department of materials science and engineering college of

College of Engineering
National Taiwan University
Micro-joints with Different Substrates
102 6

(ACPL)
Sean
ACPL

II
ACKNOWLEDGEMENT
This doctoral thesis presents the research results obtained by Dr. Yu-Jen Chen
during his study on the Ph.D. degree. Yu-Jen would like to express his sincere gratitude
to his advisor, Professor C. Robert Kao, for all his patience, inspiration and guidance in
advising him on the researches. Besides, Yu-Jen would also like to thank the Prof. C. M.
Chen, Prof. Albert Wu, Prof. Y. W. Yen, and Prof. S. K. Lin for their suggestions on
generating this thesis.
Yu-Jen would like to appreciate his colleagues and good friends in Advanced Chip
Packaging Laboratory (ACPL) at NTU: Dr. H. Y. Chuang, Dr. C. K. Chung, Dr. W. M.
Chen, J. H. Ke, T. L. Yang, C. C. Li, W. L. Shih, Y. F. Lin, T. C. Huang, M. S. Kuo, J.
J. Yu, H. Y. Kang, M. H. Chen, K. Y. Huang, Z. X. Zhu, Sean Yang, C. K. Chung, and H.
T. Chen.
In particular, Yu-Jen thanks the love and support that his family has been giving to
him. Without them, he would not be able to attentively concentrate on his studies. He
would like his girlfriend, Emilie Chiang, for her unlimited encouragement and support.
Yu-Jen Chen
National Taiwan University
10003D IC

ENEPIG-- 3D IC

Cu (UBM) Cu
(Au,Ni)Sn4(Pd,Ni)Sn4 (Au,Pd,Ni)Sn4

V
ABSTRACT
Current development of solder joints in flip-chip microelectronic packages has
range about 100 μm in diameter. However, the diameter of a typical micro-joint is only
about 10 μm in three-dimensional integrated circuits (3D IC). The volume of a solder
joint shrinks by a factor of 1000 during the transition from flip-chip to 3D IC between
generations. Due to such large reduction in volume, there are some unique requirements
and challenge for 3D IC. One such requirement is that solder joints are confined to a
much smaller space compared to solder joints in outer levels of packages. As a result,
some new issues may be arisen due to such a large miniaturization of the solder volume.
This research is divided into two parts. The first part studies the mechanical
properties and interfacial reactions of 3D IC micro-joints. In micro-joints, two kinds of
microstructures can be present in Cu/Sn/Cu sandwiches during assembly bonding
process or under a long-term aging process: (1) joints with unreacted remaining solder
and (2) joints fully occupied by intermetallic compounds (IMCs). According to the
shear test results, the shear strength decreases with aging time in both cases. This study
points out that the strength of the joints fully occupied by IMCs is much higher than the
joints with remaining solder. Finally, this study generalizes a series of failure
mechanisms for different microstructures.
The strength of the micro-joint is very sensitive to the microstructure of IMCs. For
VI
the joints with remaining solder, the main contributing factor was the planarization of
the Cu6Sn5 morphology through the aging process. For the entire joint occupied by
IMCs, characterization of microstructure indicates that the void formation at Cu3Sn/Cu
interface is the main factor to deteriorate the mechanical properties of the joints.
The second part of this study is to investigate the interfacial reactions and
mechanical properties of 3D IC micro-joints with different surface finishes under space
confinement, including ENIG (Ni/Sn-Au/Ni), ENIP (Ni/Sn-Pd/Ni), and ENEPIG
(Ni/Sn-Au-Pd/Ni). In the 3D IC packaging, the size of solder joints becomes relatively
smaller, which can result in Au or/and Pd embrittlement as the Au or/and Pd content in
the solder is too high. The phenomenon of Au or/and Pd embrittlement caused by such
space-confined region in 3D IC can severely influence the microstructure of solder
joints.
The formation of intermetallic compound of (Au,Ni)Sn4, (Pd,Ni)Sn4 or
(Au,Pd,Ni)Sn4 will appear to be a continuous layer across entire solder joints, even if
the Au or Pd layer is very thin. This brittle layer presented in the joints could
significantly deteriorate the reliability of the solder joints. Furthermore, the effect of Au
and Pd on interfacial reaction is different. The Pd embrittlement phenomenon is much
severe than the Au embrittlement when the effective Pd/Au concentration is the same
(Sn0.8Pd/Sn0.8Au). The key reason is that more Ni can diffuse into the (Pd,Ni)Sn4 IMC,
VII
so the (Pd,Ni)Sn4 layer is much thicker than (Au,Ni)Sn4.
To prevent such serious Au/Pd brittle effect under space-confined 3D IC
construction, it is suggested that with addition of Cu into the reaction by applying Cu
under bump metallization instead of using Cu added solders can successfully inhibit the
formation of (Au,Ni)Sn4, (Pd,Ni)Sn4 or (Au,Pd,Ni)Sn4, and eliminates the risk of Au
or/and Pd embrittlement.
property; Au embrittlement; Pd embrittlement.
VIII
CONTENTS
1.3 Introduction of Surface Finishes 9
1.4 Intercacial Reactions and Machanical Properties of Cu/Sn/Cu Sandwiches under
Space Confinement 18
1.5 Intercacial Reactions between Lead-free Solder and ENIG, ENIP, ENEPIG
Surface Finishes (Au and Pd Embrittltment Issues) 23
1.6 Aims of the Thesis 28
II. EXPERIMENTAL PROCEDURES
2.1 Cu Pillar Solder Bump (Cu Pillar Bump/Solder Cap) 29
2.2 Ni/Sn-Au/Ni Sandwich (ENIG) and Ni/Sn-Pd/Ni Sandwich (ENIP) 33
2.3 Ni/Sn-Au-Pd/Ni Sandwich (ENEPIG) 38
2.4 Mechanical Properties of Micro-joints Measured by Nanoindentor 40
III. RESULTS AND DISCUSSION
3.1 Interfacial Reaction and Shear Strength of Cu Pillar Solder Bump 41
3.1.1 Interfacial Reaction of Cu Pillar Solder Bump 41
3.1.2 Kinetics of IMCs Growth 45
IX
3.1.3 Shear Strength of the Joints with Remaining Solder 46
3.1.4 Shear Strength of the Joints Fully Occupied by IMCs 55
3.2 Interfacial Reaction of Ni/Sn-Au/Ni (ENIG) under Space Confinement 63
3.3 Interfacial Reaction of Ni/Sn-Pd/Ni (ENIP) under Space Confinement 71
3.4 Interfacial Reaction of Ni/Sn-Au-Pd/Ni (ENEPIG) under Space Confinement76
3.5 Characterization of Au and Pd in Interfacial Reactions 80
3.6 Mechanical Properties of Au or/and Pd Embrittlement Micro-joints 83
3.7 Solutions of Au or/and Pd Embrittlement in the Space Confined Soldering
Reaction for 3D IC Packaging 90
IV. CONCLUSIONS 95
Figure 1-1 Packaging roadmap. [1] ............................................................................... 2
Figure 1-2 Schematic drawings showing the concept of 3D IC. ................................... 2
Figure 1-3 Typical schemes for establishing 3D IC (a) Wire bonding and (b) TSV
micro-bumping. [6] ........................................................................................................ 4
Figure 1-4 TSV 3D IC packaging. [7] ........................................................................... 4
Figure 1-5 Comparison of wire bonding and TSV micro-bumping for achieving 3D IC. [6]
....................................................................................................................................... 5
Figure 1-6 Solder volumes in different packaging levels. [3] ....................................... 6
Figure 1-7 Micrographs showing the typical morphology of (a) 3D IC micro-bump and (b)
flip-chip joint. ................................................................................................................ 8
Figure 1-8 Configuration of OSP surface finish. [12] ................................................. 11
Figure 1-9 Configuration of ENIG surface finish. [12] ............................................... 12
Figure 1-10 Configuration of ENIP surface finish. [12] ............................................. 14
Figure 1-11 Configuration of ENEPIG surface finish. [12] ........................................ 17
Figure 1-12 SEM images taken from the polished cross-sections of the Cu/Sn/Cu
samples reflowed at 340 o C for (a) 10 min, (b) 40 min, and (c) 90 min. [4] ............... 19
Figure 1-13 Two kinds of microstructures can be present in the micro-joints: (a) joint
with remaining solder and (b) joint fully occupied by IMCs. ..................................... 20
Figure 1-14 Cross-sectional SEM-BSE images of Cu pillar bumps after bonding. [25]21
Figure 1-15 Die shear test force as a function of annealing time. [25] ....................... 21
Figure 1-16 BSE micrograph showing the (Au0.4Ni0.6)Sn4 layer formed in a solder joint
XI
that had been aged at 160°C for (a) 25 h, (b) 100 h, (c) 500 h, and (d) 1000 h. The solder
was eutectic SnPb, and the soldering pad was 1 mm Au over Ni layer. [37] .............. 24
Figure 1-17 Fracture surfaces for solder side and pad side of the same solder joint that
had been aged at 160°C for 500 h. The (Au,Ni)Sn4/Ni3Sn4 interface is very weak. [35]
..................................................................................................................................... 24
Figure 1-18 BSE micrographs showing the reaction between Sn3.5Ag balls and Au/Ni
substrates of different Au thicknesses at 160°C for 1,000 h. (a) 0 μm Au, (b) 0.2 μm Au,
(c) 1.0 μm Au, (d) 2.0 μm Au, and (e) over view of Fig. 1.14(d). [42] ....................... 25
Figure 1-19 BSE micrographs showing the Sn–xPd/Ni interfaces after reaction at 250 o C
for 2 min. The x were (a) 0.05, (b) 0.1, and (c) 0.2, respectively. [43] ........................ 26
Figure 1-20 Cross-sectional view showing the fracture of a Sn–0.2Pd/Ni joint after
shear test. [43] ............................................................................................................. 26
Figure 2-1 The flow chart of the Cu pillar bump with solder cap fabrication. ............ 30
Figure 2-2 As-fabricated Cu pillar solder bump structure. .......................................... 31
Figure 2-3 Schematic drawings showing the different shear position (a) shearing at solder
matrix and (b) shearing at IMCs. ................................................................................. 32
Figure 2-4 Schematic drawings showing (a) the simplified solder joint and (b) the
sandwich configuration................................................................................................ 34
Figure 2-5 Schematic drawings showing the sample preparation via hot-pressing
bonding. ....................................................................................................................... 37
Figure 2-6 As-fabricated sandwich (a) Ni/Sn-Au/Ni and (b) Ni/Sn-Pd/Ni ................. 37
Figure 2-7 As-fabricated Ni/Sn-Ag/ENEPIG structure. .............................................. 38
Figure 2-8 HysitronTI950 TriboIndenter. .................................................................... 40
Figure 3-1 FIB-SIM images of the cross-sectioned Cu pillar solder bump that had been
aged at 180°C for (a) 0 h, (b) 150 h, (c) 300 h, (d) 500 h, and (e) 750 h. ................... 43
XII
Figure 3-2 Magnified images of rectangle regions of Figs. 3.1 (c)–(e). Voids were formed
at the Cu3Sn/Cu interface. ........................................................................................... 44
