Creep and Failure of Lead-free Solder Alloys · Creep and Failure of Lead-free Solder Alloys ... solder joint using a miniature tensile ... A more capricious failure occurred along

Creep and Failure ofLead-free Solder Alloys

J.G.A. Theevenreport number MT02.03

Master’s thesis

Supervisor: prof.dr.ir. M.G.D. GeersCoach (TU/e): dr.ir. W.P. VellingaCoach (Philips): dr. J.W.C. de Vries

Eindhoven University of TechnologyFaculty of Mechanical EngineeringMaterials Technology Group

Eindhoven, March 2002

Contents

Abstract iii

1 Introduction 1

2 Thermodynamics 3

2.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2.2 Extremum principles and evolution . . . . . . . . . . . . . . . . . . . . . 4

2.3 Diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

3 Plastic Deformation Mechanisms 11

3.1 Diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

3.2 Deformation mechanics . . . . . . . . . . . . . . . . . . . . . . . . . . . 12

4 Phase Diagrams 15

4.1 Thermodynamics of phase equilibria . . . . . . . . . . . . . . . . . . . . 15

4.2 The Sn-Ag-Cu System . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

4.3 The Sn-Bi-Ag system . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

4.4 The Sn-Zn-Bi system . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

5 Experiments 22

5.1 Shear tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22

5.2 Digital Image Correlation . . . . . . . . . . . . . . . . . . . . . . . . . . 24

5.3 Contrast . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

5.4 Creep and coarsening . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

i

Contents ii

6 Experimental Results 28

6.1 Microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

6.2 Solder A (SnPb36Ag2) . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

6.3 Solder B (SnAg3.8Cu0.7) . . . . . . . . . . . . . . . . . . . . . . . . . . 34

6.4 Solder C (SnAg3.3Bi3.82) . . . . . . . . . . . . . . . . . . . . . . . . . . 38

6.5 Solder D (SnZn8Bi3) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

7 Conclusions and Recommendations 46

Bibliography 48

A Pure element transition data 50

B Lattice stabilities 51

C Thermodynamic data for the Sn-Zn-Bi system 54

D Binary Ag-Cu phase diagram 55

E Binary Sn-Cu phase diagram 56

F Binary Sn-Ag phase diagram 57

G Binary Ag-Bi phase diagram 58

H Binary Sn-Bi phase diagram 59

I Binary Zn-Bi phase diagram 60

J Binary Sn-Zn phase diagram 61

K EDX Analysis 62

L Shear deformation of gold grid 66

M ”Hall of Fame” 67

Samenvatting 70

Acknowledgements 72

Abstract

Increasing environmental awareness regarding lead-based solder alloys has caused majorefforts to develop a lead-free soldering technology. The RIPOSTE project of TU/e andPhilips CFT aims to construct a tool able to predict the reliability of miniature ICpackage soldered interconnects, subjected to (thermo-)mechanical stress, taking intoaccount the relevant microstructure and chemistry of the solder. This thesis is partof this project and deals mainly with shear loading, creep, failure and deformationbehaviour of three proposed lead-free solder alloys, i.e. SnAg3.8Cu0.7 (also called SAC),SnAg3.3Bi3.82 and SnZn8Bi3 and SnPb36Ag2 as a reference.

Due to high homologous operating temperatures of solders used in electronics, defor-mation of solder joints is always a thermo-mechanical process. This requires an under-standing of equilibrium as well as non-equilibrium thermodynamic principles. Duringdeformation there is a constant flow of energy toward the system, implying that we arenot at all dealing with a system in thermodynamic equilibrium.

If the temperature of an electronic system rises, the printed circuit board expands morethan the component, due to the difference in coefficients of thermal expansion, resultingin the solder connection being subjected to shear stress. Shear tests have been performedwith special samples to investigate the shear strength and maximum shear strain of thesolder joint using a miniature tensile stage. With the aid of digital image correlation(DIC) techniques, deformation behaviour at the micron level could be observed andcompared to the shear test data.

Intermetallic compounds at the solder-substrate interface have shown to play a crucialrole in solder joint deformation for SnAg3.8Cu0.7 and SnZn8Bi3, where inhomogeneousand highly local strains could be seen at the interface. SnAg3.3Bi3.82 showed a muchmore homogeneous deformation leading to a superior shear strength, but brittle fractureas well. SnAg3.8Cu0.7 showed the largest plastic deformation, whereas SnZn8Bi3 wassubject to peculiar ageing effects.

Shear stress in combination with high homologous temperatures also induces creep andpossibly coarsening. A creep test setup has been designed to test samples inside a stove.It appeared that SnZn8Bi3 has the highest creep rate of the three lead-free solder alloys,followed by SnAg3.3Bi3.82. SAC is the most creep-resistant alloy with about the samecreep rate as SnAg3.3Bi3.82 at 100˚C, but better resistance to creep at 150˚C.

iii

Abstract iv

ESEM and light microscopy revealed the failure mechanisms of the creep samples. InSAC solder joints, cracks parallel to the interface and perpendicular to these cracks(from one interface to the other), along large colonies could be seen and failure occurredalong cracks parallel to the interface throughout the bulk material. SnAg3.3Bi3.82 alsoshowed this, but failed along the intermetallics at the interface. SnZn8Bi3 showed colonyboundary sliding again, however, colonies in this material are much smaller and differ-ently shaped, so cracks could be seen throughout the whole solder joint in any direction.A more capricious failure occurred along these cracks throughout the bulk material andnot just along the interface. Also, EDX-analysis showed large diffusion of copper to thesurface of the solder joint, forming brass with zinc.

In this thesis, baseline thermodynamic data from literature is presented, together withthermo-mechanical data for the three lead-free solder alloys considered in the RIPOSTEproject. Combining these data with data from other research areas within this projectwill lead to a better understanding of the properties of these materials and a goodreplacement for lead-tin based solder alloys.

Chapter 1

Introduction

Lead-tin based solders have long been the most popular materials for electronic pack-aging because of their low cost and superior properties required for interconnectingelectronic components. However, the toxic nature of lead and the increasing awarenessof its adverse effect on environment and health have given rise to the need for develop-ment of lead-free solders in recent years. Several alloy compositions have been proposed,however there is a general lack of engineering information and there is also significantdisparity in the information available on these alloys. Nontoxic substitute materialsshould satisfy a number of criteria if they are to serve as an effective replacement forlead. A major factor affecting alloy selection is the melting point of the alloy, since thishas a major impact on the other polymeric materials used in microelectronic assemblyand encapsulation. Other important manufacturing issues are cost, availability and wet-ting characteristics. Reliability related properties include mechanical strength, fatigueresistance, coefficient of thermal expansion and intermetallic compound formation.

When an electronic device is in operation, the solder connections are subjected to me-chanical strains. The primary cause of these strains arise from the fact that the electroniccomponent and the board have different coefficients of thermal expansion. An exampleof how these strains are generated, between a silicon die and the substrate, is shownin figure 1.1. If the temperature of the system rises, the board expands more than thecomponent, resulting in the solder connection being subjected to shear strain. As thesystem is switched on and off, it is subjected to thermal cycling, resulting in the solderconnection being subjected to cyclic shear stresses.

Microstructure size has a great influence on the plastic deformation kinetics of solderjoints. Diffusion-driven separation of phases, also known as coarsening, has been wellexamined for binary alloys, like eutectic Sn-Pb solder. Experiments have shown thatdiffusion processes can be considerably accelerated by combination of high homologoustemperatures with mechanical stresses [1]. This research will investigate coarsening phe-nomena in ternary lead-free solder alloys, if present.

Constant thermo-mechanical loading will also induce creep. Creep is the most commonand important micromechanical deformation mechanism in solder joints, thus creep

1

Chapter 1. Introduction 2

Figure 1.1: Solder joints subjected to shear strain during thermal cycling due to CTE mismatchbetween die, solder and substrate [2]

resistance is an important mechanical property. Accelerated thermo-mechanical testingprovides a useful insight in the creep resistance of the different kinds of solder alloysinvestigated.

The three main goals of this research are (1) to provide baseline thermodynamic dataand theory on equilibrium and non-equilibrium thermodynamics regarding solder jointsfor further numerical modeling, (2) to develop experimental tools necessary to testcreep, shear and deformation properties and (3) to provide thermo-mechanical data onthe aforementioned properties. By characterizing the deformation behaviour of solderjoints with the use of digital image correlation (DIC) techniques and identifying themicromechanisms of their degradation, a much better understanding of the deformationmechanism can be achieved.

This research is part of the RIPOSTE project (Reliability Improvement with Pb-freesolders to Outlive in Shock and high Temperature Environments) of Philips CFT andTU/e. The aim of RIPOSTE is to construct a tool able to predict the reliability ofminiature IC package soldered interconnects, subjected to (thermo-)mechanical stress,taking into account the relevant microstructure and chemistry of the solder. Adding upthe results of this research will lead to a better insight in the mechanical properties ofthe proposed lead-free solder alloys. Combining this research with the other researchactivities of RIPOSTE will hopefully lead to one ”winning” lead-free solder alloy, whichwill eventually lead to better performance and a cleaner environment.

Chapter 2

Thermodynamics

2.1 Introduction

Due to the high homologous temperature, in deformation of solder alloys one is alwaysdealing with a thermo-mechanical process. This implies that there is a coupling betweenthermal and mechanical processes. This requires an understanding of equilibrium as wellas non-equilibrium thermodynamic principles, as well as of the specific nature of theinteractions leading to the energy terms involved in this process.

A major conceptual problem is to relate deformation mechanisms operating on severallevels to thermodynamic quantities. This is crucial, because during deformation there isa continuous flow of energy toward the system, so we are not at all dealing with a systemin thermodynamic equilibrium. In fact, it is not even certain that we are dealing with asystem that evolves towards a thermodynamic equilibrium. In fact, we can envisage thesolder alloy being deformed as a closed system, and the solder alloy plus deformationmodes as an isolated system. For a closed system, that exchanges energy not matter,the first and second law read

dU = dQ + dW (2.1)

dS = deS + diS (2.2)

where

deS =dQ

Tand diS ≥ 0 (2.3)

and deS is defined to be the change of the system’s entropy due to exchange of energy anddiS the change in entropy production due to irreversible processes within the system.

