Available online at www.sciencedirect.com
Acta Materialia 57 (2009) 5196–5206
www.elsevier.com/locate/actamat
Growth mechanism of Ni3Sn4 in a Sn/Ni liquid/solid interfacial reaction
J. Shen a, Y.C. Chan a,*, S.Y. Liu b
a
Department of Electronic Engineering, City University of Hong Kong, Tat Chee Avenue, Kowloon Tong, Hong Kong
b
Department of Mechanical Engineering, University of Hong Kong, Pokflum Road, Hong Kong
Received 1 November 2008; received in revised form 6 July 2009; accepted 13 July 2009
Available online 12 August 2009
Abstract
The chemical interfacial reaction of Ni plates with eutectic Sn–3.5Ag lead-free solder was studied by microstructural observations and
mathematical calculations. Compared with the Sn–3.5Ag–0.75Ni/Ni interfacial reaction, based on a simple model of the growth of the
liquid/solid chemical compound layer, the growth mechanism of Ni3Sn4 in the Sn–3.5Ag/Ni interfacial reaction is discussed and presented. The growth process of Ni3Sn4 in the Sn/Ni liquid/solid reaction interface involves the net effect of several interrelated phenomena, such as volume diffusion, grain boundary diffusion, grain boundary grooving, grain coarsening, and dissolution into the molten
solder. The growth time exponent n and morphology of Ni3Sn4 were found to be dependent on these factors.
Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Sn/Ni; Interfacial reaction; Microstructure; Growth mechanism
1. Introduction
Solders, substrate materials and their interfacial reaction
products play crucial roles in the reliability of joint assemblies in microelectronic packages because they provide electrical, thermal and mechanical continuity in electronic
assemblies [1,2]. The sustained trend towards miniaturization and functional density enhancement in electronic
devices, the use of increasingly dense arrays and fine-pitch
interconnections for microelectronic packaging, and the
development of lead-free solders in response to the strict
legislation of a ban on the use of lead-based solders have
posed several new challenges to the microelectronic packaging industry [3]. Whether in the first-level (chip-to-module) or in the second-level (module-to-board) packaging
technologies for advanced electronic applications, the most
important issue in packaging is that the molten solders flow
or spread on the substrate surfaces to form a proper metallic bond and thus achieve a perfect joint [1–3]. Hence a
thin, continuous and uniform intermetallic compound
*
Corresponding author. Tel.: +852 2788 7130; fax: +852 2788 8803.
E-mail address: [email protected] (Y.C. Chan).
(IMC) layer formed between a solder and the substrate
material is an essential requirement for good metallurgical
bonding. However, due to the inherently brittle nature of
IMC layers and their tendency to generate structural
defects, IMC layers which are too thick at the solder/substrate material interface may degrade the fatigue and fracture strengths of solder joints, leading to poor reliability of
electronic devices [2]. Copper has been the most widely
used solderable metal substrate in the under bump metallization (UBM) for flip-chip and in the bond pad for ball
grid array (BGA) applications [1–5]. Because Cu dissolves
into molten Sn-rich solders very quickly during the soldering process and the Sn–Cu IMC layers grow at a very high
rate to become thick during thermal aging, this excessive
growth of Sn–Cu IMC layers may have a deleterious effect
on the reliability of solder joints when electronic devices are
used in service at high temperatures [5–8]. Often, Ni and
Ni-based alloys, which are also solderable metals, are considered to be excellent alternatives for Cu substrates. The
rate of dissolution of Ni in Sn-based solders is very low
at the soldering temperature so that only a very thin
IMC layer is generally observed between Ni and the
Sn-based solders [9–11]. This is the reason that Ni/Au
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
doi:10.1016/j.actamat.2009.07.021
J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
and Ni/Pd metallization schemes are widely applied as substrates for solder joints in advanced microelectronics
packaging.
Because of the prime importance of the interfacial reaction between solder alloys and metal substrates, the morphologies and the growth kinetics of the IMC layer in the
Sn/Ni system have been widely researched and reported
[9,10,12–26]. However, it is difficult to achieve a clear
image regarding the phases, their morphologies and
growth kinetics (in particular, the time exponent) of the
IMC layer in the Sn/Ni interfacial reaction system from
the literature because there are ambiguous results
reported and some results are in conflict. It is widely
reported that Ni3Sn4 is the only phase formed that can
be detected in the interfacial reaction layer between molten pure Sn or Sn–X (X is one of the chemically inert
metallic elements with Ni, such as Pb, Ag, etc.) alloys
and Ni [9,15,17,22,25]. However, Ni3Sn and/or Ni3Sn2
phases, which are thermodynamically stable phases, have
also been reported to exist in the Sn/Ni couple reaction
interface [13,14]. In particular, the morphology of the
Ni3Sn4 phase has been described differently, such as whisker-like [13], scallop-like with round and smooth surfaces
(a non-faceted structure) [14,15,22], scallop-like with
cusps (a faceted structure) [9,22], a continuous thin layer
[17,22], a nonuniform and fractured layer [17], a ‘‘chunk”
type [10] and a faceted rod type [10], etc. Generally, alloying elements, which are not chemically inert with Ni,
influence the phases of the IMC layers formed between
Sn-based solders and Ni. For example, a minor Cu addition to a Sn–Ag solder changed the interfacial reaction
between molten solder and Ni dramatically to form ternary (Cu,Ni)6Sn5 or/and (Ni,Cu)3Sn4 IMCs phases in
the IMC layers [19,20,23]. The effect of minor Zn additions in a Sn–Ag solder is to nucleate a Ni5Zn21 IMC
layer and to suppress the formation of Ni–Sn IMC layers
[24]. It is interesting that, although Bi reacts with Ni to
form both soft and brittle NiBi3 IMC particles in the
Ni/Bi couple, except for Ni3Sn4, none of the other Ni–
Sn IMCs or Ni–Bi IMCs was observed when molten
58Bi–42Sn solder reacted with Ni [22]. It is easy to understand that the concentration of Sn in Sn–X (X is one of
the chemically inert metallic elements with Ni) solders
influences the thickness of the Ni3Sn4 layer because a larger amount of Sn atoms react with the Ni atoms, and of
course, this results in a thicker Ni3Sn4 layer [16,25]. The
most ambiguous description of Ni3Sn4 is its growth kinetics during thermal aging. According to the literature, various growth kinetics of the Ni3Sn4 layer have been
reported, which could be roughly classified as parabolic
kinetics [14,15,25,26], linear and parabolic kinetics (in different thermal aging stages) [22] and nonparabolic growth
kinetics [9,12,13,18].
