Characterization of the Solvus Transformation in Ternary and Quaternary Bismuth Containing Lead-Free Solder

Alloys

by

Ivan Matijevic

A thesis submitted in conformity with the requirements for the degree of Masters of Applied Science

Material Science and Engineering University of Toronto

© Copyright by Ivan Matijevic 2017

ii

Characterization of the Solvus Transformation in Ternary and

Quaternary Bismuth Containing Lead-Free Solder Alloys

Ivan Matijevic

Masters of Applied Science

Material Science and Engineering

University of Toronto

2017

Abstract

Compliance with RoHS requirements has necessitated the development of high reliability

alternatives to SAC305 lead-free solder. Bismuth containing ternary and quaternary alloys are

emerging as an attractive lead-free high reliability solder. A particular point of interest is the

solvus temperature – the point above which Bi will form a solid solution, and below which Bi

will precipitate out of the matrix. The solvus temperature of these solders was determined using a

novel differential scanning calorimetry (DSC) method. A study was conducted to characterize

the effects of aging above- and below-solvus on microstructure and mechanical properties.

Metallographic analysis showed that above-solvus aging improved bismuth distribution as well

as reduced grain size. It was concluded that above-solvus aging could be used as a restorative

treatment for in service electronics. Nano-indentation was used to collect creep and hardness

data, and it was found that above-solvus aging had a positive effect on mechanical properties.

iii

Acknowledgments

I would like to thank Professor Doug D. Perovic and Dr. Polina Snugovsky for their continued

guidance and support. Without their input this thesis would have no scope or direction. I would

also like to thank the ReMAP group for all the direction they gave me and for providing an

invaluable industry perspective. ReMAP is an industry led Canadian Centre of Excellence

dedicated to enabling research to be brought to market rapidly.

I would like to thank Professor Donald W. Kirk of Chemical Engineering for giving me access to

his group’s differential scanning calorimeter as well as for his advice regarding use of the

instrument. I would like to thank Professor Chandra V. Singh for giving me access to his groups

nano-indenter, and Matt Daly for his assistance in training me, his advice on use of the

instrument (amongst other things) and his help in keeping the machine running. I would like to

also thank Dr. Leonid Snugovsky and Professor Zhirui Wang for giving me access to their

metallographic equipment and ovens without which I would have had difficulty preparing my

samples, as well as giving me valuable advice regarding sample preparation.

I would like to thank Sal Boccia for all of his help surrounding electron microscopy, including

training and discussions on sample preparation. I would also like to thank Dr. Dan Grozea for his

assistance in finding equipment as well as his advice.

Finally, I would like to thank André Delhaise for the many discussions we had regarding scope

and direction, as well as his help in preliminary testing.

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Table of Contents

Acknowledgments.......................................................................................................................... iii

Table of Contents ........................................................................................................................... iv

List of Tables ................................................................................................................................. vi

List of Figures ............................................................................................................................... vii

Chapter 1 Background .....................................................................................................................1

Introduction .................................................................................................................................1

1.1 Causes of Solder Failure ......................................................................................................3

1.2 Current Generation SAC Alloys ..........................................................................................5

1.2.1 SAC Microstructure .................................................................................................5

1.2.2 Limitations of SAC ..................................................................................................8

1.3 Ternary and Quaternary Bismuth Containing Alloys ........................................................11

1.3.1 Advantages of Bismuth ..........................................................................................11

1.3.2 Alloy Selection.......................................................................................................13

Objectives ..................................................................................................................................15

Chapter 2 Determination of Solvus................................................................................................17

Methods .....................................................................................................................................17

1.1 Experimental Design ..........................................................................................................17

1.2 Sample Preparation ............................................................................................................18

DSC Results ..............................................................................................................................19

Verification of Experimentally Determined Solvus ..................................................................21

Discussion .................................................................................................................................23

Chapter 3 Aging Study...................................................................................................................24

Methods .....................................................................................................................................24

v

1.1 Experimental Design ..........................................................................................................24

1.2 Sample Preparation ............................................................................................................26

Results .......................................................................................................................................27

2.1 Microstructural Evaluation ................................................................................................27

2.1.1 Bi Particle Analysis................................................................................................32

2.1.2 Grain Size Analysis................................................................................................34

2.2 Mechanical Evaluation.......................................................................................................37

2.2.1 Creep Study ............................................................................................................37

2.2.2 Hardness .................................................................................................................40

Discussion .................................................................................................................................41

Chapter 4 Conclusions ...................................................................................................................48

Summary of Conclusions ..........................................................................................................48

Future Work ..............................................................................................................................50

References ......................................................................................................................................51

vi

List of Tables

Table 1 - Typical operating temperatures for various electronic grades [16] ................................. 5

Table 2 - Results of initial DSC testing ........................................................................................ 20

Table 3 – Results of secondary DSC testing ................................................................................. 21

Table 4 - Plan for aging study ....................................................................................................... 24

vii

List of Figures

Figure 1 - Typical ball grid array (BGA) solder joint (left), typical gullwing quad flat package

(QFP) solder joint (right) [2]........................................................................................................... 1

Figure 2 - Solder joints loaded in shear due to CTE mismatch between package, solder and PCB

[12] .................................................................................................................................................. 4

Figure 3 - SAC305 as-cast microstructure; (A) points to dendritic areas, (B) points to

interdendritic areas and IMC .......................................................................................................... 6

Figure 4 - Beachball SAC305 grain structure (left), view of {101} cyclic twin nucleus faceted

on {110} planes with arrows representing the preferred growth directions (right) [27] ................ 8

Figure 5 - Cross-sectional view of pad cratering (a), plan view of a pad crater (b) [6] .................. 9

Figure 6 – Excessive copper dissolution of a BGA pad (left) and a plated through hole knee

(right) [30] ..................................................................................................................................... 10

Figure 7 - Binary Sn-Bi phase diagram showing predicted solvus temperatures for Senju (A),

Violet (B) and Sunflower (C) [56] ................................................................................................ 14

Figure 8 - Graphite crucible used to cast 400mg DSC samples and a prepared DSC sample

adjacent ......................................................................................................................................... 18

Figure 9 - A representative successful DSC run (above) and a representative DSC run that

yielded no result (below) .............................................................................................................. 19

Figure 10 – Left: Sunflower aged at 90°C for 5 days (above), Sunflower aged at 70°C for 5 days

(below), Center: Violet aged at 100°C for 5 days (above), Violet aged at 80°C for 5 days

(below), Right: Senju aged at 30°C for 5 days (above) and Senju aged at 50°C for 5 days (below)

....................................................................................................................................................... 22

Figure 11 - Representative images of SAC305 microstructure taken at all aging timepoints: A)

As-Cast, B) Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at 120°C, D)

Aged 300 hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C ................... 27

viii

Figure 12 - Representative images taken of Violet microstructure at all aging timepoints: A) As-

Cast, B) Aged 300 hours at 70°C .................................................................................................. 28

Figure 13 - Lamellar structure of bismuth in as-cast Violet with evidence of grain boundary

sliding ............................................................................................................................................ 31

Figure 14 - Evidence of Ostwald ripening in sub-solvus aged Violet .......................................... 32

Figure 15 - Histogram of all timepoint particle area distributions overlaid (all particle areas

greater that 100 pixels were cropped for legibility) ...................................................................... 33

Figure 16 - Plot of standard deviation and average bismuth particle size with respect to timepoint

....................................................................................................................................................... 34

Figure 17 – Representative EBSD Euler maps of SAC305 taken at all timepoints: A) As-Cast, B)

Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at 120°C, D) Aged 300

hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C .................................... 35

Figure 18 - Representative EBSD Euler maps of Violet taken at all timepoints: A) As-Cast, B)

Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at 120°C, D) Aged 300

hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C .................................... 36

