Multimethod cross-sectional characterization
approaching the limits of diamond monophase
multilayers

M. Meindlhumer, J. Grau, H. Sternschulte, Y. Halahovets, P. Siffalovic, M.

Burghammer, D. Steinmüller-Nethl, J. Keckes

► To cite this version:

M. Meindlhumer, J. Grau, H. Sternschulte, Y. Halahovets, P. Siffalovic, et al. Multimethod cross-sectional characterization approaching the limits of diamond monophase multilayers. Materials Characterization, 2024, 212, pp.113973. 10.1016/j.matchar.2024.113973. hal-05099492

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# Multimethod cross-sectional characterization approaching the limits of diamond monophase multilayers

M. Meindlhumer<sup>a, *</sup>, J. Grau<sup>b</sup>, H. Sternschulte<sup>b</sup>, Y. Halahovets<sup>c</sup>, P. Siffalovic<sup>c</sup>, M. Burghammer<sup>d</sup>, D. Steinmüller-Nethl<sup>e</sup>, J. Keckes<sup>a</sup>

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

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## ARTICLE INFO

**Keywords:**

Hot-filament chemical vapour deposition

Diamond multilayers

Raman spectroscopy

Cross-sectional X-ray nanodiffraction

## ABSTRACT

Chemical vapour deposited diamond draws significant scientific interest due to its wide range of mechanical and functional properties, which can be easily tuned by the applied deposition conditions. Here, hot filament chemical vapour deposition was used to synthesize four monophase multilayer diamond thin films with individual layer thickness down to ~ 300 nm. Extensive structural characterization by scanning electron microscopy and Raman spectroscopy on the cross-section confirmed the different diamond morphologies originating from the different applied process parameters of the individual sublayers, where Raman spectroscopy could give a clear distinction between nano- and microcrystalline diamond down to a layer thickness of ~ 1 µm. Cross-sectional X-ray nanodiffraction identified exclusively the diamond crystal structure throughout the cross-section of the multilayered diamond thin films. Cross-sectional phase and FWHM analysis allowed to discriminate between the diamond morphologies down to ~ 500 nm, thus identifying the limitations of monophase diamond multilayers. While the microcrystalline diamond sublayers all exhibit pronounced structural gradients expressed in (i) increasing intensity of the 111 Debye-Scherrer ring, (ii) (110) fibre texture, and periodic variations of (iii) grain size and (iv) residual stress, the nanocrystalline diamond sublayers showed no pronounced cross-sectional variations. In summary, the experimental results provide insights into the cross-sectional evolution of microstructure and residual stress and the limitations of monophase multilayered diamond thin films.

## 1. Introduction

Diamond thin films have attracted significant scientific and industrial interest during the last decades, primarily due to their outstanding chemical, thermal, mechanical, and biocompatible properties [1,2]. For the synthesis of diamond films with thicknesses ranging from several nanometers up to > 100 µm, chemical vapour deposition (CVD) is commonly applied [3,4]. Similarly, the morphology of CVD diamond thin films can be easily tuned by the applied deposition conditions, where the resulting grain size can range from some nanometers up to around hundreds of micrometers [5,6].

Depending on the applied deposition conditions, highly different cross-sectional gradients of microstructure, texture and residual stress are formed within the diamond films as a consequence of the

evolutionary non-equilibrium deposition process [2,5]. These include formation of a distinct crystallographic texture [7–9] and increasing crystallite size over the film thickness in case of microcrystalline diamond [10] (MCD), while in case of nanocrystalline diamond (NCD), texture and grain size are uniform across the thickness [2,11]. In both MCD and NCD film types, the film growth is coupled with the formation of intrinsic stress [12,13], which is superimposed by thermal stress originating from the mismatch of coefficients of thermal expansion (CTE) of the film and substrate [13–15].

Cross-sectional X-ray nanodiffraction [16,17] (CSnanoXRD) with a resolution down to ~ 25 nm has had already significant success in uncovering gradual evolution of microstructure, texture and residual stress in a 22.5 µm thick bilayer film prepared by microwave plasma-assisted CVD [18], as well as in uncovering gradients in monolithic

<footnote><sup>a</sup> Department of Materials Science, Montanuniversität Leoben, Franz Josef Straße 18, 8700 Leoben, Austria</footnote>

<footnote><sup>b</sup> Fakultät für Geistes- und Naturwissenschaften, Technische Hochschule Augsburg, Postfach 110605, 86031 Augsburg, Germany</footnote>

<footnote><sup>c</sup> Institute of Physics, Slovak Academy of Sciences, Dúbravská cesta 9, 845 11 Bratislava 45, Slovak Republic</footnote>

<footnote><sup>d</sup> ESRF, The European Synchrotron, 71 Avenue des Martyrs, CS40220, 38043 Grenoble Cedex 9, France</footnote>

<footnote><sup>e</sup> CarbonCompetence GmbH, Weisstraße 9, 6112 Wattens, Austria</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: michael.meindlhumer@unileoben.ac.at (M. Meindlhumer).</footnote>
nano- (NCD) and microcrystalline (MCD) sublayers prepared by hot filament (HF) CVD process [19]. In addition to CSnanoXRD, where discrimination between micro- and nanocrystalline diamond is drawn via peak broadening and texture analysis [18,19], also Raman spectroscopy can give insight into the C—C bond states and the morphology of the diamond microstructure, which highly differ between micro- and nanocrystalline diamond [20–26].

Mechanical properties of CVD diamond depend strongly on the morphology, as extensively summarized by Hess [27]. With increasing grain size of the diamond thin films, Young's modulus and hardness increase from 400 to 1200 GPa and ~ 50 to ~ 110 GPa, respectively, while fracture strength decreases from ~ 9 to ~ 1 GPa, whereas the latter is also dependent on sample size [27]. The relationship between grain size and fracture toughness is less clear, though a trend to higher fracture toughness with increasing grain size can be seen in literature [27–32]. The advantage of NCD and ultra-nanocrystalline diamond thin films resides in decoupling the diamond grain size and the film thickness, resulting in isotropic film properties [2,27].

Temporal variability of deposition conditions during the CVD process allows for enabling the synthesis of diamond multilayers with a cross-sectional variation of mechanical and functional properties [18,33–37]. These multi-layered diamond thin films have various applications including biomedical implants [38,39], sensors for MEMS devices [40] as well as wear-resistant hard coatings for microelectronics and high-end milling tools [35,36,41–43]. As examples, Takeuchi et al. [33,34] analysed diamond multilayer thin films and revealed increased bend strength and increased adhesion, with concomitant slightly enhanced wear rate compared to monolithic MCD films. Similarly, Salgueiredo et al. [43] revealed decreased surface roughness and increased adhesion strength of multilayered diamond, while Silva et al. [35] showed additionally improved machining capability. Hinzman et al. [36] investigated multilayered diamond with sublayer thicknesses of ~ 1 µm using sub-µm diffraction, yet, the identification of the individual diamond layers by FWHM, intensity and stress analyses may be shadowed by the high roughness of the WC-Co substrate. However, Hinzmann et al. [36] discovered higher compressive stress in MCD compared to NCD in agreement with other literature [18,19]. However, despite the well-established advantages of diamond multilayers, the extent of cross-sectional gradients of texture, morphology, and residual stress in sub-µm thick sublayers has not been thoroughly investigated up to now.

In the current study, four diamond multilayers were deposited, each with a total thickness ranging from approximately 10.5 to 4.7 µm, where the multilayer architecture was achieved by alternating the parameter sets to create nanocrystalline and microcrystalline diamond sublayers during hot filament CVD. Scanning electron microscopy, cross-sectional Raman spectroscopy and cross-sectional X-ray nanodiffraction confirmed the different diamond morphologies of the individual layers. For individual layer thicknesses as thin as 750 nm, the characterization techniques allowed for a clear differentiation between microcrystalline diamond (MCD) and nanocrystalline diamond (NCD). However, in the multilayer sample with sublayers of MCD and NCD with sublayer thicknesses between 500 and 300 nm, transitional states between MCD and NCD were observed. Overall, the reported nanoscale microstructure and residual stress gradients provide unprecedented insights into the microstructure and stress evolution of monophase multilayered diamond thin films.

## 2. Experimental

### 2.1. CVD thin film deposition

The four multilayered diamond thin films were deposited on single-crystalline Si (100) wafers with a diameter of 100 mm and a thickness of 530 µm by modified HF-CVD [44]. Ta filaments with a diameter of 3.6 mm were electrically heated to 2400 °C for thermal dissociation of the precursor gases hydrogen and methane. The Ta filament temperature

was monitored by a Senotherm Metis M316 pyrometer. Before the thin film deposition, the Si wafers were seeded by the application of detonation diamond nanoparticles with an average size of 5 nm, which were dissolved in isopropyl alcohol in an ultrasonic bath. For deposition of the multilayered films, the temperature of the Si substrate of 800–850 °C was monitored by a thermocouple type K and a total gas pressure of 5 mbar was applied, while the filament-to-substrate distance was maintained at 50 mm, while CH₄/H₂ ratios of ~ 0.1 and ~ 2% were used, respectively. The applied CH₄/H₂ ratios of ~ 0.1 and ~ 2% have already been proven to successfully synthesize MCD and NCD morphologies in monolithic thin films, resulting in grain sizes in the range of ~ 10–15 nm and in the micrometer range, respectively [19].

Schematic morphologies of the two trilayer (further denoted as TL1 and TL2) and the two multilayer (further denoted as ML1 and ML2) films are given in Fig. 1a, the deposition time per sublayer was given by the growth rates evaluated from monolithic MCD and NCD films [19] deposited using similar conditions. For all layers, the gas flows are adjusted for the individual sublayers using a 180 s ramp, detailed CH₄ content over deposition time is shown for each multilayer film in Figs. 3a, 4a, 5a and 6a.

### 2.2. SEM characterization

A characterization of the TL1, TL2, ML1 and ML2 films cross-sectional morphology was performed in a ZEISS LEO 1525 scanning electron microscope (SEM) operated at an acceleration voltage of 2 kV and an aperture size of 20 µm.

