InxGa1−x N is a strategically important material for electronic devices given its tunable bandgap, modulated by the In/Ga ratio. However, current applications are hindered by defects caused by strain relaxation and phase separation in the material. Here, we demonstrate growth of homogeneous InxGa1−x N films with 0.3 < x < 0.8 up to ∼30 nm using atomic layer deposition (ALD) with a supercycle approach, switching between InN and GaN deposition. The composition is uniform along and across the films, without signs of In segregation. The InxGa1−x N films show higher In-content than the value predicted by the supercycle model. A more pronounced reduction of GPCInN than GPCGaN during the growth processes of InN and GaN bilayers is concluded based on our analysis. The intermixing between InN and GaN bilayers is suggested to explain the enhanced overall In-content. Our results show the advantage of ALD to prepare high-quality InxGa1−x N films, particularly with high In-content, which is difficult to achieve with other growth methods.

Nanometer-thin InxGa1−x N films are known as very promising optoelectronic materials for efficient blue and green light-emitting diodes (LED). A monolithic full color LED has been demonstrated by stacking InxGa1−x N layers with various x within a complete LED structure.1 The emission energy of InxGa1−x N is determined by several factors with the value of x considered as the most important. Due to the miscibility gap between InN and GaN and very high critical temperature for phase separation (1250 °C),2 a homogeneous InxGa1−x N film is regarded to be metastable. The most common techniques to prepare InxGa1−x N and LED structures are metal-organic chemical vapor deposition (MOCVD) and molecular beam epitaxy (MBE) by which the In-content (x) is mainly controlled by the growth temperature and III/V ratio. Phase separation to InN and GaN is typically seen for x > 0.3,3 even in InxGa1−x N quantum well structures that are just a few nanometers thick.4 

Recently, atomic-layer-deposition (ALD) has emerged as a promising method for the growth of III-nitrides, such as binary AlN,5,6 GaN,7,8 InN,9,10 and their alloys.11–13 The ALD approach is characterized by the alternating pulses of metal and nitrogen precursors, purged by inert gas between pulses. This makes the film growth fully dependent on surface chemical reactions. In our previous work, we showed that the gas exchange dynamics combined with plasma activation of the nitrogen precursor (NH3) at temperatures below 350 °C was essential to achieve high structural quality epitaxial InN layers for just 5 nm thick.14,15 Further extension of the ALD technology to ternary III-nitrides can open a new path for integration of ALD with existing LED and high-electron mobility transistors.

Unlike ternary oxides,16 ternary nitrides are much less explored by ALD. Recently, we explored the approach of co-evaporating precursors for InxGa1−x N, by mixing In- and Ga-triazenide precursors in the evaporator and introducing them simultaneously into the reactor.17 More common for ternary materials is to use a supercycle ALD approach based on the combination of growing respective binary compounds as a multilayer. This has been demonstrated to some extent for nitrides.11–13 The growth of InxGa1−x N is accomplished by a repeating cycle that comprises of a number (m) of InN ALD cycles (UInN) followed by a number (n) of GaN cycles (UGaN). The complete growth of In1−xGaxN can be formulated as (mUInN+nUGaN)×k, where k is the total number of super-cycles. According to the rule of mixtures,18 x in InxGa1−x N is determined not only by the respective number of unit cycles but also the growth per ALD cycle (GPC) as

(1)

m and n are the number of unit cycles of InN and GaN, respectively. GPCInN and GPCGaN are the GPC of InN and GaN in the growth of In1−xGax N, bearing in mind that GPCInN(GPCGaN) may not be the same as GPCInN (GPCGaN) determined from the binary InN (GaN) process. ρInN and ρGaN are the atomic densities of InN and GaN. Although varying the ratio of m/n has been shown effective in tuning In-content,11,12 the impact of GPC on the In content has not been discussed to our knowledge.

