Materials Sciences and Applicatio n, 2011, 2, 899-910
doi:10.4236/msa.2011.27120 Published Online July 2011 (http://www.SciRP.org/journal/msa)
899
Precipitation Kinetics and Mechanism in Cu-7
wt% Ag Alloy
Djamel Hamana1, Mohamed Hachouf2, Leila Boumaza1, Zine El Abidine Biskri1
1Phase Transformations Laboratory, Mentouri University of Constantine-Algeria; 2Nuclear Research Centre of Birine MB, Ain
Oussera, Djelfa, Algeria.
Received March 3rd, 2011; revised April 2nd, 2011; accepted April 14th, 2011.
ABSTRACT
The discontinuous precipitation kinetics and mechanism of the  (Ag-rich) phase in Cu-7 wt% Ag alloy has been inves-
tigated using dilatometric and calorimetric anisothermal analysis, optical microscopy, scanning and transmission elec-
tron microscopy and X-ray diffraction. Dilatometric and calorimetric curves present at ~ 500˚C an important effect
related to the
 (Ag-rich) phase formation and consequently the matrix β (Cu-rich) depletion. The nucleation and
growth of the precipitated phase show cells formation at initial grain boundaries; a fine lamellar structure is detected
by SEM and TEM and consists of alternate lamellar of the α (Ag-rich) and β (Cu-rich)-solid solutions. Cellular pre-
cipitation leads to the simultaneous appearance of two diffraction peaks and occurs apparently according to the
Fournelle and Clark’s mechanism. Obtained results give an Avrami exponent n = 2.0 ± 0.2 in agreement with an inter-
facial controlled process having an activation energy Ea equals to 99 ± 7 kJ/mol obtained from anisothermal analysis
by using different isoconversion methods. This activation energy expresses the discrepancy between isoconversion
methods and the analytical diffusive model. Moreover, the supersaturation rate has an effect on the lamella spacing of
the precipitated cells.
Keywords: Cu-Ag alloy, Dilatometry, Calorimetry, Scanning Electron Microscopy, Discontinuous Precipitation
1. Introduction
Discontinuous precipitation (DP) is not a rare phenome-
non; rather it occurs in many commercial binary and ter-
nary alloys, e.g. based on Ag, Al, Co, Cu, Fr, Ni or Pb.
Thus, many investigations into this phenomenon have
been conducted in the past and accordingly, many theo-
ries have been developed to account for the initiation and
growth mechanisms of DP. For example, Hirth and Gott-
stein [1] found that the occurrence of the pucker mecha-
nism in Al-Ag-Ga (and other alloys) gives evidence that
the driving force for DP may be derived from various
sources.
Discontinuous reactions are solid state moving bound-
ary phase transitions characterised by a discontinuous or
abrupt change in orientation and composition between
the matrix phase in the reactant and product aggregate
across the migrating boundary or reaction front that pro-
vides a short circuit path of diffusion. The reactions in-
clude discontinuous precipitation, discontinuous coars-
ening, discontinuous dissolution, and diffusion induced
grain boundary migration. All these reactions may ac-
count for substantial change in microstructure, composi-
tion, and material properties, and hence, deserve ade-
quate scientific attention for a better understanding [2].
As is well known, the Cu-Ag alloy system is a phase
separation type with eutectic reaction and there is a lim-
ited solubility range at both ends of the phase diagram.
The solubility limit of Cu in Ag and Ag in Cu is 14.1 and
4.9, respectively, in at% at the eutectic temperature
779˚C. With decreases in the temperature, the values
decrease rapidly to 0.35% Cu for the former and become
negligible (less than 0.06% Ag) for the latter at 200˚C [3].
The Cu-rich terminal solid solutions containing 2 - 4.8
at% Ag transform into a lamellar structure consisting of
depleted
 and  (Ag-rich) phases by discontinuous pre-
cipitation [4].
Concerning the previous studies related to discontinu-
ous precipitation in Cu-Ag, there are several publications.
For example, Wirth and Gleiter [5] have studied the dis-
continuous precipitation reaction in the Cu-5 wt% Ag
alloy and have concluded that the observations suggest
that a discontinuous precipitation reaction can occur only
at a pre-existing (static) grain boundary if the structure of

