Materials Science and Engineering B 182 (2014) 29– 36
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Materials Science and Engineering B
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hort communication
late-like cell growth during directional solidification of an–20wt%Sn high-temperature lead-free solder alloy
ashington L.R. Santosa, Crystopher Britoa, José M.V. Quaresmab, José E. Spinelli c,∗,mauri Garciaa
Department of Materials Engineering, University of Campinas, UNICAMP, P.O. Box 6122, 13083-970 Campinas, SP, BrazilFederal University of Pará, UFPA, Augusto Correa 1, 66075-110 Belém, PA, BrazilDepartment of Materials Engineering, Federal University of São Carlos – UFSCar, 13565-905 São Carlos, São Paulo, Brazil
r t i c l e i n f o
rticle history:eceived 21 June 2013eceived in revised form 16 October 2013ccepted 18 November 2013vailable online 2 December 2013
a b s t r a c t
Although Zn–Sn alloys have suitable features for high temperature solders, as for example the absence ofintermetallic compounds (IMCs) and relatively high melting temperatures, the control of the scale of themicrostructure by adequate pre-programming of the solidification thermal parameters remains still atask to be accomplished. The present study focuses on the interrelation among hardness, microstructurefeatures/segregation and solidification thermal parameters. An upward directional transient solidifica-
eywords:icrostructure
b free solder alloysn–Sn solderolidification
tion apparatus was used in order to permit samples along a range of cooling rates to be obtained for suchevaluation. The entire Zn–20wt%Sn alloy casting is characterized by a two-phase alternated structure,which resembles the morphology of a lamellar eutectic. Experimental growth laws having −1/2 and −1/4exponents are proposed relating the interphase spacing to the growth rate and the cooling rate, respec-tively. The morphology and size of the Zn-rich plate-like cells, as well as the macrosegregation pattern
ardne
ardness are shown to affect the h
. Introduction
Special attention has been given by industry for the develop-ent of lead-free high-temperature solders, which are considered
s a challenge for example in the aerospace and automotive indus-ries. These materials are considered essential for a variety ofomponents and connections, with two main concerns of being reli-ble and free from lead. Possible applications include assemblingptical components, automobile circuit boards and circuit modulesor step soldering.
Alloys from three major candidate metallic systems can be con-idered as alternatives to Pb-rich high temperature solders: Zn–Sn,u–Sn, and Bi-based alloys [1]. Each has its own superior charac-
eristics as well as some drawbacks. Lead-containing solders, suchs 63Sn–37Pb, Pb–10Sn, and Pb–3Sn (all in wt%), have been usedn various microelectronic applications along years [1]. Nowadays,trong research and development efforts with a view to replacing Pbontaining solders with Pb-free solders are being made due to bothompetitive market pressures and environmental issues [2–9]. The
ervice temperatures may exceed 120 ◦C for regular applicationsith Pb-containing solders. In addition, industry needs the melting
range to be defined from 270 to 350 ◦C so that process control couldbe ensured [10].
Even though less volume of flux may be used for Au–Sn alloysduring soldering and good creep resistance could be obtained, bothbrittle Au–Sn intermetallic compounds (IMCs) and costs regardingthe price of gold are considered the major problems of these candi-date alloys. In addition, development of Bi–Ag solder alloys is stillunder way mainly due to its inferior thermal and electrical con-ductivity as well as poor workability. Therefore, mainly due to costconcerns, hypoeutectic Zn–Sn alloys (20, 30 and 40wt%Sn) seem tobecome interesting high-temperature lead-free solder alternatives[10]. In this case, IMCs are unable to form in the state of equilibriumin the whole composition range, since the Zn–Sn system is charac-terized by a simple binary eutectic. These alloys also show a muchimproved ductility compared with other Zn-based alloys, excellentelectrical properties and oxidation resistance at high temperatureand humidity [11]. The microstructures of Zn–Sn alloys are formedby a primary �-Zn phase surrounded by the eutectic mixture con-sisting of fine �-Zn platelets and a Sn-rich �-phase. As the alloy Zncontent decreases, the fraction of the primary �-Zn phase decreaseswhile the fraction of the eutectic Sn–Zn mixture increases. Anotheraspect important to be highlighted concerning the Zn–Sn system
is the relatively wide solidification ranges of the alloys of interest,varying from 150 to 175 K.
