# Influence of the nucleation surface on the growth of epitaxial Al2O3 thermal CVD films deposited on cemented carbides

Maoxiang Zhu, Sofiane Achache, Mélanie Emo, Alejandro Borroto Ramírez, Jean-François Pierson, Frédéric Sanchette

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## ► To cite this version:

Maoxiang Zhu, Sofiane Achache, Mélanie Emo, Alejandro Borroto Ramírez, Jean-François Pierson, et al. Influence of the nucleation surface on the growth of epitaxial Al2O3 thermal CVD films deposited on cemented carbides. Materials & Design, 2022, 216, pp.110601. 10.1016/j.matdes.2022.110601. hal-04218360

HAL Id: hal-04218360

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Submitted on 26 Sep 2023

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# Influence of the nucleation surface on the growth of epitaxial Al₂O₃ thermal CVD films deposited on cemented carbides

<figure><img src="image_4.png" /></figure>

Maoxiang Zhu a,b,c,*, Sofiane Achache a,b, Mélanie Emo c, Alejandro Borroto Ramírez c, Jean-François Pierson c, Frédéric Sanchette a,b

## HIGHLIGHTS

* The nucleation of α- and κ-Al₂O₃ can be controlled with specific nucleation treatments.

* α-Al₂O₃ is the preferred alumina phase to be deposited without any nucleation treatment.

* α-Al₂O₃ is grown epitaxially on rutile (TiO₂), with (120)α-Al₂O₃//(101)rutile, (003)α-Al₂O₃//(010)rutile and [210]α-Al₂O₃//[101]rutile.

* κ-Al₂O₃ is grown epitaxially on Ti(C, N), with (100)κ-Al₂O₃//(111)Ti(C,N), (013)κ-Al₂O₃//(220)Ti(C,N) and [031]κ-Al₂O₃//[112]Ti(C,N).

## ARTICLE INFO

**Article history:**
Received 12 November 2021
Revised 17 February 2022
Accepted 25 March 2022
Available online 17 March 2021

**Keywords:**
Thermal chemical vapor deposition (CVD)
Alumina
Epitaxial growth
Bonding layer

## GRAPHICAL ABSTRACT

<figure><img src="image_5.png" /></figure>

## ABSTRACT

This work aims at understanding the nucleation and growth of alumina films grown on Ti(C,N)-based layers using an industrial-scale CVD system. Firstly, Al₂O₃ layer was deposited on Ti(C,N)-based layers without any nucleation treatment, pure α-Al₂O₃ is obtained, whereas no orientation relationship at the Ti(C, N)/α-Al₂O₃ layers interface can be observed. Secondly, Al₂O₃ layer was deposited on the bonding layer consisting of rutile TiO₂, which is obtained by oxidizing the uppermost part of Ti(C,N)-based layers. Herein, α-Al₂O₃ single-phased layer is obtained, and the epitaxial growth of α-Al₂O₃ on rutile is observed. The orientation relationships can be found as: (120)α-Al₂O₃//(101)rutile, (003)α-Al₂O₃//(010)rutile and [210]α-Al₂O₃//[101]rutile. Finally, Al₂O₃ layer was deposited on the bonding layer, produced from a TiCl₄-H₂-N₂-CH₄-CO-AlCl₃ gas mixture. Regarding this intermediate layer, despite additions of CO and AlCl₃, no evidence for oxide phases, e.g. TiO₂ and Al₂O₃ can be found, while Ti(C,N) needle-shaped grains develop. In this case, κ-Al₂O₃ is epitaxially grown on Ti(C,N), with the orientation relationships found as: (013)κ-Al₂O₃//(220)Ti(C,N), (100)κ-Al₂O₃//(111)Ti(C,N) and [031]κ-Al₂O₃//[112]Ti(C,N). Since processing parameters for the alumina depositions were always the same, it is revealed that the nucleation of α-Al₂O₃ and κ-Al₂O₃ can be accurately controlled with deposition of specific bonding layers.

© 2022 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).

## 1. Introduction

Coated cutting tools are widely used in the metal cutting industry, in order to improve productivity and process reliability [1,2].

