Influence of Homogenization Annealing of AZ91 on Mechanical Properties and Corrosion Behavior

DOI: 10.1002/adem.200700234

Influence of Homogenization Annealing of AZ91 on MechanicalProperties and Corrosion Behavior**By Ming-Chun Zhao, Ming Liu, Guang-Ling Song and Andrej Atrens*

Magnesium alloys are a class of structural materials withincreasing industrial interest in automobile service due totheir good strength to weight ratio and low density,[1–3] al-though corrosion is an issue.[4–7] AZ91 is one of the most pop-ular magnesium alloys, with the nominal composition Mg-9 % Al-1 % Zn (wt-%). The microstructure of as-cast AZ91typically consists of a matrix of a-Mg grains and a divorcedeutectic distributed along the a-Mg grain boundaries;[8,9] thedivorced eutectic consists of secondary a-Mg plus the b phase(the intermetallic Mg17Al12). For slow cooling the secondarya-Mg can precipitate b phase. For commercial utilization, themechanical properties and corrosion behavior are importantand it is well established that the mechanical properties andthe corrosion behavior of Mg alloys can be significantly influ-enced by heat treatment.[10–14] Prior heat treatments aimed atdeveloping improved mechanical properties have consistedof a solution heat treatment followed by an ageing heat treat-ment;[10,11,13,14] hereafter, named solution heat treatment plusageing (STA). STA is also designated as the T6 temper (solu-tion heat treatment plus artificially aged).[15] Recently, it wasreported[16,17] that a homogenization annealing (HA) heattreatment improved the machinability of the Mg alloysAZ31B and ZK60. HA is different from STA. HA only needsreheating up to a temperature for several hours followed byair cooling rather than the two heat treatment steps of STA,which has (i) first a solution heat treatment at a higher tem-perature followed by quick cooling and (ii) second an ageingheat treatment at a lower temperature. As a consequence, HAhas potentially a cost lower than STA. However, the mecha-

nism by which HA changes the mechanical properties is notknown nor is the influence of microstructure variation on cor-rosion behavior. Furthermore, HA has not been attempted forAZ91. This research investigates the influence of HA of AZ91on mechanical properties and corrosion behavior.

Experimental Procedure

Specimens, cut from an as-cast ingot of AZ91 with nominalcomposition Mg – 9 wt-% Al – 1 wt-% Zn, were subjected toHA at 380 °C and 410 °C for times from 1 h to 25 h. For com-parison a specimen was also subjected to a HA of 25 h at410 °C followed by an ageing treatment of 25 h at 200 °C. Themicrostructure of the samples was examined by scanningelectron microscopy (SEM) after metallographic preparation:mechanical grinding, diamond polishing and etching in a 3 %nital solution. The hardness was measured using the Vickershardness (HV) test with a load of 50 M. Tensile tests usingcylindrical tensile specimens with a gage diameter of 5 mmand a gage length of 20 mm were carried out with a cross-head speed of 0.6 mm/min at room temperature. Fracturesurfaces of the tensile specimens were examined by SEM.

Corrosion behavior was characterized by immersion tests1 M NaCl solution at room temperature. Coupons were en-capsulated in epoxy resin so that a surface with the dimen-sion of 18 mm × 27 mm was exposed to the solution. The sur-face was ground using 240, 400, 600 and 1,200 grit SiC paper,washed with distilled water and dried with warm flowingair. The working surface was horizontally immersed in1,500 ml of 1 M NaCl, which was made with analytical gradereagent and distilled water. The corrosion rate was evalu-ated[4–6,18,19] by measuring the volume of evolved hydrogen,collected in a burette above the corroding sample. The overallmagnesium corrosion reaction, Equation 1, indicates that onemolecule of hydrogen is evolved for each atom of corrodedmagnesium.

