ADHESION AND RELIABILITY OF SOLAR MODULE MATERIALS A …td803kh9320/All-Chapters-v21-E... · then...

ADHESION AND RELIABILITY OF SOLAR MODULE MATERIALS

A DISSERTATION

SUBMITTED TO THE DEPARTMENT OF

MATERIALS SCIENCE AND ENGINEERING

AND THE COMMITTEE ON GRADUATE STUDIES

OF STANFORD UNIVERSITY

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS

FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

FERNANDO NOVOA

JUNE 2015

http://creativecommons.org/licenses/by-nc/3.0/us/

This dissertation is online at: http://purl.stanford.edu/td803kh9320

© 2015 by Fernando Daniel Novoa Perez. All Rights Reserved.

Re-distributed by Stanford University under license with the author.

This work is licensed under a Creative Commons Attribution-Noncommercial 3.0 United States License.

ii



http://purl.stanford.edu/td803kh9320

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

Reinhold Dauskardt, Primary Adviser


Michael McGehee


Alberto Salleo

Approved for the Stanford University Committee on Graduate Studies.

Patricia J. Gumport, Vice Provost for Graduate Education

This signature page was generated electronically upon submission of this dissertation in electronic format. An original signed hard copy of the signature page is on file inUniversity Archives.

iii

iv

ABSTRACT

Debonding of solar module laminates (encapsulations, backsheets and frontsheets) is the

least understood failure mode in the PV industry. It is well known, however, that long-

term exposure to environmental stressors, including moisture, temperature and UV light,

often causes laminate degradation and loss of adhesion. Low levels of mechanical stress

(thermal, wind, or handling) can then lead to module debonding, component corrosion

and failure. Although debonding of solar module materials is frequently reported, it is

still not well-understood or quantified.

Solar module adhesion has been previously studied using only peeling-based techniques,

however, peel-strength is not an interfacial property and cannot be used to compare

between materials. In particular, peel-strength results are strongly convoluted with the

macroscopic plastic deformation of the peeled material, which is difficult to quantify and

unrelated to the interfacial debonding processes observed in-field.

In this dissertation, I report a series of fracture mechanics-based techniques to accurately

quantify interfacial adhesion of encapsulations, backsheets and frontsheets. Cantilever-

beam and cantilever–plate techniques were developed for full-size modules and small

lab-made specimens. These metrologies allow the use of an absolute scale (J/m2) to

quantify degradation of module materials after field or simulated exposures.

To measure backsheet debond energy, a novel single cantilever beam method using

acrylate beams was developed. To measure encapsulation debond energy, a square-corner

metrology—where debonding occurs at a constant force—was designed. To measure

frontsheet debonding, a double cantilver beam technique was used in an environment of

controlled temperature (T), relative humidity (RH), and UV light. The techniques

presented here can be easily adapted to characterize adhesion of other module laminates

and interfaces.

To measure debond growth kinetics, a load-relaxation technique was applied to the single

and double cantilever beam specimens. The debonding rates (as low as 10-9 m/s) of

selected module materials were measured as a function of T, RH, and the applied

mechanical stress (e.g. thermal, wind or handling). The debond growth rates increased up

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to 1000 times with small increases of T (10°C) and RH (15%). To understand the

physical mechanisms of debonding, a debond-reaction kinetics framework is introduced.

In this framework, the value of the strain energy release rate, G, is analyzed as the

driving-force for the debonding reaction, (similar to the Gibbs free energy for a chemical

reaction). In the debonding reaction, the molecular relaxation and bond-rupture kinetics

processes—which depend on temperature and relative humidity—determine the value of

the debond growth rate. The model can used to predict module lifetime using a defect-

tolerant approach.

Selected module materials and interfaces were characterized. The adhesion of a

Tedlar/PET/EVA backsheet decreased dramatically from 1000 to 27 J/m2 within the first

750 hrs of exposure to hot (85C) and humid (85% RH) environments. Adhesion of a

poly(ethylene-co-vinyl acetate) (EVA) encapsulation decreased slightly from 2.15 to 1.75

kJm-2 when T increased from 25 to 50C, but decreased precipitously to ~0.35 kJ/m2 at

T>55C. The debond growth rates of the encapsulations and backsheets increased up to

1000 times with small increases of T (10C) and RH (20%). Frontsheet debond growth

rate increased up to four orders of magnitude under low levels (~5mW/m2) of UV light.

The metrologies and models presented here can be used to improve and develop

encapsulation materials and interfaces, for example by optimizing surface preparation

procedures or by improving lamination methods. The techniques can also be used to

quantify loss of adhesion after field or accelerated exposures—directly on full-size

modules or on small encapsulation specimens. The debond-kinetics models can be used

to develop accelerated testing protocols in order to estimate module lifetime.

vi

PREFACE AND ACKNOWLEDGEMENTS

I am thankful to all my professors, at Stanford and ITESM, who had a deep influence in

my education. I am especially thankful to Professor Reinhold Dauskardt—my Ph.D.

adviser, mentor and boss. His scientific expertise and keen precision to forecast

technology have changed my academic style, enriched my educational experience, and

directed my contribution at Stanford. I am also thankful to Professors Michael McGehee

and Alberto Salleo for educating me in photovoltaics and thermodynamics and for their

input in this document. I am also in debt with Dr. Evelin Sullivan, my instructor from the

Technical Communication Program, who edited some of my papers and introduced me to

the details of technical writing. I also want to acknowledge Professor William Nix, who

served as my faculty resource adviser for the Stanford Graduate Fellowship (DARE) that

I received. I would also like to thank Professors Sheri Sheppard, Bruce Clemens, Bob

Sinclair, Yi Cui, and Adrian Bejan (Duke) for their advice and advocacy during the Ph.D.

program. I also appreciate the mentorship of Dr. Alexander Tumanov and Professors

Charles Cox and Leslie Field-Barth during the Ph.D. program. The earlier mentorship of

then graduate student Kanzan Inoue, and Professors Ray H. Baughman, Doroteo

Mendoza, Arturo Barba, Pedro Ponce and Arturo Ponce are also much appreciated.

A large fraction of my Ph.D. work was done in collaboration with David Miller and Nick

Bosco from NREL. They provided a number of important research questions, materials

systems, processing and feedback. Due to the applied nature of the research presented

here, collaborations with the solar materials industry, including SunPower, Dupont, and

3M, were also fundamental to obtain modules, materials and funding. I am especially

thankful to President David Ranhoff for recruiting me into SunEdison.

My research was also heavily influenced by most members of the Dauskardt group,

especially by former graduate student Ani Kamer. The younger graduate students that I

trained (mostly Scott and Warren) also enriched my educational experience at Stanford.

My research interns, Ricardo Corona, Scott Takahashi, Karl Bayer, Olgaby Martinez and

Vi Le, have helped me clarify my hypothesis and improve my teaching. They also

partially contributed with some of the research findings presented in this dissertation.

vii

The mechanical properties of interfaces are complex and intriguing. Ensuring the long-

term mechanical stability of solar module materials and interfaces is a more intriguing

challenge. The research presented in this Ph.D. dissertation resulted in the development

of metrologies and models that can be used to characterize the long-term reliability of

solar modules. Because solar cells are a truly fantastic energy asset only if they last for

long periods of time, I truly hope this dissertation can inspire, guide or complement the

reader’s efforts in industry or academia to design long-lasting solar energy technologies.

viii

TABLE OF CONTENTS

ABSTRACT ...................................................................................................................... IV

PREFACE AND ACKNOWLEDGEMENTS .................................................................. VI

TABLE OF CONTENTS ............................................................................................... VIII

LIST OF FIGURES .......................................................................................................... XI

CHAPTER 1: INTRODUCTION ....................................................................................... 1

1.1 Solar module materials ........................................................................................... 1

Failure modes ............................................................................................. 2

Materials selection ...................................................................................... 3

Loss of adhesion .......................................................................................... 4

1.2 Dissertation outline ................................................................................................. 6

1.3 References ............................................................................................................... 7

CHAPTER 2: THEORETICAL BACKGROUND: FRACTURE MECHANICS ........... 10

2.1 Griffith Fracture Theory ....................................................................................... 10

2.2 Strain Energy Release Rate, G .............................................................................. 12

Stress intensity factor, K ........................................................................... 12

Crack tip plasticity .................................................................................... 13

Equivalence of G and K ............................................................................ 14

2.3 Interfacial Fracture Energy ................................................................................... 14

DCB specimen mechanics ......................................................................... 15

Kanninen calibration ................................................................................ 17

Penado calibration.................................................................................... 17

2.4 Debond Growth Kinetics ...................................................................................... 18

Bond rupture kinetics ................................................................................ 18

Viscoelastic debond growth ...................................................................... 19

Debond growth metrology ........................................................................ 19

2.5 References ............................................................................................................. 22

CHAPTER 3: ENCAPSULATION ADHESION AND DEBONDING .......................... 23

3.1 Chapter Summary ................................................................................................. 23

3.2 Introduction ........................................................................................................... 23

3.3 Materials and Methods .......................................................................................... 26

Encapsulation SCB debond specimens ..................................................... 26

SCB debond test metrology ....................................................................... 28

SCB debond test displacement rate ........................................................... 30

Debond growth rate kinetics ..................................................................... 31

Tensile specimens...................................................................................... 32

ix

3.4 Results and Discussion ......................................................................................... 33

Environmental sensitivity of EVA modulus ............................................... 33

Mechanical compliance of SCB debond specimen ................................... 34

Encapsulation debond energy ................................................................... 35

Environmental debonding kinetics ............................................................ 37

Debond Kinetics Model............................................................................. 41

3.5 Conclusions ........................................................................................................... 45

3.6 References ............................................................................................................. 47

CHAPTER 4: SQUARE-CORNER ADHESION METROLOGY .................................. 52

4.1 Chapter Summary ................................................................................................. 52

4.2 Introduction ........................................................................................................... 52

4.3 Experimental Setup ............................................................................................... 54

Materials and adhesion coupons .............................................................. 54

Adhesion Metrology .................................................................................. 55


4.5 Metrology variations ............................................................................................. 61

Double Cantilever beam ........................................................................... 62

Single substrate configurations ................................................................ 62

4.6 Conclusions ........................................................................................................... 63

4.7 References ............................................................................................................. 65

CHAPTER 5: BACKSHEET ADHESION AND DEBONDING .................................... 68

5.1 Chapter Summary ................................................................................................. 68

5.2 Introduction ........................................................................................................... 68

5.3 Materials and Methods .......................................................................................... 70

Backsheet specimens and materials .......................................................... 70

Aging treatment methodology ................................................................... 71

Single cantilever beam test metrology ...................................................... 72

Debond energy experiment ....................................................................... 73

PVF deformation behavior ....................................................................... 73

Environmental debond growth .................................................................. 73


Backsheet debond energy .......................................................................... 74

PVF deformation behavior ....................................................................... 76

Environmental debonding ......................................................................... 77

Fracture Kinetics Model ........................................................................... 79

5.5 Conclusions ........................................................................................................... 83

5.6 References ............................................................................................................. 85

CHAPTER 6: FRONT-SHEET DEBONDING UNDER UV LIGHT ............................. 89

6.1 Chapter Summary ................................................................................................. 89

6.2 Introduction ........................................................................................................... 89

x

6.3 Experimental Procedures ...................................................................................... 92

Materials and Coatings............................................................................. 92

Crack and debond growth testing ............................................................. 93

UV and Environment accelerated debonding ........................................... 93

Failure path characterization ................................................................... 94


Moisture Sensitivity ................................................................................... 94

UV sensitivity ............................................................................................ 97

UV-Crosslinking ..................................................................................... 101

6.5 Conclusions ......................................................................................................... 102

6.6 References ........................................................................................................... 104

xi

LIST OF FIGURES

Figure 1-1: Solar module structure. .................................................................................... 1

Figure 1-2: Debonding of module laminates. ..................................................................... 2

Figure 1-3: Front sheet mechanical failure. ........................................................................ 3

Figure 2-1: Infinitely wide plate ....................................................................................... 11

Figure 2-2: Stress-state ahead of a crack tip ..................................................................... 13

Figure 2-3: Three characteristic modes of loading ........................................................... 13

Figure 2-4: Schematic diagram of double cantilever beam specimen geometry .............. 15

Figure 2-5. Representative DCB data ............................................................................... 17

Figure 2-6. Crack velocity as a function of applied crack driving force. ......................... 20

Figure 2-7. Load versus time for a subcritical cracking test ............................................. 20

Figure 2-8. Crack length versus time for a subcritical cracking test ................................ 21

Figure 3-1: Single Cantilever Beam Specimen (SCB) for encapsulant debonding. ......... 27

Figure 3-2: Debond-tip opening displacement rate, ......................................................... 31

Figure 3-3: Elastic (E) and storage (E’) modulus of the EVA encapsulation ................... 33

Figure 3-4: Mechanical compliance, C, of EVA and PVB specimens (SCB), ................. 35

Figure 3-5: Debond Energy of EVA/glass and PVB/Ti interfaces as a function of T. ..... 37

Figure 3-6: Debond growth rate of EVA/Glass and PVB/Ti ............................................ 38

Figure 3-7: Debond growth rate of EVA/Ti ..................................................................... 39

Figure 3-8: Debond growth rate of EVA/glass subject to an 8°C T gradient, .................. 40

Figure 4-1: Adhesion coupon loaded in tension by a corner. ........................................... 55

Figure 4-2: Mechanical compliance of the adhesion coupon ........................................... 57

Figure 4-3: Glass substrate stiffening factor, f. ................................................................. 58

Figure 4-4: Load-displacement curve for EVA/glass debonding ..................................... 59

Figure 4-5: Adhesion (Gc) of the EVA encapsulation bonded to glass ............................ 61

Figure 4-6: Adhesion metrology variations. ..................................................................... 64

Figure 5-1: Backsheet-Encapsulant-Glass specimen structure. ........................................ 71

Figure 5-2: Backsheet debond energy vs. aging treatment temperature ........................... 74

Figure 5-3: PVF-PET debond path characterization. ........................................................ 75

Figure 5-4: Backsheet debond energy vs. duration of aging treatment ............................ 76

Figure 5-5: Stress-strain curves of the PVF film .............................................................. 77

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Figure 5-6: Thermally-activated backsheet debond growth rate ...................................... 78

Figure 5-7: Moisture-assisted backsheet debond growth rate........................................... 79

Figure 6-1: Schematics and photograph of debond growth experiment ........................... 92

Figure 6-2: Adhesive debond growth rate of a polysiloxane/borosilicate interface ......... 94

Figure 6-3: ATR-FTIR spectra showing Si-O-Si, Si-OH and Si-CH3 .............................. 95

Figure 6-4: Cohesive crack growth rate of polysiloxane .................................................. 97

Figure 6-5: Crack growth rate in the polysiloxane film .................................................... 98

Figure 6-6: Debond growth rate of a polysiloxane/glass interface ................................... 99

Figure 6-7: Debond growth rate of a polysiloxane/quartz interface ............................... 100

Figure 6-8: Cohesive crack growth rate of in polysiloxane ............................................ 101

Figure 6-9: FITR spectra of polysiloxane film ............................................................... 102

1

1 CHAPTER 1: INTRODUCTION

Solar modules are designed to protect the solar cell from field environments and

mechanical stress[1,2]. However, preserving adhesion of module materials has been

challenging, especially in hot and humid operating environments[3,4]. Long-term

exposure (>20 years) to environmental stressors, including moisture, temperature and UV

light, often causes degradation of materials and interfaces. Low levels of mechanical

stress (thermal, wind, or handling) can then lead to module debonding, component

corrosion, and failure. Although debonding of solar module materials is frequently

reported, it has yet to be characterized, understood or quantified.

Figure 1-1: Solar module structure. The cell array is encapsulated between two elastomer layers.

A glass sheet and an insulating backsheet are bonded to each side of the encapsulation. Solar

modules protect the solar cell from the environment and mechanical stresses. Image courtesy of

Dupont.

1.1 Solar module materials

The solar module structure consists of laminated layers of moisture-impermeable, load-

bearing, and discharge-insulating materials. Inside the module, the active solar cell array

is bonded between two elastomeric encapsulation layers, as shown in Fig. 1-1. The front

encapsulation layer is also bonded to glass, or alternatively, to a moisture-barrier film

consisting of intercalated organic/inorganic layers. The back encapsulation layer is

bonded to a multilayered backsheet, and the backsheet is bonded to a junction box. All

Frame

Glass

Encapsulation

Solar Cells

Encapsulation

Backsheet

Junction Box

2

the module layers are then sealed around the edges and glued to a metallic frame. The

main function of the solar module is to isolate the solar cells from the field environment

in order to prevent corrosion. Mechanically, the module also protects the solar cell from

wind, handling, thermal and impact stresses[2]. Without the module, the solar cell would

rapidly corrode and fracture.

Failure modes

Solar module adhesive failure has been extensively reported [21]. In the presence of

aggressive environmental stressors, the encapsulation layer often debonds from the glass

substrate, as shown in Fig. 1-2a. Debonding of multilayered PV backsheets is also

frequently observed. In the most common backsheet structure (Tedlar-PET-EVA), failure

occurs at the interface between the Tedlar(PVF) and PET layers. In hot environments,

backsheet blistering can occur due to acetylene gas formation during encapsulation

(EVA) thermal degradation (Fig. 1-2b).

2a 2b

Figure 1-2: Debonding of module laminates. Debonding of glass/encapsulation (2a) and

backsheet/encapsulation (2b) can occur in the presence of active environments and mechanical

stress.

Glass fracture is also a common failure mode due to handling, wind, thermal or impact

stresses (Fig. 1-3a). When glass is replaced with a multilayered moisture-barrier film[22],

the film frequently delaminates internally at very low values of mechanical stress (Fig. 1-

3b). Front sheet mechanical failure (glass or barrier) often leads to faster pathways for

moisture diffusion, increased series resistance, and eventually loss of module

function[23]. Preserving front sheet adhesion is important not only to guarantee the long-

term performance of the leading silicon technologies, but also to allow for the

3

implementation of thin-film and organic cells that are much more sensitive to moisture

ingress.

3a 3b

Figure 1-3: Front sheet mechanical failure. PV glass fracture (3a) and barrier delamination (3b).

Both failure modes allow moisture ingress, which corrodes the active cell.

Materials selection

Silicon solar modules currently dominate the solar energy industry[5] but other cell

technologies, including thin-film organic[6] and inorganic[7], represent potential cheaper

substitutes in current and evolving markets. To house newer cell technologies, however,

improved module design and materials-selection are still required [8]. Ensuring the

interfacial adhesion of potential candidate materials is necessary to minimize production

failure and to improve field reliability. Even in conventional cell technologies, improving

module materials is necessary to further reduce module costs and to determine optimal

module fabrication procedures. In the near future, module materials will be selected to

match field operating conditions, in order to provide the most cost-effective—yet

reliable—solar energy supply.

Due to the long-term operating requirements, module materials selection is by itself

complicated, but determining the reliability of the interfaces created between new

materials is an even more significant challenge. While the solar material manufacturers

(Dupont, Dow, 3M, etc) do optimize the bulk properties of their module materials, the

interfaces created during module lamination are often completely new, especially when

the module structure has been redesigned. Building integrated photovoltaics (BIPV), for

example, requires the use of completely new substrates, where interfacial strength has

never been studied[15].

4

Most current module materials have been grandfathered-in from the 70s[9]. The most

commonly used solar module encapsulation, EVA, lacks moisture impermeability and

resistance to yellowing, and undergoes a phase transformation at operating field

temperatures[11]. Although several novel module materials have already been developed,

very few have been adopted in large-scale field applications due to reliability concerns.

Although the current rate of failure in Si-modules (~0.5 to 0.75% per year) is not by itself

a show-stopper for the solar industry, unpredictable failure is still a large source of

inefficiency. For example, inadequate glass sheet transportation has been reported to

cause 3% field-failure in several batches of a CdTe module manufacturer. Unpredictable

module failure also increases operating costs, complicates risk management[12], and

reduces the bankability of the solar project. The performance and reliability of current Si-

module materials has been mostly characterized in terms of failure occurrences[10], but

the physical mechanisms behind failure are still not well understood. Specifically, the

mechanisms leading to loss of adhesion and debonding in solar modules are still

completely unknown.

Loss of adhesion

Environmental loss of adhesion of module materials is a complex process. Bond rupture,

polymer chain scission, and phase segregation can occur in the bulk of the material[17].

Moisture ingress, lattice mismatch and interface-selective corrosion can weaken the

interfacial bond[18]. The multilayered structures used in solar modules often require

specialized bonding protocols, surface treatments, proprietary epoxies, and even the use

of intermediate “tie” layers to bond dissimilar materials. The hybrid organic/inorganic

bond—such as the one formed between the encapsulation layer and the glass (or Si)

substrate—has always been a source of major instability[19]. Hybrid interfaces are

usually susceptible to both the degrading stressors of organic and inorganic materials. In

organic materials, moisture and temperature, together with cycling loading, leads to

interfacial degradation. In inorganic materials, moisture-assisted cracking and brittle

fracture cause interfacial failure[20]. In other semiconductor technologies, ensuring the

longevity of the hybrid organic/inorganic interfaces has been a serious challenge. In solar

modules, adhesion of hybrid interfaces has yet to be quantified.

5

The adhesion of module materials is a function of both the strength of the interfacial bond

and the bulk material processes (deformation, cavitation, etc.) that occur during

debonding. Optimizing the strength of the interfacial bond is pointless when the bulk

properties degrade rapidly in the operating environment. Conversely, poor interfacial

bonding reduces the potential contribution of the bulk processes to the overall adhesion

value[16]. While the available mechanics tests (peel or shear tests) can be related the bulk

properties of the material, there is currently not a mechanics-based technique to measure

the adhesive strength of the interfacial bond. Currently available metrologies that

measure either a debonding force or a debonding stress are fundamentally incorrect.

Adhesion can only be characterized by measuring the critical strain energy release rate of

debonding, Gc (J/m2). The research presented in this Ph.D. dissertation resulted in the

development of metrologies and protocols that allow the accurate measurement of Gc in

solar module materials.

Solar module materials are exposed to very aggressive environments. Operating

conditions include cycles of temperature and humidity, and changing levels of

mechanical stress due to wind[13], handling or thermal mismatch[14]. The effect of the

environment on module material adhesion has been mostly studied by replicating tests

and sequences from the module qualification protocols[24,25]. The UV exposure, thermal

cycling, and humidity freeze tests simulate a brief field-exposure (months to a few years

at best), whereas the conditions applied in damp heat test greatly exceed those

encountered in the application[10]. The qualification tests at best query failure modes

related to module infant mortality; they do not investigate the rate of degradation. There

is currently not a single physics or chemistry-based argument that ensures that the

qualification protocol is time-scalable. Predicting lifetime and ensuring long-term

performance module materials is still a serious challenge. The currently available aging

treatments, as well as the available adhesion tests, provide limited insight on the kinetic

processes of debonding and their dependence on mechanical stress, moisture and

temperature. Kinetics-based testing protocols are therefore urgently needed not only to

assess encapsulation durability in field environments but also to define objectives for

encapsulation formulation and manufacturing. The research presented in this Ph.D.

dissertation resulted in the development of metrologies and models to study module

6

debonding-kinetics. The metrologies can be used to estimate defect-tolerant module

lifetimes, and the models constitute the basis to design time-scalable, accelerated-aging

qualification protocols.