Figure 3-3 The thickness of IMCs formed at 180°C against the square root of reaction
time. ............................................................................................................................. 45
Figure 3-4 The shear strength as a function of aging time at 180°C. .......................... 47
Figure 3-5 Ball shear failure modes (a) Ductile mode, (b) Quasi-ductile mode, (c)
Quasi-brittle mode, and (d) Brittle mode. [51] ............................................................ 48
Figure 3-6 The fracture surfaces of solder joints that have been aged at 180°C for (a) 0 h,
(b) 150 h, (c) 300 h, (d) 500 h, and (e) 750 h.. ............................................................ 51
Figure 3-7 The phase percentages of fracture surface against aging time. .................. 52
Figure 3-8 Schematic illustration showing the fracture mechanisms during aging at
180°C from 0-750 h. .................................................................................................... 54
Figure 3-9 The shear strength as a function of aging time at 180°C (a) the joints fully
occupied by IMCs, and (b) the joints with remaining solder. ..................................... 56
Figure 3-10 The fracture surface of solder joint that had been aged at 180 for (a) 0 h, (b)
150 h, (c) 300 h, (d) 500 h, and (e) 750 h. ................................................................... 59
Figure 3-11 The percentages of the different phases on the fracture surface against the
aging time. ................................................................................................................... 60
Figure 3-12 Schematic illustration showing the failure mechanism during aging at 180
for 0-750 h. .................................................................................................................. 62
Figure 3-13 Micrographs showing Ni/Sn0.8Au/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 24 h, (c) 96 h, (d) 192 h, and (e) 240 h . ................................................................ 65
(b) 24 h, (c) 72 h, (d) 168 h, and (e) 192 h . ................................................................ 68
XIII
(b) 24 h, (c) 96 h, (d) 144 h, and (e) 168 h . ................................................................ 69
(b) 24 h, (c) 48 h, (d) 72 h, and (e) 96 h . .................................................................... 70
Figure 3-17 Micrographs showing Ni/Sn0.8Pd/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 6 h, (c) 9 h, (d) 12 h, and (e) 16 h . ........................................................................ 74
Figure 3-18 Micrographs showing Ni/Sn0.29Pd/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 6 h, (c) 24 h, (d) 36 h, and (e) 54 h . ...................................................................... 75
Figure 3-19 Micrographs showing Ni/Sn1.8Ag/ENEPIG sandwiches aged at 200 o C for
(a) 0 h, (b) 50 h, (c) 100 h, and (d) 200 h. ................................................................... 77
Figure 3-20 Micrographs showing (a) Ni/Sn0.8Pd/Ni and (b) Ni/Sn0.8Au/Ni sandwiches
aged at 200 o C for different aging time. ...................................................................... 81
Figure 3-21 Micrographs showing (a)Ni/Sn-Au/Ni, (b)Ni/Sn-Pd/Ni, and (c)
Ni/Sn-Au-Pd/Ni sandwiches as Sn fully consumed. ................................................... 84
Figure 3-22 Fracture surfaces for the solder side and ball side of the same solder joint that
had been aged at 160oC for 500 hr. This solder joint fractured just from routine handling.
This solder joint failed along the (Au1-xNix)Sn4/Ni3Sn4 interface and the Ni3Sn4/Ni
interface, respectively. [35] ......................................................................................... 84
Figure 3-23 A schematic drawing (white line) showing the fracture locations of the solder
joint in Fig. 3.21. [35] .................................................................................................. 85
Figure 3-24 (a) Shear strength of the Sn–xPd/Ni joints as a function of Pd
concentration(x). (b) Typical fracture of a Sn–0.2Pd/Ni joint after the shear test....... 86
Figure 3-25 Representative load–depth curves from nanoindentation lateral to different
phases .......................................................................................................................... 89
sandwiches aged at 200 o C for 24 h. ............................................................................ 91
XIV
Figure 3-27 Micrographs showing Cu/Sn-1.8Ag/ENEPIG sandwiches aged at 200°C for
(a) 0 h, (b) 50 h, (c) 100 h, (d) 200 h, and (e) EPMA line scan across the reaction zone of
the sample shown in (d). .............................................................................................. 94
XV
LIST OF TABLES
Table 1-1 Process Summary of Organic Solderability Preservative (OSP). [12] ........ 11
Table 1-2 Process Summary of Electroless Nickel/Immersion Gold (ENIG). [12] .... 13
Table 1-3 Process Summary of Electroless Nickel/Immersion Palladium (ENIP). [12]15
Table 1-4 Process Summary of Electroless Nickel/Electroless Palladium/Immersion Gold
(ENEPIG). [12]. ........................................................................................................... 17
Table 2-1 The experimental thickness of Au/Pd layer and the effective Solder
compositions that Au/Pd layer dissolved into.............................................................. 35
Table 3-1 Quantitative analysis of the (Au,Pd,Ni)Sn4 through EPMA. ...................... 79
Table 3-2 Quantitative analysis of (a) (Pd,Ni)Sn4 and (b) (Au,Ni)Sn4 through EPMA.82
Table 3-3 Young’s modulus and hardness of Ni, Ni3Sn4, (Au,Ni)Sn4, (Pd,Ni)Sn4, and
(Au,Pd,Ni)Sn4 IMCs obtained by nanoidenter. ........................................................... 88
1
INTRODUCTION
1.1 Background Introduction
The electronic industry is rapidly expanding and the main trends in consumer
electronic products are directed toward miniaturization, portability, and high
performance. For decades, we have been followed Moore’s law in which the number
of transistors inside an IC doubles every 18 months and semiconductor manufacturers
have been effectively shrinking transistor size of integrated circuit (IC) to achieve
higher speed and performance. In recent years, imminent ending of Moore's law is the
most critical issue threatening the continuing development of semiconductor industry.
As the limitation of shrinkage reaches its critical dimension in ICs, it’ll become
increasingly more difficult and challenged. Hence, one of the strategies to resolve its
limitation in dimension shrinkage for continuous development is to apply 3D IC.
For pursuing the higher packaging density and performance, the packaging
technology advances from 2D to 3D [1-5], as shown in Fig.1.1. The basic concept of
3D IC package is to integrate two or more layers of active electronic components,
such as logic and memory chip, vertically into a single circuit, as showed in Fig. 1.2.
The benefits and advantages of 3D IC are higher functionality, lower power
requirement and consumption, and significantly reduction in size as compared to the
conventional 2D IC packaging. To overcome packaging limitation and obtain better
2
quality of ICs, the 3D IC architecture will be highly accepted and favoured by most
major semiconductor companies, including Intel, TSMC, IBM, and Samsung.
Fig. 1.1 Packaging roadmap. [1]
Fig. 1.2 Schematic drawings showing the concept of 3D IC.
3
Two different methods were used for 3D IC packaging, including wire-bonding
and through silicon via (TSV) micro-bumping [1, 6], as shown in Fig. 1.3. The wire
bonding is a common approach for 3D IC packaging, in which Au or Cu wires
connect the individual die in a stack. Au or Cu wire serves as individual connected
signal pathway between the board and each chip and the board or the chip carrier acts
as communication interface between chips. In order to have communication between
each chip, Au or Cu wire plays an important role in transferring signal from chip to
the board to other chips. However, it is possible to bond from chip to chip in the stack
without interconnected by Au or Cu wire and the board. Furthermore, this approach is
limited by the resolution of wire bonders and becomes more and more difficult as the
input/output (I/O) density in the chip stack increases.
In contrast, TSV micro-bumping is a better strategy to approach 3D IC. This
technology has the potential to offer the greatest interconnect density. Assembly
occurs at the wafer level, placing a second wafer face down on the first wafer
(face-to-face) and subsequent wafers face down (face-to-back) as the number of tiers
grows. This approach used an intermediate material, solder, on the surface of die to
make connections. A commercial TSV 3D IC packaging sample is shown in Fig. 1.4
[7]. The micro-bumps typically have a pitch of 50 to 200 μm, and the bumps in such a
configuration typically have a height of several microns. This approach offers a much
4
greater vertical interconnect density than the wire bonding approach. TSV
micro-bumping also enables the use of one or more chips, from the same or from
different fabrication technologies, in each layer of the stack [8-11]. The pros and cons
of the 3D IC architecture accomplished by wire bonding and TSV micro-bumping
with can be seen in Fig. 1.5 [6].
Fig. 1.3 Typical schemes for establishing 3D IC (a) Wire bonding and (b) TSV
micro-bumping. [6]
5
Fig. 1.5 Comparison of wire bonding and TSV micro-bumping for achieving 3D IC.
[6]
The solder volume of micro-joints is much smaller than conventional solder
joints, e.g., ball grid array (BGA) or flip chip, especially for 3D IC applications.
Fig.1.6 shows solder volume in decreasing order of BGA, flip-chip and 3D IC
micro-joints [3]. The diameter of conventional BGA and flip-chip scale solder joints
is about 760 μm and 100 μm, respectively. The diameter of typical micro-joint in a 3D
IC under development today is 10 μm, only 1/1000 the volume of a typical flip chip
scale solder joint. Therefore, some issues concerning joint properties may be arisen
because of reducing the solder volume.
6
7
Soldering reactions under space confinement has become increasingly important
due to its application for chip stacking in three-dimensional integrated circuits. The
interfacial reactions occur under a severe space confinement, and some issues
concerning joint properties may be arisen in 3D IC micro-joints. The space
confinement behavior is defined as the products of interfacial reaction, e.g.,
precipitates and/or IMCs, will occupy a large fraction of a solder joint. As shown in
Fig. 1.7, a conventional flip-chip joint compared with a 3D IC micro-joint. The
micro-joint has a pad diameter of 10 μm and a solder layer with 5 μm thickness, as
showed in Fig. 1.7 (a). For a conventional flip-chip scale solder joint, showed in Fig.
