Plastic Deformation of Al and AA5754 Between 4.2K and 295K

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Materials Science and Engineering A 491 (2008) 88–102

Plastic deformation of Al and AA5754 between 4.2 K and 295 K

Dong-Yeob Park, Marek Niewczas ∗Materials Science and Engineering, McMaster University, JHE-357, 1280 Main St. West, Hamilton, ON L8S 4L7, Canada

Received 2 April 2007; accepted 22 January 2008

bstract

Plastic deformation and work-hardening behaviour of high-purity Al and AA5754 have been studied by a combination of measurements ofechanical response, electrical resistivity and TEM. The results show that materials deformed at very low temperatures exhibit an unprecedented

evel of strength and unusual work-hardening behaviour. Strain rate sensitivity measurements suggest that deformation of high-purity Al is governedy dislocation–dislocation interactions in a broad range of temperatures, whereas Al alloys exhibit a larger thermal component of flow stress dueo the presence of solute atoms in the matrix. The electrical resistivity data provide information about the evolution of both the dislocation density
nd the dislocation mean free path and establish the contributions of dislocation storage and dynamic recovery during plastic flow of materials at.2 K. The results suggest that fracture is initiated by the collapse of the dislocation network at places where dislocations develop a critical spacingor spontaneous annihilation. This spacing is estimated as ∼8 nm at 4.2 K and ∼12 nm at 78 K.
2008 Elsevier B.V. All rights reserved.

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eywords: Aluminum; AA5754 Al alloys; Plastic deformation; Electrical resis

. Introduction

Aluminum alloys of the 5000 (Al–Mg) series have been con-idered for structural applications in the automotive industry1]. The non-heat-treatable AA5754 (Al–3 wt.% Mg) have beeneceiving considerable scientific and technological attentionecause they show little or no progressive damage accumulationrior to fracture [2].

Previous studies have shown that the formability and ductil-ty of these materials are significantly influenced by impurities,specially iron-rich intermetallic particles whereas the shearocalization is the major mechanism responsible for failure2–4]. Deformed 5xxx Al–Mg alloys are characterized by aelatively uniform distribution of dislocations because of theolute drag effect which prevents dislocations from undergo-ng further rearrangement to lower energy structures [5,6].he solute drag also increases the work-hardening rate in Allloys. In case of Al–Mg systems it has been shown that
g atoms exert a larger effect on work hardening than on
he solution strengthening [7]. Decreasing the temperatureas a similar effect; it has been observed that the work-

∗ Corresponding author. Tel.: +1 905 525 9140; fax: +1 905 521 2773.E-mail address: [email protected] (M. Niewczas).

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921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2008.01.065

; Work hardening; Strain rate sensitivity; Dislocation density; Fracture

ardening behaviour of low solute alloys is more sensitive tohe temperature changes than the work-hardening behaviour ofighly alloyed systems. This reflects the competition betweenhe dynamic recovery and the solute drag processes takinglace during plastic flow of these materials [7]. In general,he ductility of Al–Mg alloys decreases with increasing the

g content [8]. The torsion experiments reveal that Al–1%g alloy exhibits a well-defined stage IV of work harden-

ng in a broad temperature range between 78 K and 373 K9].

In contrary to Al alloys, deformed Al develops a three-imensional dislocation cell structure, the scale of whichepends upon the amount of deformation and the temperature.arker et al. [10] reported that the size of the cell structure

s reduced from 2 �m to 1.3 �m when Al is strained from0% to 30% at room temperature. Further it was shown thatith increasing strain, the homogeneous rotation of the individ-al cells within cell blocks becomes gradually more difficult,nd thus the size of cell blocks decreases. Chu and Morris11] studied the influence of the microstructure in pure Aln the work-hardening behaviour at 77 K and concluded that
re-existing subgrains may hinder the formation of disloca-ion cells. As a consequence, a high rate of work hardening isetained at high stresses leading to the improved combination oftrength and elongation. Generally, it is known that the ductility
mailto:[email protected]

dx.doi.org/10.1016/j.msea.2008.01.065

D.-Y. Park, M. Niewczas / Materials Science and Engineering A 491 (2008) 88–102 89

Table 1Composition of the AA5754 Al alloy (in wt.%)

Mg Mn Si Fe Cr Cu Ni V Zn Al

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C 3.21 0.24 0.05 0.09 –C 3.11 0.25 0.06 0.21 0.

f Al increases as temperature decreases due to the enhancedork-hardening capacity of this metal at a lower temperature

12].Most studies of the plastic deformation of Al-based alloy sys-

ems available in the literature have been carried out at room orlevated temperatures. However, there is an interest in under-tanding work-hardening behaviours of commercial alloys atower temperatures not only because of the fundamental naturef such studies but also from the practical side as many of theseystems are restricted from being used at elevated temperatures13–15].

The objective of this paper is to study mechanical proper-ies and deformation behaviour of commercial strip cast (SC),irect chill cast (DC) AA5754 Al alloys and of pure Al between.2 K and 295 K, a region of temperatures rarely explored in

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ig. 1. True stress vs. true strain characteristics at the initial strain rates of (a) 1.6 × 1or DC alloy (DC) at temperatures 4.2 K, 78 K and 295 K.

– – – – Balance0.01 0.01 0.01 0.02 Balance

he literature. Combined measurements of mechanical response,hanges in electrical resistivity during plastic deformation andEM studies of the dislocation substructure were employed toain a better understanding of the processes which control plas-ic deformation and processes which lead to failure in these

aterials.

. Experimental procedure

The experiments were carried out on SC and DC AA5754l alloys with a composition given in Table 1 and on high-
urity (99.9995%) aluminum. The as-received material was inhe form of cold-rolled sheets, 1 mm thick SC alloy and 7 mmhick pure Al. Al sheets were further cold-rolled in a laboratoryo a thickness of 3.7 mm. The as-received DC alloy was in a
0−4 and (b) 1.6 × 10−5 for SC alloy (SC) and pure Al (Al) and (c) 1.6 × 10−4

9 ence and Engineering A 491 (2008) 88–102

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0 D.-Y. Park, M. Niewczas / Materials Sci

mm thick sheet form in an annealed state with the grain sizef 23 �m.

Tensile test samples with gauge dimensions ofmm × 3.5 mm × 60 mm were machined from the as-receivedlloys parallel to the rolling direction. Samples made of Alad dimensions of 2 mm × 3.5 mm × 60 mm. Machined Al andC specimens were annealed at 350 ◦C for 30 min and 2 h,espectively, followed by air cooling, to produce recrystallizedaterials. The average grain size of annealed samples was

00 �m and 14 �m for pure Al and SC alloy. Some SC samplesere annealed at 450 ◦C for 2 h to study the effect of grain size

nd annealing temperature on mechanical properties; the finalrain size of these samples was 18 �m.

