Zinc blende structure group-III nitrides

Abstract

A method of manufacturing a semiconductor structure including a (001)-oriented zinc blende structure group III nitride layer, such as GaN, is disclosed. The layer is formed on a 3C-SiC layer on a silicon substrate. A nucleation layer is formed, recrystallization is carried out, and then a zinc blende structure group III nitride layer with a thickness of at least 0.5 μm is formed by MOVPE at a T3 temperature in the range of 750 - 1000 °C. A corresponding semiconductor structure including a zinc blende structure group III nitride layer is also disclosed, which, when characterized by XRD, is formed mostly or entirely of zinc blende structure group III nitride, preferentially over wurtzite structure group III nitride.

Description

[0001] This application is a divisional application of a Chinese application with the application number 201880036363.7 and the invention title "Zinc Blende Structure Group III Nitrides". This application is the application of the PCT international patent application PCT / EP2018 / 058250 entering the Chinese national phase. The application date is March 29, 2018, and it claims the priority of PCT / EP2017 / 057764, with the priority date being March 31, 2017. Technical Field

[0002] The present invention relates to the formation of zinc blende structure group III nitride layers, such as GaN, AlGaN, InGaN, InAlN, and more generally In x Al y Ga 1-x-y N. The characterization of these layers and their formation methods are disclosed herein. These materials have specific but not necessarily exclusive applications in the field of semiconductor structures and devices, such as light-emitting applications, such as LEDs, lasers, and other devices such as transistors, diodes, sensors, etc. Background Art

[0003] Group III nitride semiconductors can be used in a wide range of optoelectronic applications, such as multi-quantum well (MQW) LEDs and laser diodes that emit in the blue and green spectral regions. Such devices typically grow along the c-axis of the hexagonal wurtzite phase, where a strong internal polarization field passing through the quantum well causes a decrease in the emission wavelength related to the radiative recombination rate and current density [Miller et al (1985); Fiorentini et al (1999); Hammersley et al (2015)]. Although these effects are somewhat mitigated by using thin QW layers (usually 2 - 4 nm thick), for green light-emitting structures, longer radiative recombination lifetimes and relatively low internal quantum efficiencies can be observed [Nippert et al (2016); Hammersley et al (2016)]. Currently, QW structures have been designed to grow along the non-polar axes (such as the a-plane and m-plane) of the wurtzite GaN phase to avoid the polarization field and related limitations. Although the radiative lifetimes of non-polar wurtzite devices are very short et al (2013); Dawson et al (2016)] and the wavelength characteristics are independent of the current density [Detchprohm et al (2010)], and its quantum efficiency never exceeds that of the polar c-plane structure [Dawson et al (2016)]. A possible explanation for the poor performance of non-polar wurtzite devices with green emission is that in the absence of the quantum-confined Stark effect (QCSE) related to the polarization field, an indium-rich quantum well is required to achieve green emission [Fiorentini et al (1999)]. In addition to the increased interfacial strain between the GaN buffer layer and the barrier layer, the indium incorporation efficiency of the non-polar growth plane is lower compared to the polar growth plane, and a lower process temperature is required for the growth of the indium-rich layer [Zhao et al (2012)]. This potentially leads to a high density of impurities and point defects, which act as non-radiative recombination centers and further reduce the efficiency of the non-polar wurtzite MQW [Chichibu et al (2005)].

[0004] Therefore, after attracting strong attention in the mid-1990s, the GaN-related structures in the cubic zinc blende phase reappeared as a promising method to improve the efficiency of green-wavelength LEDs.

[0005] The cubic zinc blende InGaN / GaN MQW structure grown on the (001) plane may not have a polarization field because in the zinc blende phase, these fields are only caused by shear stress, and there is no shear stress in the (001)-oriented film [Hanada (2009)]. Therefore, compared with the c-plane hexagonal structure, the electron-hole wave function overlap increases, which will lead to an increase in the radiative recombination rate. In addition, for a given indium content, the energy gap of InGaN in the cubic phase is narrower than that in the hexagonal phase et al (2006); Compeán García et al (2015)], enabling the achievement of green-wavelength emission at an indium content lower than that of the non-polar wurtzite structure. However, since zinc blende GaN and InGaN are metastable under most growth conditions, the zinc blende film may contain more stable wurtzite polytype inclusions, and thus contains a large number of wurtzite-like stacking faults and platelets [Trampert et al (1997); Shen et al (2003); Wu et al (1997); Yang et al (1996)].

[0006] US2016 / 0247967 discloses suitable substrates for growing cubic GaN layers. These substrates include single-crystalline silicon wafers having spaced-apart single-crystalline silicon carbide layers formed on the silicon wafers, and amorphous or polycrystalline silicon carbide layers formed between the single-crystalline silicon carbide layers. The single-crystalline SiC is in the form of the 3C-SiC polytype. GaN is formed above the SiC layer, and epitaxy allows single-crystalline GaN to form above single-crystalline SiC, while polycrystalline GaN forms above polycrystalline / amorphous SiC. The effect of this is to release stress for the single-crystalline GaN region. Other methods may be employed to control stress in the GaN layer, for example, introducing lattice mismatch layers such as compositionally graded AlGaN, which cause compressive stress during growth to counteract the tensile stress caused by the thermal expansion mismatch. This method has been used to grow conventional GaN on Si structures [Zhu et al (2013)]. US2016 / 0247967 does not provide substantial details on how to grow GaN, or does not explain or quantify the content of the wurtzite-like regions of GaN in single-crystalline GaN. In this regard, the disclosure of US2016 / 0247967 is limited to stating that a GaN layer can be formed at a temperature below 1000 °C (preferably between 800 °C and 950 °C) using a metalorganic vapor phase epitaxy (MOVPE) process. Summary of the Invention

[0007] In order to obtain a single-phase epitaxial film with reasonable crystal quality, the growth process must be supported by powerful structural characterization techniques.

[0008] To solve at least one of the above problems, the inventors have designed the present invention. Preferably, the present invention reduces, improves, avoids, or overcomes at least one of the above problems.

[0009] Accordingly, in a first aspect, the present invention provides a method of manufacturing a semiconductor structure including a zinc blende structure Group III nitride layer having a substantially (001) orientation, the method comprising the steps of:

[0010] Providing a silicon substrate;

[0011] Providing a 3C-SiC layer on the silicon substrate;

[0012] Growing a Group III nitride nucleation layer;

[0013] Performing a nucleation layer recrystallization step; and

[0014] Depositing and growing a zinc blende structure Group III nitride layer having a thickness of at least 0.3 μm at a T3 temperature in the range of 750 - 1000 °C by MOVPE.

[0015] The present inventors have found that these steps provide a zinc blende structured group III nitride layer with improved crystallization quality, particularly in terms of reducing the formation of wurtzite structured group III nitride inclusions.

[0016] In the present disclosure, some numerical ranges are expressed as open-ended ranges with upper or lower limits or as closed-ended ranges with upper and lower limits. It is expressly stated herein that preferred ranges are disclosed, which are combinations of the upper and / or lower limits of different ranges of the same parameter.

[0017] Preferably, before growing the group III nitride nucleation layer, the 3C-SiC layer is subjected to a nitridation step at a T1 temperature in the range of 800 - 1100 °C. This step is beneficial for ensuring that there is sufficient available N for forming the subsequent group III nitride. Using a temperature T1 outside this range will reduce the PL NBE peak intensity and broaden the emission FWHM.

[0018] Now we consider the deposition and growth conditions of the group III nitride nucleation layer. Preferably, the group III nitride nucleation layer is grown at a T2 temperature in the range of 500 °C to 700 °C. More preferably, the temperature T2 is in the range of 550 - 650 °C. The growth rate can be at least 0.1 nm / s. The growth rate can be as high as 1 nm / s. The thickness of the nucleation layer can be at least 3 nm, but more preferably greater than 3 nm. More preferably, the thickness of the NL can be at least 10 nm. The thickness of the nucleation layer (NL) can be as high as 100 nm. Preferably, the thickness of the NL can be as high as 50 nm. More preferably, the thickness of the NL can be as high as 40 nm. Preferably, the selected T2 temperature is about 40 - 60 °C higher than the temperature at which the growth rate deviates from a constant value to a lower value, thereby entering a state where the ammonia flow determines the growth rate.

[0019] After the growth of the nucleation layer, a nucleation layer recrystallization step is carried out. In this step, preferably, the temperature is raised at a rate of 0.1 - 10 °C / second. More preferably, the temperature is raised at a rate of 0.5 - 5 °C / second. This method is suitable for achieving satisfactory recrystallization of the nucleation layer, allowing for subsequent deposition of high-quality epitaxial layers.

[0020] The group III nitride nucleation layer is preferably a zinc blende structured group III nitride nucleation layer.

[0021] In the step of depositing and growing a zinc blende structured group III nitride layer on the recrystallized nucleation layer, the reactor pressure is preferably not greater than 500 Torr. More preferably, the reactor pressure is not greater than 300 Torr. More preferably, the reactor pressure is not greater than 100 Torr.