3

Chapter 2. Thermodynamics 4

2.2 Extremum principles and evolution

In systems that are allowed to evolve ”spontaneously” or isolated systems, the final,equilibrium situation is characterized by an extremum value of some thermodynamicpotential, for example the Helmholtz free energy F , the Gibbs free energy G or theenthalpy H. Processes leading to a lowering of the thermodynamic potential may occurspontaneously. An equivalent, but more general point of view is that all processes forwhich dSi ≥ 0 may occur and drive the system toward equilibrium. This is true regard-less of the specific thermodynamic restrictions such as constant pressure or volume.For the aforementioned thermodynamic potentials it follows that for any spontaneouslyoccurring process:

dF = −TdSi ≤ 0 (T, V = constant) (2.4)

dG = −TdSi ≤ 0 (T, p = constant) (2.5)

dH = −TdSi ≤ 0 (S, p = constant) (2.6)

Taking the Gibbs free energy as an example:The Gibbs free energy is defined as [3]:

G = U + pV − TS = H − TS (2.7)

Using equation (2.2), one can write the change in Gibbs free energy as

dG = dU + pdV + V dp − TdS − SdT (2.8)

dG = dU + pdV + V dp − TdeS − TdiS − SdT (2.9)

dG = dQ − pdV + pdV + V dp − TdeS − TdiS − SdT (2.10)

At constant p and T this reduces to

dG = −TdiS (2.11)

So at constant p and T , G is the appropriate potential. In thermodynamic equilibrium,the value of some thermodynamic potential is at a minimum and that of the entropy ata maximum. Phase diagrams are a way of presenting the various equilibrium phases andtheir coexistence, at a certain T and p. Since they represent the ultimate state to whicha system may evolve, they form an important basis. So in the particular case of solders,we are interested in explicit values for the Gibbs free energy. In conclusion, the systemmay evolve toward equilibrium by a sequence of steps that lower G monotonically.


For the Gibbs free energy G we find for a pure compound a, assuming the simplestpossible pairwise interaction:

Ga =Nz

2Ea (2.12)

where N is the mole number, z the coordination number and Ea the interaction energy.For a mixture of two phases with fractions xa and xb = 1 − xa one assumes

G =Nz

2(xaEa + xbEb) + TSmix + Gmix (2.13)

with Nz2 (xaEa + xbEb) the interaction energy between like atoms, Gmix the interaction

energy between unlike atoms and TSmix the term due to mixing entropy Smix. For theterm Smix we find the following. Consider the example of a crystal with a total of Nsites available for the occupation of atoms or molecules, n of which are occupied byA atoms/molecules and (N − n) are occupied by B atoms/molecules. It is clear thatthe number of configurations the system can adopt depends on the number of possiblepositions W in which the B atom or molecule can place itself.

Smix = k lnW (2.14)

where

W =[

N !n!(N − n)!

](2.15)

Using Stirling’s approximation this becomes

Smix = k[N lnN − n lnn − (N − n) ln(N − n)] (2.16)

As the mole fractions of A and B are given by xa = (N −n)/N and xb = n/N , equation(2.16) reduces to

Smix = −NAk(xa ln xb + xb lnxb) = −R(xa lnxb + xb lnxb) (2.17)

as R = NAk. This then defines the ideal entropy change on mixing. If there are norepulsive or attractive interactions between atoms A and B the solution is called idealand the Gibbs energy of mixing is given by

G = RT (xa lnxb + xb lnxb) (2.18)

For the energy in the presence of the mixed terms one could define


G =Nz

2xa((1 − xb)Ea + xbEab) +

Nz

2xb((1 − xa)Eb + xaEab) (2.19)

with Eab an interaction energy between a and b atoms.This leads to

G =Nz

2[xaEa + xbEb + xaxb(2Eab − Ea − Eb)] (2.20)

So for a binary mixture the term 2Eab −Ea −Eb constitutes the mixing term Gmix anddetermines whether the ab-bonds are preferred or not.(Note that this can be written as

G =Nz

2[xaEa + xbEb + xaxb((Eab − Ea) − (Eb − Eab))] (2.21)

preluding a second derivative term in the Gibbs free energy in case atoms a and bseparated by a straight interface.)For more components, still restricting the interaction to pairs one finds:

G =Nz

2

m∑i

xi((1 −m∑

j =i

xj)Ei +m∑

j =i

xjEij), (2.22)

G =Nz

2

m∑i

(xiEi − xiEi

m∑j =i

xj + xi

m∑j =i

xjEij), (2.23)

G =Nz

2

m∑

i

xiEi −m∑i

(xiEi

m∑j =i

xj − xi

m∑j =i

xjEij)

, (2.24)

G =Nz

2

m∑

i

xiEi +m∑i

(m∑

j =i

xixjEij −m∑

j =i

xixjEi)

. (2.25)

So for a ternary system:

G =Nz

2(xaEa + xbEb + xcEc + xaxb(2Eab − Ea − Eb)

+ xaxc(2Eac − Ea − Ec) + xbxc(2Ebc − Eb − Ec)) (2.26)

In practice the situation is slightly more complex as the mixing terms are found tobe dependent on temperature and composition. One often encounters an approximatedescription of these effects proposed by Muggianu as


Gmix =∑

i

∑j>i

xixj

∑v

Ωvij(xi − xj)v with v = 1 or 2 (2.27)

This leads for example, in the case of a ternary system to the somewhat more elaborateexpression

Gmix = x1x2(Ω012 + Ω1

12(x1 − x2))

+ x2x3(Ω023 + Ω1

23(x2 − x3))

+ x1x3(Ω013 + Ω1

13(x1 − x3)) (2.28)

with the Ωvij as a + bT .

G =∑

i

xiG0i + RT

∑i

xi lnxi +∑

i

∑j>i

xixj

∑v

Ωvij(xi − xj)v (2.29)

where Ωvij is a binary interaction parameter dependent on the value of v. The above

equation for Gmix becomes regular when v = 0 and sub-regular when v = 1. In practicethe value of v does not rise above 2. Equation (2.29) assumes that ternary interactionsare small in comparison to those which arise from binary terms. This may not alwaysbe the case and where there is evidence for higher-order interactions these can be takeninto account by a further term of the type Gijk = xixjxkLijk, where Lijk is an excessternary interaction parameter. Equation (2.29) is normally used in metallic systemsfor substitutional phases such as liquid, FCC, BCC, etc. However, for phases such asinterstitial solutions, ordered intermetallics, ceramic compounds, slags, ionic liquids andaqueous solutions, simple substitutional models are generally not adequate.

The integral Gibbs energy, G0i , of a pure species is given simply by the equation

G0i = H − TS (2.30)

where H and S are the enthalpy and entropy as a function of temperature and pres-sure. Thermodynamic information is usually held in databases using some polynomialfunction for the Gibbs energy which, for the case of the Scientific Group ThermodataEurope (SGTE), is of the form

Gm[T ] − HSERm = a + bT + cT ln(T ) +

n∑2

dnTn (2.31)

The left-hand side of the equation is defined as the Gibbs energy relative to a StandardElement Reference state (SER) where HSER

m is the enthalpy of the element or substancein its defined reference state at 298.15 K, a, b, c and dn are coefficients and n representsa set of integers, typically taking the values 2, 3 and -1. From equation (2.31), further


thermodynamic properties can be obtained. Thermodynamic data for pure elements canbe found in [4], where equation (2.31) is used. Data for pure elements in the form ofequation (2.33) can be found in appendix B.

S = −b − c − c ln(T ) −∑

ndnTn−1 (2.32)

H = a − cT −∑

(n − 1)dnTn (2.33)

Cp = −c −∑

n(n − 1)dnTn−1 (2.34)

Relation to chemical potential

From the definition of the internal energy for a homogeneous system U = TS − pV +∑k µkNk it follows that

G =∑k

µkNk (2.35)

for the molar Gibbs energy

Gm =∑k

µkxk (2.36)

and for

(dGm)p,T =∑k

µkdxk (2.37)

Moreover we find

dG = V dp − SdT +∑k

µkdNk (2.38)

so

(∂G

∂p

)T,Nk

= V ,

(∂G

∂T

)p,Nk

= −S ,

(∂G

∂Nk

)T,p

= µk (2.39)


2.3 Diffusion

Consider a system consisting of two parts of equal T ; one with chemical potential µ1

and mole number N1, the other with chemical potential µ2 and mole number N2. Theflow of particles from one part to the other can be associated with a single number ofeach species involved:

−dN1k = dN2

k = dξk (2.40)

For the entropy production we find, summing over all species

diS =∑k

−(

µ1k − µ2

k

T

)dξk (2.41)

One can generalize to a continuously changing chemical potential along 1 direction bydefining the driving force across an infinitesimal distance dx as ∂µ

∂xdx. Now,

diS(x)dx =∑k

(−∂µk

∂xdx

)dξk

Tand (2.42)

diS(x)dx

dt=

∑k

(−∂µk

∂xdx

)1T

dξk

dt(2.43)

Now introducing Jk for the flux dξkdt and furthermore concentrating on terms per unit

length, we get

diS(x)dt

=∑k

(−∂µk

∂x

)1T

Jk (2.44)

At this point the assumption is made that the fluxes Jk are linearly dependent on thedriving force:

Jk = −Lk1T

∂µk

∂x(2.45)

In terms of concentration gradients:

Jk = −Lk1T

∂µk

∂nk

∂nk

∂x(2.46)

and depending on the sign of ∂µk∂nk

∂nk∂x we have Fick type diffusion or ”uphill” diffusion.

