Localized recrystallization and cracking of lead-free solder interconnections under thermal cycling

Localized recrystallization and cracking of lead-free solder

interconnections under thermal cycling

Hongtao ChenDepartment of Electronic, Aalto University School of Science and Technology, FIN-00076 Aalto, Finland; and

Shenzhen Graduate School, Harbin Institute of Technology, Shenzhen 518055, China

Maik MuellerElectronics Packaging Laboratory (IAVT), Technische Universität Dresden, 01062 Dresden, Germany

Tonu Tuomas Mattilaa) and Jue LiDepartment of Electronic, Aalto University School of Science and Technology, FIN-00076 Aalto, Finland

Xuwen LiuLaboratory of Materials Science, Department of Materials Science and Engineering, Aalto University School of

Chemical Technology, FIN-00076 Aalto, Finland

Klaus-Juergen WolterElectronics Packaging Laboratory (IAVT), Technische Universität Dresden, 01062 Dresden, Germany

Mervi Paulasto-KröckelDepartment of Electronic, Aalto University School of Science and Technology, FIN-00076 Aalto, Finland

(Received 21 January 2011; accepted 18 May 2011)

The failure mechanism of lead-free solder interconnections of chip scale package–sized Ball Grid

Array (BGA) component boards under thermal cycling was studied by employing cross-polarized

light microscopy, scanning electronic microscopy, electron backscatter diffraction, and nano-

indentation. It was determined that the critical solder interconnections were located underneath the

chip corners, instead of the corner most interconnections of the package, and the highest strains and

stresses were concentrated at the outer neck regions on the component side of the interconnections.

Observations of the failure modes were in good agreement with the finite element results. The failure

of the interconnections was associated with changes of microstructures by recrystallization in the

strain concentration regions of the solder interconnections. Coarsening of intermetallic particles and

the disappearance of the boundaries between the primary Sn cells were observed in both cases. The

nanoindentation results showed lower hardness of the recrystallized grains compared with the non-

recrystallized regions of the same interconnection. The results show that failure modes are dependent

on the localized microstructural changes in the strain concentration regions of the interconnections

and the crack paths follow the networks of grain boundaries produced by recrystallization.

I. INTRODUCTION

Solder interconnections still remain the predominant

method to form electrical and mechanical connections

between electronic components. Solder interconnection

reliability has received an increasing attention during the

recent years due to the miniaturization trend of electronic

devices as well as the adoption of the lead-free solder

alloys.1–6 Numerous studies have been carried out to identify

the root causes of failures in electronic devices under different

operation conditions. Most often electrical failures are caused

by the mismatch in the coefficient of thermal expansion

(CTE) between adjoining materials as a consequence of

changes of the device temperature due to either internally

generated heat dissipation of integrated circuits or chances of

external operation temperature. The evolution and growth of

interfacial intermetallic compound (IMC) layers during oper-

ation has often been assumed to dictate the performance of

solder interconnections due to their inherently brittle nature

and tendency to form defects. However, conditions under

which strains are concentrated on them are quite limited;

cracking of intermetallic layers appears primarily under

mechanical shock loading or fast, high-amplitude bending/

vibration. However, under thermal cycling conditions, the

solder interconnections undergo continuous cyclic plastic

deformation that leads to cracking of the solder bulk instead

of in the interfacial IMC layers of solder interconnections, see,

e.g., Refs. 7–15. Therefore, to have a better understanding of

the failure mechanism of the solder interconnections, empha-

sis should be placed, in addition to the interfacial IMC layers,

on (i) thorough microstructural characterization of “bulk”

solder interconnections, (ii) the investigation of changes of

a)Address all correspondence to this author.e-mail: [email protected]

DOI: 10.1557/jmr.2011.197

J. Mater. Res., Vol. 26, No. 16, Aug 28, 2011 �Materials Research Society 2011 2103

microstructures during operation, and (iii) the associated

changes of mechanical properties of the solder bulk.

From microstructural perspective, damage of solder

interconnections is a result of the accumulation of internal

energy in the form of increased dislocation densities in the

plastically deformed regions of the solder interconnections.

The fact that the commonoperation temperatures of electronic

devices are relatively high with respect to the melting temper-

ature ranges of most Sn-rich alloys [for Sn the homologous

temperature (T/Tmp in K) at room temperature is about 0.6]

makes the time-dependent plastic deformation (i.e., creep) of

solder interconnections significant even at low stress levels,

such as those produced under commonly used accelerated

thermal cycling test conditions. The internal energy stored in

the solder interconnections during plastic deformation pro-

vides the driving force for the restoration processes, namely

the recovery and recrystallization. However, the mechanisms

by which the internal energy is decreased by the two

processes are different, and the activation energy of recovery

is lower than that of recrystallization. Recovery is the

annihilation of lattice defects, mainly dislocations, by their

movement to the grain boundaries where they disappear and

their rearrangement to form subgrain networks by polygoni-

zation. Recovery takes place effectively in materials that have

high stacking fault energy, such as pure Sn,12,13 but the

efficiency of recovery is diminished by impurities, second

phase particles, and tangling of dislocations in the regions

of high dislocation density. Therefore, in principle,

recovery can decrease the driving force so that the

recrystallization is initiated only under restricted conditions.

