Influence of predeformation on ageing in an Al–Zn–Mg alloy—I. Microstructure evolution and...

INFLUENCE OF PREDEFORMATION ON AGEING IN AN

Al±Zn±Mg ALLOYÐI. MICROSTRUCTURE EVOLUTION

AND MECHANICAL PROPERTIES

A. DESCHAMPS{{, F. LIVET and Y. BREÂ CHET

LTPCM/ENSEEG, Domaine Universitaire, BP 75, 38402 St Martin d'HeÁ res Cedex, France

(Received 20 February 1998; accepted 22 August 1998)

AbstractÐWe investigate the ageing kinetics of a ternary Al±Zn±Mg alloy, and the coupling between pre-straining and subsequent ageing. The precipitation sequence is investigated using DSC. The microstructureis followed using TEM and Small-Angle X-ray Scattering, and the hardening resulting from precipitationstudied by hardness measurements. The e�ect of prestraining depends signi®cantly on the parameters ofthe thermomechanical treatment, and has to be considered in terms of competition between homogeneousprecipitation and heterogeneous precipitation on dislocations. The e�ect of prestraining on peak hardeningis controlled by the nucleation of the hardening Z' phase, and therefore depends critically on the heatingrate to the ageing temperature. The e�ect of prestraining on overageing is an increase in overall coarseningkinetics due to the combination of faster precipitate coarsening on dislocations and dissolution of homo-geneous precipitates followed by solute di�usion towards dislocations. # 1998 Acta Metallurgica Inc. Pub-lished by Elsevier Science Ltd. All rights reserved.

1. INTRODUCTION

Precipitation from a solid solution and the related

hardening of alloys is one of the most investigatedareas of physical metallurgy, both from the exper-imental and from the theoretical viewpoint [1±3].

However, most of the existing models focus on sim-pli®ed situations. For instance, in the case of homo-geneous precipitation, a single precipitate phase is

assumed. For heterogeneous precipitation on struc-tural defects (grain boundary, dislocations), homo-geneous precipitation is assumed to be absent. In

addition, most of the existing work is concernedwith isothermal treatments.In many systems, these simplifying assumptions

appear somewhat unrealistic: precipitation occurs

through a sequence of metastable phases, and het-erogeneous and homogeneous precipitation arecompeting phenomena. Moreover, industrial pro-

cesses often involve non-isothermal heat treatments,and/or complex thermomechanical treatments lead-ing to the interaction between precipitation and dis-

locations.In view of the variety of mechanisms involved, it

is hopeless to develop models for such situations at

the same level of details as the academic models. Inorder to understand and to model these complexsituations [4], one has ®rst to qualitatively identifythe relevant microstructure evolutions, to quantify

these evolutions, and then to use from classical

modelling the functional form of the laws, leavingas adjustable parameters (whose physical meaning

and order of magnitude are known) some of theunknown constants.The present two papers are concerned with such

a problem for a relatively simple system, AlZnMg,

which however is the base of a whole family ofaeronautic alloys. Paper I will present experimentalinvestigations of microstructural evolution during

ageing treatments in situations re¯ecting some ofthe industrial processes. Paper II [5] will present asimple model, relying on the observations of paper

I, and allowing to describe precipitation kineticsand the yield stress resulting from these thermo-mechanical treatments.The Al±Zn±Mg±(Cu) alloys family provides some

of the most interesting alloys for light-weight struc-tural applications, in particular in the aircraftindustry. Their properties are obtained through a

complex thermomechanical treatment, which endswith a solution treatment, a quench, a plastic defor-mation to relieve quench stresses and a multi-step

ageing treatment, with the usual precipitationsequence [6]:

Solid solution ÿ4 GP zones ÿ4 metastable Z'

ÿ4 stable Z�MgZn2�:

The plastic deformation step (generally of the order

of 2%) aims at relieving quench-related internalstresses which would alter dimensional stabilityduring later machining of the alloy. It introduces in

the material a large density of dislocations which

Acta mater. Vol. 47, No. 1, pp. 281±292, 1999# 1998 Acta Metallurgica Inc.