In this work, InxGa1−x N films with various x were prepared based on the supercycle ALD approach, in which m and n were chosen to obtain respective binary monolayers. We find that InxGa1−x N with 0.2 < x < 0.8 can be deposited without phase separation. We also find a low GPCInN, which should decrease x (In content) but it is, on the contrary, found notably higher than the predicted values. This is ascribed to the effective intermixing of In into GaN. We believe that our results provide insight into the control of both the thickness and composition of InxGa1−x N by ALD and shows the importance of the ALD technique for potential InxGa1−x N devices.

The growth of InxGa1−x N was done using a Picosun R-200 plasma ALD system. Trimethyl indium (TMI), TEG, and (NH3 + Ar + N2) plasma were used to provide In, Ga, and N, respectively. The InxGa1−x N was grown at 320 °C and 6 mbar. Si (100) substrates were chemically cleaned using standard RCA cleaning procedure before loading into the reactor. After temperature stabilization at 320 °C for 60 min and prior to the growth of InxGa1−x N, the Si substrate was subjected to (Ar + N2) plasma treatment for 2 min with the intention to remove residual surface oxide. After the plasma treatment, the system was purged with 100 SCCM N2 for 10 s followed by the first TMI pulse of the InN ALD cycle. The details of the InN8 and GaN19 ALD processes can be found in our previous reports. The GPC of InN8 and GaN19 was determined to be ∼0.45 Å/cycle when grown as binary materials on silicon (100) substrates. Considering the preferred growth direction along <0001> for III-N wurtzite crystals,187UInN and 5UGaN are set for a monolayer of InN and GaN.

In some experiment that was done to show the heteroepitaxy, on-axis Si-face-CMP-polished 4H-SiC substrates was used. SiC substrate was placed next to Si (100) substrate in the same growth run. The SiC substrate was subjected for the same RCA cleaning procedure as the Si substrate prior to loading into the reactor.

The surface morphologies of the InN samples were studied by a high-resolution LEO 1550 Gemini field emission scanning electron microscope. The crystalline quality, thickness, and the macroscale roughness were characterized by using x-ray diffraction (XRD) (PANalytical X'Pert Pro with a Cu-anode x-ray tube). Grazing-incidence XRD (GIXRD) and symmetric 2θ-ω and ω-scans were used to study the crystalline quality. The film thickness was determined by analyzing the x-ray reflectivity (XRR) results. The XRR data were fitted by software PANalytical X'Pert reflectivity.

Elemental composition of the films was obtained using x-ray photoelectron spectroscopy (XPS) (Kratos AXIS ULTRA DLD with Al-anode excitation) and Rutherford backscattering spectrometry (RBS). Gaussian–Laurentius functions and Shirley background were used to fit the experimental XPS data. RBS spectra were fitted by SIMNRA 7.02 code with an 1% statistic uncertainty to determine elemental compositions.

Scanning transmission electron microscopy (STEM) and selective area electron diffraction characterization were performed using the Linköping double Cs corrected FEI Titan3 60–300 operated at 300 kV.

Spectroscopic ellipsometry was used to determine the absorption coefficient and, thus, to obtain their optical bandgap by using Cody plots (see supplementary material28 for more description about the model).

Figure 1 shows the GIXRD results of the InxGa1−x N films grown by varying the number of GaN unit cycles (n) while keeping all other growth parameters unaltered. We observe only one peak between 20° and 40°, in which III-nitrides show their crystal planes with low Miller index such as (101¯0), (0002), and (101¯1). In addition to the dominant peak between 20° and 40°, another relative weak peak ∼60° is seen (see supplementary material28 for the result in Fig. S1). Both GIXRD peaks shift toward lower 2θ values with decreasing n, revealing that the spacing between corresponding planes is enlarged and consequently the In-content (x) is enhanced accordingly. Unlike MOCVD and MBE, ALD is probably the most straight forward technique to tune x in a broad range (between x = 0.2 and 0.8) as evidenced by GIXRD results of InxGa1−x N films made with decreasing number of GaN cycles (n) in the growth process.