Our recent investigations have shown that SnAgCu solder

slabs compressed up to 50% of the original height do not

recrystallize when annealed at 100 °C immediately after the

deformation.14 However, several investigations have

demonstrated that near-eutectic SnAgCu interconnections

do undergo a discontinuous change of microstructures

under thermal cycling as well as power cycling condi-

tions.7–11,13–21 Cracks are very often observed to propagate

in the regions where the recrystallization has most signif-

icantly changed the microstructure.7,9,11,15,21

In this article, the failuremodes andmechanisms of Sn-rich

solder interconnections under thermal cycling conditions are

investigated. Particular attention is paid on themicrostructural

changes preceding and taking place during the crack initiation

and propagation. The development of solder interconnections

microstructures and the formation of the cracks under the

thermal cycling conditions are studied by component boards

with different solder interconnection metallurgies.

II. MATERIALS AND METHODS

Figure 1 shows the structure and dimensions of the

Integrated Circuit (IC) package used in the study. The

package was a FR4-based chip scale packaged (CSP) Ball

Grid Array with 12 mm � 12 mm � 1 mm package

dimensions, 144 Sn3.0Ag0.5Cu or Sn1.2Ag0.5Cu 0.05Ni

bumps with the pitch of 0.8 mm and the bump diameter of

0.5 mm. No underbump metallization was applied on the

copper bump pads of the component. Electroless Ni with

immersion Au (Ni(P)|Au) or organic solderability preserva-

tive (Cu|OSP) was coated on the copper pads of the printed

wiring boards (PWBs). Three after-reflow compositions

of solder interconnections, Sn3.1Ag0.5Cu (SnAgCu),

Sn3.1Ag0.52Cu0.24Bi (SnAgCuBi), and Sn-1.1Ag-0.52

Cu-0.1Ni (SnAgCuNi), were formed by reflow soldering

the packages with different commercial solder pastes

(“SnAgCu”: Sn3.0Ag0.5Cu bumps and Sn3.8Ag0.7Cu

solder paste, “SnAgCuBi”: Sn3.0Ag0.5Cu bumps and

Sn3.0Ag0.7Cu3.0Bi solder paste, and “SnAgCuNi”:

Sn1.2Ag0.5Cu 0.05Ni bumps and Sn0.7Cu0.1Ni solder

paste). It should be noted that the above-reported compo-

sitions are calculated approximations determined based on

the nominal compositions of the solder bump and paste

alloys, volume ratios calculated from the nominal bump

dimensions and paste particle contents, and the assumption

of perfect mixing of the alloys. It should also be noted that

the compositions of the interconnections changes slightly

during the reflow soldering because of the dissolution of Cu

or Ni from the soldering pads of the PWB (for estimation of

the effects of dissolution, see, e.g., Ref. 7).

Reflow soldering of the assembled component boards was

carried out in a commercial in-line forced convection reflow

soldering oven (EWOS 5.1 N2, EPM/Heraeus, Hanau,

Germany) under air atmosphere. The temperature on the

top of the PWB underneath the CSP component was

FIG. 1. Dimensional drawing of the component.

H. Chen et al.: Localized recrystallization and cracking of lead-free solder interconnections under thermal cycling

J. Mater. Res., Vol. 26, No. 16, Aug 28, 20112104

measured with thermocouples (Super M.O.L.E. Gold, ECD,

Milwaukie, OR). The peak temperatures varied between 240

and 245 °C, and the time above 217 °C was about 40–43 s

depending on the location of the package on the board.

The component boards were thermally cycled

(ENX12-7.5CWL, ESPEC, Hudsonville, MI) between

�40 °C and +125 °C by allowing 15 min dwell times at

each temperature extreme and adjusting the temperature

ramp time to 7 min in both directions. Samples were

removed from the oven every 10 days to observe the

microstructural evolution.

Cross sections of the thermally cycled samples were

prepared along the diagonal of the component. The samples

were first cut further away from the diagonal line with the

help of a water-cooled, low-speed diamond saw to be sure

that the cutting would not disturb the material to be

inspected. The diagonal cross section of the interconnections

was reached by careful grinding. To reveal the microstruc-

ture of the critical solder interconnections, the cross-sections

were polished by standard metallographic methods. In the

minimum 30 samples of each material combination were

inspected. Each sample included 8�12 cross sections of the

interconnections depending on the orientation of the cross

section with respect to the package. Optical microscopy with

cross-polarized light (BX60, Olympus, Center Valley, PA)

was used to inspect differences in crystal orientations. Crystal

orientations and relative misorientations were quantified with

the help of orientation imaging microscopy (OIM) by

employing the electron backscatter diffraction (EBSD). The

scanning electron microscopy (SEM, JEOL6330F, JEOL,

Tokyo, Japan) was used to study the distributions of the IMC

particles of the bulk solder.