<footnote>a LASMIS, Université de Technologie de Troyes, Antenne de Nogent, Pôle Technologique de Sud Champagne, 52800 Nogent, France</footnote>

<footnote>b Nogent International Centre for CVD Innovation (NICCI), LRC CEA-LASMIS, UTT, Antenne de Nogent, Pôle Technologique de Sud Champagne, 52800 Nogent, France</footnote>

<footnote>c Institut Jean Lamour, UMR CNRS 7198, Université de Lorraine, 54000 Nancy, France</footnote>

<footnote>* Corresponding author at: LASMIS, Université de Technologie de Troyes, Antenne de Nogent, Pôle Technologique de Sud Champagne, 52800 Nogent, France.
E-mail address: maoxiang.zhu@utt.fr (M. Zhu).</footnote>
this respect, CVD Al₂O₃ coatings are considered as an ideal candidate, owing to the beneficial combination of good chemical inertness and hardness at high temperature [3–7]. CVD Al₂O₃ coatings are commonly deposited from a AlCl₃-CO₂-H₂-H₂S-HCl gas mixture [8–12]. The global process can be described by the following reaction:

$$2\text{AlCl}_3\text{(g)} + 3\text{H}_2\text{(g)} + 3\text{CO}_2\text{(g)} \rightarrow \text{Al}_2\text{O}_3\text{(s)} + 3\text{CO}\text{(g)} + 6\text{HCl}\text{(g)}$$

This overall reaction involves both homogeneous gas phase reactions and heterogeneous surface reactions [13]. Additionally, the introduction of catalyst H₂S may enhance the overall reaction rate and suppress the “dog-bone” effect to gain a better thickness uniformity [13–15].

Accordingly, κ-Al₂O₃ and α-Al₂O₃ are desired for cutting applications [5,6]. The metastable κ-Al₂O₃ has an orthorhombic structure (space group Pna2₁) with lattice parameters a = 4.84 Å, b = 8.33 Å and c = 8.95 Å, showing an ABAC- stacking sequence of oxygen close-packed planes along the c-axis, Al-ions occupy both tetrahedral and octahedral interstitial sites at a 1:3 ratio [16]. The stable α-Al₂O₃ has a trigonal crystal structure (space group R3̅c) with lattice parameters a = b = 4.75 Å and c = 12.97 Å [17]. The structure consists of an ABAB- stacking sequence of oxygen close-packed planes with Al-ions in octahedral interstices. Importantly, κ-Al₂O₃ may transform irreversibly to α-Al₂O₃ resulting in volume contraction of about 8%, due to the heat generated during the CVD process as well as metal cutting [18]. Furthermore, it is found that mechanical properties of CVD α-Al₂O₃ coatings can be enhanced by choosing an optimized texture for the α-Al₂O₃ phase [12,13]. Thus, CVD Al₂O₃ coatings have continuously been optimized with respect to the phase and film texture control [12,19,20].

Because of adhesion problems between Ti(C,N)-based layers and Al₂O₃ layers, the use of bonding layers is suggested to enhance the toughness of tools and achieve the growth control of alumina [21–25]. Halvarsson et al show that the nucleation of κ-Al₂O₃ is favorably occurring on non-oxidized surfaces of Ti(C,N) [26,27]. Despite α-Al₂O₃ shows excellent thermal stability compared to κ-Al₂O₃, it is found to be difficult to be deposited directly on Ti (C,N)-based layers [21–24]. It is also reported that deposition of the bonding layer consisting of titanium oxides, e.g. Ti₂O₃ or Ti₄O₇, between Ti(C,N)-based and Al₂O₃ layers can promote the growth of α-Al₂O₃ [5,6]. However, the basic mechanisms explaining why the nucleation of α-Al₂O₃ is favored by different oxidizing treatment steps and bonding layers are not well known.

In this work, three Al₂O₃ coatings without and with bonding layers corresponding to samples: A, B and C were performed. Two other coatings corresponding to samples: B1 and C1 were stopped after synthesis of the bonding layers B and C. Firstly, X-ray diffraction was used for phase identification. The surface morphologies were observed by scanning electron microscopy (SEM), and the film cross-sectional microstructures were investigated by transmission electron microscopy (TEM). The orientation relationships at the Ti(C,N)/Al₂O₃ or bonding/Al₂O₃ layers interface were studied by high resolution transmission electron microscopy (HRTEM).

## 2. Materials and methods

### 2.1. Sample preparation

All coatings were deposited on cemented carbide substrates (WC-Co 6 wt%) mirror polished (1 µm diamond slurry), using an industrial-scale low-pressure CVD system (Ionbond Bernex™ BPXpro 325S). The global process has been described in a previous

work [28]. The device is mainly composed of five systems: reactor chamber, gaseous precursors supply system, liquid precursors (TiCl₄ and CH₃CN) supply system, AlCl₃ generator and by-products treatment system, note that the in-situ chlorination of Al-pellets generates almost only AlCl₃ and its dimer Al₂Cl₆ under experimental conditions used [28]. The detailed parameters are given in Table 1.

