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


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  • 1. 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 examination showed that the Sc addition greatly refined the grain size, especially in the Zn-containing alloy where large amounts of subgrains and fine grains were formed. Meanwhile, a number of large primary 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 from 275 °C to 350 °C and at strain rates of 0.015 to 1.5 sÀ1. The test results showed that, in the alloys with the single addition of Zn or Sc, Zn slightly increased the flow stress but decreased the ductility and Sc worsened both the flow stress and the ductility. However, in the alloy with the combined addition of Zn and Sc, the flow stress was significantly increased at almost all testing conditions and the ductility was also increased at a higher temperature of 350 °C and lower strain rate of 0.015 sÀ1. The results of superplastic tensile tests and biaxial stretch tests dem- onstrated that the alloy with both Zn and Sc additions exhibited good high-temperature formability. The effect of Zn and Sc additions on the microstructure and the hot formability is discussed. DOI: 10.1007/s11661-008-9728-6 Ó The Minerals, Metals & Materials Society and ASM International 2009 I. INTRODUCTION grain size during the high-temperature deformation, the presence of second-phase particles at the grain bound- WITH the widespread application of lightweight aries is required to achieve superplasticity.[3] If thealuminum alloy products in the automotive industry, deformation temperature is decreased, the strain ratealloys with a combination of good formability, high increases, or, in coarse-grained materials, dislocationperformance, and low cost are required. Especially for slip becomes important and deformation becomesapplications where conditions of high temperature or concentrated within the grains.[3] If second-phase par-stress are involved, superior creep resistance or high ticles are present within the grains, they can retardstrength is an essential requirement. Because the dislocation motion, causing strengthening. Grainincrease of strength is always accompanied by a boundaries can also act as barriers to deformation atsubstantial decrease in ductility, the ambient tempera- low temperatures; hence, a small grain size will causeture ductility and formability of these alloys is always an increase in strength and a decrease in ductility.very limited, causing the conventional manufacturing Overall, grain size has an opposite effect on the flowoperations such as rolling, forging, or drawing to be stress at high and low temperatures: the smaller thedifficult, and thereby restricting their use.[1] As a result, grain size, the higher the flow stress at lower temper-forming technologies at intermediate and high temper- ature and the lower the flow stress at higher temper-atures such as warm forming, hot forging, and super- ature. Therefore, a crossover temperature Tc can beplastic forming appear to be more promising routes for defined where the effect is reversed.[4] If the formingfabricating structural components using these alloys. temperature is higher and the service temperature is At elevated temperatures, the deformation mecha- lower than Tc, a small grain size is beneficial for bothnisms occurring during plastic flow can be divided into forming and service performance. Also, if a second-three distinct classes depending upon whether they are phase particle can be precipitated at a temperatureintragranular dislocation movement, grain boundary lower than the forming temperature, it may besliding (GBS), or diffusional flow.[2] Generally, at high beneficial for both high-temperature formability anddeformation temperatures and low strain rates, GBS is low-temperature service strength. Precipitates in thethe predominant mechanism for fine-grained superplas- microstructure of aluminum alloys are such second-tic materials (grain size:<10 lm).