1.2 Dissertation outline

The rest of this dissertation is organized as follows. Chapter two, (theoretical

background), briefly describes the fracture mechanics concepts and methods to

characterize interfacial adhesion and debonding-kinetics. In this chapter, debond-growth

is rationalized using a reaction-kinetics framework. Chapter three (encapsulation) extends

the methods presented in chapter two to account for the elastomeric behavior of the

encapsulation materials, the applied loading rates, and the temperature and relative

humidity of the environment. Chapter four (square-corner test) describes a simplified

version of the metrologies developed in chapter three. Chapter five (backsheets) describes

a newly-developed single cantilever beam method to quantify backsheet adhesion and

debonding kinetics, which can be used directly on full-size solar modules and on small

specimens. In this chapter, a viscoelastic fracture-mechanics model is derived to elucidate

the physical mechanisms of backsheet debonding. Chapter six (frontsheets) uses the

methods derived in the previous chapters to characterize debonding kinetics in the

presence of UV light. The effect of the environment is investigated and modeled using a

bond-rupture kinetics approach that includes dependences on temperature, moisture and

UV. The techniques and models presented in this Ph.D. dissertation can be used to

develop accelerated testing protocols, evaluate the adoption of new encapsulations, and

make long-term reliability predictions.

7

1.3 References

[1] P. Sánchez‐Friera, Analysis of degradation mechanisms of crystalline silicon PV modules

after 12 years of operation in Southern Europe, Prog. …. (2011) 658–666.

doi:10.1002/pip.

[2] M. Quintana, D. King, T. McMahon, Commonly observed degradation in field-aged

photovoltaic modules, Photovolt. Spec. Conf. 2002. Conf. Rec. Twenty-Ninth IEEE.

(2002) 1436–1439. http://ieeexplore.ieee.org/xpls/abs_all.jsp?arnumber=1190879

(accessed April 4, 2012).

[3] G. Jorgensen, K. Terwilliger, J. Delcueto, S. Glick, M. Kempe, J. Pankow, et al., Moisture

transport, adhesion, and corrosion protection of PV module packaging materials, Sol.

Energy Mater. Sol. Cells. 90 (2006) 2739–2775. doi:10.1016/j.solmat.2006.04.003.

[4] F.J. Pern, S.H. Glick, Adhesion Strength Study of EVA Encapsulants on Glass Substrates,

Natl. Cent. Photovoltaics Sol. Progr. Rev. Meet. (2003).

[5] I. Photovoltaic, S. Cells, Inorganic Photovoltaic Solar Cells: Silicon and Beyond,

Electrochem. Soc. Interface. 17 (2008) 30–35.

http://scholar.google.com/scholar?hl=en&btnG=Search&q=intitle:Inorganic+Photovoltaic

+Solar+Cells:+Silicon+and+Beyond#0.

[6] T.D. Nielsen, C. Cruickshank, S. Foged, J. Thorsen, F.C. Krebs, Business, market and

intellectual property analysis of polymer solar cells, Sol. Energy Mater. Sol. Cells. 94

(2010) 1553–1571. doi:10.1016/j.solmat.2010.04.074.

[7] Anon, Market trend on solar cells, Kino Zair. 30 (2010) 45–50.

[8] R. Tipnis, J. Bernkopf, S. Jia, J. Krieg, S. Li, M. Storch, et al., Large-area organic

photovoltaic module-Fabrication and performance, Sol. Energy Mater. Sol. Cells. 93

(2009) 442–446. doi:10.1016/j.solmat.2008.11.018.

[9] A. Czanderna, Encapsulation of PV modules using ethylene vinyl acetate copolymer as a

pottant: A critical review, Sol. Energy Mater. Sol. Cells. 43 (1996) 101–181.

http://www.sciencedirect.com/science/article/pii/0927024895001506 (accessed April 4,

2012).

[10] J.H. Wohlgemuth, M.D. Kempe, Equating Damp Heat Testing with Field Failures of PV

Modules, in: Proc. IEEE PVSC, 2013.

[11] F. Pern, A. Czanderna, K. Emery, Weathering degradation of EVA encapsulant and the

effect of its yellowing on solar cell efficiency, Photovolt. Spec. Conf. 1991., Conf. Rec.

8

Twenty Second IEEE. (1991) 557–561.

http://ieeexplore.ieee.org/xpls/abs_all.jsp?arnumber=169275 (accessed April 5, 2012).

[12] M. Elisabeth Pate-Cornell, R.L. Dillon, Success factors and future challenges in the

management of faster-better-cheaper projects: Lessons learned from NASA, IEEE Trans.

Eng. Manag. 48 (2001) 25–35. doi:10.1109/17.913163.

[13] J. Cao, A. Yoshida, P.K. Saha, Y. Tamura, Wind loading characteristics of solar arrays

mounted on flat roofs, J. Wind Eng. Ind. Aerodyn. 123 (2013) 214–225.

doi:10.1016/j.jweia.2013.08.014.

[14] Y. Lee, A.A.O. Tay, Finite Element Thermal Analysis of a Solar Photovoltaic Module,

Energy Procedia. 15 (2012) 413–420. doi:10.1016/j.egypro.2012.02.050.

[15] B. Norton, P.C. Eames, T.K. Mallick, M.J. Huang, S.J. McCormack, J.D. Mondol, et al.,

Enhancing the performance of building integrated photovoltaics, Sol. Energy. In Press,

(2010) -. doi:10.1016/j.solener.2009.10.004.

[16] F.D. Novoa, R.H. Dauskard, Debonding Kinetics of Photovoltaic Encapsulation in Moist

Environment, Rev. Prog. Photovoltaics. (2015).

[17] M. Lane, R. Dauskardt, Plasticity contributions to interface adhesion in thin-film

interconnect structures, J. Mater. Res. 15 (2000) 2758–2769.

http://journals.cambridge.org/production/action/cjoGetFulltext?fulltextid=7998189


[18] R.P. Birringer, R. Shaviv, P.R. Besser, R.H. Dauskardt, Environmentally assisted

debonding of copper/barrier interfaces, Acta Mater. 60 (2012) 2219–2228.

doi:10.1016/j.actamat.2012.01.007.

[19] M. Giachino, G. Dubois, R.H. Dauskardt, Heterogeneous solution deposition of high-

performance adhesive hybrid films, ACS Appl. Mater. Interfaces. 5 (2013) 9891–9895.

doi:10.1021/am403032v.

[20] E.P. Guyer, R.H. Dauskardt, Fracture of nanoporous thin-film glasses., Nat. Mater. 3

(2004) 53–7. doi:10.1038/nmat1037.

[21] P. Klemchuk, M. Ezrin, G. Lavigne, W. Holley, J. Galica, S. Agro, Investigation of the

degradation and stabilization of EVA-based encapsulant in field-aged solar energy

modules, Polym. Degrad. Stab. 55 (1997) 347–365.

[22] C. Charton, N. Schiller, M. Fahland, A. Holländer, A. Wedel, K. Noller, Development of

high barrier films on flexible polymer substrates, in: Thin Solid Films, 2006: pp. 99–103.

doi:10.1016/j.tsf.2005.07.253.

9

[23] D.J. Coyle, Life prediction for CIGS solar modules part 1: Modeling moisture ingress and

degradation, Prog. Photovoltaics Res. Appl. 21 (2013) 156–172. doi:10.1002/pip.1172.

[24] IEC 61646, Thin-film terrestrial photovoltaic (PV) modules - Design qualification and

type approval, Int. Electrotech. Comm. Geneva, 1–81. (2008).

[25] IEC 61215, Crystalline silicon terrestrial photovoltaic (PV) modules—Design

qualification and type approval ", Int. Electrotech. Comm. Geneva, 1–93. (2005).

10

2 CHAPTER 2: THEORETICAL BACKGROUND: FRACTURE

MECHANICS

Fracture mechanics is the field of mechanics concerned with the study of crack

propagation and failure predictions of materials and structures that contain flaws. In

particular, linear elastic fracture mechanics (LEFM) forms the underlying basis employed

to quantify both adhesive debonding and cohesive cracking through the interfaces of solar

module structures. The theoretical background to characterize adhesion debonding using

LEFM is presented in this section.

2.1 Griffith Fracture Theory

The Griffith fracture theory introduced a fundamental connection between three key

variables in fracture mechanics: applied stress, flaw (crack) size, and material fracture

toughness[1]. Griffith developed a thermodynamic approach to explain fracture

phenomena observed in brittle glass materials whereby he considered that the total energy

in a system with a crack may be partitioned into mechanical strain energy supplied by

applied loading, U, and work required to create the crack, W. He proposed that the crack

could form only if such a process caused the total energy of the system to decrease or

remain constant, which is expressed mathematically by

𝑑𝑈

𝑑𝐴+

𝑑𝑊

𝑑𝐴≤ 0

(2-1)

where dA represents an incremental increase in surface created by crack formation and

propagation. In an infinitely wide plate (Fig. 2-1) with thickness, B, containing a through

crack with length 2a and height 2b, U is given by the Inglis stress analysis[2] showing

𝑈 = 𝑈0 −𝜋𝜎2𝑎2𝐵

𝐸 (2-2)

where U0 is the potential energy of the uncracked plane, σ is the applied stress, and E is

the plane strain Young’s modulus of the plate. Since the formation of the crack results in

the creation of two surfaces, W is given by

11

𝑊 = 4𝑎𝑏𝛾 (2-3)

where γ is the surface energy of the plate. By substituting Eq. (2-2) and Eq. (2-3) into Eq.

(2-1) under equilibrium condition and solving for the fracture stress, σc, we obtain

𝜎𝑐 = √2𝐸𝛾

𝜋𝑎

(2-4)

which is the well-known Griffith relation, which relates fracture stress to material

properties and flaw size. Using this relation, Griffith was able to demonstrate good

agreement for measured fracture stresses in glass fibers with various crack sizes, however

the relation severely underestimated the fracture stress of ductile engineering metals and

polymers. Orowan and Irwin independently modified Griffith’s relation to include an

energy dissipation term to account for materials capable of plastic flow, given by

𝜎𝑐 = √2𝐸(𝛾 + 𝛾𝑝)

𝜋𝑎

(2-5)

where γp, is the plastic work per unit area of surface created, which is typically much

larger than γ .

Figure 2-1: Infinitely wide plate with thickness B containing a through crack

having characteristic dimensions of 2a (width) and 2b (height).

2a

2b

σ

12

2.2 Strain Energy Release Rate, G

In 1956, Irwin proposed an energy approach for fracture that is based on the

pioneering work of Griffith. The strain energy release rate, G, is defined as a measure of

energy available for an increment of crack extension, represented mathematically by

𝐺 = −𝑑U

𝑑𝐴

(2-6)

The term rate does not refer to a derivative with respect to time, but rather the rate of

change in potential energy with the crack area. The strain energy release rate can also be

related to load, P, compliance, C, and crack length, a, which are easily quantifiable

through experimental methods, resulting in

𝐺 =𝑃2

2𝐵

𝑑𝐶

𝑑𝑎

(2-7)

The strain energy release rate is unrelated to the details of the crack tip, but rather is

concerned with the energy of the system as a whole[3]. As such, if information of the

crack tip stress state is not desired, the strain energy release rate is a convenient and

intuitive measure to describe and predict failure phenomena. In this dissertation, Eq. (2-7)

plays a central role; it is used repeatedly to develop new adhesion metrologies, to

determine optimal loading rates, and to calculate the values of debond-energy and

debond-driving forces in selected solar module interfaces.

Stress intensity factor, K

The strain energy release rate provides an analytical tool to characterize the global

behavior of a cracked plate, but doesn’t offer any information about the stress and

displacement fields near the crack tip. According to elastic theory, the stresses near the

crack tip (Fig. 2.2) can be defined by

𝜎𝑖𝑗 =𝐾

√2𝜋𝑟𝑓𝑖𝑗(𝜃)

(2-8)

where K is the stress intensity factor and fij(θ) is an angular function that depends on the

stress component and loading mode.

13

Figure 2-2: Stress-state ahead of a crack tip at any given point (r,θ).

Eq. (2-8) implies that the stress goes to infinity at the crack tip and zero away from the

crack tip[4]. The amplitude of the stress singularity is determined by the stress intensity

factor for each loading mode, KI, KII and KIII, corresponding to opening,

in-plane shear and out-of-plane shear, as shown in Fig. 2.3.

Figure 2-3: Three characteristic modes of loading that result in crack propagation. Mode I is

opening loading orthogonal to crack; Mode II is in-plane shear; Mode III is out-of-plane shear.

Crack tip plasticity

According to Equation (2-8), stresses at the crack tip are infinitely high. In real

materials, yielding will occur at some finite stress defined as the yield strength, σys,

resulting in a region near the crack tip that undergoes plastic deformation. The size of the

plastic zone that results can be estimated by calculating the distance away from the crack

tip, ry, where the singular stress is equivalent to the yield stress of the material.

y

r

xθ

14

𝜎𝑦𝑠 =𝐾𝐼

√2𝜋𝑟𝑦

(2-9)

√𝑟𝑦 =𝐾𝐼

𝜎𝑦𝑠√2𝜋

(2-10)

𝑟𝑦 =1

2𝜋(

𝐾𝐼

𝜎𝑦𝑠)

2

(2-11)

Equation (2-11) provides a good first order approximation of the plastic zone size, but

does not take into account that stresses in the plastic zone cannot exceed the material

yield stress. A more accurate second order approximation involving force balances yields

a plastic zone size, rp, which is twice as large as ry:

𝑟𝑝 =1

𝜋(

𝐾𝐼

𝜎𝑦𝑠)

2

(2-12)

Equivalence of G and K

The strain energy release rate quantifies the net change in potential energy that

accompanies an increment of crack extension while the stress intensity factor

characterizes the stresses, strains and displacements near the crack tip[3]. The former

describes global behavior while the latter is a local parameter. For isotropic linear elastic

solids, K and G are uniquely related by

𝐺 =𝐾𝐼

2

𝐸′+

𝐾𝐼𝐼2

𝐸′+

𝐾𝐼𝐼𝐼2

2𝜇

(2-13)

When describing fracture, it is often convenient to describe toughness in terms of either

energy or stress, depending on the mechanism. The fact that K and G are unique related

means that they can effectively be used interchangeably.

2.3 Interfacial Fracture Energy

In this dissertation, debonding is studied as an interfacial crack propagating between two

materials[5]. The most common specimen configuration to measure interfacial fracture

15

energy is the double cantilever beam specimen (DCB) described below. Other specimen

configurations are described in detail in Chapter 3 through 6.

DCB specimen mechanics

Using the definition of strain energy release rate (Equation (2-7)), the value of G for the

DCB geometry shown in Fig. 2.1 can be calculated using simple beam theory[6].

Figure 2-4: Schematic diagram of double cantilever beam specimen geometry. The solar module

material is sandwiched between two elastic beams (Si, PMMA, metal, etc).

Each cantilever is treated as a fixed beam under a point load with a length equal to the

crack length,

3

2 3

Pa

EI

(2-14)

where is the total DCB deflection, P is the applied load, a is the crack length, E is the

Young’s modulus of the beam, and I is the moment of inertia of the beam given by

3

12

BhI (2-15)

where B is the specimen thickness and h is the beam height as depicted in Fig. 2.1.

The total compliance is then,

3

3

8Beam

aC

P Bh E

(2-16)

Thus according to Equation (2-7),

16

2 2

2 3

12Beam

P aG

B h E (2-17)

where C=/P is calculated from the slope of the load-displacement experimental data.

Multiple measurements of the fracture energy can be made on a single DCB specimen by

repeatedly loading and unloading the beam while extending the crack slightly during each

cycle. A representative plot of load versus displacement for a DCB test is shown in

Figure 2-5. The initial portion of the loading curve is linear and corresponds to elastic

deformation of the “cantilever” sections of the silicon substrates. The nonlinear portion

corresponds to debonding of the film. The compliance is the inverse of the slope of the

loading line and can be used to calculate the crack length according to the following

equation:

31

3

8

'

bhCEa (2-18)

where C is the compliance measured from the load versus displacement data. The

critical load cP is taken to be the load at which the loading curve begins to deviate from

linearity. This deviation from linearity implies crack extension in the thin film stack, and

the critical value of applied G is taken to be the fracture energy [5]cG . Upon unloading,

the compliance will have increased, corresponding to an increase in crack length. In the

DCB geometry, the critical load and compliance must be measured separately for each

loading cycle.

17

Figure 2-5. Representative DCB data showing multiple loading and unloading cycles. Each cycle

further extends the crack and allows a measurement of Gc to be made.

Kanninen calibration

In 1971, Kanninen modified the beam theory solution to account for the portion of the

beam ahead of the crack tip[7]. In his augmented DCB model, he considered each half of

the DCB to be partially free in the cracked region and partially supported by an elastic

foundation in the uncracked portion. To do this, he added a Winkler foundation to the

beam theory problem. An important parameter that was introduced in this analysis is the

foundation modulus, k,

AEBk

h

(2-19)

where A is a constant equal to 2, E is the Young’s modulus of the beam, B is the beam

thickness, and h is the height of a single beam. The resulting modified compliance

relationship is

2 33

3

81 1.92 1.22 0.39Kann

a h h hC

EBh a a a

(2-20)

and the solution for the strain energy release rate is

22 2

2 3

121 0.64Kann

P a hG

EB h a

(2-21)

The crack length calculated from compliance techniques will differ from the beam theory

crack length in Equation (2-18) by 0.64h.

1/33

0.648

Kann

CEBha h

(2-22)

The optically-measured crack lengths are in good agreement with the Kanninen

predictions.

Penado calibration

Penado expanded Kanninen’s analysis to account for the presence of a thick, low-

stiffnessnes film in the DCB specimen[8]. The modified foundation compliance can be

18

simply calculated by adding the compliance value of the adhesive layer and substrate.

Reciprocally, the equivalent foundation modulus, kPen , is given by

1

1 1Pen

a s

k

k k

(2-23)

where ka and ks are the foundation modulus of the adhesive and substrate, respectively.

The foundation modulus of the adhesive includes a dependence on the layer half

thickness, t.

21

aa

a

EBk

t

(2-24)

where E and are the Young’s modulus and Poisson’s ratio of the adhesive layer and B

is the specimen thickness. The final solution for G from this analysis is given by

22 2

2 3 0.25 2 0.5

12 2 11

8

sPen

s s

EP a h hG

E B h h a h a

(2-25)

where s is the shear modulus of the substrate and the parameter is given by

43 Pen

s

k

E B (2-26)

2.4 Debond Growth Kinetics

Debonding is studied as the propagation of an interfacial crack between to materials. In

active environments, crack propagation can occur at values of G below cG .

Bond rupture kinetics

Environmentally-assisted crack growth is common, particularly in the materials and

interfaces exposed to high levels of relative humidity that contain Si-O bonds[9]. The

chemical reaction describing environmentally-assisted debonding, can be written as:

HOSiSiOHOHSiOSi 2

19

where bridging siloxane bonds are cleaved into a pair of terminal silanol groups in the

presence of water. Several solar module interfaces contain Si-O bonds.

Viscoelastic debond growth

Alternatively, interfacial crack propagation can be controlled by the viscoelastic

relaxation processes at the debonding-tip [10]. Viscoelastic stress relaxation occurs when

a constant strain o is imposed and maintained on a polymeric material. The stress is

monitored as a function of time and the response function is called the stress relaxation

modulus (E(t)):

ottE /)()( (2-27)

The value of E(t) is used in chapters 4 through 6 to understand and model the kinetics of

debonding of backsheet and encapsulation interfaces.

Debond growth metrology

Regadless of the underlying crack-growth mechanisms, the interfacial crack growth rate

depends on the applied value of G. A plot of the crack velocity versus the applied crack

driving force is called a v-G curve, which is shown in Fig. 2-3.

There are four main regions of a v-G curve. At the lowest applied values of G, there is a

threshold value below which crack propagation does not occur[11]. This value of the

applied crack driving force is denoted thG . Above

thG is the reaction region, where the

kinetics of the chemical reaction are responsible for cleaving bonds at the crack tip. As

the applied crack driving force increases further, a transport region may be encountered

where the interfacial crack velocity is determined by the transport rate of active species to

the crack tip. In the transport region, interfacial cracking is roughly independent of the

applied crack driving force. Lastly, in the critical region, the crack velocity rises

asymptotically as G approaches cG . This critical crack behavior was discussed in prior

sections about the Griffith fracture criteria.

A v-G curve can be obtained using a DCB specimen configuration. At the peak of a

DCB loading cycle, right under the critical fracture load, the position of the loading grips

are fixed, and the interfacial crack is allowed to propagate. The load will progressively

20

decrease throughout the experiment. The decreasing load causes a corresponding load

relaxation and corresponding decrease in G . As the applied G decreases from cG to

thG , the crack length, and hence the crack velocity, can be calculated at any time from the

load. A representative load versus time curve is shown in Figure 2-7, and a

representative crack length versus time curve is shown in Figure 2-8. From this

information, the v-G curve shown in Figure 2-6. Crack velocity as a function of applied

crack driving force. can be derived.

Figure 2-6. Crack velocity as a function of applied crack driving force. The four main regions of

the v-G curve (critical, transport, reaction, and threshold) are indicated.

Figure 2-7. Load versus time for a subcritical cracking test beginning at time t0.

21

Figure 2-8. Crack length versus time for a subcritical cracking test beginning at time t0.

22

2.5 References

[1] J.R. Willis, A comparison of the fracture criteria of griffith and barenblatt, J. Mech. Phys.

Solids. 15 (1967) 151–162. doi:10.1016/0022-5096(67)90029-4.

[2] B. Cotterell, The past, present, and future of fracture mechanics, Eng. Fract. Mech. 69

(2002) 533–553. doi:10.1016/S0013-7944(01)00101-1.

[3] T.L. Anderson, Fracture Mechanics, fundamentals and applications, CRC Press, 2005.

[4] J.W. Hutchinson, A.G. Evans, Mechanics of materials: top-down approaches to fracture,

Acta Mater. 48 (2000) 125–135. doi:10.1016/S1359-6454(99)00291-8.

[5] R.H. Dauskardt, M. Lane, Q. Ma, N. Krishna, Adhesion and debonding of multi-layer thin

film structures, Eng. Fract. Mech. 61 (1998) 141–162.



doi:10.1016/j.actamat.2012.01.007.

[7] M.F. Kanninen, An augmented double cantilever beam model for studying crack

propagation and arrest, Int. J. Fract. 9 (1973) 83–92. doi:10.1007/BF00035958.

[8] F.E. Penado, A Closed Form Solution for the Energy Release Rate of the Double

Cantilever Beam Specimen with an Adhesive Layer, J. Compos. Mater. 27 (1993) 383–

407. doi:10.1177/002199839302700403.

[9] M.W. Lane, J.M. Snodgrass, R.H. Dauskardt, Environmental effects on interfacial

adhesion, Microelectron. Reliab. 41 (2001) 1615–1624. <Go to ISI>://000171384900058.



[11] C.S. Litteken, R.H. Dauskardt, Adhesion of polymer thin-films and patterned lines, Int. J.

Fract. 119/120 (2003) 475–485. doi:10.1023/A:1024940132299.