2Mg � 2H� � 2H2O � 2Mg2� � 2OH� � 2H2 �1�

The corrosion morphology after immersion for 1 h was ex-amined using optical microscopy. The same surface prepara-tion was used as was used in the corrosion exposures withhydrogen collection, with the necessary consequence that thequality of the optical micrographs could not be expected toequal that of a metallographic prepared sample. To obtaingood quality micrographs the corrosion exposures wererepeated with samples that were metallographic polished but

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–[*] Dr. M.-C. Zhao,+ M. Liu,+ Dr. G.-L. Song,

Prof. A. Atrens+

Division of MaterialsThe University of QueenslandBrisbane, Qld 4072, AustraliaE-mail: [email protected]

[+] also with:Swiss Federal Laboratories for Materials Science andTechnologyEMPA, Dept 136, Überlandstrasse 129CH-8600 Dubendorf, Switzerland

[**] This work was supported by the ARC Center of Excellence,Design of Light Alloys. Atrens, Zhao and Liu would like tothank EMPA for support that allowed them to spend time atEMPA.

were not etched. The immersion time was 18 h, essentiallythe same as for the samples polished to 1,200 SiC grit. Thecorrosion morphologies for these metallographic polishedsamples were essentially the same as for the samples sub-jected to only grinding down to 1,200 SiC paper and thesecorrosion morphologies are presented in Section 3.5.

Results and Discussion

Microstructure

Figure 1 shows the microstructure evolution of AZ91caused by HA. The as-cast microstructure consisted of theprimary a-Mg matrix with the dendrite boundaries decoratednet-like by the continuous eutectic micro-constituent consist-ing of fine (a+b) in a lamellar arrangement plus coarse b parti-cles. The fine lamellar (a+b) arrangement was nearly alwaysassociated with coarse b particles (Fig. 1(a) and 1(b)) althoughthere are instances in Figure 1(a) where (i) there was the finelamellar (a+b) arrangement apparently not associated withany b particle, and (ii) there were b particles with apparentlyno associated fine lamellar (a+b) arrangement. The formation

of the fine lamellar (a+b) arrangement is beyond the scope ofthe present paper. However, Figure 1(a) and 1(b) does pro-vide instances where it appears that the fine lamellar (a+b)arrangement has nucleated some distance away from thelarge b particle and has grown towards the b particle. Thereare cases where it appears that there is a thin sheet of b sur-rounding the fine lamellar (a+b) arrangement and this thinsheet of b might have caused the nucleation of the fine lamel-lar (a+b) arrangement in a mechanism that might have somesimilarities to the nucleation of pearlite in steels.[20,21]

After HA, the volume fraction of the second b-phase parti-cles decreased and their distribution became more homoge-neous. Figure 1(c) shows that, after HA for 5 h at 380 °C, thefine (a+b) in the lamellar arrangement had dissolved and themicrostructure consisted of the a-Mg matrix and isolatedcoarse b particles. Figure 1(d) shows that, after HA for 10 h at410 °C, the coarse b particles had partly dissolved, leaving rel-atively fine b particles homogeneously distributed through-out the matrix. This indicates that a higher temperature and/or a longer time produced a more homogeneous distributionof isolated b particles throughout the microstructure.

Zhao et al./Influence of Homogenization Annealing of AZ91 on Mechanical Properties

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(a) (b)

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Fig. 1. SEM micrographs of the microstructure of AZ91 (a) and (b) as-cast, (c) after HA for 5 h at 380 °C and (d) after HA for 10 h at 410 °C.

Higher magnification micrographs, presented in Fig-ures 2(a) and 2(b), indicated that there were few small(∼ 100 nm) b precipitates in the a-Mg matrix after HA for10 h at 380 °C and 10 h at 410 °C, indicating that significantprecipitation of b did not occur from the a-Mg matrix duringHA. In marked contrast, Figure 2(c) shows that there were alarge number of fine b precipitates with size ∼ 100 nm distrib-uted throughout the matrix of AZ91 subjected to HA for 25 hat 410 °C followed by an ageing heat treatment of 25 h at200 °C.

Mechanical Properties

Figure 3(a) presents engineering stress-strain curves mea-sured at room temperature for AZ91 as-cast, after HA for 5 hfor 380 °C (designated 3805) and 10 h for 410 °C (designated4110) and for AZ91 subjected to HA of 25 h at 410 °C followedby an ageing heat treatment of 25 h at 200 °C (designated412025); these curves showed no sharp yield point but contin-uous yielding. Figure 3(b) presents the dependence of the


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(a) 380ºC-10h (b) 410ºC-10h

(c) 410ºC-25h+200ºC–25h

Fig. 2. SEM micrographs of the microstructure of AZ91 (a) after HA for 10 h at380 °C and (b) after HA for 10 h at 410 °C, and (c) after HA for 25 h at 410 °C plusaging for 25 at 200 °C.

a)

