# MBE growth of GaSb on Ge-based templates grown on Si for mid-infrared photonics [Invited]

Milan Silvestre, Sébastien Lacaze, Audrey Gilbert, S Calcaterra, A de Cerdeira Oliveira, D Impelluso, G Isella, J-B Rodriguez, Eric Tournié

► To cite this version:

Milan Silvestre, Sébastien Lacaze, Audrey Gilbert, S Calcaterra, A de Cerdeira Oliveira, et al., MBE growth of GaSb on Ge-based templates grown on Si for mid-infrared photonics [Invited]. Optical Materials Express, 2025, 15 (10), pp.2499. 10.1364/ome.569432. hal-05283375

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# MBE growth of GaSb on Ge-based templates grown on Si for mid-infrared photonics [Invited]

M. SILVESTRE,<fnref n="1" />,<fnref n="2" /> S. LACAZE,<fnref n="1" /> A. GILBERT,<fnref n="1" /> S. CALCATERRA,<fnref n="3" /> A. DE CERDEIRA OLIVEIRA,<fnref n="3" /> D. IMPELLUSO,<fnref n="3" /> G. ISELLA,<fnref n="3" /> J.-B. RODRIGUEZ,<fnref n="1" /> AND E. TOURNIÉ<fnref n="1" />,<fnref n="4" />,*

**Abstract:** We study the molecular beam epitaxy of GaSb layers on Ge-based templates grown on Si substrates by low-energy plasma-enhanced chemical vapor deposition. First, we study the growth mechanism to obtain a GaSb surface free of emerging antiphase boundaries on Ge substrates, and then we transfer this mechanism to Ge-based templates. Our investigations reveal that this mechanism is hindered by the crosshatch pattern formed during the Ge-based template growth. We demonstrate that using low-miscut (1°) Si substrates or chemical-mechanical polishing mitigates the effects of the crosshatch pattern, enabling the growth of GaSb layers free of emerging antiphase boundaries. These results represent a further step towards the monolithic integration of III-V on mid-infrared photonic platforms.

Published by Optica Publishing Group under the terms of the [Creative Commons Attribution 4.0 License](https://creativecommons.org/licenses/by/4.0/). Further distribution of this work must maintain attribution to the author(s) and the published article's title, journal citation, and DOI.

## 1. Introduction

Over the past two decades, silicon photonics research has enabled the development of low-cost, chip-scale photonic integrated circuits (PICs) that meet the demands of the telecommunications and data processing markets [1]. These applications focus on the near infrared, since silicon absorbs above ~8 µm [2]. The introduction of germanium-based materials is a natural way to extend the transparency window of silicon photonic devices into the mid-infrared (MIR) region [2]. These MIR photonic devices open the door to applications in various fields, including chemical hazard detection [3], astrophotonics [4] and spectroscopy [5,6]. For these applications, however, germanium-based passive devices [7-9] must be hybridized with efficient MIR light sources such as III-V lasers [10]. However, this technology brick is still missing.

Several techniques have been developed for hybridizing III-V lasers on group IV semiconductor platforms, including wafer bonding [11], transfer printing [12] and direct growth [13]. The latter option is an early-stage technology which will enable dense integration without incurring the cost and size limitations imposed by III-V substrates. However, several challenges must be overcome to achieve high quality material. III-V growth on group IV substrates is plagued by the formation of a high density of defects such as antiphase boundaries (APBs) [14], threading dislocations [15] and thermal cracks [16]. APBs appear at the interface between two polar domains formed during nucleation of III-V compounds on group IV substrates [17]. These 2D defects consist in III-III or V-V bonds that act as electrically charged paths, creating shortcuts that can kill devices. H. Kroemer [14] initially proposed to use bi-atomic stepped surfaces of group IV substrates in order to promote the growth of one of the two III-V polar domains and avoid APB formation. One way to obtain these surfaces is to use substrates that are highly misoriented (>4° miscut angle). However, these substrates are not compatible with the silicon industry. Recently, a new

<footnote n="1">¹IES, University of Montpellier, CNRS, F - 34000 Montpellier, France</footnote>

