Nanoindentation study on the creep resistance of SnBi solder alloy ...

Nanoindentation study on the creep resistance of SnBi solder alloy ...

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

article info

Article history:

Received 24 August 2012

Received in revised form

22 October 2012

Accepted 25 October 2012

Available online 10 November 2012






1. Introduction


SnBi, with its superior yield strength, fracture resistance and

comparable solderability with SnPb alloy [1], 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:, (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:

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 [4], carbon nano-tubes (CNTs) [6,7] and metallic fillers

including Cu [2,8,9], Ni [9], Sb[10] and Mn [11], etc., demonstrated

good potential in improving the mechanical properties of

solder alloys. Lin et al. [8] has reported a 30–40% increase in

hardness of the eutectic SnPb with less than 5% of Cu addition.

Marshall et al. [12] 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. [9] 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 [13]. 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 [14].

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. Experimental

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

displacement, h:

_e i ¼ k1U dh


U 1


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.

Table 1

Evolution of inter phase spacing (IPS) in SnBi composite alloys with increasing

filler percentage.

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

matrix [17]:

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 [18] is used as Ep,

while the literature modulus value of 134 GPa for Ni4Sn3 is used

[19] 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.

Table 2

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.

Table 3

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. [20] 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.

4. Conclusions


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.


[1] F. Hua, Z. Mei, J. Glazer, in: 48th Electronic Components and Technology

Conference, 227 (1998).

[2] F. Guo, J. Mater. Sci. Mater. Electron. 18 (2007) 129–145.

[3] J. Shen, Y.C. Chan, Microelectron. Reliab. 49 (2009) 223–234.

[4] H. Mavoori, S. Jin, J. Electron. Mater. 27 (1998) 1216–1222.

[5] X.L. Zhong, M. Gupta, J. Phys. D: Appl. Phys. 41 (2008).

[6] K. Kumar, V. Kripesh, L. Shen, A. Tay, Thin Solid Films 504 (2006) 371–378.

[7] 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

[8] D. Lin, G.X. Wang, T.S. Srivatsan, M. Al-Hajri, M. Petraroli, Mater. Lett. 53

(2002) 333–338.

[9] 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.

[10] F. Guo, G. Xu, H. He, J. Mater. Sci. 44 (2009) 5595–5601.

[11] K.S. Kim, S.H. Huh, K. Suganuma, Microelectron. Reliab. 43 (2003) 259–267.

[12] J.L. Marshall, J. Calderon, J. Sees, G. Lucey, J.S. Hwang, IEEE Trans. Compon.,

Hybrids, Manuf. Technol. 14 (1991) 698–702.

[13] L.E. Felton, C.H. Raeder, D.B. Knorr, J. Miner. Met. Mater. Soc. 45 (1993)


[14] B.N. Lucas, W.C. Oliver, Metall. Mater. Trans. A 30 (1999) 601–610.

[15] M. He, Z. Chen, G.J. Qi, C.C. Wong, S.G. Mhaisalkar, Thin Solid Films 462

(2004) 363–369.

[16] M. He, Z. Chen, G.J. Qi, Metall. Mater. Trans. A 36A (2005) 65–75.

[17] T.H. Courtney, Mechanical Behavior of Materials, second ed., McGraw Hill,


[18] J.H. Westbrook, R.L. Fleischer, Structural Applications of Intermetallic Compounds,

Wiley, New York, 2000.

[19] Z. Chen, M. He, B. Balakrisnan, C.C. Chum, Mater. Sci. Eng., A 423 (2006)


[20] J. Rosler, R. Joos, E. Arzt, Metall. Trans. A 23 (1992) 1521–1539.

More magazines by this user
Similar magazines