Intermetallic compounds in solar cell interconnections: Microstructure and growth kinetics

Solar Energy Materials & Solar Cells 159 (2017) 370–388

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Solar Energy Materials & Solar Cells journal homepage: www.elsevier.com/locate/solmat

Intermetallic compounds in solar cell interconnections: Microstructure and growth kinetics Torsten Geipel a,n, Monja Moeller b, Johann Walter a, Achim Kraft a, Ulrich Eitner a a b

Fraunhofer Institute for Solar Energy Systems ISE, Heidenhofstrasse 2, 79110 Freiburg, Germany Christian-Albrechts-University Kiel, Christian-Albrechts-Platz 4, 24118 Kiel, Germany

art ic l e i nf o

a b s t r a c t

Article history: Received 27 April 2016 Received in revised form 26 August 2016 Accepted 29 August 2016

The microstructure and intermetallic phase growth in solder joints of photovoltaic modules are investigated because of their significance for interconnection reliability. Interconnector ribbons with Sn60Pb40, Sn62Pb36Ag2, Sn43Bi57, Sn41Bi57Ag2 and Sn91Zn9 are soldered on the front busbars of industrial crystalline silicon solar cells. Cross sections are inspected using microscopy, SEM and EDX. The interconnections are isothermally aged, whereas the intermetallic layer thickness is determined successively. The microstructural changes in the bonds are characterized. Grain coarsening, volume increase of intermetallic compounds, Sn-penetration into the metallization and growth of grain boundaries between the phases are found. The composition of the intermetallic phases within Sn91Zn9-bonds is discussed. A diffusion model is used to simulate the intermetallic layer growth after 3000 h at 85 °C and thermal cycling from
40 °C to 85 °C for 600 cycles. A prognosis of the phase growth within the photovoltaic module after 25 years at the location Freiburg in Germany is made. It is found that the Ag3Snphase using Sn43Bi57 extends to 6.6 mm after 3000 h at 85 °C as compared to 2.6 mm for Sn60Pb40. After 25 years in Freiburg the Ag3Sn-layer within Sn43Bi57 joints is predicted to be 2.4 mm whereas within Sn60Pb40-bonds only 2.2 mm. & 2016 Elsevier B.V. All rights reserved.

Keywords: Photovoltaic module Interconnection Soldering Intermetallic compounds

1. Introduction Crystalline silicon photovoltaic module manufacturing consists of two main processes: the series connection of solar cells to strings and the lamination of the cell matrix into glass and encapsulation materials [1]. The first is realized by soldering, whereby ribbon interconnectors consisting of a copper core and a pre-tinned solder alloy coating are soldered on both electrodes (busbars) of the solar cell, which, in the standard H-pattern design, are positioned on front and rear side respectively. The copper core and the solder coating in the as-supplied state is shown in Fig. 1. The copper core usually varies from 0.7 mm to 1.5 mm in width and from 0.12 mm to 0.22 mm in thickness. The chosen dimensions depend on the size of the solar cell and the number of ribbons per cell. The solder alloy coating has a thickness of 10 mm–30 mm. The busbars are made of screen-printed and sintered thick-film paste with silver as the main component, inorganic binders and an organic vehicle [2,3]. The busbars are shaped as continuous or quasi-continuous tracks on the sunny side and usually interrupted pads on the rear n

Corresponding author. E-mail address: [email protected] (T. Geipel).

http://dx.doi.org/10.1016/j.solmat.2016.08.039 0927-0248/& 2016 Elsevier B.V. All rights reserved.

side. A cross section of the solar cell and a detail of the front busbar is shown in Fig. 2. The process of series interconnection involves applying flux on the busbars or ribbons and aligning the interconnectors on the front and rear side of neighboring cells. Then, while holding down, heat is transferred into the stack to melt the solder. Upon cooling the solder solidifies and forms metallurgical bonds to the busbars. Finally, the interconnected cells are cooled down in a controlled manner to allow the solder to relax thermo-mechanical stress. The whole process is automated in tabber-stringers [1]. In many respects, the design and quality of the interconnection plays a critical role for mechanical yield in production as well as conversion efficiency and long-term stability of the photovoltaic module. Coarsening of the solder and dislocation of the solder compounds as well as fatigue cracks within the interconnection have been found in photovoltaic modules after long-term outdoor exposure [4,5]. A particular concern in solder bonds, which is often linked to those failure modes, is the existence and growth of intermetallic compounds (IMC) at the interfaces between substrates and solder [6–8]. The brittleness of intermetallic layers in combination with thermal and mechanical stress is assumed to cause crack formation and propagation within the layer or at the interface. The growth kinetics of IMCs and its effects on adhesion

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Fig. 1. Cross section of an interconnector ribbon with a Sn60Pb40 solder coating in the condition as supplied. (a)–(c) Increasing level of detail of the microstructure.

Fig. 2. a) Cross section of a solar cell in the condition as supplied, b) Detail of the screen printed and sintered Ag front busbar of the solar cell in the unsoldered state.

properties within solar cell interconnections have gained attention recently. The necessity to investigate and control the IMC-growth is stressed by Nieland et. al. [9]. Moreover, Schmitt et. al. [10] correlated the thickness of intermetallic layers with adhesion strength of Sn96.5Ag3.5-soldered solar cell interconnections. Sndiffusion into the busbar is assumed to cause a loss of adhesion of cell metallization to the wafer. Jung et. al. characterized the kinetics of intermetallic phase growth of Sn60Pb40-soldered solar cell ribbons on thermally evaporated Al/Ni : V/Ag-stacks. An activation energy for Cu6Sn5 and Ni3Sn-phases respectively are derived in the study. Additionally, the IMC thickness is estimated using realistic temperature histograms [11]. A similar work is made by Kumm et. al. in which Ni : V is replaced by Ti and TiN [12]. A particular focus on the growth kinetics of the Ag3Sn intermetallic phase at the busbar metallization is placed by Yang et. al. [13]. Three common solders (Sn63Pb37, Sn62Pb36Ag2 and Sn96.5Ag3Cu0.5) are investigated during isothermal aging. Activation energies are reported for each solder as well as a prognosis of the long-term growth is made. Faes et. al. analyzed the IMC formation between In48Sn52-coated thin wires which are used to interconnect busbar-less solar cells [14]. A study of the existing intermetallic phases using less typical solders based on SnxZny is done by Chen et. al. [15]. A starting point for us to improve the long-term stability of interconnections in photovoltaic modules is to increase and ascertain the understanding of the solder microstructure initially and after isothermal aging as well as to characterize existing intermetallic phases and their growth kinetics. In this regard, we recognize that kinetic parameters to describe growth kinetics are either missing (as in [10]) or limited to one specific interconnection stack or to an individual IMC (such as in [11] or [13]). And,

from our knowledge, the IMC-growth for low-temperature solders on busbars, relevant for the interconnection of temperature-sensitive solar cells, has been widely omitted so far. Consequently, the aim of our work is to characterize the solder microstructure initially and after isothermal aging using optical microscopy, scanning electron microscopy (SEM) and energy dispersive x-ray spectroscopy (EDX). We extend the analysis to a total of five solders including leaded standard solders Sn60Pb40 and Sn62Pb36Ag2, low-temperature compatible unleaded solders Sn43Bi57 and Sn41Bi57Ag2 as well as the unleaded Sn91Zn9, which could be a replacement for Sn60Pb40 in the future. We report kinetic parameters for all solders and estimate phase growth under different conditions.