Hence, a question arises: what is the intrinsic growth
mechanism of Ni3Sn4 in the Sn/Ni liquid/solid interfacial
reaction system? Because of the significance of the Sn/Ni
interfacial reaction in advanced microelectronics packaging
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technology, there is a particular need to develop a comprehensive understanding of this system. In this paper, a eutectic Sn–3.5Ag solder alloy, which is regarded as the most
recommended lead-free solder candidate to replace leadbased solder [1–4], was used to investigate the interfacial
reaction between liquid solder and Ni plates. Moreover,
minor amounts of Ni were introduced into the Sn–3.5Ag
solder to assist in clarifying the growth mechanism of
Ni3Sn4. This paper aims to provide a clear picture of the
growth mechanism of Ni3Sn4 in the Sn/Ni interfacial reaction system.
2. Experiments
In order to avoid the influence of impurities in commercial solder alloy bars/pastes on the interfacial reaction, Sn–
3.5Ag and Sn–3.5Ag–0.75Ni (mass%, hereafter) solder
alloys were prepared from bulk rods of pure Sn, Ag and
Ni (their purities were all above 99.99%). After weighing
the individual pure metals, they were mixed and melted
in a vacuum arc furnace under a high-purity argon atmosphere to produce button-like specimens with a diameter
of about 3.5 cm. In order to get a homogeneous composition, all ingots were remelted four times. Finally they were
solidified in a water-cooled copper mold with a cooling rate
of about 20 K s1. Pure Ni plates with a thickness 0.15 mm
(the purity was above 99.99%) were adopted as substrates
for interfacial reaction tests. The pure Ni plates were cut
to 1 1 cm specimens and were polished lightly with diamond powder and degreased in a solution (99 vol.%
C2H5OH + 1 vol.% HCl) so as to remove surface oxides
and contaminants. Both multiple reflow tests and thermal
aging experiments were carried out in an air environment.
In the multiple reflow tests, Sn–3.5Ag and Sn–3.5Ag–
0.75Ni alloys were machined into /10 mm 5 mm size
samples and placed on the prefluxed Ni plates (using
water-soluble EP9301 flux) to be reflowed 1, 3, 5, and 7
times in a five-zone forced convection reflow oven (BTU
VIP-70 N) with the highest temperature of 523 K for
1 min. In the thermal aging experiments, the machined
Sn–3.5Ag and Sn–3.5Ag–0.75Ni alloy samples were placed
on the prefluxed Ni plates and subjected to high-temperature aging at 523 K and 553 K for times up to 10 h.
After multiple reflows and thermal aging tests, the specimens were sectioned carefully using a slow speed diamond
saw and mounted in an epoxy. The cross-sections of the solder/Ni interfaces were prepared using standard metallographic procedures (grinding and polishing). The
microstructures were characterized by scanning electron
microscopy (Philips, Inc. XL 40 FEG SEM) in the back-scattered electron (BSE) mode and transmission electron microscopy (Philips Tecnai G2 20 S-TWIN). Energy dispersive Xray spectroscopy (EDX) (OXFORD, Inc. ISIS300), using a
standard atomic number, absorption, fluorescence (ZAF)
correction, and X-ray diffraction (Siemens D500 XRD),
were used to determine the phase composition and the crystal
structure of the IMC layers.
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J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
3. Results
3.1. Microstructural evolution of the interfacial reaction
product
Both in the multiple reflow samples and in the thermal
aging test samples, interfacial reaction product layers were
clearly observed (see Figs. 1–3). EDS analysis results show
that these IMC layers are a composite of Sn and Ni elements.
Furthermore, XRD tests were performed in the as-reflowed
sample and a sample after 10 h thermal aging (they are the
representative of other samples) (see Fig. 4). It was found
in our study that only Ni3Sn4 was detected as the reaction
product. This result is in accord with some experimental
results reported in the literature [9,15,17,22,25], while it is
inconsistent with other test results in the literature [13,14],
which reported that Ni3Sn and/or Ni3Sn2 phases were also
formed by the Sn/Ni interfacial reaction.
