In Situ Optical Monitoring of New Pathways in the Metal-Induced

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In situ optical monitoring of new pathways in the metal induced crystallization of amorphous Ge D. Pelati, G. Patriarche, O. Mauguin, L. Largeau, F. Brisset, F. Glas, and F. Oehler Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.7b00799 • Publication Date (Web): 25 Sep 2017 Downloaded from http://pubs.acs.org on September 29, 2017

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Crystal Growth & Design

COVER PAGE Authors: D. Pelati,†,‡, G. Patriarche,‡ O., Mauguin,‡ L. Largeau,‡ F. Brisset,§ F. Glas,‡ F. Oehler*,‡ †RIBER SA, 31 rue Casimir Périer, 95870 Bezons, France ‡Centre for Nanoscience and Nanotechnology, CNRS, Université Paris-Sud, Université ParisSaclay, Route de Nozay, 91460 Marcoussis, France §Institut des Chimie Moléculaire et des Matériaux d'Orsay, CNRS, Université Paris-Sud, Université Paris-Saclay, 91405 Orsay, France ǁInstitut Photovoltaïque d’Ile-de-France (IPVF), 92160 Antony, France ABSTRACT

We use high-resolution optical microscopy to characterize in situ the processes at play during the Al-induced crystallization of amorphous Ge. In addition to the well-established aluminuminduced layer exchange (ALILE) process, we demonstrate the existence of another crystallization mechanism with different kinetics and spatial extension using in situ monitoring. Further ex situ characterizations show that both processes are active in our samples. The ALILE process is found to create a single Ge layer and 111-oriented crystallites in our growth conditions, while the other crystallization process yields a double Ge layer with a mixed 111 and 100 orientations in the bottom layer while the top layer remain amorphous. This work underlines the importance of in situ monitoring for the understanding and modeling of metal induced crystallization.

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* Fabrice Oehler - Researcher C2N - Centre for Nanoscience and Nanotechnology UMR9001 Site Marcoussis -Route de Nozay 91460 Marcoussis, France http://www.c2n.universite-paris-saclay.fr/ Phone : +33 (0) 1 69 63 63 76 / Fax : +33 (0) 1 69 63 60 06


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In situ optical monitoring of new pathways in the metal induced crystallization of amorphous Ge D. Pelati,†,‡,ǁ G. Patriarche,‡ O., Mauguin,‡ L. Largeau,‡ F. Brisset,§ F. Glas,‡ F. Oehler*,‡ †RIBER SA, 31 rue Casimir Périer, 95870 Bezons, France ‡Centre for Nanoscience and Nanotechnology, CNRS, Université Paris-Sud, Université ParisSaclay, Route de Nozay, 91460 Marcoussis, France §Institut des Chimie Moléculaire et des Matériaux d'Orsay, CNRS, Université Paris-Sud, Université Paris-Saclay, 91405 Orsay, France ǁInstitut Photovoltaïque d’Ile-de-France (IPVF), 92160 Antony, France ABSTRACT

We use high-resolution optical microscopy to characterize in situ the processes at play during the Al-induced crystallization of amorphous Ge. In addition to the well-established aluminuminduced layer exchange (ALILE) process, we demonstrate the existence of another crystallization mechanism with different kinetics and spatial extension using in situ monitoring. Further, ex situ characterizations show that both processes are active in our samples. The ALILE process is found to create a single Ge layer and 111-oriented crystallites in our growth conditions, while the other crystallization process yields a double Ge layer with a mixed 111 and 100 orientations in


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the bottom layer while the top layer remain amorphous. This work underlines the importance of in situ monitoring for the understanding and modeling of metal induced crystallization.

INTRODUCTION Metal induced crystallization (MIC) is the crystallization of an amorphous material at low temperature in contact with a metal. Although MIC is considered a pure solid-solid reaction, its existence is closely related to that of eutectics or defined compounds1. In particular Knaepen et al.2 report that crystallization temperature of the MIC between various metal/Ge couples is typically two thirds of that of the lowest temperature germanide or eutectic in the metal-Ge binary phase diagram. The particular case of Al-induced crystallization (AIC) has attracted much interest due to the low cost of Al, the reduced temperatures of the AIC of amorphous Ge and Si with respect to other crystallization techniques and its potential compatibility with existing CMOS processes. Both AIC-Si and AIC-Ge have been investigated as potential candidates for various devices, including solar cell absorbers3,4, transistors5 or buffer layers for epitaxy6,7. AIC is also referred to as the ALILE process, for aluminum-induced layer exchange. This follows from the work of Nast et al.8 on AIC-Si, which shows that the Al and amorphous Si layers exchange positions grain by grain as the Si crystallizes. Typically, the Si (Ge) atoms first diffuse and wet the Al grain boundaries. Above a critical concentration, Si (Ge) crystals nucleate and then grow laterally. This creates stress and causes the outward diffusion of the Al away from the crystalline Si (Ge). Overall the free energy of the system decreases due to two phenomena: the crystallization of the amorphous Si (Ge) material and the creation of new interfaces9–11. Thermodynamics may also explain the very different kinetics between AIC-Ge and AIC-Si, with


