softening behaviour of cold rolled continuous cast and ingot cast aluminum alloy aa5754

Materials Science and Engineering A 421 (2006) 276–285

Softening behaviour of cold rolled continuous cast and ingotcast aluminum alloy AA5754

S. Sarkar ∗, M.A. Wells, W.J. PooleDepartment of Materials Engineering, The University of British Columbia, 309-6350 Stores Road, Vancouver, BC V6T 1Z4, Canada

Received 20 September 2005; accepted 23 January 2006

Abstract

The recovery and recrystallization behaviour of ingot cast (IC) and continuous cast (CC) AA5754 cold rolled material has been compared.Included in the comparison was the effect of a homogenization treatment as well as hot rolling prior to cold rolling for the CC material. Withoutthe homogenization process, the softening kinetics between the CC and IC material were significantly different with the CC material exhibitingconsiderably slower softening kinetics. However, when the CC material was homogenized prior to cold rolling, the softening kinetics were similarto that of ingot cast (IC) material. The differences in the softening behaviour between the CC material without a homogenization treatment and ICmaterial can be related to solute drag and concurrent precipitation which occurred during annealing of the CC material resulting in the retardedgti©

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1

miaia(msAtnlth[e

0d

rowth of recrystallized grains. This was supported by measuring the evolution in the resistivity and microstructure during the annealing heatreatment. Differences in the recrystallized texture between the CC and IC material were also made during this investigation and the resultsndicated that although overall the textures were weak, the IC material consistently had a higher cube texture volume fraction.

2006 Elsevier B.V. All rights reserved.

eywords: Recovery and recrystallization; Al–Mg alloy; Homogenization; Continuous and ingot cast; Resistivity; Precipitation

. Introduction

Increasingly light weight materials such as aluminum andagnesium are receiving interest as alternate materials to steel

n automotive sheet applications. A current disadvantage in usingluminum sheet for automotive applications is that it costs signif-cantly more than steel sheet on a per pound basis. Traditionally,luminum sheet alloys have been produced via the ingot castingIC) route. In this manufacturing process, large rectangular alu-inum ingots (400–500 mm thick) are produced which are then

calped, homogenized, hot rolled, cold rolling and annealed.luminum sheet produced in this manner typically costs 4–5

imes more than steel sheets on a per pound basis [1]. Alter-atively, the continuous casting (CC) route offers a potentiallyess expensive processing route, as the aluminum is cast intohin slabs of 25 mm or less in thickness which can then be eitherot and/or cold rolled to the final desired thickness and annealed2]. In this manufacturing process, the costly stages of homog-nization and potentially hot rolling are eliminated [3].

There are some obvious differences between the continuouscast and the ingot cast processing routes for the production ofsheet aluminum. These include: (i) the solidification rate expe-rienced during casting, which influences the grain size as wellas the size and distribution of the constituent particles [4–7], (ii)the absence of the homogenization/preheat treatment stage formaterial produced via continuous casting [2,4,8], and (iii) theamount of hot and cold deformation applied to the material aftercasting [8]. These differences can affect the way the aluminumsheet responds to the final annealing treatment, i.e. the final sheetmicrostructure and mechanical properties.

Previous work on continuous and ingot cast Al–Mg(AA5xxx) alloys has indicated that the softening kinetics aregenerally faster for ingot cast material as compared to continu-ous cast material for the same level of cold deformation [8–10].The difference in the softening behaviour has been attributed to:(i) the differences in the size and distribution of the constituentparticles (formed during casting), the level of solute dissolvedin the matrix and (iii) the formation of dispersoids in the matrixeither prior to or during annealing. Typically continuous castmaterial has a much higher level of solute (particularly Mn) in

∗ Corresponding author. Tel.: +1 604 822 7886; fax: +1 604 822 3619.E-mail address: [email protected] (S. Sarkar).

solution viz. ingot cast material since continuous cast materialdoes not typically undergo a homogenization treatment where

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dx.doi.org/10.1016/j.msea.2006.01.069