Experiments with different numbers of reflow cycles or
periods of thermal aging gave Ni3Sn4 with different morphologies. The microstructure of Ni3Sn4, produced in the
as-reflowed Sn–3.5Ag/Ni sample, was loose-like, including
a thin main layer at the interface and some discontinuous
Ni3Sn4 particles embedded inside the Ni plate (see
Fig. 1a). By contrast, a continuous and uniform Ni3Sn4
layer was formed in the as-reflowed Sn–3.5Ag–0.75Ni/Ni
sample (see Fig. 1c). The microstructures of Ni3Sn4 in
Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples after being
reflowed for seven cycles were representative of all multiple
reflows samples. As seen in Fig. 1b and d, they show almost
the same microstructure compared with the corresponding
as-reflowed samples (note the thickness of the Ni3Sn4 layers
did not increase dramatically).
In the thermally aged test samples, between the solder
matrices and Ni plates, Ni3Sn4 layers were clearly
observed, showing different morphologies and thicknesses.
Obviously, both in the Sn–3.5Ag/Ni and in Sn–3.5Ag–
0.75Ni/Ni samples, the thickness of the Ni3Sn4 layers
increased with an increase in the thermal aging time. However, the evolution of morphology of these Ni3Sn4 layers
was different in the Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/
Ni samples. A thin and continuous Ni3Sn4 layer was
formed in the Sn–3.5Ag/Ni sample after aging for 5 min
at 523 K (see Fig. 2a). After thermal aging for several
hours, some isolated Ni3Sn4 particles were found to have
departed from the main Ni3Sn4 layer and entered into the
solder matrices (see Fig. 2b and c). The interfaces of the
solder/Ni3Sn4 layers were not flat and some bulges
appeared with rounded edges (this bulging morphology
of Ni3Sn4 usually is described as a ‘‘scallop” structure).
The Ni3Sn4 layers formed in the Sn–3.5Ag–0.75Ni/Ni samples were thicker than those in the Sn–3.5Ag/Ni samples
for the same thermal aging time. As seen in Fig. 2d, some
Ni3Sn4 particles, originating from the solidification of the
solder alloy, were found to be distributed in a wide region
of the solder matrices showing faceted structures in the Sn–
3.5Ag–0.75Ni/Ni sample thermally aged for 5 min. However, a few solidified Ni3Sn4 particles appeared in the solder
matrices of Sn–3.5Ag–0.75Ni/Ni samples which had been
subjected to thermal aging for several hours (see Fig. 2e
and f). The surfaces of the Ni3Sn4 layers in Sn–3.5Ag–
0.75Ni/Ni samples were also scallop-like but with faceted
Fig. 1. SEM micrographs of cross-sections of IMC layers formed in reflowed samples: (a) Sn–3.5Ag/Ni as-reflowed solder joint, (b) Sn–3.5Ag/Ni solder
joint reflowed for seven cycles, (c) Sn–3.5Ag–0.75Ni/Ni as-reflowed solder joint, and (d) Sn–3.5Ag–0.75Ni/Ni solder joint reflowed for seven cycles.
J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
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Fig. 2. SEM micrographs of cross-sections of IMC layers formed in samples thermally aged at 523 K: (a) Sn–3.5Ag/Ni solder joint thermally aged for
5 min, (b) Sn–3.5Ag/Ni solder joint thermally aged for 6 h, (c) Sn–3.5Ag/Ni solder joint thermally aged for 10 h, (d) Sn–3.5Ag–0.75Ni/Ni solder joint
thermally aged for 5 min, (e) Sn–3.5Ag–0.75Ni/Ni solder joint thermally aged for 6 h, and (f) Sn–3.5Ag–0.75Ni/Ni solder joint thermally aged for 10 h.
edges. Similar microstructures of Ni3Sn4 layers were found
in the Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples
which were thermally aged at 553 K (see Fig. 3). It is worth
noting that by careful observation, in Sn–3.5Ag/Ni samples
after thermal aging for 10 h at 523 K and 553 K, both faceted surfaces and round edge surface Ni3Sn4 scallops
appeared to a greater or lesser extent (see Figs. 2c and
3c). This phenomenon should be regarded as a transition
from round edge surface Ni3Sn4 particles to faceted surface
Ni3Sn4 particles during their growth.
3.2. Growth of Ni3Sn4 layers
Generally, there are three methods to obtain the average
thickness of the IMC layers in solder joints. The first
method (method I) is to determine the average thickness
of IMC layers by calculating the arithmetical mean thickness through measuring the layer thickness at several
equally spaced points. This method was not applied in
our tests since it will result in a relatively large error due
to the ragged edge of the solder/Ni3Sn4 interface. The second method (method II) to achieve the average thickness of
IMC layers is by using software to measure the integrated
area and the length of IMC layers first. Then, the average
thickness of IMC layers may be calculated by dividing the
integrated area by the length of the IMC layers. Since the
interfaces of the Ni/Ni3Sn4 were flatter than the interfaces
of the solder/Ni3Sn4 in the IMC layers (see Figs. 1–3), a
third method (method III) was developed to obtain the
average thickness of the IMC layers by measuring the consumption of the Ni plates. From the volume transformation during the interfacial reaction, the average thickness
of the Ni3Sn4 layers may be calculated as follows:
hNi3 Sn4 ¼
1 qNi
hNi
fNi qNi3 Sn4
ð1Þ
where hNi3 Sn4 is the average thickness of the Ni3Sn4 layers,
hNi is the average thickness of the Ni plates consumed, fNi
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J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
Fig. 3. SEM micrographs of cross-sections of IMC layers formed in samples thermally aged at 553 K: (a) Sn–3.5Ag/Ni solder joint thermally aged for
5 min, (b) Sn–3.5Ag/Ni solder joint thermally aged for 6 h, (c) Sn–3.5Ag/Ni solder joint thermally aged for 10 h, (d) Sn–3.5Ag–0.75Ni/Ni solder joint
thermally aged for 5 min, (e) Sn–3.5Ag–0.75Ni/Ni solder joint thermally aged for 6 h, and (d) Sn–3.5Ag–0.75Ni/Ni solder joint thermally aged for 10 h.