<footnote n="2">²RIBER SA, F - 95870 Bezons, France</footnote>

<footnote n="3">³L-NESS, Politecnico di Milano Dipartimento di Fisica, I - 22100 Como, Italy</footnote>

<footnote n="4">⁴Institut Universitaire de France (IUF), F-75005 Paris, France</footnote>

<footnote>*eric.tournie@umontpellier.fr</footnote>
perspective on the formation and annihilation of III-V APBs on low-miscut silicon substrates was proposed [18]. Experimental demonstrations were performed using molecular beam epitaxy (MBE) to grow GaSb [19] and GaAs [20] on Si with a miscut 0.18-0.5° oriented toward a [110] direction. In brief, the silicon substrate is first deoxidized and annealed at high temperature to obtain a homogeneous surface in which all steps are parallel to each other and perpendicular to the miscut direction. This organization is then transferred to the III-V layer during growth, which is the key point. One polarity, the main phase domain (MPD), grows faster than the other, the anti-phase domains (APD) and progressively overgrows and buries the APDs after a few hundred nm. This leaves a single-domain surface that is free of emerging APBs and suitable for growing devices.

In this work, we study the MBE growth of GaSb layers on Ge substrates and on various Ge-based templates grown on Si substrates using low energy plasma-enhanced chemical vapor deposition (LEPCVD). This is the first step toward hybridizing MIR antimonide-based lasers with Ge-based MIR PICs through direct growth. First, we determine the necessary template specifications needed to enable APD burying. Then, we demonstrate the growth of thin GaSb layers free of emerging APBs on these Ge-based photonic platforms. Interestingly, we confirm that the APD burying mechanism is the same on Ge than on Si.

## 2. Experiments

GaSb growths were performed on (001)-oriented Ge substrates with a miscut of 1° or 0.5° toward a [110] direction, as well as on Ge-based templates grown by LEPECVD [21] on silicon substrates with a miscut of 0.2°, 0.5° or 1° toward a [110] direction. The GaSb layers were grown in a solid-source RIBER C21E III-V MBE chamber equipped with valved cracker cells for group V elements. This chamber is connected to an automated cluster that enables the transfer of samples from a preparation chamber to the MBE under ultra-high vacuum conditions (~1 × 10⁻⁹ Torr). All GaSb growths were performed at 530°C (thermocouple), and the V/III growth rate ratio was kept constant around 2. We assessed the surface morphologies using a Brucker atomic force microscope (AFM) Dimension 3100 in tapping mode. AFM images of 5 × 5 µm² and 20 × 20 µm² were analyzed using Gwyddion software to determine roughness root mean square (RMS), APD surface coverage and MPD/APD height difference following the approach described in Ref. [19]. We determined threading dislocation densities (TDDs) using electron channeling contrast imaging (ECCI) in a JEOL JSM-IT800 scanning electron microscope (SEM) equipped with a standard annular pole piece mounted backscatter electron (BSE) detector. High resolution X-ray diffraction (HR-XRD) measurements were performed using a PANalytical X'Pert3 MRD instrument with a PIXcel¹D linear detector, a four-bounce asymmetric Ge(220) monochromator, and an X-ray tube that provides Cu Kα₁ radiation.

## 3. Results and discussion

We studied the growth mechanisms on Ge substrates before turning to the growth on Ge-based templates.

### 3.1. GaSb growth on Ge substrates

The first issue is substrate annealing prior to III-V nucleation. The temperature should be high enough to enable deoxidation and surface organization, but low enough to avoid surface degradation. Figure 1 shows AFM topographic images of 500-nm thick GaSb films grown on Ge substrates with a 1° miscut in the [110] direction and annealed for one hour at three different temperatures (thermocouple values). Based on previous AFM studies related to the growth of GaSb on Si, the black areas seen in the images are attributed to APDs [19]. The high APD surface coverage observed in the case of a Ge substrate annealed at 750°C before GaSb growth
(Fig. 1(a)) is attributed to insufficient thermal energy to promote surface organization. This affects the APD burying mechanism during GaSb growth. Conversely, annealing at 825°C is too high (Fig. 1(c)), likely due to thermal roughening of the surface. The optimal preparation temperature is then 790 °C, resulting in the smallest surface covered by APDs.

