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Direct Imaging of Dopant Distribution in Polycrystalline ZnO Films Fanni Lorenzo,*,† A. Brian Aebersold,‡ Monica Morales-Masis,† Martin Ledinský,§ Stéphane Escrig,∥ Aliaksei Vetushka,§ Duncan T. L. Alexander,‡ Aïcha Hessler-Wyser,† Antonín Fejfar,§ Cécile Hébert,‡ Sylvain Nicolay,⊥ and Christophe Ballif†,⊥ †

Ecole Polytechnique Fédérale de Lausanne (EPFL), Institute of Microengineering, Photovoltaics and Thin-Film Electronics Laboratory, Rue de la Maladière 71B, CH-2000 Neuchâtel, Switzerland ‡ Ecole Polytechnique Fédérale de Lausanne (EPFL), Interdisciplinary Centre for Electron Microscopy (CIME), Station 12, CH-1015 Lausanne, Switzerland § Laboratory of Nanostructures and Nanomaterials, Institute of Physics ASCR, Cukrovarnická 10, 162 00 Prague 6, Czech Republic ∥ Ecole Polytechnique Fédérale de Lausanne (EPFL), Laboratory for Biological Geochemistry, Station 2, CH-1015 Lausanne, Switzerland ⊥ Centre Suisse d’Electronique et Microtechnique, PV-Center, rue Jacques-Droz 1, CH-2002 Neuchâtel, Switzerland S Supporting Information *

ABSTRACT: Two fundamental requirements of transparent conductive oxides are high conductivity and low optical absorptance, properties strongly dependent on the free-carrier concentration of the film. The freecarrier concentration is usually tuned by the addition of dopant atoms; which are commonly assumed to be uniformly distributed in the films or partially segregated at grain boundaries. Here, the combination of secondary ion mass spectroscopy at the nanometric scale (NanoSIMS) and Kelvin probe force microscopy (KPFM) allows direct imaging of boron-dopant distribution in polycrystalline zinc oxide (ZnO) films. This work demonstrates that the boron atoms have a bimodal spatial distribution within each grain of the ZnO films. NanoSIMS analysis shows that boron atoms are preferentially incorporated into one of the two sides of each ZnO grain. KPFM measurements confirm that boron atoms are electrically active, locally increasing the free-carrier concentration in the film. The proposed cause of this nonuniform dopant distribution is the different sticking coefficient of Zn adatoms on the two distinct surface terminations of the ZnO grains. The higher sticking coefficient of Zn on the c+ surface restricts the boron incorporation on this side of the grains, resulting in preferential boron incorporation on the c− side and causing the bimodal distribution. KEYWORDS: film polarity, polycrystalline film, dopant distribution, grain boundaries, zinc oxide, NanoSIMS activated11 and, furthermore, that grain boundaries act as carrier traps.12,13 Seto compared the dependence of carrier mobility on dopant concentration in mono- and polycrystalline silicon, describing two distinct regimes.13 He ascribed the carrier mobility limitation at low doping level to a potential barrier at grain boundaries, and the carrier mobility limitation at high doping level to ionized impurities. The Seto model qualitatively accounts for the dependence of μe on temperature (for T larger than 0 K) and on Nd in polycrystalline films. The model has been updated several times,14−16 and is widely used to describe the mechanisms limiting carrier mobility in polycrystalline TCOs.17−19 The model shows that, in polycrystalline

1. INTRODUCTION Transparent conductive oxides (TCOs) are highly doped degenerated wide-bandgap semiconductors that play a major role in optoelectronics. In applications, such as lighting, displays, and solar cells, the optoelectronic properties of the TCOs are instrumental for better device performance.1−6 Keeping high conductivity while reducing the absorptance in TCOs is a known challenge which requires in-depth understanding of their elemental composition and microstructure, and how these features govern the optoelectronic properties of TCOs. Electrical conductivity and optical absorptance in TCOs are commonly tuned by controlling the dopant concentration.7−9 The influence of dopant concentration (Nd) on carrier mobility (μe) and carrier concentration (Ne) in semiconductors has been investigated since the 1950s, especially for polycrystalline silicon.10 It was observed that, in such polycrystalline samples, the carrier mobility is thermally © XXXX American Chemical Society

