# Rubidium Fluoride Absorber Treatment for Wide-Gap and Bifacial Ag(In,Ga)Se2 Solar Cells

Jan Keller, Sapna Mudgal, Natalia Martin, Olivier Donzel-Gargand, Marika Edoff

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Jan Keller, Sapna Mudgal, Natalia Martin, Olivier Donzel-Gargand, Marika Edoff. Rubidium Fluoride Absorber Treatment for Wide-Gap and Bifacial Ag(In,Ga)Se2 Solar Cells. Solar RRL, In press, 10.1002/solr.202500423. [hal-05212655](https://HAL Id: hal-05212655)

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# Rubidium Fluoride Absorber Treatment for Wide-Gap and Bifacial Ag(In,Ga)Se₂ Solar Cells

Jan Keller | Sapna Mudgal | Natalia M. Martin | Olivier Donzel-Gargand | Marika Edoff

Ångström Laboratory, Division of Solar Cell Technology, Uppsala University, Uppsala, Sweden

**Correspondence:** Jan Keller (jan.keller@angstrom.uu.se)

**Received:** 4 June 2025 | **Revised:** 17 June 2025 | **Accepted:** 20 June 2025

**Funding:** European union programme HORIZON, Grant/Award Numbers: HORIZON-CL5-2021-D3–02 101075626; Swedish Research Council, Grant/Award Number: 2020-04065; MAX IV

**Keywords:** AIGS | bifacial solar cell | CIGS | RbF-PDT | wide-gap chalcopyrite

## ABSTRACT

This study investigates the impact of a RbF post-deposition treatment (PDT) on the properties of Cu-free, wide-gap ($E_G$ = 1.5 eV) Ag(In,Ga)Se₂ (AIGS) solar cells with either opaque Mo or transparent In₂O₃:W back contacts. The formation of a Rb-In-Se surface phase and Cd in-diffusion from the buffer into this layer are detected. In addition, Rb is mainly located in Ag-depleted and In-enriched grain boundaries and at the back interface. The RbF-PDT leads to decreased fill factor (FF) values, without providing any gain in open-circuit voltage ($V_{OC}$). This FF deterioration is more pronounced when a In₂O₃:W back contact is used. It is suggested that transport barriers form at the Rb-In-Se/AIGS front and at the AIGS/GaOₓ/In₂O₃:W back interfaces after RbF-PDT. In addition, the GaOₓ thickness at the rear electrode increases. The absence of a $V_{OC}$ boost may be explained by the missing doping increase after the RbF-PDT. Lastly, the AIGS solar cells from this study show slightly lower efficiencies as compared to Cu-containing devices with the same $E_G$ but higher Ga content. However, the collection at rear illumination is marginally improved, reaching up to 72% bifaciality in short-circuit current for the best cell with an In₂O₃:W back contact.

## 1 Introduction

Today, the incorporation of the heavy alkali (HA) elements potassium, rubidium, or cesium into low-gap (Agₓ,Cu₁₋ₓ)(Inᵧ,Ga₁₋ᵧ)(Sz, Se₁₋z)₂-based absorbers is standard for solar cell applications, since highest efficiencies ($η$ > 22%) cannot be reached if only sodium (mandatory for decent performance) is added [1–3]. In most cases, the HA species is supplied in the form of evaporated (HA)F in a post absorber deposition treatment (HA-PDT), with the first report of the beneficial effect of a potassium fluoride-post-deposition treatment (KF-PDT) dating back to 2013 [4]. It was observed that the HA-PDT allows growing substantially thinner CdS buffer layers without losses in open-circuit voltage ($V_{OC}$) or fill factor (FF), resulting in higher short-circuit current densities ($J_{SC}$). Possible explanations are a facilitated Cd in-diffusion (creating type inversion) [4, 5] and an earlier formation of a fully covering

CdS layer during chemical bath deposition (CBD) [6, 7]. Typically, the HA-PDT also leads to higher $V_{OC}$ values, but the underlying physical mechanism is still under debate [3, 8–11]. Prevalent suggestions are that, after HA-PDT, the recombination rates are reduced i) *at the interface to the buffer* by the formation of a very thin, wide-gap (Alkali,Cu)-(In,Ga)-Se surface layer (possibly without Ga and Cu) [12–17], ii) *at grain boundaries* via a reduction/ passivation of charged defects [18–20], or iii) *in the grain interiors* due to a reduced concentration of deep defects [21, 22] (deemed unlikely by others [18]). Alternative explanations are an increased doping density [23–25] (but lower doping was observed as well [26]) and reduced potential fluctuations [27].

Comparably few HA-PDT studies were conducted on chalcopyrite absorbers with bandgap energies ($E_G$) > 1.4 eV, as required for a top cell in a tandem PV device [28, 29]. In case of cells based

<footnote>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.</footnote>
on a pure ternary CuGaSe₂ (CGS) absorber, a HA-PDT did not lead to a significant increase in performance [15, 30, 31]. It was proposed that the absence of a HA-In-Se (no indium in CuGaSe₂) surface phase [15] or the formation of more detrimental interface defects [31] prevent any beneficial effect. However, a small V<sub>OC</sub> gain was observed when RbF was added during the final stage of absorber formation and not as a PDT [32]. Zahedi-Azad et al. showed that a HA-PDT can lead to an improved V<sub>OC</sub> and efficiency for Cu(In,Ga)Se₂ (CIGS) solar cells with a [Ga]/([Ga]+[In]) ratio (GGI) up to 0.8 [33, 34]. They claim that while recombination at the CIGS/CdS interface is still limiting, it can be significantly mitigated by the HA-PDT. However, for too large amounts of (HA)F a transport barrier at the front contact was claimed to lead to FF losses [33]. More recently, we tested a RbF-PDT for (Ag,Cu) (In,Ga)Se₂ (ACIGS) solar cells with a GGI ≈ 0.7 and a [Ag]/([Ag]+[Cu]) ratio (AAC) of 0.5–0.7 and found a statistically significant V<sub>OC</sub> increase, leading to a top efficiency of 16.3% [35]. Photoelectron-spectroscopy measurements indicate that CIGS surfaces with a high GGI are less modified by a HA-PDT and no Rb-In-Se surface layer was detected after RbF-PDT on CIGS with GGI = 0.9 [36]. All aforementioned works are summarized in Table 1.

Overall, it can be concluded that, according to the previous investigations, the absorber should contain a sufficient (minimum) amount of indium to utilize any beneficial effect of a HA-PDT. In addition, similar to low-gap absorbers [37–40], the (HA)F dose can also become too large for wide-gap CIGS films, leading to FF deterioration [30, 33, 35].

In the present work, we investigate the effect of a RbF-PDT for Cu-free Ag(In,Ga)Se₂ (AIGS, i.e., AAC = 1) solar cells with a GGI of 0.5, leading to E<sub>G</sub> ≈ 1.5 eV. As compared to previous works on wide-gap

ACIGS solar cells [35, 41–50], the Cu was removed altogether to reach the same bandgap, but with a lower Ga content [51]. The intention is that the lower GGI improves the absorber bulk properties (i.e., electron lifetime), as observed for Ag-free CIGS solar cells [52–63]. Since a positive conduction-band offset (CBO) of about +0.35 eV is predicted at the CdS/AIGS interface for the used absorber composition [42], interface recombination is expected to play only a minor role, but a slightly detrimental effect on FF may occur [64]. Up to this point, the total exclusion of Cu to process wide-gap AIGS solar cells was barely tested before and, to the best of our knowledge, only for GGI ≥ 0.8 and without HA-PDT [65–69].

To use a wide-gap A(C)IGS device as a top cell in a tandem structure, the back-contact material needs to be transparent. In the past, different transparent conductive oxides were tested and a concise summary of most works can be found in [48, 49]. A good compromise between a sufficiently low sheet resistance (R<sub>sheet</sub>) and a high near-infrared (NIR) transparency in the final device was achieved by using an In₂O₃:W (IOW) back contact [49, 50]. Thus, in this work we compare the effect of a RbF-PDT for AIGS solar cells with standard, opaque Mo to devices with transparent IOW back contacts, both subjected to the exact same absorber deposition runs. In addition, a reduction of the standard CdS thickness was evaluated for each sample, too.

## 2 | Results and Discussion

### 2.1 | Impact of the RbF-PDT on Solar Cell Performance

Four individual Ag(In,Ga)Se₂ deposition runs, with different RbF-PDT thicknesses (0 to 20 nm, in Se atmosphere), were conducted.