There are some drawbacks for Zn based alloys used as hightemperature solders. Firstly Zn is a highly corrosive metal and
Fig. 1. Schematic representation of the experimental setup: (1) rotameter; (2) heat
0 W.L.R. Santos et al. / Materials Scie
econdly, it exhibits poor wetting behavior due to the high oxygenffinity. Lastly, the microstructure of the Zn alloy is very sensitiveo soldering temperature and relatively unstable compared withhe high lead content solders [12]. The control of microstructures the main highlighted area of future work which is required for
more comprehensive understanding of the lead-free Zn-basedolders performance [10]. Distinct microstructural arrays can beormed during soldering, which are intimately associated withhe corresponding cooling rate. It is well known that the resulting
icrostructure, unsoundness, strength, and corrosion behaviorf solders depend on the thermal processing variables imposeduring soldering or thermal cycling [2,5,6,13–16]. Studies empha-izing the effects of solidification thermal parameters such as theooling rate and the tip growth rate on the final Zn–Sn soldericrostructure have not been found in the literature.According to the literature [1,17], the liquid fraction during sol-
ering operations with the Zn–Sn alloys should be kept less than0% in volume. Thus, integrity and shape of a solder filet could bereserved [1,11,17]. Suganuma et al. [1] state that no movementf liquid within the mushy zone is expected to happen in the casef a Zn–30wt%Sn alloy at 260 ◦C with liquid fraction lower than5 vol.%. Under such condition, the solid portion may support theie-attachment structure in the multiple reflow treatment.
The ductility of the Zn–Sn alloys is considered their greatdvantage [1,10,11,17]. Higher elongation-to-fracture values cane obtained with increasing tin content. A maximum elongationalue of 23% is reported for the Zn–40wt%Sn alloy [1]. Althoughlongation varies, the ultimate tensile strength becomes reason-bly constant (around 68 MPa) with increasing tin content. Softerolders are able to preserve the joint structure by relaxation ofhermal stresses.
The aim of the present study is to carry out transient directionalolidification experiments with a Zn–20wt%Sn alloy with a view toermit the effects of the tip cooling rate on the final microstructurerrangement (morphology and scale) to be examined. The influ-nces of the tip cooling rate, on the segregation pattern and onardness profile along the directionally solidified casting length arelso aimed.
. Experimental procedure
Preceding studies give complete information about the experi-ental set-up used in the directional solidification trials [18–20].
he casting assembly used in solidification experiments is shown inig. 1. Heat is directionally extracted only through a water-cooledottom, promoting vertical upward directional solidification. Atainless steel mold was used, having an internal diameter of6 mm, a height of 150 mm and a wall thickness of 10 mm. The innerertical surface was covered with a layer of insulating alumina toinimize radial heat losses, and a top cover made of an insulatingaterial was used to reduce heat losses from the metal/air sur-
ace. The bottom part of the mold was closed with a 3 mm polishedarbon steel sheet. The Zn–20wt%Sn alloy was melted in situ byontrolling the power of the lateral electric heaters. To start solid-fication, the electric heaters were disconnected and at the sameime the controlled water flow was initiated. In order to parameter-ze the time for the onset of cooling, a melt superheat of 10% abovehe liquidus temperature was adopted. Continuous temperature
easurements in the casting were monitored during solidificationia the output of a bank of fine type J thermocouples sheathed in.6 mm outside diameter (O.D.) stainless steel tubes, and positioned
t 5, 10, 15, 30, 40, 50 and 70 mm from the heat-extracting surfacet the bottom of the casting. All thermocouples were connected byoaxial cables to a data logger interfaced with a computer and theemperature data were acquired automatically.
extracting bottom; (3) thermocouples; (4) computer and data acquisition software;(5) data logger; (6) casting; (7) mold; (8) temperature controller; (9) electric heaters;(10) insulating ceramic shielding.