The multilayered coatings consist of a 0.5 µm-thick TiN inner layer, followed by a 2.5 µm-thick medium temperature Ti(C,N) layer, a 0.5 µm-thick high temperature Ti(C,N) layer, and a 4 µm-thick Al₂O₃ layer. The so-called medium temperature Ti(C,N) (MT-Ti(C,N)) layer was deposited at 850 °C from a TiCl₄-CH₃CN-H₂-N₂ gas mixture, whereas the high temperature Ti(C,N) (HT-Ti (C,N)) layer was deposited at 1005 °C from a TiCl₄-CH₄-H₂-N₂ gas mixture.

In sample A, Al₂O₃ top layer was deposited on Ti(C,N)-based layers without any nucleation treatment. On the contrary, Al₂O₃ layer in sample B was deposited on the bonding layer B, which was obtained by oxidizing the uppermost part of the HT-Ti(C,N) layer using CO₂. Al₂O₃ layer in sample C was deposited on the bonding layer C, which was produced from a TiCl₄-H₂-N₂-CH₄-CO-AlCl₃ gas mixture. Two other coatings (samples B1 and C1) were stopped after synthesis of the bonding layers B and C.

### 2.2. Characterization

X-ray diffraction (XRD) studies were carried out using a Bruker D8 Advance diffractometer operating at 40 kV/40 mA with Kα radiation of Cu (λ Kα₁ = 1.54059 Å and λ Kα₂ = 1.54443 Å), phase identifications were conducted in Bragg-Brentano geometry with a step size of 0.02° for a 2θ-range from 20° to 60°. The peak positions were determined by fitting the diffraction peaks with Pseudo-Voigt function after Kα₂ stripping. The software MAUD was used, the instrumental effect has been considered. For clarity, the standard material® 640d is used as reference to determine the instrumental contribution, as the crystallite size of the standard material is known.

The surface morphologies were observed by scanning electron microscopy (SEM) using a Zeiss Gemini 500. Transmission electron microscopy (TEM) analyses were performed in a JEOL JEM-ARM 200F Cold FEG TEM/STEM equipped with a spherical aberration (Cs) probe corrector operating at 200 kV. STEM micrographs were obtained using high angle annular dark field (HAADF) detectors. Electron energy loss spectroscopy (EELS) analyses were performed for differently grown layers. The orientation relationships at the Ti (C,N)/Al₂O₃ layers interface were investigated by high-resolution transmission electron microscopy (HRTEM). Prior to TEM investigations, all cross-sectional thin foils were prepared using a FIB/SEM FEI Helios NanoLab 600i equipped with platinum gas injection system. FIB column used a gallium liquid metal ion source operating at an accelerating voltage up to 30 kV, and a Pt-layer was deposited to protect the sample surface, using the commercial FEI MeCpPtMe3 precursor.

## 3. Results and discussion

### 3.1. X-ray diffraction

The X-ray diffractograms of samples A, B, B1, C and C1 are shown in Fig. 1 (a) – (e). As for sample B1, it was stopped after synthesis of the bonding layer B. Since the bonding layer is thin (approx. 100 nm), peaks corresponding to titanium oxides in the 2θ region investigated exhibit insufficient intensity. To highlight the growth of rutile, sample B1 was investigated by TEM, the SAED pattern is presented in Supplementary Fig. S1.
<caption>**Table 1**</caption>

The detailed processing parameters are given, a series of coatings were performed: sample A: TiN/MT-Ti(C,N)/HT-Ti(C,N)/Al₂O₃ sample B: TiN/MT-Ti(C,N)/HT-Ti(C,N)/bonding layer B/Al₂O₃ sample B1: TiN/MT-Ti(C,N)/HT-Ti(C,N)/bonding layer B sample C: TiN/MT-Ti(C,N)/HT-Ti(C,N)/bonding layer C/Al₂O₃ sample C1: TiN/MT-Ti(C,N)/HT-Ti(C,N)/bonding layer C.