### 2.3. Raman spectroscopy

In order to obtain information on the different types of C—C bonds and their relative volume fractions in the MCD and NCD films, Raman spectroscopy was performed on the films surface and cross-section using a Renishaw System 1000 Raman spectrometer and WITec alpha300 R confocal Raman spectrometer, respectively. Therefore, cross-sections of the diamond films were mechanically mirror-polished to minimize any influence of surface roughness. The Renishaw System 1000 Raman spectrometer was operated at 25 mW laser power using a green Ar laser with a wavelength of 514.5 nm. Spectral resolution of the Renishaw System 1000 was ~ 1–2 cm⁻¹ using a 50 times magnifying objective with a numerical aperture NA = 0.25 resulting in a spot size of ~ 1.5 µm. For surface analysis, three spectra were recorded for each sample at room temperature between wavenumbers of 400 and 2000 cm⁻¹ using a measurement time of 30 s.

For the cross-sectional Raman spectroscopy, the WITec alpha300 R confocal Raman spectrometer used a solid-state laser operated at an excitation wavelength of 405 nm, a laser power of 1 mW. A step size of 200 nm was used for cross-sectional Raman analysis, where every spectrum was collected over a time of 40 s. A hundred-times magnifying objective with NA = 0.9 was used at a working distance of 0.3 mm resulting in a laser spot of ~ 450 nm in diameter given by the diffraction limit, while spectral resolution was ~ 1.3 cm⁻¹. In order to get a better insight into the MCD and NCD distributions at the cross-section of the films, the diamond Raman peak area fraction was fitted to the Raman spectra obtained at the cross-section.

### 2.4. CSnanoXRD characterization

For the CSnanoXRD experiments, thin lamellae (~ 10 × 1 × 0.53 mm³) were cut from the four multilayered diamond samples. The lamellae were then mechanically polished to a final thickness of ~ 250 µm. According to Stefenelli et al. [45], residual stress changes can be avoided when the lamella thickness is at least 5 times larger than the corresponding film thickness. Afterwards, a cross-sectional trench with a thickness of ~ 50 µm and a width of 400 µm was prepared in a ZEISS Auriga Crossbeam workstation equipped with a femtosecond laser,
<figure><img src="image_5.png" /><figcaption>Fig. 1. The schematics of the investigated monophase diamond multilayers presented in (a) shows the intended layer stacking. While the stacking is shown in scale for TL1, TL2 and ML1, the layer thicknesses of the ML2 sample are not to scale. In (b) a schematics of the X-ray nanodiffraction experiment is shown, displaying the approximate sample dimensions, the X-ray detector, and the beam stop (BS).</figcaption></figure>

keeping the above stated requirements. The reason to measure ~ 50 µm thick synchrotron lamellae is that the representative sample volume remains polycrystalline (as documented by a continuous Debye-Scherrer ring, cf. Fig. 1b) and therefore methodology of powder diffraction can be used to evaluate phases, texture and stresses [16,17]. The applied laser parameters were a pulse duration of 500 fs, a wavelength of 515 nm and a minimal focal diameter at the sample surface of ~ 20 µm. Further details on the laser system can be found elsewhere [46].

The CSnanoXRD experiments were performed at the Nanofocus extension of the ID13 beamline of the European Synchrotron (ESRF) in Grenoble, France [47–49]. The experiments were performed using a circular X-ray beam focused by multi-layer Laue lenses (MLLs) at a photon energy of 13.0 keV. The TL1, TL2 and ML1 thin film samples were scanned using a beam with dimensions of 75 × 75 nm² [47,48], while the ML2 thin film sample demanded a higher resolution and was scanned with a beam focussed to 25 × 35 nm² (Fig. 1) [49]. Before conducting the CSnanoXRD experiments, the alignment of the interface between the Si substrate and the diamond film was achieved by ensuring it was parallel to the incident beam. This alignment was accomplished by optimizing the absorption contrast between the diamond and the Si substrate through a series of absorption scans using an X-ray point detector at various sample tilts. The samples were scanned perpendicular to the Si-diamond interfaces with 50 nm steps and the 2D diffraction signal was recorded by an Eiger 4M detector at each measurement position using an exposure time of 50 ms. The exact detector geometry was calibrated using a NIST corundum powder, where the sample-to-detector distances were determined to be 99.987 mm [47], 109.848 mm [48] and 104.675 mm [49] for the experiments carried out on the TL1, TL2 and ML1, and the ML2 thin film samples, respectively. For

further experimental details, Ref. [16] should be considered.

The collected 2D X-ray diffraction data were processed using the Python package PyFAI [50,51]. In the observed 2θ region, which is determined by the relationship between the photon energy and the sample-to-detector distance, only the 111 Debye-Scherrer (DS) ring of diamond was covered. By means of a segmented azimuthal (angle δ) integration, depth-resolved plots of the diffracted intensity I(2θ, z) (radial intensity plot), the so-called phase plot, were generated for every sample depth z, where θ represent the Bragg’s angle. Furthermore, the azimuthal distributions of the diffracted intensities I¹¹¹(δ, z) and the full width at half maxima FWHM(δ, z) of the diffracted 111 DS rings were evaluated to give further insight into the multilayered microstructure and the thus introduced cross-sectional gradients originating from the particular deposition conditions.

In order to quantify the residual stress gradients in the samples, the ellipticities of the individual diamond 111 DS rings were evaluated. Therefore, for each DS ring, the 24 radial data sets I(2θ; δ, z) were obtained by azimuthal integration within 15-degree wide sections of δ. These data sets were then subsequently fitted using Pseudo-Voigt functions to determine the exact 2θ¹¹¹(δ, z) dependencies for each DS ring. In following, the 2θ¹¹¹(δ, z) distribution was used to calculate a set of lattice spacing d¹¹¹(δ, z) applying Bragg’s law (d¹¹¹(δ, z) = λ/2sinθ¹¹¹(δ, z)). Using the evaluation approach derived by Keckes et al. [16,17] considering the geometry of the CSnanoXRD experiment and an equiaxial residual stress state, the in-plane residual stress σᵣ(z) is calculated as follows

$$ \frac{\partial d^{111}(\delta, z)}{\partial \sin^2 \delta} = \sigma_r(z) \cdot \frac{1}{2} S_2^{111} \cdot d_0^{111}(z), \quad (1) $$
where the stress free lattice parameter $d_0^{11}(z)$ was calculated according to Ref. [45] and the X-ray elastic constant $\frac{1}{2}\sigma_2^{11} = 8.97 \cdot 10^{-13} \text{ Pa}^{-1}$ were adopted from literature data [52] applying the Hill grain interaction model.

### 3. Results & discussion

#### 3.1. Diamond trilayer TL1 MCD-NCD-MCD

The cross-sectional microstructure of the trilayer film depicted in Fig. 3b reveals a pattern of MCD interrupted by NCD. As competitive growth commences, the grain size progressively increases, reaching a sublayer thickness of approximately 3.5 µm. In SEM image (Fig. 3b), the transition to NCD above a thickness of about 3.5 µm appears sudden and distinct, with no observable signs of cross-sectional microstructure evolution up to a thickness of approximately 7 µm. This suggests that the sublayer was grown with a high secondary nucleation rate [2,11]. The transition back to microcrystalline diamond (MCD) appears less abrupt, primarily due to the gradual development of a columnar grain microstructure resulting from the resumed competitive growth (Fig. 3b).

Raman spectra obtained from the film surface and cross-section are shown in Fig. 2 and Suppl. Fig. 1, respectively. The Raman spectrum retrieved from the TL1 film surface is typical for MCD and comparable with literature [26,53,54,56,57,60,61] as it consists of a distinctive diamond Raman peak at $1332 \text{ cm}^{-1}$ and peaks originating from the D- and G-bands at 1350 and $1580 \text{ cm}^{-1}$, respectively, of $sp^2$ hybridised C located primarily at the grain boundaries [26,56,57,61]. The Raman signal of $sp^2$ bonded carbon is ~ 50 times stronger than the signal of $sp^3$

<figure><img src="image_6.png" /><figcaption>**Fig. 2.** Raman spectra with excitation perpendicular to the surface of the multilayer diamond films were obtained at room temperature using the Renishaw 1000 system at an excitation wavelength of 514.5 nm. While the TL1 and TL2 samples show clear evidence for MCD and NCD similar to literature [2,11,21–26,53–60], respectively, the Raman spectra obtained from the ML1 and ML2 films indicate an overlapped signal originating from the alternating MCD and NCD layers.</figcaption></figure>

bonded carbon, thus the spectra presented in Fig. 2 fits to a poly-crystalline diamond with high $sp^3$ to $sp^2$ C ratio [26] and confirms the presence of high-quality diamond at the film surface. No signal of the NCD layer below the 3.5 µm thick MCD film is visible. Similar Raman spectra were recorded at the cross-section from the MCD bottom and top sublayers (Suppl. Fig. 1b and c) located at film thicknesses of ~ 1.5 and ~ 9.0 µm, respectively. Contrary, in the Raman signal retrieved from the NCD interlayer (between the thicknesses of 3.3 and 7.0 µm), no distinctive diamond Raman peak is visible (cf. Suppl. Fig. 1b and c).

Generally, the presence of the diamond Raman peak in the Raman spectra is dependent on the excitation wavelength and the grain size of the film. The minimum grain size for its presence is ~ 50 nm, as identified by transmission electron microscopy, when using a HeNe laser [61] (632.8 nm), while the minimum size of coherently diffracting (SCDD) domains was identified to be ~ 12 nm by XRD for the presence of the diamond Raman peak when using an Ar laser [24] (514 nm).

Additionally, Sun et al. [58] identified a minimum SCDD of ~ 20 nm for the presence of the diamond Raman peak, in agreement with the SCDD evaluated by CSnanoXRD (Fig. 3f) discussed below. Additionally, strong peaks originating from the D and G-band of $sp^2$ hybridised C located at 1350 and $1580 \text{ cm}^{-1}$, are present in the spectrum of the NCD interlayer presented in Suppl. Fig. 1. Together with the concomitant peaks originating from the *trans*-polyacetylene (t-PA) $\nu_1$ and $\nu_3$ bands visible at wavenumbers of 1150 and $1450 \text{ cm}^{-1}$ the Raman spectrum obtained from the NCD interlayer can be interpreted as an indication of nanocrystallinity, in agreement with the literature [2,22,55,61,62]. The area fraction of the diamond Raman peak is displayed in Fig. 3c and allows to distinguish between the individual MCD and NCD layers in agreement with the morphology of the SEM image and the Raman spectra presented in Fig. 3b and Suppl. Fig. 1, respectively. Here, it is emphasized that the diamond Raman peak is present exclusively in the MCD layers, respectively, as shown in Fig. 3c. Contrary, the NCD layer exhibits a negligible diamond area fraction, in agreement with NCD spectra presented in Suppl. Fig. 1 and literature [2,22,55,61,62].

