Combined effect of non-equilibrium solidification and thermalannealing on microstructure evolution and hardness behavior of AZ91magnesium alloy

Z.Z. Zhou, W. Yang n, S.H. Chen, H. Yu, Z.F. XuNational Defence Key Discipline Laboratory of Light Alloy Processing Science and Technology, Nanchang Hangkong University, Nanchang 330063, PR China

a r t i c l e i n f o

Article history:Received 1 January 2014Received in revised form25 February 2014Accepted 26 February 2014Available online 4 March 2014

Keywords:Non-equilibrium solidificationHeat treatmentMicrostructureMetals and alloys

a b s t r a c t

Non-equilibrium solidification of commercial AZ91 magnesium alloy was performed by copper moldspray-casting technique and the thermal stability property of as-formed meta-stable microstructure wasinvestigated by subsequent annealing at different temperatures and times. Remarkable grain refinementappears with increasing cooling rate during solidification process, which is accompanied by a visiblecellular/dendrite transition for the grain morphology of primary phase. Moreover, the non-equilibriumsolidified alloy exhibits obvious precipitation hardening effect upon annealing at 200 1C, and theprecipitation mode of β-Mg17Al12 phase changes from discontinuous to continuous with extendingisothermal time from 4 h to 16 h, which generates an increase of resultant micro-hardness value. Aftersolid solution treatment at the elevated temperature of 420 1C, the volume fraction of β-Mg17Al12 phasedecreases and a notable grain growth phenomenon occurs, which give rise to a reduction of hardness incomparison with that of as-quenched alloy.

& 2014 Elsevier B.V. All rights reserved.

1. Introduction

Magnesium alloy has been developed as one of the mostattractive engineering materials due to its satisfactory energysaving, magnetic shield characteristic, environmental issues andabundant resources [1–3]. Unfortunately, both the poor formabil-ity and strength arising from the h.c.p. crystal lattice and the lowcorrosion resistance owing to the absence of a protective oxidefilm are major shortcomings for its wide application in aerospaceand automotive industries [4].

In recent years, rapid solidification (henceforth referred to asRS) has received quite extensive attention because of its ability toproduce structural refinement, solid solubility extension andmeta-stable phase, which in turn provide approaches for improv-ing the properties of materials [5–9]. Since the pioneering workconducted by Olsen and Hultgren [10], who investigated the effectof cooling rate on the homogeneity of solid solution, a great deal ofeffort has been spent in this field with broad scientific interest andtechnological importance [11–13]. Up to now, a series of high-performance magnesium alloys have been fabricated by differentRS methods, such as melt-spinning [14], atomization [15], surface

melting [16] and spray-casting [17]. As expected, the formedproperties are far superior to those of conventional as-cast alloys.

However, it should be noted that the products formed by theavailable RS methods are mostly low-dimensional with theshape of powder, ribbon or ring so as to achieve the best with-drawing capacity for heat from the melts. Then, a subsequentconsolidation process is needed to form a bulk ingot for commer-cial application, as reported by the foil metallurgy technique forconsolidation of rapidly solidified ribbons [18–20] and the powermetallurgy technique by hot press bonding [21]. In such a case, theprecipitation of super-saturation solid solution and the growth oforiginal refined grain are prone to occur at elevated temperaturebecause of the activated atomic diffusion behavior, which defi-nitely changes the initial non-equilibrium solidification character-istics [22–25]. Broadly speaking, the mechanical properties ofpolycrystalline materials depend critically on the average grainsize, and often improve with the reduction of average grain sizeaccording to Hall–Petch relationship [26]. From this viewpoint,understanding the thermal-stability behavior of as-quenchedmicrostructure to retain the benefits of rapid solidification proces-sing becomes the major problem for developing high-performancemagnesium alloy.