<figure><img src="image_2.png" /><figcaption>**Fig. 1.** 10 × 10 µm² AFM images of 500 nm of GaSb layers grown on Ge substrates annealed under UHV at a) 750°C, b) 790°C and c) 825°C. Black and white domains respectively correspond to antiphase domains and main phase domains. 790°C is the optimal annealing temperature</figcaption></figure>

Next, GaSb layers of different thicknesses were grown on 0.5° miscut germanium substrates to investigate whether the burying mechanism is similar to that observed on silicon, where progressive burying occurred due to different MPD and APD growth rates [19]. Figure 2(a) shows AFM images taken of samples with GaSb layers in the 250- to 500-nm thickness range. Progressive burying was also observed, with only a few APDs visible after 500 nm of GaSb was grown.

<figure><img src="image_3.png" /><figcaption>**Fig. 2.** a) 20 × 20 µm AFM images of 250, 300, 400 and 500 nm GaSb sample grown on Ge substrate. b) Mean height difference between APD and MPD in function of the GaSb deposited thickness.</figcaption></figure>
The difference in mean height between the MPD and the APD was calculated using the procedure described in Ref. [19]. This difference increases linearly with the nominal thickness of GaSb (Fig. 2(b)). This confirms that the MPD and the APD grow at different rates, as was observed in the growth of GaSb on Si. The slope of the curve in Fig. 2(b) indicates that the growth rate imbalance is ~4%, meaning the MPD grows 4% faster than the APD. This result aligns with the value obtained for GaSb grown on Si [19]. This indicates that the APD burying mechanism is independent of the group IV substrate. This finding is consistent with the burying model framework, which is based on GaSb domain growth kinetics rather than substrate composition.

This study demonstrates that an optimized annealing of Ge substrates and dedicated nucleation strategies enable APD burying during molecular-beam epitaxy of GaSb layers, and that the APD-burying mechanism is the same on Si and Ge substrates, as expected from the theoretical model. This paves the way for growing emerging APD-free GaSb layers on more complex Ge-based templates prepared on Si substrates.

## 3.2. GaSb growth on Ge-based templates

### 3.2.1. GaSb growth on as-grown Ge/SiGe/Si templates

The next stage of the study was to examine the MBE growth of GaSb layers on Ge-based templates. These were grown by LEPECVD on (100)-oriented Si substrates with a 0.5° or 1° miscut toward a [110] direction. These templates begin with a gradual composition layer that evolves progressively from Si to Ge at rates of 16.7%Ge/µm or 7% Ge/µm. This layer is then covered by a pure Ge layer, either 100 nm or 1 µm thick. The total thickness of the templates is therefore 6 µm or 15 µm. Similar templates can be used to process low-loss SiGe waveguides for MIR photonics [9]. Growing a GaSb buffer layer free of emerging APDs on such templates is the first step toward the direct integration of MIR laser on Ge-based MIR photonic platforms. Figure 3 shows the template structures and AFM images of the surfaces before and after the growth of 500 nm of GaSb.

<figure><img src="image_4.png" /><figcaption>Fig. 3. i) epitaxial stack of Ge-based templates, ii) 20 × 20 µm AFM images of the Ge-based template before GaSb growth, iii) 20 × 20 µm AFM images after the growth of 500 nm of GaSb.</figcaption></figure>

The Ge layers on the surface of the templates exhibit crosshatch patterns (Fig. 3), which are attributed to the strain relaxation of the SiGe layers grown on Si [22], [23]. Despite the differences in layer composition, the two templates have comparable roughness RMS. After growing 500 nm of GaSb, Fig. 3(a)) iii) shows a surface with morphological patterns similar to
those observed on the Ge-based template, and APDs are visible. In contrast, the sample shown in Fig. 3(b) iii) exhibits a smooth surface with very few emerging APDs.

To understand the different behaviors, the AFM images of both templates were further analyzed. Figure 4 shows that the crosshatch pattern on the template’s surface creates random local slopes despite the template’s low overall roughness.

