ong>Nanoindentationong> ong>studyong> on the creep resistance of SnBi solder alloy
with reactive nano-metallic fillers
Lu Shen a,b , Zheng Yu Tan b , Zhong Chen b,n
a Institute of Materials Research and Engineering, A n STAR (Agency for Science, Technology and Research), 3 Research Link, Singapore 117602, Singapore
b School of Materials Science and Engineering, Nanyang Technological University, Nanyang Avenue, Singapore 639798, Singapore
Received 24 August 2012
Received in revised form
22 October 2012
Accepted 25 October 2012
Available online 10 November 2012
SnBi, with its superior yield strength, fracture resistance and
comparable solderability with SnPb alloy , has drawn great
research focus as potential replacement for Pb-containing solders
in the microelectronic industry. Its low melting point (T m¼139 1C
at eutectic concentration) could effectively reduce the thermal
stress built-up at the electronic joints during the multiple reflow
processes. However, the high homologous temperature at service
environment or even at room temperature causes large creep
deformation in the alloy. As the creep properties are closely
related to fatigue life which is a function of the accumulated
plastic deformation, enhancing the creep resistance of the low Tm
solder becomes one of the chief objectives of the works carried
out with such alloys. Composite solder, as one of the effective
ways to enhance the hardness and strength of the alloy, has been
investigated extensively. However, relatively less work has been
done for the creep enhancement. In what have been reported so
far, it was suggested that creep resistance is enhanced by
stabilizing the fine-grained microstructure so that homogenized
deformation is realized [2,3].
To serve as suitable filler materials, similar density of the filler
as the constituent materials of the base solder is required as it
n Corresponding author. Tel.: þ65 6790 4256; fax: þ65 6790 9081.
E-mail addresses: ASZChen@ntu.edu.sg, ZhongChen2006@gmail.com (Z. Chen).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 561 (2013) 232–238
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A
journal homepage: www.elsevier.com/locate/msea
SnBi alloy is an attractive soldering material for temperature-sensitive electronic devices. With its
excellent yield strength and fracture resistance, SnBi alloy has become one of the promising candidates
to replace Pb-based solders. However, due to the low melting temperatures of this alloy, the prominent
time-dependent deformation at service temperatures hinders its wide applications. In this ong>studyong>, low
concentration (no more than 4 wt%) of reactive nano-metallic fillers, i.e., Cu and Ni, have been added
into the Sn–58Bi alloy aiming to enhance its creep resistance. The elastic, plastic and creep properties
are characterized by nanoindentation constant strain rate (CSR) technique. The addition of the fillers
has refined the microstructure of the solder matrix leading to moderate strengthening and hardening.
The creep resistance of the Sn–58Bi alloy has been improved with the filler addition. Two regions of
stress-dependent creep rates were found in the alloys with and without fillers. An optimum filler
concentration for creep resistance enhancement is identified at which there is a balance between the
effects of particle pinning and microstructure refinement.
& 2012 Elsevier B.V. All rights reserved.
promises a good distribution of the filler in the matrix. Moreover,
it requires good wetting between filler and base solder as well as
a similar Tm and solderability between the composite alloy and its
non-modified counter-part. Ceramic particles, such as Al 2O 3 [4,5]
and TiO2 , carbon nano-tubes (CNTs) [6,7] and metallic fillers
including Cu [2,8,9], Ni , Sb and Mn , etc., demonstrated
good potential in improving the mechanical properties of
solder alloys. Lin et al.  has reported a 30–40% increase in
hardness of the eutectic SnPb with less than 5% of Cu addition.