Note that the mechanism of diffusion represented here by the constant Lk does not inany way follow from thermodynamical considerations, and has to be put in, either from


empirical sources or theoretical considerations based on atomistics. For the change inconcentration over time one finds

∂n(x)∂t

= −∂Jk

∂x(2.47)

or

∂n(x)∂t

=Lk

T

∂2µk

∂x2(2.48)

Remembering the relations between G and the chemical potentials, the link betweenthe equilibrium phase diagram, the driving forces in non-equilibrium situations and thefluxes leading to equilibrium is clear. Input from outside thermodynamics is necessaryfor the determination of the reaction/diffusion mechanisms and speeds.

Chapter 3

Plastic Deformation Mechanisms

3.1 Diffusion

In solids, atomic movements are restricted due to bonding to equilibrium positions.However, thermal vibrations occurring in solids do allow some atoms to move. Diffusionof atoms in metals and alloys is particularly important since most solid-state reactionsinvolve atomic movements. There are two main mechanisms of diffusion of atoms ina crystalline lattice, i.e. the vacancy or substitutional mechanism and the interstitialmechanism [5].

Vacancy or substitutional diffusion mechanismAtoms can move in crystal lattices from one atomic site to another if there is sufficientactivation energy present provided by the thermal vibration of the atoms and if there arevacancies or other crystal defect for atoms to move into (point defects). As the tempera-ture of the metal increases, more vacancies are present to enable substitutional diffusionof atoms to take place. This process is called self-diffusion, which activation energy isequal to the sum of the activation energy to form a vacancy and the activation energyto move the vacancy. In general as the melting point of the metal is increased, the ac-tivation energy is also. This relationship exists because the higher-melting-temperaturemetals tend to have stronger bonding energies between their atoms. Diffusion can alsooccur by the vacancy mechanism in solid solution. Atomic size differences and bondingenergy differences between the atoms are factors which affect the diffusion rate.The combination of enthalpy H and entropy S in G explains the fact that the presenceof point defects can be stable inside the metal. Although defects lead to an increaseof the internal energy U (in the surrounding area of the defect, the atoms are out ofposition and every deviation from that position will lead to an increase of the poten-tial energy), this increase of U will be compensated for in the free enthalpy G by thesimultaneous increase of the entropy S (G = H −TS). In total, G will decrease becauseof these defects, so dG < 0; exactly what nature strives for. The expression for thefree enthalpy also clarifies why the solubility of alien atoms in a metal increases withincreasing temperature. With equal increase of entropy (caused by the point defects)

11

Chapter 3. Plastic Deformation Mechanisms 12

the contribution of this to the decrease of G at higher temperature is greater, because Sappears in a product with T . The system can therefore allow more increase of internalenergy without losing equilibrium [6].

Interstitial diffusion mechanismThe interstitial diffusion of atoms in crystal lattices takes place when atoms move fromone interstitial site to another neighboring interstitial site without permanently displac-ing any of the atoms in the matrix crystal lattice. For this mechanism to take place,the size of the diffusing atoms must be relatively small compared to the matrix atoms.Small atoms such as hydrogen, oxygen, nitrogen and carbon can diffuse interstitially insome metallic crystal lattices.

3.2 Deformation mechanics

Plastic flow is a kinetic process. In general, the strength of the solid depends on strain,strain rate and temperature. It is determined by the kinetics of the processes occur-ring on the atomic scale, i.e. the glide-motion of dislocation lines, their coupled glideand climb, the diffusive flow of individual atoms, the relative displacement of grainsby grain boundary sliding (involving diffusion and defect motion in the boundaries),mechanical twinning (by the motion of twinning dislocations) and so forth [7]. Theseare the underlying atomic processes which cause flow. But it is more convenient toto describe polycrystal plasticity in terms of the mechanisms to which the atomisticprocesses contribute. We therefore consider the following deformation mechanisms:

• Collapse at the ideal strengthFlow when the ideal shear strength is exceeded.

• Low-temperature plasticity by dislocation glide(a) Limited by a lattice resistance, (b) limited by discrete obstacles, (c) limitedby phonon or other drags and (d) influenced by adiabatic heating.

• Low-temperature plasticity by twinning.

• Power-law creep by dislocation glide and climb(a) Limited by glide processes, (b) limited by lattice-diffusion controlled climb(high-temperature creep), (c) limited by core-diffusion controlled climb (low-temperaturecreep), (d) power-law breakdown (the transition from climb-plus-glide to glidealone), (e) Harper-Dorn creep and (f) creep accompanied by dynamic recrystalli-sation.

• Diffusional flow(a) Limited by lattice diffusion (Nabarro-Herring creep), (b) limited by grainboundary diffusion (Coble creep) and (c) interface-reaction controlled diffusionalflow.


Power-law creep by climb and glideAt high temperatures, dislocations acquire a new degree of freedom; they can climb aswell as glide (figure 3.1).

Figure 3.1: Power-law creep involving cell-formation by climb. Power-law creep limited by glideprocesses alone is also possible.

If a gliding dislocation is held up by discrete obstacles, e.g. intermetallic compounds(IMCs), a little climb may release it, allowing it to glide to the next set of obstacleswhere the process is repeated. The glide step is responsible for almost all of the strain,although its average velocity is determined by the climb step. The important featurewhich distinguishes this mechanism from others is that the rate-controlling process, atan atomic level, is the diffusive motion of single ions or vacancies to or from the climbingdislocation, rather than the activated glide of the dislocation itself.

Creep deformation kinetics

The steady-state creep deformation kinetics of fine-grained metals and alloys such ascold worked-and-annealed eutectic Sn-Pb alloys, generally exhibit the characteristicbehaviour given in figure 3.2. In such a log-log plot of the shear creep rate vs theapplied stress, four regions can be identified [8].

The first three consist of straight line segments with slopes of about 3, 2 and 3-7,respectively, and the fourth with an increasing slope greater than 10. For the eutecticPb-Sn alloy, regions I and II are grain size dependent, whereas regions III and IV areindependent of phase size, in keeping with the behavior of metals and alloys in general.Steady-state creep is generally expressed by the Weertman-Dorn equation of the form

dγs

dt=

AGb

kT

(b

d

)p (τ

G

)n

D0 exp (−Q/kT ) (3.1)

where dγs/dt is the steady-state strain rate, G the shear modulus, b the Burgers vector,k Bolzmann’s constant, T the absolute temperature, d the grain (phase) size, τ theapplied stress, D0 the frequency factor, Q the activation energy for the deformationprocess, n the stress exponent, p the grain size exponent and A a constant. Their values


Figure 3.2: Schematic of strain rate vs appliedstress showing four stages, each described bythe Weertman-Dorn equation

Figure 3.3: Strain rate vs stress for the stressrelaxation of bulk Sn

and the mechanism(s) which have been proposed for each of the four regions of figure3.2 are listed in table 3.2 for the eutectic Sn-Pb solder [9].

Table 3.1: Parameters of the Weertman-Dorn equation and the proposed rate-controlling mech-anism for the high-temperature plastic deformation of eutectic Sn-Pb solder

region n p ∆H (kJ/mole) A proposed mechanismI 1.7-3 2.3 84 1015 grain boundary slidingII 1.6-2.4 1.6-2.3 57 103-105 grain boundary slidingIII 3-7 0 80-100 1015 power-law creepIV >7 0 80-100 - power-law creep breakdown

The deformation in region III should be attributed to dislocation climb and glide and isalso called matrix creep. Grain boundary sliding results in intergranular failure while theother mechanisms lead to grain deformation. In the case of multi-phase materials withlarge volume fractions of phases, interphase boundary sliding may act similar to grainboundary sliding. As the strain rate is increased or the temperature is lowered, stress-strain behaviour becomes increasingly less dependent on thermally activated processes[10].The distinct stages with constant slope as indicated in figure 3.2 for worked and annealedmicrostructures do not in general occur for reflowed solder joints. Rather, such plotsusually exhibit a continuous curvature with increasing slope, such as in figure 3.3.

Chapter 4

Phase Diagrams

Phase diagrams are a useful way of presenting the various equilibrium phases and theircoexistence, at certain compositions and temperatures. In this chapter, the ternaryphase diagrams of solder alloys B, C and D are reviewed (see table 4.1). Solder typeA is not included here, because this lead-containing alloy was merely added to thisresearch to serve as a reference for shear and creep tests. All of the phase diagrams weretaken from literature and have been calculated using the CALPHAD method, which isshortly discussed in section 4.1.

Table 4.1: The four types of solder alloys investigated and their melting points

type composition melting point (˚C)A SnPb36Ag2 179B SnAg3.8Cu0.7 217C SnAg3.3Bi3.82 210-216D SnZn8Bi3 189-199

4.1 Thermodynamics of phase equilibria

Figure 4.1 shows one of the simplest forms of phase diagrams, a system with a miscibilitygap. It is characterized by a high-temperature, single-phase field of α which separatesinto a two-phase field between α1 and α2 below a critical temperature of 900 K. Thisoccurs because of the repulsive interactions between A and B (note that these A andB are not the same A and B as in table 4.1). While a single minimum exists, the Gibbsenergy of the alloy is always at its lowest as a single phase. However, below 900 K thesystem has further possibilities to lower its Gibbs energy. Figure 4.2 shows the G vs xdiagram at 600 K. If an alloy of composition x0 were single phase it would have a Gibbsenergy G0. However, if it could form a mixture of two phases, one with composition x′

α1

and the other x′α2

, it could lower its total Gibbs energy to G′, where G′ is defined by

15

Chapter 4. Phase Diagrams 16

Figure 4.1: (a) Phase diagram for an A−B system showing a miscibility gap and (b) respectiveG vs x curves at various temperatures

G′ =x0 − x′

α1

x′α2

− x′α1

G′α2

+x′

α2− x0

x′α2

− x′α1

G′α1

(4.1)

and G′α1

and G′α2

are respectively the Gibbs energies of α at composition x′α1

and x′α2

.The equation is formed using the lever-rule [5]. A further separation to compositionsx′′

α1and x′′

α2sees a further reduction of the Gibbs energy of the two-phase mixture to

G′′.