Published by Elsevier Science Ltd. All rights reservedPrinted in Great Britain

1359-6454/99 $19.00+0.00PII: S1359-6454(98)00293-6

{To whom all correspondence should be addressed{Also at Pechiney-Centre de Recherches de Voreppe,

BP 27, 38340 Voreppe, France..

281

can alter the precipitation sequence in a number of

ways, either by heterogeneous precipitation on dis-

locations or by modi®cation of bulk precipitation

kinetics:

ÐDislocations are favorable nucleation sites

for precipitates [7], and short-circuit di�usion

paths for solutes [8±10]. This results in faster and

coarser precipitation on dislocations. Moreover,

precipitation on structural defects naturally pro-

motes the stable rather than the metastable

phases.

ÐOn the other hand, the interaction between

dislocations and solutes results in a solute ¯ux to

dislocations [11±13]. Therefore a lower solute

fraction is available for bulk precipitation.

Dislocations act also as vacancy sinks [14]. Due

to the strong in¯uence of the vacancy concen-

tration on the precipitation process in these

alloys [15±17], the presence of dislocations may

in¯uence a large volume of material around

them.

Given this variety of possible interactions, the e�ect

of the presence of dislocations on precipitation and

mechanical properties strongly depends on the alloy

family and ageing sequence. In many alloys the pre-

sence of dislocations may increase mechanical prop-

erties, and often changes the precipitation sequence.

This is the case for Al±Li [18, 19], Al±Mg±Si [20],

Al±Cu±Li [21, 22], Al±Li±Cu±Mg±(Ag) [23], for

instance. In other cases, especially when nucleation

of small-size homogeneous precipitation is easily

acheived, the presence of dislocations tends to

decrease mechanical properties. This is the case for

Al±Cu±Mg±Ag, for instance [22].

In Al±Zn±Mg±(Cu) alloys, the e�ect of predefor-

mation on the precipitation process and mechanical

properties after ageing has been studied both from

the microscopic and macroscopic point of views:

ÐOn a microscopic scale, the precipitation

process on dislocations has been studied in

some detail by Ryum [24], Allen and Vander

Sande [25, 26], and Gomiero et al. [27].

Precipitation on dislocations appears directly as

the equilibrium phase Z. The growth of precipi-

tates on dislocations is faster than in the bulk,

and is partly controlled by the angle between the

largest dimension of the precipitate and the dislo-

cation line [26]. It corresponds in the early stages

of ageing to a Cottrell±Bilby type of law, with

an activation energy of 1.25 eV, showing that the

limiting mechanism for precipitate growth is bulk

di�usion towards dislocations, rather than pipe

di�usion along dislocations (mechanism for

which the activation energy is close to

0.4 eV [28]).

ÐOn a macroscopic scale, the in¯uence of dis-

locations on the evolution of mechanical proper-

ties during ageing has also been investigated:

concerning natural ageing at room temperature,

a slowing down of precipitation kinetics has been

observed in the presence of dislocations, due to

progressive annihilation of quenched-in vacancies

on dislocations [29, 30]. Concerning ageing at

higher temperatures (120±1608C), the presence of

dislocations generally results in an acceleration

of the global precipitation kinetics [27, 29±33].

The in¯uence of dislocations on mechanical

properties is often a decrease in the hardening

response [31], but an increase in hardening po-

tential has also been reported, especially when

predeformation is carried out on a naturally aged

material [32, 34±37].

The aim of this study is to relate these two levels of

understanding by considering the following points:

ÐPrecipitation in a predeformed structure is a

competitive process between bulk precipitation

and precipitation on dislocations. Thus one

needs to consider separately the two families of

precipitates and their coupled evolution during

ageing.