<table><caption>**TABLE 1** | Compilation of works which studied the effect of heavy-alkali incorporation into wide-gap (E<sub>G</sub> = 1.4 – 1.7 eV in all cases) chalcograph absorbers for solar cells with CdS buffer layers. The effect on V<sub>OC</sub> is listed together with a note, summarizing the, for this work, most important conclusions.</caption><thead><tr><th>Study</th><th>Absorber</th><th>GGI</th><th>Alkali type</th><th>V<sub>OC</sub> change</th><th>Note</th></tr></thead><tbody><tr><td>Ishizuka 2018 [30]</td><td>CGS</td><td>1</td><td>RbF-PDT</td><td>~ +10 mV</td><td>FF decreases slightly and J<sub>SC</sub> increases. Only very moderate efficiency gain.</td></tr><tr><td>Taguchi 2018 [15]</td><td>CGS</td><td>1</td><td>RbF-PDT</td><td>No change</td><td>Suggestion: Absence of Rb-In-Se phase (no In) reason for unchanged V<sub>OC</sub>.</td></tr><tr><td>Ishizuka 2020 [31]</td><td>CGS</td><td>1</td><td>CsF-PDT</td><td>~ -50 mV</td><td>Suggestion: More detrimental defects at CdS interface with CsF-PDT.</td></tr><tr><td>Ishizuka 2022 [32]</td><td>CGS</td><td>1</td><td>RbF in-situ (3rd stage)</td><td>~ +20 mV</td><td>Suggestion: Rb reduces interface recombination.</td></tr><tr><td>Zahedi-Azad 2019 [34]</td><td>CIGS</td><td>0.6–0.8</td><td>CsF-, RbF-, aKF-PDT</td><td>> +100 mV</td><td>Claim: HA-PDT reduces interface recombination. EQE also improves.</td></tr><tr><td>Zahedi-Azad 2020 [33]</td><td>CIGS</td><td>0.8</td><td>KF-PDT</td><td>Up to +66 mV</td><td>Too much KF leads to barrier at front contact and corresponding FF losses</td></tr><tr><td>Keller 2022 [35]</td><td>ACIGS (AAC = 0.5–0.7)</td><td>0.7</td><td>RbF-PDT</td><td>~ +25 mV</td><td>Slightly reduced FF values for most samples</td></tr><tr><td>Both 2025 [36]</td><td>CIGS</td><td>0.9</td><td>RbF-PDT</td><td>+30 mV</td><td>Surfaces with high GGI less modified by HA-PDT than with low GGI. No evidence of Rb-In-Se surface phase.</td></tr></tbody></table>
In each run, soda-lime glass (SLG) substrates coated with Mo or IOW were added and sodium was provided by a NaF-precursor film on top of the back contact. In the case of the transparent back contact (TBC), the NaF addition was chosen to be lower (7 nm) as compared to Mo (15 nm), since earlier works suggest that a NaF thickness of ≥10 nm leads to excessive $GaO_x$ formation at the absorber/TBC interface, resulting in FF losses [48]. To study a potential reduction of the CdS thickness without performance losses, as enabled by the HA-PDT for low-gap (A)CIGS solar cells, two different CdS buffer layer thicknesses, a standard ≈65 nm thick ('std') and a ≈35 nm thick ('thin') film, are compared for each sample.

An overview of all samples is provided in Table 2, which lists the corresponding integral compositions (I/III = [Ag]/([Ga]+[In])) and absorber layer thicknesses ($d_{abs}$), as deduced from X-ray fluorescence (XRF) measurements, as well as the bandgap energies, as extracted from external quantum efficiency (EQE) measurements (inflection point method). As is evident from the table, the absorber properties are very similar in terms of composition and thickness for all samples, allowing a fair evaluation of the effect of the RbF-PDT. Intentionally, very close-stoichiometric AIGS compositions were targeted, since for absorbers with AAC ≥ 0.5, I/III values >0.90 lead to a significant widening of the space charge region, resulting in adequate carrier collection, despite the low diffusion length in wide-gap chalcopyrite material [35, 41, 44, 70, 71]. A $d_{abs}$ value of ≈1.1 µm was previously identified to be optimal for maximizing the efficiency under rear illumination, without sacrificing too much of the performance level at front illumination [49, 50].

Figure 1 shows the current-density versus voltage (JV) characteristics (a, b) and EQE spectra (c, d) of the best cell for each sample with a Mo and IOW back contact, respectively. Here, we only show the results for the thinner CdS, since the samples with standard-thick buffers show very similar trends (see Figure S1 in supporting information). All corresponding JV parameters can be found in Table 2, including the short-circuit current densities as derived from integration of the EQE spectra ($J_{SC,qe}$) and the corresponding, spectral-mismatch and probe-shading corrected efficiency values ($\eta_{qe} = \frac{J_{SC,qe}}{J_{SC,iv}} \cdot \eta_{iv}$).

It is obvious that the efficiency is monotonously decreasing with increasing RbF dose for all samples. The most pronounced loss is a reduction in FF, indicating the creation of a transport barrier. It was suggested that the excessive formation of an HA-In-Se phase at the absorber surface (like $RbInSe_2$) may be responsible for FF losses upon too intense HA-PDTs [8, 39, 40, 72]. Since the exclusion of Cu in the used AIGS film leads to a significant downshift of the conduction band minimum [42], an even larger CBO would be expected at a potential $RbInSe_2$/AIGS interface, leading to a more detrimental transport barrier. Alternatively, it was proposed that when too many Na atoms are replaced by heavier alkalis at the back contact interface, a significant back contact barrier establishes, explaining the FF losses [26, 73].

It is evident that the FF reduction by the RbF-PDT is more pronounced when IOW is used instead of Mo. Strong current blocking was observed before for chalcopyrite solar cells with TBCs (not if based on $SnO_2$ [48]), which is often explained by the formation of $GaO_x$ [74] that is facilitated by the presence of Na(F) [48, 75–77]. It may be speculated that the in-diffusion of Rb also

<table><caption>TABLE 2 | Absorber properties and corresponding JV parameters at front illumination (best cells) for all samples processed with varying RbF-PDT doses. The $J_{SC,bif}$ value expresses the percentage of $J_{SC,qe}$ reached at rear in relation to front illumination. The parameters marked with (*) are not reliable, due to strong current blocking in JV ($J_{SC,qe}$ much higher than $J_{SC,iv}$).</caption><thead><tr><th>d<sub>RbF</sub><br>(nm)</th><th>GGI</th><th>I/III</th><th>d<sub>abs</sub><br>(µm)</th><th>E<sub>G,qe</sub><br>(eV)</th><th>BC<br>material</th><th>CdS</th><th>V<sub>OC</sub><br>(mV)</th><th>FF<br>(%)</th><th>J<sub>SC,iv</sub><br>(mA/cm<sup>2</sup>)</th><th>J<sub>SC,qe</sub><br>(mA/cm<sup>2</sup>)</th><th>η<sub>iv</sub><br>(%)</th><th>η<sub>qe</sub><br>(%)</th><th>J<sub>SC,bif</sub><br>(%)</th></tr></thead><tbody><tr><td rowspan="3">0</td><td rowspan="3">0.50</td><td rowspan="3">0.92</td><td rowspan="3">1.05</td><td rowspan="3">1.47</td><td rowspan="3">IOW</td><td>thin</td><td>804</td><td>62.5</td><td>19.3</td><td>23.1</td><td>9.7</td><td>11.6</td><td>72</td></tr><tr><td>std</td><td>796</td><td>63.0</td><td>18.2</td><td>19.1</td><td>9.1</td><td>9.6</td><td>50</td></tr><tr><td>Mo</td><td>thin</td><td>858</td><td>74.4</td><td>19.4</td><td>23.0</td><td>12.4</td><td>14.7</td><td>-</td></tr><tr><td rowspan="3">5</td><td rowspan="3">0.49</td><td rowspan="3">0.92</td><td rowspan="3">1.17</td><td rowspan="3">1.46</td><td rowspan="3">IOW</td><td>thin</td><td>777</td><td>47.0</td><td>20.7</td><td>23.6</td><td>7.5</td><td>8.6</td><td>53</td></tr><tr><td>std</td><td>714</td><td>41.3</td><td>18.7</td><td>20.9</td><td>5.5</td><td>6.2</td><td>54</td></tr><tr><td>Mo</td><td>thin</td><td>857</td><td>69.1</td><td>19.9</td><td>23.6</td><td>11.8</td><td>14.0</td><td>-</td></tr><tr><td rowspan="3">10</td><td rowspan="3">0.50</td><td rowspan="3">0.95</td><td rowspan="3">1.13</td><td rowspan="3">1.47</td><td rowspan="3">IOW</td><td>thin</td><td>758</td><td>42.2</td><td>19.1</td><td>22.1</td><td>6.1</td><td>7.1</td><td>63</td></tr><tr><td>std</td><td>763</td><td>39.0</td><td>18.3</td><td>20.6</td><td>5.4</td><td>6.1</td><td>58</td></tr><tr><td>Mo</td><td>thin</td><td>859</td><td>64.7</td><td>18.9</td><td>22.4</td><td>10.5</td><td>12.4</td><td>-</td></tr><tr><td rowspan="3">20</td><td rowspan="3">0.51</td><td rowspan="3">0.92</td><td rowspan="3">1.08</td><td rowspan="3">1.49</td><td rowspan="3">IOW</td><td>thin</td><td>537</td><td>15.8</td><td>9.4</td><td>18.6</td><td>0.8</td><td>1.6*</td><td>86*</td></tr><tr><td>std</td><td>540</td><td>13.3</td><td>8.3</td><td>17.2</td><td>0.6</td><td>1.2*</td><td>89*</td></tr><tr><td>Mo</td><td>thin</td><td>817</td><td>58.7</td><td>18.6</td><td>22.4</td><td>8.9</td><td>10.7</td><td>-</td></tr><tr><td></td><td></td><td></td><td></td><td></td><td></td><td>std</td><td>652</td><td>63.6</td><td>16.9</td><td>20.9</td><td>7.0</td><td>8.7</td><td>-</td></tr></tbody></table>
<figure><img src="image_3.png" /><figcaption>**FIGURE 1** | JV characteristics for the best cells with Mo (a) and IOW (b) back contacts for varying RbF-PDT doses (given as RbF thickness) and with the thin (35 nm) CdS buffer layer. (c,d) Corresponding external quantum efficiency spectra. The reflection loss (1-R), as measured on the sample with Mo and 10 nm RbF, is indicated by the dashed blue curve. All corresponding results for samples with the standard-thick (65 nm) CdS layer are shown in Figure S1, supporting information.</figcaption></figure>