The “intercept” method was employed in order to measure theinterphase spacing (�) on longitudinal sections of the direction-ally solidified castings [21]. Image processing systems Neophot 32(Carl Zeiss, Esslingen, Germany) and Leica Quantimet 500 MC (LeicaImaging systems Ltd, Cambridge, England) were used to measure� and its distribution range. At least 40 measurements were per-formed for each selected position along the casting length, withthe average taken to be the local spacing. Microstructural charac-terization was also performed using a Field Emission Gun (FEG) –Scanning Electron Microscope (SEM) FEI (Inspect S 50).
X-ray diffraction (XRD) measurements were carried out in orderto determine the constitution of the phases forming the samplesmicrostructure. XRD patterns were obtained utilizing a XRD-7000Shimadzu with a 2-theta range from 20◦ to 90◦, Cu-K� radiationwith a wavelength, �, of 0.15406 nm. The segregation samples wereunderwent a fluorescence spectrometer, model Shimadzu EDX-720to estimate its average concentration through an area of 100 mm2
probe. Hardness tests were performed on the longitudinal sectionsof the samples by using a test load of 1000 g and a dwell time of10 s. A Shimadzu HMV-2 model hardness measuring test device wasused supported with a Computer Assisted hardness System (CAMS– Newage Testing Instruments, Inc.). The adopted Vickers hardnesswas the average of at least 20 measurements on each sample.
3. Results and discussion
The thermal data from the cooling curves (Fig. 2a) were usedin the determination of solidification thermal parameters, i.e., tipgrowth rate (VL) and cooling rate (T) along the casting length (seeFigs. 2b and c). These parameters were found out by considering thethermal data recorded immediately after the passage of the liquidusisotherm by each thermocouple. The time (t) intervals correspond-ing to the passage of the liquidus isotherm by each thermocoupleand the beginning of water-cooling launch (when the melt tem-perature was about 420 ◦C, i.e., t = 0) have been determined. A plot
of position of each thermocouple against these time intervals hasbeen generated. Tip growth rates were then determined based onthe time-derivative of the fitting function representing such exper-imental plot. Higher VL and T values are associated with positions
W.L.R. Santos et al. / Materials Science and Engineering B 182 (2014) 29– 36 31
F ate an
castlam
t(vcisnrifsamidcgdc
ig. 2. (a) Experimental cooling curves; and experimental values of (b) tip growth r
lose to the bottom of the Zn–20wt%Sn alloy casting decreasinglong the casting length as can be seen in Figs. 2b and c. It can beeen in Fig. 2c that the water-cooled experimental set-up permit-ed a wide range of cooling rates to be attained along the castingength (in the range 1–45 K/s). That will certainly be conducive to
corresponding significant change in the scale of the as-solidifiedicrostructure.Although the Zn–20wt%Sn alloy is placed very far from
he eutectic composition expected for the Zn–Sn systemZn–91.1wt%Sn), the Zn-rich primary phase does not exhibit a con-entional microstructural morphology, like, for instance, regularells [22] or a dendritic arrangement [20], along the entire cast-ng length. Fig. 3 depicts the macrostructure of the directionallyolidified Zn–20wt%Sn alloy casting. It can be seen that a colum-ar macrostructure prevailed along the entire casting, and someepresentative longitudinal microstructures have been includedn Fig. 3 as typical examples of the morphology of the phasesorming the microstructure of the Zn–20wt%Sn alloy. It can beeen that an alternate pattern of elongated �-Zn plate-like cellsnd an intercellular eutectic mixture characterizes the as-solidifiedicrostructure, resembling a eutectic-like lamellar arrangement
n morphology. Such cells grow essentially aligned in the verticalirection according to the heat extraction path. Very fine �-Zn cells
an be observed close to the cooled bottom of the casting, with aradual coarsening effect occurring toward the top of the castingue to the decrease in the cooling rate. Despite the solidificationonditions, which are very far from equilibrium, these plate-like
d (c) cooling rate for the Zn–20wt%Sn alloy. R2 is the coefficient of determination.
cells seem to grow side-by-side and in a coupled way with theeutectic phase during the directional solidification, with no evi-dence of regular cells for the range of cooling rates experimentallyexamined (Fig. 2c).