Fig. 4. X-ray diffractograms of the interfacial reaction product layers in
the (a) as-reflowed and (b) thermally aged Sn–3.5Ag/Ni samples (553 K,
10 h). (Note the IMC layers are so narrow that when the X-ray beam is
put over the layers, it takes in data from around the layers, such as the
solder matrices.)
is the weight fraction of Ni in the Ni3Sn4 IMC, qNi and
qNi3 Sn4 are the densities of pure Ni (8.9 g cm3) and Ni3Sn4
(8.64 g cm3 [26]), respectively.
It should be stressed that in this equation, the solubility
of Ni in the eutectic Sn–3.5Ag solder was neglected (since
the solubilities of Ni in the eutectic Sn–3.5Ag solder at
523 K and 553 K are very small; compared with the consumption by the interfacial reaction, it is negligible) and
the total Ni3Sn4 layer formed in a solder joint was regarded
as arising from the consumption of the Ni substrate. So, in
theory, the average thickness of the Ni3Sn4 layers calculated by methods II and III should be the same in the
Sn–3.5Ag/Ni samples. However, because the Ni atoms
contained in the solder alloy reacted with Sn atoms to form
more Ni3Sn4 which adhered to the Ni3Sn4 layer and
increased the thickness of Ni3Sn4 layer, there must be a difference between the value of average thicknesses of Ni3Sn4
layers in Sn–3.5Ag–0.75Ni/Ni samples determined by
methods II and III.
J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
Using methods II and III, the average thickness of Ni3Sn4
layers in Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples
were determined and are given in Fig. 5. Fig. 5a shows plots
of the Ni3Sn4 layer thicknesses against the number of reflow
cycles in Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples. It
can be seen that, calculated by method II, the thickness of
Ni3Sn4 layers in the Sn–3.5Ag–0.75Ni/Ni sample was larger
than that in the Sn–3.5Ag/Ni sample regardless of the number of reflow cycles, and their thicknesses did not change dramatically as a function of the number of reflow cycles. Using
method III, the values of the thicknesses of Ni3Sn4 layers in
the Sn–3.5Ag–0.75Ni/Ni samples are close to those of the
Sn–3.5Ag/Ni samples.
Fig. 5b shows plots of the thickness of Ni3Sn4 layers in
Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples during
thermal aging at two different temperatures as a function
of the aging time (the values of thickness were calculated
by method II). It can be seen that in both the Sn–3.5Ag/
Ni and Sn–3.5Ag–0.75Ni/Ni samples, the thickness of
Ni3Sn4 layers increased with an increase of the thermal
aging time at the two test temperatures. Moreover, the
5201
thickness of Ni3Sn4 layers in the Sn–3.5Ag–0.75Ni/Ni samples was always larger than that for the Sn–3.5Ag/Ni samples. This is because the Ni atoms in the solder alloy
reacted with Sn atoms directly at the interface of the solder/Ni3Sn4 layer without a requirement of long distance
diffusion and formed additional Ni3Sn4, which contributed
to the value of the thickness of the Ni3Sn4 layer when measured by method II. However, when measured by method
III, the plots of the thickness of Ni3Sn4 layers at two different temperatures as a function of aging time indicated that
the thickness of Ni3Sn4 layers in the Sn–3.5Ag–0.75Ni/Ni
sample was only slightly smaller than that in the Sn–
3.5Ag/Ni sample (as shown in Fig. 5c).
In order to clarify the growth kinetics of the Ni3Sn4 layers in the Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni samples
during thermal aging, an empirical power-law relationship
was used as follows:
ht h0 ¼ ktn
ð2Þ
where ht is the average thickness of the Ni3Sn4 layer at time
t, h0 is the initial thickness (i.e., at the aging time of t = 0),
Fig. 5. Variation of thickness of Ni3Sn4 layers as a function of (a) number of reflow cycles, (b) thermal aging times (the values of thickness were calculated
by method II) and (c) thermal aging times (the values of thickness were calculated by method III).
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J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
k is the growth rate constant, and n is the time exponent.
The values of k and n, for a particular aging temperature,
can be obtained by means of multivariable linear regression
analysis. In our study, a linear regression analysis of the
average thickness of the Ni3Sn4 layer as a function of thermal aging time was conducted to determine the best-fit n
values for the two aging temperatures and the equation
was rewritten as:
log ðht h0 Þ ¼ log k þ n log t
ð3Þ
where the time exponent n actually equals the slope of the
log ðht h0 Þ, vs log t plot for each temperature.
The calculated values of the thicknesses of Ni3Sn4 layers
by method III reflect the actual growth process of the
Ni3Sn4 layers in Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni
samples by an interfacial diffusion reaction and only the
data from Fig. 5c was used for log plotting (see Fig. 6).