Marshall et al.  reported superior mechanical properties
(including Young’s modulus, yield and tensile strength) of the
Cu6Sn5 powder modified SnPb (60/40) alloy for up to 40% of the
filler concentration. Guo et al.  has investigated the creep
resistance of Sn–3.5Ag alloys with low concentration of Ni and Cu
addition. Cu was reported to moderately enhance the creep
resistance of SnAg alloy at both room temperature and elevated
temperature. Ni composite alloy showed much higher creep
resistance than Cu composite at room temperature, but insignificant
effect at elevated temperatures. The differences of the Ni
and Cu fillers were attributed to their solubility and diffusively in
the base alloys. 1 wt% Cu addition was reported as a microstructure
stabilizer for SnBi alloy in which insignificant grain size
variation was found in the matrix alloy after 30 days of annealing
at 80 1C . Such fine grain microstructure is beneficial in
strengthening the alloy matrix. However, it is not ideal for creep
resistance as it promotes grain boundary sliding especially at low
strain rate. Therefore, it is expected that an optimum filler
concentration exist for each compatible filler-solder matrix system,
in which the effect of pinning and restricting of the grain
boundary motion is maximized by the filler addition.
In the present ong>studyong>, nano-Cu and Ni particles were added into
the Sn–58Bi alloy. The creep behaviour of the composite alloys were
studied by nanoindentation constant strain rate (CSR) method .
The stress and strain rate relation is obtained by capturing the stress
responses at eight controlled indentation strain rates. The indentation
creep behaviour and the effect of filler distribution and the
microstructural change are discussed. The current work contributes
to the limited literatures on the effect of filler type, loading, size, etc.
on the creep behaviour of low-melting SnBi solders.
2.1. Materials and sample preparation
Sn-Bi solder paste with the eutectic composition of 42%Sn–
58%Bi and nano-Cu and Ni particles were obtained from ESL
Europe s and US Research Nanomaterials, Inc., respectively. Two
types of composition mixture were prepared, i.e., Sn–Bi solder
paste blended with nano-copper powder (average particle size
100 nm) and Sn–Bi solder paste blended with nano-nickel powder
(average particle size o70 nm). The composite solders with
several weight percentages of the nano-fillers ranging from
0.5 wt% to 4 wt% were prepared. The mixture was homogenously
blended in an aluminum dish first before melt in a furnace
maintained at temperature of 300 1C for an hour. The mixture
composition was then air cooled to room temperature.
The solidified samples were prepared for subsequent morphology
ong>studyong> and mechanical testing. The samples were moulded
using mixture of epoxy and hardener. Sample preparation
involved grinding using progressively finer grade of silicon
carbide papers with copious amounts of water. After which, fine
polishing was carried out using 6, 3 and 1 mm diamond paste with
polishing cloth. Mirror finishing of the sample surfaces was
achieved using silica suspension as the final polishing step.
2.2. Scanning electron microscope (SEM) and energy dispersive
X-ray (EDX) analysis
The polished and indented samples were examined using JOEL
SEM 6360 for the microstructure ong>studyong>. The accelerating voltage
of 10 kV was used for the secondary electron (SEI) and backscattered
electron images (BEI) collection. Elemental analysis of
the samples was done with an Oxford EDX instrument, attached
to JEOL FESEM 7600 F system. The accelerating voltage of 15 kV
was used to examine the nano-particle distribution in the alloys.
2.3. Indentation constant strain rate (CSR) test
The indentation strain rate, _e i, is defined as the instantaneous
displacement rate of the indenter, dh/dt, divided by the instantaneous
_e i ¼ k1U dh
where k1 is a constant. The indentation stress (s) is defined as
instantaneous load divided by projected contact area. The stress
and strain rate follows the phenomenological description as
Eq. (2) when the steady state deformation process is achieved.
_e i ¼ C0sn ð2Þ
where C0 is a material constant including temperature factor and
activation energy for the creep deformation. n is the creep stress
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238 233
exponent which can be derived as the slope of the linear fit of a
series of stress and strain rate pairs plotted in logarithm scale.