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 900
the static boundary is converted into a structure of high
mobility by a structural transformation. Gust et al. [6]
have investigated the diffusion along a migrating random
grain boundary of Cu-3.8 at% Ag bicrystal by discon-
tinuous precipitation and dissolution. They concluded
that the diffusivities of migrating, sliding and stationary
grain boundaries in the Cu-Ag system are of the same
order of magnitude. This conclusion is in accord with the
general behaviour of other binary systems. Gupta [7] has
studied the kinetics of discontinuous precipitation and
dissolution of the cellular precipitate in Cu-3 at% Ag and
Cu-4 at% Ag alloys. The Cu-Ag alloys were observed to
decompose into a lamellar structure consisting of alter-
nate lamellae of the
  (Cu-rich) and  (Ag-rich) phases
when a solid solution of the alloy was aged below the
solvus temperature. The primary cell growth data were
analysed using the theories of Cahn, Hillert, Sundquist,
Turnbull and Petermann and Hornbogen. From the diffu-
sivity values, it has been shown that the growth of pri-
mary cells occurs by the diffusion of Ag along the grain
boundaries. Hamana and Choutri [8] studied the effect of
plastic deformation on the kinetics and mechanism of
cellular precipitation in Cu-6.5 wt% Ag alloy; they show
that the discontinuous precipitation rate is decreased by
small amounts of prior plastic deformation. This decrease
can be explained by the difficulty of some heterogene-
ously nucleated precipitates to follow the grain boundary
migration.
Last years, differential dilatometry has been exten-
sively used to study the precipitation reactions in differ-
ent alloys [9-15]; it has been shown that this method is
very sensitive to this kind of phase transformations.
Differential dilatometry can show effects of precipita-
tion and dissolution by the appearance of various anoma-
lies in the experimental curves, and generally, depending
on the lattice parameter and the specific volume varia-
tions, these anomalies can be expansions or contractions
[9-15]. However the shape of the dilatometric curve can
change if the second phase is a solid solution, which jus-
tifies this investigation.
A very interesting phase transformation during non
isothermal heating has been detected in Ag-8 wt% Cu
alloy [15]. This transformation, which occurred at rela-
tively low temperature (T < 300˚C), is identical to an
allotropic one. Contrary to all expectations, a region with
a new structure is formed in both sides of grain bounda-
ries apparently without lamellar precipitates. However at
higher ageing temperature (400˚C) a lamellar morphol-
ogy appears. Thus it is interesting to follow the behav-
iour of a Cu-7 wt% Ag alloy of the same system (and
phase diagram).
This work aims to study the precipitation of the second
solid solution
  (Ag-rich) phase in Cu-7 wt% Ag alloy,
using differential dilatometry (DD), differential scanning
calorimetry (DSC), optical microscopy (OM), scanning
electron microscopy (SEM), transmission electron mi-
croscopy (TEM) and X-ray diffraction (XRD). Moreover
the activation energy of the
 (Ag-rich) phase is esti-
mated from the DD and DSC data by using Kissinger and
Starink methods. The Avrami exponent n is calculated
from Johnson Mehl Avrami Kolomogorov (JMAK)
method.
2. Experimental Details
The Cu-7 wt% Ag alloy was prepared from electrolytic
copper and metallic silver of purity  99, 99% by melting
under vacuum (10–6 Torr). Cylindrical samples of 3 mm
diameter and 3 mm length were used for DSC analysis
and cylindrical samples of 4 mm diameter and 20 mm
length were used for dilatometric analysis. Square sam-
ples of 3 mm thickness and about 10 mm length were
used for OM, SEM and XRD.
The dilatometric analysis was carried out under argon
using a DI24 Adamel Lhomargy dilatometer connected
to a microcomputer. Suitable software (Logidil) allows
analysing the results. The DSC measurements were also
carried out under argon with a Setaram DSC SETsys
Evolution 1500.
The applied thermal cycle consisted of heating from
room temperature until 760˚C for both DSC and DD,
holding of 1min at this temperature and finally a cooling
until room temperature.
Two optical microscopes LEITZ MM6 and OLYM-
PUS BX51M, a scanning electron microscope FEG
JSM-7600F were employed to observe the microstruc-
tures. The compositions in the
 matrix and precipitates
phase structure were determined by an energy-dispersive
spectrometer of X-ray system (EDS) operating at 10 kV.
Moreover a transmission electron microscope (TEM)
Tecnai G2 F20 X-Twin was used to observe the cellular
precipitation and determine the lamellae composition.
The XRD diagrams were obtained using a Siemens
(Bruker) D8 Advance diffractometer with Cu K radia-
tion (V = 40 kV and A = 35 mA) and Debye-Scherrer
back-plane films (V = 40 kV and A = 20 mA) X-rays
profiles have been recorder around the (311) and (420)
diffractions peaks.
3. Results and Discussion
3.1. Dilatometric Study
Dilatometric curve of homogenised 2 h at 760˚C and
quenched in water Cu-7 wt% Ag alloy is shown in Fig-
ure 1. The derivative curve of heating segment (Figure
1(b)) presents two important anomalies:
The first is an important contraction extending in the