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Fig. 3. (a) engineering stress-strain curves measured at room temperature for AZ91 as-cast, after HA for 5 h for 380 °C (designated 3805) and HA 10 h for 410 °C (designated4110) and for AZ91 subjected to HA of 25 h at 410 °C followed by an ageing heat treat-ment of 25 h at 200 °C (designated 412025). (b) the dependence of the mechanical prop-erties {Vickers hardness (HV), yield strength (YS), ultimate tensile strength (UTS)and elongation to failure (EL)} on HA time (t) for HA at 380 °C and 410 °C.

mechanical properties {Vickers hardness (HV), yield strength(YS), ultimate tensile strength (UTS) and elongation to failure(EL)} on HA time (t) for HA at 380 °C and 410 °C. For eachtemperature, the hardness increased with increasing time upto a maximum and then slightly decreased for longer times.Peak-hardness occurred at 15 h for 380 °C and 10 h for 410 °C.The ultimate tensile strength also showed a similar trend, i.e.the ultimate tensile strength increased with increasing timeup to a maximum and then slightly decreased for longertimes. The yield strength did not change much, but wasslightly higher than that for the as-cast condition. The elonga-tion to failure was higher than that in the as-cast condition.The results showed that the mechanical properties of AZ91were improved by HA. Especially interesting were the peakstrength conditions (15 h for 380 °C or 10 h for 410 °C) forwhich the yield strength was increased by about 5 %, the ulti-mate tensile strength by about 60 %, and the elongation tofailure by about 100 % compared to the as-cast condition.

HA changed the mechanical properties of AZ91 by chang-ing the microstructure. The as-cast microstructure is charac-terized by the eutectic micro-constituent continuously distrib-uted along the boundaries of a-Mg matrix, i.e. a continuousfine lamellar a+b plus coarse b particles arranged net-likealong and adjacent to the a-Mg grain boundaries (Fig. 1(a)).This means that there was a continuous easy crack path of thebrittle b-phase along the a-Mg grain boundaries. The b-phasecan easily crack during plastic deformation. As a consequencethere was a resulting low elongation to failure and concomi-tantly a low strength: an elongation to failure of 3 % and anultimate tensile strength of only 125 MPa (Figs. 3(a) and3(b)). Fractography after the tensile tests revealed a brittletransgranular fracture mode (Fig. 4(a)). This characteristicpattern is formed when fracture reinitiates in a different ori-entation at the grain boundary net-like decorated by the con-tinuous fine a+b in the lamellar arrangement plus coarse bparticles.

HA produced microstructures consisting of the a-Mg ma-trix with isolated b particles relatively homogeneously dis-tributed throughout the a-Mg matrix (Figs. 1(c) and 1(d)).The better ductility compared with the as-cast microstructureis directly attributable to the absence of the continuous brittleb-phase at the a-Mg grain boundaries, and moreover thehigher ductility correlated with the higher ultimate tensilestrength as is evident from the tensile stress-strain curves(Fig. 3(a)). The ductility and UTS after HA for 10 h at 410 °Cwere greater than after 5 h at 380 °C; this correlated with thesmaller and more isolated b-phase particles after HA for 10 hat 410 °C. The fractography after the tensile tests of the sam-ple HA for 10 h at 410 °C (Fig. 4(b)) revealed transgranularfracture accompanied with some ductile dimples congruentwith the higher ductility compared with the as-cast micro-structure (Figs. 3(a) and 3(b)). The increase in the hardnessand the ultimate tensile strength and the small increase inyield strength (Fig. 3(b)) are attributed to solid solutionstrengthening due to an increase in Al content in the a-Mg

matrix that is caused by the dissolution of the fine b particlesin the fine eutectic a+b lamellar arrangement and some of thecoarse b particles. There is not much precipitation hardeningproduced by HA because, as indicated by Figures 2(a) and2(b), HA produced few relatively fine (∼ 100 nm) b precipi-tates in the a-Mg matrix, which would be sufficiently small toblock the movement of dislocations by mechanisms such asOrowan looping.[22,23] Figure 3 also includes data for HA for25 h at 410 °C followed by ageing for 25 h at 200 °C. These ex-periments were carried out in order to clarify the differencebetween HA and STA. The ageing heat treatment caused pre-cipitation of a fine distribution of a large number of fine bprecipitates with size of ∼ 100 nm distributed throughout thematrix, Figure 2(c) which caused an obvious increase instrengthening associated with a significant increase in yieldstrength and at the same time limited the ductility. This rein-forces the above conclusions that the main influences of HAare (i) an increase in the Al concentration in the a-Mg matrix,(ii) no continuous crack path through the matrix (due to thedissolution of the fine b particles in the fine eutectic a+blamellar arrangement and some of the coarse b particles) and(iii) a microstructure consisting of a relatively homogeneousdistribution of isolated b particles in the a-Mg matrix.

