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Progress in Materials Science 55 (2010) 428–475

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Progress in Materials Science

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Failure mechanisms of solder interconnects under currentstressing in advanced electronic packages

Y.C. Chan *, D. YangEPA Centre, Department of Electronic Engineering, City University of Hong Kong, Tat Chee Avenue, Kowloon Tong, Hong Kong

a r t i c l e i n f o

Article history:Received 18 November 2008Received in revised form 5 November 2009Accepted 6 January 2010

0079-6425/$ - see front matter � 2010 Elsevier Ltdoi:10.1016/j.pmatsci.2010.01.001

* Corresponding author. Tel.: +852 2788 7130; fE-mail address: [email protected] (Y.C. Ch

a b s t r a c t

The pursuit of greater performance in microelectronic devices hasled to a shrinkage of bump size and a significant increase in electri-cal current. This has resulted in a high current density and accom-panying Joule heating in solder interconnects, which places greatchallenges on the reliability of advanced electronic packages. Areview of current stressing-induced failures of solder interconnectsis thus timely. This review is devoted to five types of physical fail-ure mechanisms occurring in high current density applications,which include electromigration (EM), Joule heating-induced fail-ures, interfacial reactions, stress-related damage, and thermomi-gration (TM). In practice, some of these failure mechanisms aremixed together so that the real root cause cannot be easilydetected and understood. Reliability designers need to be wellinformed to evaluate the electrical characteristics, thermal charac-teristics and mechanical strength for solder interconnects inadvance. This review summarizes recent progress and presents acritical overview of the basis of atomic transport, diffusion kinetics,morphological evolution, and numerical simulation. Specialemphasis is on the understanding of the interactions of EM withother failure mechanisms. Aside from the review of the current sta-tus of knowledge, the remaining challenges as well as future direc-tions are also discussed.

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Y.C. Chan, D. Yang / Progress in Materials Science 55 (2010) 428–475 429

Contents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 429
1.1. Solder interconnects for advanced electronic packaging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4291.2. Challenges of high current density applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4291.3. The scope of this review . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 432
2. EM in solder interconnects under an electron wind force. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 432
2.1. Phase separation and atomic transport under current stressing . . . . . . . . . . . . . . . . . . . . . . . . 4322.2. Current stressing-enhanced phase coarsening. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4342.3. Nucleation and growth of voids at the interface. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4362.4. Lifetime statistics and the reliability of EM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 439
3. Joule heating-enhanced dissolution of UBM and the diffusion of on-chip metal interconnects . . . . . . 441
3.1. Effect of Joule heating due to current stressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4413.2. Dissolution of UBM layers and possible solutions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4423.3. Melting of solder interconnects due to Al diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 445
4. Effect of current stressing on the formation of IMCs and kinetic analysis . . . . . . . . . . . . . . . . . . . . . . . 448
4.1. Polarity effect of current stressing and enhanced growth of IMCs at the anode. . . . . . . . . . . . 4484.2. Dynamic equilibrium of IMCs at the cathode . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4504.3. Abnormal polarity effect of current stressing on the formation of IMCs . . . . . . . . . . . . . . . . . . 451
5. Stress-related degradation of solder interconnects under EM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453
5.1. Morphological evolution due to EM and a back stress in solder interconnects . . . . . . . . . . . . 4535.2. Mechanical deformation and degradation under current stressing . . . . . . . . . . . . . . . . . . . . . . 458
6. TM behavior in solder interconnects under a thermal gradient . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 462
6.1. TM in tin–lead solder interconnects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4626.2. TM in Sn-based lead-free solder interconnects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 468
7. Concluding remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 470
7.1. Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4707.2. Further problems which need to be addressed in the near future. . . . . . . . . . . . . . . . . . . . . . . 471 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 472References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 472
1. Introduction

1.1. Solder interconnects for advanced electronic packaging

With the trend towards higher integration and further miniaturization of Si-based devices, elec-tronic packaging is successively requiring a higher input/output (I/O) density, smaller feature size,and better performance. Concurrently, the flip chip solder interconnect has established its leadershiprole for high current density packages as thousands of solder bumps are fabricated onto one singlechip. To meet an even higher demand for device performance, the I/O number is expected to increase,while the dimensions of each individual bump will accordingly need to shrink.

According to the 2003 International Technology Roadmap for Semiconductors (ITRS), a significantdownsizing in flip chip packaging is anticipated [1]. Fig. 1 shows the anticipated variation in pad diam-eter, pad pitch and line width.

Also the bump size is expected to be reduced along with the pad size and pitch. At present, thediameter of a solder bump in use is about 100 lm or less [2]. In 2007, the diameter of micro-bumpshad been decreased to 20 lm [3]. This relentless scaling-down will inevitably place severe challengeson the reliability of micro-devices, as will be discussed later.

1.2. Challenges of high current density applications

At present, in the microelectronic industry, each solder joint is designed to carry 0.2 A, and this willbe doubled in the near future [2]. This means that the average current density through a 50 lm diam-eter solder joint may approach 104 A/cm2. This demands a reduced cross-section of the conductivelines and solder interconnects, whilst on the other hand, they are expected to conduct such a high

2002 2004 2006 2008 2010 2012 2014 2016 2018 20200

20

40

60

80

100

120

140

160

Size

(mic

rom

etre

)

Year

pad pitch pad diameter line width

Fig. 1. A downsizing in flip chip packaging, based on 2003 ITRS edition.

430 Y.C. Chan, D. Yang / Progress in Materials Science 55 (2010) 428–475

current density. Meanwhile, since Joule heating is proportional to the square of the current density,the local temperature of conductive lines and solder bumps will rise substantially. Also, during fieldservice, the solder joints will experience a temperature rise of at least 100 �C, to approximately 82%and 76% of the melting temperatures of eutectic SnPb and SnAgCu, respectively. As a consequence, un-der the combined effect of a high current density and a high homologous temperature, easy diffusionof atoms in the lattice is anticipated [4]. This renders electromigration (EM) a serious reliability issuein the application of high current density packages.

EM means a diffusion controlled mass transport phenomenon due to the application of electricalcurrent. In 1961, Huntington and Grone proposed that thermally-activated metal ion becomes essen-tially free in the lattice and is acted upon by two opposing forces (a direct force and an electron windforce) in a metal [5]. Also, they identified the electron wind force as the primary driving force respon-sible for the EM failure observed in interconnects. The electron wind force is further explained as oneforce experienced by a metal ion in the direction of the electron flow due to the momentum exchangebetween the moving electrons and the ion. Therefore, the phenomenological equation for the atomicflux due to EM (Jem) is described as:

Jem ¼ Cm ¼ CDFkT¼ C

DkT

Z�eE ¼ CDkT

Z�eqj; ð1Þ

where Z� is a dimensionless quantity known as the effective charge or the effective valence that re-flects the direction and the magnitude of the momentum exchange, e is the electron charge, E is theelectric field, q is the resistivity, j is the current density, C is the concentration of diffusing atoms, vis the drift velocity of these atoms, D is the thermally-activated diffusivity, and kT is the average ther-mal energy.

As regards the EM of tin and lead in solder alloys, the first published work was given by Branden-burg and Yeh in 1998 [6]. EM causes the net atom transport of solder material along the direction ofthe electron flow. Since 2002, ITRS has started to include this reliability problem for industrial atten-tion [7]. Table 1 lists the near-term reliability challenges requiring concern in current assembly andpackaging techniques [3]. According to the 2007 ITRS, EM will become a more limiting factor of highcurrent density packages, such as wafer-level packaging (WLP) for micro-electro-mechanical systems(MEMS). It is suggested that physical failure mechanisms such as EM, and thermal migration in com-bination with mechanical stresses should be understood and modeled for practical life assessment. In

Table 1Assembly and packaging – difficult challenges [3].

Difficult challenges Summary of issues

High current densitypackages

Electromigration will become a more limiting factor. It must be addressed through changes inmaterials together with thermal/mechanical reliability modelingWhisker growthThermal dissipation


particular, solder and the under bump metallurgy (UBM) need to be well designed to support a highcurrent density and minimize or avoid EM.

Also, thermal dissipation is addressed as a critical factor of reliability considering the large Jouleheating generated by the on-chip metal interconnects. One should appreciate that the cross-sectionalarea of conductive lines on the chip has been decreased significantly with the trend towards miniatur-ization, as shown in Fig. 1. It is true to say that the on-chip interconnects used in electronic packagesmay be based on Al or Cu materials. However, most concern is with the interconnects based on Al,where Joule heating is more pronounced (the resistivity of Cu is about 60% of that of Al). Fig. 2 providesa schematic diagram of a flip chip interconnect. Normally the electrical resistance of the Al traces is atleast one order of magnitude higher than that of the solder joints and Cu conductors, and thus the Altraces are the primary heat source. A significant Joule heating will accelerate the EM process in theneighboring solder joints, but also result in the degradation of UBM layers and even the Al trace itself.

With the passage of current in service, the momentum transfer from electrons to atoms may alsoplay a significant role in interfacial reactions. The formation of an intermetallic compound (IMC) layerat the interface is crucial for the good adhesion between a solder bump and a UBM, of which the driv-ing force comes from the chemical potential gradient between different contacting materials. How-ever, a rapid interaction may occur to form a substantial amount of IMC, which is known to bebrittle and has a negative effect on the mechanical integrity of interconnects. Therefore, the growthof IMC particles/layer accelerated by current stressing and the role of EM in interfacial reactions isbecoming an important and challenging problem in solder reliability.

Another concern is the formation of compression and tension regions inside a solder joint, whenatoms are driven from the cathode to the anode by the electron wind force. Stress generation andrelaxation are issues under exploration, and stress-related damage under a current density shouldbe paid considerable attention. In addition, the mechanical properties are direct indicators of strengthand long-term durability. It is understandable that EM would exert a certain effect on the mechanicaltransition of solder interconnects. As an important reliability factor, the mechanical behavior of solderinterconnects for high current density applications also needs to be carefully considered.

Moreover, owing to the significant heat accumulation, atomic migration process, thermomigration(TM), may be triggered and influence the reliability of packages. Due to differences in electrical resis-tance and thermal dissipation of individual parts within the flip chip interconnect structure (seeFig. 2), it is predicted that the heat accumulated at the chip side will be larger than that at the sub-strate side. This will inevitably lead to a considerable temperature gradient across solder joints, whichcan provide a driving force for atomic diffusion to trigger TM. More exactly, the driving force of TMcomes from the energy transported by the moving atoms and the interactions with the usual heat car-riers in the lattice [8]. The phenomenological equation for the atomic flux due to TM (Jtm) is [9]:

Fig. 2. A schematic diagram of a flip chip interconnect including Al traces, solder joints and Cu conductors.


Jtm ¼ Cm ¼ CDFkT¼ C

DQ �=N

kT2 ð�DTÞ; ð2Þ

where Q� is the heat of transport and is the heat flow per mole that must be supplied to maintain unitmolar flow in the steady state, N is Avogadro’s number, and DT is the thermal gradient. The otherterms have been defined under Eq. (1). Being a potential reliability concern for flip chip solder inter-connects, TM induced-void or pore formation has also been introduced in the 2007 ITRS [3].

1.3. The scope of this review

During field service, all the factors discussed are supposed to combine and act concurrently, whichwill further complicate the failure processes. Therefore, a quantitative understanding of the physicsand mechanics of each failure mechanism will bring about a good appreciation which will help inthe design and life prediction in flip chip solder interconnects, which should be of particular interestto both those in industry and in academe.

