Document [original]

RESEARCH ARTICLE

Decay mechanisms in CdS-buffered Cu(In,Ga)Se

thin-film

solar cells after exposure to thermal stress: Understanding the

role of Na

Hasan A. Yetkin

1,2

| Tim Kodalle

| Tobias Bertram

Alejandra Villanueva-Tovar

| Marin Rusu

| Reiner Klenk

| Bernd Szyszka

Rutger Schlatmann

1,4

| Christian A. Kaufmann

Competence Centre Photovoltaics Berlin

(PVcomB), Helmholtz-Zentrum Berlin für

Materialien und Energie, Berlin, Germany

Technology for Thin-Film Devices,

Technische Universität Berlin, Berlin, Germany

Department Structure and Dynamics of

Energy Materials, Helmholtz-Zentrum Berlin

für Materialien und Energie, Berlin, Germany

School of Engineering - Energy and

Information, Hochschule für Technik und

Wirtschaft, Berlin, Germany

Correspondence

Hasan A. Yetkin, PVcomB, Helmholtz-Zentrum

Berlin für Materialien und Energie,

Schwarzschildstr. 3, 12489 Berlin, Germany.

Email: hasan.yetkin@helmholtz-berlin.de

Funding information

Ministry of National Education of the Republic

of Turkey

Abstract

Due to their tunable bandgap energy, Cu(In,Ga)Se

(CIGSe) thin-film solar cells are an

attractive option for use as bottom devices in tandem configurations. In monolithic

tandem devices, the thermal stability of the bottom device is paramount for reliable

application. Ideally, it will permit the processing of a top device at the required opti-

mum process temperature. Here, we investigate the degradation behavior of chemi-

cal bath deposited (CBD) CdS-buffered CIGSe thin-film solar cells with and without

Na incorporation under thermal stress in ambient air and vacuum with the aim to gain

a more detailed understanding of their degradation mechanisms. For the devices

studied, we observe severe degradation after annealing at 300C independent of the

atmosphere. The electrical and compositional properties of the samples before and

after a defined application of thermal stress are studied. In good agreement with liter-

ature reports, we find pronounced Cd diffusion into the CIGS absorber layer. In addi-

tion, for Na-containing samples, the observed degradation can be mainly explained

by the formation of Na-induced acceptor states in the TCO front contact and a back

contact barrier formation due to the out-diffusion of Na. Supported by numerical

device simulation using SCAPS-1D, various possible degradation models are dis-

cussed and correlated with our findings.

KEYWORDS

CdS buffer layer, CIGSe, degradation mechanism, elemental interdiffusion, Na, SCAPS,

simulation, thermal stress

1|INTRODUCTION

Recently, Cu(In,Ga)Se

(CIGSe) thin-film solar cells have shown an

efficiency boost up to 23.35% by alkali post-deposition treatments

(PDTs).

According to the theoretical limit of about 33% at a bandgap

energy of 1.15 eV

for single junction devices, there is still open room

to further increase their efficiency. One approach to achieve this could

be to decrease the absorption losses in the transparent conductive

[Correction added on 11 June 2021 after first online publication: Tables 1 and 2 have been corrected in this version.]

Received: 16 March 2021 Accepted: 30 April 2021

DOI: 10.1002/pip.3438

This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium,

provided the original work is properly cited.

1034 Prog Photovolt Res Appl. 2021;29:1034–1053.wileyonlinelibrary.com/journal/pip

oxide (TCO) used as front contact material by increasing its transpar-

ency. For this, it would be of great advantage to be able to work at

elevated deposition temperatures and/or to apply a post-deposition

annealing.

3–5

Another possibility to increase device efficiencies even

past the single junction Shockley–Queisser limit is to fabricate multi-

junction solar cells by utilizing different wavelength ranges of the inci-

dent light spectrum in a stack of two or more solar cells with different

bandgap energies. Recently, an efficiency of 24.1% for a perovskite/

CIGSe tandem solar cell has been reported.

Here, a CdS-buffered CIGSe

device was used as a bottom device with the deposition temperature of

the perovskite device not exceeding 100C. However, to broaden the

choice of absorber materials that can potentially be used as top devices

in such tandem stacks, the thermal stability of the bottom device is

paramount for reliable application. Ideally, it will permit processing of

the top device at the required optimum process temperature. For a

chalcopyrite-based top device, this would have to exceed 400C.

The impact of thermal exposure on CIGSe solar cells under differ-

ent environmental conditions was formerly investigated by several

groups. Ramanathan et al. have demonstrated that heat treatments in

air (“aging”) at 250C of CIGSe/CdS/i-ZnO junctions led to the forma-

tion of a wide bandgap secondary phase as seen in the corresponding

external quantum efficiency (EQE).

Kijima and Nakada have con-

ducted vacuum annealing for 30 min on CdS and ZnS(O,OH)-buffered

CIGSe solar cells. They found that excess Cd and Zn diffuse into the

CIGSe absorber layer leading to degradation of the device perfor-

mance.

Similar observations were confirmed by other groups.

10–12

They have however not reported the effect of Na on the cell degrada-

tion. Besides this, Kazmerski et al. have observed Cu

S formation at

the CIGSe/CdS interface using X-ray photoelectron spectroscopy

(XPS) analysis leading to 50–75% loss in photovoltaic cell perfor-

mance.

Looking at potential-induced degradation (PID) of CIGSe

solar cells and modules, it has been shown that also alkali elements

play a crucial role in the deterioration of the investigated device per-

formance.

13–16

In our study here, the CIGSe solar cell devices with

and without Na incorporation are used to investigate decay mecha-

nisms of those devices after thermal stress. Accordingly, a comprehen-

sive model is presented, which describes the degradation mechanisms

in place, and it will be seen that it is indeed Na that plays a key role in

the observed, thermally induced degradation. For the sake of possible

guidance, Figure 11 might be seen while reading Sections 3 and 5.

2|EXPERIMENTAL PROCEDURES AND

NUMERICAL SIMULATION

For device fabrication, an 800-nm-thick molybdenum layer is depos-

ited by DC sputtering on top of 50 50 2-mm

-sized soda lime

glass substrates as a back contact. Na-free samples contain a 150-nm-

thick SiO

diffusion barrier that is deposited before the Mo back

contact. The 2.1-μm-thick CIGSe absorber layer is co-evaporated in

a three-stage-like process at a maximum nominal substrate tempera-

ture of 530C, showing a compositional in-depth Ga gradient and a

final molar fraction ratio Cu/(Ga +In) (CGI) of 0.90. Details of the

adapted three-stage process can be found in Heinemann et al.

After

washing the CIGSe absorber layer in 10% NH

(aq), a CdS buffer layer

is applied by chemical bath deposition (CBD). On top of the CdS, a

bilayer consisting of intrinsic ZnO (i-ZnO) and a doped ZnO:Al (AZO)

with a total thickness of approximately 190 nm is sputter deposited at

room temperature (RT) and at 150C, respectively. Finally, using a

shadow mask, Ni/Al/Ni contact grids are deposited by e-beam evapo-

ration in order to facilitate current collection. It should be noted that

one dedicated reference sample (“as-deposited”) will be shown for

each set of samples together with the respective results in order to be

able to compare the obtained results, since the sets were deposited at

different points in time. Particularly for the sets of samples used in

Section 3.2, we term the respective reference devices as, for example,

“as-deposited with Na for air annealing”or “as-deposited with Na for

vacuum annealing.”

Concerning the annealing procedures, the systems (a simple hot

plate for air annealing and a heater in a vacuum chamber for vacuum

annealing) are preheated to the desired annealing temperature. After

the temperature has stabilized, the samples are exposed to the

desired thermal stress. The temperature of the hot plate was mea-

sured with a temperature sensor showing a variation of ΔT¼ 2C.

With regard to the vacuum annealing, a 5 5-cm

sample holder is

used to handle the samples inside the vacuum system. Temperature

calibration in the vacuum system is achieved using temperature

stickers. Nevertheless, a comparatively slow heat-up rate that is cau-

sed by the sample holder increasing the thermal mass and a possibly

laterally inhomogeneous heat distribution must be taken into account.

Therefore, the accuracy of the annealing temperature in vacuum is

estimated to be in the range of ΔT¼ 5C, even though nominally

identical annealing procedures have been used.

The annealing procedures are applied either to complete devices

or separately after individual layer deposition steps, as depicted in

Figure 1. Hence, annealing in air or under vacuum is carried out after

CIGSe (just before CdS deposition), CdS, i-ZnO, and AZO deposition.

After annealing, unfinished devices are completed with the remaining

layers. In the course of this study, the front contact layers might be

seen as a main reason for the observed degradations. We therefore

employed etching procedures to selectively remove the front contact

layers and to rebuild them freshly for the purpose to reveal the effect

of front contact layers on the observed degradation. To that end,

some of the completed samples that had been annealed after CdS and

AZO during preparation are separately etched (1) for 5 min in 10% cit-

ric acid in order to remove the ZnO bilayer or (2) for 2 min in 10% HCl

or in 10% HCl and 10% KCN (complementary experiment) in order to

remove the CdS and ZnO bilayer. Afterwards, respectively, either the

ZnO bilayer or both the CdS and the ZnO bilayer are freshly

redeposited. The KCN etch is performed in an experiment, which

investigated the possible presence of Cu

(S,Se) phases at the

absorber/buffer interface as indicated by the work of Bér et al.

Their

presence can, however, be excluded here (see Figure S5). KCN etching

did not further affect the properties of the device completed with the

HCl-etched CIGS surface. Therefore, the KCN etch is not mentioned

further below; it was however performed on all the samples from

YETKIN ET AL.1035

vacuum annealing experiments. An NaF PDT consisting of the evapora-

tion of nominally 3-nm NaF at 250C was implemented for Na-

containing samples after HCl etching prior to the redeposition of

CdS/i-ZnO/AZO. For one sample without Na, air annealing at 300C

for 20 min was performed on the HCl-etched sample as well. Note that

all the samples were stored in N

-filled desiccator before and after the

deposition, annealing, and analysis. After the completion of the

annealing treatments, the analysis of the samples was performed

mostly in the first 6–8 weeks. However, it should be noted that the

temperature-dependent current density–voltage measurement in some

cases has been carried out after up to 9 months. In this case, how-

ever, it was ensured that the current density–voltage measurements of

these samples did not show any instability after the long storage time.

Current density–voltage (JV) measurements were conducted

under standard test conditions (AM1.5, 1000 W m

2

,25

C) using a

WACOM A+solar simulator. In this study, each sample contains at

least seven, mostly 15 solar cells with a nominal area of 0.97 cm

Therefore, we use box plots to visualize the JV measurements, where

25% of the data above and below the median are located within the

box. Temperature- and illumination-dependent current–voltage (JV–T)

measurements were conducted under vacuum in order to avoid con-

densation of water on the sample in a liquid N

-cooled cryostat

(CryoVac) using a Keithley 2601A source measurement unit in 4-point

contact configuration and a LED solar simulator (Oriel VeraSol) simu-

lating an AM1.5 solar spectrum with a light intensity of 1000 W m

2

The temperature ranges from 320 up to 90 K with a step size of 10 K.