[1] To maintain the fine phase particles. Therefore, fine grains and precipitates are important microstructural characteristics allowing HANLIANG ZHU, Research Fellow, and ARNE K. DAHLE, the possible combination of high service strength andProfessor, are with Materials Engineering, ARC CoE for Design in good high-temperature formability.Light Metals, The University of Queensland, Brisbane, QLD 4072, Superplastic aluminum alloys having a fine, equiaxedAustralia. Contact e-mail: AMIT K. GHOSH, microstructure with a grain size of approximately 10 lmProfessor, is with the Department of Materials Science and Engineer-ing, The University of Michigan, Ann Arbor, MI 48109. are typical of these materials.[5] However, superplastic Manuscript submitted March 31, 2008. aluminum alloys are always produced using complicated Article published online January 14, 2009 thermomechanical processes or powder metallurgy598—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 2. technology, which greatly increases costs. Also, super- conditions, allowing both solid solution strengtheningplasticity normally occurs at high-temperature (0.75 T/ and precipitation strengthening. However, the additionTm) and low strain rates (near 10À4 sÀ1).[6] Hence, the of trace amounts of Zn and Sc in combination in Al-superplastic forming process (SPF) places high demands Mg–based alloys has not been investigated. In thison forming equipment, shortens the lives of tools, and study, Zn and Sc are added to an Al-Mg alloylowers production rates, leading to a high cost of singularly or simultaneously. The microstructure andmanufactured components. All of these drawbacks limit hot formability of these alloys are compared, and thethe applications of superplastic materials and SPF role of Zn and Sc additions is discussed. The aim of thisprocesses in industry. investigation is to develop a new high-strength alumi- Recent research has found that several Al-Mg–based num alloy with good formability at a relatively loweralloys obtained by casting as well as simple thermome- temperature and a higher strain rate.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 materials II. MATERIALS AND EXPERIMENTALinto intricate sheet components in industrial practices.[7] PROCEDUREThe strength of these Al-Mg alloys can be improvedfurther by the addition of certain elements that refine the Four Al-Mg sheet alloys with various concentrationsgrain size and form additional precipitates in the of Sc and Zn were used in this study, and theirmicrostructure of the alloys. chemical compositions are listed in Table I. The as- Scandium is one such element added to aluminum received (hot-rolled) alloy sheets had a thickness ofalloys.[8,9] When the concentration of Sc in the Al alloys 8.0 mm. They were warm rolled at 180 °C to a finalexceeds a critical level, coarse primary Al3Sc particles thickness of 1.0 mm, giving a reduction ratio ofmay form. These primary Al3Sc particles can act as 87.5 pct. The rolled sheets were then recrystallized atpotent nucleation sites for aluminum grains, effectively 450 °C for 60 minutes.refining the grain size during casting.[10] On the other Tensile specimens were cut along the rolling andhand, a large supersaturation of Sc in a-Al can be transverse directions. Based on the expected tempera-achieved if the cooling rate during casting is sufficiently ture regime for warm forming operations, warm tensilehigh. Because the solid solubility of Sc decreases rapidly test temperatures were selected as 275 °C, 300 °C, andwith temperature,[11] fine Al3Sc dispersoids can preci- 350 °C. Warm tensile tests were carried out at strainpitate from the supersaturated solid solution when aged rates of 0.015, 0.15, and 1.5 sÀ1. Superplastic tensileat lower temperatures, providing an improvement in tests were conducted at 520 °C and 0.002 sÀ1. All tensilestrength due to precipitation strengthening. The Al3Sc tests were performed on an INSTRON* 4505 testingdispersoids can control the grain growth during ther-momechanical treatments, further refining the grain *INSTRON is a trademark of Instron Company, Norwood, MA.size. However, fine Al3Sc precipitates are not stable athigh temperatures due to the similarity of the latticeparameter and structure of Al3Sc to that of the Al machine with the data acquisition obtained by a digitalmatrix.