Corrosion Behavior

Corrosion behavior was evaluated for the following micro-structure conditions: (i) as-cast, (ii) peak strength (HA for10 h at 410 °C) and (iii) moderate strength (HA for 5 h at380 °C). These three conditions represent three types of micro-structures: (i) a + eutectic (coarse b + fine a + b in the lamellararrangement), (ii) a + coarse b and (iii) a + fine b (Fig. 1). Fig-ure 5 presents the corrosion behavior as characterized by thehydrogen evolution volumes for these three microstructures.The moderate strength condition (HA for 5 h at 380 °C) gavethe highest hydrogen evolution volume, whereas the peakstrength condition (HA for 10 h at 410 °C) and the as-cast con-dition gave similar and lower hydrogen evolution volumes.This means that it is possible to improve both the mechanicalproperties and the corrosion properties of AZ91 throughappropriate HA, i.e. HA for 10 h at 410 °C.

The corrosion behavior can also be related to the micro-structure. It appears that the eutectic micro-constituent of finea+b in a lamellar arrangement is beneficial to corrosion. Thisis proposed because the as-cast AZ91, with a microstructurecontaining the fine a+b in a lamellar arrangement, had a hy-drogen evolution rate lower than that the moderate strengthAZ91 (HA for 5 h at 380 °C) without the lamellar eutectic a+bin the microstructure. The existence of an Al rich a-phase inAZ91 and the different Al contents in the a-Mg matrix, thea+b in the fine lamellar arrangement and the b-phase particleshas been widely discussed.[4,5,24–27] The a+b in the fine lamel-lar arrangement may behave as an independent electro-chem-ical entity. However, this micro-constituent as a separate elec-tro-chemical entity has not been previously presented as a


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Fig. 4. (a) Fractography of the as-cast tensile tests sample revealed a brittle transgranular fracture mode. (b) The fractography after thetensile tests of the sample HA for 10 h at 410 °C revealed transgranular fracture accompanied with some ductile dimples congruent withthe higher ductility compared with the as-cast microstructure (Figs. 3(a) and 3(b)).

possible mechanism, which may influence the corrosion ofAZ91.

Corrosion Morphology

The influence of the microstructure on the corrosion be-havior was studied by observing the micro-corrosionmorphologies of a set of samples with the above three micro-structures exposed for different time periods to the test solu-tion. After a few minutes immersion in the test solution, somesmall hydrogen bubbles were evolved from all over thewhole working surface. After a longer immersion time (1 h),several sites suffered corrosion in the form of localized corro-sion. The corrosion in these local areas enlarged in areaslowly and a large number of hydrogen bubbles wereevolved from these local areas of corrosion. Figure 6 presentsthe micro-corrosion morphology of each microstructure con-dition after immersion in the test solution for 1 h.

In the as-cast condition, the corrosion localized in the a-Mgmatrix adjacent to coarse b particles which were not sur-rounded by the a+b in the fine lamellar arrangement(Fig. 6(a)). This indicates that corrosion initiated preferential-ly, and that the corrosion did not initiate in the vicinity of thea+b in the fine lamellar arrangement.

In the moderate strength condition (HA for 5 h at 380 °C)and the peak strength condition (HA for 10 h at 410 °C), pref-erential corrosion also occurred and localized in the a-Mg ma-trix adjacent to b particles (Figs. 6(b) and 6(c)), although thedistribution and size of the b particles were much different inthese two microstructure conditions (isolated coarse b parti-cles in the moderate strength condition and isolated fine bparticles in the peak strength condition).