Measurements of the EQE with light bias at 0 voltage bias and the

capacitance–voltage (CV) in dark were performed using in-house built

setups. Elemental depth profiles are determined by glow discharge

optical emission spectrometry (GD-OES) using a Spectruma GDA

650 tool.

For better comparison, the Na and Cd depth profiles are

aligned at the respective Cd onset. Please note that the Na depth

profiles are smoothed in order to ease viewability. Sheet resistance

measurements are done by 4-point measurement on the AZO layers

incorporated in the solar cell structure.

The 1D-Solar Cell Capacitance Simulator Software (SCAPS-1D

3.3.07) was used to simulate the obtained experimental results.

The

bandgap energy (E

), electron affinity (χ), and thickness of the CIGSe

absorber are determined according to the Ga/(Ga +In) (GGI) ratio

measured by GD-OES and integrated into the SCAPS simulation pro-

gram using optical absorption parameters of experimentally deposited

CuInSe

and CuGaSe

and further employing a corresponding interpo-

lation algorithm within SCAPS based on Burgelman and Marlein.

The

optical absorption parameters of the other layers, except for CdS,

were experimentally determined. In order to implement parasitic

absorption losses within the CdS layer and ensure that light absorption

in the CdS layer does not contribute to photocurrent, we used neutral

midgap defects in the CdS layer along with the embedded absorption

model in SCAPS. Note that the carrier concentration extracted from

the CV measurement is called as NCV , whereas the corresponding

input parameter for the simulation is referred to as NCIGSe

3|EXPERIMENTAL RESULTS

3.1 |Determining typically harmful thermal stress

conditions

As a starting point, we compared the JV measurements of complete

solar cells with and without Na incorporation obtained before and

FIGURE 1 Schematic illustration of

the sample processing—thermal

annealing after each deposition step and

etching with citric acid, HCl, or HCl

+NaF PDT as well as HCl +AA. It

should be noted that since there are two

types of devices, that is, those with and

those without Na, a glass substrate only

covered in half by a SiO

barrier

indicates that BOTH, devices with AND

without barrier, have been processed.

The colored frames refer to the etching

procedures. For instance, the red frame

indicates that a device without Na that is

annealed directly after CdS deposition is

etched with HCl and rebuilt. In this case,

a complete SiO

barrier is indicated,

because this is applied only to the device

without Na

1036 YETKIN ET AL.

after thermal annealing in air for 20 min to be able to define an overall

critical annealing temperature. Figure 2 displays box plots of the

JVparameters, recorded on devices, which have been annealed at the

given temperature. Clearly, the efficiency decreases with increasing

annealing temperature. Even for relatively low temperatures ≤250C,

slight deteriorations in VOC and FF yielding lower efficiencies for the

devices with and without Na are evident. In the case of solar cells with

Na incorporation, the losses in open-circuit voltage (VOC ), fill factor

(FF), and efficiency (η) are more pronounced for temperatures >250C.

At the highest annealing temperature of 300C for 20 min, the VOC

and FF decreased considerably resulting in an efficiency below 4%. In

the case of the solar cells without Na incorporation, similar trends are

observed, except for overall lower VOC and FF values, which are due

to the absence of Na in the CIGSe absorber. Again, solar cells

degraded considerably after annealing at 300C for 20 min.

Figure 3 shows illuminated (left) and dark (right) JV curves of com-

plete solar cells with Na subjected to air annealing at 300C with an

annealing duration increasing from 30 s to 20 min. Within the first

2 min of annealing, very little changes in VOC and FF can be seen.

There are also no relevant changes in JSC. From then on, VOC and FF

continue to decrease significantly, along with increasing current

losses. In addition, the formation of a mixture of kink and rollover

anomalies in the illuminated JV curves becomes apparent. The current

densities measured in the dark are increasingly blocked with longer

annealing duration. In the light of the results presented in Figures 2

and 3, the total heating flux that the device is exposed to seems more

relevant than the annealing temperature, but only once the annealing

temperature is higher than the critical threshold temperature. For the

devices investigated in this study this threshold lies 250C, as can

be seen in Figure 2. Above the critical annealing temperature, the

FIGURE 2 Results of JV measurements of solar cells with (left) and without (right) Na incorporation air annealed at up to 300C for 20 min; a

fresh device was used for each annealing temperature

FIGURE 3 Illuminated (left) and dark (right) JV curves for the time-resolved annealed solar cells with Na incorporation at 300C in air

YETKIN ET AL.1037

degradation of the devices hinges on the total heating flux (see

Figure 3). This observation is in line with the results shown by

Flammini et al.

It is obvious that annealing at 300C for 20 min is

extremely damaging for the complete solar cells of both kinds. There-

fore, for this study, these experimental stress conditions were chosen,

to study and determine the causes why the various devices degrade

under thermal stress.

3.2 |Comparing thermal stress in air and vacuum

on CIGSe devices with and without Na incorporation

In order to identify the degradation mechanisms in the solar cells after

exposure to thermal stress in more detail, we administered the

annealing procedures as presented in Figure 1. Annealing at 300C

was carried out for 20 min in air or under vacuum after CIGSe (just

before CdS deposition), CdS, i-ZnO, and AZO deposition individually.

For better understanding, the samples are named as shown in Table 1.

The corresponding JV results for each type of sample are shown in

Figure 4.

Again the as-deposited samples with and without Na for air and

vacuum annealing demonstrate different PV parameters, especially

VOC and FF. We attribute the variation in VOC and FF, which is seen

for nominally identical samples, to the general variability often seen in

our CIGSe devices over longer time periods, due to slight variations in

one of the deposited components. In case of solar cells with Na incor-

poration, air annealing after CIGSe (AA-CIGSe-wNa) results in slightly

decreased VOC and FF, whereas vacuum annealing (VA-CIGSe-wNa)

only leads to a slight VOC decrease. Air and vacuum annealing after

CdS, i-ZnO, and AZO deposition show a similarly detrimental effect

on VOC . However, in addition, air annealing leads to a significantly

enhanced decrease in FF, accompanied by a small drop of around

2mAcm

2

in JSC , all in all resulting in a strongly degraded solar cell

performance. Vacuum annealing leads to less deterioration in FF and

no observable change in JSC.

Without Na present, air annealing of the CIGSe absorber layer

(AA-CIGSe-w/oNa) facilitates an increased VOC and FFcompared to

the as-deposited sample (AA-Asdepo-w/oNa). Apart from this, similar

trends to the case with Na present can be observed in the deteriora-

tion of the JV parameters for air and vacuum annealing. As also seen

in Figure 5a,c for the air-annealed devices with and without Na, a

drop in JSC when annealing after CdS, i-ZnO, and strongest after AZO

is noted.

We conclude that the overall degradation behaviors are—in

principle—independent of the annealing atmosphere, noting that this

is not true when annealing bare as-deposited CIGSe absorber layers.

Figure 5a–d displays the JV curves of the air- and vacuum-annealed

solar cells with and without Na that showed the best efficiencies each.

The air-annealed samples after CdS, i-ZnO, and AZO exhibit strong

rollover and kink anomalies, especially in the case of devices with Na

incorporation. On the other hand, only slight rollover behavior can be

seen in case of vacuum-annealed samples after CdS and i-ZnO. Fur-

ther, EQE data are given in Figure 5e–h. There are no relevant changes

in the EQE when annealing after CIGSe in any case. Once CdS has

been deposited prior to air or vacuum annealing of the devices both

with and without Na incorporation, the EQE decreases in the blue por-

tion of the spectrum, with stronger deterioration visible when

annealing samples without Na and more so when annealing after

i-ZnO deposition (AA-i-ZnO-w/oNa). Annealing after AZO (AA-AZO-

wNa and AA-AZO-w/oNa) only shows a detrimental effect on the

EQE when done in air, independent of the Na content. In particular,

sample AA-AZO-wNa exhibits a pathological behavior by giving an

EQE > 1 at around 400 nm. We attribute this to a light-sensitive

behavior in the window and absorber layers, which causes an

TABLE 1 List of the individual sample investigated in this study

Air annealing Vacuum annealing

With Na Without Na With Na Without Na

As-deposited AA-Asdepo-wNa AA-Asdepo-w/oNa VA-Asdepo-wNa VA-Asdepo-w/oNa

Annealed after

CIGSe AA-CIGSe-wNa AA-CIGSe-w/oNa VA-CIGSe-wNa VA-CIGSe-w/oNa

CdS AA-CdS-wNa AA-CdS-w/oNa VA-CdS-wNa VA-CdS-w/oNa

CdS, then HCl etched —AA-CdS-HCl-w/oNa ——

i-ZnO AA-i-ZnO-wNa AA-i-ZnO-w/oNa VA-i-ZnO-wNa VA-i-ZnO-w/oNa

AZO AA-AZO-wNa AA-AZO-w/oNa VA-AZO-wNa VA-AZO-w/oNa

AZO, then citric acid etched AA-AZO-citric-wNa —VA-AZO-citric-wNa VA-AZO-citric-w/oNa

AZO, then HCl etched AA-AZO-HCl-wNa —VA-AZO-HCl-wNa VA-AZO-HCl-w/oNa

AZO, then HCl +NaF PDT AA-AZO-HCl-NaF-wNa —VA-AZO-HCl-wNa —

AZO, then HCl +air annealing ———VA-AZO-HCl-AA-w/oNa

Note: It should be noted that there are four different types of experiments: (1) air annealing (AA) with Na (wNa), (2) air annealing (AA) without Na (w/oNa),

(3) vacuum annealing (VA) with Na (wNa), and (4) vacuum annealing (VA) without Na (w/oNa). For each annealing experiment, there is one reference

sample named here “as-deposited.”Accordingly, as-deposited references are identified by, for example, AA-Asdepo-wNa and VA-Asdepo-wNa. The

remaining of the experiment details are indicated between the prefix “AA or VA”and suffix “wNa or w/oNa.”

1038 YETKIN ET AL.

additional flow of injected charge carriers towards the front contact

during EQE measurement similar to an apparent quantum efficiency as

suggested by Scheer and Schock

or photoconductivity effects

as described by Phillips and Roy.