ÐThese alloys experience a complex ageing

sequence, therefore one needs to consider the

dynamics of phase transformations in order to

understand the evolution of microstructure and

mechanical properties. More speci®cally, it is im-

portant to understand how the in¯uence of pre-

deformation depends on the parameters of the

ageing treatment. Among these parameters, we

will consider the e�ect of the amount of defor-

mation and of the heating rate to the ageing tem-

perature. Changing this heating rate will change

the ability for ®ne-scale precipitation in the

bulk [38] and therefore change the balance

between bulk precipitation and precipitation on

dislocations.

ÐA complete understanding can only be

achieved by considering in detail the relation-

ships between microstructure and mechanical

properties all along the ageing treatment.

On the experimental viewpoint, this study was car-

ried out on a model Al±Zn±Mg alloy using a com-

bination of experimental techniques, which enables

to follow simultaneously microstructure and mech-

anical properties on the various salient scales:

Ðon a microscopic scale, local information on

precipitation was obtained by Transmission

Electron Microscopy, and average information

was obtained by Small-Angle X-rays Scattering.

Ðon a macroscopic scale, information on

phase transformations was obtained by

Di�erential Scanning Calorimetry, and mechan-

ical properties were studied by microhardness.

A model describing competitively the kinetics of

bulk precipitation and precipitation on dislocations,

DESCHAMPS et al.: INFLUENCE OF PREDEFORMATION ON AGINGÐI282

as well as the resulting mechanical properties, ispresented in a companion paper (part II) [5].

2. MATERIAL AND METHODS

The alloy is an Al±6.1wt%Zn±2.35wt%Mg

model alloy containing also 0.1% Zr. It was pro-vided by Pechiney±Centre de Recherches deVoreppe as 3 mm thick sheets, in the cold rolledstate. The thermomechanical treatment studied in

this work is the following:

Ð1 h solution treatment at 4758C followed bya drop quench in cold water, in order to retain a

complete solid solution. The resulting grainstructure was partially recrystallized, recrystal-lized grains having a diameter of approx. 20 mm;

Ðtensile deformation, performed at a strainrate of 10ÿ4 sÿ1;

Ð3 days natural ageing;

Ðageing at 1608C, with two di�erent heatingrates: slow heating at 308C/h, and fast heating byintroducing the samples in the hot furnace. Theslow heating rate is close to heating rates used in

industrial heat treatments and will be the ``favor-able case'' for precipitation since it promotes®ne-scale precipitation of the hardening phase Z'.The fast heating rate favours reversion of GPzones formed during natural ageing and thereforewill be the ``unfavorable case''.

. Samples for TEM observations were cut with alow-speed cutting saw and mechanically thinneddown to 120 mm. The 3 mm disks were obtained

by ultrasonic cutting, and electropolished in adouble-jet Tenupol by a 33% nitric acid solutionin methanol maintained at ÿ208C and 15 V.

Samples were observed using Jeol 200 CX andJeol 3010 microscopes.

. SAXS measurements were performed in situ

during heating by recording the intensity dif-fracted by 80 mm thick samples from a CuKasource (l = 1.541 AÊ ) with a position sensitivedetector measuring scattering vectors q ranging

from 0.02 to 0.5 AÊ ÿ1. Spectra were correctedfor background, ¯uorescence and absorptione�ects. This experiment is described in detail

elsewhere [39].

The particle dimensions can be calculated in twoways [40]. In the low-q region one may use the

Guinier approximation, which gives the gyrationradius of the particles. However, the validity of theGuinier law depends on the precipitate size distri-

bution. When the microstructure consists of a singlefamily of homogeneously nucleated precipitates, theGuinier plot (ln(I) vs q2) shows a straight line in a

large range of di�usion vectors (typically in therange 0.8 < q.R< 2, where R is the precipitatemean size and q the di�usion vector). In the case ofmore complex microstructures consisting of several

families of precipitates having separate sizes andmorphologies, it is generally not possible to deter-

mine separately the Guinier radii of the di�erentprecipitate families. However, one can have anothermeasurement of the precipitate radius using Porod's

limit: regardless of the precipitate size distribution,the large-angle scattering is proportional to 1/q4,and the ratio between this proportionality factor

and the integrated intensity can be related to theratio between the total volume and the total surfaceof the precipitates, thus providing an estimate of

their Porod radius Rp:

Rp � 3V

S� 3

p�1ÿ fv�limqÿ41�I�q�:q4��1

0

I�q�q2 dq

. DSC measurements were performed on 275 mgsamples, with a reference sample of pure Al (5 N)

of identical weight, on a SETARAM calorimeter.The heating rate was varied from 30 to 3008C/h.