Plate-like cellular structures have also been reported in the liter-ature for other Zn-based alloys for high range of growth velocities.Ma et al. reported the formation of plate-like cells, with cells paral-lel to the growth direction, for growth velocities > 2.64 mm/s (andthermal gradient of 15 K/mm during Bridgman growth, i.e., cool-ing rates > 39.6 K/s) for Zn–Cu alloys in the compositional rangefrom 2.17 to 4.94 wt% Cu [23,24]. These authors affirmed that thecell configuration is determined by the crystallographic orienta-tion for strongly hexagonal systems such as that of the Zn-richphase, and that regular and plate-like cells grow in different orien-tations during directional solidification [24]. Xu et al. investigatedthe Bridgman growth of a wide range of Zn–Ag alloys composi-tions (up to 9.0 at% Ag) over a range of growth velocities from0.42 to 4.82 mm/s and a under a thermal gradient of 15 K/mm [25].These authors reported a microstructural evolution characterizedby plate-like cells for velocities higher than a critical growth veloc-ity, and a morphology typified by cells with non-regular secondarybranches in the lower range of experimental growth rates, calledby these authors as cellular/dendrites. They reported that as the
growth velocity increased during Bridgman growth of Zn–Ag alloysfrom the limit of constitutional supercooling to the limit of absolutestability, the microstructure of the Zn-rich phase was subjected tofour transitions: low velocity plane front; low velocity plate-like
32 W.L.R. Santos et al. / Materials Science and Engineering B 182 (2014) 29– 36
F and (bS utectic
ch
ctsdso2sHdZtwmc
iareafsptf
proposed by Jackson and Hunt for the lamellar growth of eutectics:�2v = constant [29]. Thus, such relationship is also shown to beable to encompass the coupled growth of alternated Zn-rich
ig. 3. (a) Directionally solidified macrostructure of the Zn–20wt%Sn alloy casting
EM image at position 10 mm revealing the presence of fine �-Zn platelets in the e
ells; regular cells or dendrites; high velocity plate-like cells andigh velocity plane front [26].
Matsugi et al. [27] and Mahmudi and Eslami [28] reported aontinuous presence of primary �-Zn surrounded by the eutec-ic mixture throughout the microstructure of a Zn–20wt%Sn alloyolidified against a carbon steel mold in air, and cast into 16 mmiameter bars, respectively. In a study by Lee et al. [17] micro-tructures of a Zn–20wt%Sn alloy are exhibited, which werebtained under two different cooling rates: 0.05 K/s (3 ◦C/min) and8 K/s. The latter one is quite similar to the longitudinal micro-tructures obtained in the present study and shown in Fig. 3.owever, in the present investigation no evidence of �-Zn den-ritic growth has been observed for the directionally solidifiedn–20wt%Sn alloy. Except for microstructures associated withhe lower regime of cooling rates experimentally examined, inhich non-regular lateral branches indicated a cellular/dendriticicrostructure (Fig. 4), plate-like cells prevailed along the entire
asting length.Plots of the corresponding microstructure parameter, i.e., the
nterphase spacing, �, can be seen in Fig. 5. In this figure, pointsre experimental average results of interphase-spacings and linesepresent empirical fits to the experimental scatter, with � beingxpressed as a power function of both the tip growth rate (Fig. 5a)nd the cooling rate (Fig. 5b), and the error bars indicate the dif-erence between minimum and maximum values of experimental
pacings. A large spacing variation with the experimental thermalarameters can be observed in Figs. 5a and b. The exponents ofhe experimental growth laws were found to be −1/2 and −1/4or VL and T , respectively. In the case of the dependence of �
) longitudinal as-cast microstructures for different positions with insertion of (c) a mixture. P is the position from the casting cooled surface.
on the growth rate, this is in agreement with the relationship
Fig. 4. Microstructure corresponding to the lower value of cooling rate experimen-tally examined (of about 0.3 K/s) indicating the occurrence of non-regular secondarybranches.