2. Theoretical background 2.1. Solder bonds of solar cells Soldering means that two or more materials are joint using a third interjacent material which is able to melt and solidify at significantly lower temperatures than the substrate melting temperatures [16]. The formation of the solder joint may be described by the consecutive stages [16,17]: 1. Heating of surfaces and solder until the melting temperature of the solder is exceeded. 2. Removing of oxides by a flux and hence permitting direct contact of the molten solder to the substrate. 3. Wetting and spreading of the molten solder on the surfaces. 4. Dissolving of the first atomic layers of the substrates into the

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Fig. 3. a) Interconnection of a solar cell and a Sn60Pb40 coated ribbon after soldering. b) The same interconnection thermally aged for 500 h at 130 °C in nitrogen atmosphere.

solder and diffusion of solder into the substrate. 5. Nucleation and solidification of intermetallic phases at the interfaces between substrates and solder as well as within the solder matrix. 6. Cool down and solidification of the molten solder. Fig. 3a shows a solar cell interconnection with Sn60Pb40 right after the soldering process in a cross sectional view using an optical microscope. The solder matrix consists of Pb-rich (dark) and Sn-rich phases (bright). The smaller bright spots within the Pbrich phases are secondary Sn-rich crystals precipitating during the cool down from the eutectic temperature to room temperature [16]. The internal structure of the solder matrix after the interconnection process strongly depends on the cooling rate. A fast cooling results in a fine grain structure with high interfacial areas and non-equilibrium solute concentrations. The structure forms larger grains already in the first 48 h at room temperature after the soldering process [18]. Generally, the solder microstructure is a result of the soldering and the cooling process as well as time and temperature of subsequent aging. The formation of intermetallic compounds using SnPb and SnBi-based solders is briefly described as follows. The reaction of the molten solder at the interface to Cu involves dissolution of Cu into the molten solder. The rate of dissolution is ≈1 μm min−1 at 200 °C. Once the local equilibrium solubility of Cu is exceeded in the near-field solder, the IMC starts to form [19]. The result is an approximately 0.5 mm thick Cu6Sn5 compound layer ( η′-phase) of scallop-type morphology [20]. Ag is dissolved into molten solder at a considerably higher rate of approximately 6 μm min
1 at 200 °C.1 The interface of solder to sintered silver busbar contains a 0.5–1.2 mm thick Ag3Sn compound layer. Excessive soldering temperature and time can result in a blocky morphology with needle-like or plate-like protrusions [21,22]. The formation of thin, uniform and continuous intermetallic layers at the interfaces are an indication for a proper metallurgical bond [17,23]. However, intermetallic compounds grow during prolonged soldering times and storage at elevated or even room temperature [24]. To explain the commonly observed changes

during thermal aging the same solar cell interconnection is shown in Fig. 3b after isothermal aging at 130 °C for 500 h. The aged interconnection is characterized by a coarsening of the solder matrix resulting in fewer and larger grains [18]. The driving force is a reduction of the Gibb's free energy by reducing the overall surface area of the grains. As can be seen, the IMC-layers grow to a thickness of 1.5 mm Cu6Sn5 and 10 mm Ag3Sn. However, the specific extent of phase growth depends on the solder type. An additional phase of Cu3Sn ( ϵ -phase) is visible in between Cu6Sn5 and Curibbon with a thickness of 1 mm. The growth of the Cu3Sn-phase can occur at the expense of the Cu6Sn5-phase [24]. Kirkendall voids can sometimes be observed within the Cu3Sn-phase [8]. 2.2. Growth kinetics of intermetallic compound layers Growth and ripening of intermetallic layers during solid-state aging is determined by diffusion of atoms across the compound layer and by the time needed to rearrange the atomic structure at the substrate-IMC interface. The flux of atoms during diffusion is slowed down with time due to the increasing IMC thickness resulting in a characteristic parabolic shape of the thickness versus time curve. If the compounds are sufficiently thick, a control of layer growth kinetics solely by diffusion is assumed [25]. According to a random walk treatment of solid state diffusion starting from the one dimensional solution of Fick's second law, the mean square diffusion distance x 2 of atoms after time t is calculated with Eq. (1) [26].

x 2 = 2Dt

(1)

D is the temperature dependent diffusion coefficient, for which the empirical Arrhenius relationship holds as given in Eq. (2).

⎛ Q ⎞ D = D0exp⎜ − ⎟ ⎝ RT ⎠

(2)

with D0 as the species specific pre-exponential factor, Q the activation energy, R is the gas constant and T the absolute temperature. Assuming that the intermetallic layer thickness x equals to the root mean square diffusion distance, Eq. (1) is rearranged to yield x as:

x=

2Dt + x 0

(3)

1

Due to the high dissolution rate of Ag into molten solder “solder scavenging” or “solder leaching” – the detachment of thick-film conductors during soldering – is a known concern in hybrid thick-film technology [2,16]

The symbol x0 describes the initial layer thickness. Different approaches exist to describe the intermetallic phase

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growth. Regularly, a simplified growth model of the form Kt is used, where K is a temperature dependent growth constant [16].2 Sometimes, a generalized power law of the form x = At n + B with variable time exponent n is used [28]. Hence, care must be taken if parameters from different authors need to be compared as they depend on the underlying growth model. The combination of Eqs. (2) and (3) gives:

x=

⎛ Q ⎞ 2·D0 ·exp⎜ − ⎟·t + x 0 ⎝ RT ⎠

(4)

In order to obtain the activation energy Q, Eq. (4) is linearized by taking the natural logarithm which results in Eq. (5).

⎞ ⎛ ⎞ 1 ⎛ Q 1 · ln⎜ x − x 0⎟ = ln⎜ 2D0t ⎟ − ⎝ ⎠ 2 ⎝ ⎠ 2R T

(5)

We approximate the expression ln( x − x0) over 1/T by a straight line with the slope m so that:

Q = − 2mR

(6)

Although Eq. (5) would allow the extraction of D0, we experience a large error using this method. Hence, we decide to perform a non-linear fit of the measured intermetallic layer thickness to Eq. (4) instead using the activation energy from Eq. (6) as a constant parameter. This approach leads to more accurate kinetic doublets. Once having the activation energy the determination of the time t1 for aging at elevated temperatures T1 needed to accelerate the phase growth within a time t2 at the temperature T2 as given by Eq. (7) is possible.

⎡Q ⎛ 1 1 ⎞⎤ t1 = t2exp⎢ ·⎜ − ⎟⎥ T2 ⎠⎦ ⎣ R ⎝ T1

(7)

Finally, a measure to quantify the quality of the modeled data with respect to the observed data is needed. We use the mean absolute error (MAE) for this purpose [29]. N

MAE =

∑i = 1

()

x ti − x^i si

N

(8)

x^i is the measured thickness at time step i. The symbol x( ti) is the modeled thickness at the time ti, N is the total number of measurements, and si is the standard deviation of the measurements at time step i. A MAE close to 0 mm denotes a good fit of the model to the experimental data since the absolute errors between observed thickness to calculated thickness are small. 2.3. Properties of solder alloys, intermetallic compounds and constitutive elements Since, the solder alloys used in the experimental study are Sn60Pb40, Sn62Pb36Ag2, Sn43Bi57, Sn41Bi57Ag2, and Sn91Zn9 the characteristics of each solder alloy and its intermetallic compounds are briefly reviewed in the following section. 2.3.1. Sn60Pb40 and Sn62Pb36Ag2 Although the “Restriction of the use of certain hazardous substances (RoHS)” bans lead from electronic equipment in the European Union, which effectively promoted the transition to lead2 It can easily be seen that K = 2D such that both approaches are compatible. Nevertheless, very often the symbol D is mistakenly used to express the growth constant such as in [6,24,27] but in order to avoid confusion it should be reserved to express the diffusion coefficient. Additionally, the treatment of x0 remains unclear in many instances.