The indentation tests were performed with the MTS Nano
Indenter s XP (TN, USA) system using a sharp Berkovich tip. Load
was applied at eight constant strain rates from 0.0005 s 1 to
0.1 s 1 to a depth of 4 mm. During loading, a small oscillated force
with known frequency and amplitude was imposed on the
nominal load to extract a series of contact stiffness (S) along
indentation loading. Continuous modulus (E) and hardness (H)
profiles can thus be calculated in respect of indentation depth
from Eqs. (3) and (4) with a single loading/unloading cycle.
E p 1
¼ pffiffiffiffiffiS ð3Þ
1 n2 2 Ac
H ¼ Pmax
where n is Possion’s ratio and set to be 0.35 for current analysis.
Ac is the projected contact area at the maximum load Pmax.
The reported modulus and hardness were averaged values in
the depth range from 2 4 mm where a constant values of the
properties are achieved. Five indentations were made at each
strain rate. The indenter area function describing the actual
contact area as a function of indentation depth was calibrated
on a standard SiO2 sample.
3. Results and discussion
3.1. Microstructure ong>studyong>
Figs. 1 and 2 show the SEM images on the cross-section area of
Sn–58Bi alloys with and without the nano-particles. The dark
regions represent the Sn-rich phase while the bright regions
represent the Bi phase (note that by the Sn–Bi phase diagram,
at 25 1C there is approximately 3 wt% of Bi dissolved in Sn
forming the Sn-rich phase, but there is no Sn dissolution in Bi
phase). The pristine alloy demonstrates a lamellar structure
where the Sn-rich phase and Bi phase inter-locked with each
other (Fig. 1(a)). The average inter phase spacing of the alternate
layered arrangement of this alloy is approximately 2.62 mm. With
small amount of nano-fillers addition, the microstructure of the
composite solder becomes finer (Fig. 1(b)–(d) and Fig. 2(a)–(d)),
which is likely to be caused by heterogeneous nucleation on the
inter metallic compounds (IMC) formed by Cu and Ni fillers with
the Sn in the matrix (elaborated later). The existence of the IMCs
may also play the role of preventing the growth of the lamellar
structures. As the weight percentage of the nano-fillers increases,
the size of inter phase spacing (IPS) generally decreases. Table 1
summarizes the evolution of IPS in SnBi composite alloys with
increased filler concentration. It is noticed that the filler addition
as low as 0.5 wt% could cause great reduction of IPS in SnBi alloy.
Further increasing the filler concentration up to 4 wt% only gently
refined the microstructures. EDX was used to verify the distribution
of the nano-fillers in the alloy matrix. Figs. 3 and 4 show the
elemental mapping of Cu- and Ni-added SnBi alloys, respectively.
Based on existing literatures [2,8,9,15,16], Cu and Ni form Cu 6Sn 5
and Ni3Sn4 IMCs with Sn upon solder melting. Therefore in the
subsequent analyses we assume the presence of these two IMCs.
It is noticed that Cu element seems to be slightly segregated in
the Sn-rich phase (Fig. 3(d)), which indicates that the Cu–Sn IMC
phase is mainly distributed in Sn-rich phase. Ni distribution, on
the other hand, appears to be quite homogeneous across both
phases (Fig. 4(d)). Further clarification is necessary but beyond
the scope of this paper. More detailed analyses using TEM and
electron diffraction are under the way in our lab.
Fig. 1. Morphology of the constituent phase distribution of (a) Sn–58Bi, and the composite alloy with (b) 1%, (c) 2%, and (d) 4% nano-Cu.
Fig. 2. Morphology of the constituent phase distribution of the SnBi composite alloy (a) with (a) 0.5%, (b) 1%, (c) 2%, and (d) 4% nano-Ni.