[0022] In the step of depositing and growing a zinc blende structured group III nitride layer on a recrystallized nucleation layer, the ratio of V to III is preferably in the range of 10 - 300. More preferably, the ratio of V to III is in the range of 20 - 150. Even more preferably, the ratio of V to III is in the range of 50 - 100. During this step, the growth rate is preferably in the range of 0.1 - 1 nm / second. For example, a growth rate of about 0.5 nm / second is suitable. Carefully selecting the ratio of V to III within the preferred range can improve the surface morphology, zinc blende phase purity, and XRD rocking curve peak width.

[0023] In the step of depositing and growing a zinc blende structured group III nitride layer on a recrystallized nucleation layer, the temperature T3 is preferably in the range of 800 - 920 °C. More preferably, the temperature T3 is at least 810 °C, even more preferably at least 820 °C, even more preferably at least 830 °C. The temperature T3 is preferably at most 910 °C, at most 900 °C or at most 890 °C. A particularly suitable range for T3 is 845 - 880 °C. Carefully selecting the temperature T3 within the preferred range can improve the Nomarski image surface morphology, XRD rocking curve peak width, and PL. For example, samples grown in the range of 860 - 880 °C show a relatively smooth surface, and the corresponding NBE PL peak is the strongest, although it is also the broadest in the data shown here. At higher growth temperatures, the surface becomes rough, the PL NBE peak significantly narrows, but the intensity of the yellow band increases.

[0024] The change in surface roughness with growth temperature can be confirmed by AFM. The phase purity determined by X-ray diffraction indicates that when T3 is 900 °C or lower, the amount of wurtzite inclusions can be greatly reduced. When T3 is higher than 900 °C, XRD analysis shows that the reflection contribution attributed to the wurtzite lattice increases, thus indicating the incorporation of hexagonal inclusions into the cubic zinc blende matrix.

[0025] The inventors have found in this work that by performing epitaxial layer growth at a relatively low pressure, the preferred conditions for temperature T3 and III-V ratio can be broadened. In a set of exemplary conditions for growing a zinc blende GaN epitaxial layer by MOVPE at a constant pressure of 100 Torr, T3 can be in the range of 850 to 890 °C, and the V / III ratio ranges from 38 to 150, resulting in a relatively smooth thin film with less than 1% wurtzite contamination. The preferred thickness of the NL is in the range of 10 - 50 nm, for example, about 22 nm.

[0026] Preferably, the group III nitride layer is based on In x Al y Ga 1-x-y N layer, where 0 ≤ x ≤ 1, 0 ≤ y ≤ 1.

[0027] The diameter of the silicon substrate is at least 100 mm. There may be different substrate diameters. It is noted that the growth process described herein can be easily extended to substrates of any suitable size, such as at least 150 mm, at least 200 mm or at least 300 mm.

[0028] In a second aspect, the present invention provides a semiconductor structure comprising a zinc blende structured Group III nitride layer, wherein:

[0029] the thickness of the Group III nitride layer is at least 0.5 μm; and

[0030] the Group III nitride layer is a single crystal zinc blende structured Group III nitride, and when the Group III nitride layer is characterized by XRD, the intensity I 10-11 attributed to the 10-11 reflection of the wurtzite structured Group III nitride and the intensity I 002 attributed to the 002 reflection of the zinc blende structured Group III nitride satisfy the following relationship:

[0031]

[0032] Unless the context otherwise requires, the features of the first aspect of the present invention can be combined individually or in any combination with the features of the second aspect of the present invention.

[0033] Preferably, at least one of the following relationships is applicable:

[0034]

[0035] More preferably, the intensity I 10-11 attributed to the 10-11 reflection of the wurtzite structured Group III nitride and the intensity I 002 attributed to the 002 reflection of the zinc blende structured Group III nitride satisfy the following relationship:

[0036]

[0037] The intensity I 10-11 attributed to the 10-11 reflection of the wurtzite structured Group III nitride and the intensity I 002 attributed to the 002 reflection of the zinc blende structured Group III nitride can be determined by two-dimensional reciprocal space mapping to form a measured reciprocal space map, which contains the expected reflections of the zinc blende structured Group III nitride 002 and the wurtzite structured Group III nitride 10-11. The reciprocal space map is a technique well known to those skilled in the art and can effectively capture a large amount of data to indicate the crystal phases present in the film.

[0038] The intensity I 10-11It may be caused by stacking faults formed on the {111} plane of the zinc blende structure group III nitride. The elongated stripes between the reflection of the 002 of the zinc blende structure group III nitride and the reflection of the expected wurtzite structure group III nitride 10-11 in the measured reciprocal space map can prove this. In this way, there may be measurable reflected X-ray intensity at defined positions in the reciprocal space, but this does not necessarily mean the existence of wurtzite structure group III nitride inclusions. Instead, the X-ray intensity can be provided by the reflection of the stacking faults. Compared with wurtzite structure inclusions, the hexagonal stacking faults have less impact on the properties of the group III nitride layer.

[0039] In a third aspect, the present invention provides a zinc blende structure group III nitride layer, wherein:

[0040] The thickness of the group III nitride layer is at least 0.5 μm; and

[0041] The group III nitride layer is a single crystal zinc blende structure group III nitride. When the group III nitride layer is characterized by XRD, the volume V zb of the zinc blende structure group III nitride and the volume V wz of the wurtzite structure group III nitride satisfy the following relationship:

[0042]

[0043] wherein V wz is evaluated based on the 1-103 reflection of the wurtzite structure group III nitride, and V zb is evaluated based on the 113 reflection of the zinc blende structure group III nitride, and the evaluation is based on:

[0044]

[0045] wherein:

[0046] V uczb is the volume of the unit cell of the zinc blende structure group III nitride,

[0047] V ucwz is the volume of the unit cell of the wurtzite structure group III nitride,

[0048] F 113 is the structure amplitude of the 113 reflection of the zinc blende structure group III nitride,

[0049] F 1-13 is the structure amplitude of the 1-103 reflection of the wurtzite structure group III nitride,

[0050] 2θ 113 is the 2θ angle of the 113 reflection of the zinc blende structure group III nitride,

[0051] 2θ 1-13 is the 2θ angle of the 1-103 reflection of the wurtzite-structured group III nitride,

[0052] I 113 is the integrated intensity of the 113 reflection of the zinc blende-structured group III nitride,

[0053] I 1-13 is the integrated intensity of the 1-103 reflection of the wurtzite-structured group III nitride.

[0054] Unless the context otherwise requires, the features of the first aspect of the present invention and / or the features of the second aspect of the present invention may be combined with the features of the third aspect of the present invention individually or in any combination.

[0055] Preferably, at least one of the following relationships is applicable:

[0056]

[0057]

[0058] Preferably, the zinc blende-structured group III nitride layer of the semiconductor structure of the second and / or third aspect is substantially (001)-oriented.

[0059] The thickness of the zinc blende-structured group III nitride layer of the semiconductor structure of the second and / or third aspect may be at least 0.3 μm.

[0060] In any one of the first, second or third aspects, the thickness of the zinc blende-structured group III nitride layer of the semiconductor structure may be at least 0.4 μm, at least 0.6 μm, at least 0.8 μm, at least 1 μm, at least 1.5 μm, or at least 2 μm.

[0061] The zinc blende-structured group III nitride layer may have a reflective layer interposed between the zinc blende-structured group III nitride layer and the substrate. This is useful for structures used as devices, particularly in light-emitting devices.

[0062] The zinc blende-structured group III nitride layer may have dimensions of at least 1 mm in two directions orthogonal to each other and to the thickness direction. More preferably, these dimensions may be at least 2 mm, at least 3 mm, at least 4 mm or at least 5 mm.

[0063] However, it should be understood that for some useful devices, the dimensions of the zinc blende-structured group III nitride layer may be smaller than the above dimensions. For example, the structure is cut for a specific device.

[0064] The present invention also provides a semiconductor device incorporating a semiconductor structure according to the second or third aspect, wherein the semiconductor device is selected from: a light emitting diode (LED), a laser, a diode, a transistor, a sensor.

[0065] Other optional features of the present invention are set forth below. BRIEF DESCRIPTION OF THE DRAWINGS

[0066] Embodiments of the present invention will now be described by way of example with reference to the accompanying drawings, in which:

[0067] Figure 1 Schematically illustrates the optical path and different goniometer movements in X-ray diffraction characterization.

[0068] Figure 2 Schematically explains the measurement geometry for mapping texture in reciprocal space and its projection onto a two-dimensional plot.

[0069] Figure 3 、 Figure 6 、 Figure 9 、 Figure 12 and Figure 15 Provide Nomarski microscope images showing the effect of nitridation temperature on the surface morphology of the zinc blende GaN film.

[0070] Figure 4 、 Figure 5 、 Figure 7 、 Figure 8 、 Figure 10 、 Figure 11 、 Figure 13 、 Figure 14 、 Figure 16 and Figure 17 Show the variation of photoluminescence (PL) optical properties with nitridation temperature.

[0071] Figure 18 Illustrates the growth rate of the nucleation layer (NL) as a function of growth temperature and V / III ratio.

[0072] Figures 19 - 27 : Figure 19 、 Figure 22 and Figure 25 Show the variation of the surface morphology observed from Nomarski microscope images with growth temperature (845 °C, 860 °C and 880 °C). Figure 20 、 Figure 21 、 Figure 23 、 Figure 24 、 Figure 26 and Figure 27 Show the variation of PL optical properties with growth temperature (845 °C, 860 °C and 880 °C).