373

free solders in electronics, photovoltaic modules are still excluded from this legislative even in the latest revision [30]. Regardless of lead's toxicity, Sn60Pb40 and Sn62Pb36Ag2 solder alloy coatings are almost always used in today's photovoltaic module production. This is because the inclusion of lead in solders has many technical and economical advantages such as improved wetting of surfaces, mechanical ductility, prevention of tin-pest and low costs [31]. These leaded alloys have a melting temperature range of 183– 189 °C and 179 °C respectively. An addition of Ag lowers the dissolution rate of silver substrate material in the molten solder [32]. However, owing to the need of further cost reductions Sn62Pb36Ag2 is practically suppressed by the cheaper Sn60Pb40 solder nowadays. 2.3.2. Sn43Bi57 and Sn41Bi57Ag2 The Sn43Bi57 and Sn41Bi57Ag2 solder alloys possess a lower melting point than SnPb(Ag) solders of 139 °C. Due to this fact, they serve as a means to reduce the buildup of thermo-mechanical stress during and after the soldering process of solar cells with a thickness of below 180 mm [33]. Since wafer thickness did not decrease as expected in the past, the motivation to switch to this material was not huge. However, the emergence of amorphous silicon/crystalline silicon heterojunction solar cells with efficiencies of above 24% could make the development of low temperature interconnection methods necessary because they contain temperature sensitive layers that degrade if exposed to temperatures above 200 °C [34]. In this regard Sn43Bi57 and Sn41Bi57Ag2 are promising candidates for low temperature soldering. Nevertheless, serious technical concerns exist to use bismuth containing solders in photovoltaic applications. Bi is a relatively brittle material compared to Sn or Pb [35]. Due to the lower melting temperature, SnBi(Ag) must operate at a high homologous temperature Th, which is defined as Th = To/Tm with To being the operating temperature and Tm being the melting temperature of the solder. A high homologous temperature accelerates segregation and the growth of large Bi-phases. These large and brittle Biphases can generate early fracture at high strain rates such as thermal shocks or mechanical impact. The observed strain-rate dependent brittleness can, however, be reduced significantly by the addition of Ag, that leads to a finer grain structure during solidification [36–38]. Moreover, Bi shows a negative thermal expansion when melted (and expands when solidified), thus it is speculated that Sn43Bi57 matrices are susceptible to extended grain boundary growth and eventually cracking. In fact, Sn43Bi57 has shown a lower fatigue life as compared to Sn63Pb37 solders [39]. Regarding the processing of the material, concerns exist because SnBi(Ag) remelts during typical lamination cycles of photovoltaic module production, which require temperatures of more than 150 °C for a few minutes. Although tests by Lalaguna et. al. showed no detrimental effect on electrical bond properties, this concern is not sufficiently investigated yet [33]. The use of Sn43Bi57 has recently gained more attention in the effort to interconnect solar cells using the “Smart Wire Technology”. In this field, these solders shall replace more expensive In58Sn42 solder alloy coatings [14]. 2.3.3. Sn91Zn9 The eutectic composition of Sn91Zn9 is sometimes considered as a replacement for lead-based solders mainly due to the low melting temperature of 198 °C, which is close to that of Sn60Pb40 and due to the low price of the material. It also possess excellent tensile, fatigue and creep properties [40]. As already mentioned in the introduction, first tests of photovoltaic interconnector ribbons with SnxZny alloy coating have been done highlighting the good peel strength and low series resistance [15].

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Table 1 Material properties of constitutive elements and IMCs. Denotion of symbols: E… Young's modulus, σ… electrical conductivity, ρ… density.

Sn Pb Bi Zn Ag Ag-busbar Cu Si Ag3Sn Cu6Sn5 Cu3Sn Cu5Zn8 AgZn Ag5Zn8

E [MPa]

s[1  106 S/m] ρ[g/cm3] CTE[10-6K-1] Refs.

41.6 14 32 96.5 76 5.74–44.25 110 112.4 78.9–99 112.3–134.7 134.2–160 170 96 97

8.7 4.8 0.89 16.9 64.5 33.3 58.8 – 0.56–9.71 5.7 11.2 – – –

5.76 11.35 9.8 7.1 10.49 8.58 7.76 2.3 – 8.28 8.9 – – –

23.8 29.1 13.3 31.2 19.6 10.4 16.4 2.5 – 16.3 19.0 – – –

[49] [49] [49] [49] [49] [50,51] [49] [49] [42,43,46,47] [46–48] [46–48] [52] [52] [52]

However, severe drawbacks exist such as poor wettability [41] and low oxidation resistance, which is attributed to the diffusion of Zn to grain boundaries and forming ZnO in particular under hot and humid environments [42]. Some of its properties can be improved by the addition of additional elements such as Al, Bi or Ag [43]. The intermetallic compounds and interdiffusion zones of Sn91Zn9 formed within the solder matrix and at the interfaces to the substrates when bonded to Cu and Ag are of greater complexity than the before mentioned systems because of the high number of possible alloy compositions that can exist simultaneously at room temperature. The binary system of Cu and Zn describes the following solid alloys at the relevant temperature with approximate elemental concentrations [44]: α -Cu consisting mainly of Cu with up to 37 at% dissolved Zn, β′-CuZn with a relatively well-defined stoichiometric ratio of 45–50 at% Zn and 50– 55 at% Cu, γ -CuZn (Cu5Zn8), for which 59 at% Zn to 69 at% Zn is possible, ϵ -CuZn with a Zn-content between 78 at% and 87 at% Zn, η -Zn which is mainly Zn with a few percent dissolved Cu. The solid phases in the Ag-Zn-system are no less complex [45]. Possible phases are: α -Ag consisting mainly of Ag with 0–40.2 at% dissolved Zn, ζ -AgZn with 37–51.2 at% Zn, γ -AgZn (Ag5Zn8) with 58.5–64.7 at% Zn, ϵ -AgZn (AgZn3) with a Zn-content between 66.2 at% and 89 at% and η -Zn with possibly 5 at% dissolved Ag. Owing to this complexity of solid phases a quantitative analysis of the actual composition with EDX leaves room for interpretation. Here, we briefly summarize the observations from other authors using Sn91Zn9 solder alloy on Cu and Ag substrates. At the interface to the Cu, primarily a planar γ -Cu5Zn8 intermetallic layer is formed. This takes place instead of the formation of Cu6Sn5, which is caused by the lower Gibb's free energy of CuZn-IMCs compared to CuSn-IMCs and the higher diffusivity of Zn in the molten solder [46,47]. In the presence of silver, additional ϵ -AgZn3 particles agglomerate at the interface [48]. It is observed that the solder-silver interface comprises of a bilayer stack of γ -Ag5Zn8 and, at the substrate side, of the ζ -AgZn phase. Additionally, ϵ -AgZn3 is claimed to exist non-uniformly at the interface as well as in the near-field solder matrix [48]. 2.3.4. Physical properties Table 1 summarizes physical properties of the constitutive elements and intermetallic phases of solder alloys used in this study taken from previous research.

Fig. 4. Semi-automatic soldering station with solar cell and ribbon on the heating chuck.

Table 2 Dimensions of the ribbons and processing conditions. Denotion of symbols: Tsold… set soldering head temperature, Tchuck…set heating chuck temperature, tsold…soldering time, tdown…down holding time. Sn60Pb40 Sn62Pb36Ag2 Sn43Bi57 Sn41Bi57Ag2 Sn91Zn9 Dimensions

1.5  0.2

[mm2] Tsold [°C] Tchuck [°C] tsold [s] tdown [s]

1.5  0.2

1.5  0.18

1.5  0.18

1.2 4.0

1.5  0.2

250

230

200

200

300

175

170

135

135

195

1.2 4.0

1.2 4.0

1.2 4.0

3.0 4.0

3. Experimental approach In order to investigate the microstructural properties of solar cell interconnections and their changes during thermal aging, solar cells are first interconnected using ribbons with different solder alloy coatings on the front busbars. Afterwards, the solar cells are trimmed and the pieces are embedded as well as polished to form cross section samples. After characterizing the initial microstructure, the cross section samples are isothermally aged and the thickness of the intermetallic layers is monitored successively. The so obtained data is processed according to the approach in Section 2.2 to extract the kinetic parameters. With the help of the kinetic parameters, predictions of the intermetallic layer growth during environmental chamber testing and under real ambient conditions during 25 years module life are made. 3.1. Materials and methods 3.1.1. Soldering process The soldering of the ribbons onto the solar cells is realized with a semi-automated soldering station from the company Somont as shown in Fig. 4. The soldering station consists of a hot plate and three rows of soldering heads. Each row has twelve soldering irons that transfer heat into the ribbons and cell by direct contact. The ribbons are fluxed with Kester 952S manually and positioned on the busbars of the cell. After 30 s warm-up time on the hotplate the soldering heads and downholders are lowered automatically for a defined time. For this study, we use the same row of soldering heads for each soldered busbar to enhance the reproducibility and uniformity of the soldering joints throughout the specimen. The ribbon dimensions and processing parameters are given in Table 2. The solar cells are commercial available mono-crystalline AlBSF cells with a thickness of 190 mm and three continuous busbars of 1.5 mm width and ≈20 μm height on the front side and six soldering pads per row on the rear side.

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aged at 85 °C, 100 °C, 115 °C, 130 °C due to the lower melting point. Isothermal aging is done in a furnace under nitrogen atmosphere. Microscopic images are taken in the initial state which is 24 h after the soldering process and after 15 h, 85 h, 155 h of isothermal aging. We age the embedded cross section samples and therefore we use a single cross section sample for each series of aging times. This technique allows the measurement of the phase growth at almost the same position. Nevertheless, cleaning and polishing of the cross section samples is needed after each aging step.

Fig. 5. Laser confocal microscope image of a cross section of an interconnection with Ag3Sn-IMC and the measuring lines for thickness determination.

