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IEEE TRANSACTIONS ON COMPONENTS, PACKAGING AND MANUFACTURING TECHNOLOGY, VOL. 3, NO. 6, JUNE 2013
High Temperature Interconnect and Die Attach
Technology: Au–Sn SLID Bonding
Torleif André Tollefsen, Andreas Larsson, Ole Martin Løvvik, and Knut E. Aasmundtveit
Abstract— Au–Sn solid–liquid interdiffusion (SLID) bonding
is a novel and promising interconnect and die attach technology
for high temperature (HT) applications. In combination with
silicon carbide (SiC), Au–Sn SLID has the potential to be a
key technology for the next generation of HT electronic devices.
However, limited knowledge about Au–Sn SLID bonding for
HT applications is a major restriction to fully realizing the HT
potential of SiC devices. Two different processing techniques—
electroplating of Au/Sn layers and sandwiching of eutectic Au–
Sn preform between electroplated Au layers—have been studied in a simplified metallization system. The latter process
was further investigated in two different Cu/Si3 N4 /Cu/Ni–P/Au–
Sn/Ni/Ni2 Si/SiC systems (different Au-layer thickness). Die shear
tests and cross-sections have been performed on as-bonded,
thermally cycled, and thermally aged samples to characterize
the bonding properties associated with the different processing
techniques, metallization schemes, and environmental stress tests.
A uniform Au-rich bond interface was produced (the ζ phase with
a melting point of 522 °C). The importance of excess Au on both
substrate and chip side in the final bond is demonstrated. It is
shown that Au–Sn SLID can absorb thermo-mechanical stresses
induced by large coefficient of thermal expansion mismatches (up
to 12 ppm/K) in a packaging system during HT thermal cycling.
The bonding strength of Au–Sn SLID is shown to be superb,
exceeding 78 MPa. However, after HT thermal ageing, the ζ phase
was first converted into the more Au-rich β phase. This created
physical contact between the Sn and Ni atoms, resulting in brittle
Ni x Sn y phases, reducing the bond strength. Density functional
theory calculations have been performed to demonstrate that the
formation of Ni x Sn y in preference to the Au-rich Au–Sn phases
is energetically favorable.
Index Terms— Au–Sn solid–liquid interdiffusion (SLID)
bonding, density functional theory (DFT), die attach, high
temperature (HT), interconnect technology.
Manuscript received April 3, 2012; revised December 13, 2012; accepted
March 9, 2013. Date of publication May 3, 2013; date of current version
May 29, 2013. This work was supported in part by the HTPEP Project, the
Research Council of Norway under Project 193108/S60, Badger, SmartMotor,
Fairchild, Roxar, and Norbitech. Recommended for publication by Associate
Editor T.-C. Chiu upon evaluation of reviewers’ comments.
T. A. Tollefsen is with SINTEF ICT Instrumentation, Oslo 0373, Norway,
and also with Vestfold University College, Institute for Micro and Nanosystems Technology, Borre 3184, Norway (e-mail: torleif.tollefsen@sintef.no).
A. Larsson is with SINTEF ICT Instrumentation, Oslo 0373, Norway
(e-mail: andreas.larsson@sintef.no).
O. M. Løvvik is with SINTEF Materials and Chemistry, Oslo 0314, Norway,
and also with the Department of Physics, University of Oslo, Oslo 0318,
Norway (e-mail: olemartin.lovvik@sintef.no).
K. E. Aasmundtveit is with Vestfold University College, Institute
for Micro and Nanosystems Technology, Borre 3184, Norway (e-mail:
knut.aasmundtveit@hive.no).
Color versions of one or more of the figures in this paper are available
online at http://ieeexplore.ieee.org.
Digital Object Identifier 10.1109/TCPMT.2013.2253353
I. I NTRODUCTION
M
ICROELECTRONIC packaging plays a vital role in
electronic devices. It serves the purposes of electrical interconnection, heat dissipation, mechanical support, and
physical protection [1]. The electrical performance, size, cost,
and reliability is to a large degree governed by the package,
which is often referred to as the bottleneck of microsystem
industry [2]. The choice of interconnection, i.e., the conductive
path required to achieve connection from a circuit element to
the rest of the circuit, is therefore of the utmost importance.
Commonly used interconnect techniques include solders
and conductive adhesives [3]. However, for high temperature
(HT) applications like automotives, drilling and well intervention systems, aerospace, space exploration, and nuclear
environments, the standard interconnect materials do not meet
the requirements regarding, e.g., HT stability [4]. Currently,
there is no clear definition of the temperature range of a HT
electronic system. In this paper, it is considered to be above
200 °C.
There is a limited range of HT interconnect techniques
[4]–[6]. One alternative is sintered nano-particle Ag, which
has good electrical and thermal conductivity [7], [8]. A nanoparticle Ag joint has a high melting point (960 °C) compared
to the low processing temperature (<300 °C) [7]. However,
Ag migration is reported to be a problem in HT applications
(particularly in combination with high power), limiting the
lifetime of the joint [9], [10].
