Effect of Cold Forging and Static Recrystallization on Microstructure

and Mechanical Property of Magnesium Alloy AZ31

Jianzhong Liu1;2, Tianmo Liu1;2;*, Hanqi Yuan1;2, Xiuling Shi1;2 and Zhongchang Wang3;*

1College of Materials Science and Engineering, Chongqing University, Chongqing 400044, P. R. China2National Engineering Research Center for Magnesium Alloys, Chongqing University, Chongqing 400044, P. R. China3World Premier International Research Center, Advanced Institute for Materials Research, Tohoku University,Sendai 980-8577, Japan

We perform a cold forging and annealing technique on Mg alloy AZ31 and investigate the impact on microstructure evolution andmechanical property. In particular, we focus on how twin and static recrystallization behave during the forging process followed by theannealing. Interestingly, we find that a large number of thick-lenticular {10�112} twins emerge at initial stage of deformation and subsequentlyevolve into narrow-band {10�111} and {30�332} twins. The transformable twins are found to be crucial for inducing static recrystallization viaproviding recrystallization sites and refined grains. Moreover, the alloy forged at room temperature and annealed at 623K is found to haveoptimal microstructure due to the complete recrystallization and therefore exhibit the highest micro-hardness, largest compressive strength, andmost significantly enhanced compressive ratio. The improved mechanical properties are comparable or even superior to those of the alloydeformed using other techniques, rendering the cold forging a promising way for further tailing properties of Mg alloy.[doi:10.2320/matertrans.M2009288]

(Received August 17, 2009; Accepted November 6, 2009; Published December 23, 2009)

Keywords: magnesium alloy AZ31, cold forging, static recrystallization, microstructure, mechanical property

1. Introduction

The abundant natural resource of magnesium is stimulat-ing extensive investigations focusing on its alloy in order toimprove the physical and mechanical properties. In fact, theinterest toward Mg alloy is twofold. On one hand, it is anattractive metal with excellent intrinsic characteristics suchas good heat dissipation, damping, electromagnetic shielding,and high specific resistance.1–3) On the other hand, it is alow-density structure material, demonstrating the ability tofunction in conditions where weight savings directly lowercost and increase efficiency. However, the large-scale useof Mg alloy as a structure material (e.g., AZ31B sheet) iscurrently restrained by its low formability at room temper-ature due to the restricted number of slip systems inhexagonal close packed structure of Mg alloy. To improvethe poor plastic workability, the modification of crystalstructures via doping Li,4,5) rare earth,6) or other foreignatoms7) has been attempted. The problem behind theseapproaches is that the workability can only be enhanced in alimited fashion because they rely critically on the chemicalcomposition of alloy.

Another well-known way to modify the microstructure ofMg alloy so as to improve its mechanical properties isthrough plastic deformation. Such deformation can producea wrought Mg alloy that has more superior properties thanthe alloy subjected to the die casting.8,9) Of all examineddeformation methods, the severe plastic deformation (SPD)has attracted the broadest attention, largely because it canoptimize the microstructures of Mg alloy and consequentlyimprove the ductility, strength, and superplasticity.10,11) Theequal-channel angular extrusion (ECAE),12,13) change-chan-

nel angular extrusion (CCAE),14) C-shape equal-channelreciprocating extrusion (CECRE),15) and accumulative roll-bonding (ARB)16,17) are just a few important SPD examples.Although these SPD techniques have been confirmed to beable to refine grains in Mg alloy via dynamic (DRX) or staticrecrystallization (SRX), they are in general complicated.Recently, there appear several simple deformation methodsthat are also effective for improving mechanical propertiesof Mg alloy, which involve roll,18) stamping,19) forwardextrusion,20) and forging.21) Among them, the techniquecombining forging with annealing has a number of advan-tages especially in workability and strengthening of Mgalloy. For this reason, much effort has been put on thistechnique,22,23) in particular, the hot forging with temperatureup to 573K.24) However, little is known about the formabilityof Mg alloy below 573K, namely, under cold-forgingcondition. A detailed understanding of the cold forgingtechnique requires a comprehensive examination of micro-structure and its impact on property, and this has so far beenlacking.

In this paper, a combined method of cold forging withannealing has been applied on Mg alloy AZ31 in order toinvestigate how microstructure evolves during the deforma-tion. The alloy forged under a strain of 10%was subsequentlyannealed at various temperatures. In addition to characteringmicrostructure, the goal of this study is to systematicallyexamine the effect of deformation and SRX on twin, grainsize, and mechanical property. The data presented in thiswork are also compared with available results obtained usingother deformation techniques.