3.1.2. Metallographic preparation, microscopy and spectroscopy In order to inspect the solder joints a part of the soldered busbar ( ≈1.5 cm × 1.0 cm ) is cut by a precision saw, embedded into a nickel-filled epoxy, ground and polished [53]. Optical and laser scanning confocal microscope images are made with a LEXT OLS4000 from Olympus. The laser image is used to measure the thickness of the IMC-layers. Three images are taken of each cross section sample. The thickness of the IMCs is measured in the image every 3 mm with the aid of parallel guides generated by the computer software of the microscope to a total of 30 measurements per cross section as indicated in Fig. 5. The arithmetic mean of these 30 measurements per time and temperature step is used as a predictor of the IMC-growth. The SEM images are obtained with a cross beam workstation Auriga 60 from Zeiss, and EDX spectroscopy is done in the same tool using a Bruker Quantax XFlash 6|60 detector. 3.1.3. Aging parameters The IMC-growth is examined for each solder at four different temperatures. The Pb and Zn containing solders are isothermally aged at 100 °C, 115 °C, 130 °C, 150 °C. The Bi-containing material is

Fig. 6. Time-temperature curve of a photovoltaic module in Freiburg, Germany over the course of one year from June 01, 2012 to May 15, 2013. The red curve shows a moving average for easier readability.

3.1.4. Temperature data for phase growth modeling The necessary mathematical theory for modeling the phase growth is described in Section 2.2, which can easily be applied to predict phase growth under isothermal conditions. Since we want to make a projection of the phase growth in a photovoltaic module under more realistic outdoor conditions for the life time of 25 years, a time-temperature curve from a real photovoltaic module installed in Freiburg, Germany is considered (see Fig. 6). This one year record starts on June 01, 2012 and ends on May 15, 2013. The hourly averaged temperature for the course of one year is reiterated 25 times and used as an input for a numerical solver algorithm to calculate the resulting layer thickness.

4. Results 4.1. Microstructural changes induced by the soldering process Interconnector ribbons are pre-tinned by the ribbon manufacturer before they are supplied to the PV module producer usually by dipping the Cu-core into molten solder. A certain microstructure of the solder matrix and initial IMC morphology at the interface from solder to Cu-core is the result of this pre-tinning process. Especially, the duration of the dip tinning process as well as the cooling phase afterwards influence the initial microstructure. The solder coating is remelted during the interconnection process of the solar cells such that the microstructure becomes entirely different from the as-delivered state of the ribbon. Furthermore, due to the elevated temperature during the process, already existing IMCs expand and new intermetallic layers or interdiffusion zones are formed at the interfaces from solder to the metallization of the solar cell according to the mechanisms described in Section 2.1. The changes of the solder alloy coating from the as-delivered condition to the state after the soldering process of the solar cell are investigated for Sn62Pb36Ag2, Sn41Bi57Ag2 and Sn91Zn9. Respective SEM images of the cross sections before and after the soldering process are given in Fig. 7. In the as-supplied state, the Sn62Pb36Ag2 solder matrix is characterized by a homogeneous distribution of Pb-rich regions with a diameter of 1–10 mm within a Sn-phase. Cavities with a diameter in the range of ≈1 μm are observed in the as-supplied Sn62Pb36Ag2 matrix, that lead to a porosity of the solder coating. The Ag-content is bonded as Ag3Snparticles, which are preferentially found at the interface to the Cucore. Occasionally, they are also observed within the matrix. After the soldering process, the microstructure of the Sn62Ag36Ag2 solder appears significantly refined. The Pb-rich phases have a diameter of 1–3 mm. The refinement of the solder matrix is caused by the remelting of the as-delivered solder coating during the interconnection and the relatively fast cooling rate after the joining process. In the case of the Sn41Bi57Ag2 solder, the pre-tinned coating of the ribbon shows a uniform distribution of Sn-rich and Bi-rich phases. As opposed to the Sn62Pb36Ag2 solder, whose initial Snrich and Pb-rich regions tend to exist as isolated areas, the Sn-rich and Bi-rich phases in the as-delivered Sn41Bi57Ag2 solder form

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Fig. 7. SEM images of cross sections of ribbons with solder coating in the as-delivered state and after the soldering process on the front side busbar of a solar cell.

large conjunct regions. A 0.5 mm thick and uniform Cu6Sn5-layer is observed in the as-supplied condition, whose morphology obtains the typical scallop-like structure after the soldering process. A particularity is the agglomeration of Ag3Sn-particles at the interface to the Cu-core. The particles have diameters of up to 5 mm. The solder matrix of Sn41Bi57Ag2 is refined after the soldering process likewise the Sn62Pb36Ag2 solder. The result is a solder

layer with a few large conjunct Bi-areas and 0.5–2 mm isolated Biregions in the Sn-phase. The diameters of the Ag3Sn-particles, attached to the Cu-core, increase slightly. The Sn91Zn9 solder matrix in the as-delivered condition has Zn-rich particles with a diameter of usually below 1 mm. Occasionally, they possess a needle-like shape as reported in Ref. [15]. In between Cu-core and solder matrix, a 1–2 mm Zn-rich layer is

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Fig. 8. Metallographic cross sections of thermally aged soldered solar cell interconnections.

found, which is assumed to be a ZnO-layer in regard to the observations described in Section 4.3.3. The ZnO-layer may be the result of a non-optimal pre-tinning process since Sn91Zn9 is no standard material for a ribbon manufacturer.

The microstructure of the Sn91Zn9 solder matrix and the structure of the intermetallic layers strongly changes after the soldering process. Ag3Sn(Zn)-particles are observed in the solder layer, which are a result of the partial dissolution of the Ag-busbar

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during prolonged soldering. A complex structure of IMCs and interdiffusion zones, including β′-CuZn, Cu5Zn8 and Ag-Sn-Zn is found at the interface to the Cu-core. Moreover, a ≈2 μm thick Ag3Sn(Zn)-layer is observed at the interface to the busbar.3 A detailed analysis of the intermetallic compounds and interdiffusion zones is found in Section 4.3.3. As a conclusion from these investigations we want to highlight that the remelting of the pre-tinned solder coating during the interconnection process leads to a change of the microstructure and grain distribution within the solder matrix. In this regard, soldering time and cooling rate during the interconnection process mainly influence the resulting microstructure. Furthermore, the partial dissolution of the Ag-busbar during soldering can affect the microstructure of the solder layer since the growth of Ag3Sn-particles at the interface to the Cu-core (in particular Sn41Bi57Ag2) or within the solder matrix (Sn91Zn9) is observed. In the following sections, whose focus is on the effects of isothermal aging, the initial condition always refers to the as-soldered state. 4.2. Characterization of the microstructural evolution during isothermal aging with optical microscopy Fig. 8 shows metallographic cross sections of the soldered interconnections in the initial state and after isothermal aging at 130 °C at 15 h, 85 h, 155 h. In the following subsections the microstructure and its observed changes are described. 4.2.1. Sn60Pb40 and Sn62Pb36Ag2 In the initial state the Sn60Pb40 solder matrix consists of evenly distributed Pb-rich and Sn-rich phases of 1–3 mm in diameter. The intermetallic layers at both interfaces have a thickness of 0.3 mm and 0.7 mm. Locally, Sn encircles cavities and lead-glass phases within the busbar, which is perceivable as gray areas around the dark glass phases in the busbar. After 15 h of isothermal aging the total number of grains inside the solder matrix starts to decrease, and larger phases with a diameter of 5–10 mm are formed. This process of grain coarsening proceeds upon further aging such that isolated grains of 10–30 mm and mostly interlinked phases are formed. We observe hairline cracks at 85 h and 155 h of aging between individual grains within the solder. The Cu6Sn5 IMC has a uniform, scallop-type appearance and tends to grow homogeneously. Cu3Sn is observable at 85 h and 155 h with a thickness of ≈0.3 μm and ≈1 μm respectively. Ag3Sn at the busbar interface grows more rapidly whereas its thickness is 3–4 mm after 85 h and 5–6 mm after 155 h. It consists of a uniform region close to the solder and local extensions almost reaching the wafer surface. Differently from the Sn60Pb40 solder alloy, the Sn62Pb36Ag2 has a finer grain structure in the initial state. Occasionally, 1–2 mm Ag3Sn-particles are located within the solder matrix but most often Ag3Sn-particles of 3–5 mm are sparsely distributed at the Cu6Sn5 interface. Basically, the observed changes during isothermal aging are similar to the ones seen in Sn60Pb40. The extent of intermetallic layer growth is equivalent to Sn60Pb40. The difference is, that a finer grain structure within the solder matrix is preserved even after prolonged thermal aging. 4.2.2. Sn43Bi57 and Sn41Bi57Ag2 The Sn43Bi57 solder has a lamellar structure of uniformly distributed Sn and Bi-phases in the initial state. Note, that the 3 We use the notation Ag3Sn(Zn) to indicate the existence of the Ag3Sncompound with traces of Zn (below 10 at%). The notation Ag-Sn-Zn is to describe an interdiffusion zone of the elements Ag, Sn and Zn with significant amounts of each ( >10 at% ) and no observable, fixed stoichiometry.