Figure 19 - Grain size of Violet for all aging treatments .............................................................. 37

Figure 20 - Optical micrographs taken at 1000x of nano-indentation sites of: A) As-Cast

SAC305, B) Aged 300 hours at 70°C SAC305, C) Aged 300 hours at 70°C and aged 12 hours at

120°C SAC305, D) As-Cast Violet, E) Aged 300 hours at 70°C Violet, F) Aged 300 hours at

70°C and aged 12 hours at 120°C Violet, G) Aged 300 hours at 70°C and aged 24 hours at 120°C

SAC305, H) Aged 300 hours at 120°C SAC305, I) Aged 300 hours at 70°C and aged 24 hours at

120°C Violet, J) Aged 300 hours at 120°C Violet ........................................................................ 38

Figure 21 - Creep rate after aging for both SAC305 and Violet ................................................... 39

Figure 22 - Hardness of SAC305 and Violet alloys with respect to aging ................................... 40

Figure 23 – As-cast Violet (left), below solvus aged Violet (right) ............................................. 42

ix

Figure 24 - The effect of otherwise invisible scratches on bismuth precipitation on the observed

surface (arrows indicate a scratch) ................................................................................................ 43

Figure 25 - Deformation introduced by the growth of a second phase particle [68] .................... 44

Figure 26 – Average grain size plotted against creep rate for all timepoints ............................... 46

1

Chapter 1 Background

Introduction

Soldering has been a fundamental process in the assembly of all electronic products for as long

as the electronics industry has existed, and is predicted to remain the primary joining technology

in electronics for the foreseeable future [1]. Soldering is the process by which two or more metal

substrates are joined by having a filler metal with a lower melting point flowed into the joint. In

the context of electronic assembly, solders are used to permanently join two substrates, typically

made of copper (Cu). These interconnects are relied upon to provide thermal, mechanical and

electrical bonds between electronic components and the printed circuit boards (PCB) they are

mounted on, and as such are critical to the longevity and reliability of electronic products.

Should a solder joint fail in service or form improperly during assembly, a given product will not

function regardless of how robust the other components are [1].

There are many geometries of solder joints based on the component being attached. Two of the

most common component solder joints are ball grid arrays (BGA) and gullwing quad flat

packages (QFP), as seen in Figure 1.

Figure 1 - Typical ball grid array (BGA) solder joint (left), typical gullwing quad flat package (QFP) solder

joint (right) [2]

2

Tin (Sn) is a major alloying element in solder materials for a number of reasons. It is valued for

its fairly low melting point, its propensity for forming intermetallic compounds and its relatively

low cost [3]. A solder joint consists of three major areas: the component side interface, the bulk

solder and the PCB side interface. This can be most easily seen in Figure 1 on the BGA solder

joint. The two interfaces are similar, but may vary based on the finish of the conductive

substrate. The copper substrate is eroded while the solder is melted and will bond chemically

with tin (Sn) to form a layer of intermetallic compounds (IMC) at the interfaces during melting

and solidification. This interfacial layer is more brittle than the bulk solder, and as such is the

primary location of most drop or shock related solder failures [4].

Eutectic or near-eutectic tin-lead (SnPb) solder has been the gold standard for the electronics

manufacturing services (EMS) industry from its birth some 50 years ago. SnPb solder has proven

to have a combination of mechanical, thermal and chemical properties that provide a robust and

reliable solder joint. However, while its mechanical properties and reliability are excellent, lead

is a hazardous substance with a significant environmental impact. In 2006, the Restriction of

Hazardous Substances Directive, or RoHS, was enacted throughout the European Union. This

legislation restricted the use of six substances deemed hazardous, and necessitated the

replacement of heavy metals, including lead, with safer alternatives [5]. This forced an industry

wide shift within the electronics manufacturing sector to use lead-free solder and soldering

processes for consumer products. Due to the fact that most products are destined for global

markets, this change impacted all EMS companies whether their target market was Europe or

not. In response to this legislation, the EMS industry adopted tin-silver-copper (Sn-Ag-Cu)

alloys (SAC alloys) [6] but many sectors, such as aerospace and defense, maintain exemptions

from RoHS and continue to use the thoroughly researched and well understood SnPb solders [7].

3

In 2011, the European Union revised the RoHS legislation to examine the exemptions

maintained by certain industries. This has driven research into high reliability lead-free solders

for use in aerospace, defense and medical industries [8].

1.1 Causes of Solder Failure

Solder joints are critical to the functionality of electronic assemblies. As such, it is important to

have an understanding of how and why they fail, and what loading conditions lead to these

failures. This is of particular interest when designing experiments to study the mechanical

properties of these alloys. There are three primary in-service causes of failure that are observed

in solder joints: overload, creep and fatigue [9].

Electronics are often handled roughly, either while being transported or assembled or while in

use. Mobile phones are expected to withstand numerous drops or other high strain rate events

over the course of their service life. While a given assembly may survive a single drop, the

cumulative damage from successive events will often lead to failure. The bulk solder is quite

ductile and often able to absorb the energy of a shock event but the IMC layer formed at the

interfaces of the solder joint is quite brittle. Brittle fracture in the IMC layer is the most common

failure mode found in drop shock testing on modern solder alloys [6], [9], [10], [11].

The loading conditions that contribute to fatigue and creep are both related to the thermal

properties of the assembly, namely the coefficient of thermal expansion (CTE). As can be seen in

Figure 2, a component and its PCB substrate will respond differently to changes in temperature,

loading the solder in shear [9], [12], [13].

4

Figure 2 - Solder joints loaded in shear due to CTE mismatch between package, solder and

PCB [12]

During assembly this CTE mismatch can introduce warpage into the assembly. While this

warping typically means the solder joints are stress free, when the assembly is placed into a

fixture and flattened the solder joints are loaded. Due to the low melting point of solders they are

at a high homologous temperature (TH) even at ambient conditions. Consequently, creep is very

active even at low stress which could lead to cracks developing during assembly or shortly

thereafter [9].

While in service, a given electronic assembly will be exposed to a large range of temperature

with respect to time. While not in use, the assembly will dwell at ambient conditions but when

powered on or the load is increased, the components heat up. This exposes the solder joints to

cyclic shear stresses, fatiguing them [9], [14], [15]. Typical temperature ranges that electronics

can be exposed to can be seen in Table 1.

5

Table 1 - Typical operating temperatures for various electronic grades [16]

Temperature Grade Temperature Range

Commercial 0°C to 85°C

Industrial -40°C to 100°C

Extended -40°C to 125°C

Military -55°C to 125°C

Automotive -40°C to 125°C

Many assemblies can be exposed to vibrations due to the nature of their application. As an

example, electronics in vehicles or industrial machinery can be exposed to not only extreme

thermal cycling but vibration as well [11], [17]. Over the service life of a given assembly, both

thermal cycling fatigue and vibration fatigue can contribute to the failure of solder joints.

1.2 Current Generation SAC Alloys

As previously mentioned, SAC alloys have become the industry standard for lead-free soldering

in consumer products. This however has not been without its own issues and was met with an

initial wave of resistance. The replacement of abundant and inexpensive lead with silver and

copper has driven intrinsic manufacturing costs up in addition to the incidental costs of

developing new lead-free soldering processes. Although increasing costs are a significant issue,

one of the main concerns the industry has with the SAC alloys is that of reliability; SAC has

been demonstrated to have poorer reliability and a consequentially shorter lifespan when

compared to its SnPb predecessor [1], [6], [18]–[20].

For the purposes of this thesis, SAC305 (Sn-3.0%Ag-0.5%Cu) will be the exemplar SAC alloy.