Figure 7 presents the micro-corrosion morphology of eachmicrostructure condition after immersion in the test solutionfor 18 h. In the as-cast condition, optical microscope examina-tion of the samples exposed for 16 h to 1 M NaCl solution

after grinding with 1200 SiC grit led to the conclusion that thea+b in the fine lamellar arrangement and the coarse b parti-cles acted as corrosion barriers; the adjacent a-Mg matrix hadundergone the majority of the corrosion, with no significantcorrosion of the a+b in the fine lamellar arrangement and thecoarse b particles. Optical microscope examination of the me-tallographic polished samples exposed for 18 h to 1 M NaClare presented in Figure 7(a). Figure 7(a) presents the follow-ing three different cases: (i) corrosion of the adjacent a-Mgwith no significant corrosion of the a+b in the fine lamellararrangement and the coarse b particles (ii) corrosion of thea-Mg up to the a+b in the fine lamellar arrangement, with theappearance of the corrosion of the a-Mg stopped at the inter-face of the a+b in the fine lamellar arrangement, and (iii) cor-rosion proceeding into the a+b in the fine lamellar arrange-ment, with both phases having been corroded to a similarextent.

For the moderate strength condition (HA for 5 h at 380 °C)and the peak strength condition (HA for 10 h at 410 °C), someisolated b particles were in the forefront of the corrosion ad-vance (Figs. 7(b) and 7(c)). However, the distance betweenthe b particles in these two microstructures was large, and thecorrosion advance is not expected to be effectively blocked bythe b phase, although the b particles were themselves notsignificantly corroded.

Corrosion Mechanism

AZ91 is a multiphase alloy, in which different micro-con-stituents may form micro-galvanic cells. Therefore the corro-sion behavior and morphology can be explained on the basisof the microstructures and the corresponding corrosion prop-erties of the a-Mg matrix, the a+b in the fine lamellar arrange-ment, and the coarse and the fine b particles. When the sam-ples were immersed in the test solution, intensive hydrogenevolution preferentially takes place at certain sites and thecorrosion gradually advances.

In the as-cast condition, Figures 6(a) and 7(a) show thatcorrosion initiated preferentially in the a-Mg matrix adjacentto coarse b particles which are not surrounded by the a+b inthe fine lamellar arrangement. Subsequently the corrosionoccurred in the a-Mg matrix, with essentially no corrosion inthe b particles and only some in the a+b in the fine lamellararrangement. The implication is that the corrosion initiatesand gradually advances by the dissolution of the a-Mg matrixadjacent to the a+b in the fine lamellar arrangement and the bparticles. When the propagation of the corrosion reaches thea+b in the fine lamellar arrangement, the corrosion is retardedto some extent. The implication is that the a+b in the finelamellar arrangement also retards the corrosion of the a-Mgmatrix in the depth direction. A schematic illustration for thiscase is depicted in Figure 8(a).

In the moderate strength condition (HA for 5 h at 380 °C)and the peak strength condition (HA for 10 h at 410 °C), iso-lated b particles are distributed throughout the microstruc-ture and the distance between the isolated b particles is large,


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Fig. 5. Corrosion behavior, measured in terms of the hydrogen evolution, for the threemicrostructure conditions: as-cast, peak strength (HA for 10 h at 410 °C) and moderatestrength (HA for 5 h at 380 °C).


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Fig. 6. The micro-corrosion morphology of each microstructure condition after immersion in 1 M NaCl solution for 1 h: (a) as-cast, (b) moderatestrength (HA for 5 h at 380 °C) and (c) peak strength (HA for 10 h at 410 °C). (a) corrosion initiated at isolated b particles as indicated by redarrows. (b) & (c) corrosion initiated at isolated b particles.

so that the corrosion attack is not effectively blocked by thistype of isolated b particles. A schematic illustration for thiscase is depicted in Figure 8(b). This case is different from thatof the as-cast microstructure.

In addition, the corrosion also involves a contribution ofmicro-galvanic corrosion between the different micro-constit-

uents in the microstructure because they have different elec-trochemical properties.[5,9] Micro-galvanic corrosion in addi-tion to spontaneous corrosion can occur between (i) the a-Mgmatrix and the a+b in the fine lamellar arrangement plus thecoarse b particles in the as-cast condition and (ii) the a-Mgmatrix and the isolated coarse b particles in the moderate


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Fig. 7. The micro-corrosion morphology of each microstructure condition after immersion in 1 M NaCl solution for 18 h: (a) as-cast, (b) moderate strength (HA for5 h at 380 °C) and (c) peak strength (HA for 10 h at 410 °C). (a) In the as-cast condition, the a+b in the fine lamellar arrangement and the coarse b particles acted tosome extent as corrosion barriers. (b) and (c) the distance between the b particles was large, and the corrosion advance is not expected to be effectively blocked by the bphase, although the b particles were themselves not significantly corroded.

strength condition (HA for 5 h at 380 °C) and (iii) the a-Mgmatrix and the fine b particles in the peak strength condition(HA for 10 h at 410 °C). Therefore, the a+b in the fine lamellararrangement and the b particles may also act as galvanic cath-odes and accelerate the corrosion of the a-Mg matrix in addi-tion to acting as an anodic barrier against the overall corro-sion propagation as mentioned above. The a+b in the finelamellar arrangement and the b particles as electro-chemicaland microstructure entities are involved in these two roles,which can be described as micro-galvanic corrosion and cor-rosion barrier.