The specimens were deformed in a screw-driven tensile-esting machine at temperatures of 4.2 K, 78 K, and 295 K atconstant cross-head velocity corresponding to an initial strain

ate of 1.6 × 10−4 s−1 and 1.6 × 10−5 s−1. The strain rate sen-itivity (SRS) measurements were carried at 4.2 K, 78 K and95 K for pure Al and at 78 K for the alloys by repeating thenstantaneous rate change between 10−4 s−1 and 10−5 s−1 dur-ng deformation.

To investigate the evolution of deformation-induced defects,he electrical resistivity measurements were carried out in situuring deformation of the samples at 4.2 K. Employed was theour-point method for the measurement of potential drop acrosshe sample with reversing the current flowing through the circuit.he resistivity ratio, RR = r295 K/r4.2 K, i.e. ratio of the resistancet 295 K over the resistance at 4.2 K of pure Al was about 2000,ndicating a very low residual content of impurities and latticeefects in the material. In contrast, the alloy had a resistivityatio of about 2. Measurements of resistivity change due toeformation-induced defects in such material presents a chal-enge; the apparatus must be able to measure potential changesf the sample smaller than 10−9 V. Other experimental detailsre available elsewhere [16].

TEM studies of the dislocation substructure were performedn selected samples with a Philips CM12 STEM electron micro-cope operating at 120 kV. Thin foils were cut in the sectionarallel to the wider face of the samples and were prepared bystandard thinning procedure. The final polishing was carriedut in a solution of 30% nitric acid in high-purity methanolHPLC) at −35 ◦C with Tenupol-5 electropolisher. Fracture sur-ace observations in samples deformed to failure were carriedut using a scanning electron microscope (SEM) and an opticalicroscope.

. Results

.1. Deformation behaviour and fracture

Fig. 1 shows true stress–true strain characteristics of SC andC alloys and pure Al deformed at 4.2 K, 78 K, and 295 K. Ten-

ile curves of the alloys exhibit extensive serrations at 295 K,
ssociated with the strain aging known as Portevin–LeChatelierPLC) effect [17,18]. The stress drop of SC alloy, �σ, duringLC instabilities increases as the flow stress increases, frompproximately 6 MPa at the early stage of deformation to about
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ens heat-treated for 2 h at 350 ◦C and 450 ◦C, respectively. Tensile tests wereonducted at 295 K, with a strain rate of 1.6 × 10−4 s−1.

MPa at the later stages of deformation. DC alloy show stressrops in the range 2–7 MPa (Fig. 1(c)). The PLC effect is sup-ressed at 78 K and samples exhibit stable plastic flows withlong onset of uniform deformation followed by the fracture,

s shown in Fig. 1. At 4.2 K, the material shows flow instabili-ies arising from the adiabatic shear deformation, of a differentature than the ones observed at room temperature. The adia-atic deformation is accompanied by intensive acoustic emissionccurring with a higher frequency and louder acoustic clickshan those observed during the PLC deformation. The sampleeformed with a lower strain rate exhibits a lower frequency ofdiabatic instabilities, as shown in Fig. 1(b). Mechanical charac-eristics reveal that pure Al shows substantially lower flow stresshan SC alloys in the entire deformation range. The Al samplesxhibit a stable plastic flow up to the point of fracture, which isnitiated at the stress of approximately 50 MPa, 150 MPa, and10 MPa at deformation temperatures of 295 K, 78 K and 4.2 K,espectively. The fracture of the alloys occurs at much highertresses, approximately 290 MPa, 500 MPa, and 700 MPa at cor-esponding temperatures of 295 K, 78 K and 4.2 K, respectivelyFig. 1(a) and (c)). The alloy samples deformed at 78 K showhe highest uniform deformation before the fracture, which isomparable to the one observed in pure Al. Compared to SClloys, DC alloys show very short or no Luders deformationfter yielding, as seen in Fig. 1(c).

Fig. 2 shows the effect of annealing temperature on mechani-al properties of SC materials. It is seen that specimens annealedt 350 ◦C and 450 ◦C exhibit qualitatively similar character of theow curve. However, they show some differences in the yieldtress, the uniform deformation, and the stress at which frac-ure occurs. For example, the sample annealed at 350 ◦C has theield stress of about 130 MPa compared to 115 MPa for speci-ens annealed at 450 ◦C. Similarly, the fracture of the sample

nnealed at 450 ◦C occurs at the stress level of about 275 MPa
ompared to 290 MPa for the sample annealed at 350 ◦C.
Fig. 3 shows fracture surface observations of SC sampleseformed at three temperatures. The micrographs reveal an elon-


SC al

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Fig. 3. SEM observations of the fracture surface of

ated dimple structure characteristic of ductile failure. Little oro particles were found on the fracture surface and within theimples indicating that the failure is not a particle stimulatedrocess in SC alloys. One of the characteristic features of theracture surface in this material is the temperature dependence ofhe dimple size; one observes the formation of smaller dimplesn samples deformed at lower temperatures. Pure Al exhibitshe same trend, as shown in Fig. 4. It is seen that, at room tem-erature, a few very large dimples dominate its fracture surfacehereas, at 4.2 K, small dimples are evenly distributed across

he fracture surface. Fracture surface observations (Figs. 3 and 4)ndicate that the dimples formed in pure Al are much larger thanhese observed in Al alloys at the same deformation temperature,uggesting the higher ductility of Al, in agreement with tensileharacteristics of these materials.

Fig. 5 shows the normalized work-hardening rate, θ/μ, plotted
s a function of the normalized effective flow stress, (σ − σy)/μ,or both pure Al and Al alloys; the yield stress (σy) is chosent the onset of the Luders deformation for the alloys. The initialork-hardening rate of the alloy, equal approximately to μ/10,
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loys deformed at (a) 295 K, (b) 78 K, and (c) 4.2 K.

s a factor of almost two higher than the work-hardening ratef pure Al, of the order of μ/20. This is being observed for allemperatures and within the broad range of the flow stresses,ndicating that more effective dislocation storage is taking placeuring alloy deformation compared to pure Al. For pure Al, theork hardening decreases very fast at 295 K and 78 K indicating

hat the intensive dynamic recovery is taking place during plas-ic flow. At 4.2 K, dynamic recovery is suppressed; the rate ofork hardening decrease is much slower than at higher tempera-

ures, indicating that more effective dislocation storage occurs inhe substructure. Fig. 5 shows that the alloy exhibits the qualita-ively similar behaviour of the work hardening with respect to theeformation temperature, i.e. the work-hardening rate decreasesaster as temperature increases, suggesting that dynamic recov-ry processes operate more effectively at higher temperatureslso in SC material. At 4.2 K the work-hardening rate of the alloy
emains at an unusually high level above μ/20 for an extendedange of plastic flow. It is seen however that the hardening rateecreases faster in SC alloy than in pure Al, suggesting thatore intensive dynamic recovery operates in SC material. It has

92 D.-Y. Park, M. Niewczas / Materials Science and Engineering A 491 (2008) 88–102

Fig. 4. SEM observations of the fracture surface of pure Al deformed at (a)295 K and (b) 4.2 K.