Other prospective HT interconnect techniques include
liquid-based solder joints [11], composite solder joints [12],
bismuth-based solder joints [13], and solid–liquid interdiffusion (SLID) joints [14], [15]. Of these techniques,
SLID bonding—also called transient liquid phase bonding
[16], [17], isothermal solidification [18], or off-eutectic bonding [19]—has shown great potential [9], [19]–[21]. SLID
bonding utilizes a binary system with one HT melting metal
and one low temperature melting metal (general principles
are shown in Fig. 1). The applied processing temperature
is higher than the melting point of the low melting point
metal, and new intermetallic compounds (IMCs) are formed.
The solidification is isothermal, and the final joint has a
higher melting point than the processing temperature. This
opens a window for new subsequent manufacturing steps
without the need for ever decreasing process temperatures
for each step [14], [15]. Another important advantage with
SLID bonding is that often a processing temperature in close
proximity to the final application temperature can be applied
2156-3950/$31.00 © 2013 IEEE
TOLLEFSEN et al.: HIGH TEMPERATURE INTERCONNECT AND DIE ATTACH TECHNOLOGY
Fig. 2.
Fig. 1. Schematic illustration of SLID bonding (TB bonding temperature, Tm
melting temperature, RT room temperature, IMC intermetallic compound).
(since the melting point of the final joint is much higher
than the processing temperature). This can help reduce the
thermo-mechanical stresses induced by coefficient of thermal
expansion (CTE) mismatches in the final package.
SLID bonding was performed in various metal systems.
Examples include Ag/In [14], [15], Ag/Sn [22], Au/In [14],
[23], Au/Sn [6], [9], [19]– [21], [23]–[25], Cu/Sn [26]–[29],
and Ni/Sn [28], [29]. Ag/In, which is among the first SLID
systems [14], [15], was investigated for HT applications in
several studies [9], [30]–[32]. Here, a HT stable joint can be
achieved (stable up to 700 °C [30]) using a processing temperature of only 210 °C, followed by annealing at 150 °C [32].
However, the HT lifetime of the joint is reported to be
limited (especially in combination with high power), due to Ag
migration [9], [10].
Au/Sn is a promising SLID system for HT applications [9],
[19]–[21], [33], [34]. Based on the Au–Sn phase (shown in
Fig. 2), several Au–Sn phases can be appropriate for HT
applications. However, when long time stability is taken into
account, the ζ phase is the most promising. The final bond
structure is reported to be layered [21]—Au/ζ /Au—where
the ζ phase undergoes a phase transition to the ζ phase
at 190 °C [35], [36]. The ζ phase has a melting point of
522 °C [35], [36], making it desirable for HT applications.
Thorough investigations of the Au–Sn phase diagram suggest
that the actual layered bond structure probably is Au/ζ /Au,
since the ζ phase is stable down to −5 °C (depending on Au
concentration). Another possible bond structure is Au/β/Au.
The β phase has an even higher melting point than the ζ
phase, i.e., 532 °C [35], [36]. However, the β phase is reported
to be more brittle than the ζ phase [37], indicating that a
Au/ζ /Au bond probably is more reliable than a Au/β/Au bond.
Au/Sn SLID was already shown good HT stability and thermal
cycling abilities in studies performed by Johnson et al. [9],
[19].
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Au–Sn phase diagram [35], [36].
Silicon carbide (SiC), a wide bandgap semiconductor, is
commonly considered as the best alternative for the next
generation of innovative, high performance, cost-effective, and
environmental-friendly drilling and well-intervention systems
for the oil industry [38]. SiC has a high breakdown field
strength, a high thermal conductivity, and offers excellent
performance in HT (up to 600 °C) and high power applications [39]–[41]. However, lack of qualified HT packaging
technology is a major limitation to fully realize the potential
of SiC.
This paper presents a study of Au–Sn SLID bonding in
a package utilizing a bipolar junction transistor (BJT) SiC
chip. First, two different processing techniques—electroplating
of Au/Sn layers and sandwiching a eutectic Au–Sn preform
between electroplated Au layers—are investigated in a simplified metallization system to find the best technique. The latter
process is further investigated for metallized Si3 N4 substrates,
in two different Cu/Si3 N4 /Cu/Ni–P/Au–Sn/Ni/Ni2 Si/SiC systems (different Au-layer thickness). Die shear tests and crosssections are performed on the different samples to characterize
the bonding properties associated with the different processing
techniques, metallization schemes, and environmental stress
tests. Furthermore, density functional theory (DFT) calculations were performed to investigate the relative thermodynamic
stability of different Au–Sn and Ni–Sn phases, to enhance the
understanding of how Ni acts as a diffusion barrier between
substrate/chip metallization and a AuSn SLID bond.
II. M ETHODOLOGY
A. Test Assemblies
1) SLID 1a & b Samples: Oxidized Si wafers with sputtered
Ti-W (60 nm)/Au (100 nm) adhesion/seed layers were used as
both substrate (diced in 4.3 × 6.6 mm2 after plating) and chip
(diced in 2 × 2 mm2 after plating) in the simplified bonding
samples, hereby referred to as SLID 1 samples. These were
then electroplated with a uniform Au layer (5 μm). The Au
electroplating was performed in a gold cyanide solution at
a temperature range 60 °C–65 °C, with a current density of
5.4 mA/cm2 . Two different types of SLID 1 samples were
manufactured.
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IEEE TRANSACTIONS ON COMPONENTS, PACKAGING AND MANUFACTURING TECHNOLOGY, VOL. 3, NO. 6, JUNE 2013
(a)
(b)
Fig. 3. SLID 1a samples. (a) Layers as plated. (b) Expected structure after
bonding.
