Eloïse Rahier

Sijia Ran

Nicolas Ratel Ramond

Shuangying Ma

Lionel Calmels

Sabyasachi Saha

Cristian Mocuta

Daniel Benoit

Yannick Le Friec

Minh Anh Luong
email: minh-anh.luong@cemes.fr

Alain Claverie
email: alain.claverie@cemes.fr

la diffusion de documents scientifiques de niveau

Keywords: phase change materials, synchrotron X-ray diffraction, Ge-rich GST, crystallization, segregation, diffusion, nucleation dominated, decomposition

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Introduction

Phase change materials (PCMs) are materials showing very different physical properties, notably optical and electrical, depending on their physical state, amorphous or crystalline [START_REF] Wuttig | Phase-Change Materials for Rewriteable Data Storage[END_REF][START_REF] Lee | Phase -Change Technology and the Future of Main Memory[END_REF][START_REF] Raoux | Phase Change Materials and Their Application to Nonvolatile Memories[END_REF][START_REF] Noé | Phase-Change Materials for Non-Volatile Memory Devices: From Technological Challenges to Materials Science Issues[END_REF][START_REF] Guo | A Review of Germanium-Antimony-Telluride Phase Change Materials for Non-Volatile Memories and Optical Modulators[END_REF] . The "canonical" alloy Ge 2 Sb 2 Te 5 (GST-225) allows today the development of a new generation of electronic memories called "e-PCM" (Electronic Phase Change Memories), where the information bit is coded by two very distinct resistive states, corresponding to the amorphous state (high resistivi ty) and the crystalline state (low resistivity) [START_REF] Bruns | Nanosecond Switching in GeTe Phase Change Memory Cells[END_REF][START_REF] Loke | Breaking the Speed Limits of Phase-Change Memory[END_REF][START_REF] Lai | Current status of the phase change memory and its future[END_REF][START_REF] Arnaud | Truly Innovative 28nm FDSOI Technology for Automotive Micro-Controller Applications Embedding 16MB Phase Change Memory[END_REF][START_REF] Cappelletti | Phase Change Memory for Automotive Grade Embedded NVM Applications[END_REF] . Beyond that, the ability to quasi-continuously tune their resistance states by applying electrical pulses makes PCMs excellent candidates for playing the role of artificial synapses in brain-inspired computers, where a network of artificial neurons is used to mimic the response of biological neurons to stimuli [START_REF] Ielmini | In-Memory Computing with Resistive Switching Devices[END_REF][START_REF] Islam | Device and Materials Requirements for Neuromorphic Computing[END_REF][START_REF] Sabina | Memristive Devices for Brain-Inspired Computing[END_REF] . However, the low crystallization temperature of GST-225 (between 120 and 180°C) is at the origin of the low thermal stability of the different electrical states offered by the corresponding devices and is thus an obstacle to many possible applications, especially in the field of embedded technologies where operating temperatures are well above 200°C. To improve performances and offer stable characteristics at higher temperatures, at least for digital devices, the enrichment of canonical GST-225 with Germanium (Ge-rich GST or GGST) is one of the solutions promoted by the industry [START_REF] Cappelletti | Phase Change Memory for Automotive Grade Embedded NVM Applications[END_REF][START_REF] Zuliani | Overcoming Temperature Limitations in Phase Change Memories with Optimized GexSbyTez[END_REF][START_REF] Kiouseloglou | A Novel Programming Technique to Boost Low-Resistance State Performance in Ge-Rich GST Phase Change Memory[END_REF][START_REF] Palumbo | Forming Operation in Ge -Rich GexSbyTez Phase Change Memories[END_REF][START_REF] Lee | Composition Segregation of Ge-Rich GST and Its Effect on Reliability[END_REF][START_REF] Redaelli | Material and Process Engineering Challenges in Ge -Rich GST for Embedded PCM[END_REF] . Indeed, GGST has a much higher crystallization temperature (> 300°C), which guarantees the integrity of the encoded information at high temperature Still, one key and limiting characteristic of PCMs is the thermal stability of the amorphous phase, which is associated with the high resistivity state of the devices. While being fundamental to device operation and understanding, the mechanisms by which such alloys crystallize and switch from one state to the other are still controversial. When this class of material was introduced [START_REF] Cheng | A Thermally Robust Phase Change Memory by Engineering the Ge/N Concentration in (Ge, N)XSbyTez Phase Change Material[END_REF][START_REF] Cheng | A High Performance Phase Change Memory with Fast Switching Speed and High Temperature Retention by Engineering the GexSbyTez Phase Change Material[END_REF] , it was thought that the different properties of GGST were resulting from a single but unknown "golden", Ge-rich, crystalline phase. Following this initial belief, recent theoretical papers have investigated the relative stability of different crystalline GST phases, in an attempt to find possible decomposition paths from "very Gerich" to "less Ge-rich" GGST alloys during annealing [START_REF] El Kheir | Density Functional Simulations of Decomposition Pathways of Ge-Rich GeSbTe Alloys for Phase Change Memories[END_REF][START_REF] El Kheir | High-Throughput Calculations on the Decomposition Reactions of Off-Stoichiometry GeSbTe Alloys for Embedded Memories[END_REF][START_REF] Cecchi | Crystallization and Electrical Properties of Ge -Rich GeSbTe Alloys[END_REF] . However, one may question the existence of many of the considered alloys (e.g., Ge 18 Sb 2 Te 5 , Ge 5 Sb 2 Te 3 , Ge 3 Sb 2 Te 3 or Ge 2 Sb 2 Te 1 ), as there is no experimental evidence that GST crystalline phases can accommodate Ge concentrations larger than that of Te in a single unit cell. In fact, the existence of such exotic Ge-rich GST alloys is usually incorrectly claimed from the chemical analysis of relatively large regions, by Energy-dispersive X-ray spectroscopy (EDX) or Electron energy loss spectroscopy (EELS), where the composition of different small, chemically different, grains can be averaged along the thickness of the sample. Alternatively, there are experimental evidences that the overall crystallization of such Ge-rich amorphous layers results in the formation of a two-phase material, where both the pure Ge and GST-225 crystalline phases coexist [START_REF] Zuliani | Overcoming Temperature Limitations in Phase Change Memories with Optimized GexSbyTez[END_REF][START_REF] Privitera | Crystallization Properties of Melt-Quenched Ge-Rich GeSbTe Thin Films for Phase Change Memory Applications[END_REF][START_REF] Agati | Chemical Phase Segregation during the Crystallization of Ge-Rich GeSbTe Alloys[END_REF][START_REF] Navarro | Trade-off between SET and Data Retention Performance Thanks to Innovative Materials for Phase-Change Memory[END_REF][START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF] . Still, the details of this crystallization process remain unclear, as well as the evolutions of the microstructure and local composition of the GGST layer, which are responsible for the observed resistivity changes. A clear understanding of these two characteristics of the GGST material is however required, if we want to be able to finely tune the resistivity of the PCM device by continuously switching the GGST material in intermediate resistivity states, as needed for advanced and neuromorphic applications.

For some authors Ge crystallizes first [START_REF] Privitera | Crystallization Properties of Melt-Quenched Ge-Rich GeSbTe Thin Films for Phase Change Memory Applications[END_REF][START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF] while for the others crystallisation starts with the formation of GST grains [START_REF] Prazakova | Temperature Driven Structural Evolution of Ge-Rich GeSbTe Alloys and Role of N-Doping[END_REF] . Moreover, the observation of the crystallization of the Ge cubic phase at temperatures as low as 310-330°C, well below the temperature at which the homogeneous crystallisation of Ge is known to occur (400°C), is puzzling. Actually, this controversy and the lack of understanding of the details of the crystallization scenario result from the difficulty to infer the mechanisms leading to the overall crystallization of GGST layers from the sole characterization of samples ex situ annealed at different temperatures and durations. Clearly, a detailed characterization of the mechanisms at work during isothermal annealing and of their relative kinetics is missing.

In this work, we report the results of our investigations of the crystallization of Ge-rich GST alloys using isothermal annealing during in situ X-ray diffraction (XRD) performed at SOLEIL synchrotron (France). These results are complemented by transmission electron microscopy (TEM) observations. We show that, prior to crystallization, the initially homogeneous material undergoes phase separation in the amorphous phase. After some incubation time, a transient GeTe phase appears which triggers the heterogeneous nucleation of nanocrystals of the Ge cubic phase. Then, this Ge crystalline phase grows through the addition of new small grains while the Ge content in the remaining amorphous GST matrix concomitantly decreases until regions finally matching the GST-225 stoichiometry are generated and spontaneously crystallize. The transient phase responsible for the observed low temperature crystallization of Ge is formally identified as the Pnma phase of GeTe. Density functional theory (DFT) calculations attest that this phase is more stable than the common R3m phase for very small size crystals. These results unveil the intriguing characteristics of the three-step crystallization of Ge-rich alloys and close a long lasting controversy concerning the order and the scenario of crystallization in these alloys.

Samples and experimental methods

Samples used in this study consisted in 100 nm-thick GGST layers deposited by physical vapor deposition (PVD) using only one target of the same GGST composition onto naturally oxidized Si(100) wafers and covered by a 20 nm thick TiN cap layer using industrial tools. In these samples the Ge concentration is such that after phase separation and crystallization, the Ge and GST-225 phases weight each about 50% of the total volume. Diffraction experiments have been performed in situ and under N 2 gas flow, using an Anton Paar DHS 1100 furnace available on the Diffabs beamline at the synchrotron SOLEIL. The X-ray energy was monochromatized to 10975 eV using a Si(111) double crystal monochromator (E/E ~10 -4 ). The photon beam was focussed vertically (using Rh-coated mirrors) and horizontally (sagittal bended second crystal of the monochromator) into a spot of ~ 250 × 250 µm [START_REF] Lee | Phase -Change Technology and the Future of Main Memory[END_REF] (full width at half maximum, FWHM). The incidence angle of the X-ray beam on the specimen was adjusted to 10°, corresponding to the value at which the diffracted signal was optimized and free from artefacts. The diffracted signals were acquired through the acquisition of two CirPAD detector frames at two different angular positions, in order to fill the gap between detector modules [START_REF] Desjardins | The CirPAD, a Circular 1.4 M Hybrid Pixel Detector Dedicated to X-Ray Diffraction Measurements at Synchrotron SOLEIL[END_REF] . One single XRD profile was acquired every 30 seconds. Quantitative information was extracted from the XRD profiles by Rietveld refinement. Scanning transmission electron microscopy (STEM) using high angle annular dark field imaging (HAADF) was used to check material inhomogeneity through Z-contrast imaging in a probe corrected JEOL JEM ARM 200F. As our previous ex situ XRD studies have shown that the incubation time, i.e. the time after which the onset of crystallization is observed, is of about several tens of minutes at 330°C [START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF] , we have explored the 310-330°C temperature range in detail, in an effort to slow down and decompose the different mechanisms which intervene in the crystallization process.