Fig. 5. Normalized by temperature dependent shear modulus work-hardeningcharacteristics as a function of reduced flow stress. SC and Al denote SCalloy and pure Al, respectively. NA and AQ denote the natural aged and theas-quenched Al–Zn–Mg alloys, respectively, from the literature [19] for com-parison.

Fig. 6. Haasen plot characteristics obtained from strain rate sensitivity testswith rate changes between 10−4 s−1 and 10−5 s−1. SC and SC4 denote SC alloysannealed for 2 h at 350 ◦C and 450 ◦C, respectively; DC denotes direct cast alloyand Al denotes pure Al. For comparison with the present data, strain rate sensi-tms

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ivity results for AA6111 are reported from Ref. [21]. A6(1 h) represent AA6111aterial aged for 1 h at 180 ◦C, and A6(SS) represents the supersaturated solid

olution AA6111.

een reported in the literature that Al–Mg–Zn alloys deformed at.2 K exhibit an abnormally high hardening rate [19]. For com-arison with present results, Fig. 5 also shows work-hardeningehaviours of naturally aged and solution-treated and quenchedl–Mg–Zn alloys from Ref. [19]. It is seen that SC alloy exhibits

n even higher work-hardening rate than the Al–Mg–Zn systemoth in aged and quenched states within a broad range of flowtress.

.2. Strain rate sensitivity measurements

The experimental data of strain rate sensitivity measurementsre analyzed by plotting (�σ/T � ln ε)|T,Σ against (σ − σy), asuggested by Haasen [20]. Here �σ is the change in stress dueo the instantaneous change of strain rate, and σ, σy, ε, T, and Σ

epresent the applied stress, yield stress, strain rate, temperature,nd structure, respectively. It is well known that two importantarameters determined from the above analysis are: (i) the ther-odynamic SRS of a material obtained as a slope of a given

haracteristic and (ii) a thermal component of the flow stressiven by the intercept of a plot with the vertical axis. Fig. 6hows Haasen plots for both pure Al and SC alloys obtained inhis work; for comparison, results for supersaturated solid solu-ion (SSS) and precipitation-hardened AA6111 alloys availablen the literature [21] are also included in the graph. It is seen thatC alloys exhibit strain rate sensitivity similar to other classes ofA6111 materials with the value of about 1.4 × 10−4 K−1. The

lloy annealed at 350 ◦C has a slightly higher positive intercepthan the one annealed at 450 ◦C. The intercept of SC alloy islose to that of SSS AA6111, whereas aged AA6111 exhibits a
igher intercept. The graph also presents data for DC AA5754lloy tested in this work, with the same nominal compositions SC alloys but higher contents of iron, i.e. 0.2 wt.% Fe in DCaterial versus 0.09 wt.% Fe in SC alloy (Table 1). It is seen that

D.-Y. Park, M. Niewczas / Materials Science

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ofAriitptieuutiseen that pure Al accumulates in average about 7 × 10 mdislocation density before the fracture. The total dislocationdensity stored in SC alloy before the fracture is higher andis of the order of 1.5–1.6 × 1016 m−2. DC material accumu-

ig. 7. Activation volume as a function of the effective flow stress on a log scaleor SC alloy and pure Al. Activation are is given on the right axis.

C data shows the lowest Haasen intercept among all systemsonsidered, which also has positive value.

Strain rate sensitivity of pure Al increases as tempera-ure decreases; the SRS values are equal to 2.98 × 10−4 K−1,.88 × 10−4 K−1, and 0.29 × 10−4 K−1 at 4.2 K, 78 K, and95 K, respectively. Fig. 6 shows that the intercept of pure Alata on the Haasen plot is equal to zero at room temperature,hereas it is slightly negative at lower temperatures. The reasonsehind this behaviour will be discussed later.

During the plastic flow of a material, dislocations overcom-ng the obstacles move from a stable state (static equilibrium)o an unstable state (second equilibrium). The apparent activa-ion area, �a, swept by a dislocation segment during the glideetween the unstable and stable equilibrium states at a giventress is evaluated from the following equation [22,23]:

a = MkT

b

∂ ln ε

∂σ

∣∣∣∣T,Σ

= MkT

mb(σ − σy)(1)

here k is the Boltzmann constant, M is the Taylor factor, b ishe Burger’s vector, and m is the strain rate sensitivity in thengineering definition.

Fig. 7 shows the activation volume, V′ = b �a, for SC alloynd pure Al as a function of the effective flow stress on a logcale, where b is the Burgers vector (∼0.28 nm) and �a is thectivation area. It is seen that pure Al exhibits an activationrea with a strong temperature dependence. For example, at theffective flow stress of 25 MPa, the activation areas have beenetermined to be approximately 10 nm2, 30 nm2, and 175 nm2

t 4.2 K, 78 K, and 295 K, respectively. For a given flow stress,he activation area increases as temperature increases due tohe higher contribution from the thermal energy. At a constantemperature, the activation area decreases as the flow stressncreases (Fig. 7). The presence of solute atoms reduces the
ctivation area of a material; the activation area of SC alloy at8 K decreases from approximately 9 nm2 to 2.5 nm2, close tohe values observed for pure Al at 4.2 K (Fig. 7).
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and Engineering A 491 (2008) 88–102 93

.3. Electrical resistivity

Fig. 8(a) shows raw data of deformation-induced resistivityver the residual resistivity value as a function of the true strainor Al, SC and DC alloys deformed at 4.2 K. The super-purel used in this work enabled us to measure change in electrical

esistivity in this material during deformation at 78 K, which isncluded in the graph. It is seen that, in every case, the resistivityncreases parabolically with the strain, indicating that deforma-ion is homogeneous on the macroscopic scale. Compared toure Al, the alloys show higher scattering in resistivity data dueo the contribution from solute and impurity scattering. Assum-ng that all deformation-induced defects are dislocations, thelectrical resistivity is recalculated into dislocation densities,sing the value of 1.8 × 10−25 � m3 for specific resistivity pernit length of dislocation in Al [24]. The data in Fig. 8(a) canhus be interpreted as the evolution of dislocation density dur-ng plastic deformation for both pure Al and Al alloys. It is

15 −2

ig. 8. Evolution of defect-induced resistivity change during deformation as aunction of (a) true strain and (b) the effective flow stress for SC alloy and purel.