<table><thead><tr><th></th><th>TiN</th><th>MT-Ti(C,N)</th><th>HT-Ti(C,N)</th><th>bonding layer B</th><th>bonding layer C</th><th>Al<sub>2</sub>O<sub>3</sub></th></tr></thead><tbody><tr><td>Temperature (°C)</td><td>900</td><td>850</td><td>1005</td><td>1005</td><td>1005</td><td>1005</td></tr><tr><td>Pressure (mbar)</td><td>70</td><td>70</td><td>70</td><td>70</td><td>70</td><td>70</td></tr><tr><td>Time (min)</td><td>40</td><td>60</td><td>20</td><td>20</td><td>20</td><td>150</td></tr><tr><td>H<sub>2</sub> (vol.%)</td><td>balance</td><td></td><td></td><td></td><td></td><td></td></tr><tr><td>N<sub>2</sub> (vol.%)</td><td>39.6</td><td>20.3</td><td>18.4</td><td>-</td><td>9.2</td><td>-</td></tr><tr><td>TiCl<sub>4</sub> (vol.%)</td><td>1.3</td><td>2.4</td><td>1.4</td><td>-</td><td>1.2</td><td>-</td></tr><tr><td>CH<sub>3</sub>CN (vol.%)</td><td>-</td><td>0.8</td><td>-</td><td>-</td><td>-</td><td>-</td></tr><tr><td>CH<sub>4</sub> (vol.%)</td><td>-</td><td>-</td><td>3.4</td><td>-</td><td>2.2</td><td>-</td></tr><tr><td>CO<sub>2</sub> (vol.%)</td><td>-</td><td>-</td><td>-</td><td>3.8</td><td>-</td><td>4.0</td></tr><tr><td>CO (vol.%)</td><td>-</td><td>-</td><td>-</td><td>-</td><td>0.9</td><td>-</td></tr><tr><td>AlCl<sub>3</sub> (vol.%)</td><td>-</td><td>-</td><td>-</td><td>-</td><td>0.4</td><td>1.6</td></tr><tr><td>HCl (vol.%)</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>2.0</td></tr><tr><td>H<sub>2</sub>S (vol.%)</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>0.3</td></tr><tr><td>Total gas flow (L.min<sup>-1</sup>)</td><td>33.8</td><td>26.1</td><td>32.5</td><td>26.0</td><td>32.5</td><td>25.0</td></tr></tbody></table>

<figure><img src="image_6.png" /><figcaption>**Fig. 1.** X-ray diffractograms of samples A, B, B1, C and C1.</figcaption></figure>

As shown in Fig. 1, the WC phase and the Ti(C,N) phase are indexed by the ICDD powder diffraction files 01-073-9900 and 01-071-6059, respectively. Ti(C,N) (111) and (200) peaks neighbour closely two other reflections belonging to the TiN layer underneath, minor overlaps near the baseline among these reflections are observed. Additionally, since the main advantage of the MT-CVD Ti(C,N) process is the suppression of interfacial η-phase, e.g. W₃Co₃C or W₆Co₆C, which is not detected, consistent with earlier studies [29–31].

Without any nucleation treatment, pure α-Al₂O₃ is obtained in sample A (see Fig. 1 (a)). It is shown that the deposition of bonding layer B also leads to the growth of α-Al₂O₃ (see Fig. 1 (b)). Further-

more, it is revealed that the deposition of bonding layer C promotes the nucleation of κ-Al₂O₃ in sample C (see Fig. 1 (d)). Sample C1 was stopped after synthesis of the bonding layer C, no evidence for oxide phases can be found herein. The α-Al₂O₃ and κ-Al₂O₃ phases are indexed by the ICDD powder diffraction files 00-046-1212 and 00-052-0803, respectively.

## 3.2. Surface morphology

Fig. 2 (a) – (c) show the surface morphologies of samples A, B and C, samples A and B exhibit similar morphologies composed of small facetted grains. Previous studies found that the κ- to α-
<figure><img src="image_7.png" /><figcaption>**Fig. 2.** SEM micrographs (a)–(c) showing surface morphologies of samples A, B and C. Some twin-related defects (black arrows) of κ-Al₂O₃ are observed in the SEM micrograph (c).</figcaption></figure>

Al₂O₃ phase transformation often occurs, this process depends strongly on the deposition time and temperature [2,22,32]. It is also observed that the α-Al₂O₃ obtained through the phase transformation exhibits relatively large equiaxed grains [2]. In addition, thermal stability of κ-Al₂O₃ coatings was studied previously by heat treatment at 1000 °C, while argon (Ar) was used as a carrier gas, it is found that the phase transformation started at 70 min [33]. However, neither diffraction peak originating from κ-Al₂O₃ (see Fig. 1 (a) and (b)), nor large equiaxed grains can be observed (see Fig. 2 (a) and (b)). Hence, no phase transformation can be expected in samples A and B. Meanwhile, twin-related defects typically present in κ-Al₂O₃ are visible in the SEM micrograph of sample C (see Fig. 2 (c)), in agreement with an earlier study [19].