Despite many experimental investigations reported on the RSmagnesium alloys to improve their strength, workability andcorrosion resistance [14–21], there is still a lack of systematicresearch on rapid solidification and the relevant thermal stability

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Physica B

http://dx.doi.org/10.1016/j.physb.2014.02.0520921-4526 & 2014 Elsevier B.V. All rights reserved.

n Corresponding author. Tel.: þ86 791 86453167; fax: þ86 791 86453167.E-mail addresses: weiyang@mail.nwpu.edu.cn, nchkyw@163.com (W. Yang).

Physica B 443 (2014) 35–42

behavior, especially for the comparisons of microstructure evolu-tion and mechanical properties of RS magnesium alloys withdifferent heat-treatment routes. In this paper, RS of AZ91magnesium alloy with different cooling rates were fabricated bychanging the inner diameter of copper mold. Then, the micro-structure characterization and annealing response of as-chilledalloy with different temperatures and isothermal times wereelucidated so as to guide the development of advancedmagnesium alloy.

2. Experimental procedures

Commercial AZ91 alloy was used in this study. The nominalcomposition of this alloy (in mass percentage) is 9.04 Al, 0.65 Zn,0.33 Mn, 0.03 Si, 0.0048 Cu, 0.0003 Ni, 0.0015 Be and Mg balance.Firstly, the magnesium blocks with the weight of �20 g weremelted in a quartz crucible by a high frequency induction heatingfacility. During the heating process, the melts were kept under theprotection of Ar atmosphere to ensure that the vaporization andoxidation of Mg was adequately controlled. Then, the molten alloywas pressurized under a purified Ar atmosphere and injectedthrough a quartz nozzle into the copper mold. To obtain differentcooling rates, the inner diameter of the copper mold was designedas 3 mm, 6 mm and 8 mm separately.

Heat-treatment experiment was carried out for the as-quenchedingot with the diameter of 8 mm and the samples were fabricatedinto rods with an height of 5 mm. The annealing procedure wasperformed in a heat-resistance furnace with the protection offlowing Ar atmosphere. To compare the thermal stability behaviorfor different temperatures, the isothermal annealing was main-tained at 200 1C and 420 1C, which correspond to aging and solidsolution treatment separately. Considering the dependence ofatomic diffusion capacity on temperature, the aging times werechosen as 4 h, 8 h, 12 h and 16 h, while that for solid solution theywere chosen as 2 h, 4 h, 6 h and 8 h.

As-prepared samples were mounted, polished and etched with5% citric acid according to the standard metallographic prepara-tion techniques. Then, microstructures of etched specimens werecharacterized by optical microscopy (VHX-600, KEYENCE) andscanning electron microscope (Quantan 200) equipped withenergy dispersive spectroscopy (INCA). The micro-hardness mea-surements were conducted by using a 200 g load and a dwell timeof 3 s with a Vickers indenter (DuraScan). For each of the samples,five measurements were carried out at different areas and theaverage value was used in the present work.

3. Results and discussion

3.1. Effect of cooling rate on the microstructure of AZ91 alloy

Fig. 1 shows the optical microstructure changes of AZ91magnesium alloy prepared at different cooling rates. The typicalmicrostructure of as-cast alloy is presented in Fig. 1a, where well-developed primary α-Mg dendrites surrounded by β-Mg17Al12phase in divorced form are clearly visible. In this condition, theresultant grain size is quite large and its distribution is non-uniform due to slower cooling rate. As reported by Ref. [27], thecooling rate in chill casting changes reversely with the innerdiameter d of copper mold. Apparently, a gradual grain refinementphenomenon appears for increasing cooling rate and the corre-sponding grain distribution becomes more homogeneous, asevidenced in Fig. 1b–d by decreasing d from 8 mm to 6 mmand finally 3 mm. The aforementioned microstructure evolutioncan be interpreted by the enhanced nucleation sites for elevatedundercooling and insufficient grain growth for limited solidifica-tion time at higher cooling rate [28]. In addition, a transition ofprimary grain morphology from cellular to dendrite can beidentified because of the strengthened interface unstable effectarising from the increase of growth rate for enhanced non-equilibrium effect [29].