2. Experimental Procedure

The material we used was a commercially available Mgalloy AZ31 with the following composition (in mass%): 3%

*Corresponding author, E-mail: tmliu@cqu.edu.cn; zcwang@wpi-aimr.

tohoku.ac.jp

Materials Transactions, Vol. 51, No. 2 (2010) pp. 341 to 346#2010 The Japan Institute of Metals

Al, 1% Zn, 0.3% Mn, and Mg (balance), which was suppliedin form of cast ingot. The ingot was first machined intocylindrical samples with a gage length of 30mm and a cross-section dimension of 15mm and then homogenized asstarting materials for the forging. The cold forging wascarried out by a set of home-made hammers at roomtemperature. To better control the strain, the cylindricalsample was placed inside a steel ring with internal diameterof 20mm during the deformation. The strains we tested were2.5%, 5%, 7.5%, 10%, and 15%, respectively, within whichthe sample forged under the strain of 10% was selected andannealed at temperatures ranging from 573 to 773K for 3 hin order to reveal microstructure evolution during the SRX.Finally, we measured micro-hardness and compressiveproperties of the annealed samples.

Microstructures were observed using optical microscopy(OM) and transmission electron microscopy (TEM). For theOM, the specimens were prepared by cutting, mechanicalpolishing, and subsequent etching in a chemical solution. Themean grain sizes of annealed samples were determined fromthe OM following the procedure given in the ASTM standardE112-95. As for the TEM, the samples were first cut fromforged billets along normal to forging direction and thenmechanically polished to thin foils. The foils were finallyelectropolished at 228K to make electron transparent. Thespecimens were characterized using the Philips TECNAI-20electron microscope operated at 200 kV. Compressive ex-periments were carried out on annealed samples with 7mmin diameter and 14mm in length, which were obtained bycutting along longitude of forged billets. The Vickershardness (HV ) was measured using a micro-hardness indenterat a load of 0.49N for 20 s.

3. Results and Discussion

3.1 Microstructure evolution during cold forgingFigure 1 show optical microstructures of as-cast and

homogenized specimens before forging. Secondary phasesemerge around grain boundaries in the as-cast case(Fig. 1(a)) but disappear largely after the homogenization(Fig. 1(b)). As a result of annealing, the grains are somewhatcoarsened with their sizes ranging from 200 mm to 250 mm,which are large enough to form twins during forging at roomtemperature.25) In fact, as we will see later, there indeedappear several types of twins during cold forging, includingthick-lenticular, secondary (double), and narrow-band twins.

Figure 2 show representative optical images of the alloysforged under various strains. From this figure, both amountand morphology of twins are observed to alter substantiallywith the rise of strain. In the strain case of 2.5%, the majorityof twins are of thick-lenticular shape, although secondarytwins are also visible. However, these thick-lenticular twinstransform, to a large extent, to narrow-band ones when thestrain is beyond 2.5% (Figs. 2(b)–2(e)). The disappearanceof thick-lenticular twins may not be attributed to the DRX,as mentioned in previous reports,26) because the alloy isdeformed at low temperature with high speed so that time isnot sufficient for the DRX. By careful analysis using electronbackscattered diffraction (EBSD), we further identify thethick-lenticular twins as {10�112} twins and narrow-band twinsas {10�111} and {30�332} ones, in good agreement with theresults obtained using rolled AZ31 sheets and as-extrudedbars.25–27)

In general, there are two typical twin modes for Mg alloy:f10�112gh10�111i tension twin and f10�111gh10�112i compressiontwin.28) Since the latter can only be generated under a largecritical resolved shear stress (CRSS) (3MPa for the formerand 112MPa for the latter),29,30) it is not surprising toobserve first the thick-lenticular twins during deformation(Fig. 2(a)). Subsequent increase of strain may accumulatestress in grains, which could eventually result in theformation of narrow-band twins (Figs. 2(b)–2(e)). This issimilar to the observations by Chino et al.26) showing thatthe narrow-band twins appear in the alloys compressedunder a strain of 11%. The main difference is that the strainused in their study is higher than ours (5%), which isattributable to the higher forging velocity and largerhomogenized grains we used.