solder depot in the shown sample is significantly lower which influences the microstructural changes during aging. The Cu6Sn5phase is 0.5 mm thin after the soldering process. Additionally, a slightly thicker 0.8–1.1 mm Ag3Sn layer is located at the solderbusbar interface. It becomes obvious that the local, non-uniform penetration of Sn into the busbar is more pronounced in the Sn43Bi57 solder. Already after 15 h of isothermal aging it is observable that Sn diffuses significantly deep into the busbar, reaching the wafer surface locally. A rapid segregation of Bi in the solder matrix takes place, which is more pronounced within samples having a low solder depot. Moreover, the Cu6Sn5-phase grows faster than within the before mentioned solders. It reaches a thickness of 1.5– 3 mm after 15 h of thermal aging. Inspecting the solder joint after 85 h reveals a coarse solder matrix. It is frequently observed that it can be entirely desaturated from Sn if the solder layer is thin. As before, a double-layer of Cu3Sn and Cu6Sn5 is clearly distinguished at the interface to the Cu. At 155 h no further growth of Cu6Sn5 is observed. Rather Cu3Sn grows at the expense of the Cu6Sn5-layer. Moreover, no additional growth of Ag3Sn takes place within this sample. The retarded growth of Cu6Sn5 and Ag3Sn is due to the depletion of Sn in the solder depot. The Sn41Bi57Ag2 solder is initially made of a moderately-fine lamellar structure of the two main components similar to Sn43Bi57. In contrast to the former solder, Bi-rich phases of ≈1 μm diameter are dispersed within the matrix. Ag3Sn-particles of 3– 5 mm in diameter attached to the Cu6Sn5-layer are observed. The local ingress of Sn into the busbar after soldering is equivalently observed in Sn41Bi57Ag2. After 15 h of thermal aging, the grains have coarsened but the smaller Bi-rich particles have significantly increased in quantity and even decreased in size to below 1 mm. Although an accumulation of Bi takes place similarly to Sn43Bi57, still a large quantity of small Bi-rich crystals remains in the Sn-rich phase after extended thermal aging. The layer thickness of Cu3Sn and Cu6Sn5 in the Sn41Bi57Ag2joint after 15 h and also in subsequent aging steps is similar to the ones in Sn43Bi57-bonds. Notably, the Ag3Sn-particles, which are attached to the Cu6Sn5-layer, extend. We assume that Ag successively diffuses from within the solder matrix and particularly the cell metallization to the interface and participates at the growth process. The growth of the Ag3Sn layer within the busbar is equal to Sn43Bi57 after 15 h. The Ag3Sn thickness reaches 2–3 mm after 15 h. After that, we observe an almost entire consumption of the busbar by Ag3Sn (10–20 mm). The layer thickness when the Ag3Snphase reaches the busbar is not taken into account during the later calculations as it does not describe the potential phase growth under the condition of an infinitely large substrate. 4.2.3. Sn91Zn9 The initial Sn91Zn9 solder bond consists of fine Zn and Ag grains dispersed within the Sn main component. Stacks of intermetallic layers and interdiffusion zones with a total thickness of 1– 2 mm exist at the interface to Cu. A thick (2–3 mm) initial intermetallic layer exists at the busbar-to-solder interface. We speculate that the thicker intermetallic layers are a result of the longer soldering times. During isothermal aging at 130 °C for 85 h an agglomeration of Ag and Zn within the solder takes place to form phases of 1–3 mm diameter. A bi-layered structure of intermetallics at the Cu-interface becomes obvious. Along with this, a limited phase growth at the busbar is observed. Even after 155 h the total IMC thickness within the busbar is only 3–4 mm. After 155 h at 130 °C, the Agand Zn-rich phases in the solder matrix as well as the intermetallic compounds at the Cu-interface change to appear darker under the

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Fig. 9. a) SEM image of the Cu3Sn and Cu6Sn5 intermetallic layers of an thermally aged Sn41Bi57Ag2 solder bond. b) EDX line scan across the interface of Cu to solder (indicated by the dashed arrow). c) SEM image of a thermally aged Ag3Sn intermetallic compound at the solder-busbar interface. d) EDX line scan across this interface.

optical microscope indicating a change in elemental composition. Finally, we observe many locations where the solder has not wetted the busbar during sample preparation, which is an indication of the poor wettability of the solder. 4.3. Scanning electron microscopy and electron dispersive x-ray spectroscopy 4.3.1. Details of the Cu3Sn, Cu6Sn5 and Ag3Sn intermetallic layers A detailed sight on the Cu3Sn and Cu6Sn5 intermetallic layers as formed within an Sn41Bi57Ag2 solder joint thermally aged for 155 h at 100 °C is given in Fig. 9a along with an EDX line scan in

Fig. 9b across this interface. The SEM image illustrates clearly the scallop-type shape of the Cu6Sn5 IMC. The EDX scan shows the normalized atomic ratios of Cu and Sn. It is difficult to observe a distinct plateau where the atomic ratios would match the predicted values from theory as 3:1 and 6:5 respectively due to the limited resolution of the measurement. An Ag3Sn intermetallic layer at the busbar-to-solder interface (Sn43Bi57) after thermal aging at 115 °C for 155 h is shown in the SEM image in Fig. 9c. Again, it is supplemented with an EDX line scan across this interface in Fig. 9d. The Ag3Sn layer forms a relatively uniform plateau in the EDX line scan with an approximate atomic ratio of Ag to Sn of 3:1. Furthermore, Sn is detected in the

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Cu6Sn5

Cu Sn

1 m

Cu6Sn5

Pb

Ag3Sn

Ag3Sn busbar

2 m

(a)

(b)

Fig. 10. a) Tilted perspective on a thermally aged and intentionally unpolished Sn62Pb36Ag2 bond after 85 h at 100 °C, b) Detail of the same Sn62Pb36Ag2 solder bond at the interface to Cu showing an overly grown Ag3Sn particle.

bulk of the busbar close to a cavity. A content of 90 at% Ag and 10 at% Sn is measured which is either primary Ag with dissolved Sn or probably a stable ζ -(Ag) phase [54]. 4.3.2. Grain coarsening and grain boundaries The changes of the microstructure of an Sn62Pb36Ag2 solder joint after 85 h at 130 °C are highlighted using the tilted SEM images in Fig. 10a. Intentionally, the specimen is not polished before taking the pictures in order to reveal the significant structural and volumetric changes, which the solder joint is subjected to during thermal aging. It can be seen that height differences are formed between solder layer and intermetallics as well as the busbar. Fig. 10b shows a 3–5 mm Ag3Sn-particle at the solder-Cu interface bulging out of the sample. Fig. 11a unveils the coarse structure of Sn-phases and Bi-phases within a thermally aged Sn41Bi57Ag2 solder joint. Occasionally, Bi-rich areas of below 1 mm in diameter are seen within the rocky agglomerated Sn and Bi-phases. A high magnification image between both phases in Fig. 11b exposes assumedly 10–20 nm wide grain boundaries within the solder. 4.3.3. Composition of the intermetallics in Sn91Zn9 Fig. 12a shows the interface from Cu-ribbon to solder within a Sn91Zn9 solder joint. The dashed arrow indicates the corresponding quantitative EDX line scan in Fig. 12e. From the EDX data it can be seen that Zn agglomerates within 1.5–2 mm thick

intermetallic layers or interdiffusion zones. The atomic ratio of Cu to Zn is approximately 5:8, indicating the existence of a Cu5Zn8phase. The SEM image shows a thin 0.2–0.3 mm intermediate phase between Cu5Zn8 and Cu, which, according to the quantitative EDX analysis, has the atomic ratio of Cu and Zn of 1:1. This suggests the presence of β′-CuZn. A significant feature at the Cu-solder interface is a chain-like arrangement of Ag-rich particles. During a number of EDX point scans at different locations, we also find significant amounts of Sn (15–20 at%) and Zn (10–20 at%) within these intermetallic compounds, whereby the atomic ratio of Ag and Sn often corresponds to 3:1, which indicates Ag3Sn. Zn may be dissolved into the highly Ag-rich regions since its solubility in solid Ag can be up to 40 at%. Fig. 12b shows a detail of the Sn91Zn9 solder matrix, in which a large amount of Ag3Sn-particles of 0.2–0.5 mm diameter with Znadditions are dispersed. Since the initial solder composition does not contain any Ag, the existence of these particles is the result of an Ag-diffusion process. It is well-known that Ag dissolves quickly into molten solder, hence we assume that the presence of Ag within the solder matrix and also at the interface to the Cu-ribbon is the result of Ag-dissolution into molten solder during the soldering process. Fig. 12b also shows primary Zn-areas, which appear darker as compared to the Sn matrix. After thermal aging the interface from solder to copper is structurally rearranged. In Fig. 12c it is observable that Sn diffuses non-uniformly through the Cu5Zn8 layer. Cu5Zn8 and β′-CuZn