[0073] Figure 28 Shows the variation of surface roughness (hollow circle symbols) and hexagonal-cubic ratio (solid circle symbols) observed by atomic force microscopy (AFM) with growth temperature.

[0074] Figure 29 Shows the variation of the FWHM of the XRD 002 and 004 rocking curves with growth temperature.

[0075] Figures 30 - 41 : Figure 30 、 Figure 33 、 Figure 36 and Figure 39 Shows the variation of the surface morphology of the Nomarski image with the V-III ratio (in the range of 23 to 152) at a constant growth temperature of 880 °C. Figure 31 、 Figure 32 、 Figure 34 、 Figure 35 、 Figure 37 、 Figure 38 、 Figure 40 and Figure 41 Shows the variation of the PL optical properties with the V-III ratio (in the range of 23 to 152) at a constant growth temperature of 880 °C.

[0076] Figures 42 - 50 : Figure 42 、 Figure 45 and Figure 48 Shows the variation of the surface morphology of the Nomarski image with the growth pressure (100, 300, and 500 Torr). Figure 43 、 Figure 44 、 Figure 46 、 Figure 47 、 Figure 49 and Figure 50 Shows the variation of the PL optical properties with the growth pressure (100, 300, and 500 Torr).

[0077] Figures 51 - 56 : Figure 51 and Figure 54 Shows the variation of the surface morphology of the Nomarski image with the thickness of the layer grown under the conditions of T = 860 °C, P = 300 Torr, and V / III = 80 (example layer thicknesses are 600 and 750 nm). Figure 52 、 Figure 53 、 Figure 55 and Figure 56 Shows the variation of the PL optical properties with the layer thickness (example layer thicknesses are 600 and 750 nm).

[0078] Figures 57 - 62Shows the variation of the surface morphology of the Nomarski image with the silane flow rate for growing a 500 nm thick film at T = 860 °C, P = 300 Torr, and V / III = 80.

[0079] Figures 63 - 68 Shows the variation of the PL optical properties with the silane flow rate.

[0080] Figure 69 Shows the Si concentration in the GaN film measured by SIMS based on the silane flow rate.

[0081] Figure 70 Shows the morphology of the Nomarski image of a typical green QW structure according to an embodiment of the present invention, taken at the center of a single mesa.

[0082] Figure 71 and Figure 72 respectively show Figure 70 the PL optical properties of the green QW structure in

[0083] Figures 73 - 75 Shows the pole figures of different wurtzite and zinc blende reflections of GaN grown on a (001)-oriented 3C-SiC / Si template.

[0084] Figure 76 Shows the crystal arrangement of the wurtzite and zinc blende GaN phases.

[0085] Figure 77 and Figure 78 Shows the two-dimensional reciprocal space map (RSM) of a zinc blende GaN sample, where the sample was grown under non-optimal conditions promoting the formation of wurtzite inclusions ( Figure 77 ) and under improved conditions yielding nearly 100% pure zinc blende GaN ( Figure 78 ), with stacking fault fringes (SF), detector fringes (DS), crystal truncation rods (CTR), and Bragg rings (BR).

[0086] Figure 79 Shows the XRD peak width plot of an optimized zinc blende GaN sample in a conventional Williamson-Hall plot, where Lorentzian (n = 1) and Gaussian-shaped peaks (n = 2) are fitted.

[0087] Figure 80 Shows the extrapolated peak width β hkl ·|G hkl | versus the polar angle c and the scattering vector magnitude |G hkl | estimated by a series of symmetric ω-scans with tilting. The circles show the measured reflections used for extrapolation.

[0088] Figure 81 It shows that for wurtzite GaN grown on relevant substrates 3C-SiC, GaAs, MgO, and Si, the XRD ω-linewidth (FWHM) decreases with the increase in film thickness.

[0089] Figure 82 and Figure 83 It shows an example of XRD wafer curvature analysis of a 4” 3C-SiC / Si template, showing a convex curvature of -51.5 km -1 .

[0090] Figure 84 It shows a graph of the relationship between the PL emission wavelength and the PL integrated intensity and the QW thickness.

[0091] Figure 85 It shows a process flow chart for forming a GaN layer according to an embodiment of the present invention.

[0092] Figure 86 It shows a schematic cross-sectional view of forming a GaN layer on SiC / Si.

[0093] Figure 87 It shows a schematic cross-sectional view of transferring a GaN layer onto another substrate for use as a semiconductor device (such as an LED).

[0094] Figures 88 - 93 It shows Nomarski optical micrographs showing the surface of a GaN epitaxial layer grown at a constant V / III of 76 and a pressure of 100 Torr in the temperature range of 850 to 910 °C. The temperature used is marked on each image.

[0095] Figure 94 It shows the variation of the wurtzite fraction of a GaN epitaxial layer obtained from XRD with temperature, where the epitaxial layer is grown at a constant V / III of 76 and a pressure of 100 Torr.

[0096] Figures 95 - 102 It shows Nomarski optical micrographs showing the surface of a GaN epitaxial layer grown at a temperature of 875 °C with a V / III of 15 - 1200 and a pressure of 100 Torr. The V / III ratio used is marked on each image.

[0097] Figure 103 It shows the variation of the wurtzite fraction of a GaN epitaxial layer obtained from XRD with the V / III ratio, where the epitaxial layer is grown at a constant temperature of 875 °C and a pressure of 100 Torr.

[0098] Figures 104 - 108Shows Nomarski optical microscopy images showing the relationship between the GaN epitaxial layer grown at a temperature of 875 °C with a V / III of 76 and a pressure of 100 Torr and the thickness of the GaN nucleation layer (NL). The thickness of the nucleation layer used is marked on each image.

[0099] Figure 109 Shows the change in the wurtzite fraction of the GaN epitaxial layer obtained from XRD as a function of the thickness of the GaN nucleation layer (NL), where the epitaxial layer is grown at a temperature of 875 °C, a V / III of 76, and a pressure of 100 Torr.

[0100] Figure 110 Shows the relationship between the integrated intensity of the XRD 002 rocking curve of the GaN epitaxial layer grown at 875 °C, 100 Torr, and a V / III of 76 and the thickness of the GaN nucleation layer (NL).

[0101] Figure 111 and Figure 112 : The change in the wurtzite fraction in the zinc blende GaN epitaxial layer grown at reaction pressures of 100 Torr and 300 Torr as determined by XRD as a function of temperature ( Figure 111 ) and V / III ratio ( Figure 112 ). Figure 111 The temperature-dependent samples in [[]] are grown with a constant V / III of 76. Figure 112 The V / III-dependent samples in [[]] are grown at 875 °C (100 Torr) and 880 °C (300 Torr). Detailed Description

[0102] Solving the green gap problem is a key challenge for the future development of LED-based lighting systems. A promising approach to achieving higher LED efficiency in the green spectral region is to grow group III nitrides in the cubic zinc blende phase. However, the metastability of zinc blende GaN and the crystal growth process typically result in phase mixing with the wurtzite phase, high mosaicity, high density of extended defects and point defects, and strain, all of which can degrade the performance of light-emitting devices. X-ray diffraction (XRD) is the primary characterization technique for analyzing these device-related structural properties because it is very inexpensive and provides rapid feedback compared to other techniques. In the present disclosure, we present various XRD techniques to identify the phase purity, primarily in zinc blende GaN thin films, to analyze their mosaicity, strain state, and wafer curvature. Different techniques are employed and illustrated on samples grown on 4" SiC / Si wafers by MOVPE or MOCVD (metalorganic vapor deposition).

[0103] X-ray diffraction (XRD) is a method suitable for the above purposes, as it is non-destructive, well-established, and can quickly provide detailed information about the structural characteristics of crystalline materials. In this paper, we illustrate how the XRD technique can be used to identify texture, phase purity, crystal orientation, and quantify the mosaicity of cubic zinc blende thin films. We use the proposed technique to characterize the properties of epitaxial GaN thin films grown on low-cost, large-area (001) cubic 3C-SiC / Si templates in order to provide rapid feedback for further growth optimization.

[0104] Crystallographic Properties of Group III Nitrides

[0105] Group III nitride materials AlN, GaN, InN and their alloys, Al x Ga 1-x N, In y Ga 1-y N, In x Al y Ga 1-x-y N (0 ≤ x ≤ 1, 0 ≤ y ≤ 1) can crystallize in wurtzite, zinc blende, and rock salt structures, where the first two are the most common phases in epitaxial thin films [Ambacher (1998); Hanada (2009)]. The hexagonal wurtzite and cubic zinc blende phases of GaN-based semiconductors are two different polymorphs of the same material. In both structures, the bonds between metal ions and nitrogen ions are tetrahedrally coordinated, and the interionic distances within the close-packed planes are approximately the same. The main difference between these two structures is the stacking order of these planes. In the wurtzite structure, for the (0001) plane, it is …AaBbAaBbAaBb…; while in the zinc blende structure, for the (111) plane, it is …AaBbCcAaBbCc…, where Aa, Bb, and Cc represent different metal-N bilayers. The interplanar distances in the wurtzite structure are

[0106]

[0107] in the zinc blende structure are

[0108]

[0109] [Cullity (1978)]. Here a and c are the respective lattice parameters of each structure, and h, k, and l are the Miller-Bravais indices of the crystal planes. The crystallographic similarity of the two polymorphs, as well as the similar formation energies of the two phases [Yeh (1992)], allows for the formation of a certain fraction of the two polymorphs during the growth of materials with defects.

