FORM 2 THE PATENTS ACT, 1970 (39 of 1970) & THE PATENTS RULES, 2003 COMPLETE SPECIFICATION (See section 10, rule 13) “ZINCBLENDE STRUCTURE GROUP III-NITRIDE” ZINCBLENDE STRUCTURE GROUP MI-NITRIDE BACKGROUND TO THE INVENTION Field of the invention The present invention relates to the formation of layers of zincblende structure group Ill-nitride, e.g. GaN, AIGaN, InGaN, InAIN and more generally ln,Al,Ga1-x-yN. Disclosed herein are characterizations of such layers and methods for their formation. These materials have particular, but not necessarily exclusive application, in the field of semiconductor structures and devices, such as light emission applications, e.g. LEDs, lasers, and other devices such as transistors, diodes, sensors. Related art The group-Ill nitride semiconductors offer a wide range of optoelectronic applications such as multiple quantum well (MQW) LEDs and laser diodes emitting in the blue and green spectral region. Such devices are commonly grown along the c-direction of the hexagonal wurtzite phase, in which strong internal electrical polarisation fields across the quantum wells result in a reduction of the radiative recombination rates, and current density-dependent emission wavelength [Miller et al (1985); Fiorentini et al (1999); Hammersley et al (2015)]. Although these effects are somewhat mitigated by the use of thin QW layers (typically 2-4 nm thick) long radiative recombination lifetimes and relatively low internal quantum efficiencies are observed for green light emitting structures [Nippert et al (2016); Hammersley et al (2016)]. QW structures grown along non-polar axes of the wurtzite GaN phase, e.g., the a-plane and m-plane, have been designed to avoid the polarisation fields and related limitations. Although the non-polar wurtzite devices show very short radiative lifetimes [Marcinkevicius et al (2013); Dawson et al (2016)] and wavelength characteristics independent of current density [Detchprohm et al (2010)], their quantum efficiencies have never surpassed those of polar c-plane structures [Dawson et al (2016)]. A possible explanation for the underachievement of green-emitting non-polar wurtzite devices could be the fact that indium-richer QWs are required to achieve green emission, in the absence of the polarisation field-related quantum confined Stark effect (QCSE) [Fiorentini et al (1999)]. Apart from the increased interfacial strain with the GaN buffer and barrier layers, the lower indium incorporation efficiency of non-polar growth planes compared to the polar growth plane requires low process temperatures for indium-rich layer growth [Zhao et al (2012)]. This potentially results in high densities of impurities and point defects, which act as non-radiative recombination centres and further decrease the efficiency of non-polar wurtzite MQWs [Chichibu et al (2005)]. Hence, GaN-related structures in the cubic zincblende phase have re-emerged as a promising approach to achieving improved efficiencies for green-wavelength LEDs after strong initial interest in the mid 1990s. Cubic zincblende InGaN/GaN MQW structures grown on the (001) plane may be polarisation field free, as in the zincblende phase these fields are only induced by shear stresses [Hanada (2009)], which are not present in (001) oriented films. Therefore, compared with c-plane hexagonal structures, the electron-hole wavefunction overlap is increased, which should lead to increases in the radiative recombination rate. Furthermore, InGaN has a narrower bandgap in the cubic phase than in the hexagonal phase for a given indium content [Schormann et al (2006); Compean Garcfa et al (2015)], allowing green-wavelength emission to be achieved at lower indium contents than in non-polar wurtzite structures. However, as zincblende GaN and InGaN are metastable under most growth conditions, it is possible that zincblende films will contain inclusions of the more stable wurtzite polytype, and therefore have a substantial content of wurtzite-like stacking faults and lamellae [Trampert et al (1997); Shen et al (2003); Wu et al (1997); Yang et al (1996)]. US 2016/0247967 discloses suitable substrates for growing cubic GaN layers. These substrates comprise a single crystal silicon wafer with spaced-apart monocrystalline silicon carbide layers formed on the silicon wafer, with amorphous or polycrystalline silicon carbide layers formed between the monocrystalline silicon carbide layers. The monocrystalline SiC is in the form of the 3C-SiC polymorph. GaN is formed over the SiC layer, with epitaxy allowing monocrystalline GaN to form over the monocrystalline SiC and polycrystalline GaN forming over the polycrystalline/amorphous SiC. The effect of this is to provide stress relief for the monocrystalline GaN regions. Other approaches to control stress in the GaN layer are possible, for example the introduction of lattice mismatched layers, such as compositionally graded AIGaN, which induce compressive stress during growth to counteract the tensile stress introduced by the thermal expansion mismatch. This approach has been used for the growth of conventional GaN on Si structures [Zhu et al (2013)]. US 2016/0247967 does not provide substantial detail of how the GaN should be grown, or any explanation or quantification of the content of wurtzite-like regions of GaN in the monocrystalline GaN. In this regard, the disclosure of US 2016/0247967 is limited to stating that the GaN layers can be formed using a Metal Organic Vapour Phase Epitaxy (MOVPE) process using a temperature less than 1000°C, preferably between 800°C and 950°C. SUMMARY OF THE INVENTION In order to achieve single phase epitaxial films with reasonable crystal quality, it is necessary to support the growth process by a powerful structural characterisation technique. The present invention has been devised in order to address at least one of the above problems. Preferably, the present invention reduces, ameliorates, avoids or overcomes at least one of the above problems. Accordingly, in a first aspect, the present invention provides a method of manufacturing a semiconductor structure comprising a substantially (001) oriented zincblende structure group Ill-nitride layer, the method including the steps: providing a silicon substrate; providing a 3C-SiC layer on the silicon substrate; growing a group Ill-nitride nucleation layer; carrying out a nucleation layer recrystallization step; and depositing and growing the zincblende structure group Ill-nitride layer by MOVPE at temperature T3 in the range 750-1000 °C, to a thickness of at least 0.3um. The present inventors have found that these steps provide a zincblende structure group Ill-nitride layer of improved crystalline quality, in particular with respect to the reduction of the formation of wurtzite structure group Ill-nitride inclusions. In this disclosure, some numerical ranges are expressed in terms of open ended ranges with upper or lower