[12,13] The high-temperature instability of the interface board using a specialized computer program.Al3Sc dispersoids is even worse in Al-Mg alloys, where The original raw data were then processed to calculateMg increases the lattice parameter of the Al matrix and true stress/true strain relations and major mechanicalprovides a better match with Al3Sc, further decreasing properties.the driving force for coarsening of the Al3Sc parti- For the biaxial stretch tests, the sheets were cut tocles.[13] However, zirconium, when added to Al-Mg rectangular blank samples of a size of L 9 W =alloys containing Sc, can partially substitute for Sc, 200 mm 9 140 mm with L in rolling direction. Boronforming coherent L12 Al3(Sc, Zr) precipitates. The nitride powder was used as the lubricant, and itmuch lower diffusivity of Zr compared to Sc in was sprayed on the blanks and baked to a dry condition.aluminum helps to stabilize the Al3(Sc, Zr) precipitates Forming tests were performed on a heated rectangularup to at least 350 °C.[11–14] Therefore, by adding Zr and die-punch device, which was mounted on an InstronSc in combination, a dense and homogeneous distribu- 1116 testing machine with 250 kN capacity. The cross-tion of thermally stable Al3(Sc, Zr) precipitates is sectional area was 110 mm 9 50 mm for the die cavityobtained during annealing.[11] Zinc has good solid and 100 mm 9 40 mm for the punch. Both the diesolubility and can form precipitates under appropriate edge and the punch had a radius of about 5 mm. Table I. Chemical Composition of Al-Mg Sheet Alloys in this Study, Weight PercentElement Fe Si Mg Mn Cr Zr Cu Sc Zn AlAlloy 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 balMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—599
  • 3. Biaxial stretch tests were conducted at 350 °C and a analysis were polished electrolytically at 20 V forfixed crosshead speed of 10 mm/s. Load vs punch 10 seconds in STRUERS  A2 solution.displacement curves were recorded using an X-Y datarecorder, and the data were used to obtain the depth of a  formed part at peak load (where necking occurred on STRUERS is a trademark of Struers Inc., Cleveland, OH.the sheet). This part depth was used as a measure ofbiaxial formability. III. RESULTS To observe the grain structures clearly, all thespecimens for optical microscopy (OM) were aged at A. Microstructure160 °C for 20 hours after the recrystallization treatment The grain structures of the four aluminum alloys areand then electrolytically polished in a perchloric acid illustrated in Figure 1. It is evident that the additions ofsolution. Specimens were also prepared for electron Zn or Sc have a great effect on the shape and size of thebackscatter diffraction (EBSD) and backscatter electron grains. As shown in Figure 1(a), in alloy 1 without Zn orsignals (BES), which was done using a JSM** 6460LA Sc addition, most grains exhibit an equiaxed shape, indicating a fully recrystallized state. However, in alloy **JSM is a trademark of Japan Electron Optics Ltd., Tokyo, Japan. 3, where Zn is the only addition (Figure 1(c)), the grains are slightly elongated and have a smaller grain size than alloy 1. This is consistent with previous findings that thescanning electron microscope (SEM). Those specimens addition of Zn can refine the grain size of rolled Al-Mgfor the investigation of the intermetallic particles were alloys.[15] In contrast, the microstructure of alloy 2,etched in 0.5 pct HF for 3 seconds, and those for EBSD 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.600—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 4. 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.which has both Zn and Sc additions (Figure 1(d)), the the backscattered SEM images of alloys 3 and 4 ofgrains are severely elongated in the rolling direction, Figure 3. Two kinds of second-phase particles arewhich is typical of rolled aluminum alloys.[16] These observed: coarse cuboidal shapes with a size up toresults demonstrate that Sc has an antirecrystallization 10 lm and fine plate shapes with a size up to 2 lm.effect in aluminum alloys.