Figure 5i–l provides information about the charge carrier distribu-

tions within the absorber after the various treatments. Air annealing

of the bare CIGSe absorbers with and without Na (AA-CIGSe-wNa

and AA-CIGSe-w/oNa) brings about an increase in NCV , more

FIGURE 4 Effect of air and vacuum annealing on the JV parameters of solar cells with (left) and without (right) Na incorporation. Please note

that in total, there are four different as-deposited samples (one for each set of air- and vacuum-annealed samples with and without Na, since the

sets were processed at different times)

FIGURE 5 JV (a–d), EQE (e–h), and NCV (i–l) of the best solar cells with and without Na incorporation exposed to air annealing (AA) and

vacuum annealing (VA) after each layer deposition at 300C for 20min

YETKIN ET AL.1039

pronounced for the sample without Na. Vacuum annealing either

leads to a slight decrease in NCV in the case with Na (VA-CIGSe-wNa

sample) or does not affect the NCV for the sample without Na

(VA-CIGSe-w/oNa sample). However, the annealing procedures

applied after CdS, i-ZnO, and AZO regardless of the environment have

striking repercussions on the NCV profiles for both cases with and

without Na. The common observation for all of these samples is a

shift of their NCV profiles to larger dSCR . While the samples with Na

demonstrate decreased NCV , the samples without Na exhibit even

increased NCV along with the entire shift of their NCV profiles. These

observations therefore stress that although the annealing atmosphere

did hardly alter the trend observed in the PV parameters (see

Figure 4), it does crucially affect the manifestation of the degradation

mechanisms in place. To reveal the underlying degradation mecha-

nisms individually, further electrical characterization of all devices and

elemental in-depth profiles measured on the as-deposited and

annealed devices will be analyzed below.

Table 2 shows the correlation between VOC and NCV as deter-

mined by CV for the best devices as observed when annealing after

the various processing steps. The absorber layers for each set of sam-

ples (with and without Na for air and vacuum annealing) are deposited

in separate, nominally equal co-evaporation processes. Na supply is

varied simply by the use of substrates with and without Na diffusion

barrier. The as-deposited samples with Na generally show higher VOC,

which is attributed to the presence of Na during CIGSe deposition. In

Table 2, the section ΔVOC compares experimentally observed values

with those derived from the relation

ΔVCV

OC ≈

kBT

qln NCV

N∗

,ð1Þ

with k

denoting the Boltzmann constant, temperature T, the elemen-

tal charge q, and the charge carrier concentration N∗

CVof the reference

device.

It should be noted that this relation is only true, if the domi-

nant recombination pathway is in the absorber bulk. Hence, values for

ΔVCV

OC are only shown for those devices, for which—as will be indi-

cated below via JV–Tanalysis—the main recombination pathway is

assumed within the absorber bulk. All other values for devices that—

according to the JV–Tanalysis—exhibit the main recombination taken

place at the interface are not calculated (denoted as not applicable

[n.a.]). The value shown for ΔVJV

OC for the as-deposited reference

device for vacuum-annealed set demonstrates that the observed

differences in VOC can be explained by the difference in NCV when Na

is present. Without Na present during growth, this is not the case.

TABLE 2 VOC,NCV,ΔVOC, and RSheet values of the best solar cells with and without Na before and after annealing at 300C for 20 min

Na As-deposited Annealed in

Annealed after Annealed after AZO, then etched and rebuilt

CIGSe CdS i-ZnO AZO Citric acid HCl HCl +NaF PDT

VOC(mV) Yes 652 Air 626 468 481 471 467 528 548

Yes 632 Vacuum 616 395 465 469 467 495 554

No 511 Air 564 342 328 290 —407

—

No 449 Vacuum 458 307 288 342 320 354 537 (AA)

NCV (10

3

) Yes 110 Air 160 4.5 4.1 7.8 4.6 5.0 6.5

Yes 55 Vacuum 31 7.9 11.3 8 9.9 6.0 5.4

No 5.4 Air 74 90 100 340 —6.2

—

No 4.3 Vacuum 4.3 39.6 19.7 37.9 9.8 19.9 30.5 (AA)

ΔVJV

OC /ΔVCV

OC Yes 0/0 Air 26/+10 184/

n.a.

171/

n.a.

181/

n.a.

4/

n.a.

+57/

n.a.

+77/n.a.

(mV) Yes 20/18 Vacuum 36/33 257/

n.a.

187/

n.a.

183/

n.a.

2/n.a. +26/

n.a.

+85/

n.a.

No 141/

n.a.

Air 88/10 310/

n.a.

324/

n.a.

362/

n.a.

—+65/

n.a.

—

No 203/

n.a.

Vacuum 194/

n.a.

345/

n.a.

364/

n.a.

310/

n.a.

22/n.a. +12/

n.a.

+95/n.a.

(AA)

RSheet (Ω/□) Yes 94 Air 107 94 85 959 107 115 101

Yes 65 Vacuum 64 74 73 113 76 76 66

No 95 Air 87 78 80 420 —95

—

No 61 Vacuum 60 71 68 91 70 71 70 (AA)

Note:ΔVOC: for comparability, the values for the “as-deposited and annealed after”devices refer to the “as-deposited”device in the “air-annealed”sample

set with Na, and the values in the section “etched and rebuilt”refer to the “annealed after AZO”in their individual sample set; AA notifies “air annealing.”

The “from JV”values are calculated using ΔVJV

OC ¼V∗

OC VOC, while the “from NCV”values, ΔVCV

OC, are calculated according to Equation (1). AA: after

etching, this sample was only air annealed at 300C for 20 min as alternative to a NaF PDT.

This sample has been air annealed after CdS deposition. After annealing, it was etched with HCl and rebuilt with the respective layers.

1040 YETKIN ET AL.

Only part of ΔVJV

OC can be attributed to a doping effect of the

CIGSe absorber.

The JV–Tmeasurements that were performed on the various

devices with and without Na are shown in Figure 6a–d. From this

measurement, the difference at the linear extrapolation of VOC to 0 K

from high T(gives activation energy [E

]) and from low T(VOC

saturation) can be interpreted as barrier height at the back contact.

Samples grown on substrates with and without Na differ in the

determined values for the E

and the back-contact barrier height (ΦBC

). The AA-Asdepo-wNa and VA-Asdepo-wNa samples show a good

agreement of E

and the bandgap energy extracted from EQE

measurement (EEQE

g), while a ΦBC of about 0.14 eV initially appears for

these samples. On the other hand, the AA-Asdepo-w/oNa and

VA-Asdepo-w/oNa samples indicate E

along with a higher ΦBCof

about 0.20 eV. Interestingly, air annealing of the sample without Na

after CIGSe (AA-CIGSe-w/oNa) raises the E

to a value that is

equal to its EEQE

g. Apart from this, any annealing after CdS deposition

leads to severe decrease in E

smaller than EEQE

g. Furthermore, both air

annealing and vacuum annealing also increase ΦBC of the Na-

containing sample to the level of the back barrier height in the

samples without Na. In contrast to these findings, however, the latter

is not significantly altered by annealing of the samples in any

atmosphere.

Table 2 also lists values for the sheet resistance Rsheet of the TCO

front contact as measured on the as-deposited and annealed devices.

Due to maintenance that was performed on the TCO deposition tool

during the experiments of this work, there is some variability in the

values for the various sets of samples that is well reproduced on glass

references. Due to the roughness of the surface of the CIGSe com-

pared to the glass surface, the TCO resistance measured on the

devices front contact is slightly higher than the reference values on

glass. The one notable effect of the annealing steps on Rsheet is a clear

FIGURE 6 The VOC(T) characteristics under one sun illumination at temperatures from 90 to 320 K for the respective CIGSe solar cells with

and without Na. Each set of sample is separately shown for with Na in (a), for vacuum annealing with Na in (b), for air annealing without Na in (c),

and for vacuum annealing without Na in (d). It should be noted that the table shown in each graph includes the bandgap energy extracted from

EQE measurement air annealing (EEQE

g), activation energy (E

), and resulting back-contact barrier height (ΦBC)

YETKIN ET AL.1041

increase when the complete device is annealed in air (AA-AZO-wNa

and AA-AZO-w/oNa), while there is only a slight increase in Rsheet

after vacuum annealing.

Finally, Figure 7 shows the Cd and Na depth profiles (for ease of

comparison aligned at the point of Cd onset in all devices) of the air-

and vacuum-annealed solar cells with and without Na as measured by

GD-OES. The as-deposited Na profiles in Figure 7b,d display an accu-

mulation of Na mainly at the Mo/CIGSe and less so at the CIGSe/CdS

interface. The VA-Asdepo-wNa sample—even though nominally

identical to sample AA-Asdepo-wNa—exhibits a lower Na content at

both interfaces and in the entire absorber layer compared to the AA-

Asdepo-wNa sample. The lower Na content of the VA-Asdepo-wNa

might substantiate the lower VOC . The main effect of air annealing

after CIGSe on Na is a flattening of the Na depth profile towards the

back contact (see Figure 7b), while the vacuum annealing after CIGSe

only gives rise to a rather slight decrease of the Na signal in the

absorber layer. Annealing after CdS, i-ZnO, and AZO leads to signifi-

cant changes in the Na depth profiles of all Na-containing samples. Na

FIGURE 7 GD-OES Cd depth profiles of the solar cells with (a, c) and without (e, f) Na as well as the Na depth profiles of the solar cells with

Na (b, d) annealed in air and under vacuum at 300C for 20 min

1042 YETKIN ET AL.

diffuses into the CdS and ZnO bilayer and with air annealing after

AZO also towards the surface of the AZO layer (see Figure 7b). On

the other hand, vacuum annealing only provokes Na out-diffusion into

the CdS layer, not into the ZnO bilayer as seen in Figure 7d. Addition-

ally, air annealing after CdS, i-ZnO, and AZO leads to significant Cd

diffusion into the CIGSe absorber with and without Na up to a depth

of 200–400 nm (see Figure 7a,e). In contrast, Cd diffusion into the

CIGSe can hardly be seen after vacuum annealing in either case (see

Figure 7c,f).

3.3 |Partial mitigation of annealing damage

In order to test a possible way to mitigate damage and for closer

investigation of the effects of thermal stress on the window layers,

those samples that were annealed after AZO were etched by HCl

(as indicated in Table 2) to completely remove all layers from the

CIGSe absorber or alternatively by citric acid in order to selectively

remove the ZnO bilayer only (see also Figure 1). In addition to the HCl

etch, a NaF PDT was carried out for both the air- and vacuum-

annealed Na-containing samples, and an air-annealing step at 300C

for 20 min for the Na-free sample directly after the etching process

instead of the NaF PDT.