Fig. 11. a) Sn41Bi57Ag2 solder joint after 155 h at 100 °C, b) High magnification detail of a grain boundary between Sn-rich and Bi-rich phases.

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Fig. 12. a) Detail of the copper-solder interface within a Sn91Zn9 solder joint. The dashed arrow indicates the line of the quantitative EDX measurement. b) Detail of the Sn91Zn9 solder matrix. c) The thermally aged interface from copper to solder with Sn91Zn9. d) The thermally aged interface from solder to busbar with Sn91Zn9. e) EDX line scan across the interface from copper to solder as given by the dashed line in (a).

appear to be partly disintegrated and “flooded” with Sn-rich regions instead of forming solid and uniform layers. Moreover, the agglomeration of ZnO in cavities at the solderbusbar interface can be seen in Fig. 12d as a result of the thermal aging process. Also the interface from solder to ribbon is afflicted with these ZnO- regions of diameters ranging from 1 mm to 20 mm. EDX point and area scans within the busbar region close to the interface to the solder reveal the presence of a continuous Ag3Sn intermetallic phase with minor additions of Zn between 5 at% and 10 at%. 4.4. Modeling the phase growth 4.4.1. Parameter determination The IMC-layer thickness is measured after each time step and at the defined aging temperatures using the laser confocal microscope. The measured IMC thickness using 30 measurements per cross section along with a linear fit of the averaged measurements to the square root of time is shown Fig. 13. The adjust- R2 of the fit is almost always in the range of 0.91–0.99, deviating only if layer thickness is still comparably small at short aging times. The

linearity of the fit implies a diffusion-controlled process. We decide to sum the Cu3Sn and Cu6Sn5-phase growth for Bicontaining solders because here the Cu3Sn phase grows notably in expense of Cu6Sn5. Moreover, we sometimes observe large variations in the initial layer thickness when using the Sn91Zn9 solder alloy. This may be explained by the poor wettability of the solder on the copper core of the ribbon during initial tinning which leads to larger inhomogeneities of the initial IMC thickness. The standard deviation of the Cu6Sn5 IMC thickness is between 0.1–0.2 mm. It shows notably higher values for Ag3Sn of 0.2–0.5 mm. Particularly, the thickness of Ag3Sn using Sn60Pb40 is subject to large variations of 2–4 mm after extended aging. Using the experimental data and Eq. (5) we generate Arrhenius plots, which enable the extraction of the activation energy of the layer growth using the slope of the linear fit of ln(x − x0) and 1/T according to Eq. (6). Fig. 14 shows one of the Arrhenius plots for the case of Ag3Sn-growth using Sn60Pb40-solder, and the linear fit as a dashed line. Sometimes outliers, mainly resulting from measurement uncertainties must be removed, such as the data point indicated with a cross in the Fig. 14, that would else lead to an erroneous averaging of the slopes. The measured intermetallic layer thickness of the Cu5Zn8 with

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Fig. 13. Measured IMC thickness and linear fit to the square root of time after isothermal aging. Sn62Pb36Ag2 is omitted due to similarity with Sn60Pb40.

our chosen aging temperatures does not result in an appropriate linear fit in the Arrhenius chart which is why we discard the data in the following analysis. The resulting kinetic parameters for all solders and the investigated IMC are cataloged in Table 3. These values are useful for making forecasts of the intermetallic layer thickness under specified temperature conditions. The value SEQ is the average of the standard errors, that are determined during the linear fit of each individual time series. Respectively, the value SE D0 means the

average of the standard errors which are calculated during the non-linear fit to Eq. (4). The MAE as defined in Eq. (8) of the parameterized model is almost always below 1 mm indicating its good predictability. 4.5. Application of the model The model is applied to predict the phase growth that we assume to occur during photovoltaic module testing at a constant

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temperature and slows down in winter. When the temperature rises in the following year the IMC-growth sets in again. The simulated intermetallic layer thickness of selected solders and IMC types after 25 years in Freiburg, Germany is shown in Fig. 16b. Due to the sequential growth and stagnation during one year a typical step-wise increase over the course of 25 years is calculated. Globally, the curve follows the parabolic shape as consequential from the theoretical considerations. We calculate the Ag3Sn intermetallic phase using standard Sn60Pb40 solder to grow to a maximum thickness of 1.2 mm. Moreover, a resulting thickness of the Ag3Sn-phase with Sn43Bi57 solder alloy of 2.4 mm is predicted. Using Sn91Zn9 the Ag3Sn reaches a thickness of 2.8 mm after 25 years. The results prognosticate that the intermetallic phase growth after 25 years in the location Freiburg is negligible. The reason for the limited phase growth is the low module temperature that rarely exceeds 60 °C. IMCs will grow more in arid regions with higher module temperature.

5. Discussion Fig. 14. Arrhenius plot of the Ag3Sn IMC-growth using Sn60Pb40-solder. The error bars indicate the SE of the data point, which originates from the SE of repeated measurements of the layer thickness. The data points have been slightly shifted in x-direction to avoid an overlapping of the error bars. Outliers are omitted from the linear fit such is the case for the data point shown as a cross in the figure.

temperature of 85 °C, as is present in damp heat, and during thermal cycling tests between
40 °C and 85 °C [55]. The simulated intermetallic layer growth during 3000 h at a constant temperature of 85 °C is shown in Fig. 15a. The unlikely influence of humidity on the phase growth is not included in the simulation. It is calculated that the intermetallic layers reach a considerable thickness already after 1000 h and 3000 h respectively. Especially the Ag3Sn-phase when using low melting point solders is prone to excessive growth up to values between 3.7 mm and 6.6 mm after 3000 h at 85 °C. The Ag3Sn in Sn60Pb40 solder joints is limited to 1.8 mm and 2.6 mm after 1000 h and 3000 h respectively. The results for 600 cycles from
40 °C to 85 °C with a cycle time of ≈2.8 h is shown in Fig. 15b. After the thermal cycling test the Ag3Sn thickness of the lead-based solders and Sn41Bi57Ag2 is predicted to increase to 1.2–1.5 mm, which is almost double the initial thickness. Particularly, Sn43Bi57 shows an extended Ag3Sn growth to a total thickness of 2.2 mm. The growth of the phases at the copper ribbon for all solders and the growth of Ag3Sn for Sn91Zn9 during thermal cycling tests can be regarded as negligible. Fig. 16a shows a progression of the Ag3Sn layer growth for Sn60Pb40 during the first year in the outdoor location Freiburg. During the first year the Ag3Sn layer grows 0.1 mm. The IMCgrowth proceeds fast during summer due to the elevated module

5.1. Discussion of the observed microstructural changes In Sections 4.2 and 4.3.2, substantial structural changes within the solder matrix and at the interfaces are described. Along with the growth of intermetallic layers at the interfaces, segregation of the constitutive elements in the solder in particular prevalent in the Bi-containing solders is observed. The segregation process is generally driven by the alloy to reduce the overall energy of the system and reach equilibrium solute concentrations [18]. From the respective phase diagrams in references [16,56] it becomes obvious that neither Sn in Bi nor Bi in Sn, as well as Sn in Pb nor Pb in Sn is soluble in significant amounts. Hence, the alloy tends to separate into larger grains. As we have identified, the phases are demarcated by pronounced grain boundaries. The effects on the mechanical properties of the interconnection may be tremendous since the overall capability of the solder to allow deformation and to resist fatigue is influenced by the granular structure within the solder. The common opinion is that a uniform and fine granular structure evenly distributes the mechanical load which leads to enhanced fatigue life [23]. It is, however, well known that Bi is brittle and cannot withstand large deformations. With regard to the large extent of segregation in Bi-solders and the inherent brittleness it is very likely that Bi-containing solders are less resistant to mechanical stresses and fatigue which could limit their usefulness in photovoltaic applications. The addition of Ag is claimed to increase the elongation to failure of SnBi-solder by a refinement of the microstructure [36,38]. In fact, we also observe a finer microstructure after