23

3 CHAPTER 3: ENCAPSULATION ADHESION AND DEBONDING

3.1 Chapter Summary

Debonding of photovoltaic (PV) encapsulation in moist environments is frequently

reported but presently not well-understood or quantified. Temperature cycling, moisture,

and mechanical loads often cause loss of encapsulation adhesion and interfacial

debonding, initially facilitating back-reflectance and reduced electrical current, but

ultimately leading to internal corrosion and loss of module functionality. To investigate

the effects of temperature (T) and relative humidity (RH) on the kinetics of encapsulation

debonding, we developed a mechanics-based technique to measure encapsulation debond

energy and debond growth rates in a chamber of controlled environment. The debond

energy decreased from 2.15 to 1.75 kJm-2 in poly(ethylene-co-vinyl acetate) (EVA) and

from 0.67 to 0.52 kJm-2 in polyvinyl butyral (PVB) when T increased from 25 to 50C

and 20 to 40C, respectively. The debond growth rates of EVA increased up to 1000-fold

with small increases of T (10C) and RH (15%). To elucidate the mechanisms of

environmental debonding, we developed a fracture-kinetics model, where the viscoelastic

relaxation processes at the debonding-tip are used to predict debond growth. The model

and techniques constitute the fundamental basis for developing accelerated aging tests

and long-term reliability predictions for PV encapsulation.

3.2 Introduction

Photovoltaic (PV) encapsulation is arguably the most important structural component of

PV modules. It holds the module together and protects it from aggressive field

environments that include moisture, external mechanical loads, and temperature

cycles[1,2]. Guaranteeing the long-term (up to 25 years) interfacial stability of PV

encapsulation is currently a remarkable challenge; preventing interfacial debonding of

similar film technologies in other industries has so far proven difficult and costly[3].

Upon encapsulation debonding, additional optical interfaces are generated within the

module, resulting in local back-reflection and reduced electrical current. Following

debonding, several corrosion processes of the module components can accelerate

24

dramatically due to direct contact with the environment[4,5]. PV encapsulation

debonding can therefore be a crucial limiting factor for the yield and reliability of PV

modules, eventually risking loss of function[6,7]. Although the effect of hot and humid

field environments on encapsulation degradation is often reported[4,6,8], the fundamental

kinetic processes that lead to debonding have not yet been characterized. Debonding in

similar structures, however, has been shown to depend strongly on moisture[9,10],

temperature[11,12], mechanical stress[13], and their synergistic interactions[14]. The

effects of these variables on PV encapsulation debonding are important not only to assess

long-term encapsulation reliability, which may limit module lifetime, but also to develop

tests and protocols to improve encapsulation formulation(s) and manufacturing (surface-

preparation) procedures.

Encapsulation debonding occurs when the strain energy release rate, G (J·m-2), or debond

driving force, is equal to the critical debond energy (adhesion) at the encapsulation

interface, Gc. However, debonding in the presence of active environments can occur at

debond-driving forces below Gc, due to the effects of moisture and temperature on the

kinetic processes that occur at the debonding-tip, such as viscoelastic relaxation,

chemical reactions, and water transport.

Encapsulation adhesion has been previously studied using peeling-based techniques

including the 180° [15–17] , 90° “T” [1,18,19], and load-controlled 45°peel tests [20].

These techniques presently dominate the study of encapsulation and backsheet adhesion,

but the resulting peel-strength is not an interfacial property and cannot be used to

compare between dissimilar materials. Peel-strength results include macroscopic plastic

deformation of the encapsulation material and the applied handle layer (tape used to

perform the peel-test), which are difficult to quantify[21–23], unrelated to the adhesive

mechanism, and irrelevant to in-field debonding. Practically, deviation of peel angle from

its nominal value as well as cavitation (due to the large separation angle) also reduce

comparability between materials. More importantly, debonding of viscoelastic films is

strongly affected by the debonding-tip deformation processes [9,10,24], which are not

well represented by the various peeling techniques.

25

Other mechanical methods that have been applied to study PV encapsulation attachment

interfaces include the overlap shear test[2,25], rotational torque[26], compressive shear

test[27], wedge test[28–30], and the blister test[31,32]. The former methods examine the

strength of the attachment layer rather than interfacial adhesion to the relevant substrate.

The overlap shear test has long been known to be limited by the mechanical compliance

(displacement per unit load) of the specimen handles[33–35]. The compressive shear test

or double-lap shear[33] specimen geometries can be used to avoid this limitation. While

insight may be gained by analogy using the former methods, the results may not be

directly related to the fundamental mechanic and kinetic processes of debonding. The

wedge and blister methods do instead make use of fundamental fracture mechanics

principles, but they have not yet been demonstrated using PV materials. The wedge test,

however, may not be capable of determining the initial (unaged) Gc in strongly adhered

systems, like PV encapsulation.

The effect of simulated environments on encapsulation adhesion has been mostly studied

by replicating tests and sequences from the module qualification protocols[36,37]. The

“UV exposure”, “thermal cycling”, and “humidity freeze” tests simulate a brief field

exposure (months to a few years at best), whereas the conditions applied in “damp heat”

test greatly exceed those encountered in the application[38]. The qualification tests at

best query infant mortality failure modes, and do not of themselves investigate the rate of

degradation. The currently available aging treatments, as well as the available adhesion

tests, provide limited insight on the kinetic processes of debonding and their dependence

on mechanical stress, moisture and T. For example, softening of the encapsulant (lower

E) can occur near the transition or melting temperatures, which can significantly reduce

encapsulation adhesion and accelerate debonding. Kinetics-based testing protocols are

therefore urgently needed not only to assess encapsulation reliability and durability in

field environments, but also to define objectives for encapsulation formulation and

manufacturing (surface preparation) procedures.

In this study, we develop a quantitative characterization technique to measure the effects

of mechanical stress, T, and moisture on the debond energy and debond kinetics of two

encapsulation materials used in PV modules: poly(ethylene-co-vinyl acetate) (EVA) and

polyvinyl butyral (PVB). To measure debond energy, we fabricated single cantilever

26

beam debond (SCB) specimens (section 3.3.1) and modified a well-known SCB test

metrology (sections 3.3.2 and 3.3.3) ,where debonding is studied as the propagation of an

interfacial crack[39]. To ensure debonding of the glass/EVA interface, we describe a

novel temperature gradient method. To measure encapsulation debond growth rate

(section 3.4), da/dt, we applied a load-relaxation technique to the SCB debond specimens

in an environment of controlled temperature, T, and relative humidity, RH. The

technique allowed debond growth rates, da/dt, to be accurately quantified (~10-4 ms-1 to

10-8) in terms of the debond-driving force, G, and the viscoelastic relaxation processes at

the debonding-tip that depend on temperature (T) and relative humidity (RH). The

synergistic effect of UV-light is not examined in this study. Tensile specimens were also

fabricated (section 3.3.5) to quantify the effect of temperature on encapsulation stiffness,

which we used to rationalize the debond-growth data. Debonding kinetics was modeled

using viscoelastic fracture-kinetics (section 3.4.6), including the softening effect of T and

RH on the encapsulation modulus. The model provides a more fundamental

understanding of the effects of the environment on encapsulation debonding. The

techniques and model can be used to develop accelerated testing protocols, evaluate the

adoption of new encapsulations, and make long-term reliability predictions.

3.3 Materials and Methods

Encapsulation SCB debond specimens

In order to fabricate representative encapsulation specimens, a 0.45-mm layer of PV

encapsulation (EVA or PVB) was laminated between a 5.5-mm-thick glass substrate (low

iron-containing, soda-lime float glass, as in the PV industry) and a 3.17-mm-thick

titanium plate (6 Al-4 V, Grade-5), as shown in Fig. 1a. Prior to lamination, the surfaces

of the Ti plate and the glass (Sn-deficient side) were carefully cleaned, including: buffing

with pumice powder; washing with an aqueous detergent (Billco, Billco Manufacturing

Inc.); rinsing with deionized water; and rinsing with isopropyl alcohol. The lamination

step was performed in a commercial instrument (Model LM-404, Astropower Inc.) at

145C for 8 minutes under applied pressure of 1 atmosphere. During lamination, the

specimen components were fixed relative to one another using Kapton tape and were

27

surrounded with a glass frame to improve thickness uniformity at the edges. Two parallel

cuts (12-mm apart), were made through the full thickness of the Ti plate and the

encapsulation layer with a high-speed tungsten carbide saw and a razor blade,

respectively (Fig. 3-1b).

1a

1b

Figure 3-1: Single Cantilever Beam Specimen (SCB) for encapsulant debonding. a) A 0.45-mm

sheet of encapsulation was thermally bonded between a 3.17-mm Ti plate and a 5.5-mm soda

lime glass substrate (figure). b) Two parallel cuts, 12-mm apart, were made through the Ti plate

and encapsulation to fabricate the SCB specimen. The glass substrate was rigidly fixed to a

testing table and mechanical load was applied to one end of the Ti beam.

The elastic beam material (Ti here) was chosen to maximize elastic cantilever deflection,

as given by:

222

93

2yy

c

yyG

l

h

l

(3-1)

where l, h, y and y are the length, thickness, yield stress and yield strain of the beam,

respectively, and Gc is the encapsulation debond energy, which is unknown before the

debond measurement, but is independent of beam material. To maximize y , we

28

calculated the product 𝜎𝑦휀𝑦2 for several candidate beam materials (in MPa), including

steel (0.023), Al (0.019), PMMA (0.039) and Ti (0.047). We used Ti beams in this study

but some polymer beams, such as PMMA, can be sometimes easier to bond to the organic

PV encapsulation[24].

SCB debond test metrology

The SCB specimen described in section 0 was fixed to a rigid testing table with screw

clamps. A loading tab with a ruby bearing was bonded to one end of the Ti beam, and the

tab was then connected to an adhesion test system (Delaminator, DTS, Menlo Park, CA)

consisting of a linear actuator (displacement-controlled, 0.1 m res.) in series with a

high-resolution mechanical load cell fixed to a high-stiffness frame. The displacement of

the linear actuator and the mechanical load were monitored every 0.1 seconds.

The actuator was retracted (developing a tensile force, P) to initiate encapsulation

debonding, until the debond length (distance from the loading tab to the debonding-tip)

was ~15mm. The Ti beam in this configuration can be treated as a single cantilever beam

fixed at the debonding-tip[39]. The mechanical compliance of the SCB specimen (40 to

60 m/N) was much larger than that of the testing system (~0.5 m/N), which allowed

for accurate debond energy measurements. A minimum beam thickness (~3 mm here)

was required to avoid plastic deformation of the beam due to bending stress during the

debonding experiment.

The SCB specimen was mechanically loaded at a constant displacement rate to propagate

debonding for a few additional millimeters and then immediately unloaded to prevent

further debonding. The loading-unloading process was repeated for several cycles at

increasing debond-lengths, adjusting the actuator displacement rate to keep the local

deformation rate (at the debonding-tip) constant (see section 3.6). At each debonding

cycle, the mechanical compliance (displacement per unit load) of the SCB specimen, C,

and the mechanical load at the onset of debonding, 𝑃𝑐 , (where the load-displacement

curve becomes non-linear) were recorded.

29

The debond driving-force, G (J/m-2), can be than calculated with the following

expression[40]:

𝐺 =

𝑃2

2𝑏

𝑑𝐶

𝑑𝑎

(3-2)

where b is the width of the Ti beam and a is the debond-length. The encapsulation

debond energy, Gc, is then defined as the value of G at the onset of debonding, where

𝑃 = 𝑃𝑐.

The value of C for the SCB specimen is also a function of the encapsulation stiffness,

which depends on T and the mechanical loading-rate [41]. Using the DTS system

described in section 3.3, the value of C was measured at selected values of a at 20 and

50°C in a chamber of controlled RH (38%). The compliance data was then fitted to the

the compliance equation for a cantilever beam with elastic foundation [42]:

𝐶 =

4

𝐸𝑏(

𝑎

ℎ)

3

{1 +3

𝐵ℎ0.25(

ℎ

𝑎) + 3 [

1

𝐵2ℎ0.5+

𝐸

8𝐺] (

ℎ

𝑎)

2

+3

2𝐵3ℎ0.75(

ℎ

𝑎)

3

} (3-3)

where E, h and G are the Young’s modulus, thickness and shear modulus of the Ti beam;

and B is an encapsulation-dependent parameter given by

𝐵 = (ℎ

12+

𝐸 ∙ 𝑡(1 − 𝜐2)

3𝐸𝐸)

−14

(3-4)

where t, and 𝐸𝐸 are the thickness, poisson’s ratio and modulus of the encapsulation,

respectively. In the case of very large a or very thin elastic beams (ℎ/𝑎 ≪ 1), Eq. (3-3)

takes the form of the compliance equation for the elastic beam in cantilever without

encapsulation[43] present, where 𝐶 = 𝑎3

3𝐸𝐼⁄

Maximizing elastic beam deflection (thinnest beam) reduces the contribution of the

encapsulation layer to the overall SCB compliance, improves the sensitivity of the

debond experiment, and reduces the size of the damage zone ahead of the debond-tip. We

therefore selected the smallest possible value of h (3.17 mm here) for the Ti beam—

corresponding to a maximum bending stress of ~70% the yield strength of Ti.

30

SCB debond test displacement rate

The debond-tip displacement rate, ��, which characterizes the deformation rates at the

debond-tip, was kept constant at every debonding cycle to minimize rate-effects in the

debond energy measurement. To keep �� constant, we derived an equation for the required

actuator displacement rate (∆) as a function of a as follows.

The debond-tip opening displacement, 𝛿, is given by:

𝛿(𝑡) =

𝐺(𝑡)

휀𝑦𝐸𝐸

(3-5)

where 휀𝑦 is the yield strain of the encapsulation. (At debonding, 𝛿 = 𝛿𝑐 and 𝐺 = 𝐺𝑐.)

Substituting Eq.(3-2) and (3-3) into the time-derivative of Eq. (3-5), and neglecting

second and higher order terms, the actuator displacement rate required to keep �� constant

(at debonding) at increasing debond-lengths can be approximated as:

∆= (𝑎

𝑎0)

2

[𝐵 + 3ℎ0.75

𝑎⁄

𝐵 + 3ℎ0.75𝑎0

⁄] ∆0

(3-6)

where ∆0 is the reference actuator displacement rate (at a reference debond-lenth 𝑎0). To

illustrate the use of this equation, we calculated the value of �� (at debonding) as a

function of a, using Eqs. (3-2), (3-3) and (3-5) with a representative SCB specimen (Fig.

2). The values of �� and a were normalized to the initial values, 𝛿0 and a0=15mm

respectively. The value of �� decreases rapidly with a when the SCB is loaded at a

constant displacement rate (∆0) but remains approximately constant (��/𝛿0 ≈ 1) when

loaded at the corrected displacement rate (∆) given by Eq. (3-6).

In thinner and stiffer encapsulation layers, where 𝑡𝐸/𝐸𝐸 ≪ h , , or when the debond

experiment starts at large initial debond-lengths (ℎ0.75/𝑎0 ≪ 1), Eq. (3-6) reduces to:

∆= (

𝑎

𝑎0)

2

∆0 (3-7)

31

which has been previously used to determine the actuator displacement rates in debond

experiments with thinner and stiffer films[10,13].

Figure 3-2: Debond-tip opening displacement rate, ��, of the SCB specimen loaded at constant

(∆0) and corrected (∆) displacement rates, as a function of debond length, a, normalized to the

initial values, 𝛿0 and a0. The value of �� decreases rapidly with a when loaded at ∆0 and remains

roughly constant when loaded at ∆. B is an encapsulant parameter and h is the Ti beam thickness

Debond growth rate kinetics

A detailed explanation of the load-relaxation method for debond growth measurement is

provided elsewhere[12,39]; a brief summary is given here. To measure the encapsulation

debond growth rate, da/dt, as a function of G, the SCB specimens were mechanically

loaded in the DTS adhesion test system (section 0) up to a value of G close to the

encapsulation debond energy. The load relaxation was then recorded as a function of time

for about 24 hrs at a fixed actuator displacement. Post-experiment analysis of the load-

relaxation data determined the debond growth rates, da/dt, over the range of ~10-8 to

~10-4 ms-1 as a function of G. The experiments were performed on new encapsulation

specimens in an environmental chamber at selected values of RH (10, 25, 40, 55 and

70%) and T (20, 30, 40 and 50°C). The load cell and actuator were allowed to equilibrate

in the chamber environment for 6-8 h before each experiment.

32

The debond path in the debond growth experiments described above shifted occasionally

from the EVA/glass to the EVA/Ti interface (and vice versa), which limited the yield of

the debond growth experiments described above. In a separate set of experiments, a

temperature gradient was established on the EVA layer, which successfully stabilized the

debond path at the warmer EVA/glass interface. To establish the T-gradient, the glass

side of the SCB specimen was heated on a “hot-plate” at a fixed temperature (48, 53 or

58°C) and the Ti side was exposed to the colder air of an environmental chamber (20, 24

or 30°C). The heat-flow required to maintain this temperature gradient (4.1 kJ∙m-2) was

~4 times that of solar illumination (1 kJ∙m-2) as calculated with a 1-dimension Fourier

heat flow model using thermal conductivities of 22, 1, and 0.23 W∙m-1K-1 for Ti, glass

and EVA[44], respectively. The purpose of this experiment was not to emulate field-

conditions, but to investigate the effect of a temperature gradient on stabilizing the

debonding path to the EVA/glass interface. The temperature at the glass/EVA (28.4, 32.5,

and 38°C) and EVA/Ti (20.3, 24.4, and 30.3°C) were also calculated with this model.

The surface of the hotplate and the sides of the SCB specimen were covered with a 5-mm

sheet of polystyrene foam to minimize heat loss. The hot plate, SCB specimen, and the

delaminator system were allowed to stabilize in the environmental chamber for one hour

to establish the T-gradient across the glass/EVA/Ti structure. Once the gradient of ~8°C

was established, the debond growth experiment described above was conducted.

Tensile specimens

The SCB mechanical compliance and optimal SCB testing displacement rate (employed

to characterize da/dt and Gc) depend on the value of the encapsulation modulus (E) as

described in sections 3.2 and 3.3. In order to fabricate tensile experiments for module

measurement, a 0.45-mm layer of EVA was laminated between two glass substrates.

Several 3.5-mm square specimens were then cut from the laminated structure with a high-

speed precision saw to make glass/EVA/glass tensile specimens (3.5 × 3.5 × 10.5 𝑚𝑚).

To measure the value of E of the EVA layer, a loading tab was attached to each glass

substrate and a monotonic load was applied from -0.5 to 0.5 MPa in an environmental

chamber of controlled relative humidity (38%). The modulus was measured at selected

values of T in recently-laminated specimens (10 days) and in specimens aged for one year

33

in a laboratory-environment (23°C, fluorescent indoor illumination of ~200 lux, and RH

ranging from 25 to 45%). The cross-section area of the tensile specimen corresponded to

a mixed state of plane-stress (edges) and plane-strain (center) in the EVA layer.

Figure 3-3: Elastic (E) and storage (E’) modulus of the EVA encapsulation as a function of

temperature (T) in new and 1-year aged (20°C/40%RH) specimens. The value of E increased 10-

fold with the one-year age and decreased exponentially with T in the 30 to 60°C range. E was

measured using a 3.5 x 3.5 mm cross-section tensile specimen, subject to monotonic loading from

-0.5 to 0.5 MPa. E’ was measured with a DMA system at 32 mHz.

3.4 Results and Discussion

Environmental sensitivity of EVA modulus

The elastic modulus of the EVA encapsulation (E) as a function of testing temperature, T,

is shown in Fig. 3-3 for recently-laminated (10 days) and laboratory-aged (1 year)

specimens. Aging was conducted at 23°C, under fluorescent indoor illumination of ~200

lux, and 25 to 45% RH. The storage modulus of EVA, EDMA, measured with dynamic

mechanical analysis (DMA) in a previous experiment by one of the authors [45], is also

included. The value of E and EDMA decreased exponentially with T in the entropy-elastic

temperature range (30-60°C) below the melting temperature (65°C), Tm. The value of E

increased ~10-fold with aging, which was attributed to thermal and possible

EV

A M

od

ulu

s (M

pa)

Temperature (°C)

34

photochemical cross-linking of EVA[4] although the aged specimens were exposed only

to indoor illumination (fluorescent). To explain, recently laminated EVA encapsulation

may still contain as much as 40% of the initial peroxide used to facilitate cross-linking

[47,48]. In fielded PV modules, the combination of residual peroxide and UV radiation

explains the cross-linking of EVA to up to ~90% gel content.

The Arrhenius equation was fitted to the inverse of the measured moduli (E-1 and EDMA -1)

as shown in Fig.3 in dotted lines (T<Tm), with an activation energy of ~37 𝑘𝐽 ∙ 𝑚𝑜𝑙−1 in

the new specimens and 40 𝑘𝐽 ∙ 𝑚𝑜𝑙−1 in the aged specimens, respectively. The

similarity between the activation energies suggests that aging does not affect significantly

the temperature-dependence of the compliance over the one year age. Similarly, the

melting temperature, Tm, (65°C)[41] did not change significantly. At T>Tm, where the

polymer is in the melt state, the values of EDMA (1.1 MPa) and E (0.7 and 4 MPa) were

independent of T.

Mechanical compliance of SCB debond specimen

The mechanical compliance, C, of the SCB specimen as a function of debond length (a)

is shown in Fig. 3-4 for EVA (circle data points) and PVB (square data points)

specimens, at two selected values of T. The compliance of the bare Ti cantilever beam is

also shown for reference as a solid line. The value of C increased with a in the SCB

specimens and in the bare Ti beam. The specimen with the lowest encapsulation storage

modulus, PVB (G’~0.4 to 40 MPa)[49], exhibited the largest SCB compliance, followed

by the specimen with EVA (G’~20 to 1000 MPa)[41], followed by the bare Ti beam.

The value of C increased with T as a consequence of the temperature-dependence of the

encapsulation modulus; this dependence was stronger in the more compliant

encapsulation (PVB).

To obtain an analytical expression of C as a function of a, we fitted Eq. (3-3) to the

compliance data of the SCB specimens with EVA (dashed-line) and PVB (dotted-line), as

shown in Fig. 3-4. The fitted equations were then used to calculate the encapsulation

debond energy (Eq. (3-2)) as described in section 0.

35

Figure 3-4: Mechanical compliance, C, of EVA (□) and PVB (○) specimens (SCB), as a function

of cubed debond length, a3, at two selected environmental temperatures. The value of C

increased with T and a. A compliance model (Eq. 3), shown in dashed (EVA) and dotted (PVB)

lines, was fitted to the data. The compliance of the bare Ti beam in cantilever is shown for

reference as a solid line.

Encapsulation debond energy

The measured debond energy of the EVA/glass interface, Gc, decreased from 2.15 to 1.8

kJm-2 in the 25 to 50°C temperature range (Fig. 5). At temperatures close to Tm (65°C),

the value of Gc decreased six-fold to 0.3 kJm-2. The value of Gc remained constant in the

melt state at any T, similar to the modulus in Fig. 3.

The debond energy of the PVB/Ti interface, shown in Fig. 3-5 (circular data-points), also

decreased with T in the 20 to 40°C temperature range. The values of Gc for PVB/Ti (0.67

to 0.52 kJm-2) were much lower than those of EVA/glass in the same temperature range,

which can be attributed to the lower yield stress of PVB and to the nature of the

attachment at the interface. EVA is typically bonded to glass covalently, as enabled

through a silane chemistry (primer) typically present in the EVA formulation[27]. In

contrast, PVB is more weakly bonded via hydrogen bonding[27].