W.L.R. Santos et al. / Materials Science and Engineering B 182 (2014) 29– 36 33
100
101
102
Zn-20wt. %Sn
Inter
phas
e spa
cing,
λ(μ
m)
Tip gro wth rate , VL (mm /s)
λ.=23 .4VL -1/2 - R2=0.94
(a)
10-1 100 101
101
102
Zn-20w t. %Sn
Inter
phas
e spa
cing,
λ(μ
m)
Cooli ng rate, T (K /s)
λ.=38 T -1/4 -R2=0.93
(b)
F for a Zl struct
psati
ig. 5. Interphase spacing as a function of (a) tip growth rate and (b) cooling rate
ongitudinal SEM and optical images, respectively, and arrows correlate each micro
late-like cells/eutectic mixture during the transient directional
olidification of the Zn–20wt%Sn solder alloy. A −1/4 exponentllows a fair to good agreement between the empirical fit andhe experimental points corresponding to the dependence of thenterphase spacing on the cooling rate (Fig. 5b). Such exponent is
n–20wt%Sn alloy. Details of microstructure are given by (a) cross-section and (b)ure with the corresponding � value.
supported by analytical heat flow expressions relating VL and T ,
where T is shown to be given by a constant × VL
2 for unidirectionaltransient solidification conditions [30].
Fig. 6 shows the experimental XRF (X-Ray FluorescenceSpectrometer) concentration profile along the length of the
34 W.L.R. Santos et al. / Materials Science and Engineering B 182 (2014) 29– 36
10090807060504030201008
10
12
14
16
18
20
22
24
Zn-20wt%Sn - XRF Zn-20wt%Sn - 2 D Image anal ysis
Tin
cont
ent (
wt.%
Sn)
Position (mm )
FZ
ZicsStirftcasttpip
ecstvtf
(�frt
wvTcli
pAet
50454035300
100200300400500600700800900
5000
10000
15000
20000
25000
30000
Inte
nsity
(a.u
.)
Angle (2θθ)
Sn Zn
Zn-20wt.%Sn
Position =5mm
5045403530200300400500600700800900
100 0110 0
500 0
1000 0
1500 0
2000 0
25000
3000 0
Position =10mm
Inte
nsity
(a.u
.)
Angle (2θθ)
Sn Zn
Zn-20wt.%Sn
5045403530400500600700800900
1000110012001300
5000
10000
15000
20000
25000
30000
Position =40mm
Inte
nsity
(a.u
.)
Ang le (2θθ)
Sn Zn
Zn-20 wt.%Sn
ig. 6. Macrosegregation profile along the length of the directionally solidifiedn–20wt%Sn alloy casting.
n–20wt%Sn alloy casting. Sn is rejected at the solidificationnterface during the progress of solidification since the partitionoefficient (k) is lower than 1.0. The solute-rich liquid ahead theolidification front tends to move toward the casting bottom, sincen is slightly denser than Zn, while the solidification progresses ver-ically upwards. The final positive macrosegregation of Sn, shownn Fig. 6, can be further explained by the mechanism of inverse seg-egation which involves volumetric contraction on solidificationollowed by interdendritic flow of the Sn-enriched melt. This leadso a definitive solute enriched zone at the bottom of solidified Zn–Snasting. A large solidification range characterizes the Zn–20wt%Snlloy, i.e., 185 ◦C. This feature is known as a cause of long-rangeegregation. Three extra points were plotted within Fig. 6 in ordero confirm the trend found by XRF analyses. These average Sn con-ent values were determined by counting the area fractions of eachhase using an image processing software (Image J). At least 10
mages were examined to yield the value corresponding to eachoint inserted in Fig. 6 (white squares).
According to Fig. 6, the composition is rather constant consid-ring the range of relative positions from 10 to 50 mm. This rangeorresponds to cooling rates from 1.0 to 12.5 K/s and interphasepacings from 40 to 20 �m. This means that after the drop of tin con-ent along the first Sn-enriched region, the subsequent quite largeariations in cooling rate and interphase spacing were not enougho change local tin concentrations, which tended to diminish onlyor cooling rates lower than 1.0 K/s.
The X-ray diffractograms determined for the Zn–20wt%Sn alloyFig. 7) depict the presence of peaks associated with both �-Sn and-Zn phases with lower intensities of � peaks found for positions
arther from the cooled bottom of the casting, i.e., lower coolingates. This feature found by XRD analysis can be rather connectedo the macrosegregation profile shown in Fig. 6.