### 3.3. Microstructural analyses

Fig. 3 (a) is a cross-section STEM HAADF micrograph of sample A, and Fig. 3 (b) shows EELS mapping acquired from the dashed square indicated in (a). Since acetonitrile (CH₃CN) is considered as the single carbon and nitrogen feeding source, it is generally accepted that the C/(C + N) ratio of MT-Ti(C,N) coatings is almost independent from deposition parameters and it is typically around 0.5 [34,35]. However, the chemical composition of HT-Ti(C,N) coatings is highly reliant on CH₄ and N₂ contents in the gas mixture [36]. Herein, a relatively low CH₄/N₂ ratio was used for the deposi-

tion of the HT-Ti(C,N) layer. Therefore, the MT-Ti(C,N) layer is carbon-rich and the HT-Ti(C,N) layer is nitrogen-rich, as shown in Fig. 3.

Fig. 4 (a) shows a cross-section TEM BF image of sample A, and Fig. 4 (b) is a TEM BF image with lower magnification to present the entire Al₂O₃ layer, which is approximately 4 µm-thick. Fig. 4 (b.1) is a SAED pattern acquired from the pink circle area indicated in Fig. 4 (b). This diffraction pattern evidences the formation of α-Al₂O₃. Fig. 4 (c) shows a TEM BF image acquired from the red circle area indicated in Fig. 4 (a), at the HT-Ti(C,N)/α-Al₂O₃ layers interface. Fig. 4 (c.1) and (c.2) present SAED patterns acquired from the pink and green circle areas indicated in Fig. 4 (c), α-Al₂O₃ is grown on the HT-Ti(C,N) layer. Furthermore, detailed analysis was carried out using high resolution transmission electron microscopy (HRTEM), but no orientation relationship at the HT-Ti(C,N)/α-Al₂O₃ layers interface can be highlighted.

As mentioned beforehand, the metastable κ-Al₂O₃ phase should be the preferred alumina phase grown on non-oxidized surfaces of Ti(C,N), with the epitaxial growth of {001} κ-Al₂O₃ close-packed planes on {111} Ti(C,N) [19,21]. Apparently, it is not always the case. It is also reported that the κ- to α-Al₂O₃ phase transformation initiated at free surfaces, e.g. interfacial voids, often occurs in a CVD reactor resulting in pronounced cracking [2]. However, no visible cracking along the grain boundaries can be observed. Thereby, the phase transformation is ruled out, and it can be speculated that

<figure><img src="image_8.png" /><figcaption>**Fig. 3.** (a) STEM HAADF micrograph showing a cross-section of sample A. (b) EELS mapping obtained from the black dashed square indicated in (a).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 4.** (a) TEM BF image showing a cross-section overview of sample A. (b) TEM BF image with lower magnification to present the entire Al₂O₃ layer. (c) TEM BF image acquired from the red circle area at the interface indicated in (a). (b.1), (c.1) and (c.2) are selected area electron diffraction (SAED) patterns corresponding to differently grown layers. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

since α-Al₂O₃ is the most thermodynamically stable phase, it is favorably occurring with the deposition conditions used in the current study.

Fig. 5 (a) is a cross-section STEM HAADF micrograph of sample B, and Fig. 5 (b) shows EELS mapping acquired from the black dashed square indicated in Fig. 5 (a). From EELS analyses, the bonding layer B consisting of TiO₂ is highlighted between the HT-Ti(C,N) and Al₂O₃ layers. As the deposition parameters were always the same for Ti(C,N)-based layers, the carbon-rich MT-Ti(C,N) and nitrogen-rich HT-Ti(C,N) layers are evidenced.

Fig. 6 (a) shows a cross-section TEM BF image of sample B, almost the entire Al₂O₃ layer is composed of large columnar grains. It is reported that the grain refinement at the start of α-Al₂O₃ growth can be achieved by the oxidizing treatment using a gas mixture of CO₂ and CO [15]. Herein, the uppermost part of the HT-Ti(C,N) layer was oxidized without CO, few small grains can thus be observed, whereas some interfacial voids are found at the bonding/α-Al₂O₃ layers interface indicated by black arrows in Fig. 6 (a). It is assumed that the volume contraction of the order of 25–30% associated with the phase transformation of titanium

<figure><img src="image_10.png" /><figcaption>**Fig. 5.** (a) STEM HAADF image showing a cross-section of sample B. (b) EELS mapping acquired from the black dashed square indicated in (a).</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>**Fig. 6.** (a) TEM BF image showing a cross-section overview of sample B. Some porosity at the interface are observed (indicated by arrows). (b) and (d) TEM BF images showing the α-Al₂O₃ and HT-Ti(C,N) layers. (c) TEM BF image acquired from the red area at the bonding/α-Al₂O₃ layers interface indicated in (a). (b.1)-(d.1) SAED patterns corresponding to different layers. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

oxides, due to the interdiffusion of the bonding and Ti(C,N)-based layers [6,15].