To better understand the effect of temperature on debond energy, Eq.(3-5) can be solved

for Gc [9]:

SC

B C

om

pli

ance

, C

(m

m/N

)

Cubed Debond Length, a3(m3)

36

𝐺𝑐 = 𝛿𝑐휀𝑦𝐸(𝑡, 𝑇) (3-8)

where 𝐸(𝑡, 𝑇) is the temperature-dependent modulus of the material at the debonding-tip.

In the EVA/glass interface, the extent of plastic deformation during debonding is

determined entirely by the weaker EVA encapsulation (𝜎𝑦 ≈ 4.5 𝑀𝑃𝑎) and not by the

stronger glass substrate. We therefore assume that the debonding process of EVA/glass

can be represented by considering only the mechanical behavior of the EVA. In

particular, 𝐸(𝑡, 𝑇) = 𝐸𝐸𝑉𝐴(𝑡, 𝑇), where 𝐸𝐸𝑉𝐴 is the modulus of EVA.

In other polymer studies[50,51], the value of c has been assumed independent of T,

however, the strong viscoelastic behavior of EVA and other PV encapsulation materials

suggests that 𝛿𝑐 depends strongly on T. This dependence can be demonstrated by

substituting the measured values of Gc and 𝐸 from Fig. 3-5 and 3-2, respectively, into

Eq(3-8) and considering that the value of 휀𝑦 is independent of T in most polymers[50,51]

including EVA[52]. The resulting values of 𝛿𝑐 for the EVA/glass interface can then be

fitted to an empirical equation valid in the 20 to 50°C interval:

𝛿𝑐 ≈

3.78 ∙ 10−4 ∙ 𝑇2 − 2.01 ∙ 10−1 ∙ 𝑇 + 27

103 ∙ 휀𝑦

(3-9)

The value of 𝛿𝑐 increases significantly with T .

The temperature dependence of Gc (Fig. 3-5) can therefore be rationalized as the

competing effects of lower E and larger 𝛿𝑐 at higher T. The implication is that measuring

tensile and yield properties of the encapsulation, such as E or 𝜎𝑦, is not sufficient to

characterize the debond energy in encapsulation/glass systems. The stress state at the

debond-tip, represented by 𝛿𝑐, can drastically affect the value of the interfacial debond

energy Gc , as demonstrated by Eq.(3-8) and (3-9).

It is important to highlight that only fracture-mechanics based experiments, like the SCB

technique presented here, can capture the effects of debond-tip geometry. The absence of

debond-tips in the peel test experiments[17–19,53] suggests that the peel-strength may be

strongly correlated to the uniaxial properties of the encapsulation (E, 𝜎𝑦), which can yield

37

results irrelevant to the encapsulation/substrate system as describe above. For example,

considering only the value of E of EVA, which decreases continuously through Tm (Fig.

3-3), would not provide any information on the sharp discontinuity of the debond energy

at T=Tm (Fig. 3-5). The sudden decrease of debond energy can only be understood by

considering the inability of EVA to maintain the debond-tip geometry at temperatures

near Tm.

Figure 3-5: Debond Energy of EVA/glass and PVB/Ti interfaces as a function of T. The debond

energy of EVA/Glass decreased gradually with T in the 20 to 50°C range, decreased precipitously

near the melting temperature of EVA (65°C), and remained constant at higher temperatures. The

debond energy of PVB/Ti was much lower than that of EVA/glass. The debond experiments were

conducted using the single cantilever beam technique in a chamber of controlled environment at

38% relative humidity.

Environmental debonding kinetics

The encapsulation debond growth rate, da/dt, as a function of debond-driving force, G, is

shown in Fig. 3-6 at selected values of T for the EVA/glass (T=20, 30,50°C) and PVB/Ti

(T=20, 40 °C) interfaces. The debond growth rate increased with G, displaying a

concave-down da/dt-G debond growth curve, which shifted to lower values of G at

higher T. Conversely, the debond growth rate increased with T at any given value of G,

demonstrating that encapsulation debonding is a thermally activated process[9].

Deb

on

dE

ner

gy,

Gc

(J/m

2)

Temperature (°C)

38

Debonding of PVB/Ti occurred at much lower values of G than in EVA/glass. Debond

driving-force thresholds, Gth, operationally defined at ~10-8 m/s, were much lower for the

PVB/Ti (Gth =200 J/m2) compared to the EVA/glass interface. The values of da/dt were

more sensitive to temperature changes at lower (compared to high) test temperatures, for

example: In EVA/glass, higher values of T shifted the value of G (corresponding to

da/dt=10m/s) by 50 𝐽 ∙ 𝑚−2°C−1 at low T (20 to 30°C ), but only 35 𝐽 ∙ 𝑚−2°C−1 at

high T (30 to 50°C). Modulus stiffening due to aging was considered negligible for most

data in Fig.6, including the curves at low temperature (20 and 30°C) and the high

debond-growth portion of the 50°C data (10-6 to 10-4 m/s), which was obtained in the first

two hours of the experiment. The low debond-growth portion of the 50°C data (collected

over two days) may include stiffening effects due to temperature cross-linking.

Figure 3-6: Debond growth rate of EVA/Glass (□) and PVB/Ti (○) as a function of debond

driving-force (G) at selected temperatures (T). Debonding of PVB/Ti occurred at much lower

values of G than in EVA/Glass. In both interfaces, the debond growth rates were shifted to lower

values of G with increasing T. The data was collected using a load-relaxation technique in a

chamber of controlled environment. A Fracture-kinetics model (dashed lines, Eq. 21) was fitted

to the data.

The debond growth rate of EVA/Ti, measured at five selected values of relative humidity,

RH, is shown in Fig. 3-7 as a function of G. The debond growth curves were shifted to

Deb

on

dG

row

th R

ate,

da/

dt

(m/s

)

Debond Driving Force, G(J/m2)

39

lower values of G at higher values of RH. Conversely, the debond growth rate increased

with RH at any given value of G, demonstrating that encapsulation debonding is a

moisture-activated process. Similar to the effect of T on da/dt, debonding was more

sensitive to RH changes in dryer environments: the value of G corresponding to da/dt

=100 m/s decreased 250 Jm-2 from 10 to 25% RH, but only 150 Jm-2 from 40 to 55%.

The similarities between the effects of T and RH on debond growth rate suggest a

common debonding mechanism, which we propose is determined by the viscoelastic

relaxation of the material ahead of the debond-tip. To better understand the effect of T, G

and RH on debond growth, a viscoelastic fracture-kinetics model was developed in

section 0.

Figure 3-7: Debond growth rate of EVA/Ti as a function of debond driving-force (G), at selected

values of relative humidity, (RH). The debond growth rates were shifted to lower values of G

with increasing RH. The data was collected using a load-relaxation technique in a chamber of

controlled environment. A fracture-kinetics model (dashed lines, Eq. 21), which includes the

plasticizing effect of moisture on EVA, was fitted to the data.

The values of da/dt in the presence of a T-gradient (across the EVA layer) are shown in

Fig. 3-8 as a function G at three selected temperatures of the EVA/glass interface.

Debonding in the presence of a T-gradient occurred at much lower values of G (400 to

1000 Jm-2) than those at uniform T (G~ 1200 to 2500 Jm-2). Although the maximum

Deb

ond

Gro

wth

Rat

e, d

a/dt

(m/s

)


40

temperature difference across the encapsulation thickness was only a few degrees Celsius

(~ 8°C), the low values of G (or conversely, increased values of da/dt) in the presence of

a T-gradient can be attributed to the size and location of the debonding-tip. The

viscoelastic deformation rates that control debond growth are faster near the warmer glass

substrate than in the bulk of EVA. The resulting asymmetrical mass transport shifts the

debond-tip towards EVA/glass interface. The stiffer glass substrate then restricts the

amount of material that undergoes viscoelastic relaxation near the debond tip, which

effectively weakens the interface. For example, exactly at the EVA/glass interface, the

debonding-tip is adjacent to only half as much viscoelastic material than in the bulk of

the encapsulation layer, which roughly corresponds to the ratio of G values in Fig. 3-6

and 3-8 for equivalent values of T.

Figure 3-8: Debond growth rate of EVA/glass subject to an 8°C T gradient, as a function of

debond driving-force (G), at three selected EVA/Glass temperatures. The debond path was

stabilized at the warmer EVA/glass interface in the presence of the T-gradient. The debond

growth rates were shifted to lower values of G at higher values of T. A Fracture-kinetics model

(dashed lines, Eq. 21) was fitted to the debond growth data.

Deb

on

dG

row

th R

ate,

da/

dt

(m/s

)


41

Debond Kinetics Model

We recently derived a fracture kinetics model for viscoelastic debonding in PV

backsheets[9]. The core assumption of the model is that the debond growth rate (da/dt)

can be approximated by the rate of formation of the viscoelastic zone at the debonding-

tip. The viscoelastic zone is modeled akin to a plastic zone, rp, and forms in a time

approximated by the viscoelastic modulus relaxation time, 𝑡𝑇[50,51] (time at 63% stress

reduction due to relaxation). The debond growth rate can then be described by:

𝑑𝑎

𝑑𝑡=

𝑟𝑝

𝑡𝑇

(3-10)

Where rp is given by:

To include the effect of T on da/dt, the time-temperature superposition principle[50] was

applied to a power-law equation describing the time dependence of the backsheet

modulus, 𝐸[51,54]:

𝐸1𝑡−𝑛 = 𝐸1 (

𝑡𝑇

𝑎𝑇)

−𝑛

(3-12)

where 𝐸1 is the modulus coefficient, n is a measure of the rate sensitivity of the material

(n=0 for perfectly elastic solids), and 𝑎𝑇 is the Arrhenius time-temperature “shift” factor

described by:

1

𝑎𝑇= 𝑒

−𝐸𝑎𝑅

(1𝑇

−1𝑇𝑟

)

(3-13)

where Ea is the activation energy of the viscoelastic relaxation process, Tr is a reference

temperature, and R is the gas constant. It is important to note that the temperature

dependence of the modulus (measured in section 4.1) also exhibited an Arrhenius-type

dependence (as shown in Fig. 3-3), which validates the use of Eq. 13 in this model.

The resulting debond growth rate in the original model was written as:

𝑟𝑝 =

𝜋

8

𝛿𝑐

휀𝑦

(3-11)

42

𝑑𝑎

𝑑𝑡=

𝜋

8

𝛿𝑐

휀𝑦(𝛿𝑐 휀𝑦𝐸1)1

𝑛⁄(𝐺)

1𝑛 ∙ 𝑒

−𝐸𝑎𝑅

(1𝑇

−1𝑇𝑟

)

(3-14)

Note that as previously discussed in section 4.3, the value of 휀𝑦 is largely independent of

temperature in most polymers[50,51] including EVA[52]. Also, the value of 𝛿𝑐 can be

approximated with Eq. (3-9).

The model will now be extended to explicitly account for RH and to consider the

presence of the encapsulant/glass interface as follows. Like in section 4.3, we assume

that the deformation related to EVA/glass debonding is determined exclusively by the

mechanical properties of the EVA layer and not by the stronger and stiffer glass

substrate. We then model the plasticizing effect of water on the EVA glass transition

temperature, Tg, with the Gordon-Taylor equation:

𝑇𝑔 =

𝑤𝑑𝑇𝑔𝑑 + 𝑘𝑤𝑤𝑇𝑔𝑤

𝑤𝑑 + 𝑘𝑤𝑤

(3-15)

where k is a proportionality constant, wd and ww are the weight ratios of dry EVA and

water, respectively, 𝑇𝑔𝑑 is the glass transition temperature of dry EVA, and 𝑇𝑔𝑤 =

−135°𝐶 is a reference value used for aqueous plasticization[55]. Due to the low

saturation limit of water in EVA (~0.01 ≪ 1) [56], the shift of 𝑇𝑔 with increasing 𝑤𝑤 can

be linearized as:

𝑇𝑔 = 𝑇𝑔𝑑 + ∆ ∙ 𝑤𝑤 (3-16)

where ∆= 𝛼 ∙ 𝑘 ∙ 𝑇𝑔𝑤, where 𝛼 is a constant.

We then assume that the shift of the entire viscoelastic E-T curve (due to aqueous

plasticization) can be approximated with the shift of 𝑇𝑔:

𝐸(𝑇, 𝑤𝑤) ≈ 𝐸(𝑇 + ∆ ∙ 𝑤𝑤, 0) (3-17)

We further assume that the value of 𝑤𝑤 in the plastic zone at the debond tip is

proportional to the adsorbed water at the EVA surface. In a previous study[57], the

adsorbed weight-ratio of water at the EVA surface was shown to linearly depend on the

partial pressure of water vapor, 𝑝𝐻20 , which corresponds to a Henry’s law-type

43

adsorption isotherm. The value of of 𝑤𝑤 at the debond-tip can then be approximated

with:

𝑤𝑤 = 𝑚 ∙ 𝛽 ∙ 𝑝𝐻20 (3-18)

where m is a proportionality constant and 𝛽(𝑇) is the Henry’s law adsorption coefficient

that can be expressed with the van’t Hoff equation using data of EVA [57] in the 20 to

55°C range:

𝛽(𝑇) ≈ 𝛽(𝑇𝜃)𝑒

𝐶(1𝑇

−1

𝑇𝜃) (3-19)

where 𝑇𝜃 = 298𝐾, C=4200 and 𝛽(𝑇𝜃) = 0.03 𝑚𝑏𝑎𝑟−1

To relate 𝑝𝐻20 to RH, we use the well-known equation:

𝑝𝐻20(𝑚𝑏𝑎𝑟) = 𝜑 ∙ 𝑅𝐻 (3-20)

where 𝜑 is the saturation vapor pressure of H20, which can be described with a fit of

atmospheric data values: 𝜑(𝑚𝑏𝑎𝑟) = 0.0744 ∙ 𝑇2 − 42.531 ∙ 𝑇 + 6101.4

Using Eqs. (3-17) through (3-20), the encapsulation debond growth rate with explicit G,

RH and T dependence can be finally written as:

𝑑𝑎

𝑑𝑡= 𝛼 ∙ 𝛿𝑐

1−1𝑛⁄

∙ (𝐺)1

𝑛⁄ ∙ 𝑒−

𝐸𝑎𝑅

(1

𝑇+𝛾𝑅𝐻 −

1𝑇𝑟

)

(3-21)

where 𝛼 =𝜋

8∙ 𝐸1,𝑑𝑟𝑦

−1𝑛⁄ ∙ 휀𝑦

−1−1𝑛⁄

, where 𝐸1,𝑑𝑟𝑦 is the encapsulation modulus

coefficient in dry environments (𝑤2 = 0) and the coefficient (𝛾), in the exponent of Eq.

(3-21) is given by:

𝛾 = 𝑚∆ ∙ 𝛽(𝑇𝜃)𝑒

𝐶(1𝑇

−1

𝑇𝜃)

∙ 𝜑(𝑇) (3-22)

The parameters 𝛼, 𝑛, 𝐸𝑎 in Eq. (3-21) and 𝑚∆ in Eq. (3-22) were fitted to the data in

Figs. 3-6 through 3-8. The resulting value of Ea was 956 kJ∙mol-1 for the EVA/glass

interface, 357 kJ∙mol-1for the EVA/Ti interface, and 1560 kJ∙mol-1 for the EVA/glass

interface with a T-gradient. The values of n=0.043 and 𝑚∆= 97 °C represent EVA

parameters that are common to both the EVA/glass and EVA/Ti interfaces. The value of

44

𝛿𝑐 was approximated with Eq. (3-9). The model curves, shown as dotted lines in Figs. 6,

3-7 and 3-8, are in close agreement with the debond growth rate data. Note that the values

of 𝛾 𝛼 and Ea for the PVB/Ti interface cannot be uniquely defined without further

temperature data not available in the present study.

The expression 𝑇 + 𝛾𝑅𝐻 in the exponent of Eq. (3-21) accounts for the combined effects

of T and RH on debond growth: increasing RH is equivalent to increasing T, and the

value of 𝛾 can then be interpreted as the RH to T equivalence in debonding kinetics. The

value of 𝛾 for EVA ranged from 113 to 156 °C in the T=20°C to 60°C interval:

increasing RH by 1% was equivalent to increasing T by 1.13 to 1.56 degrees Celsius.

The debonding mechanism of the EVA/glass interface changes at values of T close to Tm,

as reported in section 4.3. Close to Tm, the debonding-tip becomes unstable and the value

of 𝛿𝑐 decreases abruptly which led to lower values of G, and accounts for the increased

values of da/dt in Fig. 6 at 50°C, as compared to those predicted by Eq. (3-21). This

mechanism transition is particularly relevant to PV applications because roof-mounted

modules have been verified to achieve up to 105 C in desert locations[58]. Likewise, the

high-RH data (55% and 70%) in Fig. 7 was higher than predicted by Eq.(3-21),

suggesting that the mechanism transition at high-RH is similar to that at high-T.

Comparing Figs. 3-6 and 3-7, the change in debonding mechanism occurred only when

the quantity 𝑇 + 𝛾𝑅𝐻 reached a common value (~103°C). We therefore suggest that,

there is a critical value of 𝑇 + 𝛾𝑅𝐻 at which the debond-tip becomes unstable and

debonding may accelerate dramatically. This mechanism transition is therefore expected

to occur at much lower temperatures in tropical locations (high RH) than in desert

locations (low RH). The RH/T conditions for catastrophic (increased da/dt) debonding in

EVA encapsulation may then be approximated with the equation: 𝑇 + 𝛾𝑅𝐻 ≈ 103°C.

We now propose that the conditions for catastrophic debonding in other elastomeric

materials that undergo water and temperature plasticization may also be approximated

with: 𝑇 + 𝛾𝑅𝐻 ≈𝑇𝑚

𝑓+ 𝛾𝑅𝐻𝑚, where Tm is the melting temperature of the encapsulation

material as measured at 𝑅𝐻𝑚 , and 𝑓 > 1 is an empirical constant. The following

condition may be satisfied to avoid catastrophic debonding:

45

𝑇 + 𝛾𝑅𝐻 <

𝑇𝑚

𝑓+ 𝛾𝑅𝐻𝑚

(3-23)

which can be used to evaluate the adoption of new elastomeric encapsulation materials

given the values of RH and T from the environment and the values of 𝑇𝑚, 𝑓 and 𝛾 from

the encapsulation material. Satisfying Eq. (3-23) implies that the operating T and RH are

sufficiently lower than the debond-tip “softening” point, as characterized by Tm and RHm.

Eq (3-21) can then be used to estimate encapsulation lifetime, which may be defined as

the time required for the growth of an initial debond defect (typically due to

manufacturing or installation), to a final critical size, where loss of adhesion is

anticipated. Changes in modulus due to UV light and thermally activated cross-linking

(see Fig. 3-3) should also be considered when using Eq (3-21).

In this study, higher modulus (E) corresponded to higher debond energies and lower

debond growth rates. Higher values of E are usually correlated to higher polymer cross-

link densities, which are desirable to hinder viscoelastic-controlled debonding. However,

very high cross-link densities may also sharpen the debonding-tip, reducing the ability of

the material to accommodate stress and leading to weaker interfaces and faster debond

rates. To reduce in-field debonding, modulus stiffening via cross-linking may therefore

be optimized between these two competing regimes to efficiently improve module

durability.

3.5 Conclusions

The combined effects of mechanical stress, moisture and temperature on the loss of

encapsulation adhesion and debond growth kinetics were studied in two types of PV

encapsulation materials: EVA and PVB. A crack-growth technique was modified to study

interfacial encapsulation debonding. The technique accounts for the rate and temperature-

dependence of the encapsulation mechanical properties at the debonding-tip. The debond

energy of EVA and PVB decreased with testing T, which was rationalized as the result of

the competing effects of lower encapsulation modulus and larger debond-tip opening

displacements at higher T. Debond kinetics were quantified as a function of mechanical

46

driving-force, T and RH. The debond growth rates of EVA increased up to 1000-fold with

small increases of, T (10C) and RH (15%).

To elucidate the mechanisms of environmental debonding, a fracture-kinetics model was

developed based on the viscoelastic relaxation and plasticization processes at the

debonding-tip. The model accurately describes the debond growth rates as a function of

encapsulation material, applied mechanical driving-force, T, and RH in the environment.

The RH and T-dependences of the encapsulation modulus were used to successfully

predict debond growth. The model suggests that increasing RH is equivalent to increasing

T. An encapsulation property (𝛾), which quantifies the relative effects of RH and T on

debond growth was also introduced. The model importantly suggests that a transition to

rapid debonding may occur as a function of T and/or RH. The techniques and models can

be used to evaluate the adoption of new encapsulation materials, to develop accelerated

testing protocols, and to make long-term reliability predictions.

47

3.6 References




[2] G.J. Jorgensen, K.M. Terwilliger, M.D. Kempe, T.J. McMahon, Testing of packaging

materials for improved PV module reliability, Conf. Rec. Thirty First IEEE Photovolt.

Spec. Conf. 2005. (2005) 499–502. doi:10.1109/PVSC.2005.1488179.

[3] S.-Y. Kook, R.H. Dauskardt, Moisture-Assisted Subcritical Debonding of a

Polymer/Metal Interface, J. Appl. Phys. 91 (2002) 1293. doi:10.1063/1.1427139.





doi:10.1002/pip.



(2002) 1436–1439.

[7] J.E. Granata, W.E. Boyson, J. a. Kratochvil, M. a. Quintana, Long-term performance and

reliability assessment of 8 PV arrays at Sandia National Laboratories, 2009 34th IEEE

Photovolt. Spec. Conf. (2009) 001486–001491. doi:10.1109/PVSC.2009.5411336.

[8] D.L. King, M. a. Quintana, J. a. Kratochvil, D.E. Ellibee, B.R. Hansen, Photovoltaic

module performance and durability following long-term field exposure, AIP Conf. Proc.

565 (1999) 565–571. doi:10.1063/1.58001.

[9] F.D. Novoa, D.C. Miller, R.H. Dauskardt, Environmental mechanisms of debonding in

photovoltaic backsheets, Sol. Energy Mater. Sol. Cells. (2013) 1–7.

doi:10.1016/j.solmat.2013.08.020.

[10] S.R. Dupont, F. Novoa, E. Voroshazi, R.H. Dauskardt, Decohesion Kinetics of

PEDOT:PSS Conducting Polymer Films, Adv. Funct. Mater. (2013) n/a–n/a.

doi:10.1002/adfm.201302174.

[11] B. Lawn, An atomistic model of kinetic crack growth in brittle solids, J. Mater. Sci. 10

(1975) 469–480.

48



doi:10.1016/j.actamat.2012.01.007.

[13] S.R. Dupont, M. Oliver, F.C. Krebs, R.H. Dauskardt, Interlayer adhesion in roll-to-roll

processed flexible inverted polymer solar cells, Sol. Energy Mater. Sol. Cells. 97 (2012)

171–175. doi:10.1016/j.solmat.2011.10.012.

[14] S.-Y. Kook, J.M. Snodgrass, A. Kirtikar, R.H. Dauskardt, Adhesion and Reliability of

Polymer/Inorganic Interfaces, J. Electron. Packag. 120 (1998) 328.

doi:10.1115/1.2792642.

[15] ISO 8510-2, Adhesives—Peel test for a flexible-bonded-to-rigid test specimen

assembly—Part 2: 180 degree peel, Int. Electrotech. Comm. Geneva, 1–5. (1990).