[0110] It should be noted that, as is well known, the cubic structure can be described by three Miller index symbols (h, k, l), while the hexagonal structure can be described by Bravais-Miller index symbols (h, k, i, l). However, in the Bravais-Miller index symbols, i = -(h + k), so it is also entirely possible to describe the crystal planes and X-ray reflections of the hexagonal structure by three Miller index symbols (h, k, l).

[0111] Experimental Description and Basics of X - ray Diffraction

[0112] X-ray diffraction (XRD) is one of the most commonly used methods for characterizing crystalline samples. This method is based on the measurement of X-ray reflections, and the pattern of X-ray reflections represents the Fourier transform image of the crystal structure in reciprocal space. The diffraction angle (2θ) of the hkl reflection and the interplanar spacing d of the (hkl) plane hkl are interrelated by Bragg's law:

[0113] 2 hkl ·sinθ = λ (Equation 3).

[0114] For X-ray characterization of cubic zinc blende GaN thin films, we used two different standard laboratory diffractometers and a Cu-Kα1 light source High-resolution measurements were carried out on a Philips X'Pert diffractometer, where the radiation from the X-ray tube was filtered by an asymmetric four-crystal Bartels monochromator. Thereafter, an adjustable crossed-slit collimator further reduced the size of the beam and decreased the divergence angle to a few arcseconds, after which it impinged on the sample at the incident angle ω. Then the X-rays scattered from the sample at 2θ were measured, which could be measured directly using a gas proportional point detector (open detector configuration) or after passing through an additional monochromator (triaxial configuration) for high-resolution analysis. The sample was mounted on an Eulerian cradle, which could rotate about the sample normal (φ) and tilt the sample relative to the beam path plane (bpp) (χ). By changing the x- and y-positions of the sample stage and reducing the irradiation area, different regions on the sample surface could be focused and analyzed. Figure 1 Illustrates the beam geometry and the different movements of the sample stage.

[0115] Reciprocal space maps were measured using a PANalytical Empyrean diffractometer equipped with a 2-bounce hybrid monochromator, 1 / 4° slits, an Eulerian cradle, and a PIXcel solid-state area detector. This configuration ensured high intensity and allowed for rapid and precise measurement of large maps in reciprocal space.

[0116] Since correct alignment is an important factor for accurate evaluation of Bragg reflections and lattice properties, the goniometer 2θ angle was calibrated for the primary beam before each measurement. Then, following the best practice method recommended by Fewster and Andrew (1995), the sample was moved into the primary beam path (z movement) until half of the intensity was blocked.

[0117] Measurements were mainly carried out on wurtzite GaN thin films grown by metalorganic vapor phase epitaxy (MOVPE) on a 4'(001) cubic 3C-SiC template deposited on a Si substrate, which will be discussed in detail below. To relieve the strain in the large-area template, the 3 μm to 8 μm thick SiC layer was etched with a polycrystalline square grid to form mesa structures with millimeter-scale lengths. The present invention is not limited to using substrates of this size, but it should be noted that embodiments of the present invention allow the use of large-area substrates, which allows for efficient growth and device processing.

[0118] Suitable SiC / Si substrates are disclosed in US2016 / 0247967.

[0119] The SiC layer was polished by chemical mechanical polishing (CMP) to reduce the surface roughness (measured by AFM) from about 5 - 10 nm to less than about 1 nm.

[0120] Thin - Film Preparation

[0121] The growth of wurtzite GaN and InGaN using MOVPE is described below. The so-called two-step growth method was used, which includes performing substrate cleaning and a nitridation step at high temperature (800 °C < T1 < 1100 °C), growing a thin GaN nucleation layer at low temperature (500 °C < T2

[0122] < 700 °C), a nucleation layer recrystallization step, and depositing a GaN layer at high temperature (750 °C < T3 < 1000 °C).

[0123] The temperatures cited here are high-temperature measurement values corrected for emissivity measured using the in-situ monitoring tool EpiTT provided by LayTec AG. This value is calibrated relative to a Si / Al eutectic wafer.

[0124] Figure 85 A process flow diagram for forming a GaN layer according to an embodiment of the present invention is shown.

[0125] Figure 86A schematic cross-sectional view showing the formation of a GaN layer and active devices on SiC / Si is presented. The thickness of the layers is not drawn to scale. The silicon substrate 100 has a 3C-SiC layer 102. A GaN nucleation layer 104 is formed on the 3C-SiC layer 102. A cubic GaN buffer layer 106 is formed on the GaN nucleation layer 104. Then, an active device layer 108 is formed.

[0126] Figure 87 Shows Figure 86 A schematic cross-sectional view of the transfer of the GaN layer formed in to another substrate for use as a semiconductor device (e.g., an LED) is shown. First, a reflective p-contact 110 is formed on the active device layer 108. At this time, if necessary, a buried n-contact layer (not shown) can also be formed. Then, the device structure is bonded to a silicon processing wafer 114 using a bonding layer 112 (including an optional diffusion barrier). After bonding, the original Si / SiC substrate is removed to prevent light absorption. The nucleation layer 104 and the cubic GaN buffer layer 106 can be left in place as shown and used as part of any light extraction structure, or they can also be removed, leaving only the active device layer 108 bonded to the silicon processing wafer 114 for further device processing.

[0127] The thin film is characterized by well-known photoluminescence (PL) techniques, in which the thin film is irradiated above the bandgap of the GaN material with a laser (266 nm Q-switched). This promotes the formation of a large number of electron-hole pairs, which recombine to emit light. The emission spectrum is captured. It should be noted that the bandgap of cubic GaN is approximately 3.2 eV, and the bandgap of hexagonal GaN is approximately 3.4 eV.

[0128] Substrate Nitridation:

[0129] The substrate is exposed to a flowing mixture of ammonia and hydrogen (ratio: 3 / 17) for 360 seconds at a reactor pressure of 100 Torr and a temperature of 960 °C. Using a temperature below or above 960 °C reduces the PL NBE peak intensity and broadens the emission FWHM, as Figure 4 , Figure 5 , Figure 7 , Figure 8 , Figure 10 , Figure 11 , Figure 13 , Figure 14 , Figure 16 , Figure 17 shown.

[0130] Nitriding of the surface can partially remove the oxide from the SiC surface and fill the surface with N atoms when Ga is available, preparing for the formation of GaN.

[0131] In an alternative embodiment, the oxide can be completely removed by a thermal method, but this requires a processing temperature of about 1400 °C.

[0132] Nucleation - Layer Deposition :

[0133] Deposit about 40 nm of GaN using TMG at a flow rate of 93 μmol / min, ammonia at a flow rate of 0.15 slm, at a reactor pressure of 500 Torr and a temperature of 575 °C. The growth rate is about 0.3 nm / s. The selected growth temperature is about 40 - 60 °C higher than the temperature at which the growth rate deviates from a constant value to a lower value, thus entering a state where the ammonia flow determines the growth rate. Figure 18 Shows the relationship between the growth rate of the nucleation layer (NL) and the growth temperature and V / III ratio.

[0134] Due to the low temperature, it is considered that the nucleation layer formation step forms relatively small GaN nuclei.

[0135] It should be noted that, if desired, a buffer layer can be provided between the SiC layer and the GaN layer to manage the thermal expansion mismatch between these layers. This buffer layer can consist of an AlN layer or an Al x Ga 1-x N layer (where the composition changes gradually or is continuously graded) or consist of a combination of AlN / Al x Ga 1-x N layers.

[0136] Recrystallization :

[0137] At a pressure of 100 Torr and an ammonia flow rate of 0.5 slm, a hydrogen flow rate of 20 slm, the temperature is raised to the epitaxial layer growth temperature at a rate of 1 °C / sec. The residence time is 30 seconds. Optionally, a higher ammonia flow rate can be used, which helps to prevent or reduce roughening of the nucleation layer.

[0138] The recrystallization step aims to improve the crystal quality of the nucleation layer by smoothing the layer surface and preventing the formation of surface crystal planes that may lead to wurtzite-like inclusions. In the recrystallization step, N (from NH3) is provided on the surface of the GaN nucleation layer to reduce or avoid decomposition of the GaN nucleation layer at high temperatures.

[0139] Epitaxial - Layer Growth :

[0140] Perform GaN deposition using TMG at a flow rate of 140 μmol / min, ammonia at a flow rate of 0.25 slm, at a reactor pressure of 300 Torr and a temperature of 860 °C, with a growth rate of about 0.5 nm / s. When changing the growth temperature between 845 and 880 °C at a fixed V / III ratio, the surface morphology and PL changes of the observed Nomarski image are shown (see Figures 19 - 27) indicates that the optimal growth temperature is approximately 860 °C; the samples grown at 860 °C show a relatively smooth surface, and the corresponding NBE PL peak is the strongest but also the broadest. At higher growth temperatures, the surface becomes rough, the PL NBE peak becomes significantly narrower, but the intensity of the yellow band (YB) increases. The cited temperature is based on in-situ real-time measurements using an emissivity-corrected pyrometer. The pyrometer can be calibrated with reference to the Al / Si eutectic temperature or a calibration light source such as the AbsoluT system provided by LayTec AG. Those skilled in the MOVPE field should be aware of the calibration procedure for such an emissivity-corrected pyrometer system.