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ates in average 1.1 × 1016 m−2 dislocation density before theracture.

Fig. 8(b) shows deformation-induced resistivity as a functionf effective flow stress for all materials studied. For pure Al,esistivity increases roughly parabolically with the flow stressp to the fracture. The resistivity characteristics for alloys showppreciable scattering of data points; the scattering increases ashe flow stress increases. This reflects the unstable deformationue to adiabatic shearing of the lattice, which contributes tohe scattering of the flow stress at the same dislocation content.ig. 8(b) shows that the resistivity evolution as a function oftress for the SC alloy follows closely the characteristic of purel. It is seen that, for a given flow stress level, resistivity of DC

lloy is lower than that of pure Al or SC material in a broadange of stresses. As previously shown, the resistivity data cane recalculated into the dislocation densities. This data showshat, at the same flow stress level, Al deformed at 78 K produceshe highest dislocation densities among all materials studied.

.4. TEM observations

Fig. 9 shows the microstructure of SC alloy beforend after plastic deformation at different temperatures. Thennealed SC material contains non-homogeneously distributedne particles smaller than 1 �m in diameter, produced dur-

ng strip casting of the alloy [25]. These particles arebserved both in the form of strings aligned along rollingirection as shown in Fig. 9(a) and also in the form of ran-om clusters distributed non-homogeneously in the Al-richatrix.Fig. 9(b) shows TEM observations of the microstructure in

he SC material deformed 10% at room temperature. Disloca-ions are stored in the entire volume of the material and onean recognize a formation of characteristic cell structure withhe higher dislocation density accumulated in the cell wall andomewhat less dense in the cell interior. The average cell sizef this structure is of the order of 0.5–0.7 �m. It is seen thathe formation of these features is not associated with the secondhase particles as no particles are found in this area. The alloyeformed 17% at 4.2 K exhibits on average much higher disloca-ion densities. In some areas, as is shown in Fig. 9(c), dislocationsorm insulated cells with a size of the order of 0.2–0.5 �m. Inighly deformed samples (Fig. 9(d)) the dislocations are storeduite homogeneously in the volume of the specimen and theres no clear separation into higher and lower dislocation densityreas as it was seen in the sample deformed at room temperature.

Fig. 10 shows examples of the dislocation substructure inure Al deformed till fracture, i.e. 35%, 30%, and 43% of straint 295 K, 78 K, and 4.2 K, respectively. In comparison with theC alloy, the substructure of pure Al is markedly different. In allases, the micrographs show well-developed substructures withislocation sub-boundaries separating areas of recovered cell
nteriors. There are substantially less defects found inside theubgrains than was observed in SC alloy. Smaller cells are pro-uced at lower temperatures. The scale of the subgrains rangesetween 1 �m and 3 �m in material deformed at 295 K, 0.6 �m
dmsr

and Engineering A 491 (2008) 88–102

nd 1.7 �m in Al deformed at 78 K, and 0.2 �m and 0.7 �m inl deformed at 4.2 K.TEM observations (Figs. 9 and 10) show that the disloca-

ion content and the nature of dislocations stored in SC materialre different from these found in pure Al. Pure Al exhibits atrong tendency towards grain subdivision and development ofow misoriented sub-boundaries within larger grains. This hasot been observed in SC alloy, which tends to form a rela-ively homogeneous network of dislocations, particularly at lowemperatures.

. Discussion

.1. Deformation and work-hardening behaviour

Strip cast aluminum alloys deformed at 4.2 K and 295 Kxhibit plastic flow instabilities associated with two differentrocesses operating during tensile deformation of the material.n the subsequent section, we discuss briefly these processes,o develop a better understanding of the hardening behaviour in

aterials studied in this work.At room temperature, PLC deformation is responsible for

he serrated plastic flow of SC alloys. The magnitude of theerrations increases as strain increases. The PLC instabilities arelassified into five different types A–E [18]. The SC alloy shows-type serrations through most of the plastic flow at 295 K, withsmall onset of type A instabilities observed at the early stage ofeformation (Figs. 1 and 2). As the strain rate decreases, moreolute atoms diffuse to dislocations and the magnitude of stressrops (oscillation) increases (Fig. 1). Present results agree withther studies of the Portevin–LeChatelier effect in Al alloys,ublished in the literature, e.g. [17,26].

Load instabilities at 4.2 K arise from the thermomechanicalnstability in the substructure occurring during adiabatic shear-ng of the lattice. These processes have been studied in detail forarious systems in the 1950s and 1960s including pure metals27,28] and Al alloys [29–31]. When shear deformation initiatest low temperatures due to the low thermal conductivity of theaterial, the heat generated during this process is localized insmall volume of the deformed specimen. Basinski [31] mea-

ured that during adiabatic deformation the temperature of pureopper increases locally to 60 K. One can expect that the localncrease of the temperature in the SC sample during adiabatichearing can be at least of this order of magnitude or higher. Theocalized heating produces softening of the substructure alonghe deformation path and leads to the shear localization on theength scale crossing a number of grain boundaries. The adi-batic deformation produces load drop instabilities associatedith characteristic sound clicks and affects the work-hardeningehaviour of the material at 4.2 K. The adiabatic shear localiza-ion can also be initiated at a very high strain rate where the timeor the thermal diffusivity is short and the generated heat cannote dissipated in the volume of the material [32,33]. In samples
eformed at a low strain rate, the heat per unit time generated byoving dislocations is lower and the local temperature increases
lowly. Hence, the alloy deformed at 4.2 K with a lower strainate exhibits less frequent stress drops, whereas it shows exten-


Fig. 9. TEM observations of the microstructure of SC alloys: (a) prior to deformation, (b) the substructure developed after 10% of deformation at 295 K at thefl formas f defot

st

m

ow stress of 130 MPa, (c) dislocation substructure developed after 17% of deubstructure developed in the SC sample deformed till fracture, i.e. after 33% oensile direction).

ive serrations with a high frequency of load drop at 295 K dueo PLC deformation.

In comparison with Al alloys, pure Al shows a stable defor-ation at 4.2 K because of the better thermal conductivity, which

aTma

tion at 4.2 K corresponding to the flow stress of 320 MPa, and (d) dislocationrmation at 4.2 K corresponding to the flow stress of 520 MPa (TA denotes the

llows fast dissipation of the heat produced during the slip event.his is also reflected in work-hardening characteristics of bothaterials (Fig. 5), showing that the work-hardening rate of SC

lloy at 4.2 K decreases faster than that of pure Al, suggesting

96 D.-Y. Park, M. Niewczas / Materials Science and Engineering A 491 (2008) 88–102

F re atfl MPa,

tbisSt

a

ig. 10. TEM observations of the microstructure of pure Al deformed till fractuow stress of 50 MPa, (b) 30% of deformation at 78 K at the flow stress of 150

hat the recovery rate is higher in alloy than pure metal. Suchehaviour of the work-hardening results from the adiabatic local-
zed heating causing softening of the substructure and inducinghear localization, which increases the apparent recovery rate inC alloy. In the absence of adiabatic deformation one can expect
hat the recovery rate in SC material should be lower.

tate

different temperatures, i.e. (a) 35% of deformation at 295 K corresponding to a(c) 43% of deformation at 4.2 K corresponding to a flow stress of 415 MPa.