(a)
(b)
Fig. 5. SLID 2a & b samples. Sketch of expected layer structure after
bonding. (a) SLID 2a. (b) SLID 2b.
(a)
(b)
Fig. 4. SLID 1b samples. (a) Layers as plated. (b) Expected structure after
bonding.
a) SLID 1a: Sn (2 μm)/Au (0.1 μm) layers were electroplated on the chip side using a tin sulfate solution at
room temperature, with a current density of 10 mA/cm2 .
The thin Au layer was applied to minimize oxidation,
making fluxless bonding possible [24]. A Sn/Au-plated
chip was then bonded to a substrate (see Fig. 3 for
illustration).
b) SLID 1b: A eutectic Au 80 wt.% Sn 20 wt.% preform
(7.5 μm) was sandwiched between a chip and substrate
to make the joint. The preform was purchased from
Micro Joining KB (see Fig. 4 for illustration).
The bonding was performed in two steps; first, a flip chip
bonder was used to pick and place at a moderate temperature
(120 °C) applying a force of 35 N for 30 s. Secondly, the
positioned samples were bonded using a hotplate in a vacuum
chamber and a clamping force to ensure intimate contact. The
applied force translated to a pressure 2.5 MPa on the surface
to be bonded. First, the samples were heated to 250 °C and
held there for 5 min (to bake out any residual moisture and
to assure a uniform temperature distribution in the bonding
layers). Then, the samples were heated to 350 °C, and kept
there for 20 min to ensure that the desired phases were created.
2) SLID 2a & b Samples: Commercialy purchased Si3 N4
substrates (from Denka Chemicals) with active metal bonded
Cu (150 μm) and plated Ni–P (7 wt% P) were used as
substrates for both SLID 2a & b samples. The substrates had
symmetrical metallization (Cu/Ni–P layers on both top and
backside) to minimize warpage of the substrate due to CTE
mismatches between Si3 N4 and Cu. An additional Au layer (3
or 5 μm) was electroplated on the substrates in a gold cyanide
solution at a temperature range 60 °C–65 °C, with a current
Fig. 6. Picture of manually aligned SLID 2a & b samples on a hot plate
fastened with a clamping force.
density 2.7 mA/cm2 . The substrate was diced in 6 × 6 mm2
samples after plating.
The BJT SiC dummy chips, delivered from Fairchild, had
sputtered Ni2 Si (140 nm)/Ni (300 nm)/Au (100 nm) metallization. The chips were electroplated with a uniform Au layer
(5 μm), and diced in 1.855 × 3.4 mm2 samples. Two different
types of SLID 2 samples were produced.
a) SLID 2a: As for SLID 1b samples, an eutectic Au
80 wt.% Sn 20 wt.% preform (7.5 μm) was sandwiched
between the chip and the substrate to make the joint.
For SLID 2a samples, the electroplated Au layer on the
substrate was only 3 μm. This means that there would
be no excess Au left on the substrate side of the joint
[see Fig. 5(a) for illustration].
b) SLID 2b: Same as SLID 2a samples, but with 5 μm
electroplated Au on both sides, resulting in excess Au
on both the substrate and the chip side in the final joint
[see Fig. 5(b) for illustration].
The bonding was performed in two steps: first, the substrate,
the preform, and chip were aligned manually on a hot plate
and fastened with a clamping force (see Fig. 6). Secondly,
the samples were bonded using the hotplate in a vacuum
chamber. The same bonding profile as for SLID 1 samples
were used. However, the final bonding time was reduced from
20 to 10 min based on work published in [20].
TOLLEFSEN et al.: HIGH TEMPERATURE INTERCONNECT AND DIE ATTACH TECHNOLOGY
100
150
100
50
0
0
50
100
Time (min)
Fig. 7. Cycling profile for thermal cycling. The temperature in the joint was
measured by a k-type thermocouple and an Agilent 34970A data acquisition
unit.
B. Test Methods and Equipment
Thermal cycling tests were peformed in a Heraeus HT
7012S2 thermal cycling chamber. Both SLID 1a & b and
SLID 2a & b samples were cycled between 0 °C and 200 °C,
with a gradient 10 °C/min, and a dwell time 15 min at
the temperature extremes (see Fig. 7 for cycling profile—the
temperature in the joint was measured by a k-type thermocouple and an Agilent 34970A data acquisition unit). Two levels
of cycling were performed (500 and 1000 cycles). The cycled
samples were shear tested in a Dage 2400A shear tester with
a 50 kgf load cartridge, and the results were compared to
measurements on as-bonded samples.
SLID 1a & b and 2a & b samples were aged in air at 250 °C
for 6 months in a Binder laboratory oven. The aged samples
were shear tested in a Dage 2400A shear tester with a 50 kgf
load cartridge, and compared to measurements on as-bonded
samples.
Cross-sectioning was performed on all groups of samples.
The samples to be cross-sectioned were embedded in epoxy
resin prior to grinding, and then ground (on SiC paper grade
320 through 4000, using water cooling) and polished (using
6 μm diamond particles and an alcohol-based lubricant prior
to fine polishing using 3 and 1 μm diamond particles with a
water- and oil-based lubricant). Note that SiC is a very hard
material compared to the Au and the Au–Sn phases, making
it difficult to prepare planar cross-sections.