A phase plot was generated from the CSnanoXRD data acquired from the multilayer film through azimuthal integration over 360 degrees, as illustrated in Fig. 3d. According to the chosen experimental geometry (Sec. 2.4) described elsewhere [18], exclusively the 111 DS ring of the diamond phase at $2\theta = 26.84$ deg. is present in the diffraction signal. Additionally, no other crystalline phases were detected for both the MCD and the NCD sublayers (Fig. 3d). Furthermore, a rather sharp diamond 111 diffraction peak was observed in the MCD sublayers, indicating a high portion of coarse crystalline diamond, in agreement with the SEM and Raman data presented in Fig. 3b and c, respectively. Contrary, between ~ 3.3 µm < z < 7 µm, low intensity and high apparent FWHM of the diamond 111 peak are observed together with a concomitantly enhanced background signal indicating the presence of amorphous phases (Fig. 3d), in agreement with the presence of the *trans*-polyacetylene peaks in the Raman signal (Suppl. Fig. 1) and the absent diamond peak (Fig. 3c) of the NCD sublayer.

Diffraction data obtained from the multilayer film was furthermore used for a qualitative texture analysis, the cross-sectional evolution of the azimuthal intensity distributions along DS 111 rings $I^{111}(\delta,z)$ is presented in Fig. 3e. Beginning with the randomly oriented seed layer (between z = 0 and 0.3 µm) a rapid transition to a $\langle 110 \rangle$ fibre texture was revealed, indicated by azimuthal maxima of the 111 DS ring at $\delta = \pm 90$ and $\pm 35.3$ deg., where $\delta = \pm 0$ deg. corresponds to the film's out-of-plane orientation (film growth direction). The further narrowing of the $I^{111}(\delta,z)$ distributions up to z ≈ 3.3 µm indicates further sharpening of the $\langle 110 \rangle$ fibre texture during the competitive thin film growth in agreement with literature [9,19]. Upon transition to the NCD sublayer, the $\langle 110 \rangle$ fibre texture is greatly diminished, nearly a random distribution of azimuthal intensities is displayed, despite the interruptions given by the segmented CCD detector. However, azimuthal maxima of the 111 DS ring are still present at $\delta = \pm 90$ and $\pm 35.3$ deg., in agreement with literature [7,19]. At z ≈ 7.2 µm, again a rapid transition
<figure><img src="image_7.png" /><figcaption>Fig. 3. Cross-sectional characterization of a 10.5 µm thick diamond multilayer thin film composed of two MCD layers interrupted by a NCD interlayer. In (a) the applied methane concentration is shown, while in (b), (c) and (d-g) results retrieved by SEM, Raman spectroscopy and CSnanoXRD are presented, respectively. In detail, two layers of MCD exhibiting a gradual evolution of microstructure due to competitive grain growth are interrupted by a NCD layer with no recognizable microstructure (b). The difference between the NCD and the MCD layers are further highlighted by the fitted diamond Raman peak area fractions (c). The bottom and top MCD layers exhibit a distinctive diamond Raman intensity, while the interrupting NCD layer does not (c). In (d), the CSnanoXRD phase plot exclusively the diamond 111 reflection is present, albeit the intensity and apparent peak width clearly distinguish the MCD and NCD layers. Azimuthal intensity distributions reveal a gradually intensifying (110) fibre texture in the MCD layers, which is highly diminished in the interrupting NCD layer (e). The cross-sectional FWHM variation reveals a gradual decrease of the oscillatory FHWM with the increasing MCD sublayer thickness, while the FWHM gradually rises to ~ 0.6 deg. in the middle of the NCD interlayer (f). The residual stress evaluated from the 111 Debye-Scherrer Ring shows an oscillatory variation around a mean value of ~ 0.25 GPa in the MCD layers, while the interrupting NCD layer exhibits a rather constant tensile stress of + 0.6 GPa (g). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

to a sharp (110) fibre texture is observed, further reflected by the SEM image in Fig. 3b, where pronounced 111 facets of the diamond crystals are present at the surface oriented ≈ 35 deg. from the films surface normal, in agreement with literature [9].

FWHM of the 111 DS ring was evaluated to get qualitative insights

into thin film morphologies and is presented in Fig. 3f. In the TL1 film directly at the substrate-thin film interface a FWHM of ~ 0.18 deg. was evaluated which rapidly decreases to ~ 0.14 deg. at z = 0.5 µm for both out-of-plane and in-plane diffraction vector orientations. From z = 0.5 up to z ~ 3.0 µm, the average in-plane FWHM slightly decreases to ~
0.13 deg., whereby oscillations around the mean value are observed (Fig. 3f), similar to a monolithic MCD film [19]. The out-of-plane FWHM shows similar oscillations without further decrease of the average FWHM, which may be attributed to the reduced intensity in out-of-plane orientation (δ = 0 deg., cf. Fig. 3e). Here, the evaluated FWHM 0.18–0.11 deg. values correspond to sizes of coherently diffracting domains (SCDD) of ~60–40 nm, apparently significantly smaller than the grain size observed by SEM (Fig. 3b). Furthermore, the combination of elevated in-plane and diminished out-of-plane FWHM obtained from the MCD film disagrees with the columnar-grained thin film morphology detected by SEM indicating a significant amount of lattice defects in the MCD. Above z ~ 3.0 µm, the FWHM rapidly increases to a peak value of ~ 0.6 deg., where the in-plane FWHM is slightly lower than the out-of-plane FWHM, indicating a plate-like morphology, similar to investigations carried out on a monolithic NCD thin film [19]. The evaluated FWHM values correspond to SCDDs of ~10–15 nm, in agreement with the Raman data retrieved from NCD [24,58,61] as discussed above and cross-sectional data from literature [18,19]. Though constant deposition conditions were applied, FWHM analysed from the TL1 film already started to decrease at z ~ 6.0 µm, despite negligible diamond Raman signal up to z ~ 7.0 µm (cf. Fig. 3c). Between z ≈ 7.0 and 10.5 µm the FWHM is similar compared to the values obtained from the first MCD layer (0 µm < z < 3.0 µm) further indicating similar morphology, as found by the SEM and Raman analysis (Fig. 3b,c, Supp. Fig. 1).

At the interface of the TL1 film, a residual stress στ~0.25 GPa was evaluated from the DS ring which decreased to ~-0.25 GPa at z ~ 1 µm. Up to z ~ 3.0 µm, the στ magnitudes oscillated between -0.3 and +0.6 GPa with an average value of ~ 0.25 GPa (Fig. 3g). The oscillatory στ agrees with a monolithic MCD thin film prepared using similar deposition conditions [19], though the average stress in the monolithic film was evaluated with ~ -0.5 GPa, which represents quite a significant stress increase. In the NCD sublayer (between z ~ 3-7.2 µm), the stress profile was constant at a value of ~ 0.6 GPa, again shifted to more tensile stress of ~ 0.35 GPa compared to the monolithic NCD film deposited with similar conditions [19]. In the top MCD layer, the oscillatory stress remerged, beginning with a decrease to ~0 and oscillation around a mean value of ~ 0.25 GPa. Generally, the residual stress magnitudes evaluated from the MCD and NCD sublayers of TL1 film are shifted towards more tensile stress by ~ 0.75 and 0.35 GPa, respectively, compared to their monolithic counterparts investigated in [19], despite the sublayer's similar morphology [19]. Given the CTEs of MCD and NCD [14], their mismatch to the Si substrate [63] and the deposition temperature (Sec. 2.1), a thermal stress component of ~ -0.6 to -0.9 GPa and ~ -0.1 to ~0.4 GPa can be calculated for MCD and NCD, respectively, in agreement with literature [14,18,64]. Concomitantly, tensile intrinsic stress components of ~ 0.95 to 1.15 GPa and ~ 0.7 to 1.0 GPa are evaluated from the residual stress data (cf. Fig. 3g) of the MCD and NCD sublayers, respectively, which in agreement with the στ magnitudes are shifted to more tensile stress compared to their monolithic counterparts [19]. This may be attributed to a coupled effect of changed thermal and intrinsic stresses, since the layers are not isolated in contrast to their monolithic counterparts. One explanation may be that the NCD and MCD layers at z = 3-7.2 µm and z = 7.2-10.5 µm, respectively, are grown on diamond of different morphologies, which may influence the intrinsic stress component. However, the MCD layer at z = 0-3.0 µm, is grown directly on the Si substrate (similar to a monolithic MCD film) and also exhibits a higher tensile στ magnitude (compared to the monolithic MCD film in Ref. [19] deposited with similar conditions) indicating a rather complex residual stress distribution of multilayered diamond thin films. Notably, Gruber et al. [18] investigated similar cross-sectional stress gradients in a bilayer diamond deposited by plasma-assisted CVD. However, a direct comparison is rather difficult, since no monolithic counterparts were investigated in this study, which would allow clarifying the influence of the multilayer thin film system on the residual stress [18].

3.2. Diamond trilayer TL2 NCD-MCD-NCD

A cross-sectional microstructure of the trilayer TL2 film is presented in Fig. 4b and shows two ~ 1.5 µm thick NCD sublayers interrupted by a ~ 3 µm thick MCD sublayer. The NCD diamond sublayers show no signs of cross-sectional microstructure evolution, with uniform morphology for both NCD sublayers (Fig. 4b). After changing the deposition conditions in order to deposit the MCD interlayer, immediately competitive growth starts, leading to a gradually increasing grain size up to a sublayer thickness of ~ 3.0 µm and a columnar grain morphology (Fig. 4b). In the SEM, the transition above a thickness of ~ 4.8 µm to NCD appears instantaneous and sharp, indicating NCD growth with a high secondary nucleation rate [2,11].