Fig. 2 presents the scanning electron microscopy images of theAZ91 alloy fabricated at different cooling rates. It can be deducedthat the average grain size of primary α-Mg reduces from largerthan 200 μm for as-cast state (Fig. 2a) to smaller than 25 μm foringot as-quenched into copper mold with d¼8 mm (Fig. 2b).Continuous refinement still prevails with the further decrease ofd to 3 mm (Fig. 2c). It has been reported [30] that the equilibriumpartition coefficient (the ratio of solute concentration in the solidto that in the liquid) of solute Al in primary α-Mg is merely 0.3,which means that more and more Al atoms are rejected accom-panying the growth of α-Mg. In terms of the relatively widetemperature interval between the liquids and the solids curves,the solute segregation becomes more serious with the process oftransformation. Therefore, divorced eutectic β-Mg17Al12 phaseforms in the interdendritic region because of the enrichment ofAl element, as indicated by the EDS result shown in Fig. 2a.

With the increase of cooling rate, the non-equilibrium effect isenhanced and the migration rate of solid/liquid interface isaccelerated, which may be comparable or larger than thatfor atom diffusion across the interface. In such a case, the soluteatoms seem to be caught by the growing phase and the segrega-tion coefficient will deviate from the initial 0.3 to unity, asdescribed by the solute trapping theory [29]. Accordingly, remark-able solute super-saturation effect appears due to the adoption of

Fig. 1. Optical microstructures of AZ91 magnesium alloy with different cooling rates. (a) As-cast; (b) spray-casting, d¼8 mm; (c) spray-casting, d¼6 mm; and (d) spray-casting, d¼3 mm.

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non-equilibrium solidification, which can be supported by the EDSmeasurements shown in Fig. 3. As expected, increase of solutecontent can be seen for larger cooling rate. It should be noted that

the cooling rate herein was calculated in relation to inner diameterof copper mold, as mentioned in Ref. [27].

3.2. Effect of aging on the microstructure of RS AZ91 alloy

The optical microstructure changes of RS magnesium alloy afteraging at 200 1C for different isothermal times are shown in Fig. 4.As can be seen clearly, the average size of the primary grain is notmuch influenced by extending the holding time during agingtreatment (Fig. 4a–d). Moreover, the dendrite structure has beenmaintained after aging treatment, and its morphology appearsmore clearly, especially after aged for 16 h, as illustrated in Fig. 4d.According to thermodynamic calculation, the equilibrium concen-tration of solid solubility of Al in magnesium at 200 1C is merely2.9 wt% [30], which is much lower than the super-saturation valueformed in RS process (Fig. 3). Therefore, the precipitation processappears and the formation of β-phase precipitates occurs. Gen-erally speaking, precipitation occurring in supersaturated alloy canbe conducted either discontinuously or continuously. The former ischaracterized by the cellular growth of secondary phase adheringto grain boundary and the formation of a lamellar structure behinda moving grain boundary, while the latter proceeds in all the

Fig. 2. Scanning electron microscopy images of the AZ91 alloy with different cooling rate. (a) As-cast; (b) spray-casting, d¼8 mm; and (c) spray-casting, d¼3 mm.

Fig. 3. Variation of Al content in primary grain as a function of cooling rate duringsolidification.

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remaining regions of the supersaturated matrix and forms largenumbers of fine dispersive particles.