The cross-sectioned samples and the fracture surfaces of
shear tested samples were investigated by optical microscopy
(Neophot 32), scanning electron microscopy (SEM: JEOL
JSM-5900LV) and energy-dispersive spectroscopy (EDS:
Oxford X-MAX 50).
C. Ab Initio Calculations
Periodic DFT calculations were performed using the
Vienna ab initio simulation package (VASP) [42], [43]
within the generalized gradient approximation according to
Perdew et al. [44]. The projector augmented wave (PAW)
method was used to treat the core regions within an allelectron framework [45]. The plane wave basis set cut-off
energy was 450 eV, and the density of the k-point grid
was at least 0.25 Å−1 for all the structures. Standard PAW
potentials were used, which implied that the following valence
Die shear strength (MPa)
Temperature (°C)
200
907
As bonded
500 cycles
1000 cycles
80
60
40
20
0
SLID 1a
SLID 1b
Sample group
Fig. 8. Die shear strength of SLID 1a & b samples as a function of thermal
cycles (0 °C–200 °C, 10 °C/min, dwell time of 15 min). The bars show the
standard deviation for the different groups.
Fig. 9. Optical microscopy image of a cross-section of an as-bonded SLID 1b
sample showing the challenges associated with co-planarity during bonding.
orbitals were used in: 5d10 6s1 (Au), 5s2 5p2 (Sn), and 3d8 4s2
(Ni). The criterion for self-consistency was that the energy
difference between two consecutive iterations was <10−6 eV.
All structures were relaxed with a quasi-Newton method (the
residual minimization scheme with direct inversion in the
iterative subspace) using a force convergence criterion of
0.05 eV/Å. The atomic positions as well as the unit cell size
and shape were relaxed simultaneously. Structural relaxation
was performed twice, before a final high-accuracy calculation
of the total energy. Experimental structures from a database
(Inorganic Crystal Structure Database) were used as input for
the structural relaxations [46].
III. R ESULTS AND D ISCUSSION
A. Thermal Cycling
1) SLID 1a & b Samples: After 500 and 1000 thermal
cycles, the die shear strength was determined and compared
to the as-bonded strength. Five samples from each group were
tested. The results have some clear trends (shown in Fig. 8).
a) Variations: There are large variations (standard deviation) in the shear strength. The reason for this most
likely originate from the sample manufacturing. A clamp
force was used to ensure sufficient contact between the
bonding surfaces (see Fig. 6). The clamp force is only
in contact with a restricted part of the surface area of
the chip and the substrate. During the bonding process,
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IEEE TRANSACTIONS ON COMPONENTS, PACKAGING AND MANUFACTURING TECHNOLOGY, VOL. 3, NO. 6, JUNE 2013
b)
c)
d)
e)
parts of the bonding materials will be liquified, meaning
that if the clamp force is not symmetrically placed
on the chip/substrate, problems with co-planarity will
occur. This was confirmed by cross-section pictures (see
Fig. 9).
Die shear strength: The shear strength has a substantial
increase after 500 cycles. The shear strength after 1000
cycles also increased compared to the as-bonded samples, but it had decreased compared to the strength after
500 cycles. The explanation for the relative increase in
shear strength probably stems from the challenges associated with co-planarization. During thermal cycling,
the samples are regulary heated and kept at 200 °C,
increasing the diffusion rate between the bonding partners, i.e., making a stronger joint. The reduction in the
shear strength after 1000 cycles compared to 500 cycles
indicates that there is a reduction in fatigue lifetime
during thermal cycling. However, due to the limited
amount of samples, and the large standard deviation,
this should be further investigated.
SLID 1a versus 1b: The SLID 1b samples has higher die
shear strength than the SLID 1a samples after thermal
cycling. There are no obvious reason for this trend, but
careful investigations of the cross-sections of 1a and
1b samples revealed that there was a through crack in
all SLID 1a samples (Fig. 10). There were no cracks
in the SLID 1b samples, so the through crack in the
1a samples probably explains the lower strength. The
cracks, located at the interface between the electroplated
Au and Sn layers on the chip side, probably originated
from contaminatons during Sn plating (e.g. on the Au
surface, not pure enough Sn, etc.). Hieber et al. [47]
experienced problems regarding the pureness of the
electroplated Sn in a CuSn SLID joint. They reported
that only e-beam evaporated Sn is pure enough for CuSn
SLID bonding. This shows the importance of the Sn
quality. However, strong and uniform CuSn SLID joints
have been produced using electroplated Sn [26].
Bond configuration: The bond interface of the the good
samples (without co-planarity problems) was uniform
with regard to IMC formation. Fig. 11 shows a optical
microscopy and a SEM image of the cross-section of a
SLID 1b sample. EDS was used to identify the bonding
phase as most probably the ζ phase. This supports our
assumption that a AuSn SLID bond most probably has
a Au/ζ /Au structure when the overall Au:Sn ratio in the
bondline is designed appropriately.
Fracture surfaces: The fracture surfaces of samples
with high-bond strength were located at the interfaces
between the substrate/chip and Ti-W. The fracture surfaces of samples with low-bond strength were located
in the actual bond layer. This indicates that the samples with low die shear strength fail because of coplanarization issues. However, the samples with high die
shear strength fail in the chip/substrate metallization,
indicating that a good AuSn SLID bond has higher
bond strength than 80 MPa (the highest measured bond
strength). This is also supported by previous work by
Fig. 10. Optical microscopy image of a cross-section of a SLID 1a sample
showing the through crack in the electroplated Au/Sn interface.