The Raman spectra obtained from the TL2 films surface is presented in Suppl. Fig. 1 and shows a spectrum exclusively related to NCD [2,22,55,61,62]. In detail, in the Raman signal retrieved from the TL2 films surface (cf. Fig. 2), no distinctive diamond Raman peak is visible. Additionally, strong peaks originating from the D and G-band of sp² hybridised C located at 1350 and 1580 cm⁻¹, respectively, are present in the spectrum presented in Fig. 2. Together with the concomitant the peaks originating from the *trans*-polyacetylene v₁ and v₃ bands visible at wavenumbers of 1150 and 1450 cm⁻¹, respectively, the Raman spectrum obtained from the NCD interlayer is similar to the NCD layer in the TL1 sample (Suppl. Fig. 1a and b) and interpreted as an indication of nanocrystallinity [2,22,55,61,62]. Notably, given the thickness of the top NCD layer of ~ 1.5 µm, it is reasonable to conclude that the maximum information depth of the Raman signal using an excitation wavelength of 514.5 nm is below 1.5 µm for NCD. Similar Raman spectra have been obtained at the cross-section from the NCD bottom and top sublayers (Suppl. Fig. 2) located at z = 0.6 and 5.8 µm, respectively, which are devoid of a diamond Raman peak. Together with the concomitant peaks originating from the v₁ and v₃ t-PA bands visible at wavenumbers of 1150 and 1450 cm⁻¹, respectively, the Raman spectra obtained at z = 0.6 and 5.8 µm can be interpreted as NCD, in agreement with literature [2,22,55,61,62]. The absence of the diamond Raman peak in the Raman spectra depends on both the grain size and the excitation wavelength. In this case, the calculated SCDD of approximately 15 nm from the CSnanoXRD analysis (as seen in Fig. 4f) falls well below the minimum grain size required for the presence of a diamond Raman peak, as reported in the literature [24,58,61]. Contrary, the Raman spectrum obtained at z = 3.2 µm exhibits a distinctive Raman peak at 1332 cm⁻¹ and a comparably low intensity originating from the D and G-bands at 1350 and 1580 cm⁻¹, respectively, of sp² hybridised C located primarily at the grain boundaries [26,56,57,61].

Given the fact that the Raman signal of sp² bonded carbon is approximately 50 times stronger than the signal of sp³ bonded carbon, the spectra obtained at z = 3.2 µm fits to a polycrystalline diamond with high sp³ to sp² C ratio [26]. Furthermore, the area fraction of the diamond Raman peak was evaluated from the spectra obtained at the cross-section of the TL 2 film (cf. Suppl Fig. 2) and is presented in Fig. 4c. It allows to clearly distinguish the individual MCD and NCD layers, since the diamond signal is exclusively present in the MCD sublayer, while on the contrary, in the NCD sublayers diamond signal vanishes (cf. Fig. 4c). The area fraction of the diamond Raman peak is in agreement with the morphology of the SEM image presented in Fig. 4b.

The phase plot derived from azimuthal integration of the CSnanoXRD data is presented in Fig. 4d. Besides diamond, there were no other crystalline phases detected in either the MCD or the NCD sublayers (as shown in Fig. 4d). These can be distinguished by their distinct intensity and the apparent width of the 111 diamond DS ring, as well as the background intensity. The enhanced background detected in the NCD layers (z ≈ 0-1.6 µm and z ≈ 4.7-6.3 µm in Fig. 4d) can be related to the higher amount of amorphous sp² hybridised C and the presence of t-PA in the NCD layers, in agreement with the Raman data [2,22,55,61,62] discussed above (cf. Suppl. Fig. 2). The transition to the MCD layer appears sharp, the maximum intensity of the diamond 111
<figure><img src="image_8.png" /><figcaption>Fig. 4. Cross-sectional characterization of a 6.3 µm thick diamond multilayer thin film composed of two NCD layers interrupted by a MCD interlayer. In (a) the applied methane concentration is shown, while in (b), (c) and (d-g) results retrieved by SEM, Raman spectroscopy and CSnanoXRD are presented, respectively. In detail, two layers of NCD exhibit no apparent cross-sectional evolution of microstructure due to growth conditions favouring ongoing secondary nucleation of grains, while competitive growth forms elongated columnar grains in the MCD interlayer (b). The difference between the NCD sublayers and the MCD interlayer are further highlighted by the fitted diamond Raman peak area fractions retrieved from the cross-section presented in (c). The bottom and top NCD layers exhibit negligible diamond Raman area fractions, while in the interrupting MCD layer the diamond Raman signal is maximized (c). In (d) the CSnanoXRD phase plot exclusively the diamond 111 reflection is present, albeit the intensity and apparent peak width clearly distinguish the MCD and NCD layers. Azimuthal intensity distributions reveal a weak (110) fibre texture in the NCD layers, while the interrupting MCD layer exhibits immediately a strong (110) fibre texture which intensifies with increasing thickness (e). The cross-sectional FWHM variation reveals a rapid increase from the seed layer to a plateau of ~0.4–0.5 deg., while in the MCD layer oscillatory FWHM with an average value of ~0.14 deg. was evaluated. In the top NCD layer again a plateau value of ~0.4–0.5 deg. was found (f). The residual stress evaluated from the 111 Debye-Scherrer Ring increases abruptly from 0 to 1.35 GPa during the transition from the seed layer to the CVD grown film (z ≈ 0–0.25 µm), while during growth of the NCD layer σ<sub>r</sub> decreased slightly. An oscillatory variation with decreasing average stress from 1 to 0 GPa was detected in the NCD interlayer, while again a rapid increase to σ<sub>r</sub> ~ 1 GPa was evaluated from the top NCD layer (g). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
DS ring is reached within 200–300 nm and remains steady up to $z \approx 4.7$ µm, indicating a high portion of crystalline diamond, again in agreement with the Raman data presented in Suppl. Fig. 2 and the literature [26,53,54,56,57,60,61]. Although the transition from MCD to NCD appears sharp in the SEM image (Fig. 4b), in the CSnanoXRD phase plot, the transition from high to low intensity (i.e., from MCD to NCD) exhibits a width of ~400 nm (Fig. 4d). It is important to highlight that the interface roughness of the MCD sublayer, resulting from faceted growth, as illustrated in Fig. 4b, could potentially account for the observed broadening in the transition from MCD to NCD in the CSnanoXRD phase plot, as depicted in Fig. 4d.

Azimuthal intensity distributions $I^{111}(\delta, z)$ used for qualitative texture analysis are presented in Fig. 4e. Between $z \approx 0-1.6$ µm the azimuthal intensity distributions show a rather continuous distribution along $\delta$, with only slightly enhanced signal of the 111 DS ring at $\delta = \pm 90$ and $\pm 35.3$ deg., indicating a non-pronounced (110) fibre texture in the NCD layer, in agreement with the literature [7,18]. At $z \approx 1.6$ µm, a rapid transition sets in, a strongly pronounced (110) fibre texture is developed within 200–300 nm ($z \approx 1.6$ µm). The further narrowing of $I^{111}(\delta, z)$ at $\delta = \pm 90$ and $\pm 35.3$ deg. indicates further sharpening of the (110) fibre texture until $z \approx 4.7$ µm. Above $z \approx 4.7$ µm, again a rapid transition to a non-pronounced (110) fibre texture expressed by rather continuous $I^{111}(\delta, z)$ is observed, which is finished at $z \approx 5.0-5.1$ µm. The comparably broader transition width of ~300–400 nm from MCD to NCD, in comparison to 200–300 nm at the transition from NCD to MCD ($z \approx 1.6-1.8$ µm) can be related to the intrinsic roughness induced by the faceted µm-sized diamond grains (cf. Fig. 4b). However, even taking into account some additional broadening of the transition between MCD and NCD in the TL1 film due to experimental conditions, as discussed above, the transitions between the different diamond morphologies is significantly narrower in the TL2 film (cf. Fig. 3e and Fig. 4e).

Cross-sectional qualitative diamond 111 peak width analysis of the TL2 film is presented in Fig. 4f. Directly at the substrate-thin film interface in-plane and out-of-plane FWHM of 0.23 and 0.27 deg. were evaluated from the diffraction data, which rapidly rise to plateau values of ~ 0.42 and 0.47 deg. for in-plane and out-of-plane directions, respectively, at $z \approx 0.3$ µm. Up to a film thickness of $z \approx 1.5$ µm, the in-plane FWHM remained constant at ~0.4 deg., while out-of-plane FWHM slightly increased to a peak value of ~ 0.49 deg., corresponding to a SCDD of ~15–17 nm (Fig. 4f), indicating a plate-like morphology, in agreement with literature [19]. While the transition from MCD to NCD was determined between $z \approx 1.6$ and $1.8$ µm in the phase (Fig. 4d) and texture (Fig. 4e) plots, the FWHM evaluated both in-plane and out-of-plane direction starts to decrease slightly already at $z \approx 1.5$ µm, before a steep decrease to a value of ~ 0.16 deg. is observed at $z \approx 1.6$ µm (Fig. 4f). Between $z \approx 1.8$ µm and $z \approx 4.7$ µm an oscillatory FWHM was evaluated for both in-plane and out-of-plane directions varying around the mean value of ~0.16 deg. Between $z \approx 4.7$ µm and $5.1$ µm a rapid increase of the FWHM was detected (Fig. 4f) together with the concomitant decrease of peak intensity (Fig. 4d) and more continuous $I^{111}(\delta, z)$ distributions (Fig. 4e), further confirming the rapid transition between MCD and NCD in the TL2 film. Again, FWHM values of ~ 0.42 and 0.47 deg. were evaluated for in-plane and out-of-plane directions, respectively. Despite similar applied process parameters, the plateau FWHM values, obtained from the NCD layers of the TL2 film differ remarkably from the ones obtained for the TL1 film, resulting in a ~ 50% enlarged SCDD (cf. Figs. 3e and 4e). However, the Raman signal obtained from the surface (Fig. 2) and cross section (Suppl. Fig. 2, z = 0.6 and $5.8$ µm) of the TL2 film confirm C bonding states typical for NCD, in agreement with literature [2,22,55,61,62].