In order to observe the distribution of precipitates in annealedRS alloy, the scanning electron microscopy images for the samplesafter aged for 4 h and 16 h are displayed in Fig. 5, which representthe shortest and longest time for aging treatment. Apparently,both the morphology and volume fraction of phases (Fig. 5a and c)are similar to those in as-quenched state (Fig. 2b). However, asseen from enlarged micrographs in Fig. 5b, flakes of intermetalliccompounds with distinct anisotropic morphology can be identifiedalong the grain boundary and the maximum length grown intothe α-Mg matrix is nearly 2 μm, which indicates discontinuousprecipitation behavior of β phase and is consistent with theprevious investigation [19]. With increasing isothermal time, manyfine particles are found in the interior of matrix, which can beascribed to the occurrence of continuous precipitation (Fig. 5d).It is well known that the heterogeneous nucleation number forcontinuous precipitates mainly depends on the crystal defectswithin the matrix. As for RS alloy, various nucleation sites, such asvacancies, dislocations, and stacking faults are formed during the

non-equilibrium transformation process, which further favor theoccurrence of continuous precipitates as shown in Fig. 5d. More-over, the evenly dispersed fine compound particles can act asobstacles to suppress the grain boundary sliding [30]. So, it can beexpected to retard the grain growth effectively, as evidenced inFig. 4.

From the preceding discussion, the volume fraction of precipi-tated β grains increases with extending the time for aging process,According to mass conservation, the formation of solute-rich βphase (�43.95 wt% Al) will consume a certain amount of Al atomsand reduce the super-saturation extent in initial non-equilibriumstate. This is also verified by the EDS analysis shown in Fig. 6,where continuous decrease of solute content of Al in primary grainis observed with increasing the aging time during this solid–solidtransformation process.

3.3. Effect of solid solution on the microstructure of RS AZ91 alloy

Fig. 7 shows the optical microstructure changes of RS AZ91alloy after solid solution treatment for different times, which is

Fig. 4. Optical microstructures of RS magnesium alloy after aging at 200 1C for different times. (a) 4 h; (b) 8 h; (c) 12 h; and (d) 16 h.

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much different from those shown in Fig. 4. Obviously, increasingthe annealing temperature diminishes the dendrite form of grainsand the whole microstructure transforms into polygon-like shape.In comparison with the aging treatment, the solid solution at420 1C is capable to dissolve the β phase into the magnesiummatrix completely (Fig. 7a). This is due to the relatively hightemperature used for the dissolution of divorced β phase in grain

boundary. According to the previous study [31], the time requiredto obtain the homogeneous microstructure for solid solution ofas-cast AZ91 alloy is at least �24 h because of the sluggishdiffusion rate of Al in magnesium solid state. However, the timefor solid solution to obtain homogeneous microstructure incurrent study is merely 2 h, which can be ascribed to the reduceddiffusion distance owing to refined microstructure (Fig. 1) and theincreased solute content arising from RS process (Fig. 3).

Another feature observed in the samples after solid solutiontreatment at 420 1C is the coarsening of primary fine grains withincreasing the annealing time (Fig. 7a–d). The grain size, whichwas calculated using an image analysis software, is shown in Fig. 8,which indicates the prevalent grain growth behavior. With theincrease of holding time from 2 h to 8 h, the average grain size hasincreased by a factor of 2–3, i.e., changing from �35 μm to�85 μm.

Generally speaking, polycrystalline materials tend to reduce thesystem free energy by reducing the amount of grain boundaryarea. It has been reported that the kinetic equation of grain growthfor many metallic materials during isothermal annealing can begiven by [32]

Dn�Dn0 ¼ kt ð1Þ

where D is the average grain size after annealing for time t, D0 isthe initial grain size, n is the grain growth exponent, and k thetemperature dependent rate constant in an Arrhenius relationship.Previous work [32] showed that n equals 2 as for single phasesystem without precipitates, where the grain growth kinetics is

Fig. 5. Scanning electron microscopy images of the AZ91 alloy after aging at 200 1C for different times. (a) and (b) 4 h; (c) and (d) 16 h.

Fig. 6. Variation of Al content in primary grain as a function of holding time afteraging at 200 1C.