catalyzes the $GaO_x$ formation or that the greater thermal budget by the PDT step leads to thicker $GaO_x$. However, the latter point is deemed unlikely, since the RbF-PDT is done at a temperature of 350°C (duration of ≈14 min for 20 nm RbF), while $GaO_x$ was found to form only for temperatures ≥390°C [78, 79]. Generally, it cannot be excluded that the total alkali concentration is just too high after HA-PDT when an IOW back contact is used, since it allows for only very minor alkali diffusion from the absorber back into the SLG (in stark contrast to Mo). To clarify some of these questions, the Rb distribution and interface modifications after RbF-PDT are illustrated and discussed later in Section 2.2.

Only in the case of 20 nm RbF and for an IOW back contact does the PDT-induced photocurrent blocking also result in a decreased $J_{SC}$, expressed by a constant reduction of the EQE level (see Figure 1d). For all other samples, the EQE is not significantly affected by the RbF-PDT or the choice of the back contact. The reflectance (R) spectra for the samples with 10 nm RbF and Mo are added in Figure 1c,d and S1(c,d) for full devices with thin and standard CdS layers, respectively. Very good carrier collection is observed for all samples, facilitated by the electrical field in the fully depleted absorber films (see capacitance-voltage (CV) measurements in Figure S2 in supporting information). Remaining losses can be mainly attributed to parasitic absorption in the ZnO:Al front electrode and, to a lesser extent, to transmission losses by incomplete absorption in the AIGS film as well as absorption in ordered vacancy compounds (OVCs) at the absorber surface [80]. When comparing the EQE spectra of samples with thin and standard CdS (Figure 1c,d vs Figure S1(c,d)) it may appear as if the collection in the NIR region is reduced in the latter case. However, the differences in EQE in the NIR region are mainly

arising from differences in reflection losses (1-R) and R-corrected EQE spectra, as presented in Figure S3 in supporting information for the sample with Mo and 10 nm RbF, only deviate in the spectral region where CdS absorbs.

Figure 2 illustrates the trends in JV parameters with increasing RbF dose for the best cells of each sample (values in Table 1). It is obvious that $V_{OC}$ is not improved by the RbF-PDT for any of the devices. In the case of a Mo back contact, it stays at the reference level up to 10 nm RbF and is then reduced for the highest dose. When IOW is used, $V_{OC}$ deteriorates already for the smallest RbF dose and then monotonously decreases further. We speculate that the loss in $V_{OC}$ and the creation of the transport barrier share the same origin. The trends in FF are similar to the ones in $V_{OC}$, while the situation for $J_{SC,qe}$ was already discussed earlier. Overall, this leads to a monotonous decline in efficiency with increasing RbF dose for all samples. The efficiency of the best cells with an IOW back contact decline from ≈12% to ≈2% when applying the most intense RbF-PDT. Clearly, a RbF-PDT is not suitable for AIGS solar cells, in contrast to Cu-containing ACIGS devices with a similar bandgap [35]. This indicates that the presence of Cu in the absorber plays a crucial role for the effectiveness of an HA-PDT. However, the results also suggest that an HA-PDT may generally be detrimental when a TBC is used in combination with a wide-gap absorber, which needs to be confirmed in future works. In contrast, for low-gap CIGS solar cells grown at temperatures ≤400°C (i.e., forming very little to no $GaO_x$) with an $In_2O_3:Sn$ TBC, the application of a RbF-PDT was shown to provide very good performance [78].

Notably, thinning down the CdS from our standard 65 nm to 35 nm did not deteriorate any JV parameter. On the contrary,
<figure><img src="image_4.png" /><figcaption>**FIGURE 2** | Trends in (a) open-circuit voltage, (b) short-circuit current density, (c) fill factor, and (d) efficiency with increasing RbF-PDT dose (given as RbF thickness), respectively. The results are shown for the best cells with Mo and IOW back contacts and for thin (35 nm) and standard (65 nm) CdS buffer layer thicknesses.</figcaption></figure>

$V_{OC}$ and FF rather improved slightly with the thinner CdS buffer, while $J_{SC}$ even increased by about 1–2 mA/cm², due to the reduced parasitic absorption. In a previous works, we found that our CdS deposition process leads to closed films on low-gap CIGS absorbers for average layer thicknesses ≥30 nm, leading to no more $V_{OC}$, but a continuous FF increase for thicker films up to the standard thickness [81]. At this point, we cannot explain why the thinner buffer layer is clearly advantageous for the performance of the AIGS solar cells analyzed in this study. In the future, we plan to further reduce the CdS thickness to identify the optimum configuration.

Overall, the performance level of the AIGS cells with a Mo back contact and without HA-PDT is similar to samples which include Cu and have the same bandgap [41, 48, 49]. Table 3 compares the best cells from this study (AAC = 1, GGI = 0.5) with the best Cu-containing cell from our earlier work (AAC = 0.8, GGI = 0.6) [49]. For the standard CdS thickness, the sample with Cu shows a higher FF, but very comparable $V_{OC}$ and $J_{SC}$ values. The reason for the lower FF for the AIGS sample in this study may be the more

pronounced CBO at the CdS interface (≈0.35 eV vs ≈0.25 eV [42]). In summary, reducing the Ga content from GGI = 0.6 to GGI = 0.5 for a fixed bandgap of 1.46–1.47 eV by increasing the AAC to 1 did not result in an improved cell performance. This questions the hypothesis that the deep Ga₁ defect is responsible for the large $V_{OC}$ deficit of wide-gap ACIGS solar cells [55–61].

Light soaking is often seen to improve the performance of (HA)F-treated ACIGS solar cells [1, 82, 83]. However, none of the samples in this work (with or without PDT) improved by light soaking, but instead even showed $V_{OC}$ degradation after 4 weeks of storage in ambient conditions. It is therefore advised to limit the AAC value to ≤0.5 and accept a slightly higher GGI value for a given bandgap, which was also shown to minimize metastability issues [70, 71]. On the contrary, the thermodynamical stability of the absorber (i.e., also of the final cell) is predicted to decrease with increasing Ga content and AAC values close to 0.5 [84]. Thus, an AAC ≈ 0.4 may be optimal, still allowing for a sufficiently wide space charge region for close-stoichiometric absorbers, while being more stable and potentially gaining from a RbF-PDT.

<table><caption>**TABLE 3** | Absorber properties and JV parameters of the best Cu-free AIGS cells with Mo back contacts from this work as compared to the corresponding parameters of the best Cu-containing ACIGS cell from our previous work [49]. Compositional depth grading and device structure are the same for all devices and for none of them a HA-PDT was applied.</caption><thead><tr><th>Sample</th><th>AAC</th><th>GGI</th><th>d<sub>abs</sub><br>(μm)</th><th>CdS</th><th>E<sub>G</sub><br>(eV)</th><th>V<sub>OC</sub><br>(mV)</th><th>FF<br>(%)</th><th>J<sub>SC,qe</sub><br>(mA/cm<sup>2</sup>)</th><th>η<sub>qe</sub><br>(%)</th></tr></thead><tbody><tr><td>This work, No RbF</td><td>1.00</td><td>0.50</td><td>1.05</td><td>Thin</td><td>1.47</td><td>858</td><td>74.4</td><td>23.0</td><td>14.7</td></tr><tr><td>This work, No RbF</td><td>1.00</td><td>0.50</td><td>1.05</td><td>Std</td><td>1.47</td><td>851</td><td>73.1</td><td>20.9</td><td>13.0</td></tr><tr><td>Best cell from [49],<br>No RbF</td><td>0.81</td><td>0.59</td><td>1.00</td><td>Std</td><td>1.46</td><td>854</td><td>77.1</td><td>21.5</td><td>14.2</td></tr></tbody></table>
When using the wide-gap device as a top cell in a tandem configuration, it would be desirable to ‘flip’ the entire structure upside down and illuminate through the rear side (SLG/IOW/ etc.). In this way, a glass sheet could be saved. However, due to the low diffusion length of the wide-gap chalcopyrite material, the carrier collection is much worse at rear as compared to front illumination. As a consequence, the $J_{SC}$ bifaciality factor ($J_{SC,bif} = \frac{J_{\text{rear,SC,abs}}}{J_{\text{front,SC,qe}}}$) was found to be limited to ≈60% for wide-gap ACIGS solar cells with $d_{\text{abs}} \ge 1$ µm [49, 50]. Figure 3a shows the EQE spectra at front and rear illumination for all samples with a TBC from this study (only best cells). The corresponding $J_{SC,bif}$ values can be found in Table 2. Excluding the cell without RbF-PDT and thin CdS, the collection at rear illumination improves with RbF dose. This would indicate an increased diffusion length by the Rb incorporation. However, since the best $J_{SC}$, bif value of 72% is measured for just this sample without RbF and thin CdS, evidence for such a trend cannot be claimed. When comparing Cu-containing (AAC ≈ 0.8) cells from our previous work to the AIGS samples with similar $E_G$ and the same device structure from this work, the bifaciality is roughly on the same level (see Figure 3b). However, the best $J_{SC,bif}$ of 72% from this work is outstanding for an absorber thickness >1 µm. Samples which show strong blocking in JV that reduces the $J_{SC,qe}^{\text{front}}$ value (like for 20 nm RbF-PDT in this study) are excluded in Figure 3b.