At the substrate-thin film interface of the TL2 film a residual stress $\sigma_r$ ~ 0 GPa was evaluated from the 111 DS ring which increased rapidly to a peak value ~ 1.27 GPa at $z \sim 0.3$ µm during the transition from the seed layer to the CVD grown film (cf. Fig. 4g). From $z \sim 0.3$ µm to $z \sim 1.8$ µm the residual stress continually decreases to ~ 0.73 GPa, together with a slight increase in the out-of-plane FWHM (cf. Fig. 4f). The rapid

increase of the residual stress between $z = 0$ and ~ 0.3 µm may be interpreted as the transition from the seed layer to the steady film growth, where the increase of the FWHM goes hand in hand with a concomitant decrease of the mismatch of CTEs [14]. Given the mismatch of the CTEs of NCD and Si, again a thermal stress component of ~ -0.1 to -0.4 GPa can be estimated for the NCD layer and in following the intrinsic stress component ranges from ~ 1.37 to ~ 1.67 GPa and ~ 0.83 to ~ 1.13 GPa for $z \sim 0.3$ and ~ 1.8 µm, respectively, significantly higher than in the TL1 film or in a monolithic NCD film [19].

After the transition to the MCD layer, an oscillatory stress profile was detected, which on average showed a gradient of average values from ~ 1 to ~ 0 GPa between $z \sim 1.8$ and $z \sim 4.7$ µm, respectively (cf. Fig. 4g). Here, peak values of residual stress were observed at similar z-positions compared with the oscillatory maxima and minima of the in-plane FWHM (cf. Fig. 4e and f). Furthermore, the gradual decrease of residual stress may be associated with an increasing mismatch of the CTEs of MCD with the Si substrate due to a further increase of grain size [14] (cf. Fig. 4b), which in turn further increases the compressive thermal stress component and reduces the overall $\sigma_r$ magnitudes. However, given the sharp transition between NCD and MCD at $z \sim 1.6-1.8$ µm confirmed by SEM (Fig. 4b) and the CSnanoXRD data (Fig. 4d-f), a gradual change of the intrinsic stress component is more likely. Supposing a similar thermal stress component as in the TL1 film (~ -0.6 to -0.9 GPa) given by the same deposition conditions (cf. Sec. 2.1) and thin film morphology (compare Figs. 3 and 4), the intrinsic stress component gradually decreases from ~ 1.6 to 1.9 GPa to ~ 0.6 to 0.9 GPa at $z \sim 1.8$ and $z \sim 4.7$ µm, respectively. Concomitant with the FWHM increase between $z \sim 4.7$ and $5.1$ µm, also $\sigma_r$ rises from 0 to ~ 1 GPa, similar to the rising $\sigma_r$ magnitudes from the beginning of the film growth. The total stress values evaluated from the TL2 film (Fig. 4f) are shifted further to more tensile $\sigma_r$ magnitudes compared to the TL1 film (Fig. 3f) and the monolithic counterparts of the MCD and NCD sublayers [19], which could indicate a significant influence of the (i) layer stacking, (ii) individual sublayer thickness and (iii) the abruptness of the transition between MCD and NCD on the intrinsic stress component of $\sigma_r$. This possibility will be further supported by the comparable $\sigma_r$ magnitudes evaluated from the ML1 and ML2 films presented in Sec. 3.3 and 3.4, respectively, where the former shows also rapid transitions between the MCD and NCD sublayers.

### 3.3. Diamond multilayer ML1

The cross-sectional microstructure of the ML1 film is presented in Fig. 5b. It shows four MCD sublayers interrupted by three sublayers NCD (Fig. 5b). The morphology of the first and second MCD, as well as the NCD layers with an individual sublayer of ~ 2 µm appear similar compared to the trilayer films TL1 and TL2 investigated prior. For the NCD layers with reduced thickness also no apparent morphological change compared to TL1 and TL2 could be detected, which can be related to the high secondary nucleation rate [2,11], decoupling the microstructure from the film thickness. In the top two MCD layers, narrowing of the columnar grains may be interpreted by the reduced layer thickness, which obviously hinders the gradual evolution of columnar grains.

The Raman spectrum obtained at the surface of the ML1 film presented in Fig. 2 shows a distinct diamond Raman peak at $1332 \text{ cm}^{-1}$ accompanied by a pronounced signal from the D and G-bands at 1350 and $1580 \text{ cm}^{-1}$, respectively, of sp$^2$ hybridised C located primarily at the grain boundaries [26,54,55,57,60]. Additionally, the presence of signal originating from the t-PA $\nu_1$ and $\nu_3$ bands visible at wavenumbers of $1150$ and $1450 \text{ cm}^{-1}$, respectively, indicates presence of signal originating from NCD [21–23,58], confirming overlapping Raman spectra from the top MCD and NCD layers (Fig. 2, Fig. 5b). Details about Raman spectra obtained from MCD and NCD sublayers have been already discussed in detail in Sec. 3.1. The Raman spectra presented in this study are furthermore in good agreement with literature
<figure><img src="image_9.png" /><figcaption>**Fig. 5.** Cross-sectional characterization of the 9 µm thick diamond multilayer thin film ML1. In (a) the applied methane concentration during deposition (Fig. 1) is shown, while in (b), (c) and (d-g) results retrieved by SEM, Raman spectroscopy and CSnanoXRD are presented, respectively. In detail, SEM identifies the four MCD layers by three layers of NCD (b). The difference between the NCD layer and the MCD layers are highlighted by the fitted diamond Raman peak area fractions, where high and negligible fractions indicate MCD and NCD, respectively (c). The red arrow in (c) indicates that in the topmost NCD layer the area fraction of the diamond Raman peak is not reduced to zero, in agreement with the Raman spectra presented in Suppl. Fig. 3. In (d) the CSnanoXRD phase plot exclusively the diamond 111 reflection is present, albeit the intensity and apparent peak width clearly distinguish the MCD and NCD layers of the ML 1 film. Azimuthal intensity distributions reveal a gradually intensifying (110) fibre texture in the MCD layers, which is highly diminished in the interrupting NCD layer (e). The cross-sectional FWHM variation reveals a gradual decrease of the oscillatory FHWM with the decreasing NCD sublayer thickness, while the FWHM in the MCD sublayers oscillates around a mean value of ~ 0.15 deg., independent of layer thickness except for the growth from the seed layer (f). The residual stress evaluated from the 111 Debye-Scherrer Ring is presented in (g) and shows an oscillatory increase to $\sigma_r$ ~ 1.0 GPa up to $z \approx 2.1$ µm, followed by constant $\sigma_r$ of 1.0 GPa up to $z \approx 3.7$ µm in NCD. The oscillation of $\sigma_r$ remerges in the next MCD layer with a concomitant decrease of the absolute $\sigma_r$ magnitude ($z \le 5.8$ µm), while tensile stress is built up and reduced thereafter in NCD and MCD layers, respectively (5.8 µm $\le z \le 9.0$ µm) (g). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
[21-24,54,56-58,60-62]. The Raman spectra obtained from individual positions along the cross-section of the ML1 are presented in Suppl. Fig. 3. There, at $z \approx 1.25, 5.25, 7.5$ and $8.75$ µm, in the centre of the respective MCD sublayers, Raman spectra with a distinctive diamond Raman peak at $1332 \text{ cm}^{-1}$ were found. Additionally, these spectra were composed of signal originating from the D-and G-band at 1350 and $1580 \text{ cm}^{-1}$, respectively, and as discussed above, the spectra are a clear indication for a presence of diamond with a high amount of $sp^3$ hybridised C [26,53,54,56,57,60,61]. Contrary, the spectra obtained at cross-sectional positions of $z \approx 3.25$ and $6.5$ µm are devoid of a distinctive diamond Raman peak, while exhibiting signal originating from the t-PA $\nu_1$ and $\nu_3$ bands visible at wavenumbers of $1150 \text{ cm}^{-1}$ and $1450 \text{ cm}^{-1}$, respectively, thus clearly indicating NCD diamond, in agreement with literature [21-23]. However, in the topmost (supposed) NCD layer, located between ~ 7.6 and 8.4 µm according to the Raman intensity of diamond presented in Suppl. Fig. 3, a slight diamond peak is detected in addition to the t-PA $\nu_1$ and $\nu_3$ bands. This indicates that the grain size increased beyond the threshold of the detection of the diamond Raman signal, while still a significant amount of t-PA is present at the grain boundaries, similar to Mohr et al. [24]. However, the SCDD will be discussed below, where only a slight increase in the topmost NCD layer (compared to the other NCD layers) was revealed (cf. Fig. 5f). Thus, the presence of the diamond Raman signal may be related to a combination of the finite size of the laser spot of 450 nm (Sec. 2.2), which coupled with the intrinsic roughness of the film (Fig. 5b) may also indicate that the spatial limit of the resolution of the Raman spectrometer is approached. However, the topmost NCD layer is ~ 800 nm thick and a thorough analysis of the SEM image (Fig. 5b) suggest a roughness of ~ 150-250 nm, thus at least one Raman spectrum should be obtained exclusively within the topmost NCD layer. In addition, the area fraction of the diamond Raman peak was evaluated from the spectra measured at the cross-section (Suppl. Fig. 3) and is presented in Fig. 5c.

Up to the third MCD layer ($z \approx 6.8-7.6$ µm), the area fraction of the diamond Raman peak directly gives the presence and absence of the diamond Raman peak, which can be interpreted as MCD and NCD, respectively (Fig. 5c, Suppl. Fig. 3). In the topmost NCD ($z \approx 7.6-8.4$ µm) layer, the diamond area fraction is not reduced to zero, indicating again the presence of an intermediate state between MCD and NCD.