Received: November 9, 2016 Accepted: February 2, 2017 Published: February 2, 2017 A

DOI: 10.1021/acsami.6b14350 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces films, the barrier height and depletion width at grain boundaries depend on the dopant spatial distribution (dipole approximation) and it assumes uniform spatial distribution of dopants in the film. Indeed, most of the works on TCOs and polycrystalline semiconductors assume homogeneous doping throughout the grains, with possible modifications only at the grain boundaries.20,21 In this work we investigate the validity of this assumption and show that it does not hold for borondoped zinc oxide (ZnO:B) films deposited using the low pressure metalorganic chemical vapor deposition (LPMOCVD) technique. We report that a-textured polycrystalline ZnO:B films with double-faced grains show a structural difference between the two sides of the grains which, using NanoSIMS technique, is found to impact the dopant incorporation inside the film. Further, we use Kelvin probe force microscopy (KPFM) to show that this dopant distribution affects the electronic properties of the film: the contact potential on the surface of doped samples follows the same nonuniform distribution as the dopant atoms. We attribute the observed bimodal distribution of the dopant atoms to the different type of surface termination on each side of the grain. These findings play a crucial role in the optimization of polycrystalline LP-MOCVD ZnO:B films, widely applied as electrodes in copper indium gallium selenide22,23 and silicon heterojunction solar cells,24,25 both technologies with proven conversion efficiencies of more than 22%.26−28 Indeed, simulations (not included in this paper) showed that LP-MOCVD ZnO films characterized by a bimodal dopant distribution are likely to reduce the resistive losses with respect to those having a uniform dopant distribution.29

with a scanning probe microscope (Bruker, Dimension Icon) used in Kelvin probe force configuration.34 Bruker PFQNE-AL probes (silicon pyramidal tip on a silicon nitride cantilever) were used. In order to reduce roughness-related artifacts before NanoSIMS and contact potential measurements, the ZnO film surface was polished with the chemical mechanical procedure35 shown in Figure S1.

3. RESULTS 3.1. Structural Difference between the Two Sides of a Grain. The grains composing polycrystalline zinc oxide films deposited by LP-MOCVD appear as wedges on the film surface, as shown in Figure 1a. These films are a-textured,

Figure 1. SEM micrographs of a-textured ZnO:B films illustrating the structural differences between the two sides of the grain: (a) as deposited, tilted view; (b) as deposited, top view; (c) after polishing, top view; and (d) after HCl-etching, top view. Film thickness ∼7 μm.

2. EXPERIMENTAL SECTION 2.1. Sample Preparation. ZnO thin films were deposited using LP-MOCVD on 4 × 4 cm2 0.5 mm-thick borosilicate glass substrates (Schott AF32). Diethylzinc ((C2H5)2Zn; DEZ) and water vapor (H2O) were used as precursors for zinc and oxygen, respectively. The intentionally doped films were made by adding diborane (B2H6, diluted at 2% in argon) in the gas phase as precursor for B, which acts as an n-type dopant in ZnO. The gases were injected into the chamber through a showerhead facing the hot plate (30 × 30 cm2) where the substrates are placed. A detailed description of the deposition system can be found elsewhere.30 During deposition the hot plate temperature was 180 °C, the total gas flow was 200 sccm, the H2O/DEZ ratio was 1 and for the doped films the B2H6/DEZ ratio was 0.01. This deposition condition leads to the a-textured rough films31 normally used as front electrodes in solar cells.22,25 2.2. Sample Characterization and Analysis. The thickness of the films was 7 μm, as measured by a stylus-profilometer (Ambios, XP200). The film morphology was assessed by a scanning electron microscope (JEOL, JSM-7500 TFE) in secondary electron mode. TEM samples were prepared by tripod wedge polishing using diamond lapping films and finishing with a colloidal silica suspension. The (scanning) TEM work was carried out on a FEI Tecnai Osiris operated at 200 kV. The scanning TEM image was acquired with a high angle annular dark field detector using a 115 mm camera length, corresponding to inner and outer collection angles of 65 and 200 mrad, respectively. CBED patterns were recorded along the [21̅1̅0] zone axis with a beam semiconvergence angle of 2.1 mrad. Blochwave CBED simulations were performed with JEMS.32 The dopant distribution in the film was assessed by NanoSIMS (Cameca, NanoSIMS 50L) using a < 100 nm-wide cesium ion (Cs+) beam to sputter the surface. Seven different isotopes were detected: BO−, BO2−, 64ZnO−, 66ZnO−, O−, CN− and C2−, as shown in Figure S3. Further details on the system are available elsewhere.33 The contact potential differences were measured in ambient atmosphere