This process can continue but is limited to a critical point where the compositionscorrespond to xE

α1and xE

α2where any further fluctuation in the compositions causes

the Gibbs energy to rise. This point is then a critical point and the phases α1 and α2

with compositions xEα1

and xEα2

respectively are defined as being in equilibrium witheach other. At this point it is convenient to define the fraction of each phase using theequations

Nα1 =xE

α2− x0

xEα2

− xEα1

and Nα2 =x0 − xE

α1

xEα2

− xEα1

(4.2)

where Nα1 and Nα2 are the number of moles of α1 and α2 respectively. The criticalposition can be defined as follows. The system A−B with composition x0 has reachedan equilibrium where its Gibbs energy is at a minimum. Equation (4.1) can then beused in combination with a Newton-Raphson technique to perform a Gibbs energyminimalisation with respect to the composition of either A or B. The calculation of thephase diagram is then achieved by calculating phase equilibria at various temperaturesbelow 900 K and plotting the phase boundaries for each temperature.


Figure 4.2: G vs x diagram at 600 K for an A − B system shown in figure 4.1(a) showingseparation of a single-phase structure into a mixture of two phases

4.2 The Sn-Ag-Cu System

The thermodynamic assessment of the ternary Sn-Ag-Cu system is based on the binarysystems Ag-Cu, Sn-Cu and Sn-Ag as shown in appendices D, E and F, respectively [11].The liquidus projection of the ternary system is shown in figure 4.3. Concerning thesubstitute material for the Pb-Sn eutectic solder alloy whose melting temperature is183˚C, one should pay attention to the Sn-rich corner of the liquidus surface diagram[12]. Figure 4.4 shows the partial liquidus surface projection focused on the ternaryeutectic reaction in table 4.2 1.

Table 4.2: Eutectic equilibrium of the Sn-Ag-Cu system

reaction phase mass % Ag mass % Cu mass % SnL 3.73 0.83 95.42

L Ag3Sn + Cu6Sn5 + (Sn) Ag3Sn 73.17 0 26.83(216.9˚C) Cu6Sn5 0 39.07 60.93

(Sn) 0.07 0 99.93

4.3 The Sn-Bi-Ag system

Another group of candidate alloys for lead-free solder alternatives contains Sn, Bi andAg. The phase diagram provides a basic road map for the evaluation of propertiesof alloys in this system. Unfortunately, an experimentally established ternary phase

1values are taken from http://www.metallurgy.nist.gov/phase/solder/solder.html


Figure 4.3: Liquidus projection of the Sn-Ag-Cu system

Figure 4.4: Liquidus surface in the Sn-richcorner

diagram appears to be unknown. Nevertheless, the three binary systems Sn-Ag, Ag-Biand Sn-Bi are well documented and can be found in appendices F, G and H, respectively.The Sn-Bi and Ag-Bi systems are simple eutectic systems with limited solubilities intheir terminal solid solutions [13]. The Ag-Sn system has two intermetallic phases,both of which form peritectically, and also a eutectic equilibrium. When the binarysystems are known, thermodynamic calculation provides an extremely useful tool toobtain quantitative information about the ternary system. From the obtained ternaryphase diagram, the melting temperature range for solders and the solidification path,as well as the susceptibility to intermetallic formation with various substrates can befound.

Extrapolation to the ternary system

In order to extrapolate the binary systems to a ternary system, some assumptions weremade. The phases L, (Sn) and (Ag) were treated in all three binary systems as solutionphases. For the ternary extrapolation of the Gibbs energies of these phases, no ternaryinteractions are used for the solution phases. For the (Bi) phase, solubility was onlyconsidered for Sn and, therefore, (Bi) is treated as a semi-stoichiometric compound withno solubility for Ag. Negligible solubility of Bi in (ζAg) and Ag3Sn was assumed. Thus,(ζAg) is treated as semi-stoichiometric compound and Ag3Sn is treated as an ordinarystoichiometric compound in the calculation of the ternary system. The temperaturesand compositions of the phases for the invariant reaction are given in Table 4.3.

The predicted liquidus projection is shown in figure 4.5. Figure 4.6 shows the liquidussurface in the Sn-rich corner.


Table 4.3: Eutectic equilibrium of the Sn-Ag-Bi system

reaction phase mass % Ag mass % Bi mass % SnL 0.68 55.85 43.47

L Ag3Sn + (Sn) + (Bi) Ag3Sn 73.17 0 26.83(137.1˚C) (Bi) 0 99.89 0.11

(Sn) 0.03 20.60 79.37

Figure 4.5: Liquidus projection of the Sn-Ag-Bi system

Figure 4.6: Liquidus surface in the Sn-richcorner

4.4 The Sn-Zn-Bi system

Thermodynamics

Due to a lack of data on thermodynamic properties of terminal solid solutions as wellas on their phase boundaries, only binary contributions were utilized, and the ternaryterms were ignored. In the case of the liquid phase, the availability of experimental dataallows the inclusion of in principle a ternary term. The Gibbs free energy of the liquidphase is described using the following equation:

GLm =

∑i=Bi,Sn,Zn

0GLi xL

i + RT∑

i=Bi,Sn,Zn

xLi lnxL

i

+ xLBi xL

Sn LLBi,Sn + xL

Bi xLZn LL

Bi,Zn

+ xLSn xL

Zn LLSn,Zn + ∆exGtern (4.3)

where


LLi,j =

n∑m=0

mLLi,j · (xi − xj)m (4.4)

∆exGtern = xLBi x

LSn xL

Zn(xLBi

0LLBiSnZn + xL

Sn1LL

BiSnZn + xLZn

2LLBiSnZn) (4.5)

and where the coefficient mLLi,j is the parameter in the sub-binary system and nLL

BiSnZn

may be temperature dependent. SGTE data for a selection of elements can be found inappendix A. Additional SGTE data for Bi, Sn and Zn can be found in appendices Band C. Binary phase diagrams of the Sn-Bi, Zn-Bi and Sn-Zn systems can be found inappendices H, I and J, respectively.Figure 4.8 shows calculated isothermal sections at 135, 170 and 250˚C, respectively.The characteristic features of the phase diagram in this ternary system are (1) thesolubilities of Bi and Sn in (Zn) are negligibly small, (2) the solubility of Zn in (Bi) alsosmall, and (3) (Zn) directly equilibrates with liquid or (Bi) and (Sn). Figure 4.7 showsthe projection of the liquidus surface in the Sn-rich corner.

Figure 4.7: Liquidus projection in the Sn-rich corner

In the Sn-Zn-Bi system, no experimental data for the isothermal sections and verylimited information for the liquidus surface are available. Only a calculated liquidussurface in the entire composition range was presented by Pelton et al. (1977), where thecalculated eutectic reaction occurs at about 137˚C and 2.5 at.% Zn and 54 at.% Sn. Inthe calculation of Malakhov et al. [14], the eutectic reaction takes place at T = 130˚Cwith compositions of four equilibrium phases given in table 4.4.


Figure 4.8: Calculated isothermal sections at (a) 135˚C, (b) 170˚C and (c) 250˚C

Table 4.4: Eutectic equilibrium of the Sn-Zn-Bi system

reaction phase Bi (wt.%) Sn (wt.%) Zn (wt.%)liquid 54.54 42.75 2.71

L (Bi) + (Sn) + (Zn) (Bi) 97.75 1.99 0.26130˚C (Sn) 22.37 77.35 0.28

(Zn) 1.93 · 10−3 7.68 · 10−4 99.997

Chapter 5

Experiments

Different experiments have been done with the four solder alloys. This chapter willfirst describe what kind of sample has been used and how it should be polished inorder to expose the solder’s microstructure. Next, a concise description of the digitalimage correlation technique used to measure strain fields will be presented, followed bya description of the coarsening and creep test setup.

5.1 Shear tests

When an electronic device is in operation, the solder connections are subjected to me-chanical stresses and strains. The primary cause of these stresses and strains arise fromthe fact that the electronic component and the board have different coefficients of ther-mal expansion, as shown in figure 1.1. If the temperature of the system rises, the boardexpands more than the component, resulting in the solder connection being subjectedto shear stress. In order to investigate the shear strains caused by these shear stressesin more detail, it is necessary to perform shear tests on the four different kinds of solderalloys. A special shear sample has been designed for this purpose, which is shown isfigure 5.1.

Figure 5.1: Shear sample

The copper plates are 25 mm long, 9 mm wide and 1 mm thick and are soldered togetherby a 10 mm long and 0.5 mm wide solder strip. These measurements were chosen to make

22

Chapter 5. Experiments 23

it suitable for the miniature tensile stage. The tensile stage can be mounted under theoptical microscope as well as in the ESEM, so pictures can be taken during deformation.These pictures can then be stored and analyzed later.

The samples were used for three important reasons. First of all, important data aboutthe shear properties of the different kinds of alloys can be obtained. Stress-strain curvescan be generated, which give valuable information on the maximum elastic and plasticstrains. Another reason is the investigation of the deformation behaviour of the solderalloys; using DIC techniques (see section 5.2), strain fields inside the solder strip can bevisualized. And finally, microstructure coarsening under thermo-mechanical loading, aswell as creep, can be observed using a special test setup which is described in the nextsection.