To shed light on how twins behave during forging, wepresent in Fig. 3 images of the specimen forged under thestrain of 15%. From Fig. 3(a), one can see that the narrow-band twins traverse grain boundaries and change theirorientations, in good agreement with the observations byKoike.25) This is due to the formation of step-like ledgeswhen basal slip dislocations inside the f10�111g twins head fortheir neighboring grains.31) Further TEM observations con-firm the presence of slip dislocations both around (Fig. 3(b))and within twins (Fig. 3(c)). In addition, the imagingsuggests that slip and twin interact during forging in eithera repulsive or an attractive manner, depending on theirmode types. Specifically, since the interaction between slipand {10�112} twin is repulsive,28) dislocations may pile up at

(a) (b)

Fig. 1 Optical microstructures of AZ31 alloy before cold forging: (a) as-cast and (b) homogenized at 673K.

342 J. Liu, T. Liu, H. Yuan, X. Shi and Z. Wang

(b)(a)

(c) (d)

(e)

Fig. 2 Optical micrographs of the alloy forged under the strains of (a) 2.5%, (b) 5%, (c) 7.5%, (d) 10%, and (e) 15%.

(a) (b)

(c)

Fig. 3 Two types of twin in the alloy forged under strain of 15% at room temperature: (a) narrow-band {10�111} twin observed by the OM

and (b) thick-lenticular {10�112} twin by the TEM. (c) Bright-field TEM image of interactions between thick-lenticular twins. The inset

shows a selected area diffraction pattern of {10�112} twins.

Effect of Cold Forging and Static Recrystallization on Microstructure and Mechanical Property of Magnesium Alloy AZ31 343

twin boundaries, which can therefore result in high concen-tration of local stress. Apart from this interaction, twinsthemselves may interact without altering their originalorientations (Fig. 3(c)), which infers that the growth of twinsmight be impeded by other twins at boundaries.

Although twins are known to contribute less to strain thanslips, they are critical for improving deformability of Mgalloy via tuning orientations so as to reactivate other slipsystems.27) At the initial stage of forging, only the tension{10�112} twin and basal slip are important for the deform-ability because of their low CRSS, despite that other types ofslip and twin may also contribute. Further deformationrequires more strain than what the twinning alone can offer,which causes crystallographic slips and dislocations. Thedislocations can pile up at grain or twin boundaries and thusinduce large stress, which could eventually form compres-sion {10�111} twins. These high-CRSS twins may rotatelattice, activate basal slip, and facilitate dislocation move-ment, thereby making further forging of Mg alloy under largestrain likely.32)

3.2 Microstructure characteristics after SRXTo investigate the role of twins on recrystallization, we

show in Fig. 4 microstructures of the alloys forged under

strain of 10% and annealed at various temperatures for 3 h.Clearly, the twins are largely substituted by well equiaxedgrains, indicating that the SRX takes place and that grainsstart to grow. The mean grain sizes for the alloys annealed attemperatures ranging from 573K to 773K by a step of 50Kare determined to be 9, 15, 23, 26, and 33 mm, respectively.Although the grains are coarsened with the rise of temper-ature, they all are refined after deformation.

Figure 5 shows microstructure of the alloy forged underthe strain of 10% and annealed at 623K for 2 h. From thisfigure, twins are found to distribute where the SRX occurs orfine grain appears. We thus conclude that the twin SRX takesplace during annealing. This is because severe deformationenables the formation and subsequent accumulation ofdislocations and twins especially inside coarse grains. Thesedislocations can arrange themselves to form cell structuresand walls and then evolve into sub-grains, thereby increasinggrain-boundary energy dramatically. Such grains withnonequilibrium boundaries could store sufficient energy toundergo the SRX during annealing. Finally, a comparisonbetween Fig. 4 and 5 reveals that as the annealing continues,low-angle grains are inclined to evolve into high-angleequiaxed ones, suggestive of the important role of twin andSRX on grain refining.

(a) (b)

(c) (d)

(e)

Fig. 4 Optical micrographs of the alloy forged under strain of 10% and annealed at (a) 573K, (b) 623K, (c) 673K, (d) 723K, and

(e) 773K for 3 h.

344 J. Liu, T. Liu, H. Yuan, X. Shi and Z. Wang

3.3 Mechanical properties after SRXTo investigate how microstructure evolution affects

mechanical properties, we first measured micro-hardness(HV ) for the AZ31 alloy forged under strain of 10% andannealed at various temperatures. As seen in Fig. 6, the HV

decreases gradually with the rise of annealing temperature,which can be ascribed to the grain coarsening (Fig. 4).The fast coarsening can be understood by considering anexponential relationship between diffusion coefficient (D)of grain boundary and annealing temperature (T), as ex-pressed by D ¼ D0 expð�Q=KTÞ. Evidently, grain boundarymigrates swiftly as the temperature increases, which even-tually results in quick growth of neonatal grains. Further,we note in this figure that the HV and grain size (d)follow approximately the Hall-Petch relationship given byHV ¼ H0 þ KHd

�1=2,33) where H0 and KH are materialconstants and calculated to be 17.3 and 239.3, respectively,for the annealed alloy. This means that strengthening is adirect consequence of gain refining.