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different locations for the initial nucleation site: inside Al grain boundaries, at the Al/amorphousGe(Al/a-Ge) or the Al/substrate interface10,12. The local environment of the crystal nuclei may also promote a preferential crystal orientation of the final Si or Ge layer13. There is now a growing body of literature aimed at obtaining (111)-oriented Ge layers using AIC. The main difference between standard (randomly oriented) AIC-Ge and oriented AIC-Ge is the presence of a carefully engineered diffusion barrier, usually an Al or Ge oxide (AlGeOx), between the Al and Ge layers. Through various optimizations of layer thickness14, temperature15, diffusion barrier16, substrate17, and more recently Al-grain size18, near complete(111)-oriented layers can be obtained. Additional work by Park et al.19 on Au-induced crystallization of Ge shows that the chemical nature of the lower interface, Au-silica or Au-alumina, may promote Ge (100) orientation instead of (111). Modeling highlights the central role of the diffusion barrier or interlayer in the formation of oriented crystallites by AIC11,20,21. The ability to monitor and quantify the nuclei density or the crystallization rate in situ during annealing is unmatched compared to time series in which thermal treatments are stopped and resumed at regular intervals. For the optimization of (111)-oriented AIC-Ge, for which annealing durations can reach several tens of hours15, in situ monitoring becomes a critical requirement. However, the technical difficulties in observing hot samples (200-350°C) in vacuum or nitrogen ambient at high magnification have limited the use of in situ optical microscopy. Some recent works on AIC-Si are still based on low-resolution imaging22–24. Only last year, Tutashkonko et al.25 have reported better resolution (~5-10µm) which allowed the authors to better characterize the processes at play during the AIC of Si. The present work reports the in situ optical monitoring of the AIC of amorphous Ge with an even higher spatial resolution (~1µm), close to the optical limit.


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EXPERIMENTAL

Figure 1. Schematic of the in situ optical monitoring experimental setup with the programmable hotplate, the nitrogen flux, the optical system and the motorized autofocus system. We use a long working distance microscope 50x objective (Mitutoyo Plan Apo 378-805-3), with an optical resolution limit of 500 nm, associated with a programmable hot plate (HARRY GESTIGKEIT) which lid has been extensively modified to provide a uniform nitrogen supply and an optical viewport made of fused silica (Figure 1). After damping of the mechanical vibrations, the optical resolution is ultimately limited by the turbulent air flow between the relatively hot fused silica viewport and the cooler surface of the microscope objective. At such a high magnification, we find that a motorized automatic focus is necessary to maintain sharp images, given the different thermal expansions of various elements during the 8-15 h of annealing.


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Figure 2. Main steps of sample preparation and annealing for the AIC of amorphous Ge (a-Ge) leading to a final crystalline Ge layer (c-Ge). A thin (1 nm) amorphous Ge is used as underlayer. The AlGeOx diffusion barrier is obtained by air oxidation of another a-Ge layer. The fabrication steps of the samples are shown in Figure 2. All materials are deposited in the same e-beam evaporator (Plassys MEB 550SL), which allows us to switch material sources without exposing the sample to atmosphere. All thicknesses and deposition rates are calibrated using a quartz scale inside the evaporator. Starting from a clean silica surface, we first deposit a 1 nm Ge underlayer26 (at a rate of 0.01 nm/s), followed by 20 nm Al (1-2 nm/s) which is coated with a 2 nm thin Ge layer (0.1 nm/s). This top Ge layer is exposed to air (9 min) to provide an AlGeOx diffusion barrier27, before the deposition of the final 20 nm Ge layer (1 nm/s). When a transparent substrate is required, the silica surface is that of a 2’’ bulk silica wafer (JGS2, UniversityWafer Inc.). Otherwise, we use thermal silica obtained by dry oxidation, (1050°C 15 min) from a standard 2’’ Si (100) wafer (Siltronix). All silica surfaces are HF treated (at 5% for 3 min) prior to Al and Ge deposition in order to remove potential surface contamination.