S. Sarkar et al. / Materials Science and Engineering A 421 (2006) 276–285 277

Table 1Chemical composition of AA5754 material used in this study (in wt.%)

AA5754 Mg Mn Fe Si Al

Continuous cast 3.11 0.20 0.16 0.090 ∼BalanceIngot cast 3.07 0.24 0.17 0.057 ∼Balance

much of this solute can precipitate out in the form of disper-soids [8,11]. The presence of high levels of solute can cause astrong solute-dislocation interaction during the annealing heattreatment and thereby retard the kinetics of both recovery andrecrystallization [10]. Precipitation of fine dispersoids duringthe annealing treatment may also affect the softening behaviourof the cold rolled aluminum alloys, as has been observed inAA5xxx and AA3xxx alloys [12,13]. Research has also shownthat when the hot rolled CC material is given a high-temperaturepreheat treatment (∼474 ◦C for 3 h) prior to cold rolling, the soft-ening kinetics are enhanced [14]. Studies have indicated thatthere is an increase in the measured conductivity during homog-enization treatment (∼560 ◦C) on various AA5xxx CC materialswhich offers indirect evidence for the precipitation of Mn andMg rich particles [15].

Although previous research has shown that there is a differ-ence in softening kinetics for aluminum sheet processed via ICand CC routes and that this difference is related to the homoge-nization heat treatment there is still a lack of understanding onthe reasons why aluminum sheet produced via the CC and ICproduction routes behaves so differently. This lack of knowledgeincludes: (1) a more fundamental insight into the evolution ofthe microstructure during the recrystallization process and theeffect the CC and IC route has on this and (2) a more preciseunderstanding of the change of solute in solid solution duringboth the homogenization as well as annealing heat treatmentsptefob(octd

2

2

NTtotb(

Table 2Thermal-mechanical processing route for continuous and ingot cast AA5754used in this study

Values given in the brackets are the % reduction. Names highlighted in bold arethe starting material received from Novelis and subsequently heat treated anddeformed at UBC.

facilities at Novelis. To avoid the complication of the retainedstrain on subsequent processing in the CC material which hadbeen hot rolled, this material was annealed for 30 min at 500 ◦Cin a salt bath which resulted in a fully recrystallized microstruc-ture. To investigate the effect of homogenization process on thesoftening kinetics, some of the as-cast CC AA5754 was given ahomogenization treatment at 500 ◦C of 24 h in a salt bath priorto cold rolling at UBC.

In total, four different processing histories for AA5754 werecompared, i.e. (i) ingot cast material processed through the tra-ditional route of casting, homogenizing, hot rolling and coldrolling (IC-HR), (ii) continuous cast material hot rolled andthen cold rolled (CC-HR), (iii) continuous cast material directlycold rolled (CC-DCR), and (iv) continuous cast material given ahomogenization treatment prior to cold rolling (CC-H). Detailsof the thermal-mechanical processing history for both the CCand IC materials are given in Table 2. Micrographs of the start-ing materials prior to cold deformation are shown in Fig. 1.

2.2. Annealing experiments

To study the recrystallization and recovery behaviour ofthese materials, a series of isothermal annealing experimentswere conducted using a salt bath in the temperature range of200–325 ◦C. The composition of the salt bath was 60% potas-sium nitrite + 40% sodium nitrate and the typical measured heat-ubb

2

tna

rior to and after cold rolling. This research attempts to addresshese issues by doing a systematic investigation of the differ-nces in the material response to an annealing heat treatmentor aluminum sheet produced using ingot casting and continu-us casting technologies. This includes an investigation of: (i)oth recovery and recrystallization kinetics during annealing,ii) using resistivity measurements to de-convolute the influencef the change in dislocation density from the change in soluteontent, and (iii) a detailed microstructure characterization ofhe evolution in the number of recrystallized grains over timeuring annealing of IC and CC material.