meaning that the grains have the a-axis perpendicular to the substrate. The lateral surfaces of each wedge are characterized by straight lines which are roughly parallel to the ridge of the wedge (cf., Figure 1a and 1b). We believe that these lines correspond to basal plane stacking faults observed with high density by transmission electron microscopy (TEM) imaging, and so indicate the direction of the c-planes.36,37 Noteworthy, these lines were assumed to correspond to faceted surface steps on the wedges,38 thus exposing the a-planes parallel to the substrate and the c-planes perpendicular to the substrate. We note, however, that in this case the fineness of the steps is such that they have not been successfully resolved either by atomic force microscopy or by TEM cross section imaging. Scanning electron microscopy (SEM) observations of asdeposited (Figure 1a and 1b) and treated films (mechanochemical polished in Figure 1c, and HCl-etched films in Figure 1d) underline that each wedge consists of two distinct sides. The main differences between these two sides are (i) the slope of the faces (Figure 1a and 1b); (ii) the surface features, one side shows little bumps along the stacking faults lines, while the other shows elongated features roughly perpendicular to the stacking faults lines (Figure 1b); (iii) when polished, there appears a “middle line” dividing the grain into a brighter and a darker side (Figure 1c); and (iv) when immersed in a HClbased solution, one of the two surfaces is etched slower than the other, forming triangular spikes roughly perpendicular to the stacking fault lines (Figure 1d). B


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the patterns measured at the points 3 and 4 for the same orientation of c-axis relative to grain morphology but for a greater local thickness of 46 nm. 3.2. Direct Imaging of Boron Distribution. We have demonstrated that the different kinds of surface termination impact the growth rate and the microstructure of the grains. They could therefore also affect the way in which the dopant atoms are incorporated into the film. To verify the latter hypothesis we assess the distribution of the B atoms in atextured (i.e., with the majority of the grains having the a-axis oriented perpendicular to the substrate) polycrystalline ZnO films using NanoSIMS. There are three challenges to these measurements. First, for the LP-MOCVD ZnO:B film thickness normally used in solar cell applications (1−2 μm; note that Inbased TCOs are usually much thinner49−51), the grain size of an a-textured film is around 100 nm.52 These grain sizes are similar to the NanoSIMS spatial resolution, thus it is not possible to distinguish composition within an individual grain. However, for a-textured films, the grain size increases roughly linearly with the film thickness.19 Thus, for this study we prepared films with a thickness of 7 μm, corresponding to an average lateral grain size of about 2 μm at the surface (Figure 1a and 1b). Second, the roughness of the film just after deposition (shown in Figure 1a) can generate artifacts during SIMS measurements. Therefore, the film was polished using a chemical mechanical polishing process35 described in the Experimental Section and sketched in Figure S1. After polishing, the surface becomes suitably flat, as shown in Figure 1c. Third, when using a positively charged primary ion beam (Cs+), it is only possible to collect negatively charged secondary ions, making it difficult to detect species such as B that cannot easily be negatively ionized. To increase the secondary ion emission detection, it is preferable to setup the detection for BO− and BO2−.41 To assess the reliability of these procedures, we performed the measurements on a moderately doped sample and on a nonintentionally doped (nid) sample as a reference. The results of the NanoSIMS emission maps acquired on the polished ZnO films are shown in Figure 3. In the secondary electron maps (Figure 3a and 3b), the bright/dark contrast of each marked grain is readily visible for both nid and moderately doped samples. The map of the BO2 distribution (Figure 3c and 3d) is different between the two samples: while the nid sample does not contain any BO2, in the moderately doped sample the BO2 is preferentially segregated to the regions of the film corresponding to the bright sides in the secondary electron maps. Specifically, the bright regions in the secondary electron map contain two to three times higher amount of BO2 than that present in the darker regions of the grains (cf. Figure 3h). Moreover, Figure 3d and 3h show that, within the bright sides of the grain, the BO2 concentration is roughly uniform. We therefore describe the dopant distribution within the grains of the film as bimodal. As further evidence, in contrast, the ZnO distribution (Figure 3e and 3f) is more uniform than the BO2 distribution. In the nid sample the relative variation is below 5%, while for the moderately doped sample the relative variation is up to 10%. In addition, in the doped sample, the regions presenting lower values apparently correlate with higher values of BO2 (cf., Figure 3h); we attribute this correlation to the substitution of Zn by B atoms. A similar correlation for substitutional dopants was already observed in aluminum doped ZnO analyzed by atom probe tomography20 and energy dispersive X-ray spectroscopy.21 Finally, we note that, as well as