Five types of failure mechanism are presented in this review. Section 2 addresses the dominantphysical mechanisms during EM in solder interconnects, including phase separation, phase coarsen-ing, and void formation. Also, from an engineering point-of-view, this section gives a summary ofthe lifetime statistics and reliability evaluation of the EM of solders. Section 3 reviews the dissolutionof the UBM due to an accelerated interstitial diffusion, and the diffusion of on-chip Al trace under Jouleheating, then introduces research results on the time-dependant melting behavior of solder intercon-nects under current stressing. Section 4 presents a critical literature review related to interfacial reac-tions. The kinetics dominating in these interfacial reactions are investigated and discussed. Asummary of stress-related degradation in solder interconnects is presented in Section 5. The morpho-logical evolution due to EM, and the back stress induced are described. Also some mechanical defor-mation and degradation mechanisms under current stressing are summarized as a part of an overallunderstanding of the mechanical behavior. Section 6 discusses the reliability concerns of TM. The ther-mo-transports of Pb, Sn, Cu and Bi in solder interconnects under a thermal gradient are introduced.Lastly but importantly, some issues that need to be clarified in the near future are proposed inSection 7.

2. EM in solder interconnects under an electron wind force

2.1. Phase separation and atomic transport under current stressing

In eutectic two-phase solder joints, phase separation is likely to occur under EM due to the differ-ent atomic diffusivities over a range of operational temperatures. It was found that in a eutecticSn37Pb solder, Sn and Pb are the dominant diffusing species during EM at room temperature andabove 100 �C, respectively. This is in agreement with a study of tracer diffusion, where the resultshowed that Sn has a larger diffusivity than Pb below 100 �C whereas Pb is the faster moving speciesabove this temperature [10]. Brandenburg and Yeh observed Pb migration with the electron flowabove 100 �C accompanied by phase separation in tin–lead solder [6]. This was also supported by laterfindings in EM experiments above 100 �C [11,12]. In the case of EM at 100 �C, Agarwal et al. examinedthe EM behavior and identified that Pb is still the dominant diffusion species [13]. Compared to theresults above, EM tests were carried out at room temperature by Liu et al. [14]. It was found that nearroom temperature Sn is the dominant diffusing atom, and interfaces serve as the fastest kinetic path ofmass transport. However, a recent study through X-ray fluorescence spectroscopy has shown that Snatoms move faster than Pb in the initial stages even above 100 �C [15]. Further investigations on themovement of Sn and Pb undergoing EM over a wide temperature range need to be developed to iden-tify the in situ atomic transport. On the whole, this variation of migration behavior imposes a seriouslimitation on an understanding of the physics of failure, and on conducting and interpreting acceler-ated tests of EM in tin–lead solder.

Since eutectic SnBi is very close to eutectic SnPb in its microstructure and constitution, SnBi solderexhibits similar EM characteristic [16,17]. Fig. 3 illustrates the mass accumulation of Bi in a Sn58Bi

Fig. 3. SEM image of mass accumulation of Bi in a Sn58Bi solder joint under a current density of 5 � 103 A/cm2 at 75 �C.


solder joint under a current density of 5 � 103 A/cm2 at 75 �C. As the dominant effective charge carrier,Bi migrated along with the electron flow and substantially accumulated at the anode side. Also, as aresult of the large concentration of Bi, a phase segregation between Bi and Sn was evident, which be-haved like the Pb segregation with Sn in Sn37Pb solder.

For most lead-free solders, e.g., Sn3.5Ag, Sn0.7Cu, and Sn4Ag0.5Cu, since the amount of elementssubstituted is not as high as that of Bi in eutectic Sn58Bi, the migration of the substitutional elements(Ag, Cu) is insignificant as compared with the self diffusion of Sn and thus the phase separation may benegligible. Instead, a fast interstitial diffusion of Cu, Ni and Ag – induced dissolution of UBM and IMCplays a dominant role in the electrical failure of the lead-free solder systems above, as will be men-tioned in Section 3.2.

To understand the migration mechanism, the atomic flux due to EM was estimated from the phasedisplacement, and the product of diffusivity and effective charge number was also calculated. Accord-ing to Eq. (1), the product of the diffusivity and the effective charge number, DZ�, should be:

Table 2DZ� val

Test

35 �C55 �C75 �C

DZ� ¼ JemkT

eqjC: ð3Þ

The total volume of atomic transport during the time in operation can be obtained approximatelyfrom the product of the width and the cross-sectional area of the solder joint. According to the re-search by Lee et al. [18], the accumulated width of atoms increased linearly with time during EM.In the case considered by us, the product DZ� was calculated to be 6.5 � 10�11 cm2/s for the migrationof Pb under a current density of 2.0 � 104 A/cm2 at 100 �C [19]. This is the same order of magnitude asthe results from other groups [18]. To measure the displacement driven by EM, inert particles or nano-indentations have been used as surface markers, which moved in the opposite direction to the EM fluxas expected. Also, the product for the migration of Bi was investigated [20]. Table 2 lists the DZ� valuesobtained for a Sn58Bi solder under different test conditions. However, since the diffusivity is stronglyrelated to the microstructural changes in the EM process, there exists a large uncertainty in determin-ing the value of the effective charge number. A few studies have been conducted to unravel the diffu-sivity and explore the effective charge number [18,21].

On the other hand, for a SnAgCu solder with a higher melting temperature, under similar experi-mental conditions, the marker motion was found to be less significant, which indicates that the rateof EM is lower in this lead-free solder [21,22].

ues for a Sn58Bi solder under different test conditions [20].

conditions Accumulation rate (cm/s) Atomic flux (atoms/cm2 s) Product DZ� (cm2/s)

/5 � 103 A/cm2 3.02 � 10�10 4.86 � 1012 6.60 � 10�11

/5 � 103 A/cm2 4.70 � 10�10 7.56 � 1012 1.09 � 10�10

/5 � 103 A/cm2 1.16 � 10�9 1.87 � 1013 2.84 � 10�10


2.2. Current stressing-enhanced phase coarsening

During current stressing, another noticeable phenomenon of the microstructural evolution in sol-der joints is phase coarsening. Fig. 4 illustrates the effect of current stressing on phase coarsening inSn37Pb and Sn3.5Ag0.5Cu solder joints under a moderate current density of 6 � 102 A/cm2 at 125 �C[23]. Fig. 4a and b shows the typical microstructures of Sn37Pb and Sn3.5Ag0.5Cu solder joints as-re-flowed, in which a fine sized Pb-rich phase or Ag-rich IMC (Ag3Sn) is dispersed in a Sn-rich matrix,respectively. After being stressed for 600 h, the coarsening of the Pb and Ag-rich phase was remark-able, as shown in Fig. 4c and d. Using electrothermal simulation and an infrared thermal technique,it was found that the current stressing induced a substantial Joule heating in the solder interconnects.Hence, it was proposed that the phase coarsening was enhanced by the current stressing as a result ofenhanced diffusion related to an elevated temperature and atomic stimulation [23,24].

Also, Fig. 5a presents the percentage distribution versus the particle size after current stressing forAg-rich phases. With an increase in stressing time, the heights of the peaks decreased and the curveextended to a larger size with a wider range, which indicates a general tendency of overall coarsening.The average size as a function of time is plotted in Fig. 5b. One recalls the typical equation of phase sizecoarsening (Ostwald ripening) is [25]:

Fig. 4.phase c

dn � dn0 ¼ At ¼ A0e

�QkT t; ð4Þ

where d is the phase particle size after being annealed at a temperature of T for a time t, d0 is the initialphase particle size, A0 is a pre-exponential term, Q is the rate-controlling activation energy, and n isthe phase size exponent related to a specific atomic transportation mechanism.

(a) Typical microstructure of a Sn37Pb solder joint as-reflowed, (b) a Sn3.5Ag0.5Cu solder joint as-reflowed, (c) Pb-richoarsening in a Sn37Pb solder joint, and (d) Ag-rich IMC coarsening in a Sn3.5Ag0.5Cu solder joint [23].

0

5

10

15

20

25

30

35

40

0 0.4 0.8 1.2 1.6 2 2.4 2.8 3.2 3.6

Ag-rich Particle Size (micrometer)

Perc

enta

ge (

%)

Stressed for 600 hours(412 particles)

Stressed for 200 hours(302 particles)

As reflowed(540 particles)

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0 100 200 300 400 500 600 700

Aging or Stressing Time (hour)

Ave

rage

Pha

se S

ize

(mic

rom

eter

)

Stressed at 125 degree Celsius

Aged at 125 degree Celsius

(a)

(b)

Fig. 5. (a) The percentage distribution versus the particle size after current stressing for Ag-rich phases, and (b) the average sizeas a function of time [23].


Then the derivative form of this equation is:

lndddt¼ ln

A0

n� Q

kTþ ð1� nÞ ln d: ð5Þ

Through linear regression, the value of n for Ag-rich phase growth was estimated to be 2.96, cor-responding to the controlling kinetics of volume diffusion. Also, the n for Pb-rich phase growth in ourstudy was approximately 3.28, which corresponded to the co-kinetics of grain boundary and volumediffusion [26]. Similar to eutectic SnPb solder, current stressing-enhanced Bi coarsening was also de-tected in SnBi solder joints in our group [20].

Moreover, a phase-coarsening model including the influence of the current density was developedby Ye et al. based on an experimental study of the coarsening of the Pb phase in eutectic Sn37Pb flipchip solder joints [27]. It was found that the electric current had a greater influence on the Pb phase


growth than the temperature. Thus, a modified equation was proposed by adding a current densityterm into Burke and Turnbull’s equation:

dn � dn0 ¼ Ajmt; ð6Þ

where j is the current density, and m is the exponent for the current density. The n was found to be 5.5for the Pb-rich phase growth, and the current density exponent m was about 3. However, a theoreticalexplanation for the m is not yet clear.

It is understandable that a solder with a phase-coarsening microstructure is susceptible to damageby EM. Solder normally operates at a high homologous temperature (TH > 0.5), where both lattice andgrain boundary diffusivities play important roles. The net contribution of these two diffusion mecha-nisms in solder is dependent on the grain size. Therefore, measures need to be taken to convolute thearrangement of grain boundaries so that the process of diffusion is dominated by lattice rather thangrain boundary diffusion. For a better EM-resistant performance, the microstructure with triple pointsor intersecting grain boundaries is preferred, since these intersections interrupt the diffusion path andforce the atoms to transport laterally following the grain boundaries or migrate into the blockinggrains [28].

Chen et al. found that doping with Cu could retard the phase coarsening of a solder under currentstressing [29]. The coarsening rate was significantly reduced from 4.6 to 1.4 lm3/h when 1 wt.% Cuwas added. They proposed that the Cu added reacted with Sn, and small Cu6Sn5 IMC precipitates wereformed inside the solder. These precipitates then acted as roadblocks which retarded the movement ofgrain boundaries, leading to a reduced grain coarsening rate. Also, they investigated the effect of a0.5 wt.% Ag addition on the EM behavior. It was found that some plate-like Ag3Sn precipitates wereformed inside the solder, and behaved as obstacles that intercepted the atomic migration [30].

2.3. Nucleation and growth of voids at the interface

During EM, atomic diffusion-induced microstructural evolution includes not only phase separationand coarsening, but also void creation at interfaces. Fig. 6 illustrates the typical EM failure of a Sn37Pbsolder joint under a current density of 2.0 � 104 A/cm2 after 100 h at 100 �C [31]. A tilting effect of thecurrent [32–34] in such a line-to-bump structure was evident during the EM process. The Pb-richphase migrated towards and accumulated at the anode side corresponding to the entry direction ofthe electron flow. Moreover, it is noticeable that these voids were primarily generated near the currentcrowding regions because of the concentrated flux divergence and more serious UBM consumption.

A three-dimensional finite-element simulation was performed to demonstrate the current densitydistribution in a solder interconnect with a 1.6 A current applied. As shown in Fig. 7, the local currentdensity at the entry location reached 7 � 109 A/m2 (i.e., 7 � 105 A/cm2), at least one order of magni-tude larger than the average. Near this current crowding region, the atomic flux divergence may causevacancy accumulation and hence voids, as can be seen in Fig. 6.