Polishing

When trying to expose and visualize a metal’s microstructure, grinding and polishingis needed. For this purpose, a grinding table and a polishing machine (Struers DAP-U with a Struers Pedemin 2 specimen mover) have been used. Because of the smallthickness of the samples (1 mm), it was impossible to press them onto the grindingpaper or polishing cloth manually. Therefore, the samples had to be glued onto smallsteel cylinders which fit in the three holes of the polishing machine’s specimen holderplate. The glue used is based on cyanoacrylate and can be dissolved by immersing it inacetone during approximately half a day (depending on the amount of glue used).The polishing method is based on the Metalog Guide from Struers [15] for a type Ametal (very soft). Some adjustments have been made to optimize the polishing resultsand at the same time reduce the total preparation time and reduce wear of the polishingcloths. An overview of the correct preparation method for solder alloys is given belowin table 5.1. All grinding and polishing steps are done with a rotation speed of 150 rpmand the given force is the force per sample. After each step, the samples must be cleanedin an ultrasonic bath for 1-2 minutes to remove any remaining abrasive particles.

Table 5.1: Sample preparation method

cloth force (N) time (min.) lubricant abrasive sizeplane grinding SiC paper 25 until plane water grit 320fine grinding MD Largo 30 4 blue 9 µm

MD Dur 25 3 blue 6 µmpolishing MD Mol 20 2 red 3 µm

MD Chem 10 1 OP-S -


5.2 Digital Image Correlation

Introduction

Digital Image Correlation (DIC) is a non-contact technique, used to measure strainfields on specimens being subjected to deformation. The technique compares two digitalimages, one taken before deformation and one taken after. The digital images are madeup of a rectangular array of grey pixels. These pixels are assigned an 8 bit value, sothat they can assume a grey level from 0 (white) to 255 (black). The grey level of thepixel represents the light intensity, usually received from a camera. Computer software(Aramis) uses correlation algorithms and displacement fields to match regions of the twoimages to each other. From the differences in positions of the features on the specimenthe strain can then be calculated. This technique will be particularly useful in analyzingthe inhomogeneous strain that can occur at or below the micron level [16].

Aramis

Mapping of the facetsThe software initially places a square grid of equally spaced points over the sourceimage. Centered on each of these points is one square facet (figure 5.2).

Figure 5.2: Source grid and facets mapped to destination grid and facets

The software tries to match the facet from its position on the source image to its positionon the destination image. The software must find the source facet’s deformation beforeit can find the best position of the match on the destination image. Further analysis ofthe algorithm is beyond the scope of this report and can be found in [16].


Ideal imagesThe best images for the software to work with contain small, finely distributed featureswith a high contrast. The software finds it easier to match this type of image as thesurfaces produced by the grey levels are identifiable. A ’speckled’ image can also beprovided artificially on low contrast samples, as described in section 5.3. The softwarecan then more easily map corresponding points from source to destination image.

5.3 Contrast

In order to gain more insight in the deformation behaviour of the solder alloys, DIC tech-niques can be used to visualize the strain field. However, the DIC software Aramis needshigh-contrast images to be able to calculate the strain field accurately. Unfortunately,the solder alloys didn’t show sufficient contrast, so this had to be created artificially.To achieve this on the shear samples, microsieves from Aquamarijn (a spin-off companyfrom the University of Twente 1) have been used. A sketch of these microsieves is shownin figure 5.3.

Figure 5.3: Microsieve

The black area is the silicon support for the nine membranes with a thickness of 1 µm(grey). Inside these membranes are areas of 500×500 µm (white) with circular holesof 5 µm diameter and 10 µm apart. If this microsieve is placed on the solder joint ofthe shear sample and gold is sputtered on the sample, the gold will pass through thesmall holes onto the solder. Before doing this, the sample was painted dark grey, thuscreating a good contrast between the dark paint and light gold speckles. The microsievehad to be pressed onto the microsieve, without braking it (the silicon support is very

1http://www.microsieve.com


brittle). The design for a clamping device is shown in figure 5.4 and a close-up of thebottom side of the microsieve holder is shown in figure 5.5. The rectangular pocket inthe microsieve holder is 0.3 mm deep, while the microsieve itself is 0.5 mm thick. Thismeans that the microsieve will stick out 0.2 mm, thus ensuring a good contact with theshear sample. To make the handling a bit more easy, the microsieves were ’glued’ in themicrosieve holder, using a small amount of silicon gel on the edges.

5.4 Creep and coarsening

Microstructure size has a great influence on the plastic deformation kinetics of solderjoints. Since room temperature for eutectic Sn-Pb alloys is already 0.65 Tm (Tm is themelting temperature in K), phase coarsening by diffusion can be expected at highertemperatures. When subjected to additional mechanical loading, this process will beaccelerated [1][17]. Although coarsening has only been investigated in binary solderslike Sn-Pb, this research will also examine coarsening in ternary alloys, if present.Creep can be expected as well, which will ultimately lead to failure of the solder joint.To investigate these two phenomena, the test setup shown in figure 5.6 was used. Thedevice can be hung on a special rack and placed in a stove. Weights can be attached tothe bottom hook and because of the symmetrical design, the force will always be verticaland no bending moment will act on the samples. This also speeds up the testing, becausetwo samples can be used at the same time.The samples were taken out of the furnace briefly every day and pictures were taken.Because just half of the solder strip was painted grey, the evolution of the microstructurecould still be examined. The samples used for this experiment were also provided witha gold grid, this time to measure creep. The coarsening/creep tests were performed attwo temperatures and two loads.

The results of the experiments described in this chapter will be presented in Chapter 6.


Figure 5.4: Microsieve clamping device Figure 5.5: Microsieve holder

Figure 5.6: Creep test setup

Chapter 6

Experimental Results

In order to gain better insight in the deformation behaviour of the proposed lead-freesolder alloys, one must look at the microstructural evolution of the solder joint at dif-ferent scales and different modes. First of all, shear tests were performed to determinethe solder’s shear strength, maximum shear strain and crack formation. One level down,DIC-measurements were done on all types of alloys to produce strain field images ofthe deforming solder joint. Furthermore, creep tests have been performed at two tem-peratures and two loads, in order to determine creep rates for all solder types (table6.1). Creep experiments have started with solders B and C at 100˚C. Due to a verylow creep rate, experiments with solder D have been done using higher loads. SolderA showed a very high creep rate, so the load for these experiments was lowered. Shearstrain was measured with the gold grids as well, so the samples had to be taken out ofthe stove briefly every day to take pictures. These creep samples were then analyzed inthe ESEM to determine the failure mechanism, so this could be compared to the resultsof the normal shear tests and DIC-measurements.

Table 6.1: Creep schedule

100˚C 150˚C1 MPa AJ1 AJ2 -

1.5 MPA - A10 A132 MPa AJ3 AJ4 -3 MPa - A11 A14

1.5 MPa C5 C7 C14 C162.5 MPa - C12 C133 MPa C8 C10 -

1.5 MPa B4 B9 BJ1 BJ22.5 MPa - B18 B193 MPa B7 B14 -

1.5 MPa - D9 D162 MPa D10 D13 -3 MPa D6 D11 D14 D15

28

Chapter 6. Experimental Results 29

6.1 Microstructures

The microstructures of the four types of solder alloys in undeformed state have been ex-amined in order to better understand differences that occur in these microstructures inthe case of thermo-mechanical loading. Photographs taken with an optical microscopeare shown in figures 6.1 through 6.4, respectively.

Solder A shows a eutectic structure; the dark areas are Pb, the light phase is Sn. Silverin the form of Ag3Sn can also be observed, only not clearly in figure 6.1. Figure K.15in appendix K shows this more clearly. The needle shaped structure on the left is theintermetallic compound (IMC) Cu6Sn5.Solder B clearly shows two hexagonal Cu6Sn5 particles. Here, the black areas are Snwith Ag3Sn and Sn (eutectic) in between.

Figure 6.1: Microstructure of solder A (brightfield)

Figure 6.2: Microstructure of solder B (darkfield)

Solder C again shows Cu6Sn5 particles (H-shaped). The black phase is Sn with a eutecticmixture of Ag3Sn and Sn in between.Solder D clearly shows the bright lamellar structures of Zn in a black Sn matrix withBi dissolved in the Sn.EDX-analysis of all solder types has been done as well to determine the presence andlocation of detected elements. Pictures can be seen in appendix K.


Figure 6.3: Microstructure of solder C (darkfield)

Figure 6.4: Microstructure of solder D (darkfield)

During soldering, reactions occur between solders and conductors and IMCs may nu-cleate and grow at the solder/conductor interface. The presence of these IMCs is anindication of good metallurgical bonding and a thin and continuous layer is an essentialrequirement for good wetting and bonding. It also produces distinct improvements inmechanical properties of joint. However, due to their brittle nature, too thick of an IMClayer at the solder/conductor interface may degrade the reliability of the solder joint[18]. In the case of alloys A, B and C, Cu6Sn5 will form at the interface and in the caseof alloy D, CuZn (γ brass) will form. This can also be seen in appendix K.


6.2 Solder A (SnPb36Ag2)

Shear data for solder joints made of this ”classic” lead-containing alloy is shown in figure6.5 and is corrected for the stretching of the 690N loadcell [19]. It is obvious that thismaterial shows a low maximum shear strength of about 33 MPa and has a relativelylow maximum shear strain. It should be noted that this data represents the solder jointand not the material itself. As can also be seen from the DIC-measurements, strainconcentrates at the intermetallics at the interfaces, so it’s the combination of solderalloy and interface that determines the mechanical properties of the joint.

0 0.5 1 1.50

5

10

15

20

25

30

35

40

Shear strain

She

ar s

tres

s [M

Pa]

Shear strain vs shear stress for solder A

Figure 6.5: Shear strain vs shear stress curves for type A solders

Post failure analysis of the shear samples revealed that almost all samples showed crackinitiation along both interfaces (figure 6.6).

Figure 6.6: Crack initiation along both interfaces


Strain field images (figures 6.7 through 6.10) confirm crack formation along the interface.High shear strains can be seen in this region, whereas deformation in the rest of thesolder joint is relatively low. The four markers on the x-axis show the macroscopicstrains where the four DIC pictures were taken.

