# Thermochemical treatment of sputtered-AlN/2D MoS2 seed layers: a new elaboration process of highly c-axis AlN films

J Patouillard, R Gassilloud, Frédéric Mercier, A Dussaigne, B Hyot, M

Bernard, S Cadot, N Gauthier, N Vaxelaire, N Bernier, et al.

► To cite this version:

J Patouillard, R Gassilloud, Frédéric Mercier, A Dussaigne, B Hyot, et al. Thermochemical treatment of sputtered-AlN/2D MoS2 seed layers: a new elaboration process of highly c-axis AlN films. Journal of Physics: Materials, 2025, 8 (2), pp.025014. 10.1088/2515-7639/adc3cb hal-05396537

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# Thermochemical treatment of sputtered-AIN/2D MoS₂ seed layers: a new elaboration process of highly c-axis AIN films

To cite this article: J Patouillard et al 2025 J. Phys. Mater. 8 025014

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# Thermochemical treatment of sputtered-AIN/2D MoS₂ seed layers: a new elaboration process of highly c-axis AlN films

J Patouillard¹,²,³,* ID, R Gassilloud¹, F Mercier² ID, A Dussaigne¹, B Hyot¹ ID, M Bernard¹, S Cadot¹, N Gauthier¹, N Vaxelaire¹, N Bernier¹, F Martin¹, C Raynaud¹, F Gianesello³, A Mantoux² and E Blanquet²

¹ Université Grenoble Alpes, CEA, LETI, F-38000, Grenoble, France

² University Grenoble Alpes CNRS, Grenoble INP, SIMaP, 38000 Grenoble, France

³ STMicroelectronics, 850 rue Jean Monnet, F-38926 Crolles, France

* Author to whom any correspondence should be addressed.

E-mail: julien.patouillard2@cea.fr

**Keywords:** sputtering, MoS₂, AlN, c-axis, NH₃

Supplementary material for this article is available [online](#)

## Abstract

AlN-based acoustic filters are key devices in radio frequency communications. The AlN material crystal quality on silicon substrates limits the current performances of these devices. Atomic layer deposition -grown 2D-MoS₂ thin film can be used as a template on silicon to improve AlN crystal quality. However, after deposition of thick sputtered AlN films (≥200 nm), a systematic delamination of the AlN/MoS₂ stack appears, drastically limiting the integration of these materials in RF devices. We propose a new elaboration process to grow thick AlN films with improved crystal quality while avoiding the delamination issue. This process is based on the chemical conversion of the MoS₂ layer through a very thin 5 nm AlN seed using a 1000 °C NH₃ heat treatment at 400 mbar. We show a trade-off between the AlN thickness and the reactive annealing conditions to allow the diffusion of nitrogen towards the underlaying MoS₂ and its conversion into Mo(Oₓ)Nᵧ. Then, we observe that after the heat treatment, it is still possible to grow thick AlN epitaxial films without any delamination on the substrate while still improving the AlN crystal quality. Using this new route, a 1 µm thick sputtered AlN layer with a record AlN (002) rocking curve value of 0.71 °C on a silicon-based substrate was obtained without any cracks or delamination.

## 1. Introduction

Surface and bulk acoustic wave filters (in the so-called BAW technologies) are widely used in modern mobile telecommunication. Their operating principles are based on the propagation of surface or vertical acoustic waves within a piezoelectric material. AlN shows its highest piezoelectric response along the c-axis (d₃₃ = 5.53 pm.V⁻¹) and is used as the active layer in BAW devices. It displays an interesting tradeoff with a reasonable electromechanical coupling coefficient (k²ₐ₈) of around 7% and a high quality factor Q compared to ZnO, PZT, LiNbO₃ or LiTaO₃ [1]. Meanwhile, the rapid development of communication standards such as 5 G results in new requirements in filtering technology. The filters must operate at a higher frequency with a larger bandwidth (better k²ₐ₈). The current AlN crystalline quality of AlN-based BAW devices is insufficient to address these new requirements. The traditional route to elaborate c-axis textured AlN thin films is a low-temperature reactive sputtering process on molybdenum electrodes deposited on silicon substrates. The resulting AlN (002) rocking curve (RC) FWHMs are typically above 1 °C [2].