Fig. 8 displays the progress of void growth in Sn3.5Ag1.0Cu solder joints under a current density of1.5 � 104 A/cm2 at 125 �C [19]. Fig. 8a shows the typical morphology of the interface before the exper-iment. After a stressing time of 75 h, as shown in Fig. 8b, voids were initiated from the upper-rightcorner, and gradually displaced the current to the surrounding areas which resulted in a lateralgrowth. Since the growth of voids induced the redistribution of the current, it is also reasonable to findthat the voids were developed towards the periphery of the UBM opening, where the current density isoriginally low. This experimental finding verifies the finite-element simulation by Liang et al. [35].

Fig. 8c shows the microstructural development after 280 h. It is evident that the voids continuouslyextended from the right-hand to the left-hand regions. Fig. 8d displays further void growth after 425 h.The propagation of voids decreased the effective contact area of the current path and induced a moreserious current crowding, and thus accelerated the void growth along the interface. This process con-tinued till the voids finally spread across the complete contact window at 515 h, as shown in Fig. 8e.

From Fig. 8, it is believed that the first void nucleation took less than 14% of the failure time. Thefailure time was then more dependent on the void growth than the void nucleation. Likewise, Chiu andChen monitored void formation and propagation in Sn37Pb solder joints under a current density of

Fig. 7. Current density distribution of a Sn37Pb solder joint stressed with a 1.6 A current (the current crowding effect isapparent at the entry direction of the current) [31].

Fig. 6. (a) SEM image of void formation at the current crowding region (the upper-right corner) of a Sn37Pb solder joint under acurrent density of 2.0 � 104 A/cm2 after 100 h at 100 �C, and (b) local magnified micrograph [31].


6.5 � 103 A/cm2 at 150 �C [36]. They found that voids started to form at approximately 10% of the fail-ure time and they grew for the remaining of 90% of the failure time. This is different from the EMbehavior of Al and Cu interconnects. In Al and Cu interconnects, the failure is basically controlledby the nucleation of voids, and the growth becomes very rapid once the voids are produced. By con-trast, in the study by Yeh et al., they proposed that it took 88% of the failure time to initiate the first

Fig. 8. SEM images of the morphological evolution in Sn3.5Ag1.0Cu solder joints under a current density of 1.5 � 104 A/cm2 at125 �C (a) before the experiment, the time point A, (b) after 75 h, 14% of the failure time, the time point B, (c) after 280 h, 53% ofthe failure time, the time point C, (d) after 425 h, 81% of the failure time, the time point D, and (e) after 515 h, 98% of the failuretime, the time point E.


few voids, whereas only 12% of the failure time was spent in void propagation until the final open fail-ure [32]. According to their study, the incubation time for void nucleation was relatively long.

Based on the results from the time points A to E in Fig. 8, the growth rate was found to be about0.32 lm/h for the whole process, which matches well with the result by Chiu and Chen [36]. Theyfound a void growth rate of 0.3 lm/h in the later stages in Sn37Pb solder joints. However, this is dif-ferent from the investigation conducted by Zhang et al. [37]. They reported a void growth rate of4.4 lm/h in Sn4.0Ag0.5Cu solder joints which experienced EM under a current density of3.7 � 103 A/cm2 at 146 �C. A theoretical value was also calculated under a continuity condition accord-ing to the kinetic model they proposed, which was in accord with the experimental result. However,similar to research in Al interconnects, thin-film test structures should be prepared to directly mea-sure the material depletion in solder over an EM period, so that the atomic drift velocity may be pre-cisely obtained. This method has been utilized to explore some EM parameters in Sn3.5Ag solder [38].

The growth of voids has been understood in relation to a variation in electrical resistance. A Kelvinstructure was designed and employed to monitor the resistance variation of a solder joint with thepropagation of voids [39,40]. A change in bump resistance as small as 0.01 mX could be detected inthis way, and it is effective to monitor how the void growth induced the resistance change in a singlesolder joint. It was found that, when the percentage depletions of the contact opening were about 50%and 80%, the maximum resistance increases could reach 70% and 250% of its initial value, respectively.

0 100 200 300 400 500 600

0.15

0.20

0.25

0.30

0.35

0.40

0.45

Vol

tage

(V

)

Time (h)

A B C DE

Fig. 9. The voltage as a function of time when an interconnect was stressed by 1.2 A at 125 �C (note this data arises from thesame experiments as Fig. 8, and the time points A–E refer to the same times in both figures).


Three-dimensional simulations for different stages of void propagation by Liang et al. also fitted wellwith these results [35].

Fig. 9 shows the typical variation in the voltage as a function of time. Solder interconnects experi-encing EM mostly exhibit such a characteristic of the variation of resistance [32,41,42]. There is a longincubation time with very little resistance change below 90% of the stressing time. The solder jointsonly contribute a minor part to the resistance of the whole interconnect structure as compared withthe Al traces and the Cu conductors, thus the effect of void propagation on the resistance of an inter-connect was less significant before the UBM was completely detached from the solder. This is why sol-der interconnects mostly retain a low electrical resistance in the early stages of the stressing time,although void accumulation has occurred in the solder joints. After that, the resistance rose abruptlyto an open circuit, since the expanding voids led to the final failure in the contact. It is also suggestedthat the cause of the rapid increase of resistance in the later stages may involve Al degradation, whichwill be emphasized in more detail in Section 3.3.

2.4. Lifetime statistics and the reliability of EM

From an engineering point-of-view, a mean-time-to-failure (MTTF) estimation of solder intercon-nects is of great interest, and a systematic reliability evaluation on EM is needed. For Al interconnects,it is reported that the EM lifetime mostly followed a log-normal distribution. Recently, Black’s modelhas been introduced to describe the EM lifetime of a solder based on the assumption that the failure iscontrolled by void damage, and a log-normal function has been applied frequently [42–44]. However,the reason for such a log-normal distribution has not been clarified. On the other hand, a Weibull anal-ysis has been performed on time-to-failure (TTF) data in other research studies [6,45]. In our case, theEM failure of Sn3.5Ag1.0Cu solder joints followed Weibull statistics very well with current densitiesranging from 1 � 104 to 2 � 104 A/cm2 at 100, 125, and 150 �C [19]. Fig. 10 illustrates the distributionfunction under various current densities at 125 �C. As expected, the reliability of EM degraded with anincrease of current density.

Nevertheless, it is noticed that the predicted lifetimes did not match well with the measured ones,particularly under a higher current density [46]. As a result, it is suggested that the model should be

Fig. 10. Weibull cumulative distribution under various current densities at 125 �C (A: 2.0 � 104 A/cm2, B: 1.5 � 104 A/cm2, C:1.0 � 104 A/cm2).


modified to include the effect of current crowding and Joule heating. A modified Black’s equation wasproposed by inserting a multiplying factor (c) of the current density and a temperature increment (DT)into Black’s equation [34,46]:

MTTF ¼ A1ðcjÞm

expQ

kðT þ DTÞ

� �; ð7Þ

where A is a constant, j is the current density in the solder, m is an exponent for current density, Q isthe activation energy for EM, k is Boltzmann’s constant, and T is the average temperature.

By adding the current factor and a constant temperature increment, Choi et al. obtained activationenergies of 0.5 and 0.8 eV for SnPb and SnAgCu solders, respectively [46]. Chae et al. also consideredthe effect of Joule heating, and the activation energy calculated was 0.86–0.94 eV and the current den-sity exponent 2.1–2.2 for SnAg solder joints [42]. Also, by virtue of a numerical simulation and tem-perature coefficient resistance method (TCR ¼ DR

R0

1DT, R0 is the resistance of the Al trace at T0, DR is the

resistance variation, and DT is the temperature rise), we deduced the c and the DT (depending on theapplied current and ambient temperature), respectively. The average activation energy obtained wasabout 0.62 eV, and the current density exponent ranged from 1.46 to 1.89 [19]. Table 3 summarizesthe EM reliability parameters for Sn-based lead-free solder interconnects from accelerated life testsin different research groups [19,34,42,46].

More recently, Chiang et al. have compared the predicted values based on Black’s equation, the pre-dicted values based on the modified equation, and the measured MTTFs under test conditions of

Table 3Statistics of EM reliability parameters of lead-free solders [19,34,42,46].

Research group Test conditions,T (�C), j (A/cm2)

Activation energy,Q (eV)

Current densityexponent, m

Univ. of Texas at Austin (Sn3.5Ag) [42] 115–150,4.12–5.67 � 104

0.86–0.94 2.1–2.2

Univ. of California, LA (SnAgCu) [46] 125–160,2.75–3 � 104

0.8 –

National Tsing Hua Univ. (Sn3Ag0.5Cu) [34] 125–165,0.74–1.68 � 104

0.88 2.11

EPA, City Univ. of HK (Sn3.5Ag1.0Cu) [19] 100–150,1.0–2.0 � 104

0.58–0.66 1.46–1.89


different current densities and temperatures. They have found that the deviation of predicted valuesfrom the experimental results have been reduced based on the modified equation [34]. Indeed, furthereffort is necessary to identify the actual current density and bump temperature in solder intercon-nects. Also, the physics of failure after accelerated life tests needs to be established to confirm its con-sistency with the proposed failure mechanism. Otherwise, the EM reliability would be incorrectlyevaluated.

3. Joule heating-enhanced dissolution of UBM and the diffusion of on-chip metal interconnects

3.1. Effect of Joule heating due to current stressing

In high current density packages, heat accumulation cannot be ignored since Joule heating is pro-portional to the square of the current density. As mentioned in Section 2.3, the current crowding effectinevitably leads to a local temperature rise which in turn accelerates the nucleation and growth ofvoids inside the solder joint. More significantly, as the foremost heat source, Joule heating from theon-chip metal interconnects is of particular concern. This has been verified with thermal infrared mea-surements [47]. Fig. 11 shows the temperature distribution in a flip chip interconnect when stressed

Fig. 11. Thermal infrared measurement for the chip side, with the Al trace exhibiting the highest temperature [47].

Fig. 12. Temperature distribution within a solder interconnect when stressed by 3.7 � 104 A/cm2 at 75 �C (hot region occurredin the Al trace) [48].


by 104 A/cm2 at an ambient temperature of 70 �C. The temperature in the middle of the Al traces wasmuch higher than that at the circular Al pads. The edges of the UBM (marked point A) and the passiv-ation openings (marked point B) also exhibited higher temperatures than the Al pads above the solderjoints.

Also, Fig. 12 shows a numerical simulation result of the temperature distribution within a solderinterconnect, including an Al trace, the half-bump solder and Cu conductor, when stressed by3.7 � 104 A/cm2 at 75 �C [48]. It is apparent that the highest temperature occurred in the Al trace. Sim-ilar results were found by other research groups [47,49].

Since the Al trace is the dominant heat source together with local Joule heating inside the solderitself, it is expected that hot spots should occur where the Al traces enter the solder joint. Near thishot spot region, atomic diffusion will be thermally accelerated so that the UBM layer will be damaged.Also, lattice diffusion of Al atoms will possibly be initiated because of the local high current densityitself. These mechanisms may be combined and will be discussed later.

3.2. Dissolution of UBM layers and possible solutions

The dissolution of a Cu UBM in a eutectic Sn37Pb solder joint under current stressing has been de-tected [50,51]. Under a current density of 103 A/cm2 at 150 �C for 0.5 h, the solder joint failed with anopen circuit, as one of the corners of the Cu UBM was dissolved and replaced by solder according to themicrostructural analysis. Hu and co-workers also reported the rapid, asymmetrical, and localized dis-solution of a Cu UBM at the cathode side [52,53]. The average dissolution rate was 1 lm/min when thecurrent density through the eutectic Sn37Pb solder joint was 2.5 � 104 A/cm2 at 100 �C. From the loca-tion and geometry of the dissolved Cu, it is believed that current crowding played a critical role in thisrapid dissolution. When the current density was increased to 4 � 104 A/cm2, extensive dissolution ofthe Cu UBM occurred even at an ambient temperature of 30 �C.