It is known that ZnO crystallizes in a noncentrosymmetric wurtzite crystal structure, and, consequently, the surfaces formed perpendicular to the c-axis are terminated by Zn atoms in the c+ direction and by O atoms in the c− direction39,40 as shown in Figure 2b. These two types of surface

Figure 2. Definition of the crystallographic directions within the ZnO grain. (a) Crystallographic directions with respect to a grain (SEM micrograph, tilted view). (b) Arrangement of zinc and oxygen atoms in the wurtzite structure showing the two preferential surface terminations perpendicular to the c-axis: O-termination in the c− direction and Zn-termination along the c+ direction. (c) High angle annular dark-field scanning TEM image of a plan-view TEM sample. Convergent beam electron diffraction patterns along the [211̅ 0] zone axis of a selected grain with matching simulations, showing that throughout the entire grain the c-axis is perpendicular to the middle line, and directed from the upper to lower side of the analyzed grain, without an inversion across the middle line (cf., 1 vs 2 and 3 vs 4).

terminations have different properties: in particular different growth rates and etching behavior are often reported for ZnO41−44 as well as in other wurtzite-based films.45 In our case, the difference in slope between the two faces can be related to the difference in growth rate between the two sides. Combined with the different etching behavior, this suggests that the two lateral surfaces possess a different polarity with a different type of surface termination that causes the observed asymmetries. This hypothesis requires that the sense of the c-axis does not change within a grain and in particular not across the middle line, that would otherwise suggest the presence of an inversion boundary commonly found in ZnO varistor ceramics.46 We verified that the c-axis direction does indeed remain the same within a grain by performing a TEM analysis using convergent beam electron diffraction (CBED) at different locations within a grain, as shown in Figure 2. The contrast in CBED patterns arising from dynamical scattering is commonly used to determine the sample thickness and polarity of crystals by comparison to simulated CBED patterns.47,48 Here, CBED patterns are taken along the [211̅ 0̅ ] zone axis of ZnO. These patterns are not centrosymmetric, a fact which can be used to determine the sense of the c-axis. CBED patterns within a single grain were recorded closer to the thin specimen edge (points 1 and 2) and further away from it (points 3 and 4). Figure 2c shows that the patterns recorded at the points 1 and 2 match a pattern simulated for a local thickness of 25 nm of the TEM lamella, and with an invariant c-axis going from the bright toward the dark side. A similarly good match is found between C


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Fermi level and the bottom of the conduction band (EF − Ec) of the film increases34 VCPD =

χf − (E F − Ec) − φt e

(2)

where χf is the electron affinity of the film surface. Equation 2 is valid assuming that the film surface is representative of the bulk and that, for example, no Fermi level pinning takes place at the surface.35,36 Considering that in degenerately doped semiconductors the value EF − Ec increases with increasing Ne,40 Equation 2 shows that VCPD decreases when increasing the doping level in the film. Figure S2 schematically explains how the contact potential difference is generated in our measurement setup, which is composed of a n-doped Si tip and a ZnO:B film. In principle, from eq 2 it should be possible to calculate the Ne of the film.55−58 In reality, because of the nonuniformity of the electric field at the tip surface, the measured values provided by the KPFM setup are influenced by the tip geometry and tip− sample distance.59,60 This generates an undesired additional interaction of the cantilever with the electric field, leading to an underestimation of VCPD, which makes quantitative analysis problematic. Therefore, we use the relation stated in eq 2 in a qualitative fashion. Figure 4 compares the surface height signal with the VCPD signal for nid and moderately doped films. In both samples, we observe a difference in height, within 5−10 nm, between the two sides of the grain (Figure 4a and 4b). Note that this height

