!!Effect of sc and zn additions on microstructure and hot formability of al mg sheet alloys

Effect of Sc and Zn Additions on Microstructureand Hot Formability of Al-Mg Sheet Alloys

HANLIANG ZHU, ARNE K. DAHLE, and AMIT K. GHOSH

Trace elements of Sc or Zn were added to Al-Mg sheet alloys. Microstructural examinationshowed that the Sc addition greatly refined the grain size, especially in the Zn-containing alloywhere large amounts of subgrains and fine grains were formed. Meanwhile, a number of largeprimary intermetallic particles formed in the Sc-adding alloys. In order to evaluate the high-temperature formability, warm tensile tests were carried out at temperatures ranging from275 �C to 350 �C and at strain rates of 0.015 to 1.5 s�1. The test results showed that, in thealloys with the single addition of Zn or Sc, Zn slightly increased the flow stress but decreased theductility and Sc worsened both the flow stress and the ductility. However, in the alloy with thecombined addition of Zn and Sc, the flow stress was significantly increased at almost all testingconditions and the ductility was also increased at a higher temperature of 350 �C and lowerstrain rate of 0.015 s�1. The results of superplastic tensile tests and biaxial stretch tests dem-onstrated that the alloy with both Zn and Sc additions exhibited good high-temperatureformability. The effect of Zn and Sc additions on the microstructure and the hot formability isdiscussed.

DOI: 10.1007/s11661-008-9728-6� The Minerals, Metals & Materials Society and ASM International 2009

I. INTRODUCTION

WITH the widespread application of lightweightaluminum alloy products in the automotive industry,alloys with a combination of good formability, highperformance, and low cost are required. Especially forapplications where conditions of high temperature orstress are involved, superior creep resistance or highstrength is an essential requirement. Because theincrease of strength is always accompanied by asubstantial decrease in ductility, the ambient tempera-ture ductility and formability of these alloys is alwaysvery limited, causing the conventional manufacturingoperations such as rolling, forging, or drawing to bedifficult, and thereby restricting their use.[1] As a result,forming technologies at intermediate and high temper-atures such as warm forming, hot forging, and super-plastic forming appear to be more promising routes forfabricating structural components using these alloys.

At elevated temperatures, the deformation mecha-nisms occurring during plastic flow can be divided intothree distinct classes depending upon whether they areintragranular dislocation movement, grain boundarysliding (GBS), or diffusional flow.[2] Generally, at highdeformation temperatures and low strain rates, GBS isthe predominant mechanism for fine-grained superplas-tic materials (grain size:<10 lm).[1] To maintain the fine

grain size during the high-temperature deformation, thepresence of second-phase particles at the grain bound-aries is required to achieve superplasticity.[3] If thedeformation temperature is decreased, the strain rateincreases, or, in coarse-grained materials, dislocationslip becomes important and deformation becomesconcentrated within the grains.[3] If second-phase par-ticles are present within the grains, they can retarddislocation motion, causing strengthening. Grainboundaries can also act as barriers to deformation atlow temperatures; hence, a small grain size will causean increase in strength and a decrease in ductility.Overall, grain size has an opposite effect on the flowstress at high and low temperatures: the smaller thegrain size, the higher the flow stress at lower temper-ature and the lower the flow stress at higher temper-ature. Therefore, a crossover temperature Tc can bedefined where the effect is reversed.[4] If the formingtemperature is higher and the service temperature islower than Tc, a small grain size is beneficial for bothforming and service performance. Also, if a second-phase particle can be precipitated at a temperaturelower than the forming temperature, it may bebeneficial for both high-temperature formability andlow-temperature service strength. Precipitates in themicrostructure of aluminum alloys are such second-phase particles. Therefore, fine grains and precipitatesare important microstructural characteristics allowingthe possible combination of high service strength andgood high-temperature formability.Superplastic aluminum alloys having a fine, equiaxed

microstructure with a grain size of approximately 10 lmare typical of these materials.[5] However, superplasticaluminum alloys are always produced using complicatedthermomechanical processes or powder metallurgy

HANLIANG ZHU, Research Fellow, and ARNE K. DAHLE,Professor, are with Materials Engineering, ARC CoE for Design inLight Metals, The University of Queensland, Brisbane, QLD 4072,Australia. Contact e-mail: [email protected] AMIT K. GHOSH,Professor, is with the Department of Materials Science and Engineer-ing, The University of Michigan, Ann Arbor, MI 48109.