Next, we investigated room-temperature compressiveproperties for the forged alloy before and after annealing,as shown in Fig. 7. The stress in both cases increases sharplyat initial stage of deformation but gradually afterwards untilthe maximum point. A key difference is that the fracturestrain for the alloy before annealing (4.75%) is much lowerthan that for the annealed alloy having a value of 18.75%. Inaddition, the ultimate compressive strength (UCS) in theunannealed case (248MPa) is also smaller than that in theannealed case (367MPa), although the yield strength for the

unannealed alloy is higher (205 vs. 167MPa). Interestingly,the hardening exponent in the unannealed case is found to bevery small, which can be ascribed to pinning of dislocations.As well known, dislocations may heavily tangle andaccumulate either along boundaries or in interior of twinsand grains during the forging, thus causing hardening.However, the tangled dislocations are relieved significantlyafter annealing due to the SRX, which explains theaforementioned property change.

Finally, we presented in Fig. 8 yield strength, UCS, andcompressive ratio for the alloy annealed at various temper-atures. From this figure, we see that they all have the optimalvalues at 623K. A notable difference is that the yield strength

(a) (b)

Fig. 5 (a) Optical micrographs of the alloy forged under strain of 10% and annealed at 623K for 2 h. (b) Magnified plot of (a).

550 600 650 700 750 800

50

55

60

65

d -1/2/mm-1/2

Temperature, T / K

0.14 0.16 0.18 0.20

Har

dnes

s, H

V /

HV

Fig. 6 Micro-hardness of the alloy after SRX as a function of annealing

temperature (bottom x axis) and d�1=2 (top x axis), where d is grain size

obtained from Fig. 4.

0 4 8 12 16 200

100200300400

Strain, ε / %

Annealed

(b)0

100200300400500

Stre

ss, σ

/ MP

a

Unannealed

(a)

Fig. 7 Compressive stress-strain characteristics of the alloy forged under

strain of 10% (a) before annealing and (b) after annealing at 623K for 3 h.

550 600 650 700 750 800

120

180

240

300

360

Com

pres

sive

Rat

io (

%)

Temperature, T / K

Stre

ss, σ

/ MP

a

Yield Strength

Compressive Ratio

UCS

14

16

18

20

Fig. 8 Yield strength, ultimate compressive strength (UCS), and compres-

sive ratio of the alloy forged under strain of 10% and annealed at various

temperatures for 3 h. Compressive ratio is represented by right axis.

Effect of Cold Forging and Static Recrystallization on Microstructure and Mechanical Property of Magnesium Alloy AZ31 345

at 573K is larger than that at 675K, in contrast to the casesof UCS and compressive ratio. This can be attributed itsresidual stress and crystallographic texture as well as itsincomplete SRX. However, the SRX at 623K is completed,which refines grains and are responsible for the optimalproperties. Finally, these values are comparable to the dataobtained using the hot-extrusion technique.25)

4. Conclusions

We have applied a cold forging and annealing method toMg alloy AZ31 and examined systematically the micro-structures and mechanical properties. We have found that thethick-lenticular {10�112} twins are formed at initial stage ofdeformation but subsequently transformed to narrow-band{10�111} and {30�332} twins. Further annealing of the alloyreveals that static recrystallization takes place in twininteriors and that twin is of central importance to grainrefining. Moreover, the micro-hardness and mean grain sizeare found to follow approximately the Hall-Petch relation-ship, which means that grain refining is crucial for alloyhardening. By measuring several mechanical quantities, wehave found that the combined deformation technique iseffective in improving the mechanical properties of AZ31alloy in a way comparable to those of the processed alloy byhot extrusion.

Acknowledgements

This work was supported in part by a Grant-in-Aid forScientific Research on Priority Area, ‘‘Atomic Scale Mod-ification (Grant No. 474)’’, from MEXT of Japan, and aNational 973 Major Project of China, ‘‘The Key FundamentalProblem of Processing and Preparation for High PerformanceMagnesium Alloy’’ under Grant No. 2007CB613700.

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