A variation concerning the intensities of the peaks associatedith the Sn-rich phase is observed in Fig. 7 as the cooling rate is
aried. Each XRD spectrum corresponds to a particular cooling rate.his means that � peaks of lower intensity are connected to lowerooling rate values, i.e., higher relative positions along the castingength. This is in agreement with the preceding XRF results depictedn Fig. 6.
The alloy hardness increases with the decrease in the inter-
hase spacing for the Zn–20wt%Sn alloy as can be seen in Fig. 8.
typical residual hardness impression is depicted in Fig. 8 cov-ring a considerable area of the microstructure. This guaranteeshat the measured hardness is representative of the ensemble
Fig. 7. Typical X-ray diffraction (XRD) patterns of the Zn–20wt%Sn solder alloysamples for different positions along the casting length.
Zn-rich plate-like cells/eutectic mixture. A modified Hall–Petchtype correlation (HV = H0 − K1�−1/2 + K2�−1) fits adequately theexperimental scatter. Smaller interphase spacing means a betterdistribution and alternation of the reinforcing phase, i.e., theZn-rich phase since Zn has a higher hardness than Sn. However,the interphase spacing is one of the factors affecting the hardness
profile. The local Sn composition (which varies along the castinglength according to Fig. 6) is also affecting the local hardness. Whilesmaller � at positions close to the casting surface contributes toincrease the local hardness, the Sn content higher than that of the
W.L.R. Santos et al. / Materials Science and Engineering B 182 (2014) 29– 36 35
0.210.180.150.12 0.330.300.270.24
26
28
30
32
34
36
38
40
42
44 H
ardn
ess
HV
λ-1/2 ( μm-1/2
)
Zn-20wt%Sn HV = 31.8-25 (λ)
-1/2+130 (λ)-1
inter
atmoframa
4
amZlambwlofot
ohttp
[[
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[
[
Fig. 8. Hardness evolution as a function of the
lloy nominal composition acts in the opposite way contributinghe decrease it, since hardness is an interactive property. The
odified Hall–Petch approach, which relates hardness to the scalef the microstructure in the present investigation, differs slightlyrom previous Hall–Petch type equations proposed recently toelate hardness and the microstructural scale of Al–Fe and Zn–Culloys [31,32]. This is caused by the fact that tin is not only accu-ulated in the first positions of the Zn–20wt%Sn alloy casting but
lso dispersed for positions higher than 60 mm (see Fig. 6).
. Conclusions
An alternate pattern of elongated �-Zn plate-like cells andn intercellular eutectic mixture characterizes the as-solidifiedicrostructure along the length of the directionally solidified
n–20wt%Sn solder alloy casting, resembling a eutectic-like lamel-ar arrangement in morphology. The cooling and growth ratesttained during solidification were enough to promote develop-ent of high velocity plate-like cells. The relationship proposed
y Jackson and Hunt for the growth of lamellar eutectics (�2v = C)as able to encompass the alternated growth of the Zn-rich plate-
ike cells and the eutectic mixture during the directional growthf the Zn–20wt%Sn alloy. In addition, a Hall–Petch type statistical-ormula is proposed in order to represent the hardness dependencyn both the microstructure scale and long-term segregation alonghe casting length.
The experimental interrelations for the interphase spacingbtained through the present study can be used coupled with the
ardness Hall–Petch type expressions allowing for the first time inhe case of Zn–Sn solder alloys the preprogramming of the scale ofhe solidification microstructure during soldering in terms of somearticular level of strength which is desired to be achieved.
[
[
phase spacing for a Zn–20wt%Sn solder alloy.
Acknowledgments
The authors acknowledge the financial support provided byFAPESP - São Paulo Research Foundation, Brazil (grant 2012/08494-0) and CNPq (The Brazilian Research Council).