Results and determination of the crystallization scenario

Figure 1 is a map showing the time-evolution of the XRD profiles during in situ annealing of the GGST specimen at 310°C. During the first 7 hours, the GGST layer stays in the amorphous phase as evidenced by the featureless shape of the diffused profile. After 9 hours, two main phases are present which are easily identified as the Ge cubic and GST-225 cubic phases, as previously observed after annealing at higher temperatures [START_REF] Zuliani | Overcoming Temperature Limitations in Phase Change Memories with Optimized GexSbyTez[END_REF][START_REF] Privitera | Crystallization Properties of Melt-Quenched Ge-Rich GeSbTe Thin Films for Phase Change Memory Applications[END_REF][START_REF] Agati | Chemical Phase Segregation during the Crystallization of Ge-Rich GeSbTe Alloys[END_REF][START_REF] Navarro | Trade-off between SET and Data Retention Performance Thanks to Innovative Materials for Phase-Change Memory[END_REF][START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF] . Since we are looking for the chain of events leading to the crystallization of these phases, the evolution of the diffraction profiles during this lap time is of interest. Figure 2 shows the variations of the integrated intensities of the two main characteristic peaks of the Ge and GST-225 phases during isothermal annealing.

Fig. 2: Variations of the integrated diffracted intensities of the two main peaks characteristic of the Ge and GST-225 cubic phases during annealing. The GST peaks rise about 20-30 minutes after the Ge peaks have almost reached their maximum values.

This result evidences that the Ge peaks do appear before those characteristics of the GST-225 phase.

The diffracted intensity of the Ge peaks dramatically increases during the first 30 minutes reaching about 70-80 % of the final values before the GST phase starts to nucleate. This is the first clear experimental evidence that GGST alloys crystallize by first nucleating the Ge phase and that the GST-225 phase crystallizes only after a large proportion of the Ge phase has already crystallized. Moreover, it should be noted that, in comparison to Ge, the GST-225 phase builds up progressively, more slowly, needing up to 90 min to diffract with an intensity of about 70 % of the final values.

Figure 3 shows the time-evolution of the mean size of the Ge and GST-225 crystallites during annealing, as extracted from the widths of the diffraction peaks and after Rietveld refinement. This figure shows that the size of the Ge crystallites remains quite constant over time, with small variations between 4 and 6 nm. Combining this information with the dramatic intensity variations observed in Figure 2, we can deduce that the crystallization of the Ge phase is nucleation-dominated, i.e. proceeds through the accumulation of new small crystallites all of about the same size. The small size increase noted from 9 to 11 hours of annealing probably results from a limited growth of existing crystals at the end of the phase separation process. While nucleated later, the GST crystallites are relatively much bigger, even at the onset of crystallization, in the 15-18 nm range. They do not evolve too much during annealing. At first, they tend to decrease in size during the "nucleation period" lasting 90-120 minutes and during which the corresponding diffracted intensities increase (Fig. 2). Then, they increase again for longer annealing time, probably as a result the slight growth of the existing grains, again at the end of the phase separation process. It is important to note that these sizes are surprisingly large and well above any reported crystallite sizes observed during the nucleation-limited crystallization of pure GST-225 [START_REF] Luong | Effect of Nitrogen Doping on the Crystallization Kinetics of Ge2Sb2Te5[END_REF][START_REF] Kooi | Situ Transmission Electron Microscopy Study of the Crystallization of Ge2Sb2Te5[END_REF] .

To set up and validate a scenario able to explain all the observed characteristics of the crystallization of GGST alloys, several questions must be answered. First of all, why is Ge crystallization observed as such a low temperature? The homogeneous crystallization of pure amorphous Ge is usually observed above 400°C whilst there are reports showing this mechanism can be activated at temperatures down to 380°C [START_REF] Germain | Crystallization in Amorphous Germanium[END_REF] . However, heterogeneous crystallization of Ge may occur at much lower te mperatures. As an extremely favourable case, the solid phase epitaxy (SPE) of Ge, i.e. the crystallization of amorphous Ge in contact with a Ge crystalline (001) surface, can be activated at temperatures down to 310°C, still using a reasonable timescale [START_REF] Claverie | Recrystallization and End of Range Defects in Germanium[END_REF] .

Coming back to Figure 1 and focussing on the onset of crystallization, one can note that there are a number of distinct features in the XRD profiles which appear slightly before the appearance of the peaks corresponding to the Ge phase. In particular, a peak is clearly visible on the right side of the GST peak, at about 22°. Another weaker one is seen as a shoulder of the Ge peak, at about 20°. This evidences that a transient phase, which nucleates first then quickly disappears, is at the origin of the heterogeneous nucleation of Ge at low temperature. The quite strong peak observed at about 22° has been recently detected by Navarro's group after ex-situ annealing in the 350°C range, as a shoulder on the right side of the (200) GT-225 peak [START_REF] Prazakova | Temperature Driven Structural Evolution of Ge-Rich GeSbTe Alloys and Role of N-Doping[END_REF] . However, due to the lack of complementary peaks, the authors could not formally identify the phase responsible for this feature. Instead, they proposed that a "G + GST" unknown phase forms and initiates the crystallisation of the GST-225 phase. The sensitivity of the synchrotron XRD allied to the acquisition of profiles all along the in-situ isothermal annealing allows us to clarify this point. The broadening of the diffraction peaks corresponding to the GeTe phase is similar to that of the Ge phase. From this qualitative observation, we can infer that GeTe particles sizes are in the range of few nms. Figure 4 shows the XRD profile obtained after 30 000 s annealing i.e., during the short period when all the three phases are present. Beyond the peak at 22°, six additional peaks could be evidenced, rendering phase identification unambiguous. The transient phase, the first phase to appear at the onset of crystallisation, is identified as the dense GeTe phase (PDF card 00-044-0819), sometimes referred as type II GeTe, of orthorhombic structure (space group Pnma), and known to be the stable bulk phase under high pressure conditions [START_REF] Shimada | Crystallization of Amorphous Germanium Telluride under Pressure[END_REF] .

Fig. 4: XRD profile recorded after 30 000 s annealing at 310°C. The three phases are observed. The GeTe Pnma phase is formally identified through its seven characteristic peaks.

While the exhaustive study of the relative stability of facetted nanocrystals goes far beyond the scope of this paper, our DFT calculations show that not only pressure but size and termination may stabilize the orthorhombic Pnma structure at the expense of the rhombohedric R3m structure. Indeed, we have calculated the ground state energy of several slabs of the R3m and Pnma GeTe phases (with the densities of 49.75 and 52.94 Å 3 /formula unit (f.u.) respectively computed for the bulk phases), terminated by different surfaces and as function of the slab thickness. For that, we have used the firstprinciples software VASP [START_REF] Kresse | Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set[END_REF] (exchange-correlation functional PBE0sol 36 , cut-off energy of 500 eV, kmesh corresponding to 𝛿𝑘 𝑥 = 𝛿𝑘 𝑦 = 0.188 Å -1 , ground state energy convergence below 10 -5 eV/f.u., force convergence of 10 -3 eV Å -1 , GeTe surfaces of periodically repeated slabs separated by 20 Å of vacuum). For each of the GeTe phases and crystal orientations that we considered, we found that the ground state energy 𝐸 𝑠𝑙𝑎𝑏 depends linearly on the slab thickness. This dependence can be modelled by the simple equation 𝐸 𝑠𝑙𝑎𝑏 = 2𝐸 𝑠𝑢𝑟𝑓 + 𝑛𝐸 𝑏𝑢𝑙𝑘 , where 𝐸 𝑠𝑢𝑟𝑓 and 𝐸 𝑏𝑢𝑙𝑘 are the surface and bulk energy, respectively per GeTe surface and bulk formula unit and 𝑛 is the number of GeTe atomic layers stacked in the slab.

Figure 5 shows the energy differences between R3m and Pnma GeTe stacks terminated by different surfaces, as a function of the number of GeTe atomic layers stacked in the slabs. This figure suggests that the (001) terminated Pnma phase can be more stable than the R3m phase when less than 10 GeTe layers are involved, i.e. for crystallites thinner than 3 nm. This is due to the lower energy of the GeTe (001) Pnma surfaces, the presence of which counterbalances the lowest energy of the bulk R3m phase.

To check whether a modification of the inter-bilayer bonds of the slabs also contribute to the results shown in Fig. 5 beside the contribution of the surface energy [START_REF] Ronneberger | Changes of Structure and Bonding with Thickness in Chalcogenide Thin Films[END_REF] , we have studied the variations of the inter-bilayer spacing for all the slabs considered in this figure. Noticeable variations of the inter-bilayer distance only occur in the vicinity of the first 3-4 layers the closest from the surfaces, which means that mainly surface effects (including atomic layer relaxation near the surface) contribute to the energy difference shown in Figure 5. We thus believe that the lower energy of the Pnma phase is the driving force for the stabilization of thin nanocrystals of this phase; of course, a more rigorous calculation would have required to compute the interface energy betwee n thin GeTe slabs and an amorphous phase of the Ge-rich GST alloys, instead of the free surfaces that we considered here . Finally, the scenario suggested by the evolution of the XRD profiles must be completed by information on the state of the material during the "incubation time" which lasts for more than 7 hours before the onset of crystallisation occurs. Figure 6 compares STEM-HAADF images of the material after deposition and after 4 hours annealing at 310°C. While the image of the as-deposited material is uniform in contrast, the annealed layer shows contrast inhomogeneity which can be ascribed to the decomposition of the layer into Ge-rich and Ge poor-regions. In other words, the material is undergoing progressive phase separation during this lap time, the matrix tempting to expulse Ge and forcing it to segregate. Interestingly, the Ge-rich regions have a diameter of a few nanometers. Putting all these results together, a scenario able to describe all the characteristics of the crystallisation of GGST alloys can now be proposed. It is illustrated in figure 7. This diagram shows the evolution of the system during thermal annealing and leading to the crystallization of a GGST alloy. In the amorphous state, the material is relatively homogeneous in composition but in a metastable state. As soon as sufficient thermal energy is supplied, the system decomposes by expelling and segregating the Ge and depleting the rest of the matrix. Compositional fluctuations then appear at the nanometric scale. The kinetics of the phenomenon is only limited by the capacity of the atomic Ge to diffuse in the surrounding matrix, amorphous and still rich in Ge.