The rapid dissolution of Cu UBM is attributed to an interstitial diffusion of noble and near-noblemetals enhanced by Joule heating. It is well known that the interstitial diffusion of dilute elementsin tin is significant [54–56]. A series of fundamental studies on diffusion and EM of Cu, Ni, Ag andAu in lead–tin alloys have been developed since the 1980s [57–59]. As the lattice constants of aand b in tin are much larger than that of the c axis, the open structure along the c axis facilitates fasterinterstitial diffusion than along the other orthogonal directions. Taking Ni for example, the diffusivityof Ni along the tetragonal c axis is about 7 � 104 times than that at right angles at 120 �C [56], and theEM is relatively very fast.

Therefore, it is not difficult to understand why Ni was consumed during EM experiments too,although Ni is used as a diffusion barrier in UBM application. Fig. 13 illustrates the effect of EM on

Fig. 13. Elemental mapping at the UBM/IMC interface in a Sn3Ag1.5Cu solder joint after 1967 h under a current density of1 � 104 A/cm2 at 150 �C [60].

Fig. 14. (a) Schematic diagram of a Cu pillar bump with a solder cap, and (b) Focused Ion Beam (FIB) image of Cu pillar bumpswith a height of 80 lm [3].


a multi-layer UBM film of Ti (0.2–0.5 lm)/Ni(V) (0.325 lm)/Cu (0.5–1.0 lm) [60]. It was found that,after experiencing a downward electron flow, the Ni and Cu constituents in the UBM began to spreadinto the solder, and the UBM was gradually consumed. In this case, voids formed at the UBM/IMCinterface due to UBM consumption under the combined effect of interstitial diffusion and large Jouleheating.

Recently, it has been proposed that a possible solution to the effects of EM in solder joints would bea thick Cu pillar. The thick Cu pillar could be fabricated as the UBM, and a thin cap layer of solder


would be required for the bump, as shown in Fig. 14 [3]. An additional electroplated Ni layer has beensuggested to suppress Cu diffusion into the solder body, thus practically inhibiting IMC formation andKirkendall voiding [61]. It is expected that this will be effective for dealing with the problem of UBMdissolution and void accumulation, since the thick Cu pillar is designed to spread the current from thecontact to an approximately uniform and low density. This proposal has been supported by experi-mental studies and numerical simulation [62,63]. However, a substantial formation of IMC at theinterface is becoming an issue. Also, TM may be initiated since a large temperature gradient is gener-ated across the shallow solder interconnects.

It was noted that the rotation of b-Sn grains occurred in Sn-based solder under current stressingbecause of their anisotropic properties [64,65]. This re-orientation resulted in a realignment of Sngrains along with the current flow, thereby reducing the resistance of the solder. It is also known thatthe diffusion of Ni/Cu in UBM was much enhanced along the c axis of Sn crystals, which contributed tothe dramatic dissolution of the UBM. Therefore, one should say that the orientation of Sn grains isclosely related to the reliability of Sn-base lead-free solders. Recently, Lu et al. have investigated

Fig. 15. Elemental mapping of the interface between a solder and the UBM (a) unfailed after 1000 h current stressing, and (b)failed after 1711 h current stressing [68].


the effect of Ag in a Sn-based solder and concluded that the grain re-orientation of Sn was blocked dueto the presence of cyclic twinning and a stable Ag3Sn IMC network, in turn the dissolution of UBM wascomparatively mitigated [66]. They further explored the effect of an additional 0.6 wt.% Zn in Sn1.0Agsolder and obtained a positive result [67]. The Zn doping stabilized the Ag3Sn and Cu6Sn5 IMC net-works, and suppressed the formation of Cu3Sn IMC. More importantly, polycrystalline-like structureswere found to form at the solder/UBM interface. Although it seems that Zn doping could not controlthe grain orientation in bulk solder, the strong binding with Cu effectively slowed down the Cu diffu-sion, and thus stabilized the solder microstructure. This is a creative study to explore the doping effecton the microstructural evolution and thus the enhancement of EM resistance. A further nano-dopinginto solder is also anticipated to support a higher current density and attenuate the EM damage.

3.3. Melting of solder interconnects due to Al diffusion

For flip chip solder joints with an Al/Ni(V)/Cu UBM, if the Ni layer is consumed completely theadhesion of the UBM to the solder will be degraded. Also, Al diffusion in the Al trace will be triggeredas a result of the high current density and local heating. Liu and Lin reported an Al flux-induced failureat the cathode side of a Sn97Pb and Sn37Pb composite solder joint with a downward electron flow[68]. Fig. 15a shows that the location of the Ni layer matched well with that of the Cu layer. In thedownward electron flow case, the Ni(V) layer would be gradually consumed over a prolonged periodof time. As Fig. 15b shows, the Ni completely diffused into the solder and the V layer was also dam-aged. Furthermore, Al started to spread within the solder joint. EM and the accompanying Joule heat-ing drove the Al away from the Al trace and pushed it into the solder. The diffusion of Al into a Sn3.5Agsolder was also detected by Shao et al. [69]. They found that the solder filled in where a Ti/Cr–Cu/CuUBM had been located and some CuAl2 IMC formed in the region where the Al pad had been situated.

Using an infrared microscope Liang et al. detected the fracture of an Al trace while the current den-sity through the Al trace was about 1.2 � 106 A/cm2 [70]. They speculated that EM damage had alsooccurred in the Al trace, and that the degradation of the Al trace may be responsible for an abrupt tem-perature rise. Additionally, their thermoelectric simulations supported this. It was the degradation ofthe Al trace, instead of void formation, that contributed to the formation of a hot spot.

It has been proposed that solder melting under current stressing is a time-dependent phenomenon[28,71]. According to previous research, the principal reason for an incubation time was attributed tothe process of void generation and propagation, and solder lifetime was explained and approachedthrough modeling void accumulation [37,53]. However, Ouyang et al. have observed the melting ofeutectic Sn37Pb solder joints due to the Joule heating of the Al traces [71]. They suggested that Al dis-solution expedited the rise of electrical resistance of a solder interconnect and hence led to the finalmelting of the solder. Since the resistance change of the Al trace was dependent on the dissolution rateof Al into the solder, an incubation period was required for a temperature rise which could providesufficient heat to melt the solder joint. In this way they explained why the melting of the solder exhib-ited a time-dependent characteristic.

Recently, the melting failure of Sn3.5Ag1.0Cu solder interconnects has been studied under a cur-rent density of 2.3 � 104 A/cm2 at 125 �C [72]. A new failure mechanism involving the combined effectof solder EM and Al diffusion was proposed. Fig. 16 shows typical stages of the morphological evolu-tion. Firstly, with a downward electron flow, voids occurred at the interface between the Cu–Sn IMClayer and the solder, especially in the current crowding region (see Fig. 16a). Secondly, as Fig. 16bshows, the voids gradually extended to the surrounding areas due to vacancy super-saturation.Thirdly, the creation of pancake-type voids decreased the effective contact area, which led to moreserious current crowding. Meanwhile, the Joule heating due to current crowding was enhanced be-cause of poor heat dissipation around the voids. Under such accumulated effects, the atomic diffusionof Ni(V) in the UBM was accelerated, and the barrier that prevents the dissolution of Al into the solderno longer existed. Therefore, the diffusion of Al was triggered and some voids were found in the Al pad,as demonstrated in Fig. 16c. Also, from the local magnified micrograph shown in Fig. 16f, it is noticedthat the Ni(V) layer previously attached to the Al pad had disappeared as compared with Fig. 16e. Niatoms were dissolved and consumed to form a Cu–Ni–Sn ternary IMC, and the V layer above the voidsextruded and begun to lose its structural integrity, so that the dissolution of Al through this layer was

Fig. 16. SEM images of different stages of the morphological evolution in Sn3.5Ag1.0Cu solder joints under a current density of2.3 � 104 A/cm2 at 125 �C after (a) 92 h, 25% of the failure time, the time point A, (b) 245 h, 66%, B, (c) 295 h, 80%, C, (d) 361 h,98%, D, (e) local magnified micrograph of the interface at the time point B (the dotted region), and (f) C [72].


more rapid. Fourthly, with the progress of Al dissolution, the EM in connecting Al trace was initiatedand expedited [19,72], so that further melting failure of solder interconnects was produced, as shownin Fig. 16d.

A finite-element simulation was applied to understand the current density distribution in the flipchip interconnects. Fig. 17 displays the evolution of the current density in Al interconnects alone (seenfrom underneath). The current density reached more than 106 A/cm2, which is sufficiently high totrigger the EM of the Al. According to Fig. 17a, the current density at the exit location of the Al padranged from 1.2 � 1010 to 1.4 � 1010 A/m2 (i.e., from 1.2 � 106 to 1.4 � 106 A/cm2) before voids weredeveloped. The modeled maximum value occurred at the connecting corner of the Al pad and the Al

Fig. 17. Current density distribution in the Al interconnect alone (seen from underneath) (a) before void growth (the currentdensity at the exit location was 1.2 � 1010–1.4 � 1010 A/m2), and (b) after void growth (the current density at the exit locationwas 1.5 � 1010–1.7 � 1010 A/m2).


trace. By contrast, when the voids propagated, the location of the maximum current density wastransferred to the exit location of the Al pad, and it reached 1.7 � 1010 A/m2 (i.e., 1.7 � 106 A/cm2),as shown in Fig. 17b. This simulation indicates that the current density through the Al pad tendedto be enhanced due to the decrease of contacting area at the interface, and this result supports theexperiments well.

The total incubation time for melting the solder was found to be dependent on the rates of voidgrowth and Al diffusion in this case. Therefore, the solder melting exhibited a unique time-dependentcharacteristic. In the initial stages, the rate of void growth varied from 0.24 to 0.53 lm/h. This was re-lated to the void nucleation and propagation. In the later stage, before the final failure, the depletion ofthe Al also exhibited a linear relationship with time, which was ascribed to the EM of the Alinterconnect.

It has been known that the change in the trace resistance is a linear function of the atomic driftvelocity [73]. In this case, the relationship between the rates of resistance change of the traces (o(DR/R)/ot) and material depletion (o(DL)/ot) may be described as:

@ðDR=RÞ@t

� qrSAl

qAlSr� 1

� �1L@ðDLÞ@t

/ @ðDLÞ@t

¼ td; ð8Þ

where the subscripts r and Al refer to the under-layer and Al trace, respectively, q is the electricalresistivity, S is the cross-sectional area of the specific layer, R is the initial trace resistance, L is the ini-tial trace length, and md is the atomic drift velocity. Based on an electrical characteristic, the rate ofresistance change was estimated to be 0.9% h�1. This rate of change then represents the drift of Alatoms in the later stage.


4. Effect of current stressing on the formation of IMCs and kinetic analysis

The interactions at the anode and the cathode are different due to the polarity of the electric cur-rent. Generally, it is found that the growth of IMCs is enhanced at the anode while inhibited at thecathode, resulting in an excessive growth of the IMC layer at the anode and a dissolution of theIMC layer at the cathode under current stressing. Either of these effects may put the reliability of sol-der joints at risk and lead to the failure of the whole circuit in electronic devices.

4.1. Polarity effect of current stressing and enhanced growth of IMCs at the anode

Fig. 18 illustrates the typical interfacial evolution of Cu/Sn3.8Ag0.7Cu/Cu interconnects under cur-rent stressing [74]. Taking the IMC thickness as the total value (including both the Cu6Sn5 and Cu3Snlayers), the thickness change as a function of time under various current densities at 120 �C has beenplotted, and is shown in Fig. 19. It is apparent that the growth rates of the IMC layer were verydifferent at the anode and the cathode. The IMC layer at the anode was always thicker than that at

Fig. 18. SEM images of IMC evolution at interfaces in Sn3.8Ag0.7Cu solder joints under a current density of 3.2 � 104 A/cm2

[74].

Fig. 19. Change of the IMC thickness with time under various current densities at 120 �C [74].


the cathode after the same time of current stressing. This result is in good agreement with the researchby Chen and co-workers in the Sn/Cu, Sn/Ag, and Sn/Ni systems [75–79]. Also, it is noted that the va-cancy flux became supersaturated and condensed to form voids at the cathode, as shown in Fig. 18h.Void formation at interfaces has already been covered in Section 2.3.