To address this challenge, different research axes are investigated. Doping AlN with scandium enhances the electromechanical coupling coefficient k²ₐ₈ but at the expense of the quality factor Q [3]. Another straightforward alternative is to improve AlN crystal quality using a high-temperature (≈1000 °C) deposition method such as metal organic chemical vapor deposition (MOCVD) usually using sapphire [4, 5] or silicon carbide [6, 7] substrates. However, due to the diameter size and the subsequent cost of the wafers combined with the material and these substrates, the MOCVD is not a cost-effective solution for large scale
production. Recently, with the emergence of 2D-TMDs material, in particular MoS₂ [8] and WS₂ [9], some studies have been published focusing on the use of these materials for the growth of *c*-axis textured AlN layers by magnetron sputtering [10], helicon sputtering [11, 12] or MOCVD [13]. However, to the best of our knowledge, no AlN-based RF devices using 2D-TMDs as seed layers have been reported in the literature. In addition, starting from a critical AlN thickness, the strong predominance of van der waals (VdW) interactions at the MoS₂/SiO₂ interface and between the MoS₂ atomic plans in the out-of-plane direction cannot handle anymore the AlN residual stress present in the AlN films and lead to a detrimental and systematic delamination of the AlN/MoS₂ stack [12]. This delamination issue, which is a major hindrance to the integration of these materials in RF devices, has been studied in detail. To not overload the introduction of this article, these results are presented in supplementary information (see figure S1 S3 and table S1).

In a previous work [8], we proposed an innovative material approach relying on the use of an atomic layer deposition (ALD)-grown 2D-MoS₂ nanofilm as a template to improve the crystal quality of a 100 nm sputtered AlN layer on a silicon-based substrate. The strong initial mosaïcity of MoS₂ is transmitted to the AlN layer thanks to a local VdW epitaxy of AlN on the MoS₂ grains. The 100 nm thick AlN layer exhibits a strong reduction in the AlN (002) RC of 0.5 °C compared to that directly sputtered on SiO₂ (RC of 4.0 °C). However, we observe an onset of delamination of the stack at the MoS₂/SiO₂ interface for AlN thickness above 200 nm. The VdW interactions between the MoS₂ plans and at the MoS₂/SiO₂ interface can not handle the AlN residual stress formed during the sputtering deposition of the metal nitride for thicknesses higher than 200 nm. This delamination issue is a major hindrance to the integration of these materials in RF devices.

In this article, we will first rapidly describe the delamination issue. Then, we will detail a new elaboration process of AlN films relying on a NH₃ heat treatment of a thin AlN/MoS₂ stack solving the detrimental delamination of the stack while improving the AlN *c*-axis textured growth. This AlN seed layer has two roles. First, it is thin enough to allow the diffusion of nitrogen species towards the underlying MoS₂ and its conversion into Mo(Oₓ)Nᵧ. Second, it is thick enough to ensure the epitaxial regrowth of a thick AlN layer grown on the seed.

## 2. Experimental methods

### 2.1. MoS₂ and AlN growth

MoS₂ nanofilms were deposited by ALD on Si/SiO₂ (500 nm) substrate at 90 °C using the process described by Cadot *et al* [13]. Then, a crystallisation annealing under a N₂ flow (2 l min⁻¹) is achieved at 900 °C for 30 s at atmospheric pressure. The MoS₂ preparation leads to the formation of a well-controlled uniform three monolayers of MoS₂ thin film. Thin AlN layers were subsequently deposited on top of MoS₂ with a deposition time varying between 4 s and 10.4 s by reactive sputtering in a SPTS (Trikon) Sigma fxP 200 mm magnetron physical vapour deposition. For each deposition time, the corresponding thickness has been measured by XRR and is presented in figure S4. As the layer thicknesses deposited are very thin (≤ 5 nm), we will use the deposition time instead of the real thickness in this paper. At this scale, we need to consider the substrate roughness to avoid any formation of AlN islands or uneven films. You can refer to our previous article [8], especially the cross-sectional TEM images showing the smoothness of the initial substrate. The deposition parameters were initially tuned to achieve the best texture of the AlN films on MoS₂ (3 monolayers)/SiO₂ (500 nm)/Si(100) wafers. The temperature was set to 350 °C with a standard target power of 2 kW. The argon and nitrogen flow rates (for the plasma) are set to 10 sccm and 50 sccm, respectively. As the substrate bias is suspected to degrade MoS₂, no bias was applied on the wafers to ensure soft sputtering conditions and preserve the MoS₂ surface.

