Inhibition of Cracking in Cu6Sn5 Intermetallic Compounds at Sn-Cu Lead-Free
Solders and Cu Substrate Interfaces
K. Nogita1, S. Suenaga2, S. D. McDonald1, H. Tsukamoto1, J. Read1 and T. Nishimura2
Materials Engineering, The University of Queensland, St. Lucia, Brisbane QLD, 4072 Australia
2
Nihon Superior Co. Ltd., NS Bldg., 1-16-15 Esaka-Cho, Suita City, Osaka, 564-0063 Japan

1

Abstract
The authors have successfully developed a method of limiting cracking in the Cu6Sn5 intermetallic
compounds, which are formed at the interface between lead free solders and their Cu substrates. In thispaper, it is shown that the level of cracking in the (Cu,Ni)6Sn5 reaction layer and the morphology of theintermetallics varies with the Ni content. To explore the mechanisms associated with the differences injoint integrity, detailed synchrotron X-ray powder diffraction (XRD) has been used with Rietveld analysis.
The results show Ni stabilises a high temperature allotrope of the Cu6Sn5 phase avoiding stresses induced
by a volumetric change that would otherwise occur on phase transformation .

Introduction

Cracking measurements

The nature of cracking in the reaction layer present in
two types of samples, ball grid array (BGA) soldered anddipped Cu plate interfaces, has been investigated. Two
compositions of lead-free solder were used, Sn-0.7wt%Cu(SC) and Sn-0.7wt%Cu-0.05wt%Ni (SCN). BGA typesamples were obtained from 500µm diameter solder ballson organic solderability protection (OSP) Cu boards with acommon reflow process. Figure 1 shows the reflow profile
for the BGA process. This cycle was conducted twice foreach BGA specimen.
Temperature , T (oC)In the last few decades significant research has been
directed at developing lead-free solders for the electronics
industry(1). One attractive group of lead-free solder alloys
are those based on the Sn-Cu-Ni system which have foundapplications in wave, dip and iron soldering processes andare cheaper than most other candidate alloys due to theomission of expensive elements such as silver or rareearths(2). The intermetallic Cu6Sn5 that forms between
Sn-Cu lead-free solder alloys and Cu substrates influences
the soldering process and the subsequent joint properties.
Cu6Sn5 exists in two crystal structures with an allotropictransformation from monoclinic η'-Cu6Sn5 at temperatures
lower than 186ºC to hexagonal η-Cu6Sn5. In conventional
lead-free solder alloys the brittle Cu6Sn5 intermetallic isoften associated with the presence of a large number ofmicrocracks which may form in response tothermo-mechanically induced stresses associated with thephase transformation. The authors have successfully
developed a method of limiting cracking in the Cu6Sn5intermetallic compounds (IMC) by stabilising the hightemperature hexagonal Cu6Sn5 phase with Ni additions(3).
The phase stability induced by Ni being present in the
Cu6Sn5 phase and its possible implications are discussedwith reference to the experimental results.

Experimental procedure

250
200
150
100
50
0
0

50

100

150

200

250

300

350

Time, t (sec)Figure 1. Temperature profile of a reflow process.