Evolution of inter phase spacing (IPS) in SnBi composite alloys with increasing
Filler concentration (wt %) Nano-Cu filler (lm) Nano-Ni filler (lm)
0 (Sn–58Bi) 2.6270.18 2.6270.18
0.5 1.7370.26 1.7770.14
1.0 1.5370.13 1.7470.36
2.0 1.5170.23 1.2870.23
3.0 1.3970.13 1.2570.10
4.0 1.2170.19 1.1870.10
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238
3.2. Elastic modulus and hardness
Fig. 5(a) and (b) shows the modulus (E) and hardness (H) of the
Sn–58Bi and its nano-particle composites with filler concentration
up to 4 wt%. The values of E and H are listed in Table 2. The
pristine Sn–58Bi is shown to possess E and H of 38.43 GPa and
287.3 MPa, respectively. The two types of particles demonstrate
similar strengthening effect, in which 4 wt% filler addition
increases the elastic modulus of the SnBi solder slightly
(8 12%). The increment is believed to be attributed to the stiff
nature of IMCs particles in the composite. The moduli of the
composite can be expressed by the empirical Eq. (5), which is the
modification from the iso-strain arrangement of the fillers and
Ec ¼ V mEm þkV pEp
where Ep,m and Vp,m are the IMC particle and matrix modulus and
volume fraction, respectively. k is empirically determined falling
in the range of 0 to 1. The higher the value of k, the greater the
enhancement effect from the IMC particles. The slope of the linear
fit, s, of the composite modulus vs. IMC volume fraction (Vp) can
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238 235
Fig. 3. Elemental mappings of cross-section of SnBi with 4 wt% nano-Cu addition. (a) Low angle backscatter electron (LABE) image and EDX mappings of (b) Sn, (c) Bi,
and (d) Cu.
Fig. 4. Elemental mappings of cross-section of SnBi with 4 wt% nano-Ni addition. (a) Low angle backscatter electron (LABE) image and EDX mappings of (b) Sn, (c) Bi,
and (d) Ni.
be expressed as Eq. (6):
s ¼ k Ep 38:43 ð6Þ
For Cu6Sn5, the reported modulus of 90 GPa  is used as Ep,
while the literature modulus value of 134 GPa for Ni4Sn3 is used
 for Ni composite prediction. As such, the k values of the
nano-Cu and nano-Ni modified SnBi alloy are both found to be
approximately 0.3. The composite modulus of the SnBi(Cu) and
SnBi(Ni) alloys is thus predicted by following equation provided
Fig. 5. (a) Modulus and (b) hardness of Sn–58Bi alloys with and without nano-Cu
and Ni additions.
Modulus and hardness of Sn–58Bi and its composites with nano-Cu and nano-Ni
the solder materials consist of continuous phases without defects.
Ec ¼ 38:43 V SnBi þ0:3V IMCEIMC ð7Þ
Fig. 5(b) shows the hardness of the Sn–58Bi alloy steadily
increases with nano-Cu and Ni concentration. The enhancement
in hardness is more significant, in which 4 wt% of Cu and Ni
addition has increased the hardness of the pristine Sn–58Bi by
16% and 23%, respectively. The hardening effect from the metal
fillers is believed not only coming from the hard IMC phase, but
also the strengthening of the base alloy due to refinement of the
microstructure. The hardness behaviour of the composite alloy is
thus not able to be predicted by simple relationship between the
properties of filler and the matrix.
3.3. Creep behaviour
Modulus (GPa) Hardness (MPa)
Sn–58Bi 38.4370.57 287.3720.0
SnBi–0.5%Cu 38.2971.75 299.9725.9
SnBi–1%Cu 40.3470.74 317.4711.9
SnBi–2%Cu 39.4371.98 311.9717.0
SnBi–3%Cu 39.1270.64 327.7717.0
SnBi–4%Cu 41.6770.98 333.8734.1
SnBi–0.5%Ni 38.8170.26 298.7712.8
SnBi-1% Ni 39.1971.01 297.5710.3
SnBi–2% Ni 38.8971.31 326.578.0
SnBi–3% Ni 39.6972.30 318.3730.5
SnBi–4% Ni 43.2371.08 354.177.6
The addition of filler increases the creep resistance of the unmodified
sample by shifting the strain rate–stress curve towards
higher stress level at the same strain rates. Fig. 6(a) and (b) shows
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238
Fig. 6. Strain rate and stress relation of Sn–58Bi solder comparing with its
composite solders with (a) 4 wt% Cu addition, and (b) 4 wt% Ni addition.