. Experimental

.1. Initial material

The materials used in this investigation were provided byovelis Global Technology Centre (Kingston, Ontario, Canada).he chemical composition for both materials is very similar and

he details are given in Table 1. The starting materials consistedf: (i) ingot cast (IC) hot band produced in industry (fully recrys-allized), (ii) continuous cast (CC) AA5754 hot band producedy pilot scale facilities at Novelis (partially recrystallized), andiii) continuously as-cast AA5754 produced using pilot scale

p rate for the samples was 25 ◦C/s. The holding times in the saltath were varied from 30 s to 15 h to investigate the softeningehaviour of the material.

.3. Material characterization

The softening behaviour of the annealed materials was quan-ified by measuring the change in the mechanical properties,amely the yield strength, as a function of annealing timend temperature. Limited optical metallography was also done

278 S. Sarkar et al. / Materials Science and Engineering A 421 (2006) 276–285

Fig. 1. Micrographs of the AA5754 prior to cold rolling with thermal-mechanical histories as outlined in Table 2: (a) CC (recrystallized hot band), (b) CC (as-cast),(c) CC (as-cast + homogenization), and (d) IC (hot band). All the pictures were taken through the thickness of the material at the 1/4 thickness position.

to determine the initial (as deformed) and final (fully recrys-tallized) microstructure of the materials. Some measurementson partial recrystallized structures were also done to assessthe validity of indirectly estimating the fraction recrystallizedthrough the mechanical property measurements. The grain sizewas measured according to the ASTM E112-88 standard (Jef-fries’ method). For this method, grain sizes were calculatedfrom micrographs taken at approximately 350× magnificationsuch that each micrograph contained at least 250 grains. Frac-tion recrystallized at various stages during annealing was esti-mated according to the ASTM E562-89 standard point countingmethod. To some extent, all the materials exhibited a variation ingrain size and microstructure from the surface to the center of thesheet. All measurements reported in this work were measuredat the 1/4 thickness position.

Resistivity measurements were done to study the microstruc-ture changes during the annealing and homogenizing heat treat-ments on continuous cast materials in as-cast, deformed, par-tially softened and fully softened conditions. A Verimet M4900Cconductivity meter was used to measure the conductivity of thematerials, which was then converted to corresponding resistiv-ity values. Each sample thickness was kept at least 1.5 mm toavoid any kind of thickness effect on the measurement. Resis-tivity measurements for the CC-HR and IC-HR were not doneas the final thickness of these materials was too thin.

In addition, evolution of the microstructure including the sizeafihAl

quantitative texture measurements were done on each of the fullyrecrystallized samples using the electron backscattered diffrac-tion (EBSD) system.

3. Results

3.1. Softening kinetics

The starting point for the annealing experiments was thecold rolled materials. The yield stress after 80% cold rollingshowed little variation, ranging from 335 MPa for the CCrecrystallized hot band (CC-HR) to 352 MPa for the CC-DCRmaterial (see Table 3). Upon annealing, there was a significantreduction in the yield stress, first due to static recovery and forlonger times at higher temperatures due to static recovery andrecrystallization. Fig. 2 summarizes the comparative softeningbehaviour of the IC and CC AA5754 materials after annealingat 200 ◦C for various lengths of time. Referring to Fig. 2, it canbe seen that the yield stress of the material gradually decreases

Table 3Flow stress for 80% cold rolled continuous and ingot cast AA5754 at threedifferent conditions

Material conditionprior to coldrolling

As-rolled yieldstress (MPa)

Recrystallizedstart stress (MPa),annealed at 300 ◦C

Fullyrecrystallizedstress (MPa)

CCCI

E

nd number of recrystallized grains per unit area were quanti-ed at various stages of annealing to understand the effect ofomogenization process on the CC-DCR and CC-H materials.t least 500 grains were studied for the calculation of recrystal-

ized grain density on partially recrystallized materials. Finally

C-HR 335 227 117C-DCR 352 236 118C-H 342 221 116

C-HR 343 219 112

rror associated with the stress measurement was ±5 MPa.


Fig. 2. Isothermal recovery kinetics of AA5754 at 200 ◦C after 80% cold workwith different processing histories.

over time, typical of what is observed during recovery [16,17].The IC material exhibits the fastest recovery kinetics followedby the CC material that had received a homogenization heattreatment (CC-H), the CC hot band (CC-HR) and finally theCC material that was directly cold rolled (CC-DCR).