. Vickers Microhardness was carried out with aweight of 500 g. Each data point represents an

average of ®ve measurements, and the scatter wasapproximately 2 HV.

3. TEM RESULTS: NATURE AND DISTRIBUTIONOF PRECIPITATES

TEM observations provide information on the

nature, size and spatial distribution of precipitateson a local and microscopic scale. We will investi-gate how these characteristics are modi®ed in the

presence of dislocations, depending on the par-ameters of the ageing treatment.The very nature of the precipitates present after

the heat treatment is in¯uenced by the presence ofdislocations [41].

. In the undeformed material, a classical evolution

of the microstructure is observed. GP zones arereplaced in the early stages of ageing by themetastable Z' phase, which is responsible for thepeak hardening of the alloy [42, 43]: after 1 h at

1608C most precipitates are disk-shaped Z'.During the later stages of precipitation, Z' under-goes a transformation in the stable phase Z,mostly in the Z1 and Z4 orientation, some Z2being yet present.

. When the alloy is deformed before ageing, the

precipitation sequence is quite di�erent. After 1 hat 1608C one ®nds mostly Z' precipitates in thebulk and Z2 precipitates on dislocations, with a

very small fraction of Z1 and Z4. At longer ageingtimes, mainly Z2 and Z4 precipitates are observed,with a small residual fraction of Z1, whichdiminishes during the ageing.

The size and spatial distribution of precipitatesduring ageing is also modi®ed by the presence ofdislocations.

DESCHAMPS et al.: INFLUENCE OF PREDEFORMATION ON AGINGÐI 283

. In the undeformed material, precipitates remain

homogeneously distributed in space during the

ageing treatment [Figs 1(a)±3(a)], except for a

small precipitate-free zone close to grain bound-

aries. After 1 h at 1608C, precipitates are disk-

shaped and approximately 20 AÊ in radius. After

700 h at 1608C, precipitates are more rounded

and still show a strain contrast corresponding to

partial coherency. Their mean radius is approxi-mately 120 AÊ .

. At short ageing times (here 1 h at 1608C), the in-¯uence of dislocations on the size and spatial dis-

tribution of precipitates depends strongly on the

heating rate. After a slow heating rate of 308C/h,the material does not seem much perturbed by

the presence of dislocations [Fig. 1(c)]. Although

some coarse precipitates lying on dislocations arepresent, precipitation in the bulk occurs similarly

as in the undeformed material. This can be re-

lated to the favorable conditions for nucleation

of Z' precipitates on GP zones (see below). Afterrapid heating, the situation is quite di�erent

[Fig. 1(b)]: the microstructure is characterized by

a strong spatial gradient in the precipitates sizes.Precipitates on dislocations are coarse, of radius

often larger than 100 AÊ . Far from dislocations,

precipitates show a homogeneous distributionand a size identical to the undeformed material

Fig. 1. TEM micrographs of the material aged 1 h at1608C in di�erent conditions. (a) No predeformation,rapid heating; (b) 10% predeformation, rapid heating; (c)

10% predeformation, slow heating.

Fig. 2. TEM micrographs of the material aged 50 h at1608C after a rapid heating. (a) Undeformed material; (b)

material predeformed 10%.


(radius of 20 AÊ ). In between, a continuous size

gradient can be observed. Moreover, one can ob-serve a small precipitate-free zone close to dislo-

cations.