1.2.1 SAC Microstructure

The as-cast SAC305 microstructure is best described by its typical solidification pathway. First,

tin-beta (solid solution tin) will nucleate and solidify to form primary dendrites. Tin-beta

incorporates no Cu and very little Ag, and as such these are segregated to the liquid around these

6

dendrites. These interdendritic areas grow into a Cu6Sn5, Ag3Sn and beta-tin ternary eutectic.

The distinction between these areas can be see in Figure 3. It has been observed however that in

a BGA grid, there will be some solder joints with an unusual microstructure, such as forming one

or two large dendrites or a binary Ag3Sn-Sn phase instead of a ternary eutectic. Forming very

large tin dendrites is dangerous mechanically due to the anisotropy of tin; if the growth direction

happens to be unfavorable in shear then the solder joint will exhibit significantly worse

mechanical properties, and the CTE mismatch of tin in different orientations introduces internal

stresses during thermal cycling.

Ag3Sn forms a network of thin plates while Cu6Sn5 has multiple morphologies. It appears as

continuous scalloped grains in interfacial IMC and may form hexagonal or acicular faceted

whiskers on top. It appears as either hexagonal or acicular whiskers in the bulk matrix. Cu3Sn

appears as a planar thin layer and is only present in interfacial IMCs [21].

Figure 3 - SAC305 as-cast microstructure; (A) points to dendritic areas, (B) points to

interdendritic areas and IMC

A

B

7

Unusual SAC microstructures are a combination of factors. SAC alloys have been shown to

undercool to temperatures of 20°C or more below liquidus readily. It has been shown that Ag3Sn

and Cu6Sn5 are poor catalysts for beta-tin nucleation, and as such significant undercooling is

required to start freezing. Under some conditions, tin-beta will be the first phase to nucleate

while under others Ag3Sn will be formed as the primary phase. The resultant microstructures are

very different; beta-tin nucleation will give what is considered a typical microstructure seen in

Figure 3, while Ag3Sn nucleation leads to the formation of large plates of that brittle phase [22],

[23]. It should be noted that the cooling rate plays a significant role in microstructure as well; it

has been shown that cooling rate corresponds to the volume fraction of beta-tin dendrites formed

during freezing. The volume fraction increases from 5% with a cooling rate of 1°C/s to 65% with

a cooling rate of 100°C/s. Thermal connections in the form of copper traces and the release of

latent heat from adjacent solder joints nucleating earlier or later during the cooling process can

produce a wide range of cooling rates in a BGA component’s solder joints. This complex and

dynamic system means that there is a significant range of microstructure produced in the solder

joints of a single component during one processing window [24].

An example of one such unusual microstructure observed in SAC alloys is the so called “Kara’s

Beachball” which can be seen in Figure 4. Tin is a highly anisotropic material. The tetragonal

unit cell of tin has different mechanical properties, coefficients of thermal expansion and growth

rates from a molten state based on direction. The anisotripic growth of Sn dendrites at high

undercooling results in a “beachball” grain structure caused by cyclic twinning growth from a

silver or copper atom nucleation site. The resultant grain structure is very coarse and if it

solidifies in an unfavorable orientation this can lead to significantly reduced mechanical

8

properties when compared to interlaced dendrites which form at lower undercooling as other

directions of growth become more favorable [25]–[27].

Figure 4 - Beachball SAC305 grain structure (left), view of {101} cyclic twin nucleus

faceted on {110} planes with arrows representing the preferred growth directions (right)

[27]

1.2.2 Limitations of SAC

One of the most pressing issues with SAC alloys is their high melting point and consequentially

high processing temperatures. SAC305 has a melting range of 217-220°C which is a significant

increase from eutectic SnPb which melts at 183°C. This increased processing temperature

increases intrinsic assembly costs due to an increased energy requirement during processing,

while also increasing the carbon footprint of electronics assembly [6].

Both the PCB substrate and the components being soldered are typically sensitive to temperature.

As a result, while SnPb solders were routinely processed at 40°C above their melting point, SAC

alloys see a much lower superheat. This may also cause issues with not passing out of the “pasty

range” prior to solidification, the area of a phase diagram between solidus and liquidus

temperatures at non-eutectic compositions. This necessitates a significantly higher degree of

process control, especially when dealing with complex populations and temperature sensitive

components as well as reducing the wettability of the solder during processing [1], [26], [28].

9

The higher processing temperatures also have forced a shift in substrate material. The PCB

laminates that were used with SnPb solders have been replaced by new materials with higher

glass transition temperatures (Tg) to allow the substrate to survive processing. These new

materials, however, are stiffer due to the inclusion of ceramic particle fillers and the use of new

curing agents. They are consequently more susceptible to warpage and delamination, and are

susceptible to a new failure mode known as pad cratering. Pad cratering refers a failure in the

laminate of a PCB below the copper substrate as can be seen in Figure 5. This may sever internal

copper traces, causing electrical failure [6], [20], [28]–[30].

Figure 5 - Cross-sectional view of pad cratering (a), plan view of a pad crater (b) [6]

SAC305 has been studied at length, and while the as-cast mechanical properties even outperform

SnPb solders, the microstructural changes they experience while in service are significant and

detrimental. Intermetallics present in the solder will coarsen as the solder joints age, which is

both an issue in the bulk solder and at the interfacial IMC layers. This increases the susceptibility

of these solder joints to brittle fracture caused by drop and shock related events. Grain growth is

very active in these systems at the operating temperatures, so even solder joints with a relatively

fine, interlaced dendritic grain structure may develop into coarse grains in unfavorable

orientations [6], [29]–[36].

10

Silver is present in SAC alloys to reduce the liquidus temperature, which both reduces process

temperature and improves wetting. High amounts of silver (such as in SAC305) also improves

thermal fatigue resistance. However, when present in this quantity, Ag3Sn intermetallics formed

in the bulk solder form brittle platelets that cause poor performance in high strain rate events,

and add substantial cost to the solder [6], [37].

SAC solders demonstrate inferior wetting behavior when compared to SnPb solder, even with a

high amount of silver present. While it is adequate to achieve a good solder joint, the surface

tension of the melt prevents flow on a copper pad. As such, there may be exposed copped in the

completed assembly which is susceptible to corrosion. This also makes wave soldering of high

aspect ratio through hole components more difficult as they are more likely to have inadequate

barrel fill or to have excessive voiding. This requires reworks of the assembly which presents yet

another issue. Dissolution of the copper substrate occurs quite quickly, which means that rework

may entirely dissolve corners of plated through holes and thin BGA pads as can be seen in

Figure 6 [6], [30].

Figure 6 – Excessive copper dissolution of a BGA pad (left) and a plated through hole knee

(right) [30]

As SAC solder joints are fatigued in service, the alloy has a tendency to recrystallize ahead of the

fatigue crack. This introduces a new mode of plastic deformation in the form of grain boundary

11

sliding. Once this happens, cracking at grain boundaries and voiding at triple junctions will allow

fatigue cracks to grow in an intergranular fashion [13], [34], [38].

Tin whiskers are spontaneous growths of thin tin structures from the surface of a solder joint.

These are a significant reliability concern as they can bridge leads and cause short circuits, cause

issues with high-frequency circuits and create loose conductive debris. Lead is well known to

suppress tin whisker formation, but SAC alloys are susceptible and indeed there has been an

increase in whisker related failures from the time RoHS has taken effect [36], [39]–[41].

1.3 Ternary and Quaternary Bismuth Containing Alloys

Bismuth containing solders were first considered by the industry when RoHS was being

introduced. They were found to have good properties that would make them a very viable

candidate to replace SnPb solders with one critical shortcoming [20]. As the industry was just

moving away from lead containing solders and finishes, components could be expected to

occasionally be pre-tinned with leaded solder. The Sn-Pb-Bi system forms a ternary eutectic at

95.3C which is within the operating range of some assemblies, and as such is not acceptable

[42]. For this reason, they were not considered for many years and SAC alloys became the

industry standard. Now that RoHS has been in effect for a decade, the industry has been largely

purged of leaded components and as such this risk is mitigated [6], [30]. The alloys being

considered for this study all have copper for its positive effects in IMC formation and prevention

of copper pad dissolution, and some have silver to lower melting temperature, improve wetting

and improve performance in thermal cycling.