## 2.2 | Material Characterization

To understand why the RbF-PDT does not lead to an improvement of any photovoltaic parameter when applied to our wide-gap AIGS absorber, the treatment-induced chemical and structural changes are presented and discussed in this section. In the first part 2.2.1., surface modifications after RbF-PDT

are illustrated and the second part 2.2.2. analyzes the effect of the Rb incorporation underneath the absorber surface (i.e., front interface, bulk, grain boundaries, and back interface).

### 2.2.1 | Absorber Surface Modifications upon RbF-PDT

Figure 4 shows scanning electron microscopy (SEM) images of the AIGS surfaces with 20 nm RbF thickness (Mo back contact), before and after rinsing in deionized water for 140 s. Corresponding images with lower magnification can be found in Figure S4, supporting information.

The water rinse mimics the chemical processes occurring during the initial phase of the CdS-CBD, when water-soluble phases are dissolved in a $\text{NH}_3$ solution before CdS starts growing [12, 85]. It is therefore expected that the morphology and elemental

<figure><img src="image_5.png" /><figcaption>**FIGURE 4** | Scanning electron microscopy (SEM) images of the AIGS surface subjected to 20 nm RbF-PDT before and after 140 s water rinse.</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>**FIGURE 3** | (a) External quantum efficiency spectra at front and rear illumination for all samples (best cells only). The pale and bold lines correspond to the cells with 65 nm and 35 nm thick CdS layers, respectively. (b) $J_{sc}$ bifaciality factors as a function of absorber thickness for Cu-containing ACIGS samples from [50] and for the AIGS samples from this study (excluding the ones with 20 nm RbF). The ACIGS devices from [50] are in the compositional range of AAC = 0.7–0.8 and GGI = 0.5–0.6 with $E_G$ = 1.4–1.5 eV (values of AIGS in Table 2).</figcaption></figure>
distributions of the H₂O-rinsed absorber surfaces are very similar in the final device (i.e., after CdS deposition). Isolated and agglomerated particles as well as bigger patches are observed before the water rinse, similar to what is reported for low-gap CIGS [9, 12, 23, 30] and wide-gap ACIGS [35] surfaces. After all those, obviously water-soluble, particles are removed by the water rinse (or in the CBD solution), a very thin porous layer with openings or trenches is left behind.

To chemically analyze the surface before the water rinse, synchrotron-based X-ray photoemission electron microscopy (XPEEM) measurements were conducted on the sample with 20 nm RbF-PDT (same sample as in Figure 4). This technique allows to probe the chemical species in only the upper ≈2 nm of the sample surface (excitation photon energies = 100–455 eV) and to map the elemental composition with a lateral resolution in the order of tens of nanometers. Figure 5a shows again an SEM image of the sample surface for comparison, while Figure 5b illustrates the XPEEM maps for the absorber elements (Ga3d, In4d, Se3d, Ag3d) and Rb (3d) for a selected region of the measurement (different area as compared to SEM image). The entire scanned area can be found, together with a corresponding low-energy electron microscopy image, in Figure S5 in the supporting information. Due to differences in the photoelectron kinetic energies of 1s lines, fluorine and oxygen could not be measured in the emission mode at the same probing depth. Clearly,

regions of locally increased elemental concentrations are observed, very likely indicating the presence of the particles visible in SEM. Interestingly, the Ga and In signal intensities are anti-correlated, as highlighted in the overlay image in Figure 5b and in the 2-dimensional histogram in Figure 5c. A comparison of the photoemission spectra of photoelectrons originating from a Ga-rich area (region B in Figure 5b) with Ga-poor regions, as presented in Figure S6 (a,b) in the supporting information, reveals that the Ga agglomerates do not contain any Indium at all (In4d peak not detected). In contrast, the regions without Ga-rich particles contain mainly In with only a minor amount of Ga. We suggest that the Ga-enriched areas are either GaF, Rb-Ga-F, or GaOₓ, maybe with contributions of Se. Rubidium can be found everywhere on the surface, suggesting that beside the particles observed in SEM, a (possibly very thin) Rb-compound with strongly reduced Ga content formed, see also Figure S6(c,d) in the supporting information. We propose that these are regions where a thin Rb-In-Se surface layer is exposed, which is supported by complementing analysis later in the manuscript.

The regions with highest Rb content also contain the highest Se concentration (e.g., region A in Figure 5b), leading to a clear correlation in the 2-dimensional histogram in Figure 5d. These regions may point to Rb-Se-X (X = F, Ga, O or nothing) agglomerates at the surface, suggested to also form for RbF-treated low-gap CIGS [86]. Finally, smaller Ag-rich particles are observed

<figure><img src="image_7.png" /><figcaption>**FIGURE 5** | (a) Secondary electron microscopy image of the AIGS absorber surface subjected to 20 nm RbF-PDT (Mo back contact), before water rinse. (b) Corresponding surface-sensitive, elemental intensity maps of the absorber elements and Rb, as deduced from X-ray photoemission electron microscopy measurements (different area as compared to (a)). An overlay of the Ga and In intensities is added, too. (c) Ga versus In intensity, plotted as a pixel density distribution map (2-dimensional histogram). (d) Same 2-dimensional histogram for Rb versus Se.</figcaption></figure>
(e.g., region C in Figure 5b). No correlation with any other detectable element was found, suggesting that either metallic silver, AgF or AgOₓ is present at the surface. An overlay of the Ga, In, Ag, and Se maps in Figure S7 in the supporting information (slightly larger region) confirms that the Ga-, Ag-, and Rb-rich regions are different, spatially separated particles and that the indium signal stems from the regions beside those particles.

To also quantify overall concentration changes of the absorber elements as well as of fluorine and rubidium after the water rinse, conventional X-ray photoelectron spectroscopy (XPS) was performed on the samples shown in Figure 4 and 5. Figure 6 illustrates the corresponding XPS spectra for the AIGS sample with 20 nm RbF-PDT (Mo back contact) before and after water rinse.

Generally, the changes in surface chemistry by the water rinse resemble what we observed before for Cu-containing, wide-gap ACIGS absorbers [35]. The fluorine and sodium signals vanish, the Rb intensity is strongly reduced and a peak at the binding energy (*E*<sub>b</sub>) of ≈21 eV disappears. All other signals, stemming from the chalcopyrite AIGS absorber, are increased after the water rinse. This indicates that before the water rinse the measurement is

dominated by the presence of water-soluble Rb- and F- (and to a minor extent Na-) containing compounds. In addition, the vanishing of the peak at *E*<sub>b</sub> ≈ 21 eV, which can be assigned to a X-Ga-F compound (with X being Rb and/or O, C, H, Na or nothing), suggests that at least some of the water-soluble particles also contain a substantial concentration of Ga. This is in line with the detection of the Ga-enriched and In-free particles in the XPEEM maps before the water rinse (see Figure 5b) and should result in a Ga depletion at the very surface of the absorber in the final device, that is, after CdS deposition. On the first sight, it may appear less coherent that the XPS measurements do not reproduce the Rb-Se correlation observed in the XPEEM maps, which suggest the presence of water-soluble Rb-Se (-X) compounds. In fact, the Se3d XPS-signal is significantly increasing after the water rinse. A possible explanation may be that the non-rinsed surface is mainly covered by Se-free (or Se-poor) particles like GaF and RbF (or corresponding mixed compounds) and less so by Rb-Se (-X) particles. This would indeed lead to the observed increase of the Se (as well as In and Ag) XPS-signal after the water rinse, mainly stemming from the Rb-In-Se surface phase and the underlying (or locally exposed) AIGS absorber. Generally, the differences in probing depth (≈5 nm vs. ≈2 nm) and beam intensity between

<figure><img src="image_8.png" /><figcaption>**FIGURE 6** | Photoemission spectra of selective core level electrons for the absorber elements as well as for Rb, Na and F, as deduced from XPS measurements on the bare AIGS surface subjected to 20 nm RbF-PDT (Mo back contact), before and after 140s water rinse (i.e., same surfaces as shown in Figure 4). The measurement settings were exactly the same for both samples, allowing for a direct comparison of signal intensities.</figcaption></figure>
the in-house XPS and synchrotron-based XPEEM measurements
make a direct comparison of the emitted photoelectron spectra
challenging.

