessing route, gives rise to di�erent corrosion behaviours for materialsprepared by di�erent processing routes.

The results presented above reveal considerable in¯uence of this microstructuralheterogeneity especially the volume fraction and nature of b phase, and coring(leading to unequal distribution of aluminium and zinc) on corrosion behaviour ofcast AZ91D alloy. A detailed discussion of the results is given below.

Fig. 12. Potentiodynamic polarisation curves for AZ91D ingot and die-cast in 3.5% NaCl solution.


4.1. E�ect of corrosive environment

A solution of 3.5% NaCl was found to be an aggressive environment for thepresent alloys (Table 2). The drastic increase in corrosion rate (4 times) observedfor die-cast and ingot in 3.5% NaCl solution (Fig. 4) may be interpreted as due tothe participation of chloride ions in the dissolution reaction. Chloride ions areaggressive for both magnesium and aluminium. The adsorption of chloride ions tomagnesium surface, transforms Mg(OH)2 to easily soluble MgCl2 [10]. The criticalconcentration of chloride ions required for pitting in uninhibited NaCl wasreported to be about 0.002 to 0.02 M NaCl [10]. This value increases with increasein pH of the solution.

The e�ect of pH on the corrosion behaviour is in agreement with the Eh±pHdiagram of magnesium [7]. In highly alkaline solution, the corrosion rate for die-cast and ingot was found to be comparatively low due to the stability of Mg(OH)2®lm. The e�ect of these parameters was also found to depend on themicrostructure. A more detailed paper on this subject will be published separately.

4.2. E�ect of microconstitutents on corrosion behaviour

4.2.1. b phaseFor both materials, b phase showed maximum resistance to corrosion

environment. This phase did not undergo corrosion in 3.5% NaCl except that itwas undermined due to corrosion at the interface. However, the in¯uence of thisphase on the overall corrosion behaviour was di�erent in ingot and die-castbecause of the di�erence in morphology. The lower corrosion rate and betterpassivation of die-cast owes to ®ne distribution of b phase. Earlier investigationson the corrosion behaviour of AZ91 alloys proved that the b phase could play adual role in the dissolution behaviour [3]. It can act as either a galvanic cathodeor a kinetic barrier to dissolution [3]. Lunder et al. [4] reported that the corrosionpotential of b phase in 5% NaCl saturated with Mg(OH)2 is about 490 and 420mV cathodic to pure Mg and AZ91 alloy, respectively. Since this phase is highlycathodic to the a phase, hydrogen evolution preferentially takes place on b phaseto make it an e�ective cathode. This has been observed in the present studyduring in-situ corrosion observation. On the other hand if the a grains are ®neand b fraction is not too low, the oxide ®lm formed on b phase is nearlycontinuous [3]. It has been reported that [4,11] the ®lm formed on b phase is morestable because of the presence of higher amount of aluminium. The continuous

Table 2

Icorr, Ecorr and bc values for die-cast and ingot in 3.5% NaCl solution

Material Icorr (mA cmÿ2) Ecorr (mV) bc(mV)

Die-cast 0.0063 ÿ1461.0 169

Ingot 0.011 ÿ1441.5 182


stable oxide ®lm acts as barrier so that the dissolution of a phase is inhibited. Forthis to occur the interparticle distance between b phase should not be too large. Indie-cast this was possible due to the close net work of b phase. On the other hand,in ingot, the interparticle distance was very high so that such a networking wasimpossible. In contrast, the undermining of b phase occurred due to the corrosionof a phase (more detailed discussion given later) leads to reduction of b fractionon the surface which in turn causes an increase in corrosion rate. This is exactlywhat was observed in the case of ingot above an exposure time of 6 h (Fig. 2).The XRD data (Fig. 10(a)) also showed that the undermining of b phase from thesurface of this material was quite high. On the other hand on die-cast, bundermining was found to be comparatively low (Fig. 10(b)), the fraction of b onsurface increased with time, ®nally leading to a network of b phases that reducesthe corrosion rate with exposure time. Fig. 5(d) shows the b network in die-cast inand around the corroded area.