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In addition to optical microscopy, grazing incidence x-ray diffraction (GIXRD, incident angle = 0.3 °) is used to assess the crystalline properties of the layer (Rigaku SmartLab equipped with a Cu rotating anode delivering Cu K-alpha doublet radiation). Compared to conventional geometries (Bragg-Brentano), GIXRD is able to extract signal from very thin layers but only the planes perpendicular to the sample normal are allowed to diffract. Therefore, a strong (111) texture is not characterized by a strong 111 diffraction peak, but by the presence of strong 110 and 112 diffraction lines, which are perpendicular to the [111] direction. In addition to XRD, electron backscatter diffraction (EBSD) is used to assess the crystal orientation of the grains at the micron scale (ZEISS Supra 55 VP, Hikari/OIM TSL EDAX). Top-view chemical analysis is performed in a scanning electron microscope (SEM, FEI Magellan) by energy dispersive x-ray spectroscopy (EDXS, Oxford INCA). Cross-sectional Transmission Electron Microscopy (TEM, FEI Titan) with chemical analysis by EDXS (Bruker Super-X) is used for high resolution chemical mapping, using focused ion beam (FEI Scios) to extract zones of interest from the sample.

RESULTS AND DISCUSSION


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Figure 3. In situ optical microscopy of the back of the sample (Al-side) during the AIC annealing at 247 °C. (i) and (j) are zoomed views of the areas indicated by rectangles in (d) and (f), respectively. Scale bar for images (a) to (h) is 40 µm. Black arrows indicate the MIC-1 nucleation and growth. Red arrows indicate the onset of the MIC-2 process.


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Figure 3 illustrates the optical monitoring of a sample during annealing. After an initial temperature ramp from ambient to 238 °C (~1 h), the sample is maintained at a fixed temperature (238 ± 1 °C) for 16 hours while optical images are acquired every 30 s. We use a transparent substrate (see experimental section) to monitor the optical changes in the Al layer at the bottom interface. After 0.05 h (Figure 3a), we observe the formation of small spots which enlarge in a dendritic fashion over several tens of microns (arrows, Figure 3b). We refer to this process as MIC-1 in the rest of the text. After about 2.9 h of annealing, a new darker contrast appears near the previously formed dendrites (red arrows, Figure 3e and Figure 3i). This feature then slowly expands (Figure 3e and Figure 3f), encircling the previous dendrite and merging with other darker fronts (Figure 3f and Figure 3g) until it finally fills the field of view in approximately 14 h (Figure 3h). This second process is described as MIC-2 in the rest of the manuscript. The complete movie, cropped to a selected area to limit file size, can be found in the Supporting Information. High magnification images (Figure 3i and Figure 3j) show that both processes are actually dendritic. MIC-1 (Figure 3i, gray contrast) produces long dendritic branches which are typically 20-30 microns long and a couple micron wide. MIC-2 (Figure 3j, darker contrast) consists of a dense collection of small dendrites, with a typical scale of 1-2 micron. Additional in situ optical monitoring on a sample without Ge underlayer shows the same evolution with the MIC-1 and MIC-2 processes occurring on the surface in a similar manner (see Supplementary Information S1).


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Figure 4. Spatial extension of the MIC-1 (a) and MIC-2 (b) processes as a function of time. In (a), letters A, B, C refer to three different branches of MIC-1 dendrites (inserts show images taken at times indicated by dotted lines, scale bar is 40 µm). (b) Spatial extension of the MIC-2 process using length measurements, L1 and L2, (see inserts, scale bar is 40 µm). To better appreciate the different kinetics of the two phenomena at the same temperature (238 ± 1 °C), we now quantify in Figure 4 the spatial extension rate of the MIC-1 and MIC-2 from the movie shown in Figure 3. In Figure 4a, we measure the spatial extension of large dendrites created by the MIC-1 process. We find that the ‘arms’ of the MIC-1 dendrites all expand at a


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fixed rate of 10.5 ± 0.3 nm/s, until they stop completely. The stop time can vary between dendrites and between the arms of the same dendrite. All the observed MIC-1 dendrites stop growing after 1 h or 1.5 h of annealing and terminate well before reaching another MIC-1 dendrite. The situation is very different for the MIC-2 process (Figure 4b), with a slower but steady growth rate of 2.1 nm/s, measured over more than 7 h. The crystallization front progresses in an isotropic manner, with the same velocity, measured directly from length. The MIC-2 process only stops when the whole surface of the sample has been covered or if the sample is cooled down. The kinetics of the MIC-1 and MIC-2 processes can be monitored from the top or the back of the sample, see Supporting Information S2, as the Ge layer is thin enough (20 nm) to be semi-transparent in the visible range. The kinetics measured from the top are comparable to that measured from the bottom. They differ only by the small difference in the annealing temperature due to the different substrates used in the experiment (fused silica versus oxidized silicon wafers).