Nanoindentation measurements were performed on the

polished cross sections of the solder interconnection with

the help of the TriboIndenter (Hysitron Inc., Minneapolis,

MN) nanomechanical testing instrument.With an ultrahigh-

resolution load and displacement control, the device is able

to determine various material properties, such as hardness

and elastic modulus, from surfaces in submicron scales.

Each indentation cycle included a 15-s holding period

before unloading to reduce the effects of viscoplasticity

on the measured properties. The Berkovich indenter with

a 130-nm tip curvature was used in the measurement.

Stress–strain analysis of the solder interconnections was

carried out with the help of the commercial finite element

software ANSYS v. 11.0. The symmetrical structure of the

component assembly allowed the employment of the

quarter model. The model composed of a global model

with coarse meshes and a local model with fine meshes.

The constraint equations were used to tie together the

meshes in the global and local models, and the displace-

ments were transferred along the boundary between the

two models. Each solder interconnection was roughly

meshed with 96 elements in the global model, while the

finer meshes with 3060 elements were applied for the

diagonal solder interconnections in the local model. The

local model had realistic shape of the solder interconnec-

tion and detailed pad design (solder-masked-defined pad

on the component side and non-solder-masked-defined

pad on the PWB side).

The elastic material properties for thermomechanical sim-

ulation are listed in Table I.22–24 The solder material was

modeled as viscoplastic using Anand’s constitutive equations

with parameters provided by Reinikainen.25,26 The Anand

constitutive model is composed of a flow equation and three

evolution equations that describe strain hardening or soften-

ing during the primary and secondary creep stages.25

The PWB and the interposer layer of the package were

orthotropic. All the other materials in the model, including

solder, were assumed isotropic. The three-dimensional

element type SOLID45 was assigned to all the materials

except the solder interconnections. The solder intercon-

nections were modeled with the element type VISCO107,

which is a viscoplastic solid element for the Anand

constitutive model and is designed to solve rate-dependent

large strain plasticity problems. The symmetry boundary

conditions were applied as mechanical constraints, and the

center node at the bottom of PWB was fixed to prevent

free rigid body motion.

III. RESULTS AND DISCUSSION

The finite element method was used to characterize the

strain and stress concentrations in different interconnec-

tions of the package as well as in the interconnections with

the highest stress and strain. The results were correlated

with the observations from the physical failure analyses.

A. FEM results and analysis

The simulation results of the global model showed that

the solder interconnections with highest strains are located

at the corner locations of the inner periphery underneath

the Si chip instead of the corner most interconnections of

the package (Fig. 1). This result is in agreement with the

observations from the physical failure analyses, and similar

results have been presented in Ref. 27. Compared with the

TABLE I. Elastic material properties for thermal-mechanical

simulation.

Materials

Young’s

modulus (GPa) CTE (ppm/°C)

Poisson’s

ratio

Silicon 131 2.7 0.3

FR-4 In-plane 17.7 16 0.28

Out-of-plane 7.72 84

Substrate In-plane 18.8 15 0.39

Out-of-plane 8.27 57

Mold compound 15.8 15.5 0.3

SnAgCu 49 � 0.07T 21.3 + 0.017T 0.35

Copper 117 17 0.38


J. Mater. Res., Vol. 26, No. 16, Aug 28, 2011 2105

corner most solder interconnections of the package, the

critical solder interconnections have higher viscoplastic

strain energy densities and stresses. A possible reason for

this distribution of strains and stresses is related to the

structure of the package and the choice of its materials. The

fact that the silicon chip, whose CTE is low (Table I), is

attached to the FR4 interposer of the package (with

relatively high CTE) influences the total CTE of this area

and causes a mismatch in thermal expansion between

component and PWB. While the solder joints in the center

region underneath the diemay assumingly experience lower

stress than the corner most interconnections of the package,

the strains and stresses increase toward the edges of the rigid

Si die. The strains/stresses in the interconnections along the

outermost periphery of the package (which is not covered

by the die) remain low due to the diminished influence of

the silicon chip. In the edge region of the package, the

CTE mismatch is dictated by the interposer/mold com-

pound against the PWB. The maximum viscoplastic strain

energy density, the von Mises stress, and the normal stress

of the critical solder interconnections are 0.415, 45.13, and

41.73 MPa, respectively. The corresponding values of the

corner most interconnections of the package are 0.034,

20.49, and 17.95 MPa, respectively. Figure 2 shows the

contour plots of the viscoplastic strain energy density and

the normal stress on the diagonal cross section of the critical

interconnections at the low-temperature portion of the

thermal cycle. The von Mises stress has similar distribution

as the viscoplastic strain energy density. As shown in

Figs. 2(a) and 2(b), the maximum normal stress and the

maximum viscoplastic strain energy density are highly

localized at the outer corner regions on the component side

of the solder interconnections.