Since the occurrence of voids brings complexity into the analysis of the evolution of the IMC layer,it is more reasonable to probe the kinetics that dominate the IMC growth at the anode. According toFig. 19, the growth of IMCs at the anode has a parabolic dependence on time since the square of thick-ness increased linearly with stressing time, which indicates that the formation of the IMC layer wasmainly controlled by a diffusion mechanism. It is suggested that a back stress resulting from theEM, which will be discussed in detail in Section 5.1, was responsible for this growth behavior. AsCu is the dominant diffusing species in Cu–Sn IMC formation [80], the IMC growth rate at the interface(d(Dx)/dt) based on the transport of the Cu flux was expressed and further simplified to [74]:

dðDxÞdt¼ G D

DCg

Dx� Cg

DkT

DrXDx

� �� GJ0em � G0Jem

¼ aþ a�

Dxþ b ða�; b P 0; anode; a�; b 6 0; cathodeÞ; ð9Þ

where D is the diffusivity of Cu, Cg is the Cu concentration in the IMC, r is the stress, X is the atomicvolume, J0em is the Cu EM flux in the IMC, Jem is the Cu EM flux in the solder, and G and G’ are knownconstants related to Cu concentrations in the Cu anode, the IMC and the solder.

From Eq. (9), it is evident that the (a + a�)/Dx term contributes to parabolic growth, while the bterm contributes to linear growth. Without electric current (a� = 0, b = 0), the IMC growth would fol-low a parabolic growth law. In addition, since a� is positive, the growth rate of the IMC layer should berelatively large at the anode. This is in accord with the experimental results above.

The parabolic law for the growth of the total IMC layer is also applicable to the evolution of a singleIMC layer. In Sn/Cu/Sn/Cu/Sn sandwich-type structures, uniform layers of Cu6Sn5 and Cu3Sn were ob-served at the interfaces where electrons flew from the Sn side to the Cu side at 170 �C [78]. Moreover,the growth of both the Cu6Sn5 and Cu3Sn IMC layers exhibited a parabolic trend with the passage oftime.

Xu et al. reported the effect of EM on the IMC growth in Ni/Sn3.8Ag0.7Cu/Ni interconnects at dif-ferent ambient temperatures [81]. They also found that the growth of the Ni–Cu–Sn IMC layer at theanode was faster than that at the cathode, and the growth at the interface had a parabolic dependence


on the stressing time. However, both were faster than that due to isothermal aging because of theJoule heating effect, which is distinct from that in the cases discussed above.

By contrast, Gurov and Gusak developed a model where the electric current contributes to the lin-ear term in the growth of the IMC layer, suggesting a reaction-controlled mechanism for EM [82]. Theysuggested that the phase growth changes gradually from parabolic to linear at a constant rate whenthe electric current favors the growth. An investigation by Chae et al. showed a similar result in Cu/SnAg/Cu and Ni/SnAg/Ni interconnects under a current density of 5.2 � 104 A/cm2 [42]. Also, a kineticmodel was formulated to verify the linear growth of the IMC layer when an EM driving force domi-nates the chemical interdiffusion [83,84].

It is well known that without current stressing, the growth of the IMC layer follows a parabolicbehavior [85,86]. However, exactly how the current influences the IMC growth is obviously quite com-plicated from the results cited above. More attempts need to be made to probe the underlying kineticswhich dominate the growth of the IMC layer under current stressing.

4.2. Dynamic equilibrium of IMCs at the cathode

As shown in Fig. 18b, d, f and h, the IMC growth at the cathode is retarded, in comparison with thatat the anode. This growth is accompanied by the consumption of Cu from the cathode side. The follow-ing discussion will be developed in the light of the flux of Cu.

Due to a larger electron scattering cross section, the Cu atoms experience a larger EM force than theSn atoms. In addition, the Cu solute atoms in solder have a larger diffusivity compared to that in theCu–Sn IMC and in the Cu substrate [87]. As a consequence, EM fluxes of Cu in the IMC and in the Cusubstrate are insignificant and thus negligible. IMC dissolution, and the Cu EM flux in the solder be-come the two dominant mechanisms in the process of Cu consumption. It is well known that boththe IMC dissolution flux and the Cu EM flux in the solder are dependent on the temperature. Therefore,a critical temperature (Tc) can be determined by balancing the two fluxes [88]:

CSnDo expð�Q Cu=Solder=kTÞ

kTZ�Cu=SoldereE ¼ CCuA exp � Q

kT

� �; ð10Þ

where CSn and CCu are the atomic concentrations of Sn and Cu, respectively, QCu/Solder and Q are theactivation energies for Cu diffusion in the solder and Cu consumption, respectively, Z�Cu=Solder is theeffective charge number for Cu diffusion in the solder, A is a constant, E is the electric field, T is theEM temperature.

If define:

Y ¼ CSnDoZ�Cu=SoldereE=CCukA: ð11Þ

Then, Eq. (10) can be expressed as:

YT

exp �Q Cu=Solder

kT

� �¼ exp � Q

kT

� �: ð12Þ

If QCu/Solder = Q, the critical temperature is equivalent to Y. If QCu/Solder – Q, then the critical temper-ature needs to be resolved based on Eq. (12). More importantly, this indicates that when the EM tem-perature is above the critical temperature, Cu consumption is controlled by the IMC dissolutionprocess and is typically linear with time. On the other hand, when the temperature is below the crit-ical temperature, the EM flux of Cu in the solder is dominant, and thus the Cu consumption has a par-abolic relationship with time. Under this condition, a constant Cu–Sn IMC thickness can be obtainedby balancing the Cu chemical flux and the Cu EM flux in the solder, which suggests that a dynamicequilibrium state is reached [2,89]. Liu et al. observed the interfacial evolution at the cathode in Snbumps under a current density of about 5 � 103 A/cm2 at 155 �C, and noticed that the IMC layer thick-ness approached a certain value in the later stage of current stressing [88]. A systematic experimentalinvestigation of this dynamic equilibrium was conducted by Tu [2]. He found that with a furtherreduction of the current density, the IMC layer thickness at the cathode was stabilized after a periodof 55 h stressing time in the case of 5 � 103 A/cm2, which is shown in Fig. 20.

Fig. 20. Change of Cu6Sn5 IMC thickness at the cathode in Sn3.8Ag0.7Cu solder joints under various current densities at 150 �C(a plateau was found after 55 h in the case of 5 � 103 A/cm2) [2].


4.3. Abnormal polarity effect of current stressing on the formation of IMCs

As mentioned in Section 2.1, under current stressing, Pb atoms will migrate with the electron flowand be accumulated substantially at the anode side in SnPb solder joints. It is suggested that due to theobstruction by the Pb barrier, the IMC layer growth at the interface would exhibit an abnormal char-acteristic. The IMC evolution in SnPb solder under a current density of 6.2 � 102 A/cm2 at 125 �C wasinvestigated by Wu et al. [90]. However, since the applied current was relatively moderate, the retar-dant effect on the evolution of the IMC layer due to the segregated Pb layer was insignificant.

As eutectic SnBi has been considered as a preferential solder for low temperature applications, thebehavior of SnBi solder was studied to verify the effect of current stressing on the IMC formation inthis type of two-phase solder joints. Fig. 21 shows the typical interfacial evolution in Cu/Sn58Bi/Cuinterconnects under a current density of 5 � 103 A/cm2 after different stressing times at 75 �C [91].It was found that the growth of the IMC layers at both sides (anode and cathode) was enhanced bythe electric current, and the IMC layer at the cathode grew faster than that at the anode. This abnormalphenomenon is different from what was discussed above in Sections 4.1 and 4.2. Also, systematic EMexperiments on this solder system at several ambient temperatures (35 �C, 55 �C and 75 �C) were con-ducted to establish the kinetic parameters. It is interesting to note that although the growth behaviorat the two sides showed the abnormal polarity effect, the IMC layer thickness at the cathode followeda parabolic growth law, as shown in Fig. 22.

According to an empirical equation on interfacial reactions during solid-state aging [92], the IMClayer thickness as a function of time and temperature can be described by:

x ¼ dffiffitpþ xo ¼ Do exp

�QkT

� � ffiffitpþ xo; ð13Þ

where x is the IMC layer thickness at aging time t, x0 is the initial IMC layer thickness, Do is the intrin-sic diffusivity of Cu in the IMC layer, Q is the activation energy for Cu diffusion, and T is thetemperature.

Based on Eq. (13), the intrinsic diffusivity was estimated to be 9.9 � 10�5 m2/s. Also, the activationenergy for Cu diffusion in the Cu–Sn IMC was about 0.92 eV, which agrees well with those reported forthe growth of Cu–Sn IMC [85].

Fig. 21. SEM images of the interfacial evolution in Cu/Sn58Bi/Cu interconnects under a current density of 5 � 103 A/cm2 at75 �C for (a) 192 h at the anode side, (b) 192 h at the cathode side, (c) 384 h at the anode side, (d) 384 h at the cathode side, (e)576 h at the anode side, and (f) 576 h at the cathode side [91].


A possible explanation for the abnormal polarity effect in this case is as follows. At the anode, thegrowth of the IMC layer was initially enhanced by the current stressing. However, due to the Bi EM-induced back stress inside the solder, the diffusion of Cu and Sn to the anode was inhibited. FromFig. 21, some Cu–Sn IMC precipitates were detected within the Bi-rich layer, which may be attributedto a Cu flux in the opposite direction to the electron flow. Also, with a long period of EM time, the Bi-rich layer became thick enough to block the transfer of Sn atoms. Hence, further development of theIMC layer was limited due to the lack of a Sn source at the anode. On the other hand, at the cathode,the electron flow was in the same direction as the Cu chemical diffusion, so that the atomic flux of Cuinto the IMC was accelerated. Meanwhile, the outward Cu flux was significantly inhibited due to the

Fig. 22. Variation of the Cu–Sn IMC thickness at the cathode side of Cu/Sn58Bi/Cu interconnects with the square root of time atvarious temperatures [91].


back stress and the obstruction by the Bi layer. Therefore, the IMC layer at the cathode exhibited pref-erential growth as the inward atomic fluxes were larger than those flowing out.

A back stress-induced abnormal phenomenon was also observed in eutectic SnZn solder. Because ofits low melting point, excellent mechanical properties and low cost, SnZn solder has been developedas a substitute for SnPb. Zhang et al. detected an abnormal interfacial variation in eutectic Sn9Zn/Cuinterconnects under a current density of 4.3 � 103 A/cm2 at 140 �C [93]. They found that the Cu–ZnIMC layer at the cathode was about 2.3 lm thicker than that at the anode after 166 h. Also, uponincreasing the temperature to 185 �C, the Cu–Zn IMC was replaced by Cu–Sn IMC at the anode andby a mixture of Cu–Zn and Cu–Sn IMCs at the cathode. They speculated that under the effect of theEM-induced back stress gradient, more Zn atoms migrated to the cathode. Then the Sn atoms closeto the anode reacted with Cu atoms to form a new Cu–Sn IMC layer. At longer times, the Zn migrationto the cathode was accelerated so that the initial Cu–Zn IMC was completely replaced by a Cu–Sn IMClayer. The investigation by Kuo and Lin supported this finding [94]. They noticed a similar abnormalpolarity effect of Cu–Zn IMC, i.e., the growth rate at the cathode was a little faster than that at the an-ode. However, the rationale behind the experimental findings should be further verified.

5. Stress-related degradation of solder interconnects under EM

5.1. Morphological evolution due to EM and a back stress in solder interconnects

When atoms are driven from the cathode to the anode by the electron wind force, the latter will bein compression and the former in tension. In a cross-sectioned solder joint for an in situ observation, itis expected that the compressive stress will be released from the free surface causing hillocks or whis-kers to occur at the anode. Using thin film solder strips, Liu et al. first investigated the formation ofatomic hillocks in pure tin under a current density of 105 A/cm2 at room temperature [14]. As canbe seen in Fig. 23, with a prolonged current stressing, Sn grains were extruded out as hillocks. Anexplanation was given in terms of a stress relief mechanism that a hillock or whisker grows fromthe surface under compression [95].