### 2.2. NH₃ heat treatments

We chose NH₃ as a chemical reagent since it is prone to decompose at high temperature, forming NH radicals, which are expected to lead to nitrogen species diffusion through AlN and react with the Mo metal present in MoS₂. One must notice that N₂ is a very stable gas that does not decompose below 1000 °C and oxygen-based reagents are eliminated to avoid oxidation of the AlN seed layer forming Al–O amorphous layer. Two different equipments were used to perform the nitridation process. First, the NH₃ heat treatments were carried out on AlN/MoS₂ stacks in a home-made furnace shown in figure 1, at a temperature of 1000 °C for 60 min. In this furnace, the deposition time of the AlN layer within the AlN/MoS₂ stacks is 4 s, 5 s, 6 s, 7 s, and 8 s, corresponding to film thicknesses varying between 2 nm and 4 nm.

The temperature of the sample was monitored by an optical pyrometer pointing at the sample surface. The NH₃ gas line opened only during the 1000 °C temperature plateau with a flow rate set at 50 sccm. The NH₃ pressure was set at 50 mbar using a throttle valve. The furnace is equipped with a LayTec EpiTT system to monitor in situ reflectances with three different wavelengths at 405 nm, 633 nm and 949 nm called R405, R633 and R949, respectively. The reflectance is calculated from the ratio (value between 0 and 1) of the
<figure><img src="image_5.png" /><figcaption>**Figure 1.** Home-made furnace equipped with a LayTec EpiTT system to monitor *in situ* sample reflectance.</figcaption></figure>

reflected luminous flux to the incident luminous flux. In this work, we focus on the wavelength at 405 nm (R405). For each sample, the R405 reflectance acquisitions were recorded during the overall $NH_3$ heat treatments but only the acquisitions during the temperature plateau have been exploited.

In parallel, a nitridation process was implemented in an industrial Aixtron MOVPE reactor used as $NH_3$ furnace only (no deposition was achieved in the MOVPE reactor). This equipment enables to use higher $NH_3$ pressures and assists the $MoS_2$ nitridation as has been investigated by Hyot *et al* [12]. The temperature was set at 1000 °C for 10 min on AlN/$MoS_2$ samples with an AlN sputtering deposition time of 10.4 s (≈5 nm). The $NH_3$ flow rate and pressure were set at 3000 sccm and 400 mbar, respectively. The AlN/$MoS_2$ samples and thermal heat treatment conditions investigated in each piece of equipment are reported in Supplementary Information table S2.

## 2.3. Characterizations

Raman spectra were collected in the backscattering geometry using a Renishaw in-via spectrometer equipped with a Peltier cooled charge coupled device detector. The light was focused onto the sample surface using a x100 (0.85 numerical aperture) short working objective lens (0.27 mm). The resulting spot diameter was around 0.7 µm. A 532 nm wavelength laser diode was used as light source with a typical laser power of 1 mW. Raman spectra experimental line shapes were fitted with Gaussian–Lorentzian components on WiRE software.

The x-ray photoelectron spectroscopy (XPS) spectra were recorded with a PHI 5000 VersaProbe II spectrometer (Physical Electronics) using monochromatic Al $K_α$ radiation (1486.6 eV) and with an overall energy resolution of 0.6 eV (with silver as reference). The analysis angle was set at 45 °C. The experimental line shapes were fitted with Gaussian–Lorentzian (G-L (70–30)) components on CasaXPS software.