Therefore, the corrosion behavior of the three microstruc-ture conditions, i.e. (i) the as-cast condition, (ii) the peakstrength condition (HA for 10 h at 410 °C) and (iii) the moder-

ate strength condition (HA for 5 h at 380 °C), may be summa-rized as follows. For the peak strength condition (HA for 10 hat 410 °C) isolated fine b particles cannot effectively block thecorrosion propagation because of the large distance betweenthe isolated fine b particles; the mechanism is schematically il-lustrated in Figure 8(b) and the micro-corrosion morphologyis presented in Figure 7(c). Furthermore, galvanic corrosionprinciples[28–32] indicate that a small cathode (the isolated fineb particles) causes negligible galvanic corrosion whencoupled with a large anode (the a-Mg matrix). As a conse-quence, corrosion in the peak strength condition (HA for 10 hat 410 °C) mainly involves the spontaneous dissolution of thea-Mg. For the moderate strength condition (HA for 5 h at380 °C), the isolated coarse b particles in the microstructure



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(a) (b)Fig. 8. (a) For the as-cast condition, the a-Mg matrix adjacent to coarse b particles which lack the surrounding fine lamellar a+b is corroded initially due to spontaneousdissolution of the a-Mg and due to micro-galvanic coupling between the a-Mg matrix and the continuous fine lamellar a+b plus coarse b particles. When the corrosionpropagation reaches the region of the a+b in the fine lamellar arrangement, the corrosion is retarded due to the corrosion barrier effect. (b) For the peak strength condition(HA for 10 h at 410 °C) the isolated fine b particles cannot effectively block the corrosion propagation because of the large distance between the isolated fine b particles.Furthermore the small isolated fine b particles cause negligible galvanic corrosion when coupled with a large anode (the a-Mg matrix). As a consequence, corrosion in thepeak strength condition (HA for 10 h at 410 °C) mainly involves the spontaneous corrosion of the a-Mg. For the moderate strength condition (HA for 5 h at 380 °C), theisolated coarse b particles in the microstructure can also not cause effective blocking of the corrosion propagation because of the large distance between the isolated coarse bparticles, however, galvanic corrosion is significant and the moderate strength condition gives the highest hydrogen evolution volume.

can also not cause effective blocking of the corrosion propa-gation because of the large distance between the isolatedcoarse b particles; the mechanism is schematically illustratedin Figure 8(b) and the micro-corrosion morphology is pre-sented in Figure 7(b). These coarse b particles cause micro-galvanic acceleration of the adjacent a-Mg matrix. The corro-sion in the moderate strength condition (HA for 5 h at 380 °C)is (i) micro-galvanic corrosion of the a-Mg matrix caused bycoupling with the isolated coarse b particles and (ii) the spon-taneous dissolution of the a-Mg. This is reflected in theimmersion testing results in which the moderate strengthsamples (HA for 5 h at 380 °C) show hydrogen evolution vol-umes higher than for the peak strength sample (HA for 10 hat 410 °C). For the as-cast condition, the a-Mg matrix adjacentto coarse b particles which lack the surrounding fine lamellara+b is dissolved initially due to spontaneous dissolution ofthe a-Mg and due to micro-galvanic corrosion between thea-Mg matrix and the continuous fine lamellar a+b plus coarseb particles as schematically illustrated in Figure 8(a). Whenthe corrosion propagation reaches the region of the a+b in thefine lamellar arrangement, the corrosion is retarded to someextent due to the corrosion barrier role, which is schematical-ly illustrated in Figure 8(a) and the micro-corrosion morphol-ogy is presented in Figures 6(a) and 7(a). This is also support-ed by the fact that the hydrogen evolution volume of the as-cast condition is initially slightly higher than that for the peakstrength sample (HA for 5 h at 410 °C), but becomes slightlylower with the advance of the corrosion attack.