Figure 6.7: Deformation at 19% shear strain Figure 6.8: Deformation at 26% shear strain


Creep tests for this solder were performed to serve as a reference for the other soldertypes. However, this material showed very poor resistance to creep, so many samplesfailed before pictures could be taken for strain measurements. To provide an indicationof the creep rates, table 6.2 shows the time-to-failure for a number of samples. Aninteresting phenomenon, is that failure of the creep samples occurred completely alongone interface, so no crack initiation along both interfaces. So, the cracks along the grainor colony boundaries inside the solder joint didn’t seem to have a major influence on


final failure.

Table 6.2: Time-to-failure for solder A samples (hours)

100˚C (0.83Tm) 150˚C (0.93Tm)1.5 MPA - A10 A13 ∼16h2 MPa AJ3 AJ4 ∼48h -3 MPa A16 A19 <24h A11 A14 <16h

ESEM-analysis has been performed on the creep samples to check the deformationmechanism. Figure 6.11 shows a BSE-image of a deformed solder joint. Crack formationcan be seen clearly, as well as large deformations along the interface. Figure 6.12 shows aclose-up of the same sample, where can be seen that cracks form along colony boundaries.Failure however occurred along the interface for every creep sample.

Figure 6.11: Sample AJ1, crack formationalong colony boundaries and interface (BSEdiode A-B image)

Figure 6.12: Sample AJ1, close-up of crackformation along colony boundaries (normalBSE image)


6.3 Solder B (SnAg3.8Cu0.7)

Shear data for solder B joints (also called SAC) is presented in figure 6.13 and the datahas been corrected for the stretching of the loadcell. As can be seen, the shear strengthof this alloy is comparable to that of solder A. There is one big difference though, andthat is the relatively high maximum shear strain. So the rather low maximum shearstrength of this alloy is compensated by its enormous plastic shear deformation, whichcan take on values of more than 2.

0 0.5 1 1.5 2 2.5 30

5

10

15

20

25

30

35

40

Shear strain

She

ar s

tres

s [M

Pa]

Shear strain vs shear stress for solder B

Figure 6.13: Shear strain vs shear stress curves for type B solders

When looking at the failure mode and crack formation during normal shear testing, wecan observe crack formation at the interface region, like we’ve also seen in solder A. Thiscan also be illustrated by looking at the strain field images of the deforming solder joint(figures 6.14 through 6.17, respectively). It is obvious that shear deformation localizesat the interface(s), ultimately leading to failure in this region. Like with solder A, cracksalso initiated from both sides interface which could be observed for all samples. Thefour markers on the x-axis of figure 6.13 show the macroscopic strains where the fourDIC pictures were taken.




Creep tests were performed according to table 6.1 and show a very high resistance tocreep, as displayed in figure 6.18. Figure 6.19 shows a log creep rate vs log shear stressplot for SAC, taken from literature to serve as a reference [20]. It can be seen thatthe creep rate from own experiments is higher than which can be seen in figure 6.19.ESEM-pictures have been taken to investigate the microstructural evolution of the alloywhen undergoing constant thermo-mechanical shear loading.


100

101

10−10

10−9

10−8

10−7

10−6

10−5

10−4

log shear stress [MPa]

log

shea

r st

rain

rat

e [1

/s]

Average creep rate vs shear stress (solder B)

100 C (0.76Tm)150 C (0.86Tm)

Figure 6.18: Log creep rate vs log shear stressfor solder B (own results)

Figure 6.19: Log creep rate vs log stress forsolder B, taken from literature

Figures 6.20 and 6.21 show an optical microscope photograph and an ESEM-photographof creep-induced crack formation along colony boundaries. Note the cracks formingparallel to the interface as well as perpedicular to the interface and the deformationalong the bottom interface in figure 6.21. The cracks parallel to the interface ultimatelylead to failure of the solder joint, whereas failure during normal shear tests occurs alongthe interfaces.

Figure 6.20: Sample B19, crack formationalong colony boundaries (dark field)

Figure 6.21: Sample B19, crack formationalong colony boundaries (BSE diode A-B)


To complete the microstructural analysis, an EDX-map has been made to determinethe presence and location of the alloy’s elements. Figures 6.22 through 6.25 show thediffusion of copper into the bulk material and the finely dispersed silver particles, mostof which is actually Ag3Sn. Copper and tin can be found in the interface; this is theintermetallic compound Cu6Sn5.

Figure 6.22: SE-image Figure 6.23: Silver

Figure 6.24: Copper Figure 6.25: Tin


6.4 Solder C (SnAg3.3Bi3.82)

Shear data for solder C joints is plotted in figure 6.26, where the data has been correctedfor the stretching of the loadcell. The first feature that can be noticed is the very highshear strength, unlike any other solder type investigated. A disadvantage however, isthat this material fractures in a brittle way. This could be observed for all samples,except for sample C4 (curve with the highest maximum shear strain). Not all samplesthat were investigated are shown in this plot, because of a problem with the displace-ment data. These samples did however show brittle fracture and about the same shearstrength. Failure occurred along one interface as well as both interfaces (crossing overof the crack, like solders A and B) for some samples.

0 0.5 1 1.50

10

20

30

40

50

60

70

Shear strain

She

ar s

tres

s [M

Pa]

Shear strain vs shear stress for solder C

Figure 6.26: Shear strain vs shear stress curves for type C solders

The strain field images clearly show a different pattern in the shear strain than thosefor solders A and B. As can be seen in figures 6.27 through 6.30, deformation is muchmore homogeneous and no large shear strains along the interface can be seen, exceptfor at the very last moment (figure 6.30 was taken 4 seconds prior to fracture). The fourmarkers on the x-axis show the macroscopic strains where the four DIC pictures weretaken.




As for solder B, this solder also shows good resistance to creep, as can be seen in figure6.31. Creep tests were performed according to table 6.1. Closer examination of thesamples using light microscopy and the ESEM reveals that this material is also subjectto colony boundary sliding during constant thermo-mechanical loading.


100

101

10−10

10−9

10−8

10−7

10−6

10−5

10−4


log

shea

r st

rain

rat

e [1

/s]

Average creep rate vs shear stress (solder C)

100 C (0.77Tm)150 C (0.87Tm)

Figure 6.31: log creep rate vs log shear stress for solder C

Figure 6.32 shows cracks forming along colony boundaries. A remarkable feature of thisalloy is, that failure occurred along the interface and not along the colony boundarycracks. Figure 6.33 shows a close-up of the interface, where sliding along the intermetalliclayer (Cu6Sn5) can be seen.

Figure 6.32: Sample C10, crack formationalong colony boundaries (BSE diode A-B)

Figure 6.33: Sample C10, sliding of grainsalong intermetallics at the interface


Again, an EDX-map has been made to show the colonies in a different way. Figures6.34 through 6.37 show the distribution of copper, silver and tin. Bismuth was leftout, because this was finely dispersed throughout the whole region and no particularconcentrations could be found anywhere. It can be seen from the silver map that theorientation of the grains is different for both sides of the small crack, thus indicating acolony boundary.

Figure 6.34: SE-image Figure 6.35: Silver

Figure 6.36: Copper Figure 6.37: Tin


6.5 Solder D (SnZn8Bi3)

Shear data for solder D joints can be seen in figure 6.38. Again, data has been cor-rected for the stretching of the loadcell. As can be observed, this material shows greatinconsistency in the shear strength. This can however be explained by looking at themicrostructure in the course of time. The appearance of air bubbles could be detectedafter a period of time. Immediately after polishing, no air bubbles could be detected,but the ”older” the samples got, the greater the amount of holes became. The two shearstress-shear strain curves with the lowest shear strength were the oldest and showed alarge amount of holes, leading to weakening of the solder joint. The two curves with thehighest shear strength were the youngest. This probably has something to do with theeasy oxidation of zinc.

0 0.5 1 1.50

10

20

30

40

50

60

Shear strain

She

ar s

tres

s [M

Pa]

Shear strain vs shear stress for solder D

Figure 6.38: Shear strain vs shear stress curves for type D solders

When we look at the shear deformation of this alloy, figures 6.39 through 6.42 showcrack formation along the interface, but also through the bulk material. However, DICmeasurements of old samples (with air bubbles) showed crack formation along the in-terface, where most of the holes appeared. The four markers on the x-axis show themacroscopic strains where the four DIC pictures were taken.




Creep data was collected again according to table 6.1. It can be seen that this alloy hasa higher creep rate than solder types B and C, which makes this alloy, in combinationwith the formation of air bubbles, more unreliable.


100

101

10−10

10−9

10−8

10−7

10−6

10−5

10−4


log

shea

r st

rain

rat

e [1

/s]

Average creep rate vs shear stress (solder D)

100 degrees150 degrees

Figure 6.43: log creep rate vs log shear stress for solder D

Again, light microscopy and ESEM microscopy were utilized to examine the microstruc-tural evolution of this material. Figures 6.44 and 6.45 very clearly show crack formationalong colony boundaries. It can be seen that colonies are differently shaped and smallerthan the colonies that could be seen in solder types B and C. This ultimately leads tointergranular failure, rather than failure along the interface. This could also be observedfrom the normal shear experiments and DIC measurements.

Figure 6.44: Sample D11, crack formationalong colony boundaries, (bright field)

Figure 6.45: Sample D11, crack formationalong colony boundaries (BSE diode A-B)


An interesting feature of the thermal aging of this solder type is the vast diffusion ofcopper to the surface of the joint, forming brass with zinc. EDX-analysis of sample D16clearly shows this phenomenon (figures 6.46 through 6.49).