The value of the time exponent n of the Sn–3.5Ag/Ni sample at 523 K was 0.32 ± 0.013, while it was 0.43 ± 0.011 at
553 K. The value of the time exponent n of the Sn–3.5Ag–
0.75Ni/Ni sample was 0.30 ± 0.020 at 523 K and was
0.42 ± 0.020 at 553 K, which are slightly smaller than that
in the Sn–3.5Ag/Ni sample at the corresponding temperatures. Meanwhile, the activation energy, Q, for the growth
of Ni3Sn4 in the Sn–3.5Ag/Ni and Sn–3.5Ag–0.75Ni/Ni
samples, was calculated from an Arrhenius relationship
of the growth rate:
k ¼ k 0 eðQ=RT Þ
ð4Þ
where R is the gas constant (8.314 J mol1 K1) and k 0 is a
pre-exponential coefficient. The activation energies were
calculated from the slope of the Arrhenius plots using a linear regression model and the values found were
10.7 kJ mol1 (Sn–3.5Ag/Ni sample) and 11.6 kJ mol1
(Sn–3.5Ag–0.75Ni/Ni sample), which are close to the values reported by others [9,15].
Fig. 6. Log plot of the growth of Ni3Sn4 layers in Sn–3.5Ag/Ni and Sn–
3.5Ag–0.75Ni/Ni samples during thermal aging tests (the thicknesses of
Ni3Sn4 layers were measured by method III).
4. Discussion
In a simple model, analysis of the growth kinetics of the
Ni3Sn4 layer starts with the simplest case, that is, the
growth of a solid layer between the molten solder and plate
which forms only one IMC phase (Ni3Sn4) and the interfacial reaction takes place in accordance with the reaction
equation: 3Ni + 4Sn = Ni3Sn4. On the assumption that
the Ni3Sn4 formed is a parallel-plate layer whose thickness
is the same over the entire surface of contact of the Sn/
Ni3Sn4/Ni and the length of the Ni3Sn4 layer in the direction normal to the direction of diffusion of Sn/Ni interface
is considerably greater than its thickness, so that the edge
effects on the growth of the Ni3Sn4 layer can be neglected,
and the growth of Ni3Sn4 can be described by a kinetic
equation [27]:
dh
k cSn
k cNi
þ
¼
dt 1 þ kkcSn h 1 þ kkcNi h
dSn
ð5Þ
dNi
where k c is a chemical constant and k d is a diffusional constant. This equation representing the growth of the Ni3Sn4
layer covers both the growth at the Sn/Ni3Sn4 interface and
the growth of the Ni3Sn4/Ni interface independently, and
the growth at each interface includes two alternative processes – the Sn and Ni atoms diffusing through the Ni3Sn4
layer (a physical process) to react with the Ni and Sn atoms
(a chemical process) to form the Ni3Sn4 layer. It is obvious
that in the very initial period the interaction of the Sn and
Ni when the Ni3Sn4 layer is very thin (several atoms in
thickness), the diffusion of Sn and Ni atoms to react with
Ni and Sn atoms can be regarded as almost instantaneous
due to the very short diffusion distance. This corresponds
to fulfilling the condition that k c k d =h in Eq. (5). Therefore, since the chemical reaction takes place at an almost
constant rate, the growth of the Ni3Sn4 layer is usually referred to as having linear kinetics (n = 1). Similarly, it is
easy to understand that after a continuous increase of the
Ni3Sn4 layer, the growth of the Ni3Sn4 layer becomes more
and more dependent on the rate volume diffusion of Sn and
Ni atoms through the bulk of the Ni3Sn4 layer, whereas the
effect of the rate of growth by the chemical reaction gradually decreases and eventually becomes insignificant. At
this stage, k c k d =h should be fulfilled in Eq. (5) and the
growth rate of the Ni3Sn4 layer by diffusion is proportional
to its existing thickness. In practice, the time dependence of
the total thickness increase of the Ni3Sn4 layer is described
by parabolic growth kinetics (n = 0.5).
The reflow test results revealed that even after only one
reflow cycle, Ni3Sn4 layers with thicknesses of about
1.33 lm were formed in both the Sn–3.5Ag/Ni and Sn–
3.5Ag–0.75Ni/Ni samples and the thicknesses of these
Ni3Sn4 layers increased very slowly during multiple reflows
(see Fig. 1). This indicates that in the very initial period of
the interaction of the molten solders and Ni plates, the Sn
and Ni atoms reacted directly without long distance diffusion to form Ni3Sn4 layers quickly following linear kinetics
J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
(although this growth process was not observed directly in
our tests because this linear kinetic growth regime of the
Ni3Sn4 layer is of a short duration at the high aging temperature due to the rapid chemical reaction of Sn and Ni
atoms. However, the research results regarding the linear
kinetic growth of Cu6Sn5 layers in a copper–tin reaction
couple which was reported by Tu and Thompson support
the growth mechanism proposed here [28]). After the initial
period, the Ni3Sn4 layer formed grew slowly with parabolic
kinetics because the diffusion rate of Sn and Ni atoms
became slower due to the thicker Ni3Sn4 layer. In this
stage, in order to observe the variation of the thickness
of the Ni3Sn4 layer easily, a relatively long high-temperature thermal aging must be performed.
The actual growth of the Ni3Sn4 layer by the Sn/Ni
interfacial reaction during soldering is more complex than
the description in the simple model above because it
involves the net effect of several interrelated phenomena,
such as diffusion through the layer via bulk (volume diffusion) and grain boundary diffusion, grain boundary grooving, grain coarsening, and dissolution into the molten
solder. For example, in the Sn–3.5Ag/Ni samples, the grain
boundaries in the Ni plate act as rapid paths for the Sn
atoms to pass through and react with Ni atoms to form
Ni3Sn4 particles, which became embedded in the Ni plate.