Fig. 3 shows the softening data for CC-HR material at threedifferent annealing temperatures (i.e. 275, 300 and 325 ◦C).Optical metallography confirmed that the final plateau inthe yield stress values corresponded to a fully recrystallizedmicrostructure. Detailed optical metallography was used to char-acterize the onset of recrystallization and this point is marked inFig. 3 by the vertical arrows. As can be observed in Fig. 3, theannealing temperature did not affect the yield stress at whichthe recrystallization process started or the fully recrystallizedyield stress (at least for the conditions studied in this work).The recrystallization start and finish yield stresses are given inTable 3 for each of the materials investigated.

Assuming the results that the yield stress for the initiationof recrystallization and final recrystallized yield stress wereessentially independent of processing history and annealing tem-

FAi

Fig. 4. Comparison of measured recrystallization kinetics for 80% cold rolledCC-HR at 300 ◦C based on mechanical property measurements and quantitativemetallography.

perature, it is possible to estimate the fraction recrystallizedbased on the softening data.

FR = σRxi − σt

σRxi − σRx(1)

where FR is the fraction recrystallized, σRxi the yield stress atthe onset of recrystallization, σRx the yield stress in the fullyrecrystallized state and σt is the material yield stress as a functionof time (t) during the recrystallization process, i.e. ranging fromσRxi to σRx. This estimate for the fraction recrystallized was veri-fied against quantitative metallography for the case of annealingat 300 ◦C. As shown in Fig. 4, an estimation of the fractionrecrystallization using optical metallographic techniques givesgood agreement with recrystallization kinetics determined usingmechanical test data and thus, for this alloy system this approachcan be used to characterize recrystallization behaviour.

Fig. 5 compares the recrystallization behaviour for boththe continuous and ingot cast materials at 300 ◦C. Similarto the recovery kinetics the IC material exhibited the fastestrecrystallization kinetics, followed by the CC material with thehomogenization treatment, CC material which had been hotrolled and CC material which was directly cold rolled. The timeneeded to obtain 50% recrystallization is approximately 325 sfor the IC material as compared to approximately 480 s for theCC material with the homogenization treatment, 800 s for theCC material hot band and 1875 s for the CC material whichw

ccmCeiTasg

ig. 3. Isothermal softening kinetics for 80% cold rolled CC-HR material.rrows indicate recrystallization start temperature for the three different anneal-

ng temperatures.

as directly cold rolled.Fig. 6 shows the fully recrystallized grain structure of both

ontinuous and ingot cast materials after recrystallization isomplete. It can be observed from Fig. 6 that all the recrystallizedaterials had a mostly equiaxed grain structure. However, theC materials which did not undergo a homogenization treatmentxhibited some elongated grains parallel to the rolling directionn the fully recrystallized structure as pointed out in the figure.he average recrystallized grain size values for all four materialsre presented in Table 4. As can be observed from Table 4 andupported by Fig. 6, the continuous cast materials exhibit a finerrain size as compared to the ingot cast material.


Fig. 5. Isothermal recrystallization kinetics of AA5754 at 300 ◦C after 80% coldwork with different processing histories.

Table 4Average grain size of recrystallized 80% cold rolled continuous and ingot castAA5754

Material conditionprior to cold rolling

Initial grainsize (�m)

Recrystallized grainsize (�m)

CC-HR 29.0 12.0CC-DCR 60.0 9.5CC-H 65.0 10.6IC-HR 32.0 14.0

Error associated with the grain size measurement was ±1.0 �m.

Fig. 7. Change in resistivity of CC AA5754 in the as-cast condition duringhomogenization at 500 ◦C for 30 h.