1

Dipped Cu type samples were obtained by dipping a Cu
plate (C1220P), 10mm x 30mm x 0.3mm, with common
flax into a molten solder held at in 250°C, 10 sec. Selected
samples were subsequently annealed at 120°C for 1000
hours. All samples were embedded in resin and polished incross-sectioned to observe the solder joint IMC using
SEM/EDS. The microstructures of the samples werestudied using SEM (JEOL JSM-6460LA, Japan) with
20kV in secondary image mode. The crack length andcrack number of IMC was measured by digital imaging
using commercial SEM image analysis software, JEOL
AnalysisStation. The crack number and crack length werenormalised by the total measured length of the solder joint,
which was typically close to 300µm. EDS point elemental
analysis of Sn, Cu and Ni in the IMC was conducted by
the average of five measurements.
Phase stabilisation
To obtain Cu6Sn5 and (Cu,Ni)6Sn5 intermetallic particles
for detailed experimentation, two types of samples (as castsample ‘A’ and annealed ‘B’) were produced using a direct
alloying method. High purity copper, tin (and nickel)
ingots were placed in a clay-graphite crucible and heated
to 700ºC in an electric resistance furnace under Ar gas
flow. The melt was held at this temperature for aminimum of two hours. Using a stainless steel mould of
15mm diameter and 10mm height, approximately 15g of
the alloy was cast from the melt and allowed to solidify
while cooling in air. The solid sample was then annealed
in a heat treatment furnace at 400 ºC for 528 hours andcooled to room temperature (designated series ‘A’) andselected samples were then subjected to a subsequent
treatment of 180 ºC for 168 hours (designated series ‘B’).
Chemical analysis confirmed we obtained Cu6Sn5 andCu5.5Ni0.5Sn5 samples for both A and B series. The
chemical analysis data of Cu6Sn5 and (Cu,Ni)6Sn5 areshown in Table 1.
Table 1. Chemical composition of XRD samples.
No.

Samples

A1 Cu6Sn5

Sample annealing
conditions
400°C 528hr

crystal structure and lattice parameters were carried out at
the BL15 Kyushu synchrotron, SAGA-LS. The XRD
spectra for the Rietveld structure refinement were obtainedin the 2θ range of 20 to 75 degree after scanning from film
exposed in a Debye-Scherrer type XRD machine.
Rietveld analysis of the XRD signal peak intensities, peak
profiles and peak positions were carried out using the
RIETAN-2000(4) computer package. Standard silicon(NIST640C) powder, National Institute of Standards, wasused for the calibration of the Rietveld analysis withidentical experimental conditions used for the samples.
For the refinement the parameters were optimized tominimize the residual Rwp, Rp and S factors(4).
Throughout the determination of the lattice parameters
using structure refinement of the samples, V and W, which
are both machine dependent parameters, and B, anisotropic displacement parameter, were fixed with thevalues obtained from the standard silicon powder asshown in Table2
Table 2 Refined Rietveld parameters for standard Si
Gaussian FWHM parameter, U
Gaussian FWHM parameter, V
Gaussian FWHM parameter, W
Scherrer coefficient for Gaussian broadening, P
Lorentzian Scherrer broadening, X
Anisotropic Scherrer broadening, Xe
Strain broadening, Y
Anisotropic strain broadening, Ye
Preferred-orientation parameter, rLattice parameter, aLattice parameter, bLattice parameter, cLattice parameter, αLattice parameter, βLattice parameter, γOverall isotropic displacement parameter, Q
Silicon atom Occupation factor, gFractional coordinate, x
Fractional coordinate, yFractional coordinate, z
Isotropic displacement parameter, B
R wp
Rp
Rb
S

1.54E-03-8.67E-04
3.31E-04
0
1.77E-02
0
7.69E-03015.431195.431195.43119
90
90
90
0
1
0
0
0
0.58
12.74
8.00
2.85
1.09

Results and discussion

Cmpostions (at%)
Cu
Ni
Sn
54.40
0.00 45.60

A2 Cu5.5Ni0 .5 Sn6B1 Cu6Sn5

400°C 528hr

50.12

4.58

45.30

400°C 528hr + 180°C 168hr

54.40

0.00

45.60

B2 Cu5.5Ni0 .5 Sn6

400°C 528hr + 180°C 168hr

50.12

4.58

45.30

The sample was crushed in an agate mortar to obtainpowder for the XRD experiments. XRD studies on the

Observation and quantification of cracking at thesoldered interface
Figure 2 to 4 show SEM micrographs of cross-sectioned
BGA, dipped Cu type samples and dipped Cu samples
with subsequent 120°C annealing of (a) SC and (b) SCN,
respectively. Visible cracks can be seen in the intermetallic

2

compound (IMC) layer between the solder alloy and theCu substrate on (a) SC alloy but not or few in (b) SCN
alloys where Ni is present.
(a)