FIG. 1.

GIXRD results of InxGa1−x N films on Si substrates grown with various number of unit cycles for GaN. The peak intensity of every measured curve was normalized to 1 for visual clarity of peak position.

FIG. 1.

GIXRD results of InxGa1−x N films on Si substrates grown with various number of unit cycles for GaN. The peak intensity of every measured curve was normalized to 1 for visual clarity of peak position.

Close modal

The observation of a single, dominant peak in Fig. 1 indicates that the studied InxGa1−x N films is rather homogeneous. Such homogeneity can diminish as evidenced by showing multiple and broader XRD peaks when small m and n were applied (see supplementary material28 for the experimental result in Fig. S2). Considering that the <0002> is the natural preferred growth direction of wurtzite III-nitrides on substrates without epitaxial relationship18 and corroborated by our own observations,8,19 we believe the XRD peaks arise from (0002) InxGa1−x N. Assuming strainfree InxGa1−x N films, the In-content can be tuned from x = 0.87 for n = 2 to x = 0.33 for n = 25 based on Vegard's law. However, accurate determination of the In-content by XRD peak position is very challenging as it is influenced by both In-content and strain condition. The strain originates not only from the lattice mismatch between the substrate and the film, but also the In/Ga ratio and its distribution within the InxGa1−x N films. Haider et al. reported substantial deviations of In-content obtained from different characterization techniques such as GIXRD, energy dispersive x-ray spectroscopy (EDX), and XPS.11 To further investigate this issue, RBS and XPS are employed in our study. RBS is a powerful, nondestructive, and standardfree ion-beam-based technique for a fast quantitative analysis with high accuracy that is not sensitive to the strain in the target material. In addition, the high uncertainties in quantification by XPS, mainly coming from sputter-etching artifacts, are not applicable in RBS.20 From our XPS analysis, the carbon level is not detectable, indicating that it is below 1 at. %, and the oxygen level is in the range of 2–7 at. %. The detected oxygen level is found to increase with the In content. We would like to emphasize that the studied InGaN films were very thin (<30 nm) and without any protective/anti-oxidation measures. Postoxidation and surface absorbents are inevitable and will disturb the accuracy of the results significantly. In addition, the N content of our InGaN films cannot be extracted accurately. First, the excitation source of the Al Kα line for our XPS facility renders spectral overlap between the auger peaks (LMM) of Ga and N1s. Second, the detection of light elements in RBS, such as B, C, N, and O, is hindered by the substrate with heavy element.20 (see supplementary material28 for the discussion of metal/nitrogen ratio extracted from RBS in Sec. III). As the main focus of this work is about In tunability and crystalline quality, we will not further discuss about impurities and also metal/nitrogen ratios of our InGaN films. Table I summarizes the In-content and thickness of InxGa1−x N thin films obtained by different techniques, together with their optical bandgaps. The In-content determined by all analytic techniques is similar within a range of a few atomic percent, indicating that the In1−xGax N films are rather homogeneous. Constant overestimation of In-content by GIXRD is likely due to residual strain within the InxGa1−x N films and in line with previous observations.11 

TABLE I.

In-content and thickness of InxGa1−x N films on Si substrates by various combinations of k, m, and n.