Fig. 6 (b) presents a TEM BF image of Al₂O₃ layer near the coating surface, and Fig. 6 (b.1) is a SAED pattern acquired from the pink circle area indicated in Fig. 6 (b). Fig. 6 (c) shows a TEM image at the bonding/α-Al₂O₃ layers interface acquired from the red circle area indicated in Fig. 6 (a). As shown in Fig. 6 (c.1) and (c.2), SAED patterns reveal the growth of α-Al₂O₃ on the rutile bonding layer. Titanium dioxide (TiO₂) crystallizes primarily in three different phases, rutile, anatase and brookite, among which rutile state (tetragonal, space group P4₂/mnm, with lattice parameters a = b = 4.59 Å and c = 2.96 Å) is the most stable one [37]. Herein, sample B1 was stopped after synthesis of the bonding layer B, and the growth of rutile is evidenced (see Supplementary Fig. S1).

Fig. 7 (a) shows a TEM image at the bonding/α-Al₂O₃ layers interface, the rutile bonding layer is approximately 107 nm-thick. Fig. 7 (b) is a HRTEM image obtained from the yellow square indicated in Fig. 7 (a), where the epitaxial growth of α-Al₂O₃ on rutile is observed. As α-Al₂O₃ and TiO₂ rutile structures can be both considered as being composed of close-packed oxygen layers with cations distributed in octahedral sites, the atomic arrangement are supposed to play a critical role in controlling the properties of the stack α-Al₂O₃/rutile during the initial growth stages of α-Al₂O₃ [38]. Fig. 7 (c) and (d) illustrate the epitaxial relationships at the rutile/α-Al₂O₃ interface. Similar relationships have also been reported previously, where rutile thin films are epitaxially grown on α-Al₂O₃ substrates by reactive ionized cluster beam deposition

(RICBD) or metalorganic chemical vapor deposition (MOCVD) [38-40]. Although different deposition techniques are used, the epitaxial growth at the α-Al₂O₃/rutile interface is always highlighted.

Firstly, the lattice mismatch for (101) rutile and (120) α-Al₂O₃ with the in-plane orientations [101] rutile and [210] α-Al₂O₃ is low, less than a few percent [38]. Meanwhile, according to the structural mismatch schema at the rutile/α-Al₂O₃ interface proposed by Daneu et al, an ideal fitting region is shown, while an A- (or B-) type oxygen layer of α-Al₂O₃ faces an A- (or B-) type oxygen layer of rutile [41]. Secondly, it is found that rutile film consists of three equal domains rotated by 120° [37]. The reason behind this particular arrangement is that a tetragonal unit cell with a 2-fold symmetry matches with a hexagonal unit cell with a 3-fold symmetry in the stack rutile/α-Al₂O₃. It is shown that [101] direction of (010) rutile can orient with any of the three equivalent [210] directions of (001) α-Al₂O₃ [40]. Thus, the orientation relationships experimentally found can now be explained.

Fig. 8 (a) shows a cross-section STEM HAADF micrograph of sample C, Fig. 8 (b) presents EELS mapping acquired from the black square indicated in Fig. 8 (a). As shown in Fig. 8 (a), no evidence for oxide phases can be found at the bonding/Al₂O₃ layers interface, while Ti(C,N) needle-shaped grains are observed. It is reported that CO addition results in the incorporation of oxygen into the Ti(C,N) phase leading to the microstructure refinement, but the C/N ratio and the lattice parameter of the Ti(C,N) phase will not be affected [1]. However, EELS analysis shows that no evidence for oxide
<figure><img src="image_12.png" /><figcaption>**Fig. 7.** (a) TEM BF image showing the bonding/α-Al₂O₃ layers interface of sample B. (b) HRTEM image acquired from the yellow square indicated in (a). (c) and (d) Fast Fourier Transform (FFT) images obtained from the blue (bonding layer) and red (α-Al₂O₃ layer) dashed squares indicated in (b), the orientation relationships are illustrated by yellow dashed lines. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 8.** (a) STEM HAADF image showing a cross-section of sample C. (b) EELS mapping acquired from the black square indicated in (a).</figcaption></figure>

phases can be detected in the bonding layer. Hence, it is believed that the microstructure of bonding layer C is more likely dependent on the deposition temperature and the CH₄/N₂ ratio of the gas mixture used, according to Cheng et al [36].