B. Observations on the evolution of

microstructures during thermal cycling

The inspections of the as-soldered microstructures of

the interconnections revealed differences in microstruc-

tures between the compositions under investigations.

Figure 3 shows the microstructures of all the combina-

tions of experimental variables as inspected with the help

of the cross-polarized light microscopy and SEM. As

inspected under the cross-polarized light, differences in

the number of crystal orientations were observed be-

tween the compositions of the interconnections but not

between the interconnections formed on the different

protective coatings. In principle, the primary phase of

solidification in near-eutectic SnAgCu solder alloys can

be b-Sn, Cu6Sn5, or Ag3Sn depending on the composi-

tion, but the nominal after-reflow compositions of the

interconnections in this study are in the primary b-Snphase region. Thus, the solidification structures of the

interconnections are dictated by the solidification of

b-Sn. High degree of undercooling as caused by the

suppressed nucleation of primary Sn during solidification

has been reported in the literature.28 The high under-

cooling results in a cellular (or dendritic) structure of the

SnAgCu solder joints.

In the as-reflowed conditions, the solder interconnections

are composed of relatively few solidification colonies that

enclose the uniformly oriented Sn cells. (The term “solid-

ification colony” is used here to emphasize the fact that

FIG. 2. The left-hand side edge of the interconnection faces the center,

and the right faces the corner of the package: (a) viscoplastic strain

energy density distribution in the local model, (b) normal stress

distribution in y-direction in the local model, and (c) quarter view of

the interconnections of the package shows the distribution of von Mises

stresses in different interconnections.



under the reflow conditions employed, a cellular structure is

generated in which the difference of crystal orientations

between individual cells is very small, below 3°. The

difference in crystal orientation between the colonies is

high, above about 10–15°. The use of this term helps us also

to make a verbal distinction between the as-solidified

microstructures and the recrystallized grains.) The cellular

structure of the b-Sn is clearly distinguishable as they are

surrounded by regions composed of small Cu6Sn5 and

Ag3Sn IMC particles uniformly dispersed in b-Sn. The topimages in Fig. 3 show an example of a cross section of each

interconnection composition in the as-reflowed state as

imaged by employing the cross-polarized microscopy. In

the after reflow conditions, the interconnections consist of

relatively few solidification colonies of uniformly oriented

b-Sn cells distinguished by high-angle boundaries. The

cross sections of the SnAgCu and the SnAgCuBi intercon-

nections showed distinctly fewer number of crystal orienta-

tions (typically only one or two) than the SnAgCuNi

interconnections (typically three to eight). Similar results

have been reported in other studies.29,30 The larger number

of crystal orientations in SnAgCuNi interconnections is

most likely a consequence of the low Ag content in the

solder bulk, which is also shown in the comparison of the

binary SnAg and SnCu solder interconnections on different

underbump metallization reported in the literature.31,32

As inspected with the help of the SEM with higher

magnification, differences were observed also between the

interconnections on the soldering pads with different

protective coatings, namely in the amount and distributions

of Cu6Sn5 and Ag3Sn intermetallic phases. The SEM

images in Fig. 3 illustrate that the cellular structure of the

b-Sn in the SnAgCu and the SnAgCuBi interconnections

was clearly distinguishable [as they were surrounded by

thick regions composed of uniformly distributed Cu6Sn5and Ag3Sn in the b-Sn matrix; Figs. 3(a), 3(b), 3(c), 3(d),

and 3(e)]. However, the SnAgCuNi interconnections did

not show a clear cellular structure and the intermetallic

FIG. 3. Scanning electron microscopy (SEM) (top row) and polarized light (bottom row) images of the as-solidified solder interconnections with

different PWB protective coatings: (a) SnAgCu (SAC)/OSP, (b) SnAgCuBi (SACB)/OSP, (c) SnAgCuNi (SACN)/OSP, (d) SAC/Ni(P)|Au, (e)

SACB/Ni(P)|Au, and (f) SACN/Ni(P)|Au.



particles were muchmore evenly distributed over the whole

volume of the solder interconnections [Figs. 3(c) and 3(f)].