1.3.1 Advantages of Bismuth

Many bismuth containing ternary and quaternary alloys have been shown to have melting

temperatures lower than SAC305 by as much as 10°C. This would be enough to use lower Tg

12

PCB materials as with leaded solders, as well as widening process windows considerably,

allowing a higher superheat of the solder which among other things improves wetting. Bismuth

as an additive has also been show to improve solder wetting [20], [29], [43], [44].

The typically lower percentage of silver used in these solders reduces the amount of brittle IMC

formed in the matrix, which improves drop shock performance [20]. The lower silver content

also helps reduce the cost of these alloys. Bismuth has been shown to inhibit the growth of

interfacial and bulk IMC during aging, which would help preserve mechanical properties with

time in service [45].

Bismuth, being soluble in tin, has a solid solution strengthening effect on the solder. Bismuth

precipitates form particles that pin grain boundaries and act as barriers to dislocation motion. The

addition of bismuth should serve to stabilize grain size in bismuth containing alloys and improve

its mechanical properties [44], [46], [47].

It has been shown that the presence of bismuth acts to stabilize the microstructure in aging

through a number of mechanisms. Bismuth containing alloys have been shown to outperform

SAC alloys in mechanical testing after aging and during accelerated thermal cycling (ATC) [46],

[48], [49]. It has also been demonstrated that bismuth plays a role similar to lead in the

suppression of tin whiskering by stress relaxation [6], [39].

It has been shown that bismuth containing solders exhibit increased creep resistance compared to

SAC equivalents, primarily by way of Zener pinning [25], [50], [51].

Lead-free solder alloys exposed to low temperatures may experience a phase transformation

from beta-tin (body centered tetragonal) to alpha-tin (diamond cubic). This phase transformation

causes a 26% increase in volume, which could initiate considerable cracking and have a very

13

detrimental impact on solder joint reliability. This effect was suppressed by the presence of lead,

which is a soluble impurity. Silver and copper, however, are insoluble which has made SAC

alloys susceptible to this transformation. Bismuth, like lead, is a soluble impurity, which may

make bismuth containing alloys more attractive for low temperature applications [52].

Finally, one of the most interesting features of bismuth containing alloys is the fact that the

solvus line for bismuth lies within the range of temperatures one could expect the assembly to

experience in service. This solvus line represents the temperature above which bismuth will tend

to dissolve into solid solution and below which it will precipitate out of solid solution to form

bismuth particles. It is expected, and has been observed, that the alloys experience significant

microstructural modification when passing this temperature [6], [47], [48]. Thus it is important to

characterize the exact nature of these changes to best understand how they impact solder

reliability.

1.3.2 Alloy Selection

The alloys selected for investigation in this work are: Violet (Sn-2.25%Ag-0.5%Cu-6.0%Bi),

Sunflower (Sn-0.7%Cu-7%Bi) and Senju (Sn-2.0%Ag-0.5%Cu-3.0%Bi). Violet is a variant of

“Paul”, an alloy proposed by Vianco [20], [53]. It has an excellent pasty range, a lower silver

content such that Ag3Sn plates do not form which may improve drop and shock performance and

have a higher bismuth content to help mitigate whisker formation and growth [19], [20], [53].

Sunflower has an excellent pasty range and features no silver whatsoever, which may again

improve mechanical performance and reduces cost significantly. Senju was selected due to its

low bismuth content and because of the significantly different solvus temperature it has as a

result.

14

A binary Sn-Bi phase diagram can be used as a rough approximation of these ternary and

quaternary systems’ solvus temperatures. The addition of alloying elements such as silver or

copper to the binary Sn-Bi system will reduce the solubility of bismuth in tin, and as such one

expects an increase in the solvus temperature as the percentage of alloying elements increases.

As seen in Figure 7, the solvus temperature of many of these candidate ternary and quaternary

alloys is in a range that an assembly could conceivably be subjected to during use. This

necessitates the understanding of how these alloys behave as they age above and below their

respective solvus temperatures, and of the resultant microstructures. There has been some

literature investigating the phase diagrams of these ternary and quaternary systems, but little

study has been devoted to experimental verification of the solvus temperature [54], [55].

Figure 7 - Binary Sn-Bi phase diagram showing predicted solvus temperatures for Senju

(A), Violet (B) and Sunflower (C) [56]

15

Objectives

SAC alloys have significant limitations in high reliability applications. With RoHS exemptions

being revisited, Bi containing solder alloys are becoming a very attractive solution to this

problem. While it has been demonstrated that these alloys can outperform SAC in reliability

testing, the majority of characterization has been from the perspective of existing standards

established for current generation solder alloy systems.

To better understand the reasons for which bismuth containing solders are outperforming current

SAC alloy, it is important to characterize the behavior of bismuth in these systems beyond

nebulous statements regarding solid solution strengthening and grain boundary pinning. Current

solder systems are already dynamic, but the addition of bismuth introduces a solid-state phase

transformation that lies within the feasible range of service temperatures. Little work has been

done to examine the changes these systems experience when aged above and below the solvus

temperature yet conventional knowledge suggests that aging in these two regimes will have

significantly different effects upon the microstructure and consequently mechanical behavior of

these alloys. To this end, this thesis aims to experimentally determine the solvus temperatures for

the aforementioned bismuth containing ternary and quaternary solder alloys. Once this value is

established, this thesis aims to characterize the effects of aging at above and below solvus

temperatures. The microstructure of aged bismuth containing solders will be characterized by

analysis of bismuth particle uniformity as well as grain size analysis using electron backscatter

diffraction (EBSD). To examine the effects of these microstructural changes on the mechanical

properties of these solders, the creep properties and hardness will be tested using

nanoindentation. This thesis hopes to contribute in these ways to the fundamental understanding

16

of such a dynamic alloy system with the goal of helping optimize these alloys for reliability

applications.

17

Chapter 2 Determination of Solvus

Methods

1.1 Experimental Design

This project aimed to use differential scanning calorimetry (DSC) to experimentally quantify the

solvus temperature, to be then verified using traditional metallographic techniques. The difficulty

of using DSC for such an investigation is that the solid-solid phase transformation generates very

little signal. As such, a DSC methodology involving a significantly increased sample mass and a

high heating rate was devised to maximize the signal produced [57].

The heat flow signal in DSC is not an exact representation of the true heat flow. The reported

value is delayed in time by the heat capacity of the pan and of the sample being measured which

smears the true curve to the right in heating. It stands to reason that samples with a higher mass

or using a high heating rate will experience a more significant degree of thermal lag, as is the

case in the aforementioned tests. Unfortunately, due to the nature of the peak being characterized

a lower mass would have been impractical to use with the available instrumentation. The DSC

instrumentation used in these tests does however automatically compensate for thermal lag by

modeling the smear it introduces in real time. Considering that this thesis is concerned only with

the peak temperature, and that the proposed methodology would be using a significantly lower

heating rate than has been used in studies requiring a higher degree of precision [58], it was

decided that the effects of thermal lag could be neglected and that farther study lies outside the

scope of this paper [58]–[60]. Additionally, there are no other expected thermal events near the

solvus temperature and thus the expected effect is a few degree shift of peak temperature at worst

as opposed to an overlap of events [57].