References
[1] K. Suganuma, S. Kim, K. Kim, JOM 61 (2009) 64–71.[2] W.R. Osório, L.C. Peixoto, L.R. Garcia, N. Mangelinck-Noël, A. Garcia, J. Alloys
Compd. 572 (2013) 97–106.[3] J. Chen, J. Shen, D. Min, C.F. Peng, J. Mater. Sci.: Mater. Electron. 20 (2009)
1112–1117.[4] Y. Li, K.-S. Moon, C.P. Wong, Science 308 (2005) 1419–1420.[5] E. C adırlı, U. Böyük, S. Engin, H. Kaya, N. Maraslı, A. Ülgen, J. Alloys Compd. 486
(2009) 199–206.[6] L.R. Garcia, W.R. Osório, L.C. Peixoto, A. Garcia, J. Electron. Mater. 38 (2009)
2405–2414.[7] A. Ahmido, A. Sabbar, H. Zouihri, K. Dakhsi, F. Guedira, M. Serghini-Idrissi, S. El
Hajjaji, Mater. Sci. Eng. B 176 (2011) 1032–1036.[8] I. Wang, J. Duh, C. Cheng, J. Wang, Mater. Sci. Eng. B 177 (2012) 278–282.[9] A. Sharif, Y.C. Chan, Mater. Sci. Eng. B 106 (2004) 126–131.10] G. Zeng, S. McDonald, K. Nogita, Microelectron. Reliab. 52 (2012) 1306–1322.11] J.E. Lee, K.S. Kim, K. Suganuma, J. Takenaka, K. Hagio, Mater. Trans. (2005)
2413–2418.12] V. Chidambaram, J. Hattel, J. Hald, Microelectron. Eng. 88 (2011) 981–989.13] F.X. Che, W.H. Zhu, E.S.W. Poh, X.W. Zhang, X.R. Zhang, J. Alloys Compd. 507
(2010) 215–224.14] M. Wang, J. Wang, H. Feng, W. Ke, Corros. Sci. 63 (2012) 20–28.15] G. Montesperelli, M. Rapone, F. Nanni, P. Travaglia, P. Riani, R. Marazza, Mater.
Corros. 59 (2008) 662–669.16] X. Liu, M. Huang, Y. Zhao, C.M.L. Wu, L. Wang, J. Alloys Compd. 492 (2010)
433–438.17] J.E. Lee, K.S. Kim, K. Suganuma, M. Inoue, G. Izuta, Mater. Trans. 48 (2007)
584–593.18] I.L. Ferreira, J.E. Spinelli, J.C. Pires, A. Garcia, Mater. Sci. Eng. A 408 (2005)
317–325.19] M.V. Canté, J.E. Spinelli, I.L. Ferreira, N. Cheung, A. Garcia, Metall. Mater. Trans.
20] D.M. Rosa, J.E. Spinelli, I.L. Ferreira, A. Garcia, Metall. Mater. Trans. A 39 (2008)2161–2174.
21] M. Gunduz, E. C ardili, Mater. Sci. Eng. A 327 (2002) 167–185.22] P.R. Goulart, K.S. Cruz, J.E. Spinelli, I.L. Ferreira, N. Cheung, A. Garcia, J. Alloys
Compd. 470 (2009) 589–599.23] D. Ma, Y. Li, S.C. NG, H. Jones, Acta Mater. 48 (2000) 419–431.24] D. Ma, Y. Li, S.C. N.G., H. Jones, Sci. Technol. Adv. Mater. 2 (2001) 127–130.25] W. Xu, Y.P. Feng, Y. Li, G.D. Zhang, Z.Y. Li, Acta Mater. 50 (2002) 183–193.26] W. Xu, Y.P. Feng, Y. Li, Z.Y. Li, Mater. Sci. Eng. A 373 (2004) 139–145.
[
[[
d Engineering B 182 (2014) 29– 36
27] K. Matsugi, G. Sasaki, O. Yanagisawa, Y. Kumagai, K. Fujii, Mater. Trans. 48 (2007)1105–1112.
28] R. Mahmudi, M. Eslami, J. Mater. Sci.: Mater. Electron. 22 (2011) 1168–1172.
30] I.L. Ferreira, C.A. Santos, V.R. Voller, A. Garcia, Metall. Mater. Trans. B 35 (2004)285–297.
31] B. Silva, A. Garcia, J.E. Spinelli, Mater. Lett. 89 (2012) 291–295.32] C. Brito, C.A. Siqueira, J.E. Spinelli, A. Garcia, Mater. Lett. 80 (2012) 106–109.