[16] J. Wohlgemuth, D. Cunningham, P. Monus, J. Miller, A. Nguyen, Long Term Reliability

of Photovoltaic Modules, 2006 IEEE 4th World Conf. Photovolt. Energy Conf. (2006)

2050–2053. doi:10.1109/WCPEC.2006.279905.





[19] A.J. Kinloch, C.C. Lau, J.G. Williams, The peeling of flexible laminates, Int. J. Fract. 66

(1994) 45–70. doi:10.1007/BF00012635.

[20] N.E. Shephard, J.P. Wightman, A Simple Device for Measuring Adhesive Failure in

Sealant Joints, Sci. Technol. Build. Seals, Sealants Weather. Seventh Vol. ASTM STP

1334. (1998).

[21] K.-S. Kim, N. Aravas, Elastoplastic analysis of the peel test, Int. J. Solids Struct. 24

(1988) 417–435. doi:http://dx.doi.org/10.1016/0020-7683(88)90071-6.

[22] Y. Wei, J.W. Hutchinson, Interface strength , work of adhesion and plasticity in the peel

test, Int. J. Fract. 93 (1999) 315–333. doi:10.1023/A:1007545200315.

[23] J.L. Gardon, Peel adhesion. I. Some phenomenological aspects of the test, J. Appl. Polym.

Sci. 7 (1963) 625–641. doi:10.1002/app.1963.070070219.

[24] A. Kamer, K. Larson-Smith, L.S.C. Pingree, R.H. Dauskardt, Adhesion and degradation

of hard coatings on poly (methyl methacrylate) substrates, Thin Solid Films. 519 (2011)

1907–1913. doi:10.1016/j.tsf.2010.08.116.

[25] M.D. Kempe, Accelerated UV Test Methods for Encapsulants of Photovoltaic Modules,

in: 33rd IEEE, 2008: p. 5.

49

[26] G. Jorgensen, T. McMahon, Accelerated and outdoor aging effects on photovoltaic

module interfacial adhesion properties, Prog. Photovoltaics …. (2008) 519–527.

doi:10.1002/pip.

[27] V. Chapuis, S. Pélisset, M. Raeis-Barnéoud, C. Li, H.-Y., Ballif, L.-E. Perret-Aebi,

Compressive-shear adhesion characterization of polyvinyl-butyral and ethylene-vinyl

acetate at different curing times before and after exposure to damp-heat conditions, Prog.

Photovoltaics Res. Appl. 22 (2014) 405–414. doi:10.1002/pip.2270.

[28] J.A. Marceau, Y. Moji, J.C. McMillan, A Wedge test for Evaluating Adhesive-Bonded

Surface Durability, Adhes. Age. (1977) 28–34.

[29] ISO 10354 “Adhesives -- Characterization of durability of structural-adhesive-bonded

assemblies -- Wedge rupture test, 1992.

[30] ASTM D3762 - Standard Test Method for Adhesive-Bonded Surface Durability of

Aluminum (Wedge Test), West Conshohocken, PA, 2010.

[31] K.-K. Liu, K. Wan, Multi-scale mechanical characterization of a freestanding polymer

film using indentation, Int. J. Mater. Res. 99 (2008) 862–864. doi:10.3139/146.101712.

[32] L.B. Freund, S. Suresh, Thin Film Materials: Stress, Defect Formation and Surface

Evolution, Cambridge University Press, 2008.

[33] D.-A. Mendels, S.A. Page, Y. Leterrier, J.-A.E. Manson, A Modified Double Lap-Shear

Test as a Mean to Measure Intrinsic Properties of Adhesive Joints, in: Proc. Euro. Conf.

Comp. Mats., 2000.

[34] O. Volkersen, Luftfahrtforschung, in: 15, 1938: p. 41.

[35] M. Goland, E. Reissner, No Title, J. Appl. Mech. 11 (1944) A17–27.


type approval, Int. Electrotech. Comm. Geneva, 1–81. (2008).


qualification and type approval ", Int. Electrotech. Comm. Geneva, 1–93. (2005).






[41] D.C. Miller, M.D. Kempe, S.H. Glick, S.R. Kurtz, Creep in Photovoltaic Modules:

Examining the Stability of Polymeric Materials and Components, Proc. IEEE PVSC.

(2010) 262–268.

50



407. doi:10.1177/002199839302700403.



407. doi:10.1177/002199839302700403.

[44] S. Ghose, K.A. Watson, D.C. Working, J.W. Connell, J.G. Smith, Y.P. Sun, Thermal

conductivity of ethylene vinyl acetate copolymer/nanofiller blends, Compos. Sci. Technol.

68 (2008) 1843–1853. doi:10.1016/j.compscitech.2008.01.016.



(2010) 262–268.

[46] S. Turri, M. Levi, M. Cristini, A. Sanguineti, Dynamic and Thermo-Mechanical Properties

of Some Specialty Fluoroelastomers for Low Tg Seal Materials, J. Polym. Res. 14 (2007)

141–145. doi:10.1007/s10965-006-9093-3.




[48] F. Pern, Ethylene‐vinyl acetate (EVA) encapsulants for photovoltaic modules:

Degradation and discoloration mechanisms and formulation modifications for improved,

Die Angew. Makromol. Chemie. 252 (1997) 195–216.

[49] A. DasGupta, D. David, A. Misra, Modification of polyvinyl butyral rheological

properties via blending with ionomeric polyvinyl butyral, Polym. Bull. 25 (1991) 657–

660. doi:10.1007/BF01032661.

[50] G. Marshall, L. Coutts, J. Williams, Temperature effects in the fracture of PMMA, J.

Mater. Sci. 9 (1974) 1409–1419.

[51] R. Schapery, A method for predicting crack growth in nonhomogeneous viscoelastic

media, Int. J. Fract. 14 (1978) 293–309.

[52] S. Hobeika, Y. Men, G. Strobl, Temperature and Strain Rate Independence of Critical

Strains in Polyethylene and Poly(ethylene- c o -vinyl acetate), Macromolecules. 33 (2000)

1827–1833. doi:10.1021/ma9910484.



51

[54] W. Bradley, W. Cantwell, H. Kausch, Viscoelastic creep crack growth: a review of

fracture mechanical analyses, Mech. Time-Dependent Mater. I (1997) 241–268.

[55] Y. Roos, Water Activity and Plasticization, in: Food Shelf Life Stab., CRC Press, 2000.

doi:doi:10.1201/9781420036657.ch1.

[56] M.D. Kempe, Control of moisture ingress into photovoltaic modules, Conf. Rec. Thirty-

First IEEE Photovolt. Spec. Conf. 2005. (2005). doi:10.1109/PVSC.2005.1488180.

[57] G.R. Mon, L.C. Wen, R.S. Sugimura, R.G.. J. Ross, Reliability studies of photovoltaic

module insulation systems, Proc. 19th Electr. Electron. Insul. Conf. (1989).

doi:10.1109/EEIC.1989.208251.

[58] M.D. Kempe, D.C. Miller, J.H. Wohlgemuth, S.R. Kurtz, J.M. Moseley, Q. (Annie) Shah,

et al., Field Testing of Thermoplastic Encapsulants in High Temperature Installations, n.d.

52

4 CHAPTER 4: SQUARE-CORNER ADHESION METROLOGY

4.1 Chapter Summary

Long-term solar module exposure to environmental stressors, including moisture,

temperature and UV, leads to encapsulation loss of adhesion, which together with

mechanical stress (thermal, wind, handling) can lead to module debonding, component

corrosion and eventually loss of function. To quantify glass/encapsulation adhesion

(J/m2), we laminated a Ti/glass/EVA/glass/Ti square coupon, initiated a debond defect at

the EVA/glass interface, and applied two opposing tensile forces at one corner of the

coupon. In this loading configuration, steady-state debonding occurs at a constant force

(Pc) and the energy of adhesion (J/m2) can be calculated with Gc=Pc2/(4B), where B is the

cross-section bending stiffness of the bilayer (Ti/glass) substrates. We provide graphical

and analytical methods to calculate B even in the presence of glass surface cracking. We

used this metrology to quantify the increase of EVA/glass adhesion (0.48 to 1.5 kJ/m2) in

the first 9 min of post-lamination, 140°C, cross-linking treatment. Adhesion then

decreased with further exposure to the 140°C environment. We extend this technique to

easily quantify adhesion of other module laminates and interfaces (e.g. encapsulation/Si)

using both lab-made coupons and full-size solar modules. The metrology presented here

can be used to quantify loss of adhesion after field or simulated exposures using an

absolute scale (J/m2), in order to improve encapsulation materials and to optimize surface

preparation procedures.

4.2 Introduction

Solar module encapsulation materials are exposed to aggressive field environments for

long periods of time (>20 years) [1,2]. These materials and their interfaces are

significantly affected by temperature [3,4], moisture [5,6] and UV light, which together

with field, thermal and wind stresses, can lead to debonding, component corrosion [7,8]

and eventually loss of function [9,10]. Quantifying encapsulation adhesion is important in

order to improve encapsulation materials, optimize surface preparation procedures, and

design accelerated-exposure protocols for module lifetime estimation. The currently

53

available lifetime estimation methods are largely based on curve-fitting and extrapolation

of past-performance; very little is understood about the fundamental mechanisms of

degradation, particularly those related to component loss of adhesion. The metrology

presented here can be used to quantify the loss of encapsulation adhesion after field or

simulated exposures, which constitutes the basis to design meaningful accelerated-

exposure protocols for encapsulation lifetime estimation.

Debonding of encapsulation occurs when the strain-energy release rate G (J/m2), which is

a function of the mechanical stress applied on the module (e.g. thermal, wind or

handling), is greater than the debond energy of the interface, Gc. The magnitude of Gc is a

function of not only the interfacial bond strength but also of the debond-front

deformation processes, which can be affected by the environment, including moisture and

temperature.

The strength of attachment of encapsulation has been mostly studied using peel-based

techniques, including the 180° [11–13] , 90° “T” [1,14,15], and load-controlled 45°peel

tests [16]. Peel strength, however, is mostly determined by the macroscopic yield and

tensile properties of the bulk of the encapsulation layer [17–19] rather than by the

interfacial debonding processes, especially in strongly adhered interfaces. Additionally,

the debonding-tip stresses and related local deformation processes observed in field

debonding are not well represented in the various peeling metrologies [5,6,20]. In order

to quantify interfacial adhesion in solar modules, our group developed several mechanics-

based metrologies for selected module components [5,6,21–23]. These metrologies can

be used to accurately quantify interfacial adhesion, but they require active crack-length

monitoring and specimen machining.

We have now developed a simplified, mechanics-based adhesion metrology that is both

crack-length independent and requires minimal specimen preparation. To quantify

encapsulation adhesion (J/m2), we fabricated a Ti/glass/EVA/glass/Ti square adhesion

coupon, initiated a debond-defect between the EVA and glass at a corner of the coupon,

and applied two opposing tensile forces at that corner to propagate debonding. When

loaded by its corner, the coupon debonds at a constant mechanical force (Pc) and the

energy of adhesion (J/m2) can be easily calculated with Gc=Pc2/(4B), where B is the

54

cross-section bending stiffness of the bilayer (Ti/glass) substrate. We provide analytical

and graphical methods to calculate B even in the presence of glass surface cracking

during testing. We used this metrology to quantify the rapid increase of EVA/glass

adhesion (0.48 to 1.5 kJ/m2) during the first 9 min of cross-linking treatment at 140°C

(post-lamination). Adhesion then decreased with further exposure to the 140°C

environment. The technique is also extended to easily quantify adhesion in other module

laminates using both lab-made coupons and full-size solar modules. Finally, the

technique can be used to quantify loss of module adhesion due to environmental stressors

after field or accelerated exposures (moisture, temperature, UV), which provides the basis

for developing accelerated tests and protocols for solar module lifetime predictions.

4.3 Experimental Setup

Materials and adhesion coupons

To fabricate bilayer square substrates (Ti/glass), a 150m-thick cover glass

(Fisherbrand) was bonded to a 830 m-thick Ti substrate (Grade 5 alloy) with a high-

strength epoxy (Locite 420) under a spring-clamp pressure of ~4 kPa. To cure the epoxy,

the clamped substrates were then annealed for twenty-five minutes at 70°C. We selected

the thinnest possible Ti substrate (limited by the yield strength of Ti) to maximize the

elastic substrate deflection (and curvature) during the adhesion test. Large deflections are

necessary to deform the EVA past its rupture strain, and Ti provides the maximum elastic

deflection and elastic energy density of all available materials as described in the next

subsection. We also recommend the use of thin glass substrates (<150m) to avoid

internal severing of the glass (cohesive glass fracture) during the adhesion test.

To fabricate Ti/glass/EVA/glass/Ti adhesion coupons (Fig.4-1), a 500 m layer of 33%

wt. vinyl acetate-EVA encapsulation (the most common encapsulation in the PV

industry) was laminated between the glass surfaces of the bilayer (Ti/glass) substrates

described above, for 10 minutes at 120°C in a vacuum bag at 0.75 atm of barometric

pressure. Thermal stresses in the bilayer substrate can be neglected due to the similar

thermal expansion coefficients of Ti and glass (8.6 and 8.5 m·m-1K-1). Before

55

lamination, the glass surfaces were rinsed three times with acetone and isopropyl alcohol

and treated with 254nm UV light of 4 mW∙cm-2 intensity. To introduce a debond-defect,

a piece of Teflon tape was bonded to one of the glass substrates at a selected corner of

before lamination. Alternatively, a razor blade can be wedged at the debonding corner

after lamination.

Figure 4-1: Adhesion coupon loaded in tension by a corner. A sheet of encapsulation (EVA,

shown in white) was thermally bonded between two glass (150m) square substrates bonded to

secondary titanium (830m) substrates as shown. Debonding may be initiated by introducing a

debond defect in the debonding corner during the lamination step, or alternatively, by wedging a

razor blade after lamination. The titanium and glass substrates were bonded with a high-strength

epoxy.

The adhesion coupons were then annealed for 2, 9, 24 and 124 min at 140°C under no

applied pressure to determine the effect of EVA cross-linking on EVA/glass adhesion.

The 140°C heat treatment is commonly used in the PV industry for at least 10 minutes to

ensure complete cross-linking of the encapsulation layer.

Adhesion Metrology

To load the adhesion coupon by a corner, aluminum tabs were bonded to both Ti

substrates at the corner where the debond-defect was introduced. The loading tabs were

then connected to an adhesion test system (Delaminator, DTS, Menlo Park, CA)

consisting of a linear actuator (displacement-controlled, 0.1 um res.) in series with a high-

resolution mechanical load cell. The actuator was then retracted at a constant

56

displacement rate (developing a tensile force, P) to open up the corner and to propagate

debonding. The experiments were conducted in a chamber of controlled temperature

(60°C) and humidity (50%).

To quantify adhesion, the strain energy release rate, or debond driving force, G (J/m2),

can be calculated with the following expression[24]:

𝐺 =

𝑃2

2𝑏

𝑑𝐶

𝑑𝑎

(4-1)

where a is the debond-radius (distance from the loading corner to the propagating debond

line), b is the arc-length of the debond line, and C is the mechanical compliance

(displacement per unit force) of the square adhesion coupon loaded by a corner.

The value of C as a function of a was calculated using finite element analysis (Fig. 4-2,

square data points) and modeled as two 1-D cantilever triangular beams of length a (Fig.

4-2, solid line). The 1-D approximation was remarkably accurate in describing the value

of C. Using this approximation, the value of G for the adhesion coupon loaded by a

corner can be written as

𝐺 =

𝑃2

4𝐵

(4-2)

where B is the cross-section bending stiffness of the bilayer (Ti/glass) substrate given by

where Eg and ET are the Young’s moduli of glass and Ti, respectively, and t and h are the

glass and Ti substrate thicknesses, respectively. It is important to highlight that the value

of G in Eq. (4-2) does not depend on the value of a, unlike most other specimen

configurations that use Eq. (3-2). Debonding is therefore expected to occur at a constant

of value of P. When the glass substrate is much thinner than the Ti substrate, (𝑡/ℎ)2 << 1,

as is the case here, the value of B can be approximated with:

𝐵 =

𝐸𝑔2𝑡4 + 𝐸𝑇

2ℎ4 + 2𝐸𝑔𝐸𝑇𝑡ℎ(2𝑡2 + 3𝑡ℎ + 2ℎ2)

12(𝐸𝑔𝑡 + 𝐸𝑇ℎ)

(4-3)

𝐵 ≈ 𝐸𝑇𝐼𝑇 (1 + 𝑓

𝐸𝑔𝑡

𝐸𝑇ℎ)

(4-4)

57

where 𝑓 = 3 and IT is the bending stiffness (per unit width) of the Ti cross-section.

Figure 4-2: Mechanical compliance of the adhesion coupon loaded in tension by a corner. The 1-

D triangular beam model (—) accurately describes the 2-D finite element analysis data (□). The

1-D model consists of two triangular beams firmly fixed at the debond line. The line equation is

C=36 ∙a2/B, where B is the Ti+glass bending stiffness.

The glass substrate breaks during loading (surface cracking, parallel to the propagating

debond line) but remains perfectly bonded to the Ti substrate. Glass fracture is negligible

for the overall energy balance but can significantly reduce the value of B. To account for

this effect, the value of 𝑓 in Eq. (4-4) was calculated (with FEA) as a function of crack

spacing, s, normalized to the thickness of the glass substrate, t, as shown in Fig. 4-3. The

value of s is relatively uniform across the debonded substrates; the loaded coupon is

subject to uniform bending moments during testing. At large crack spacings (s/t >10), the

value of 𝑓 may still be approximated with the thin-substrate (no-crack) value of 𝑓 = 3.

At smaller crack spacings (s/t <10), the value of 𝑓 can be obtained from Fig. 3 using the

average crack spacing of the debonded glass substrate, measured with a micrometer or a

microscope. Even at very closely spaced cracks (i.e. s/t =5), the value of B, as calculated

with Eq. (4-4) and Fig. 4-3 (B=77), was similar to the no-crack value of Eq. (4-3) (B=79).

Alternatively, the value of B (at any crack spacing) can be calculated using the load-

displacement curve and the coupon dimensions as described below.

58

Figure 4-3: Glass substrate stiffening factor, f. The value of f describes the contribution of the

glass substrate to the glass/Ti bilayer bending stiffness, as derived in in Eq. (4-4). The value of f

is shown as a function of glass substrate crack spacing (normalized to glass substrate thickness, t)

for t=100 m (■) and t=150 m (□). The value of f depends strongly on crack spacing in

abundantly cracked glass (s/t<10), and may be approximated with the un-cracked thin-substrate

assumption f=3 at s/t>15. The fractured glass remained perfectly bonded to the Ti substrate in

the adhesion experiments.

A representative load-displacement curve for the debond test is shown in Fig. 4-3. The

value of P increased upon initial loading, plateaued during steady-state debonding at Pc,

and decreased abruptly when the debond line reached the middle of the coupon as

expected. To graphically calculate the value of B, we located the post-plateau inflection

point (Pc , c) as shown in Fig. 4-3. We then measured the post-debond corner aperture

(distance between the two glass substrates) with a micrometer and subtracted the initial

encapsulation thickness to calculate the zero load displacement,. The value of B can

then be calculated as

where l is the side-length of the coupon. In this study, l=50 mm. Alternatively, the value

of can be calculated by extrapolating the unloading load-displacement curve to P=0 N.

𝐵 =

𝑃𝑐

4(∆𝑐 − 𝛿)∙ 𝑙2

(4-5)

59

For example, the value of B for the adhesion experiment in Fig. 4-5 was 77 J.m when

calculated using 𝛿 =1.7mm in Eq. (4-5) and was 79 J.m when approximated using f = 5

and Fig. 4-3 in Eq.(4-4).

Figure 4-4: Load-displacement curve for EVA/glass debonding obtained by loading the square

adhesion coupon by a corner. The value of P increased with initial loading, plateaued during

steady-state debonding at Pc, and decreased abruptly at c when the debond line reached the

middle of the coupon. Adhesion (J/m2) can be calculated with Gc=Pc2/(4B), where B is the

bending stiffness of the bilayer substrate (Ti+glass). Glass fracture events during initial loading

are shown as sharp load-drops that do not affect the value of Gc. The value of B depends on crack

spacing as described in Eq. (4-4) and Fig. 3.

The encapsulation adhesion, Gc (J/m2) can finally be reported as the value of the debond-

driving force (G) where 𝑃 = 𝑃𝑐, namely,

𝐺𝑐 =

𝑃𝑐2

4𝐵 (4-6)

The value of Gc of interfaces of viscoelastic materials is often sensitive to the debonding

rate, and the actuator displacement rate should be reported when characterizing

encapsulation adhesion. Furthermore, when the coupon is loaded at a constant

displacement rate (at the corner aperture), the debonding rate (at the debond line)

decreases gradually as debonding propagates, which explains the slightly negative slope

of the Pc plateau in Fig. 4-5. To maintain the debonding rate constant (and to achieve a

60

flatter Pc), the corner-aperture displacement, ∆, may be increased with time using the

following quadratic equation:

∆= 𝑘 ∙ 𝑡2 (4-7)

where k is a proportionality constant and t is time. However, as shown in Fig. 4-4, using

constant displacement rates, ∆= 𝑘 ∙ 𝑡, may be sufficient to obtain accurate values of Pc.


The value of Gc (J/m2) of the EVA/glass interface is shown in Fig. 4-5 as a function of

heat treatment time at T=140°C. The initial value of Gc increased from 0.55 to 1.7 kJ/m2

(~300%) during the first 9 min of exposure to the 140°C treatment. This treatment,

together with peroxide additives[25,26], is commonly used (post-lamination) to promote

cross-linking of the EVA layer. Higher cross-link densities promote stiffening of the

encapsulation layer [27], which delays viscoelastic debonding [22], and increases the

value of Gc, as shown in Fig. 4-5 at t < 9 min. At longer exposures to the 140°C

environment, the value of Gc (J/m2) decayed gradually due to EVA-Glass interfacial bond

degradation. Similar thermally-activated bond degradation kinetics have been measured

in other PV laminates and polymer layers [5].

The value of Gc of the EVA/Glass interface depends on the strength of the covalent bond

(glass-EVA) but also on the deformation properties of EVA. In particular, the higher-

energy deformation processes of the EVA layer—including cavitation ahead of the

debonding-tip, and crazing and bridging behind the debonding-tip—can only occur

because of the strong covalent bonds of the glass-EVA interface. For example, the value

of Gc of the PVB/glass interface (~0.2 kJ/m2) reported elsewhere [22] is much lower than

that of EVA/glass because PVB/glass is bonded only via van der Waal interactions. It is

important to note that these processes are simply not present in the currently-available

peeling techniques.

61

Figure 4-5: Adhesion (Gc) of the EVA encapsulation bonded to glass as a function of heat

treatment time (140°C, post-lamination), measured with a square coupon loaded by a corner (Fig.

1). The value of Gc increases during the first 9 min of treatment due to EVA cross-linking and

decreases with further exposure to the 140°C environment. All specimens were previously

laminated at 120°C for 10 min.

4.5 Metrology variations

Using the same mechanics principles presented above, the square-corner metrology can

be adapted to measure adhesion of other module interfaces. For example, to measure

EVA/Si adhesion, a Si substrate (t<150m) can be used instead of glass to fabricate

Ti/Si/EVA/Si/Ti square coupons. To measure adhesion of encapsulation on metals or

other materials that exhibit large elastic deflections (e.g. PMMA or PET), a simpler

adhesion coupon can be fabricated, for example, Ti/Encapsulation/Ti, where the value of

B=IT in Eq. (4-5). Similarly, when depositing a thin silica film on the Ti substrate to

fabricate Ti/silica/EVA/silica/Ti coupons, the value of B can be approximated with IT; the

debonded silica/EVA interface can be representative of glass/EVA while avoiding glass

substrate cracking.