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy901
Temperature (˚C)
Temperature (˚C)
Temperature (˚C)
Figure 1. Dilatometric curves of the complete cycle (a) and
derivative heating (b) and cooling (c) curves of homogenised
2h at 760˚C and water quenched Cu-7wt% Ag alloy (heat-
ing and cooling rate b = 5˚C min–1).
temperature interval [374 - 516˚C] with a peak of the
derivative curve at 450˚C, related to the precipitation
process of the  (Ag-rich) phase (solid solution).
The second is an expansion in the temperature interval
[516 - 760˚C] with a peak of the derivative curve at
687˚C, which can be attributed only to the dissolution of
 (Ag-rich) phase (for the composition of our alloy the
solubility limit is about 740˚C).
The fact that the precipitation gives a contraction can
be explained by the reduction of the crystalline parameter
of the β (Cu-rich) solid solution attributed to its Ag de-
pletion (as rAg = 1.44 Å > rCu= 1.28 Å); its effect is more
important than that effect of the apparition of the α
(Ag-rich) solid solution which must give an expansion.
Moreover the observed expansion is explained by the
dissolution process which leads to an increase of the lat-
tice parameter due to the entrance of Ag atoms in the β
(Cu based) solid solution.
In conclusion, during the precipitation process of

(Ag-rich) phase Ag atoms diffuse toward the grain
boundaries, decreasing the alloy volume and giving a
contraction detected by dilatometric analysis in the heat-
ing segment. This effect is followed by an expansion
related to the solution treatment.
However, the cooling segment presents an expansion
in the temperature interval [680 - 520˚C] with a peak of
the derivative curve at 609˚C (Figure 1(c)), which is
related to the precipitation of the α (Ag-rich) solid solu-
tion at high temperature after solution treatment. This is
in accordance with the results obtained by Gupta [7]
where de maximum growth rate is observed at this tem-
perature range. In this temperature range, the supersatu-
ration is weak (2.2 wt% Ag) and consequently, during
cooling, the reduction of crystalline parameter of the β
(Cu-rich) solid solution is less important than the in-
crease of the specific volume due to the precipitation of α
(Ag-rich) solid solution, which leads to an expansion in
the dilatometric curve. Thus one can conclude that in the
cooling segment the precipitation reaction effect corre-
sponds to the maximum growth rate of precipitated cells
measured during isothermal treatment [7].
3.2. Calorimetric Study
In order to confirm and understand the dilatometric
anomalies, we have carried out a calorimetric study. DSC
curve recorded during the heating of the homogenised 2h
at 760˚C and quenched in water sample (Figure 2(a)),
shows one exothermic peak in the temperature interval
[378 - 505˚C] with a maximum situated at 449˚C, corre-
sponding certainly to the formation of
 (Ag-rich) phase
particles, as observed in the dilatometric curve (Figure
1).
This exothermic peak is followed by an endothermic
one with a minimum situated at 641˚C, corresponding to
the dissolution of the α (Ag-rich) phase. The first endo-
thermic effect (at law temperatures) is due to the machine
artefact.
In conclusion, differential scanning calorimetry confirms

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 902
Temperature (˚C)
Temperature (˚C)
Figure 2. DSC heating curves and dilatometric derivative
heating curve (heating rate b = 5˚C min–1) (a, b) of Cu-7
wt% Ag alloy homogenised 2h at 760˚C and water
quenched.
the previous results observed on dilatometric curves; the
same temperature range is observed for the tow reactions
effects (Figure 2(b)).
3.3. Study of the Precipitation Kinetics
The calorimetric analysis can also be used for calculating
the activation energies according to the method devel-
oped by Kissinger [16]:
2
1
ln( /)/C
Kispap
YbTERT (1)
and Starink [17]