Manuscript submitted March 31, 2008.Article published online January 14, 2009

598—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A

technology, which greatly increases costs. Also, super-plasticity normally occurs at high-temperature (0.75 T/Tm) and low strain rates (near 10�4 s�1).[6] Hence, thesuperplastic forming process (SPF) places high demandson forming equipment, shortens the lives of tools, andlowers production rates, leading to a high cost ofmanufactured components. All of these drawbacks limitthe applications of superplastic materials and SPFprocesses in industry.

Recent research has found that several Al-Mg–basedalloys obtained by casting as well as simple thermome-chanical processing have tensile elongations of greaterthan 100 pct at higher strain rates and lower tempera-tures than typical superplastic materials. This elongationgives sufficient ductility to manufacture the materialsinto intricate sheet components in industrial practices.[7]

The strength of these Al-Mg alloys can be improvedfurther by the addition of certain elements that refine thegrain size and form additional precipitates in themicrostructure of the alloys.

Scandium is one such element added to aluminumalloys.[8,9] When the concentration of Sc in the Al alloysexceeds a critical level, coarse primary Al3Sc particlesmay form. These primary Al3Sc particles can act aspotent nucleation sites for aluminum grains, effectivelyrefining the grain size during casting.[10] On the otherhand, a large supersaturation of Sc in a-Al can beachieved if the cooling rate during casting is sufficientlyhigh. Because the solid solubility of Sc decreases rapidlywith temperature,[11] fine Al3Sc dispersoids can preci-pitate from the supersaturated solid solution when agedat lower temperatures, providing an improvement instrength due to precipitation strengthening. The Al3Scdispersoids can control the grain growth during ther-momechanical treatments, further refining the grainsize. However, fine Al3Sc precipitates are not stable athigh temperatures due to the similarity of the latticeparameter and structure of Al3Sc to that of the Almatrix.[12,13] The high-temperature instability of theAl3Sc dispersoids is even worse in Al-Mg alloys, whereMg increases the lattice parameter of the Al matrix andprovides a better match with Al3Sc, further decreasingthe driving force for coarsening of the Al3Sc parti-cles.[13] However, zirconium, when added to Al-Mgalloys containing Sc, can partially substitute for Sc,forming coherent L12 Al3(Sc, Zr) precipitates. Themuch lower diffusivity of Zr compared to Sc inaluminum helps to stabilize the Al3(Sc, Zr) precipitatesup to at least 350 �C.[11–14] Therefore, by adding Zr andSc in combination, a dense and homogeneous distribu-tion of thermally stable Al3(Sc, Zr) precipitates isobtained during annealing.[11] Zinc has good solidsolubility and can form precipitates under appropriate

conditions, allowing both solid solution strengtheningand precipitation strengthening. However, the additionof trace amounts of Zn and Sc in combination in Al-Mg–based alloys has not been investigated. In thisstudy, Zn and Sc are added to an Al-Mg alloysingularly or simultaneously. The microstructure andhot formability of these alloys are compared, and therole of Zn and Sc additions is discussed. The aim of thisinvestigation is to develop a new high-strength alumi-num alloy with good formability at a relatively lowertemperature and a higher strain rate.

II. MATERIALS AND EXPERIMENTALPROCEDURE

Four Al-Mg sheet alloys with various concentrationsof Sc and Zn were used in this study, and theirchemical compositions are listed in Table I. The as-received (hot-rolled) alloy sheets had a thickness of8.0 mm. They were warm rolled at 180 �C to a finalthickness of 1.0 mm, giving a reduction ratio of87.5 pct. The rolled sheets were then recrystallized at450 �C for 60 minutes.Tensile specimens were cut along the rolling and

transverse directions. Based on the expected tempera-ture regime for warm forming operations, warm tensiletest temperatures were selected as 275 �C, 300 �C, and350 �C. Warm tensile tests were carried out at strainrates of 0.015, 0.15, and 1.5 s�1. Superplastic tensiletests were conducted at 520 �C and 0.002 s�1. All tensiletests were performed on an INSTRON* 4505 testing

machine with the data acquisition obtained by a digitalinterface board using a specialized computer program.The original raw data were then processed to calculatetrue stress/true strain relations and major mechanicalproperties.For the biaxial stretch tests, the sheets were cut to

rectangular blank samples of a size of L 9 W =200 mm 9 140 mm with L in rolling direction. Boronnitride powder was used as the lubricant, and itwas sprayed on the blanks and baked to a dry condition.Forming tests were performed on a heated rectangulardie-punch device, which was mounted on an Instron1116 testing machine with 250 kN capacity. The cross-sectional area was 110 mm 9 50 mm for the die cavityand 100 mm 9 40 mm for the punch. Both the dieedge and the punch had a radius of about 5 mm.