<figure><img src="image_5.png" /><figcaption>Fig. 4. Profile lines taken in three directions on the crosshatched surface. Slopes are created by local roughness in all directions. Steepness of these slopes are characterized by θ<sub>local</sub>, the angle between the vector normal to the surface and the vector normal to the substrate.</figcaption></figure>

At each point on the surface, the local angle, θ<sub>local</sub>, can be defined as the angle between the vector normal to the pattern and the vector normal to the nominal substrate surface (see Fig. 4). In other words, θ<sub>local</sub> represents the surface’s local inclination due to roughness. The normalized distribution of θ<sub>local</sub> can be extracted from the entire AFM image using Gwyddion software [24] and compared to the miscut angle of the substrate. To quantify the fraction of the surface with a local inclination greater than the miscut value (θ<sub>miscut</sub>), we calculate I(θ<sub>miscut</sub>) using the following equation:

$$I(\theta_{\text{miscut}}) = \int_{\theta_{\text{miscut}}}^{90^{\circ}} \text{Surface angle distribution } d\theta$$

Figure 5 shows the surface inclination distributions of the two templates presented in Fig. 3.

Figure 5 shows that, in the first case, 33.1% of the initial Ge surface has a local inclination greater than that induced by the substrate miscut. Consequently, GaSb growth in these areas follows the local inclination rather than the miscut direction, transferring the initial crosshatch pattern to the GaSb surface, as illustrated in Fig. 3(a)(iii). This prevents the APD burying process, which requires a surface on which GaSb steps can develop in only one direction imposed by the miscut [19]. In the second case, when θ<sub>miscut</sub> = 1°, only 0.7% of the initial Ge surface has local slopes greater than θ<sub>miscut</sub>. This value is low enough to permit unidirectional GaSb step-flow across the entire surface of the sample. Figure 3(b) iii) shows the morphology of the surface after the GaSb growth, with only 0.36% of the surface still covered by APD, indicating a successful burying process. These results demonstrate the impact of crosshatch patterns on Ge-based templates on the APD burying mechanism during GaSb growth. However, a miscut angle of 1° in the [110] direction mitigates this issue by providing a local inclination greater than that of the crosshatch pattern.

### 3.2.2. GaSb growth on Ge/Si templates

The crosshatch pattern is particularly pronounced after SiGe growth on Si. Therefore, pure Ge layers grown on Si may be seen as an alternative, and such templates have already been used
<figure><img src="image_6.png" /><figcaption>**Fig. 5.** Normalized histogram of θ<sub>local</sub> values extracted from all pixels of AFM images of the crosshatch patterns on the surface of the templates.</figcaption></figure>

for MIR photonics circuits [8]. Thus, we grew two 4-µm-thick pure-Ge layers on Si substrates with miscuts of 0.27° and 1° toward a [110] direction. The template grown on the 0.27° miscut substrate underwent an additional chemical-mechanical polishing (CMP) step to remove any remaining surface roughness. Figure 6 shows the template structures and AFM images of the surfaces before and after growing 500 nm of GaSb.

<figure><img src="image_7.png" /><figcaption>**Fig. 6.** i) epitaxial stack of Ge-based templates, ii) 20 × 20 µm AFM images of the Ge-based template before GaSb growth, iii) 20 × 20 µm AFM images after the growth of 500 nm of GaSb.</figcaption></figure>

Figure 6(a) i) shows that if a crosshatch pattern is visible on the non-polished 1° template, its amplitude is much lower than that of Ge on SiGe templates. Conversely, the polished template is perfectly flat. In both cases, the roughness RMS is drastically reduced compared to SiGe-containing templates, and APDs are almost buried after 500 nm GaSb growth. Image Fig. 6(b) iii) shows that the remaining APDs are perpendicularly aligned, which could be reminiscent of a
crosshatch morphology at the template level. However, since no crosshatch pattern was observed on the surface of the polished template, we propose that a faint crosshatch pattern reappeared during sample annealing prior to III-V growth, as we have observed on other samples (not shown). Figure 7 shows the surface inclination distributions of the templates presented in Fig. 6.