Present results reveal that SC alloy deformed at 4.2 K exhibitsremarkably high strength of the order of 700 MPa just before

he fracture (Fig. 1). This is comparable to the strength levelchieved in precipitation-hardened materials [19,21]. Comparedo those age-hardenable Al–Mg and Al–Zn–Mg alloys, the hard-ning rate of SC alloy at 4.2 K is much higher in the wide

D.-Y. Park, M. Niewczas / Materials Science

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ig. 11. The effect of dislocation storage on the strain hardening of SC alloynd pure Al, respectively. θo is defined as the slope of a straight line though therigin in the region of linearity (see text for details).

ange of flow stresses (e.g. Fig. 5). It has been suggested that,n non-heat-treatable alloys, solute atoms lower stacking faultnergy, therefore reducing dynamic recovery and enhancingork-hardening capacity [34]. The strong hardening of SC

lloy must be attributed to the influence of alloying elementsresent in the solid solution and it is evident, from the resultshown in Fig. 5, that this influence is particularly effective atow temperatures. Lloyd [7] studied mechanical properties andork-hardening behaviours of a 5xxx series of Al–Mg alloysown to 85 K and concluded that out of all elements present inhese systems Mg inserts the strongest effect on the hardeningehaviour through the solute drag contribution, which inhibitsislocation rearrangement and increases hardening rate.

With the exception of 4.2 K, which will be discussed later,he work hardening decreases approximately at the same rateor both materials at 78 K and slightly faster in Al than SC at98 K (Fig. 5), indicating that the solute atoms and impurities doot change substantially the kinetics of dynamic recovery pro-esses at this temperature range. Thus, these processes must beoverned by dislocation–dislocation interactions. As discussedreviously, the higher apparent recovery rate in SC materialt 4.2 K compared to pure Al is a direct consequence of thediabatic shear deformation, which produces substantial soften-ng of the microstructure on the macroscopic scale and affectshe work-hardening behaviour of the alloy. To provide a moreetailed picture of the storage mechanisms operating duringeformation of SC alloy in comparison with pure Al, Fig. 11hows characteristic of (σ − σy) dσ/dε plotted as a function offfective flow stress, σ − σy. In the Kocks–Mecking descrip-ion [35], the strain-hardening rate is expressed in terms of twoompeting mechanisms: athermal storage of dislocations andynamic recovery:

= θo − θr(γ, T ) = θo − kτ (2)

here θ is the macroscopic-hardening rate, θo is the athermaltorage term of the strain-hardening rate, θr is the dynamic recov-ry term, γ is the shear rate, T is the temperature, k is the constant

r5td


nd τ is the shear stress. The rate of the work hardening decreases proportional to k representing the rate of dynamic recovery.n the case of deformation of polycrystalline materials, one canrite by application of the Taylor factor that [35]:

σ − σy)θ = (σ − σy)[θo − θr(σ, ε, T )] (3)

here σ is the true tensile stress, σy is the yield stress, and ε ishe strain rate. The initial slope of the relationship (3) representshe athermal storage of dislocations, whereas deviation from lin-arity occurring at higher stresses indicates the onset of dynamicecovery process. In the present experiments θo for Al is equalo about 1.3 GPa, whereas for SC alloy θo is found to be equal topproximately 2.2 GPa. θo is thus almost exactly equal to ∼E/50or pure Al and ∼E/32 for SC alloy. The athermal storage term,o, is larger for SC alloy than pure Al. This indicates that SClloy more effectively accumulates dislocations in the structure.

The athermal-storage contribution θo is related to the dislo-ation mean free path, l, and dislocation density, ρ, through theollowing equation [35]:

o = αμ

2l√

ρ(4)

here μ is the shear modulus and α is the constant. Refer-ing back to Fig. 11 and data at 4.2 K, SC characteristic deviaterom the linearity at the effective flow stress of approximately0–80 MPa whereas pure Al at about 80–100 MPa, correspond-ng to average stored dislocation densities of approximately× 1013 m−2 and 3 × 1014 m−2, respectively (Fig. 8(b)). TEMbservations reveal independently that at this deformation stageoth materials develop a relatively homogeneous dislocationubstructure. One can estimate that the average dislocation meanree path calculated based on Eq. (4) is equal approximatelyo 0.3 �m for pure Al and about 0.6 �m for SC alloy. Thiss much smaller than the initial grain size of these materials600 �m Al and 14 �m SC alloy) and much larger than the aver-ge distance between solute atoms in the glide plane in SC alloy1.5 nm. In case of pure Al, the dislocation mean free path

orresponds rather closely to the size of the cell structure pro-uced after large deformation. More detailed discussion of thisspect of the present work will be provided later in Section 4.3.owever, based on the above estimate one can conclude that

he dislocation–dislocation interactions (not dislocation–grainoundaries or dislocation–particle or dislocation–solute interac-ions) are important in the dislocation storage in these materials.he dislocation mean free path corresponds relatively closely to

he size of the observed cell structure in Al, which indicates thatislocation storage occurs on cell walls and therefore processesccurring within the cell walls determine the rate of the dynamicecovery in this metal at least at the low temperature regime.

After a certain amount of plastic deformation, (σ – σy) dσ/dε

urves (Fig. 11) develop a negative curvature. At 4.2 K, theownward curvature occurs at approximately the same flowtress of about 350 MPa for both pure Al and SC alloy, cor-
esponding to average dislocation density of approximately× 1015 m−2 (Fig. 8). Although not confirmed, it is unlikely that
he point at which these curves turn down is associated with theevelopment of any form of damage process because it occurs

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elatively far below the ultimate tensile strength of these materi-ls. The alternate possibility is that some accelerated recovery isaking place in the substructure, initiated at a dislocation densityf the order of 4 × 1015 m−2. It is noted that at higher tempera-ures (e.g. 78 K) dynamic recovery occurs much earlier in pure Alhan SC alloy. This implies that SC material is more resistant tohis process, which intensifies after more advanced deformationhan in pure Al.