Figure 3. Maps of emitted signals obtained using the NanoSIMS technique. Comparison between nonintentionally and moderately doped samples for (a, b) secondary electrons, (c, d) boron dioxide, (e, f) zinc oxide, and (g, h) line profiles for the three previous types of signals. Film thickness ∼7 μm.

the NanoSIMS measurements, we have attempted to confirm the observation of a bimodal B distribution using TEM-based electron energy-loss spectroscopy. However, boron is notoriously volatile in TEM53,54 (i.e., disappears from the thin TEM lamella from the vacuum and/or from electron beam illumination) and unfortunately its identification by this method has not been possible. 3.3. Spatially Resolved Electronic Properties. To verify whether the BO2 detected by NanoSIMS corresponds to the active B dopants in the ZnO films, that is, increasing Ne on one side of the grain, we performed KPFM scans on both nid and moderately doped samples. To correlate the spatial dopant distribution with the electronic potential, we measured the contact potential difference (VCPD) on a polished a-textured film similar to the one analyzed by NanoSIMS. VCPD depends on the film and tip workfunctions, respectively, φf and φt33 eVCPD = φf − φt

Figure 4. Scanning probe microscope measurements on polished samples. Comparison between nonintentionally doped and moderately doped samples for (a, b) surface height, (c, d) contact potential difference (VCPD), and (e, f) line profiles referring to the two previous parameters.

(1)

where e is the elementary positive charge. By assuming that the tip workfunction is constant during the scanning, it is possible to show that VCPD decreases when the difference between the D


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ACS Applied Materials & Interfaces difference has little effect on the determination of VCPD, since no corresponding difference in VCPD is observed in the nid film, as shown in Figure 4c. This minimal effect of the height difference on the determination of VCPD is further supported by the homogeneous local potential distribution over each side of a grain and the similar VCPD on the bright side of different grains in the doped sample (Figure 4d). Figure 4e and 4f show the line profile of the contact potential difference, where the higher side of the grain shows a lower VCPD, that is, a higher Ne than the lower side (eq 2). This demonstrates that the B segregation on one side of the grain, as indicated by the BO2 mapping, provides active B dopants, locally increasing the Ne in the ZnO:B films. The measured VCPD difference between the two sides of the grain in the doped film is around 15 meV. According to the Burstein−Moss shift, corrected for the band gap renormalization, the difference in film workfunction between a carrier concentration of ∼3 × 1019 cm−3 (side containing less boron) and ∼1.2 × 1020 cm−3 (side containing more boron) should be around 80−90 meV.61 The discrepancy between the measured and the expected difference in workfunction can be explained considering the technical aspects of the standard KPFM technique.34 Namely, the measured value strongly depends on the geometry of the cantilever and the tip itself, as well as on the tip’s actual condition (its shape, sharpness and contamination).59,60 It has been shown that this limitation, which is difficult to quantify, is introduced by the interaction between the whole cantilever (not only the tip) and the sample surface. This interaction is an undesired contribution which generally reduces the range of surface potential changes compared to expected values (contrast and absolute value of the CPD). This effect is well illustrated, for example, by Maturova et al.62 In this paper a mixture of two polymers with a large contrast in workfunction (1.4 V) was probed by KPFM. The measured contrast of the CPD was below 100 mV, several times smaller than the real difference.

in the direction parallel to the substrate (c-axis). As shown in Figure 5b, a film growing along the c+ direction offers three

Figure 5. Cross sections of a single grain: (a) in a TEM micrograph and (b) in a drawing showing that the lateral surfaces of the wedge exposed to the gas phase surface terminations having a different number of bonds at the cation site: one along the c− direction and three along the c+ direction. Note that the vertical line contrast in the TEM image corresponds to multiple, dense basal plane stacking faults as mentioned in the main text.