Two additional specimen configurations, the double cantiever beam and the single-sided

test can also be used to quantify adhesion. In the double cantilever beam specimen, the

62

debonding load decreases with the value of a, which allows for debond growth rate

monitoring. In the single-sided test, adhesion measurements can be done directly on full-

size modules with the square or beam substates. Both methods are described below.

Double Cantilever beam

Using a high-speed tungsten carbide saw, the adhesion coupon described in section 4.3.1

can be cut into several double cantilever beam (DCB) specimens, as shown in Fig. 6a.

The resulting DCB specimen requires active monitoring of the value of a to calculate the

value of the strain energy release rate as follows:

𝐺 =

𝑃2𝑎2

𝑤𝐵

(4-8)

where w is the width of the beam. The value of a can be measured optically, or calculated

with the slope (s) of a load-displacement curve using

𝑎 = √3𝐵

2𝑠

3

(4-9)

As in the square-corner test, G=Gc when P=Pc at debonding. In the DCB test, the value

of Pc decreases with a, as can be inferred from Eq. (4-8). To reduce the value of Pc at

debonding, the encapsulation material between the beams can be removed with a razor

blade to increase the value of a before the adhesion experiment. The DCB specimen can

also be used to measure the debonding kinetics of the EVA/glass interface: loading the

specimen under the value of Pc, fixing the displacement, and monitoring the load-

relaxation due to debond propagation.

Single substrate configurations

To measure backsheet adhesion directly on full-size modules, a Ti substrate (square or

beam) can be bonded to the solar module backsheet. The resulting single-sided specimen

(Ti/backsheet/module) can be loaded in tension as shown in Fig. 4-6.b. Single-sided

specimen fabrication is described in detail in [5] and is briefly summarized here: The Ti

substrates were bonded to the backsheet with a high-strength epoxy. In order to confine

63

debonding to the section of the backsheet directly under the Ti substrates, an incision was

made with a blade at the contour of the substrates (through the full thickness of the

backsheet and underlying encapsulant). The module was then firmly fixed to a testing

table, and the substrates were loaded in tension as shown in Fig. 4-6b. The value of G can

be calculated as half of that given by Eq. (4-2) and (4-8) for the square and beam

substrates, respectively. A bilayer (Ti/glass) beam can also be used in the single-sided

specimen configuration. The value of B does not depend on the specimen configuration.

The value of a in the single sided cantilever beam is given by 𝑎 = (3𝐵/𝑠)1/3. The Ti

substrate can be replaced by other materials that exhibit large elastic deflections,

including PMMA, which is sometimes easier to bond to other organic materials [5].

4.6 Conclusions

Encapsulation adhesion (J/m2) was quantified using a Ti/glass/EVA/glass/Ti square

coupon loaded by a corner. Debonding of the EVA encapsulation occurred at a constant

mechanical force (Pc), which eliminated the need for debond-length monitoring and

simplified specimen preparation and data analysis. Adhesion, Gc (J/m2), was calculated

with the simple equation Gc=Pc2/(4B), where B is the composite (Ti+glass) bending

stiffness, which can be calculated or measured directly. We used this metrology to

quantify the rapid gain of EVA/glass adhesion (0.55 to 1.7 kJ/m2) in the first 9 min of

140°C cross-linking (post-lamination). Similarly, we quantified the ensuing loss of

adhesion with further exposure to the 140°C environment. We extend this technique to

easily quantify adhesion of other module laminates and interfaces (i.e. encapsulation/Si)

using both lab-made coupons and full-size solar modules. The metrology presented here

can be used to compare encapsulation materials using an absolute (J/m2) scale and to

quantify loss of adhesion after field or simulated exposures, which provides the

framework to guarantee the structural stability of the solar module.

64

6a

6b

Figure 4-6: Adhesion metrology variations. a) Double cantilever beam (DCB) with bilayer

(Ti/glass) substrates. Using cantilever beams, load relaxation (at a fixed displacement) can be

monitored to calculate debonding rates (m/s) as a function of applied G (J/m2). b) Single-

substrate configurations of square and cantilever beam substrates for full-size module adhesion

testing. The Ti substrates were bonded to the backsheet, the module was firmly fixed to a testing

table, and the substrates were loaded in tension as shown. The backsheet and encapsulation at the

contour of the substrates was removed blade before testing with a razor blade—to isolate the

layers directly under the substrates.

65

4.7 References





materials for improved PV module reliability, Conf. Rec. Thirty First IEEE Photovolt.

Spec. Conf. 2005. (2005) 499–502. doi:10.1109/PVSC.2005.1488179.


(1975) 469–480. http://www.springerlink.com/index/Q5000TJ5V8830210.pdf (accessed

August 1, 2013).



doi:10.1016/j.actamat.2012.01.007.

[5] F.D. Novoa, D.C. Miller, R.H. Dauskardt, Environmental mechanisms of debonding in

photovoltaic backsheets, Sol. Energy Mater. Sol. Cells. (2013) 1–7.

doi:10.1016/j.solmat.2013.08.020.

[6] S.R. Dupont, F. Novoa, E. Voroshazi, R.H. Dauskardt, Decohesion Kinetics of

PEDOT:PSS Conducting Polymer Films, Adv. Funct. Mater. (2013) n/a–n/a.

doi:10.1002/adfm.201302174.




2012).



doi:10.1002/pip.







Photovolt. Spec. Conf. (2009) 001486–001491. doi:10.1109/PVSC.2009.5411336.

66




of Photovoltaic Modules, 2006 IEEE 4th World Conf. Photovolt. Energy Conf. (2006)

2050–2053. doi:10.1109/WCPEC.2006.279905.





[15] A.J. Kinloch, C.C. Lau, J.G. Williams, The peeling of flexible laminates, Int. J. Fract. 66

(1994) 45–70. doi:10.1007/BF00012635.

[16] N.E. Shephard, J.P. Wightman, A Simple Device for Measuring Adhesive Failure in

Sealant Joints, Sci. Technol. Build. Seals, Sealants Weather. Seventh Vol. ASTM STP

1334. (1998).

[17] K.-S. Kim, N. Aravas, Elastoplastic analysis of the peel test, Int. J. Solids Struct. 24

(1988) 417–435. doi:http://dx.doi.org/10.1016/0020-7683(88)90071-6.


test, Int. J. Fract. 93 (1999) 315–333. doi:10.1023/A:1007545200315.

[19] J.L. Gardon, Peel adhesion. I. Some phenomenological aspects of the test, J. Appl. Polym.

Sci. 7 (1963) 625–641. doi:10.1002/app.1963.070070219.



1907–1913. doi:10.1016/j.tsf.2010.08.116.


processed flexible inverted polymer solar cells, Sol. Energy Mater. Sol. Cells. 97 (2012)

171–175. doi:10.1016/j.solmat.2011.10.012.



[23] V. Brand, C. Bruner, R.H. Dauskardt, Cohesion and device reliability in organic bulk

heterojunction photovoltaic cells, Sol. Energy Mater. Sol. Cells. 99 (2012) 182–189.

doi:10.1016/j.solmat.2011.11.035.


67







http://onlinelibrary.wiley.com/doi/10.1002/apmc.1997.052520114/abstract (accessed April

5, 2012).

[27] F. Pern, S. Glick, EVA encapsulants for pv modules: Reliability issues and current

R&D status at NREL, Renew. Energy. (1996) 367–370.


2012).

68

5 CHAPTER 5: BACKSHEET ADHESION AND DEBONDING

5.1 Chapter Summary

The backsheets used in photovoltaic modules are exposed to aggressive field

environments that may include combined temperature cycles, moisture, and mechanical

loads. The effects of the field environment on backsheet debonding, which can lead to

module degradation (corrosion) and loss of function, are still not well understood or

quantified. Employing a newly developed quantitative mechanics technique, we report

the effect of aging on backsheet debond energy, including the separate effect of

temperature, mechanical stress and relative humidity on debond growth rate. The debond

energy of the backsheet decreased dramatically from 1000 to 27 J/m2 within the first 750

hrs of exposure to hot (85C) and humid (85% RH) aging treatments. The debond

growth rate increased up to 500-fold with small changes of temperature (10C) and

relative humidity (20%). To elucidate the mechanisms of environmental debonding, we

developed a fracture-kinetics model, where the molecular relaxation processes at the

debond front are used to predict debond growth. The model and techniques form the

fundamental basis to develop accelerated aging tests and long-term reliability predictions

for photovoltaic backsheets.

5.2 Introduction

Photovoltaic (PV) backsheets are exposed to aggressive field environments that may

include temperature cycles, moisture and complex mechanical loads. The field

environments present a significant challenge for the yield and reliability of PV modules,

eventually leading to loss of function [1–3]. The effect of the environment on backsheet

debonding is frequently reported[4,5] but is still not well understood or quantified. It is

well known, however, that the physical and chemical degradation mechanisms that

govern debonding in other structures are affected by moisture[6–8], temperature[9],

mechanical stress[10], and their synergistic interactions.[10–13] Characterizing the role

of these variables on PV backsheet debonding is important not only to predict backsheet

69

lifetime and to assess long-term module reliability but also to design certification tests for

aggressive operating environments.

Backsheet debonding occurs when the debond-driving force, G (J·m-2), which is a

function of the applied mechanical stress, is equal to the critical debond energy of the

backsheet, Gc. However, debonding in the presence of active environments can occur at

debond-driving forces below the critical debond energy, due to the effects of temperature

and the environment on the kinetic processes at the debond front, such as molecular

relaxation, chemical reactions, and water transport.

Backsheet debonding has been previously studied using techniques including: the 180°

peel[14–16] , 90° “T” peel tests[17–19], 45°peel test (load-controlled) [20], overlap

shear[21], and rotary shear[22,23] tests. These techniques have been used to characterize

the effect of aging treatments, such as damp-heat[21,24] or UV exposure[24], on the

adhesive properties of the backsheet. The peel tests, that presently dominate the study of

backsheet adhesion, may be subject to variability, repeatability, and interpretation of the

results. In some techniques, e.g., overlap shear, the applied mixed-mode load[25] can

render results irrelevant to the application [26]. In other cases, e.g., 180° peel, the applied

load may favor cohesive rather than interfacial failure. In most cases, the peeling process

involves not only interfacial failure, but also uniaxial plastic deformation of the peeling

tape, which is hard to quantify[27–29], and is irrelevant to the backsheet application.

Environmental conditions applied in the literature typically replicate the test sequences

from module qualification protocols[30,31]. The “UV exposure”, “thermal cycling”, and

“humidity freeze” tests simulate a brief field exposure (months to a few years at best); the

conditions applied in “damp heat” test greatly exceed those encountered in the

application [32]. These sequences at best query infant mortality, and do not of themselves

investigate the rate of degradation. The currently available adhesion tests, as well as the

aging treatments applied in the literature, provide limited insight on the kinetic processes

of debond growth and their dependence on mechanical stress, moisture and temperature.

We developed a quantitative characterization technique that may examine the effects of

mechanical stress, temperature, and moisture on the adhesion of backsheet and the

kinetics of backsheet debonding which we describe with a fracture kinetics model.

70

Employing a newly developed single cantilever beam (SCB) testing method[9], where

debonding is studied as a crack propagating between polymer layers[33–35], we

investigated the effects of indoor aging treatment on the debond energy of the backsheet

as a function of treatment temperature, humidity and duration. To investigate the kinetics

of debonding, we measured the backsheet debond growth rate, da/dt, using a load

relaxation technique in an environment of controlled temperature, T, and relative

humidity, RH. The technique allows debond growth rates, da/dt, to be accurately

quantified (10-8 to 10-4 m/s) in terms of the debond-driving force, G, and related to

molecular relaxation processes at the debond front. Debond growth was interpreted using

a fracture-kinetics model, which provides a more fundamental understanding of the

mechanisms of debond propagation and can be used to develop accelerated testing

techniques and to make long-term reliability predictions.

5.3 Materials and Methods

Backsheet specimens and materials

In order to fabricate representative backsheet specimens, a 0.45-mm layer of encapsulant

(Ethylene-vinyl-acetate, EVA) was laminated between float glass and a photovoltaic

backsheet, consisting of a layer of polyvinyl fluoride (PVF), a layer of polyethylene

terephthalate (PET), and a “seed” layer of EVA, which improves adhesion to the EVA

(Fig. 5-1a). Prior to lamination, the glass surface (Sn-deficient side) was carefully

cleaned, including: buffing with pumice powder; washing with an aqueous detergent

(Billco, Billco Manufacturing Inc.); rinsing with deionized water; and rinsing with

isopropyl alcohol. The lamination step was performed in a commercial instrument

(Model LM-404, Astropower Inc.) at 145C for 8 minutes under an applied pressure of 1

atmosphere. During lamination, the specimen components were fixed relative to one

another using Kapton tape and were surrounded with a glass frame to improve their

thickness uniformity at the edges.

71

1a

1b

1c

Figure 5-1: Backsheet-Encapsulant-Glass specimen structure. a) A 0.45-mm layer of Ethylene-

vinyl-acetate encapsulant (EVA) was laminated between a polyvinyl fluoride-polyester

backsheet with EVA seed layer and a 3-mm soda lime glass substrate. b) Single Cantilever beam

(SCB) specimen. Mechanical load is applied to one end of a PMMA beam (5x10x100mm),

bonded to the backsheet. c) Side view of the SCB specimen. Dotted lines represent 1mm-wide

cuts at the contour of the PMMA beam through the backsheet and encapsulant. Layer thicknesses

in the stacks are not drawn to scale.

Aging treatment methodology

Indoor aging of the backsheet specimens was performed in environmental chambers

maintained at the temperature and relative humidity conditions of 65°C and 0%RH,

Polyvinyl fluoride

PET

EVA Seed

EVA Encapsulant

Tempered Glass

a)

72

85°C and 0%RH, and 45°C and 85%RH for a treatment duration of 1000 hours, and at

85°C and 85%RH for treatment durations of 500, 750 and 1000 hrs. Although the aging

conditions and durations do not necessarily correspond to the operating environments of

PV applications, they are intended to invoke degradation of the polymeric materials or

represent worst-case environments. In particular, the “damp heat” treatment (85°C and

85%RH for 1000 hours) is specifically required during the protocols used to qualify

module products for sale [36,37].

Single cantilever beam test metrology

A single cantilever beam (SCB) testing metrology, based on the well-known double

cantilever beam method[38,39], was developed to quantify the debond energy of the

backsheet. The glass substrate of the backsheet specimens was fixed to a testing table

using a C-clamp. A Poly-methyl-methacrylate (PMMA) beam, 5-mm wide, 10-mm thick

and 100-mm long, was then bonded to the backsheet specimen with super glue (ethyl 2-

cyanoacrylate) and an accelerant (sodium bicarbonate) as shown in Fig. 1b. In order to

confine debonding to the section of the backsheet directly under the PMMA beam, an

incision was made with a blade at the edges of the PMMA beam (through the full

thickness of the backsheet and underlying encapsulant), as shown in Fig. 5-1c.

A loading tab with a ruby bearing was then bonded to one end of the beam to facilitate

mechanical loading. The loading tab was connected to an adhesion test system

(Delaminator, DTS, Menlo Park, CA) consisting of a linear actuator (displacement-

controlled) in series with a high-resolution mechanical load cell fixed to a high-stiffness

frame. The displacement of the linear actuator (perpendicular to the backsheet plane) and

the corresponding mechanical load were monitored every 0.1 seconds using a LabView

interface. To initiate backsheet debonding, a tensile force was applied on the loading tab

until the debond length (distance from the loading tab to the debond front) was ~15mm.

The PMMA beam in this configuration can be treated as a single cantilever beam (SCB)

fixed at the debond front. The stiffness of the SCB was much smaller than that of the

testing system, which allowed for accurate measurements. A minimum beam thickness

was required to avoid plastic deformation of the beam (due to bending stress) during the

debond energy experiment.

73

Debond energy experiment

The SCB specimen was mechanically loaded with the adhesion test system (described

above) in a laboratory environment at a fixed displacement rate of 10m/s. The applied

load, P, was recorded as a function of displacement to obtain a load-displacement curve.

The debond driving force, G, was calculated with the SCB expression[38]

𝐺 =

6𝑃2𝑎2

𝑏2𝐸′ℎ3

(5-1)

where b is the beam width, h the beam thickness, E’ the plane-strain elastic modulus of

PMMA and a the debond length, which can be optically measured or calculated with the

load-displacement curve as described elsewhere[9]. Finally, the critical backsheet

debond energy, Gc, is reported as the value of G at the onset of debonding, where the

load-displacement curve starts to deviate from linearity. The PMMA beam was not

plastically deformed during the experiment; the maximum bending stress in the beam was

much smaller than the yield stress of PMMA.

PVF deformation behavior

In order to measure the environmental sensitivity of the deformation behavior of the

polyvinyl fluoride film, microtensile specimens, 50 m-thick, 4 mm-wide and 60 mm-

long, were prepared. The specimens were attached to a displacement-controlled linear

actuator in series with a high resolution load cell to measure their stress-strain curves.

The experiments were conducted at two strain rates (2x10-6 s-1 and 2x10-4 s-1) and at four

values of RH (20, 40, 60 and 80%) in an environmental chamber. The specimens and

testing system were allowed to equilibrate for 10 hrs in the chamber environment before

testing.

Environmental debond growth

A detailed explanation of the load-relaxation method for debond growth is provided

elsewhere[9,38]; a brief summary is given here: To measure the backsheet debond

growth rate, da/dt, as a function of G, the SCB specimens were mechanically loaded in

the adhesion test system to a value of G close to the backsheet debond energy. The load

74

relaxation inherent in environmental debonding was then recorded as a function of time

for about 24 hrs. Post-experiment analysis of the load relaxation data determined the

debond growth rates, da/dt, over the range of ~10-4 to ~10-8 m/s as a function of G. The

debond growth experiments were performed on unaged encapsulation specimens in an

environmental chamber at selected values of relative humidity (20, 40, 60 and 80%) and

temperature (20, 40, 60 and 80°C). The load cell and actuator were allowed to equilibrate

in the chamber environment for 6-8 h before each experiment.


Backsheet debond energy

The backsheet debond energy, Gc, as a function of aging treatment temperature in

specimens aged for 1000 hrs in humid (RH = 85%) and dry (RH = 0%) environments is

shown in Fig. 5-2. The value of Gc was highest in the unaged specimens and decreased

with treatment temperature and relative humidity.

Figure 5-2: Backsheet debond energy vs. aging treatment temperature in dry (0%RH) and humid

(85% RH) environments. The debond energy decreased (indicating increased propensity for

failure) with elevated temperature and greater relative humidity. The arrow in the inset indicates

debond path location: Adhesive at the PVF-PET interface for the aging treatment at 85°C and

85% RH with partially cohesive failure occurring in the PVF layer for all the other treatments.

Each data point is the average of at least 10 measurements in ambient laboratory conditions using

the single cantilever beam technique. Layer thicknesses in the inset are not drawn to scale.

75

The backsheet debonded between the PVF and PET films in all specimens. In the

specimens aged at the highest temperature (85°C), the debond path was located strictly at

the PVF-PET interface. In contrast, in the unaged specimens and in the specimens aged

at low (45°C) and intermediate (65°C) temperatures, the debond path was partially

cohesive in the PVF film (remnants of the PVF film were left on the debonded PET

surface), as shown in Fig. 5-3a. The high values of debond energy in these specimens

can be explained by the extensive deformation required to leave the film remnants.

Accordingly, in the specimens with low debond energy (aged at 85°C) the debonded PET

surfaces were free of PVF (Fig. 5-3b).

Figure 5-3: PVF-PET debond path characterization. a) Photograph of the debonded PET surface,

representative of the unaged and moderately aged specimens (low aging time, temperature or

humidity). In these specimens, remnants of the PVF film (white spots) were left on the PET

surface after debonding. The extensive deformation of the PVF film required to produce this

fracture morphology corresponds to the high measured values of Gc. b) Photograph of the

debonded PET surface in aggressively aged specimens (85°C /85%RH for 1000 hours). Absence

of PVF residue corresponds to strictly interfacial failure at the PET-PVF interface and low

measured values of debond energy.

The value of Gc as a function of aging treatment duration in specimens aged at 85°C and

85% RH (damp-heat treatment) is shown in Fig. 5-4. The debond energy decreased

dramatically (indicating increased propensity for failure) with treatment duration within

the first 750 hrs from 1000 to 27 J/m2 and remained constant during the last 250 hrs. As

in the previous experiment involving aging temperature, the loss in debond energy

corresponded to a change in the backsheet debond path, which was partially cohesive in

the PVF in the specimens aged for 0 and 500 hrs and was strictly adhesive at the PVF-

PET interface in the specimens aged for 750 and 1000 hrs. Interestingly, the same

76

treatment (85°C and 85% RH), but at more prolonged exposures (2000 hours), has been

often reported to hydrolyze the PET layer[40,41].

PVF deformation behavior

The stress-strain curves of the PVF film, measured at two strain rates and at four values

of RH, are shown in Fig. 5-5. At the low strain rate (2x10-6 s-1) the yield stress and the

stress during plastic deformation decreased with RH as shown by the solid lines in Fig. 5-

5. The Young’s modulus, EPVF, also decreased with RH, but only at intermediate

humidity levels, from 2150 MPa at 40% RH to 1850 MPa at 60% RH. The value of EPVF

did not change significantly from 20% to 40% or from 60 to 80% RH. At the high strain

rate (2x10-4 s-1), the stress-strain curves were not affected by the value of RH, as shown

by the dotted lines in Fig. 5-5. To clarify, the softening of the PVF, demonstrating here

the effect of water-induced plasticization, is importantly most pronounced at the lesser

applied strain rate. Noticeably, the lesser applied strain rate is more common in the PV

application environment, e.g., following from thermal transient (~1 to 2·10-3 s-1) or

moisture swelling (~1·10-5 to 1·10-8 s-1) [42].

Figure 5-4: Backsheet debond energy vs. duration of aging treatment at 85°C and 85% RH. The

debond energy decreased to very low values (~27 J/m2) within the first 750 hrs of treatment.

Arrow in stacking diagram indicates debond path location: partially cohesive in PVF after 0 and

500 hrs of treatment and strictly adhesive at the PVF-PET interface after 750 and 1000 hrs. Each

data point is the average of at least 10 measurements in laboratory air conditions with the single

cantilever beam technique. Curve describing trend is best-fit of a straight line. The layer

thickness in the insets is not drawn to scale.

77

Figure 5-5: Stress-strain curves of the PVF film measured at two strain rates and at four values of

relative humidity, RH. At the lesser strain rate (solid lines) the yield and plastic deformation

stresses decreased with RH. The Young’s Modulus of the film decreased from 2150 MPa at low

RH (20 & 40%) to 1850 MPa at high RH (60 & 80%). At the high strain rate (dotted lines), the

stress-strain curves were not affected by the value of RH. The experiments were conducted on

microtension test specimens attached to a linear actuator in series with a high resolution load cell.