limits, or in terms of closed ended ranges with upper and lower limits. It is expressly stated here that preferred ranges are disclosed herein that are combinations of upper and/or lower limits from different ranges for the same parameter. Preferably, before growing the group Ill-nitride nucleation layer, the 3C-SiC layer is subjected to a nitridation step at a temperature T1 in the range 800-1100 °C. This step is considered to be advantageous in terms of ensuring the plentiful availability of N for subsequent group Ill-nitride formation. The use of temperature T1 outside this range is considered to reduce the PL NBE peak intensity and broadens the emission FWHM. We now consider the conditions for deposition and growth of the group Ill-nitride nucleation layer. Preferably, the group Ill-nitride nucleation layer is grown at temperature T2 in the range 500-700 °C. More preferably, the temperature T2 is in the range 550-650 °C. The growth rate may be at least 0.1 nm/s. The growth rate may be up to 1 nm/s. The thickness of the nucleation layer may be at least 3 nm, but is more preferably greater than 3 nm. More preferably, the thickness of the NL may be at least 10nm. The thickness of the nucleation layer (NL) may be up to 100nm. Preferably, the thickness of the NL may be up to 50nm. More preferably, the thickness of the NL may be up to 40nm. Preferably, the chosen temperature for T2 lies about 40-60 °C above the temperature where the growth rate deviates from a constant value to a lower value, entering a regime in which the ammonia flow determines the growth rate. Following the growth of the nucleation layer, there is the nucleation layer recrystallization step. In this step, preferably the temperature is ramped up at a rate of between 0.1-10 °C/second. More preferably the temperature is ramped up at a rate of between 0.5-5 °C/second. This is found to be a suitable approach for satisfactory nucleation layer recrystallization, permitting subsequent high quality epilayer deposition. The group Ill-nitride nucleation layer is preferably a zincblende structure group Ill-nitride nucleation layer. In the step of depositing and growing the zincblende structure group Ill-nitride layer on the recrystallized nucleation layer, the reactor pressure is preferably not more than 500 Torr. More preferably, the reactor pressure is not more than 300 Torr. More preferably, the reactor pressure is not more than 100 Torr. In the step of depositing and growing the zincblende structure group Ill-nitride layer on the recrystallized nucleation layer, the V-to-lll ratio is preferably in the range 10-300. More preferably, the V-to-lll ratio is in the range 20-150. More preferably the V-to-lll ratio is is in the range 50-100. During this step, the growth rate is preferably in the range 0.1-1 nm/second. A growth rate of about 0.5 nm/second has been found to be suitable, for example. Careful selection of the V-to-lll ratio within the preferred ranges shows improvements in the surface morphology, the zincblende phase purity and XRD rocking curve peak widths. In the step of depositing and growing the zincblende structure group Ill-nitride layer on the recrystallized nucleation layer temperature T3 is preferably in the range 800-920 °C. More preferably, temperature T3 is at least 810 °C, more preferably at least 820 °C, more preferably at least 830 °C. Temperature T3 is preferably at most 910 °C, at most 900 °C, or at most 890 °C. A particularly suitable range for T3 is found to be 845-880 °C. Careful selection of temperature T3 within the preferred ranges shows improvements in the surface morphology from Nomarski images, XRD rocking curve peak width and PL. For example, samples grown in the range 860 -880 °C shows a relatively smooth surface and corresponding NBE PL peak is the strongest, although also the broadest in the data presented here. At higher growth temperature, it is found that the surface roughens, the PL NBE peak narrows significantly but also the yellow band increases in intensity. Changes in surface roughness with growth temperature were confirmed from AFM. The phase purity, as determined by XRD shows that it is possible to very substantially reduce the amount of wurtzite inclusions when T3 is 900 °C or lower. When T3 is higher than 900 °C the XRD analysis shows an increasing contribution of reflections due to the wurtzite lattice, indicating incorporation of hexagonal inclusions in the cubic zincblende matrix. It has been found in this work that it is possible to widen the preferred conditions of temperature T3 and lll-V ratio by carrying out epilayer growth at relatively low pressure. In one exemplary set of conditions for growing zincblende GaN epilayers by MOVPE at a constant pressure of 100 Torr, T3 can be in the range 850 and 890 °C, with a V/lll ratio of 38 to 150. This results in a relatively smooth film with a wurtzite contamination of less than 1%. The preferred thickness of the NL is in the range 10-50 nm, for example about 22nm. Preferably, the group Ill-nitride layer is an lnxAlyGa1-x-yN based layer, where 0 {111} facets in zincblende GaN. This unit cell arrangement is illustrated in Fig. 76, and it is unsurprising, since the closed-packed planes of each structure are parallel to each other and differ only in the stacking sequence. Other groups observed a similar arrangement of the two phases by XRD [Qu et al (2001); Tsuchiya et al (1998)] and transmission electron microscope measurements [Trampert et al (1997)], and it was also found that the zincblende and wurtzite phases arrange as alternating zb-wz-lamellae [Wu et al (1997)]. As our example in Fig. 75 indicates, (0001) wurtzite GaN is not necessarily formed with equal probability on each of the four inequivalent {111}Zb facets. Table 1: Atomic positions in the ideal wurtzite and zincblende unit cell. wurtzite structure zincblende structure Ga3+ 0, 0, 0; 1/3,2/3, V2; 0, 0, 0; V2, 0, 1/2; 1/2, 1/2, 0; 0, V2,1/2; N3- 0, 0, 1/3,2/3,7/8; 3/4, 3/4, 3/4; V4, 3/4, 1/4; 3/8; 1/4, 1/4, 3/4; 3/4, 1/4, 1/4; wurtzite structure zincblende structure Ga3+ 0,0,0; 1/3,2/3, V2; 0, 0, 0; V2, 0, 1/2; 1/2, 1/2, 0; 0, V2,1/2; N3- 0,0, 1/3,2/3,7/8; 3/4,3/4,3/4; V4,3/4,1/4; 3/8; 1/4, 1/4, 3/4; 3/4, 1/4, 1/4; In samples grown under optimised conditions to maximise the zincblende phase, the X-ray intensity in the expected positions of the wurtzite phase reflections is negligible, i.e. only slightly above the background noise level. In these cases it is very likely that the signal originates not from diffraction by hexagonal wurtzite inclusions in the epilayers, but from diffuse scattering on planar defects, such as stacking faults. This can be shown by measuring a two-dimensional reciprocal space map as described later. Besides wurtzite inclusions, zincblende GaN thin films could also contain twinned zincblende regions. Similar to stacking faults, these are introduced by stacking errors of a single (111) plane, but in contrast to stacking faults, the zincblende matrix continues with a different stacking sequence ...AaCcBbAaCcBb... . With respect to the surrounding GaN matrix, the zincblende twins are tilted by approximately 70.4° around the [1-10] axis, so that twin and matrix have the relation (111)twin || (115)mathx [Tsuchiya et al (1998)]. Such zincblende twins, and possibly twinned wurtzite-like material with a similar relation, cause the weak 111 reflections at x of about 15° (circled in Fig. 73), and possibly at x of about 83° (out of range in Fig. 73). Their volume portion is in the low percentage range. Reciprocal space maps for texture analysis Two-dimensional reciprocal space maps (RSMs) - a combination of oo-26-scans with a stepwise change of the co-angle after each scan - of suitable zincblende and wurtzite phase reflections can be used to analyse the phase purity of GaN samples, as well as several other structural properties. Suitable reflections include 002ZB and 10-11wz, as shown in Figs. 77 and 78 for two different samples with and without hexagonal inclusions. In both reciprocal space maps the 002 reflections of zincblende GaN and 3C-SiC are clearly visible by their high intensities. The low intensity streaks running along <111 > through the 002 reflections are caused by diffuse scattering from {111} stacking faults in the structure, where diffracted X-rays suffer an additional phase shift between both sides of a stacking fault. Stacking faults may also lead to a small shift of the GaN reflections out of the ideal position. Another feature passing through the 3C-SiC reflection on a 26-arc is the detector streak (DS), caused by the instrument function of the diffractometer. The streak intersecting the 002zb GaN reflection normal to the surface is the so called (X-ray) crystal truncation rod (CTR), whose shape is influenced by surface structures in agreement with observations from atomic force microscopy. The partially visible Bragg ring in the RSMs (with 26 of about 35.6°) stems from polycrystalline SiC deposited on the etched grid of the 3C-SiC/Si templates, and is neither related to the SiC mesa regions nor to the GaN epilayer. As the GaN epilayer is much thinner than the SiC template, similar Bragg rings from GaN grown on the etched grid are much weaker and often not visible. In the presence of wurtzite-like inclusions in the zincblende GaN thin film (Fig. 67) two additional 10-11wz reflections of the wurtzite GaN phase appear, which are absent in the sample without these inclusions (Fig. 78). Since the stacking fault streak overlaps with the wurtzite phase reflections, these streaks can be easily misinterpreted as signal of small amounts of hexagonal inclusions in texture maps. In comparison to texture mapping, discussed in the section above, reciprocal space maps are much faster to perform even with high integration times by using a CCD detector. This increases the signal to background noise ratio, and hence allows the quantification of even relatively low proportions of wurtzite-like GaN inclusions. However this method presumes a fixed epitaxial relation and provides no additional information on the presence of cubic twins. Mosaicitv analysis Due to the lack of suitable homo-substrates, cubic zincblende GaN-based nitrides are typically grown heteroepitaxially on foreign cubic substrates, such as GaAs [As et al (2000); Yang et al (1996); Shen et al (2003); Qu et al (2001); Tsuchiya et al (1998)], SiC [Wu et al (1997); Chichibu et al (2003)], Si [Lei et al (1991)], and various other cubic materials (e.g. GaP [Cheng et al (1995)], MgO [Compean Garcfa et al (2015)]). The lattice mismatch between the different materials results in a high mosaicity and the formation of defects at grain boundaries in the epilayers. In general, mosaicity should be avoided as it negatively influences the physical properties of the sample, e.g. causes high electrical resistances at grain boundaries [Fujii et al (2010)]. And hence, mosaicity should preferably be quantified for crystal growth optimisation. In a simplified model originating from powder diffractometry, the thin film consists of mosaic blocks (grains), which differ slightly in their finite size and orientation relative to each other. The spread in size, tilt, and twist, together with microstrain and compositional inhomogenities (in the case of alloys) lead to broadening of the X-ray reflections in reciprocal space. Mosaic tilts lead to an angular broadening of reflections perpendicular to the surface, while twists result in an azimuthal spread around the surface normal. Hence for both mosaic tilt and twist, the absolute broadening in reciprocal space AGhki increases linearly with the magnitude of the scattering vector |Ghki|. The finite lateral size of mosaic grains causes a broadening parallel to the interface, being inversely proportional to the average real space size L and independent of the scattering vector magnitude (AGhki = 2TT/L). The effects of tilt, twist and finite grain size convolute to the spread of a reflection hkl measured by co-scans in skew-symmetry as follows: [Lee et al (2005)]. Here 3 denotes the integrated breadth, and the exponent n takes values between 1 and 2 depending on the Gaussian n and Lorentzian (1 - n) contribution to a Pseudo-Voigt fit (n = 1 + n,2) (See appendix in Srikant et al (1997)]). The measured peak broadening is then a combination of this mosaic broadening of the sample and the instrument function (without a sample). The latter one can be neglected as long as it is much narrower than the mosaic broadening. Experimentally, the peak broadening effect due to lateral size and tilt can be separated by measuring a series of co-scans of different order symmetric reflections 00/, and plotting 3n-Gn versus Gn in a modified Williamson-Hall plot (not shown). The slope of the line is related to the tilt component (3tiitn) and the ordinate offset is related to the average grain size ((2Tr/L)n). Unfortunately only the symmetric 002 and 004 zincblende GaN reflections are accessible with the commonly used Cu-Ka radiation, which significantly limits the accuracy, especially of the finite size determination. Fig. 79 shows the linear behaviour of traditional Williamson-Hall plot for which typically Lorentzian broadenings (n = 1) are taken, even though Lorentzian profiles often do not fit very well to measured profiles. In comparison with the curve for Gaussian fits (n = 2), this results in a much larger value for the finite size L, as also pointed out by Lee et al (2005). As X-ray intensity profiles can be empirically described as a convolution of Gaussian and Lorentzian functions, the real lateral finite size is within these two extremes, depending on the portion of both profiles. Usually one gets the Gaussian and Lorentzian ratio from profile fitting, but they might vary for different reflections of a series. However, as the