Overall, we suggest that the RbF-PDT results in the formation of water-soluble Rb-, F-, and Ga-containing particles/agglomerates at the surface. After water rinse (or CBD), those particles are gone and only a very thin Rb-containing phase is left on the surface, which is probably depleted in Ga. The SEM image in Figure 4 suggests that this phase either has local openings or varies lat- erally in thickness. It should be mentioned here that the Rb3d XPS-signal is very weak after the water rinse, but the correspond- ing XPEEM measurement (presented in Figure S8 in supporting information), being more surface-sensitive and conducted under higher excitation flux, clearly shows a Rb signal across the entire surface (albeit with significantly reduced intensity as compared to before the water rinse). In line with the XPS analysis, the Ga- rich particles (leading to Ga-In anti-correlation before water rinse) vanished in the XPEEM maps after the water rinse.

## 2.2.2 | Chemical and Structural Modifications underneath the Absorber Surface

Figure 7a,b shows the scanning transmission electron micros-
copy (STEM) cross-section images of the samples with standard
CdS and TBC without and with 20 nm RbF-PDT, respectively.
The dark spots are Ag precipitations, forming during lamella
preparation, as reported before for ACIGS solar cells [35, 44].
No significant structural changes are observed on the micrometer
scale after RbF-PDT. An open question raised in Section 2.1 was
whether the RbF-PDT leads to a thicker GaOₓ film at the AIGS/
IOW interface, potentially explaining the stronger FF deteriora-
tion for cells with IOW as compared to Mo. Indeed, as can be seen
in Figure 7c,d, which highly magnifies the back contact region, a
thicker GaOₓ layer (3.0 nm on average) is found after RbF-PDT as
compared to the sample without PDT (1.7 nm on average). The
latter value is in good agreement with our earlier work on ACIGS
with an IOW TBC [49]. However, a GaOₓ thickness of 3.0 nm is

<figure><img src="image_9.png" /><figcaption>**FIGURE 7** | Scanning transmission electron microscopy (STEM) bright-field images of the entire cross-section for the samples with IOW and standard thick CdS without PDT (a) and subjected to 20 nm RbF-PDT (b). (c,d) Corresponding high-magnification images of the AIGS/GaOₓ/IOW back contact regions.</figcaption></figure>

still thinner as observed in other studies on (A)CIGS devices
with a TBC, which show no significant transport barrier and
high FF values for GaOₓ thicknesses up to at least 4 nm
[76–79]. Thus, it is deemed unlikely that the rather moderate
increase in GaOₓ extension from 1.7 to 3.0 nm alone explains
the strong photocurrent blocking in JV for the RbF-treated sam-
ples in this work.

Energy-dispersive X-ray spectroscopy (EDS) in STEM mode was
performed at the back contact region of the sample without RbF,
see Figure S9 in supporting information. Besides confirming the
presence of the GaOₓ layer, the analysis shows that the grain
boundaries close to the back contact are Ag-depleted and In-
enriched (same is true for majority of analyzed grain boundaries,
independent of their location, not shown here). A depletion of
group-I elements (Ag or Cu) is frequently observed in ACIGS
random, high-angle grain boundaries [18, 35, 87, 88] and claimed
to make them electrically more benign [89]. It was found that
with increasing Ga and for GGI > 0.4, grain boundaries in
CIGS are more likely to be group-I enriched and thereby poten-
tially more detrimental. However, no such Ag-rich grain bound-
aries were identified in this study.

The same STEM-EDS analysis was conducted for the correspond-
ing sample with 20 nm RbF-PDT, as illustrated in Figure 8a,b.
Again, the grain boundaries are Ag-depleted and In-enriched,
but in this case also decorated by rubidium. It was suggested that
the introduction of Rb leads to a slightly more pronounced (or fre-
quent) group-I depletion in CIGS grain boundaries [18, 87, 88],
which cannot be confirmed nor excluded in the present study.
We further find that Rb is also located at the AIGS/GaOₓ/IOW
interface of our samples. Extracted absolute and normalized
([0,1]) EDS line-scans across the back interface, presented in
Figure 8c,d, suggest that Rb may even be present inside the very
thin GaOₓ layer, potentially modifying its properties. Even the for-
mation of a Rb-Ga-O compound, such as RbGa₅O₁₁, may be think-
able [90].

In any case, it can be speculated that the excessive presence of Rb
at the back interface alters the electronic properties of one or sev-
eral of the involved layers (e.g., doping and band alignment) and
creates a back-contact barrier, explaining the strong FF losses
with increasing RbF dose. Even for CIGS solar cells with Mo back
contacts, the excessive replacement of Na by Rb atoms was sug-
gested to create a back-contact barrier [26, 73]. Indeed, we also
observe such an ion-exchange mechanism, resulting in a decreas-
ing Na content in the absorber bulk and at its interfaces with
increasing RbF dose (see alkali depth profiles in Figure S10, sup-
porting information). We conclude that an excessive concentra-
tion of Rb at the back-contact interface may be even more
detrimental when IOW is used instead of Mo.

From the EDS maps and the STEM image in Figure 8, it seems as
if GaOₓ locally grows spherically into the absorber. However, a
closer look reveals that those bright spots in the STEM image are
small voids at the interface, with GaOₓ at their edges and presum-
ably form when the NaF precursor layer was locally too thick, as
occasionally observed before when using a TBC [76, 77]. More
statistics are needed to determine whether such voids are more
frequent when a RbF-PDT is applied.
<figure><img src="image_10.png" /><figcaption>**FIGURE 8** | (a) STEM bright-field image of the back-contact region for the sample with IOW and standard 65 nm thick CdS, subjected to 20 nm RbF-PDT. (b) Corresponding elemental distribution maps, as deduced from energy-dispersive X-ray spectroscopy. (c,d) Extracted elemental line scan across the back interface(s) at position “LS3” in (b) in absolute and normalized ([0,1]) atomic concentrations, respectively. In (d) the normalized profiles are shown in a narrower region as compared to (c) to highlight the changes across the interface.</figcaption></figure>

Finally, RbF-PDT-induced modifications of the heterojunction are presented. Figure 9a shows a STEM image of the AIGS/CdS/ZnO:Al interface region of the sample with 20 nm

RbF-PDT and the standard 65 nm CdS thickness with an IOW back contact. The corresponding elemental distribution maps, as measured by EDS, are presented in Figure 9b. Clearly, also

<figure><img src="image_11.png" /><figcaption>**FIGURE 9** | (a) STEM bright-field image of the heterojunction region for the sample with IOW and standard 65 nm thick CdS, subjected to 20 nm RbF-PDT. (b) Corresponding elemental distribution maps, as deduced from energy-dispersive X-ray spectroscopy. (c,d) Extracted elemental line scans across the heterojunction at position “LS1” in (a) and across a grain boundary at position “LS2” in (a), respectively.</figcaption></figure>
the grain boundaries at the front are strongly depleted in Ag, as well as In- and Rb-enriched. Rubidium is also present at the AIGS/CdS interface. This is true for the majority of the analyzed areas/positions (see also Figure S11, supporting information, for full cross-section), pointing to an (almost) complete Rb-coverage of the AIGS surface after the PDT. This is in line with the XPEEM results, showing Rb across the entire surface, even after H₂O rinse. To detect chemical changes across the ACIGS/CdS junction, an EDS line scan was extracted from the maps at the position indicated as 'LS1' in the STEM image. The results are shown in Figure 9c. Two notable features are detected. Firstly, a thin (<10 nm) Rb-rich film is present between the AIGS chalcopyrite absorber and the CdS buffer layer. This film mainly consists of In, Se and Rb and maybe small amounts of Ga, but is fully depleted in Ag. This strongly suggests the formation of a RbInSe₂ layer (maybe with traces of Ga) at the AIGS surface after RbF-PDT. Alternatively, a Rb-In-Se layer with a stoichiometry different from 1:1:2, such as 1:3:5 (RbIn₃Se₅) or 1:5:8 (RbIn₅Se₈), as observed for OVCs, or even Rb-enriched In₂Se₃ may be thinkable. However, at the position of the Rb peak, the stoichiometry extracted from EDS already leads to a RbIn₂.₇Se₅.₇ compound. Considering that the layer is very thin and no Rb plateau is reached, it can be assumed that the real Rb content is significantly higher in the film. Thus, we propose that a RbInSe₂ film is most likely.