Fig. 24 illustrates EM in the form of hillocks at the anode in a eutectic Sn37Pb solder joint under acurrent density of 4.4 � 104 A/cm2 after 1050 h. Here again, the surface morphology became rough be-cause of a non-uniform stress distribution. Interestingly, it can be seen that a dimple occurred at thecathode side, while a bulge was formed at the anode side. This is because the accumulated stress had

Fig. 24. Hillocks at the anode in a eutectic Sn37Pb solder joint under a current density of 4.4 � 104 A/cm2 after 1050 h.

Fig. 23. Hillocks at the anode side of a pure Sn strip under a current density of 105 A/cm2 after 80 h at room temperature [14].


been relaxed in the matrix by the dimple and bulge of the solder surface. In addition, the hillocks anddimples were observed in Cu/Sn9Zn/Cu solder interconnects under current stressing [96]. The forma-tion of hillocks in the middle of the bulk solder was ascribed to the compressive stress driven by EM.The average growth rate of the hillocks was measured to be approximately 1.3 � 10�8 cm3/h.

Synchrotron X-ray microscopy has been applied to provide information regarding a depth profilefor the accumulated stresses [97,98]. Fig. 25 demonstrates the hillock and valley formation in a eutec-tic tin–lead solder joint under a current density of 104 A/cm2 after 72 h. The depth profile obtainedwith confocal laser microscopy for the joint, which is shown next to the micrograph, indicates thatthe maximum height of the hillocks near the anode was about 16 lm, and the depth of the valley nearthe cathode was about 34 lm. It can also be seen that the surface features of the hillock region exhib-ited rows of striations with a spacing of several micrometres. These striations were on the side of thehillock facing the anode, and initiated from the anode end then propagated in the direction opposite tothe electron flow during current stressing. Such markings can be considered as an indication of thematerial that was being squeezed out as a result of the compressive stress.

Also, whiskers in thin film solder strips have been detected under prolonged current stressing[14,99]. Lin et al. investigated Sn whisker growth under current densities ranging from 4.5 � 104 A/cm2 to 3.6 � 105 A/cm2 [99]. It was found that a higher current density accelerated the growth ofSn whiskers. In addition, the current crowding effect played an important role in the growth of whis-kers as expected.

Recently, Ouyang et al. have reported the formation of whiskers in solder joints [100]. Fig. 26a andb shows the growth of whiskers at the upper-right corner (anode side) in eutectic SnPb and SnAgCu

Fig. 26. Whisker growth at the anode (chip side) (a) a eutectic SnPb solder joint under a current density of 104 A/cm2 after 48 h,and (b) SnAgCu solder joint under a current density of 1.4 � 104 A/cm2 after 248.5 h [100].

Fig. 25. Hillock and valley formation in a eutectic tin–lead solder joint under a current density of 104 A/cm2 after 72 h (thedepth profile is shown to the left) [98].


solder joints under a current density of above 104 A/cm2 after 48 h and 248.5 h, respectively. The com-position of the whiskers was confirmed as 93 wt.% Sn by EDX. These whiskers were initiated from thecracked surface at the chip side. When Pb atoms were pushed towards the anode, a compressive stresson Sn-rich grains was produced and then Sn whiskers were squeezed out. Moreover, it was noticedthat the cross-sectioned surface of the SnPb solder exhibited a dimple and bulge structure after EM,while the surface of SnAgCu solder remained flat. This phenomenon suggests that the rate of EM inSnAgCu was smaller than that in SnPb.


On the basis of the Nabarro-Herring model of the equilibrium vacancy concentration, more vacan-cies are generated in the tensile region, while less vacancies occur in the compressive region, so thereis a vacancy concentration gradient decreasing from the cathode to the anode [101]. The atomic fluxunder a combined electrical and mechanical force can be expressed as:

Fig. 27.(no ma

J ¼ CDkT

Z�eqj� CDkT

Xdrdx

; ð14Þ

where r is the hydrostatic stress, dr/dx is the stress gradient, and X is the atomic volume. The otherterms have been defined before. The first part represents the flux due to EM, whereas the second partstands for the opposite flux due to back stress.

A plausible explanation for the above difference between the rates of EM of SnAgCu and SnPb isthen given in terms of the back stress [100]. Since the elastic modulus or stiffness of the SnAgCu solderis larger than that of the SnPb, the back stress gradient in SnAgCu could be higher. Hence, the effect ofthe back stress on the retardation of EM was relatively larger for the SnAgCu solder.

If the stress balances with the wind force at a critical length, there should be no net atomic flux(J = 0), which has been well known as the Blech condition [102]. According to Eq. (14), the criticallength (Xc) can be obtained:

Xc ¼rcX

Z�eqj: ð15Þ

The effect of back stress and the critical length in a solder was investigated by Wei and Chen [103].Eutectic SnPb solder strips with lengths ranging from 5 lm to 200 lm were prepared for the study anda length-dependent EM behavior was identified. Fig. 27 shows the microstructural characteristics ofvarious solder strips under a current density of 2 � 104 A/cm2 after 490 h at 100 �C. No material deple-tion or voids could be detected for the 5 lm and 10 lm long strips. By taking the critical compressiveyield stress (rc) of SnPb solder (27 MPa), the critical length was estimated to be 11 lm under suchexperimental conditions. This value agrees well with the experimental results.

The effect of the back stress was further studied by an area array of nanoindentation markers onthe cross-section of solder joints by Xu et al. [22]. Most markers moved against the EM-induced atomicflux, indicating that the effect of the electron wind force was larger than that of the back stress in theircase. After 360 h of current stressing, the average marker movement from the cathode to the anodewas plotted and is shown in Fig. 28a.

Also, the atomic flux can be calculated as follows:

J ¼ VXðStÞ ¼

uXt

; ð16Þ

where V is the total volume of atomic transport, u is the marker displacement, X is the atomic volume,S is the cross-sectional area, and t is the operation time.

Microstructural characteristics of various solder strips under a current density of 2 � 104 A/cm2 after 490 h at 100 �Cterial depletion could be detected for the 5 lm and 10 lm long strips) [103].

Fig. 28. (a) Marker movement at different locations in a solder joint under current stressing after 360 h, and (b) the stressgradient as a function of the location [22].


By combining Eqs. (14) and (16), and assuming that the effect of the back stress gradient onthe marker movement could be neglected when the marker was far enough from the anode, the stressgradient as a function of marker displacement can be described as (when defining K = (1/C)(kt/D)(1/X2t)):

drdx¼ Kðuo � uÞ; ð17Þ

where uo is the marker displacement near the cathode, and K is a constant for a given temperature andtime.

Hence, the stress at any location is:

r ¼Z x

0Kðuo � uÞdx: ð18Þ

Substituting u(x) and the boundary conditions into Eq. (18), the stress gradient at any locationcould be determined, which is shown in Fig. 28b. The stress gradient near the anode was 97 kPa/lm, and it decreased gradually to zero with distance.

In addition, the damage under the combined effect of the electron wind force and the thermo-mechanical stress has been investigated using a coupled-field simulation [104]. It has been found thata substantial thermal stress accumulated around the interface at the chip side, especially in the NiUBM and (Ni,Cu)3Sn4 IMC layers. As shown in Fig. 29a, the maximum stress was 138 MPa. Also, thestress variation in direction y with the distance from point A is plotted in Fig. 29b, where a stressgradient of 1.67 � 1013 Pa/m was found. Therefore, the thermal mismatch-induced force for Ni atomicdiffusion was estimated to be 1.82 � 10�16 N, which was comparable to the electron wind force(2.82 � 10�16 N) in this case. At the initial stage, if the stress is not released effectively, stress migra-tion would interact with EM and influence the reliability of solder joints, which needs more experi-mental studies.

Fig. 30. U field fringe of a Sn4Ag0.5Cu solder joint (a) during current stressing, and (b) after the current was terminated [105].

Fig. 29. Thermal stress distribution in Ni UBM and (Ni,Cu)3Sn4 layers (at the chip side) (a) stress distribution, and (b) variationof the stress with the distance from point A in the direction y [104].


5.2. Mechanical deformation and degradation under current stressing

To detect the EM-induced mechanical damage, a Moiré interferometric technique was used to ob-tain the in situ displacement evolution of solder joints under electric current stressing [105]. Largedeformations may be observed in solder joints under a current density of 104 A/cm2. Figs. 30a and31a display the U field and V field fringes in a Sn4Ag0.5Cu solder joint after 1500 h of current stressing,

Fig. 31. V field fringe of a Sn4Ag0.5Cu solder joint (a) during current stressing, and (b) after the current was terminated [105].


respectively. The U field fringes were predominantly in the vertical direction with concentrations onboth vertical edges, indicating that a large normal deformation was developed in the horizontaldirection. Instead, the V field fringes were predominantly in the horizontal direction suggesting a large

0 200 400 600 800 1000 1200

42

44

46

48

50

52

54

56

58

60

62

Mod

ulus

(G

Pa)

Time (h)

anodic joint cathodic joint

Fig. 32. Variation of modulus of Sn3.5Ag1.0Cu solder joints under a current density of 2.0 � 104 A/cm2 after different times at125 �C.

Fig. 33. SEM images of solder joints after mechanical shear testing under a current density of 2.55 � 104 A/cm2 after 10 h at140 �C (a) chip side, and (b) substrate side [106].


normal deformation in the vertical direction. Also, Figs. 30b and 31b show the field fringes after thecurrent was switched off. Although the fringes became less clear, there were little changes in bothU and V field fringes. This means that the deformations created by the high current density were irre-versible, which is attributed to the re-arrangement of defects and atoms in the material and accom-panying local volumetric change.

Nano-indentation tests were conducted to explore the mechanical behavior of solder joints afterEM [19]. Fig. 32 illustrates the variation of modulus of Sn3.5Ag1.0Cu solder joints under a current den-sity of 2.0 � 104 A/cm2 after different times at 125 �C. It is apparent that the modulus of solder jointstended to decrease after current stressing (except for that of the cathodic joint after 305 h), and themechanical properties were degraded as compared with the original values. When interfacial voidsare initiated due to flux divergence under EM, bond damage occurs in solder joints. From a physical

Fig. 34. Typical fractographs of Sn3.5Ag1.0Cu solder joints (substrate side) (a) as reflowed, (b) under current stressing after144 h, and (c) under current stressing after 288 h [48].


perspective, the modulus is directly related to the atomic bonds. Hence, it is understandable that themodulus decreased with the passage of time under current stressing.

The effect of EM on the shear behavior of flip chip solder joints was studied by Nah et al. [106]. Itwas found that the mode of shear failure changed after EM and depended on the direction of electronflow. Originally, shear-induced fracture occurred in the bulk of the solder without current stressing.However, under a current density of 2.55 � 104 A/cm2 after 10 h at 140 �C, as shown in Fig. 33, fractureoccurred instead at the cathode interfaces between the solder and IMCs. This is because EM dissolvedand drove Cu or Ni atoms from the UBM or bond pad into the solder, and this resulted in the largegrowth of brittle Sn-based IMCs at the cathode side. Therefore, shear failure occurred predominantlyat the cathode interface.

Also, fractographs of solder joints before and after EM were examined for a comparison [24]. Fig. 34shows the typical sheared fracture surfaces of Sn3.5Ag1.0Cu solder joints (substrate side) under acurrent density of 2.1 � 104 A/cm2 at room temperature. For the case without current stressing, thefracture mode was in the bulk solder cutting through the region just near the (Cu,Ni)6Sn5 IMC layer,and the fracture surface exhibited large amounts of ductile deformation with big dimples. At longerstressing times, the interface became brittle and less plastic deformation was observed. Thismechanical deterioration with stressing time is attributed to void formation and stress accumulationat the interface.