Lorentzian portion is often very small in such fits, mosaic block sizes estimated from pure Gaussian fits give a relatively good estimate. The azimuthal spread around the surface normal due to mosaic twist can be determined from off-axis reflections with large polar angles x measured in screw-symmetric geometry. Ideally one would use one of the in-plane reflections (x of about 90°), but those often exhibit only very low intensities and are generally difficult to measure. For (001) oriented zincblende GaN films the 331 reflection (x of about 76.7°) may be better used instead. Alternatively, the integrated breadths of a series of different off-axis reflections extracted from skew-symmetric co-scans can be extrapolated with Eq. (9) to determine the twist component. Fig. 80 shows such an extrapolation, for which we converted Eq. (9) to 3hkin-|Ghki|n, and fit this function (for n = 2) to the measured peak broadenings of an optimised cubic GaN sample by using the tilt and finite size values from the Williamson-Hall plot in Fig. 79. The circles mark the measured reflections, and the contours are provided to indicate the extrapolated peak width in reciprocal space as function of polar angle x and scattering vector magnitude |Ghki|. The contours represent constant peak width. The profile at x= 0° was already shown before (Fig. 79 ) and is only influenced by the tilt and the finite size of the mosaic blocks. With increasing polar angle x the broadening gradually increases, revealing a slightly higher mosaic twist (x = 90°) of 0.864° than tilt (x = 0°) of 0.755°. This trend is not much pronounced as tilt and twist are very similar, but it becomes more apparent for larger scattering vectors |Ghki| as the contribution of the finite size to the peak broadening decreases. Defect densities In general mosaic tilt and twist are assumed to be associated with the formation of threading dislocations at grain boundaries in the thin film. Hence the XRD peak broadening is sometimes used to estimate the defect density in a thin film, by following different mosaic tilt models discussed in literature. According to these models the threading dislocation density DTD in a well oriented mosaic film is proportional to [Fewster (1989)], while in poorly oriented films with randomly oriented grains and strictly statistically spread Burgers and line vectors, the threading dislocation density is proportional to P2tiit/twist [Dunn and Koch (1957)]. Here the parameter L is the average lateral finite size of the grains, while denotes the Burgers vector of the dislocation with a value of in the case of perfect dislocations in zincblende GaN. In contrast to wurtzite GaN material, where the threading dislocation line vector propagates predominantly along the [0001] c-direction, the threading dislocations in zincblende GaN run along multiple <110> directions. Thus, the equations above do not allow separating between edge-type, mixed-type, or screw type dislocations in zincblende GaN. However it is well known that, the dominating threading dislocation type in zincblende like structures are 60°-type perfect dislocations [Blumenau et al (2000)]. In an intensive comparative study using XRD and transmission electron microscopy (TEM) to estimate the defect densities in wurtzite GaN films, Metzger et al (1998) found a good match with the random distribution model (Eq. xy), even though the assumptions of the model are not fulfilled at all in oriented epitaxial thin films. Contrary to expectations, the model for oriented mosaic films revealed threading dislocation densities which were more than a magnitude lower than the values estimated from TEM. Lee et al (2005) came to similar conclusions and noted that large differences in the measured dislocation densities between TEM and XRD are common. In general XRD seems to slightly overestimate the threading dislocation density when the twist component is used, and underestimate the density when the broadening due to tilt is used [Lee et al (2005)]. Furthermore it should be noted that for very thin films XRD also samples the tilts associated with misfit dislocations at the GaN/SiC interface. If the Burger's vectors of the dislocations are randomly oriented, the associated strain fields will tend to cancel out as the film thickness increases, however if the Burger's vectors are not random the tilts can persist. This discussion shows that there are still some limitations on the current understanding of even the measurements of the more widely studied wurtzite GaN, and that defect densities estimated by XRD need be handled with care. This is especially the case when comparing samples of different layer thickness. The influence of layer thickness Typically the intensity spread of X-ray reflections is not constant but decreases with increasing film thickness, as shown in Fig. 81 by the full width at half maximum (FWHM) of the 002 reflection in co-scans. It is also evident that zincblende GaN (lattice constant GaN: a = 4.50597 A) grown on low lattice mismatch substrates, like 3C-SiC (lattice constant SiC: a = 4.3596 A, therefore 3.4 % compressive) and MgO (lattice constant MgO: a = 4.213 A, therefore 7.0 % compressive), exhibit a lower mosaicity than similar thick cubic GaN films grown with much larger mismatch on Si (lattice constant Si: a = 5.4311 A, therefore -17.0 % tensile) or GaAs (lattice constant GaAs: a = 5.65352 A, therefore -20.3% tensile). Moreover, Fig. 81 shows that MOVPE grown zincblende GaN (our data) compares well with state-of-the-art cubic GaN films grown by MBE [Kemper et al (2015); Martinez-Guerrero et al (2002)]. The decreasing intensity spread of the reflections with increasing film thickness is commonly associated with an overall reduction in defect density and an improvement in material quality for thicker epitaxial films. Transmission electron microscopy investigations reveal a strong reduction in stacking fault density with increasing layer thickness by reactions between pairs of stacking faults under formation of perfect edge dislocations or partial threading dislocation. Martinez-Guerrero et al (2002) observe a nearly exponential decay of the stacking fault density from 5 x 106 cm-2 to 3 x 105 cm-2 in the first 500 nm of zincblende GaN growth. In our MOVPE grown cubic GaN films TEM measurements have revealed that the stacking fault density reduces from 107 cm-2 directly at the template interface to 3 x 104 cm-2 close to the surface of 1200nm thick films. However, the stacking fault density affects foremost the shape and intensity profile along the SF-streak in reciprocal space (as shown in Barchuk et al) for basal plane stacking faults in wurtzite GaN and in Dupraz et al (2015) for stacking faults in face-centered cubic (fee) nano-crystals), but co-scans of the symmetric 002ZB reflection have almost no overlap with the stacking fault streak profile. Hence the observed peak narrowing with increasing layer thickness