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mainly controlled by grain-boundary migration driven by grain-boundary energy. However, when grain growth is prevented byimpurities or fine intermetallic particles existing on grain-bound-ary, n42 is frequently reported, which implies sluggish growthrate according to the Zener pinning theory. As seen from Fig. 7, theβ phase in grain boundary has disappeared completely after solidsolution for 2 h. Thus, the grain growth exponent n will definitelybe reduced toward 2. According to Eq. (1), notable grain growthbehavior occurs because of the weakened pinning effect.

3.4. Measured hardness of as-prepared AZ91 alloy

Fig. 9 presents the measured hardness of as-prepared AZ91Dalloy. Generally speaking, both the microstructure refinement andformation of over-saturation solid through rapid solidification (RS)processing can result in substantial improvement of mechanicalproperties. For these reasons, the hardness increases continuouslywith increasing the cooling rate (Fig. 9a). It increases from 68 HVfor as-cast specimen to 104 HV for RS alloy with the innerdiameter of 3 mm.

Fig. 7. Optical microstructures of RS magnesium alloy after solid solution at 420 1C at different times. (a) 2 h; (b) 4 h; (c) 6 h; and (d) 8 h.

Fig. 8. Variation of grain size for primary phase as a function of isothermal timeafter solid solution at 420 1C.

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Moreover, the Vickers hardness of samples, as chilled by thecopper mold with d¼8 mm and isothermal annealing at 200 1C fordifferent times, is given in Fig. 9b. As can be seen clearly, a furtherincrease of hardness prevails, which can be attributed to theoccurrence of fine-scale and uniform distribution of precipitatesin the alloy (Fig. 5). From the above discussion, it has beenmentioned that the formation of precipitates reduces the solutecontent in super-saturation of primary grain (Fig. 6). So, it can be

inferred that the precipitation hardening plays a stronger role thansolid solution strengthening. However, Vickers hardness decreasesgradually after annealing at 420 1C with extending the holdingtime, as presented in Fig. 9c. This can be attributed to thedissolution of β phase in accordance with the accelerated atomicdiffusion rate. In addition, the notable grain growth of matrixshould also account for the degradation of hardness.

4. Conclusions

Non-equilibrium solidification of commercial AZ91 magnesiumalloy was fabricated successfully via copper mold spray-castingtechnique and the thermal stability was investigated by annealingwith different temperatures and isothermal times. In comparisonwith the conventional as-cast alloy, striking grain refinementoccurs in the as-quenched alloy and a transition of primary grainmorphology from cellular to dendrite can be observed clearly.Moreover, the extent of solute super-saturation is enhanced withincreasing the cooling rate because of solute trapping effect. As forthe samples aged at 200 1C, the average grain size of primaryphase is not much varied with extending isothermal time, exceptfor the transition from discontinuous precipitation to continuousprecipitation favored by the enhanced crystal defects arising fromRS process. After solid solution of as-quenched AZ91 alloy at420 1C, the initial dendrite morphology changes into polygon-like shape and a single-phase solid solution microstructure can beobtained. The disappearance of β phase gives rise to the weakenedpinning effect for grain-boundary migration and a notable graingrowth of primary α-Mg phase. Enhancement of Vickers hardnessoccurs with increasing cooling rate during solidification process, aswell as for the aged samples with extending isothermal time. Incontrast, slight reduction appears after solid solution because ofthe dissolving of β phase and the obvious grain growth behavior.

Acknowledgments

The work is supported by the fund of the National NaturalScience Foundation of China (Grant no. 51164028), NationalDefense Key Disciplines Laboratory of Light Alloy ProcessingScience and Technology in Nanchang Hangkong University(GF201201003). The authors thank the Instrumental AnalysisCenter of Nanchang Hangkong University and Prof. Y.L. Ai forproviding the necessary testing equipments. Great appreciation isalso given to the anonymous reviewer for the valuable suggestionsand comments on the manuscript.

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