In addition, Ren et al. examined the tensile behavior of Cu/Sn3.8Ag0.7Cu/Cu solder interconnectsunder current densities ranging from 1 � 103 A/cm2 to 5 � 103 A/cm2 at 100 �C [107]. They observedthe ductile-to-brittle transition in which the fracture migrated from the middle to the cathode inter-face of solder joints with increasing current density. Fig. 35 illustrates this movement of rupture posi-tion under a creep stress of 7 MPa. Likewise, the transition is explained by the polarity effect of EM,

Fig. 35. SEM images of Cu/Sn3.8Ag0.7Cu/Cu solder interconnects after tensile testing (a) without EM, (b) under a currentdensity of 3.3 � 103 A/cm2, and (c) under a current density of 5 � 103 A/cm2 [107].


especially the accumulation of vacancies at the cathode interface. As a whole, these EM-inducedmechanical property changes of flip chip solder joints would strongly impact on the reliability of flipchip technology.

6. TM behavior in solder interconnects under a thermal gradient

With the trend towards greater integration and further miniaturization in the microelectronicsindustry, the cross-sectional area of conductive lines on chips has been decreased significantly. Thishas led to a dramatic accumulation of Joule heating in first-level interconnects, as discussed in Section3.1. Then a considerable thermal gradient will possibly be built up across solder joints, which can pro-vide a driving force for atomic diffusion to trigger TM.

6.1. TM in tin–lead solder interconnects

The earliest report regarding the combined effects of EM and TM in solder joints was given by Yeet al. [108]. With microstructural observations and marker measurements, they found that TM in flipchip solder joints may assist or counter EM, depending on the direction of the thermal gradient andelectric field.

The individual contribution of TM to the failure of solder joints was described by Huang et al. [109].They proposed the design of a test structure of flip chip solder joints which can be applied to conductTM without EM. Generally, in the interconnect structure, the Al traces are the primary heat sourcesbecause of their larger resistance. Hence, it is believed that a certain thermal gradient exists in thepowered solder joints due to the temperature difference. Moreover, owing to the excellent thermalconductivity of the silicon chip, a similar thermal gradient is also formed across the adjacent un-pow-ered joints. These un-powered solder joints are thus investigated for a TM study since no current isapplied to them.

Fig. 36 shows the typical TM phenomenon of tin–lead composite solder joints (Sn97Pb and Sn37Pb)under a temperature gradient of above 1000 �C/cm after 5 h at 150 �C. The arrows indicate the direc-tion of the electron flow. According to the cross-sectional observations after current stressing, the ef-fect of TM was clearly visible across the un-powered solder joints, since in both of them Sn (darkregions) moved to the chip side, the higher temperature end, and Pb (light regions) to the substrateside, the lower temperature end. The advantage of using a composite solder is that there is an inho-mogeneous compositional distribution in such samples, and the phase redistribution by TM can beeasily recognized. Also, this set of samples was used to conduct experiments in situ by detectingchanges on the cross-sectioned surfaces directly during current stressing. Likewise, the phase redistri-bution of Sn and Pb was recognized in the neighboring un-powered solder joints [109].

Morphological evolution due to TM has also been detected in eutectic tin–lead solder joints.Fig. 37a shows SEM images of a row of tin–lead solder joints after 50 h at 150 �C. Fig. 37b demon-strates the detailed microstructure of joint 4 at a higher magnification [110]. According to thesemicrographs, it is believed that Pb migrated to the substrate side under the temperature gradientacross the un-powered solder joints. This was supported by the EDX analysis of local regions. As

Fig. 36. SEM images of TM in un-powered tin–lead composite solder joints (Sn97Pb and Sn37Pb) under a temperature gradientafter 5 h at 150 �C [109].

Fig. 37. (a) SEM images of a row of solder joints from 1 to 12, with solder joints from 5 to 8 under current stressing after 50 h at150 �C (Pb accumulation in the un-powered solder joints 4 and 9), and (b) the detailed microstructure of joint 4 at a highermagnification [110].


shown in Fig. 37b, the average concentration of accumulated Pb at the substrate side was about65.2 at.%, and the concentration of Sn at the chip side approached 86.3 at.%. The width of the accumu-lated Pb-rich phase band reached approximately 15 lm, i.e., half of the joint standing height. This re-sult agrees well with TM in tin–lead composite flip chip solder joints [109].

A possible explanation for the above is as follows. The flow of atoms under a thermal gradient de-pends on the heat of transport (Q�), defined as the difference between the heat carried by the movingatoms and the heat of the atoms in the initial state (the hot end or the cold end) [111]. For the atomswhich move from the hot end to the cold end, the Q� is negative since they lose heat. For atoms movingfrom the cold to the hot end, then the Q� is positive. Pb atoms are the dominant diffusing species witha higher diffusivity in eutectic tin–lead solder above 120 �C [10,112]. Therefore, based on the micro-structural evolution as shown in Figs. 36 and 37, it is speculated that with a negative Q�, Pb atomsmigrated from the higher temperature side to the lower temperature side under the temperaturegradient. Meanwhile, Sn atoms moved slowly and replenished the vacancies due to the depletion ofPb atoms. Macroscopically, the Pb-rich phase migrated to one side and the Sn-rich phase was‘‘pushed” towards the opposite side on the basis of a constant volume process. However, the


mechanism of the reversed Sn flux during TM remains unclear, and the sign of Q� for Sn cannot be con-firmed as yet.

It is important to obtain the temperature distribution of first-level solder interconnects to under-stand the TM behavior of solder joints. Finite element modeling and simulation were applied to pre-dict the electrothermal characteristics in a TM study. Fig. 38a shows the temperature distribution of aflip chip test structure with a 1.8 A current applied at 150 �C [110]. A temperature difference betweenthe chip side and the substrate side was visible. Significantly, Fig. 38b describes the detailed temper-ature distribution of an un-powered joint in the interconnect, where there existed a temperature dif-ference above 3 �C between the chip side and the substrate side. This means that a temperaturegradient greater than 1000 �C/cm (3 �C/30 lm � 1000 �C/cm) was built up across the un-poweredjoint due to the Joule heating from the neighboring Al traces. A simulation by Ye et al. also showedthe existence of a thermal gradient of about 1500 �C/cm in their flip chip test structures [108]. In com-bination with the simulation, thermocouples, and the temperature coefficient of resistance methodwere used to verify real temperatures in solder interconnects [31]. However, due to the unique inter-connect structure and the limitation of these measurement methods, it is difficult to pinpoint the

Fig. 38. (a) Temperature distribution of a quarter of the flip chip test structure, and (b) an un-powered solder joint (joint 9)[110].


temperature gradient across solder joints. Recently, a thermal infrared technique has been employedby Hsiao and Chen to obtain the thermal gradient directly in cross-sectioned solder joints [113]. As canbe seen from Fig. 39a, a uniform temperature distribution occurred in the solder bump before currentstressing. Then Fig. 39b shows the temperature distribution of a solder joint under an alternating cur-rent density of 9.2 � 104 A/cm2. Since there is no EM effect under alternating current stressing, and thealternating current produces a similar Joule heating as does the direct current, the alternating currentwas applied to independently investigate the TM behavior in solder joints. Fig. 39c shows the temper-ature profile along the dashed line in Fig. 39b, in which the average temperature at the chip side was16.0 �C higher than that at the substrate side. The thermal gradient was calculated to be approxi-mately 2143 �C/cm. This trial is significant since it verified the existence of a large thermal gradientacross real flip chip solder joints with experimental data.

Fig. 39. (a) Temperature distribution of a solder joint before current stressing, (b) under an alternating current density of9.2 � 104 A/cm2, and (c) the temperature profile along the dashed line in (b) (the thermal gradient was estimated to be 2143 �C/cm) [113].


Also, from Fig. 37b it is noticeable that the Pb-rich phase accumulated at the lower left side, i.e., thelower temperature region of solder joint 4. Likewise, the Pb redistribution in solder joint 9 showed asimilar tendency, i.e., Pb migrated to the lower right side (the lower temperature region), as Fig. 37ashows. Taking solder joint 4 for example, since its right side was closer to the heat source, it is possiblethat a temperature gradient was established laterally from the right side to the left side. Thus the Pb-rich phase not only migrated to the substrate side under the vertical temperature gradient, but alsomoved to the lower temperature region driven by the lateral temperature gradient across this solderjoint. This lateral TM was also observed in composite solder joints [109,114]. As Fig. 36 shows, in theun-powered joints near to the powered joints, e.g., joint 9, it is evident that the Sn redistribution wastilted towards the powered joints.

Morphological variations at different cross-sectional planes of a solder joint during TM are involvedbecause of different thermal dissipations [31]. Fig. 40a shows the obvious TM of Pb at the periphery oftwo solder joints, which is similar to that in Fig. 37. Then a stepwise cross-sectional analysis was con-ducted by gradually grinding the solder joints to the center of the passivation opening. Fig. 40b dem-onstrates the cross sections of solder joints 4 and 9 after re-polishing by about 50 lm. It is noted thatthe TM of Pb was not as apparent as that in Fig. 40a. Pb-rich phases were uniformly distributed in sol-der joint 4. Pb accumulation in solder joint 9 was also slight and only the Pb-rich phase at the periph-ery of the solder exhibited a TM characteristic. This means that TM of the inner solder region was lesssignificant than that of the outer solder (the surface layer). One can understand that the temperaturedistribution inside the center of a solder joint was more uniform. By contrast, it is easier to build up alarge thermal gradient across the surface layer of solder joints where a substantial heat dissipation isachieved, since the outer solder is close to the ambient environment.

Fig. 40. SEM images of two cross-sectional planes for un-powered solder joints 4 and 9 (a) after the first polishing (outersolder), and (b) after re-polishing 50 lm or so (inner solder) [31].

Fig. 41. (a) SEM image of TM in a eutectic tin–lead solder joint under a thermal gradient after 27 h at 100 �C, and (b) localmagnified micrograph (the finer lamellar microstructure occurred after TM) [116].


It is worth mentioning that during the TM, the Pb grains were even more uniformly dispersed in thetin-matrix, although the bulk of the Pb had moved to the substrate side, as shown in Fig. 37b. Thismeans that the lamellar microstructure became much finer after the TM process. This has also beendetected in Sn58Bi solder joints under a TM-enhanced effect in our group [115]. Ouyang et al. founda similar phenomenon, which is shown in Fig. 41 [116]. They suggested that the formation of this finerlamellar structure created a more disordered higher entropy state [2,116]. Also, according to their esti-mate, entropy production by heat propagation was many orders of magnitude larger than that byatomic migration, and it is thus conceivable that entropy production in TM could affect the micro-structure substantially.

In order to understand the mechanism of atomic transport, the atomic flux and the heat of trans-port during the TM process were estimated. Taking a central displacement (Dx) of 7.5 lm in Fig. 37b,the total volume of atomic transport (Vtm) during the operation time (t) can be approximately obtainedfrom the product of the displacement and the cross-sectional area (S) of the solder joint. Therefore,taking the atomic volume of Pb (X) as 3.0 � 10�23 cm3, the atomic flux of Pb due to TM (Jtm) can becalculated as follows:

Jtm ¼Vtm

XðStÞ ¼ADxXðStÞ ¼

7:5� 10�4

3:0� 10�23 � 50� 3600� 1:4� 1014 ðatoms=cm2 sÞ: ð19Þ

Also, the atomic flux due to TM can be expressed as Eq. (2), which has been given in Section 1.2.Hence, taking a predicted temperature gradient of 1100 �C/cm, an atomic diffusivity of Pb of4.0 � 10�13 cm2/s [10], we substituted these values into Eq. (2), and obtained the molar heat of trans-port as about 27.2 kJ/mole. Compared to the result reported by Ouyang et al. (25.3 kJ/mole) [116], themolar heat of transport of Pb is slightly different.

In addition, Chuang and Liu estimated the molar heat of transport of Pb as 22.2 kJ/mole under athermal gradient of 1010 �C/cm, by measuring the displacement of artificial markers [117]. Signifi-cantly, they found that the average displacement of atoms increased almost linearly with time duringTM. More recently, markers fabricated by a FIB method have been employed to measure the rate of TMby Hsiao and Chen [113]. With a thermal gradient of 2143 �C/cm, a molar heat of transport of 26.8 kJ/mole has been obtained for the transport of Pb.