Out-of-plane measurements (Rocking curve, $2θ/θ$ scans and XRR) have been performed using Panalytical® Xpert 4 circles goniometer equipped with a Cu anode ($λ = 1.5406 Å$) as source. The beam is monochromatized and collimated with a multilayer curved mirror and a parallel plate collimator is mounted in the diffracted beam path.

In-plane measurements were performed on a Rigaku, SmartLab diffractometer equipped with Cu $K_α$ radiation ($λ = 1.5418 Å$). The copper anode operated at 30 kV and 15 mA.

For TEM analysis, the sample was prepared by $Ga^+$ focused ion beam milling. The thin foil was observed in STEM mode at 200 kV using a convergence semi-angle of ≈18 mrad for the incident electron probe in a probe-corrected ThermoFisher Titan Themis microscope equipped with the Super-X detector system for energy dispersive x-ray (EDX) spectrometry. EDX hypermaps were acquired with a pixel size of 50 pm, a pixel dwell time of 18 µs and a total acquisition time of ≈4 min. Hypermaps were processed in the Bruker Esprit v2.2 software using standard TEM recipes for background subtraction and Gaussian peak deconvolution. STEM High-angle annular dark-field (HAADF) images were acquired with a camera length of 86 mm corresponding to inner and outer collection angles of the HAADF detector (fischione model M3000) of ≈78 and 230 mrad. Gun lens and spot size values were selected to provide a probe current of approximately 20 pA. We paid attention to verify that there was no electron beam induced damage by comparing STEM images before and after each STEM/EDX acquisition.
<figure><img src="image_6.png" /><figcaption>**Figure 2.** Evolution of the normalized reflectance R405 measured during the 1000 °C temperature plateau under NH₃ on different AlN/MoS₂ samples. The AlN thickness is varied between ≈2 nm (4 s deposition time) and ≈4 nm (8 s deposition time).</figcaption></figure>

### 3. Results

We started to assess the diffusion of nitrogen into AlN. A first set of experiments was performed in the home-made furnace equipped with the *in situ* reflectance monitoring system on the AlN/MoS₂ stacks described in experimental methods. Figure 2 shows R405 signal recorded during 1000 °C plateau under NH₃ ambience for various AlN thicknesses on MoS₂. A complete acquisition curve is shown and explained in figure S5. An experimental issue occurred during the measurement of the AlN 7 s/MoS₂ stack at the end of the thermal treatment, but the beginning and the first half can be exploited.

We clearly observe a progressive increase in R405 followed by a stabilization due to a change in color of the stack with time. The required time to see a color change depends on the thickness of the AlN layer. One must remember that the AlN thickness does not affect the signal, since it is fully transparent at this wavelength [14]. The thicker the AlN film, the slower the reaction kinetics. AlN is a rather good diffusion barrier against nitrogen at 1000 °C, since 8 s of AlN deposition corresponds to roughly 4 nm of AlN. As AlN is highly transparent at the working wavelength, we assume that the color change is mainly due to the transformation of the MoS₂ sub-layer.

To confirm the changes of the 2D-MoS₂ through the AlN seed layer, Raman spectroscopy was performed before and after heat treatment and the results are reported in figure 3. It shows that the MoS₂ Raman modes $E_{2g}^1$ and $A_{1g}$ located at ≈382 cm⁻¹ and ≈407 cm⁻¹, respectively, completely vanished after NH₃ exposure. It confirms a clear modification of MoS₂ below the AlN seed layer. A second set of experiments under NH₃ at 1000 °C was performed in the Aixtron equipment. The inset of figure 3 presents similar results obtained in the Aixtron equipment. As shown in the inset of figure 3, both $E_{2g}^1$ and $A_{1g}$ MoS₂ Raman modes completely vanished on the AlN 10.4 s MoS₂ sample confirming the efficiency of this reactive heat treatment.