Stress exponents of Sn–58Bi and its composites with nano-Cu and nano-Ni fillers.
Filler concentration (wt %) Cu filler Ni filler
n1 n2 n1 n2
0 (Sn–58Bi) 5.20 2.35 5.20 2.35
0.5 5.58 2.42 6.20 3.73
1.0 7.04 2.51 7.45 3.85
2.0 7.23 3.24 5.89 3.60
3.0 7.92 3.31 5.29 2.95
4.0 6.11 2.14 5.16 2.93
a representative stress-dependent strain rate behaviour of the
Sn–58Bi alloy with 4 wt% of metal fillers. The indentation creep
test results show two stress regions in the strain rate–stress
curves of Sn–58Bi alloys with and without metallic fillers. The
transition stress is approximately unchanged at 170 MPa in spite
of the moderate enhancement of the creep resistance by metal
filler addition. Table 3 lists the stress exponents of the samples
measured at the two stress regions. n1 is defined as the slope of
strain rate–stress curved at higher stress region (4170 MPa);
while n2 is the one measured at stress region lower than 170 MPa.
n 1 for the two types of metal filler samples is found in the range
from 5.16 to 7.45, which implies the dominant creep mechanism
in this stress range is related to dislocation movement in the bulk.
To compare the creep resistance among various samples with
different filler loadings, the strain rates at a particular stress,
selected at 200 MPa and 100 MPa to represent the high stress and
low stress regions respectively, are shown in Fig. 7(a)–(d). The
pristine SnBi alloy shows the highest strain rate (lowest creep
resistance) at both stress levels. With addition of nano-Cu or Ni,
the creep rate reduces significantly and achieved best creep
resistance when filler concentration reaches 3 wt% for Cu filler
composite and 1 wt% for Ni filler composite.
Fig. 8(a) shows the SEM image of SnBi with 4 wt% Ni deformed
at 354 MPa. The indentation imprint displayed a regular triangle
with evaginated boundaries indicating discernible pile-up of the
material during the indentation process. It is believed that the
indentation induced deformation is accomplished by the bulk
deformation inside Sn-rich and Bi phases. As a result, the
hardened individual phases push the surrounding material outwards
causing significant pile-up around the indented area.
However, when material is deformed at a lower stress of
116 MPa where a lower strain rate is applied, significant number
of pores are formed at the boundaries of the adjacent phases as
shown in Fig. 7(b). Prominent delamination at phase boundaries
is formed at the surrounding area of the indent. It is believed that
at a lower stress condition, the deformation of the material is
mainly channelled through slipping or shifting at the phase
boundaries between Sn-rich and Bi phases. Comparing to the
indentation creep results, the stress exponent n 2 at lower stress
region (o170 MPa) is in the range from 2.42 to 3.85, which
corresponds well with the phase boundary sliding mechanism.
As shown above, the creep resistance reaches a maximum at
3 wt% Cu and 1 wt% Ni filler concentration, and then starts
decrease when more filler is added. With the nano-filler incorporation,
it is believed there are two major strengthening
mechanisms from the stiff metal particles. One comes from the
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238 237
Fig. 7. Strain rates of SnBi(Cu) alloys deformed at (a) 200 MPa and (b) 100 MPa. Strain rates of SnBi(Ni) alloys deformed at (c) 200 MPa and (d) 100 MPa.
pinning of the dislocation movement through effective particle–
dislocation interactions. Rosler et al.  demonstrated that such
interactions are primarily responsible for the excellent creep
resistance enhancement of carbide dispersion-strengthened aluminum
alloys. Secondly, the obvious microstructural refinement
due to filler addition strengthens the matrix due to effect such as
Hall–Petch hardening. However, such microstructure refinement
promotes diffusional creep (e.g. Nabarro-Herring, Coble) and
grain boundary sliding, adversely affecting creep resistance especially
at lower stress regions. With the two effects functioning
simultaneously, it is believed that at certain low filler concentrations,
the metal filler pinning effect dominants the creep deformation
until an optimum filler concentration is reached, after that
creep rate starts to increase due to significant amount of diffusional
creep and phase boundary sliding is involved, which causes
the stress exponent shift to a lower value.