Table I. Chemical Composition of Al-Mg Sheet Alloys in this Study, Weight Percent

Element Fe Si Mg Mn Cr Zr Cu Sc Zn Al

Alloy 1 0.19 0.13 3.29 0.18 0.19 0.18 0.49 — — balAlloy 2 0.15 0.11 3.48 0.19 0.19 0.13 0.52 0.32 — balAlloy 3 0.15 0.1 3.48 0.19 0.21 0.18 0.50 <0.01 0.64 balAlloy 4 0.16 0.11 3.39 0.18 0.21 0.17 0.48 0.33 0.76 bal

*INSTRON is a trademark of Instron Company, Norwood, MA.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—599

Biaxial stretch tests were conducted at 350 �C and afixed crosshead speed of 10 mm/s. Load vs punchdisplacement curves were recorded using an X-Y datarecorder, and the data were used to obtain the depth of aformed part at peak load (where necking occurred onthe sheet). This part depth was used as a measure ofbiaxial formability.

To observe the grain structures clearly, all thespecimens for optical microscopy (OM) were aged at160 �C for 20 hours after the recrystallization treatmentand then electrolytically polished in a perchloric acidsolution. Specimens were also prepared for electronbackscatter diffraction (EBSD) and backscatter electronsignals (BES), which was done using a JSM** 6460LA

scanning electron microscope (SEM). Those specimensfor the investigation of the intermetallic particles wereetched in 0.5 pct HF for 3 seconds, and those for EBSD

analysis were polished electrolytically at 20 V for10 seconds in STRUERS� A2 solution.

III. RESULTS

A. Microstructure

The grain structures of the four aluminum alloys areillustrated in Figure 1. It is evident that the additions ofZn or Sc have a great effect on the shape and size of thegrains. As shown in Figure 1(a), in alloy 1 without Zn orSc addition, most grains exhibit an equiaxed shape,indicating a fully recrystallized state. However, in alloy3, where Zn is the only addition (Figure 1(c)), the grainsare slightly elongated and have a smaller grain size thanalloy 1. This is consistent with previous findings that theaddition of Zn can refine the grain size of rolled Al-Mgalloys.[15] In contrast, the microstructure of alloy 2,where Sc is the only addition (Figure 1(b)), and alloy 4,

Fig. 1—Microstructure in the thickness direction of various sheets: (a) alloy 1, (b) alloy 2, (c) alloy 3, and (d) alloy 4. Arrows in (d) indicatelarge second-phase particles.

**JSM is a trademark of Japan Electron Optics Ltd., Tokyo, Japan.

�STRUERS is a trademark of Struers Inc., Cleveland, OH.


which has both Zn and Sc additions (Figure 1(d)), thegrains are severely elongated in the rolling direction,which is typical of rolled aluminum alloys.[16] Theseresults demonstrate that Sc has an antirecrystallizationeffect in aluminum alloys.[9] Furthermore, the width ofthe grains is much smaller in alloy 4 than in the otherthree alloys.

In order to clearly reveal the details of the differencein the grain structure of the alloys with or without Scaddition, EBSD was conducted on alloys 3 and 4. Thegrain mappings of the two alloys are shown in Figure 2.The microstructure of the two alloys consists of elon-gated and equiaxed grains. The microstructure of alloy 4shows large amounts of fine grains and subgrains with asize of less than 2 lm. These fine grains and subgrainsare not uniformly distributed in the microstructure butsurround larger grains (Figure 2(b)).