TEM observations reveal that, at the flow stress of00–350 MPa, Al develops a well-defined cell structure whereasC produces a relatively homogeneous, Taylor-like dislocationrrangement. We estimate that the relative volume fraction ofhe cell walls to the cell interiors in pure Al is approximately0–70%, whereas the density of dislocations stored in the wallselative to the dislocation density inside the cell is in the ratio0:1 (Fig. 10(c)). Assuming that the simple rule of mixturepplies, i.e.:

T = fρw + (1 − f )ρi (5)

here ρT is the total dislocation density stored in the mate-ial, ρw is the dislocation density stored in the wall, ρi ishe dislocation density stored in the cell interior and f is theolume fraction occupied by the cell walls, one can estimatehat the downward curvature of (σ − σy) dσ/dε curves or theccelerated dynamic recovery occurs in Al when the disloca-ion density in the wall and the cell interior are approximately:w = 1.1 × 1016 m−2 and ρi = 1.1 × 1015 m−2. As mentionedreviously for SC alloys this process occurs at the average dislo-ation densities ∼4 × 1015 m−2. These figures correspond to theverage dislocation spacing in SC alloy of the order of 16 nm andpproximately 10 nm and 30 nm in the wall and inside the cell ofure Al, respectively. It is seen that dynamic recovery intensifieshen the spacing between dislocations approaches 10–15 nm,

ndicating that similar fundamental processes responsible forislocation annihilation operate in Al and SC alloy.

It will be of interest to apply the same approach to estimatehe average dislocation spacing in both materials at the point ofracture. In this case, the dislocation densities are approximatelyw = 1.9 × 1016 m−2 and ρi = 1.9 × 1015 m−2 in the cell wall and

he cell interior of pure Al and ∼1.5 × 1016 m−2 in SC alloy.his represents average dislocation spacing in the SC alloy of

he order of 8 nm and approximately 7 nm and 23 nm in the wallnd inside the cell of pure Al, respectively. It is striking that theverage dislocation spacing in SC alloy and in highly dislocatedalls of Al correspond very well to each other. This may indicate

hat, independent of the composition of the system, the fracturenitiates by some form of dislocation collapse occurring in thereas where the dislocation densities approach a critical levelr a critical spacing for spontaneous annihilation. We estimatehis spacing for Al and SC alloy to be between 7 nm and 8 nmt 4.2 K.

In pure Al deformed at 78 K the total dislocation density
t the fracture is ∼1.8 × 1015 m−2 (Fig. 8). The volume frac-ion of the cell walls to the cell interiors is approximately0–80% and we estimate the similar 10:1 ratio for the den-ity of dislocations stored in the wall relative to cell interior
irtt


n Al deformed at 78 K (Fig. 10(b)). The dislocation densi-ies, ρw and ρi, are thus approximately: ρw = 6.4 × 1015 m−2

nd ρi = 6.4 × 1014 m−2, which represents average dislocationpacing of 12 nm and 39 nm in the wall and inside the cell ofure Al at 78 K. We do not attempt to evaluate these figures forl deformed at 295 K or for alloys, as they would be subjected

o a large error.

.2. Mechanism of plastic deformation

Haasen characteristics show that one mechanism of plas-ic deformation operates during deformation of pure metalnd also in Al alloys studied in this work (Fig. 6). Haasenlot of pure Al deformed at room temperature passeshrough the origin, suggesting that the flow stress is deter-

ined by dislocation–dislocation interactions and obeys theottrell–Stokes law, which is also confirmed in Fig. 7. A good

inear relationship between the activation volume and the effec-ive flow stress (Fig. 7) indicates that Cottrell–Stokes law holds atow temperatures as well as room temperature because it implieshat τeV′, where τe is the effective flow stress and V′ is thectivation volume, is independent of strain, which is known ashe modified Cottrell–Stokes law [36]. The deviation from theinearity for the alloy in Fig. 7 is usually attributed to the impu-ity effect on the short-range stress which gives rise to a higholute–dislocation component [37]. Al shows a slightly nega-ive intercept at lower temperatures in the Haasen plot (Fig. 6),hich again suggests that its substructure provides a more strain

ate sensitive component for the flow stress than forest dislo-ation interactions. This athermal component may arise fromhe nature of the dislocation substructure produced at low tem-eratures characterized by the higher volume fraction of theell walls comprising the higher dislocation content and largerisorientations across the walls.The y-axis intercepts of the Haasen plot for alloys are always

igher than for pure Al, due to the presence of solute atoms inhe aluminum matrix, which provide a larger thermal compo-ent to the flow stress. If more solutes precipitate as in the casef DC alloy, the thermal contribution decreases and the value ofintercept is lower. This will also suppress or eliminate Luderseformation as observed experimentally in Fig. 1(c). Therefore,t is not surprising that SC alloy, containing more alloying con-ents in the solid solution than DC alloy due to the effect ofhe thermomechanical process, shows the higher positive inter-ept on the Haasen plot. At the same time, SC alloy annealedt 450 ◦C, indicated as SC4 in Fig. 6, shows a lower thermalomponent than our reference SC material annealed at 350 ◦C.t also shows the lower yield stress, but the work-hardeningehaviour of these two specimens is the same (Fig. 2). Thiss attributed to two combined factors: one related to the grainize difference and the second to the composition of the matrix.he strain rate sensitivity of the material is very sensitive to theompositional change; small changes in the solid solution are
mmediately visible on the Haasen characteristic of the mate-ial [22]. Optical observations reveal that, with the exception ofhe grain size, there is no substantial difference in the nature,he size or distribution of precipitates present in SC samples

ence and Engineering A 491 (2008) 88–102 99

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Table 2Values of the fitting parameters n, k1, and k2 to Eqs. (6) and (7)

Materials

Al (78 K) Al (4.2 K) SC (4.2 K) DC (4.2 K)

n values 0.54 0.5 0.34 0.29k1 (×108) 2.52 4.22 8.12 7.07k

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D.-Y. Park, M. Niewczas / Materials Sci

nnealed at 450 ◦C and 350 ◦C. This suggests that the largerrain size of the sample annealed at a higher temperature affectshe yield stress and should lower the thermal component of theow stress, as observed experimentally. It is also interesting toote that the y-axis intercept of SC alloys is close to the one ofolution treated AA6111 [21]. This indicates that both materialsre characterized by a similar strength of obstacles, which pro-ides the similar thermal component to the flow stress duringhermal activation. When fine precipitates are formed in agedA6111, the thermal component is even larger than observed in

olution treated material [21].