The CSnanoXRD phase plot of the ML 1 film is presented in Fig. 5d, where exclusively the 111 DS ring of the diamond phase at $2\theta = 26.84$ deg. is present in the diffraction signal. Additionally, no other crystalline phases were detected for both the MCD and the NCD sublayers (Fig. 5d). Furthermore, a rather sharp diamond 111 diffraction peak was observed in the MCD sublayers, indicating a high portion of coarse crystalline diamond, in agreement with the SEM and Raman data presented in Fig. 5b and c, respectively. Contrary, low intensity of the diamond 111 reflection, an apparent high FWHM and an enhanced background signal was found when investigating the NCD sublayers at $z \approx 2.1-3.7$, $\approx 5.7-6.6$ and $\approx 7.6-8.4$ µm (Fig. 5d), in agreement with the presence of the trans-polyacetylene peaks and the partly absent diamond peak in the Raman signal (Fig. 5c, Suppl. Fig. 3) of the NCD sublayer.

The cross-sectional evolution of the azimuthal intensity distribution $I^{111}(\delta,z)$ retrieved from the ML1 film is presented in Fig. 5e. From the randomly oriented seed layer (between $z = 0$ and $0.3$ µm) a rapid transition to a (110) fibre texture was revealed, indicated by azimuthal maxima of the 111 DS ring at $\delta = \pm 90$ and $\pm 35.3$ deg. The further narrowing of the $I^{111}(\delta,z)$ distributions up to $z \approx 2.1$ µm indicates further sharpening of the (110) fibre texture during the competitive thin film growth [9]. Upon transition to the first NCD sublayer ($z \approx 2.1-3.7$ µm), the (110) fibre texture is greatly diminished, nearly a random distribution of azimuthal intensities is displayed, despite the interruptions given by the segmented CCD detector (Fig. 5e). However, a preferential orientation of the crystallites along the (110) directions can be deduced from the $I^{111}(\delta,z)$ distributions, in agreement with literature [7]. At $z \approx 3.7$ µm, again a rapid transition to a sharp (110) fibre texture is observed, similarly between $z \approx 5.7$ and $6.6$ µm, the (110) fibre texture

is again greatly diminished. However, in the thinner MCD layers between $z \approx 6.6-7.4$ and $8.4-9$ µm, apparently the sublayer thickness is too small for the evolution of a pronounced fibre texture (Fig. 5e). Especially in the topmost MCD sublayer, a rather random orientation of the diamond crystallites is detected (Fig. 5e).

Cross-sectional evolution of the FWHM and SCDD of the ML1 film is presented in Fig. 5f. Directly at the substrate/thin film interface, in the first MCD sublayer, a FWHM of ~ 0.25 deg. was evaluated from the 111 DS ring which decreases for both in-plane and out-of-plane orientations to ~ 0.15 deg. up to $z \approx 0.5$ µm and reflects the transition from the seed layer to the CVD grown MCD sublayer [19]. Again, in agreement to the results from the TL1 and TL2 films presented in Figs. 3f and 4f, respectively, as well as monolithic films [19] an oscillatory FWHM was detected for all four MCD sublayers (Fig. 5f). From $z \approx 2.1$ to $2.4$ µm the FWHM rises rapidly, reaching peak values of ~ 0.46 and 0.50 deg. for the in-plane and out-of-plane orientations, confirming a plate-like morphology in the NCD layer. Subsequently, between $z \approx 3.7-4.0$ µm, the FWHM rapidly decreases and oscillation around a mean value of ~ 0.15 deg. reappear in the adjacent MCD layer (Fig. 5f). Both the width of the plateau and the maximum FWHM are reduced to ~ 200 nm and ~ 0.39 deg., respectively, in the following NCD layer between $z \approx 5.7-6.6$ µm, which can be reasoned by the significantly reduced NCD sublayer thickness of ~ 0.9 µm. However, given the Raman spectrum evaluated at $z = 6.5$ µm (cf. Suppl. Fig. 3), this NCD sublayer still shows all characteristics associated with nanocrystallinity in literature [21-24,54,56-58,60-62]. It should be further mentioned that the apparent thickness of the NCD sublayer may be partly overshadowed by the interface roughness due to the faceted growth of the MCD layers (cf. Fig. 5b). In the adjacent MCD sublayer ($z \approx 6.6-7.6$ µm), no apparent changes of the FWHM compared to the layers located at $z \approx 0-2.1$ and $\approx 3.7-5.7$ µm could be detected (Fig. 5f). Finally, in the topmost NCD layer ($z \approx 7.6-8.4$ µm), no plateau width was discernible, and the peak width further decreased to ~ 0.35 deg. and 0.39 deg. for in-plane and out-of-plane orientations, respectively, with the concomitant presence of a diamond Raman peak in the spectrum recorded at $z = 8.0$ µm (cf. Suppl. Fig. 3). Thus, the minimum SCDD was evaluated with 18 and 16 nm for in-plane and out-of-plane orientation, respectively, for the presence of a distinctive diamond Raman peak (compare Fig. 5f and Suppl. 3). However, the CSnanoXRD qualitative peak width analysis may not fully capture the increase in grain size, as it doesn't solely account for grain size but also takes into consideration lattice defects expressed as strains of 2^nd^ and 3^rd^ order. Additionally, considering the MCD sublayers of the ML1 film at $z \approx 0-2.1$, $\approx 3.7-5.7$, $\approx 6.6-7.6$ and $\approx 8.4-9$ µm, (i) the in-plane and out-of-plane FWHM evolutions and magnitudes are almost identical and (ii) despite a small growth zone at the beginning of each MCD sublayer, the average FWHM is rather constant (Fig. 5f). This finding is contradictory to the columnar grain morphology observed by SEM (Fig. 5b), where a decreasing sublayer thickness goes hand in hand with a concomitant decrease in grain size. Thus, in agreement with the data presented for the TL1 and TL2 films presented in Fig. 3f and 4f, respectively, it can be concluded that the major contributions to the FWHM in the MCD sublayers down to a thickness of ~ 750 nm may originate in structural defects resulting in strains of 2^nd^ and 3^rd^ order, since the morphology is only partly reflected by the FWHM distributions presented in Fig. 5f.

In-plane residual stress evaluated from the ML1 film shows a rapid increase from 0 to ~ 1 GPa during transition from seed layer to CVD grown film between $z = 0$ and $0.5$ µm, respectively, with a further increase of $\sigma_r$ up to ~ 1.25 GPa at $z \approx 1.0$ µm (Fig. 5g). The rapid stress increase and the peak value is similar to $\sigma_r$ evaluated in the TL2 film from the NCD between 0 and $1.6$ µm (cf. Fig. 4g) and can be associated with a tremendous gradient intrinsic stress built up from to ~ 1.85 to 2.15 GPa at $z \approx 1.0$ µm, given the prior determined thermal stress magnitudes (cf. Sec. 3.1 and 3.2) of ~ -0.6 to -0.9 GPa for MCD. Up to $z \approx 2.1$ µm an oscillating $\sigma_r$ around an average of ~ 1 GPa was evaluated from the MCD layer, while the stress oscillation ceased throughout the
first NCD layer between $z \approx 2.1$ and 3.7 µm and the stress magnitude slightly decreased from ~ 1.0 to ~ 0.75 GPa, respectively. Given the prior determined thermal stress magnitudes (cf. Sec. 3.1 and 3.2) of ~ -0.1 to -0.4 GPa for NCD, the decrease of $\sigma_r$ amounts to a concomitant decrease of intrinsic stress from ~ 1.1-1.4 GPa to ~ 0.85-1.15 GPa, in good agreement with the stress magnitudes evaluated from the NCD layers of TL1 (Fig. 3g, Sec. 3.1) and TL2 films (Fig. 4g, Sec. 3.2). Contrary, the observed $\sigma_r$ levels in NCD between $z \approx 0$ and 2.1 µm contradict the findings from the TL1 and TL2 films, as well as literature [18,36], where a more compressive stress state was found for MCD in comparison with NCD. During growth of the second MCD layer, the oscillations of $\sigma_r$ remerged accompanied by a rapid stress decrease from ~ 1.25 to ~1 GPa between $z \approx 4$ to 4.7 µm. The stress decrease is located at similar z-positions, where strong oscillations of the FWHM were detected, indicating some changes in thin film morphology or density of lattice defects (cf. Fig. 5f). However, the $\sigma_r$ magnitude shifted towards 0 GPa at $z \approx 5.7$ µm. During the transition to the top two NCD layers at $z \approx 5.7-6.6$ and $\approx 7.6-8.4$ µm a rapid increase of $\sigma_r$ from 0 to ~ 1 GPa was evaluated from the (cf. Fig. 5g), which computes to and intrinsic stress level ~ 1.1-1.4 GPa (using calculated thermal stress magnitudes of ~ -0.1 to -0.4 GPa for NCD), similar to the TL2 film (cf. Fig. 4g). This behaviour was followed by an oscillatory decrease in the subsequent MCD layer ($z \approx 6.6-7.6$ µm) and an oscillatory stress profile in the top MCD layer ($z \approx 8.4-9.0$ µm). For the top layers (3.7 µm $< z < 9.0$ µm), more compressive and more tensile stress magnitudes were evaluated from the MCD and NCD, respectively, in agreement with data from monolithic MCD and NCD films [19], the TL1 (Fig. 3g) and TL2 (Fig. 4g) films and other multilayer films in literature [18,36]. Here, at least down to a sublayer thickness of ~ 750 nm, it appears that the absolute $\sigma_r$ and intrinsic magnitudes are highly dependent on the sublayer thickness and layer stacking (cf. Figs. 3g, 4g and 5g).

### 3.4. Diamond multilayer ML2

The cross-sectional microstructure of the ML2 film is presented in Fig. 6b. It shows four MCD layers interrupted by three NCD layers (Fig. 6b). The morphology of the first MCD sublayer with an individual layer thickness of ~ 2 µm appears similar compared to the trilayer films TL1 (Fig. 3b) and TL2 (Fig. 4b) as well as the multilayer film ML1 (Fig. 5b) investigated prior. In detail, a columnar grain morphology with increasing grain size was found up to a thickness of ~ 2.4 µm. Apparently, the change from MCD to NCD at a thickness of ~ 2.4 µm appears again to be instantaneous, like the other transitions from MCD to NCD at ~ 3.5 and 4.2 µm film thickness, in agreement with the other films presented in this study (cf. Figs. 3b, 4b and 5b). Thus, a high secondary nucleation rate [2,11], decoupling the microstructure from the film thickness can be assumed for the NCD sublayers. Evidently, the SEM confirms a multi-layered morphology, yet with highly diminished grain size in the MCD layers, whose thickness is reduced to below 500 nm.