as shown in Fig. 71 cannot be directly related to the stacking fault reduction. Several reports in the literature (see the references indicated in Fig. 81) suggest that the trend in Fig. 71 is due to a reduction in threading dislocation density with increasing film thickness, as a result of threading dislocation reaction, but TEM evidence for this assertion is scarce [As (2010); Kemper (2015); Rusing (2016); Lischka (1997)]. Theoretical models predict that the threading dislocation density is inversely proportional to the film thickness t [Ayers (1995)]. Combining these models with the reflection broadening due to mosaicity reveals a decrease in the intensity spread by a factor oft-1 or t"1/2, depending on whether an oriented thin film (Eq. (10)) or a powder sample (Eq. (11)) is assumed. As one can see from the dashed lines in Fig. 81 the experimental data do not follow the predicted trend. Instead the observed decay is much weaker, following approximately a t"1/3 dependency. This may be explained by the generation of new threading dislocations, when stacking faults react with each other, and which to the best of our knowledge is not considered in current models. Furthermore one should consider that the models predict a threading dislocation density reduction after a certain thickness, while XRD is an integration method, which provides a weighted average value over the whole layer thickness. It should also not be forgotten that there is a natural reduction in the width of the X-ray reflections with increasing layer thickness, as the number of scattering atoms increases. This all makes a comparison of the material quality of samples with different thickness difficult. Material parameter for the strain analysis The lattice parameters for zincblende Ill-nitrides are not yet well established experimentally, since such thin films suffer from stacking disorder, undoubtably a high density of line defects and wurtzite inclusions, resulting in local strain variations and relatively broad reflections. Furthermore, most X-ray diffraction experiments on predominantly zincblende GaN films are focussed on phase purity analysis, rather than high resolution lattice parameter measurements. Our own measurements of the zincblende GaN lattice parameter using high-resolution 26-ω-scans of 8 on- and off-axis reflections and a least squares fit give a value of (4.50597± 0.00038) A, which is in good agreement with experimental data by Novikov et al (2010), and can be used as reference data for strain analyses in zincblende GaN thin films. Table 2 -Lattice parameters and elastic constants of wurtzite and zincblende GaN, InN, and AIN GaN InN AIN wurtzite 3wz (A) 3.18940 [Paszkowicz 3.5446 [Strite et al 3.11197 [Paszkowicz et et al (2004)] (1992)] al (2004)] Cwz (A) 5.18614 [Paszkowicz 5.7034 [Strite et al 4.98089 [Paszkowicz et et al (2004)] (1992)] al (2004)] zincblende azb (A) 4.4913 *1 4.9393 *1 4.3136 *1 4.5105 *2 5.0128 *2 4.4010 *2 4.5041 *3 4.9882 *3 4.3717 *3 4.510 ±0.005 5.01 ±0.01 4.373 ± 0.002 [As et al [Novikov(2010)]*4 [Schormann et al (2006)]*4 (2010)]*4 4.50597± 0.00038*5 5.02 ± 0.005 [Compean Garcfa et al (2015)] *4 — Cn (GPa) 293*6 187*6 304*6 C12 (GPa) 159 *6 125*6 160*6 *4: experiment *5: this work (experiment) *6: recommended by Vurgaftman and Meyer (2003) But, to the best of our knowledge, no experimentally determined accurate lattice parameters are mentioned in the literature for zincblende InN and AIN. In these cases it is therefore necessary to derive the values from the well-established wurtzite lattice parameters awz and cwz- However, in wurtzite-like Ill-nitrides strong internal electric fields lead to a distortion of the unit cell from the ideal shape, with 1.633. In reality Cwz is typically smaller and awz slightly larger than in the ideal case, and hence the estimated zincblende parameter of nominally unstrained Ill-nitrides can differ substantially, as one can see from the values in Table 2. As awz is less affected by the wurtzite unit cell distortion than cwz, this parameter gives reasonably good values for the natural lattice constant of the zincblende phase azb- Alternatively, the lattice parameters derived from the unit cell volumes can be used. Presumably, the natural unstrained lattice constants of zincblende nitrides are somewhere between these theoretical values, and this assumption is in good agreement with the experimental data known so far. Table 2 also contains the elastic constants Cn and C12 for the zincblende Ill-nitrides as stated in Vurgaftman and Meyer (2003), and which can be used for stress and strain calculations. Strain During the growth of thin films on foreign substrates and during the heterostructure growth of alloys with different compositions, the films are subjected to varying stresses, which often result in an elastic deformation of the crystal lattice. Such lattice strains have a significant impact on the physical properties and the performance of semiconductor devices. Hence, the understanding and monitoring of these strains during device development are of high importance. In the following sections we will discuss the different sources of strain, and describe how the strain in a thin zincblende GaN film can be measured. Lattice mismatch strain In epitaxial thin films the lattice mismatch between the thin film and the underlying template produces biaxial in-plane strains, when the size of both lattices are forced to match each other. Three different states are commonly used to describe the thin film deformation. The thin film is fully strained when its lattice matches the dimensions of the template lattice at the common interface, while the film is fully relaxed when its lattice is undistorted and has its natural dimensions. The state between both extremes is called partially relaxed. In reciprocal space the lattice mismatch strain results in a shift of the reciprocal lattice points (RLPs) of the GaN thin film from their expected position and with respect to the RLPs of the substrate. The relative separation between layer and buffer peaks can be either measured in several individual ω-26-scans, or more commonly by collecting a reciprocal space map in asymmetric geometry. The latter generally gives a better overview of the relationship between the X-ray reflections of the different layers, but a correction of the sample miscut by a second scan is required for the lattice mismatch strain evaluation. Moreover, one should take into account that the layer used as reference may be affected by the substrate as well, which lowers the accuracy of this method. Ideally one should use a substrate peak as reference, but there can be a large separation in reciprocal space for systems with a large mismatch. The strain in a certain direction of the thin film is then given as follows: where ao is the natural lattice constant, and a, is the measured constant in the same direction. Due to the