.3. Electrical resistivity and TEM

Analysis of the electrical resistivity data for pure Al and Allloys show that the following relationship between the effectiveow stress and the total dislocation density holds up to largetresses:

− σy = MαEbρn (6)

here E is the Young modulus, b is the Burgers vector, M is theaylor factor ∼3.01, α is a constant and n is another constant.ig. 12 plots log of the effective flow stress as a function of

he log of the dislocation density, i.e. f(log(σ − σy)) = log(�ρ)or different materials. The slope of this figure gives a value ofhe exponent, n, in Eq. (6), which varies between pure Al, Allloys and deformation temperature (Fig. 12). The n values deter-ined from the linear regression analysis are given in Table 2.l deformed at 4.2 K shows that the flow stress is proportional to

he square root of the total dislocation density, i.e. σ ∼ ρ1/2 (e.g.
38]). SC and DC alloys exhibit substantial deviations from thiselationship with the n exponent equal to 0.29 and 0.34 for DCnd SC alloy, respectively (Table 2). It is important to emphasizehat Eq. (6) reflects the relationship between the total disloca-
ig. 12. Flow stress as a function of the total dislocation density for Al, SC andC alloy. Slopes of these relationships give values of the exponent n in rela-

ionship given in Eq. (6). The n values are 0.54, 0.5, 0.34, 0.29 for Al deformedt 78 K and 4.2 K, SC alloy deformed at 4.2 K and DC alloy deformed at 4.2 K,espectively (see the text for details). Linear regression lines are also includedn the figure.

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2 2.24 1.10 0.40 0.10

ion density stored in the materials during deformation at 4.2 Knd the flow stress. It has been shown on copper that total dis-ocation density produced during deformation at 4.2 K has twoomponents: one contributes to the flow stress and resistivityunrecoverable defects) and the other contributes only to thelectrical resistivity (recoverable defects) [16]. Our preliminarytudies of this effect suggest that it is similarly true for the casef low temperature deformation of pure Al. One can thus expecthat, after accounting for these two dislocation density compo-ents, n values corresponding to the relationship between theensity of unrecoverable defects and the flow stress will be dif-erent from those given above. Basinski and Basinski [39], afterompiling various data available in the literature, reported forure copper the relationship: τ ∼ ρ0.43 between the flow stressnd the dislocation density, which is attributed to the effect ofther obstacles such as impurities on the flow stress.

Let us now discuss in more detail the evolution of electricalesistivity (total dislocation density) in pure Al. For a given flowtress, the resistivity change of Al at 4.2 K is lower than at 78 KFig. 8(b)). This indicates that, in order to support a given stress, aigher dislocation density must be stored in the material at higheremperatures (see also Fig. 12). This reflects the strength of thebstacle structure that a given dislocation distribution exerts onhe mobile dislocations during plastic flow. It is clear that theislocation network produced at 78 K provides weaker obstacleshan the network developed at 4.2 K, so that more defects areequired to resist the glide of mobile dislocations at the sametress level.

The strength of the obstacle structure is reflected in the valuesf the activation area shown in Fig. 7. It is seen that the activationrea at 78 K is larger than at 4.2 K indicating that dislocationsan more easily overcome the obstacles. On the other hand thectivation area of SC alloy at 78 K resides close to the range ofhe one of pure Al at 4.2 K, which suggests that, at the same tem-erature, SC alloy develops stronger and more stable obstacletructure than Al, in agreement with mechanical characteristics.

The evolution of the dislocation mean free path in Al and Allloys is obtained from the variation of dislocation density withtrain (Fig. 8(a)) according to the expression l = M/b(dρ/dε),here M is the Taylor factor ∼3.01, b is the Burgers vector, ρ is

he density of dislocations and ε is the true strain [35]. Fig. 13hows the dislocation mean free path as a function of effectiveow stress. For materials deformed at 4.2 K the mean free pathecreases exponentially and for DC and SC alloy it stabilizes
t the value of 0.1 �m at the fracture. In case of Al deformedt 4.2 K and 78 K the mean free path at the point of fracture isqual to 0.4 �m and 1.5 �m, respectively. It is interesting to note

100 D.-Y. Park, M. Niewczas / Materials Science

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ture. In all materials studied, fracture occurs reproducibly withina narrow range of strength, work-hardening rate, dislocationdensity and the dislocation mean free path characteristic fora given material and deformation temperature. The fracture

ig. 13. Evolution of the dislocation mean free path for SC and DC alloyseformed at 4.2 K and Al deformed at 4.2 K and 78 K.

hat at 78 K the mean free path decreases roughly linearly withhe stress, what suggest that plastic deformation of aluminum at8 K is controlled by different kinetics of storage and recoveryhan at 4.2 K.

In all above cases, the mean free path of dislocations deter-ined from electrical resistivity measurements shows a very

ood agreement with the scale of the dislocation substructureroduced during deformation as seen by TEM. In the case of purel deformed at 4.2 K, the mean free path of 0.4 �m at the fracture

orresponds to the size of the cell structure, which indicates thatell boundaries provide significant rates of dislocation storage.n Al alloys, on the other hand the dislocations are stored quasi-omogeneously in the volume and the mean free path reflectshe spacing between obstacles, i.e. forest dislocations within theislocation network.

As the mean free path decreases and the density of dislocationncreases, dynamic recovery plays a more important role duringlastic deformation. This process has been discussed in Section.1 based on the behaviour of the work-hardening rate and it wasoncluded that at temperatures of 78 K and 298 K the kineticsf dynamic recovery of Al and Al alloy is very similar and iteems to be controlled by the dislocation annihilation occurringithin highly dislocated areas of the substructure. The electrical

esistivity results can provide insight into these processes occur-ing at 4.2 K. As known, the evolution of the dislocation densityith strain has commonly been assumed to be represented by

he sum of two contributions. One contribution is proportionalo the square root of the dislocation density associated with theardening storage rate and another contribution is proportionalo the dislocation density associated with the recovery rate. Theate of dislocation density change with strain is thus expresseds

dρ √
dε
= k1 ρ − k2ρ (7)

here k1 is the constant representing the work-hardening rate,hereas k2 is another constant proportional to the rate of

Fott


ynamic recovery. In the present work these contributions haveeen determined by fitting the relationship of dρ/dε = f(ρ) with1√

ρ − k2ρ. The experimental data points and fits are shownn Fig. 14, and constants k1 and k2 are given in Table 2. Forure Al, the constant k1 increases as the temperature decreasesrom 78 K to 4.2 K and is higher for SC and DC alloys than purel. These results, in addition to the hardening characteristicsiscussed in Section 4.1, give direct evidence that the work-ardening component of the flow stress is larger in Al alloyshan in pure Al. On the other hand, Table 2 shows that for purel, the k2 value is larger at 78 K than at 4.2 K and k2 in Al alloys

s lower than in pure metal. This indicates that annihilation ofislocations occurs with a higher rate at higher temperatureshere the thermal energy is available, whereas these processes

re slowed down in Al alloys. This is attributed to the effect ofolute atoms, which increases the dislocation storage capacitynd suppresses dynamic recovery possibly through their influ-nce on the stacking fault energy, affecting the nature of theislocation structure produced in these materials. It should bemphasized here that the above analysis reflects changes occur-ing entirely within the dislocation substructure without regardo the macroscopic deformation behaviour of the materials. Asiscussed in Section 4.1, the work-hardening characteristic ofC alloy at 4.2 K (Fig. 5) gives a higher apparent recovery rate

han in pure Al, as a result of adiabatic deformation. One can seehat electrical resistivity data are insensitive to this effect, as theignal arises from the integrated defect content in the sample,hich does not decrease during adiabatic shearing.