The corresponding JV curves are displayed in Figure 5a–d. As it

was described above, the JVcurves of the air-annealed solar cells

exhibit a mixture of strong rollover and kink anomalies (see Section 4

for more details). After etching of the window layers and rebuilding

them, these anomalies are partially recovered. Also, etching and

redeposition of the TCO restore Rsheet to values in the range of the

original values seen in Table 2. In particular, the following effects

occur: (1) removing only the TCO (not the CdS) of the Na-containing

sample (AA-AZO-citric-wNa) gives rise to a partial recovery of the

current loss, while the VOC is hardly affected. For the corresponding

vacuum-annealed samples (VA-AZO-citric-wNa and VA-AZO-citric-

w/oNa), the citric acid etch shows no recovery effect. (2) HCl etching

and rebuilding the air-annealed solar cells with Na (AA-AZO-HCl-

wNa) with new CdS and TCO leads to a disappearance of the kink

anomaly along with the complete recovery of the JSC loss, while the

rollover is hardly affected. Additionally, removal and rebuilding of

the complete window layer improves the VOC of the device by more

than 60 mV compared to the AA-AZO-wNa sample. A similar observa-

tion is made for the sample that was vacuum annealed after AZO,

then HCl etched (VA-AZO-HCl-wNa) and rebuilt. Without Na, sample

AA-CdS-HCl-w/oNa, which was air annealed after CdS, then HCl

etched and rebuilt also shows a partial improvement in VOC by 65 mV,

while the sample that was vacuum annealed after AZO, then HCl

etched and rebuilt (VA-AZO-HCl-w/oNa) only shows a rather weak

gain in VOC . (3) When the HCl-etched air-annealed sample, AA-AZO-

HCl-NaF-wNa, is exposed to NaF PDT prior to CdS redeposition, an

additional increase in VOC is notable in conjunction with a remaining

strongly pronounced rollover. (4) Exposing the HCl-etched, vacuum-

annealed sample to a NaF PDT (VA-AZO-HCl-NaF-wNa) improves its

VOC by more than 80 mV and reduces the slight rollover behavior

present before etching. (5) A general observation concerning the JV–T

results from Figure 6a–d is that an HCl etch and NaF PDT raise E

while citric acid etch does not further change the corresponding E

(6) The air-annealing treatment for the sample without Na that was

vacuum annealed after AZO (VA-AZO-HCl-AA-w/oNa) instead of the

NaF PDT leads to an increase in VOC by almost 200 mV along with a

remarkable increase in NCV and fully alleviates the slight rollover

behavior seen in the JV characteristic of its as-deposited reference

(VA-Asdepo-w/oNa). With respect to the JV–T-result in Figure 6d, it

strikingly brings the E

to the same level of EEQE

In summary, this indicates that there is a loss in JSC that can be

mitigated when “rebuilding”the AZO after air annealing. An almost

complete recovery of the JSC loss can only be realized when the CdS

layer is etched off together with the TCO and rebuilt. The annealing-

induced VOC loss can also be partially recovered by rebuilding the

nside of the junction. However, a large part of the total VOC loss

remains in the CIGSe absorber, even though the rebuilding of the

CIGSe/CdS improves the VOC. This again coincides with the observed

evident diffusion of Na through the window layers during annealing, if

present.

4|PRINCIPLE SCAPS-1D MODELS FOR

POSSIBLE DEGRADATION

To establish an understanding for the degradation mechanisms that

may be caused by exposure of the CIGSe solar cell to thermal stress,

some principle models for experimentally observed degradation

mechanisms including compositional changes near the interfaces are

introduced. First, a basic model of the as-deposited CIGSe devices,

which reproduces the JV and CV measurements is presented using

SCAPS-1D.

The device properties that are utilized for the basic model for

an as-deposited, Na-containing device are listed in the first

section of Table S1 and are based on Gloeckler et al.

In Figure 8,

the measured and simulated JV,CV, and EQE curves as well as the

corresponding energy band diagram of the as-deposited CIGSe solar

cell with Na (AA-Asdepo-wNa) are displayed. The conduction band

offset between CIGSe and CdS as well as CdS and ZnO are

assumed to be ΔECIGSe=CdS

C=0.1 eV (Spike) and ΔECdS=ZnO

C¼0:2eV

(cliff), respectively, as proposed by Sozzi et al.

The minority carrier

lifetime is set to τ

=23 ns with bulk recombination via neutral

defects placed at 0.6 eV above the valence band maximum (VBM).

Note that this neutral defect level is used to represent the overall bulk

recombination in the devices. The real distribution of the bulk defects

is known to be more complicated.

An acceptor defect is placed at

the CIGSe/CdS interface at 0.47 eV above the VBM of CIGSe

absorber. This is motivated by the exposure of samples after CIGSe

deposition to air and daylight prior to the CdS processing step

and

enables the simultaneous simulations of the JV and CV characteristics,

which under absence of these defects could not be achieved. In our

model, they show rather weak influence on the corresponding JV

results revealing that the largely dominant recombination mechanism

YETKIN ET AL.1043

is present within the absorber, not at the interfaces (see Figure S1a),

depending on the chosen energetic position and defect density. More-

over, based on the JV–Tmeasurement (see Figure 6a), a back-contact

barrier with a height of 140 meV is added to the basic model.

Characteristic phenomena that are observed in the experimental

JV curves (Figure 5) are mostly the kink (I) and rollover (II) anomalies.

In the literature, these anomalies are often connected with experi-

mental findings, which will be shortly reviewed here. With this,

numerical models will be tailored to fit the experimental findings

observed after the CIGSe solar cells have been exposed to thermal

stress.

4.1 |Kink anomaly

The kink anomaly is generally due to a charge carrier extraction

barrier for the photocurrent resulting in a voltage-dependent pho-

tocurrent collection.

Possible causes for this anomaly are forma-

tion of a thin p

layer in the near CIGSe absorber surface close to

the CdS buffer,

a highly positive conduction band offset (spike)

at the CIGSe/CdS interface,

or deep acceptor trap states in the

CdS layer.

30,31

Here, for our devices that are exposed to thermal

annealing, as based on the JV–Tmeasurement results, an increased

spike formation seems unlikely.

4.2 |Rollover anomaly

The rollover anomaly is described to be a current saturation in the first

quadrant of the JV curve, that is, under forward bias, revealing a

charge carrier injection barrier.

23,32

The first possible cause for that

behavior is the presence of a back-contact barrier representing a hole

injection barrier at the CIGSe/Mo interface.

23,32–38

Another possible

cause is acceptor states at the CdS/ZnO front interface.

However, it

should be noted that these types of defects at the CdS/ZnO interface

can also induce a kink anomaly resulting in lower photocurrent under

forward bias and FF as suggested by Nguyen et al. and Urbaniak and

Igalson.

39,40

In addition to these, Villanueva-Tovar et al. have pro-

posed that a strong cliff at the CdS/ZnO interface can also cause an

injection barrier leading to a rollover anomaly.

FIGURE 8 The measured (a) JV and (b) CV curves of the as-deposited CIGSe solar cell with Na for air annealing (AA-Asdepo-wNa) shown in

Section 3.2 (open circles) and the corresponding simulated JV and CV curves (solid line). (c) Corresponding simulated energy band diagram of the

device under illumination at 0 V including the energy defect levels in the absorber bulk and at the CIGSe/CdS interface. (d) The measured and

fitted EQE curves of the as-deposited CIGSe solar cell with Na incorporation

1044 YETKIN ET AL.

Regarding the integration of the experimentally observed findings

from Section 3.2 into the device simulation, these can be related to

the following possible effects:

•Back contact barrier: the presence of Na during CIGSe absorber

deposition is crucial to catalyze MoSe

formation creating an

ohmic contact between the Mo and the CIGSe layer.

42,43

The

experimentally observed flattening of the Na depth profiles after

all annealing procedures could be a sign for an increased back-

contact barrier. Therefore, in the case of Na-free solar cells, a

higher back-contact barrier is initially assumed. Figure 6a–d

reveals the back-contact barrier height for the corresponding

devices.

•Acceptor states in the CdS and at the CdS/i-ZnO interface:

annealing after CdS, i-ZnO, and AZO leads to Na diffusion espe-

cially into the CdS and the ZnO, as seen in Figure 7b,d. Na atoms

are known to create deep acceptor states in CdS and ZnO at an

energy of around 0.3 eV above the VBM of each layer.

44–46

Since

Na accumulation in the vicinity of the CIGSe absorber surface is

not detected, the formation of a p

layer can be excluded.

Furthermore, a high positive conduction band offset (increased

spike formation) at CIGSe/CdS can be excluded for the devices

annealed after CdS, i-ZnO, and AZO as based on the JV–Tmea-

surement results (see Figure 6). Accordingly, such deep acceptor

states are included into the SCAPS model and used to fit the

Na-containing devices. If no Na is present in the solar cells, no

acceptor trap states are implemented in the corresponding

simulations.

•n-type surface: another experimental observation is strong Cd

diffusion into the CIGSe absorber layer after air annealing. This can

modify the absorber material close to the junction from ptype to

ntype via the formation of CdCu defects.

12,47–50

As it was seen in

Figure 7a, air annealing-induced Cd diffusion into CIGSe cannot be

removed by HCl etching. Therefore, in all cases that the samples

were annealed together with CdS layer on top of CIGSe, the effect

of the Cd diffusion as an n-type surface layer without changing the

whole absorber bandgap grading is used.

•Negative band offset at the CIGSe/CdS interface (cliff formation):

based on JV–Tmeasurements (see Figure 6a–d), a cliff-like energy

band alignment at the CIGSe/CdS interface is assumed when

annealing devices with CdS layer present on top of the CIGSe

absorber, since the conditions Ea<EEQE

gare interpreted as dominant

recombination to occur at the interface. This is considered true

also after etching since—as mentioned above—once it has taken

place, Cd diffusion is permanent and the CIGSe absorber layers are

Na poor. For simulation, the magnitude of the cliff-like band align-

ment at CIGSe/CdS is adjusted by changing the electron affinity of

the CdS layer.

Figure 9 exemplarily displays how the above mentioned possible

degradation mechanisms individually influence the simulated JV and

EQE curves of the basic device model. The corresponding simulated

energy band diagrams are shown as well. If a back-contact barrier of

ΦBC ¼0:4 eV is present at the CIGSe/Mo interface while keeping car-

rier concentration of the CIGSe (NCIGSe

A) constant in the simulation, a

rollover anomaly forms (see Figure 9a,b). However, this manifestation

of the back-contact barrier depends strongly on the carrier concentra-

tion in the absorber layer: for example, decreasing the NCIGSe

Ain the

example gives rise to an increased VOC loss rather than a rollover due

to increased back-contact recombination. This can be clearly seen in

the related band diagram as well.

The effect of acceptor states in the CdS layer is shown in

Figure 9c,d indicating the formation of a kink anomaly. It is apparent

that this leads to reduced FF and JSC. With increasing acceptor defect

density, the severity of the kink anomaly becomes more visible, since

the extraction barrier for generated electrons is more pronounced due

to the compensation of the n-type conductivity within CdS layer by

acceptor states accounting for a reduced FF. Furthermore, the slight

drop in JSC most likely arises from the reduced dSCR as shown in the

respective energy band diagram in Figure 9d. Even though acceptor

states at the CdS/i-ZnO interface show virtually the same effect, even

more pronounced, they additionally lead to the formation of an injec-

tion barrier and therefore of a rollover (see Figure 9e,f). This behavior

is rather similar to the case of above mentioned possible cause of

strong cliff at CdS/ZnO interface for a rollover, as proposed by

Villanueva-Tovar et al.