To evaluate the composition of the ‘converted MoS₂’ material, we performed XPS on the samples before and after heat treatments (see table S3 in SI for details). The Mo3d and S2p regions for the AlN 4 s MoS2 sample before and after heat treatment are visible in figures 4(a) and (b), respectively (Mo3p spectra were also acquired and gave similar results to Mo3d). N1s region is not shown since the nitrogen contribution is mainly due to the presence of Al and N bonds from the top AlN seed layer. Moreover, the similar binding energies between Mo-N and Al-N make deconvolution difficult.

Before NH₃ heat treatment, Mo3d deconvolution displays two main peaks at 228.2 eV and 231.2 eV corresponding to Mo⁴⁺3d⁵/₂ and Mo⁴⁺3d³/₂, respectively, which are attributed to Mo-S bonds. Mo3d⁵/₂ and Mo3d³/₂ peaks at 232.2 eV and 235.0 eV are attributed to Mo⁶⁺ state corresponding to a small amount of MoO₃ on the top surface. We also observe a slight peak at 225.4 eV attributed to S2s due to the overlapping of Mo3d with S2s spectrum. After NH₃ heat treatment, we clearly observe a change in the XPS signal. First, we notice the absence of S2s signal while Mo3d is still visible with multiple environments. The disappearance of the S2s signal is also confirmed in S2p region visible in figure 4(b) and in table S3 where no sulfur is detected in the NH₃-treated samples. Second, the Mo3d FWHM is larger after annealing, indicating a
<figure><img src="image_7.png" /><figcaption>**Figure 3.** Raman spectra of MoS₂ before heat treatment (in black) and AlN 8 s MoS2 after 1000 °C NH₃ heat treatment (in red). The inset presents the Raman spectrum of AlN 10.4 s MoS2 after NH₃ heat treatment in aixtron equipment. The peak at ≈520 cm⁻¹ is attributed to the Si substrate.</figcaption></figure>

a)

<figure><img src="image_8.png" /></figure>

b)

<figure><img src="image_9.png" /><figcaption>**Figure 4.** (a) Mo3d and (b) S2p XPS spectra of AlN 4 s MoS2, before and after 1000 °C NH3 heat treatment. (a) Reprinted with permission from [12]. Copyright (2024) American Chemical Society.</figcaption></figure>

noticeable evolution in the Mo bonding environment. By deconvoluting Mo3d spectra, we can attribute the Mo3d5/2 and Mo3d3/2 peaks at 228.0 eV and 231.0 eV, respectively, to Mo-N bonds [15, 16] which is the main part of the spectra. Peaks located at 228.6 eV and 231.7 eV are attributed to molybdenum oxynitride which is probably a mixture of MoO2, MoO3 and Mo2N [17]. Third, the Mo3d5/2 and Mo3d3/2 peaks at 232.2 eV and 235.3 eV are assigned to a slight amount of MoOx. Finally, a 1000 °C heat treatment under ammonia also seems to modify the bottom SiO2. Indeed, a shoulder in the Si2p signal is visible after heat treatment due to the appearance of a new environment around 101 eV and 102 eV and attributed to SiOxNy [18, 19] (see figure S6). The presence of SiOxNy also confirms the diffusion of N through the stack.

Time-of-flight secondary ion mass spectrometry (ToF-SIMS) was performed and no sulphur was detected in the stack or within the substrate suggesting that sulphur atoms have rather counter-diffused through the AlN seed layer (see figure S7).