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Figure 5. Ex situ characterization of samples at different stages of annealing. (a) optical microscopy. Capital letters ‘A’, ‘B’ and ‘C’ refers to samples before ‘A’ and after ‘B’ the MIC-1 appearance, while sample ‘C’ is fully annealed. Scale bars are 10 microns. (b) GIXRD of the previous samples. The top axis indicates the position of Ge-related diffracted intensities. Red label marks the intensities characteristic of a perfect (111) texture in GIXRD. Al-related diffracted intensities (111, 220, 311, 222) are indicated by arrows. The letter ‘R’ refers to a reference Ge powder diffractogram. The baseline for the analysis of each diffractogram is indicated by a solid black line.

Table 1. Calculated Lotgering factors ( ) for the 111 (100) texture of samples (B) and (C) reported in Figure 5.

(B)

0.08

0.12

(C)

0.84

0.74

Using in situ optical microscopy, we determine precisely the moment when to cool the sample abruptly to study specific configurations. In Figure 5, we compare the GIXRD signals of the reference sample A prior to the annealing, of sample B with only MIC-1 dendrites and of a fully annealed sample C with MIC-1 and MIC-2 covering the whole surface. Optical micrographs and the corresponding diffractograms are presented in Figure 5a and Figure 5b. Given the large difference of GIXRD signal between the samples, the count rates are plotted in log scale in Figure 5b. The baseline of each diffractogram is indicated by a solid black line. The “reference” sample A only shows amorphous material, while “MIC-1 only” sample B exhibits a small


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diffraction signal and “MIC-1-2” sample C shows the most intense peaks. Comparison to the reference Ge powder diffractogram (R) (ICDD card 00-004-0545) shows that the 220, 422 and 400 diffracted intensities are more intense in our sample, which indicates a preferential crystal orientation of the crystallites. Simple texture analysis is performed using a Lotgering factor28,29 adapted to GIXRD. The Lotgering factor L varies between two extreme values, L=0 and L=1, respectively corresponding to perfect random orientation and perfect texture. It is computed from the normalization of parameters p and p0 using the following formula: =

− 1 −

In the following, we note the parameter p0 for the (111) texture as in the case of

conventional XRD (i.e. Bragg–Brentano geometry or similar). It is computed by summing all the (111)-related intensities of a perfect randomly-oriented sample divided by the sum of all the diffracted intensities. If no such sample is available for the experimental determination of , reference powder data from the literature can be used. =

+ + + ⋯ ∑

Parameter is computed in the same way using the diffractogram of the sample of interest. If the sample is perfectly (111)-textured only the Ixxx intensities are present and we find

= 1. Alternatively, if the sample has a perfect random orientation we obtain =

. The corresponding Lotgering factor simply rescales the previously determined

parameter so that = 0 for perfect random orientation and = 1 for a perfect (111) texture.


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Here we adapt the Lotgering factor to the GIXRD geometry . To obtain we now sum over the diffracted intensities which are perpendicular to the [111] directions, i.e. the [110] and [211] families. For our experimental diffractogram (see Figure 5), we use the following formula:

=

+ ! + !! ∑

The parameter is computed in the same manner from reference powder data.

Similarly, the Lotgering factor of the (100) texture can be computed by summing the 220, 400 and 440 diffracted intensities, which are all perpendicular to [100]. Note that the constructed GIXRD Lotgering factor is less discriminating than the standard XRD Lotgering factor due to overlapping reflection families, as seen above the diffracted intensities and !!

contributes to both and .

In Table 1, we compare the values of and . We observe that sample (B) with the MIC-1 dendrite has a small preference toward the (100) texture while the fully annealed sample (C) is more textured along the (111) direction.


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Figure 6. EBSD analysis of the Ge layer after complete AIC and removal of Al by wet etching. The normal of the Ge crystallites is shown projected along the vertical direction in (a), i.e. perpendicular to the substrate, and in the plane in (b), i.e. parallel to the substrate.