[0141] The change in surface roughness with growth temperature was confirmed by AFM (see Figure 28 ), and the phase purity was determined by the hexagonal-cubic XRD peak ratio, indicating that there are no obvious hexagonal inclusions in the cubic films grown at 860 °C or lower temperatures. At growth temperatures above 860 °C, under the growth conditions reported herein, XRD analysis shows an increase in the contribution of the reflections attributed to the wurtzite lattice, thus indicating that hexagonal inclusions are incorporated into the cubic zinc blende matrix.

[0142] As the growth temperature increases, the PL NBE peak becomes narrower, which may be related to the same trend shown by the FWHM values of the 002 and 004 XRD rocking curves (see Figure 29 ).

[0143] When changing the ratio of the ammonia input flow rate to the Ga precursor TMG input flow rate (V / III ratio, see Figures 30 to 41 ), the changes in surface morphology and PL spectra were observed, and the changes indicate that higher ammonia, as well as the increase in the PL NBE peak intensity and its FWHM, will improve the surface morphology while reducing the intensity of YB.

[0144] When changing the growth pressure from 100 Torr to 300 Torr to 500 Torr ( Figures 42 - 50 ), the observed changes in surface morphology and PL are similar in effect (from the perspective of surface morphology and PL) to a corresponding increase in growth temperature of approximately 15 to 30 °C.

[0145] When increasing the layer thickness, the surface morphology and PL spectra show that the surface becomes rough, but the PL NBE peak intensity increases and the FWHM decreases. Figures 51 - 56 Examples of 600 and 750 nm thick layers are shown in

[0146] N - type Doping:

[0147] When incorporated into the GaN lattice, Si acts as an electron donor, making GaN n-type. We used silane (SiH4) as the Si precursor, diluted to 50 ppm in hydrogen. Other n-type doping sources are also possible, such as disilane (Si2H6), germane (GeH4), digermane (Ge2H6), or oxygen-containing precursors. N-type conductivity was observed in the Si-doped layer, and the existence of high-quality ohmic metal contacts was also demonstrated on Si-doped GaN. For silane flow rates up to 70 sccm, the surface morphology of the layer, as well as the PL NBE peak and yellow luminescence band, were hardly affected (see Figures 57 - 61 ). At higher flow rates, surface pits appeared (see Figure 62 ). The optical properties of Si-doped GaN are shown in Figures 63 - 68 , and the silane flow rate is shown.

[0148] The silane input flow rate is linearly proportional to the Si concentration, where the Si concentration was measured in the bulk using SIMS (see Figure 69 ).

[0149] InGaN / GaN Quantum Wells:

[0150] InGaN deposition was carried out using TMG at a flow rate of 8.2 μmol / min, TMI at a flow rate of 9.7 μmol / min, ammonia at a flow rate of 446 mmol / min, at a reactor pressure of 300 Torr and a temperature of 700 to 800 °C, with a growth rate of approximately 0.8 nm / min. For GaN barrier growth, the same conditions were used except for the TMI flow.

[0151] For a nominal quantum well (QW) width of 2 nm, the growth temperature of InGaN that achieves a PL peak of 450 nm is very similar to the growth temperature used for a standard wurtzite c-plane structure with the same well width. However, increasing the QW width to 10 nm allows the emission wavelength to be extended to 540 nm at the same growth temperature.

[0152] Figure 70 An optical micrograph of a typical green QW structure is shown, while Figure 71 and Figure 72 show the PL spectra at high and low laser excitation powers, respectively. As the excitation density increases, the strong, broad InGaN peak at 521 nm shifts to 540 nm. The structure includes 5 quantum wells with a nominal thickness of 10 nm formed at 300 Torr. It can be observed from Figure 84 that for a constant laser excitation power, when saturation begins, the peak emission wavelength increases approximately linearly as the well width increases up to a thickness of 8 nm. As the well width increases, the PL peak intensity shows a small decrease, but due to the green gap, it is much smaller than expected in a wurtzite GaN Q well.

[0153] Texture Analysis

[0154] Polymorph Identification

[0155] The phases and orientations present in GaN films can be identified by XRD texture analysis, where different selection rules are used for the reflections that occur in the zinc blende and wurtzite phases. For certain diffraction angles 2θ, the reflections of the two phases overlap, but there are other particularly suitable diffraction angles where the reflections of only one of the two phases can be observed at a time. For example, for 2θ of approximately 34.5°, the wurtzite 0002 and zinc blende 111 reflections that both occur are not suitable, nor are the wurtzite 11-20 and zinc blende 220 reflections (2θ of approximately 57.8°). Herres et al. (1999) proposed using the cubic 200 (2θ of approximately 40.0°) and hexagonal 10-12 (2θ of approximately 48.1°) reflections for films that are mainly (111) zb and (0001) wz orientations and demonstrated that reasonable results could be obtained. However, for cubic films that are mainly (001) oriented, different zinc blende reflections should be preferred because the {100} reflections from the crystal planes are usually very weak and are superimposed by surface scattering effects, making it difficult to determine the in-plane relationships of the films. There are several other suitable reflection combinations that we can use for texture analysis, such as the 113 zb reflection (2θ of approximately 69.0°) and the 1-103 wz reflection (2θ of approximately 63.4°) because they are well separated in reciprocal space and the characteristic diffraction patterns are relatively easy to interpret.

[0156] Texture Maps

[0157] For XRD texture analysis, the angular distribution of selected reflections in reciprocal space is measured by plotting the surface of a hemisphere with a radius given by a specific Bragg condition (see Figure 2 ). For this purpose, the sample is rotated about its surface normal (φ scan) and is gradually tilted (χ step) towards the beam path plane after each scan. The measured intensities are shown as a polar projection (radius = χ) or a stereographic projection (radius = tan(χ / 2)). In the so-called pole figure, the center represents the direction of the surface normal, while the poles at the edge of the figure (χ = 90°) represent the directions within the surface plane.

[0158] To determine the phase purity, main orientation, and crystallographic relationships between different textures of nitride films, at least two texture maps - one for each phase - must be measured.

[0159] Figure 73 、 Figure 74 andFigure 75 show the pole figures collected at 2θ = 34.5°, 2θ = 68.9°, and 2θ = 63.4° for the GaN epilayer grown on a 3C-SiC / Si template under non-optimized conditions. In Figure 73 , four strong reflections at χ of approximately 57° can be clearly seen. From the four-fold symmetry, it can be seen that these reflections most likely represent the 111 zb reflection of the zinc blende phase. However, this result alone does not prove the absence of the wurtzite phase. Even if there is no cubic phase at all, the reflections measured in Figure 73 can represent the 0002 wz peak, which originates from four different twins of the hexagonal wurtzite component grown on the {111} plane of 3C-SiC. In the more likely case, a mixture of two phases has been grown, which will contribute to the measurement of reflections from both phases.

[0160] For further examination, we measured the distribution of the zinc blende 113 zb reflection ( Figure 74 ) and the wurtzite 1-103 wz reflection ( Figure 75 ), which do not overlap with the reflections of the other phase. The pattern of the 113 zb reflection ( Figure 74 ) shows four-fold symmetry, with three reflections (2 at χ of approximately 72° and 1 at χ of approximately 25°) arranged around a common 111 zb pole. This can clearly confirm that the main orientation of the cubic film is (001), corresponding to the SiC / Si template orientation, which is indicated by the weak 004Si reflection at the center of the pole figure. In Figure 75 , the {1-103} wz plane of the wurtzite phase causes many reflections, which form a distorted hexagonal pattern around their common center 0001 pole (invisible). The results show that the hexagonal wurtzite phase also exists in the GaN film, but it is the minority phase, with much weaker reflection intensity compared to the cubic zinc blende reflections.