As the phase separation intensifies during the incubation time, the compositional fluctuations and the low solubility of Sb in Ge [START_REF] Okamoto | Germanium-Antimony)[END_REF][START_REF] Bordas | Phase Diagram of the Ternary System Ge-Sb-Te[END_REF] allow the formation of increasingly Ge-rich regions, until they approach the GeTe stoichiometry. As this separation is activated above 300°C, i.e. above the crystallization temperature of GeTe (around 230°C), thin GeTe grains can crystallize with the Pnma phase as soon as they are formed. Upon contact with areas richer in Ge, these GeTe crystallites trigger the heterogeneous crystallization of Ge from 310°C onwards. During annealing, the Ge phase develops, via the nucleation of new small grains and, as a consequence, the GGST matrix becomes "purified", being progressively depleted in Ge, while remaining in the amorphous phase. When the Ge content is such that the stoichiometry is locally close to that of GST-225, these regions spontaneously crystallize in the cubic GST-225 phase since the annealing temperature is much higher than the crystallization temperature of this phase (120-180°C).

This scenario well explains why GST-225 crystals appear later and are larger than observed during the crystallization of pure GST-225 [START_REF] Luong | Effect of Nitrogen Doping on the Crystallization Kinetics of Ge2Sb2Te5[END_REF][START_REF] Kooi | Situ Transmission Electron Microscopy Study of the Crystallization of Ge2Sb2Te5[END_REF] . In a GGST alloy, the crystallization of GST-225 is obtained only when very depleted Ge regions appear, i.e. when the phase separation is well advanced, which is finally obtained within larger domains. Thus, the crystallization of the GST-225 phase does not solely depend on temperature but also on a diffusion length L, i.e. on a characteristic time that depends on the temperature (L=Dt with D=D 0 exp(-E/kT)), this is the time that the excess Ge needs to reach the (crystalline or not) Ge phase.

Finally, our findings suggest that the change of resistivity of a GGST cell submitted to some thermal processing should probably not be solely ascribed to the reduction of the size of the amorphous GGST region in the device, but to the progressive decomposition of the GGST layer into two phases, leading to the possibility of charge carrier percolation through this composite material.

Conclusions

In summary, our XRD and TEM results unambiguously show that the crystallization of GGST involves its decomposition/separation into "pure" phases that crystallize successively through three steps, GeTe (fugitive), cubic Ge and finally cubic GST-225. This phase separation can be thermally activated above 310°C. The crystallization of GeTe is homogeneous and results from compositional fluctuations generated by the separation of Sb and Ge (non-miscible). These crystal seeds allow the heterogeneous crystallization of Ge at low temperature (from 310°C). The growth of the Ge phase, through the generation of small grains with a diameter of about 3-5 nm, allows the dissolution of the GeTe grains and the purification of the surviving amorphous matrix until its stoichiometry goes down to that of GST-225 where it can crystallize as large regions (> 15 nm). These phenomena are limited by the diffusion of Ge within the amorphous GGST matrix during its depletion in Ge.

Fig. 1 :

 1 Fig. 1: XRD map obtained during in situ isothermal annealing at 310°C of an amorphous GGST layer in the synchrotron (E=10975 eV). The Ge cubic phase appears after an incubation time of nearly 7 hours, followed by the GST-225 cubic phase several minutes later. A transient phase is observed to nucleate first, just before the onset of Ge crystallization. The vertical lines show the structure factors of the GeTe Pnma structure (PDF card 00-044-0819). The blue horizontal between 41000 and 43000 seconds is due to a momentary beam loss in the synchrotron ring.

Fig. 3 :

 3 Fig. 3: Variations of the Ge and GST crystallite size during annealing. The GST grains are about 3 times larger than the Ge grains.

Fig. 5 :

 5 Fig. 5: Energy difference 𝐸 𝑠𝑙𝑎𝑏 -𝐸 𝑠𝑙𝑎𝑏 {𝑃𝑛𝑚𝑎, (001)}, versus the number of GeTe atomic layers in the GeTe slabs. For thicknesses smaller than 10 GeTe layers, the Pnma phase is the most stable. Dark lines are guides to the eyes.

Fig. 6 :

 6 Fig.6: STEM-HAADF images of the material, as deposited (left) and after 4 hours at 310°C (right). The dark regions seen on the image of the annealed sample are of lower "mean Z", probably due to their enrichment in Ge. These contrasts are not seen before annealing.

Fig. 7 :

 7 Fig.7: Evolution of the GGST during thermal annealing. On the abscissa, the thermal energy given to the GGST material; on the ordinate, the phase composition. Circles without/with a star respectively indicate amorphous/crystalline phases. During phase separation, regions of GeTe and Ge stoichiometries are generated. In parallel, the amorphous matrix gets depleted in Ge. The crystallization of GeTe is observed after 7 hours at 310°C and triggers that of Ge through heterogeneous nucleation. The crystallization of GST-225 is observed at the end of the phase separation, when the amorphous matrix goes down to the required GST-225 stoichiometry.

Acknowledgements:

This work is part of the "Ô-GST Project" and partially funded by the nano2022 (IPCEI) initiative. It was granted access to the HPC resources of CALMIP supercomputing center under allocation 2020/2021-[P17025]. The authors acknowledge the synchrotron SOLEIL for allocating the beam time, and more particularly the support group of Diffabs beamline for keeping the beamline up and running .

Declaration of Competing Interest

The authors declare that they have no known competing financial interests.

Authors


C P Saini

S Bhowmick

A Barman

N Kumar

A Das

S A Khan

A Claverie

D Kanjilal

R N Mahato

K Singh

A Kanjilal

Defect engineered blue photoluminescence in ZnO:Al/TiO2 heterostructures

Keywords: Ion irradiation, transition metal oxides, heterostructure, photoluminescence, defects

Tailoring the blue photoluminescence (PL) in Al-doped ZnO (AZO)/TiO2 heterostructures is demonstrated by controlled induction of shallow defect centres by 50 keV Ar + -ions. This is established by a combination of temperature dependent PL and electron paramagnetic resonance spectroscopy. The dominant blue-violet PL in as-grown sample comprises of a near band-edge emission, along with a peak associated with radiative recombination of the electrons in shallow donor levels (Zn interstitials) and the holes from the valence band. However, the evolution of an additional yellow-green PL band at a fluence of 110 15 ions/cm 2 is governed by the deep donor levels, particularly ionized oxygen vacancies. The irradiation at 110 16 ions/cm 2 further leads the formation of Zn vacancies (shallow acceptors) owing to the development of O-rich surface. The structural modifications of these samples have been investigated by FE-SEM, TEM, and Rutherford back-scattering. While small micro-cracks are found at a fluence of 210 16 ions/cm 2 , the formation of graded-layers is obtained at the highest fluence of 510 16 ions/cm 2 owing to 2 ballistic intermixing and diffusion of the constituents. Detailed investigation suggests that a significant amount of Ti atoms is diffused in AZO by complete deterioration of the AZO/TiO2 matrix at the highest fluence.

INTRODUCTION

Successful fabrication of low dimensional transition metal oxide devices involves in-depth knowledge of defects, especially oxygen vacancies (Vo). Among various metal oxides, both Aldoped ZnO (AZO) and titanium oxide (TiO2) are attractive candidates for exploring Vo for wide range of technological applications, such as light emitting diodes, [1][2][3] solar cells, [4][5][6] memory storage, 7,8 etc. In particular, the growth of an AZO film on TiO2 leads to the formation of a type-II heterostructure owing to their close matching of bandgaps with an extraordinary charge separation capability by lowering recombination rate. 3,9 This bilayer structure is now considered to be promising even in the field of water splitting 10 and photocatalysis, [START_REF] Zhang | Materials Science Forum[END_REF] though the charge transport is significantly influenced by the point defects, especially at their interface. The chemical stability and the tolerance to nonstoichiometric phases of the individual material also make them flexible for tailoring interstitials/vacancies in a desired way without compromising the material properties. Previously we reported the role of Vo in nanostructured TiO2 for controlling the functionality of the resistive switching devices, [START_REF] Barman | [END_REF] and for tuning the wettability and optical bandgap. 13,14 In case of nonstoichiometric ZnO (i.e. Zn-rich or O-rich ZnO), previous theoretical investigation revealed that the formation energy of defects, calculated as a function of respective Fermi level, usually dictates them to be charged. 15,16 However, it is difficult to precisely determine the formation energy of defects as the position of Fermi level in ZnO strongly depends on the doping type and its concentration. 16 For instance, in n-type ZnO, Vo and Zn interstitial (Zni) are deep and shallow donors and are having high formation energies, whereas Zn vacancies (VZn) are deep acceptors with low formation energies and in turn act as compensating centres. 15 Oba et al. 16 on the other hand, revealed that both Vo and Zni have very low formation energies and thus serve as source of hole compensation centres in p-type ZnO, provided Fermi level is low.

In this context, photoluminescence (PL) spectroscopy is known to be a very sensitive and powerful technique to track the existence of defects in metal oxides like AZO and TiO2, and to understand the nature of sub-band gap states generated either by perturbing the crystal structure, 17 doping with foreign elements, 18,19 or by directly introducing Vo and/or interstitials in the matrix. [20][21][22] For example, Lin et al. reported an intense blue and green, and faint yelloworange luminescence band in ZnO nanostructures, and attributed this in the light of deep level defects, 20 particularly Zni, ionized Vo (𝑉 𝑜 + /𝑉 𝑜 ++ ) and oxygen interstitials (Oi). In a similar fashion, anatase dominated TiO2 nanostructures also provide a smooth PL signal, positioned at about 416 nm, where the involved mechanism is discussed in terms of the charge-transfer process from the conduction band to the localized self-trapped exciton states, in the presence of Vo. 21 Although previous studies have tried to understand the nature of defects and their corresponding impact on PL characteristics, detailed investigation of Vo and/or impurity driven localized electronic states in both TiO2 and AZO, in conjunction with in-depth understanding of the PL properties of the AZO/TiO2 heterostructure are still scarce. Using a bilayer structure adds more complexity, particularly at the interface due to the possible diffusion and intermixing of the constituent elements in each layer as a result of annealing or radiation exposure. In this respect, low-energy ion bombardment is an ubiquitous surface science technique due to its capability to control the evolution of Vo in a predetermined depth, and also their spatial distributions by proper selection of ion species, energy and fluence. Samples can be systematically examined by low temperature PL spectroscopy immediately after bombardment.