1.7 (b), the diameter and solder height are about 100 μm. If these two joints were
subjected to a very similar thermal aging condition, the IMC thickness in each joint is
nearly the same. In 3D IC micro-joints, the interfacial reaction proceeds under space
confinement. Therefore, IMCs will occupy a large portion of the joint volume or
become an all-IMCs joint. In addition, IMC grains growing from the opposite
interfaces impinge on each other in the joints. This behavior does not occur in
conventional BGA or flip-chip scaled solder joint. The key reason of the space
confinement behavior is due to the extremely low solder volume in an extremely
small joint. In such a small joints, the mechanical properties dominate by IMCs. The
8
IMCs tends to be hard and brittle, so that joints containing a large portion of IMC
might perform peculiar behaviors and properties. Most of the important issues remain
uncertain because there is very limited information regarding the space confined
interfacial reactions in the literatures. In the following chapter, interfacial reactions
and mechanical properties of 3D IC micro-joints with different surface finishes under
a severe space confinement will be proposed and discussed.
Figure 1.7 Micrographs showing the typical morphology of (a) 3D IC micro-bump
and (b) flip-chip joint. (Photos provided by ITRI and J. H. Ke)
9
1.3 Introduction of Surface Finishes
Copper base metal on the soldering pads of micro-joints has to be coated with a
surface finish to preserve the wetting property during the storage period before
assembly. These surface finishes are usually consisted by multi-layers in different
levels of assembly sequence. However, to use the lead-free solder requires higher
assembly temperature and places increased demands on the surface finish if it can be
survived after multiple reflow process. The surface finishes can serve several
functions as following purpose:
Providing a solderable coating layer to form a strong solder joint
Offering a medium that is suitable for certain bonding condition during assembly
of components
Protecting the copper pad from oxidation or chemical corrosion during assembly
process
Providing a barrier layer to minimize copper dissolution during assembly process
Four kinds of surface finishes are widely used in PCB fabrication, including:
Organic Solderability Preservative (OSP), Electroless Nickel/Immersion Gold (ENIG),
Electroless Nickel/Immersion Palladium (ENIP), and Electroless Nickel/Electroless
Palladium/Immersion Gold (ENEPIG).
The following is a brief introduction of surface finishes, including configurations,
deposit processes, and key issues:
Organic Solderability Preservative (OSP)
With the increased use of surface mount devices on PCB, OSP have become the
most widely adopted method of surface finishing. Generally, OSP are low-cost,
simple to apply and offer excellent planarity. OSP consists of water-based organic
compound that selectively bonds to Cu and provides an organometallic layer to
protect the Cu pad from oxidation during soldering. The thickness of OSP layer will
typically be in the range of 20-50 nm, as shown in Fig. 1.8 [12]. The OSP layer can
maintain the solderability of the underlying copper until soldering can occur, and the
decomposition temperatures for OSP developed for Pb-free solder assembly is
approximate 260°C. At the early stage of reflow, the OSP will be completely
decomposed and the solder then react with the exposed Cu to form Cu-Sn IMCs. An
overview of the benefits and concerns of the OSP surface finish is summarized in
table 1.1 [12].
Table 1.1 Process Summary of Organic Solderability Preservative (OSP). [12]
12
Electroless Nickel/Immersion Gold (ENIG)
A surface finish consisted of electroless Ni/immersion Au (ENIG) was widely
used for hi-end applications. The electroless Ni is deposited on the exposed, activated
Cu surface, and has a thickness ranged from 3-6 μm. The Au layer is typically
deposited by immersion process on the Ni surface, and has a thickness ranged from
50-100 nm, as shown in Fig. 1.9 [12]. The Au layer is to provide good wettability for
soldering, and the Ni layer serve as a diffusion barrier to avoid the excess interaction
between solder and Cu. At soldering, the Au layer is rapidly dissolved into molten
solder forming the AuSn4 in solder matrix, and the solder then react with the exposed
N layer to form a Ni-Sn IMC. An overview of the benefits and concerns of the ENIG
surface finish is presented in table 1.2 [12].
Fig. 1.9 Configuration of ENIG surface finish. [12]
13
Table 1.2 Process Summary of Electroless Nickel/Immersion Gold (ENIG). [12]
14
Electroless Nickel/Immersion Palladium (ENIP)
ENIP surface finish consisted of electroless Ni/immersion Pd. The electroless Ni
is deposited on the activated Cu surface, and has a thickness ranged from 3-6 μm. The
Pd layer is deposited by immersion process on the Ni surface, and the thickness is
about 50-200 nm, as presented in Fig. 1.10 [12]. The Pd layer provides good wetting
property, and the Ni layer is a diffusion barrier. At the early stage of reflow, the Pd
layer will be completely dissolved into the molten solder precipitating the PdSn4
within the solder matrix, and the solder then react with the exposed Ni layer to form a
Ni-Sn IMC. An overview of the benefits and concerns of the ENIP surface finish is
shown in table 1.3 [12].
Fig. 1.10 Configuration of ENIP surface finish. [12]
15
Table 1.3 Process Summary of Electroless Nickel/Immersion Palladium (ENIP). [12]
16
Electroless Nickel/Electroless Palladium/Immersion Gold (ENEPIG)
One of the most popular surface finishes under consideration is ENIG. However,
the deposition process of ENIG meets a serious reliability concern knows as black pad
phenomenon [13-15]. This defect is a result of the Ni surface corrosion during the
immersion Au process. The black pad defect could cause poor wettability and the
resulting deterioration of mechanical properties of solder joints [16]. To solve such
reliability concerns, the ENEPIG tri-layer surface finish was proposed to replace
ENIG [17-18]. The sequence is very similar to the previously described ENIG process,
with the exception of the additional palladium deposition step by electroless process.
The typically thickness of Au layer of the surface finish is about 20-50 nm, and the Pd
layer is about 50-200 nm thickness, as shown in Fig. 1.11 [12]. The Au layer is to
provide an excellent wettability for solders; the Pd layer is to protect the Ni layer from
corrosion, and the Ni layer serve as a diffusion barrier. At soldering, the Au and Pd
layer could be totally dissolved into a molten solder forming the (Au,Pd)Sn4 in solder
matrix, and the solder then react with the exposed Ni layer to form a Ni-Sn IMC.
Table 1.4 briefly summarized the benefits and concerns of the ENEPIG surface finish
[12].
17
Gold (ENEPIG). [12]
under Space Confinement
Cu is a common material used as a metallization pad in electron packaging. Li et
al. have studied the microstructure feature of 3D IC scaled micro-joints. Fig. 1.12
shows the microstructure of Cu/Sn(25 μm)/Cu sandwiches were thermal aged at 340
o C for 10-90 min [4]. Fig. 1.12 (a) shows that Cu6Sn5 and Cu3Sn forming at each
interface at 10 min. When the reaction time extended 40 min, as showed in Fig. 1.12
(b), Cu6Sn5 grains growing from the opposite interfaces started to impinge on each
other. The original Sn was no longer in the form of a continuous layer, but it was
separated by the vertically impinged Cu6Sn5 grains. After aging for 90 min, Sn was
totally consumed and the joint was entirely occupied by IMCs of Cu6Sn5 and Cu3Sn,
as shown in Fig. 1.12 (c). According to these results, 3D IC micro-joint is
demonstrated having a chance to become an all-IMCs joint. Consequently, the
mechanical properties of IMCs and other reaction-induced microstructure features
play significant roles.
19
Fig. 1.12 SEM images taken from the polished cross-sections of the Cu/Sn/Cu
samples reflowed at 340 o C for (a) 10 min, (b) 40 min, and (c) 90 min. [4]
As the size and pitch of solder joints decrease rapidly, the reliability of the
package is more and more dependent on the mechanical properties of the solder joint.
Two kinds of microstructures can be present in the micro-joints during assembly
process or after long term storage at high temperature [19-21]. One is the joints with
remaining solder, and the other one is the joints fully occupied by IMCs, as shown in
Fig.1.13 (a) and (b), respectively. Under such a condition, the mechanical properties
of a solder joint is no longer dominated solely by the properties of solders. Now the
mechanical properties of the IMCs and other reaction-induced microstructure features
play dominant roles. As IMCs tend to be hard and brittle, a solder joint composed
mainly of IMCs might have peculiar mechanical properties.
20
Fig. 1.13 Two kinds of microstructures can be present in the micro-joints: (a) joint
with remaining solder and (b) joint fully occupied by IMCs.
Die shear test has been widely used to determine joint strength [22-24]. Most
studies on the mechanical properties of Cu pillar solder bumps have been reported in
the whole package [25-27]. However, the die shear test includes the effects of solder
joints, underfill, and the joints geometry. Therefore, the standard deviation of the ball
shear strength is pretty large.
Kim et al. used die shear test to investigate the effect of IMC growth on the
mechanical bonding strength of Cu pillar bumps [25]. The mechanical properties of
Cu pillar solder bumps in this case were measured in a package. The height and
diameter of the Cu pillars were 50 μm and 80 μm, respectively. The thickness of the
solder was 8 μm, as presented in Fig. 1.14. The die shear test results of the solder
joints that were aged at 150°C for 0–300 h were presented in Fig. 1.15. The results
revealed that the die shear force decreased with aging time, but the correlation
21
between the microstructure and bonding strength were not discussed clearly in
literatures [25-27].
Fig. 1.14 Cross-sectional SEM-BSE images of Cu pillar bumps after bonding. [25]
Fig. 1.15 Die shear force as a function of annealing time. [25]
22
In present study, a ball shear test is conducted to determine the joint strength of a
single Cu pillar solder bump, and the standard deviation of the ball shear test data is
found to be relatively small. The advantage of shearing one single solder joint at a
time is that the properties of the solder joint can be directly evaluated.
The mechanical properties of the 3D IC scaled micro-joints are very sensitive to
the solder reaction and the microstructure of IMCs [28]. In addition, the evolution of
IMCs can also affect the failure modes. Therefore, the purpose of this study is to
investigate the effects of IMCs on the single-joint shear strength of micro-joint and
uncover which microstructure, joints with remaining solder or joints fully occupied by
IMCs, has better mechanical properties. Moreover, this study also generalizes a series
of failure mechanisms for different microstructures.