Fig. 1a and b shows microstructures where some Ni3Sn4
particles were separated from the main Ni3Sn4 layer and
were embedded in the grain boundaries of Ni plates. However, in the Sn–3.5Ag–0.75Ni/Ni samples, because the Ni
atoms contained in the molten Sn–3.5Ag–0.75Ni solder
reacted with Sn atoms directly and rapidly in the initial period of interaction of Sn and Ni, these formed Ni3Sn4 particles which were nucleated adjacent to the surface of the
Ni plate with a match in their crystallographic orientation
relationships, and thus impeded the Sn atoms passing
through grain boundaries to form discrete Ni3Sn4 particles
embedded in the Ni plate. The microstructure of the Sn–
3.5Ag–0.75Ni/Ni samples show that no Ni3Sn4 particles
were observed to be embedded in the Ni plates, and therefore the Ni3Sn4 layer was more uniform and ‘‘compact”
than that in Sn–3.5Ag/Ni sample (compare Fig. 1c and d).
Whether or not the Ni3Sn and/or Ni3Sn2 phases form in
the Sn/Ni interfacial reaction system is difficult to determine, although they are also thermodynamically stable
phases in the Sn–Ni binary system [17]. Based on the simple
model above, suppose the reactivity of the Ni plate with
regard to the Sn atoms remains constant, the flux of Sn
atoms through the Ni3Sn4 layer continuously decreases as
the thickness of the Ni3Sn4 layer increases with the passage
of aging time [27]. This will result in the formation of some
areas lacking in Sn atoms near the Ni plate surface and
multiphase layers may form following the chemical reaction equations: 2Sn + 3Ni = Ni3Sn2 and Sn + 3Ni = Ni3Sn. However, although many tests have been carried
out on the Sn/Ni interfacial reaction, very few reports have
been made of the existence of the Ni3Sn2 and/or Ni3Sn
phases (in our test, no Ni3Sn2 and/or Ni3Sn were detected
5203
by X-ray analysis, see Fig. 4). The variation between the
theoretical predictions and experimental results is likely
to be attributed to the kinetics of the actual growth of
the IMC layer at the Sn/Ni interface being different from
that described in the simple model. Microstructural observations show that the IMC layers formed in the Sn/Ni
interface are not whole single crystals but are in a polycrystalline form with grain boundaries (see Figs. 2 and 3). The
effect of the anisotropic growth of grains and grain coarsening led to the IMC layers showing scallop-like morphologies, where the grain boundaries and grain boundary
grooves acted as rapid diffusion paths for the Sn and Ni
atoms to react with each other. Hence, the areas supposedly depleted in Sn atoms in the simple model hardly
existed in the actual Sn/Ni reaction interface to form the
Ni3Sn2 or Ni3Sn phases. So these phases could not be
resolved by our analytical techniques (including the use
of EDX) or perhaps were not formed because of kinetic
constraints.
Since the actual Sn/Ni interfacial reaction involves several processes, the time exponent n of IMC growth kinetics
would not be expected to be a simple 0.5 following parabolic growth kinetics. Although many researchers have
reported that both the thickening and grain growth kinetics
of Ni3Sn4 in the Sn/Ni interfacial reaction followed a parabolic law in their tests [14,15,25,26], this only means the
relationship between the values of thickness and time can
be fitted to a parabolic law, but does not necessarily mean
that the growth of Ni3Sn4 layer is totally controlled by volume diffusion. In fact, more precise tests results showed
that different values of n were achieved in the growth of
the Ni3Sn4 layer. Kang and Ramachandran [12] reported
that in the temperature range 573–786 K, in the initial stage
of thermal aging, the value of n was 0.54; at an intermediate stage of thermal aging, the value of n was 0.12; and in
the final stage of thermal aging, the value of n was 0.63.
Ghosh [9] calculated the value of n by multivariant linear
regression analysis of experimental data and obtained values of n which were always slightly smaller than 1/3.
Although he stressed that, due to the limited amount of
experimental data, the confidence limit of his kinetic
parameters could not be given, a time exponent n 6 1/3 is
in good agreement with the growth model of Sn–Cu IMC
based on the effect of grain boundary diffusion, grain
boundary grooving and grain coarsening [29]. Tao et al.
have reported that the growth of the Ni3Sn4 layer at lower
thermal aging temperatures (453 K, 513 K and 573 K) gave
parabolic growth kinetics (n = 0.5), while at a high temperature (693 K), due to the two-phase mixture of
Ni3Sn4 + solder at the interface, the growth of the Ni3Sn4
layer gave linear growth kinetics (n = 1) [22].