In-situ corrosion studies further con®rm this fact where the corrosion sites weresubsided much faster on die-cast compared to that on ingot. Scratching of thesurface also showed better passivation of die-cast. The reason for this may besame as that explained before.

Corrosion product collected from the solution for ingot showed presence of bphase (Fig. 11(a)). As the XRD analysis was carried out on corrosion productcollected after 6 days of exposure, the result indicated the stability of b phase inaqueous solutions. b phase is reported to be stable in a wide pH range (4±14) [4].The corrosion potential of this phase is close to pure Al in neutral and mildlyacidic solutions where as in highly acidic conditions, it approaches to the value forMg [4]. Corrosion properties of this phase are reported [4] to be superior to eitherof its components in mildly alkaline environments.

4.2.2. a phase: e�ect of coringThe morphology of the corroded surface given in Fig. 5 together with EDX

analysis given in Figs. 8 and 9 convincingly proved that apart from the presence bphase, coring play a bigger role in corrosion behaviour. This is especially true inthe case of ingot due to a wider width of zone with higher aluminium content(Fig. 8). In this material, corrosion was nucleated at regions with aluminiumcontent R8±6%. The e�ect of aluminium on the corrosion behaviour of Mgalloys has also been subject of many investigations [3,5,11±13]. It has beengenerally agreed that the presence of aluminium was bene®cial in improving thecorrosion behaviour of magnesium. Later, it was found that for e�ectiveprotection an optimum level of aluminium was required. Lunder et al. [4] reportedthat this quantity was about 8% where as Warner et al. [12] using TEM studieson rapidly solidi®ed ribbons of composition Mg-9%Al, arrived at the conclusionthat it is >5%. Hehmann et al. [13] studied corrosion behaviour of rapidlysolidi®ed splates and ribbons with aluminium concentration varying between 9and 62.3 wt%. They observed a decrease in the corrosion rate and Ecorr, from 9.6to 23.4 wt% Al in aerated 0.01 M sodium chloride solution. But none of these


studies reported the e�ect of the concentration di�erence of aluminium in themicrostructure near to the b phase on corrosion behaviour of Mg±Al alloys.

The present study shows that the concentration of aluminium decreases (Fig. 8)as we go away from the b phase, to a value of06% with in the primary a matrix.As explained earlier in the section on results, for ingot the variation was found tobe in the range of035% on b surface to06% in the primary a phase. Comparingthis concentration pro®le with the corrosion morphology (Fig. 5), it is clear thatthe region preferentially corroded was the one with less than 8% aluminium. Thevalue is in agreement with the value reported by Lunder et al. [4]. This shows thatwhether it is in a single-phase alloy or multiphase alloy, the amount of aluminiumneeded for considerable protection is more than 8%. In multiphase alloys,variation of concentration of aluminium with in the microstructure leads tocorrosion of regions where aluminium content is less than 8%.

An analysis of the aluminium concentration in ingot near the b phase at variouslocations led to the conclusion that (Fig. 8) the width of the region with higheramount of aluminium (>8%) varies from point to point. As shown in Fig. 8, itcan be as high as 80±100 mm to few microns. For ingot, at several locationscorrosion was initiated just adjacent to b boundary ®nally leading undermining ofthis phase. It is believed that, these are the regions where aluminiumconcentration just adjacent to b phase was <8%. However, it is important to

Fig. 13. Ecorr values for Mg±Al alloys as a function of Al and Zn concentration. (Values from Refs. [4]

and [11]: Lunder et al.: in 5% NaCl saturated with Mg(OH)2 and Daloz et al.: in ASTM D1384 water).


note that with exposure time or at lower pH conditions [14], the regions withaluminium content >8% was also corroded.

In the case of die-cast a similar feature can be seen on careful observation(Fig. 9). As explained earlier, concentration pro®ling was di�clt due tointerference. The width of the region surrounding this phase with higheraluminium content was found to be comparatively small. Fig. 5(d) revealed thatthe corrosion was preferentially started within the grain surrounded by the bphase ®nally leading to a cellular structure with a network of undissolved b phase.