This difference is most likely a consequence of lower

Ag content of the SnAgCuNi interconnections compared to

the Ag content of the SnAgCu and the SnAgCuBi compo-

sitions. Furthermore, the microstructures of the intercon-

nections on the Cu|OSP-coated soldering pads were more

heterogeneous than those of the interconnections on the Ni

(P)|Au due to the larger amounts of primary Cu6Sn5particles. This is caused by the several orders of magnitude

higher dissolution rate of Cu into the molten Sn compared

with that of Ni and the increased Cu content of the melt of

the Cu|OSP interconnections.33 In one of our earlier studies,

the dissolution rate of Cu in Sn0.5Ag0.5Cu (wt%) was

found to be about 0.07 lm/s (at;250 °C), whereas that of

Ni is only 0.003 lm/s.7

The development of microstructures was studied from

samples that were taken out in the course of the thermal

cycling test and cross sectioned for inspection. During the

test, the as-solidified microstructures transformed into

more or less equiaxed grain structures. Typically only

a fraction of the solder interconnection cross section

participates in cyclic deformation because the stress/

strain distribution inside the barrel-shaped solder inter-

connections concentrated on the edges of the solder

interconnections close to both interfacial regions as

indicated by the results of the finite element analysis.

Figure 4 shows a typical example of a failed intercon-

nection, where cracking of the solder interconnection is

accompanied by the distinct change of microstructures by

recrystallization. The changes of microstructures were first

observed in the strain concentration regions and, after the

initiation of recrystallization, the recrystallized volume

gradually expanded from the edges toward the center,

across the interconnections in the interfacial regions of the

interconnections. Comparison of the images shows that

the propagation path of the crack [Fig. 4(a)] is enclosed

entirely within the recrystallized region of the intercon-

nection [Fig. 4(b)]. The incubation time of recrystalliza-

tion varied significantly from one interconnection to

another. The first indications of recrystallization were

observed in the samples that were taken out after 500

thermal cycles, but it took about 2000 cycles until

recrystallization was consistently observed in every

interconnection on the corner regions of the packages.

The incubation time of recrystallization in the SnAgCuNi

interconnections was the shortest. They began to recrys-

tallize after about 500 cycles. The incubation time of the

SnAgCu was in the range of 1000–1500 cycles and that of

the SnAgCuBi interconnections was about 1000–2000

cycles. Differences between the incubation times can be

attributed to the fact that the strength of the SnAgCuNi

alloy is lower as compared with the SnAgCu and SnAg-

CuBi alloys. Due to the lower amount of Ag, higher

amount of plastic work per thermal cycle is stored in the

material and the accumulation of the stored energy in the

low-silver SnAgCuNi interconnections is therefore faster.

Furthermore, the onset of recrystallization on the inter-

connections on the Cu|OSP-coated PWBs was faster than

that on the interconnections on the Ni(P)|Au regardless of

the solder composition.

Figures 4(c)–4(f) show the OIM results obtained from the

EBSDmeasurement of the cross section in Figs. 4(a) and 4(b).

Figure 4(d) shows the orientation in normal direction ([001]

inverse pole figure) and allows a comparison with the cross-

polarized light micrograph in Fig. 4(b). Both pictures are in

good agreement, and the majority of the b-Sn grains were

already observable with cross-polarized light microscopy.

Themisorientation between adjoiningb-Sn grains is shown inFigs. 4(c) and 4(e). Both figures show that most of the

misorientations are below 5°. It is likely that those boundaries

are subgrains or low-angle grain boundaries formed by

dislocations. The misorientation distribution in Fig. 4(c) also

shows only a small amount of high-angle grain boundaries

that are mostly gathered in the crack area [Figs. 4(e) and 5(a)].

No special boundaries such as twins (57.2° and 62.8° rotation

around [100] axis with twin planes K1 {101} and {301},

respectively) have been found in the present solder intercon-

nections. This result differs from the findings of Lehman

et al.34 This finding points out that there may be significant

differences in the microstructures of solder interconnections

from one component board to another. At the time of writing,

the reasons for these differences are not fully understood but it

is well known that solder composition, choice of PWB

protective coating, choice of underbump metallization, and

the temperature–time profile of the reflow soldering process

all influence the microstructures of solder interconnections.

Because of such differences in the microstructures, it is also

possible that the failure modes and lifetimes differ from one

board assembly to another.

Figure 4(d) shows that the center region of the solder

joint consists of several larger grains with gradual change in

orientation, which is indicated by the inserted unit cells.

Those grain boundaries correlate in general with the mis-

orientation peaks at 4°, 7°, and 9° [Figs. 4(c) and 4(e)].