To further confirm the MoS₂ phase conversion, XRD measurements were carried out before and after NH₃ heat treatment. It is worth noting that the initial MoS₂ thin film is polycrystalline with a random in-plane orientation and a preferential out-of-plane orientation along *c*-axis [20]. Figures 5(a) and (b) show
a)

<figure><img src="image_10.png" /></figure>

b)

<figure><img src="image_11.png" /><figcaption>**Figure 5.** In-plane XRD diffractograms of AlN/MoS₂ samples (a) before and (b) after 1000 °C NH₃ heat treatment. (a) and (b) Reprinted with permission from [12]. Copyright (2024) American Chemical Society.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>**Figure 6.** Cross-sectional TEM images at different magnifications of AlN 10.4 s MoS2 sample after 1000 °C NH3 heat treatment. Reprinted with permission from [12]. Copyright (2024) American Chemical Society.</figcaption></figure>

the in-plane XRD spectra before and after NH₃ heat treatment, respectively. Before heat treatment, in-plane measurements display three distinct peaks. Two of them located at 2θ≈33.2 °C and 2θ≈59.5 °C are attributed to AlN (100) and AlN(110) planes (JCPDS N°C00-025-1133). It shows that, as the MoS₂ thin film, the AlN seed layers present a random in-plane orientation. The last peak at 58.3 °C corresponds to the MoS₂ (110) plane diffraction. After heat treatment, AlN planes are still visible but the MoS₂ (110) contribution has vanished. This first observation agrees with the previous Raman and XPS characterizations. We also observe two new slight peaks located at 2θ≈35 °C and 2θ≈63 °C after heat treatment. Their attribution is uncertain but they could be attributed to a MoN₀.₄₆ or a Mo₂N (JCPDS N°C01-076-3074 and N°C04-016-6273, respectively). In any case, whatever the precise structure formed after ammonia treatment, these analyses confirm that MoS₂ has been well modified and the quality of the AlN seed layer seems to be preserved.

To further confirm that the AlN seed quality has not been degraded, STEM cross-sections of AlN 10.4 s MoS2 after NH3 heat treatment at different magnifications are presented in figure 6. The inset shows a high-resolution TEM observation of the AlN/MoS2 interface of the reference stack (sputtered AlN (100 nm)/MoS2) presented in our previous work [8] without any treatment. A homogenous and continuous MoS2 film is observed with its VdW structure. The Al and N atoms appear to stack with a remarkable lattice matching on the MoS2 surface suggesting a local VdW epitaxy of the AlN grains on the MoS2 grains. After ammonia treatment, the cross-sectional STEM images of AlN 10.4 s MoS2 show a discontinuous interface with the substrate instead of the expected MoS2 lamellar structure. A structural reorganization is observed with the formation of nodules up to several tens of nanometers. These nodules appear recessed within the SiO2 material. The magnification at atomic resolution in the nodules shows that the initial VdW plans of MoS2 disappear after heat treatment. Moreover, even if the AlN seed layer is very thin, it appears well crystallized and displays distinct planes parallel to the growth direction suggesting a preferential c-axis
<figure><img src="image_13.png" /><figcaption>**Figure 7.** Energy dispersion mappings of the AlN/MoS₂/SiO₂ interfaces after 1000 °C NH₃ heat treatment. Reprinted with permission from [12]. Copyright (2024) American Chemical Society.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Figure 8.** (a) AlN (002) RC of 1 µm thick AlN layer sputtered on AlN 10.4 s MoS2 stack after 1000 °C NH₃ heat treatment; (b) pictures of the tape test made on the 1 µm thick AlN layer; (c) picture of a tape test made on a 100 nm thick AlN layer on untreated AlN 10.4 s MoS2 stack.</figcaption></figure>

<figure><img src="image_15.png" /></figure>

textured growth. Finally, the AlN layer is lying on the SiO₂ layer after heat treatment which could improve its adhesion on the substrate.

STEM images were coupled with EDX analysis. The measurements are presented in figure 7. Oxygen (O) contribution in green comes from the SiO₂ substrate and the top of the AlN seed layer indicating that the seed layer is significantly oxidized possibly due to sample reventing. Silicon (Si) in cyan is predominant at the bottom of the image and comes from the SiO₂ substrate. The Mo element in yellow is only visible in the bright areas of the HAADF image below the AlN seed layer and confirms the structural reorganization of the MoS₂ layer during the heat treatment. Indeed, as presented in the inset of figure 6 and in line with our previous work [8], the initial MoS₂ film is continuous with three monolayers before heat treatment. The N mapping in red indicates that N is mainly associated with Al, but is also localized in the Mo nodules. Finally, no sulphur was detected during EDX analysis which confirmed the results obtained by XPS (figure 4(b)) and ToF-SIMS (figure S7).