To clarify the texture analysis obtained by GIXRD, we have performed EBSD on a fully annealed sample. For this particular sample, the top layer was etched chemically (H3PO4) to expose only the bottom Ge layer. Figure 6a presents a map of crystal normal projected along the vertical direction. We remark that Ge crystals form isolated islands (30-40 µm diameter), mostly (111)-textured, organized around a central dendritic zone with typical (100) texture. As a example a (111)-oriented island is limited by a dotted line in Figure 6a, while the corresponding central dendrite is highlighted by a full line. The in-plane projection of the crystallite normal is shown in Figure 6b. There, the central 100 dendrite are marked by black lines. We observe that each large (111)-textured island is actually composed of different grains, with random in-plane orientations, organized around the dendritic (100)-textured core.


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Figure 7. Top view SEM image (a) and associated EDXS map (b) showing the Al (red) and Ge (light blue) signals of a partially annealed sample. The label R indicates the reference area, M1 the dendrite created by MIC-1 and M2 the surface created by MIC-2. The dashed white line indicates the FIB cut used to extract a TEM sample containing the interfaces R/M2 (pink rectangle) and M1/M2 (blue rectangle). (c-e) TEM-EDXS maps of the Ge (light blue), Al (red) and O (green) signals from R/M2. (f-h) TEM-EDXS map from M1/M2. Scale bars in (c-h) are 100 nm. In the next experiment, we stopped the annealing at the start of the MIC-2 phase when MIC-1, MIC-2 and reference area coexist. This sample was first analyzed by EDXS in top-view SEM. Figure 7a shows the secondary electron (SE) image, in which the signal from the reference area


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(R, dark gray) appears very different from that of the large MIC-1 dendrite and MIC-2 area (M1 and M2, light gray). However, there is no clear SE contrast between the MIC-1 and MIC-2 area. The corresponding chemical image (Figure 7b) shows the opposite: the upper MIC-1 dendrites, are now clearly visible with a Ge signal and very little Al signal. To help visualizing the MIC-1 area in the SE image (Figure 7a), we report the shape visualized in EDXS (Figure 7b), using a thin white line. Contrary to the clear MIC1-MIC2 different composition, we note that the reference area and the MIC-2 area appear similar, with both Al and Ge signals. This particular structure was further investigated by cross-sectional TEM using a focused ion beam to extract a slice along the dotted line of Figure 7a and Figure 7b. Two areas are of particular interest, the R/M2 interface at the crystallization front between MIC-2 and the reference area (Figure 7c Figure 7e), and region M2/M1 at the interface between MIC-1 and MIC-2 (Figure 7f - Figure 7h). The R/M2 region shows the inversion of the Ge layer, initially at the top (R) and then at the bottom of the stack (M2), after the MIC-2 process, with a mixed region at the growth front (i). The region M2/M1 shows a different structure with the inverted layering in the M2 area (Ge below Al) and a double layer of Ge in the MIC-1 area, without any Al signal (Figure 7f and Figure 7g). Both areas R/M2 and M2/M1 yield similar oxygen signals, in which we distinguish the silica substrate, the AlGeOx interlayer and top surface oxidation (white arrows, Figure 7e and Figure 7h).


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Figure 8. EBSD analysis of sample stopped just after MIC-1. (a) Analysis of the top Ge layer from an SEM image with partial chemical signal and selected EBSD pattern. (b) Similar analysis of the bottom layer, exposed after wet etching. (c-d) EBSD analysis of the bottom Ge dendritic layer, showing the projection of crystallite normal along the vertical direction (c) and in the plane (d).

To further specify the crystal orientation of the Ge double layer in the M1 region, we have performed additional EBSD characterization. In Figure 8a, we first investigate the crystallinity of the top Ge layer. For this particular SEM observation we obtain a partial chemical contrast using 20 kV acceleration and a strong negative polarization of the SE detector to select backscattered electron. These conditions allow to reveal the MIC-1 dendrite shape, despite the presence of uniform top Ge layer. The EBSD pattern of the top Ge layer inside and outside of the dendritic area show no feature, corresponding to amorphous Ge material. We then investigate the bottom


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layer using wet etching (H2O2) to remove the top amorphous Ge layer. Figure 8b shows the SEM image (20 kV, standard imaging conditions for secondary electrons) along with the EBSD patterns of the Al layer (outside the dendrites area) and Ge area (inside the dendrites). The Al layer presents weak diffraction patterns, while dendrite Ge displays strong diffraction patterns. All EBSD patterns in Figure 8a-8b are acquired in the same experimental conditions and can be directly compared. Further EBSD analysis of the crystalline Ge area is shown in Figure 8c and 8d. The crystalline Ge dendrites exhibits a mixed (100) and (111) texture in the vertical direction, Figure 8c, and are composed of smaller grains with a random in-plane orientation, Figure 8d.