23
1.5 Interfacial Reactions between Lead-free Solder and ENIG, ENIP, ENEPIG
Surface Finishes (Au and Pd Embrittlement Issues)
Popular surface finishes under consideration include Au/Ni (ENIG), Pd/Ni
(ENIP), and Au/Pd/Ni (ENEPIG). During soldering, the Au or/and Pd layers could be
dissolved into solder rapidly [29–35], exposing the Ni layer below. Previous results
indicate that the role of Au and Pd is very similar in soldering reactions [36].
In SnPb solder joints, the Ni layer reacts with Sn and a continuous Ni3Sn4 layer
forms at the interface between SnPb and Ni. During the solidification stage of the
reflow, the dissolved Au atoms precipitate out as many needle-like AuSn4 particles
dispersed in the SnPb matrix. During the subsequent aging, a large amount of AuSn4
particles in solder matrix migrated to the SnPb/Ni3Sn4 interface and formed a
continuous (Au,Ni)Sn4 layer after thermal aging [37], as shown in Fig. 1.16. The
mechanism of (Au,Ni)Sn4 resettlement is to seek Ni for a lower Gibbs-free energy
[33]. From previous studies, the Au embrittlement phenomenon will happen if the
effective Au concentration in solder is higher than 0.1 wt% in SnPb solder joints [37].
The presence of continuous AuSn4 is a key reason causing the so-called Au
embrittlement effect. According to the ball shear teat results, the interface between
continuous (Au,Ni)Sn4 and Ni3Sn4 is very weak [35], as presented in Fig. 1.17. It
could deteriorate the mechanical properties of solder joints significantly.
24
Fig. 1.16 BSE micrograph showing the (Au0.4Ni0.6)Sn4 layer formed in a solder joint
that had been aged at 160°C for (a) 25 h, (b) 100 h, (c) 500 h, and (d) 1000 h. The
solder was eutectic SnPb, and the soldering pad was 1 mm Au over Ni layer. [37]
Fig. 1.17 Fracture surfaces for solder side and pad side of the same solder joint that
had been aged at 160°C for 500 h. The (Au,Ni)Sn4/Ni3Sn4 interface is very weak. [35]
25
Although the Au embrittlement seems no longer a problem since the lead-free
(high-Sn) solders was used in soldering [38-41], Ho et al. proposed that the Au
embrittlement is still a potential concern for lead-free solder system if the Au
concentration is high enough [42]. In lead-free solder joints, a continuous (Au,Ni)Sn4
layer could form if the Au concentration in solder was higher than 0.8 wt.%, and a
continuous (Au,Ni)Sn4 layer would not form if the Au concentration was lower than
0.4 wt.% [42], as shown in Fig. 1.18.
Fig. 1.18 BSE micrographs showing the reaction between Sn3.5Ag balls and Au/Ni
substrates of different Au thicknesses at 160°C for 1,000 h. (a) 0 μm Au, (b) 0.2 μm
Au, (c) 1.0 μm Au, (d) 2.0 μm Au, and (e) over view of Fig. 1.14(d). [42]
26
Ho et al. proposed that the Pd embrittlement is also a critical issue in lead-free
solder if the Pd concentration is high enough. In lead-free solder joints, a continuous
(Pd,Ni)Sn4 layer could form if the Pd concentration in solder was higher than 0.2
wt.% [43], as shown in Fig. 1.19. As presence of continuous (Pd,Ni)Sn4 layer in the
joint, the (Pd,Ni)Sn4/Ni3Sn4 interface is the weakest place of the joint [43], as
presented in Fig. 1.20. In general flip-chip and BGA size, the effective Au or/and Pd
concentrations is below this threshold and will not cause any embrittlement problem.
Fig. 1.19 BSE micrographs showing the Sn–xPd/Ni interfaces after reaction at 250 o C
for 2 min. The x were (a) 0.05, (b) 0.1, and (c) 0.2, respectively. [43]
Fig. 1.20 Cross-sectional view showing the fracture of a Sn–0.2Pd/Ni joint after shear
test. [43]
27
In the BGA scale packaging, the Au or/and Pd layer is often deposited by an
electroplating process, and has a thickness as high as 1 μm. This Au or/and Pd
embrittlement problems can be resolved in the short term by reducing the Au or/and
Pd thickness, so that the effective Au or/and Pd concentration in solder is lowered. In
many of three-dimensional integrated circuits, 3D ICs, packaging schemes under
development today, micro solder bumping is a very promising technology enabling
chip-to-chip and chip-to-wafer stacking. One most well-perceived character is that the
solder volume of the micro solder joint is extremely small, about 1000 times small
than that of a flip-chip joint. Due to this solder volume reduction, the effective Au
or/and Pd concentration in such micro solder joints can become very high. From this
point of view, the Au or/and Pd embrittlement issues may become relevant again in
the 3D IC scaled packaging. The objective of the present study is to further
investigate the Au or/and Pd embrittlement problem in micro-joints. The extent and
effects of this issue will be studied and discussed.
28
This research investigates the interfacial reactions between 3D IC micro-joints
and different substrates and related mechanical properties. It is divided into two parts.
The first part studies the mechanical properties and interfacial reactions of 3D IC
micro-joints. Two kinds of microstructures can be present in Cu/Sn/Cu sandwiches
during assembly process or after long term storage: (1) joints with remaining solder
and (2) joints fully occupied by IMCs. The objective is to estimate the mechanical
properties of micro-joints with those two kinds of microstructures, and discuss the
correlations between the shear strength and the evolution of IMCs during the aging
process. Finally, this study generalizes a series of failure mechanisms for different
aging times.
The second part studies the interfacial reactions and mechanical properties of 3D
IC micro-joints with different surface finishes under space confinement, including I.
ENIG (Ni/Sn-Au/Ni), II. ENIP (Ni/Sn-Pd/Ni), III. ENEPIG (Ni/Sn-Au-Pd/Ni). In
3D IC packaging, the size of solder joints becomes smaller, that can result in Au
or/and Pd embrittlement as the Au or/and Pd concentration in the solder is too high.
The purpose of this study is to investigate the Au or/and Pd embrittlement issues in
space confined soldering reaction for 3D IC applications. We also want to explore
new effective approaches to eliminate the Au or/and Pd embrittlement problem.
29
2.1 Cu Pillar Solder Bump (Cu Pillar Bump/Solder Cap)
The flow chart of the Cu pillar solder bump fabrication process used in this study
is shown in Fig. 2.1. A 4-inch (111) Si wafer was sputtered with a 300-Å -thick Cr
layer, and a 3000-Å -thick Cu layer was deposited to serve as a seed layer (Fig. 2.1 (a)).
Thick negative photoresist (KMPR 1000) was spin coated on the wafer as a plating
mask. The thickness of the photoresist (PR) was 50 μm and the diameter of the pattern
was 50 μm (Figs. 2.1 (b)–(c)). Then, the Cu pillar bump with a thickness of 35 m
was electroplated at room temperature using a commercial CuSO4 electroplating
solution with a commercial additive (Fig. 2.1 (d)). After the plating, the PR was
stripped off using acetone at 55°C (Fig. 2.1 (e)). Next, the stencil printing process was
used to deposit the solder ball. The SAC405 solder ball with a 50 m diameter was
supplied by Senju Metal Industry Co., Ltd. The stencils acted as a mask to print flux
and deposit the solder ball. A mildly activated rosin flux was applied on the surface of
the Cu pillar bump using a commercial screen printer, before depositing the solder
ball on the top surface of the Cu pillar bump (Figs. 2.1 (f)–(g)). The constant volume
of the micro-ball ensured the bump height was uniform, and the composition of minor
elements in the micro-ball was accurately controlled. The reflow process was carried
out in a vacuum quartz furnace at a peak temperature of 260°C for 10 min in a
30
nitrogen atmosphere (Fig. 2.1 (h)).
Fig. 2.2 shows the as-fabricated Cu pillar solder bump structure. To avoid
artifacts forming during the mechanical polishing, all the cross-sectional simples were
prepared by focus ion beam (FIB), and the inset shows the central plane of the
samples at which the FIB was directed. The diameter and height of both the Cu pillar
bump and the solder cap were 50 m and 35 m, respectively.
Fig. 2.1 The flow chart of the Cu pillar bump with solder cap fabrication.
31
Fig. 2.2 As-fabricated Cu pillar solder bump structure.
The specimens were aged at 180°C for different time durations (0, 150, 300, 500,
750 h). The ball shear test was performed using a DAGE4000 micrometer. The shear
loading speed was 100 μm/s; the test load was 250 g, and the maximum load was 5 kg.
Shear tests for two different shear heights, 55 μm (position A) and 38 μm (position B),
were carried out in order to simulate the strengths of the joints with Sn remained and
the joints fully occupied by IMCs, respectively, as presented in Fig. 2.3. Each mean
value of shear strength was averaged from 50 specimens.
32
Fig. 2.3 Schematic drawings showing the different shear position (a) shearing at
solder matrix and (b) shearing at IMCs.
The cross-sectional microstructure was observed using a focused ion beam
scanning microscope (FIB-SIM), and the fracture surface of solder joints after shear
test was observed by scanning electron microscopy (SEM) with a back-scattered
electron image (BEI) detector. The composition of the IMCs was analyzed
quantitatively using an energy dispersive spectrometer (EDS). The Cu-Sn IMC area
was obtained from image analysis and the IMC thickness was obtained by dividing
this area with the interfacial length between the IMC and Cu pillar bump.
33
2.2 Ni/Sn-Au/Ni Sandwich (ENIG) and Ni/Sn-Pd/Ni Sandwich (ENIP)
In the 3D IC packaging, the solder volume of the micro-joints is extremely small.
Due to this solder volume reduction, the effective Au or/and Pd concentration in
micro-joints can become very high. It is assumed that the Au or/and Pd layer is
completely dissolved into the molten solder. The effective Au and Pd concentration in
a solder joint can be calculated with the following equations
CAu(wt. %) = 100 MAu
MSn + MPd
where CAu and CPd are the effective Au and Pd concentration. MAu, MPd, and MSn
represent the mass of Au, Pd, and Sn in a solder joint, respectively. For easy
understanding, the solder joint is simplified as a cylindrical structure, as shown in Fig.