<figure><img src="image_8.png" /><figcaption>**Fig. 7.** Comparison of local angle distribution of a Ge surface after chemical-mechanical polishing (CMP) and an as grown Ge surface with crosshatch patterns</figcaption></figure>

Figure 7 shows that the faint crosshatch pattern on the unpolished template creates a local inclination much lower than the 1° miscut value. Therefore, it should not interfere with GaSb growth. This is confirmed by the smooth and APD-free GaSb surface obtained after growth (Fig. 6(a)) iii). Furthermore, Fig. 7 shows that after polishing any local inclination on the Ge surface is lower than the initial, very low, 0.27° miscut. Again, an APD-free GaSb surface is obtained (Fig. 6(b)) iii). These results show that GaSb growth on Ge-on-Si templates is facilitated by reducing the magnitude of the crosshatch pattern. CMP can further reduce local inclination and thus enable GaSb growth on templates with a very low miscut.

### 3.3. Comparison of APB-free GaSb layers grown on Ge-based photonic platforms and GaSb grown on Si or Ge substrates

In this section, we compare the quality of the best GaSb layers grown on Ge-on-SiGe-on-Si templates (see Section 3.2.1) and Ge-on-Si templates (see Section 3.2.2), with that of GaSb layers grown on Si and Ge substrates. To evaluate the material quality, AFM, electron channeling contrast imaging (ECCI) and high-resolution X-ray diffraction (HR-XRD) measurements of these samples were carried out and the results are summarized in Table 1.

The four samples exhibit very few emerging APD (<1%) and low roughness RMS values, below 2 nm. Sample d), grown on a Ge/SiGe/Si template exhibits the highest roughness. This is due to the ripples visible in Fig. 3(b) iii) which are related to the crosshatch pattern present on the initial surface of the template.

To further investigate the crystalline quality of the samples, HRXRD and ECCI measurements were carried out. Figure 8 shows the vertically shifted ω-2θ scans of the four samples.

Figures 8(a) & (b) correspond to GaSb layers grown on Si and Ge substrates with 0.5° miscut. The Si and Ge substrate peaks are aligned with their theoretical Bragg angles. The positions of the GaSb peaks are close to the theoretical value, indicating that the GaSb layers are nearly
<table><caption>Table 1. Material characteristics of GaSb layers grown on different templates</caption><thead><tr><th>Sample (500 nm of GaSb)</th><th>Miscut angle (°)</th><th>APD coverage (%)</th><th>RMS (nm)</th><th>FWHM (arcsec)</th><th>TDD (cm<sup>-2</sup>, estimated from XRD)</th><th>TDD (cm<sup>-2</sup>, estimated from ECCI)</th></tr></thead><tbody><tr><td>a) GaSb/Si</td><td>0.5</td><td>0.02</td><td>1.1</td><td>529</td><td>8.1 × 10<sup>8</sup></td><td>1.6 × 10<sup>9</sup></td></tr><tr><td>b) GaSb/Ge</td><td>0.5</td><td>0.05</td><td>1.5</td><td>567</td><td>9.3 × 10<sup>8</sup></td><td>1.9 × 10<sup>9</sup></td></tr><tr><td>c) GaSb/(CMP)Ge/Si</td><td>0.27</td><td>0.63</td><td>1.4</td><td>554</td><td>8.9 × 10<sup>8</sup></td><td>2.0 × 10<sup>9</sup></td></tr><tr><td>d) GaSb/Ge/SiGe/Si</td><td>1</td><td>0.36</td><td>2.0</td><td>1096</td><td>3.5 × 10<sup>9</sup></td><td>1.9 × 10<sup>9</sup></td></tr></tbody></table>

<figure><img src="image_9.png" /><figcaption>Fig. 8. XRD ω-2θ scans of GaSb grown on different templates. Cf. Table 1 for epitaxial stacks.</figcaption></figure>

fully relaxed. The slight shift of the GaSb peak in both cases can be attributed to the residual strain induced by a mismatch of the thermal expansion coefficients during cooling of the sample. Sample c)'s XRD record shows the three peaks of Si, Ge and GaSb (Fig. 8(c)), while Fig. 8(d) clearly confirms the presence of the graded SiGe layer with a composition evolving from pure Si to pure Ge. Figure 9 shows the ω-scan of the (004) peaks of the four GaSb layers plotted in relative arcseconds. The full widths at half maximum (FWHM) of these peaks are reported in Table 1.