It is meaningless to judge the precise values of the time
exponent n in the growth of Ni3Sn4 in different reports
since it is dependent on the test conditions, measuring
methods and, most importantly, it is a dynamically changing value. But it is worth clarifying the relationship
between the influencing factors and time exponent n in
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J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
the growth of Ni3Sn4. In our study, the time exponent n of
growth of Ni3Sn4 in the Sn–3.5Ag/Ni sample is
0.32 ± 0.013 at 523 K and 0.43 ± 0.011 at 553 K. The value
of n of 0.32 ± 0.013 at 523 K is in good agreement with a
mathematical model based on grain boundary diffusion
controlling the mass transport and grain coarsening reducing the availability of grain boundaries as diffusion paths as
the IMC layer grows thicker [29]. While the value of n
increased to 0.43 ± 0.011 at 553 K, this is because during
long-time high-temperature thermal aging, due to the driving force of the Gibbs–Thomson effect [30], some Ni3Sn4
particles coarsened and became separated from the main
Ni3Sn4 layer and were dispersed into the solder matrices
(see Figs. 2 and 3). This led to the situation that the main
Ni3Sn4 layer thinned out considerably and in some areas of
the Ni3Sn4/Ni surface, the Ni3Sn4 layer was too thin to
provide ‘‘protection” to the surface of the Ni plate. Hence,
Sn and Ni atoms passed through the very thin layer to react
with each other, and therefore, this increased the rate of
growth of the Ni3Sn4 layer markedly. Tao et al. have
reported a linear rate of growth of the Ni3Sn4 layer
(n = 1) at a high temperature (693 K) by this mechanism
[22]. Still, it should be stressed that if the growth kinetics
of the Ni3Sn4 layer follow a linear law this does not mean
that the growth of Ni3Sn4 is totally controlled by the chemical reaction without an influence from diffusion after longtime high-temperature thermal aging. On the contrary, this
only means that the relationship between the values of the
thickness and the time fit a linear law (with a value of n
close to 1), and that the chemical reaction dominated the
growth process (in fact, the value of n cannot be equal to
1 except in the very initial stages of thermal aging,
theoretically).
In our study, the experimental results of nonparabolic
growth kinetics of the Ni3Sn4 layer are in good agreement
with the model of the growth of the IMC layer based on
grain boundary diffusion, grain boundary grooving, grain
coarsening and dissolution into the molten solder [29].
However, the effects of these factors on the growth of the
Ni3Sn4 layer could not be observed separately, directly
and dynamically. So the addition of minor amounts of
Ni into the solder alloy was used to assist in verifying the
controlling mechanisms. When Ni was introduced into
the solder alloy, Ni3Sn4 particles formed by the direct reaction nucleated and adhered to the Ni3Sn4 layer and experienced growth to thicken the whole Ni3Sn4 layer due to the
match in their crystallographic orientation relationships.
Hence, when measured by method II, the thickness of the
whole Ni3Sn4 layer in the Sn–3.5Ag–0.75Ni/Ni sample
was larger than that in the Sn–3.5AgiNi sample (see
Fig. 5b). However, when method III was used to measure
the thickness of Ni3Sn4 layers formed only by the diffusion
reaction, it was found that the thickness of the Ni3Sn4 layer
in the Sn–3.5Ag–0.75Ni/Ni sample was slightly smaller
than that in the Sn–3.5Ag/Ni sample (see Fig. 5c). This
means the addition of minor amounts of Ni into the Sn–
3.5Ag solder suppressed the growth of the Ni3Sn4 layer
which formed by the diffusion reaction. The reason for this
result is that the Ni3Sn4 formed by the direct reaction accelerated the grain coarsening of Ni3Sn4 (see Figs. 2 and 3: relatively large Ni3Sn4 gains appeared in the Sn–3.5Ag–
0.75Ni/Ni sample after long-time thermal aging), and thus
reduced the availability of grain boundaries and grain
grooves as diffusion paths so as to decrease the growth rate
of the Ni3Sn4 layer formed by the diffusion reaction.
Hence, in this way, the effect of the influencing factors on
the growth of the Ni3Sn4 layer was proved indirectly.
Obviously, during the growth of Ni3Sn4 by the Sn/Ni
liquid/solid interfacial reaction, some of the Ni3Sn4 layer
also dissolved into the molten solder. Fig. 7 gives a TEM
bright-field image and a selected area diffraction pattern
of a Ni3Sn4 particle, which formed in the solder matrix
by solidification, showing a faceted crystal structure. The
faceted Ni3Sn4 particles formed both in the Sn–3.5Ag–
0.75Ni solder matrix and in the Sn–3.5Ag–0.75Ni/Ni interfaces. Hence, because the microstructures of Sn–3.5Ag/Ni
samples after long-time high-temperature thermal aging
show both faceted and round edged Ni3Sn4 particles (see
Figs. 2c and 3c), it is easy to understand that the Ni3Sn4
particles with faceted surfaces originated from the those
Ni atoms which were dissolved from the Ni3Sn4 layer
and re-reacted with Sn to solidify as particles. However,
it is difficult to analyze the effect of the dissolution of
Ni3Sn4 on the growth of the Ni3Sn4 layer quantitatively,
although the dissolution of Ni3Sn4 into molten solder
was proved by microstructural observations. According
to Dybkov’s research regarding the kinetics of dissolution
of a solid in a liquid by a solid/liquid interfacial reaction,
Fig. 7. Bright-field TEM image and selected area diffraction pattern of a
Ni3Sn4 particle formed in the Sn–3.5Ag–0.75Ni solder matrix showing a
faceted structure.