[9] Furthermore, the width of Electron dispersive spectroscopy (EDS) examination ofthe grains is much smaller in alloy 4 than in the other the second-phase particles shows that the stochiometricthree alloys. ratio of Al and Zr in alloy 3 and Al and (Sc + Zr + In order to clearly reveal the details of the difference Mg + Zn) in alloy 4 are close to 3:1. It is likely thatin the grain structure of the alloys with or without Sc the coarse cuboidal particles in alloy 3 are Al3Zraddition, EBSD was conducted on alloys 3 and 4. The (Figure 4(a)) and, in alloy 4, Zn- and Mg-enrichedgrain mappings of the two alloys are shown in Figure 2. Al3(Sc, Zr) (Figure 4(c)). The amount of coarse particlesThe microstructure of the two alloys consists of elon- found in the Sc-containing alloys is much larger thangated and equiaxed grains. The microstructure of alloy 4 those in the Sc-free alloys. The fine plate phases areshows large amounts of fine grains and subgrains with a Fe-rich intermetallic particles (Figures 4(b) and (d)).size of less than 2 lm. These fine grains and subgrainsare not uniformly distributed in the microstructure but B. Warm Tensile Test Resultssurround larger grains (Figure 2(b)). Large second-phase particles can be seen in the OM True stress/true strain curves of the investigated alloysmicrostructure in Figure 1. Most of the second-phase tested at warm temperatures ranging from 275 °C toparticles are located at the grain boundaries, as seen in 350 °C and strain rates of 0.015 to 1.5 sÀ1 are presentedMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—601
  • 5. Fig. 4—Chemical compositions of large intermetallic phases: (a) and (b) in alloy 3 and (c) and (d) in alloy Figures 5 and 6. It is found that the flow stress and strength of the alloy but can decrease the ductility atfracture strain are dependent on temperature and strain warm test temperatures. Alloy 3 (containing Zn) hasrate. At lower temperatures of 275 °C and 300 °C or lower elongations than alloy 1 for all test conditions;higher strain rates of 0.15 and 1.5 sÀ1, the curves however, the flow stress levels of alloy 3 change from aexhibit a peak in the yield stress that increases with lower value than those of alloy 1 to a higher value as thestrain rate. However, at 350 °C and 0.015 sÀ1, alloy 1 test temperature increases. These results show that, with(Figure 5(e)) exhibits an initial stress transient—the increasing test temperature, the strengthening effect offlow stress increases rapidly, reaches a maximum at a Zn in the Al-Mg sheet alloy becomes pronounced but isstrain of about 0.03, and then rapidly decreases to a accompanied by a dramatic decrease in the ductility.steady-state value before decreasing prior to fracture. Compared to the other three alloys, alloy 4 exhibits aWhereas, in alloy 4 (Figure 6(f)), the flow stress higher peak stress and a lower elongation at almost allincreases generally and achieves the maximum value test conditions except at 350 °C and 0.015 sÀ1, whichat a strain of 0.2, indicating that the initial stress result in a larger elongation of 130 pct. Therefore, thetransient is absent. For alloys 2 (Figure 5(f)) and 3 combination addition of Zn and Sc in the Al-Mg(Figure 6(e)), the flow stress quickly reaches a maxi- sheet alloy cannot only significantly increase themum value at a small strain (less than 0.05) and the strength but also enhance the ductility at an appropriatestress-strain behavior of alloys 2 and 3 are similar to deformation condition.that of alloy 1. The variation in elongation with the testconditions is illustrated in Figure 7. It is evident that C. Superplastic Tensile and Biaxial Stretch Test Resultselongation increases with temperature but decreaseswith strain rate. The superplastic tensile test and biaxial stretch test At the test temperature of 275 °C, almost all the stress were carried out to investigate hot formability of thelevels and elongations of alloy 2 are lower than those of alloys at higher temperatures and complex stress condi-alloy 1 (except at 0.015 sÀ1), as shown in Figures 5 and tions. The test results for the superplastic elongation