[0161] The phase purity of the GaN mixture can be estimated by integrating the reflection intensities of each phase and determining the ratio of these values. As suggested by Herres et al. (1999), this may be sufficient to provide rapid feedback for optimizing crystal growth activities. However, for a more accurate quantification of the volume fractions of the zinc blende and wurtzite GaN phases, additional corrections considering the different scattering efficiencies of the two structures and their crystal planes are required. Since the zinc blende and wurtzite phases have the same absorption coefficient, only the different structure factors F hkl , unit cell volumes V UC, and a geometric correction known as the Lorentz polarization (LP) factor is sufficient. Ignoring smaller correction factors (such as absorption correction and temperature correction), the integrated intensity of a single reflection hkl is proportional to the volume amount V of the material phase is given by the following formula:

[0162]

[0163] where the Lorentz polarization factor is

[0164]

[0165] The structure amplitude is

[0166] [Cullity(1956)],

[0167] The ψ factor depends on the diffraction angle of the powder sample (ψ = sin -1 (θ), while it is a constant for single crystal [Reynolds(1986)]. The coordinates x j , y j , z j are the positions of each atom in the unit cell and are listed in Table 1. The atomic scattering factors f of Ga and N j are proportional to the number of electrons of each atom and also have a complex dependence on the diffraction angle and wavelength. The relevant details are described in the literature: [Cullity(1956); International Tables for Crystallography, Volume C(2004); Waasmaier and Kirfel(1995)] and the online databases [Cromer-Mann coefficients:

[0168] http: / / www.ruppweb.org / Xray / comp / scatfac.htm and

[0169] http: / / www.ruppweb.org / new_comp / scattering_factors.htm and DABAX library (ESRF)

[0170] http: / / tx.technion.ac.il / ~katrin / f0_CromerMann.txt]. With these corrections, Figure 74 and Figure 75 the {113} in zb and {1-103} wzThe reflection reveals that the volume fraction of wurtzite GaN in the sample is approximately 69 vol%. The detection accuracy for the data collection here is a few volume percentages and is mainly affected by the large variations in the integrated intensities of the reflections in different crystal directions.

[0171] Therefore, when performing XRD characterization on the group-III nitride layer, the volume V zb of the wurtzite-structured group-III nitride and the volume V wz of the zinc-blende-structured group-III nitride have a relative volume ratio of:

[0172]

[0173] where V wz is evaluated based on the reflection of the wurtzite-structured group-III nitride 1-103 (i.e., 1-13 in the conventional Miller index notation), and V zb is evaluated based on the reflection of the zinc-blende-structured group-III nitride 113.

[0174] Rearranging the equations already given:

[0175]

[0176] where:

[0177] V uczb is the volume of the unit cell of the zinc-blende-structured group-III nitride,

[0178] V ucwz is the volume of the unit cell of the wurtzite-structured group-III nitride,

[0179] F 113 is the structure amplitude of the 113 reflection of the zinc-blende-structured group-III nitride,

[0180] F 1-13 is the structure amplitude of the 1-103 reflection of the wurtzite-structured group-III nitride,

[0181] I 113 is the integrated intensity of the 113 reflection of the zinc-blende-structured group-III nitride,

[0182] I 1-13 is the integrated intensity of the 1-103 reflection of the wurtzite-structured group-III nitride.

[0183] The crystallographic relationship between the zinc-blende and wurtzite phases can be obtained by combining the Figures 63 - 65 pole figures, and this relationship is (111) zb ||(0001) wz , where two of the four unequal {111} planes in zinc-blende gallium nitride are [11-2]zb ||[-1010] wz and [-110] zb ||[1-210] wz 。This unit cell arrangement is shown in Figure 76 and this is not surprising since the close-packed planes of each structure are parallel to each other and differ only in the stacking sequence. Other families have observed similar arrangements of the two phases by XRD [Qu et al (2001); Tsuchiya et al (1998)] and by transmission electron microscopy measurements [Trampert et al (1997)], and it has also been found that the zinc blende and wurtzite phases are arranged in alternating zb-wz-lamellae [Wu et al (1997)]. As we pointed out in the example in Figure 75 , it is not necessarily formed with equal probability on each of the four non-equivalent {111} zb planes. The (0001) wurtzite GaN is formed.

[0184] Table 1: Atomic positions in ideal wurtzite and zinc blende unit cells

[0185]

[0186] Samples are grown under optimal conditions to maximize the zinc blende phase, and the X-ray intensity at the expected position of the wurtzite phase reflection can be neglected, i.e., only slightly above the background noise level. In these cases, it is very likely that the signal is not caused by the diffraction of hexagonal wurtzite inclusions in the epitaxial layer, but rather originates from diffuse scattering on plane defects (such as stacking faults). This can be illustrated by measuring two-dimensional reciprocal space maps, as described below.

[0187] In addition to wurtzite inclusions, the zinc blende GaN film may also contain twinned zinc blende regions. Similar to stacking faults, these are introduced by stacking errors on a single (111) plane, but in contrast to stacking faults, the zinc blende matrix continues with a different stacking sequence...AaCcBbAaCcBb... Relative to the surrounding GaN matrix, the zinc blende twin is tilted by approximately 70.4° about the [1-10] axis, and thus, the relationship between the twin and the matrix is (111) twin ||(115) matrix [Tsuchiya et al (1998)]. Such zinc blende twins and possibly wurtzite-like twinned materials with similar relationships are circled at χ of approximately 15° in Figure 73 and may cause weak 111 reflections at χ of approximately 83° (outside the range in Figure 73 ). Their volume fraction is in the low percentage range.

[0188] Reciprocal - Space Maps for Texture Analysis

[0189] Appropriate two-dimensional reciprocal space maps (RSMs) of wurtzite and zinc blende back-reflections—combined with ω-2θ scans and a stepwise change in the ω angle after each scan—can be used to analyze the phase purity and several other structural properties of GaN samples. Appropriate reflections include 002 ZB and 10-11 WZ , as Figure 77 and Figure 78 shown, for two different samples with and without hexagonal inclusions, respectively. In both reciprocal space maps, the 002 reflections of high-intensity zinc blende GaN and 3C-SiC are clearly visible. The low-intensity streaks along <111> passing through the 002 reflection are caused by diffuse scattering due to {111} stacking faults in the structure, where the diffracted X-rays undergo an additional phase shift between the two sides of the stacking fault. Stacking faults may also cause a slight displacement of the GaN reflections from their ideal positions. Another feature across the 3C-SiC reflections on the 2θ arc is the detector streak (DS), which is caused by the instrumental function of the diffractometer. The streaks intersecting the 002 zb GaN reflection perpendicular to the surface are the so-called (X-ray) crystal truncation rods (CTRs), whose shape is affected by the surface structure and is consistent with atomic force microscopy observations. The partially visible Bragg rings (2θ ≈ 35.6°) in the RSM originate from polycrystalline SiC deposited on the etched grating of the 3C-SiC / Si template and are unrelated to both the SiC mesa region and the GaN epitaxial layer. Since the GaN epitaxial layer is much thinner than the SiC template, the similar Bragg rings generated by the GaN grown on the etched grating are much weaker and usually not visible. The presence of wurtzite-like inclusions ( Figure 67 ) in the zinc blende GaN film results in two additional 10-11 wz reflections for the wurtzite GaN phase, such reflections being absent in samples without these inclusions ( Figure 78 ). Since the stacking fault streaks overlap with the wurtzite reflections, these streaks are easily misinterpreted as signals from a small amount of hexagonal inclusions in the texture map.

[0190] Compared with the texture maps described above, reciprocal space mapping can be performed much faster even with a CCD detector and even at high integration times. This increases the signal-to-noise ratio, thus allowing quantification of the lower proportion of wurtzite-type GaN inclusions. However, this method assumes fixed epitaxial relationships and does not provide additional information about the presence of cubic twins.

[0191] Mosaicity Analysis

[0192] Due to the lack of a suitable homogeneous substrate, cubic zinc blende GaN-based nitrides are usually heteroepitaxially grown on a hetero-cubic substrate, such as GaAs [As et al (2000); Yang et al (1996); Shen et al (2003); Qu et al (2001);

[0193] Tsuchiya et al (1998)], SiC [Wu et al (1997); Chichibu et al (2003)], Si [Lei et al (1991)] and various other cubic materials (such as GaP [Cheng et al (1995)], MgO [Compeán García et al (2015)]). The lattice mismatch between different materials leads to a high mosaicity and the formation of defects at the grain boundaries of the epitaxial layer. Generally, mosaicity should be avoided because it has a negative impact on the physical properties of the sample, for example, generating a high resistance at the grain boundaries [Fujii et al (2010)]. Therefore, it is preferably to quantify the mosaicity to optimize crystal growth.

[0194] In a simplified model derived from powder diffraction method, the thin film consists of mosaic blocks (grains) with finite sizes and directions that are slightly different from each other. The spread of size, tilt, and twist, as well as microstrain and compositional inhomogeneity (for alloys) cause the broadening of X-ray reflections in reciprocal space. Mosaic tilt causes the reflection angle perpendicular to the surface to expand, while twist causes the azimuthal expansion around the surface normal. Therefore, for mosaic tilt and twist, the absolute broadening ΔG hkl in reciprocal space increases linearly with the magnitude of the scattering vector |G hkl |. The finite lateral size of the mosaic grains causes broadening parallel to the interface, which is inversely proportional to the average real-space size L and independent of the magnitude of the scattering vector (ΔG hkl = 2π / L). The effects of tilt, twist, and finite grain size are convoluted with the spread of the reflection hkl, and the spread of the reflection hkl is measured by the skew-symmetry ω-scan as follows:

[0195]

[0196] [Lee et al (2005)]. Here β represents the integral width, and the exponent n takes values between 1 and 2, depending on the contributions of the Gaussian η and Lorentzian (1 - η) to the Pseudo-Voigt fit (n = 1 + η 2 ) (see Appendix of Srikant et al (1997)).