J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
the rate of dissolution, dc
; of the Ni3Sn4 layer in the molten
dt
solders during aging may be expressed by the equation [27]:
dc
S
¼ k ðcs cÞ
dt
V
ð6Þ
where C is the concentration of Ni in the molten solders
measured at time t, C s is the saturation concentration of
Ni in the molten solder at the aging temperature, k is the
dissolution rate constant of Ni3Sn4, S is the interfacial area
of the Ni3Sn4 layer in contact with the molten solders, and
V is the volume of the molten solder. In our tests, the rate
, of Ni3Sn4 in the molten solders could not
of dissolution, dc
dt
be calculated to evaluate the effect of dissolution on the
growth of the Ni3Sn4 layer quantitatively unless the unknown dissolution rate constant, k, was given. However,
5205
at least from the microstructural observations that no faceted Ni3Sn4 was formed in the Sn–3.5Ag solder matrix (see
Figs. 2 and 3), we can assume that during a long period of
time of thermal aging, the dissolution of the Ni3Sn4 layer
into molten solders is slow and the concentration of Ni
in the molten solders is far from the saturation concentration of Ni in these molten solders at the aging temperature.
Notwithstanding this, according to the above Eq. (6), compared with the Sn–3.5Ag solder, the higher concentration
of Ni in the Sn–3.5Ag–0.75Ni solder should decrease the
rate of dissolution of the Ni3Sn4 layer into the molten solder and increase the rate of growth of the Ni3Sn4 layer in
the Sn–3.5Ag–0.75Ni/Ni sample during aging. However,
using method III, the thickness of the Ni3Sn4 layer in the
Sn–3.5Ag–0.75Ni/Ni sample was slightly smaller than that
Fig. 8. Schematic diagrams of the growth behavior of the Ni3Sn4 layer with different values of the time exponent n with different conditions.
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J. Shen et al. / Acta Materialia 57 (2009) 5196–5206
in the Sn–3.5Ag/Ni sample. Hence, one can suppose that
without the effect of dissolution of the Ni3Sn4 layer on
the growth of the Ni3Sn4 layer, the thickness of the Ni3Sn4
layer in the Sn–3.5Ag–0.75Ni/Ni sample should be much
smaller than that in the Sn–3.5Ag/Ni sample. This proves
that grain coarsening again suppressed the growth of
Ni3Sn4.
Based on the discussion above, a clear image of the
growth mechanism of the Ni3Sn4 layer in the Sn/Ni reaction interface is given in the schematic diagram of Fig. 8.
In the very initial period of interaction of Sn and Ni,
because the Ni3Sn4 layer formed was very thin, the Sn
and Ni atoms reacted with each other instantaneously
without long range diffusion. This growth behavior of the
Ni3Sn4 layer gave linear kinetics and a time exponent
n = 1 (see Fig. 8a). After several reflow cycles or thermal
aging for several minutes (these conditions are close to
the procedures used in the microelectronic packaging in
industry), if a relatively continuous and even Ni3Sn4 layer
formed at the interface of the Sn/Ni couple, the growth
behavior of the Ni3Sn4 layer gave parabolic kinetics and
a time exponent n close to 0.5 due to volume diffusion dominating the diffusion process of Sn and Ni atoms (see
Fig. 8b). However, usually, this is not the case in the Sn/
Ni reaction interface since grain boundary diffusion, grain
boundary grooving, grain coarsening and dissolution into
the molten solder also occurred during the growth of
Ni3Sn4. When grain boundary diffusion and grain boundary grooving dominated the diffusion process of Sn and
Ni atoms, the growth rate of the Ni3Sn4 layer increased
and the time exponent n increased with a value between
0.5 and 1 (see Fig. 8c). By contrast, the grain coarsening
of Ni3Sn4 should reduce the availability of grain boundaries and grooves as diffusion paths so that it should slow
down the growth of the Ni3Sn4 layer (see Fig. 8d). With
these conditions, according to the mathematic model [29],
the time exponent n should decrease to be near 1/3 (the
value of the time exponent achieved in our tests by thermal
aging the Sn–3.5Ag/Ni sample at 523 K is 0.32 ± 0.013,
which is good agreement with the model). After long-time
high-temperature thermal aging, due to grain coarsening
and the Gibbs–Thomson effect, some Ni3Sn4 particles
became separated from the main Ni3Sn4 layer, and therefore this resulted in the formation of some bare interfaces
at the main Ni3Sn4 layer. In this situation, Ni3Sn4 grew
rapidly in these bare areas and the overall thickness of
the Ni3Sn4 layer increased sharply. Then, the time exponent n increased to be close to 1. In practice, some researchers have called this linear growth kinetics.
5. Conclusions
With the help of a classical simple model of the liquid/
solid interfacial reaction, the growth mechanisms of Ni3Sn4
in actual Sn–3.5Ag/Ni couples were investigated and dis-
cussed. An introduction of a minor addition of Ni into
the solder alloy changed the growth process of Ni3Sn4 at
the Sn–3.5Ag–0.75Ni/Ni reaction interface and proved
the actual growth of Ni3Sn4 involves the net effect of several interrelated phenomena, such as diffusion through
the layer via bulk (volume) diffusion and grain boundary
diffusion, grain boundary grooving, grain coarsening, and
dissolution into the molten solder. The time exponent n
and morphology of Ni3Sn4 were dependent on the relative
values of these factors, i.e., which factor/factors dominated
the growth process. Notwithstanding this, reflow tests
showed that the slow growth rate of Ni3Sn4 in industrial
soldering applications is desirable since a brittle and thick
Ni3Sn4 layer will severely weaken the solder joints.
Acknowledgments
The authors would like to acknowledge the financial
support provided by an RGC Competitive Earmarked Research Grant (Project No. 9041222, CityU.111307). Special
thanks to Prof. B. Ralph in Brunel University for his cooperation in this study.
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