and7. It can also be seen that as the test temperature cup depth are show in Figure 8. All four alloys achieveincreases from 300 °C to 350 °C, the stress levels of alloy an elongation of over 150 pct at the superplastic testing2 approach that of alloy 1, but the elongations of alloy 2 condition. Alloy 4 reaches an elongation of 302.3 pct,remain much lower than those of alloy 1. These results reflecting a superplastic characteristic, and exhibits theindicate that the addition of Sc does not increase the best formability among the four investigated alloys. For602—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 6. 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.METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—603
  • 7. 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 other three alloys, elongation decreases from alloy 1 Under the biaxial stretch condition, the anisotropy ofto alloy 2 and to alloy 3, which is consistent with the microstructure between the rolling and transverse directionsuniaxial tensile test results at 350 °C and 0.015 sÀ1. should have an effect on the deformation behavior.[17]604—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 8. Fig. 7—Elongation of various alloys: (a) alloy 1, (b) alloy 2, (c) alloy 3, and (d) alloy 4. The plastic strain ratio of width to thickness, R, was measured after superplastic testing. The R values for alloys 1 and 4 are 0.917 and 0.430, respectively, indicating that alloy 4 has more anisotropic ductility than alloy 1. This anisotropic ductility results in a slightly lower cup depth in alloy 4 than in alloy 1, though it has better uniaxial ductility. However, the biaxial stretch formability of alloy 4 (with both Zn and Sc 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 ofFig. 8—Superplastic tensile and biaxial stretch test results of investi- about 0.7 wt pct, most of the zinc is in solid solution ingated aluminum alloys. a-Al according to the Al-Zn-Mg phase diagram.[18]METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—605
  • 9. The Zn in solid solution has an antirecrystallization trace amounts of Sc are added in addition to Zn, as inability due to the solute drag effect on grain boundary alloy 4, the flow stress is higher and the ductility is lowermigration.[19] Furthermore, no Zn-rich second-phase (Figures 6(b), (d), and (f)) than the other three alloys atparticles were identified in alloy 3 (Al-Mg-Zn). Hence, almost all deformation conditions (the exception beingthe only effect of Zn is solute drag on grain boundary 350 °C and 0.015 sÀ1). This is because the change in themigration during thermal exposure. This effect prevents lattice parameter of a-Al due to Zn addition becomesthe grain growth and recrystallization process during more pronounced with the addition of trace amounts ofannealing, resulting in the finer and more elongated Sc, greatly increasing the solid solution strengthening.grains in alloy 3 (Figure 1(c)) compared to those in alloy Second-phase particles have a strong influence on1 (Figure 1(a)). formability of Al sheet alloys. The second-phase parti- The solid solubility of Sc in a-Al of the binary Al-Sc cles can nucleate voids or cracks and influence thesystem is considered to be only approximately growth and coalescence of these voids or cracks. The0.025 wt pct at 430 °C[9] and 0.07 wt pct at 500 °C.[13] most common modes of initiation are the complete orThe solid solution of Sc can also prevent the grain partial interfacial decohesion of the second-phase par-growth and recrystallization process during the thermo- ticle and the surrounding matrix, or the fracture ofmechanical treatments. Moreover, Mg and Sc mutually second-phase particles.[23] If the second-phase particlereduce each other’s solid solubility in the Al alloy; has a rounded shape or high fracture strength, crackshence, the equilibrium concentration of Sc is expected to usually initiate at the particle/matrix interface.[24] In thedecrease as temperature decreases or with increasing Mg investigated alloys, the Fe-rich intermetallic particlescontent.