Corrosion of the Fine Lamellar (a+b) Arrangement

Figure 7(a) presents two cases that were unexpected fromprior research on Mg corrosion:[4,5,9] (a) corrosion of the a-Mgup to the a+b in the fine lamellar arrangement, with theappearance of the corrosion of the a-Mg stopped at the inter-face of the a+b in the fine lamellar arrangement, and (b) corro-sion proceeding into the a+b in the fine lamellar arrangement,with both phases having been corroded to a similar extent.

Prior research on Mg corrosion mechanisms, particularlythe work by Song, Atrens and Dargush,[9] leads to the expec-tation that the b phase causes micro-galvanic acceleration ofthe adjacent a phase. It is thus surprising that the a+b in thefine lamellar arrangement can act as a corrosion barrier. Apossible reason for a barrier effect would be provided by theb phase if there is a continuous sheet of the b phase at theedge of the a+b in the fine lamellar arrangement.

The expectation that the b phase causes micro-galvanicacceleration of the adjacent a phase means that it is surprisingthat the corrosion of the a+b in the fine lamellar arrangementoccurred in a manner in which both phases had been cor-roded to a similar extent. Examination of Figure 7(a) indi-cated that the corrosion was not simply corrosion of thea-phase causing undermining (and the falling out of) the bphase particles. The corrosion at the same extent of thea-phase and the b-phase might profit from additionalresearch.

Proposed HA Heat Treatment

The HA heat treatments studied in this work changed themicrostructure by causing the dissolution of the continuousfine lamellae a+b and partial dissolution of the coarse b parti-cles; the aluminium of the b phase dissolves in the a-Magne-sium. The resulting microstructure, with a relatively homoge-neous distribution of isolated b particles, has bettermechanical properties. In terms of corrosion, a superior HAfor 10 h at 410 °C is proposed; this particular HA produces amicrostructure with (i) peak strength (ii) improved ductilityand (iii) no loss of corrosion properties compared with the as-cast condition. This HA is a novel heat treatment producingproperty enhancement of AZ91 and is economically feasiblefor industrial production to AZ91 manufacturers. This HAonly requires a simple isothermal heat treatment, whichcauses dissolution into solid solution of the continuous finelamellar a+b and some dissolution of the coarse b particlesproducing a microstructure of isolated fine b particles in thea-matrix. In Mg alloy technology, this enhanced HA providesa prospect for producing AZ91 with improved mechanicalproperties without the loss of corrosion properties.

Conclusions

Property enhancement of as-cast AZ91 was produced by asimple homogenization annealing. HA for 10 h at 410 °Ccauses an improvement in hardness, ultimate tensile strength,and elongation to failure without a loss of corrosion proper-ties. The microstructure consisted of fine b particles homoge-neously distributed throughout the a-Mg matrix. The im-provement of the mechanical properties, compared with theas-cast condition, is due to the absence of a continuous easycrack path. The isolated fine b particles do not lead to an ob-vious loss of corrosion properties because they act as a smallcathode connected to a large anode (the a-Mg matrix) so thatthe peak aged condition (HA for 10 h at 410 °C) has negligiblemicro-galvanic corrosion

The moderate strength condition (HA for 5 h at 380 °C)gave the highest corrosion rates as measured by the highesthydrogen evolution volume. The peak strength condition(HA for 10 h at 410 °C) and the as-cast condition give similarand lower hydrogen evolution volumes. The corrosion behav-ior after HA can be interpreted as a result of the competitionbetween microstructure features acting as a corrosion barrieror as a galvanic cathode. The fine lamellar a+b in the as-castcondition act as a corrosion barrier to some extent. When thecorrosion attack reaches the region of the fine lamellar a+b,the corrosion is retarded to a certain extent. HA for 5 h at380 °C changes the microstructure causing the dissolution ofthe fine lamellar a+b leaving isolated coarse b particles; galva-nic corrosion becomes significant with the result that themoderate strength condition gives the highest hydrogen evo-lution volume. With refinement of b particles, the small cath-ode (the isolated fine b particles) cause negligible galvaniccorrosion of the relatively large anode (the a-Mg matrix),


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which explains the no obvious loss of corrosion properties forthe peak strength condition (HA for 10 h at 410 °C) comparedto the as-cast condition.

Received: September 11, 2007Final version: October 04, 2007

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Influence of Homogenization Annealing of AZ91 on Mechanical Properties and Corrosion Behavior

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