.4. Fracture

Let us finally discuss the way our materials approach frac-

ig. 14. Rate of the dislocation density change with the strain dρ/dε as a functionf the dislocation density ρ. The characteristics are fitted to the rate equa-ion: dρ/dε = k1

√ρ − k2ρ with k1 and k2 constants given in Table 2. Both

he experimental data points and the fitting curves are shown in the graph.

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urface observations reveal that all materials undergo ductileracture by the nucleation and growth of voids. In the case ofl alloys failure is not particle-stimulated and particles seem

o have weak influence on ductility. From the work-hardeningharacteristics (Fig. 5), it can be deduced that, at 4.2 K, fracturen Al alloys occurs far before the work-hardening capacity isxhausted, i.e. before Considere criterion is meet. Therefore,here must be some mechanism that triggers the nucleationf voids and leads to the material failure. In Section 3.1, itas shown that SC samples exhibit the largest elongation at8 K where both Portevin–LeChatelier and adiabatic deforma-ion are suppressed. This suggests that these processes facilitatehe premature necking and fracture of samples. The litera-ure provides a considerable amount of data suggesting thatLC deformation has a detrimental effect on ductility, e.g.40–42].

There is much less data available regarding the effect ofdiabatic deformation on deformation behaviour and frac-ure. Present results reveal that the adiabatic deformationffects the macroscopic work-hardening rate of the alloys and,lthough it may aid in nucleating strain localization, adiabatichear is not a primary cause of the fracture. This is con-luded from the fact that the homogeneous elongation of theample deformed with a higher strain rate, where the adia-atic deformation operates very intensively, is not lower thanhe elongation of samples deformed at a lower strain rateFig. 1).

TEM observations show that there are significant differencesn the nature of the dislocation substructure developed in Alnd Al alloys. The intriguing feature is that the average spacingetween dislocations within highly dislocated areas at the pointf fracture is represented by the same figure in both Al and Allloys. Thus, it seems that the failure occurs when a certain dis-ocation density limit or a critical dislocation spacing is achievedn the substructure to allow for spontaneous annihilation of dis-ocations and global collapse of the dislocation network in thesereas. We estimate a critical dislocation spacing of about 8 nm inl and Al alloys at 4.2 K and ∼12 nm in pure Al at 78 K. Thesegures correspond well to Brown’s theoretical estimations ofritical height for the athermal collapse of dislocation dipoles inopper [43].

Destabilization of the dislocation substructure at stresses,s these developed before the fracture, will cause persistentow localization in softened volumes of the material and will

ead to nucleation of voids, their growth and eventually neck-ng and failure. Assuming that voids form in highly dislocatedreas (e.g. cell walls in Al), this should produce a fractureurface with a high density of small dimples at low temper-tures as there are many places where voids will nucleate.ontrarily, the coarser scale of the substructure developed atigher temperatures provides less potential sites for the voiducleation. Thus, there will be fewer voids formed; conse-uently, a few voids with the large size eventually dominate
n the fracture surface. This is reflected in the nature ofhe fracture surface produced in Al and Al alloys at 295 Khich is filled up predominantly with a few large dimples
Figs. 3 and 4).aE


. Summary and conclusions

Deformation behaviours during tensile tests of strip andirect chill cast AA5754 Al alloys and of high-purity Al haveeen studied. The focus is on the low temperature regimehere mechanical properties of these materials have not been

ufficiently explored. Al alloys show two kinds of flow insta-ilities, at 295 K resulting from dynamic strain aging knowns Portevin–LeChatelier effect and at 4.2 K from adiabaticeformation. At 78 K, the flow instability is suppressed andhese materials exhibit a homogeneous deformation. Plas-ic deformation of high-purity Al is governed mainly byislocation–dislocation interactions, and a stronger athermalomponent of the flow stress observed at lower temperaturesrises from the nature of the dislocation cell structure producedn this material. Al alloys exhibit a larger thermal component ofhe flow stress due to the presence of solute atoms in the matrix.

Electrical resistivity measurements reveal that all materialsxhibit different kinetics of dislocation storage and recovery,hich is reflected both in the constitutive relationship betweenow stress and the total dislocation density and in the way theseaterials approach fracture. The results reveal that SC alloy

eformed at 4.2 K exhibits a remarkably high work-hardeningapacity, with the work-hardening rate of the order of μ/20 forost of the plastic flow, attaining eventually the strength of

00 MPa just before the fracture. The work-hardening capacityf SC alloy at 4.2 K is more effective than in precipitation-ardened materials probably due to a higher dynamic recoveryomponent in heat-treatable alloys.

The evolution of the dislocation mean free path shows a goodgreement with TEM observations of the substructure and sug-ests that dislocations are accumulated within the cell walls inure Al and quasi-homogeneously in the volume of Al alloys.articles have small or no influence both in the process of dis-

ocation storage and during the fracture.Direct measurements of the evolution of defect content during

eformation at 4.2 K suggest that the rate of dynamic recoveryccurring within a dislocation network is slower in Al alloyshan in pure Al. However, adiabatic deformation, which oper-tes during plastic flow, affects the macroscopic work-hardeningehaviour of alloys and gives the higher apparent recovery ratehan observed in pure metal.

Present results show that fracture occurs when dislocationensity approaches a critical level locally in highly dislocatedreas. Dislocation density figures give an estimate for the criticalpacing between dislocations in these areas to trigger sponta-eous annihilation. At 4.2 K this figure is ∼8 nm for all materialshereas at 78 K it has been estimated only for Al and is ∼12 nm.he weakening of the substructure, arising from the collapse of

he dislocation network under the high stresses, forces the flowocalization, the nucleation of voids and fracture.

cknowledgements

The authors wish to thank the Centre for Automotive Materi-ls and Manufacturing (CAMM) for financial support. Dr. Olafngler of Hydro Aluminium GmbH is gratefully acknowledged

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or providing high-purity aluminum used in this work. AA5754lloys were kindly supplied by Novelis. We wish to thank Dr.jorn Holmedal for comments on the manuscript.

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Plastic Deformation of Al and AA5754 Between 4.2K and 295K

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