[9] Consequently, the addition of about have a rounded shape (Figure 3) and the Al3(Sc, Zr) or0.3 wt pct Sc in the Al-Mg–based alloys can result in Al3Zr intermetallics have a high fracture strength; hence,the formation of coarse primary intermetallic particles interface decohesion is likely to be the major mode ofduring casting. Such particles, with a size of up to crack nucleation. Furthermore, voids usually nucleate10 lm, are seen in Figures 1(d) and 3(b) and are likely to first on larger particles that contain small internal orbe Al3(Sc, Zr) enriched in Zn and Mg (Figure 4(c)). interfacial defects. Larger particles also induce more According to previous studies, these coarse particles rapid decohesion with increasing strain, because thehave a complex effect on the microstructure. First, constrained zone of plasticity at the interface becomesAl3(Sc, Zr) intermetallic particles formed in the melt can larger.[25] Thus, coarser second-phase particles are moreact as potent nucleation sites for aluminum grains,[9] deleterious to the mechanical properties than finer ones.refining the grain structure during casting. Second, The addition of Sc in the investigated alloy producesthermomechanical treatments result in large amounts much coarser and larger number of Al3(Sc, Zr) particlesof deformation-induced stored energy surrounding par- and no significant effect on the morphology andticles with a size above 1 lm. This increased stored distribution of the Fe-rich intermetallic particles. Con-energy greatly promotes the nucleation of recrystallized sequently, it is believed that the larger number of coarsegrains; hence, coarse intermetallics can act as efficient Al3(Sc, Zr) particles nucleate voids early and promotenuclei for recrystallization (also known as particle- the growth and coalescence of the voids, contributingstimulated nucleation (PSN)[20]). This PSN effect can greatly to the decrease in both flow stress and ductility inextend for more than ten particle diameters.[21,22] Con- the Sc-containing alloys. This is the likely cause for alloysequently, fine grains and subgrains are formed in zones 2 having inferior mechanical properties at the elevatedaround the particles. These zones are readily seen in the test temperatures. However, the flow stress at lower testEBSD mapping, as shown in Figure 2(b). Third, the temperatures or higher strain rates and the ductilityprimary intermetallic particles can also limit grain at 350 °C and 0.015 sÀ1 for alloy 4 increase. This isgrowth during thermomechanical treatments. Combined because the contributions from solid solution strength-with the solute drag effect due to Sc and Zn on grain ening and other mechanisms compensate the loss of theboundary migration, much more elongated and thinner flow stress and ductility due to large second-phasegrains are formed in the microstructure of alloy 4 than particles.that of alloy 1. Of the three elevated temperature deformation mech- anisms, GBS, slip creep, and diffusion creep, the contribution from diffusion creep is considered to beB. Effect of Zn and Sc Additions on Hot Formability negligible.[26] At relevantly low temperature, i.e., T < Tc, The Zn in a-Al increases the flow stress and decreases the grain-boundary bond is strong and slip creep is athe ductility of alloy 3 compared to alloy 1 due to solid predominant deformation mode. The slip creep processsolution strengthening; this effect is more pronounced at is found to be solute drag creep (SDC) in Al-Mg–basedthe elevated test temperatures. As mentioned in Section alloys.[1] However, at higher temperature, i.e., T > Tc,A, the solid solubility of Sc in a-Al of the binary Al-Sc the grain-boundary bonds are weakened, and yet shearsystem is only approximately 0.025 wt pct at 430 °C,[9] and normal stresses must be transferred across at the warm forming temperatures of and below Thus, inelastic displacement in grain-boundary regions350 °C, the amount of Sc in solid solution is expected to must be larger than those within the grain interior,be lower than this value for alloys 2 and 4. When Sc is which may even remain elastic. This view is essentiallysingularly added to alloy 2, the solid solution strength- the same as the grain mantle vs core deformationening is not significant, and the flow stress and elonga- model.