[0197] Then, the measured peak broadening is the combination of the mosaic broadening of the sample and the instrumental function (without sample). As long as the latter is much narrower than the mosaic broadening, it can be neglected. Experimentally, it is possible to measure the ω-scans of a series of symmetric reflections 00l of different orders and plot β in a modified Williamson-Hall plot (not shown). n ·G n versus G n to distinguish the peak broadening effects due to the lateral size and tilt. The slope of the line is related to the tilt component (β tilt n ), and the vertical offset is related to the average grain size ((2π / L) n ). Unfortunately, the commonly used Cu-Kα radiation can only obtain the symmetric 002 and 004 wurtzite GaN reflections, which greatly limits the accuracy, especially the accuracy of determining the finite size. Figure 79 shows the linear behavior of β hkl n in a traditional Williamson-Hall plot, for which the Lorentzian broadening (n = 1) is usually adopted, although the Lorentzian curve usually does not fit the measured curve very well. The value of the finite size L is much larger compared to the Gaussian fitting curve (n = 2), as pointed out by Lee et al. (2005). Since the X-ray intensity curve can be empirically described as the convolution of a Gaussian function and a Lorentzian function, the actual lateral finite size is within these two limiting values, depending on the parts of these two curves. Usually, the Gaussian and Lorentzian ratios can be obtained from curve fitting, but they may vary for a series of different reflections. However, since the Lorentzian part in such fittings is usually small, a relatively good estimate of the mosaic block size can be obtained from the pure Gaussian fitting.

[0198] The azimuthal angles scattered around the surface normal due to the mosaic distortion can be determined from the off-axis reflections with large polar angles χ measured in a helical symmetric geometry. Ideally, one of the in-plane reflections (χ ≈ 90°) would be used, but these reflections usually only exhibit very low intensities and are usually difficult to measure. For a (001)-oriented wurtzite GaN film, the 331 reflection (χ ≈ 76.7°) can be better used. Alternatively, the integral widths of a series of different off-axis reflections extracted from the asymmetric ω-scans can be extrapolated using Equation (9) to determine the distortion component.

[0199] Figure 80 shows such an extrapolation, where we transform Equation (9) to β hkl n ·|G hkl | n , by using Figure 79The slope and the finite size value of the Williamson-Hall plot, fitting this function (for n = 2) to the measured peak broadening of the best cubic GaN samples. The circles mark the measured reflections, and the contour shows the extrapolated peak width in the reciprocal space β hkl ·|G hkl | as a function of the polar angle χ and the magnitude of the scattering vector |G hkl |. The contours represent constant peak widths. The curve at χ = 0° has been shown previously ( Figure 79 ), and it is only affected by the tilt and finite size of the mosaic blocks. As the polar angle χ increases, the broadening gradually increases, showing that the mosaic distortion of 0.864° (χ = 90°) is slightly higher than the tilt of 0.755° (χ = 0°). Since the tilt and distortion are very similar, this trend is not obvious, but for larger scattering vectors |G hkl |, this trend becomes more obvious as the contribution of the finite size to the peak broadening decreases.

[0200] Defect Density

[0201] Typically, mosaic tilt and distortion are considered to be related to the formation of threading dislocations at the grain boundaries of the thin film. Therefore, by following different mosaic tilt models discussed in the literature, XRD peak broadening is sometimes used to estimate the defect density in the thin film. According to these models, the threading dislocation density D TD in a well-oriented mosaic film is proportional to β tilt / twist :

[0202]

[0203] [Fewster(1989)], but in a film with a large misorientation, the grain orientations are random and the Burgers vectors and line vectors are strictly distributed, and its threading dislocation density is proportional to β 2 tilt / twist :

[0204]

[0205] [Dunn and Koch(1957)]. The parameter L here is the average lateral finite size of the grains, and b TD represents the value of the Burgers vector of the dislocation in the case of complete dislocation in zinc blende GaN, which is

[0206] In contrast to wurtzite GaN materials, where threading dislocation line vectors mainly propagate along the

[0001] c direction, the threading dislocations in zinc blende GaN extend along multiple <110> directions. Therefore, the above equation cannot distinguish between edge, mixed, or screw dislocations in zinc blende GaN. However, it is well known that the main type of threading dislocation in the zinc blende structure is the 60° perfect dislocation [Blumenau et al (2000)].

[0207] In an in-depth comparative study of using XRD and transmission electron microscopy (TEM) to estimate the defect density in wurtzite GaN films, Metzger et al (1998) found that it could well match the random distribution model (Equation xy), even though the assumptions of this model were not satisfied at all in the case of epitaxial films with an orientation. Contrary to expectations, the model for the oriented mosaic films showed that the threading dislocation density was one order of magnitude lower than the value estimated by TEM. Lee et al (2005) reached a similar conclusion and pointed out that it was common for there to be a large difference in the dislocation density measured between TEM and XRD. Generally, when using the distortion component, XRD seems to slightly overestimate the threading dislocation density, while when using the broadening due to tilt, XRD underestimates this density [Lee et al (2005)]. In addition, it should be noted that for very thin films, XRD also samples the tilt related to the misfit dislocations at the GaN / SiC interface. If the Burgers vectors of the dislocations are randomly oriented, then as the film thickness increases, the associated strain fields will tend to cancel each other out. However, if the Burgers vectors are not random, the tilt will persist. The discussion shows that even for the measurements of the more widely studied wurtzite GaN, there are still some limitations in the current understanding, and the defect density estimated by XRD needs to be handled with caution. This is especially true when comparing samples with different layer thicknesses.

[0208] Effect of Layer Thickness

[0209] Generally, the intensity distribution of X-ray reflection is not constant but decreases with the increase of film thickness, the full width at half maximum (FWHM) of the 002 reflection in the ω scan, as Figure 81 shown. It is also obvious that for zinc blende GaN (lattice constant GaN: ) grown on low lattice mismatch substrates such as 3C-SiC (lattice constant of SiC: with a compressive rate of 3.4%) and MgO (lattice constant of MgO: ) and the zinc blende GaN grown on more mismatched Si (lattice constant of Si: with a tensile rate of -17.0%) or GaAs (lattice constant of GaAs: Therefore, it has a lower mosaicity compared to cubic GaN films with similar thickness grown at a tensile rate of -20.3%. In addition, Figure 81 shows that wurtzite GaN grown by MOVPE (our data) can be comparable to cubic GaN films grown by MBE in the prior art [Kemper et al (2015); Martinez-Guerrero et al (2002)]. The decrease in the reflection spreading intensity with the increase in film thickness is generally related to the overall reduction in the defect density of thicker epitaxial films and the improvement in material quality. Transmission electron microscopy studies show that in the case of the formation of complete edge dislocations or partial penetration dislocations, the reaction between paired stacking faults causes the stacking fault density to decrease significantly with the increase in layer thickness. Martinez-Guerrero et al (2002) observed that in the first 500 nm of wurtzite GaN growth, the stacking fault density decreased from 5×10 6 cm -2 to 3×10 5 cm -2 , almost showing an exponential decay. In our MOVPE-grown cubic GaN thin films, TEM measurements show that the stacking fault density decreases from 10 7 cm -2 directly at the template interface to 3×10 4 cm -2 close to the surface of the 1200 nm thick film. However, for basal plane stacking faults in wurtzite GaN and stacking faults in face-centered cubic (fcc) nanocrystals [Dupraz et al (2015)], the stacking fault density mainly affects the shape and intensity distribution along the SF stripes in reciprocal space (as shown by Barchuk et al), but the ω-scan of the symmetric 002 ZB reflection hardly overlaps with the stacking fault distribution. Therefore, as Figure 71 shows, the narrowing of the observed peaks with the increase in layer thickness cannot be directly related to the reduction in stacking faults.

[0210] Several reports in the literature (see the references in Figure 81 ) indicate that Figure 71 this trend is due to the reaction of penetration dislocations, and the penetration dislocation density decreases with the increase in film thickness, but the TEM evidence is insufficient [As (2010); Kemper (2015); Rüsing (2016); Lischka (1997)]. Theoretical models predict that the penetration dislocation density is inversely proportional to the film thickness t [Ayers (1995)]. Combining these models with the reflection broadening caused by mosaicity, it can be found that the reduction amplitude of the intensity distribution is t -1 or t -1 / 2, which depends on whether it is an oriented film (Equation (10)) or a powder sample (Equation (11)). As can be seen from the dashed line in Figure 81 , the experimental data do not follow the predicted trend. Instead, the observed decay is much weaker, following approximately t -1 / 3 dependency. This can be explained by the fact that when stacking faults react with each other, new threading dislocations can be generated, which, to our knowledge, is not taken into account in the current model. In addition, it should be taken into account that the model predicts a decrease in the threading dislocation density after a certain thickness, and XRD is an integral method that provides a weighted average over the entire layer thickness. It should also be taken into account that as the number of scattering atoms increases, the width of the X-ray reflection naturally decreases with increasing layer thickness. All these make it difficult to compare the material quality of samples with different thicknesses.

[0211] Material Parameters for Strain Analysis

[0212] The lattice parameters of zinc blende type III nitrides cannot be well determined experimentally because such films exhibit stacking disorder, undoubtedly a high density of line defects and wurtzite inclusions, resulting in local strain variations and relatively broad reflections. In addition, most X-ray diffraction experiments on mainly zinc blende GaN films have focused on phase purity analysis rather than high-resolution lattice parameter measurements.