Finally, to confirm that the AlN seed layer is not degraded by the heat treatment, we performed the sputtering deposition of a thick AlN layer. Figure 8(a) shows the AlN (002) RC measurement of a 1 µm thick AlN layer sputtered on AlN 10.4 s MoS2 after ammonia treatment. It displays a RC value of 0.71 °C highlighting a pronounced c-axis orientation. To the best of our knowledge, such a low RC value has never been reported for a 1 µm thick AlN sputtered layer on silicon-based substrates. Also, it proves that the detrimental delamination issue inherent in the initial AlN/MoS₂ stack has been solved. This is qualitatively visible in figures 8(b) and (c) where we achieved tape tests. Considering that the adhesion between the tape and the AlN surface is always the same, we observed in figure 8 (c) the delamination of a 100 nm thick AlN
<figure><img src="image_16.png" /></figure>

layer on an untreated AlN/MoS₂ stack, while 1 µm thick AlN film on the annealed stack remains on the surface (see figure 8(b)).

## 4. Discussion

The previous experimental characterizations demonstrate the chemical conversion under NH₃ heat treatment of the lamellar structure of MoS₂ into another structure, without VdW interactions. This new structure is composed of nodules of Mo(Oₓ)Nᵧ. The heat treatment induces a kind of de-wetting of the MoNₓ-based layer at the interface between SiO₂ and AlN. The precise crystalline identification of the new compound is uncertain. Meanwhile, the in-plane XRD measurements and cross-sectional STEM confirm that MoS₂ was completely transformed without significant degradation of the AlN seed layer. Interestingly, the Mo(Oₓ)Nᵧ nodules appear recessed within the SiO₂ substrate, suggesting that the oxide participates in their round shapes through surface energy minimization. At 1000 °C, SiO₂ starts to soften and affects the surface cohesion energy of the Mo(Oₓ)Nᵧ layer leading to round shapes. This phenomenon improves the AlN adhesion on the SiO₂ substrate.

The color change of the AlN/MoS₂ stack during NH₃ heat treatment is exploited to analyze the diffusion behavior of nitrogen in AlN. The study of the diffusion process requires the use of a well-chosen model. An interesting model is the diffusion-limited aggregation model developed by Witten and Sander [20]. However, in this model, the diffusion and formation of fractal structures are limited by the availability/access to growth sites which does not seem to be the case in this work. Therefore, we based our diffusion study on applying and comparing the use of linear and parabolic laws in diffusion processes formalized and generalized by Wagner et al [21]. The linear law is described as follows: Y = k₁.t (where Y is the thickness of AlN, k₁ a linear constant and t the time) while the parabolic law is Y² = kₚ.t (in this case kₚ is a parabolic constant). A limitation in the conversion reaction of nitrogen with MoS₂ by the availability of nitrogen should follow a linear trend, whereas a parabolic law described a limited N species availability at the AlN/MoS₂ interface through the AlN barrier diffusion. First, we need to determine where to extract the conversion time on the curves in figure 2 to apply a kinetic model. For each curve, we decided to choose the inflection point of the curve (dR405/dt = max). Then, the values of dR405/dt = max for each AlN thickness (determined by XRR in figure S4) are plotted as a function of the normalized time in figures 9 (a) and (b), respectively. After a linear regression, the R² coefficient appears better with the parabolic law than with the linear one (R² = 0.987 vs. R² = 0.904). As it could be expected, the MoS₂ conversion through the AlN seed layer is limited by the diffusion of nitrogen species in the AlN barrier.