Large second-phase particles can be seen in the OMmicrostructure in Figure 1. Most of the second-phaseparticles are located at the grain boundaries, as seen in

the backscattered SEM images of alloys 3 and 4 ofFigure 3. Two kinds of second-phase particles areobserved: coarse cuboidal shapes with a size up to10 lm and fine plate shapes with a size up to 2 lm.Electron dispersive spectroscopy (EDS) examination ofthe second-phase particles shows that the stochiometricratio of Al and Zr in alloy 3 and Al and (Sc+Zr+Mg+Zn) in alloy 4 are close to 3:1. It is likely thatthe coarse cuboidal particles in alloy 3 are Al3Zr(Figure 4(a)) and, in alloy 4, Zn- and Mg-enrichedAl3(Sc, Zr) (Figure 4(c)). The amount of coarse particlesfound in the Sc-containing alloys is much larger thanthose in the Sc-free alloys. The fine plate phases areFe-rich intermetallic particles (Figures 4(b) and (d)).

B. Warm Tensile Test Results

True stress/true strain curves of the investigated alloystested at warm temperatures ranging from 275 �C to350 �C and strain rates of 0.015 to 1.5 s�1 are presented

Fig. 2—Grain mapping of (a) alloy 3 and (b) alloy 4. Arrows indicate fine grain zones.

Fig. 3—Morphology of coarse primary intermetallic particles of (a) alloy 3 and (b) alloy 4. White phases are intermetallic particles.


in Figures 5 and 6. It is found that the flow stress andfracture strain are dependent on temperature and strainrate. At lower temperatures of 275 �C and 300 �C orhigher strain rates of 0.15 and 1.5 s�1, the curvesexhibit a peak in the yield stress that increases withstrain rate. However, at 350 �C and 0.015 s�1, alloy 1(Figure 5(e)) exhibits an initial stress transient—theflow stress increases rapidly, reaches a maximum at astrain of about 0.03, and then rapidly decreases to asteady-state value before decreasing prior to fracture.Whereas, in alloy 4 (Figure 6(f)), the flow stressincreases generally and achieves the maximum valueat a strain of 0.2, indicating that the initial stresstransient is absent. For alloys 2 (Figure 5(f)) and 3(Figure 6(e)), the flow stress quickly reaches a maxi-mum value at a small strain (less than 0.05) and thestress-strain behavior of alloys 2 and 3 are similar tothat of alloy 1. The variation in elongation with the testconditions is illustrated in Figure 7. It is evident thatelongation increases with temperature but decreaseswith strain rate.

At the test temperature of 275 �C, almost all the stresslevels and elongations of alloy 2 are lower than those ofalloy 1 (except at 0.015 s�1), as shown in Figures 5 and7. It can also be seen that as the test temperatureincreases from 300 �C to 350 �C, the stress levels of alloy2 approach that of alloy 1, but the elongations of alloy 2remain much lower than those of alloy 1. These resultsindicate that the addition of Sc does not increase the

strength of the alloy but can decrease the ductility atwarm test temperatures. Alloy 3 (containing Zn) haslower elongations than alloy 1 for all test conditions;however, the flow stress levels of alloy 3 change from alower value than those of alloy 1 to a higher value as thetest temperature increases. These results show that, withincreasing test temperature, the strengthening effect ofZn in the Al-Mg sheet alloy becomes pronounced but isaccompanied by a dramatic decrease in the ductility.Compared to the other three alloys, alloy 4 exhibits ahigher peak stress and a lower elongation at almost alltest conditions except at 350 �C and 0.015 s�1, whichresult in a larger elongation of 130 pct. Therefore, thecombination addition of Zn and Sc in the Al-Mgsheet alloy cannot only significantly increase thestrength but also enhance the ductility at an appropriatedeformation condition.

C. Superplastic Tensile and Biaxial Stretch Test Results

The superplastic tensile test and biaxial stretch testwere carried out to investigate hot formability of thealloys at higher temperatures and complex stress condi-tions. The test results for the superplastic elongation andcup depth are show in Figure 8. All four alloys achievean elongation of over 150 pct at the superplastic testingcondition. Alloy 4 reaches an elongation of 302.3 pct,reflecting a superplastic characteristic, and exhibits thebest formability among the four investigated alloys. For

Fig. 4—Chemical compositions of large intermetallic phases: (a) and (b) in alloy 3 and (c) and (d) in alloy 4.


Fig. 5—True stress/true strain curves of alloys 1 and 2 at different temperatures and strain rates: (a), (c), and (e) alloy 1 and (b), (d) and(f) alloy 2.


the other three alloys, elongation decreases from alloy 1to alloy 2 and to alloy 3, which is consistent with theuniaxial tensile test results at 350 �C and 0.015 s�1.