Finally, in Figure 9g,h, the effect of an n-type layer at the CIGSe

surface is shown. Since it is mainly affecting the JSC, we show the sim-

ulated EQE curves instead of the JV curves. A 50-nm-thick n-type

layer decreases the EQE in the short wavelength region depending

highly on the magnitude of the n-type doping (Nntype

D), thereby reduc-

ing JSC . Regarding the impact of increasing the thickness of n-type

layer at constant Nntype

D, the decrease in EQE is further amplified,

while VOC is slightly increased due to decreased interface recombina-

tion. Consequently, such an n-type layer at the CIGSe surface can lead

to poor blue response seen in EQE curves.

If no Na is present during growth, the basic device model, as

introduced above, has to be adjusted to be able to represent the mea-

sured JV and CV characteristics. As seen from Table 2, in Na-free

devices, NCV is lower, and—as was argued above—a higher back-

contact barrier needs to be introduced. Further, JV–Tmeasurements

imply that a cliff-like band alignment is present at the CIGSe/CdS

interface. It should also be noted that due to the fact that Na

decreases the In–Ga interdiffusion during CIGSe growth, Na-free

devices are expected to exhibit a slightly larger minimum bandgap E

when compared to Na-containing devices as also suggested by Cabal-

lero et al.

In our numerical simulation routine, this, however, is of no

concern, as for each device the experimentally measured in-depth Ga

gradient is used. Finally, in order to be able to fit the measured data,

the minority carrier lifetime of the AA-Asdepo-w/oNa sample has to

be decreased from 23 ns in the basic model with Na to 12 ns when no

Na is present to account for an additional VOC and FF loss, as pro-

posed by Zakay et al.

This is done by increasing the capture cross

section of the neutral defects in the absorber layer. It should also be

noted that within this work, a “good”or “bad”quality of the absorber

layer, as judged by the minority carrier lifetime in the bulk, is tuned by

YETKIN ET AL.1045

an adjustment of only the capture cross section of the neutral defects

in the absorber layer rather than by an increase of the neutral defect

density. Taken as a whole, this device (AA-Asdepo-w/oNa) accord-

ingly suffers mostly by interface recombination, which is justified by

its JV–Tresult, as simulated by SCAPS (see Figure S1b). The simula-

tion parameters for all solar cells with and without Na can be found in

Tables S3, S2, S5, and S4.

5|DISCUSSION—DEGRADATION

MECHANISMS IN CIGSE SOLAR CELLS

UNDER THERMAL STRESS

Obviously, thermal stress—regardless of the atmosphere—results in

severe degradation of the investigated CIGSe solar cells causing

degeneration of electrical and optoelectronic characteristics. Distinct

FIGURE 9 The effects of the various

principle causes individually applied to the

basic device model with Na present as

simulated in SCAPS-1D. The effect of

integration of (a, b) the back-contact barrier

along with NCIGSe

A, (c, d) acceptor trap states

in the CdS layer, (e, f) acceptor states at the

CdS/i-ZnO interface, and (g, h) n-type

surface along with the increasing thickness

and Nntype

Dinto the basic device model on

the JV characteristic (EQE for the n-type

surface) and energy band diagram

1046 YETKIN ET AL.

possible causes as reasons behind these losses have been defined.

Some losses can be recovered to a certain extent after removing the

front contact by the use of citric acid or HCl. Using the observed

trends and their associated effects on the annealed devices, we now

attempt to draw up a more general model for the origins of these

losses.

5.1 |Annealing after CIGSe

In the case of air annealing after CIGSe, the solar cell with Na shows a

lower VOC compared to the AA-Asdepo-wNa device, while NCV does

not show a corresponding decrease but even increases slightly. This is

attributed to both an increased bulk recombination and an increased

back-contact barrier height, as confirmed by JV–Tin Figure 6a, due to

the reduced Na content in the CIGSe absorber and near the CIGSe/

Mo back interface that was clearly seen in Figure 7b. Taken together,

using those effects along with the experimentally determined higher

value for NCIGSe

Ain the simulation, the experimental device characteris-

tics of both JV and CV of the sample AA-CIGSe-wNa can be

reproduced well (see Figures 10a and S4a, respectively). In contrast,

vacuum annealing after CIGSe with Na leads to a decreased VOC .

Here, however, a lower NCV is determined. GD-OES depth profiles

also only show a minor effect on the Na in-depth profile. The ΔVOC

values from JV and CV measurements shown in Table 2 reveal that

most of the VOC loss arises from the decreased NCV for this device.

Correspondingly, the device characteristics can again be reproduced

in SCAPS using decreased NCIGSe

Aalong with a slightly increased ΦBC

(see Figure 10b).

SampleAA-CIGSe-w/oNawithoutNathatwasairannealedafter

CIGSe, on the other hand, exhibits a relatively high VOC increase accom-

panied by an NCV increase in contrast to the AA-Asdepo-w/oNa

FIGURE 10 Simulated JV characteristics of the degraded solar cells in comparison with the experimentally derived JV curves after air (a, c)

and vacuum annealing (b, d) as well as with and without Na, respectively

YETKIN ET AL.1047

sample. Perhaps the most significant finding is that air annealing of

the CIGSe absorber without Na leads to the change of the main

recombination mechanism from the interface (AA-Asdepo-w/oNa) to

the absorber bulk (AA-CIGSe-w/oNa) according to the JV–Tresults

seen in Figure 6c. With ΔVOC ¼88/10 mV from Table 2, the VOC

calculated according to Equation (1) for the AA-CIGSe-w/oNa sample,

much of the VOC loss due to the absence of Na can be mitigated by

air annealing of the bare CIGSe absorber. In spite of the absence of

Na in the absorber, oxygen atoms seem able to passivate Se vacancies

leading to a reduction in the donor concentration as suggested by sev-

eral studies.

53–55

This may hold for both solar cells air annealed after

CIGSe with and without Na incorporation, but only for NCV. However,

the air annealing of the Na-free CIGSe cannot completely make up

the lack of Na as can be seen by the fact that the VOC of this device

(564 mV) is still substantially lower than the one of the Na-containing

AA-Asdepo-wNa reference (652 mV). It would seem to imply that

oxygen-induced passivation of the CIGSe absorber is not exclusively

enough to make the cells as efficient as Na despite the interface pas-

sivation and NCV increase. Comparing the simulated devices of the

AA-Asdepo-wNa and AA-Asdepo-w/oNa, the low VOC and efficiency

are associated with four reasons: the lack of Na leads to more severe

defect-assisted recombination in the absorber bulk (leading to a lower

), a rather low NCV, a higher ΦBC, and a cliff forms at the CIGSe/CdS

interface as is also implied by JV–Tcharacterization in Figure 6a,c.

According to the model proposed here for the AA-CIGSe-w/oNa sam-

ple, a spike at the CIGSe/CdS interface is set to act as interface pas-

sivation due to the increased interface energy bandgap; the absence

of Na on the other hand still results in severe recombination in the

absorber bulk and a higher back-contact barrier. Implementing all of

the above mentioned implications into the basic device model, a good

agreement between the measurements and simulations of the

corresponding JV and CV results (see Figure S4) could be realized as

seen in Figure 10c. We conclude that in the investigated devices, Na,

beside its doping effect, plays a crucial role reducing the severity of

the bulk defects (higher τ

), limiting the back-contact barrier height,

and providing the prominent spike-like band lineup at the CIGSe/CdS

interface. Similarly, the dopant and passivation effects of Na have

been reported by Zakay et al. and Cojocaru-Mirédin.

52,56

Here, a

question for further basic research is posed with respect to the reason

for the formation of the spike-like band lineup at the CIGSe/CdS

interface. Does it rather depend on the band bending at the CIGSe

surface or the electron affinity? In conclusion, annealing of the bare

CIGSe absorber layers seems to be less critical in terms of the degra-

dation mechanisms.

5.2 |Annealing after CdS, i-ZnO, and AZO

It is generally observed for Na-containing devices that only about one

third of the VOC decrease of the samples that are air and vacuum

annealed after CdS, i-ZnO, and AZO can be accounted for by an NCV

decrease, except for the VA-CdS-wNa sample matching only one fifth

of the VOC decrease. According to the model proposed here, the rest

of the VOC losses arise from a combination of a decreased NCIGSe

Aand

in the CIGSe absorber due to Na out-diffusion from CIGSe, an

increased ΦBC due to Na depletion at the CIGSe/Mo interface, and a

cliff formation at the CIGSe/CdS interface. On the other hand, a VOC

decrease by a decreased NCV does not directly hold for the Na-free

devices annealed after CdS deposition, which experimentally show

even higher NCV . However, the measured charge carrier profiles of

these devices could contain defect contributions arising from the front

contact layers, since the AA-CdS-HCl-w/oNa device shows still shifted,

however reduced NCV as seen in Figure 5k. This might be a reason for

the defect contribution that leads to the observed additional increase

in NCV, when annealing is applied to the devices after CdS deposition.

This effect can consistently be observed for the etched and rebuilt

devices with and without Na as well (see Figure 5i–l). Additionally, the

presence of a higher ΦBC can also cause a decrease in the NCV

profile along with the shift to larger dSCR, as revealed by SCAPS simu-

lations (see Figure S2). However, a decreased NCIGSe

Aas an input

parameter in the simulation with increasing ΦBC results in the

entire shift of the NCV profile to the right along with significantly

narrower dSCR at an elevated value of NCV . This should also be

taken into account when interpreting the experimentally

determined NCV profiles. Consequently, the shift of the experimentally

determined NCV profiles to the right can be a strong indicator for an

increased ΦBC , which has been proven by the JV–Tmeasurements in

Figure 6a–d.

Looking at the Na depth profiles of the samples with Na that were

air and vacuum annealed after CdS deposition in Figure 7b,d, a high

amount of Na is obviously located at the CdS layer coming from the

CIGSe absorber regardless of the annealing environment. A compari-

son of the Na depth profiles of the air- and vacuum-annealed devices

reveals that the nature of the CBD–CdS, which contains a consider-

able amount of water and OH



ions,

attracts mobile Na ions from

the CIGSe absorber due to its rather low electronegativity.

These

results clearly unveil that the wet nature of the CBD–CdS buffer layer

or water containing air environments poses a major problem in terms

of the Na out-diffusion from the CIGSe absorber layer. Thereby for

simulation, it is generally assumed that all the samples that are

annealed with a CdS layer present on the CIGSe absorber have

reduced τ

, caused by the absence of Na. In the same vein, comparing

the Na depth profiles of the AA-AZO-wNa and VA-AZO-wNa sam-

ples, it is possible to identify additional Na diffusion towards to the

AZO surface in case of air annealing. This also confirms the triggering

mechanism of Na diffusion by water and/or oxygen in air environ-

ment. Furthermore, Table 2 compares the sheet resistance value RSheet

of the AZO layer of each solar cell, showing that RSheet rapidly

increases upon annealing under presence of either Na or air (and

strongest in case of both). Air annealing-induced additional Na diffu-

sion into the AZO layer correlates to the strongest increase in RSheet

considerably increasing the overall series resistance of the device AA-

AZO-wNa. Interestingly, in view of the Na depth profiles and sheet

resistance values from Table 2, a comparison of the AA-AZO-wNa,

AA-AZO-w/oNa, and VA-AZO-wNa samples provides strong evidence

that Na and humid air have a detrimental effect on the TCO

1048 YETKIN ET AL.

conductivity or that Na enhances a detrimental effect of humid air on

the TCO, as a catalytic effect of Na was described in several studies

before.