Since this area experiences the lowest stress and strain

during the thermal cycling test, it is reasonable that the grain

structure in the center of the solder joint still represents the

as-solidified state. The areas close to both interfaces

experienced higher strain and, therefore, more significant

changes in grain structure. The cracked region on the

package side interface of the solder interconnections shows

primarily high misorientations of the grains. Figure 4(e)

illustrates well that the high-angle grain boundaries are

located very close to the crack path, while the misorienta-

tion becomes smaller the larger the distance from the crack

region is. Figure 4(f) shows all grain boundaries between

adjoining grains that are formed by a rotation around

a certain axis. It shows that a large amount of low-angle

boundaries of the interconnection are formed by rotations



around [100], [110], and [001], while most high-angle grain

boundaries in the cracked region and in the strain concen-

tration region on the PWB side interface do not have this

relationship. A very narrow tolerance (1°) of the rotation

axis was chosen in the analysis to distinguish between the

different rotations causing the low-angle grain boundaries

(when a higher tolerance, e.g., 5°, was chosen, more low

misorientation boundaries were highlighted and all three

rotations around [100], [110], and [001] overlapped). These

three rotations and the mentioned overlap could be

explained by the simultaneous activation of different slip

systems which could form dislocation walls and create

FIG. 4. (a) The bright light micrograph shows the crack path of a failed solder interconnection during thermal cycling, (b) a cross-polarized light image

of the same interconnection shows the change of grain structure as caused by the cyclic deformation, (c) orientation imaging microscopy misorientation

distribution between adjoining b-Sn grains found in the solder joint, (d) inverse pole figure [001] showing the orientation in normal direction, (e) electron

backscatter diffraction map of the cross section showing the misorientations and distribution of grain boundaries, and (f) grain boundaries by rotations

around certain directions (note that boundaries where red [100] and green [110] lines overlap appear olive green).



rotating subgrains. It is remarkable that rotations around

[101] and [111] are not affected significantly by this

tolerance change. Independent from the chosen tolerance,

it was shown that high-angle grain boundaries above 15° do

not show this relationship.

Figure 5(b) shows a higher magnification of the

microstructures on the PWB side of the interconnection,

where cracking has not yet taken place but similar

coarsening of the intermetallic particles and formation

of grain boundaries can be identified. A closer inspection

of the region shows that primarily low-angle misorienta-

tions are formed on the PWB side region, where the

recrystallization has not yet progressed as far as in the

vicinity of the crack path on the package side interfacial

regions of the solder interconnections. It seems that the

formation of the low-angle boundaries is a preliminary

stage before recrystallization. A similar observation has

been reported also in Ref. 35.

A correlation of the microstructures on the PWB

[Fig. 5(b)] and the package side regions [Fig. 5(a)]

indicates that the Sn cells begin to rearrange by the

gradual coalescence of the Sn cells, during which the

cellular structure disappears and the low-angle boundaries

emerge. The regions surrounding the Sn cells formed during

solidification disappear as the small intermetallic particles

coalesce and distribute sparsely mainly on the formed grain

boundaries. These experimental findings indicate that there

are in fact two different mechanisms that cause the observed

changes of microstructures. The first to take place is the

gradual rotation of small volumes in the stress-concentration

regions. The mechanism following this stage is the discon-

tinuous change of microstructures by recrystallization.

It is well known from the microstructural perspective

that damage of solder interconnections is a result of the

accumulation of internal energy in the form of increased

dislocation densities in the plastically deformed regions of

the solder interconnections (i.e., where strain hardening

takes place). The stored energy is released through re-

covery (annihilation of lattice defects and the formation of

low-angle grain boundaries due to a rearrangement of

dislocations in deformed materials with a high dislocation

density) and recrystallization (the nucleation and growth of

new defect-free grains). Thus, in principle hardness of the

FIG. 5. Higher magnification backscatter scanning electron micro-

graphs of the cross section shown in Fig. 4: (a) cracked region on the

package side and (b) left corner on the PWB side (black particles are

Cu6Sn5 intermetallics, bright and grey particles are Ag3Sn). The grain

boundaries have been inserted manually by superposition of the EBSD

results.

FIG. 6. Image under optical microscope with cross-polarized light after

indentation test.



strain-hardened regions of the solder interconnections

should be higher compared with the solidification colonies

in the center regions of the solder interconnections, while

the hardness of the recrystallized grains should be lower

than that of the strain-hardened regions. To characterize

the different regions, hardness and elastic modulus of the

as-soldered, recrystallized, and the rotated colonies were

measured with the help of the nanoindentation test. Figure

6 shows a cross-polarized micrograph of the SnAgCuNi

solder interconnection in which the indents are clearly

visible. The high-resolution scanning probe microscopy

(SPM) images show that some of the grains boundaries

were made visible by the stress field of the indentation tip

during the measurements. The measured values of hardness

of the three different regions are presented in Fig. 7. The

values are averages from 12 to 18 indents, and the indicated

range is the standard deviation of the measurements. The

hardness of the rotated regions is about 17% higher than that

of the as-soldered regions, and the hardness of the recrystal-

lized regions is about 30% lower than the rotated regions. It

should also be noted that standard deviation of the hardness

measurements from the recrystallized regions is only about

8% of the average while that of the as-solidified regions is

about 15%. It is suggested that the microstructural changes

shown in Fig. 4(b) are caused by increased dislocation

densities in the regions of stress–strain concentration and

the subsequent rotation of the solidification colonies (by

recovery) to reduce the internal stresses.