bonds to the incoming Zn adatom while a film growing along the c− direction offers only one bond, thus the desorption probability of a Zn adatom is higher along the c− direction, and in turn the Zn-sticking coefficient is lower. On the other hand, B has a lower vapor pressure in the gas phase (for moderately doped samples, the ratio between DEZ and B2H6 is roughly 100) and therefore we can assume no desorption of B adatoms from the film surface. Thus, the B-sticking coefficient does not depend on the number of bonds at the cation site, which suggests that B would be equally incorporated on both sides of the grain. However, as discussed above, Zn prefers to sit on the c+ side of the grain, therefore hindering the incorporation of B on this side, and which explains the higher concentration of B atoms measured in the side of the grain growing along the c− direction. While here we provide a model that gives one possible explanation of the incorporation difference, future work is however required to clearly establish the relation between the doping distribution and the film polarity. It is noteworthy to comment here that, by means of simulations, we identified that the bimodal dopant distribution is actually beneficial for the film conductivity compared to a uniform dopant distribution.29 The simulation outcomes show that, for the same carrier concentration, the contact resistivity at a grain boundary is lower for the bimodal case, at the interface between the B-rich side of one grain to the B-rich side of another grain. Therefore, assuming the existence of a continuous percolation path through the dopant-rich regions (a plausible assumption according to SEM observations of the LP-MOCVD ZnO:B films, as shown in Figure 6), the film conductivity would be higher for the bimodal distribution than for the uniform one.29

4. DISCUSSION We demonstrated that the dopant distribution in a-textured polycrystalline ZnO:B films is bimodal and that the two sides of each grain are characterized by different microstructural, chemical composition and electronic properties. In this section we discuss the reason leading to these differences. The large majority of literature reports show that the dopant distribution in polycrystalline films is uniform or, more rarely, segregated at grain boundaries. To our knowledge, the here observed bimodal distribution has not been previously reported in literature. An incorporation dependency on film polarity has been reported for single crystalline GaN63,64 and ZnO films.65 In particular, Mg-doped ZnO deposited by pulsed laser deposition showed higher Mg concentration on c− oriented than on c+ oriented films.66 The authors attributed this observation to the difference in sticking coefficient between Zn and Mg atoms on the O- and Zn-terminated surfaces (c− being the O-terminated surface, and c+ the Zn-terminated one; cf., Figure 2b).67 Kato et al. further suggest that, on the one hand, Zn has a high vapor pressure in the gas phase and therefore some of the adsorbed Zn adatoms can desorb from the film surface. The desorption probability of Zn adatoms depends on the number of bonds at the cation site: the more available bonds, the lower the desorption probability. In LP-MOCVD ZnO the grains grow both in the direction perpendicular to the substrate (a-axis) and E



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5. CONCLUSION In this Research Article, we demonstrated the difference in microstructure between the two sides of individual grains in atextured LP-MOCVD ZnO films. In addition, we unveiled a remarkable bimodal spatial distribution of B dopants using NanoSIMS and KPFM techniques. B atoms were found to be preferentially incorporated into one side of each ZnO grain. Contact potential measurements showed how the nonuniform dopant distribution locally affects the carrier concentration in the film. The microstructural difference and the observed bimodal distribution are attributed to the difference in surface termination between the two sides of each grain. By controlling the in-plane orientation of the grains, it should be possible to obtain continuously doped regions that would constitute low resistivity paths for charge carriers. ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b14350. Details of the mechanochemical polishing, electronic energy band alignment for the system composed by tip and film used in the KPFM, and maps of all the secondary particles (BO−, BO2−, 64ZnO−, 66ZnO−, O−, CN−, C2−) emitted by nid and doped ZnO samples (PDF)

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REFERENCES

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Figure 6. (a) Electronic band diagram at a grain boundary assuming the bimodal dopant distribution. (b) Example of an arbitrary “low resistivity” percolation path within the dopant-rich regions of the film.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +41 21 695 4347. Fax: +41 0 21 695 4201. ORCID

Fanni Lorenzo: 0000-0002-2642-2434 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Funding

This work was funded by the Swiss National Science Foundation (FNS) under the project ZONEM (Grant No. 137833) and supported by the project 16-10429J of the Czech Science Foundation. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Prof. Anders Meibom for the discussions on the interpretation of NanoSIMS measurement, Benoit̂ Delaup for the HCl etching of the samples and Dr. Quentin Jeangros, Dr. Robin Schäublin, and Dr. Stephan Gerstl for experimental trials with alternative techniques and discussions. F


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