DISCUSSION Combining the GIXRD results (Figure 5) and the cross-section TEM experiment (Figure 7), it is clear that the MIC-2 process is indeed the ALILE process, as initially reported on amorphous Si8,9 and later on amorphous Ge14,30. The Al and Ge layers are clearly seen to exchange positions in the TEM images while the Ge diffraction signal significantly increases. EBSD also shows that the ALILE/ MIC-2 process, which results in coverage of the whole sample surface, yields mostly (111)-oriented Ge crystallites (Figure 6). Preferential (111) crystal orientation is also clearly visible in the GIXRD diffractogram and quantified by the calculated Lotgering values (Table 1). This type of (111) preferential orientation of the Ge crystallites has already been reported for similar samples in which a silica substrate, a Ge underlayer26,31and an AlGeOx27,31,32diffusion barrier are combined and annealed at low temperature.


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Using only ex-situ measurements, we could analyze the MIC-1 process in a similar fashion. Reference experiments carried out at 270 °C (not shown) indicate that, in the absence of Al, the sample does not evolve over 40 h. In the presence of Al, GIXRD indicates that a small Ge volume crystallizes as the MIC-1 dendrites appear. We can thus safely conclude that MIC-1 is indeed an AIC process during which Al catalyzes the crystallization of amorphous Ge. The dendritic shape of the (100) oriented regions, visible in the EBSD maps (Figure 6), reminds of the geometry observed optically in Figure 3, with a MIC-1 dendrite in the middle of a (111) island created by the MIC-2/ALILE process. This suggests that these partially (100)oriented regions revealed by EBSD are related to the large dendrites created by the MIC-1 process. This hypothesis is also supported by the Lotgering values calculated from GIXRD (Table 1) which hint at a preferential but partial (100)-orientation of the MIC-1 crystallites. The EBSD results of Figure 6 are obtained on a fully annealed sample in which all the area has been transformed by the MIC-1 and MIC-2/ALILE processes. The relevant interface in the partially annealed sample studied by TEM is thus the M2/M1 region (Figure 7a, blue rectangle, and Figure 7f to 7h). We observe there that the standard structure of the ALILE process (MIC-2) stops on a double layer of Ge. The top-view chemical image (Figure 7b) indicates that this double layer exists not only in the M2/M1 region, but also extends to the reference area yet unaffected by the ALILE process. From the ex situ measurements, we infer that the double Ge layer is likely created before the ALILE process, in the form of large Ge dendrites. The in situ optical microscopy results lift any remaining doubts, Figure 3 shows that the only large dendrites observed on the sample are created by the MIC-1 process prior to ALILE (MIC-2) and that nothing happens after or


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concurrently. Combining ex situ and in situ results we conclude that the dendrites created by the MIC-1 process takes the form of a double layer of Ge. Further analysis of the MIC-1 dendrites requires a specific experiment in which we use the in situ optical monitoring to stop the annealing before the onset of the ALILE (MIC-2) process. This allows one to selectively investigate the crystal orientation created by the MIC-1 mechanism. In Figure 8 we observe that the top Ge layer of the sample is still amorphous while the lower Ge layer is crystalline with mixed (100) and (111) orientations. As (111) crystallites can be both created by the MIC-1 and the ALILE (MIC-2) processes, this also explains why the (100)-textured regions observed of Figure 6 only loosely define the dendrite shape. Careful reading of the literature suggests that MIC-1 might already have been observed, but not identified as a distinct process. For instance, Nakazawa et al.30 report that AIC can be problematic for Ge layers thinner than 30 nm (we use here 20 nm thick layers). Toko et al6,33 observe that double layers of crystalline Ge can form during the Ge-AIC. The top Ge layer consists of random oriented crystallites and can be removed by appropriate wet etching. Double crystallized layer have also been observed during the AIC of SiGe material34. The main difference with the current findings is that the present top Ge layer remains amorphous after the MIC-1 process. However we do observe a Ge single layer, with (100)-textured dendritic shapes, in Figure 6, after the both the ALILE (MIC-2) and wet etching has been performed. This indicates that either the ALILE or the etching was successful in removing or relocating a fraction of the top Ge layer, possibly due to its small thickness.