kmnIn% by rule of mixtureaIn% by XPSIn% by RBSIn% by XRDEstimated thickness (nm)bMeasured thickness (nm)cFitted optical bandgapd (eV)
23 25 17 29 31 35 33.12 26.9 ± 0.8 2.43 
28 20 20 32 34 40 34.02 27.2 2.393 
34 15 26 36 41 48 33.66 25.1 ± 1.8 2.39 
44 10 34 59 47 56 33.66 24.1 ± 2 2.36 
60 51 69 64 72 32.40 21.9 ± 1.4 2.12 
80 72 77 98 87 32.40 20.6 ± 1.7 1.99 
kmnIn% by rule of mixtureaIn% by XPSIn% by RBSIn% by XRDEstimated thickness (nm)bMeasured thickness (nm)cFitted optical bandgapd (eV)
23 25 17 29 31 35 33.12 26.9 ± 0.8 2.43 
28 20 20 32 34 40 34.02 27.2 2.393 
34 15 26 36 41 48 33.66 25.1 ± 1.8 2.39 
44 10 34 59 47 56 33.66 24.1 ± 2 2.36 
60 51 69 64 72 32.40 21.9 ± 1.4 2.12 
80 72 77 98 87 32.40 20.6 ± 1.7 1.99 
a

The GPC of 0.45 Å/cycle are assumed for both InN and GaN as we observed in our earlier studies.

b

Thickness=(mGPCInN+nGPCGaN)×k. GPC* of 0.45 Å/cycle for both InN and GaN is assumed.

c

The measured thickness is obtained by fitting the results of x-ray reflectivity measurement.

d

The optical bandgap is obtained by using spectroscopic ellipsometry.

According to the dimensionless mixing number (Mx) proposed to differentiate the interfacial sharpness related to the layer thickness in the ALD process,21 a distinguishable composition variation can only happen if the diffusion length of one metal in another metal nitride is smaller than the thickness of each bilayer.16 Although the diffusion length of In (Ga) in GaN (InN) under the growth environment is not known, the finite thickness of GaN in our case (ideally no thicker than 1.2 nm) should prevent the formation of separated phases.

We would like to point out that there is no sign of epitaxy between In1−xGax N and Si. The crystal grains are randomly oriented with respect to the Si substrate. Despite the hexagonal crystal grains as majority, cubic inclusion is also found. (see supplementary material28 for the TEM result of InGaN on Si in Fig. S3). To highlight the heteroepitaxy and the good homogeneity of the films, the SiC substrate is used for STEM study. Figure 2(a) shows a constant In/Ga ratio across entire InxGa1−x N (7UInN + 5UGaN) film from the energy dispersive x-ray elemental line profiles (STEM-EDX). Pronounced In segregation is not seen according to our STEM-EDX element mapping images. [Figs. 2(c)2(e)]. In addition, the heteroepitaxy of the In1−xGax N film is evidenced in the corresponding Fast Fourier Transform pattern [Fig. 2(b)]. The AB-type stacking of wurtzite InxGa1−x N and the ABAC-type stacking of the 4H–SiC substrate are well-aligned along [011¯0] and [0001] directions. The a and c lattice constants derived from the FFT are ∼3.454 and ∼5.521 Å, respectively. This corresponds to an In content (x) of 0.72, in line with the GIXRD, but can also be affected by the strain in the material. From the lattice-resolved image, the density of structural disorder, such as stacking faults, at the interface is higher than for binary InN.14 The same constant In/Ga ratio and heteroepitaxy can also be seen from a Ga-rich InxGa1−x N (7UInN + 25UGaN) film (see supplementary material28 for the TEM result in Fig. S4). However, in the Ga-rich films, there seems to be an In-rich regime close to the surface and the lattice stacking is then less coherent especially after a few layers on top of the Ga-rich InxGa1−x N film. The reason why Ga-rich InxGa1−x N film is worse than In-rich ones requires further investigation.

FIG. 2.

(a) STEM-EDX line scan and (b) HR-STEM image of an InxGa1−x N (7UInN + 5UGaN) film grown on 4-H SiC (0001) substrate and its corresponding FFT. EDX element maps of In (c), Ga (d), and Si (e).

FIG. 2.

(a) STEM-EDX line scan and (b) HR-STEM image of an InxGa1−x N (7UInN + 5UGaN) film grown on 4-H SiC (0001) substrate and its corresponding FFT. EDX element maps of In (c), Ga (d), and Si (e).