relatively low material quality often found in zincblende nitride materials there are currently no accurate reference values for the natural lattice parameter available in the literature, as discussed in the previous section. For GaN one can use the experimentally determined values we provided in Table 2. For other group-Ill nitrides we recommend using the values which are derived from the wurtzite a-parameter or from the wurtzite cell volume (see Table 2), as the wurtzite lattice parameters are well known. Assuming that the thin film is stress-free in the growth direction (commonly labelled as z), the strain of a (001) oriented film in the growth direction is given via Hooke's law as where ex and ey are the strains in the two in-plane directions, and Cn and C12 are the materials' elastic constants (see Table 2) [Dunstan (1997)]. In the case of isotropic in-plane strain (ex = ey) this can be further simplified. The strain relations for other orientations differ from the equation above, and are published elsewhere in literature, such as in Dunstan (1997). Thermal mismatch strain & growth induced strain The small strain in an epitaxial thin film originates from the thermal mismatch between the used substrate and epilayer, or is formed in the early stage of growth. It is typically much smaller than strain due to lattice mismatch, but can be large compared to the residual mismatch strain in a partially relaxed film. As GaN has a larger thermal expansion coefficient than SiC and Si [Wahab et al (1994); La Via (2012); Okada and Tokumaru (1984)] the remaining thermal strain after cooling down from the growth temperature leads to a tension in the GaN film at the interface to the substrate. For typical zincblende GaN growth temperatures in the range between 700°C and 1000°C, the theoretical thermal strain is between 1.1 * 10"3 and 1.6 x 10"3 when a Si substrate is used. Growth-induced strain occurs due to island coalescence during nucleation on the substrate in the early stages of growth. Its magnitude is given by the smallest possible gap between two islands A and the average island size in the particular in-plane direction [Hoffman (1976)]. Relative lattice parameter measurements as described above are insufficiently accurate to determine such small strains, as the resolution is often low and as the substrate itself may also be influenced by the strain. Instead analyses of very small strains require absolute measurement of the lattice parameters, using high resolution 26-oo-scans of a larger set of different reflections. Then the as measured plane spacings dj are matched by the plane spacings of a model crystal using a least-squares fit: Often it is necessary to choose a suitable coordinate system, which describes the geometry of the problem better than the natural lattice. The following example illustrates this. Table 3 lists 26 values for different reflections, which were measured from a zincblende GaN film grown on a 3C-SiC/Si template with 4° miscut in [110] direction. The Bragg angles of all reflections tilted in the miscut direction hhlare significantly smaller than the similar reflections tilted away from the miscut h-hl, indicating a difference in the lattice dimensions for these two directions. In consequence the natural crystal lattice is slightly sheared within the growth plane. This can be simplified by using a new coordinate system x', y', z' with x' (y') parallel (perpendicular) to the sample miscut, and z' pointing in the growth direction. Note that by this coordinate transformation the new Using this approach, the least squares fit (see above) together with Bragg's law give a unit cell size of x' = (6.39566 ± 6.7 x 10-4) A, y' = (6.38465 ± 5.5 x 10"4) A, and z' = (4.49236 ± 3.2 x 10"4) A in the new coordinate system, and an anisotropy of the in-plane strain of εx' = (3.65 ± 0.11) x 10"3 resp. εy' = (1.92 ± 0.09) x 10"3. Table 3 - Reflections in the natural coordinate system (hkl) and rotated coordinate system (h'kT) of a zincblende GaN-sample, with 29 and A(29) derived from high-resolution 29oo-scans. Reflection hkl in Reflection h'kT Measured FWHM natural lattice in x'y'z' lattice 29 (°) A(29) O 002 002 40.0871 0.2452 004 004 86.6495 0.5002 1-13 023 69.2284 0.3132 2-24 044 114.0130 0.6891 3-31 061 96.1246 0.4233 113 203 69.1963 0.3690 224 404 113.8499 0.7261 331 601 95.9402 0.5198 Reflection hkl in I Reflection h'kT I Measured I FWHM natural lattice in x'y'z' lattice 29 (°) A(29) (°) =00l 002 40.0871 0.2452 004 004 86.6495 0.5002 1-13 023 69.2284 0.3132 2-24 044 114.0130 0.6891 3-31 061 96.1246 0.4233 113 203 69.1963 0.3690 224 404 113.8499 0.7261 331 601 95.9402 0.5198 It is known that a substrate miscut can lead to strain relaxation in epitaxial thin films via alignment of threading dislocations [Young et al (2010); Chen et al (2007)], but then the zincblende GaN layer would be less strained in the miscut direction (ex') than in the perpendicular direction (ey'). Since we observed the opposite case, we can rule out this relaxation mechanism for this sample. Instead the results indicate that the strain anisotropy is probably due to the coalescence of islands with different size in the two in-plane directions, which has been observed in atomic force microscopy images. Wafer curvature analysis In heteroepitaxial thin films stress above a certain level can be relieved by the formation of defects, or in the case of tensile surface stress by the formation of cracks. Moreover the stress in the thin film can be lowered by bowing of the whole sample. This is often the case in thick, medium stressed epilayers, such as templates and buffer layers. Thermal strains can also lead to a significant wafer bowing. This is especially a problem with large area templates with diameters up to 8", where even small bows lead to significant deviations in uniformity during growth and processing. Hence, control and management of strain and wafer bowing are of significant interest. The bow of a wafer can be determined with XRD by measuring the incident beam angle ω for the symmetric reflection at different positions of the sample X1 In bowed samples the lattice planes are also curved with the bow of the wafer. In consequence the incident angle needs to be corrected for different positions along the diameter of the wafer. Wafer curvature and radius of bow R are then given by the relative change of ω and X1: [Inaba (2014)]. Since the reflections of the zincblende GaN epilayer are often relatively broad, it is more suitable to use a narrower symmetric reflection of the underlying template. By using larger range of measurement positions, and a beam mask to reduce the irradiated area on the sample surface, the resolution of the measurement can be further improved. Figs. 82 and 83 show such a curvature measurement for a 4" 3C-SiC/Si template, where the 002 SiC reflection was used. The curvature of the wafer along the measured direction can be