Under the biaxial stretch condition, the anisotropy ofmicrostructure between the rolling and transverse directionsshould have an effect on the deformation behavior.[17]

Fig. 6—True stress/true strain curves of alloys 3 and 4 at different temperatures and strain rates: (a), (c), and (e) alloy 3 and (b), (d), and(f) alloy 4.


The plastic strain ratio of width to thickness, R, wasmeasured after superplastic testing. The R values foralloys 1 and 4 are 0.917 and 0.430, respectively,indicating that alloy 4 has more anisotropic ductilitythan alloy 1. This anisotropic ductility results in aslightly lower cup depth in alloy 4 than in alloy 1,though it has better uniaxial ductility. However, thebiaxial stretch formability of alloy 4 (with both Zn andSc additions) is still better than those of alloys 2 and 3,which have Sc and Zn additions, respectively.

IV. DISCUSSION

A. Effect of Zn and Sc Additions on Microstructure

When Zn is added to Al-Mg sheet alloys in the level ofabout 0.7 wt pct, most of the zinc is in solid solution ina-Al according to the Al-Zn-Mg phase diagram.[18]

Fig. 7—Elongation of various alloys: (a) alloy 1, (b) alloy 2, (c) alloy 3, and (d) alloy 4.

Fig. 8—Superplastic tensile and biaxial stretch test results of investi-gated aluminum alloys.


The Zn in solid solution has an antirecrystallizationability due to the solute drag effect on grain boundarymigration.[19] Furthermore, no Zn-rich second-phaseparticles were identified in alloy 3 (Al-Mg-Zn). Hence,the only effect of Zn is solute drag on grain boundarymigration during thermal exposure. This effect preventsthe grain growth and recrystallization process duringannealing, resulting in the finer and more elongatedgrains in alloy 3 (Figure 1(c)) compared to those in alloy1 (Figure 1(a)).

The solid solubility of Sc in a-Al of the binary Al-Scsystem is considered to be only approximately0.025 wt pct at 430 �C[9] and 0.07 wt pct at 500 �C.[13]The solid solution of Sc can also prevent the graingrowth and recrystallization process during the thermo-mechanical treatments. Moreover, Mg and Sc mutuallyreduce each other’s solid solubility in the Al alloy;hence, the equilibrium concentration of Sc is expected todecrease as temperature decreases or with increasing Mgcontent.[9] Consequently, the addition of about0.3 wt pct Sc in the Al-Mg–based alloys can result inthe formation of coarse primary intermetallic particlesduring casting. Such particles, with a size of up to10 lm, are seen in Figures 1(d) and 3(b) and are likely tobe Al3(Sc, Zr) enriched in Zn and Mg (Figure 4(c)).

According to previous studies, these coarse particleshave a complex effect on the microstructure. First,Al3(Sc, Zr) intermetallic particles formed in the melt canact as potent nucleation sites for aluminum grains,[9]

refining the grain structure during casting. Second,thermomechanical treatments result in large amountsof deformation-induced stored energy surrounding par-ticles with a size above 1 lm. This increased storedenergy greatly promotes the nucleation of recrystallizedgrains; hence, coarse intermetallics can act as efficientnuclei for recrystallization (also known as particle-stimulated nucleation (PSN)[20]). This PSN effect canextend for more than ten particle diameters.[21,22] Con-sequently, fine grains and subgrains are formed in zonesaround the particles. These zones are readily seen in theEBSD mapping, as shown in Figure 2(b). Third, theprimary intermetallic particles can also limit graingrowth during thermomechanical treatments. Combinedwith the solute drag effect due to Sc and Zn on grainboundary migration, much more elongated and thinnergrains are formed in the microstructure of alloy 4 thanthat of alloy 1.

B. Effect of Zn and Sc Additions on Hot Formability

The Zn in a-Al increases the flow stress and decreasesthe ductility of alloy 3 compared to alloy 1 due to solidsolution strengthening; this effect is more pronounced atthe elevated test temperatures. As mentioned in SectionA, the solid solubility of Sc in a-Al of the binary Al-Scsystem is only approximately 0.025 wt pct at 430 �C,[9]so at the warm forming temperatures of and below350 �C, the amount of Sc in solid solution is expected tobe lower than this value for alloys 2 and 4. When Sc issingularly added to alloy 2, the solid solution strength-ening is not significant, and the flow stress and elonga-tion between alloys 1 and 2 is similar. However, when