In this article, we report the impact of 50 keV Ar + ions in forming Vo in the AZO layer grown on top of a thin amorphous TiO2 layer and their possible role behind the visible range PL emission at room temperature (RT). Further, temperature dependent PL analysis has been carried out in the range of 15 to 300 K to shed light on the involvement of ion beam induced defects, especially Zni and Vo. We also show the tolerance limit before reaching the extreme condition by developing micro-cracks via formation of voids and local evolution of a graded layer by heavy mixing of the constituent elements of AZO and TiO2 films at a fluence of 510 16 ions/cm 2 . Such ion fluence dependent Vo and VZn in AZO/TiO2 heterostructures have been found

to be singly charged according to low temperature Electron Paramagnetic Resonance (EPR) study.

EXPERIMENTAL

About 10 nm thick amorphous TiO2 layer was grown at RT on ultrasonically cleaned Si substrate (area 1×1 cm 2 ) by radio frequency (RF) magnetron sputtering (Excel Instruments) using a 100 W power supply. Highly pure (99.999%) oxygen and Ar gases were injected into the vacuum chamber (base pressure 10 -7 Torr) with a flow rate of 12 sccm and 30 sccm, respectively, to maintain the working pressure of 8 mTorr during deposition from a 2 inch diameter and 6 mm thick TiO2 target (MTI corp., purity 99.99%). The substrate-to-target distance was kept at 13 cm with a substrate rotational speed of 9 rpm to have a uniform film thickness. Subsequently, about 95 nm thick AZO film was grown on TiO2 at RT from a 2 inch diameter and 6 mm thick AZO target (MTI corp., purity 99.99%) by RF magnetron sputtering using 50 W power supply. Highly pure Ar gas (99.999%) was introduced into the vacuum chamber with a flow rate of 30 sccm, while the working pressure was maintained at 5 mTorr during entire AZO deposition process. The substrate-to-target distance was kept at 8 cm and the substrate rotational speed was 12 rpm. We should note here that TiO2 and AZO have different sputtering yields, therefore growth parameters are optimized accordingly. Moreover, no post annealing or heat treatment was carried out on the AZO/TiO2 samples. Following this AZO/TiO2 heterostructures were bombarded eventually with 50 keV Ar + -ions under normal incidence at RT (schematically shown in Fig. 1), keeping a constant beam current of ~1 A and fluences in the range of (0.1-5)10 16 ions/cm 2 .

The surface morphologies of the AZO/TiO2/Si heterostructures have been studied before and after Ar + -ion irradiation using field-emission scanning electron microscopy, FE-SEM (TESCAN MIRA-II). The chemical compositions and film thicknesses were extracted by analyzing the RBS (random) spectra, where the data were recorded by using 2 MeV He + ions produced by 1.7 MV Pelletron accelerator at the Inter-University Accelerator Centre (IUAC), New Delhi. The visible range PL measurements were performed in the temperature range of 15-300 K by mounting the sample in a closed-cycle helium cryostat, while the samples were excited by a 30 mW He-Cd laser (325 nm wavelength). Moreover, ion beam irradiation induced structural evolution along with intermixing of bilayer films were examined by cross-sectional TEM (XTEM) using a Tecnai FEI (Cs corrected). Here, all TEM specimens were prepared by Focused Ion Beam (FIB Helios dual beam) after depositing a Pt layer on the sample surface in order to protect the top surface. In addition, EPR experiments were carried out at the Dresden High Magnetic Field Laboratory (Hochfeld-Magnetlabor Dresden, HLD) of Helmholtz-Zentrum Dresden-Rossendorf (HZDR) with the help of an X-band "Bruker Elexsys E500" EPR spectrometer operating in the range of 9.35 to 9.40 GHz and equipped with "Oxford

Instruments" helium-4 flow-type cryostat with a lowest accessible temperature of 2 K. The data were recorded for a microwave power of 1 mW and 5 Oe field modulation amplitude.

RESULTS AND DISCUSSION

Prior to Ar + -ion exposure, the penetration depth of ions and vacancy distribution were estimated by Stopping and Range of Ions in Matter (SRIM 2013) calculations. 23 A typical depth dependent vacancy distributions profile in AZO film is shown in Fig. 2, revealing the formation of both Zn and O vacancies within 80 nm and dominance in the range of 20-30 nm from the surface.

Typical plan-view FE-SEM image of as-grown AZO/TiO2 heterostructure is displayed in Fig.

3(a)

, showing that the AZO surface is made of small grains with an average diameter of 15 nm.

Further, 50 keV Ar + -ion bombardment leads to the gradual increase in roughness. Micro-cracks which are visible after irradiation with a fluence of 110 16 ions/cm 2 [Fig. 3(b)] are found to be transformed to a wrinkle-like pattern at a fluence of 210 16 ions/cm 2 [Fig. 3(c)]. Further increase in ion fluence up to 510 16 ions/cm 2 leads to complete deterioration of the AZO/TiO2 heterostructure owing to preferential sputtering of surface atoms followed by the formation of large voids (70-80 nm in diameter) in Si substrate, as evidenced from Fig. 3(d).

For microstructural investigation of AZO/TiO2 structure with increasing ion fluence, detailed XTEM have been carried out of some selected samples. As can be seen in fig. 4 (a), low magnification XTEM of as-grown AZO/TiO2/Si sample revealed the smooth and conformal growth of AZO film which is located in between the Si substrate and Pt layer though it's hardly shows any signature of TiO2 film owing to very low thickness. We should note here that a thin Pt layer was grown prior to TEM sample preparation for protecting the top surface of AZO. The average thickness of AZO film, measured by imageJ software using spectra is found to be around In order to study the modification of average thickness, atomic density, and chemical composition of the bilayer structure, RBS spectra were recorded before and after ion beam exposure, as exhibited in Fig. 5. The chemical composition of the films and (areal) atomic densities of the constituent elements were extracted by fitting the RBS spectra using Rump software 24 (shown in Fig. 6), where the values are summarized in Tables I andII, respectively.

As discerned from Fig. 5, the Si edge was clearly shifted towards higher channel number with increasing ion fluence and it became severe at 5×10 16 ions/cm 2 (shown by dashed arrow). Close inspection suggests a gradual reduction of the Zn and Ti peak intensities in conjunction with a shift in Ti edge towards higher channel numbers, especially above a critical fluence of 1×10 16 ions/cm 2 . The shift of Ti peak towards Zn is equally related to the increasing growth of SiOx layer on Si substrate. This could be a signature of the systematic reduction of average film thickness of AZO from its initial value of ~101 nm to ~90 nm (using average density of 7.3 × 10 22 atoms/cm 3 ) due to sputtering loss of Zn and O atoms from the surface, in parallel to interfacial mixing of the TiO2 film with increasing ion fluence up to 2×10 16 ions/cm 2 consistent with our XRTEM results. This is though not applicable for 5×10 16 ions/cm 2 fluence as the AZO/TiO2 heterostructure was found to be entirely degraded [see Fig. 4(f) for clarity]. However, this effect is not that prominent up to a fluence of 1×10 16 ions/cm 2 owing to the relatively high bonding/cohesive energy of O in TiO2 (-9.79 eV) than that of ZnO (-7.52 eV), 25,26 and also for having almost 10 times large ZnO film thickness with respect to TiO2. Indeed, the observed broadening of the Ti peak at 5×10 16 ions/cm 2 can be explained in the light of the diffusion of segregated Ti into the remaining AZO. In order to understand the phenomenon, it is now important to check the oxygen edge with increasing ion fluence, though it is clear in Fig. 6. As discerned from Fig. 6(a), the O signals from the AZO and TiO2 films are overlapped in the RBS random spectrum and thus provide a broad O peak, though the contribution is more from the AZO film because of higher thickness. We cannot however neglect the contribution of oxygen present in the natural oxide layer on Si substrate. Interestingly, the oxygen peak width is found to decrease gradually with increasing fluence, especially above 110 16 ions/cm 2 [see Figs.

6(b)-(d)]

, and this can be attributed to the decreasing thicknesses of the oxide layers by ion beam sputtering loss of atoms from the AZO surface region. Since the oxide signal in the lower channel region is mainly lost due to the increasing fluence (Fig. 5), it is clear that the oxygen from TiO2 film is equally sputtered with AZO when the energetic Ar + ions can reach there above a fluence of 110 16 ions/cm 2 . Although ion implantation leads to modification in film stoichiometry, the variation in areal density of elements before and after ion exposure is however modest (see Table II). Beside this, the signature of Ar implanted atoms has also been witnessed with increasing ion fluence (see Table II), leading to the evolution of Ar peak at 210 16 ions/cm 2 [Fig. 6(c)] and further intensified at 510 16 ions/cm 2 , as revealed from Fig. 6(d). Comparing the atomic density of Ar in the samples irradiated with fluences of 210 16 and 510 16 ions/cm 2 , one can see that almost a factor of 4 numbers of Ar atoms were lost in the former case. However, the intensity of the sharp Si edge remains unaffected.

The RBS results are found to be consistent with the corresponding cross-sectional TEM (XTEM) results showing a decrease of the initial AZO film thickness by ~10 nm at a fluence of 110 16 ions/cm 2 in the presence of Ar. 16 ions/cm 2 , followed by a complete degradation of the AZO/TiO2 heterostructure at 510 16 ions/cm 2 good agreement with our XRTEM results.

Since the fitting program RUPM involves two important parameters like the interface position and its standard deviation (), the knowledge of variance of the second component, say surface roughness (s) and the interfacial concentration gradient (i), is very important for proper understanding, details can be found in References. 27,28 In the present bilayer structure, as we have seen the evolution wrinkle-like features at surface [see Fig. 3(c)], the effect of s cannot be neglected in the present fitting process for 210 16 ions/cm 2 . This is equally applied for 510 16 ions/cm 2 due to the formation of clusters on the Si substrate with void-like structures [see Fig. II).

As can be seen from Table I, as-grown AZO film exhibits relatively higher Zn concentration than that of O, showing the stoichiometry imbalance in the as-grown sample. This stoichiometry imbalance could be due to lack of O incorporation during the growth of AZO films, leading to the formation of O deficient AZO films. On the other hand, the increase in ion fluence apparently gives rise to a decrease in overall Zn concentration, consistent with our Energy Dispersive X-ray Spectroscopy (EDS) results (not shown). This stoichiometric restoration may appear from the preferential sputtering of Zn atoms from the surface (owing to the lower surface binding energy of Zn, compared to that of O) and diffusion, and readjustment of O atoms. Further increase of fluence up to 5×10 16 ions/cm 2 causes the complete deterioration of the AZO/TiO2 interface via intermixing of constituent elements, which in turn leads to the formation of a single graded layer structures (see Table II).