Figure 6.46: SE-image Figure 6.47: Tin

Figure 6.48: Zinc Figure 6.49: Copper

Chapter 7

Conclusions andRecommendations

All three lead-free solders have shown to have advantages over SnPb36Ag2, the mostcommon solder alloy used in the past. DIC showed that intermetallic compounds at thesolder-substrate interface play a crucial role in solder joint deformation for SnAg3.8Cu0.7(SAC) and SnZn8Bi3, where inhomogeneous and highly local strains could be seen atthe interface. SnAg3.3Bi3.82 showed a much more homogeneous deformation leading tosuperior shear strength, but brittle fracture as well. SnAg3.8Cu0.7 showed the largestplastic deformation, whereas SnZn8Bi3 was subject to aging effects. The older the sam-ples got, the more air bubbles could be detected inside the solder joint, which havedetrimental effects on shear strength. A summary of the shear test results can be seenin table 7.1.

Table 7.1: Summary of shear test results

Solder type A B C Dτmax [MPa] 32-35 32-37 66-67 58→42

γmax 1-1.3 2-2.5 0.7 0.6→1failure ductile ductile brittle ductile at first,

fracture fracture fracture brittle after ageing

The DIC measurements done with the gold grid sputtered onto the solder joints givesvery good results. Aramis had no problems to correlate subsequent facets with eachother. This method will therefore be suitable for other applications, where strain fieldshave to be calculated as well.

It appeared that SnZn8Bi3 has the highest creep rate of the three lead-free solder alloys,followed by SnAg3.3Bi3.82 as can be seen in figure 7.1. SAC is the most creep-resistantalloy with about the same creep rate as SnAg3.3Bi3.82 at 100˚C, but better resistanceto creep at 150˚C. SnPb36Ag2 showed such poor creep resistance that most of thesamples failed before creep could be measured, so the superiority of the lead-free solders

46

Chapter 7. Conclusions and Recommendations 47

100

101

10−10

10−9

10−8

10−7

10−6

10−5

10−4

Average creep rate vs shear stress for all solder types


log

shea

r st

rain

rat

e [1

/s]

solder B, 0.76Tm solder B, 0.86Tm solder C, 0.77Tm solder C, 0.87Tm solder D, 0.8Tm solder D, 0.9Tm reference line n=3

Figure 7.1: Creep data for all solder alloys

with regard to creep is evident. Unfortunately, coarsening could not be observed, so nodata on this subject could be collected.

Furthermore, one important general conclusion that can be drawn from the picturestaken with the optical microscope and ESEM, is that failure occurs along colony bound-aries. Each type of solder alloy investigated showed this behaviour, as can be seen bythe pictures presented in chapter 6. If possible, formation of colonies should be avoidedin the reflow process, as these will have a detrimental effect on solder joint reliabilityat constant thermo-mechanical loading. In SAC solder joints, cracks parallel to the in-terface and perpendicular to these cracks (from one interface to the other), along largecolonies could be seen and failure occurred along cracks parallel to to interface through-out the bulk material. SnAg3.3Bi3.82 also showed horizontal and vertical cracks alonglarge colony boundaries, but didn’t fail along these cracks. Failure did occur along theintermetallics at the interface. SnZn8Bi3 showed the same phenomenon as the othertwo lead-free solder alloys: colony boundary sliding. However, colonies in this mate-rial are much smaller and differently shaped, so cracks could be seen throughout thewhole solder joint in any direction. A more capricious failure occurred along these cracksthroughout the bulk material and not along the interface. Also, EDX-analysis showedlarge diffusion of copper to the surface of the solder joint, forming brass with zinc.

The three main objectives, as described in the introduction, were all met. Baselinethermodynamic data from literature for further numerical modeling has been given,as well as basic theory on equilibrium and non-equilibrium thermodynamic principles.Futhermore, tools have been developed to measure the shear, creep and deformationproperties of the different kinds of solder alloys and finally, data on these properties hasbeen collected to be used for the RIPOSTE project. Combining the results presented inthis thesis with results from other research areas within this project will lead to a goodreplacement for current lead-containing solders.

Bibliography

[1] W. Dreyer and W.H. Muller. Modeling Diffusional Coarsening in Eutectic Tin/LeadSolders: a Quantitative Approach. International Journal of Solids and Structures,38, 2001.

[2] M. Abtew and G. Selvaduray. Lead-Free Solders in Microelectronics. MaterialsScience and Engineering, 27, 2000.

[3] N. Saunders and A.P. Miodownik. CALPHAD, A Comprehensive Guide. PermagonMaterials Series. Elsevier Science Ltd., 1998.

[4] A.T. Dinsdale. SGTE Data for Pure Elements. CALPHAD, 15, 1991.

[5] W.F. Smith. Foundations of Materials Science and Engineering. Engineering Me-chanics Series. McGraw-Hill, Inc., 1993.

[6] D. Landheer. Metaalkunde. Lecture notes, august 2000.

[7] H.J. Frost and M.F. Ashby. Deformation Mechanism Maps. Permagon Press, 1982.

[8] H. Conrad, Z. Guo, Y. Fahmy, and D. Yang. Influence of Microstructure Size on thePlastic Deformation Kinetics, Fatigue Crack Growth Rate and Low-Cycle Fatigueof Solder Joints. Journal of Electronic Materials, 28(9), 1999.

[9] P.L. Hacke, A.F. Sprecher, and H. Conrad. Microstructure Coarsening DuringThermo-Mechanical Fatigue of Pb-Sn Solder Joints. Journal of Electronic Materi-als, 26(7), 1997.

[10] H. Mavoori, J. Chin, S. Vaynman, B. Moran, L. Keer, and M. Fine. Creep, StressRelaxation and Plastic Deformation in Sn-Ag and Sn-Zn Eutectic Solders. Journalof Electronic Materials, 26(7), 1997.

[11] K.W. Moon, W.J. Boettinger, U.R. Kattner, F.S. Biancaniello, and C.A. Handw-erker. Experimental and Thermodynamic Assessment of Sn-Ag-Cu Solder Joints.Journal of Electronic Materials, 29(10), 2000.

[12] I. Ohnuma, M. Miyashita, K. Anzai, X.J. Liu, H. Ohtani, R. Kainuma, andK. Ishida. Phase Equilibria and the Related Properties of Sn-Ag-Cu Based Pb-Free Solder Alloys. Journal of Electronic Materials, 29(10), 2000.

48

BIBLIOGRAPHY 49

[13] U.R. Kattner and W.J. Boettinger. On the Sn-Bi-Ag Ternary Phase Diagram.Journal of Electronic Materials, 23(7), 1994.

[14] D.V. Malakhov, X.J. Liu, I. Ohnuma, and K. Ishida. Thermodynamic Calculationof Phase Equilibria of the Bi-Sn-Zn System. Journal of Phase Equilibria, 21(6),2000.

[15] Struers. Metalog Guide, 2000.

[16] D. Luff. An Analysis of Digital Image Correlation Applied to Scanning ElectronMicroscope Images. Technical Report MT00.020, TU/e, 2000.

[17] L. Li and W.H. Muller. Computer Modeling of the Coarsening Process in Tin-LeadSolders. Computational Materials Science, 21, 2001.

[18] K. Zeng and J.K. Kivilahti. Use of Multicomponent Phase Diagrams for PredictingPhase Evolution in Solder/Conductor Systems. Journal of Electronic Materials,30(1), 2001.

[19] S.H.A. Boers. Private communication.

[20] A. Schubert, H. Walter, A. Gollhardt, and B. Michel. Materials Mechanics andReliability Issues of Lead-free Solder Interconnects. 2001.

Appendix A

Pure element transition data

Values below are taken from [4] and are in J/mole.

H298 − H0 S298 Ttrans ∆transH ∆transS ∆transCp

Ag FCC A1 5745 42.55 1234.93 11296.8 9.1477 1.5402Bi RHO A7 6426.624 56.735 544.55 11296.8 20.7479 0.6521Cu FCC A1 5004 33.15 1357.77 13263.28 9.7684 1.5391Pb FCC A1 6870 64.8 600.612 4773.94 7.9485 1.1867Sn BCT A5 6323 51.18 505.078 7029.12 13.9169 -1.0252Zn HCP A3 5657 41.63 692.68 7322 10.5706 1.6653

50

Appendix B

Lattice stabilities


Data for Ag relative to FCC A1

(Difference in Gibbs energy between each phase and this reference phase, according toequation (2.33).)

LIQUID11025.076 − 8.891021 T − 1.034 · 10−20 T 7 (298.15 < T < 1234.93)11508.141 − 9.301747 T − 1.412 · 1029 T−9 (1234.93 < T < 3000)

BCC A23400 − 1.05 T (298.15 < T < 3000)

HCP A3300 + 0.3 T (298.15 < T < 3000)

Data for Cu relative to FCC A1

LIQUID12964.736 − 9.511904 T − 5.849 · 10−21 T 7 (298.15 < T < 1357.77)13495.481 − 9.922344 T − 3.642 · 1029 T−9 (1357.77 < T < 3200)

BCC A24017 − 1.255 T (298.15 < T < 3200)

HCP A3600 + 0.2 T (298.15 < T < 3200)

51

Appendix B. Lattice stabilities 52

Data for Bi relative to RHOMBO A7

LIQUID11246.067 − 20.63651 T − 5.955 · 10−19 T 7 (298.15 < T < 544.55)11336.259 − 20.810418 T − 1.661 · 1025 T−9 (544.55 < T < 3000)

BCC A211297 − 13.9 T (544.55 < T < 3000)

BCT A54148.07 (544.55 < T < 3000)

FCC A19900 − 12.5 T (544.55 < T < 3000)

HCP A39900 − 11.8 T (544.55 < T < 3000)

TETRAGONAL A64148.07 (544.55 < T < 3000)

TET α14234 (544.55 < T < 3000)

Data for Sn relative to BCT A5

LIQUID7103.092 − 14.087767 T + 1.47031 · 10−18 T 7 (100 < T < 505.08)6971.587 − 13.814382 T + 1.2307 · 1025 T−9 (505.08 < T < 3000)

FCC A14150 − 5.2 T (298.15 < T < 3000)

HCP A33900 − 4.4 T (298.15 < T < 3000)

RHOMBO A72035 (298.15 < T < 3000)

BCC A24400 − 6.0 T (298.15 < T < 3000)