Sn07CuCu6Sn 5Cu
10 µm

layer may be present between the Cu substrate and the
Cu6Sn5 IMC in all samples(5), however, this must be verythin and could not recognized in the current study at thelimits of SEM magnification. Table 3 shows EDS
elemental analysis results (fine average point) of the IMC
layer. The chemical composition of the IMCs are Cu6Sn5or (Cu,Ni)6Sn5. It is obvious that the IMC in SCN samples
of BGA and dipped Cu contain about 5at% Ni to form
(Cu,Ni)6Sn5. Since the only source of Ni is in the solder for
SCN samples, it is expected that the 5at% Ni migrate fromthe solder during growth of the IMC layer.
(a)
Sn07Cu(b)
Cu6Sn5

Sn07Cu005Ni(Cu,Ni)6Sn5

Cu

Cu
10 µm
Figure 2. Cross-sectioned SEM image of BGA solder
structures of (a) SC and (b) SCN.

(b)
Sn07Cu005Ni(Cu,Ni)6Sn 5

Table 3 Chemical compositions of IMC

Sam plesCu (at%) Ni (at%) Sn (at%)
BGA-S C58
042
BGA-SCN
51
544
Dip-S C57
043
Dip-S CN52
543

It is noted that all cracks formed in IMC are nearparallel to the Cu substrate, but none are perpendicular orpresent in other directions. This tendency may be due tothe growth direction of IMC as determined by theinterface crystallography or thermal conditions combinedwith a preferred habit plain for cracking, or otherwise
related to the direction of stresses during thermal
expansion . Further investigations on the crack orientation
are required. The IMC present in this layer is commonly
known to be Cu6Sn5 and it is apparent significantly fewer
cracks are present in the Cu6Sn5 layer when Ni is present
in the Sn-0.7Cu solder. According to the literature a Cu3Sn

Cu

Figure 3. Cross-sectioned SEM image of dipped Cu
sample of (a) SC and (b) SCN.

Figure 5 to 7 indicate the normalized crack length andcrack number in IMCs of BGA, dipped Cu and dipped Cu
with subsequent annealing samples, respectively. The
IMC layers in SCN solders, which contain around 5 at%
Ni, experience less cracking than those without Ni in all
solder conditions and there are no cracks present in dipped
Cu samples of SCN solder. Therefore, by including Ni in
the solder cracking can be inhibited at the solder joint by in
the IMC layer. In the next section, we demonstrate themechanism of crack inhabitation is likely to be related tophase stabilisation by using synchrotron powder XRD for
Cu6Sn5 and (Cu,Ni)6Sn5 with 5 at%Ni samples.

3

Cu6Sn 5

Normalised crack number
and length

(a)

0.3
Crack number
Crack length

0.2

0.1

0
SC

SCN

Samples

Cu

Figure 7 Normalised crack length and number in IMC ofdipped Cu sample with annealing.

(b)

Crystallography and phase stability of IMCs

(Cu,Ni)6Sn 5

Cu

Normalised crack number
and length

Figure 4. Cross-sectioned SEM image of dipped Cu with
120C annealed sample of (a) SC and (b) SCN.
0.5
Crack number
Crack length

0.4
0.3
0.2
0.1
0
SC

Samples

SCN

Normalised crack number
and length

Figure 5 Normalised crack length and number in IMC of
BGA
0.3
Crack number
Crack length