Typical temperature dependent PL (15-300 K) spectra of the as-grown AZO/TiO2 heterostructure are displayed in Fig. 7(a). Since the penetration depth of He-Cd laser (325 nm wavelength) in ZnO is limited to 50-80 nm, 29 the majority of optical absorbance and consequent photoexcitation occur within the AZO layer of heterostructure. Since our XRTEM results clearly demonstrated that the underneath TiO2 layer was amorphous in nature, the observed PL mainly originates from the AZO film. It appears that the spectrum consists of at least five components (indicated by 1 to 5), and the peak intensity increases systematically with decreasing T down to 15 K, where all associated peaks are well resolved and distinguishable. Close inspection further reveals a broad visible range hump in the as-grown sample at RT, and it is attributed to the dominance of defect-related luminescent centers (LCs), especially Vo and interstitials in AZO, 30,31 though no signature of band-edge emission was recorded. By lowering the temperature, the gradual increase in PL intensity was observed, which is mainly attributed to reduction of the non-radiative de-excitation channels, frizzing-out phonons and concentration reduction of free carriers in the conduction band (CB). 32 In particular, the peak 1, detected at the emission wavelength of 378 nm (3.28 eV) was found to be blue shifted continuously with lowering T, whereas peak 2, detected at 397 nm (3.12 eV), was slightly blue shifted only at lowest T (15 K).

In fact, peak-1 is attributed to near band-edge (NBE) emission whereas Peak-2 is likely to be associated with the transition from neutral donor-bound exciton, free exciton and/or donor to acceptor pair. [33][34][35][36][37] Although no discernible change in position of peak 3 and 4 was found with decreasing T, the position of peak 5 however shifted towards higher wavelength region (see dashed line of peak 5) with decreasing T. These peaks are mainly associated with defect related LCs arises within the bandgap; details can be found in Refs. 9,30,38,[START_REF] Peng | White-light emitting ZnO-SiO2[END_REF] Here, the 3-5 PL peaks are found to be at 420 (2.95 eV), 434 (2.85 eV) and 460 nm (2.69 eV), in the blue-violet region, but their origin and mechanism are still controversial. 9,38 For instance, Zeng et al. have shown that violet emission in ZnO can be attributed to the electron transition from shallow donor states, located at 0.22 eV below the CB, to the valence band (VB) and related to the neutral Zni interstitial (Zni * ), whereas transitions from ionized Zni interstitials(Zni + and/or Zni ++ ) of slightly lower energies than Zni * (0.54 to 0.65 eV below the CB) are responsible for the blue emission. 40 In fact, the probability of formation of Zni interstitials is significantly higher in asgrown AZO film owing to have relatively high Zn concentration, in consistent with the RBS (see Table I). It is worth to note that Zni * is formed in ZnO due to the formation of a Frenkel pair through atomic collisions (𝑍𝑛 𝑍𝑛  Zni * + VZn * ) which are further ionized (Zni *  Zni + + e -and Zni +  Zni ++ + e -) and finally lead to the formation of extended Zni interstitial states. 41 In contrast, Peng et al. have shown that the band bending found in ZnO grains can result electron depletion in the regions where electron transition from the CB to defect levels located inside the depletion regions (0.7 eV above the VB) governs the blue emission. [START_REF] Peng | White-light emitting ZnO-SiO2[END_REF] A schematic representation of all these proposed defect related PL emissions in AZO is shown in Fig. 7(e).

On the other hand, the PL spectrum of the sample implanted with a fluence of 110 15 ions/cm 2 (see Figure 7(b)) is clearly different from the one obtained in as-grown sample, especially at the longer wavelength region. In particular, the broad hump which appears in the range of 450-600 nm, in the green-yellow-orange band, (indicated by an arrow on peak-3) becomes even more prominent at low T (15 K). It is attributed to the radiative recombination of charge carriers from the CB to deep acceptors like oxygen interstitials or associated to ionized Vo in AZO matrix. 20,30,31 In particular, green luminescence is attributed to the transition of electrons from neutral or singly ionized Vo (located 0.8-1.0 eV below the CB minima) to the VB edge whereas the transitions from CB to oxygen interstitials and doubly ionized Vo are responsible for yellow and orange emissions. 20,30 Lin et.al. has shown that yellow luminescence in ZnO is associated to the band-acceptor transition and further red shifted with decreasing temperature. 20 Interestingly, unlike what is observed in the as-grown sample, the NBE of AZO as shown by dashed line of peak-1 in Fig. 7(a) has completely disappeared. The quenching of such NBE is also previously reported and attributed to the defects like vacancies in Zinc sublattice which may evolve here during ion irradiation and eventually acts as trapping sites. 42 On the other hand, the LCs related violet emission, centered at 402 and 421 nm, are still visible and are indicated by peaks-1 and 2, respectively. The origin of the emission at 421 nm seems to be understood (discussed above), but the mechanism behind the PL emission found at 402 nm is still controversial, although sometimes associated with the donor-acceptor pair with longitudinal optical (LO) phonon replicas. [START_REF] Peng | White-light emitting ZnO-SiO2[END_REF] The trend was confirmed when observing the spectrum obtained on the sample bombarded with a fluence of 110 16 ions/cm 2 . This time, peak-2 (421 nm) is slightly blue shifted as evidenced in Fig. 7(c). Such blue shifting suggests some alteration of the environment or the nature of defect responsible for this emission, probably by the formation of shallow acceptors like VZn with different charge states. 40 In fact, irradiation with a fluence of 1×10 16 ions/cm 2 causes the preferential surface sputtering of the Zn atoms, leading to the reduction of the relative concentration of Zn (see Table I) and in turn promote the formation of VZn. Further increase of the ion fluence (i.e. 510 16 ions/cm 2 ) leads to some dramatic change in the PL spectrum, which can be ascribed to the strong intermixing of the AZO/TiO2/Si structure and the formation of voids as seen in Fig. 7(d). Interestingly, no PL signal was detected at RT on this sample and this could result from the dominance of phonons in this structure. In fact, it follows a similar trend in the as-grown sample below 250 K. As can be seen, peaks-1 in Fig. 7(a), which are assigned to near band-edge emission of AZO, was dramatically evolved again [see peak 1 in controversially discussed in the light of singly charged vacancies, such as 𝑉 𝑜 + and 𝑉 𝑍𝑛 -. 38,[44][45][46] In fact, 𝑉 𝑜 + and 𝑉 𝑍𝑛 -(generated by trapping of one electron from the conduction band and consequently positively charged with respect to lattice) were previously considered as shallow acceptor and deep donor defects, respectively. 38,[46][47][48] In the as-grown sample, the existence of 𝑉 𝑍𝑛 in ZnO is significantly less probable (due to relatively smaller density) owing to its larger defect formation energy (4 eV) than 𝑉 𝑜 + and therefore discarded. Previous EPR studies on ZnO have shown that 𝑉 𝑜 + has tendency to localized on the surface region of ZnO which in turn results the EPR signal at g = 2.004. 49,50 As can be seen from Fig. 7 (a), the value of g-factor is very close in all as-grown as well as irradiated samples and matches well with g-factor of free electron (ge = 2.0023), indicating the delocalization of surface defects. 51 Beside this, other defect states like doubly ionized Vo (𝑉 𝑜 ++ ), Zn and O in interstitial site (i.e., Zni and Oi 2-) could also be formed with ion irradiation and may coexist with the paramagnetic defects (𝑉 𝑍𝑛 -and/or 𝑉 𝑜 + ) found in the AZO layer. However, it is difficult to detect these defects with EPR as they are diamagnetic in nature. 38 Nevertheless, displacement of the Zn and O atoms from their regular lattice sites leads to the possible formation of singly ionized vacancies, where the concentration increases further under the exposure of energetic Ar + ions. Moreover, linewidth of the EPR signals is enhanced from 1.20 to 1.31 mT by exposing to an ion fluence of 110 16 ions/cm 2 , but is decreased to 1.14 mT at an ion fluence of 510 16 ions/cm 2 [see Fig. 8(b)]. The EPR linewidth depends generally on the spin-lattice and spin-spin interactions of unpaired electrons, while the latter one dominates more. In fact, a spin-spin interaction is related with mutual spin flips resulted by dipolar and exchange interactions between assembly of spins in the sample. 52 Here, the initial increase in EPR linewidth can be attributed to the increase in strain within the layer and to its effect on the crystalline environment of the defect centers. With further increase in ion fluence, the decrease in EPR linewidth is most likely associated with the reduction of heterogeneous distribution of 𝑉 𝑜 + centers by suppressing strain via evolution of voids as evidenced by FE-SEM and TEM results and so the effect on g tensor. 14

CONCLUSIONS

In conclusion, we report a compositional variation and the subsequent change of defects in AZO/TiO2 heterostructures by exposing to 50 keV Ar + ions. We have correlated these defects to the observed PL in the region of visible range as a function of temperature. As-grown AZO layer on TiO2/Si surface initially provides an emission peak in the blue-violet region, which is attributed to the transition of electrons from the near band edge and Zn interstitials to the valence band. An additional broad PL band in the green-yellow region (450-600 nm) is evidenced at a fluence of 110 15 ions/cm 2 owing to active participation of deep donors originated from the ionized Vo. Further increase in ion fluence led to the formation of microcracks with gradual reduction in Zn concentration via preferential surface sputtering of Zn atoms, suggesting the evolution of VZn in AZO layer. The highest fluence of 510 16 ions/cm 2 causes a complete deterioration of the AZO/TiO2 bilayer structure by intermixing of the constituent elements and evolution of voids which eventually leads to the formation of the chemically graded intermediate layer. 5.1510 16 3.4010 18 9.6310 16 3.3710 16

102

  nm. However, high resolution TEM (HRTEM) image of dashed circular region [Fig.4 (b)] not only confirms the crystalline nature of the AZO film but also clearly evident the existence of thin amorphous TiO2 film with average thickness of 10 nm. Closed inspection, further confirmed the formation of very thin (4.5  0.5 nm) native oxide SiOx layer at the TiO2/Si interface which originated due to the diffusion of oxygen radicals during the growth of TiO2 film.Moreover, the XTEM image of the AZO/TiO2 heterostructure after bombarding with a fluence of 11016 ions/cm 2 is displayed in Fig4(c), indicating a significant reduction in AZO film thickness (92 nm) due to sputtering of AZO layer. Close inspection further revealed a surprising enhancement in the underneath TiO2 thickness from 10 to 14 nm [see Fig.4(d)]. We believe that this could be due to ion beam induced swelling of the TiO2 film because of outdiffusion of oxygen atoms from the native oxide layer for reducing the difference in chemical potential at their interfaces.Further increase in ion fluence up to 5×1016 ions/cm 2 leads to complete degradation of the AZO/TiO2 matrix in conjunction with evolution of voids [see Fig.4(e)]. This is clearly visible in HRTEM spectra as shown by Fig.4(f), acquired at dashed circle in Fig.4(e), indicating the formation of bigger voids within the range of 40-50 nm along with small ones. This is consistent with our SEM results. In particular, the evolution of such bigger voids can be results the coalescence of smaller ones bombarded at a fluence of 5×1016 ions/cm 2 . In addition, ion irradiation causes the intermixing of AZO and TiO2 layers by formation of Ti, Zn and O rich graded layers and reduction in overall film thickness of 65  5 nm though its not uniform throughout the surface. [see Figs.4 (f)]. Close inspection of the spectra further revealed the evolution of amorphous layer near the surface of Si substrate, enriched with O up to a depth of 52  2 nm from the surface.