Fig. 13 shows the Ecorr values for Mg±Al alloys as a function of Al and Zncontent reported in the literature. Increasing aluminium content in the materialshifted the potential towards anodic side. A comparison of EDX result (Fig. 8)with Fig. 13 reveals that a variation in the electrochemical potential could beexpected with in the a phase with regions of higher Al content showing higheranodic potentials. The corrosion results showed an opposite trend. The reason forthis behaviour is not clear. However, as shown in Fig. 13, the presence of Zn mayalter the situation. In the present study aluminium rich regions found to have ahigher amount of zinc (Figs. 8 and 9). Daloz et al. [11] reported that in Mg±Alalloys, presence of Zn in¯uences the Ecorr of both matrix and precipitate and hada bene®cial e�ects on the corrosion resistance of the alloys.

4.3. Corrosion behaviour of ingot cast vs. die-cast

Overall comparison of the results reveal that, die-cast was marginally lesscorrosive than ingot. However, a comparison of the corrosion features inFig. 5(a)±(d) reveal deep pitting on die-cast. It was found that these areas (graininterior) contain comparatively less quantity of aluminium (Fig. 9). In the case ofingot, the corrosion was spread out in the comparatively large primary a phase.The eutectic a phase in die-cast was more corrosion resistant due to its higheraluminium content even though it was very narrow compared to that in ingot.The lower aluminium content of eutectic a phase in ingot compared to die-castmakes it susceptible to corrosion at longer exposure times that led to underminingof b phase which was found to be high for ingot material from XRD data.

Electrochemically die-cast was found to be more cathodic due to the presence ofhigher area fraction of b at the surface (Fig. 12). The potential of an alloy isdetermined by the potential of the constituent phases and the area fractioncovered by each phase [15]. As explained earlier, the b phase is highly cathodic toa phase in Mg alloys, so that the presence of a higher fraction of b phase on thesurface shifts the potential towards more cathodic side. The more cathodic currentdensity at all potentials may also be due to the same factor. The anodic andcathodic tafel slope remained the same for both materials, showing that theelectrochemical reactions are the same.

XRD data showed that the quantity of b phase removed from the surface dueto corrosion was high in ingot (8%) (Fig. 10(a) and (b)). However, the 2%reduction in b phase on the surface for die-cast can be substantial considering thesize of b phase in this material. The nature of corrosion products was found to be


di�erent for ingot and die-cast (Fig. 11(a) and (b)). For ingot, the compoundsidenti®ed are fairly in agreement with that reported in the literature. Earlierstudies have reported the presence of Mg(OH)2 and MgH2 in the corrosionproduct of Mg and Mg alloys [16±20] even though none of the studies havereported the presence of a mixed oxide of aluminium and magnesium, and thepresence of b phase. The reason for the amorphous structure for product for die-cast is not clear. This needs further investigation. However our work on theanalysis of corrosion product formed in di�erent environments points to the factthat the presence of more aluminium oxide in the product leads to the abovementioned result [14].

As mentioned earlier, on both materials at a few sites hydrogen evolution wasvery vigorous and severe corrosion was found underneath. This may be thelocation where AlFeMn phases may be found. Lunder et al. [4] also reported asimilar phenomena in AZ91 alloy in 5% NaCl. The AlFeMn phases act as a goodcathode for hydrogen evolution.

In summary, the corrosion behaviour of cast AZ91 alloy depends signi®cantlyon the size and distribution of b phase and distribution of aluminium and zincwith in the microstructure. Although a close network of b phase provides abarrier to dissolution, an optimum level of aluminium and zinc is essential toensure more corrosion resistance for the a phase. Further, a higher aluminiumcontent adjacent to b phase reduces the undermining of the same that help inimproving the corrosion resistance.

5. Conclusions

1. AZ91D die-cast showed higher corrosion resistance and better passivation thanAZ91D ingot due to ®ne grain structure and b phase. Electrochemically die-castwas found to be more cathodic to ingot.