18

In accordance with ASTM E793, a ramp time of two minutes before the expected thermal event

(predicted using the binary phase diagram) was used to establish an adequate baseline [61]. A

TA Instruments Q2000 DSC was programmed to equilibrate at -20°C, ramp at 30K/min to 120°C

and to hold isothermally for two minutes at the final temperature. The DSC used in this study did

not have a cooling loop, and as such testing was performed only in heating. Three samples of

each alloy were tested initially, and additionally three samples of binary Sn-Bi solder with equal

bismuth content to each alloy (Sn-7.0%Bi for Sunflower, Sn-3.0%Bi for Senju and Sn-6.0%Bi

for Violet) were tested as a control. An additional 7 samples of Violet and Sunflower were tested

to confirm the results found in the initial runs. In order to verify the results of DSC testing,

traditional metallographic techniques were employed.

1.2 Sample Preparation

All DSC samples were prepared from 99.99% purity 50 gram ingots. These ingots were cut in

two then melted in a graphite crucible on a laboratory hot plate to form smaller bars. This

process was repeated until the desired sample mass of 400 mg was reached. The samples were

then melted and drop cast into aluminum DSC pans to maximize thermal conductivity as seen in

Figure 8.

Figure 8 - Graphite crucible used to cast 400mg DSC samples and a prepared DSC sample

adjacent

19

They were kept in their liquid state for 30 seconds before being placed on a room temperature

steel block to solidify. These pans were then aged for one week below their binary equivalent

alloy’s solvus temperature to ensure significant amounts of bismuth had precipitated from the

beta-tin matrix.

DSC Results

The initial phase of DSC testing showed that while many runs yielded no results, those that did

had results that were both realistic and in line with predicted values. Representative failed and

successful runs can be seen in Figure 9.

Figure 9 - A representative successful DSC run (above) and a representative DSC run that

yielded no result (below)

20

Table 2 - Results of initial DSC testing

Alloy

Sample

1 Peak

(°C)

Sample

2 Peak

(°C)

Sample

3 Peak

(°C)

Average

(°C)

Difference

(°C)

Sunflower 83 82 82 82

3 Sn-

7.0%Bi 82 76 N/R 79

Violet N/R 86 92 89

7 Sn-

6.0%Bi 82 82 N/R 82

Senju N/R N/R 42 42

6 Sn-

3.0%Bi N/R N/R 36 36

The difference seen in Table 2 between the solvus of the ternary and quaternary alloys and their

binary equivalents were as expected. Sunflower and its binary equivalent had a smaller

difference in solvus temperature when compared to Violet and Senju, likely because has fewer

additives to the binary system. Violet and Senju, which both have comparable concentrations of

copper and silver, had a predictably similar difference in solvus with relation to their respective

binary equivalents. The experimentally determined solvus temperatures of the binary alloys align

with what is present in the calculated phase diagram [54], and the ternary and quaternary alloys

had solvus temperatures only slightly higher than those binary alloys.

The second phase of DSC testing focused on Violet and Sunflower, and excluded Senju due to

the difficulty encountered when attempting to produce a signal in such a low bismuth alloy. As

can be seen in Table 3, many of the runs again yielded no result. The runs that did have a result

had excellent agreement with the results of the initial DSC testing. From both sets of testing, a

solvus temperature of 90°C was taken for Violet, 80°C for Sunflower and 40°C for Senju.

21

Table 3 – Results of secondary DSC testing

Alloy Sample Peak Temperature (°C)

Violet

(Sn-2.25%Ag-0.5%Cu-6.0%Bi)

1 91

2 91

3 N/R

4 90

5 N/R

6 N/R

7 N/R

Average 91

Sunflower

(Sn-0.7%Cu-7.0%Bi)

1 80

2 82

3 84

4 N/R

5 85

6 N/R

7 85

Average 83

Failed runs were discarded as outliers when analyzing the data, but for this reason it was deemed

necessary to verify these predicted values and to assume significant error is present in the

experimentally determined solvus temperatures.

Verification of Experimentally Determined Solvus

Samples were prepared and aged for 5 days at 10°C both above the experimentally determined

solvus temperature and 10°C below the experimentally determined solvus temperature. Nine

images were taken at randomly selected locations across three samples for each timepoint. The

resulting SEM images were examined for evidence of aging above and below the solvus

temperature, such as penetration of bismuth particles into the beta-tin dendritic areas, particle

uniformity, number of particles and evidence of Ostwald ripening.

22

Figure 10 – Left: Sunflower aged at 90°C for 5 days (above), Sunflower aged at 70°C for 5

days (below), Center: Violet aged at 100°C for 5 days (above), Violet aged at 80°C for 5

days (below), Right: Senju aged at 30°C for 5 days (above) and Senju aged at 50°C for 5

days (below)

Figure 10 shows samples that were aged above their solvus temperature. These samples had

bismuth particles, seen as bright spots in the micrographs, present in the dendritic regions as

opposed to being confined to the eutectic regions that formed upon solidification. There is

evidence of Ostwald ripening in both sets of images, which can be seen where coarse particles

are surrounded by a bismuth depleted zone, but the below-solvus aged samples show more sites

where it has occurred. The particle size appears to be more uniform in the samples that were

aged above solvus when compared to the samples aged below solvus. All this confirms that the

solvus temperatures for Violet lies in a range between 80°C and 100°C, in a range between 70°C

and 90°C for Sunflower and in a range between 30°C and 50°C for Senju.

23

Discussion

A novel DSC methodology for experimentally determining the solvus temperature of ternary and

quaternary Bi containing solder alloys was designed and implemented.

Over the course of these experiments the absence of a peak corresponding to the solvus phase

transformation was observed in a high percentage of tests. During explorative testing conducted

prior to the experiments described in this thesis, it was found that the signal produced from

samples aged above solvus prior to DSC testing could not be detected. This aging treatment was

intended to increase the uniformity of the samples and prevent bismuth segregation, but it is

possible that little bismuth had precipitated out of solid solution in the space between the aging

and testing. In the interest of achieving better signal, aging was conducted prior to testing below

the predicted solvus temperature to promote bismuth precipitation in subsequent samples. This

increased the signal to detectable limits, indicating that a higher percentage of bismuth had

indeed precipitated out of the tin-beta matrix. This however meant that the bismuth was likely

segregating as well, which is proposed as a possible cause of no observed peaks in some

samples. Another possibility is that the failed samples did not wet to the aluminum pan substrate.

This would have increased the thermal resistance between the pan and the sample, possibly

distorting the peak across such a wide area that it was no longer detectable. Voiding at the

interface with the pan could have had a similar effect.

The experimental solvus temperature was found to be approximately 91°C for Violet and 83°C

for Sunflower after two rounds of testing, and 42°C for Senju after one round of testing. While

this value incorporates an unknown error in the form of thermal lag, the solvus temperatures can

be said to lie with certainty in the range of 80°C to 100°C for Violet and 70°C to 90°C for

Sunflower and between 30°C and 50°C for Senju.

24

Chapter 3 Aging Study

Methods

With solvus temperatures established, the effects of aging in the below and above solvus regimes

could be studied.

1.1 Experimental Design

Aging temperatures were selected to both reflect real-world service conditions and to mimic a

hypothetical in-service restorative treatment. From existing knowledge, it can be seen that above

solvus aging could serve to “restore” an over aged alloy. A high duration above solvus timepoint

was added to have as a comparison for above solvus treatments of a more realistic duration. A

plan of aging timepoints can be seen in Table 4. This aging plan was applied to both Violet and

SAC305.

Table 4 - Plan for aging study

Aging

Treatment As-Cast 300h 70°C

300h 70°C + 12h 120°C

300h 70°C + 24h 120°C

300h 120°C

Solvus

Regime N/A

Below

Solvus Mixed Mixed

Above

Solvus

As-cast samples were naturally selected as a timepoint to assess the initial properties of Violet.