easily determined graphically by linear interpolation of the measured incident beam angles. A positive (negative) slope corresponds to a concave (convex) bow of the wafer. The example in Figs. 82 and 83 gives a convex bow of -51.5 km-1 (respectively R = -19.4 m). It should be noted however that this technique measures the curvature of the substrate planes. If the substrates already contain a high density of dislocations or a grain structure then the planes in the substrates may already be curved before layer growth and therefore the measured bow may not give an accurate reflection of the residual stress in the wafer. For this reason there may also be discrepancies between the bow measured by X-ray and optical techniques. We find that for the high quality templates used in these studies the discrepancy between the curvature measured by X-ray diffraction and optical techniques is negligibly small. Further studies Following on from the experimental work reported above, further studies have been carried out to investigate the effect of reaction pressure (and other parameters and conditions) on the growth of cubic zincblende GaN films. In summary, cubic zincblende GaN films were grown by metalorganic vapour-phase epitaxy (MOVPE) on 3C-SiC/Si(001) templates and characterized using Nomarski optical microscopy and X-ray diffraction. In particular, the surface morphology and material quality were assessed as a function of the low-temperature nucleation layer thickness (between 3 and 44nm in these experiments), the epigrowth temperature (in the range of 850 to 910 °C in these experiments) and the V/lll ratio (from 15 to 1200 in these experiments) for a reaction pressure of 100 Torr. The reduction in reaction pressure from 300 to 100 Torr as presented here forms the main difference with the earlier results reported above. In this particular case, a window for particularly suitable MOVPE growth conditions was identified between 850 and 890 °C with a V/lll ratio between 38 and 150, resulting in relatively smooth, zincblende GaN films with a wurtzite impurity content of less than 1%. The effect of a reduction in the epilayer reaction pressure from 300 to 100 Torr on the zincblende phase purity is illustrated in Figs. 111 and 112. These graphs show an equally low or lower wurtzite fraction for the 100 Torr data set under similar conditions of temperature and V/lll ratio. Hence, the MOVPE growth window for the good zincblende GaN material quality is substantially widened under reduced reaction pressure conditions. All samples were grown in a Thomas Swan 6x2" close-coupled showerhead MOVPE reactor on 3-SiC/Si templates. The SiC templates consisted of an about 3 um-thick 3C-SiC layer on a Si (001) substrate with a thickness of between 0.75 mm and 1 mm and a misorientation of 4° towards [110]. For the GaN growth, trimethyl gallium (TMG) and ammonia were used as Ga- and N-precursors, respectively, while hydrogen was used as the carrier gas. The total gas flow was kept constant at 20 standard litres per min (slm). The growth procedure comprised of a high-temperature thermal anneal of the substrate followed by a low-temperature nucleation layer deposition and finally the growth of the epilayer proper at high temperature. The temperatures quoted are those recorded by a Laytec EpiTT in-situ optical monitoring system, calibrated against an Al/Si eutectic wafer. The thermal annealing step of the template took place at 960 °C in a mixture of hydrogen and 3 slm ammonia. The GaN nucleation layer (NL) was grown at 600 °C, 500 Torr, and a V/lll ratio of 720 to a thickness of 44 nm for sample sets A and B (see below). The epilayer growth pressure was maintained at 100 Torr while the epilayer thickness was kept constant at 300nm. Two sample sets were prepared in which the variables were the epilayer growth temperature (sample set A) and the V/lll ratio (sample set B). For set A, the epilayer growth temperature was varied between 850 and 910 °C at a constant V/lll ratio of 76 in the gas phase. For set B, the V/lll ratio during epilayer growth at 880 °C was varied between 15 and 1200 by changing the ammonia flow rate at a constant TMG flow rate of 145 umol/min. All GaN epilayers were doped with Si to a nominal concentration of mid-1018 cm-3 using silane (50 ppm SiH4 in H2). Finally, a third sample set (sample set C) consisted of 300nm thick GaN epilayers grown at 875 °C, 100 Torr and V/lll of 76 on low-temperature nucleation layers that varied in thickness between 3 and 44nm. XRD phase analysis was performed using a PANalytical Empyrean diffractometer equipped with a Cu-Ka1 source (A = 1.54056 A), 2-bounce hybrid monochromator, 1/4°-slit, Eulerian cradle, and a PIXcel solid-state area detector. Reciprocal space maps (RSMs) around the 113 zb-GaN and 1-103 wz-GaN reflections were measured parallel and perpendicular to the miscut direction of the substrate. The intensity profile along the SF streak between the 113 and 1-103 reflections were extracted from the RSMs and subsequently fitted with a maximum of three Pseudo-Voigt functions for the zincblende and wurtzite phases and a third, ill-defined, hence defective phase probably related to stacking faults. The integrated intensities of the fitted profiles were used to quantify the wurtzite fraction of the GaN epilayers. The residual peak intensity assigned to stacking faults was not quantified in this work. The 002 peak broadening was measured with a Philips X'Pert diffractometer equipped with an asymmetric 4-crystal Bartels monochromator (A = 1.54056 A), 5x5mm2 cross slit collimator, Eulerian cradle, and gas proportional point detector without further secondary optics. The intensity profiles of open detector co-scans were fitted with Pseudo-Voigt functions. Sample set A - temperature series The first set of samples (set A) to be discussed consisted of six samples in which the epilayer growth temperature was varied between 850 and 910 °C, in steps of 10 to 15 °C. In this series of experiments, the V/lll ratio was kept constant at 76 which represented an intermediate value within the range of values explored in sample set B (see below). The surface morphology of the samples grown below 895 °C is characterized by elongated features or striae, as shown in the Nomarski optical micrographs in Figs. 88-93. The striae are aligned along the [1-10] direction, i.e. in the direction perpendicular to the substrate's miscut indicating an in-plane anisotropy. The elongated features shorten to a few micrometers in length with increasing temperature. At 895 °C and above, the surface becomes more granular and roughens further with increasing temperature. The XRD analysis of sample set A reveals that the wurtzite fraction measured perpendicular to the miscut increases gradually with increasing growth temperature but remains below 1% as shown in Fig. 94. The measurements obtained parallel to the miscut indicated no wurtzite