2. Size and morphology of b phase and coring were found to have signi®cantin¯uence on corrosion behavior of AZ91D alloy. In both materials, corrosioninitiated at areas where aluminium concentration was less than 8%. Hydrogenevolution took place preferentially on b phase but only at few locations.

3. XRD pattern showed a higher percentage of b undermining on ingot. XRDdata also showed the presence of undermined b phase in corrosion product (foringot) in addition to Mg(OH)2 and small quantities of magnesium-aluminiumoxide and MgH2. The corrosion product for die-cast was found to be moreamorphous.

Acknowledgements

The authors would like to thank Dr. Qiu Jian Hai of NTU for his help in


conducting the electrochemical experiments and Drs. B.H. Hu and Ian Pinwill ofGintic for providing the magnesium alloys. Financial assistance from NanyangTechnological University through research grant RG79/98 is gratefullyacknowledged.

References

[1] D. Magers, Magnesium Alloys and their Applications, in: B.L. Mordike, K.U. Kainer (Eds.),

Proc. of the Conf. on Magnesium Alloys and their Applications, Werksto�-

infomationsgesellschaft, 1998, p. 105.

[2] S. Schumann, F. Friedrich, in Ref. (D. Magers, Magnesium Alloys and their Applications, in:

B.L. Mordike, K.U. Kainer (Eds.), Proc. of the Conf. on Magnesium Alloys and their

Applications, Werksto�-infomationsgesellschaft, 1998, p. 3.

[3] Guangling Song, Andrej Atrens, Matthew Dargusch, Corros. Sci 41 (1999) 249.

[4] O. Lunder, J.E. Lein, T. Kr. Aune, K. Nisancioglu, Corrosion 45 (1989) 741.

[5] O. Khaselev, J. Yahalom, Corros. Sci 40 (1998) 1149.

[6] Annual Book of ASTM Standards Parts 3 and 4, 1977, p. 722.

[7] M. Pourbaix, Atlas of Electrochemical Equilibrium Diagrams in Aqueous Solutions, NACE,

Houston, TX, 1966.

[8] JCPDS File, V.1.30, 1997, Card No. 10-0238.

[9] I.J. Polmear, in: Light Alloys: Metallurgy of the Light Metals, 2nd ed, Edward Arnold, New

York, 1989, p. 183.

[10] Heike Bommer, in Ref. (D. Magers, Magnesium Alloys and their Applications, in: B.L. Mordike,

K.U. Kainer (Eds.), Proc. of the Conf. on Magnesium Alloys and their Applications, Werksto�-

infomationsgesellschaft, 1998, p. 79.

[11] D. Daloz, P. Steinmetz, G. Michot, Corrosion 53 (1997) 944.

[12] T.J. Warner, N.A. Thorne, G. Nussbaum, W.M. Stobbs, Surf. Interf. Anal 19 (1992) 386.

[13] F. Hehmann, R.G.J. Edyvean, H. Jones, F. Sommer, in: Proc. of the Int. Conf. on Powder

Metallurgy Aerospace Materials, Lucerne, Switzerland, November 2±4, 1987, p. 46 (Shrewsbury,

UK; Metal Powder Report Publishing Services, 1987).

[14] Rajan Ambat, Naing Naing Aung and W. Zhou, Nanyang, J. Appl. Electrochem. (2000) in press.

[15] F. Mansfeld, Corrosion 27 (1971) 436.

[16] J.D. Hanawalt, C.E. Nelson, J.A. Peloubet, Trans AIME 147 (1942) 273.

[17] K. Huber, J. Electrochem, Soc 100 (1953) 376.

[18] O. Fruhwirth, G.W. Herzog, I. Hollerer, A. Rachetti, Surf. Technol 24 (1985) 301.

[19] S. Yamaguchi, J. Appl. Phys 25 (1954) 1437.

[20] S. Diplas, P. Tsakiropoulos, R.M.D. Brydson, J.F. Watts, Mat. Sci. Tech 14 (1998) 699.


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Evaluation of microstructural e•ects on corrosion ... 2000 - Corrosion Scie… · passivation compared with the ingot. In die-cast and ingot, hydrogen evolution took place preferentially