The Raman spectrum obtained at the surface of the ML2 film presented in Fig. 2 shows a diamond Raman peak at $1332 \text{ cm}^{-1}$ of low intensity (compared to the signal of the ML1 and TL1 films) accompanied by a pronounced signal from the D and G-bands at 1350 and $1580 \text{ cm}^{-1}$, respectively, of sp² hybridised C located primarily at the grain boundaries [26,54,55,57,60]. Additionally, the presence of signal originating from the t-PA $\nu_1$ and $\nu_3$ bands visible at wavenumbers of 1150 and $1450 \text{ cm}^{-1}$, respectively, indicates presence of signal originating from NCD [21–23,58], confirming overlapping Raman spectra from the top MCD and NCD layers (Fig. 2, Fig. 6b) similar to the ML1 film discussed in Sec. 3.3. For further details about Raman spectra obtained from MCD and NCD sublayers Sec. 3.1 should be considered, furthermore the Raman spectra presented in this study are in good agreement with literature [21–24,54,56–58,60–62].

The Raman spectra evaluated from the cross-section of the ML2 film are presented in Suppl. Fig. 4. Up to $z = 2.0$ µm the spectra consist of a distinctive diamond Raman peak at $1332 \text{ cm}^{-1}$ of high intensity

accompanied by signal from the D and G-bands at 1350 and $1580 \text{ cm}^{-1}$, respectively, of sp² hybridised C located primarily at the grain boundaries [26,54,55,57,60] and thus the spectra presented in Suppl. Fig. 4 at $z = 0.2, 0.8, 1.4$ and 2.0 µm fit to a polycrystalline diamond with high sp³ to sp² C ratio [26]. Contrary, the spectra obtained at $z = 2.6, 3.2, 3.8$ and 4.4 µm consist of a diamond peak of low intensity, accompanied by comparably high signal originating from D and G-bands at 1350 and $1580 \text{ cm}^{-1}$, respectively, indicating sp² hybridised C as well as peaks at 1150 and $1450 \text{ cm}^{-1}$ originating from the $\nu_1$ and $\nu_1$ t-PA, respectively. This is a clear indication for an overlapping signal originating in NCD and MCD, which can be related to the low sublayer thicknesses (cf. Fig. 6b). The area fraction the diamond Raman peak was evaluated from the spectra obtained at the cross-section of the ML2 film (cf. Suppl. Fig. 4) and is presented in Fig. 6c. In agreement with the Raman spectra obtained from the surface (Fig. 2) and the cross section (Suppl. Fig. 4), as well as the morphology uncovered by SEM (Fig. 6c) up to $z = 2.0$ µm, high diamond Raman area fractions were evaluated. Beginning at $z = 2.0$ µm the diamond Raman area fraction decreases, while at $z > 2.4$ µm the diamond Raman peak area fraction oscillates at a low average value (Fig. 6c), in agreement with the Raman spectra in Suppl. Fig. 4 and the SEM cross-section of the ML2 sample. Since the individual sublayer thicknesses are in the order of the laser spot of ~ 450 nm, only an intermediate state between NCD and MCD is shown by the diamond Raman area fraction.

Again, the CSnanoXRD phase plot exclusively shows the diamond 111 peak (cf. Fig. 6d) in case of the ML2 film. Up to $z = 2.4$ µm, a sharp diamond peak at $2\theta = 26.84$ deg. is present in the diffraction signal. At thicknesses larger than $z = 2.4$ µm, the decreased MCD and NCD sublayer thickness is particularly reflected by the phase plot presented in Fig. 6d, since the intensity and apparent FWHM are only fluctuating between $z = 2.4$ and 4.7 µm, thus no individual layers can be distinguished in contrast to the diffraction data of the other multilayer films presented in Figs. 3d, 4d and 5d.

Similarly, up to $z = 2.4$ µm a gradually evolving (110) fibre texture can be determined from the azimuthal intensity distribution $I^{111}(\delta, z)$ retrieved from the ML2 film (cf. Fig. 6e), which is expressed by intensifying azimuthal maxima of the 111 DS ring at $\delta = \pm 90$ and $\pm 35.3$ deg. with increasing film thickness. Between $z = 2.4$ and 4.7 µm, also an undulating, less pronounced (110) fibre texture can be determined from the azimuthal intensity distribution $I^{111}(\delta, z)$. In detail, the undulations are represented by a rather diminished (110) fibre texture, where more continuous $I^{111}(\delta, z)$ was evaluated between $z \approx 2.5-3.0$, $\approx 3.4-3.8$ and $\approx 4.2-4.5$ µm, while azimuthal maxima at $\delta = \pm 90$ and $\pm 35.3$ deg. are somewhat amplified between $z \approx 3.0-3.4$, $\approx 3.8-4.2$ and $\approx 4.5-4.7$ µm. However, vertical line features of different length are present throughout $z = 2.4$ and 4.7 µm in Fig. 6e at random azimuthal orientations $\delta$, which can be related to diffraction from individual crystallites. This is in clear contrast to the NCD sublayers of the other multilayer films, where no line features representing individual crystallites were found and the slight (110) fibre texture was detected from the continuous azimuthal distribution of the diffraction signal (cf. Figs. 3e, 4e and 5e).

Cross-sectional evolution of the FWHM and SCDD of the ML2 film is presented in Fig. 6f. Directly at the substrate/thin film interface, in the first MCD sublayer, a FWHM of ~ 0.25 deg. was evaluated from the 111 DS ring, which decreases for both in-plane and out-of-plane orientations of the diffraction vector to ~ 0.15 deg. up to $z \approx 0.5$ µm and reflects the transition from the seed layer to the CVD grown MCD sublayer [19]. The in-plane FWHM stays rather constant at a value of ~ 0.15 deg. up to $z \approx 2.4$ µm, while the out-of-plane FWHM starts to increase from $z \approx 1.6$ µm, which may be related to the low intensity at $\delta = 0$ deg. as a consequence of the emerging (110) fibre texture (cf. Figs. 6e). FWHM evaluated from the z-positions where a more continuous $I^{111}(\delta, z)$ was found, showed decreasing maximum in-plane and out-of-plane FWHM (cf. Fig. 6f). In detail, the maximum in-plane FWHM between $z \approx 2.5-3.0$ µm, $\approx 3.4-3.8$ µm and $\approx 4.2-4.5$ µm, was evaluated as 0.38, 0.33 and 0.26
<figure><img src="image_10.png" /><figcaption>**Fig. 6.** Cross-sectional characterization of a diamond multilayer thin film with a total film thickness of 4.7 µm composed of 4 MCD layers interrupted by 3 NCD sublayers. In (a) the applied methane concentration is shown, while in (b), (c) and (d-g) the results retrieved by SEM, Raman spectroscopy and CSnanoXRD are presented, respectively. In detail, four layers of MCD exhibiting are interrupted by three layers of NCD as evidenced by SEM (b). The difference between the MCD layer (up to *z* ≈ 2.4 µm) and the multilayered structure are highlighted by the fitted diamond Raman peak area fraction, where the MCD layer exhibits a distinctive high diamond Raman signal, while the multilayered structure shows a significantly diminished area fraction of the diamond Raman peak (c). In (d) the CSnanoXRD phase plot exclusively the diamond 111 reflection is present, here the peak intensity and apparent peak width vary, yet the differentiation is less clear compared to the other films (Figs. 2d, 3d and 4d). Azimuthal intensity distributions reveal a gradually intensifying (110) fibre texture in the MCD layer up to *z* ≈ 2.4 µm, which is diminished in the layers above. Though azimuthal gradients are still present in the 6 alternating MCD and NCD layer, no pronounced texture or cross-sectional variation thereof is present (e). The cross-sectional FWHM variation (f) reveals a gradual decrease of the oscillatory FHWM to ~ 0.15 deg. with increasing MCD sublayer thickness up to *z* ≈ 2.4 µm, while the FWHM in the 6 alternating MCD and NCD layer oscillates between 0.15 and 0.45 deg., respectively. In detail, the peak FWHM decreases with decreasing layer thickness from ~ 0.42 to ~ 0.27 deg. (f). The residual stress evaluated from the 111 Debye-Scherrer Ring is presented in (g) and shows an oscillatory variation around *σ*<sub>r</sub> ~ 0.75 GPa throughout the entire film thickness. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
<caption>**Table 1**</caption>

Local maximal FWHM<sup>max</sup> (z) and minimal FWHM<sup>min</sup> (z) evaluated from the individual diamond films, as well as the corresponding minimum and maximum SCDDs, respectively. In brackets the cross-sectional position of the observed peak stress values is given.

<table><thead><tr><th></th><th>FWHM<sup>min</sup> (z)<br>[deg]</th><th>SCDD<sup>max</sup> (z)<br>[nm]</th><th>FWHM<sup>max</sup> (z)<br>[deg]</th><th>SCDD<sup>min</sup> (z)<br>[nm]</th></tr></thead><tbody><tr><td>TL1</td><td>≈0.11–0.15<br>(7.5–10.5 µm)</td><td>≈40–60<br>(7.5–10.5 µm)</td><td>≈0.6<br>(5 µm)</td><td>≈10–15<br>(5 µm)</td></tr><tr><td>TL2</td><td>≈0.15<br>(2–5 µm)</td><td>≈40<br>(2–5 µm)</td><td>≈0.48<br>(0.5–1.5 µm)</td><td>≈15<br>(0.5–1.5 µm)</td></tr><tr><td>ML1</td><td>≈0.15<br>(4.0–5.6 µm)</td><td>≈40<br>(4.0–5.6 µm)</td><td>≈0.51–0.53<br>(2.5–3.5 µm)</td><td>≈12<br>(2.5–3.5 µm)</td></tr><tr><td>ML2</td><td>≈0.15<br>(0.3–2.0 µm)</td><td>≈40<br>(0.3–2.0 µm)</td><td>≈0.45<br>(2.5 µm)</td><td>≈14<br>(2.5 µm)</td></tr></tbody></table>

deg., respectively, while the out-of-plane FWHM at the same positions was evaluated as 0.46, 0.36 and 0.33 deg., respectively. Furthermore, in the layers with amplified texture located at z ≈ 3.0–3.4, ≈ 3.8–4.2 and ≈ 4.5–4.7 µm, the minimum in-plane FWHM at those z-positions was evaluated with 0.17, 0.19 and 0.20 deg., respectively, indicating a slight increase of the FWHM with further decreasing layer thickness. Additionally, no oscillatory FWHM was evaluated at those positions. This finding contrasts with the FWHM evaluated from the MCD layers in the ML1 film (cf. Fig. 5f), where no dependence of the FWHM was found (down to a layer thickness of ~ 750 nm, Sec. 3.3). The decreasing and increasing FWHM in the individual sublayers indicate furthermore an intermediate state between MCD and NCD, similar to the results from the CSnanoXRD phase and texture plots presented in Fig. 5d and e, respectively. However, due to the faceted growth of MCD up to a thickness of 2.4 µm induces a significant interface roughness (cf. Fig. 6b), which may corroborate the CSnanoXRD and Raman spectroscopy results.