This spatial precipitate size gradient can be attribu-ted to the precipitation process during natural age-

ing: during room temperature ageing, the kineticsof GP zones formation depends critically on the

supersaturated vacancy concentration [15±17]. Onthe other hand, the presence of dislocations pro-

gressively slows down the natural ageing kineticsbecause of progressive annihilation of quenched-in

vacancies on dislocations during naturalageing [29, 30]. Thus, after natural ageing the GP

zone structure is less developed close to dislo-cations. Nucleation of the Z' phase during ageing at

temperatures like 1608C depends critically both onthe GP zone structure [44±47] and on the heating

rate [38]: good nucleation conditions require simul-taneously, large GP zones radii so that they are

stable up to the temperature where Z' nucleationcan take place (typically 1208C), and a slow heating

rate so that the growth rate of GP zones during theheating stage is larger than the growth rate of theircritical radius until nucleation of Z' takes place.

Thus, when the GP zone structure is poorly devel-oped and the heating rate is fast, nucleation of Z' isdi�cult and results in large precipitate sizes.

. During later stages of ageing, precipitates lyingon dislocations continue to grow [Figs 2(b) and3(b)] (especially those which are connected to

several dislocations). The precipitate-free-zonearound dislocations extends, and after 700 h onecan hardly ®nd any precipitates outside dislo-

cations. One cannot observe any signi®cant di�er-ence in the materials experiencing fast and slowheating rates.

This extension of a precipitate-free-zone around dis-locations can be attributed to the elastic interactionof solute atoms and dislocations. This interactionresults in a solute ¯ux towards dislocations [11±13]

which enables the rapid growth of precipitates ondislocations and the solute depletion of the materialsurrounding dislocations.

. When strain prior to heat treatment is increased,the qualitative e�ect of dislocations remains thesame but the volume fraction of material in¯u-

enced by their presence is larger.

4. SMALL-ANGLE X-RAYS SCATTERING ANDDIFFERENTIAL SCANNING CALORIMETRY:

KINETICS OF PRECIPITATION

Whereas TEM provides qualitative informationon precipitation, SAXS measurements allow us toobtain quantitative information on the precipitate

sizes and the precipitation kinetics. Moreover, DSCinvestigations allow us to identify the precipitationsequence.

4.1. Undeformed material

Consider ®rst the in¯uence of heating rate on theprecipitation in the undeformed material. Both the

precipitation sequence and the precipitation kineticsare sensitive to this parameter.During heating at constant rate, one can observe

a series of endothermic and exothermic peaks onthe DSC thermogram [Fig. 4(a)].

. The ®rst endothermic peak (A) can be attributed

to the reversion of GP zones formed during natu-ral ageing [42, 47±50]. As expected, this reversionis of larger magnitude for the fast heating rate:

the faster the heating rate, the faster the increaseof the critical radius for GP zone dissolution. Itis worth noticing that some reversion occurs attemperatures as low as 508C.

Fig. 3. TEM micrographs of the material aged 700 h at1608C after a rapid heating. (a) Undeformed material; (b)

material predeformed 10%.


. When the temperature reaches approximately

1208C (for the slow heating rate), Z' precipitatesnucleate on existing GP zones, which results in a

strong exothermic peak (B) in the thermogram.

At higher temperatures one can observe two

other exothermic peaks (C and D), the ®rst corre-

sponding to the Z' 4 Z transformation, and the

second probably to the precipitation of T phase,

the alloy being in the (Al + Z +T) region of the

phase diagram. At the slow heating rate, the Z'and Z' 4 Z peaks are well separated. This indi-

cates that Z' nucleates su�ciently easily so that

the precipitates formed are relatively stable and

do not immediately transform into Z. This is not

the case for the fast heating rate: in this case the

peaks B and C largely overlap. It thus appears

that an increased reversion of GP zones duringheating results in a more di�cult nucleation of Z'at higher temperatures, in the same way as wasdiscussed in Section 3. Possibly Z' nucleation

occurs at a larger radius, triggering the Z' 4 Ztransformation almost immediately.

