XRR has been performed to study the growth rate of all InN films. However, only samples grown by ALD at 320 °C revealed clear X-ray intensity fringes induced by the smooth and continuous film. The thickness fringes are barely seen for samples by ALD at 400 °C and thereafter disappear completely for all other samples (supplementary information, Fig. ). We believe this is a consequence of high surface roughness observed in most of the studied samples as shown in Fig. . Meanwhile, the absence of noticeable thickness fringes also indicates that the existence of buffer layer between InN and the substrate is negligible. Crystalline, hexagonal InN can be found by XRD in the samples deposited by ALD as seen from the InN (0002) diffraction peak at 31.3° in Fig. . The satellite peaks in the vicinity of InN (0002), observed in the samples deposited at 320 °C, are thickness-correlated fringes. The presence of thickness fringe can be ascribed to heterojunctions with smooth surface and sharp interface. Such thickness fringes disappears in InN grown at higher temperatures due to higher surface roughness. Despite that the roughness increases with increasing temperature, the stacking coherence of InN crystal was not affected as shown by the -scans against the InN(0002) between 320 and 400 °C. The full-width-half-maximum of the InN (0002) peak in ALD deposited films being nearly identical to the SiC substrate, in the range of 20 to 50 arcs (Fig. ). For samples grown by ALD at 450 °C, the crystalline quality of InN is notably degraded by showing a broad FWHM of 600 arcs. The additional peaks at 30.58° and 32.93° in the 2θ-ω scan can be indexed to In2O3 and metallic In, respectively. In droplets, as observed in Fig. , have previously been reported during growth of InN under elevated temperatures and In-rich conditions. Among the samples deposited by CVD, only the sample deposited at 450 °C show XRD peaks that can be indexed to InN (Fig ). The FWHM of the InN (0002) in the -scan for the sample deposited by CVD is roughly 350 arcsec (Fig. ), which is significantly wider than for films prepared by ALD at lower temperatures in this work, but is comparable to previously reported CVD films. In Fig. , a down shift of InN (0002) peak for approximately 0.15° with respected to the strain-free InN (0002) at 31.3° is observed for all samples except the one by ALD at 320 °C (Fig. ). Such XRD peak shift is mostly likely due to strain caused by the lattice mismatch to the substrate and/or defects within the films. According to our earlier study, InN layer becomes relaxed efficiently via the formation of misfit dislocations. A strain-free InN can be indexed for samples thicker than 5nm by XRD. Instead of strain from substrate, we suggest that oxygen could play a significant role for this. Owing to the very oxyphilic nature of In, seen from the standard Gibbs free energy of formation of -831 kJ/mol for In2O3 and -105 kJ/mol for InN, any residual O in the reactor and post-oxidation from air exposure can results in InOxNy or even In2O3. Peak down-shift with respect to InN (0002) in XRD symmetric scan has been observed for InOxNy before it turns completely into In2O3, which typically forms the cubic structure that is very different from the wurtzite structure of InN. According to EDX element mapping (Fig. ), In, N, O and Si are uniformly distributed for samples deposited at 320 °C by ALD (Fig. ) while an aggregation of oxygen together with In and N can be observed for samples deposited at 400 °C by ALD (Fig. ). It has been found that the non-c plane surfaces are more favorable for oxygen incorporation in GaN. Assuming InN behaves similarly as GaN, the scattered small crystal grains seen in the sample deposited at 400 °C afford a higher surface area leading to enhanced oxygen incorporation and oxidation. We therefore suggest that the shift of the XRD peak in The results in Figs. and shows that the ALD process holds advantages over the conventional continuous CVD for deposition of InN. There are two major differences between the ALD and CVD approaches: the ALD approach uses plasma activation for the ammonia and the time-resolved supply of the precursors. The difference in the film deposition seen in Fig. is likely explained by one or both differences. When a time-resolved precursor supply was used in the CVD reactor, no film was obtained at any temperature. When the plasma was disabled in our ALD reactor and a thermal ALD process for InN was attempted, no sign of InN deposition could be seen even at 450 °C. On the other hand, with plasma activation of the NH3 pulse, InN can be deposited even at 180 °C in the ALD reactor, albeit with a reduced deposition rate and crystalline quality. This demonstrates the importance of the plasma activation of the NH3 pulse. The utilization of plasma, affording more reactive N species, has been shown to increase the growth rate as well the crystalline quality of InN in both ALD and CVD systems. The plasma activation therefore seems paramount for the InN deposition at low temperature.