3.2. Microstructure changes during homogenization

To examine why the homogenization treatment had such adramatic effect on the softening kinetics, electrical resistivitymeasurements were conducted on the CC material at varioustimes during the homogenization process and the annealingheat treatment after cold rolling. It was assumed that changesin electrical resistivity would be indirect evidence of changesin the solid solution level of the alloy. Fig. 7, illustrates thechange in the electrical resistivity of the material over timeduring a homogenization treatment at 500 ◦C. As shown inFig. 7, there is a large change in the resistivity of the CC as-castmaterial during the homogenization treatment, which would

Fig. 6. Micrographs of recrystallized AA5754 at 300 ◦C after 80% cold rolling with(c) CC-H, and (d) IC-HR. All the pictures were taken through the thickness of the str

thermal-mechanical histories as outlined in Table 2: (a) CC-HR, (b) CC-DCR,ip at the 1/4 thickness location, arrows indicate elongated grains in (a) and (b).


Fig. 8. Increase in resistivity values during cold rolling to various strains.

be consistent with a reduction in solute-supersaturation in theas-cast CC material due to the precipitation of dispersoids suchas Mn/Fe rich particles.

Further, to investigate the possible concurrent precipitationof dispersoids during annealing of the cold deformed materials,additional resistivity measurements were conducted. For thispurpose, the as-cast CC materials (with or without homoge-nization) were 80% cold rolled (true strain of approximately1.85) and resistivity measurements were conducted as a func-tion of annealing time at 300 ◦C. It was observed that there wasa significant increase in electrical resistivity of the alloy as afunction of the cold deformation (see Fig. 8). For example, theresistivity of the direct cold rolled as-cast material increasedfrom 52.3 to 53.8 n� m for an imposed strain of approximately1.85 while the cold rolled homogenized material increased from49.5 to 51.2 n� m (note that the change in resistivity with colddeformation was almost the same for the two cases). Fig. 9summarizes the subsequent evolution of the electrical resistivityduring annealing at 300 ◦C for three cases, i.e. (i) as-cast materialwithout cold rolling, (ii) as-cast material rolled to 80% reduc-

Fa(

Table 5Resistivity values for CC-DCR and CC-H materials before and after cold defor-mation and after complete recrystallization at 300 ◦C

Material Resistivity values (n� m)

Before rolling After rolling After completerecrystallization

CC-DCR 52.3 53.8 50.1CC-H 49.5 51.2 49.3

Error associated with measurement was ±0.2 n� m.

tion in thickness, and (iii) as-cast material homogenized for 24 hat 500 ◦C and then cold rolled 80%. For the case of the as-castmaterial without deformation, there is essentially no change inelectrical resistivity indicating that there is no significant changein the microstructure. In contrast, the cold rolled as-cast materialwithout homogenization shows a rapid drop in electrical resis-tivity after the first 60 s (from 53.8 to 51.4 n� m) of annealingand then a much slower drop in resistivity with annealing time(i.e. from 51.4 to 50.3 n� m). Finally, the homogenized, coldrolled material also shows a rapid initial drop in resistivity (i.e.from 51.1 to 49.5 n� m) but in this case, the subsequent drop inresistivity is less (i.e. from 49.5 to 49.3 n� m). The importantvalues of resistivity are summarized in Table 5.

Finally the evolution of the recrystallized grain structure forthe CC-DCR and the CC-H materials at various stages of anneal-ing at 300 ◦C were studied by optical metallography. The evo-lution of the recrystallized grain structure was characterized bymeasuring the evolution for the number of recrystallized grainsas a function of annealing time. Fig. 10 shows the change in graindensity (the number of grains per unit area) for the direct coldrolled CC material and the homogenized/cold rolled material asa function of annealing time at 300 ◦C. It can be observed that forthe case of the homogenized cold rolled material, the area densityof the recrystallized grains rapidly increases and then reaches asteady state value at the early stages during the recrystallizationprocess (at approximately 20% recrystallized). In contrast, ther

Faw

ig. 9. Change in resistivity during annealing at 300 ◦C for CC materials in thes-cast (without homogenization treatment) and 80% cold deformed conditionwith and without experiencing a homogenization treatment).

ecrystallized grain density in the direct cold rolled CC mate-

ig. 10. Measured evolution in recrystallized grains over time during annealingt 300 ◦C in both homogenized and non-homogenized AA5754 CC materialshich had been cold rolled to 80%.