TABLE I
D IE S HEAR S TRENGTH OF SLID 2a S AMPLES AS A F UNCTION OF N O OF
T HERMAL C YCLES (0 °C–200 °C, 10 °C/min, D WELL T IME OF 15 min)
As Bonded (MPa)
500 Cycles (MPa)
1000 Cycles (MPa)
>78
>78
>78
71.8
64.3
>78
>78
69.0
66.9
50.3
>78
>78
43.8
36.4
35.5
TABLE II
D IE S HEAR S TRENGTH OF SLID 2b S AMPLES AS A F UNCTION OF N O OF
T HERMAL C YCLES (0 °C–200 °C, 10 °C/min, D WELL T IME OF 15 min)
As Bonded (MPa)
500 Cycles (MPa)
1000 Cycles (MPa)
>78
>78
>78
>78
>78
>78
>78
>78
>78
50.9
>78
>78
>78
>78
>78
>78
>78
>78
>78
52.9
>78
>78
>78
>78
>78
>78
69.9
68.4
66.3
65.5
Johnson et al. [19], who reported a die shear strength
above 90 MPa for AuSn SLID joints.
EDS analysis indicates that a AuSn SLID bond is constructed of a layered Au/ζ /Au structure. When the joining is
well performed, high-bond strengths are achieved. Importantly,
a AuSn SLID joint can withstand thermal cycling between
0 °C and 200 °C in a package with small CTE mismatches
(Si chip and substrate).
2) SLID 2a & b Samples: After 500 and 1000 thermal
cycles, the die shear strength was determined and compared
to the as-bonded strength. Ten SLID 2b samples and five 2a
samples were tested. The results have some clear trends
(shown in Fig. 12, Tables I and II).
a) Variations: The standard deviation in the shear strength
was greatly reduced compared to SLID 1a & b samples.
The assembly was performed as for SLID 1a & b
samples, but experience assured a higher yield. Note
that the majority of the samples had a bond strength
above the equipment limit (50 kgf), indicating that there
could be a considerable variation not measureable with
the applied equipment.
TOLLEFSEN et al.: HIGH TEMPERATURE INTERCONNECT AND DIE ATTACH TECHNOLOGY
(a)
909
(b)
Fig. 11. Optical microscopy image of (a) a SLID 1a sample with (b) a magnified SEM image of a representative region. The rectangles show regions where
EDS was performed. The atomic percent for the regions is included. Note that the features on the left side of the SEM image most probably stems from
cross-section preparations.
Die shear strength (MPa)
100
As bonded
500 cycles
1000 cycles
80
60
40
20
0
SLID 2a
SLID 2b
Sample group
Fig. 12. Die shear strength of SLID 2a & b samples as a function of the
number of thermal cycles (0 °C–200 °C, 10 °C/min, dwell time 15 min). Note
that this figure is only included to visualize the main trends. Since many of
the tested samples did not fracture during testing, due to the equipment limit
(50 kgf), a proper average cannot be made.
b) Shear strength: The bond strength of SLID 2b samples
remained relatively unchanged and superb (>78 MPa)
during thermal cycling (note that the majority of the
samples had a bond strength above the equipment limit,
indicating that there could be a degradation in the bond
strength not measureable with the applied equipment).
For SLID 2a samples, there was a decrease in the bond
strength as a function of the number of thermal cycles.
In SLID 2b samples, there was excess Au on both
the substrate and the chip side (Fig. 13). In SLID 2a
samples, there was no excess Au left on the substrate
side (Fig. 14), causing formation of brittle Au–Ni–Sn
IMCs during thermal cycling. These brittle IMCs are
believed to be the primary cause of the degradation of
the bond strength.
Fig. 13. Optical microscopy image of a cross-section of a SLID 2b sample
showing a uniform bondlayer, with excess Au on both substrate and chip side.
The different phases were identified by SEM and EDS.
Fig. 14. Optical microscopy image of a cross-section of a SLID 2a sample
showing a uniform bondlayer, but with excess Au on only the chip side. On
the substrate side brittle Au–Ni–Sn phases have been created, weakening the
die shear strength. The different phases were identified by SEM and EDS.
c) Bond configuration: As for SLID 1a & b samples,
the bond interface was uniform with regard to IMC
formation, EDS was used to identify the bonding phase
as most probably the ζ phase.
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Die shear strength (MPa)
100
As bonded
80
Aged
60
40
20
0
SLID 1a
SLID 1b
Sample group
Fig. 16. Die shear strength of SLID 1a & b samples as a function of
thermal ageing (250 °C, 6 months). The bars show the standard deviation for
the different groups.
TABLE III
D IE S HEAR S TRENGTH OF AS B ONDED AND T HERMALLY A GED
(250 °C, 6 M ONTHS ) SLID 2a S AMPLES
As Bonded (MPa)
Aged (MPa)
>78
>78
>78
>78
>78
43.8
71.8
36.4
64.3
35.5
Fig. 15. SEM images of the fracture surface of SLID 2a & b samples. The
fracture surfaces were identified with optical microscopy, scanning electron
microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS).
d) Fracture surfaces: In Fig. 15, SEM images of the
fracture surfaces of thermally cycled (1000 cycles) SLID
2a & b samples are shown. The fracture surface of 2b
samples was located at the chip/Ni2 Si/Ni interface, again
indicating that a good AuSn SLID bond has higher bond
strength than 78 MPa. The fracture surface of 2a samples
was located at the ζ phase/Au–Ni–Sn IMCs interface,
confirming that these brittle IMCs are the primary cause
of the degradation of the bond strength for 2a samples.