An enhancement effect of TM on the phase coarsening of eutectic tin–lead solder has also beeninvestigated [118]. A possible explanation is that the inward atomic flux of Pb across the thermal gra-dient was larger than the outward atomic flux for any particular Pb-rich phase particles, and this accu-mulation of the atomic flux in the Pb phase was the major contributor to the enhancement ofcoarsening.

Fig. 42. (a) Original micrograph of a Sn3.5Ag solder joint and (b) micrograph of a solder joint after 800 h of TM testing with atemperature gradient of 2829 �C/cm at 100 �C (note that the Sn whiskers were present at the chip side, and the markers movedtoward the substrate side) [119].


6.2. TM in Sn-based lead-free solder interconnects

As stated above, the migration of the Sn flux during TM has been unclear and has necessitated fur-ther investigations for lead-free solders.

A more recent study by Hsiao and Chen reveals the TM characteristic of Sn in lead-free solder [119].They investigated the TM behavior of Sn in Sn3.5Ag solder joints under a temperature gradient of2829 �C/cm at 100 �C. As mentioned in Section 6.1, an alternating current was used to eliminate theEM effect, thus facilitating an independent study of TM. After 800 h of TM testing with a 0.57 A alter-nating current, it is interesting to find that hillocks were pushed out from the chip side, as shown inFig. 42b. These hillocks were generated by the mass transfer of the Sn at the hot side, providing directevidence that Sn was transported along the direction opposite to the thermal gradient. In addition, bymeasuring the marker movement, they obtained the TM flux and molar heat of transport of Sn as5.0 � 1012 atoms/cm2 s and 1.36 kJ/mole, respectively. Our studies also show the similar tendency thatSn atoms migrate towards the higher temperature side in Sn3.0Ag0.5Cu solder joints under a thermalgradient [19,120].

A TM of Cu atoms has been proposed. The microstructural evolution in Cu/Sn4Ag0.5Cu/Cu solderinterconnects has been studied under a thermal gradient of 1000–1200 �C/cm [121,122]. It has beenfound that the two major microstructural differences between TM and isothermal samples were thelack of a Cu3Sn layer at both the higher and lower temperature sides, and the thinning of the Cu6Sn5

layer at the higher temperature side for the TM samples. Supposedly, this thinning of the Cu6Sn5 layerwas due to its disintegration, during which the Cu atoms moved to the lower temperature side underthe thermal gradient. Meanwhile, the absence of the Cu3Sn layer was a result of an insufficient Cu con-centration. More recently, the TM of interstitial Cu in SnAg flip chip solder joints has been reported byChen et al. [123]. It has been suggested that the void formation at the chip side was attributed to a fastinterstitial diffusion of Cu atoms from the UBM into the Sn matrix. The driving force of Cu diffusion

Fig. 43. (a) Schematic diagram of a line-type test structure (Ni/Sn58Bi/Cu), (b) temperature distribution in the test structure,and (c) temperature distribution in the solder joint only (a thermal gradient of about 527 �C/cm existed across the solder joint)[115].


was due to a large thermal gradient accredited to Joule heating across the solder bumps. Further at-tempts need to be made to verify the real characteristic of the TM of Cu.

More than that, a specific line-type test structure (Ni/Sn58Bi/Cu) has been applied to investigatethe combined effect of EM and TM of Bi, as shown in Fig. 43a [115]. As Ni shows a higher electricalresistivity than Cu, a large temperature difference may be created at two sides of the solder joint dur-ing the current stressing (downward from the Ni to the Cu sides). By finite-element simulation, it wasshown that a thermal gradient of about 527 �C/cm existed in the solder joint under a current density of5 � 103 A/cm2 at 50 �C, as demonstrated by Fig. 43b and c. Temperature measurements using thermo-couples also supported this. By varying the direction of the electrical current, the counteracted andenhanced effects by TM were detected separately. It can be seen from Fig. 44 that the migration ofBi atoms was more pronounced when the Ni wire was used as the cathode. According to the experi-mental findings, we speculate that the Bi has a similar TM characteristic to Pb and shows a negative Q�.Then, if the direction of the thermal gradient was opposite to that of the electron flow, the TM coun-teracted the EM and retarded the diffusion of the Bi atoms (case 1). Otherwise, the TM assisted the EM,and the diffusion of the Bi atoms was enhanced (case 2). In addition, based on the results from thesetwo cases, the atomic fluxes of Bi due to EM and TM were differentiated and estimated to be1.48 � 1013 atoms/cm2 s and 5.38 � 1012 atoms/cm2 s, respectively.

Fig. 44. SEM images of Ni/Sn58Bi/Cu solder joints under a current density of 5 � 103 A/cm2 after 384 h at 50 �C (a) case 1: EMcounteracted by TM, (b) anode side in case 1, (c) cathode side in case 1, (d) case 2: EM enhanced by TM, (e) anode side in case 2,and (f) cathode side in case 2 [115].


7. Concluding remarks

7.1. Summary

In this review, we have discussed five types of physical failure mechanisms of solder interconnectsfor high current density applications. Phase separation and void growth, the essential physical pro-cesses occurring in EM, were introduced in Section 2. The pronounced effect of current crowding onEM damage is worthy of attention. Also, EM reliability parameters for Sn-based lead-free solder inter-connects in recent publications were collected for summary. One should say that the modification ofBlack’s model is significant, which exerts a large effect on the lifetime statistics and reliability evalu-ation for EM failure.


Joule heating in first-level solder interconnects is substantial. This was demonstrated in Section 3through experimental and numerical investigations. Due to interstitial diffusion enhanced by Jouleheating, the consumption of the UBM layer is noticeable so that new structural design (e.g., Cu/Ni pil-lar) and microstructural improvement need to be developed to support high current densities andminimize EM. In addition, the Al diffusion-induced damage in flip chip interconnects has attractedsome interest. We proposed a failure mechanism involving the combined effect of solder EM and Aldiffusion, and thus offered an explanation for the time-dependent behavior of melting failure of thesolder under current stressing.

Vacancies are provided by the flux opposite to that of the lattice diffusion of atoms, but also comefrom interfacial diffusion. As a result, another physical mechanism, interfacial reactions, was outlinedin Section 4. Due to the polarity effect of the electrical current, the formation of IMC layers is enhancedat the anode while retarded at the cathode. The IMC layer at the anode grows preferentially with pro-longed stressing time. On the other hand, IMC formation at the cathode can reach a dynamic equilib-rium when the electron wind force is comparable to the chemical force, and the Cu consumption ispredominated by the EM process instead of dissolution. In addition, some abnormal polarity phenom-ena which were present in SnBi and SnZn solders have been summarized and included for comparison.

We reviewed the stress-related degradation of solder interconnects under a current density in Sec-tion 5. Due to the relief of compressive stresses, morphological evolution is apparent in the form ofhillocks or whiskers near the anode. One important factor, the back stress generated, was investigatedto understand the damage due to EM. Also, the mechanical deformation was identified through aninterferometric technique. In addition, the deterioration of solder interconnects under current stress-ing was detected through a series of mechanical tests.

An attempt to explore the TM behavior of solder interconnects has been made recently, and thiswas covered in Section 6. By employing a testing method of differentiating TM from EM, the TMbehavior of Pb was understood in terms of morphological evolution, atomic transport and by numer-ical simulation. Pb shows a negative heat of transport. The TM of Sn has also been studied. On the basisof experimental findings it is speculated that Sn atoms exhibit a different TM characteristic opposite tothat of Pb atoms. More recently, the TM of interstitial Cu has been reported, which states that the Cualso has a negative heat of transport. In addition, a specific line-type test structure has been utilized toinvestigate the combined effect of EM and TM of Bi. It was revealed that Bi has a similar TM behaviorto that of Pb.

7.2. Further problems which need to be addressed in the near future

The current carrying capability of Al/UBM/solder should be considered based on the limitation dueto EM in the design rules. Routing design of Al interconnects needs to be implemented to mitigate thecurrent crowding and Joule heating. A pad structure that produces a uniform current distributionwithin the bump interconnect is recommended to avoid the dissolution of Al. In the case with EMproblems of Al, a relative enlargement of the cross-section of the Al trace is also a factor that couldbe considered. For UBM, in Section 3.2, the solution of a thick Cu pillar with a Ni electroplated layerhas been suggested, which can be expected to alleviate the effect of current crowding and accompa-nying heat accumulation at the interface. For solders, new candidates with improved current densitycapabilities and a higher operating temperature are expected. Considering the Sn-grain rotation undercurrent stressing, a further nano-doping into solder is anticipated to stabilize the microstructure as tolimit the fast interstitial diffusion of noble or near-noble metals. In addition, taking into account dif-ferent applications in industry, EM studies in Sn-based solders should be extended to other lead-freesolders, such as InAu.

Although void formation at the UBM side has been confirmed as a major failure mechanism in sol-der interconnects for high current density applications, IMC growth cannot be ignored, which has beendiscussed in Section 4. Moreover, as mentioned in Section 5.2, the extensive IMC growth at the sub-strate side exerts a great influence on the mechanical reliability. It is known that the kinetics dominat-ing interfacial reactions have not been established yet, and the laws of IMC growth (parabolic orlinear) under current stressing are still under investigation. Therefore, the interfacial reactions undercurrent stressing are important and challenging problems, particularly for the application of micro-


bumps with Cu pillars in 3D packages, such as TSV (through silicon via) bonding, where the IMCs sub-stantially form at the interface and hence a major phase transformation is induced [124,125].

Another issue of key importance is the TM of Sn. Although there have been a few studies made re-cently, the present understanding of the TM behavior of Sn is still limited. If one need differentiate TMfrom EM further, it is suggested that each end of the connecting wires (same materials) in a line-typetest structure should be put at different temperatures [126]. In this way, the effect of current stressingis completely removed and only the driving force due to TM is available. As compared with flip chipsamples, the advantages of line-type samples lie in the fact that the temperature gradient can be mea-sured more easily.

So far, few studies on the combined effect of TM and mechanical stress have been reported [127].For the mechanical characteristics, it should be mentioned that high strain rate fracture failure is asimportant as low cycle fatigue failure, which has been of major concern in recent years. Shear testsand tensile tests should be performed to evaluate the mechanical behavior of solder joints after TM.If atoms migrate from the higher to the lower temperature side by the force due to the thermal gra-dient, a reversed flux of vacancies will move to the higher temperature side. Consequently, the inter-face at the higher temperature side will become mechanically degraded. This interaction between TMand the applied stress is of considerable importance.

Acknowledgements

We would like to acknowledge Dr. M.O. Alam, Dr. B.Y. Wu and Dr. X. Gu at the EPA Centre, City Uni-versity of Hong Kong. Without their significant contributions, the overview cannot be finished. Specialthanks are extended to Prof. B. Ralph at Brunel University, UK, and Dr. J. Shen at the EPA Centre, CityU,HK, for their valuable suggestions.

We also wish to express our sincere gratitude to Prof. K.N. Tu’s group at University of California atLos Angeles (Dr. H. Gan, Dr. C.Y. Liu, Dr. J.W. Nah, Dr. F. Ren, Dr. A.T. Huang, Dr. F.Y. Ouyang), Prof. C.Basaran’s group at State University of New York at Buffalo (Dr. H. Ye), Prof. C. Chen’s group at NationalChiao Tung University (Dr. S.H. Chiu, Dr. C.C. Wei, Dr. H.Y. Hsiao), Prof. A. Lee at Michigan State Uni-versity, Prof. M.H. Ren at National Sun Yat-sen University, Dr. K.L. Lin’s group at National Cheng KungUniversity (Dr. Y.H. Liu), Dr. L.H. Xu at Intel, Chandler, AZ, for the permission in using their figures todevelop discussions in this review.

This work is supported by RGC of Hong Kong – GRF project: RGC project no. 111309 and CityU pro-ject no. 9041486.

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