2.4. Accordingly, the equation can be written as
CAu(wt. %) = 100 πr2dAuρAu
πr2dSnρSn + πr2dPdρPd
The dSn, dAu, and dPd are the thickness of Sn, Au, and Pd, respectively. And, the
density of Sn, Au and Pd are assumed as ρSn=7.3 g·cm −3
, ρAu =19.3 g·cm −3
, and ρPd
34
Fig. 2.4 Schematic drawings showing (a) the simplified solder joint and (b) the
sandwich configuration.
As ENIG was served as a surface finish, let’s consider a joint with dSn=10 μm,
but dAu=30, 50, 100 or 150 nm, which are practical conditions for 3D IC packaging.
Assuming Au is completely dissolved into solder, the resultant solder compositions
are Sn0.8Au, Sn1.3Au, Sn2.6Au and Sn3.9Au (wt.%), as shown in table 2.1.
As ENIP was served as a surface finish, let’s consider a joint with dSn=10 μm and
the surface finish dPd=10 or 30 nm. It is assumed that Pd is totally dissolved into the
molten solder, and the resultant solder compositions are Sn0.29Pd and Sn0.8Pd
(wt.%), as shown in table 2.1.
35
Table 2.1 The experimental thickness of Au/Pd layer and the effective Solder
compositions that Au/Pd layer dissolved into.
Four Sn-Au alloys (Sn0.8Au, Sn1.3Au, Sn2.6Au and Sn3.9Au (wt.%)) and two
Sn-Pd alloys (Sn0.29Pd and Sn0.8Pd (wt.%)) were aged in vacuum at 800 o C for 240
h, and then quenched in water. The alloy compositions were verified using an
inductively coupled plasma atomic emission spectroscopy, ICP-AES, to confirm the
composition uncertainty is acceptable.
Fig. 2.5 schematically displays the sandwiched samples that were prepared by
bonding process. Upper die and lower die were prepared by electroplating Ni layer on
the Si wafers with the pre-sputtered Cr (300 )/Ni (3000 ) seed layer. Each alloy
was then used to join sandwich samples at 250 o C. A bonding process was applied to
prepare the sandwiches. Before bonding, a fluxed solder sphere was placed onto the
lower die, and the upper die was then upside down placed on the solder sphere. Lower
36
die, together with the upper die and the in-between solder sphere were kept on a hot
plate at a fixed temperature. A force was loaded on the upper die pushing it toward to
the lower die when the solder is molten. The force was expected to ensure a close
contact between the upper and the lower die, but it has no effect on the joint quality
since the force was only loaded on the spacer, which was designed to determine the
solder thickness, rather than on the soldering interface. The heating time is 1 min
measured after the solder is molten, and the sample was then cooled down to room
temperature. The as-fabricated Ni/Sn-Au(10μm)/Ni and Ni/Sn-Pd(10μm)/Ni
sandwiches are shown in Fig. 2.6(a)-(b), respectively.
The sandwiches were then aged at 200 o C for microstructure characterisations.
The thickness of the solder layers was 10 m, confined with a spacer between two
chips. Ignoring the edge regions, the lateral dimensions of each sandwich were 1 mm
× 1 mm. Under this configuration, the space was confined along the vertical direction;
and along the lateral directions, the space could be considered as infinite. Thus, the
direction of the interaction fluxes can be considered limited to along the vertical
direction.
37
Fig. 2.5 Schematic drawings showing the sample preparation via hot-pressing
bonding.
38
2.3 Ni/Sn-Au-Pd/Ni Sandwich (ENEPIG)
The specimens are commercial micro-joints that were supplied by ASE, Inc. The
substrates were printed circuit boards (PCBs) with 20μm thick Cu pad. A ENEPIG
surface finish was made by sequentially depositing a layer of Au (60 nm), Pd (100
nm), and Ni(P) (6 m) on Cu substrate, schematic diagram from top to bottom. The
chip side had a Cu pillar bump structure whose top surface had been plated with Ni (3
m) to serve as a diffusion barrier. The Cu pillar and the ENEPIG were bonded
together using Sn1.8Ag solder (20m). Reflow process had a peak temperature of
255°C and lasted for 300 sec in a nitrogen atmosphere. The as-fabricated
Ni/Sn1.8Ag/ENEPIG structure is shown in Fig. 2.7. After assembly, the sandwiches
were aged at 200°C for periods of up to 200 h (0, 50, 100, and 200 h).
Fig. 2.7 As-fabricated Ni/Sn-Ag/ENEPIG structure.
39
All the sandwiches (ENIG, ENIP, and ENEPIG samples) were
metallographically polished to reveal the interfaces and the cross-sectional
microstructures of the solder joints. The cross-sectional microstructure was observed
by a scanning electron microscopy (SEM) equipped with a back-scattered electron
image (BEI) detector. The composition of the IMCs was analyzed quantitatively using
an electron microprobe analyzer (EPMA). The IMCs area was obtained from image
analysis and the IMCs thickness was obtained by dividing this area with the
interfacial length between the IMCs and UBMs.
40
2.4 Mechanical Properties of Micro-joints Measured by Nanoidentor
Young’s modulus and hardness of different IMCs and phases were measured
using a nanoindenter (TI 950 TriboIndenter, Hysitron, as shown in Fig. 2.8). The
samples were polished to reveal their cross-sectional microstructure of the solder
joints. A Berkovich diamond indenter was pressed into the specimen at a load of 1200
μN, and the diameter of the diamond tip is 100 nm. The indenter was then held at the
peak load for 5 sec before it completely withdrew from the specimen, and the
recorded data were used to compute the mechanical properties of the specimen. Each
mean value of elasticity modulus and hardness were calculated by averaging seven
data.
41
3. RESULTS AND DISCUSSION
3.1 Interfacial Reaction and Shear Strength of Cu Pillar Solder Bump
3.1.1 Interfacial Reaction of Cu Pillar Solder Bump
The FIB-SIM image of the cross-sectional joint as-reflowed is shown in Fig. 3.1
(a). Cu6Sn5 and Cu3Sn were observed at the interface of the solder and Cu. The
Cu6Sn5 exhibited typical scallop-type morphology. The discontinuous Cu3Sn was
marked by a circle and no void formations were observed at the interface between
Cu3Sn and Cu.
Fig. 3.1 (b) shows the FIB-SIM image of the cross-sectional joint aged at 180°C
for 150 h. The Cu6Sn5 grains grew through a ripening process, so the Cu6Sn5
morphology transformed from a scallop to a planar type after aging owing to a
decrease in interfacial energy [45]. A discontinuous Cu3Sn layer was observed
forming between the Cu6Sn5 and Cu. This phenomenon had been mentioned by Liu
and Yu. [46-48]. In Liu’s case, the discontinuous Cu3Sn layer was attributed to the
massive formation of Kirkendall voids at the Cu3Sn/Cu interface under particular
current stressing conditions. However, the formation of discontinuous Cu3Sn in our
case is much similar to that in Yu’s study [46]. They ascribed the morphology to S
addition in the commercial Cu electroplating bath. Unfortunately, the exact
mechanism still remains uncertain and needs further studies. The chemical additives
42
in electroplating bath are intended to improve the plating qualities such as uniformity
and surface roughness, etc., which are crucial especially for fabricating fine pitch and
high aspect ratio joints.
Figs. 3.1 (c)–(e) show FIB-SIM images of the cross-sectional joint aged at 180°C
for 300, 500, and 750 h, respectively. Cu6Sn5 displayed the planar-type morphology.
The Cu3Sn layer was discontinuous, and the layer area increased with aging time.
After over 300 h of aging, a series of micro void formations were observed along the
Cu3Sn/Cu interface and the amount of void formations increased with aging time, as
shown in Figs. 3.2 (a)–(c). Furthermore, it can be clearly observed that the multiple
fine Sn grains in the as-reflow condition coarsened into 1-3 large Sn grains after
thermal aging, as shown in Fig. 3.
43
Fig. 3.1 FIB-SIM images of the cross-sectioned Cu pillar solder bump that had been
aged at 180°C for (a) 0 h, (b) 150 h, (c) 300 h, (d) 500 h, and (e) 750 h.
44
Fig. 3.2 Magnified images of rectangle regions of Figs. 3.1 (c)–(e). Voids were
formed at the Cu3Sn/Cu interface.
45
3.1.2 Kinetics of IMCs Growth
Fig. 3.3 shows the average thickness of IMCs (Cu6Sn5, Cu3Sn) as a function of
aging time (0–750 h) at 180°C. The total thickness of Cu-Sn IMCs is the sum of the
layer thicknesses of Cu6Sn5 and Cu3Sn. The layer thicknesses of Cu6Sn5 and Cu3Sn
increased with prolonged aging time. The growth kinetics followed a parabolic
relationship. After the aging process, over 25% of the height of the solder joint was
occupied by IMCs.
Fig. 3.3 The thickness of IMCs formed at 180°C against the square root of reaction
time.
Aging time at 180 oC (hr1/2)
0
2
4
6
8
10
12
Shear Strength of Solder Joint
The shear test results of the individual solder joints that were aged at 180°C for
0–750 h are presented in Fig. 3.4. The shear strength of the solder joint as-reflowed
was 61.4 MPa. In previous studies [49-50], the shear strength of the solder joint in
BGA or at a micro-ball scale was reported to be approximately 58–65 MPa. Although
the solder volume in this study is different from that of previously reported studies,
the strength of the Cu pillar solder bumps is of the same order of magnitude.
The result shows that the shear strength decreases with aging time. The
as-reflowed solder joint showed the highest shear strength and there was a 22 MPa
(36%) decrease in shear strength at the end of the entire aging process (750 h).
Interestingly, the first 150 h of aging caused the most significant degree of shear
strength decrement, which equalled 22% degradation. A further 600 h of aging only
caused an additional 14% decrease in shear strength. Many previous studies have also
indicated that the joint strength would degrade with the aging process [25, 28].
However, the degree of weakening in the joint strength and the joint failure modes
found in this study showed dissimilarities from previous studies. The present study
shows that there is a new mechanism that can cause a significant solder joint shear
strength decrease by simply aging for 150 h.
47
Fig. 3.4 The shear strength as a function of aging time at 180°C.