16,59

It is stressed that if any of the Na-containing samples is exposed

to thermal stress, they exhibit almost the same ΦBC ≈0:20 eV as in

the case of the Na-free solar cells, which also show a higher ΦBC even

after annealing (see Figure 6a–d). The integration of the above men-

tioned observations into the simulations has generated a good corre-

lation between measurement and simulation for annealed devices

after CdS deposition (see Figure 10a–d). Besides this, the measured

JV curves of the air-annealed devices exhibit a stronger rollover

anomaly than the vacuum-annealed devices. As discussed in Section 4,

the severity of the rollover anomaly hinges on NCIGSe

Aand the ΦBC (see

Figure 9a,b). Accordingly, the strong rollover behavior of the air-

annealed samples with and without Na could only be simulated with

further increased ΦBC and NCIGSe

A(see Tables S3, S2, S5, and S4), even

though their measured NCV is lower.

All the devices annealed after CdS deposition exhibit a clear kink

behavior (voltage-dependent current loss) in the JV characteristics,

which is also discussed in Section 4. As mentioned above, after the

annealing treatment, Na atoms are mainly located within the CdS

layer and also its surroundings, that is, CdS/i-ZnO interface. Na is

known to generate deep acceptor defects in the CdS

and ZnO.

45,46

Considering the experimental observations and possible reasons for a

kink behavior as well as accordingly integrating the deep acceptor

defects into the CdS layer and at the CdS/i-ZnO interface, the

established SCAPS model shows only modest correlation between

measured and simulated JV characteristics for kink behavior. In the

case of the samples without Na, however, there is inconsistency with

this argument, since no interdiffusion apart from Cd diffusion into

CIGSe is detectable from GD-OES in-depth profiles. Therefore, kink

behavior of the annealed samples after CdS deposition without Na is

simulated in SCAPS model by an additional light-dependent shunt

resistance, which reproduces well their measured JV characteristics

along with the other observations as seen in Figure 10c,d.

Another significant aspect we assume is the formation of a cliff at

the CIGSe/CdS interface that is inferred from the difference between

and EEQE

gas determined via JV–Tmeasurements on the devices,

which were annealed after CdS deposition, as shown in Figure 6a–d.

Accordingly, Na-containing devices that are annealed after CdS depo-

sition suffer highly from recombination at the CIGSe/CdS interface. In

the case of Na-free devices, even the as-deposited case shows intrin-

sically high interface recombination at the CIGSe/CdS interface, again

interpreted as a cliff formation. Annealing of this junction with the

presence of CdS on top of CIGSe leads to increased cliff formation.

With these findings integrated into the simulations for all the samples

with and without Na that are annealed after CdS deposition, the

experimentally observed VOC losses could be well simulated.

5.3 |Current loss analysis

For current loss analysis, EQE results are consulted. There is a small

decrease in JSC due to reduced collection in the long-wavelength

region for the samples that are air annealed after CIGSe with and

without Na (AA-CIGSe-wNa and AA-CIGSe-w/oNa). Taking into

FIGURE 11 Schematic representation of the suggested thermal stress-induced degradation mechanisms of the CdS-buffered CIGSe thin-film

solar cells with (a) and without Na (b) in air and under vacuum along with the initial states of the solar cells (as-grown)

YETKIN ET AL.1049

account that an increased charge carrier density observed for these

devices (see Table 2) will cause a reduced width of the space charge

region, a smaller effective collection length is suggested to be the rea-

son for this observation.

60,61

Additionally, the EQE curves of the solar cells that are annealed

after CdS deposition with and without Na are inclined in the wave-

length range between 400 and 1000 nm, which is interpreted as a

confirmation of an n-type layer formation on top of the absorber

(cf. Figure 9g,h) due to the Cd diffusion into CIGSe via the formation

of CdCu defects that can be clearly seen in Figure 7a,e. Without Na in

the devices, this effect is stronger visible. By the aid of the measured

Cd in-depth profiles as seen in Figure 7a,b,e,f, the thickness of the n-

type surface on top of the CIGSe is subsequently chosen 200 and

50 nm for air-annealed and vacuum-annealed devices in the simula-

tions, respectively. There is no indication that the etching and rebuild-

ing of the window layers affected the n-type surface layers, which is

why they are still present in the simulations for those devices as well.

With the presence of the accordingly chosen n-type surface layer and

tailoring the Nntype

Din order to be able to reproduce the experimen-

tally observed effect in their EQE curves, the simulated device perfor-

mances correlate well with the experiment. However, it should also

be noted that—even though an HCl etch could not remove once dif-

fused Cd from the CIGSe surface region—the effectiveness of a possi-

ble n-type surface is also reduced by redeposition of fresh window

layers.

According to the EQE result for the device with Na that was

air annealed after i-ZnO, the response of the short wavelength

region that is characterized by absorption losses in the CdS may

suggest the formation of a layer with a higher bandgap than that

of CdS, such as maybe CdZnS as suggested by Ramanathan et al.

However, we tend to suppose that it is rather the same pathologic

trend as when annealing after AZO with Na (exhibiting an EQE >1)

that is already in effect and which we attribute to a light-sensitive

behavior in the window layers causing an additional flow of

injected charge carriers towards the front contact during measure-

ment of the spectral response. It has been argued that this behav-

ior (EQE > 1) happens only when applying a voltage bias.

contrast, Figure S3 implies that here it is the application of a light

bias that induces this pathologic behavior of the spectral response,

similarly proposed by Phillips and Roy,

while applying voltage bias

seems not to have an effect.

Additional current loss is observed in the JV characteristics (see

Figure 10a,c) for the devices that are air annealed after AZO with and

without Na compared to the samples that are air annealed before

AZO deposition. This clearly indicates that this additional drop in JSC

is driven by the increased AZO resistivity (see Table 2) in those

devices resulting in higher series resistance, as also discussed above.

These increases in AZO resistivity were translated into the SCAPS

simulations as an overall increase in series resistance, resulting in a

good agreement with the experimental data, since an increase of the

resistivity of the TCO cannot directly be integrated into the AZO layer

in SCAPS due to the fact that this is a two-dimensional effect, which

is not applicable to SCAPS-1D.

5.4 |Etching procedures

After removing the ZnO bilayer by etching with citric acid, there is no

significant change seen in VOC and NCV as well as in the shape of the

JVcurve in any of the samples with and without Na (AA-AZO-citric-

wNa, VA-AZO-citric-wNa, and VA-AZO-citric-w/oNa) as compared to

the corresponding samples (AA-AZO-wNa, VA-AZO-wNa, and VA-

AZO-w/oNa). On the other hand, the removal of the CdS layers along

with the ZnO bilayers by HCl etching leads to an increase in VOC of

more than 60mV for the sample AA-AZO-HCl-wNa and 25 mV for

the VA-AZO-HCl-wNa sample. The reader is reminded that citric acid

does not and HCl does remove the CdS together with the TCO front

contact during etching and that in annealed devices, the CdS layer is

extremely Na rich, as seen in Figure 7b. Additionally, we infer from

the JV–Tmeasurements that the HCl-etched sample (AA-AZO-HCl-

wNa) exhibits a smaller cliff-like CBM offset than the AA-AZO-wNa

and AA-AZO-citric-wNa samples (see Figure 6a), indicating that the

Na accumulation in the CdS might be the origin of the magnitude of

the cliff formation at the CIGSe/CdS interface. This would indeed

explain the more pronounced cliff formation at the CIGSe/CdS inter-

face, since Na is known to raise the electron affinity of the CdS

layer.

Applying this to our SCAPS model, we ascribe up to one third

of the annealing-induced VOC loss to the Na diffusion into the CdS

layer regardless of the annealing environment. With reduced electron

affinity of the CdS layer, HCl-etched samples can be fitted well to the

measured characteristics in both cases of air and vacuum annealing

(see Figure 10a–d).

Applying a NaF PDT after the HCl etch on the CIGSe absorber of

the air-annealed sample with Na (AA-AZO-HCl-NaF-wNa) recovers

the VOC by an additional 20 mV. It could nevertheless be argued that

the NaF PDT itself does not contribute to a carrier concentration

increase (see Figure 5i) and seems to be ineffective for the CIGSe bulk

as seen from its Na depth profile in Figure 7b. The effectiveness of

the NaF PDT could be attenuated by the presence of a secondary

phase such as Cu

(S,Se) which is well known to be detrimental for

CIGSe thin-film solar cells.

10,63,64

This was investigated in a separate

experiment, for which the reader may refer to Figure S5. The experi-

ment shows that secondary-phase formation can be excluded in our

case. The relative increase in VOC seen for the sample VA-AZO-HCl-

NaF-wNa is higher than for the sample AA-AZO-HCl-NaF-wNa; how-

ever, both of them have comparable VOC . Applying the NaF PDT to

the vacuum-annealed and HCl-etched sample (VA-AZO-HCl-NaF-

wNa) shows a subtle increase in its Na depth profile in the whole

absorber as can be seen in Figure 7d. Consequently, there are two

main effects of the NaF PDT in terms of the recovery of the degraded

devices: a surface or interface effect and a bulk effect. With respect

to the NaF PDT-treated sample AA-AZO-HCl-NaF-wNa, the fact that

the negative band offset (cliff) at the CIGSe/CdS interface is reduced

results in a good correlation of simulated and experimental device

characteristics as seen in Figure 10a. Hence, here the effectiveness of

the NaF PDT remains at the CdS/n-type CIGSe interface and cannot

reach the CIGSe absorber bulk possibly due to the formation of a

rather thick n-type CIGSe surface, which the Na diffusion is unable to

1050 YETKIN ET AL.

overcome, as indicated above. As a result, it can be concluded that a

NaF PDT supports better junction properties at the CIGSe/CdS inter-

face. On the other hand, the presence of both a decreased negative

band offset (cliff) and an increased τ

of the CIGSe absorber due to

the slightly increased Na content in the CIGSe (bulk effect) for the

NaF PDT-treated sample VA-AZO-HCl-NaF-wNa reproduces

the experimentally determined curves well (see Figure 10b). This coin-

cides with the fact that the assumed thickness of the n-type CIGSe

surface for vacuum-annealed devices is rather thin, which presumably

allows the NaF PDT on this device to be more effective.