[27] Thus, the elevated-temperature constitutivetion between alloys 1 and 2 is similar. However, when behavior has been described by assuming that both606—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 10. mantle and core deformation contribute to the overalldeformation, that is,[27] et ¼ em þ ec ½1Š where et is the total deformation rate; and and ec emare 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=d2 ðr À r0 Þ1þq ½2Šwhere 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 ðK þ A1 =d3 Þrn ½3Š ÀQ=RTwhere K is K1 e , and A1 is a constant relating tostress 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 mantle Fig. 9—Zener–Hollomon parameter as a function of modulus-region of the grains, and the SDC process occurs within compensated stress.the grain cores. The GBS controlling mechanism gener-ally requires a stable grain size of less than 10 lm, high surrounded by the finer grains is greatly increased.deformation temperatures (0.75 T/Tm), and low strain Moreover, according to Eq. [3], the fine grains alsorates (near 10À4 sÀ1).[1] In coarse-grained materials increase the contribution from grain cores. Conse-(grain size: 10 lm), the contribution from grain cores quently, the increased contribution from both grainis larger than that from grain mantles and, thus, SDC is mantles and grain cores increases the ductility of alloy 4the controlling deformation mechanism at various con- compared to the other three alloys. The increasedditions. The SDC may still be the controlling deforma- contribution from grain mantles is larger than thattion mechanism in fine-grained materials (grain size: from grain cores in alloy 4. This is demonstrated by the10 lm) at higher strain rates or lower temperatures. different shapes of the flow curves of alloy 4 compared The deformation temperature and strain rate can be to the other alloys at the same deformation condition ofrepresented by one parameter, the Zener–Hollomon 350 °C and 0.015 sÀ1. The flow curves for alloys 1parameter, Z=e expðQ=RTÞ, and hence deformation through 3 exhibit a sharp peak stress after reaching thebehavior can be investigated by a plot of Z values as a yield point, which is the characteristic of slip creep,[29]function of modulus-compensated stress, r/E. Plots of while strain hardening is apparent during plastic defor-the four investigated alloys are shown in Figure 9, with mation for alloy 4 at the same deformation condition.Q = 110 kJ/mol.[1] The Z values of alloys 1, 2, and 3 can The strain hardening is consistent with dynamic grainbe fit to a straight line with n = 8, while those of alloy 4 growth in GBS.[1] Moreover, the contribution from GBSare plotted at a higher stress level due to solid solution also enables this alloy to have better biaxial stretchstrengthening, and they show a significant decrease in the formability at 350 °C. When the deformation tempera-flow stress at lower Z values (350 °C, 0.015 sÀ1). ture is further increased to 520 °C, superplastic elonga- The lower m value (1/n 0.2) at the higher Z value tion of over 300 pct can be achieved. Thus, theindicates that greater contribution to the overall defor- combined addition of Sc and Zn in Al-Mg sheet alloysmation is from grain cores, and SDC is the controlling can generate an excellent combination of low-temperaturedeformation mechanism for the investigated alloys.[1] strength and high-temperature formability.However, at the lower Z value (350 °C, 0.015 sÀ1), thesefine grains in alloy 4 can easily slide due to the decreasein r0 of Eq. [2];[28] thus, they can increase the contribu- V. SUMMARYtion from grain mantle as well as from GBS duringdeformation. Also, GBS can occur in grain mantles of The additions of Sc or Zn in Al-Mg sheet alloys formlarger grains when the mobility of the larger grains solid solutions in a-Al and large Al3(Sc, Zr) intermetallicMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 40A, MARCH 2009—607