Environmental debonding

The debond growth rate, da/dt, as a function of the debond-driving force, G, measured at

four selected temperatures, is shown in Fig. 5-6. The debond growth rate increased with

G, displaying a concave-down debond growth curve. This curve was shifted to lower

values of G at higher temperatures, or conversely, the debond growth rate increased with

temperature at any given value of G, implying that backsheet debonding is a thermally

activated process. The effect of temperature on debond growth acceleration was

particularly strong at low temperatures; for example, the debond growth rate at a fixed

value of G=165 J/m2 increased 500-fold from 10°C to 20°C, but it only increased 100-

fold from 20°C to 30°C. To clarify, while PVF-PET debonding is most readily achieved

at elevated temperature, debonding is more sensitive to temperature changes at a lesser

temperature. The backsheet debonded between the PVF and PET films and exhibited

partially cohesive failure in the PVF.

Str

ess,

(M

pa)

Strain,

78

Figure 5-6: Thermally-activated backsheet debond growth rate vs. debond-driving force at four

selected environment temperatures. The debond growth rates were shifted to lower values of G

at higher temperatures. Model fits (10 and 20 °C) and model predictions (30 and 40°C) of the

debond growth rate are provided with a fracture-kinetics model using two temperature equations:

solid lines show the model using the Arrhenius equation and dotted lines show the same model

using the Williams-Landel-Ferry equation. Debonding occurred at the PVF-PET interface with

partially cohesive failure in the PVF layer. The experiments were performed with a single

cantilever beam load-relaxation technique in a chamber of controlled environment at 40% relative

humidity. Layer thicknesses in the inset are not drawn in scale.

The debond growth rate as a function of G, measured at four selected values of relative

humidity, is shown in Fig. 5-7. The debond growth curves were shifted to lower values

of G when the RH was increased from 40 to 60%. However, increasing RH from 20 to

40% or from 60 to 80% did not affect the debond curves significantly. Interestingly, the

measured value of EPVF in Fig. 5-5 had a similar dependence on RH, suggesting that

moisture-assisted debonding is controlled by the modulus of the PVF film. From the

observed behavior—where both the debond growth rate and EPVF were sensitive to RH

only in the 40-60% interval—, we expect no further changes in debond growth rate at

higher or lower RH values. As in the previous temperature experiment, the backsheet

debonded between the PVF and PET films and exhibited partially cohesive failure in the

PVF.

Deb

ond

Gro

wth

Rat

e, d

a/dt

(m/s

)


79

Figure 5-7: Moisture-assisted backsheet debond growth rate vs. debond-driving force at four

selected values of relative humidity. Solid lines show fracture-kinetics model fits (at 20 & 40%

RH) and model predictions ( at 60 & 80% RH) of the effect of moisture on debond growth rate.

The dotted line shows predictions (at 60 & 80% RH) of the fracture kinetics model, modified to

take into account the synergistic effects of moisture and debond growth rate. Debonding was

partially cohesive in the PVF layer. The experiments were performed with a single cantilever

beam load-relaxation technique in a chamber of controlled environment at 20°C. Layer

thicknesses in the inset are not drawn in scale.

Fracture Kinetics Model

We extended an existing polymer fracture kinetics model[43–45] to account for the effect

of G, temperature and relative humidity on debonding of the PVF-PET interface. The

central assumption of the original model is that crack growth in polymers occurs due to

crack-tip molecular relaxation processes, which occur at a constant crack tip opening

displacement[43,44] , c, given by:

𝛿𝑐 =

𝐺

𝜎𝑦

(5-2)

where y is the yield stress of the polymer. Rearranging Eq. 5-2 and assuming linear

viscoelastic behavior, where 𝜎𝑦 = 휀𝑦𝐸(𝑡, 𝑇), the value of G can be expressed as:

Deb

ond

Gro

wth

Rat

e, d

a/dt

(m/s

)


80

𝐺 = 𝛿𝑐휀𝑦𝐸(𝑡, 𝑇) (5-3)

where y is the yield strain of the material and 𝐸(𝑡, 𝑇) is the relaxation modulus. A

significant limitation of this model is the lack of moisture dependence. We therefore

adapted the model as follows to account for the effect of moisture on debond growth and

to include the presence of thin layers of PVF and PET.

Since the onset of plastic deformation in PVF occurs at stress values of ~30 MPa, which

are lower than in PET (~47 to 60 MPa), the deformation that occurs during debonding is

determined by the PVF film and not by the stronger PET film. This was further

supported by the debond energy experiments (section 5.3.1), where the debonded PVF

film exhibited extensive plastic deformation (Fig. 5-3a) but the PET film did not show

any visible damage. We therefore assume that the debonding process can be represented

by considering only the mechanical behavior of the PVF. In particular, we let 𝐸(𝑡, 𝑇) =

𝐸𝑃𝑉𝐹, where 𝐸𝑃𝑉𝐹 is the relaxation modulus of PVF.

The time dependence of 𝐸𝑃𝑉𝐹 at constant temperature, 𝑇 = 𝑇0, can be described with a

power-law equation, based on viscoelastic relaxation mechanics [43–45]:

𝐸𝑃𝑉𝐹 𝑇=𝑇0= 𝐸1𝑡−𝑛 (5-4)

where 𝑡 is the modulus relaxation time, E1 is the modulus at unit time, and n is a measure

of the rate sensitivity of the material (n=0 for perfectly elastic solids).

To account for the effect of temperature, we apply the time-temperature superposition

principle to Eq. 5-4 using T0 as the reference temperature:

𝐸1𝑡−𝑛 = 𝐸1 (

𝑡𝑇

𝑎𝑇)

−𝑛

(5-5)

where 𝑡𝑇 is the equivalent relaxation time that may be used in Eq. 5-4 to calculate 𝐸𝑃𝑉𝐹 at

any temperature, T, and 𝑎𝑇 is the time-temperature “shift” factor that can be described

with the Arrhenius equation:[43,46]

1

𝑎𝑇= 𝑒

−𝐸𝑎𝑅

(1𝑇

−1𝑇0

)

(5-6)

81

where Ea is the activation energy of the molecular relaxation process and R is the gas

constant.

Substituting Eqs. 5-4 and 5-5 into Eq. 3 results in:

The molecular relaxation processes that lead to Eq. 5-4 (and thus to Eq. 5-7) can also be

affected by moisture in the environment, which reduces the value of 𝐸𝑃𝑉𝐹 as shown in

Fig. 5-5. To incorporate this effect into the model, we let the value of 𝐸1 (which is

proportional to 𝐸𝑃𝑉𝐹) be a function of relative humidity, 𝐸1(𝑅𝐻), and assume that n is

independent of moisture. The value of 𝐸1(𝑅𝐻) can be estimated with the measured value

of 𝐸𝑃𝑉𝐹 as a function of relative humidity from the data in Fig. 5-5.

Similar to other fracture kinetic models[43–45], we assume the formation of a plastically

deformed zone in the PVF of size rp at the debond front:

We further assume that the yield strain, 휀𝑦, is independent of temperature and RH, and

that debond growth over a distance equal to rp occurs in the same time it takes to develop

the plastic zone, which may be approximated by the modulus relaxation time, 𝑡𝑇. [43–45]

The debond growth rate is therefore described by:

𝑑𝑎

𝑑𝑡=

𝑟𝑝

𝑡𝑇

(5-9)

Substituting 𝑡𝑇 and 𝑟𝑝 from Eq. 5-7 and 5-8, respectively, into Eq. 5-9, the debond growth

rate can be written as:

𝑑𝑎

𝑑𝑡=

𝜋

8

𝛿𝑐

휀𝑦(𝛿𝑐 휀𝑦)1

𝑛⁄(

𝐺

𝐸1(𝑅𝐻))

1𝑛 1

𝑎𝑇

(5-10)

Substituting Eq. 6 into Eq. 10, the backsheet debond growth rate with Arrhenius

temperature dependence is finally described by:

𝐺 = 𝛿𝑐휀𝑦𝐸1 (

𝑡𝑇

𝑎𝑇)

−𝑛

(5-7)

𝑟𝑝 =

𝜋

8

𝛿𝑐

휀𝑦

(5-8)

82

𝑑𝑎

𝑑𝑡 𝐴𝑟𝑟=

𝜋

8

𝛿𝑐


𝑛⁄(

𝐺

𝐸1(𝑅𝐻))

1𝑛

𝑒−

𝐸𝑎𝑅

(1𝑇

−1𝑇0

)

(5-11)

To determine the value of the kinetic variables (Ea, T0 and n), Eq. 5-11 was fitted to the

experimental data at 10 and 20°C in Fig. 5-6, and at 20 and 40% RH in Fig. 5-7. The

value of 𝐸1(𝑅𝐻) was assumed to be proportional to the moisture-dependent value of EPVF

in Fig. 5-5 (measured at the low strain rate). Eq. 11 was then used to make model

predictions (solid lines) at 30 and 40°C in Fig. 5-6, and at 60 and 80% RH in Fig. 5-7.

The model predictions are in good agreement with the experimental data, which indicates

that backsheet debond growth rate can be modeled with a thermally-activated relaxation

process, and that the effect of moisture on debond growth is determined by the modulus

of PVF. The model predictions, however, were less accurate for the high da/dt portion of

the data at 60 and 80% RH in Fig. 5-7. The deviation occurred because EPVF was

assumed to be moisture-dependent, but as shown in Fig. 5-5, it becomes moisture-

independent, and increases in value, at the high strain-rates corresponding to high values

of da/dt.

To account for the synergistic effect of strain rate and relative humidity on the value of

𝐸𝑃𝑉𝐹, we assumed that the value of 𝐸𝑃𝑉𝐹 at 60 and 80% RH increased (linearly) with

da/dt, from the moisture-dependent low-strain-rate measured value of 1850 MPa

(assumed to correspond to low da/dt~10-9 m/s) to the moisture-independent high-strain-

rate measured value of 2150 MPa (assumed to correspond to high da/dt~10-3 m/s). This

assumption implies that the effect of moisture on the relaxation processes that determine

debonding is stronger at low debond growth rates. The model predictions at 60 and 80%

RH with linearly-increasing 𝐸𝑃𝑉𝐹 improved the accuracy of the model as shown by the

dotted-line in Fig. 5-7.

A less empirical form of the debond growth rate equation can be obtained if, instead of

the Arrhenius model, the Williams-Landel-Ferry (WLF) model for time-temperature

superposition[47] is used to describe 𝑎𝑇:

𝑙𝑜𝑔(𝑎𝑇) =

−𝐶𝑎(𝑇 − 𝑇𝑔)

𝐶𝑏 + (𝑇 − 𝑇𝑔) (5-12)

83

where Ca=17.44, Cb=55.6 K-1 and Tg is the glass transition temperature of PVF.

Substituting 𝑎𝑇 from Eq. 12 into Eq. 10, the debond growth rate using the WLF equation

can then be described with

𝑑𝑎

𝑑𝑡 𝑊𝐿𝐹=

𝜋

8

𝛿𝑐


𝑛⁄(

𝐺

𝐸1(𝑅𝐻))

1𝑛

10

𝐶𝑎(𝑇−𝑇𝑔)

𝐶𝑏+(𝑇−𝑇𝑔)

(5-13)

which was fitted to the data at 20 and 30°C in Fig. 6 to obtain the value of Tg and n. The

value of Tg that gave the best data fit was -11°C, which corresponds to one of the

transition temperatures of PVF.[48,49] Eq. 5-11 was then used to make model

predictions at 10 and 40°C, which are shown as dotted lines in Fig. 5-6. The model

predictions are in good agreement with the experimental data.

The backsheet debond growth rate described here (Eq. 5-11 and 5-13) may be used to

predict backsheet lifetime, which may be defined as the time required for the growth of

an initial debond defect, typically due to manufacturing, to a final critical size, where loss

of function is anticipated.

5.5 Conclusions

A newly developed fracture-mechanics characterization technique was used to measure

the debond energy of a PV backsheet after several aging treatments. The debond energy

decreased with treatment duration, temperature and relative humidity. In the unaged and

moderately aged specimens, the debond path was partially cohesive in the PVF film. In

the aggressively aged specimens, the debond path was completely adhesive between the

PVF and PET films.

Using a load relaxation technique, the effect of mechanical stress, temperature, and

relative humidity of the testing environment, on backsheet debond growth rate was

quantified. The technique allowed the debond growth rates to be accurately determined

(10-8 to 10-4 m/s) and related to the molecular relaxation mechanisms at the debond front.

The value of the debond growth rate increased up to 500-fold with small changes in

temperature (10C) and relative humidity (20%). Backsheet debonding was accurately

modeled with fracture-kinetics as a crack propagating at the PET- PVF interface. The

84

debonding process was shown to be thermally-activated and the debond growth rate

dependence on temperature was modeled in terms of both an Arrhenius and a Williams-

Landel-Ferry based relaxation process. The debond growth dependence on moisture was

correlated to the mechanical modulus of the PVF film, including its variation with

relative humidity. For the aging conditions examined, debonding was not governed by the

PET or its hydrolysis. The techniques and models developed can be exploited to acquire

not only a fundamental understanding of debonding damage but also to develop

accelerating tests and make long term reliability predictions.

85

5.6 References


pottant: A critical review, Solar Energy Materials and Solar Cells. 43 (1996) 101–181.


module performance and durability following long-term field exposure, AIP Conference

Proceedings. 565 (1999) 565–571.


photovoltaic modules, Photovoltaic Specialists Conference, 2002. Conference Record of

the Twenty-Ninth IEEE. (2002) 1436–1439.



Photovoltaic Specialists Conference (PVSC). (2009) 001486–001491.

[5] P. Sánchez-Friera, M. Piliougine, J. Peláez, J. Carretero, M. Sidrach de Cardona, Analysis

of degradation mechanisms of crystalline silicon PV modules after 12 years of operation

in Southern Europe, Progress in Photovoltaics Research and Applications. 19 (2011) 658–

666.

[6] A. Bhatnagar, M. Hoffman, Fracture and Subcritical Crack Growth Behavior of Y-Si-Al-

O-N Glasses and Si3N4 Ceramics, Of the American Ceramic. 96 (2000) 585–596.

[7] E. Guyer, Effect of solution pH on the accelerated cracking of nanoporous thin-film

glasses, Journal of Materials Research. 20 (2005) 680–687.

[8] B. Lawn, An atomistic model of kinetic crack growth in brittle solids, Journal of Materials

Science. 10 (1975) 469–480.


debonding of copper/barrier interfaces, Acta Materialia. 60 (2012) 2219–2228.



1907–1913.

[11] M.S. Oliver, R.H. Dauskardt, Mechanical fatigue of hybrid glasses., Small (Weinheim an

Der Bergstrasse, Germany). 6 (2010) 1892–6.

[12] J.M. Snodgrass, D. Pantelidis, M.L. Jenkins, J.C. Bravman, R.H. Dauskardt, Subcritical

debonding of polymer/silica interfaces under monotonic and cyclic loading, Acta

Materialia. 50 (2002) 2395–2411.

86

[13] S. Wiederhorn, L. Bolz, Stress corrosion and static fatigue of glass, Journal of the

American Ceramic …. 1969 (1970).


assembly—Part 2: 180 degree peel, International Electrotechnical Commission: Geneva,

1–5. (1990).


of Photovoltaic Modules, 2006 IEEE 4th World Conference on Photovoltaic Energy

Conference. (2006) 2050–2053.


National Center for Photovoltaics and Solar Program Review Meeting. (2003).


transport, adhesion, and corrosion protection of PV module packaging materials, Solar

Energy Materials and Solar Cells. 90 (2006) 2739–2775.


assembly—Part 1: 90 degree peel, International Electrotechnical Commission: Geneva, 1–

6. (1990).

[19] A.J. Kinloch, C.C. Lau, J.G. Williams, The peeling of flexible laminates, International

Journal of Fracture. 66 (1994) 45–70.

[20] J.P. Shephard, N. E.; Wightman, A Simple Device for Measuring Adhesive Failure in

Sealant Joints, Science and Technology of Building Seals, Sealants and Weatherproofing:

Seventh Volume, ASTM STP 1334. (1998).


materials for improved PV module reliability, Conference Record of the Thirty First IEEE

Photovoltaic Specialists Conference 2005. (2005) 499–502.

[22] G. Jorgensen, T. McMahon, Accelerated and outdoor aging effects on photovoltaic

module interfacial adhesion properties, Progress in Photovoltaics: …. (2008) 519–527.

[23] N.G. Dhere, N.R. Raravikar, Adhesional shear strength and surface analysis of a PV

module deployed in harsh coastal climate, Solar Energy Materials Solar Cells. 67 (2001)

363–367.

[24] G. Jorgensen, K. Terwilliger, S. Glick, J. Pern, T. McMahon, Materials Testing for PV

Module Encapsulation, NREL/CP-520-33578. (2003).

[25] D.-A. Mendels, S.A. Page, Y. Leterrier, J.-A.E. Manson, A Modified Double Lap-Shear

Test as a Mean to Measure Intrinsic Properties of Adhesive Joints, in: Proc. Euro. Conf.

Comp. Mats., 2000.

87

[26] J.A. Marceau, Y. Moji, J.C. McMillan, A Wedge test for Evaluating Adhesive-Bonded

Surface Durability, Adhesives Age. (1977) 28–34.

[27] N.A. K-S. Kim, Elastoplastic analysis of the peel test, International Journal of Solids and

Structures. 24 (1988) 417–435.


test, International Journal of Fracture. 93 (1999) 315–333.

[29] J.L. Gardon, Peel adhesion. I. Some phenomenological aspects of the test, Journal of

Applied Polymer Science. 7 (1963) 625–641.


qualification and type approval ", International Electrotechnical Commission: Geneva, 1–

93. (2005).


type approval, International Electrotechnical Commission: Geneva, 1–81. (2008).




interconnect structures, Journal of Materials Research. 15 (2000) 2758–2769.

[34] M. Lane, R.H. Dauskardt, N. Krishna, I. Hashim, Adhesion and reliability of copper

interconnects with Ta and TaN barrier layers, Journal of Materials Research. 15 (2011)

203–211.


processed flexible inverted polymer solar cells, Solar Energy Materials and Solar Cells. 97

(2012) 171–175.


qualification and type approval ", International Electrotechnical Commission: Geneva, 1–

93. (2005).


type approval, International Electrotechnical Commission: Geneva, 1–81. (2008).


film structures, Engineering Fracture Mechanics. 61 (1998) 141–162.

[39] ISO 25217, Adhesives—Determination of the mode 1 adhesive fracture energy of

structural adhesive joints using double cantilever beam and tapered double cantilever

beam specimens”, International Electrotechnical Commission: Geneva, 1–24. (2009).

88

[40] M.D. Kempe, S.R. Kurtz, J.H. Wohlgemuth, D.C. Miller, M.O. Reese, A.A. Dameron,

Modeling of Damp Heat Testing Relative to Outdoor Exposure, Proc. Asian PVSEC,

Yokohama, Japan. (2011).

[41] W. McMahon, H.A. Birdsall, G.R. Johnson, C.T. Camilli, Degradation Studies of

Polyethylene Terephthalate, J. Chem. Eng. Data. 4 (1) (1959) 57–59.



(2010) 262–268.

[43] G. Marshall, L. Coutts, J. Williams, Temperature effects in the fracture of PMMA, Journal

of Materials Science. 9 (1974) 1409–1419.

[44] W. Bradley, W. Cantwell, H. Kausch, Viscoelastic creep crack growth: a review of

fracture mechanical analyses, Mechanics of Time-dependent Materials. I (1997) 241–268.

[45] R. Schapery, A method for predicting crack growth in nonhomogeneous viscoelastic

media, International Journal of Fracture. 14 (1978) 293–309.

[46] V. Gonda, K.M.B. Jansen, L.J. Ernst, J. Den Toonder, G.Q. Zhang, Micro-mechanical

testing of SiLK by nanoindentation and substrate curvature techniques, Microelectronics

Reliability. 47 (2007) 248–251.

[47] M. Williams, R. Landel, J. Ferry, The temperature dependence of relaxation mechanisms

in amorphous polymers and other glass-forming liquids, J. Am. Chem. Soc. 679 (1955).

[48] W. Reddish, Chemical structure and dielectric properties of high polymers, Pure Appl.

Chem. 5 (1962) 723–742.

[49] S. Osaki, S. Uemura, Y. Ishida, Effects of a static electric field upon dielectric properties

of poly (vinylidene fluoride) and poly (vinyl fluoride), Journal of Polymer Science Part A-

2: Polymer Physics. 9 (1971) 585–594.

89

6 CHAPTER 6: FRONT-SHEET DEBONDING UNDER UV LIGHT

6.1 Chapter Summary

Front-sheet module materials are employed in relatively hostile environments where they

are exposed to large amounts of solar UV radiation. In this chapter, we provide the first

quantitative characterization of the effect of UV on the molecular bond rupture kinetics

that determine cohesive cracking and adhesive debonding. We report the cohesive or

adhesive crack growth rate, da/dt, in terms of the applied strain energy release rate, G

(J/m2), of a transparent polysiloxane coating on quartz and borosilicate substrates. We

employ a load relaxation technique that allows debonding rates as low at 1 nm/sec to be

accurately quantified and related to the rupture of molecular bonds at the tip of the

propagating crack. To elucidate the degradation processes leading to environmental

debonding, the kinetics of crack and debond growth are interpreted using an atomistic

fracture mechanics model, which we modified to include the contribution of UV light on

bond rupture. The bond rupture processes investigated underlie the principal causes of

degradation in transparent organic materials exposed to terrestrial environments and can

be exploited to acquire, not only a fundamental understanding of damage formation and

progression, but also to develop accelerated testing techniques and make long term

reliability predictions.

6.2 Introduction

Transparent organic coatings for front-sheet module applications are employed in

relatively hostile environments where they are exposed to moist, extreme temperatures,

complex mechanical loads and even large amounts of solar UV radiation, which presents

a significant challenge for the yield and reliability of the solar module, potentially leading

to device loss of function [1–3]. Of particular interest is the effect of the environment and

UV light on the loss of interfacial adhesion and resistance to debonding in the presence of

residual stresses, thermomechanical cycling and mechanical or vibrational loading, which

has been frequently reported [1–5], but is still not well understood. Silicon solar cell

module failure has been repeatedly attributed to frontsheet degradation and debonding[6–

90

9], which occurs after extensive solar light exposure in environments of high humidity

and temperature, but which has not been fully characterized. Organic transparent films on

other devices such as aircraft windows[10], OLED displays[11], structural elements[12]

and MEMS[13] also suffer from similar degradation and debonding problems attributed

to UV light that have not been investigated yet.

Degradation of organic materials exposed to UV light has been traditionally studied by

aging the material with UV light of selected wavelengths to measure changes in such

properties as color[1,14], chemical composition[15], stiffness and viscosity[1]. While

this approach has been useful in evaluating the UV-resistance of a material, it provides

limited insight into the kinetics of the degradation processes and ignores the synergistic

interactions of multiple environmental agents such as moisture, temperature, oxygen and

mechanical stress. In particular, the interaction of UV light with the highly strained

material near the propagating crack tip that leads to crack growth and debonding is

completely unknown.