Table 3 Kinetic parameters of the intermetallic phase growth. Solder

IMC

Q[kJ mol
1]

SEQ[kJ mol
1]

D0[m2 s
1]

SE D0⎡⎣ m2s−1⎤⎦

x0[ μm]

MAE [ μm]

Sn60Pb40

Ag3Sn Cu6Sn5 Cu3Sn Ag3Sn Cu6Sn5 Cu3Sn Ag3Sn Cu6Sn5 þ Cu3Sn Ag3Sn Cu6Sn5 þ Cu3Sn Ag3Sn(Zn)

129.7 85.7 99.4 128.1 81.2 111.3 124.8 100.5 153.2 147.7 87.9

3.9 4.6 6.2 8.2 6.0 6.2 26.5 30.5 7.3 31.5 16.9

1.44 4.25  10
8 2.20  10
6 0.77 10.00  10
8 7.58  10
5 2.43 4.81  10
6 8.60  103 88.5 6.76  10
7

0.04 0.15  10
8 0.16  10
6 0.03 0.11  10
8 0.57  10
5 0.60 0.48  10
6 1.07  103 15.6 0.65  10
7

0.70 0.32 0.00 0.70 0.31 0.00 0.88 0.50 0.83 0.52 2.00

0.14 0.19 0.46 0.34 0.29 0.55 0.56 0.83 1.78 0.64 0.44

Sn62Pb36Ag2

Sn43Bi57 Sn41Bi57Ag2 Sn91Zn9

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Fig. 15. a) Simulated phase growth for 3000 h at 85 °C as would occur in damp heat tests according to IEC 61215:2005-04 [55], b) Simulated phase growth for 600 thermal cycles (cycle time E2.8 h) A: Sn60Pb40-Ag3Sn, B: Sn60Pb40-Cu6Sn5, C: Sn60Pb40-Cu3Sn, D: Sn62Pb36Ag2-Ag3Sn, E: Sn62Pb36Ag2-Cu6Sn5, F: Sn62Pb36Ag2-Cu3Sn, G: Sn43Bi57-Ag3Sn, H: Sn43Bi57-Cu6Sn5, I: Sn41Bi57Ag2-Ag3Sn, J: Sn41Bi57Ag2-Cu6Sn5, K: Sn91Zn9-Ag3Sn(Zn).

Fig. 16. a) Simulated growth of the Ag3Sn-phase of the Sn60Pb40-bond in the outdoor location Freiburg, Germany during the first year. b) Simulated IMC-growth for selected solders in Freiburg over the course of 25 years.

soldering with Sn41Bi57Ag2 which is partially retained after thermal aging. We are optimistic that the modification by Ag is able to significantly improve long-term stability of Bi-containing solders. Grain coarsening in solder joints can be a dominant failure mode and is a major reliability concern in electronics [57,31,58]. It can evidently lead to fatigue cracking in PV modules after prolonged outdoor exposure [4,59–61]. The impact of IMC-layer growth in addition (or in comparison) to grain coarsening of the solder matrix needs to be further investigated using failure analysis of field-aged PV modules and correlation of IMC-growth to mechanical properties of the joint. Furthermore, the phenomenon of Sn-penetration into the busbar metallization, which is observed regardless of using

standard leaded solders or low melting point solders, is considered a risk for PV module reliability. Significant amounts of Sn are particularly localized around lead-glass particles and voids in the busbar. We assume that the diffusivity of Sn from the solder matrix along defects, cavities and grain boundaries along lead-glass particles is higher than in solid silver. Thus, Sn agglomeration is promoted at those features in the busbar. A correlation between local Sn-penetration and deterioration of the adhesion of the screen-printed metallization to the substrate is reported by [10,62]. Diffusion swelling of the metallization and subsequent failure of the mechanical bonding of the glass frit to the wafer is assumed to be one cause of the adhesion loss [63,64]. Another conjecture is that Sn atoms induce a reduction reaction of the binder component in the metallization, which destroys the

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oxide bonding of the metallization to the wafer [65]. Although it is evident that mechanical strength of the interconnection is deteriorated due to the phenomenon of Sn-penetration a quantitative analysis about the impact on PV module reliability is still missing. A correlation of extent of Sn-penetration, adhesion loss and PV module performance in accelerated aging tests or under outdoor conditions is needed. 5.2. Discussion of the experimental approach and simulation Since grain coarsening and IMC-growth are mainly based on diffusion processes, which in turn strongly depend on temperature, it is justified to approach these phenomena with solely isothermal aging experiments. Other environmental factors such as humidity are negligible and do not influence the quantitative results of our investigation. Particularly, the understanding of the growth kinetics of the IMCs and the determination of practically useful model parameters requires systematic aging experiments at different isothermal temperatures. We chose to age the lead-based solder joints in a temperature range from 100 °C to 150 °C in order to achieve an extent of phase growth on a reasonable time scale, that can be easily measured by our instrumentation. The low melting point solders must be aged in a lower temperature range (85– 130 °C) to avoid remelting of the solder at 150 °C. The chosen temperature range is not linked to the actual module temperature under field conditions, which is usually limited to ≈80 °C . The observations on the microstructural changes in the solar cell interconnections under isothermal aging conditions must, however, be clearly differentiated from those during thermal cycling tests, damp heat tests or outdoor exposure of the PV module. Although highly relevant, a consideration of these test conditions is beyond the scope of this paper, but we want to provide some references to commonly observed failure modes: Microstructural changes during thermal cycling and outdoor exposure is dominated by crack growth within the solder joint [5,60,61]. The reason is the thermo-mechanical stress induced by the CTE-mismatch of Si-wafer and Cu-core and the limited capability of the solder layer to compensate the stress. During damp heat conditions water vapor ingresses into the PV module, which leads to the release of acetic acid from the ethylen-vinyl acetate encapsulation polymer. The acetic acid determines a corrosion and ablation of the solar cell metallization [66,67]. Likewise the experimental approach, our simulation is based on a solely thermally-driven diffusion process, whose kinetic is described by the Arrhenius relationship with a single activation energy. We find good agreement between simulation and experimental data up to 750 h at elevated temperatures above 85 °C. The MAE has a maximum of 1.8 mm. On that basis we imply validity of the proposed model for other aging temperatures and other time scales. However, the simulation lacks a cross check with experimentally determined IMC-growth after extended damp heat or thermal cycling treatment as well as after prolonged outdoor exposure. Practical problems arise regarding proper cross section preparation if encapsulated interconnections are to be investigated. Furthermore, the processes and materials must be fully known and tightly controlled in order to make a meaningful comparison between simulation and experiment. Our simulation results must be understood as a thorough prognosis under the given boundary conditions. 5.3. Interpretation of the EDX measurements of Sn91Zn9 Generally, we observe that the region of solder to copper ribbon interface is afflicted with a very complex structure of intermetallic compounds and interdiffusion zones. Recalling Fig. 12a