Close modal

Despite homogeneous In-content (x), the predicted In-content is significantly lower than experimentally determined values. Moreover, all InxGa1−x N films are thinner (determined by XRR) compared to their predicted values as indicated in Table I. Both observations indicate that the actual GPCInN and GPCGaN for In1−xGax N are smaller than those found for binary InN and GaN. According to the rule of mixture,18 the thickness (T) of an InxGa1−x N film grown via the supercycle ALD approach can be expressed by using

(2)

Considering GPCInN and GPCGaN as two variables, each InxGa1−x N film grown with different combinations of m, n, and k can be plotted as a line. Those lines shown in Fig. 3 be seen as mathematical binary linear simultaneous equations. Assuming GPCInN and GPCGaN are constant and independent on the number of m, n, and k, a single solution for GPCInN and GPCGaN is expected. However, such single solution is not available according to our experimental results. Instead, the closest approximation based on the intersections between all lines, the outcome of GPCInN0.26Å/cycle and GPCGaN0.40Å/cycle is obtained. According to our analysis, both GPCInN and GPCGaN are smaller than those estimated from binary cases, with a more pronounced reduction for GPCInN. In general, the reduction of overall GPC is often ascribed to nucleation delay when switching from one binary ALD process to another, which has been reported for ternary oxides.16,22,23

FIG. 3.

Plots of GPCs of InN and GaN for the supercycle ALD grown InxGa1−x N films using T=(mGPCInN+nGPCGaN)×k. All relevant values are from Table I.

FIG. 3.

Plots of GPCs of InN and GaN for the supercycle ALD grown InxGa1−x N films using T=(mGPCInN+nGPCGaN)×k. All relevant values are from Table I.

Close modal
FIG. 4.

Cody plots of InxGa1−x N films by various combinations of k, m, and n as indicated in Table I. Solid lines are experimental results. The fitted optical bandgaps are obtained by extrapolating the linear region of the Cody-plot to the level of zero absorption as indicated by dashed line of each experimental results.

FIG. 4.

Cody plots of InxGa1−x N films by various combinations of k, m, and n as indicated in Table I. Solid lines are experimental results. The fitted optical bandgaps are obtained by extrapolating the linear region of the Cody-plot to the level of zero absorption as indicated by dashed line of each experimental results.

Close modal

Nucleation delay is often associated with the density of reactive surface sites, exchange reaction, and persistent ligands.16,24 Those impacts are expected to fade upon increasing ALD cycles,24,25 and thus, the GPC* should increase accordingly until a steady state. This argument is further supported by seeing increasing thickness of InxGa1−x N films with increasing m and n (Fig. S4), where all lines possess identical expression of 420GPCInN+300GPCGaN If GPCInN and GPCGaN were independent of m and n, superposition lines would have been observed. According to our analysis (see supplementary material28 for the analysis in Sec. 6), the increasing thickness of In1−xGax N films with increasing m is mainly due to the enhancement of GPCInN (increasing from 0.26 to 0.36 Å/cycle) rather than GPCGaN (0.4 Å/cycle). From this perspective, the effect of nucleation delay is more pronounced for InN than GaN. A more severe nucleation delay of InN than of GaN was also found by using tris(1,3-diisopropyltriazenide)indium(III) as precursors for InxGa1−x N in our reactor.17 

Since GPCInN and GPCGaN are significantly smaller than we anticipated at the first place, the prediction of In-content shown in Table I should be revised. Although the actual GPC* is not known for accurate prediction, reduced GPCInN than GPCGaN will lead to even lower In-content prediction [Eq. (1)]. Considering that all our experimental results suggest a higher In-content than the estimated, a plausible explanation is that the “GaN” grown in the process is in fact InxGa1−x N. Unintentional In incorporation into the adjacent GaN layer is often seen in In1−xGaxN/GaN multiple quantum well structures grown by MOCVD.26 We speculate that the interdiffusion of In into GaN consumes the In-adatoms on the surface of GaN, resulting in reduced GPCInN and excessive In content seen in our experiment.