1.92
2
ln1.0008 C
Stapa p
YTb ERT  (2)
where b represents the heating rate, Ea the activation en-
ergy (kJ/mol), C1, and C2 are constants, Tp the tempera-
ture peak observed in the DSC curves and R the gas con-
stant.
Such analysis can be performed in case where heat ef-
fects are caused by a single-precipitation process [16-19].
To estimate the activation energy of  (Ag-rich) phase,
the samples were homogenised at 760˚C for 2 h and
quenched in water. The DSC experiments (Figure 3(a))
were performed with different heating rates where we
note that the peaks are shifted toward higher tempera-
tures with the increase of the heating rate.
On the basis of linear relationships between ln(b/Tp
2)
and 1000/Tp (Kissinger method) (Figure 3(b)) and be-
tween ln(Tp
1.92/b) and 1000/Tp (Starink method) (Figure
3(c)), the average values for the activation energy of

(Ag-rich) phase formation were 95 ± 4 kJ/mol and 96 ± 4
kJ/mol respectively (correlation factor is R2 ≈ 0.99)
(Figure 3).
Avrami exponent n is obtained directly from Johnson
Mehl Avrami Kolomogorov (JMAK) equation [[20]:
 
1
4
lnln 1lnCanb



 
 (3)
where a is the transformed fraction (Figure 3(d)). The fit
of the transformed fraction curves gives that it has an
exponential growth, v.s temperature T, of the form:



1expaA BkT  (4)
Temperature (˚C)
ln(b/Tp2)
1000/TP(K-1)

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 903
Temperature (˚C)
Figure 3. DSC heating curves (heating rate b = 5, 10, 15, 20
and 30˚C min–1) of Cu-7 wt% Ag alloy homogenised 2h at
760˚C and water quenched (a), Kissinger plot (b), Starink
plot (c), transformed fraction (d) and plots of ln[ln(1−a) –1] vs
ln(b) (e).
where A, B and k are constants related to the heating rate
and the alloy composition.
According to Matusita and Sakka [20] the mean value
of the exponent n obtained for 470, 480 and 490˚C (Fig-
ure 3(e)) is 2.0 ± 0.2. This value is in agreement with an
interfacial controlled process [21].
Dilatometry has been employed over the last few dec-
ades to detect phase transformations which involve
changes in lattice parameters. In dilatometry experiment,
the maximum in the reaction rate coincides with the
transformation peak in the differential thermodilatometry
ln(Tp
1.92/b)
curves (DTD) [14]. Thus dilatometry analysis can also be
used for calculating of the activation energies using dif-
ferent isoconversion methods and consequently the acti-
vation energy of the
 (Ag-rich) phase precipitation
where also confirmed by dilatometric analysis.
The peaks of the DTD curves are, as in the case of the
DSC curves, shifted to higher temperatures with the in-
crease of the heating rate, with an increase of the mini-
mum value of the Thermal Expansion Coefficient (TEC)
(Table 1). The obtained differential thermodilatometry
peaks and characteristic temperatures (Figure 4), serve
as a basis for determining the activation energies of Ag
precipitation reaction. Various rates of heating (1, 2, 5,
10 and 15˚C/min) were used for determining the activa-
tion energy, the methods of Kissinger and Starink lead to
a linear relationships for Ag precipitation (correlation
factor is R2 ≈ 0.98). The activation energy of
 (Ag-rich)
phase formation is 102 ± 9 kJ/mol. There are small dif-
ferences between calculated activation energy which can
be related to the sensitivity of DTD and DSC methods.
1000/TP(K-1)
Finally, Activation energy obtained using different
isoconversion methods with DTD and DSC data analysis
(Table 2), equals to 99 ± 7 kJ/mol. This last value re-
mains smaller than those obtained by Gupta using dis-
continuous precipitation kinetics models [7]. Therefore,
he has found that mean activation energy was practically
the same for the tow different alloy compositions: Cu-3
at% Ag and Cu-4 at% Ag; 130.72 kJ/mol and 130.42
kJ/mol respectively. However the small difference in the
composition between our studied Cu-7 wt% Ag (Cu-4.25
at% Ag) alloy and his Cu-4 at% Ag alloy, gives different
activation energy calculated by the isoconversion meth-
ods. This expresses the discrepancy between isoconver-
sion methods and the analytical diffusive models.
3.4. Microstructure Observations (OM and
SEM)
3.4.1. As-Quenched State
Figure 5 presents the microstructure of homogenised and
quenched sample; one can observe a polycrystalline
structure with twins and without precipitates.
3.4.2. Ageing States
3.4.2.1. Ageing at 250˚C
Precipitation at this temperature is very slow. The optical
micrograph of homogenised and aged for 54 h at 250˚C