Fig. 9 (a) shows a cross-section TEM BF image of sample C, and Fig. 9 (b) is a TEM BF image with lower magnification to present the Al₂O₃ layer near the coating surface. Fig. 9 (b.1) is a SAED pattern acquired from the orange circle area indicated in Fig. 9 (b). It is shown that pure κ-Al₂O₃ is obtained. Fig. 9 (c) shows a TEM BF image acquired from the red area indicated in Fig. 9 (a). Fig. 9 (c.1) and (c.2) are micro-diffraction and SAED patterns acquired

form the blue (bonding layer C) and orange (κ-Al₂O₃ layer) circle areas at the interface indicated in Fig. 9 (c), which confirm the nucleation of κ-Al₂O₃ on Ti(C,N) (bonding layer C). Since Ti(C,N) needle-shaped grains are extremely fine, the superposition of spots could falsify the result, the micro-diffraction was used.

Fig. 10 (a) shows the TEM image at the bonding/κ-Al₂O₃ layers interface, Ti(C,N) needle-shaped grains are observed. Fig. 10 (b) is the HRTEM image obtained from the yellow dashed square indicated in Fig. 10 (a). Fast Fourier Transform (FFT) images acquired from the green, blue and red dashed squares indicated in Fig. 10 (b) are shown in Fig. 10 (c)-(e), it is shown that Ti(C,N) needle-
<figure><img src="image_14.png" /><figcaption>**Fig. 9.** (a) TEM BF image showing a cross-section overview of sample C, needle-shaped grains developed at the bonding/κ-Al₂O₃ layers interface. (b) and (d) TEM BF images showing the κ-Al₂O₃ and HT-Ti(C,N) layers. (c) TEM BF image obtained from the red circle area at the bonding/κ-Al₂O₃ layers interface indicated in (a). (b).1)-(d.1) SAED patterns corresponding to differently grown layers. (c.2) Micro-diffraction image of the bonding layer. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

shaped grains consist of twin-related domains with the growth of parallel {111} planes.

Fig. 11 (a) shows the TEM BF image at the bonding/κ-Al₂O₃ layers interface. Fig. 11 (b) presents the HRTEM image acquired from the red square indicated in Fig. 11 (a). Fig. 11 (c) and (d) are FFT images acquired from the blue and orange squares indicated in Fig. 11 (b). As has been discussed by Liu *et al*, the oxygen sublattice in κ-Al₂O₃ with an -ABAC- stacking sequence has the space group P6₃/mmc, and the positions of Al-ions reduce the symmetry to Pna2₁ [16]. Several studies show that alumina deposition on TiN {111} planes yields generally the epitaxial growth of {001} κ-Al₂O₃ close-packed planes, with directions parallel in these two phases, e.g. [010]κ-Al₂O₃//[110]TiN, the misfit is about 9% [26,27]. According to Halvarsson *et al*, if oxygen ions occupy A-positions, titanium ions can be placed in interstitial B- or C-positions on the close-packed oxygen plane, TiN crystals can be formed with an -ABCABC- or -ACBACB- stacking sequence [26]. It is similar to that found in Fig. 9 (c.1) and (c.2), the orientation relationships can be expressed as: (001)κ-Al₂O₃//(111)Ti(C,N) and [010]κ-Al₂O₃//[101]Ti(C,N). The atomic arrangement schema of (001) κ-Al₂O₃ created by software VESTA is given in Supplementary Fig. S2.

Furthermore, the epitaxial growth at the lateral interface of bonding layer C is also observed, with different orientation relationships found as: (100)κ-Al₂O₃//(111)Ti(C,N), (0113)κ-Al₂O₃//(220)Ti(C,N) and [031]κ-Al₂O₃//[112]Ti(C,N). It is reported earlier that the surface morphology of bonding layer may be responsible for the phase composition of Al₂O₃ [15,42]. It shows that when the bonding layer consists of needle shaped grains, with needles almost parallel to the deposition direction, the nucleation of α-Al₂O₃ is favored. Conversely, the formation of a nano-transition needle-like phase is also proved to promote the growth of κ-Al₂O₃ [4]. Thereby, it is believed that the effect of surface morphology of bonding layers can be discarded. To verify the orientation relationships, the arrangement of O-atoms in the plane (100)κ-Al₂O₃ along the direction [031] and Ti-atoms in the plane (111)Ti(C,N) along the direction [112] are illustrated by VESTA (see Fig. 12), and the Crystallography Open Database files 1,000,442 and 1,540,253 corresponding to κ-Al₂O₃ and Ti(C,N) phases are used. It is speculated that the crystallographic arrangement may control the properties of the stack κ-Al₂O₃/Ti(C,N) during the initial growth stage of κ-Al₂O₃. The distance between every two O-atoms and Ti-atoms at the interface is calculated and pre-
<figure><img src="image_15.png" /><figcaption>**Fig. 10.** (a) TEM BF image showing the bonding layer of sample C composed of fine needle-shaped grains. (b) HRTEM image obtained from the yellow area indicated in (a), and the twin boundary can be seen from the bottom left corner to the top right corner. (c)-(e) Fast Fourier Transform (FFT) images along [112] projection, obtained from dashed squares indicated in (b). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>**Fig. 11.** (a) TEM BF image showing the bonding layer of sample C. (b) HRTEM image obtained from the red square area indicated in (a). (c) and (d) Fast Fourier Transform (FFT) images acquired from the blue (bonding layer C) and orange (κ-Al₂O₃ layer) square areas indicated in (b), the orientation relationships are illustrated by yellow dashed lines. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