Notably, the AuSn SLID joints withstand thermal cycling
between 0 °C and 200 °C for a package with large CTE
mismatches (12 ppm/K difference between SiC and the thick
Cu film). The importance of excess Au on both substrate
and chip side in the final bond is also demonstrated. The
soft Au layer is important since it absorbs thermo-mechanical
stresses in the package, induced by, e.g., CTE mismatches
between the chip and the conducting layer. Excess Au on
both sides of the final joint is also important since it acts
as a diffusion barrier between the Au–Sn phases and the
chip/substrate metallization, which tends to develop brittle
IMCs like Ni3 Sn4 [48], greatly reducing the lifetime and the
reliability of the package. Furthermore, excess Au is central
for the prediction of the properties of the joint, since it assures
stable material phases with predictable/known properties.
The superb die shear strength of a AuSn SLID joint shows
that this is one of the most promising SLID bonding material
schemes for HT applications. In addition, the ability of AuSn
SLID joints to absorb stress makes them very attractive.
However, for applications with temperatures above ∼0.5× Tm ,
creep is considered to be important for the long time reliability [2]. (Tm is the melting point given in absolute temperature.) For a ζ phase bond, 0.5 × Tm = 125 °C, making it
vulnerable for creep during long time HT applications with
constant stresses in the system (e.g., from CTE mismatches).
But, by, e.g., applying a process temperature in close range
to the operation temperature (which often is possible for
SLID bonds, since the melting point of the final joint is well
above the processing temperature), stresses induced by the
CTE mismatches in the system can be minimized. The effect
of creep in a Au–Sn SLID joint should be more carefully
investigated.
B. Thermal Ageing
1) SLID 1a & b Samples: After six months of thermal
ageing, the die shear strength was determined and compared
to the as-bonded strength. Five samples from each group were
tested, and the results are shown in Fig. 16.
The results had the same tendency as the thermally cycled
SLID 1a & b samples. There was a large standard deviation,
the shear strength increased after ageing, the bond strength of
1b samples was higher than that of 1a samples, and the fracture
surface of samples with high-bond strength was located at
the interfaces between the subsrate/chip and Ti-W, while the
fracture surface of samples with low bond strength was located
in the actual bond layer. For a more thorough discussion about
these results, see the thermal cycling test section.
TOLLEFSEN et al.: HIGH TEMPERATURE INTERCONNECT AND DIE ATTACH TECHNOLOGY
911
TABLE IV
D IE S HEAR S TRENGTH OF AS B ONDED AND T HERMALLY A GED
(250 °C, 6 M ONTHS ) SLID 2b S AMPLES
As Bonded (MPa)
Aged (MPa)
>78
>78
>78
>78
>78
>78
>78
>78
>78
50.9
>78
>78
>78
>78
>78
72.9
66.9
58.5
55.9
54.1
Die shear strength (MPa)
100
As bonded
Aged
80
60
40
20
Fig. 18. SEM image of a SLID 2a sample. The composition shown was
found by EDS analyses.
Notably, these results demonstrate that the Au/ζ /Au bond
is not stable at 250 °C. Sn continues to diffuse into the pure
Au layers/Au continues to diffuse into the ζ layer, creating a
β/ζ layer with no excess Au between the diffusion barrier,
Ni, and the bond layer. When there is a physical contact
between Sn and Ni, interdiffusion results in the formation of
the brittle Ni3 Sn2 and Ni3 Sn4 IMCs, leaving only pure Au
in the original bond (probably with some Sn). In Fig. 19, a
schematic overview over the AuSn SLID bond structure at
different stages is shown.
C. DFT Calculations
0
SLID 2a
SLID 2b
Sample group
Fig. 17. Die shear strength of SLID 2a & b samples as a function of
thermal ageing (250 °C, 6 months). Note that this figure is only included to
visualize the main trends. Since many of the tested samples did not fracture
during testing, due to the equipment limit (50 kgf), a proper average cannot
be made.
2) SLID 2a & b Samples: After six months of thermal
aging, the die shear strength was determined and compared
to the as-bonded strength. Five SLID 2a samples and 10
SLID 2b samples were tested. The results are shown in Fig. 17,
Tables III, and IV.
The results can be compared to the thermally cycled
SLID 2a & b samples. However, the bond stength of the
aged SLID 2a samples was reduced more than that of the
thermally cycled samples. Furthermore, the bond strength of
the SLID 2b samples wasalso reduced (the bond strength of
SLID 2b samples was not reduced during thermal cycling).
Cross-sections of the thermally aged samples revealed large
alterations in the bond structure compared to the as-bonded
structure. In Fig. 18, a SEM image of a cross section of an aged
SLID 2a sample is shown. It reveals that all the Sn in the AuSn
SLID bond has diffused into the diffusion barrier, Ni, creating
Ni3 Sn2 and Ni3 Sn4 , leaving only pure Au in the original bond.