Fig. 8. SEM images of Sn–58Bi with Ni 4% deformed at stress of (a) 354 MPa; and (b) 116 MPa.
The elastic, plastic and creep properties of Sn–58Bi alloy have been
studied with varying amount of nano-metal filler additions. Moderate
stiffness enhancement is found in the two types of composite alloys.
The formed hard particles and the microstructure refinement due to
nano-filler addition promote the hardening of the alloys. The creep
resistance of the Sn–58Bi is found to increase with addition of nanometal
fillers. At higher stress range (4170 MPa), the creep mechanism
is dominated by dislocation movement in the bulk; while at
lower stress range, it is dominated by phase boundary sliding. 3 wt%
Cu and 1 wt% Ni are found to be the optimum concentrations for the
creep resistance enhancement for Sn–58Bi alloy.
 F. Hua, Z. Mei, J. Glazer, in: 48th Electronic Components and Technology
Conference, 227 (1998).
 F. Guo, J. Mater. Sci. Mater. Electron. 18 (2007) 129–145.
 J. Shen, Y.C. Chan, Microelectron. Reliab. 49 (2009) 223–234.
 H. Mavoori, S. Jin, J. Electron. Mater. 27 (1998) 1216–1222.
 X.L. Zhong, M. Gupta, J. Phys. D: Appl. Phys. 41 (2008).
 K. Kumar, V. Kripesh, L. Shen, A. Tay, Thin Solid Films 504 (2006) 371–378.
 S.M.L. Nai, J. Wei, M. Gupta, J. Electron. Mater. 35 (2006) 1518–1522.
L. Shen et al. / Materials Science & Engineering A 561 (2013) 232–238
 D. Lin, G.X. Wang, T.S. Srivatsan, M. Al-Hajri, M. Petraroli, Mater. Lett. 53
 F. Guo, S. Choi, K.N. Subramanian, T.R. Bieler, J.P. Lucas, A. Achari,
M. Paruchuri, Mater. Sci. Eng., A 351 (2003) 190–199.
 F. Guo, G. Xu, H. He, J. Mater. Sci. 44 (2009) 5595–5601.
 K.S. Kim, S.H. Huh, K. Suganuma, Microelectron. Reliab. 43 (2003) 259–267.
 J.L. Marshall, J. Calderon, J. Sees, G. Lucey, J.S. Hwang, IEEE Trans. Compon.,
Hybrids, Manuf. Technol. 14 (1991) 698–702.
 L.E. Felton, C.H. Raeder, D.B. Knorr, J. Miner. Met. Mater. Soc. 45 (1993)
 B.N. Lucas, W.C. Oliver, Metall. Mater. Trans. A 30 (1999) 601–610.
 M. He, Z. Chen, G.J. Qi, C.C. Wong, S.G. Mhaisalkar, Thin Solid Films 462
 M. He, Z. Chen, G.J. Qi, Metall. Mater. Trans. A 36A (2005) 65–75.
 T.H. Courtney, Mechanical Behavior of Materials, second ed., McGraw Hill,
 J.H. Westbrook, R.L. Fleischer, Structural Applications of Intermetallic Compounds,
Wiley, New York, 2000.
 Z. Chen, M. He, B. Balakrisnan, C.C. Chum, Mater. Sci. Eng., A 423 (2006)
 J. Rosler, R. Joos, E. Arzt, Metall. Trans. A 23 (1992) 1521–1539.