sented in Fig. 12, the mismatch is about 7%, and slightly lower mismatch is obtained for (013)κ-Al₂O₃ and (220)Ti(C,N). Hence, the basic mechanism explaining why the nucleation of κ-Al₂O₃ is favored in sample C can be clearly explained.

4. Conclusions

In this work, a series of CVD Al₂O₃ coatings were deposited without (sample A) and with different bonding layers (samples B
<figure><img src="image_17.png" /><figcaption>**Fig. 12.** Schematic crystallographic arrangement of O-atoms in the plane (100)$_{\kappa-Al_2O_3}$ along the direction [031], and Ti-atoms in the plane (111)$_{Ti(C,N)}$ along the direction [112].</figcaption></figure>

and C). Two other samples B1 and C1 were stopped after synthesis of the bonding layers B and C. The interfacial microstructures were investigated in detail, and main results can be summarized as follows:

(1) The Al₂O₃ layer in sample A was deposited on the non-oxidized HT-Ti(C,N) layer, pure α-Al₂O₃ is obtained, but no orientation relationships at the interface can be observed. Thus, it can be concluded that α-Al₂O₃ is the preferred alumina phase to be deposited with the deposition conditions used in this work, as it is the most thermodynamically stable phase.

(2) The Al₂O₃ layer in sample B was deposited on the bonding layer B consisting of rutile, α-Al₂O₃ is epitaxially grown on rutile, and the orientation relationships can be expressed as: $\langle 1\bar{2}0 \rangle_{\alpha-Al_2O_3} // \langle 10\bar{1} \rangle_{\text{rutile}}, \langle 003 \rangle_{\alpha-Al_2O_3} // \langle 010 \rangle_{\text{rutile}}$ and $\langle 2\bar{1}0 \rangle_{\alpha-Al_2O_3} // \langle 101 \rangle_{\text{rutile}}$.

(3) Regarding sample C, pure κ-Al₂O₃ is obtained, due to the epitaxial growth of κ-Al₂O₃ on Ti(C,N) with the epitaxial relationships found as: $\langle 100 \rangle_{\kappa-Al_2O_3} // \langle \bar{1}1\bar{1} \rangle_{Ti(C,N)}, \langle 0\bar{1}3 \rangle_{\kappa-Al_2O_3} // \langle 2\bar{2}0 \rangle_{Ti(C,N)}$ and $\langle 031 \rangle_{\kappa-Al_2O_3} // \langle 112 \rangle_{Ti(C,N)}$. It is believed that it is necessary to satisfy the epitaxial growth of κ-Al₂O₃ on Ti(C,N) to promote the nucleation of κ-Al₂O₃, otherwise, α-Al₂O₃ is the preferred alumina phase to be deposited with the deposition conditions used.

(4) Since processing parameters for the alumina deposition were always the same, and the epitaxial growth of α-Al₂O₃ on rutile as well as that of κ-Al₂O₃ on Ti(C,N) are evidenced, it is revealed that the nucleation of α-Al₂O₃ and κ-Al₂O₃ can be accurately controlled with different treatments of the nucleation surface.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

The authors would like to acknowledge assistance from the Institut Jean Lamour and the Groupement d'Intérêt Public Haute-Marne (GIP 52) for their support. Special thanks to Sylvie Migot (CC3M -Institut Jean Lamour) for FIB/SEM investigations and Jaafar Ghanbaja (CC3M-Institut Jean Lamour) for the insightful discussions.

Appendix A. Supplementary material

Supplementary data to this article can be found online at
https://doi.org/10.1016/j.matdes.2022.110601.

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