Figure 6 (

 6 b) also suggests that the Ar atoms are mostly confined within the heterostructure at this fluence. Given the existence of Ar within the Si substrate [see Figs. 6(c) and (d)] indicates a significant reduction of the combined film thickness at a fluence of 210

  3(d) for clarity].Considering the XTEM results, it is clear that the RUMP fitting in Fig.5(d) has been affected not only by the surface roughness, but also by the mixed film/Si interface, especially the formation of a thick and uneven interfacial SiOx layer beside the development of voids and rough Si surface. We believe that the increase in interfacial SiOx layer with uneven Si surface has strong impact on the Si edge beyond the critical fluence of 11016 ions/cm 2 , as evident from Fig.5. Since the fitting the ballistic mixing effect in RBS spectra are very much dictated by the atomic density, stopping power employed for Si and residual channeling in Si substrate etc. Here we have enlisted the areal atomic densities of the constituent atoms have been determined for four different samples (see Table

Fig 7 (

 7 Fig 7(d)] where the intensity of peak-2 is significantly higher, unlike as-grown one. Close

FIG. 1 FIG. 2 FIG. 3 FIG. 4 FIG. 5 FIG. 6 FIG. 7 FIG. 8

 12345678 FIG. 1 Schematic representation of the 50 keV Ar + -ion exposure to AZO/TiO2 heterostructure

Table I :

 I Ar + ion induced thickness and compositional variation of the AZO/TiO2 heterostructures, extracted by fitting the RBS spectra using of RUMP software.

	Layer	Element(s)		Atomic percentage (0.3 at.%) for samples
	composition		As-grown	Irradiated with	Irradiated with	Irradiated with
				110 16 ions/cm 2	210 16 ions/cm 2	510 16 ions/cm 2
	TiO2	Ti	50	31.6	24.2	13.4
		O	50	68.4	75.8	-
		Zn	54.5	50.5	47.4	28.9
	AZO	O	43.5	45.5	46.3	45.3
		Al	2	2	2.1	2.1
		Ar	0	2	4.2	10.3

Table II :

 II AZO film thickness and the corresponding areal density of the constituent elements with increasing ion fluence, extracted by RUMP software.

				Areal density (atoms/cm 2 )
	Samples	AZO
		Thickness
	irradiated with	(nm  3 nm)	Ti	O	Zn	Ar
	fluence
	(ions/cm 2 )
	0	101	6.0410 16	5.3210 18	4.4510 17	0
	110 16	95	5.5410 16	5.2510 18	3.9510 17	-
	210 16	90	5.3710 16	5.1210 18	2.9310 17	5.2510 15
	510 16	65 (graded
		structures)

ACKNOWLEDGEMENTS

CPS would like to acknowledge the financial support received from the Dr. D. S. Kothari Postdoctoral fellowship (F.4-2/2006(BSR)/PH/19-20/0074), UGC. AK would like to acknowledge the financial support received from Shiv Nadar University. The help received from the scientists at Inter-University Accelerator Centre is highly acknowledged. Authors are also grateful to Dr. S. Zvyagin (High magnetic field laboratory) and Dr. S. Prucnal (Institute of Ion Beam Physics and Materials Research) form HZDR for performing the EPR and low temperature PL experiments, respectively. One of authors (DK) is thankful to Department of Atomic Energy for awarding Raja Ramanna fellowship.


Minh Anh Luong
email: minh-anh.luong@cemes.fr

Sijia Ran

Mathieu Bernard

Alain Claverie
email: claverie@cemes.fr

An Experimental Study of Ge Diffusion through Ge2Sb2Te5

Ge thermal diffusion is not only the phenomenon limiting the decomposition and crystallization of Ge-rich GeSbTe alloys but is also responsible for some of the failures of the phase change memory devices using them. However, Ge diffusion in the canonical Ge2Sb2Te5 (GST-225) or in Ge-rich GeSbTe (GGST) alloys has been little studied experimentally. For these reasons, we have designed and set up a series of experiments aimed at highlighting and studying this diffusion within the solid phase, under technologically relevant conditions. For that, dedicated GST/Ge/GST structures have been grown at purpose, and the redistribution of Ge in GST-225 layers during isothermal and isochronal annealing has been studied using scanning and transmission electron microscopy based techniques. While Ge diffusion in amorphous GST-225 is negligible at low temperatures (i.e., below 140°C), it is fast in crystalline GST-225 above 220°C and proceeds via the grain boundaries. During such annealing, Ge diffusion is only limited by the emission of Ge atoms from the solid source and by the density of the grain boundaries providing diffusion paths to Ge through the polycrystalline layer, which tends to decrease as a function of the annealing time as the GST grains coalesce. These results show that diffusion phenomena can be activated at moderate temperatures in GST alloys and can change the chemical composition and morphology of Ge/GST composite materials. This supports the hypothesis that Ge diffusion may be responsible for the observed resistance drift and loss of integrity of the Phase Change Memories based on Ge-rich GeSbTe alloys in the SET state.

Introduction and Motivation

Phase-Change Memories (PCMs) offer a number of characteristics which make them suitable to replace and extend the range of applications of the current non-volatile electronic memories. [START_REF] Burr | Phase Change Memory Technology[END_REF][START_REF] Wuttig | Phase-Change Materials for Rewriteable Data Storage[END_REF][START_REF] Fong | Phase-Change Memory-Towards a Storage-Class Memory[END_REF][START_REF] Noé | Phase-Change Materials for Non-Volatile Memory Devices: From Technological Challenges to Materials Science Issues[END_REF][START_REF] Guo | A Review of Germanium-Antimony-Telluride Phase Change Materials for Non-Volatile Memories and Optical Modulators[END_REF] Beyond the possibility to scale them down and integrate them into large arrays, they show impressive data retention and endurance characteristics. [START_REF] Arnaud | Truly Innovative 28nm FDSOI Technology for Automotive Micro-Controller Applications Embedding 16MB Phase Change Memory[END_REF][START_REF] Cappelletti | Phase Change Memory for Automotive Grade Embedded NVM Applications[END_REF] The working principle of digital PCMs relies on the phase transition between an amorphous (high resistivity) state and a crystalline (low resistivity) state, which can then be associated to the binary states of an electronic memory. Switching from one state to the other is obtained by heating of a dome of a few tens nanometers in diameter with appropriate current pulses able to locally melt then quench the material or, alternatively, recrystallize it. Furthermore, intermediate resistance states can be programmed in such memory nodes so that multilevel data storage concepts are thinkable, using PCMs. [START_REF] Ielmini | In-Memory Computing with Resistive Switching Devices[END_REF][START_REF] Islam | Device and Materials Requirements for Neuromorphic Computing[END_REF][START_REF]Memristive Devices for Brain-Inspired Computing -1st Edition[END_REF] The canonical alloy Ge2Sb2Te5 (GST-225) is the most widely used active material in PCM cells. However, it is often doped with other elements to enhance the device performances. Automotive and smart card applications, for example, require alloys offering a higher thermal stability, i.e. a crystallization temperature (Tx) higher than that of GST-225 (about 170°C). This is why Ge-rich GST (GGST) alloys which offer Tx as high as 300-350°C are receiving today increasing attention. [START_REF] Zuliani | Overcoming Temperature Limitations in Phase Change Memories With Optimized GexSbyTez[END_REF][START_REF] Palumbo | Forming Operation in Ge-Rich GexSbyTez Phase Change Memories[END_REF][13][START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF][15] Different from GST-225, a congruent material which keeps a homogeneous chemical composition during phase transition, amorphous Ge-rich GST alloys (GGST) decompose during thermal annealing, and the materials ultimately crystallize by forming a mixture of Ge and GST-225 grains. 13,[START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF][START_REF] Agati | Chemical Phase Segregation during the Crystallization of Ge-Rich GeSbTe Alloys[END_REF] The crystallization kinetics, and by extension the speed of the SET operation of the PCM device, will be limited by the rate at which Ge separates from the GST alloys, i.e. on Ge diffusivity in this amorphous glass. Moreover, a number of failures affecting the characteristics of PCM devices based on GST alloys can be attributed to segregation phenomena of the chemical elements constituting these alloys or to the inhomogeneities in the stoichiometry of the phases in presence. This is the case of the "RESET stuck" of GST-225-based PCMs and of the "SET and RESET stuck" affecting GGST-based PCMs. [START_REF] Cheng | A High Performance Phase Change Memory with Fast Switching Speed and High Temperature Retention by Engineering the GexSbyTez Phase Change Material[END_REF][START_REF] Ciocchini | Modeling Resistance Instabilities of Set and Reset States in Phase Change Memory with Ge-Rich GeSbTe[END_REF][19] In both cases, it is strongly suspected that the repetition of sequences , i.e. "melt-quench-recrystallization" (of the order of 10 11 RESET/SET cycles), results in the separation of chemical species in and around the dome region of the cell. These dynamic melting and quenching effects depend on the diffusivity of the chemical species within the PCM cells, in the liquid and super-cooled liquid phases. Another phenomenon probably involves Ge diffusivity: the SET drift observed in GGST cells, partly minimised by the nitrogen doping. [START_REF] Zuliani | Overcoming Temperature Limitations in Phase Change Memories With Optimized GexSbyTez[END_REF]20,[START_REF] Prazakova | Temperature Driven Structural Evolution of Ge-Rich GeSbTe Alloys and Role of N-Doping[END_REF] It is commonly accepted that these variations in resistivity observed at temperatures ranging from room temperature to some 200°C result from Ge diffusion via these grain boundaries and from the subsequent atomic rearrangements at the grain boundaries. [START_REF] Ciocchini | Modeling Resistance Instabilities of Set and Reset States in Phase Change Memory with Ge-Rich GeSbTe[END_REF][START_REF] Mitrofanov | Ge L 3 -Edge X-Ray Absorption near-Edge Structure Study of Structural Changes Accompanying Conductivity Drift in the Amorphous Phase of Ge2Sb2Te5[END_REF][START_REF] Singh | In Situ TEM Study of Crystallization and Chemical Changes in an Oxidized Uncapped Ge2Sb2Te5 Film[END_REF] In summary, several types of diffusivity would be needed to describe and simulate the atomic mechanisms involved in the forming, SET and RESET operations of PCM cells, as well as to predict the observed degradations (drifts or failures) of PCM cells based on GST alloys, over time. Unfortunately, the literature is poor in data relevant to the diffusion of Ge in GST alloys. Firstly, we note that only one publication (Privitera et al. [START_REF] Privitera | Atomic Diffusion in Laser Irradiated Ge Rich GeSbTe Thin Films for Phase Change Memory Applications[END_REF] ) reports an (indirect) experimental estimate of the diffusivity of Ge, and only in the liquid phase. The reason for this lack of experimental results is largely due to the difficulty of evidencing and measuring the diffusion of a species within a matrix already containing at least 22% of this element. Only significant variations in composition can then be experimentally detected. On the other hand, the design and fabrication of dedicated structures and the implementation of relevant methodologies to measure, over time and as a function of temperature, the diffusion of Ge in a GST matrix are not trivial. Most of the available data therefore come from theoretical work obtained by various approaches, at fairly variable temperatures. For example, Li et al. [START_REF] Li | Atomic-Scale Observation of Carbon Distribution in High-Performance Carbon-Doped Ge2Sb2Te5 and Its Influence on Crystallization Behavior[END_REF] and Shen et al. [START_REF] Shen | Dynamic Evolution of Thermally Induced Element Distribution in Nitrogen Modified Phase Change Materials[END_REF] calculated the diffusivity of Ge in monocrystalline GST-225 at 327°C. Li obtained a slightly higher diffusivity of Sb than Ge and Te. Privitera et al. [START_REF] Privitera | Atomic Diffusion in Laser Irradiated Ge Rich GeSbTe Thin Films for Phase Change Memory Applications[END_REF] and Akola et al. [START_REF] Akola | Structural Patterns in GeSbTe Phase-Change Materials[END_REF] calculated the diffusivity of Ge in the liquid phase. Their results are consistent and in agreement with the experimental estimate and again show a slightly higher diffusivity of Sb, compared to Ge and Te. Finally, Novielli et al. [START_REF] Novielli | Atomic Migration in Phase Change Materials[END_REF] have extracted values of these diffusivities after calibration of a model describing the segregation of species in PCM cells. The analysis of these publications shows that while the diffusivity of Ge is relatively well documented in the liquid phase, the quantities necessary for the simulation of a number of technological situations are non-existent.