385

Ag-rich intermetallic compounds in the form of particles attached to the γ -Cu5Zn8 interface are found. As opposed to Song et. al. we identify a significant amount of Sn within this type of intermetallic compound. Song et al. concluded from XRD-studies that these particles are AgZn3 and Sn would not be involved in the intermetallic formation [48]. However, judging from our EDX-measurements we think that the particles are primarily Ag3Sn with approximately 20 at% Zn dissolved in the silver. From the phase diagram of Ag-Zn, it is possible that up to 40 at% Zn can be dissolved in solid Ag [45]. The reason for the discrepancy between Song et. al. and our study is possibly rooted in the nature of the samples. Whereas the former uses bare copper and silver wires respectively which are dipped into the molten solder, we use pretinned copper ribbons prepared by the vendor at, for us, unknown conditions. Thermal aging leads to a vast reorganization of the intermetallic compound structure at the copper-solder interface. It is reported that Sn penetrates into the γ -Cu5Zn8 resulting into the formation of Cu6Sn5 [47]. In fact, we equivalently see the Sn-penetration and Cu6Sn5 formation in our microscopic investigations. Additionally, Cu6Sn5 can grow during thermal treatment due to an instability of γ -Cu5Zn8 when heated to 350 °C. These temperatures may have been present during the initial tinning of the ribbon at the vendor [68]. The existence of a β′-CuZn between γ -Cu5Zn8 and Cu along with an unidentified third intermetallic layer has been reported elsewhere [47]. Under the optical microscope this layer is characterized by a bronze visual appearance as compared to the gray dark γ -Cu5Zn8 IMC. There is disagreement concerning the elemental composition of the primary intermetallic phase at the busbar to solder interface. We mostly find the atomic ratio of Ag and Sn of 3:1 with additions of Zn of a few at% which quite clearly is not AgZn3 as reported by Chen et al. [15]. Zn rather tends to migrate into cavities at the interfaces and forms ZnO instead of allocating in the busbar and participating in the phase growth there. 5.4. Discussion of the kinetic parameters The temperature dependence of the intermetallic phase growth is described by the Arrhenius relationship. Since the activation energy Q is included in an exponential relationship it decisively shapes the extent of temperature dependence of the diffusion coefficient D over temperature. The pre-exponential factor D0 can be considered as a scaling factor merely shifting the exponential term in parallel to lower or higher diffusion coefficients. The diffusion coefficients D of the phase growth are plotted over temperature in Fig. 17a. The series for Sn62Pb36Ag2 is omitted since it practically equals the one of Sn60Pb40. The influence of a high activation energy on the phase growth can be seen in particular for the phases with Sn41Bi57Ag2 (lines 3 and 7). High activation energies lead to a steeper increase of the diffusion coefficient with temperature. Due to this fact, the diffusion coefficient of Sn41Bi57Ag2 may even fall below the one of Sn60Pb40 (line 1) at low temperatures although the former is excessively afflicted with phase growth at high temperatures. In contrast, a low activation energy leads to a lower increase of D with temperature. This may limit the phase growth at high temperatures, as for Sn91Zn9 (line 5), but may also become significant at low temperatures. We ask the question if a correlation between material properties of the solders and kinetics of the phase growth exist. First, the activation energy is plotted for each solder in Fig. 17b whereby the solders are ordered from left to right according to increasing amount of Sn. As can be seen, a relationship of increasing Sncontent to lower activation energy is present. Considering that Sn

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Fig. 17. a) Diffusion coefficients of the growth kinetics of intermetallic layers with different solders. (1) Sn60Pb40-Ag3Sn, (2) Sn43Bi57-Ag3Sn, (3) Sn41Bi57Ag2-Ag3Sn, (4) Sn91Zn9-Ag3Sn(Zn), (5) Sn60Pb40-Cu6Sn5, (6) Sn43Bi57-Cu6Sn5 þCu3Sn, (7) Sn41Bi57Ag2-Cu6Sn5 þ Cu3Sn, b) Activation energy sorted from left to right by increasing Sncontent.

is mainly involved in the phenomenon of phase growth, it seems plausible that a high Sn-content in the solder matrix lowers the energy barrier4 for the phase growth to take place because more reactive material is available and able to diffuse across the interfaces. The function of the Sn-content on the kinetics of phase growth has been investigated before by Vianco et. al. [28]. There it is concluded, that layer growth between different solders is determined by the Sn-content at equivalent homologous temperatures. Besides the Sn-content, the homologous temperature is another factor that determines the growth rate of the intermetallic layers. Recalling our phase growth data, a distinct trend of low melting point solders (Sn43Bi57 and Sn41Bi57Ag2) exists to show excessive phase growth at temperatures above 85 °C. Regarding the relatively low Sn-content, one would expect a generally reduced growth rate of these solders. However, since the homologous temperature is very high when Sn43Bi57 and Sn41Bi57Ag2 are aged at temperatures above 85 °C, this becomes the governing factor of the layer growth.

6. Conclusion The microstructure of five different solders is studied. Generally, grain coarsening during thermal aging is observed in a way that the initially fine structure of solidified crystals of sub-micrometer size gradually transforms into larger phases with a diameter of 1–10 mm and, in the later stages, 10–30 mm. Grain coarsening is particularly pronounced in Sn43Bi57 and Sn41Bi57Ag2 alloys which can lead to a complete depletion of Sn in the solder matrix and concomitant agglomeration of Bi between ribbon and busbar. The addition of Ag to SnPb and SnBi refines the grain structure 4 According to Menzinger and Wolfgang: “They (van't Hoff and Arrhenius) interpreted the activation energy as the height of the energy barrier which has to be overcome by the relative translational motion of the reactants in order for reaction to occur” [69].

initially, and these solders are less prone to grain coarsening during thermal aging. Details of the microstructural changes are inspected using SEM. We observe structural changes in the solder matrix, such as volume increase of intermetallic particles, grain boundary growth between the phases of 10–20 nm and microcracks which are all likely to have a negative impact on the mechanical stability of the solder joint. We observe the rapid growth of Ag3Sn within the busbar of 5– 10 mm depending on temperature and time of aging as well as solder type. The worst cases are Sn43Bi57 and Sn41Bi57Ag2 where Ag3Sn consumes the entire busbar after 85 h at 130 °C. The intermetallic phases at the Cu-interface grow to a total thickness of 1.5– 3 mm. Generally, the homologous temperature of the solder dominates the phase growth in such a way that low melting point solders show significantly faster phase growth at elevated temperatures. It is observed that Sn penetrates the busbar locally even in the as-soldered condition. Thereby, Sn occasionally reaches down to the wafer surface after short aging times. Sn tends to agglomerate at cavities and around lead glass particles within the busbar. The Sn-penetration is a possible risk for PV module reliability since it is linked to adhesion loss of the busbar metallization. Within the Sn91Zn9 solder joint we identify the existence of a main 2 mm thick Cu5Zn8 phase between copper and solder which is accompanied with a 0.2–0.3 mm β′-(CuZn) phase after 85 h at 130 °C or higher conditions. Attached to the Cu5Zn8 phase are particles of an Ag-Sn-Zn composition that could be interpreted as Ag3Sn with dissolved Zn. The presence of Ag in the solder matrix is the result of a dissolution of the busbar during the soldering process. In contrast to other authors, the phase within the busbar is not identified as AgZn3 but as Ag3Sn with minor additions of Zn. Instead, we observe the agglomeration of Zn and the formation of ZnO within cavities that can reach several micrometers in diameter at the interfaces to the substrates. The kinetic parameters of the phase growth are determined based on the root mean square diffusion distance of atoms and

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using the Arrhenius relationship. With the so obtained kinetic parameters absolute errors of model to experimental data of below 1 mm is achieved in most cases indicating the good predictability of the model. It is shown that the activation energy correlates with the Sn-content of the solders. A higher Sn-content lowers the activation energy. This is plausible because the existence of a large amount of reactive material in the solder matrix lowers the energy barrier for the phase growth reaction to occur. The kinetic parameters are applied to simulate the phase growth at a constant temperature of 85 °C and under thermal cycling conditions from
40 °C to 85 °C. Ag3Sn grows to a thickness of 1.8 mm with SnPb(Ag) after 1000 h at 85 °C and 2.6 mm after 3000 h. The thickness using Sn43Bi57 after 1000 h and 3000 h are 4.2 mm and 6.6 mm respectively. Sn41Bi57Ag2 yields a thickness of 2.5 mm and 3.7 mm. The Ag3Sn-phase grows to a thickness of 1.2–1.5 mm during 600 thermal cycles if Sn60Pb40, Sn62Pb36Ag2, or Sn41Bi57Ag2 is used, but extends to 2.2 mm for the case of Sn43Bi57. The growth of the Ag3Sn layer using Sn91Zn9 as well as the phases at the Cu-ribbon during thermal cycling are negligible. The application of the model for the calculation of the intermetallic layer thickness after 25 years outdoor exposure in Freiburg, Germany yields the following results. In the worst case, the Ag3Sn-phase with Sn43Bi57, which starts at 0.8 mm is projected to have a thickness of 2.4 mm after 25 years compared to 1.2 mm when the standard Sn60Pb40 is used.

Acknowledgments

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The authors would like to thank Bruker-Spaleck GmbH, Germany and Ulbrich of Austria GmbH for their kind support of this work.

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