inclusions along that direction and hence are not shown in the graph. Sample set B -V/lll ratio series The second series of samples (set B) consisted of eight samples in which the epilayer growth temperature was kept constant at 875 °C, i.e. an intermediate value of set A, while the V/lll ratio was varied from 15 to 1200, with a factor of about 2 between each value. The Nomarski optical micrographs in Figs. 95-102 show the variation in surface morphology with increasing V/lll ratio from granular to striated and again coarsening. The results of the XRD phase analysis of sample set B indicates a small increase of wurtzite inclusions from 0% at V/lll of 15 to 1% at V/lll of 300, as is shown in Fig. 103 when measured perpendicular to the miscut. At the highest V/lll values of 600 and 1200, the wurtzite fraction increases more rapidly to 3 % and finally 11%. No significant presence of the wurtzite phase was found in all samples when measured parallel to the miscut and hence is omitted from the graph. Sample set C - NL thickness series The third series of samples (sample set C) comprised of five samples in which the nucleation layer thickness was increased from 3 to 44nm by a factor of about 2 in each step, while the epilayer growth conditions was kept constant at a temperature of 875 °C, a pressure of 100 Torr and a V/lll ratio of 76. The Nomarski optical micrographs in Figs. 104-108 show the variation in surface morphology with changing NL thickness. Of the five samples, the only one that looks unusual is the one grown with the thinnest (3nm) NL showing large pits in the epilayer. Other than that, the surface striae seem to coarsen with increasing NL thickness. The XRD phase analysis of sample set C (see Fig. 109) indicates that the sample with the 3nm-thick NL has the greatest wurtzite fraction at about 3%, while the other samples have less than 1% of wurtzite phase. The integrated intensity of the 002 rocking curves for sample set C are shown in Fig. 110. The changing peak intensity indicates an improvement in material quality with increasing NL thickness and the peak intensity saturates at a NL thickness of 22nm or greater. Discussion of sample sets A, B and C The structural data of sample set A (varied temperature at a constant V/lll ratio of 76) reveal that a growth temperature below 895 °C yields relatively smooth film surfaces (see Figs 88-93) with a wurtzite-fraction of less than 1% (see Fig. 94). At the higher growth temperatures, the surface degrades marginally and the wz-GaN fraction edges towards 1%. Hence, the preferred growth temperature is below about 890 °C at a V/lll of 76 and reaction pressure of 100 Torr. A constant growth temperature of 875 °C (i.e. within the most favourable temperature window) was used for sample set B while the V/lll ratio was varied between 15 and 1200 with a factor of about 2 between each value. At the lower and higher ends of this range (see Figs. 95-102), the surfaces are more granular. The flattest surface is probably observed for the samples grown at a V/lll ratio between 38 and 150, as far as the Nomarski micrographs allow for a qualitative ranking. The wz-GaN contamination remains very small for the samples up to V/lll of 300, despite their differences in surface morphology (see Fig. 103). The fast rise in wurtzite contamination at a V/lll ratio above 300 is accompanied with a slight worsening of the surface morphology. It is worth pointing out that sample set A was grown at a V/lll ratio of 76, which falls in the range of V/lll ratios established to be most favourable from the study of sample set B. Similarly, sample set B was grown at an epilayer growth temperature of 875 °C, which falls in the temperature range established to be most preferred from the study of sample set A. Therefore, the surface and material properties of the zincblende GaN films should vary little across this MOVPE growth window. The study of the low-temperature nucleation layer thickness (sample set C) indicates that the preferred thickness is around 22nm. The use of a thinner GaN NL yields a loss of material quality while a thicker NL causes the surface morphology to coarsen. In conclusion, considering the collective data of surface morphology and phase purity of samples sets A to C, the preferred MOVPE growth conditions for zincblende GaN epilayers at a constant pressure of 100 Torr are found within a small parameter window of temperature between 850 and 890 °C and V/lll ratio of 38 to 150, resulting in a relatively smooth film with a wurtzite contamination of less than 1%. The preferred thickness of the NL is around 22nm. While the invention has been described in conjunction with the exemplary embodiments described above, many equivalent modifications and variations will be apparent to those skilled in the art when given this disclosure. Accordingly, the exemplary embodiments of the invention set forth above are considered to be illustrative and not limiting. Various changes to the described embodiments may be made without departing from the spirit and scope of the invention. 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A method according to claim 1 wherein the group Ill-nitride nucleation layer is a zincblende structure group Ill-nitride nucleation layer. 3. A method according to claim 1 or claim 2 wherein, before growing the group Ill-nitride nucleation layer, subjecting the 3C-SiC layer to a nitridation step at a temperature T1 in the range 800-1100 °C. 4. A method according to any one of claims 1 to 3 wherein the group Ill-nitride nucleation layer is grown at temperature T2 in the range 500-700 °C. 5. A method according to any one of claims 1 to 4 wherein the group Ill-nitride nucleation layer is grown to a thickness of greater than 3 nm and not more than 10Onm. 6. A method according to any one of claims 1 to 5 wherein the group Ill-nitride nucleation layer is grown to a thickness in the range 10-50nm. 7. A method according to any one of claims 1 to 6 wherein, in the step of depositing and growing the zincblende structure group Ill-nitride layer, the reactor pressure is not more than 500 Torr. 8. A method according to any one of claims 1 to 7 wherein, in the step of depositing and growing the zincblende structure group Ill-nitride layer, the reactor pressure is not more than 300 Torr. 9. A method according to any one of claims 1 to 3 wherein, in the step of depositing and growing the zincblende structure group Ill-nitride layer, the V-to-lll ratio is not more than 300. 10. A method according to any one of claims 1 to 9 wherein V-to-lll ratio is not more than 150. 11. A method according to any one of claims 1 to 10 wherein temperature T3 is in the range 800-920 °C. 12. A method according to any one of claims 1 to 11 wherein temperature T3 is in the range 820-890 °C. 13. A method according to any one of claims 1 to 12 wherein the group Ill-nitride layer is a lnxAlyGai-x-yN based layer, where 0