trace amounts of Sc are added in addition to Zn, as inalloy 4, the flow stress is higher and the ductility is lower(Figures 6(b), (d), and (f)) than the other three alloys atalmost all deformation conditions (the exception being350 �C and 0.015 s�1). This is because the change in thelattice parameter of a-Al due to Zn addition becomesmore pronounced with the addition of trace amounts ofSc, greatly increasing the solid solution strengthening.Second-phase particles have a strong influence on

formability of Al sheet alloys. The second-phase parti-cles can nucleate voids or cracks and influence thegrowth and coalescence of these voids or cracks. Themost common modes of initiation are the complete orpartial interfacial decohesion of the second-phase par-ticle and the surrounding matrix, or the fracture ofsecond-phase particles.[23] If the second-phase particlehas a rounded shape or high fracture strength, cracksusually initiate at the particle/matrix interface.[24] In theinvestigated alloys, the Fe-rich intermetallic particleshave a rounded shape (Figure 3) and the Al3(Sc, Zr) orAl3Zr intermetallics have a high fracture strength; hence,interface decohesion is likely to be the major mode ofcrack nucleation. Furthermore, voids usually nucleatefirst on larger particles that contain small internal orinterfacial defects. Larger particles also induce morerapid decohesion with increasing strain, because theconstrained zone of plasticity at the interface becomeslarger.[25] Thus, coarser second-phase particles are moredeleterious to the mechanical properties than finer ones.The addition of Sc in the investigated alloy producesmuch coarser and larger number of Al3(Sc, Zr) particlesand no significant effect on the morphology anddistribution of the Fe-rich intermetallic particles. Con-sequently, it is believed that the larger number of coarseAl3(Sc, Zr) particles nucleate voids early and promotethe growth and coalescence of the voids, contributinggreatly to the decrease in both flow stress and ductility inthe Sc-containing alloys. This is the likely cause for alloy2 having inferior mechanical properties at the elevatedtest temperatures. However, the flow stress at lower testtemperatures or higher strain rates and the ductilityat 350 �C and 0.015 s�1 for alloy 4 increase. This isbecause the contributions from solid solution strength-ening and other mechanisms compensate the loss of theflow stress and ductility due to large second-phaseparticles.Of the three elevated temperature deformation mech-

anisms, GBS, slip creep, and diffusion creep, thecontribution from diffusion creep is considered to benegligible.[26] At relevantly low temperature, i.e., T<Tc,the grain-boundary bond is strong and slip creep is apredominant deformation mode. The slip creep processis found to be solute drag creep (SDC) in Al-Mg–basedalloys.[1] However, at higher temperature, i.e., T>Tc,the grain-boundary bonds are weakened, and yet shearand normal stresses must be transferred across them.Thus, inelastic displacement in grain-boundary regionsmust be larger than those within the grain interior,which may even remain elastic. This view is essentiallythe same as the grain mantle vs core deformationmodel.[27] Thus, the elevated-temperature constitutivebehavior has been described by assuming that both


mantle and core deformation contribute to the overalldeformation, that is,[27]

et� ¼ em

� þ ec� ½1�

where et�is the total deformation rate; and em

�and ec

�

are contributions from the mantle and core deforma-tion, respectively. The mantle contribution to the ten-sile strain rate can be written as[27]

em� ¼ A=d2ðr� r0Þ1þq ½2�

where A is constant for a fixed temperature, r is tensilestress, and r0 is the tensile equivalent of s0. The value ofq may vary between 0.1 and 0.4 for a well-recrystallizedgrain structure containing a small volume fraction ofdispersoids and 0.3 and 1 for a recovered subgrainmicrostructure containing a high volume fraction ofdispersoids.

The strain rate contributed by grain core, includingaccelerated deformation at grain corners, can beexpressed as

ec� ¼ ðKþ A1=d

3Þrn ½3�

where K is K1e�Q=RT, and A1 is a constant relating to

stress concentration and enhanced dislocation creep atgrain corners. The term d3 arises due to the number ofgrain corners present per unit volume.

In the preceding equations, GBS and its accommoda-tion processes are assumed to take place in the mantleregion of the grains, and the SDC process occurs withinthe grain cores. The GBS controlling mechanism gener-ally requires a stable grain size of less than 10 lm, highdeformation temperatures (0.75 T/Tm), and low strainrates (near 10�4 s�1).[1] In coarse-grained materials(grain size: >10 lm), the contribution from grain coresis larger than that from grain mantles and, thus, SDC isthe controlling deformation mechanism at various con-ditions. The SDC may still be the controlling deforma-tion mechanism in fine-grained materials (grain size:<10 lm) at higher strain rates or lower temperatures.