Fracture Surface and Failure Mode
In previous studies, four types of failure modes were identified in the fracture
analysis for the ball shear test [51-52]. The failure modes include ductile (100% area
with solder), brittle modes (almost without solder), and two mixed failure modes, i.e.,
quasi-ductile (or <50% area with exposed pad) and quasi-brittle (or >50% area
without solder) modes, as shown in Fig. 3.5. In this investigation, only ductile,
quasi-ductile, and quasi-brittle modes were observed in the ball shear test.
0 200 400 600 800
Aging time at 180 oC (h)
30
40
50
60
70
48
Fig. 3.5 Ball shear failure modes (a) Ductile mode, (b) Quasi-ductile mode, (c)
Quasi-brittle mode, and (d) Brittle mode. [51]
49
In order to assess the elemental distribution of the fracture surface, a series of
SEM investigations with EDS elemental mapping were carried out on the samples
aged at 180°C for 0–750 h, as shown in Fig. 3.6. The EDS elemental mappings for the
elements Cu (red), Sn (green), and Ag (blue) were presented. The shear direction was
from left to right. The compositions were marked on the micrograph, which was
determined by EDS analysis. The Sn element area was marked by a dotted circle, and
the Cu element around the fracture surface was the Cu seed layer.
The fracture surface and failure modes changed with the aging process. Only Sn
was observed on the fracture surface after reflow, as shown in Fig. 3.6 (a). After 150 h
of aging, the fracture surface was occupied by Sn and Cu6Sn5, as shown in Fig. 3.6 (b).
As the aging time reached 300 h, Cu was exposed on the fracture surface alongside Sn
and Cu6Sn5, as shown in Fig. 3.6 (c). Similar fracture surfaces were observed after
aging for 500 h and 750 h, with Cu6Sn5 and Cu occupying the fracture surfaces along
the shear direction, respectively, as shown in Figs. 3.6 (d) and (e).
The phase percentages of fracture surface for different aging time were
calculated, as shown in Fig. 3.7. 100% of the fracture surface was occupied by Sn
after reflow. After 150 h of aging, approximately 60% of the fracture surface was
occupied by Sn and the rest of the fracture surface was occupied by Cu6Sn5. This
indicated that the joints experienced shear failure in a ductile mode in these two
50
conditions. After 300 h of aging, Sn retained approximately 25% of the fracture
surface, and Cu6Sn5 and Cu occupied approximately 45% and 30% of the fracture
surface, respectively. This indicated that the joints experienced shear failure in a
quasi-ductile mode. After aging for 500 h, less than 5% Sn remained on the fracture
surface, and Cu6Sn5 and Cu occupied approximately 45% and 50% of the fracture
surface, respectively. The joints experienced shear failure in a quasi-brittle mode.
Similar fracture surfaces and failure modes were observed as aging time reached 750
h. The only difference was that the area of exposed Cu on the fracture surface
increased to 60%. The area of Sn on the fracture surface decreased with aging time. In
contrast, the area of exposed Cu on the fracture surface increased with aging time.
This indicated that the failure mode transformed from ductile mode to quasi-ductile
mode and finally to quasi-brittle mode.
51
Fig. 3.6 The fracture surfaces of solder joints that have been aged at 180°C for (a) 0 h,
(b) 150 h, (c) 300 h, (d) 500 h, and (e) 750 h.
52
Fig. 3.7 The phase percentages of fracture surface against aging time.
Failure Mechanism
By combining the findings shown in Fig. 3.1 and Fig. 3.6, a schematic drawing
summarizing the results of fracture analysis is presented as Fig. 3.8. The
cross-sectional view of the Cu pillar solder bump microstructure is shown on the first
row of this figure, and the corresponding top view of the fracture surface on the pad
side is shown on the second row. By comparing the fractographs of the cross-sectional
view and the top view, the possible fracture paths were presented by a dotted line for
each condition.
After reflow, the Cu6Sn5 exhibited scallop-type morphology, and it was found
that the fracture occurred within the Sn matrix, leaving the exposed surface to be
totally occupied by Sn. As the aging time reached 150 h, the Cu6Sn5 layer became
relatively smooth. Interestingly, the fracture surface shifted to the Cu6Sn5/Sn interface.
53
Thus, the exposing surface was a combination of Cu6Sn5 and Sn. In the 300–750 h
aged samples, voids formed at the Cu3Sn/Cu interface, and the amount of voids
increased with aging time. Therefore, cracks propagated along the solder/Cu6Sn5
interface and also extended to the Cu3Sn/Cu interface. The corresponding fracture
morphology along the shear direction was made of Cu6Sn5 and Cu. As the amount of
voids increased, the percentage of the exposed Cu area on the fracture surface
increased correspondingly.
A possible reason for the higher shear strength in the joint with scallop-type
Cu6Sn5 is that this morphology could disturb the crack propagation because of the
extremely rough interfacial morphology between the compound and the multiple fine
Sn grains. Contrarily, the smooth interface between planar Cu6Sn5 and large
coarsened Sn grains was nearly parallel to the shear direction after aging, so the crack
could probably propagate along this interface. In addition, the void formation could
deteriorate the mechanical properties of the joint [53-55]. Thus, the fracture can
propagate at the interface between Cu3Sn and Cu.
The microstructure characterizations suggest that the key reason for the
degradation of joint strength is not void formations, but the planarization of the
Cu6Sn5 morphology. Hence, it is very important to preserve a scallop-type Cu6Sn5 in
the remaining solder to achieve a better joint strength.
54
Fig. 3.8 Schematic illustration showing the fracture mechanisms during aging at
180°C from 0-750 h.
55
3.1.4 Shear Strength of the Joints Fully Occupied by IMCs
Shear Strength of Solder Joint
The joints fully occupied by IMCs and the joints with remaining unreacted solder
were aged at 180°C for 0–750 hours as shown in Fig. 3.9. For the entire joint
occupied by IMCs, the shear strength decreased with aging time. The highest shear
strength of solder joint happened in the initial as-reflow condition which was 107.6
MPa. The shear strength decreased to 99.8 MPa after 150 hours of aging. After 300
hours of aging, the shear strength of the solder joint greatly dropped down to 76 MPa.
As aging time reached 500 hours and 750 hours, the shear strength of the solder joint
slightly decreased to 67.5 MPa and 65.4 MPa, respectively. There was a 42 MPa
(39%) decrease in shear strength at the end of the entire aging process (750 hours).
The first 150 hours of aging only caused 8% degradation as compared to the as-reflow
stage. In the period of 150–300 hours aging time, 22% significant reduction of shear
strength was observed. With further 450 hours of aging process, from 300–750 hours,
it only reduced further about 10% more than the previous aging time.
By comparing with the joint strengths of two cases (all-IMCs joints and joints
with remaining unreacted solder), the shear strength decreases with aging time in both
cases. The joints fully occupied by IMCs had stronger joint strength than the joints
with remaining unreacted solder, and the one with fully occupied by IMCs is as higher
56
as 1.5 times in strength than the one with remaining unreacted solder in all conditions.
It should be noted that the minimum strength of completely reacted IMCs joints,
which was 750 hours of aging process, was even higher than the maximum strength in
the joints of the remaining unreacted solder. Hence, the presence of fully IMCs in the
joints during assembly bonding process is an important determination to obtain a
better joint strength.
Fig. 3.9 The shear strength as a function of aging time at 180°C (a) the joints fully
occupied by IMCs, and (b) the joints with remaining solder.
57
Fracture Surface
Solder Joints were aged at 180°C for 0-750 hours. Shear tests were performed
with different aging times and then their fracture surfaces were analyzed by the EDS
elemental mapping as shown in Fig. 3.10. The EDS elemental mappings for the Cu
element signal (red), Sn (green), and Ag (blue) were presented. The phases were
marked on the micrographs, which were determined by EDS analysis. The Cu element
around the fracture surface was the Cu seed layer.
Shear tests are conducted by setting the shear height solely on the IMCs of the
micro-joints. We are able to determine the failure mode of solder joints are mostly
brittle mode with observation by SEM. However, even though their primary failure
mechanism was brittle mode, the morphology of fracture surface varied with different
aging samples. Shear direction of shear test was performed from left to right. Similar
fracture surfaces were observed for the as-reflow sample and the sample with 150
hours aging time. They both had Cu6Sn5 and Cu occupying on the fracture surfaces
along the shear direction as shown in Figs. 3.10 (a) and (b). For the sample with aging
over 300 hours (300-750 hours), we noticed that majority of fracture surfaces was on
the Cu region with partly Cu3Sn remaining on the fracture surfaces as shown in Fig.
3.10 (c)-(e).
58
As shown in Fig. 3.11, the phase percentages of the fracture surface for different
aging time were calculated. For as-reflow sample, approximately 71% of the fracture
surface was occupied by Cu6Sn5 and the rest of the fracture surface was occupied by
Cu. With the sample aging for 150 h, approximately 48% of the fracture surface was
Cu6Sn5 and the rest of fracture surface was Cu. For the sample with aging for 300 h,
the percentages of Cu3Sn and Cu on the fracture surface were approximately 32% and
68% respectively. Similar fracture surfaces were observed as aging time reached 500
h and 750 h and we had observed less than 10% of Cu3Sn remained on the fracture
surface with the rest occupied by Cu. As aging time increased, the presence of Cu-Sn
IMC on the fracture surface decreased. On the other hand, the Cu region on the
fracture surface increased as aging time increased.
59
Fig. 3.10 The fracture surface of solder joint that had been aged at 180 for (a) 0 h,
(b) 150 h, (c) 300 h, (d) 500 h, and (e) 750 h.
60
Fig. 3.11 The percentages of the different phases on the fracture surface against the
aging time.
Failure Mechanism
To uncover failure mechanism of fully IMCs solder joints, we compared and
analysed their cross-sectional view of the microstructure and the top view of the
fracture surface on the pad side, and schematic illustration of fracture analysis was
shown in Fig. 3.12. We proposed the potential fracture path was initiated and followed
by the white dotted line indicating on the diagram of cross-sectional view of the
microstructure for each condition.
No void was observed at the interface between Cu3Sn and Cu on the as-reflow
sample and the sample with aging for 150 h. The distribution of fracture surfaces was
similar for both of them in which Cu6Sn5 and Cu were occupied almost to the same
degree. The crack propagated transgranularly across the Cu6Sn5 grain, and finally
61
extended to Cu3Sn/Cu interface due to the very concentrated stress at the final stage of
shearing. So actually we may find this morphology occur every samples.