Applying an air-annealing treatment to the sample without Na

that was vacuum annealed after AZO and etched with HCl instead of

application of a NaF PDT (VA-AZO-HCl-AA-w/oNa) has led to a simi-

larly improved VOC as the NaF PDT treatment on the samples with Na

and even better diode behaviors (without rollover anomaly). Along

with the remarkable increase in NCV, this again suggests the previously

mentioned oxygen-induced passivation. Its JV–Tcharacteristic shown

in Figure 6d reveals that this device is not limited at the CIGSe/CdS

interface. Therefore, the JV and CVcharacteristics of this device are

well simulated by employing a positive band offset (spike) at the

CIGSe/CdS, as in the case of as-deposited devices with Na. This moti-

vates further investigations of a combination of an air-annealing treat-

ment and a NaF PDT in order to make more efficient and possibly

more thermally stable devices. In conclusion, the latter approach for

the Na-free CIGSe devices illustrates that overall degraded devices

can be recovered or even improved. However, the complete recovery

of this device might be dependent on the thickness of the n-type

CIGSe surface, as stated above. Further investigation should be car-

ried out in order to better understand whether severity of the forma-

tion of n-type CIGSe surface prevents the complete recovery of the

degraded solar cells.

A summary of the main findings and of the principal issues for the

general representation of the degradation mechanisms of the CdS-

buffered CIGSe solar cells is schematically illustrated in Figure 11. It is

suggested that the degradation mechanisms of the CdS-buffered

CIGSe solar cells are in principle independent of the thermal stress

environment in terms of the overall device performances. However,

this does not hold for the AZO degeneration and rollover anomaly. It

was found that the solar cell performances decrease markedly when

the devices are annealed after CdS, i-ZnO, and AZO deposition. Air

annealing leads to severe Cd diffusion into CIGSe, populating the

CIGSe surface with n-type doping, while vacuum annealing does not

bring about severe Cd diffusion. The major issue is Na out-diffusion

from CIGSe absorber into CdS and TCO layers regardless of the envi-

ronment, leading to an increased ΦBC , reduced NCV and τ

, cliff-like

formation at the CIGSe/CdS interface, and acceptor defects in CdS

layer and at CdS/i-ZnO interface, since the presence of water and



bonds in the CBD-deposited CdS layer attracts Na ions due to

their low electronegativity.

Moreover, Na also expedites the degra-

dation of the AZO conductivity, when annealing takes place in air

together with AZO layer. All things considered, it is alluded by help of

the etching procedures and annealing directly after CIGSe deposition

that the CIGSe absorber material is thermally stable. It is worth noting

that the conclusions and decay mechanisms for the CBD–CdS-

buffered CIGSe solar cells that are presented here can only be argued

to apply for the composition CGI 0.9, which is used here, as it repre-

sents a commonly used composition in the field of CIGSe solar cells

and modules. Most of the degradation mechanism we discussed

depend on the diffusivity of Na within the CIGSe, which, for example,

is known to be different in Cu-poor materials from stoichiometric or

Cu-rich compositions as proposed by Nishinaga et al.

Therefore, the

device's degradation mechanisms are likely to also show a depen-

dence of the Cu content of the CIGSe absorber.

6|CONCLUSION

The degradation mechanisms of the CdS-buffered CIGSe solar cells

with and without Na incorporation under thermal stress were investi-

gated. To distinguish degradation effects, annealing was performed on

unfinished layer stacks after CIGSe, CdS, i-ZnO, and AZO deposition.

In addition, the etching processes—removing degraded window com-

ponents from the annealed devices—has been performed as a mitiga-

tion strategy. In all cases, the electrical, optoelectronic, and

compositional properties of the CIGSe devices have been character-

ized before and after thermal stress was applied in air and under vac-

uum conditions. It is proposed that the degradation mechanisms of

the CdS-buffered CIGSe solar cells are in principle independent of the

thermal stress environment. This does however not hold for the AZO

degradation and the rollover anomaly. It was found that the solar cell

performance drops severely when the devices are annealed after CdS,

i-ZnO, or AZO layer deposition. Na diffusion from the CIGSe absorber

into the CdS and TCO layers is the main cause for the observed deg-

radations, since the movement of the Na ions is primarily triggered by

the presence of water and OH



bonds in the CBD-deposited CdS

layer, which both attract Na due to its low electronegativity. By the

same token, Na accelerates the degradation of the AZO conductivity

resulting in an increased device series resistance when annealing

takes place in an air environment. Na may also cause acceptor-like

defects in the CdS, TCO layers, and the respective interfaces. As JV–T

measurements have shown, any annealing procedure carried out after

CdS deposition seems to give rise to a cliff formation in the band

lineup at the CIGSe/CdS interface resulting in a severe VOC drop.

Consequently, Na out-diffusion has a very deleterious effect on the

device performance. In addition, Cd diffusion into the CIGSe absorber

is also observed after annealing, particularly when the devices are

exposed to air annealing. This may cause the formation of an n-type

surface via the formation of CdCu defects. By means of the etching

processes, it is proposed that mitigation of the observed loss seems to

be possible at least in part when redepositing fresh front contact

layers such as CdS and TCO layers. The absence of Na in CIGSe in

Na-free devices and in devices with Na after annealing, who have suf-

fered Na loss by out-diffusion from CIGSe, seems to induce a

decreased τ

and NCV , an increased back-contact barrier height, and

a lower quality junction at the CIGSe/CdS interface due to cliff forma-

tion. Towards increased thermally stable CIGSe devices, no Na or at

YETKIN ET AL.1051

least no Na diffusion from the CIGSe absorber into the buffer/TCO

layers, no Cd diffusion into the absorber, and stable back-contact

properties have to be achieved. Future research should further

develop new thermally and chemically suitable materials or material

combinations for use as buffer layers in CIGSe thin-film solar cells.

ACKNOWLEDGEMENTS

The authors thank B. Bunn, K. Mayer-Stillrich, J. Lauche,

T. Münchenberg, S. Stutzke, M. Kirsch, and I. Dorbandt for their tech-

nical support with respect to device fabrication and S. Stutzke,

D. Schewitz, Z. Salami, S. Ghaderi, and S. Knoop for their support with

device characterization. H. A. Yetkin gratefully acknowledges the

financial support of the Ministry of National Education of the Republic

of Turkey.

DATA AVAILABILITY STATEMENT

The data that support the findings of this study are available from the

corresponding author upon reasonable request.

ORCID

Hasan A. Yetkin https://orcid.org/0000-0002-1401-5866

Rutger Schlatmann https://orcid.org/0000-0002-5951-9435

REFERENCES

1. Nakamura M, Yamaguchi K, Kimoto Y, Yasaki Y, Kato T, Sugimoto H.

Cd-free Cu(In,Ga)(Se,S)

thin-film solar cell with record efficiency of

23.35%. IEEE J Photovolt. 2019;9(6):1863-1867.

2. Shockley W, Queisser HJ. Detailed balance limit of efficiency of p-n

junction solar cells. J Appl Phys. 1961;32(3):510-519.

3. Koida T, Kaneko T, Shibata H. Carrier compensation induced by

thermal annealing in Al-doped ZnO films. Materials. 2017;10(2):141.

4. Yin S, Shirolkar MM, Li J, Li M, Song X, Dong X, Wang H. Influences

of defects evolvement on the properties of sputtering deposited ZnO:

Al films upon hydrogen annealing. AIP Advances. 2016;6(6):65020.

5. Tong C, Yun J, Chen YJ, Ji D, Gan Q, Anderson WA. Thermally dif-

fused Al:ZnO thin films for broadband transparent conductor. ACS

Appl Mater Interfaces. 2016;8(6):3985-3991.

6. NREL. Best research-cell efficiency chart. National Renewable Energy

Laboratory: Golden, Colorado; 2020.

7. Siebentritt S, Avancini E, Bär M, et al. Heavy alkali treatment of

Cu(In,Ga)Se

solar cells: surface versus bulk effects. Adv Energy Mater.

2020;10(8):1903752.

8. Ramanathan K, Contreras MA, Tuttle JR, et al. Effect of heat treatments

and window layer processing on the characteristics of Cu(In,Ga)Se

thin film solar cells. In: Conference Record of the Twenty Fifth IEEE

Photovoltaic Specialists Conference IEEE; 1996:837-840.

9. Kijima S, Nakada T. High-temperature degradation mechanism of

Cu(In,Ga)Se

-based thin film solar cells. Appl Phys Express. 2008;1(7):

75002.

10. Kazmerski LL, Jamjoum O, Ireland PJ, Mickelsen RA, Chen WS. For-

mation, growth, and stability of the CdS/CuInSe

interface. J Vacuum

Sci Technol. 1982;21(2):486-490.

11. Wi JH, Kim TG, Kim JW, Lee WJ, Cho DH, Han WS, Chung YD. Pho-

tovoltaic performance and interface behaviors of Cu(In,Ga)Se

solar

cells with a sputtered-Zn(O,S) buffer layer by high-temperature

annealing. ACS Appl Mater Interfaces. 2015;7(31):17425-17432.

12. Koprek A, Cojocaru-Mirédin O, Wuerz R, Freysoldt C, Gault B,

Raabe D. Cd and impurity redistribution at the CdS/CIGS interface

after annealing of CIGS-based solar cells resolved by atom probe

tomography. IEEE J Photovolt. 2016;7(1):313-321.

13. Fjallstrom V, Salome PMP, Hultqvist A, et al. Potential-induced degra-

dation of Cu(In

1x

,Ga

)Se

thin film solar cells. IEEE J Photovolt. 2013;

3(3):1090-1094.

14. Yamaguchi S, Jonai S, Hara K, et al. Potential-induced degradation of

Cu(In,Ga)Se

photovoltaic modules. Jpn J Appl Phys. 2015;54(8S1):

8KC13.

15. Salomon O, Hempel W, Kiowski O, et al. Influence of molybdenum

back contact on the PID effect for Cu(In,Ga)Se

solar cells. Coatings.

2019;9(12):794.

16. Theelen M, Hans V, Barreau N, Steijvers H, Vroon Z, Zeman M. The

impact of alkali elements on the degradation of CIGS solar cells. Prog

Photovolt: Res Appl. 2015;23(5):537-545.

17. Heinemann MD, Mainz R, Österle F, Rodriguez-Alvarez H, Greiner D,

Kaufmann CA, Unold T. Evolution of opto-electronic properties dur-

ing film formation of complex semiconductors. Sci Reports. 2017;7:

45463.

18. Bär M, Klaer J, Weinhardt L, et al. Cu

2x

surface phases and their

impact on the electronic structure of CuInS

thin films—a hidden

parameter in solar cell optimization. Adv Energy Mater. 2013;3(6):

777-781.

19. Kodalle T, Greiner D, Brackmann V, et al. Glow discharge optical

emission spectrometry for quantitative depth profiling of CIGS thin-

films. J Anal Atom Spectrom. 2019;34(6):1233-1241.

20. Burgelman M, Nollet P, Degrave S. Modelling polycrystalline semicon-

ductor solar cells. Thin Solid Films. 2000;361:527-532.

21. Burgelman M, Marlein J. Analysis of graded band gap solar cells with

SCAPS. In: Proceedings of the 23rd European Photovoltaic Solar

Energy Conference; 2008:2151-2155.