Table 6Measured volume fraction (%) of individual texture components for both theingot cast and continuous cast AA5754 after 80% cold deformation and recrys-tallization at 325 ◦C

Texture component CC-DCR CC-H CC-HR IC-HR

Cube 3.00 4.95 5.47 6.38Brass 5.58 5.22 6.00 5.10Copper 5.55 7.03 6.58 6.42Goss 2.17 2.95 2.18 1.73S 12.39 15.42 15.13 13.11R 13.64 15.31 15.69 14.65P 2.91 3.68 4.00 4.40Q 9.94 9.08 8.90 9.76H 2.69 1.49 1.93 1.88CH 6.65 6.30 6.59 5.95R. cube 6.69 7.94 7.39 8.36

rial continues to increase during the recrystallization processand does not reach a steady-state value until recrystallizationis approximately 80% complete. Fig. 10 also shows the finalrecrystallized grains density is higher for the CC-DCR materi-als as compared to the CC-H material, which is consistent withthe finer measured grain size of the CC-DCR material.

3.3. Texture

The recrystallized texture measurement for both the continu-ous and ingot cast materials is presented in Table 6. The resultspresented in Table 6 show that although the majority of thetexture is random there is a trend in terms of the cube texturecomponent being higher in the IC materials as well as the CCmaterial which had been homogenized prior to cold rolling. Thisresult is quite similar to the finding published elsewhere for CCmaterial with or without homogenization [18].

4. Analysis and discussion

The current work is consistent with previous work in showingthat that there are significant differences in the rate of recrys-tallization for the different processing routes. The time to reach50% recrystallization can be as much as 5–6 times longer fordirect cold rolled CC material compared to material processedttstwac

acatine

First, during homogenization of the as-cast CC material,changes in electrical resistivity will be dominated by soluteeffects. The effect of Mn and Fe in solution has a particu-larly strong effect on resistivity, i.e. 29.4 and 25.6 n� m/wt.%respectively [22] and the importance of the precipitation of (Mn,Fe) dispersoids has been previously reported [12,13,15]. Asshown in Fig. 7, the change in resistivity during homogeniza-tion is approximately 3 n� m which would roughly correspondto approximately a combined 0.1 wt.% change of the Mn and Fecontent in solution or about one quarter of the total Mn/Fe con-tent of the alloy which seems reasonable. It can also be observedthat after 24 h of annealing at 500 ◦C, there is no further changein resistivity. It will be assumed that this indicates the system isnear equilibrium at this point.

Using this assumption, it is possible to examine the roleof the resistivity contribution from dislocations by consideringthe deformation of the homogenized material. Fig. 8 shows thedependence of electrical resistivity on the level of cold defor-mation. It can clearly be observed that the rate of increase ofelectrical resistivity with cold deformation decreases at largerstrains. This result arises due to the fact that: (i) rate of storageof dislocations with strain decreases at larger strain and (ii) itappears that the scattering effectiveness of dislocations is smallerat high dislocation densities is smaller (as will be discussed fur-ther in the following sections).

To estimate the contribution to resistivity from dislocationdttlma

σ

wfpwtcmtd

ρ

wiao

edcmg

hrough a traditional ingot processing route. In the following,he focus will be on the role of solute (primarily Mn) in solidolution and its precipitation either during the homogenizationreatment or concurrent with recovery and recrystallization. Weill attempt to show that it is possible to extract information,

t least in a semi-quantitative manner, on this question from aareful analysis of electrical resistivity measurements.