Understanding the physical and chemical degradation mechanisms that govern

environmental debonding under UV is necessary to assess long-term reliability, to predict

lifetime and to design certification tests for transparent organic films exposed to large

doses of UV. Characterizing the degradation mechanisms is also necessary for the

transferability of experimental results to larger systems, since the interactions of UV light

with the organic material are highly complex and vary among materials. For example, in

the presence of 02 or H2O, a UV-photoxidation reaction can degrade the bonding in some

organic materials[16,17], but UV cross-linking at specific wavelengths can actually

strengthen their inter-chain structure[18]. The effect of UV on bond stability can even be

temporary[19], where an excited molecular bond relaxes back to its original state unless

another environmental species attacks it while excited. The traditional approaches to

evaluating UV degradation in a material with such complex responses to UV would

provide very limited insight on its degradation mechanisms, especially when

characterizing coating debonding and crack growth, which depends strongly on the

synergistic interactions of all the environmental species and mechanical stresses[20,21].

91

Crack growth and adhesive debonding of films in the presence of environmental species,

such as moisture, temperature, pH has been studied using fracture mechanics, where the

crack growth rate is measured in the material or at the interface between two

materials[20–22][23,24]. This technique has been employed to characterize the effect of a

number of key parameters on crack growth and debonding, including interface chemistry

and morphology, plasticity in adjacent layers, and environmental factors such as

temperature and moisture[25–27]. However, the application of this fracture mechanics

technique to characterize the effect of UV light on debond or crack growth has not yet

been reported.

In the present study, we provide the first quantitative characterization of the effect of UV

on the molecular bond rupture kinetics that determine cohesive cracking and adhesive

debonding rate of transparent organic materials employing fracture mechanics. We report

the crack and debond growth rate, da/dt, in terms of the applied strain energy release rate,

G (J/m2), of a transparent polysiloxane coating on quartz and borosilicate substrates. We

employ a load relaxation technique that allows debonding rates as low at 1 nm/sec to be

accurately quantified and related to the rupture of molecular bonds at the tip of the

propagating crack. The effect of photon energy and light intensity on debond rate

acceleration is demonstrated, together with the existence of a well-defined debond growth

plateau rate. To elucidate the degradation processes leading to envrionmental debonding,

the kinetics of the crack and debond growth process are interpreted using an atomistic

fracture mechanics model, which we modified to include the contribution of UV light on

bond rupture. The bond rupture processes investigated underlie the principal causes of

degradation in transparent organic materials exposed to terrestrial environments and can

be exploited to acquire, not only a fundamental understanding of damage formation and

progression, but also to develop accelerated testing techniques and make long term

reliability predictions.

92

6.3 Experimental Procedures

Materials and Coatings

The transparent polysiloxane coating employed is a proprietary formulation consisting of

a Si-O-Si network, a methyl group and other organic flexible and hydrophobic end

groups. The coating was prepared using a sol-gel formulation based on tetraethoxysilane

and methyltriethoxysilane reactants, which were hydrolyzed with methanol to produce

the film precursor. The precursor was applied by flow-coating borosilicate and quartz

substrates (50x50x1 mm), and the methanol was allowed to evaporate for 90 seconds in

laboratory air. In order to prepare double cantilever beam (DCB) specimens, a blank

substrate was bonded to the deposited film by applying constant pressure at 80oC for 6

hours, forming two types of sandwich structures: glass/polysiloxane/glass and

quartz/polysiloxane/quartz. The final thickness of the polysiloxane film constrained

between the substrates was 30 m. Double cantilever beam (DCB) specimens, 5 mm in

width, 50 mm in length and 2.03 mm in total thickness (Fig. 6-1) were fabricated from

the sandwich structures with a high-speed wafer saw. Loading tabs were attached to the

end of the DCB specimens and a razor blade was inserted in between the substrates to

propagate a crack in the film for a few millimeters.

1a 2b

Figure 6-1: Schematics (a) and photograph (b) of debond growth experiment with in-situ UV

inside an environmental chamber of controlled environment. A mechanical loading system drives

the displacement ( of a double cantilever beam (DCB) to a fixed value and a sensor measures

debond propagation with in-situ UV light employing a load relaxation technique. The experiment

was conducted inside an environmental chamber at 20 deg C and 40% RH with cold-cathode UV

lamps.

93

Crack and debond growth testing

Cohesive crack growth (in the polysiloxane film) and adhesive debond growth (at the

polysiloxane-glass or polysiloxane-quartz interface) were characterized as a function of

the debond driving force, G(J/m2). A fracture mechanics load-relaxation technique was

employed, where the DCB specimens were loaded on a micromechanical test system

[DTS Mechanical Delaminator Test System, DTS Company, Menlo Park,CA] at a

constant displacement rate of 2μm/s up to a predetermined load at which the

displacement was fixed. Crack or debond growth took place for two days as the

mechanical load decreased and the DCB mechanical compliance increased. Post-

experiment analysis of the load relaxation and increasing compliance of the DCB

specimens determined the crack and debond growth rates over the range of ~10-4 to ~10-

10 m/s as a function of applied strain energy release rate, G, to produce a characteristic

da/dt vs G curve. More detailed descriptions of the general methods of crack and debond

growth measurements are described elsewhere[20].

UV and Environment accelerated debonding

Employing the DCB load relaxation technique described above, the crack/debond growth

rates were measured at 20oC and 40% RH with in-situ UV light of either 3.4eV or 4.9eV.

The light sources employed were cold-cathode quartz-tubing UV lamps (Fig. 6-1) fixed

at selected distances from the DCB specimen. The light intensity was measured with a

digital UV radiometer with interchangeable sensors [UVP Bioimaging Systems, CA].

Crack and debond growth experiments were conducted at intensities of 0.6 and 1.2

mW/cm2 with the 3.4eV UV light (365 nm wavelength) and at 2 and 6.6 mW/cm2 with

the 4.9 eV UV light (254 nm wavelength). All experiments were carried out in an

environmental chamber (Associated Environmental Systems, Ayer, MA) and the test

system was allowed to equilibrate in the chamber environment for 4-6 h before each

experiment. The temperature of the DCB specimen was continuously monitored. Heating

due to UV exposure was negligible.

94

Failure path characterization

The failure surfaces were characterized with Fourier Transform Infrared (FTIR)

spectroscopy on a Vertex 70 FTIR spectrometer (Bruker Optics Inc., Billerica, MA)

equipped with a micro ATR (A529-P MIRacle) accessory. Data collection and spectral

calculations were performed using OPUS (version 5.5) software. The spectra (2 cm-1

resolution) were obtained in the frequency range 1620-600 cm-1


Moisture Sensitivity

The adhesive debond growth rate (da/dt) of the polysiloxane/borosilicate interface is

shown in Fig. 6-2 as a function of the applied G at 1.4% and 40% RH. At high values of

da/dt, the value of da/dt exhibited an exponential dependence on the applied G. At low

values of da/dt, a debond growth threshold, Gth—under which debonding stops—was

measured at both values of RH. At high humidity (RH=40%), the debond growth curve

was shifted to lower values of G and higher values of da/dt.

Figure 6-2: Adhesive debond growth rate of a polysiloxane/borosilicate interface as a function of

debond driving force, G, at 1.4% and 40% relative humidity. The debond growth rate increased

and the debond threshold decreased with increasing relative humidity. A Maxwell-Boltzman

model was employed to calculate the H20 debonding reaction order coefficient (n=1.6). The

experiment was performed in an environmental chamber at 20°C.

95

Adhesive failure (borosilicate/polysiloxane) was confirmed with FTIR spectroscopy.

Reference spectra were obtained for the polysiloxane film and borosilicate substrate. The

polysiloxane film exhibited a distinctive peak at ~760nm corresponding to the Si-CH3

bond, and the borosilicate substrate exhibited a peak at ~950nm corresponding to the Si-

OH bond. Each side of the debonded DCB specimen exhibited only the spectra of either

polysiloxane or borosilicate, indicating perfectly adhesive failure, as shown in Fig. 6-3.

Figure 6-3: ATR-FTIR spectra showing Si-O-Si, Si-OH and Si-CH3 bond absorbances of

borosilicate and polysiloxane (references). The spectra of the debonded surfaces is also shown.

The spectra of each of the debonded surfaces matches either the polyxiloxane or the borosilicate

spectra, indicating adhesive failure.

The dependence of debond growth rate on G and RH can be described with an atomistic

bond rupture model[27] based on a Maxwell-Boltzman distribution of molecular bonds

breaking and healing along the debond front where the debond growth rate is given by

EGv

dt

da 2sinh0

(6-1)

Where v0 and are macroscopic debond growth parameters related to the chemical

environment and atomic arrangement; G is the debond driving force and 2E is the

intrinsic resistance to bond rupture of the material in the reactive environment given by:

))]ln(()[(2

2* OHABBE PRTnN

(6-2)

96

where N is the number of molecular bonds per unit area of the debond plane, R is the gas

constant, T is the absolute temperature,B is the chemical potential of the unreacted

bonds, B* is the chemical potential of the reactive complex, A* is the chemical

potential of the active species and n is the number of H20 molecules required in the H20

adsorption reaction at the crack tip:

nH2O + Si-O Si-O* Si-OH + -OH (6-3)

A detailed description of the atomistic model and RH dependence can be found

elsewhere[27,28].

The dependence of debond growth rate on G and RH is derived by combining Eq.1 and 3:

))ln()([sinh 2

0

n

OHPRTTNGv

dt

da (6-4)

which, for ideal has behavior of H20, at constant temperature and for values of G larger

than Gth, simplifies to:

GRHK

dt

da nexp (6-5)

where K is a proportionality constant. Eq. 6-5 was applied to the data in Fig. 2 and the

value of K, n and were fitted to the data at 1.4% RH and 40% RH. The best fit was

obtained with a value of n=1.6, which suggests that on average, 1.6 H20 molecules per

adhesive bond are needed for the debonding reaction (adhesive failure) to occur.

The cohesive crack growth rate in the polysiloxane film as a function of the applied G at

1.4% and 40% RH, is shown in Fig. 6-4. The crack growth curve was shifted to higher

values of da/dt in the high RH=40% environment. Eq. 6-5 can also be used to model the

crack growth data in Fig. 6-4. The value of K, n and were fitted to the data at 1.4% RH

and 40% RH. The best data fit was obtained with a value of n=1, which indicates that on

average, one H2O molecule per cohesive bond is needed for the crack growth reaction

(cohesive failure) to occur. The value of n=1 is frequently reported for other interfaces

where Si-O rupture is the limiting reaction[24].

97

Figure 6-4: Cohesive crack growth rate of polysiloxane as a function of debond driving force, G,

at RH=1.4% and 40%. The crack growth curve was shifted to higher values of da/dt and slightly

lower values of G at the high RH=40% environment. An atomistic Maxwell-Boltzman model

was employed to find the H20 debond reaction order coefficient (n=1). The experiments were

performed in an environmental chamber at 20°C

UV sensitivity

The cohesive crack growth rate of the polysiloxane film as a function of applied G,

measured with in-situ UV light of 3.4 eV and in the dark, is shown in Fig. 6-5. The

presence of UV light did not affect the crack growth curves. The polysiloxane film is

transparent to UV-light of this photon energy.

The adhesive debond growth rate of the polysiloxane/borosilicate interface is shown in

Fig. 6-6 as a function of applied G, measured with in-situ UV light of 3.4eV and in the

dark. Surprisingly, in the presence of UV light, the debond growth rate curves (adhesive

failure) were shifted to lower values of applied G, including the threshold region.

Possible mechanisms for the effect of UV on adhesive debond growth of the

polysiloxane/borosilicate interface are proposed here. First, the transport rates of the

active species (H2O and OH-) may be higher under UV light; Second, the UV absorption

of the material at the polysiloxane/borosilicate interface may change under high

mechanical stresses[24] –similar to the effect of high pressures in other materials[29];

98

And third, UV light may form radicals in the polysiloxane material that can speed up the

bond rupture reaction[26]. Regardless of the specific mechanism in which UV light

affects the transparent material, we can reformulate the atomistic model of Eq. 1 to

include the UV contribution to the debond growth process.

Figure 6-5: Crack growth rate in the polysiloxane film as a function of debond driving force, G,

with in-situ UV of 3.4eV and in the dark . No significant difference was observed in the crack

growth rates of the experiments with and without UV. The experiment was performed in an

environmental chamber at 20°C and 40% relative humidity.

We propose that, when the material is radiated with the UV photon flux, , a fraction of

the arriving photons, 2f , is absorbed near the debond front, and a bonding electron is

excited into a metastable state that decays after a lifetime, ,losing its energy by emission

or by thermal processes, similar to the absorption-decay processes in semiconductors[30].

Using the mass-action law, the area density of the excited electrons in the material,phn , is

proportional to and to the interfacial bond density, N

Nfnph 2 (6-6)

99

Figure 6-6: Debond growth rate of a polysiloxane/glass interface as a function of debond driving

force, G, with in-situ UV of 3.4eV at two selected intensities and in the dark (0 mW/cm2). The

debond growth rates increased and the debond thresholds decreased with increasing UV light

intensity (I). An atomistic Maxwell-Boltzman model (solid lines) was fitted to the data to predict

the measured debond growth rates with in-situ UV of 1.2 mW/cm2 (dashed line). All experiments

were performed in an environmental chamber at 20°C and 40% relative humidity.

The excess energy absorbed in the material is then proportional to hv , and the total energy

available for debonding can be expressed as:

hvnfE phUV 1 (6-7)

where the constant 1f is a measure of the energy absorption efficiency of the material.

Substituting Eq. 6 in Eq. 7 yields the contribution of UV to the debond driving force:

INCE hvUV

(6-8)

Where the constant 21 ffChv is a measure of the energy transfer efficiency from

the free photon to the breaking bond and I is the light intensity is defined as hvI

The UV driving force,UVE , can then be incorporated in Eq. 1 as follows:

Ehv INCGv

dt

da 2sinh0

(6-9)

100

which was applied to the experimental debond growth data in Fig. 6-6. The values of 0v ,

and E were fitted to the data in the dark and the value of

hvC was fitted to the data

with in-situ UV of 0.6 mW/cm2 (Fig. 6-6, solid lines). Eq 6-9 was then used to predict

the debond growth rates with in-situ UV at the higher intensity of 1.2 mW/cm2. The

prediction (Fig. 6-6, dashed line) is in good agreement with the measured growth rates.

Figure 6-7: Debond growth rate of a polysiloxane/quartz interface as a function of debond driving

force, G, with in-situ UV of 4.9eV at two selected intensities and in the dark (0 mW/cm2). The

debond thresholds decreased with increasing UV light intensity. The debond growth rates with

in-situ UV described a G-independent plateau at intermediate values of G. The data was fit to an

atomistic Maxwell-Boltzman model (solid lines) to predict the measured debond growth rates

with in-situ UV of 6.6 mW/cm2 (dashed line) and to explain the formation of the debond rate

plateaus. All experiments were performed in an environmental chamber at 20°C and 40%

relative humidity.

The debond growth rates of the polysiloxane/quartz interface, measured as a function of

G with in-situ UV light of 4.9 eV and in the dark are shown in Fig. 6-7. The data for the

experiment in the dark exhibited a debond threshold and an exponential debond growth

region, similar to the experiments in Figs. 6-4 to 6-6. However, under 4.9 eV

illumination, the debond growth curves also exhibited a debond plateau (values of da/dt

independent of G). The low da/dt portion of the data, including the debond threshold, was

shifted to lower values of G in the presence of UV. To model this shift, Eq. 6-9 was fitted

101

to the data in the dark and under UV light of 2 mW/cm2 (Fig. 6-6, solid lines). The

prediction at 6.6 mW/cm2 (dotted-line) is in good agreement with the low da/dt portion of

the data.

Figure 6-8: Cohesive crack growth rate of in polysiloxane as a function of debond driving force,

G, using new and pre-exposed specimens (4.9 eV). Crack growth in the pre-exposed specimens

occurred at higher values of G than in the pristine specimens due to increase in the cross-link

density. All experiments were performed in an environmental chamber at 20°C and 40% relative

humidity.

UV-Crosslinking

The cohesive crack growth rate of the polysiloxane film as a function of applied G is

shown in Fig. 6-8 for specimens pre-exposed to UV-light of 4.9eV (dose ~ 3000 kJ/m2)

and for pristine specimens. Crack growth in the pre-exposed specimens occurred at

higher values of G than in the pristine specimens, which corresponds to the higher

polysiloxane bond density (N) due to UV cross-linking, as verified with FTIR

spectroscopy (Fig. 6-9). Eq. 6-9 can be modified to account for the increasing cross-link

bond density:

102

Ehv INCG

v

dt

da2

sinh0

0

(6-10)

where = uvN /

0N is the ratio of the post-UV bond density to the initial bond density.

Eq. 6-10 was fitted to the crack growth rate data of the pre-exposed and pristine

specimens in Fig. 6-8. The value of 5.1 indicates that the bond density increased

50% over the length of this experiment.

Figure 6-9: FITR spectra of polysiloxane film as a function of exposure to UV light of 4.9 eV.

The Si-O-Si peak increases with exposure to UV. All data is normalized to the the Si-CH3 peak.

6.5 Conclusions

The first quantitative characterization of the effect of UV on the molecular bond rupture

kinetics was described. We reported the cohesive or adhesive crack growth rate, da/dt, of

a transparent polysiloxane coating on quartz and borosilicate substrates, in terms of the

applied strain energy release rate, G (J/m2). Debond growth of the polysiloxane coating

occurred at significantly lower values of G with UV light. Debond growth of

polysiloxane films that were previously exposed to 4.9eV irradiance, however, occurred

103

at higher values of G due to increased cross-linking. To elucidate the degradation

processes leading to environmental debonding, the kinetics of the crack and debond

growth process were interpreted using an atomistic fracture mechanics model, which we

modified to include the contribution of UV light and cross-linking on bond rupture.

The bond rupture processes investigated here underlie the principal causes of degradation

in transparent organic materials exposed to terrestrial environments and can be exploited

to acquire, not only a fundamental understanding of damage formation and progression,

but also to develop accelerated testing techniques and make long term reliability

predictions.

104

6.6 References




2012).


module performance and durability following long-term field exposure, AIP Conf. Proc.

565 (1999) 565–571. doi:10.1063/1.58001.





[4] M. Jørgensen, K. Norrman, F.C. Krebs, Stability/degradation of polymer solar cells, Sol.


[5] A. Hinsch, J. Kroon, R. Kern, Long‐term stability of dye‐sensitised solar cells, Prog.

Photovoltaics Res. Appl. 438 (2001) 425–438. doi:10.1002/pip.397.

[6] F. Pern, S. Glick, EVA encapsulants for pv modules: Reliability issues and current

R&D status at NREL, Renew. Energy. (1996) 367–370.


2012).



[8] W.H. Jr, S. Agro, J. Galica, Investigation into the causes of browning in EVA

encapsulated flat plate PV modules, Photovolt. Energy. (1994) 893–896.


[9] H. Wenger, C. Jennings, Carrisa Plains PV power plant performance, Photovolt. Spec.

Conf. 1990., Conf. Rec. Twenty First IEEE. 2 (1990) 844–849.




1907–1913. doi:10.1016/j.tsf.2010.08.116.

105

[11] R. Seifert, S. Scholz, B. Lussem, K. Leo, Comparison of ultraviolet- and charge-induced

degradation phenomena in blue fluorescent organic light emitting diodes, Appl. Phys. Lett.

97 (2010) 013308. doi:10.1063/1.3460285.

[12] M. a. Said, B. Dingwall, a. Gupta, a. M. Seyam, G. Mock, T. Theyson, Investigation of

ultra violet (UV) resistance for high strength fibers, Adv. Sp. Res. 37 (2006) 2052–2058.

doi:10.1016/j.asr.2005.04.098.

[13] R. Truckenmiiller, P. Henzi, D. Herrmann, V. Saile, W.K. Schomburg, Bonding of

polymer microstructures by UV irradiation and subsequent welding at low temperatures,

Microsyst. Technol. 10 (2004) 372–374.




http://onlinelibrary.wiley.com/doi/10.1002/apmc.1997.052520114/abstract (accessed April

5, 2012).

[15] F. Pern, A. Czanderna, K. Emery, Weathering degradation of EVA encapsulant and the

effect of its yellowing on solar cell efficiency, Photovolt. Spec. Conf. 1991., Conf. Rec.

Twenty Second IEEE. (1991) 557–561.


[16] A. Heller, The Role of Oxygen in Photooxidation of Organic Molecules on Semiconductor

Particles, Society. 0 (1991) 5261–5267.

[17] O. Chiantore, L. Trossarelli, Photooxidative degradation of acrylic and methacrylic

polymers, Polymer (Guildf). 41 (2000) 1657–1668.

http://www.sciencedirect.com/science/article/pii/S0032386199003493 (accessed April 5,

2012).

[18] T.-S. Kim, N. Tsuji, N. Kemeling, K. Matsushita, D. Chumakov, H. Geisler, et al., Depth

dependence of ultraviolet curing of organosilicate low-k thin films, J. Appl. Phys. 103

(2008) 064108. doi:10.1063/1.2894727.

[19] J.M. Taguenang, a. Kassu, P.B. Ruffin, C. Brantley, E. Edwards, a. Sharma, Reversible

UV degradation of PMMA plastic optical fibers, Opt. Commun. 281 (2008) 2089–2092.

doi:10.1016/j.optcom.2007.12.031.




interconnect structures, J. Mater. Res. 15 (2000) 2758–2769.

106

http://journals.cambridge.org/production/action/cjoGetFulltext?fulltextid=7998189


[22] M. Lane, R.H. Dauskardt, N. Krishna, I. Hashim, Adhesion and reliability of copper

interconnects with Ta and TaN barrier layers, J. Mater. Res. 15 (2011) 203–211.

doi:10.1557/JMR.2000.0033.

[23] S. Dill, S. Bennison, Subcritical Crack Growth Behavior of Borosilicate Glass under

Cyclic Loads: Evidence of a Mechanical Fatigue Effect, J. Am. 76 (1997) 773–776.

http://onlinelibrary.wiley.com/doi/10.1111/j.1151-2916.1997.tb02897.x/abstract (accessed

March 21, 2012).

[24] S. Wiederhorn, L. Bolz, Stress corrosion and static fatigue of glass, J. Am. Ceram. ….

1969 (1970). http://onlinelibrary.wiley.com/doi/10.1111/j.1151-

2916.1970.tb15962.x/abstract (accessed August 7, 2013).

[25] A. Bhatnagar, M. Hoffman, Fracture and Subcritical Crack Growth Behavior of Y-Si-Al-

O-N Glasses and Si3N4 Ceramics, Am. Ceram. 96 (2000) 585–596.

http://onlinelibrary.wiley.com/doi/10.1111/j.1151-2916.2000.tb01237.x/abstract (accessed

March 21, 2012).

[26] E. Guyer, Effect of solution pH on the accelerated cracking of nanoporous thin-film

glasses, J. Mater. Res. 20 (2005) 680–687. doi:10.1557/JMR.2005.0079.


(1975) 469–480. http://www.springerlink.com/index/Q5000TJ5V8830210.pdf (accessed

August 1, 2013).

[28] S. Wiederhorn, E. Fuller, Micromechanisms of crack growth in ceramics and glasses in

corrosive environments, Met. Sci. 14. (1980).

http://www.ingentaconnect.com/content/maney/msc/1980/00000014/F0020008/art00019


[29] A. Jayaraman, Diamond anvil cell and high-pressure physical investigations, Rev. Mod.

Phys. (1983). http://rmp.aps.org/abstract/RMP/v55/i1/p65_1 (accessed April 9, 2012).

[30] S. Mayer, J. Lau, Electronic Materials Science : For Integrated Circuits, Macmillan

Publishing Company, New York, 1990.

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