70°C was selected as a representative service temperature, and 300 hours was chosen as the

duration for this aging to simulate some arbitrary amount of time in service. Two above solvus

restorative treatments were then applied to samples aged for 300h at 70°C. These were chosen to

be 12 hours and 24 hours at 120°C as these are the limits of what the industry could consider

reasonable for a restorative treatment without causing significant damage to the PCB laminate

and components. A long term above solvus timepoint was also added to represent the ideal

25

microstructure after the restorative process. These values were chosen with the input of the

ReMAP team who have extensive industry experience.

SEM microscopy was conducted using a Hitachi SU3500 with EBSD capabilities. Fields of view

were randomly chosen and of equal magnification across samples. A magnification of 1000x was

chosen for bismuth particle analysis to best resolve fine particles without losing larger particles.

Nano-indentation was used to test both creep and hardness. For creep testing, an indent force of

5000 µN was used, and this force was held for one hour. The duration was selected such that

stage two steady state creep, where work hardening and annealing are in balance and well

understood creep mechanisms can be observed, was apparent [62], [63]. The actual time was

determined by adjusting the duration of the tests until a significant linear region in strain over

time charts was seen. The indentation force was selected such that the indent would have a

relatively small (roughly 15 µm across) footprint. This was done to ensure the collected data

would be only representative of the beta-tin matrix, which would be variable with time, and not

of the intermetallics surrounding the beta-tin which would contaminate the results. Care was

taken to ensure creep test indents would lie only within dendritic spaces. Hardness was measured

using 10000 µN indents with a hold time of 10 seconds. This larger footprint was selected to

encompass numerous intermetallics and be representative of the bulk material. Seven indent

locations per sample were selected for creep testing, and two locations were selected for

hardness testing.

It was decided that three samples with seven indents on each per timepoint would be adequate to

give statistically relevant results for nano-indentation creep testing. Due to the three-dimensional

stress field introduced by nano-indentation, typical creep metrics such a stress tensors were

considered to be outside of the scope of this project. While it is possible to calculate stress

26

tensors it is a very involved process [63]–[65] and with the inclusion of a SAC305 reference in

the testing plan it was deemed unnecessary. IMC coarsening has been investigated at depth in

existing literature and as such it is considered outside the scope of this project.

1.2 Sample Preparation

Samples were prepared from 99.99% purity 50 gram ingots. These ingots were cut in half and

melted repeatedly until the desired sample mass of roughly 12 grams was achieved. This was

done to prevent bismuth segregation and maintain uniformity of the samples.

Samples were placed into preheated firebrick holders before being placed into ovens for aging.

This was done to minimize the temperature fluctuations the samples would experience when the

oven door is opened to remove other timepoints from the oven.

Tin alloys are known to be difficult to etch, and after some testing this method was abandoned.

EBSD was selected as the method for characterizing grain size. Metallographic cross sections

were prepared with much attention paid to preserving the true grain structure for subsequent

EBSD imaging. Samples were ground with 400, 800, 1200 and 1200 fine grit papers for a minute

at each stage and then gently lapped with 3µm diamond polishing compound for 8 minutes. The

long polishing step was to ensure that the damage layer created by grinding had been removed.

Finally, samples were given a final polish using colloidal silica for two minutes [66], [67].

Samples were dried in vacuum and imaged as soon as possible after cross sectioning to minimize

the effects of surface migration of bismuth particles.

27

Results

2.1 Microstructural Evaluation

After aging, multiple images were taken of the cross sectioned samples. In this section,

observations and relevant micrographs will be presented.

The relatively low impact of aging on SAC305 microstructure can be seen in Figure 11. There

was some limited IMC coarsening observed but no calculations were carried out.

Figure 11 - Representative images of SAC305 microstructure taken at all aging timepoints:

A) As-Cast, B) Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at

120°C, D) Aged 300 hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C

A B

C D

E

28

As seen in Figure 12, as aging time above solvus increases, bismuth particles tend to be more

uniformly distributed, and to have a more uniform size. Bismuth penetration into dendritic

spaces increased as time aged above the solvus temperature was increased.

Figure 12 - Representative images taken of Violet microstructure at all aging timepoints:

A) As-Cast, B) Aged 300 hours at 70°C

A

B

29

Figure 12 – Representative images taken of Violet microstructure at all aging timepoints:

C) Aged 300 hours at 70C and aged 12 hours at 120C, D) Aged 300 hours at 70C and aged

24 hours at 120C

C

D

30

Figure 12 – Representative images taken of Violet microstructure at all aging timepoints:

E) Aged 300 hours at 120°C

As can be seen in Figure 13, there is evidence of particle induced strain being accommodated by

grain boundary sliding on the surface in areas with a high bismuth concentration. The dark line

surrounding the cluster of bismuth was observed under tilt and found to be shadowing from a

step. The lamellar structure of bismuth seen within the displaced grain was common throughout

as-cast samples.

E

31

Figure 13 - Lamellar structure of bismuth in as-cast Violet with evidence of grain boundary

sliding

Violet samples aged below solvus were found to have structures such as those seen in Figure 14.

Coarse bismuth particles surrounded by an area depleted of bismuth are a good indication of

Ostwald ripening, the process by which large particles grow at the expense of fine particles

which are consumed.

32

Figure 14 - Evidence of Ostwald ripening in sub-solvus aged Violet

2.1.1 Bi Particle Analysis

Three random fields of view were imaged at 1000x for bismuth particle analysis for each

timepoint. The images’ white balance was adjusted using ImageJ software and the size in pixels

of all particles was collected. A histogram of particle area was made for all timepoints which can

be seen in Figure 15. The distribution was found to be a negative exponential distribution,

indicating that there could be many fine particles potentially smaller than one pixel that were

below the imaging threshold.

33

Figure 15 - Histogram of all timepoint particle area distributions overlaid (all particle

areas greater that 100 pixels were cropped for legibility)

To better understand the difference between these distributions, the mean of particle area and

standard deviation of particle area for each timepoint were calculated as can be seen in Figure

16. Standard deviation is effectively the width of the data set; if standard deviation decreases

particle uniformity increases and if standard deviation increases this indicates that the

distribution is stretching meaning particle uniformity is decreasing.

0

500

1000

1500

2000

2500

3000

3500

1 4 7

10

13

16

19

22

25

28

31

34

37

40

43

46

49

52

55

58

61

64

67

70

73

76

79

82

85

88

91

94

97

10

0

Freq

uen

cy

Particle Area (pixels)

Histogram of Bi Particle Area for all Timepoints

AC 70C 12h Treat 24h Treat 120C

34

Figure 16 - Plot of standard deviation and average bismuth particle size with respect to

timepoint

2.1.2 Grain Size Analysis

EBSD was conducted on samples that were prepared for imaging in SEM. One sample of the

three for each timepoint was treated as such. EBSD maps were taken at three points along the

diameter of each sample: at the two extremes and one in the center. The data for SAC305 and

Violet can be seen in Figures 17 and 18 respectively. Gain size data was collected in Channel 5

HKL Tango, which can be seen in Figure 19. It was found that SAC305 had large, coarse grains

as previous work has shown. Violet was found to have finer grains than SAC305 in as-cast, and

with increasing time and temperature above solvus was found to have a reduction in grain size.