It is particularly noteworthy in Figs. 5(a) and 5(b) that

the coarsening of intermetallic particles is strong in the

heavily deformed corner regions of the interconnections,

and the coarsened intermetallic particles have (often) been

left at the boundaries of recrystallized Sn grains. The

element mapping results show that the coarsened interme-

tallic particles are mostly Ag3Sn. Similar results have been

reported in Refs. 36–38. Compared to solder interconnec-

tions with high Ag content (Sn3.1Ag0.5Cu, Sn3.1Ag0.52-

Cu0.24Bi), the coarsening of Ag3Sn in low Ag content

interconnections (Sn–1.1Ag0.52Cu0.1Ni) was not as ob-

vious as it was in the higher Ag content SnAgCu and

SnAgCuBi interconnections. Furthermore, similar coars-

ening of the Cu6Sn5 or (Cu,Ni)6Sn5 particles was not

observed, most likely due to the low concentration of

Cu and Ni (about 0.5 wt% and less than 0.1 wt%,

respectively). The difference in the coarsening rate of

the Ag3Sn particles between the “as-soldered” midsec-

tions of the interconnections and the regions undergoing

FIG. 7. Measured hardness and elastic modules of as-soldered, recrystal-

lized, and rotated microstructures (with standard deviations).

FIG. 8. Cross-polarized and SEM images of SACB|OSP solder inter-

connections after thermal cycling for 90 days: (a) cross-polarized image,

(b) a SEM micrograph and an element map of Ag of the area A, and (c) a

SEM micrograph and an element map of Ag of the area B.



cyclic deformation was significant. Thus, the primary

reason for the coarsening seems to be the increased

diffusion rates caused by the cyclic deformation rather

than the cyclically increased temperature.

There is a similarity between the recrystallized micro-

structures and the rotated regions, namely that the edges of

the different crystal orientations are clearly distinguishable,

but the disappearance of the cell boundaries is not as

FIG. 9. Cross-polarized images of solder interconnections after thermal cycling for 90 days: (a) SAC interconnections on Cu|OSP and (b) SAC

interconnections on Ni(P)|Au.

FIG. 10. Cross-polarized and SEM images of SnAgCuBi solder interconnections after thermal cycling for 90 days: (a) cross-polarized image,

(b) SEM image, (c) higher magnification of A and B areas in (b), and (d) Ag mapping of (c).



obvious in the rotated regions as in the recrystallized

regions. Differences in the number, size, and distribution

of the intermetallic particles between the changed [Fig. 8(b)]

and the as-soldered regions [Fig. 8(c)] of the interconnections

are distinguishable. The boundaries composed of the small

Cu6Sn5 and Ag3Sn intermetallic particles around the Sn cells

have faded and the Ag3Sn intermetallic particles have

coalesced and became sparsely distributed.

The nucleation mechanism of the recrystallized grains is

yet uncertain, however. Based on what is being said above it

is likely that the high-angle grain boundaries formed by the

rotation and coalescence of the subgrains act as a nucleation

sites of recrystallizing grains. However, nucleation can also

take place at the grain boundaries (preferable at triple points)

of the Sn matrix and the intermetallic particles. The stable

nuclei, formed by either of the mechanisms, then gradually

grow at the expense of the neighboring grains and consume

all the deformed volume of the interconnections. The final

outcome is an equiaxed microstructure (high-angle misor-

ientations between crystals in random orientation), where

cracks nucleate and propagate intergranularly through the

solder interconnections along the continuous network of

high-angle grain boundaries.

C. Nucleation and propagation of cracks

It is particularly noteworthy that even though there was

considerable variance in the incubation period of recrystal-

lization, the nucleation of cracks took place within relatively

narrow range between about 1000 and 1500 thermal cycles.

Thus, nucleation of cracks can take place before the change

of microstructures by recrystallization, but the propagation of

cracks without the influence of recrystallization seemed

FIG. 11. Microstructure of a SnAgCuNi interconnection after thermal cycling for 60 days: (a) cross-polarized and (b) SEM micrographs; 80 days:

(c) cross-polarized and (d) SEM micrographs; 90 days: (e) cross-polarized and (f) SEM micrographs.



sluggish judging by the measured crack lengths of samples

that were taken out of the test chamber at the same point of

time. Furthermore, recrystallization had initiated eventually

in all failed solder interconnections, and the networks of

grain boundaries had influenced the propagation of cracks in

all electrically failed interconnections observed in this study.