Experimental Challenge and Approach

Figure 1 shows the structural and chemical characteristics of a typical GGST-based PCM cell in the RESET, i.e. in its highly resistive, state. While the region far from the cell is composed of a mixture of Ge and GST-225 grains, the dome cell region is characterised by an almost pure Ge "crown" surrounding an amorphous dome much poorer in Ge than the initial GGST alloy. These compositional heterogeneities are then found, more or less pronounced, in the SET and RESET states and are characteristics of PCM cells based on GGST alloys. 13,14,16 Describing the crystallization of the amorphous dome, i.e. the SET operation and, beyond that, the thermal stability of the SET and RESET states, would require the ability to describe the diffusive behaviour of Ge in contact with and inside the GST material, both in the amorphous and polycrystalline phases. However, studying this kinetic phenomenon at this spatial scale is inaccessible. For these reasons, we have attempted to reproduce, by "planarizing" it, the situation encountered in a PCM dome cell structure in order to study in detail the behaviour, during thermal annealing, of a Ge layer in contact with a GST-225 layer, either amorphous or crystalline, as shown in Fig. 1. The structure we propose (Fig. 1b) not only allows us to "mimic" the situation encountered within a cell, but also to perform, in principle, a classical diffusion experiment from a solid source (Fig. 1c). In such an experiment, the atoms constituting the source (B) diffuse within the layers (A) and show concentration profiles as a function of distance from the source that evolve over time, typical of the diffusion mechanism activated at the temperature considered. The laws of diffusion and some mathematical manipulations allow, in principle, to deduce the diffusion coefficient at a given temperature from the fitting of concentration profiles measured over time by simple mathematical functions. [START_REF] Peterson | Grain-Boundary Diffusion in Metals[END_REF][START_REF] Sommer | Direct Determination of Grain-boundary and Dislocation Self-diffusion Coefficients in Silver from Experiments in Type-C Kinetics[END_REF][START_REF] Mehrer | Diffusion in Solids[END_REF] As diffusion is thermally activated, the measurement of diffusion coefficients obtained at different temperatures can then be used to establish a general expression for the diffusivity of the species under study. However, as it will be emphasized later, such an approach is restricted to the usual cases where the solid has a thickness larger than the diffusion length.

We have therefore fabricated such structures by physical deposition on a thermally oxidized 200 mm Si wafer. A 50 nm-thick amorphous Ge "source" layer was buried between two 115 nm thick amorphous GST-225 layers. This GST/Ge/GST stack was deposited at room temperature in an Ar atmosphere by magnetron sputtering on a Si(100) wafer of 200 mm diameter in an ultra-high vacuum industrial tool (EVATEC ClusterLine200TM) by alternating sputtering of a Ge2Sb2Te5 target and of a pure Ge target. Films were deposited on Si substrates previously cleaned by Ar sputtering. The stack was protected against oxidation by depositing in situ, immediately after fabrication, a 10 nm thick SiN capping layer. Annealing experiments were performed either in situ in a TEM, following a procedure minimizing experimental artefacts, as presented in detail by Luong et al. [START_REF] Luong | Effect of Nitrogen Doping on the Crystallization Kinetics of Ge2Sb2Te5[END_REF] , or ex situ, in a conventional annealing furnace, under nitrogen flux. XRD measurements were carried out using a Bruker D8 Discover diffractometer with a Cu -Ka1 source (λ = 1.5406 Å). Thin specimens suitable for cross-sectional transmission electron microscopy (XTEM) observations were prepared by focus ion beam (FIB) using a FEI Helios NanoLab 600 (FEI Company) operating with a 30 keV Ga ion beam, and finally polished and cleaned at 2 keV -3 pA to reduce the amount of amorphous materials on the specimen. Conventional and in-situ TEM heating were performed on a Philips CM20FEG operating at 200 kV. The EDX maps and spectra were acquired on the same microscope, which is equipped with a Microanalyzer QUANTAX XFlash detector with a 30 mm 2 active area providing an energy resolution of 127 eV. The EDX quantification was performed using Esprit 1.9 software from Bruker, employing the "Cliff and Lorimer ratio method" for quantifying constituent elements, with an absolute uncertainty of about 3% (molar fraction). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) was acquired on a probe corrected ARM JEM JEOL 200F microscope, operating at 200 kV.

Results

As-deposited GST-225/Ge/GST-225 sandwiched structure

Figure 2 shows a bright field (BF) TEM images and EDX chemical maps of the as-grown structure. The amorphousness of the GST and Ge layers is confirmed as well as the targeted 225 composition of the GST layers. While TEM imaging shows abrupt interfaces, the chemical gradients extracted from EDX extend over about 10 nm in the maps and profiles, what must be taken as the spatial resolution limit resulting from the experimental set up (electron beam probe size, specimen tilt, interfacial roughness and verticality, etc.).

Effect of temperature on Ge diffusion

Firstly, in-situ TEM heating of the specimen has been used to explore the temperatures at which different mechanisms could be activated during thermal annealing. The in-situ heating experiment was performed using a Gatan 652 double tilt heating holder which was connected to a temperature controller and acted as a furnace-type holder. During the annealing, the temperature was increased by steps of 5°C from about 100°C, using a ramping rate of 1°C/s, and held for 5 min at this temperature during which the film was imaged. These observations are summarised in Figure 3. In situ TEM heating confirms that the GST layers start to crystallize through homogeneous nucleation in the cubic GST-225 phase from about 140°C. As time passes and the temperature increases, the grains grow in size and coalesce at about 300°C when the cubic grains transit to the stable hexagonal phase. Most importantly, the thickness of the Ge layer (about 50nm) is observed to progressively decrease during annealing, notably from 250°C, and disappears after annealing at 450°C, which shows that Ge diffuses well into the GST layers at these temperatures. In order to identify more precisely the temperatures and times for which Ge diffusion could be demonstrated and studied, we "explored" the result of a panel of thermal anneals, over a wide range of annealing times and temperatures. Figure 4 shows a selection of chemical maps obtained by EDX after ex-situ annealing at different temperatures of the studied structure. BF-TEM and STEM-EDX analyses confirm that Ge diffusion can be activated above 200°C. A rapid diffusion occurs at temperatures above or equal to 300°C, for which the Ge layer disappears completely in less than an hour. Strong chemical heterogeneities can then be observed within the GST layers. It is observed that Ge first segregates at the boundaries and junctions of the GST-225 grains, and forms Ge grains. X-ray and selected area electron diffractions (XRD and SAED, respectively) show that these Ge grains develop in the amorphous state and become crystalline only for annealing temperatures equal to or higher than 300°C (shown in SI figure S2,3). Therefore, we have studied in detail the diffusion of Ge at temperatures well below 300°C, allowing the study of Ge diffusion for annealing times of the order of several hours. As a result of all these observations, two temperature ranges were identified as particularly relevant for technology and PCM cells: i) annealing at temperatures below 140°C should allow us to study the possible diffusion of Ge in the amorphous phase, and ii) annealing at 220° for increasing times (isothermal annealing), or for a given (16h) time at temperatures between 200 and 260°C (isochronal annealing) to study Ge diffusion through the GST-225 polycrystalline layer.