0

10

20

30

40

50

60

70

As-Cast 300h 70°C 120h 70°C + 12h120°C

120h 70°C + 24h120°C

300h 120°C

Mean Particle Area (pixels) and Standard Deviation

Standard Deviation

Mean Particle Area

35

Figure 17 – Representative EBSD Euler maps of SAC305 taken at all timepoints: A) As-

Cast, B) Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at 120°C,

D) Aged 300 hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C

E

C D

A B

36

Figure 18 - Representative EBSD Euler maps of Violet taken at all timepoints: A) As-Cast,

B) Aged 300 hours at 70°C, C) Aged 300 hours at 70°C and aged 12 hours at 120°C, D)

Aged 300 hours at 70°C and aged 24 hours at 120°C, E) Aged 300 hours at 120°C

A

C

E

B

D

37

Figure 19 - Grain size of Violet for all aging treatments

2.2 Mechanical Evaluation

2.2.1 Creep Study

As mentioned previously, creep testing was conducted on 3 samples for each timepoint. Some

representative optical images of indents are shown in Figure 20. It is clear that SAC305 did not

have a significant change visually between timepoints, and there was a large degree of grain

boundary sliding occurring in all samples. All SAC305 indents were of roughly the same size

and there was no clear trend. Violet samples showed evidence of grain boundary sliding in

samples aged below solvus, and the degree of that plastic deformation mechanism grew less as

time aged above solvus was increased. The size of Violet indents appears to be roughly the same

for all below solvus aged timepoints but above solvus aged timepoints had a smaller indent

footprint. These three timepoints were all similar in size as well.

0

0.5

1

1.5

2

2.5

3

3.5

4

4.5

5

As-Cast 300h 70°C 300h 70°C + 12h120°C

300h 70°C + 24h120°C

300h 120°C

d (

um

)

Average Grain Diameter of Violet for all Timepoints

38

Figure 20 - Optical micrographs taken at 1000x of nano-indentation sites of: A) As-Cast

SAC305, B) Aged 300 hours at 70°C SAC305, C) Aged 300 hours at 70°C and aged 12

hours at 120°C SAC305, D) As-Cast Violet, E) Aged 300 hours at 70°C Violet, F) Aged 300

hours at 70°C and aged 12 hours at 120°C Violet, G) Aged 300 hours at 70°C and aged 24

hours at 120°C SAC305, H) Aged 300 hours at 120°C SAC305, I) Aged 300 hours at 70°C

and aged 24 hours at 120°C Violet, J) Aged 300 hours at 120°C Violet

A B C

D E F

G H

I J

39

In Figure 21 it can be seen that Violet outperformed SAC305 consistently in creep testing. The

creep response of SAC305 followed no trends while Violet saw a definite decrease in creep rate

after aging above solvus.

Figure 21 - Creep rate after aging for both SAC305 and Violet

Statistical analysis was conducted of Figure 21. T-tests showed that as-cast and 300h 70°C

timepoints for Violet are statistically indistinguishable. 300h 70°C + 12h 120°C was statistically

distinct from both 300h 70°C and the 24 hour treatment, and the 24 hour treatment and 300h at

120°C were the same. SAC305 showed no real trend, and all points were statistically

indistinguishable. It can be seen that the first two timepoints for Violet were very similar, that the

12 hour treatment had a positive effect on creep resistance, and that the 24 hour treatment was

0

0.02

0.04

0.06

0.08

0.1

0.12

As Cast 300h 70°C 300h 70°C + 12h120°C

300h 120°C + 24h120°C

300h 120°C

Cre

ep R

ate

(nm

/s)

Nano-Indentation Creep Rate by Alloy and Aging Treatment

SAC305

Violet

40

more effective still. The 24 hour treatment appeared to be on par with the change caused by a

300 hour treatment to the creep resistance of these alloys.

2.2.2 Hardness

The hardness values collected from nano-indentation in this study seen in Figure 22 are in line

with what much of existing literature has also found. SAC305 was found to sharply drop in

hardness from as-cast to the remaining timepoints, after which it stayed roughly the same. Violet

was found to have a similar hardness across all timepoints.

Figure 22 - Hardness of SAC305 and Violet alloys with respect to aging

0

5

10

15

20

25

30

35

40

45

50

As Cast 300h 70°C 300h 70°C + 12h120°C

300h 120°C + 24h120°C

300h 120°C

Har

dn

ess

(Vic

kers

)

Hardness of SAC305 and Violet

SAC305 Violet

41

Discussion

From all of the results presented above, it can be seen that the addition of bismuth has positive

effects on the microstructure regardless of aging conditions. Violet consistently outperformed

SAC305 in mechanical testing.

A major issue that this project faced was how active this system was at room temperature. Care

was taken to ensure that samples were taken from final polishing to imaging as quickly as

possible, but scheduling conflicts meant that some samples experienced an additional day or two

of room temperature aging. This is evident in in the plot of mean particle area (Figure 16) versus

timepoint, where data followed no obvious trend and the two samples that experienced additional

room temperature aging had much higher average grain sizes than one would expect when

looking at their neighbors. These two timepoints were the 12 hour treatment and the 300 hour

above solvus treatment.

The as-cast microstructure of Violet was found to be similar to SAC305 in that there was a beta-

tin dendritic structure with interdendritic spaces where IMC and bismuth were found. Lamellar

bismuth structures were observed in as-cast Violet and can be seen in Figure 13. These structures

are the result of bismuth precipitates growing along specific crystallographic orientations due to

the anisotropy of tin. When Violet was aged below solvus, the lamellar structures had a tendency

to ripen into large bismuth particles, as can be seen in Figure 23. These large coarse particles

were examined at higher magnification, and it was found that they were surrounded by an area

depleted of bismuth, indicating the process by which they formed was Ostwald ripening.

42

Figure 23 – As-cast Violet (left), below solvus aged Violet (right)

It was found that there was some minor penetration of bismuth particles into the dendritic spaces

when as-cast and sub solvus aged samples were compared. This is likely due to local changes in

composition; in as-cast samples significant bismuth segregation was observed and as such some

areas may have had a lower solvus temperature, promoting the redistribution of bismuth instead

of coarsening. This effect was minor, especially when compared to the above solvus aged

samples.

Above solvus aging had the opposite effect; there were significantly fewer coarse particles in the

two short-duration treatments when compared to the sub solvus aging and there was significant

penetration of bismuth into the dendritic areas. The 12 hour treatment showed less bismuth

mobility than the 24 hour treatment. The 300 hour treatment was extremely uniform. This is a

result of bismuth dissolving into the beta-tin matrix when held at a temperature above solvus and

tending towards an equilibrium bismuth concentration with time. The longer the treatment, the

more time bismuth has to diffuse to areas of low bismuth concentration. When the sample was

cooled, bismuth stopped diffusing to areas of low concentration and began to precipitate out of

solid solution. This effect is best seen by the degree to which bismuth particles can be observed

in the otherwise bismuth poor dendritic areas.

43

During testing, it was found that bismuth had an affinity for free surfaces and other high energy

areas. This is best seen in micrographs of improperly prepared samples, such as in Figure 24

below. Bismuth was found to readily precipitate on scratched areas, even when the scratch itself

was so minor that it could not be seen in SE imaging.

Figure 24 - The effect of otherwise invisible scratches on bismuth precipitation on the

observed surface (arrows indicate a scratch)

Bismuth precipitates were commonly found on triple junctions of grains (another high free

energy area), and prior literature had found evidence of bismuth precipitating onto the surface of

solder joints after accelerated thermal cycling [6]. Bismuth precipitation was found to introduce

internal strain significant enough to promote grain boundary sliding on the surface. Evidence of

this can be seen in Figure 13.

44

The grain size of as-cast Violet was found to be significantly more fine than that of SAC305.

Due to the coarse nature of most bismuth particles in the as-cast condition (greater than 1

micron), it is speculated that as bismuth precipitates form immediately following solidification

they introduce strain into the lattice surrounding them. This causes particle stimulated nucleation

of recrystallized grains as can be seen in Figure 25.

Figure 25 - Deformation introduced by the growth of a second phase particle [68]

The grain size of Violet was found to be relatively stable when aging below solvus. The reason

no significant grain growth was observed is due to the Zener pinning effect bismuth particles

have on grain boundary motion [47]. When aged above solvus, the grain size of Violet was found

to decrease. The cause of this is recovery and formation of subgrains (polygonization) followed

by subgrain rotation recrystallization, all driven by the internal strain of bismuth precipitates

forming as the alloy cools from above solvus temperatures [68].