Overlays of the STEM image with elemental maps of Cd, Ag, and In or Rb are provided in Figure 10a,b, respectively and separately also in Figure S12, supporting information. Additionally, another, higher magnified, region of the heterojunction was analyzed, which is shown in Figure S13, supporting information. From these figures it becomes clear that the assumed RbInSe₂ phase is not always as 'thick' as at the position of 'LS1', but most often <5 nm. We believe that this corresponds to the porous surface structure observed in SEM after water rinse (Figure 4). Still, it is surprising that a RbInSe₂ phase of several Nanometer thickness does not result in a stronger Rb peak after the water rinse in XPS analysis, while Rb is clearly visible in the corresponding XPEEM map.

Secondly, Cd seems to diffuse into the upper few nanometers of the assumed RbInSe₂ film, as suggested by the ≈5 nm shift of the Cd onset as compared to the sulfur onset in Figure 9c. Both features, the formation of a RbInSe₂ layer and the Cd in-diffusion, are perfectly in line with earlier observations obtained from

<figure><img src="image_12.png" /><figcaption>**FIGURE 10** | Overlays of the STEM image from Figure 9 (20 nm RbF-PDT, standard-thick CdS, IOW back contact) with the elemental distributions of Cd, Ag, and In (a) or Rb (b). The pixel opacity of the elemental maps was set linear to the respective normalized atomic concentrations.</figcaption></figure>

atom-probe tomography on low-gap CIGS absorbers subjected to RbF-PDT (even the minor Ga incorporation into RbInSe₂) [87].

Another EDS line scan was extracted across a grain boundary at position 'LS2' (indicated in the STEM image). The results in Figure 9d confirm the Ag-depletion and In- and Rb-enrichment in the grain boundary, while Se and Ga do not change.

Overall, our results suggest that even when Cu is entirely replaced by Ag in the chalcopyrite absorber, as it is the case in this study, a very thin RbInSe₂ surface layer forms upon RbF-PDT. All other RbF-PDT-induced features, like the grain boundary chemistry and the overall Rb distribution are also similar for CIGS and AIGS solar cells.

Coming back to the solar cell performances presented and discussed in part 2.1, it may be speculated that the continuously reduced *FF* with increasing RbF dose arises from a thicker or more covering RbInSe₂ layer. The CBO at the RbInSe₂/AIGS interface is expected to be much larger as compared to the RbInSe₂/CIGS interface [8, 39, 40, 42, 72], potentially explaining why an optimized RbF-PDT can result in *FF* improvements for CIGS, while the *FF* only deteriorates for AIGS. In addition, we propose that the presence of excessive Rb at the back contact is more harmful when a TBC is used, presumably by modifying the GaOₓ properties and catalyzing its growth (i.e., making it thicker), explaining why those samples show a larger *FF* loss than the ones with Mo.

However, the question remains why no beneficial effect on $V_{OC}$ is observed by the RbF treatment of AIGS absorbers, in contrast to low- and wide-gap CIGS solar cells, as long as sufficient indium is present. The only difference in this work is that Cu is completely absent, indicating its crucial role for a $V_{OC}$-boost after RbF-PDT. A possible reason may be that the RbF-PDT did not lead to an increase in doping density, as reported for CIGS [23–25]. All samples show more or less complete depletion (a typical feature of high-Ag ACIGS absorbers) independent of the treatment, as exemplified for 0 and 20 nm RbF-PDT in Figure S2 in the supporting information. The physical origin may lie in the much larger bulk solubility of Na and Rb for AIGS as compared to CIGS absorbers, which determines the final concentration of group-I vacancies [91].

### 3 | Conclusions

The effect of an RbF-PDT on the performance of Cu-free, wide-gap ($E_G$ = 1.5 eV) AIGS solar cells with either opaque Mo or transparent In₂O₃:W back contacts is studied. We find that the *FF* monotonously deteriorates with increasing RbF dose for all samples without observing any gain in $V_{OC}$ or $J_{SC}$. The *FF* loss is much more pronounced when In₂O₃:W is used, probably due to a more detrimental effect of the presence of Rb at the AIGS/GaOₓ/IOW back interface, where it replaces sodium. Rubidium may even diffuse into the GaOₓ and negatively impact its electrical properties. A facilitated GaOₓ formation by the RbF-PDT is found, too. However, even for the maximum RbF dose, the GaOₓ thickness only increases from ≈1.7 nm (no PDT) to ≈3.0 nm (20 nm RbF), which is deemed unlikely to explain
the severe FF losses alone, considering the results from earlier
publications.

Advanced material characterization reveals that the RbF-PDT-
induced absorber modifications are very similar for AIGS and
CIGS absorbers: The formation of a Rb-In-Se surface phase
(likely RbInSe₂) is detected, as well as a weak Cd in-diffusion into
the same. It is suggested that the large conduction band offset at
the potential RbInSe₂/AIGS interface causes an additional loss in
fill factor. Otherwise, Rb is mainly located in group-I depleted
and In-enriched grain boundaries and at the back interface.
However, most grain boundaries are also found to be Ag-depleted
and In-enriched without RbF-PDT.

We further propose that the desired V<sub>OC</sub> boost is absent for AIGS solar cells, since the Rb incorporation does not lead to a doping increase for close-stoichiometric Cu-free absorbers, which remain fully depleted, independent of a PDT step. The possible explanation may be the higher alkali solubility in the chalcopyrite lattice. A reduced interface recombination by the RbF-PDT was not expected anyhow, due to a favorable band alignment of the untreated AIGS absorber with the CdS buffer layer. Overall, this work shows that a RbF-PDT is not suitable for wide-gap AIGS solar cells, especially when a transparent back contact is used.

Lastly, the strategy to improve the minority carrier lifetime,
by entirely replacing Cu by Ag to reduce the Ga content
(GGI = 0.5 in this work) for a given bandgap, failed. As compared
to Cu-containing ACIGS devices with the same E<sub>G</sub>, the AIGS
solar cells from this study show similar V<sub>OC</sub> and J<sub>SC</sub>, but lower
FF values. This finding indicates that the high Ga content itself is
not the main reason for the large V<sub>OC</sub> deficit of wide-gap ACIGS
solar cells. However, while no beneficial effect is observed at
front illumination, the collection at rear illumination may be
slightly improved as compared to ACIGS devices with same
bandgap, reaching up to 72% J<sub>SC</sub> bifaciality for the best cell with
an IOW back contact (η = 11.6% at front illumination).

In summary, removing Cu from ACIGS absorbers altogether does
not seem to be a viable route to enhance the efficiency of corre-
sponding wide-gap solar cells for use as a top cell in tandem devi-
ces. Additionally applying a RbF-PDT makes the situation even
worse.

4 | Experimental Section

**4.1 | Solar cell processing**

A stack of SLG/(Mo or IOW)/NaF/AIGS/CdS/ZnO:Al was used for the solar cell samples in this work. The opaque Mo back contact (320 nm, R<sub>sheet</sub> ≈ 0.7 Ω sq<sup>-1</sup>) was direct-current sputtered, while IOW (300-380 nm, R<sub>sheet</sub> ≈ 10 Ω sq<sup>-1</sup>) was radio-frequency sputtered from an In<sub>2</sub>O<sub>3</sub>-WO<sub>3</sub> compound target (nominally 1 wt.% WO<sub>3</sub>) in an Ar (40 sccm) and O<sub>2</sub> (0.4 sccm) gas mixture. Sodium in-diffusion from the SLG was allowed (no alkali diffusion-barrier), but additionally 7 nm- (for IOW) or 15 nm-thick (for Mo) NaF precursor layers were coated on top of the back electrodes via thermal evaporation. Next, a

three-stage (Ag-poor → Ag-rich → Ag-poor) co-evaporation
process was applied to grow ≈1.1 µm thick AIGS films at a
maximum substrate temperature of T = 550°C (≈3 min at
T > 500°C). Higher Ga and lower In rates in the initial growth
stages led to the formation of a bandgap grading toward the
back contact (see also Figure S10 in the supporting informa-
tion). For samples subjected to a RbF-PDT, the metal sources
were then switched off and RbF was deposited at T ≈ 350°C
for different times at a rate of ≈1 nm/min in Se-containing
atmosphere. In total, 4 AIGS depositions with different RbF
doses (0, 5, 10, and 20 nm) were done (see integral compositions
in Table 2), each containing SLG/Mo and SLG/IOW samples.
Subsequently, the samples were cut in two halves, which were
then coated with a CdS buffer layer of either the standard
(≈65 nm) or a reduced (≈35 nm) thickness in a CBD process
at 60°C for 495 s or 330 s, respectively. Slight deviations in
CdS thickness (±5 nm) between individual depositions cannot
be ruled out, potentially leading to minor variations in the
EQE spectra at short wavelengths. Finally, a ZnO:Al (270 nm,
R<sub>sheet</sub> ≈ 50 Ω/sq) window layer was sputtered on top. In the final
step, the samples with Mo were sectioned into 14 individual
solar cells (area = 0.05 cm<sup>2</sup>) by mechanical scribing, while for
samples with IOW 16 individual solar cells (area = 0.1 cm<sup>2</sup>)
were created via local, selective removal of the buffer and win-
dow layers (etching in HCl), allowed by a photolithography
masking process.