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Figure 9. Schematic of the studied AIC processes MIC-1 and ALILE (MIC-2) transforming amorphous-Ge (a-Ge) into crystalline Ge (c-Ge).

In Figure 9 we summarize the result of the MIC-1 and ALILE (MIC-2) processes. It is clear that both processes are characterized by the diffusion of amorphous Ge into the Al layer followed by its crystallization. It is important to note that the local environments of the MIC-1 and the MIC-2 (ALILE) processes are comparable. The oxygen signal is the same for both R/M2 and M2/M1 regions (Figure 7e and Figure 7h) with a clear signal from the AlGeOx diffusion barrier. We thus have two distinct AIC processes in comparable locations and yet they have very different characteristics. The crystallization rate of MIC-1 is several times that of the MIC-2 (ALILE) at the same temperature (Figure 4). The MIC-1 process also stops, for some unknown reason, well before the complete coverage of the surface, contrary to the MIC-2/ALILE process. In all our samples, the MIC-2/ALILE process always starts in the vicinity of a MIC-1 dendrite (Figure 3). This suggests that MIC-1 creates boundaries with a high free energy, possibly strained or defective. The observation of mixed (100) and (111) texture in the MIC-1 indicates that Ge crystal nucleation may occur not only at the lower Al/SiO2 interface14,19 but also elsewhere such inside Al grain boundaries10,34. This suggests that the MIC-1 process may be related to small defects in the AlGeOx barrier that we are not able to resolve in our TEM observations (Figure 7). This hypothesis could also explain the different kinetics of the MIC-1 and ALILE (MIC-2) process, as the first may proceed nucleation inside the Al grain boundaries and the latter from nucleation at the Al/SiO234 interface. Without time dependence and high-resolution in situ optical monitoring, it is easy to mistake the MIC-1 region as a poorly oriented area created by a single ALILE process or by consecutive


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ALILE processes. Our in situ optical monitoring clearly demonstrates that both hypotheses are invalid, at least in our experimental conditions. During this distinct AIC process amorphous Ge diffuses and crystallizes inside the Al layer with mixed (100) and (111) orientations while the top Ge layer remain amorphous. From the practical point of view, the MIC-1 phenomena is detrimental to the quality of the layers obtained through ALILE and should be limited. Further work is now required to better identify the active crystallization pathways in MIC-1 and its associated mechanism.

CONCLUSIONS In conclusion, we identify and characterize a new AIC process, distinct from ALILE, by using a combination of high-resolution in situ optical microscopy and ex situ analyses. This new AIC process results in the formation of a double layer of Ge in which only the bottom layer is crystalline with mixed (100) and (111) orientations while the top layer is amorphous. It also occurs before the ALILE process, at lower temperature and with a faster crystallization rate, which suggests a different formation mechanism possibly through defects in the AlGeOx diffusion barrier. In practice, this phenomenon limits the crystalline quality of (111)-oriented Ge films obtained by AIC. Its existence requires new investigations to better identify the active crystallization pathways in AIC processes and the associated mechanism. AUTHOR INFORMATION *(F.O.) Email: [email protected] SUPPORTING INFORMATIONS


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•

Movie associated to the images shown in Figure 3, cropped to a selected area. Compressed due to minimize file size.

•

S1 “In situ optical microscopy from the back with no Ge underlayer”

•

S2 “Spatial extension of MIC-1 and MIC-2 measured from the top”

Acknowledgement The authors thank Institut Photovoltaïque d'Ile-de-France (IPVF) under framework Project E-3 for financial support. We also acknowledge ANR “Investissement d’Avenir” program (TEMPOS project no. ANR-10-EQPX-50) for having funded the acquisition of the NANOTEM platform (Dual beam FIB-FEG FEI SCIOS system and TEM-STEM FEI Titan Themis) used in this work.

“The authors declare no competing financial interest.”