Appendix B. Lattice stabilities 53

Data for Zn relative to HCP A3

LIQUID7157.213 − 10.29299 T − 3.5896 · 10−19 T 7 (298.15 < T < 692.68)7450.168 − 10.737066 T − 4.7051 · 1026 T−9 (692.68 < T < 1700)

BCC A22886.96 − 2.5104 T (298.15 < T < 1700)

FCC A12969.82 − 1.56968 T (298.15 < T < 1700)

Appendix C

Thermodynamic data for theSn-Zn-Bi system


Bi-Sn 0LLBiSn = 490 + 0.97 T

1LLBiSn = −30 − 0.235 T

0LbctBiSn = 2120 − 1.44 T

1LbctBiSn = −3710.0

0LrhoBiSn = −5760 + 11.834 T

Bi-Zn 0LLBiZn = 18265.09 − 8.6763 T

1LLBiZn = −6061.21 + 0.79581 T

2LLBiZn = −6422.6 + 11.72 T

3LLBiZn = 7227.44 − 9.2905 T

4LLBiZn = 21123.07 − 27.147 T

5LLBiZn = −20747.56 + 22.0176 T

6LLBiZn = −7600.36 + 13.16 T

0LhcpBiZn = 35000

0LrhoBiZn = 10000

Sn-Zn 0LLSnZn = 12710 − 9.162 T

1LLSnZn = −5360 + 3.45 T

2LLSnZn = 835.0

0LbctSnZn = 9260

0LhcpSnZn = 40000

Bi-Sn-Zn 0LLBiSnZn = −76485.59 + 98.4963 T

1LLBiSnZn = 7048.7 − 9.25285 T

2LLBiSnZn = −265.89 − 11.10087 T

54

Appendix D

Binary Ag-Cu phase diagram

55

Appendix E

Binary Sn-Cu phase diagram

56

Appendix F

Binary Sn-Ag phase diagram

57

Appendix G

Binary Ag-Bi phase diagram

58

Appendix H

Binary Sn-Bi phase diagram

59

Appendix I

Binary Zn-Bi phase diagram

60

Appendix J

Binary Sn-Zn phase diagram

61

Appendix K

EDX Analysis

Solder type B (SnAg3.8Cu0.7)

Figure K.1: SE image Figure K.2: Tin

Figure K.3: Copper Figure K.4: Silver

62

Appendix K. EDX Analysis 63

Solder type C (SnAg3.3Bi3.82)

Figure K.5: SE image Figure K.6: Tin

Figure K.7: Silver Figure K.8: Bismuth


Solder type D (SnZn8Bi3)

Figure K.9: SE image Figure K.10: Bismuth

Figure K.11: Tin Figure K.12: Zinc


Solder type A (SnPb36Ag2)

Figure K.13: SE image Figure K.14: Lead

Figure K.15: Silver Figure K.16: Tin

Figure K.17: Copper

Appendix L

Shear deformation of gold grid

Figure L.1: Shear deformation at step 18 Figure L.2: Shear deformation at step 22

Figure L.3: Shear deformation at step 26 Figure L.4: Shear deformation at step 30

66

Appendix M

”Hall of Fame”

Figure M.1: Sample B7, fracture area in a failed solder joint after shear testing (BSE image)

67

Appendix M. ”Hall of Fame” 68

Figure M.2: Sample D16, fracture area in a failed solder joint, caused by constant thermo-mechanical loading (SE image)


Appendix M. ”Hall of Fame” 69


Figure M.5: Sample AJ1, fracture area near interface, caused by constant thermo-mechanicalloading (BSE image)

Samenvatting

Een toenemend milieubewustzijn met betrekking tot loodhoudende soldeerlegeringenheeft tot grote inspanningen geleid om een loodvrije soldeertechniek te ontwikkelen. HetRIPOSTE project van de TU/e en Philips CFT heeft als doel een tool te ontwikkelen, diede betrouwbaarheid van gesoldeerde miniatuurverbindingen in IC-packages voorspelt,onderworpen aan (thermo-)mechanische belasting, waarbij rekening wordt gehoudenmet de relevante microstructuur en chemische eigenschappen van het soldeer. Dezescriptie maakt onderdeel uit van dit project en behandelt hoofdzakelijk afschuifbelasting,kruip, faalgedrag en vervormingsgedrag van drie loodvrije soldeerlegeringen, namelijkSnAg3.8Cu0.7 (SAC), SnAg3.3Bi3.82 en SnZn8Bi3 en SnPb36Ag2 als referentie.

Als gevolg van hoge homologe bedrijfstemperaturen van soldeerverbindingen in elektron-ica, is de deformatie van deze verbindingen altijd een thermo-mechanisch proces. Ditvereist kennis van zowel evenwicht- als niet-evenwicht thermodynamica, omdat tijdensvervorming een constante flow van energie naar het systeem aanwezig is. Dit impliceertdat er geenszins sprake is van een systeem in thermodynamisch evenwicht.

Als de temperatuur van een elektronisch systeem stijgt, dan zet de PCB (printed circuitboard) meer uit dan de component als gevolg van een verschil in uitzettingscoefficient;dit resulteert in afschuifspanningen in de soldeerverbinding. Teneinde de afschuifsterkteen maximale afschuifrek te kunnen onderzoeken, is er een speciaal sample ontworpen,zodat proeven in een miniatuur-trekbank uitgevoerd konden worden. Met behulp vandigitale beeldcorrelatie (DIC) kan het vervormingsgedrag op micronschaal worden geob-serveerd en vergeleken worden met de data van de afschuifproeven.

Intermetallische verbindingen aan de soldeer-substraat interface blijken een belangrijkerol te spelen in de deformatie van de soldeerverbinding bij SnAg3.8Cu0.7 en SnZn8Bi3.Inhomogene en hoge locale rekken kunnen worden waargenomen aan de interface. Dedeformatie van SnAg3.3Bi3.82 is echter veel homogener, wat leidt tot een superieureafschuifsterkte, maar ook brosse breuk. SnAg3.8Cu0.7 heeft de grootste plastische de-formatie, terwijl SnZn8Bi3 onderhevig is aan vreemde verouderingseffecten.

Afschuifspanning in combinatie met hoge homologe temperaturen brengt ook kruip enmogelijk coarsening teweeg. Een proefopstelling om kruip te meten is ontworpen omsamples in een droogstoof te testen. Het blijkt dat SnZn8Bi3 de hoogste kruipsnel-heid heeft van de drie loodvrije soldeerverbindingen, gevolgd door SnAg3.3Bi3.82. SACis de meest kruip-bestendige soldeerlegering met ongeveer dezelfde kruipsnelheid als

70

Samenvatting 71

SnAg3.3Bi3.82 bij 100˚C, maar een betere weerstand tegen kruip bij 150˚C. Coarsen-ing kon niet worden gedetecteerd, waardoor het niet mogelijk was om data hierover teverzamelen.

ESEM- en lichtmicroscopie onthulden de faalmechanismen van de kruipsamples. In SAC-soldeerverbindingen konden parallel aan de interface en loodrecht hierop (van de eneinterface naar de andere), langs grote kolonies worden waargenomen. Het uiteindelijkefalen van de verbinding werd veroorzaakt door scheuren parallel aan de interface doorhet materiaal heen. Bij SnAg3.3Bi3.82 kon dit eveneens worden waargenomen, maarfalen gebeurde langs de intermetallische verbindingen aan de interface. Dit was ook inSnZn8Bi3 te zien, alleen zijn de kolonies in dit materiaal kleiner en hebben een anderevorm. Scheurvorming kon door de hele soldeerverbinding gezien worden in alle richtin-gen. Een wat grilliger falen langs de koloniegrenzen en dwars door het materiaal heenkon hier worden waargenomen en niet alleen langs de interface. EDX-analyse wees tevensuit dat grote hoeveelheden koper naar het oppervlak van de verbinding diffundeert endaar brons vormt met zink.

Dit onderzoek heeft thermodynamische gegevens uit de literatuur opgeleverd, alsmedethermo-mechanische gegevens voor het RIPOSTE project voor drie loodvrije soldeer-verbindingen. De combinatie van deze data met data van andere onderzoeksgebiedenbinnen dit project zal leiden tot een beter begrip van deze materialen en een goedevervanger voor loodhoudende soldeerlegeringen.

Acknowledgements

Writing this thesis has been a job I couldn’t have done without the help of many people.Their help and support is greatly appreciated and will not be forgotten. I would like tothank in particular:My supervisor, Marc Geers, who made it possible to do this project and who has alwaysbeen very supportive and understanding during difficult personal circumstances.My coach, Willem-Pier Vellinga, who’s ideas and enthusiasm have been a great contri-bution to my research.Hans de Vries, who’s expertise in the field of lead-free soldering has been very useful formy research.Jo van Gerven, who has provided me with the shear samples.Jo Sijben, who has made additional samples for me.Toon Hoeben, who has often helped me with his knowledge of equipment and materialsand who has always made sure I got the materials and consumables I needed.Toon van Gils, who has very accurately realized all the designs I made for my experi-ments.Pauline Schmit, who made it possible for me to use the sputter coater at the faculty ofChemical Technology.Jeroen Visser, who helped me with EDX-analysis and who provided me with a numberof EDX-map images.Dick Landheer, who was always prepared to help me with phase diagrams or otherquestions related to alloys and metals.My colleagues and friends from student fellowship ”3004”, who always made sure thatthe atmosphere in our room was good and who helped me with advice and creativitywhen I needed it, especially Sebastiaan Boers.My friends and family, who have stood by me and supported me throughout this project.

And last but certainly not least I would like to thank my girlfriend Manon, who’spatience and love helped me get through difficult periods and who supported me un-conditionally.

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Creep and Failure of Lead-free Solder Alloys · Creep and Failure of Lead-free Solder Alloys ... solder joint using a miniature tensile ... A more capricious failure occurred along