Diffusion in the amorphous phase: annealing at 120°C

Figure 5 shows the results obtained after thermal annealing of the structure for 7 and 13 days at 120°C. After 7 days of annealing, a few grains of cubic GST-225 embedded in the amorphous matrix can be observed. These grains grow for longer annealing times until the layer is fully crystalline after 13 days of annealing. The EDX chemical analysis of the structure does not show significant variations in the chemical composition within the layer. The small change in the concentration gradient at the Ge/GST interface can be attributed to small changes in experimental conditions and do not significant. This experiment therefore shows that Ge does not redistribute in GST-225 to any significant extent at this low temperature (120°C), both in the amorphous and crystalline phases.

Diffusion in the crystalline phase

Isothermal annealing at 220°C

Figure 6a shows TEM images of the structure after annealing at 220°C for increasing times. At this temperature, the crystallization of GST-225 is rapid (a few minutes) and the diffusion of Ge therefore takes place mainly through the GST-225 layer in the polycrystalline state. During annealing, the Ge layer becomes thinner and thinner, showing that Ge diffuses into the GST-225 layers. We have measured the variation of the width of the Ge layer over time, which characterizes the Ge "emission" towards the adjacent GST-225 layers. This emission from the source is not constant but decreases with time (shown in Fig. 6b), until it stops at long times. EDX chemical analyses of these layers (Fig. 7) reveal the resulting progressive increase in Ge concentration within the GST-225 layers. This enrichment is of the order of 10% after several days at 220°C. Although the concentration gradient at the Ge/GST interface decreases and becomes blunted over time, this reduction remains modest and is not sufficient to explain the strong slowdown of the diffusive transfer of Ge from the source to the GST layers over time. On the other hand, it is important to note that the GST-225 layers mature and that the grains become increasingly larger, almost doubling in size between the beginning and the end of the diffusion experiment (Fig. 6c). High angle annular dark field imaging using scanning TEM (HAADF-STEM), a technique where the contrast is sensitive to the atomic number of the elements, clearly shows that Ge is segregated at the boundaries between the GST-225 grains and grows there over time in the form of grains (see Figure 7). Again, SAED and XRD analyses shown in Figure S2,3 show that Ge remains in amorphous phase both in the source and at the grain boundaries. The fact that the concentration profiles remain flat during annealing evidences that the diffusion is very fast and that the diffusion length involved is very large compared to the thickness of the GST-225 layers. For this reason, classical analysis of the diffusion profiles leading to the extraction of a diffusion coefficient is impossible. It thus seems that the diffusion we are witnessing is a fast diffusion through the grain boundaries, typical of the "C-type diffusion" in Harrison's classification. [START_REF] Peterson | Grain-Boundary Diffusion in Metals[END_REF][START_REF] Sommer | Direct Determination of Grain-boundary and Dislocation Self-diffusion Coefficients in Silver from Experiments in Type-C Kinetics[END_REF][START_REF] Mehrer | Diffusion in Solids[END_REF] Its slowing down over time must be related to the growth of the GST-225 grains and the reduction of the diffusion paths offered by the grain boundaries. It is also very likely that the stability of the various grain boundaries increases over time, making the interfaces less defective and therefore less likely to support very fast diffusion.

Isochronous annealing of 16h

Figure 8 shows the chemical maps obtained by EDX on the samples after 16h annealing at temperatures between 220 and 260°C as well as the concentration profiles extracted from these maps. It can be seen that, as the temperature increases over this short temperature range, the thickness of the Ge layer dramatically decreases. As a result, the Ge concentration in the GST layers increases up to almost 20%, in excess of the regular 22% stoichiometry of the GST-225 phase. Heterogeneities of composition appear, and betray the formation of amorphous Ge grains within the GST-225 layers, at the grain boundaries. It is again important to note that no Ge gradient can be evidenced in the GST layers. The observation of some extension of the gradient over larger distances at the interface between the Ge and GST layers only results from the averaging of the composition at and close to this interface by EDX analysis, while this interface tends to become rough and lose its planarity as it shrinks. These observations show that the mechanisms allowing the emission and diffusion of Ge into the GST-225 polycrystalline layers via the grain boundaries are very strongly activated at these temperatures.

Conclusions

Studying the diffusion of an element through a compound containing that same element is complicated. In this work, we have designed a suitable structure to study the thermal diffusion of Ge through the GST-225 alloy. We have studied this diffusion in technologically relevant situations using techniques based on transmission electron microscopy. We have shown that this diffusion could not be evidenced through the amorphous phase of GST, i.e. for temperatures below 140°C beyond which GST-225 crystallizes, even for times of the order of more than ten days. On the other hand, Ge diffusion can be strongly activated at temperatures above 220°C. At these temperatures, this diffusion is very fast and takes place after the crystallization of GST-225 via the boundaries separating the crystalline grains. It follows the "type C" diffusion behaviour as depicted in Harrisson's classification. It corresponds to the situation in which the segregation coefficient is high, which is consistent with the fact that amorphous GGST alloys decompose during annealing. [START_REF] Luong | On Some Unique Specificities of Ge-Rich GeSbTe Phase-Change Material Alloys for Nonvolatile Embedded-Memory Applications[END_REF][15][START_REF] Agati | Chemical Phase Segregation during the Crystallization of Ge-Rich GeSbTe Alloys[END_REF] During annealing, Ge diffusion is only limited by the emission of Ge atoms from the solid source and by the density of the diffusion paths available through the polycrystalline solid. This density decreases with time as the GST-225 grains grow in size. However, this rapid diffusion cannot be quantified by the experiments presented in this work because the Ge concentration, while increasing with annealing time, remains independent of the distance to the source. In other words, Ge concentration gradients cannot be demonstrated because the Ge diffusion length (D.t) is much larger than the thickness of the designed GST-225 layer (115nm). However, these results show that diffusion phenomena can affect the chemical composition and morphology of Ge/GST composite materials and thus contribute to the understanding of the resistance drift and loss of integrity which affect the PCMs based on GGST alloys in the SET state, notably at temperatures above 200°C.

Figure 1 :

 1 Figure 1: a) Transmission electron microscopy (TEM) image, chemical maps and schematic illustration evidencing the typical structure and chemical distributions of Ge-Sb-Te within a GGSTN cell after the forming operation. b) scheme showing the planar sandwiched structure reproducing the interfaces of interest in a GGST cell. c) Top, classical structure used to study the diffusion of an impurity into a solid from an infinite source. Bottom, chemical profiles classically observed during diffusion experiments.

Figure 2 :

 2 Figure 2: Top left, BF-TEM image of the initial GST/Ge/GST structure. Top right, composition of the structure as measured by quantitative EDX. The composition of the amorphous GST is 225, as expected. Bottom, EDX chemical maps showing the homogeneity of Ge-Sb-Te distribution.

Figure 3 :

 3 Figure 3: A series of TEM images extracted from the movie recorded during in-situ annealing. A plot of the heating procedure is shown in the supporting information Figure S1. Further detail of the heating procedure has been presented in the Ref 32 .

Figure 4 :

 4 Figure 4: A series of chemical maps obtained by EDX showing the distribution of Ge, Sb and Te species in the structure after annealing. The GST-225 layers remain homogeneous up to temperatures of around 240°C.

Figure 5 :

 5 Figure 5: BF-TEM and SAED images of the layers, and Ge concentration profiles extracted from EDX analysis after annealing at 120°C for 7 days (a) and 13 days (b).

Figure 6 :

 6 Figure 6: a) BF-TEM images showing the progressive thinning of the Ge source layer during isothermal annealing at 220°C. b) thickness of the Ge layer that diffused into one of the GST-225 layers. c) evolution of the average grain size of GST-225 during this annealing.

Figure 7 :

 7 Figure 7: Left, time evolution of the Ge profiles during isothermal annealing at 220°C (extracted from EDX mapping). Right, HAADF STEM image taken on the sample annealed at 220°C for 81h, showing Ge segregation (in black) at the GST-225 grain boundaries and the diffusion paths.

Figure 8 :

 8 Figure 8: a) Chemical maps (EDX) of the structures annealed 16h at different temperatures and b) Geconcentration profiles extracted from these maps. At high temperature the signal appears noisy, a consequence of the compositional heterogeneities detected in the images.

Table 1

 1 summarises the values of the diffusion coefficients of Ge, Sb and Te in GST-225 alloy that can be gathered in literature.

	Authors	Method	Phase	T ( o C)	Diffusion coefficient	Ref
					(cm 2 s -1 )
	Privitera et al. Laser irradiated and	Molten phase	627	DGe/GST = 2 to 8x10 -5	Ref 24
		COMSOL
		calculations
	Akola et al.	Density functional	Molten phase	627	DGe/GST = 3.93 × 10 -5	Ref 25
		calculations			DSb/GST = 4.67 × 10 -5
					DTe/GST = 3.78 x 10 -5
	Li et al.	ab initio molecular	Crystalline	327	DGe/GST = 5.17 x 10 -6	Ref 26
		dynamics			DSb/GST = 4.28 x 10 -6
		simulations			DTe/GST = 2.35 x 10 -6
	Shen et al.	ab initio molecular	Crystalline	327	DGe/GST= 1.16 × 10 -6	Ref 27
		dynamics			DSb/GST= 1.5 × 10 -6
		simulations			DTe/GST= 1.18 × 10 -6
	Novielli et al. Finite element	Molten	620-887 Diffusivity 10 -5 to 10 -4	Ref 28
		method	Crystalline	410-620 Diffusivity 10 -9 to 10 -7
			Amorphous/	113-410 Diffusivity 10 -12 to 10 -9
			Crystalline

Table 1 :

 1 Ge-Sb-Te diffusivity values reported in literature.

Acknowledgment

This work is part of the "Ô-GST Project" and partially funded by the nano2022 (IPCEI) initiative and by the ST/CNRS Action on GST materials.

Supplementary Materials:

Figure S1 shows the typical heating profile used for in situ TEM experiment. Step jump: 5°C, ramping rate: 1°C/s. Once the set temperature is reached, 5 min of waiting time before observation. Figure S2 shows XRD patterns of samples annealed at 220°C for different time durations. The pink vertical bars represent the expected positions of GST-225 FCC diffraction peaks. The extra peaks at 38° and 44.5° are generated by the substrate on the XRD machine. Figure S3 shows XRD patterns of the as-deposited and annealed samples, at 280°C, 300 and 340°C for 1h. The Ge peak starts to appear only after annealing at 300°C/1h. Author Contributions: Investigation, writing -review and editing, Minh Anh Luong; Discussion and review the manuscript, Sijia Ran; Sample fabrication, discussion and review the manuscript, Mathieu Bernard; Project administration and funding acquisition, supervision and review manuscript, Alain Claverie. All authors have read and agreed to the published version of the manuscript.

Conflicts of Interest: The authors declare no conflict of interest.