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(10) Wang, Z. M.; Wang, J. Y.; Jeurgens, L. P. H.; Mittemeijer, E. J. Phys. Rev. B 2008, 77, 045424. (11) Sarikov, A.; Schneider, J.; Berghold, J.; Muske, M.; Sieber, I.; Gall, S.; Fuhs, W. J. Appl. Phys. 2010, 107, 114318. (12) Wang, Z. M.; Wang, J. Y.; Jeurgens, L. P. H.; Mittemeijer, E. J. Scr. Mater. 2006, 55, 987–990. (13) Kurosawa, M.; Toko, K.; Kawabata, N.; Sadoh, T.; Miyao, M. Solid-State Electron. 2011, 60, 7–12. (14) Kurosawa, M.; Kawabata, N.; Sadoh, T.; Miyao, M. ECS J. Solid State Sci. Technol. 2012, 1, 144–147.


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For Table of Contents Use Only In situ optical monitoring of new pathways in the metal induced crystallization of amorphous Ge D. Pelati, G. Patriarche, O., Mauguin, L. Largeau, F. Brisset, F. Glas, F. Oehler*

High-resolution optical microscopy is used to characterize in situ the Al-induced crystallization of amorphous Ge on glass. Combined with ex situ measurements (TEM, GIRXD, EBSD), we identify two crystallization mechanisms with separate kinetics and spatial extension: the classical ALILE process, associated to a (111)-oriented Ge layer, and a new AIC process which produces a double Ge layer in which the top layer remains amorphous and the bottom layer is crystalline with mixed (100) and (111) orientations.


29

CCD camera

Crystal Growth & Design Page 30 of 39 Motor Optical tube

1 2 light Objective 3 Window 4 Thermocouple Sample Heater 5 N26 7 8 9 ACS Paragon Plus Environment 10 11

interlayer (a-Ge)

Page 31 of 39 Crystal Growth & Design

underlayer (a-Ge) 1 2 3 4 5 6 7 8 9 10 11 12 13

SiO2

AlGeOx

Al

Al

SiO2

SiO2

deposition

c-Ge SiO2

Al

a-Ge

c-Ge

Al

SiO

SiO2

ACS Paragon Plus 2Environment

selective etching

annealing

(b) 40 μm Crystal Growth & Design Page 32 of 39

(a)

1 2 3 t=0.05 4 5 (c) 6 7 8 9 10 11 t=0.69 12 13 (e) 14 15 16 17 18 19 t=5.42 20 21 (g) 22 23 24 25 26 27 28 t=8.75 29 30 (i) 31 32 33 34 35 36

h

t=0.19 h (d)

(i) h

t=2.90 h (f) (j)

t=6.12 h

h

(h)

h

t=13.6 h

10 μm

(j)


10 μm

(a)

Page 33 of 39 Crystal C Growth & Design

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 (b) 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30

B A

ABC C C

B

B A

65 60 55 50 45 40 35 30 25 20 15 10 5 0

A

L1

L1

L2

L1 L2 ACS Paragon Plus Environment 5

6

7

8

9

10

11

12

Log Intensity (a. u.)

44 0

Page 34 of 39

42 2 51 1

40 0 33 1

31 1

CrystalBGrowth & DesignC

22 0

1 2 3 4 (b) 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23

A

11 1

(a)

(A) (B) Al

(C) (R) ACS Paragon Plus Environment

2θχ (°)

(a) (b) Page 35 of 39 Crystal Growth & Design

1 2 3 4 5 10 6 μm 7 8 9 10 {100} 11

10 μm (a)

{111}

ACS Paragon Plus Environment c-Ge

{101}

substrate

(b)

(a)

R/M2 1 2 3 4 20 5 6 7 8 9 R10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30

Page 36 of 39

M2/M1

M2/M1 μm

R/M2

M2/M1 (c)

i

(b)

R/M2 Crystal Growth & Design

M2

M2

iI

(f) M1

(d)

(g)

(e)

(h)


Ge (a) Al Ge Page 37 of 39 Crystal Growth & Design

Al

substrate

1 2 3 μm 10 4 5 6 (b) 7 8 9 10 11 12 13μm 10 14 15 (c) 16 17 18 19 20 21 22 23 24 15 μm 25

Al

{111} {100}

Al Ge substrate

Al

Al Ge substrate

(d)

{111} {100}

{101}


15 μm

{101}

AlGeOx

1 2

a-Ge Al

Al

a-Ge Crystal Growth & Design Page 38 of 39 c-Ge

Al

c-Ge


SiO2

ALILE / MIC-2

SiO2

MIC-1

Al

Al

a-Ge

Page Crystal ofc-Ge 39 Growth & Design c-Ge39 ACS Paragon Plus Environment 1 30 μm 100 2 nm

In Situ Optical Monitoring of New Pathways in the Metal-Induced

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