0.2

0.1

0
SC

SCN

SamplesFigure 6 Normalised crack length and number in IMC ofdipped Cu sample

From the reported phase diagram(5), Cu6Sn5 exists in two
crystal structures with an allotropic transformation frommonoclinic η'-Cu6Sn5 at temperatures lower than 186ºC tohexagonal η-Cu6Sn5. Ambient temperature theoreticaldensities of monoclinic and hexagonal are evaluated as
8.270 g/cm3 and 8.448 g/cm3, respectively(6). This means if
the IMC makes a phase transformation from the‘quenched’ hexagonal phase to monoclinic phase during
operation, there will theoretically be a 2.15% volume
expansion, which would conceivably be accompanied bysignificant stress in the IMC layer. We assumed the lower
cracking in the SCN compositions are due to the
avoidance of the phase transformation and the associatedvolume change when the hexagonal phase is stabilized by
Ni. Table 4 show fractional coordinates for the Rietveld
analysis of RIETAN-2000(4) of (a) hexagonal η-Cu6Sn5(P63/mmc) and (b) monoclinic η'-Cu6Sn5 (C2/c). There arevery small differences in of the X-ray diffraction peaks
between the two except for the strong peak height ratio
and the presence of several relatively weak peaks in the 2
theta range of 25-50°(6). Therefore, strong synchrotron
X-ray was required to distinguish between the two.
Table 4. Unit cell-external parameters of (a) hexagonal
η-Cu6Sn5 (P63/mmc) and (b) monoclinic η'-Cu6Sn5 (C2/c).
(a) P63/mmc
Atom
Sn1
Cu1
Cu2
(b) C2/c
Atom
Cu1
Cu2
Cu3
Cu4
Sn1
Sn2
Sn3

Site
2(c)
2(a)
2(d)

Occ
1
1
0.2

x
0.333
0.000
0.333

y
0.667
0.000
0.667

z
0.250
0.000
0.750

Site
8(f)
8(f)
4(a)
4(e)
8(f)
8(f)
4(e)

Occ
1
1
1
1
1
1
1

x
0.101
0.306
0.000
0.000
0.391
0.285
0.000

y
0.473
0.504
0.000
0.160
0.162
0.655
0.799

z
0.202
0.610
0.000
0.250
0.529
0.358
0.250

4

101

(a)

110

102

Counts (a.u.)

202

A2
C2

201

002

112

103

211

A1
C1

20

30

40

50

60

70

A2

Counts (a.u.)Counts (a.u.)

A2
B2

C1
A1

32

34

36

38

Angle (°)

40

42

201

002

112

103

202

211

B1
A1

20

30

40

A2
B2

50

60

70

002

A1
B1

30

32

34

36

38

40

42

44

Angle (°)

Figure 8. XRD peaks of (a) A1 (Cu6Sn5) and A2(Cu5.5Ni0.5Sn5), and (b) B1 (Cu6Sn5) and B2 (Cu5.5Ni0.5Sn5).
Doted area between 30-45° in 2θ is magnified.

Table 5 shows the Rietveld analysis results of annealed
Cu6Sn5 (B1) and Cu5.5Ni0.5Sn5 (B2) samples based onhexagonal (P63/mmc) and monoclinic (C2/c) structure
settings. Comparing the refinement parameters of Rwp,
Rp and S(4) of hexagonal and monoclinic settings (lower
values for these parameters indicate a better match) the
Cu6Sn5 sample (B1) is a better fit for the monoclinic(C2/c) phase than for the hexagonal phase (P63/mmc). It is
expected that not only small monoclinic peaks but theintensity ratio of each strong peak are a better fit formonoclinic refractions than that of hexagonal. In contrast
the Cu5.5Ni0.5Sn5 sample (B2) is a better fit for the
hexagonal phase. In conclusion, it is clear that the Cu6Sn5
sample with annealing at 180°, which is just below thereported phase transformation temperature, shows amonoclinic structure but does not show any of thehexagonal high temperature phase. This result is inagreement with previously reported results(3).

B1

Crystallographic
setting
Monoclinic

B1

Hexagonal

7.35

4.63

3.18

B2

Monoclinic

7.74

4.81

2.45

B2

Hexagonal

5.42

3.78

1.72

Samples30

110 102

Table 5. Rietveld analysis results of annealed Cu6Sn5 (B1)and Cu5.5Ni0.5Sn5 (B2)

002

C2

101

(b)

Counts (a.u.)