With the sample aging for 300–750 h, voids formed at the Cu3Sn/Cu interface,
and the amount of voids increased with aging time. Due to the formation of voids, the
fracture surface shifted to the Cu3Sn/Cu interface and the morphology of fracture
surface appeared to be limited Cu3Sn existed within Cu. As the amount of voids
increased with aging time, the percentage of Cu region on the fracture surface
increased correspondingly.
With the knowledge of void formation could deteriorate the mechanical
properties of the solder joints [53-55], crack could propagate at the Cu3Sn/Cu
interface after a long-term aging process. According to the shear test result, we
observed the maximum reduction in shear strength occurred on the sample of 150-300
hours aging. It may imply that the void formation could induce about 29% reduction
in shear strength. Consequently, the void formed at the Cu3Sn/Cu interface was the
main factor leading to the tremendous reduction of the joint strength.
62
Fig. 3.12 Schematic illustration showing the failure mechanism during aging at 180
for 0-750 h.
According to the microstructure characterizations, the solder joints with no void
had higher joint strength. After over 300 h of aging, void formed at the Cu3Sn/Cu
interface could cause a significant reduction of the joint strength. As a result, void
formation was the main factor to deteriorate the mechanical properties of the joints.
Thus, it is very important to further investigate and improve the reliability of joints by
inhibiting void formation. Fortunately, we are able to hinder the void formation by
introducing the alloying of Ni, Zn, etc. into soldering reaction [56-59].
63
Confinement
In 3D ICs packaging, the solder volume of the micro-joint is very small. Due to
such a small solder volume, the effective Au concentration in micro-joints can
become very high. In addition, the interfacial reaction proceeds under space
confinement, so the solder joints will fully occupy by IMCs. Therefore, the Au
embrittlement issues concerning joint properties may be arisen in 3D IC micro-joints.
To examine this issue, Sn-Au alloys with a series of Au contents, Sn0.8Au,
Sn1.3Au, Sn2.6Au, Sn3.9Au (wt.%), were prepared to react Ni substrate.
Experimentally, Ni/Sn-Au(10 μm)/Ni sandwiches were hot-pressing bonded at 250 o C
and then aged at 200 o C for different aging time.
Fig. 3.13 shows the scanning electron micrographs of Ni/Sn0.8Au(10 μm)/Ni
during the aging process. In the as-bonded condition, aged for 0 h, a thin layer of
Ni3Sn4 formed at each Ni/Sn interface, as shown in Fig. 3.13 (a). A few needles in the
white color distributed within the solder region. The needles with a averaged
composition of 79.9Sn-9.2Ni-10.9Au (at.%) formed in the solder region. The
Au-bearing phase were identified as (Au,Ni)Sn4 according to their chemical
composition analysed by electron microprobe analyzer, EPMA. After aging process,
the Ni3Sn4 layers became thickened, and (Au,Ni)Sn4 particles started to resettled at the
64
interface, as displayed in Fig. 3.13 (b)-(c). The mechanism of (Au,Ni)Sn4 resettlement
is to seek Ni for a lower Gibbs-free energy [33]. This redistribution behavior of
(Au,Ni)Sn4 is well discussed in many previous studies regarding Au embrittlement
issues [38-41]. As the aging time reached 192 h, the Ni3Sn4 layers obviously became
thickened, leaving simply a narrow region consisted of Sn and (Au,Ni)Sn4 phases that
remains in the middle of the joint, as shown in Fig. 3.13 (d). At 240 h, shown in Fig.
3.13 (e), Sn was fully consumed and the Ni3Sn4 layers from the opposite interfaces
had impinged and bridged to each other, leaving isolated (Au,Ni)Sn4 particles
discontinuously distributing in the middle of the joint. By comparing Fig. 3.13 (c), (d)
and (e), it seems that the (Au,Ni)Sn4 particles were somewhat flattened and pushed
toward the centre of the joint when the reaction proceeded. Grain ripening is
considered to be responsible for this peculiar redistribution. However, one important
factor, i.e. the joint space confinement, makes ripening along the lateral direction,
instead of becoming a spherical structure. As the reaction proceeded, (Au,Ni)Sn4
grains eventually impinged the Ni3Sn4 layers growing toward the centre of the joint.
Interaction between the two intermetallic compounds eventually resulted in the
formation of layer-like (Au,Ni)Sn4 particles.
65
Fig. 3.13 Micrographs showing Ni/Sn0.8Au/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 24 h, (c) 96 h, (d) 192 h, and (e) 240 h.
66
In Ni/Sn0.8Au/Ni, the distribution of (Au,Ni)Sn4 particles is discontinuous.
However, one can speculate that (Au,Ni)Sn4 can become a continuous layer as the
effective Au concentration in the solder becomes higher. Fig. 3.14 shows the
microstructure evolution as the original Au content was 1.3 (wt.%). As shown in Fig.
3.14 (a), some needles-like (Au,Ni)Sn4 distributed within the solder matrix. As
indicated in Fig. 3.14 (b) and (c), resettlement of (Au,Ni)Sn4 occurred as the sample
was aged for 24 h. During the aging process, it seemed that (Au,Ni)Sn4 particles were
pushed toward the middle of the joint as illustrated in Fig. 3.14 (b)-(c). At 168 h, as
showed in Fig. 3.14 (d), the (Au,Ni)Sn4 particles became disc-like or ribbon-like
particles with a lateral dimension of about 20 μm.
As the Sn fully consumed, the joints was consisted of (Au,Ni)Sn4 and Ni3Sn4.
Fig. 3.14 (e) shows that although (Au,Ni)Sn4 particles did not yet form a continuous
layer in Ni/Sn1.3Au/Ni, but the lateral dimension of the particle had reached as large
as about 20 m. This morphology, in fact, posted practical concerns of Au
embrittlement due to that the brittle (Au,Ni)Sn4 phase has a chance to become a
continuous layer throughout the solder joint. As in Ni/Sn2.6Au/Ni and Ni/Sn3.9Au/Ni,
shown in Fig. 3.15 (e) and Fig. 3.16 (e), a continuous (Au,Ni)Sn4 layer formed across
the entire of joints, and the entire solder joint was composed of
Ni3Sn4/(Au,Ni)Sn4/Ni3Sn4. It should be noticed that once (Au,Ni)Sn4 particles
67
continuously occupy the middle zone throughout the joint, it may significantly
deteriorate the mechanical strength due to a fact that cracking is found easily
propagating at the vicinity of (Au,Ni)Sn4 [35], and leads to devices failure.
In 3D IC micro solder bumping, the solder joint merely contains a pad diameter
as small as 10 m. Our results indicate that once the Au thickness exceeds 50 nm (1.3
wt.%), (Au,Ni)Sn4 has a chance to form a continuous layer throughout such joints,
and posts the risk of Au embrittlement. Previous studies [37-41] have revealed that
(Au,Ni)Sn4 in solder joint can cause sharply decrease the joint strength. Considering a
joint that contains a continuous (Au,Ni)Sn4 layer across the interfaces. The weakest
point of the joint would probably be at the (Au,Ni)Sn4/Ni3Sn4 interfaces or
(Au,Ni)Sn4 phase. Correspondingly, the safety threshold of the Au metallisation
thickness in such joint configuration is 30 nm (0.8 wt.%). This study gives a warning
that Au embrittlement problems which do not exist in the conventional lead-free
solder joint will become relevant in 3D IC packages.
68
(b) 24 h, (c) 72 h, (d) 168 h, and (e) 192 h.
69
(b) 24 h, (c) 96 h, (d) 144 h, and (e) 168 h.
70
(b) 24 h, (c) 48 h, (d) 72 h, and (e) 96 h.
71
Confinement
To examine the Pd embrittlement issue in 3D IC packaging, Sn-Pd alloys with a
series of Pd contents, Sn0.29Pd and Sn0.8Pd (wt.%), were prepared to react Ni
substrate. Experimentally, Ni/Sn-Pd(10 μm)/Ni sandwiches were hot-pressing bonded
at 250 o C and then aged at 200
o C.
Fig. 3.17 shows the scanning electron micrographs of Ni/Sn0.8Pd(10 μm)/Ni
during the aging process. In the as-bonded condition, shown in Fig. 3.17 (a), a thin
layer of Ni3Sn4 formed at each Ni/Sn interface. Massive facets in the white color
distributed within the solder matrix. The facets with a averaged composition of
80.0Sn-14.2Ni-5.8Pd (at.%) formed in the solder region. The Pd-bearing phase were
identified as (Pd,Ni)Sn4 according to their chemical composition analysed by EPMA.
After aging for 6 h and 9 h, the Ni3Sn4 layers became thickened, and (Pd,Ni)Sn4
particles started to resettled at the Ni3Sn4 interface, as shown in Fig. 3.17 (b)-(c). The
mechanism of (Pd,Ni)Sn4 resettlement is as same as (Au,Ni)Sn4, that is to seek Ni for
a lower Gibbs-free energy. This redistribution behavior of (Pd,Ni)Sn4 is discussed in
previous research regarding Pd embrittlement issues [43]. At 12 h, shown in Fig. 3.13
(d), the (Pd,Ni)Sn4 layers from the opposite interfaces had impinged and bridged to
each other, leaving isolated Sn pocket in the middle of the joint.
72
Fig. 3.18 shows the microstructure evolution as the Pd content decreased to 0.29
(wt.%). In the initial condition, some facet (Pd,Ni)Sn4 distributed within the solder
matrix, as shown in Fig. 3.18 (a). As revealed in Fig. 3.18 (b), resettlement of
(Pd,Ni)Sn4 occurred as the sample was aged for 6 h. During the aging process,
(Pd,Ni)Sn4 particles grew from the opposite interfaces and then bridged to each other,
leaving isolated Sn pocket in the joint, as illustrated in Fig. 3.18 (c)-(d).