22. Flammini MG, Debernardi N, Le Ster M, Dunne B, Bosman J,

Theelen M. The influence of heating time and temperature on the

properties of CIGSSe solar cells. International Journal of Photoenergy

2016; 2016.

23. Scheer R, Schock HW. Chalcogenide Photovoltaics: Physics, Technolo-

gies, and Thin Film Devices: John Wiley & Sons; 2011.

24. Phillips JE, Roy M. Resistive and photoconductive effects in spectral

response measurements. In: Conference Record of the Twentieth

IEEE Photovoltaic Specialists Conference IEEE; 1988:1614-1617.

25. Ott T, Schönberger F, Walter T, Hariskos D, Kiowski O, Salomon O,

Schäffler R. Verification of phototransistor model for Cu(In,Ga)Se

solar cells. Thin Solid Films. 2015;582:392-396.

26. Gloeckler M, Fahrenbruch AL, Sites JR. Numerical modeling of CIGS

and CdTe solar cells: setting the baseline. In: 3rd World Conference

on Photovoltaic Energy Conversion, 2003. Proceedings of, Vol. 1

IEEE; 2003:491-494.

27. Sozzi G, Troni F, Menozzi R. On the combined effects of

window/buffer and buffer/absorber conduction-band offsets, buffer

thickness and doping on thin-film solar cell performance. Solar Energy

Mater Solar Cells. 2014;121:126-136.

28. Karki S, Paul P, Rajan G, et al. Analysis of recombination mechanisms

in RbF-treated CIGS solar cells. IEEE J Photovolt. 2018;9(1):313-318.

29. Hölscher T, Schneider T, Maiberg M, Scheer R. Impact of air-light

exposure on the electrical properties of Cu(In,Ga)Se

solar cells. Prog

Photovolt: Res Appl. 2018;26(11):934-941.

30. Rothwarf A. The CdS/Cu

S solar cell: basic operation and anomalous

effects. Solar Cells. 1980;2(2):115-140.

31. Pudov AO, Sites JR, Contreras MA, Nakada T, Schock HW. CIGS J–V

distortion in the absence of blue photons. Thin Solid Films. 2005;480:

273-278.

32. Rockett A, Van Duren JKJ, Pudov A, Shafarman WN. First quadrant

phototransistor behavior in CuInSe

photovoltaics. Solar Energy Mater

Solar Cells. 2013;118:141-148.

1052 YETKIN ET AL.

33. Ledinek D, Donzel-Gargand O, Sköld M, Keller J, Edoff M. Effect of

different Na supply methods on thin Cu(In,Ga)Se

solar cells with

rear passivation layers. Solar Energy Mater Solar Cells. 2018;

187:160-169.

34. Neugebohrn N, Hammer MS, Neerken J, Parisi J, Riedel I. Analysis of

the back contact properties of Cu(In,Ga)Se

solar cells employing the

thermionic emission model. Thin Solid Films. 2015;582:332-335.

35. Puttnins S, Hammer MS, Neerken J, et al. Impact of sodium on the

device characteristics of low temperature-deposited Cu(In,Ga)Se

solar cells. Thin Solid Films. 2015;582:85-90.

36. Richter M, Schubbert C, Eraerds P, Parisi J, Riedel I, Dalibor T, Palm J.

Comprehensive simulation model for Cu(In,Ga)(Se,S)

solar cells. Solar

Energy Mater Solar Cells. 2015;132:162-171.

37. Koishiyev GT, Sites JR, Kulkarni SS, Dhere NG. Determination of back

contact barrier height in Cu(In,Ga)(Se,S)

and CdTe solar cells. In:

2008 33rd IEEE Photovoltaic Specialists Conference IEEE; 2008:1-3.

38. Demtsu SH, Sites JR. Effect of back-contact barrier on thin-film CdTe

solar cells. Thin Solid Films. 2006;510(1–2):320-324.

39. Nguyen Q, Orgassa K, Koetschau I, Rau U, Schock HW. Influence

of heterointerfaces on the performance of Cu(In,Ga)Se

solar cells

with CdS and In(OH

) buffer layers. Thin Solid Films. 2003;431:

330-334.

40. Urbaniak A, Igalson M. Metastable variations of the fill factor in CIGS

thin film solar cells. MRS Online Proceedings Library Archive, 1165;

2009.

41. Villanueva-Tovar A, Kodalle T, Kaufmann CA, Schlatmann R, Klenk R.

Limitation of current transport across the heterojunction in Cu(In,Ga)

solar cells prepared with alkali fluoride postdeposition treatment.

Solar RRL. 2020;4(4):1900560.

42. Caballero R, Kaufmann CA, Eisenbarth T, et al. Influence of Na on

Cu(In,Ga)Se

solar cells grown on polyimide substrates at low temper-

ature: impact on the Cu(In,Ga)Se

/Mo interface. Appl Phys Lett. 2010;

96(9):92104.

43. Caballero R, Nichterwitz M, Steigert A, Eicke A, Lauermann I,

Schock HW, Kaufmann CA. Impact of Na on MoSe

formation at the

CIGSe/Mo interface in thin-film solar cells on polyimide foil at low

process temperatures. Acta Materialia. 2014;63:54-62.

44. Varley JB, Lordi V, He X, Rockett A. First principles calculations of

point defect diffusion in CdS buffer layers: implications for Cu(In,Ga)

(Se,S)

and Cu

(Zn,Sn)(Se, S)

-based thin-film photovoltaics. J Appl

Phys. 2016;119(2):25703.

45. Meyer BK, Sann J, Zeuner A. Lithium and sodium acceptors in ZnO.

Superlattice Microst. 2005;38(4–6):344-348.

46. Meyer BK, Stehr J, Hofstaetter A, Volbers N, Zeuner A, Sann J. On the

role of group I elements in ZnO. Appl Phys A. 2007;88(1):119-123.

47. Nakada T. Nano-structural investigations on Cd-doping into Cu(In,Ga)

thin films by chemical bath deposition process. Thin Solid Films.

2000;361:346-352.

48. Liao D, Rockett A. Cd doping at the CuInSe

/CdS heterojunction.

J Appl Phys. 2003;93(11):9380-9382.

49. Persson C, Zhao YJ, Lany S, Zunger A. n-type doping of CuInSe

and

CuGaSe

.Phys Rev B. 2005;72(3):35211.

50. Cojocaru-Mirédin O, Choi P, Wuerz R, Raabe D. Exploring the p-n

junction region in Cu(In,Ga)Se

thin-film solar cells at the nanometer-

scale. Appl Phys Lett. 2012;101(18):181603.

51. Caballero R, Kaufmann CA, Eisenbarth T, et al. The effect of NaF

precursors on low temperature growth of CIGS thin film solar cells

on polyimide substrates. Phys Status Solidi (a). 2009;206(5):

1049-1053.

52. Zakay N, Stange H, Alpern H, et al. Phototransport properties of

CuInSe

thin films: the influence of Na and planar defects. Phys Rev

Appl. 2020;14(2):24005.

53. Cahen D, Noufi R. Defect chemical explanation for the effect of air

anneal on CdS/CuInSe

solar cell performance. Appl Phys Lett. 1989;

54(6):558-560.

54. Dirnstorfer I, Burkhardt W, Kriegseis W, et al. Annealing studies on

Cu(In,Ga)Se

: the influence of gallium. Thin Solid Films. 2000;361:

400-405.

55. Rau U, Braunger D, Herberholz R, Schock HW, Guillemoles JF,

Kronik L, Cahen D.. Oxygenation and air-annealing effects on the

electronic properties of Cu(In,Ga)Se

films and devices. J Appl Phys.

1999;86(1):497-505.

56. Cojocaru-Mirédin O. Role of planar defects in Cu(In,Ga)Se

thin film

solar sells. In: 2020 47th IEEE Photovoltaic Specialists Conference

(PVSC) IEEE; 2020:2623-2626.

57. Kylner A, Rockett A, Stolt L. Oxygen in solution grown CdS films for

thin film solar cells. In: Solid State Phenom, Vol. 51 Trans Tech Publi-

cations; 1996:533-540.

58. Gordy W, Thomas WJO. Electronegativities of the elements. J Chem

Phys. 1956;24(2):439-444.

59. Theelen M, Barreau N, Hans V, Steijvers H, Vroon Z, Zeman M. Deg-

radation of CIGS solar cells due to the migration of alkali-elements.

In: 2015 IEEE 42nd Photovoltaic Specialist Conference (PVSC) IEEE;

2015:1-6.

60. Rajan G, Karki S, Poudel D, et al. Study of instabilities and degradation

due to moisture ingress in the molybdenum back contact of Cu(In,Ga)

solar cells. In: 2018 IEEE 7th World Conference on Photovoltaic

Energy Conversion (WCPEC) (A Joint Conference of 45th IEEE PVSC,

28th PVSEC & 34th EU PVSEC) IEEE; 2018:3037-3039.

61. Karki S, Deitz JI, Rajan G, et al. Impact of water ingress on molybde-

num thin films and its effect on Cu(In,Ga)Se

solar cells. IEEE J Photo-

volt. 2019;10(2):696-702.

62. Yılmaz S, Atasoy Y, Tomakin M, Bacaksız E. Comparative studies of

CdS, CdS:Al, CdS:Na and CdS:(Al–Na) thin films prepared by spray

pyrolysis. Superlattice Microst. 2015;88:299-307.

63. Zweigart S, Walter T, Koble C, Sun SM, Ruhle U, Schock HW.

Sequential deposition of Cu(In,Ga)(S,Se)

. In: Proceedings of 1994

IEEE 1st World Conference on Photovoltaic Energy Conversion-

WCPEC (A Joint Conference of PVSC, PVSEC and PSEC), Vol. 1 IEEE;

1994:60-67.

64. Hsieh TP, Chuang CC, Wu CS, Chang JC, Guo JW, Chen WC. Effects

of residual copper selenide on Cu(In,Ga)Se

solar cells. Solid-State

Electron. 2011;56(1):175-178.

65. Nishinaga J, Sugaya T. Crystalline characteristics of epitaxial Cu(In,Ga)

layers on GaAs (001) substrates. In: 2020 47th IEEE Photovoltaic

Specialists Conference (PVSC) IEEE; 2020:2251-2257.

66. Ott T, Walter T, Unold T. Phototransistor effects in Cu(In,Ga)Se

solar

cells. Thin Solid Films. 2013;535:275-278.

67. Hegedus SS, Shafarman WN. Thin-film solar cells: device measure-

ments and analysis. Prog Photovolt: Res Appl. 2004;12(2–3):155-176.

SUPPORTING INFORMATION

Additional supporting information may be found online in the

Supporting Information section at the end of this article.

How to cite this article: Yetkin HA, Kodalle T, Bertram T, et al.

Decay mechanisms in CdS-buffered Cu(In,Ga)Se

thin-film

solar cells after exposure to thermal stress: Understanding the

role of Na. Prog Photovolt Res Appl. 2021;29:1034–1053.

https://doi.org/10.1002/pip.3438

YETKIN ET AL.1053