The deformation temperature and strain rate can berepresented by one parameter, the Zener–Hollomonparameter, Z=e

�expðQ=RTÞ, and hence deformation

behavior can be investigated by a plot of Z values as afunction of modulus-compensated stress, r/E. Plots ofthe four investigated alloys are shown in Figure 9, withQ = 110 kJ/mol.[1] The Z values of alloys 1, 2, and 3 canbe fit to a straight line with n = 8, while those of alloy 4are plotted at a higher stress level due to solid solutionstrengthening, and they show a significant decrease in theflow stress at lower Z values (350 �C, 0.015 s�1).

The lower m value (1/n< 0.2) at the higher Z valueindicates that greater contribution to the overall defor-mation is from grain cores, and SDC is the controllingdeformation mechanism for the investigated alloys.[1]

However, at the lower Z value (350 �C, 0.015 s�1), thesefine grains in alloy 4 can easily slide due to the decreasein r0 of Eq. [2];

[28] thus, they can increase the contribu-tion from grain mantle as well as from GBS duringdeformation. Also, GBS can occur in grain mantles oflarger grains when the mobility of the larger grains

surrounded by the finer grains is greatly increased.Moreover, according to Eq. [3], the fine grains alsoincrease the contribution from grain cores. Conse-quently, the increased contribution from both grainmantles and grain cores increases the ductility of alloy 4compared to the other three alloys. The increasedcontribution from grain mantles is larger than thatfrom grain cores in alloy 4. This is demonstrated by thedifferent shapes of the flow curves of alloy 4 comparedto the other alloys at the same deformation condition of350 �C and 0.015 s�1. The flow curves for alloys 1through 3 exhibit a sharp peak stress after reaching theyield point, which is the characteristic of slip creep,[29]

while strain hardening is apparent during plastic defor-mation for alloy 4 at the same deformation condition.The strain hardening is consistent with dynamic graingrowth in GBS.[1] Moreover, the contribution from GBSalso enables this alloy to have better biaxial stretchformability at 350 �C. When the deformation tempera-ture is further increased to 520 �C, superplastic elonga-tion of over 300 pct can be achieved. Thus, thecombined addition of Sc and Zn in Al-Mg sheet alloyscan generate an excellent combination of low-temperaturestrength and high-temperature formability.

V. SUMMARY

The additions of Sc or Zn in Al-Mg sheet alloys formsolid solutions in a-Al and large Al3(Sc, Zr) intermetallic

Fig. 9—Zener–Hollomon parameter as a function of modulus-compensated stress.


particles. The solid solution of Sc or Zn in a-Al canprevent the grain growth and recrystallization processand thereby refine the grain size in thermomechnicaltreatments prior to hot deformation. Also, large parti-cles lead to particle stimulated nucleation of recrystal-lization, resulting in the formation of numeroussubgrains and fine grains around the particles. Theselarge particles can limit grain growth during thermo-mechnical treatments, resulting in a smaller grain size.

During tensile testing at temperatures ranging from275 �C to 350 �C and strain rates of 0.015 to 1.5 s�1,solid solution strengthening of Zn or Sc increases theflow stress and decreases ductility. In the alloys con-taining Sc, large intermetallic particles cause crackinitiation during warm deformation, further decreasingthe flow stress and ductility of these alloys. Further-more, during warm deformation, SDC in grain coresand GBS in grain mantles occur simultaneously. How-ever, SDC is the controlling deformation mechanism,demonstrated by the strain rate sensitivity coefficient ofbelow 0.2, in alloys 1 through 3 and alloy 4 at lower testtemperatures and higher strain rates. With increasingdeformation temperature and decreasing strain rate, thecontribution from GBS increases. At 350 �C and0.015 s�1, the numerous fine grains in alloy 4 (contain-ing both Sc and Zn additions) further promote GBS,increasing the ductility and resulting in an elongation of130 pct. The benefit of the combined additions of Sc andZn has also been verified by the results of superplastictensile tests and biaxial stretch tests. On the basis of theexperimental results, the combined addition of Sc andZn in Al-Mg may be an important method of producinga high-strength sheet alloy with excellent high-temperatureformability.