Spectroscopic ellipsometry is employed to investigate the optical properties of InxGa1−x N films with different x. The bandgap of the InxGa1−x N film can be determined by extrapolating the linear region of respective Cody-plot as shown in Fig. 4 (see supplementary material28 for the plot of refractive index against wavelength and more information of fitting model in Sec. 7). The optical bandgaps of InxGa1−x N films obtained from our spectroscopic ellipsometry increase progressively from 1.99 to 2.43 eV with decreasing x as indicated in Table I. We would like to highlight that the bandgap estimated by our spectroscopic ellipsometry study is significantly higher than expected by considering the bandgap of InN as 0.7 eV. We believe it is due to their Fermi levels that are well above the conduction minimum, which is a known phenomenon for InN and In-rich InGaN.27 

In summary, InxGa1−x N films grown by ALD with a supercycle approach show a broad tunability in the In/Ga ratio by varying the number of GaN ALD cycles in the supercycle. The experimentally determined In-content is higher than the values predicted by the supercycle model. We ascribe this enhanced In incorporation to pronounced intermixing between InN and GaN bilayers. In addition, the actual growth of InxGa1−x N is not an ideal linear combination of GPCInN and GPCGaN found in the respective binary process. Notable reduction of GPCInN and GPCGaN is ascribed to nucleation delay when switching from one binary ALD process to another. No phase segregation is observed from an In0.6Ga0.4N film by TEM or XRD. Our results highlight the advantage of the ALD technique to prepare homogeneous, high In-content InxGa1−x N films, which could be used for extending the emission wavelength to red and near infrared in light-emitting-diodes.

This work was financially supported by the Swedish Foundation for Strategic Research through the project “Time-resolved low temperature CVD for III-nitrides” (No. SSF-RMA 15-0018) and by the Knut and Alice Wallenberg Foundation through the project “Bridging the THz gap” (No. KAW 2013.0049). The authors acknowledge the Knut and Alice Wallenberg Foundation for support of the electron microscopy laboratory in Linköping and for project funding (No. KAW 2020.0033). The Swedish Research Council and the Swedish Foundation for Strategic Research are acknowledged for access to ARTEMI, the Swedish National Infrastructure in Advanced Electron Microscopy (Grants Nos. 2021-00171 and RIF21-0026). Babak Bakhit gratefully acknowledges financial support from Swedish Research Council (VR), Grant No. 2019-00191 (for accelerator-based ion-technological centre in tandem accelerator laboratory in Uppsala University) and Grant No. 2021-00357.

The authors have no conflicts to disclose.

Chih-Wei Hsu: Conceptualization (lead); Data curation (equal); Formal analysis (equal); Investigation (lead); Methodology (lead); Validation (lead); Visualization (lead); Writing – original draft (lead). Ivan Martinovic: Data curation (equal); Formal analysis (equal); Software (equal); Writing – review and editing (supporting). Roger Magnusson: Data curation (equal); Formal analysis (equal); Validation (supporting); Writing – original draft (supporting); Writing – review and editing (supporting). Babak Bakhit: Data curation (equal); Formal analysis (equal); Writing – review and editing (supporting). Justinas Palisaitis: Data curation (equal); Formal analysis (equal); Resources (supporting); Writing – review and editing (supporting). Per. O. Å. Persson: Conceptualization (supporting); Investigation (supporting); Resources (supporting); Supervision (supporting); Writing – review and editing (supporting). Polla Rouf: Data curation (equal); Formal analysis (equal); Writing – review and editing (supporting). Henrik Pedersen: Funding acquisition (lead); Project administration (lead); Resources (equal); Writing – original draft (supporting); Writing – review and editing (lead).

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

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Supplementary Material