  • 11. particles. The solid solution of Sc or Zn in a-Al can 3. O.D. Sherby and J. Wadsworth: Prog. Mater. Sci., 1989, vol. 33,prevent the grain growth and recrystallization process pp. 169–221. 4. J. Pan, J. Tong, and M. Tian: Fundamental of Materials Science,and thereby refine the grain size in thermomechnical Qinghua University Publisher, Beijing, 2002, pp. 176–86.treatments prior to hot deformation. Also, large parti- 5. B.Z. Bai and K.A. Padmanabhan: Mater. Sci. Forum, 1997,cles lead to particle stimulated nucleation of recrystal- vols. 243–245, pp. 191–96.lization, resulting in the formation of numerous 6. L.P. Troeger and E.A. Starke, Jr.: Mater. Sci. Eng. A, 2000, vol. A277, pp. 102–13.subgrains and fine grains around the particles. These 7. E.M. Taleff, P.J. Nevland, and P.E. Krajewski: Metall. Mater.large particles can limit grain growth during thermo- Trans. A, 2001, vol. 32A, pp. 1119–30.mechnical treatments, resulting in a smaller grain size. 8. J. Royset and N. Ryum: Scripta Mater., 2005, vol. 52, pp. 1275– During tensile testing at temperatures ranging from 79.275 °C to 350 °C and strain rates of 0.015 to 1.5 sÀ1, 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,solid solution strengthening of Zn or Sc increases the vol. 49, pp. 1327–37.flow stress and decreases ductility. In the alloys con- 11. B. Forbord, H. Hallem, J. Røyset, and K. Marthinsen: Mater. Sci.taining Sc, large intermetallic particles cause crack Eng. A, 2008, vol. 475, pp. 241–48.initiation during warm deformation, further decreasing 12. M.M. Sharma, M.F. Amateau, and T.J. Eden: J. Alloys Compd., 2006, vol. 416, pp. 135–42.the flow stress and ductility of these alloys. Further- 13. K.L. Kendig and D.B. Miracle: Acta Mater., 2002, vol. 50,more, during warm deformation, SDC in grain cores pp. 4165–75.and GBS in grain mantles occur simultaneously. How- 14. R.A. Karnesky, L. Meng, and D.C. Dunand: Acta Mater., 2007,ever, SDC is the controlling deformation mechanism, vol. 55, pp. 1299–308.demonstrated by the strain rate sensitivity coefficient of 15. M.A.G. Bernal, D.H. Silva, and V.S. Rangel: J. Mater. Sci., 2007, vol. 42, pp. 3958–63.below 0.2, in alloys 1 through 3 and alloy 4 at lower test 16. J. Liu and J.G. Morris: Mater. Sci. Eng. A, 2004, vol. 385,temperatures and higher strain rates. With increasing pp. 342–51.deformation temperature and decreasing strain rate, the 17. G.E. Dieter: Mechanical Metallurgy, McGraw-Hill Bookcontribution from GBS increases. At 350 °C and Company, New York, NY, 1986. 18. H. Loffler, I. Kov, and J. Lendvai: J. Mater. Sci., 1983, vol. 18,0.015 sÀ1, the numerous fine grains in alloy 4 (contain- pp. 2215– both Sc and Zn additions) further promote GBS, 19. T. Kobayashi: Mater. Sci. Eng. A, 2000, vol. 286, pp. 333–41.increasing the ductility and resulting in an elongation of 20. T.G. Nieh, L.M. Hsiung, J. Wadsworth, and R. Kaibyshev: Acta130 pct. The benefit of the combined additions of Sc and Mater., 1998, vol. 46, pp. 2789–2800.Zn has also been verified by the results of superplastic 21. J.S. Vetrano, S.M. Bruemmer, L.M. Pawlowski, and I.M. Robertson: Mater. Sci. Eng. A, 1997, vol. 238, pp. 101–07.tensile tests and biaxial stretch tests. On the basis of the 22. R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juulexperimental results, the combined addition of Sc and Jensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen,Zn in Al-Mg may be an important method of producing and A.D. Rollett: Mater. Sci. Eng. A, 1997, vol. 238, pp. 219–74.a high-strength sheet alloy with excellent high-temperature 23. D. Lassance, D. Fabregue, F. Delannay, and T. Pardoen: Prog. Mater. Sci., 2007, vol. 52, pp. 62–129.formability. 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, REFERENCES Metalworking: Bulk Forming, ASM INTERNATIONAL, Mate- rials Park, OH, 2005, pp. 563–86.1. T.R. McNelley, K.O. Ishi, A.P. Zhilyaev, S. Swaminathan, P.E. 28. I. Charit and R.S. Mishra: Acta Mater., 2005, vol. 53, pp. 4211–23. Krajewski, and E.M. Taleff: Metall. Mater. Trans. A, 2008, 29. D.H. Bae and A.K. Ghosh: Acta Mater., 2000, vol. 48, vol. 39A, pp. 50–64. pp. 1207–24.2. T.G. London: Z. Metall., 2005, vol. 96, pp. 522–31.608—VOLUME 40A, MARCH 2009 METALLURGICAL AND MATERIALS TRANSACTIONS A