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 904
Table 1. Temperatures and Thermal Expansion Coefficient at the thermodilatometry contraction peaks for different heating
rates.
Heating Rate (˚C/min) 1 2 5 10 15
Peak temperature (˚C) 400 409 450 479 495
TEC (10-6/˚C) 7,3996 8,5245 9,3250 12,5352 12,7096
Temperature (˚C)
Figure 4. Differential thermodilatometry (DTD) curves
obtained at different heating rate (a), Kissinger plot (b) and
Starink plot (c).
Figure 5. Optical micrograph of Cu-7 wt% Ag alloy ho-
mogenised 2 h at 760˚C and quenched in water.
sample show cells formation at initial grain boundaries.
At the first stage, the mechanism has the morphology of
the single seam nodules which can indicate the occur-
rence of Fournelle and Clark’s mechanism [22]. After
cells growth, new cells have been formed on the second
side of the grain boundary (Figure 6(a)); one can not
observe the precipitate colonies especially at lower age-
ing temperature.
ln(b/Tp2)
Cells growth is produced after 86 h of ageing and the
lamellar precipitates in the cell interior are not observed.
The growth of cells 1 and 2 into the grain (1) is followed
by a new cell formation at the initial position of the grain
boundary (Figure 6(b)), which can also become a second
reaction front and permits the formation of a new cell 3
into grain (2). The double seam morphology is well re-
vealed.
1000/TP(K-1)
3.4.2.2. Ageing at 400˚C
At this temperature discontinuous precipitation is very
fast. The optical microstructures of the sample aged 15
min at 400˚C are shown in Figure 7 where we can note
the appearance of developed cells.
ln(Tp1.92/b)
The TEM analysis of the early stages confirms the
presence of lamellar precipitates at grain boundary (Fig-
ure 8). The alloy is observed to decompose into a lamel-
lar structure consisting of alternate lamellae of the

(Cu-rich) and
 (Ag-rich) phases. It seems that lamellar
precipitation was formed after the grain boundary migra-
tion, from its original position and the formation of al-
lotriomorphs, according to the mechanism of Fournelle et.
Clark [23]. TEM micrograph and the corresponding EDX
analysis of different areas show the high concentration of
Ag in O1:  (Ag-rich) phase and high concentration of
the Cu in O2:  (Cu-rich) phase in the form of lamellae
1000/TP(K-1)

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 905
Table 2. Average activation energy obtained using Kissinger and Starink methods and DTD and DSC data analysis.
Kissinger method Starink method
Obtained DSC curves activation energy (kJ/mol) 95 ± 4 96 ± 4
Obtained DTD curves activation energy (kJ/mol) 102 ± 9 102 ± 9
Average activation energy (kJ/mol) 99 ± 7 99 ± 7
Figure 6. Microstructures evolution at a grain boundary of
a sample homogenised 2 h at 760˚C, quenched then aged at
250˚C for 54 h (a) and 86 h (b).
Figure 7. Microstructures of a sample homogenised 2 h at
760˚C, quenched then aged at 400˚C for 15 min (a and b).
Figure 8. TEM micrographs showing the early stages of
discontinuous reaction with lamellar precipitates in Cu-7
wt% Ag alloy after ageing at 400˚C for 15 min. OGBP :
original grain boundary position, RF: reaction front and
GB : grain boundary.
(Figure 9). The two regions O3 and O4 which correspond
to the supersaturated matrix have the same composition.
Moreover, the SEM analysis of three regions shows that
the concentration of Ag is higher in the cell composed of
fine lamellar precipitates and lower in the supersaturated
solid solution (Figure 10(a)). A much larger magnific-

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 906
(a)
(b)
Figure 9. TEM micrograph of Cu-7 wt% Ag alloy aged at 400˚C for 15 min (a) and the corresponding analysis of the differ-
ent areas (O1, O2, O3 ,O4 ) (b).

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 907
Figure 10. SEM micrographs of Cu-7 wt% Ag alloy homogenised 2 h at 760˚C, quenched then aged at 400˚C for 15 min (X
10,000 (a)) and (X 50,000 (b)); the analysis in the indicated three regions 001, 002 and 003 gives 7.91 wt% Ag, 9.13 wt% Ag
and 7.57 wt% Ag respectively.
tion (X 50,000) shows the appearance of fine lamellar
structure in the cell, with lamella spacing about ~45 nm,
composed of the second phase  (Ag based solid solution)
and the matrix formed by the equilibrium  phase (cop-
per based solid solution) (Figure 10(b)). The lamella
spacing is smaller than that measured by Gupta at the
same temperature in Cu-6.5 wt% Ag (Cu-4 at% Ag) al-
loy [7], which reflects the supersaturation rate effect on
the lamella spacing.
Thus one can conclude that “empty” cells are really
composed of fine lamellar structure, characteristic of a
cellular precipitation.
The discontinuous precipitation is confirmed by X-ray
diffraction; two diffraction rings of the same peaks (331)
and (420) are observed during intermediate stages (Fig-
ure 11) and consequently, two lattice parameters exist
simultaneously: the lattice parameter of the still super-
saturated solid solution