**4.2 | Material characterization**

A Zeiss Merlin SEM was used to illustrate the absorber surfaces subjected to a RbF-PDT before and after rinsing it in deionized water for 140 s. STEM and EDX analyses were performed on a FEI Titan Themis 200 XFEG instrument equipped with a SuperX detector G2 and operated at 200 kV. The TEM lamellae were prepared via a Ga-based focused ion beam in a Crossbeam 550 Zeiss system, following a standard lift-out technique. A final milling step at maximum 2 kV ion accelerating voltage was performed on both sides of the lamellae, and no further electron exposure was done during focused ion beam preparation before the TEM analyses.

To laterally map the chemical structure of the AIGS surface sub-
jected to 20 nm RbF-PDT, X-ray photoemission electron micros-
copy and low-energy electron microscopy measurements were
performed at the MAXPEEM synchrotron beamline at MAX
IV [92] on samples before and after water rinse. The XPEEM
mapping images were acquired with photon energies between
100 and 455 eV, leading to a constant kinetic energy of excited
photoelectrons of ≈85 eV. Representative high-resolution core
level spectra of absorber elements Ga3d, In4d, Se3d, and Rb3d,
were extracted and used for quantitative analysis. The as-
deposited samples were packed in a N₂ atmosphere to minimize
contamination during transport to the beamline and the meas-
urements were carried out at room temperature and under vac-
uum conditions. A brief thermal annealing step (5 min at 100°C)
was conducted prior to XPEEM measurements, to remove ad-
sorbed H₂O from air exposure during sample mounting.

Additionally, the surface chemistry was also measured in an in-
house Quantera II Scanning XPS Microprobe system (Physical
Electronics) on the same samples, using monochromatic Al Kα (1.487 keV) X-ray radiation and charge compensation. In this way, also the fluorine signal could be measured.

Integral absorber compositions were determined with an XRF spectrometer Epsilon 5 (Panalytical). Glow-discharge optical emission spectroscopy was conducted in a Spectruma Analytik GDA 750HR system to measure the alkali and GGI depth-profiles of the complete solar cell stacks with Mo back contacts. While no quantification was possible for Na and Rb, the GGI was quantified by cross-calibration to the XRF-deduced compositions. The sheet resistance of the electrodes was measured in a linear 4-point probe configuration.

## 4.3 | Electro-optical characterization of solar cells

The EQE (only best cell each) and JV (all cells per sample) measurements of completed solar cells were conducted in home-built setups. To avoid any reflection from the stage surfaces, a highly absorbing, black sponge was placed underneath the solar cells with IOW. The JV characteristics were measured in forward V-sweep at a temperature of 25°C under illumination by a class AAA solar simulator from G2V (Pico), using light emitting diode (LED) arrays. The light intensity was adjusted to match the $J_{SC}$ value at 1 sun irradiance of a calibrated Si reference solar cell. Capacitance-voltage profiling was conducted from $V = -0.4$ to $+0.3$ V at 60 kHz and an amplitude of 25 mV, using an Agilent 4284A precision LCR meter and a Keithley 2401 source meter. A dielectric constant of $\epsilon_r = 10$ was assumed for the AIGS material in this study.

### Acknowledgments

This work was financially supported by the European union programme HORIZON (call: HORIZON-CL5-2021-D3-02, project ID: 101075626 (SITA)), Swedish Research Council (Dnr. 2020-04065) and the StandUp for Energy programme. We acknowledge the MAX IV Laboratory for time on the MAXPEEM beamline under proposal no 20250030. We would like to thank Dr. Alexei Zakharov and Dr. Yuran Niu for assistance during the XPEEM measurements. Research conducted at MAX IV, a Swedish national user facility, is supported by the Vetenskapsrådet (Swedish Research Council, VR) under contract 2018-07152, VINNOVA (Swedish Governmental Agency for Innovation Systems) under contract 2018-04969 and Formas under contract 2019-02496. The authors further acknowledge Myfab Uppsala for providing facilities and experimental support. Myfab is funded by the Swedish Research Council (2020-00207) as a national research infrastructure.

### Conflicts of Interest

The authors declare no conflicts of interest.

### Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

### References

1. J. Keller, K. Kiselman, O. Donzel-Gargand, et al., *Nature Energy* 9 (2024): 46.

2. M. Nakamura, K. Yamaguchi, Y. Kimoto, Y. Yasaki, T. Kato, and H. Sugimoto, *IEEE Journal of Photovoltaics* 9 (2019): 1863.

3. P. Jackson, R. Wuerz, D. Hariskos, E. Lotter, W. Witte, and M. Powalla, *Physics Status Solidi RRL* 586 (2016): 583.

4. A. Chirila, P. Reinhard, F. Pianezzi, et al., *Nature Materials* 12 (2013): 1107.

5. P. Reinhard, B. Bissig, F. Pianezzi, et al., *Chemistry of Materials : A Publication of the American Chemical Society* 27 (2015): 5755.

6. T. M. Friedlmeier, P. Jackson, D. Kreikemeyer-Lorenzo, et al., *IEEE 43rd Photovoltaic Specialists Conference (PVSC)* (IEEE, 2016), 0457–0461.

7. T. M. Friedlmeier, P. Jackson, A. Bauer, et al., *IEEE Journal of Photovoltaics* 5 (2015): 1487.

8. T. Kodalle, T. Bertram, R. Schlatmann, and C. A. Kaufmann, *IEEE Journal of Photovoltaics* 9 (2019): 1839.

9. T. Kodalle, M. D. Heinemann, D. Greiner, et al., *Solar RRL* 2 (2018): 1800156.

10. P. Reinhard, F. Pianezzi, B. Bissig, et al., *IEEE Journal of Photovoltaics* 5 (2015): 656.

11. T. Feurer, F. Fu, T. P. Weiss, et al., *Thin Solid Films* 670 (2019): 34.

12. P. Reinhard, B. Bissig, F. Pianezzi, et al., *Nano Letters* 15 (2015): 3334.

13. E. Handick, P. Reinhard, R. G. Wilks, et al., *ACS Applied Materials & Interfaces* 9 (2017): 3581.

14. E. Handick, P. Reinhard, J. H. Alsmeier, et al., *ACS Applied Materials & Interfaces* 7 (2015): 27414.

15. N. Taguchi, S. Tanaka, and S. Ishizuka, *Applied Physics Letters* 113 (2018): 113903.

16. S. Ishizuka, N. Taguchi, and P. J. Fons, *Journal of Physics Chemistry C* 123 (2019): 17757.

17. C. P. Muzzillo, J. D. Poplawsky, H. M. Tong, W. Guo, and T. Anderson, *Progress in Photovoltaics: Research and Applications*, 26 (2018): 825.

18. S. Siebentritt, E. Avancini, M. Bär, et al., *Advanced Energy Materials* 10 (2020): 1903752.

19. N. Nicoara, R. Manaligod, P. Jackson, et al., *Nature Communications* 10 (2019): 3980.

20. N. Nicoara, T. Lepetit, L. Arzel, S. Harel, N. Barreau, and S. Sadewasser, *Scientific Reports* 7 (2017): 41361.

21. C. P. Muzzillo, *Solar Energy Materials and Solar Cells* 172 (2017): 18.

22. S. Karki, P. K. Paul, G. Rajan, et al., *IEEE Journal of Photovoltaics* 7 (2017): 665.

23. I. Khatri, H. Fukai, H. Yamaguchi, M. Sugiyama, and T. Nakada, *Solar Energy Materials and Solar Cells* 155 (2016): 280.

24. A. Laemmle, R. Wuerz, and M. Powalla, *Physcis Status Solidi RRL* 7 (2013): 631.

25. O. Ramírez, M. Bertrand, A. Debot, et al., *Solar RRL* 5 (2021): 2000727.

26. F. Pianezzi, P. Reinhard, A. Chirilä, et al., *Physical Chemistry Chemical Physics* 16 (2014): 8843.

27. S. A. Jensen, S. Glynn, A. Kanevce, et al., *Journal of Applied Physics* 120 (2016): 063106.

28. A. S. Brown and M. A. Green, *Physica E* 14 (2002): 96.

29. R. K. Kothandaraman, Y. Jiang, T. Feurer, A. N. Tiwari, and F. Fu, *Small Methods* 4 (2020): 2000395.

30. S. Ishizuka, N. Taguchi, J. Nishinaga, Y. Kamikawa, S. Tanaka, and H. Shibata, *Journal of Physical Chemistry C* 122 (2018): 3809.

31. S. Ishizuka, N. Taguchi, J. Nishinaga, Y. Kamikawa, and H. Shibata, *Solar Energy* 211 (2020): 1092.
32. S. Ishizuka, R. Okamoto, and S. Ikeda, *Advanced Materials Interfaces* **9** (2022): 2201266.