The parabolic constant is evaluated as kₚ = 4.351 × 10⁻¹⁷ and is related to the diffusion coefficient (D) of nitrogen in AlN. However, the evaluation of a diffusion coefficient D requires pressure and temperature studies to elaborate a complete model. Meanwhile, we can at least try to apply a diffusion model based on Fick's second law [22]:

$$c(z, t) = \frac{N}{\sqrt{\pi Dt}} e^{-\frac{z^2}{4Dt}}, \qquad (1)$$
<figure><img src="image_17.png" /><figcaption>Figure 10. Mechanism proposal of the 2D-MoS₂ chemical modification through the AlN seed layer during 1000 °C NH₃ heat treatment.</figcaption></figure>

where C is the concentration of nitrogen species at a given time (t) and depth (z) within AlN, N is the total amount of nitrogen species that diffuses in AlN, and D is the diffusion coefficient of nitrogen in AlN. We assume that the concentration at the inflection point is C₀/2 where C₀ is the concentration leading to the total conversion of MoS₂, then we have:

$$c(z,t) = \frac{C_0}{2} \Leftrightarrow z = \sqrt{-4Dt \ln \left( \frac{C_0 \sqrt{\pi Dt}}{2N} \right)}. \qquad (2)$$

By defining N, D and C₀ as fitting parameters and applying this equation (2) to the curve in figure 9(b), we can extract a diffusion coefficient of N in AlN: D ≈ 2.42 × 10⁻²¹ m² s⁻¹. To the best of our knowledge, no publication mentions a diffusion coefficient of N in AlN, since the aluminum melts at a rather low temperature of 660 °C, and the direct thermal nitridation of Al is not possible above this temperature.

A mechanism for the chemical transformation of MoS₂ through the AlN seed layer is described in figure 10.

In the first step, during the 1000 °C heat treatment, N species cross the AlN seed layer through diffusion paths (probably through the grain boundaries), and reach the MoS₂ sub-layer. In the second step, N reacts with MoS₂ inducing the conversion of the lamellar layer into a molybdenum nitride compound following this reaction: xN + MoS₂ → MoNₓ + 2 S. Since no sulphur atoms are detected neither in the stack nor in the substrate, S probably counter-diffuses into AlN and is exhausted out of the stack probably in the form of H₂S (g). Hence, a nitridation in a reducing ambience seems mandatory to ensure the correct transformation of the 2D-sulphide. Finally, we assume that MoNₓ also reacts with the surrounding SiO₂ to form nodules of a Mo(Oₓ)Nᵧ which appear recessed into the SiO₂. This could occur at the periphery of the nodules, where dangling bounds are prompt to react during surface minimization, a temperature where SiO₂ enters a plastic mechanical regime. As the crystalline quality of the AlN seed layer is not (or nearly not) affected by the heat treatment (and since it is quite a rigid material), it locally settles directly on the SiO₂ substrate and seeds the further growth of a thick AlN layer showing improved crystalline quality (compared to an identical layer deposited directly on SiO₂).

## 5. Conclusions

In this paper, we demonstrate the proof of concept of a new elaboration process of c-axis textured AlN piezo-material using the chemical transformation of lamellar MoS₂ seed layer into Mo(Oₓ)Nᵧ nodules through a very thin AlN layer under ammonia heat treatment. It appears the N species diffuse through the thin AlN layer towards the underlying MoS₂ without affecting its crystalline quality. The heat treatment improves the adhesion of the AlN layer on the substrate and solves the systematic detrimental delamination of thick AlN films (≥200 nm) directly deposited on MoS₂. With this process, we report a 1 µm thick sputtered-AlN layer on a silicon-based substrate without any delamination with a record AlN (002) RC of 0.71 °C. This work paves the way towards the integration of AlN layers with enhanced crystal quality on silicon-based substrates.

## Data availability statement

All data that support the findings of this study are included within the article (and any supplementary files).
Acknowledgment

Part of this work was performed at the Platform for Nanocharacterization of the CEA Grenoble, MINATEC Campus, with support from the French Recherche Technologique de Base programme.

ORCID iDs

J Patouillard https://orcid.org/0000-0002-1277-9221

F Mercier https://orcid.org/0000-0002-5210-857X

B Hyot https://orcid.org/0000-0001-6021-9256

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