0 and that of the depleted solid
solution
 Figure 12 shows the corresponding micro-
structures.
X-ray diffraction spectrum (a) and optical micro-
graph(b) of Cu-7 wt% Ag alloy homogenised 2 h at
760˚C quenched, heated until 500˚C then quenched in
water (heating rate b = 5˚C·min –1) are presented in Fig-
ure 13. This temperature is corresponding to the precipi-
tation process of  (Ag-rich) phase detected by DSC and
DTD (Figures 1 and 2). One can see the diffraction
peaks of  and  phases (Figure 13(a)) and the lamellar
precipitation (Figure 13(b)).
X-ray diffraction spectrum (a) and optical micrograph
(b) of Cu-7 wt% Ag alloy homogenised 2 h at 760˚C
quenched then heated until 687˚C then quenched in water
(heating rate b = 5˚C·min–1) are presented in Figure 14.
This temperature is corresponding to the peak of the de-
rivative dilatometric curve, related to the dissolution
process of  phase (Figure 1). One can see only the dif-
fraction peaks of  phase (Figure 14(a)) and observed
fine structure results from the dissolution of the precipi-
tates formed during heating (Figure 14(b)).
Precipitated cells growth occurs by the growth mecha-
nism “double seam nodules” [22] detected by optical,
SEM and TEM observations. Moreover the initiation of
the lamellar precipitation produced by the of Fournelle
and Clark’s mechanism [23] with lamella spacing about
~45 nm.
4. Conclusions
As observed in Ag-8 wt% Cu alloy [15], practically the
same precipitation and dissolution processes of the 
(Ag-rich) phase have been detected by DTD and DSC in
Cu-7 wt% Ag alloy. The transformation, which occurred
at relatively high temperature (T ~ 500˚C), is related to
the precipitation process of  (Ag-rich) phase according
to the Fournelle and Clark’s mechanism. The decomposi-
tion of the supersaturated solid solution was very slow at
low temperature (250˚C) and the interlamellar spacing is
so thin that the cells seem to be “empty”. Discontinuous
precipitation tended to be accelerated after ageing at
higher temperature (400˚C) but always with a very fine
lamellar spacing. However a fine lamellar morphology is
only detected by SEM and TEM.
Average activation energy obtained using different
isoconversion methods with DTD and DSC data analysis
expresses the discrepancy between isoconversion meth-
ods and the analytical diffusive model. Moreover, the
cooling segment of DTD analysis permits to predict the
temperature range of the maximum isothermal growth
rate of precipitated cells.
5. Acknowledgements
The authors like to express their sincere thanks to Ir. Yuri

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 908
2θ(º)
2θ(º)
2θ(º)
Figure 11. X-rays diffraction peaks and Debye Scherrer back plane films of Cu-7 wt% Ag alloy, homogenised 2h at 760˚C
quenched (a, a’) then aged at 400˚C for 1 h (b, b’) and 15 h (c, c’).
Figure 12. Microstructures evolution of sample homogenised 2h at 760˚C quenched (a) then aged at 400˚C for 1h (b) and 15 h
(c): observation of cellular precipitation.

Precipitation Kinetics and Mechanism in Cu-7 wt% Ag Alloy 909
2θ(º)
Figure 13. X-ray diffraction spectrum (a) and optical micrograph (b) of Cu-7 wt% Ag alloy homogenised 2h at 760˚C
quenched then heated until 500˚C then quenched in water (heating rate = 5˚C·min–1).
2θ(º)
Figure 14. X-ray diffraction spectrum (a) and optical micrograph (b) of Cu-7 wt% Ag alloy homogenised 2h at 760˚C
quenched then heated until 687˚C then water quenched (heating rate = 5˚C·min–1).
Rikers, Application Specialist TEM Applications Labo-
ratory FEI Company BV, and Franck Charles, General
Manager, JEOL (Europe) SAS, for assistance with the
TEM and SEM analysis respectively.
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