In-plane residual stress evaluated from the 111 DS ring in the ML2 film is presented in Fig. 6g. Generally, throughout the cross-section of the ML2 film, an oscillatory stress profile undulating around an average $\sigma_r$ magnitude of ~ 0.75 GPa was evaluated from the CSnanoXRD data. Therefore, when using the thermal stress component calculated in Sec. 3.1 as ~ -0.6 to ~ -0.9 GPa, an intrinsic stress of ~ 1.35 to 1.65 GPa can be evaluated for the MCD up to z ≈ 2.4 µm. In detail, stress maxima with magnitudes of ~ 1.0 and 1.1 GPa were found at z ≈ 2.4 and 3.6 µm, respectively, at the approximate positions of the FWHM maxima (cf. Fig. 6f), though no such peak was found at z ≈ 4.3 µm, the position of the topmost FWHM maximum. The peaks may be attributed to the change of thermal stress component, which was calculated in Sec. 3.1 as ~ -0.6 to ~ -0.9 GPa and ~ -0.1 to ~ -0.4 GPa for MCD and NCD, respectively. In following, when excluding a change of the intrinsic stress magnitude, a shift of $\sigma_r$ of up to 0.5 GPa towards the tensile may be expected, which is in good agreement with the experimental data presented in Fig. 6g. Remarkably, the average $\sigma_r$ magnitudes of ~ 0.75 GPa differs quite notably from the TL2 (Fig. 4g) and ML1 (Fig. 5g) films, while the stress magnitudes are more close to the TL1 film (Fig. 3g), where average stresses of 0.6 and ~ 0.3 GPa were evaluated for the NCD and MCD layers, respectively.

## 4. Summary of observed cross-sectional effects

The results reported in Figs. 3–6 allow us to draw valuable conclusions on the common gradual effects occurring within individual NCD and MCD sublayers as well as on the transition phenomena occurring at the interfaces between these sublayers. It is interesting to compare common trends within the NCD and MCD sublayers, which possess different crystallinity, grain size, crystallographic texture, residual stress and crystalline diamond phase fraction.

The most obvious are changes and trends in the observed cross-sectional effects can be best interpreted by the variations of FWHM and residual stress. First it should be noted, that at the interface between

Si and the multilayer diamond thin films, the FWHM magnitude is ~ 0.2–0.25 deg. (cf. Figs. 3f, 4f, 5f and 6f), independent of the subsequent growth of MCD (TL1, ML1 and ML2) or NCD (TL2). This effect can be correlated to the growth of the multilayer diamond films from the prior applied diamond seed crystals (cf. Sec. 2.1). In case of the NCD sublayers, the FWHM decreases with decreasing layer thickness in NCD as can be seen from Figs. 3f, 4f, 5f and 6f for the TL1, TL2, ML1 and ML2 films, respectively. Additionally, the diamond 111 FWHM values and their respective cross-sectional positions are presented in Table 1. At sublayer thicknesses of 750 nm and below, the increased SCDD concurs with the presence of a distinctive diamond Raman peak in the obtained Raman spectra in agreement with literature [24,58,61].

Contrary, the FWHM of the MCD sublayers remains unaffected down to a sublayer thickness of 750 nm, as shown in Figs. 3f, 4f, 5f and 6f for the TL1, TL2, ML1 and ML2 films, respectively. Only at sublayer thicknesses of 500 nm and below, an increase of the FWHM was observed with decreasing sublayer thickness. Thus, the minimum FWHM<sup>min</sup> (z) and the respective maximum SCDD<sup>max</sup> (z) were found at ≈ 0.15 deg. and ~40–60 nm respectively. Here it should be noted that the SCDD<sup>max</sup> (z) evaluated from the CSnanoXRD data is highly divergent from the grain size observed by SEM for all investigated diamond films (cf. Figs. 3b, 4b, 5b and 6b). It can be concluded that the FWHM in the MCD is mainly governed by lattice defects expressed as strains of 2<sup>nd</sup> and 3<sup>rd</sup> order.

Overall, the residual stress analysis from all films confirmed that the stress magnitudes of MCD are shifted mostly towards more compressive stress compared to NCD films (cf. Figs. 3g, 4g, 5g and 6g) in agreement with literature [18,19,36]. The thermal stress component calculated from the CTE mismatch of diamond [14,65] and Si [63] ranges from -0.6 to -0.9 GPa in the case of MCD. Additionally, the CTE of NCD [14] and amorphous C [66,67] depends highly on the amount of sp<sup>2</sup>-hybridised C—C bonds, yielding thermal stress of ~ -0.1 to ~ -0.4 GPa and up to 1 GPa, respectively, supporting the conclusion of more compressive stress in MCD compared to NCD. However, given the experimental stress values obtained from the diffraction data of the TL1, TL2, ML1 and ML2 films in Figs. 3g, 4g, 5g and 6g, respectively, it appears that the overall residual stress magnitudes are not only influenced by the layer stacking but also the individual sublayer thickness. To further support this observations, the peak residual stress values and their respective cross-sectional positions are presented in Table 2.

<caption>**Table 2**</caption>

Localized maximal $\sigma_r^{max}(z)$ and minimal $\sigma_r^{min}(z)$ residual stress evaluated from the individual diamond films. In brackets the cross-sectional position of the observed peak stress values is given.

<table><thead><tr><th></th><th>σ<sub>r</sub><sup>max</sup>(z)<br>[GPa]</th><th>σ<sub>r</sub><sup>min</sup>(z)<br>[GPa]</th></tr></thead><tbody><tr><td>TL1</td><td>~0.6 GPa (3.0–7.2 µm)</td><td>-0.3 (1.2 µm)</td></tr><tr><td>TL2</td><td>1.31 GPa (2.5 µm)</td><td>-0.24 (4.6 µm)</td></tr><tr><td>ML1</td><td>1.35 (2.0 µm)</td><td>-1.0 (4.8 µm)</td></tr><tr><td>ML2</td><td>0.98 (2.3 µm)</td><td>0.33 (0.8 µm)</td></tr></tbody></table>
Additionally, it should be noted that independent of the residual stress evolution during growth), directly at the Si to diamond interface, the residual stress level is in a narrow range of 0.25–0.5 GPa (cf. Figs. 3g, 4g, 5g and 6g). This narrow range gives another indication that the evolution of the Si-diamond interface is dominated by the growth from the diamond seed crystals (cf. Sec. 2.1), as documented by the FWHM data discussed above.

Finally, even the higher spatial resolution (cf. Sec. 2.4) employed for the CSnanoXRD analysis of the ML2 thin film does not allow to discern fully between NCD and MCD layers alternating on a 2 µm thick MCD base layer (Fig. 6d-g), thus showing the cross-sectional limits for monophase diamond multilayers deposited by HF-CVD on Si. However, it has to be acknowledged that a NCD base layer with a lower interface and surface roughness may be beneficial for a further in-depth analysis of nanostructured diamond multilayers.

## 5. Conclusions

In this work, the gradual development of microstructure and residual stress in four multilayered diamond thin films of different sublayer thicknesses prepared by chemical vapour deposition was investigated by SEM, Raman Spectroscopy and CSnanoXRD.

The experimental results retrieved by Raman spectroscopy allow to discriminate between MCD and NCD down to an individual layer thickness of 1 µm, with overlapping spectra obtained for lower sublayer thickness. Cross-sectional X-ray nanodiffraction on the multilayer films revealed a highly different sublayer microstructure evolution during thin film growth, accompanied by different residual stress gradients, where the clear distinction between MCD and NCD is not possible below a layer thickness of 500 nm in the chosen layer stacking and experimental geometry. Furthermore, the sublayer thickness and layer stacking appear to have a critical influence on the residual stress magnitudes of monophase diamond thin films, which may influence their applicability.

## CRediT authorship contribution statement

**M. Meindlhumer:** Writing – original draft, Visualization, Validation, Investigation, Data curation, Conceptualization. **J. Grau:** Investigation, Data curation. **H. Sternschulte:** Investigation, Data curation. **Y. Halahovets:** Investigation. **P. Siffalovic:** Resources, Investigation. **M. Burghammer:** Resources, Investigation. **D. Steinmüller-Nethl:** Resources, Project administration. **J. Keckes:** Writing – review & editing, Supervision, Project administration, Funding acquisition, Conceptualization.

## 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.

## Data availability

Data will be made available on request.

## Acknowledgement

Financial support for this work was provided by Österreichische Forschungsförderungsgesellschaft mbH (FFG), project “CrossDiamond” (No. 880566).

The provision of synchrotron facilities by the ESRF (European Synchrotron Radiation Facility) is gratefully acknowledged by M. M. and J. K. J.G. and H.S. are thankful to Sven Ulrich and Steffen Mumbauer from the Institute for Applied Materials (IAM-AWP) of the Karlsruhe Institute of Technology for the possibility to use their Raman spectrometer and

their helpful assistance. The financial support by the HSA project resource “Diamant 3D” and by the Federal Ministry for Economic Affairs and Climate Action (BMWK) on the basis of a decision by the German Bundestag is gratefully acknowledged.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.matchar.2024.113973.

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