Figure 9 shows that peaks a), b) and c) have Gaussian shapes with similar FWHMs, while peak d) is about twice as broad. We estimated the threading dislocation density (TDD) using the Ayer's method [25], which is based on the FWHM of these HRXRD peaks. The results are shown in Table 1. The TDD values of samples a), b) and c) are very close, in the high $10^8$ cm$^{-2}$ range. The TDD value of sample d) is approximately four times higher, in the mid $10^9$ cm$^{-2}$ range. These values were compared to the ones measured on 6.4 µm x 4.8 µm ECCI images (Fig. 10).

ECCI yields TDDs of $1.5 \times 10^7$ cm$^{-2}$ for the two Ge-based templates, and TDDs of approximately $2 \times 10^9$ cm$^{-2}$ for the four GaSb layers. These values show that the TDDs after growing 500 nm of GaSb do not depend on the substrate's (with nominal TDDs around $10^3$ cm$^{-2}$) or template's TDDs. This can be explained by the fact that the substrate/template's TDDs are
<figure><img src="image_10.png" /><figcaption>**Fig. 9.** XRD ω scans of GaSb peaks of GaSb templates grown on different templates. Cf. Table 1 for epitaxial stacks.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**Fig. 10.** ECCI image of sample d) (GaSb grown on Ge/SiGe/Si template). Each white dot corresponds to one emerging threading dislocation.</figcaption></figure>

negligible compared to the dislocation density needed at the GaSb/Ge interface to accomodate the ~8% lattice-mismatch between GaSb and Ge. The TDDs values obtained by ECCI for the GaSb samples are comparable to those obtained from the HRXRD curves, except for sample d), which was grown on the SiGe-containing template. This means that for samples a) to c), the FWHM of the GaSb peak is mainly determined by the presence of dislocations, while
additional contributions to the peak broadening probably exist for sample d). Transmission electron microscopy – which is beyond the scope of this study – would be needed to clarify the nature of the defects responsible for the HRXRD broadening. Note that such high TDDs will not be a problem for the future integration of MIR lasers such as interband cascade lasers and quantum cascade lasers which are little sensitive to dislocations [13].

## 4. Conclusion

In summary, we studied the MBE growth of GaSb layers on various Ge-based templates. We first grew layers on bulk Ge substrates with a 0.5-1° miscut toward a [110] direction. We demonstrated that after optimizing the substrate preparation, APDs can be buried within a 500 nm GaSb layer. We also showed that the burying mechanism is similar to that previously observed during the growth of GaSb on Si substrates.

Then we extended the study to Ge-based templates grown on Si substrates, which can be used for MIR photonic integrated circuits. We demonstrated that the crosshatch pattern resulting from the relaxation of the (Si)Ge layers on Si can limit the burial of the APDs because it can induce local, random inclinations that are greater than the substrate miscut. Step flow growth occurs in all directions in this case, thereby suppressing the difference in growth rates between the main and anti-phase domains. To ensure that the miscut exceeds most local inclinations related to the crosshatch pattern, we demonstrate that one can slightly increase the substrate miscut value up to 1° or perform CMP on the template prior to III-V growth. In both cases, 500-nm-thick GaSb layers free of emerging APDs can be obtained.

These results demonstrate the feasibility of hybridizing GaSb, free of emerging APBs, on Ge-based templates grown on Si. This is a further step toward the monolithic integration of III-V semiconductor heterostructures on mid-infrared photonic platforms.

**Funding.** Agence Nationale de la Recherche (ANR-11-EQPX-0016, ANR-21-ESRE-0026); HORIZON EUROPE Excellent Science (101128598).

**Acknowledgment.** Part of this work was supported by the France 2030 program (Equipex EXTRA, ANR-11-EQPX-0016; Equipex + HYBAT, ANR-21-ESRE-0026). This project has received funding from the European Union's Horizon Europe research and innovation program (101128598 - UNISON). Views and opinions expressed are, however, those of the authors only and do not necessarily reflect those of the European Union or the European Research Council. Neither the European Union nor the granting authority can be held responsible for them.

**Disclosures.** The authors have no conflict to disclose.

**Data availability.** The data that support the findings of this study are available from the corresponding author upon reasonable request.

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