This is con®rmed by in situ SAXS data. In Fig. 4(b)we have represented the evolution of Guinier radiuswith ageing time at 1608C after slow or rapid heat-

ing. After rapid heating the initial radius of precipi-tates is much larger than after slow heating,

showing that nucleation of Z' is actually more di�-cult. For long ageing times, no di�erence between

the two heating rates can be detected: the initialdi�erence in microstructures is too small to persistthrough the coarsening process.

Fig. 4. In¯uence of the heating rate on the phase transformations during ageing at 1608C: (a) DSCthermograms performed after 3 days of natural ageing at two di�erent heating rates (the DSC signal ismultiplied by a factor of 10 in the case of the 308C/h heating rate); (b) evolution of the Guinier radiusduring ageing at 1608C as a function of the heating rate to the ageing temperature, investigated by in

situ Small-Angle X-rays Scattering.


Thus the two heating rates to the ageing tempera-ture provide two very di�erent situations: a slow

heating rate favors nucleation of Z' on GP zones,resulting in a ®ne precipitate distribution. A fastheating rate makes nucleation of Z' di�cult because

of GP zones reversion, and the resulting precipitatesare coarser. This di�erence in behaviour for di�er-ent heating rates has important consequences on

the response of the predeformed material to heattreatment.

4.2. Deformed material

In Figs 5(a) and 5(b) we have represented the

DSC thermograms for the undeformed anddeformed material, in the case of a slow heatingrate and in the case of a rapid heating rate. It

clearly appears that the introduction of dislocationsprior to ageing induces signi®cant di�erences in thetransformation sequence. Although the reversion

peak is not much modi®ed, the Z' peak of the unde-formed material is replaced by a peak of much

higher intensity. From electron microscopy results

(see Section 3) we interpret this peak by the simul-taneous precipitation of Z' in the bulk and of Z on

dislocations. Thus the Z' 4 Z transformation peak

(B), concerning only bulk precipitates, is weaklyobserved in the case of slow heating, and is not

observed at all in the case of rapid heating. In this

last case, even if the Z' 4 Z event is present, it can-not be separated from the main peak of much

higher intensity. Moreover, it can be observed thatin the case of rapid heating, the intensity of the pre-

cipitation peak increases with the amount of prede-

formation.

In situ SAXS experiments provide a quantitativestudy of the modi®cations in precipitation kinetics

induced by the presence of dislocations.

The modi®cation of the precipitation kinetics

with predeformation can be quanti®ed by in situSAXS.

Consider ®rst the favorable case of a slow heating

rate. We have seen in Section 3 that the precipitate

Fig. 5. In¯uence of predeformation on the DSC thermograms performed after 3 days natural ageing.(a) 308C/h heating rate; (b) 3008C/h heating rate.


microstructure contains two families of precipitates,

lying in the bulk and on dislocations. The smaller

the precipitates the larger the scattering vector of

their contribution to the di�racted intensity. If the

two precipitates families are su�ciently far apart,

one can obtain from a Guinier plot information on

both sizes [51]. However, in our case the two pre-

cipitate families were not separated enough and it is

not possible to derive any meaningful Guinier

radius from the di�usion spectra. Therefore we

used in this case the Porod radius (see Section 2).

Calculation of the Porod radius requires an esti-

mate of the precipitate volume fraction fv.

However, in our case fv is small (of the order of

4%) and 1ÿ fv31. It also requires the estimation of

the integrated intensity Q0 ��10 I�q�q2 dq. The evol-

ution of integrated intensity during ageing with a

heating rate of 308C/h is shown in Fig. 6(a). One

can ®rst observe that, at the beginning of the heat-

Fig. 6. In¯uence of predeformation on the in situ Small-Angle X-rays Scattering measurements duringageing at 1608C in the case of a heating rate to the ageing temperature of 308C/h. (a) Evolution of themeasured integrated intensity; (b) evolution of the Porod radius calculated from the extrapolated inten-

sity for long ageing times.


ing stage, the integrated intensity decreases with

prestrain. This corresponds to the fact that annihil-ation of vacancies during natural ageing slowsdown the GP zone formation kinetics.