As the Sn fully consumed, the joints was consisted of (Pd,Ni)Sn4 and Ni3Sn4. As
in Ni/Sn0.29Pd/Ni and Ni/Sn0.8Pd/Ni, shown in Fig. 3.18 (e) and Fig. 3.17 (e), a
continuous (Pd,Ni)Sn4 layer formed throughout the entire of joints, and the entire
solder joint was composed of Ni3Sn4/(Pd,Ni)Sn4/Ni3Sn4, and more than half of the
joint is occupied by (Pd,Ni)Sn4. It should be noticed that once (Pd,Ni)Sn4 particles
continuously occupy the middle zone across the joint, it may significantly deteriorate
the mechanical strength due to a fact that cracking is found easily propagating at the
vicinity of (Pd,Ni)Sn4 and the (Pd,Ni)Sn4/Ni3Sn4 interface [43].
In 3D IC micro-joints, the dimension of solder joint is about 10 m. The results
reveal that even the Pd layer below 10 nm (0.29 wt.%), (Pd,Ni)Sn4 can form a
continuous layer across the joints, and posts the risk of Pd embrittlement. For
practical purposes, it is nearly impossible to keep Pd metallisation thickness below 10
nm when the immersion or electroless process is used.
73
As a matter of fact, the practical Au or Pd thickness is generally greater than 10
nm. For some particular cases, Au or Pd layer can be as thick as 500~1000 nm when
the electroplating process is used. Accordingly, it is argued that the so-called Au or Pd
embrittlement indeed becomes relevant in 3D IC packaging, even with the use of
Sn-based lead-free solders which posts no Au or Pd embrittlement concern in flip-chip
and BGA size scale packaging. The main reason to be responsible for the
re-emergence of Au or Pd embrittlement is the high effective Au or Pd concentration
in the joint due to the very small solder volume.
74
Fig. 3.17 Micrographs showing Ni/Sn0.8Pd/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 6 h, (c) 9 h, (d) 12 h, and (e) 16 h.
75
Fig. 3.18 Micrographs showing Ni/Sn0.29Pd/Ni sandwiches aged at 200 o C for (a) 0 h,
(b) 6 h, (c) 24 h, (d) 36 h, and (e) 54 h.
76
Confinement
ENEPIG contains the Au and Pd, so the Au embrittlement issue should be
concerned along with Pd embrittlement in this case. To examine this issue, the
commercial micro-joints, Ni/Sn-1.8Ag(20 μm)/ENEPIG sandwiches, aged at 200°C
for 0-200 h, as shown in Fig. 3.19. After reflow, both the Au and Pd layers
disappeared from the interface, and a number of Au, Pd, Ni, Sn–bearing particles
distributed over the solder matrix, as shown in Fig. 3.19 (a). These particles were
identified as (Au,Pd,Ni)Sn4 IMCs examined by EPMA. Besides, a thin layer of Ni3Sn4
was observed at both interfaces. After aging at 200°C for 50 h, top and bottom layers
of Ni3Sn4 was seen to grow thicker, and the (Au,Pd,Ni)Sn4 particles coarsened into
larger ones, as shown in Fig. 3.19 (b). It was noted the overall amount of
(Au,Pd,Ni)Sn4 also have been increased. After 100 h of aging, as shown in Fig. 3.19
(c), (Au,Pd,Ni)Sn4 particles coarsened into even larger ones, and many of them
bridged across both interfaces. As aging time reached 200 h, as shown in Fig. 3.19 (d),
all the Sn completely consumed, and the entire solder joints were occupied by three
continuous layers in the sequence of Ni3Sn4, (Au,Pd,Ni)Sn4, Ni3Sn4. By examining
microstructure of samples, there are few cracks observed along the
(Au,Pd,Ni)Sn4/Ni3Sn4 interface and on the edge of solder joint. These cracks occur
77
despite very careful polishing, which suggests that (Au,Pd,Ni)Sn4 is very weak and
fragile. This brittle layer presented in a joint significantly deteriorates the reliability of
the solder joints [38-41]. The formation of a continuous layer of (Au,Pd,Ni)Sn4
unambiguously shows that the Au embrittlement issue re-emerge as a reliability issue
in micro joints.
Fig. 3.19 Micrographs showing Ni/Sn1.8Ag/ENEPIG sandwiches aged at 200°C for
(a) 0, (b) 50 h, (c) 100 h, and (d) 200 h.
78
Furthermore, the thickness of Ni3Sn4 was almost constant after over 100 h of
aging, but the (Au,Pd,Ni)Sn4 can grow continuously until all Sn had been converted
into (Au,Pd,Ni)Sn4. As aging time reached 200 h, solder joint became an all-IMCs
joint, and more than half of micro-joint (10.5 μm) was occupied by (Au,Pd,Ni)Sn4.
The reason to be responsible for the Ni3Sn4 slowdown is considered due to the
regroup of (Au,Pd,Ni)Sn4 back to the interface.
One of the most striking features in this experiment is that a (Au,Pd,Ni)Sn4 layer
with 10.5 μm in thickness was produced and originated from Au (60 nm) and Pd (100
nm) layers. To understand the mechanisms responsible for such behavior, it is
important to keep track of compositions of (Au,Pd,Ni)Sn4 IMCs at different stages of
thermal aging at 200 o C determined by EPMA, as presented in Table 3.1, so as to
figure out how it evolves. The concentration of Sn in (Au,Pd,Ni)Sn4 phases are nearly
constant, but the concentrations of Au and Pd decrease with time. The decrease in Au
and Pd is made up by the corresponding increase in Ni concentration. According to
the previous literature, the (Au,Pd,Ni)Sn4 phase is structurally the same as the initial
AuSn4 phase (orthorhombic), and Au, Pd, and Ni occupy the same sublattice while Sn
occupies its own sublattice [60]. The presence of Pd atoms apparently is able to
stabilize the AuSn4 phase even when the Au concentration in (Au,Pd,Ni)Sn4 is as low
as 2.1 at.%, as shown in Table 3.1.
79
80
3.5 Characterization of Au and Pd in Interfacial Reactions
According to this study, the effect of Au and Pd on interfacial reactions appears
to be different as their effective Pd/Au concentrations, Sn0.8Pd/Sn0.8Au, stay the
same. Although the crystal structures of (Au,Ni)Sn4 and (Pd,Ni)Sn4 are the same
(orthorhombic) [36], the morphologies of (Au,Ni)Sn4 and (Pd,Ni)Sn4 are different, as
shown in Fig 3.20. The morphology of (Au,Ni)Sn4 is needle-like, and that of
(Pd,Ni)Sn4 is facet-like. Microstructure characterization indicates that after reflow,
massive (Pd,Ni)Sn4 distributed within the SnPd solder matrix, while only a few
(Au,Ni)Sn4 appeared in the SnAu solder. As Sn fully consumed in both joints, the
isolated (Au,Ni)Sn4 particles distributed in the middle of the joints, but (Pd,Ni)Sn4
formed a very thick continuous layer across the entire solder joints. Therefore, the Pd
embrittlement phenomenon is much severe than the Au embrittlement when the
effective Pd/Au concentration is the same in both cases.
To uncover the characterization of interfacial reactions between Au and Pd, the
composition of (Au,Ni)Sn4 and (Pd,Ni)Sn4 IMCs at different stages of thermal aging
at 200 o C were determined by EPMA, as presented in Table 3.2. The decrease in Au
or Pd concentration is made up by the corresponding increase in Ni concentration.
Interestingly, the concentration of Ni in (Pd,Ni)Sn4 IMC is much higher than that in
(Au,Ni)Sn4. It is noted that the Ni concentration in (Pd,Ni)Sn4 is as high as 17.2 at%
81
as Sn completely consumes, and the Ni concentration in (Au,Ni)Sn4 is lower than 12
at%. In short, more Ni can diffuse into the (Pd,Ni)Sn4 IMC, hence the (Pd,Ni)Sn4
layer is much thicker than (Au,Ni)Sn4. As a result, the Pd embrittlement issue is much
severe than the Au embrittlement. Accordingly, it can be concluded that the Pd
embrittlement problem has a larger effect in 3D IC packaging.
Fig. 3.20 Micrographs showing (a) Ni/Sn0.8Pd/Ni and (b) Ni/Sn0.8Au/Ni sandwiches
aged at 200 o C for different aging time.
82
Table 3.2 Quantitative analysis of (a) (Pd,Ni)Sn4 and (b) (Au,Ni)Sn4 through EPMA
83
3.6 Mechanical Properties of Au or/and Pd Embrittlement Micro-joints
In this study, most of the Au and Pd embrittlement issues in 3D IC micro-joints
have been revealed; yet, the mechanical properties of micro-joints with continuous
(Au,Ni)Sn4 or (Pd,Ni)Sn4 layer is still uncertain. Therefore, the aim of this study is to
estimate the mechanical properties of severe Au or/and Pd embrittlement micro-joints
by nanoindenter.
In Au or/and Pd embrittlement micro-joints, the entire solder joint was
comprised of Ni3Sn4/(Au,Ni)Sn4/Ni3Sn4, Ni3Sn4/(Pd,Ni)Sn4/Ni3Sn4, or
Ni3Sn4/(Au,Pd,Ni)Sn4/Ni3Sn4 as Sn fully consumed, as presented in Fig. 3.21.
According to previous study, it is seen that as the continuous (Au,Ni)Sn4 or
(Pd,Ni)Sn4 layer formed in the joint, the (Au,Ni)Sn4/Ni3Sn4 and (Pd,Ni)Sn4/Ni3Sn4
interfaces appear to be the weakest part of the joint, thereby decreasing the joint
strength significantly [35, 43, 61], as shown in Fig. 3.22-24, and Fig. 1.20,
respectively.
84
Ni/Sn-Au-Pd/Ni sandwiches as Sn fully consumed.
Fig. 3.22 Fracture surfaces for the solder side and ball side of the same solder joint
that had been aged at 160oC for 500 hr. This solder joint fractured just from routine
handling. This solder joint failed along the (Au1-xNix)Sn4/Ni3Sn4 interface and the
Ni3Sn4/Ni interface, respectively. [35]
85
Fig. 3.23 A schematic drawing (white line) showing the fracture locations of the
solder joint in Fig. 3.21. [35]
Fig. 1.20 Cross-sectional view showing the fracture of a Sn–0.2Pd/Ni joint after shear
test. [43]
86
Fig. 3.24 (a) Shear strength of the Sn–xPd/Ni joints as a function of Pd
concentration(x). (b) Typical fracture of a Sn–0.2Pd/Ni joint after the shear test. [61]
Estimation of the hardness and Young’s modulus of different IMCs is established
to further assess the mechanical properties of Au or/and Pd embrittlement joints

department of materials science and engineering college of

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