Electrical resistivity measurements have often been used asn indirect measure of solute in solution [19]. However, therean also be a significant contribution to resistivity in aluminumlloys from dislocation scattering [20,21]. It is, thus, necessaryo carefully deconvolute these two effects in order to providensight into the changes in solute level during annealing. Fortu-ately, the current results offer a number of limiting cases whereither only the solute effect or the dislocation effect dominates.

ensity, it is necessary to estimate the dislocation density ofhe material as a function of the level of cold deformation. Ashe direct measurement of dislocation density is extremely chal-enging, we will estimate its value from the yield stress of the

aterial. The yield stress of a solid solution hardening alloy suchs AA5754 is given by

ys = σ0 + σss + σHP + αGbMρ1/2 (2)

here σ0 is the intrinsic lattice resistance to dislocation motionor aluminum, σss the solid solution contribution (in this case,redominately from Mg), σHP is the grain size contributionhich can be described by the Hall–Petch equation and the final

erm is the forest hardening contribution where α is geometriconstant (0.3), G the shear modulus of aluminum (26 GPa), b theagnitude of the Burgers vector (0.286 nm), M the Taylor fac-

or (3.06) and ρ is the dislocation density. Thus, the dislocationensity can be estimated from the change in flow stress as

=(

σCW − σRx

αGbM

)2

(3)

here σcw is the yield stress of the cold worked material. Note, its not necessary to evaluate the intrinsic strength, solid solutionnd grain size contributions as these are common to the strengthf the cold worked and fully recrystallized material.

It is now possible to plot the evolution of the increase inlectrical resistivity as a function of the estimated dislocationensity for the homogenized material as shown in Fig. 11. In thisase, the data for both deformed as-cast CC and homogenizedaterials are shown. The two have a similar dependence sug-

esting that the resistivity increase in both cases comes from the


Fig. 11. Increase in resistivity values due to cold deformation as a function ofincrease in dislocation density. Eq. (4) shows a good match with the experimentaldata.

dislocation scattering contribution (i.e. even though the as-castmaterial is supersaturated in solute, this does not change duringcold deformation). The dependence of resistivity on dislocationdensity in aluminum has been reported to be linear [21,23] andthe rate of the initial increase is consistent with this in the cur-rent work. However, at large dislocation densities, there appearsto be a drop off in the rate of increase of resistivity with dis-location density. There could be a number of reasons for thisdeviation, particularly the estimate of dislocation density fromflow stress measurement may not be sufficiently refined or theremay indeed be a deviation from linearity at large dislocationdensities when the dislocation spacings become very small. Nev-ertheless, a strong correlation is observed and for the purpose ofanalyzing the data, the dependence of resistivity increase withdislocation density can be captured by

�ρresistivity = A[1 − exp(−B�ρdislocation)] (4)

where A and B are two empirical constants which have values of10−15/m2 and 2.2 n� m, respectively. The prescribed functionsuits the current situation, i.e. it is initially linear with a slopeequal to the product of AB and for larger dislocation densities, therate of increase of resistivity with dislocation density decreases.It is of interest to compare the dependence of resistivity on dislo-cation density reported in the literature for pure aluminum andone finds that the present results show a dependency of 3–10tcMt

ttc3tes

Fig. 12. Reduction in dislocation density during annealing at 300 ◦C for CC-Hmaterial as calculated from resistivity and softening data.

achieved which offers confidence that the proposed method isself-consistent.

Finally, it is possible to examine the case of most interest,i.e. situations where concurrent softening and precipitation areoccurring. To deconvolute the two effects the following proce-dure was adopted:

1. the change in dislocation density was estimated based on theevolution of the yield stress of the alloy during annealingusing Eq. (3);

2. using these estimates for the dislocation density, the contribu-tion to the change in the electrical resistivity was calculatedusing Eq. (4);

3. the contribution from changes in solute level was then cal-culated by subtracting the dislocation contribution from thetotal resistivity change experimentally measured.

Figs. 13 and 14 show examples of the results when this proce-dure is applied to the experimental data for annealing at 200 and

Fm

imes the values for pure Al as reviewed by Nabarro [21]. Onean speculate that this deviation comes from the segregation ofg atoms to dislocations and thereby affected the efficiency of

he scattering.Having established a relationship between electrical resis-

ivity and the estimated dislocation density, it is useful to fur-her examine its applicability. Fig. 12 shows estimates for thehange in dislocation density as a function of annealing time at00 ◦C for the homogenized CC material (note: the solute con-ent of the matrix is presumed to remain constant during thisxperiment) based on yield stress and electrical resistivity mea-urements. Excellent agreement between the two estimates is

ig. 13. Contribution from dislocation and solute on resistivity for CC-DCRaterial during annealing at 200 ◦C for approximately 3 h.