Notice that the composition analysis was performed with EDS.
Due to the uncertainty in EDS analyses, and the thickenss of
the Ni–Sn IMC layers, there probably is some remaining Sn
in the bond layer. As seen from the Au–Sn phase diagram in
Fig. 2, the Au phase can contain up to 2–3 at.% Sn in solid
solution at room temperature.
DFT calculations utilizing VASP were performed to enhance
the understanding of interdiffusion at the Au–Sn/Ni interface. The enthalpy of formation for hypothetical reactions
taking place at the interface was calculated by comparing the
ground state electronic total energy of the various compounds.
Only combinations in the relevant part of the phase diagram
(Au/Ni > 0.75) were included. The results are shown in
Table V and Fig. 20, and confirm that a combination of pure
Au and intermetallic Ni–Sn is thermodynamically more stable
than a combination of Au–Sn and pure Ni. This is valid for
all combinations of Au–Sn phases (the ζ -phase Au31 Sn5 , the
β-phase Au11Sn, or Au5 Sn) and Ni–Sn phases (Ni3 Sn2 or
Ni3 Sn4 ). It is furthermore instructive to see how the enthalpy
of reaction defines a convex hull of the final combination of
Au and Ni–Sn when plotted as a function of the Au/Ni ratio in
Fig. 20. That is, depending on the effective Au/Ni ratio, both
Ni3 Sn2 and Ni3 Sn4 will be thermodynamically accessible in
the relevant part of the phase diagram. The local, effective
Au/Ni ratio will depend on a number of parameters like the
global Au/Ni ratio, the temperature profile, the morphology,
other phases present, as well as the kinetic barriers of solidstate diffusion.
The calculations above were performed at the ground state,
i.e., with no contributions from temperature, entropy, or zeropoint motion. This means that we have in reality only probed
the phase diagram at 0 K. We have also neglected the possibility of other phases in this part of the phase diagram, including
amorphous ones. Finally, the factors governing the effective
Au/Ni ratio, mentioned above, have also been disregarded.
Nevertheless, the results are consistent and conclusive enough
to rationalize the tendency of Sn to diffuse from the Au–Sn
phases into Ni when they are in physical contact. As long as
912
IEEE TRANSACTIONS ON COMPONENTS, PACKAGING AND MANUFACTURING TECHNOLOGY, VOL. 3, NO. 6, JUNE 2013
Fig. 19. Schematic overview of the AuSn SLID bond structure at different stages in a SLID 2b sample. Note that the As bonded and After HT ageing
structures are based on experimental observations, whereas the after TC structure is deduced as an intermediate state between the two aforementioned, based
on the phase diagram (Fig. 2).
TABLE V
G ROUND S TATE E NTHALPY OF R EACTION Hr OF P OTENTIAL C HEMICAL
R EACTIONS AT THE Au–Sn/Ni I NTERFACE C ALCULATED BY DFT. A
N EGATIVE E NTHALPY OF R EACTION I NDICATES T HAT THE R EACTION I S
E NERGETICALLY FAVORABLE . Hr I S G IVEN IN kJ/mol P ER Au ATOM
Potential Reaction
4Au11 Sn + 3Ni → 44Au + Ni3 Sn4
4Au31 Sn5 + 15Ni → 124Au + 5Ni3 Sn4
2Au11 Sn + 3Ni → 22Au + Ni3 Sn2
4Au5 Sn + 3Ni → 20Au + Ni3 Sn4
2Au31 Sn5 + 15Ni → 62Au + 5Ni3 Sn2
2Au5 Sn + 3Ni → 10Au + Ni3 Sn2
Au/Ni
0.94
0.89
0.88
0.87
0.81
0.77
Hr (kJ/mol)
−2.47
−3.19
−4.14
−3.02
−6.15
−6.70
0
-1
ΔHr (kJ/mol)
-2
-3
-4
-5
-6
-7
-8
0,75
0,80
0,85
0,90
Au/Ni
0,95
D. Some Final Remarks
When the DFT results are seen together with the experimental results from the ageing of SLID 2a & b samples, we
can understand how Ni–P acts as a diffusion barrier between
Cu and Au; Sn from intermetallic Au–Sn will diffuse into
Ni and form intermetallic Ni–Sn. This furthermore stresses
the importance of having a residual pure Au layer acting as
a diffusion barrier between substrate and chip metallization
as well as on both sides of the final Au–Sn SLID bond. One
then needs to take into account that the final Au–Sn SLID bond
may be a Au/ζ /Au bond. For a Au–Sn SLID bond to withstand
long-time HT exposure, the bondline should be designed with
a Sn:Au ratio corresponding to max. 8 at.% Sn, to ensure
surplus Au also after transformation to a ζ phase.
It should be mentioned here that Ti seems to act as a better
diffusion barrier than Ni for interdiffusion of Sn. The SLID 1b
samples had exactly the same bonding metallization as the
SLID 2b samples. However, the SLID 1b samples used Ti-W
as a diffusion barrier instead of Ni. After six months of thermal
ageing at 250 °C, there was no trace of Sn in the Ti-W layer.
This is consistent with the work performed by Anhock et al.
[49], where Ti was found to be a better diffusion barrier than
Cr, Ni, Pd, and Pt.