The failure analyses showed that the crack had com-

monly propagated on the package side of the solder

interconnections in the proximity of the intermetallic

layers. Inspection of the samples that were removed from

the cycling chamber sequentially after the start of the test

revealed that the cracks sometimes nucleated before

distinct recrystallization could be observed [Fig. 9(b)].

However, the propagation of cracks was influenced by the

formation of the networks of grain boundaries by recrys-

tallization [as shown in Fig. 9(a)] in all electrically failed

interconnections. As illustrated in Figs. 9(b) and 10(a),

cracks can occasionally propagate relatively far without

distinguishable recrystallization (rotation of solidification

colonies are visible, however), but they were not observed

to have propagated entirely through the interconnections

without its influence. After the initiation of recrystallization

of solder interconnections, the recrystallized volume grad-

ually expanded over the diameter of the interconnections as

the cycling was continued as illustrated in Fig. 11. After the

recrystallization had initiated, cracks were not observed to

have propagated further than the recrystallized region.

Thus, it seems evident that the networks of the newly

formed grain boundaries provide favorable sites for cracks

to propagate intergranularly with less energy than in the

solidification colonies.

Decohesions of the grain boundaries were frequently

observed as shown in Fig. 12. It is highly likely that the

cohesion between the recrystallized Sn grains is low due to

the fact that impurities as well as the intermetallic particles

are expelled from the growing recrystallized grains and

pushed into the grain boundaries. Furthermore, the for-

mation of new grain boundaries by recrystallization also

provides good preconditions for plastic deformation by

grain boundary sliding. The recrystallized regions are

therefore more easily deformed plastically than the sur-

rounding regions and the damage accumulation in the

recrystallized regions is increased. Significant grain

boundary sliding can cause cavitation along the grain

boundaries and thereby enhance the propagation of cracks

through the recrystallized regions. Finally, the thermome-

chanical anisotropy of the (recrystallized) Sn grains can

also enhance the nucleation and propagation of micro-

cracks along their boundaries as the value of the CTE of Sn

single crystal in the [001] directions is about two times that

in the [100] 5 [010] direction.39 Occasionally, when the

recrystallized grain size was larger, cracks were found to

propagate also transgranularly (Fig. 13). It seems that the

cracks tend to propagate transgranuarly only when the

stress state, orientation, and geometry of the grains are not

in favor of cracking along the grain boundaries.

FIG. 12. (a, b) Cross-polarized and (c) SEMmicrographs of the SACN

interconnections on the Ni(P)|Au after thermal cycling for 90 days.

FIG. 13. (a) Higher magnification cross-polarized and (b) bright-field

micrographs of the recrystallized region in Fig. 6.



IV. CONCLUSIONS

(1) The critical solder interconnections of the BGA

packages were under the chip corners, and the stress

concentration areas were on the component side corner

regions on the package edge side of the diagonal cross

sections of the solder interconnections.

(2) No significant differences were found between the

microstructures withNi(P)|Au and OSP protective coatings,

except that the Cu6Sn5 particles were more numerous in the

interconnections soldered on the Cu|OSP PWB pads.

(3) Higher numbers of solidification colonies were

found in the as-solidified microstructures of the SnAg-

CuNi solder interconnections compared with those of the

SnAgCu and SnAgCuBi solder interconnections, which

typically were composed of only one or two solidification

colonies with different crystal orientations. Within solid-

ification colony there is a cellular structure of b-Sn (cells

surrounded by small Ag3Sn and Cu6Sn5 particles) in

which the difference in crystal orientations between in-

dividual cells is small (below 3°). The difference in crystal

orientation between the colonies is high (above about 15°).

(4) Two different types of changes of the as-solidified

microstructures were identified, and both changes were

highly localized at the strain concentration regions of the

solder interconnections.

(i) Formation and rotation of the subgrains by recovery.

(ii) The discontinuous change of microstructures into

more or less equiaxed grain structure by recrystallization.

(5) Nanoindentation test showed decreased hardness in

the recrystallized regions compared with the nonrecrystal-

lized regions of the same cross section.

(6) Nucleation of cracks was observed to take place

before the change of microstructures by recrystallization,

but the propagation of cracks was influenced by the

formation of the networks of grain boundaries by re-

crystallization in all electrically failed interconnections.

ACKNOWLEDGMENTS

The authors would like to thank Prof. Emer. Jorma

Kivilahti for his inextinguishable enthusiasm for the

topics discussed in this article. The author would like to

thank Dr. V. Vuorinen for his valuable discussion and

help in the SEM studies. The authors would also like to

thank Mr. Jussi Hokka for his help in sample preparation

and managing the thermal cycling tests. The Academy of

Finland is acknowledged for funding this work.

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Localized recrystallization and cracking of lead-free solder interconnections under thermal cycling