During isothermal ageing at 1608C, intensity ®rstincreases, and then decreases quickly. This decreaseis caused by the loss of the scattered intensity corre-

sponding to the largest precipitates in the beamstop.This decrease occuring earlier in the predeformed

material is a ®rst indication that precipitates arecoarser in this state. In order to calculate the Porodradius all along the ageing treatment, the integrated

intensity was extrapolated for long ageing times.Figure 6(b) represents the evolution of Porod

radius during the ageing treatment with a slow

heating rate, as a function of prestrain. It appears®rst that after natural ageing the Porod radius is

identical whatever the prestrain. After natural age-ing, only the volume fraction is changed by pre-straining, and not the precipitate radius. Thus the

e�ect of the presence of dislocations on natural age-ing is mostly a kinetic one: the thermodynamics ofGP zone nucleation (re¯ected by the nucleation

radius) is unchanged, whereas the kinetics of GPzone precipitation is slowed down by the presence

of dislocations because of their decreasing thevacancy supersaturation.During the 1608C heat treatment, it clearly

appears that prestraining induces larger averageprecipitates radii. This is consistent with the TEMobservations.

In the case of rapid heating, qualitative infor-mation can be derived from Guinier plots (Fig. 7).

After a short ageing time [30 min at 1608C,Fig. 7(a)], the average Guinier radius in the unde-formed material is 20 AÊ . In the 10% deformed ma-

terial, we can observe one case where it is possibleto separate the two families of precipitates: one pre-

cipitate family is of radius 20 AÊ , i.e. the same radiusthan in the undeformed material. This is consistentwith the TEM observations that precipitates far

from dislocations had the same radius than in theundeformed material. The second precipitate familyhas an average radius of 40 AÊ , which is also consist-

ent with the TEM results.After 20 h at 1608C, it is not possible to dis-

tinguish the precipitates in the bulk and on dislo-cations [Fig. 7(b)]. It can be only qualitativelyobserved that in the predeformed material the scat-

tered intensity is higher and concentrated at lowerscattering vectors than in the undeformed material,and therefore that the mean radius is larger, as can

be observed by TEM.In this case of rapid heating, the calculation of

the Porod radius did not show any di�erencebetween the undeformed and deformed material.Since we know from both the TEM observations

and the above Guinier analysis that the mean pre-cipitate size is larger in the undeformed material,this can only result from an increase of the precipi-

tate aspect ratio due to precipitation on dislo-cations: a larger aspect ratio increases the total

surface of precipiate interfaces and thereforedecreases the calculated Porod radius.

5. CONSEQUENCES ON THE MECHANICALPROPERTIES

The in¯uence of prestraining on the microstruc-tural evolution during heat treatment induces di�er-ent mechanical properties, which result from the

superposition of the modi®cation of the precipi-tation characteristics and of the intrinsic strainhardening due to dislocations. We will here only in-

vestigate the microhardness as an indicator of yieldstress.The material shows a typical age-hardening

curve, with underaged states during the heatingstage to 1608C, peak ageing around 1 h at 1608C,and overageing at longer ageing times, correspond-ing to precipitate coarsening.

Fig. 7. In¯uence of predeformation on the in situ Small-Angle X-rays Scattering measurements during ageing at1608C in the case of a fast heating rate. (a) Guinier plotafter 30 min at 1608C; (b) Guinier plot after 20 h at

1608C.


. Consider ®rst the ``favorable'' case of a slow

heating rate [Fig. 8(a)].

The e�ect of predeformation on peak hardening is

quite complex: a prestraining of 1±10% slightly

increases the peak hardening, whereas a larger pre-

strain starts a decrease of peak hardness. This is a

result from the combination of the hardening e�ect

due to dislocations and the coarser radius of pre-

cipitates on dislocations. In any case, the e�ect of

prestraining on peak hardening is very small.