Fig. 14. Contribution from dislocation and solute on resistivity for CC-DCRmaterial during annealing at 300 ◦C for approximately 3 h.

300 ◦C, respectively. First, if we consider annealing at 200 ◦Cwhere only static recovery is observed, one can see in Fig. 13that there is considerable overlap between the depletion of solute(i.e. the precipitation reaction) and the static recovery process.The precipitation of dispersoids in the deformed structure andthe resulting pinning of the structure would explain the lowerrates of static recovery that was observed for this situation (seeFig. 2).

On the other hand, at higher temperatures such as 300 ◦Cwhere both recovery and recrystallization are observed (resultsshown in Fig. 14), one observes that the resistivity contribu-tion from the solute drops very sharply in the first 60 s ofannealing and then remains constant. Referring back to Fig. 3,it can be seen that the onset of recrystallization occurs atapproximately 120 s. It is, therefore, possible to conclude thatthe precipitation reaction at 300 ◦C again occurs concurrentwith static recovery but prior to the onset of recrystallization.Thus, the slower rate of recrystallization observed for this sit-uation is likely related to precipitation of dispersoids in thedeformed structure and on grain boundaries for which the com-bined effect is to delay recrystallization. This interaction withthe deformed structure for this case would also be consis-tent with the observation that appearance of new recrystallizedgrains occurred to a much later stage of recrystallization (seeFig. 10 and Section 3). This could be understood in terms ofthe requirement for a rearrangement of the deformed structureics

rsaaecpt

5. Conclusions

Based on the research conducted, a number of conclusionscan be made about the effect of processing route on the annealingbehaviour of AA5754 after cold rolling. The results from thisstudy indicate that although there is little difference in the macro-scopic mechanical properties and work hardening behaviour ofthe material after cold rolling, the variation in the initial solutelevel in starting materials can exert a significant influence onthe softening kinetics (recovery and recrystallization), and finalrecrystallized microstructure during further processing. Furtherspecific conclusions include:

1. Softening kinetics of the CC material without homogeniza-tion are significantly slower than IC material. When the CCmaterial is given a homogenization heat treatment similar tothat experienced by the IC material, the softening kinetics areincreased and approach those measured in the IC material.

2. Using a new approach to deconvolute the solute and disloca-tion effects upon the evolution of electrical resistivity duringannealing, it was possible to show that the precipitation ofdispersoids occurs concurrently with static recovery beforethe onset of recrystallization.

3. Recrystallized texture measurements shows that both CC andIC materials exhibit random texture and the homogeniza-

A

TtfSatlt

R

n order to nucleate new grains due to the inhibition of this pro-ess because of the precipitates that formed in the deformedtructure.

Based on the results presented, it is evident that the processingoute, specifically the application of a homogenization step has aignificant influence on the softening kinetics during annealings well as the final recrystallized microstructure of the AA5754fter cold working. It appears that the majority of the differ-nce in the rate of softening kinetics between the ingot cast andontinuous cast material is related to the role of solutes and dis-ersoids (i.e. whether they are in or out of solution) as well asheir distribution within the matrix.

tion treatment seems to play a minor role in determiningthe final texture of the materials. However application ofa homogenization treatment increases the presence of cubetexture.

cknowledgements

The authors would like to acknowledge the Novelis Globalechnology Centre for providing the material for this investiga-

ion. We would also like to acknowledge Johnson Go of UBCor providing softening data for ingot cast AA5754 material.pecial thanks go to Dr. David Lloyd, Dr. Mark Gallerneaultnd Mr. Ed Luce of the Novelis Global Technology Center forechnical assistance in material preparation. Finally we wouldike to acknowledge NSERC for providing financial support forhis research.

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softening behaviour of cold rolled continuous cast and ingot cast aluminum alloy aa5754

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