REFERENCES1. T.R. McNelley, K.O. Ishi, A.P. Zhilyaev, S. Swaminathan, P.E.

Krajewski, and E.M. Taleff: Metall. Mater. Trans. A, 2008,vol. 39A, pp. 50–64.

2. T.G. London: Z. Metall., 2005, vol. 96, pp. 522–31.

3. O.D. Sherby and J. Wadsworth: Prog. Mater. Sci., 1989, vol. 33,pp. 169–221.

4. J. Pan, J. Tong, and M. Tian: Fundamental of Materials Science,Qinghua University Publisher, Beijing, 2002, pp. 176–86.

5. B.Z. Bai and K.A. Padmanabhan: Mater. Sci. Forum, 1997,vols. 243–245, pp. 191–96.

6. L.P. Troeger and E.A. Starke, Jr.: Mater. Sci. Eng. A, 2000,vol. A277, pp. 102–13.

7. E.M. Taleff, P.J. Nevland, and P.E. Krajewski: Metall. Mater.Trans. A, 2001, vol. 32A, pp. 1119–30.

8. J. Royset and N. Ryum: Scripta Mater., 2005, vol. 52, pp. 1275–79.

9. J. Røyset and N. Ryum: Int. Mater. Rev., 2005, vol. 50, pp. 19–44.10. K.B. Hyde, A.F. Norman, and P.B. Prangnell: Acta Mater., 2001,

vol. 49, pp. 1327–37.11. B. Forbord, H. Hallem, J. Røyset, and K. Marthinsen: Mater. Sci.

Eng. A, 2008, vol. 475, pp. 241–48.12. M.M. Sharma, M.F. Amateau, and T.J. Eden: J. Alloys Compd.,

2006, vol. 416, pp. 135–42.13. K.L. Kendig and D.B. Miracle: Acta Mater., 2002, vol. 50,

pp. 4165–75.14. R.A. Karnesky, L. Meng, and D.C. Dunand: Acta Mater., 2007,

vol. 55, pp. 1299–308.15. M.A.G. Bernal, D.H. Silva, and V.S. Rangel: J. Mater. Sci., 2007,

vol. 42, pp. 3958–63.16. J. Liu and J.G. Morris: Mater. Sci. Eng. A, 2004, vol. 385,

pp. 342–51.17. G.E. Dieter: Mechanical Metallurgy, McGraw-Hill Book

Company, New York, NY, 1986.18. H. Loffler, I. Kov, and J. Lendvai: J. Mater. Sci., 1983, vol. 18,

pp. 2215–40.19. T. Kobayashi: Mater. Sci. Eng. A, 2000, vol. 286, pp. 333–41.20. T.G. Nieh, L.M. Hsiung, J. Wadsworth, and R. Kaibyshev: Acta

Mater., 1998, vol. 46, pp. 2789–2800.21. J.S. Vetrano, S.M. Bruemmer, L.M. Pawlowski, and I.M.

Robertson: Mater. Sci. Eng. A, 1997, vol. 238, pp. 101–07.22. R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juul

Jensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen,and A.D. Rollett: Mater. Sci. Eng. A, 1997, vol. 238, pp. 219–74.

23. D. Lassance, D. Fabregue, F. Delannay, and T. Pardoen: Prog.Mater. Sci., 2007, vol. 52, pp. 62–129.

24. H. Toda, T. Kobayashi, and A. Takahashi: Mater. Sci. Eng. A,2000, vol. 280, pp. 69–75.

25. D.H. Bae and A.K. Ghosh: Acta Mater., 2002, vol. 50, pp. 511–23.26. T.G. London: J Mater. Sci., 2006, vol. 41, pp. 597–609.27. A. Ghosh: ASM Handbook, vol. 14A, Constitutive Equations,

Metalworking: Bulk Forming, ASM INTERNATIONAL, Mate-rials Park, OH, 2005, pp. 563–86.

28. I. Charit and R.S. Mishra: Acta Mater., 2005, vol. 53, pp. 4211–23.29. D.H. Bae and A.K. Ghosh: Acta Mater., 2000, vol. 48,

pp. 1207–24.


!!Effect of sc and zn additions on microstructure and hot formability of al mg sheet alloys

Business

Transcript of !!Effect of sc and zn additions on microstructure and hot formability of al mg sheet alloys