The e�ect of prestraining on the overaging is

much clearer: is appears that prestraining results in

an apparent increase in the coarsening kinetics.

From the microstructural observations, it appears

clearly that this is due to the increasing fraction of

solute included in the coarse precipitates on dislo-

cations and on their faster coarsening kinetics.

. Now turn to the ``unfavorable'' case of a rapid

heating [Fig. 8(b)].

Fig. 8. In¯uence of predeformation on the evolution of microhardness during ageing at 1608C. (a)308C/h heating rate; (b) hot furnace entrance.


In this case the e�ect of predeformation on peakhardening is a drastic decrease: 10% predeforma-

tion results in a loss of 20 HV of peak hardening,i.e. approximately 60 MPa. This decrease occurs inspite of the hardening e�ect of a large dislocation

density, thus we can estimate that the loss of pre-cipitation hardening is at least of 20%. In ®rst ap-proximation peak hardening scales with the square

root of the volume fraction precipitated when theaverage radius is equal to the transition radiusbetween shearing and by-passing. Thus a decrease

of 20% in mechanical properties would mean thatthe volume fraction precipitated at peak hardeningin the predeformed material is some 35% lowerthan the equivalent for the undeformed material.

As the time for peak hardening is not very di�erentin the two materials, this does not seem to corre-spond to the experimental observations. Therefore

this decrease in yield stress is a priori surprisingwhen compared to the average characteristics ofprecipitation.

We have observed by TEM that the microstruc-ture after short ageing times is characterized by acontinuous size gradient from the bulk precipitates

to the precipitates on dislocations. We will see inthe companion paper that the reason for thedecrease in yield stress with prestraining in case ofrapid heating is to be found in the wide size distri-

bution of precipitates induced by the presence ofdislocations [5].

6. CONCLUSIONS AND DIRECTIONS FORMODELING

We have shown using di�erent experimental tech-niques that the e�ect of prestraining on precipi-tation is complex and depends on the amount of

prestraining as well as on the conditions of heattreatment. The evolution of the microstructure hasbeen quanti®ed, and the e�ect on the hardness has

been investigated. The salient features of the e�ectof prestraining are:(a) The presence of dislocations not only induces

nucleation and growth of larger precipitates at thedislocation core, but also in¯uences the materialsurrounding the dislocations. This can be describedby a growing in¯uence radius around the dislo-

cation.(b) Short ageing times: a slow heating rate

favours a ®ne and homogeneously distributed

nucleation of the hardening Z' phase. The in¯uenceof dislocations is restricted to the presence of largerprecipitates connected to the dislocations, and the

bulk material experiences unchanged precipitationcharacteristics. The related in¯uence of dislocationson the peak strength is very small.

A fast heating rate causes nucleation of the hard-ening phase Z' to be more di�cult. This e�ect isemphazised close to dislocations due to the slowerGP zones formation in these regions during natural

ageing. This results in a large precipitate size gradi-ent in the material, from unchanged precipitate

sizes far from dislocations to large precipitates closeto the dislocations and on dislocations. The relatedin¯uence of dislocations on the mechanical proper-

ties is a large decrease of the peak strength withincreasing predeformation.(c) Long ageing times: at long ageing times the

evolution of the microstructure can be described bythe increase of the in¯uence radius of dislocationsdue to progressive dissolution of precipitates close

to dislocations to the pro®t of precipitates on dislo-cations. This mechanism results in a larger averageprecipitate size with increasing predeformation, andcauses a decrease of mechanical properties in the

coarsening regime. This acceleration of coarseningis independant of the initial heating rate up to theageing temperature.

AcknowledgementsÐThe authors wish to acknowledgePechiney±Centre de Recherches de Voreppe for providingthe material and for ®nancial support. The valuable helpof David Godard for the DSC experiments is alsoappreciated.

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Influence of predeformation on ageing in an Al–Zn–Mg alloy—I. Microstructure evolution and...

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