1,00
Fig. 20. Enthalpy of reaction Hr of the tentative reactions listed in Table V,
as a function of the Au/Ni ratio. The enthalpy is given in kJ/mol per Au
atom. The dotted line designates the convex hull of the reaction products
(Ni–Sn intermetallic phases and Au) in this part of the ternary Au–Ni–Sn
phase diagram at 0 K.
the solid-state diffusivity is high enough for such transport to
take place, there is a strong thermodynamic driving force for
the formation of intermetallic Ni–Sn.
IV. C ONCLUSION
Two different AuSn SLID processing techniques were investigated, where sandwiching of a eutectic Au–Sn preform
between electroplated Au layers was found to be the preferred
scheme. Initial processing issues regarding co-planarization
and through cracks in the Au/Sn interface were solved, and
strong uniform Au/ζ /Au joints were produced.
The bond strength of a Au–Sn SLID bond is shown to be
superb, >78 MPa. However, for joints without excess Au on
TOLLEFSEN et al.: HIGH TEMPERATURE INTERCONNECT AND DIE ATTACH TECHNOLOGY
both substrate and chip side, the shear strength is reduced by
thermal cycles.
Importantly, it is demonstrated that a Au–Sn SLID joint
can absorb thermo-mechanical stresses induced by large CTE
mismatches (12 ppm/K) in a package during HT thermal
cycling. However, after long-time HT exposure, the joint
is transformed from Ni/Au/ζ /Au/Ni to Nix Sn y /Au/Nix Sn y ,
apparently through a Ni/β/Ni stage. The primary cause of the
formation of the Nix Sn y IMCs has been investigated by DFT
calculations. The results indicate that the combination of pure
Au and intermetallic Ni–Sn is energetically more favorable
than pure Ni and intermetallic Au–Sn. The Sn:Au ratio should
thus be below 8 at.%. When the insight above is implemented,
Au–Sn SLID bonding is an excellent candidate for HT die
attach and interconnect technology.
ACKNOWLEDGMENT
The authors would like to thank T. T. Luu and
Dr. K. Wang for manufacturing the SLID 1a&b samples and to
Dr. M. M. V. Taklo for proofreading.
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Torleif André Tollefsen was born in Kristiansand,
Norway, in 1978. He received the M.Sc. degree
in material physics from the University of Oslo,
Oslo, Norway, in 2008. He is currently pursuing
the Ph.D. degree in applied micro- and nanosystems
with Vestfold University College, Borre, Norway.
He was a System Engineer at FMC Technologies,
Bergen, Norway, from 2008 to 2010, working with
systems for the subsea production of oil and gas.
In 2010, he joined the Research Institute SINTEF
and Vestfold University College to pursue a Ph.D.
degree. From 2013, he has been a Research Scientist at SINTEF. He has
authored more than 25 journal and conference papers. His current interests
include packaging for micro and nanosystems, with a special focus on
materials for harsh environments.
Andreas Larsson was born in Eskilstuna, Sweden,
in 1980. He received the M.Sc. degree in engineering
physics–applied physics from Uppsala University,
Uppsala, Sweden, in 2006.
He has been a Structural Engineer with the Norwegian Oil and Gas industry since 2006, where
he performed finite element analysis and concept
development. Since 2008, he has been with the
Research Institute SINTEF, currently as a Senior
Scientist responsible for harsh environment packaging. He has authored more than 40 publications
and holds patents. His current research interests include packaging and
interconnect technology for harsh environments, with the important topics
being multiphysics simulations and concept development.
Ole Martin Løvvik was born in Norway in 1968.
He received the M.Sc. degree in theoretical physics
and the dr.scient. degree in energy physics from the
University of Oslo (UiO), Oslo, Norway, in 1993
and 1998, respectively.
He was a Post-Doctoral Researcher at UiO from
1999 to 2006 and the Institute for Energy Technology, Kjeller, Norway. He was an Associate Professor
with the Norwegian University of Life Sciences, As,
Norway, in 2001, a Consultant with the Research
Council of Norway, Oslo, in 2006, a Research Fellow with the Institute for Energy Technology, Kjeller, from 2006 to 2008,
a Visiting Scientist with the University of Hiroshima, Hiroshima, Japan, in
2008, and a Visiting Professor with Osaka University, Osaka, Japan, in 2012.
He is currently a Senior Researcher with the research institute SINTEF and
an Adjunct Professor with the University of Oslo. He has authored around
65 papers in international journals with referee, and has given more than 90
invited and contributed talks at international conferences. His current research
interests include atomic-scale modeling of materials for energy technologies,
with the important topics being electronic materials, materials for hydrogen
technology, and light weight metals.
Knut E. Aasmundtveit received the M.S. degree
in technical physics from the Norwegian Institute
of Technology, Trondheim, Norway, in 1994, and
the Ph.D. degree in materials physics from the
Norwegian University of Science and Technology,
Trondheim, in 1999.
He was with Alcatel Space Norway/Ame Space
(now Kongsberg Norspace), Horten, Norway, as a
Radio Frequency System Design Engineer, until
2004, when he joined the Department of Micro
and Nano Systems Technology, Vestfold University
College, Borre, Norway, as an Associate Professor. He has authored or coauthored four book chapters and more than 60 journal and conference papers.
His current research interests and teaching activities include packaging and
integration technology for micro and nano systems, such as intermetallic
bonding, polymer-based bonding, nanomaterials integration in microsystems,
and biomedical packaging.