From the results of the nano-indentation study, it can clearly be seen that aging above solvus has

a very positive effect on the mechanical properties while aging below solvus did not negatively

45

impact the mechanical properties of Violet. This can be attributed to both improving bismuth

particle distribution, and consequently improving the effect of Zener pinning, and the decrease in

grain size that is experienced. The solid solution strengthening and precipitation hardening

effects of bismuth mean that even as-cast and sub solvus aged Violet outperforms SAC305 in

hardness and creep as was demonstrated. It was found that the 24 hour above solvus treatment

was comparable to the 300 hour treatment in terms of effect on mechanical properties.

In Figure 26, it can be seen that the creep rate has a tendency to reduce with respect to grain size.

The relationship appears to be linear, which implies that Coble and Nabarro-Herring creep are

not the active creep mechanisms, despite the high homologous temperature of the system. This is

likely the result of significant Zener pinning preventing the motion of grain boundaries, making

these mechanisms energetically unfavorable. The apparent relationship creep rate has with

respect to grain size seen in Figure 26 is quite likely actually a function of the uniformity of

bismuth particle distribution. Lower grain size directly corresponds to longer duration aging

above solvus, which in turn is related to bismuth distribution and improved Zener pinning. This

indicates that Harper-Dorn creep is possibly the dominant creep mechanism. This mechanism

has been observed in tin systems in the past, and is dislocation climb controlled [69]. When a

pinned dislocation is unable to continue moving, it can climb out of the glide plane if a vacancy

diffuses to it and resume its motion. Harper-Dorn creep is limited by the diffusion of vacancies,

but considering the high homologous temperature, diffusive creep should be the active creep

mechanism [51], [62].

46

Figure 26 – Average grain size plotted against creep rate for all timepoints

The hardness of Violet was found to stay stable across all timepoints, and outperformed the

hardness of SAC305 significantly, which saw a major drop in hardness in all aging conditions

after as-cast.

As-cast

300h 70°C

300h 70°C + 12h 120°C

300h 120°C + 24h 120°C

300h 120°C

0

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

2 2.5 3 3.5 4 4.5 5

Cre

ep R

ate

(nm

/s)

d (µm)

Creep Rate with Respect to Average Grain Size

47

The current industry standard for testing the reliability of solders is accelerating thermal cycling

(ATC). Here, assemblies are subjected to rapid temperature fluctuations which induce significant

thermal stresses into solder joints. The typical temperatures range from below zero to over

100°C. With the knowledge that above solvus aging has a restorative effect on the microstructure

and mechanical properties of bismuth containing solder alloys, one has to consider that the

standards defined by the industry for testing current generation solders are no longer applicable.

In bismuth containing alloys, the restorative treatment is applied in every cycle, while SAC

alloys experience microstructural degradation. While ATC does indeed test the performance of

both alloys in fatigue, it decouples fatigue performance from microstructural degradation for new

generation bismuth containing solders. This explains why they have consistently been shown to

outperform SAC alloys in ATC, when the testing is unequal. This highlights the importance of

combined testing, which has lower cycling temperatures but introduces vibration to promote

earlier failure, as ATC in sub solvus temperatures would require years of testing to produce

failures and as such data on solder alloys.

48

Chapter 4 Conclusions

Summary of Conclusions

This thesis aimed to study new generation lead-free solder alloys. A major difference with this

new generation of lead-free solders is the introduction of second phase precipitation hardening

into the system. The effect of aging in above and below solvus regimes has not been studied in

much depth in existing literature but is of great importance to the microstructure of the solder

alloys, and consequently their performance.

A novel DSC methodology was developed for the detection of low energy solid-state phase

transformations. This method leveraged the increased signal from high mass samples and high

heating rates to produce a detectable peak. The thermal lag introduced by using these parameters

was minimized where possible, both by the use of built in software correction and by minimizing

thermal resistance by directly casting samples into examination pans. Using this methodology,

the solvus temperatures for three alloys were characterized. The experimental solvus

temperatures were found to be 91°C, 83°C and 42°C for Violet, Sunflower and Senju

respectively. These values were validated using metallography, and determined to lie between

80°C to 100°C for Violet and 70°C to 90°C for Sunflower and between 30°C and 50°C for

Senju.

Using the solvus temperatures that were previously determined, the microstructural changes that

these alloys experience during aging above and below solvus were characterized. Violet (Sn-

2.25%Ag-0.5%Cu-6.0%Bi) was used as an exemplar alloy. The microstructure was examined in

terms of grain size and bismuth precipitates.

49

Bismuth particle precipitation and redistribution were studied in Violet. The behavior of these

particles was found to be highly dependent on the temperature of aging. In as-cast samples,

bismuth was found to be segregated to interdendritic areas and formed large, coarse particles

with low uniformity. When samples were aged below solvus, there was significant evidence of

Ostwald ripening, and the particle size uniformity was found to decrease slightly. Bismuth

particles were not observed penetrating significantly into dendritic areas when aged below

solvus. Bismuth particles were found to dramatically increase in uniformity and distribution

when aged above solvus. This is a result of the dissolution and re-precipitation of bismuth when

the solvus temperature is passed in heating and cooling respectively. Bismuth was found to

penetrate deeper into dendritic areas with increased time spent aging above solvus.

The grain size of Violet was found to be relatively stable when aged below solvus, likely due to

the grain boundary pinning effects of bismuth acting as a barrier to grain growth. The grain size

of Violet was found to decrease when aged above solvus. The recovery and subsequent

nucleation of recrystallized grains, or the recrystallization of recovered subgrains by subgrain

rotation, is responsible for the reduction in grain size of above solvus aged Violet. The driving

force for this process is the introduction of internal strain by the precipitation of bismuth

particles from solid solution (particle stimulated nucleation).

A nano-indentation study was conducted to test the creep properties and hardness of Violet when

compared to SAC305. It was found that below solvus aging did not have a statistically

significant effect on creep resistance when compared to as-cast creep rate. The creep rate of

Violet was shown to decrease when aged above solvus. It was observed that Coble and Nabarro-

Herring creep were not the dominant mechanisms, indicated by a proposed independence from

grain size, but that Harper-Dorn dislocation climb controlled creep was likely the active

50

mechanism. The hardness of Violet was found to stay stable with both above and below solvus

aging, while the hardness of SAC305 decreased considerably with aging.

The above solvus aging treatments explored in this work were found to be effective ways of

restoring and even improving the mechanical properties of Violet aged at sub solvus

temperatures (simulating service conditions). These treatments were found to improve

microstructural uniformity, namely the redistribution of bismuth, which in turn improved

hardness and creep resistance.

Future Work

While a number of interesting properties of bismuth containing solder alloys were characterized,

there are still significant gaps in our knowledge. It would be of interest to see if the

recrystallization behavior that SAC alloys experience ahead of a fatigue crack is present in

bismuth containing alloys, and whether with the Zener pinning of bismuth this is in fact a

beneficial effect that slows the growth of these cracks. The particle size distributions revealed

that the majority of bismuth particles were smaller than what could be observed. It would be of

interest to find the true peak to be able to characterize the Zener pinning effect in more detail.

There is still significant work to be done on optimizing the restorative, above solvus aging

treatments before they could be adopted by industry. It is of interest to find the optimal duration

of the treatment such that minimal thermal damage is introduced to the assembly. It is possible

that the time spent at temperature during one such treatment could be reduced with thermal

cycling, but this thesis only examined isostatic thermal aging. It would also be of interest to

determine when “overripening” of bismuth particles occurs to optimize when to apply restorative

treatments on in-service assemblies.

51

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