To compare the effects of Ni on the crystallography, the
XRD pattern obtained from A1 (Cu6Sn5) and A2(Cu5.5Ni0.5Sn5), and Bl (Cu6Sn5) and B2 (Cu5.5Ni0.5Sn5),
samples are shown in Figure 8a and 8b. XRD peaks
clearly shows a hexagonal Cu6Sn5 structure in all Ni
containing samples present as Cu5.5Ni0.5Sn5 (A2 and B2)
but no monoclinic phase was detected, while all samples
without Ni (A1 and B1) clearly indicate small peaks,
which are expected from the monoclinic phase. The doted
area of the figures between 30 to 45 º in 2θ is magnified,and small extra peaks which are not present in hexagonaldiffractions are indicated with arrows. Intensities of weak
monoclinic peaks are higher in B1 sample than A1 sample.
This suggests sample A1 (non-annealed at 180°C sample)
may contain a mixture of hexagonal and monoclinic as aresult of being ‘quenched’ from 400°C. XRD analysis ofmonoclinic Cu6Sn5 without Ni present has been reportedby Ghosh and Asta(6) using samples annealed at 150°C for
80 days. They reported several relatively weak peaks inthe 2θ range of 25 - 50º clearly in monoclinic Cu6Sn5,
which could not be found in the present (Cu,Ni)6Sn5
samples both with and without annealing. It is concluded
that in the current study for this condition, around 5 at% of
Ni in Cu6Sn5 stabilised the high temperature hexagonalphase. Furthermore, cracking of IMCs in SC samples
shown in Figures 2a, 3a and 4a are expecting to result
from phase transformation from hexagonal to monoclinic
during cooling. Also, importantly, it was noted that thedegree of cracking in the SC samples and the potential for
more cracking during loading is likely to depend on thecooling rate related to phase transformation kinetics andtemperature, since it is expected that the transformation isincompleted.

44

R wp
6.81

Rp
4.42

S
2.95

5

Table 6. Lattice parameters of annealed Cu6Sn5 (B1) andCu5.5Ni0.5Sn5 (B2)
Lattice parameters (Å)

Samplesaa(STD)cc(STD)

B1

4.223

0.002

5.113

0.002

B2

4.197

0.002

5.103

0.001

To analyse how lattice parameters change with Nicontent, we used the hexagonal crystallographic data of
P63/mmc for all samples. As shown in Table 6, both ‘a’axis and ‘c’ axis decrease their lattice parameters with Nicontents, however, the ‘a’ axis displays a larger decrease
than that of the ‘c’ axis. Recently, Yu et. al.(7) reported thestructural and electrical properties of Cu6-xNixSn5 (withx=0, 1,2) from first-principles calculations. Assuming thatthe low temperature phase was monoclinic η'-Cu6Sn5, they
concluded that, as the Cu sites are periodically substituted
by Ni on Cu2 or 8f sites, the volume of the unit cells
shrunk and the distance between atoms shortened.
Based on energy and density of states calculations they
demonstrated that stoichiometric ternary Cu6-xNixSn5
compounds are more stable than Cu6Sn5 and that
Cu4Ni2Sn5 is the most stable phase. The measured
compositions of IMCs in Table 1 of this study show that
the IMCs are close to Cu5.5Ni0.5Sn5. Present results of
XRD show a systematic ‘shrinking’ of the lattice
parameter (or volume per unit cell), must be due to Ni
atoms in the lattice.

Conclusions
A method of limiting cracking in the Cu6Sn5 intermetallic
compounds found at the interface of lead free solder alloys
and their substrates has been successfully developed byadding small amount of Ni to a Sn-Cu lead-free solder.
The intermetallic layers in Sn-Cu-Ni solders experience
less cracking than those without Ni present. Synchrotron
XRD was used to show that the stabilisation of thehexagonal (Cu,Ni)6Sn5 occurs in when Ni-containing
solder alloys are used. This phase stabilization may
prevent volume changes that could contribute to thecracking process in the IMC layer in Ni-free alloys where
the hexagonal-monoclinic transformation occurs at
approximately 186 ºC.

Acknowledgment
This research had been conducted under an international
cooperative research program between the University of
Queensland, Australia and Nihon Superior Company,
Japan. XRD experiments were performed at the BL15 in
the SAGA Light Source (SAGA-LS) under the
Nanotechnology Support Project of the Ministry of
Education, Culture, Sports, Science and Technology,
Japan (project ID: 080625N). The author thanks Dr. T.
Yamashita at Mesaplexx Pty. Ltd. for Rietveld analysis
and Drs R. Ohtani and K. Sumitani of SAGA-LS for
valuable discussions and XRD experiments.

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