33. S. Zahedi-Azad, M. Maiberg, and R. Scheer, *Progress in Photovoltaics: Research and Applications* **28** (2020): 1146.

34. S. Zahedi-Azad, M. Maiberg, R. Clausing, and R. Scheer, *Thin Solid Films* **669** (2019): 629.

35. J. Keller, H. Aboulfadl, L. Stolt, O. Donzel-Gargand, and M. Edoff, *Solar RRL* **6** (2022): 2200044.

36. L. Both, D. Hauschild, M. Blankenship, et al., *Applied Physics Letters* **126** (2025): 021603.

37. J. Keller, F. Chalvet, J. Joel, et al., *Progress in Photovoltaics: Research and Applications* **26** (2018): 13.

38. J. Keller, O. V. Bilousov, J. Neerken, et al., *Solar RRL* **4** (2020): 2000248.

39. T. P. Weiss, S. Nishiwaki, B. Bissig, et al., *Advanced Materials Interfaces* **5** (2018): 1701007.

40. T. Y. Lin, I. Khatri, J. Matsuura, et al., *Nano Energy* **68** (2020): 104299.

41. J. Keller, P. Pearson, N. Shariati Nilsson, O. Stolt, L. Stolt, and M. Edoff, *Solar RRL* **5** (2021): 2100403.

42. J. Keller, K. V. Sopiha, O. Stolt, et al., *Progress in Photovoltaics: Research and Applications* **28** (2020): 237.

43. G. M. Hanket, J. H. Boyle, and W. N. Shafarman, *34th IEEE Photovoltaic Specialists Conference (IEEE, 2009)*, 001240–001245.

44. J. Keller, L. Stolt, K. V. Sopiha, J. K. Larsen, L. Riekehr, and M. Edoff, *Solar RRL* **4** (2020): 2000508.

45. W. Shafarman, C. Thompson, J. Boyle, et al., *35th IEEE Photovoltaic Specialists Conference (IEEE, 2010)*, 000325–000329.

46. K. Kim, S. K. Ahn, J. H. Choi, et al., *Nano Energy* **48** (2018): 345.

47. J. Keller, L. Stolt, T. Törndahl, and M. Edoff, *Solar RRL* **7** (2023): 2300208.

48. J. Keller, L. Stolt, O. Donzel-Gargand, T. Kubart, and M. Edoff, *Solar RRL* **6** (2022): 2200401.

49. J. Keller, L. Stolt, O. Donzel-Gargand, A. F. Violas, T. Kubart, and M. Edoff, *Solar RRL* **8** (2024): 2400430.

50. J. Keller, S. Mudgal, C. Hägglund, K. Kiselman, and M. Edoff, *Solar RRL* **9** (2025): 202500101.

51. J. H. Boyle, B. E. McCandless, W. N. Shafarman, and R. W. Birkmire, *Journal of Applied Physics* **115** (2014): 223504.

52. M. A. Contreras, L. M. Mansfield, B. Egaas, et al., *Progress in Photovoltaics: Research and Applications* **20** (2012): 843.

53. S. Siebentritt and U. Rau, Wide-Gap Chalcopyrites (Springer-Verlag Berlin Heidelberg, 2006).

54. S. Hegedus and W. N. Shafarman, *Progress in Photovoltaics: Research and Applications* **12** (2004): 155.

55. B. Huang, S. Chen, H. Deng, et al., *IEEE Journal of Photovoltaics* **4** (2014): 477.

56. J. Pohl and K. Albe, *Physical Review B* **87** (2013): 245203.

57. C. Spindler, D. Regesch, and S. Siebentritt, *Applied Physics Letters* **109** (2016): 032105.

58. C. Spindler, F. Babbe, M. H. Wolter, et al., *Physical Review Materials* **3** (2019): 090302.

59. S. Lany and A. Zunger, *Journal of Applied Physics* **100** (2006): 113725.

60. G. Hanna, A. Jasenek, U. Rau, and H. W. Schock, *Physics State Solid (a)* **179** (2000): R7.

61. G. Hanna, A. Jasenek, U. Rau, and H. W. Schock, *Thin Solid Films* **387** (2001): 71.

62. M. R. Balboul, H. W. Schock, S. A. Fayak, A. A. El-Aal, J. H. Werner, and A. A. Ramadan, *Applied Physics A* **92** (2008): 557.

63. M. Raghuwanshi, E. Cadel, P. Pareige, et al., *Applied Physics Letters* **105** (2014): 013902.

64. M. Gloeckler and J. R. Sites, *Thin Solid Films* **480-481** (2005): 241.

65. T. Nakada, K. Yamada, R. Arai, H. Ishizaki, and N. Yamada, *MRS Online Proceedings Library* **865** (2004): 111.

66. T. Umehara, F. Zulkifly, K. Nakada, and A. Yamada, *Japanese Journal of Applied Physics* **56** (2017): 08MC09.

67. J. K. Larsen, O. Donzel-Gargand, K. V. Sopiha, et al., *ACS Applied Energy Materials* **4** (2021): 1805.

68. K. Yamada, N. Hoshino, and T. Nakada, *Science and Technology of Advanced Materials* **7** (2006): 42.

69. Y. S. Murthy and P. J. Reddy, *Vacuum* **41** (1990): 1448.

70. P. Pearson, J. Keller, L. Stolt, O. Donzel-Gargand, and C. Platzer Björkman, *Solar RRL* **8** (2024): 2400220.

71. P. Pearson, J. Keller, L. Stolt, and C. Platzer Björkman, *Physics Status Solidi B* **260** (2023): 2300170.

72. G. Sozzi, S. D. I. Napoli, M. Enna, and R. Menozzi, *46th IEEE Photovoltaic Specialists Conference (IEEE, 2019)*, 953–956.

73. J.-H. Yoon, J.-H. Kim, W. M. Kim, et al., *Progress in Photovoltaics: Research and Applications* **22** (2014): 90.

74. T. Nakada, Y. Hirabayashi, T. Tokado, D. Ohmori, and T. Mise, *Solar Energy* **77** (2004): 739.

75. M. D. Heinemann, V. Efimova, R. Klenk, et al., *Progress in Photovoltaics: Research and Applications* **23** (2015): 1228.

76. J. Keller, N. Shariati Nilsson, A. Aijaz, et al., *Progress in Photovoltaics: Research and Applications* **26** (2018): 159.

77. J. Keller, W.-C. Chen, L. Riekehr, T. Kubart, T. Törndahl, and M. Edoff, *Progress in Photovoltaics: Research and Applications* **26** (2018): 846.

78. S. C. Yang, T. Y. Lin, M. Ochoa, et al., *Nature Energy* **8** (2022): 40.

79. A. Ali, D. Jeon, W. Kim, et al., *Advanced Energy Materials* **15** (2025): 2500899.

80. J. Keller, N. Martin, O. Donzel-Gargand, et al., *Solar RRL* **8** (2024): 2301018.

81. J. Keller, F. Gustavsson, L. Stolt, M. Edoff, and T. Törndahl, *Solar Energy Materials and Solar Cells* **159** (2017): 189.

82. R. Carron, S. Nishiwaki, S. Yang, et al., *Preprint available at Research Square* (2022), https://doi.org/10.21203/rs.3.rs-2116168/v1.

83. I. Khatri, K. Shudo, J. Matsuura, M. Sugiyama, and T. Nakada, *Progress in Photovoltaics: Research and Applications* **26** (2018): 171.

84. K. V. Sopiha, J. K. Larsen, O. Donzel-Gargand, et al., *Journal of Materials Chemistry A* **8** (2020): 8740.

85. F. Larsson, O. Donzel-Gargand, J. Keller, M. Edoff, and T. Törndahl, *Solar Energy Materials and Solar Cells* **183** (2018): 8.

86. C. K. Boumenou, H. Phirke, J. Rommelfangen, et al., *Advanced Functional Materials* **33** (2023): 2300590.

87. M. Raghuwanshi, A. Vilalta-Clemente, C. Castro, et al., *Nano Energy* **60** (2019): 103.

88. A. Vilalta-Clemente, M. Raghuwanshi, S. Duguay, et al., *Applied Physics Letters* **112** (2018): 103105.

89. M. Raghuwanshi, R. Wuerz, and O. Cojocaru-Mirédin, *Advanced Functional Materials* **30** (2020): 2001046.

90. C. A. Juillerat, V. Kocevski, V. V. Klepov, J. W. Amoroso, T. M. Besmann, and H. C. zur Loye, *Journal of the American Ceramic Society* **103** (2020): 6531.
91. H. Aboulfadl, K. V. Sopiha, J. Keller, et al., *ACS Applied Materials & Interfaces* 13 (2021): 7188.

92. Y. Niu, N. Vinogradov, A. Preobrajenski, et al., *Journal of Synchrotron Radiation* 30 (2023): 468.

**Supporting Information**

Additional supporting information can be found online in the Supporting Information section.