[0213] We measured the lattice parameters of zinc blende GaN using high-resolution 2θ-ω-scans, scanning a total of 8 in-plane and off-axis reflections, and the values obtained by least-squares fitting are in very good agreement with the experimental data of Novikov et al. (2010) and can be used as reference data for strain analysis of zinc blende GaN films.

[0214] Table 2 – Lattice parameters and elastic constants of wurtzite and zinc blende GaN, InN, and AlN

[0215]

[0216] * 1 :

[0217] * 2 :

[0218] * 3 :

[0219] * 4 : Experiment

[0220] * 5 : This work (experiment)

[0221] * 6 : Suggestions by Vurgaftman and Meyer (2003)

[0222] However, to our knowledge, the accurate lattice parameters of wurtzite InN and AlN determined experimentally are not mentioned in the literature. Therefore, in these cases, it is necessary to derive these values from the well-established wurtzite lattice parameters a wz and c wz However, in wurtzite-like group-III nitrides, a strong internal electric field causes the unit cell to deform from its ideal shape, and its c wz / a wz ratio is In fact, c wz is usually smaller than the ideal case, while a wz is slightly larger than the ideal case. Therefore, as can be seen from the values in Table 2, the estimated zinc-blende parameters of nominally unstrained group-III nitrides can vary significantly. Since a wz is less affected by the wurtzite unit cell distortion than c wz , this parameter can provide a reasonable value for the natural lattice constant of the zinc-blende phase a zb . Alternatively, lattice parameters derived from the unit cell volume can be used. It is speculated that the natural unstrained lattice constants of zinc-blende nitrides lie between these theoretical values, and this assumption is in good agreement with the experimental data known so far.

[0223] Table 2 also contains the elastic constants C 11 and C 12 of zinc-blende group-III nitrides described in Vurgaftman and Meyer (2003), which can be used for stress and strain calculations.

[0224] Strain

[0225] During the growth of thin films on a hetero-substrate and during the growth of heterostructures of alloys with different compositions, the thin films are subject to varying stresses, which usually cause elastic deformation of the lattice. Such lattice strain has a significant impact on the physical properties and performance of semiconductor devices. Therefore, it is very important to understand and monitor these strains during device development. In the following sections, we will discuss different strain sources and describe how to measure the strain in zinc-blende GaN thin films.

[0226] Lattice - Mismatch Strain

[0227] In an epitaxial film, when two lattice sizes are forced to match each other, the lattice mismatch between the film and the underlying template generates biaxial in-plane strain. Three different states are commonly used to describe the film deformation. When the lattice of the film matches the size of the template lattice at the common interface, the film will be fully deformed, while when the film lattice is undeformed and has its natural size, the film will be fully relaxed. The state between the two extremes is called partially relaxed.

[0228] In reciprocal space, the lattice mismatch strain causes the reciprocal lattice points (RLPs) of the GaN film to be displaced from their expected positions relative to the RLPs of the substrate. The relative distance between the layer and buffer peaks can be measured in several individual ω-2θ scans, or more commonly by collecting a reciprocal space map in an asymmetric geometry. The latter usually provides a better overview of the relationship between the X-ray reflections of different layers, but for lattice mismatch strain assessment, sample tilting needs to be corrected by a secondary scan. In addition, it should be taken into account that the layer used as a reference may also be affected by the substrate, which reduces the accuracy of this method. Ideally, the substrate peak should be used as a reference, but for systems with large mismatches, there may be a large gap in reciprocal space.

[0229] The strain of the film in a certain direction is as follows:

[0230]

[0231] where a0 is the natural lattice constant, and a i is the measured constant in the same direction. Since the material quality is usually relatively low in zinc-blende nitride materials, there is currently no accurate reference value for the natural lattice parameter in the literature as described in the previous section. For GaN, the experimentally determined values provided in Table 2 can be used. For other group-III nitrides, we recommend using the values derived from the wurtzite a parameter or the wurtzite unit cell volume (see Table 2) because the wurtzite lattice parameters are well known.

[0232] Assuming that the film is stress-free in the growth direction (usually labeled as z), the strain of the (001)-oriented film in the growth direction is given according to Hooke's law as:

[0233]

[0234] where ε x and ε y are the strains in two in-plane directions, C 11 and C 12 are the elastic constants of the material (see Table 2) [Dunstan (1997)]. For isotropic plane strain (ε x = ε y) The above equation can be further simplified. The strain relationships in other directions are different from the above equation and are published in other literature, such as Dunstan (1997).

[0235] Thermal - Mismatch Strain and Growth - Induced Strain

[0236] The small strain in the epitaxial film originates from the thermal mismatch between the substrate used and the epitaxial layer, or is formed in the early stage of growth. It is usually much smaller than the strain due to lattice mismatch, but it may be larger compared to the residual mismatch strain in the partially relaxed film.

[0237] Since GaN has a larger coefficient of thermal expansion than SiC and Si [Wahab et al (1994); La Via (2012); Okada and Tokumaru (1984)], after cooling down from the growth temperature, the residual thermal strain causes tensile stress in the GaN film at the interface with the substrate. For typical zinc-blende GaN with a growth temperature between 700 °C and 1000 °C, when using an Si substrate, the theoretical thermal strain is between 1.1×10 -3 and 1.6×10 -3 .

[0238] The growth-induced strain is generated due to the coalescence of islands during the nucleation process on the substrate in the early stage of growth. Its magnitude is given by the minimum gap Δ between two islands and the average size of the islands in a specific in-plane direction:

[0239]

[0240] [Hoffman (1976)].

[0241] As described above, the relative lattice parameter measurement results are not sufficient to accurately determine such a small strain because the resolution is usually low and the substrate itself may also be affected by strain. Instead, to analyze very small strains, a large set of high-resolution 2θ-ω scans of different reflections is required to make absolute measurements of the lattice parameters. Then, the measured interplanar spacing d j is matched with the interplanar spacing of the model crystal:

[0242]

[0243] To improve the accuracy of this method, sometimes it is necessary to use a weighting coefficient W j , such as 2θ j / Δ(2θ j ) [Roder etal (2006)] or d j -2 / Δ(d j-2 ) = 0.5·tan(θ j ) / Δ(2θ j )(this operation) is used to account for the inaccuracy of the measured value Δ(2θ j ).

[0244] Generally, a suitable coordinate system must be chosen that describes the geometry of the problem better than the natural lattice. The following example illustrates this. Table 3 lists the 2θ values of different reflections, which were measured from a wurtzite GaN film grown on a 3C-SiC / Si template that was obliquely cut by 4° in the

[110] direction. The Bragg angles of all reflections tilted along the oblique cut direction hhl are much smaller than those of similar reflections tilted away from the oblique cut direction h - hl, indicating that the lattice sizes in these two directions are different. Therefore, the natural lattice is slightly sheared within the growth plane. It can be simplified by using a new coordinate system x', y', z' (where x' (y') is parallel to (perpendicular to) the sample oblique cut and z' points in the growth direction). It should be noted that by transforming this coordinate, the new unit cell is larger than the unit cell in the natural lattice Using this method, the least squares fit (see above) together with Bragg's law gives the unit cell size in the new coordinate system The anisotropy of the in-plane strain is ε x ' = (3.65 ± 0.11) × 10 -3 , ε y ' = (1.92 ± 0.09) × 10 -3 .

[0245] Table 3 – Reflections of the Natural Coordinate System (hkl) and Rotated Coordinate System (h'k'l') of the Zinc - Blende GaN Sample, where 2θ and Δ(2θ) are from High - Resolution 2θω Scans

[0246]

[0247] It is well known that substrate oblique cutting can cause strain relaxation in epitaxial films due to the alignment of threading dislocations [Young etal(2010); Chen et al(2007)], but the strain in the wurtzite GaN layer in the oblique cut direction (ε x ') is less than the strain in the perpendicular direction (ε y '). Since we observed the opposite situation, the relaxation mechanism for this sample can be excluded. Instead, the results suggest that the strain anisotropy may be due to the coalescence of islands with different sizes in two in-plane directions, which can be observed in atomic force microscopy images.

[0248] Wafer Curvature Analysis

[0249] In heteroepitaxial thin films, stress above a certain level can be relieved by forming defects, or tensile surface stress can be relieved by forming cracks. Additionally, the stress in the film can be reduced by bending the entire sample. This is typically the case in thick, moderately stressed epitaxial layers (such as templates and buffer layers). Thermal strain can also cause significant wafer bending. This is especially a problem for large-area templates up to 8” in diameter, where even small bending can lead to significant deviations in uniformity during growth and processing. Therefore, it is very important to control and manage strain and wafer bending.

[0250] Wafer curvature can be determined by XRD by measuring the incident beam angle ω of symmetric reflections at different positions of the sample x j on the wafer. In a bent sample, the lattice planes also bend with the bending of the wafer. Therefore, it is necessary to correct the incident angle for different positions along the wafer diameter. Then, the wafer curvature and bending radius R are obtained from the relative changes in ω j and x j : j [Inaba(2014)]. Since the reflections of the zinc blende GaN epitaxial layer are usually relatively broad, it is more appropriate to use the narrower symmetric reflections of the underlying template. By using a larger range of measurement positions and using a beam mask to reduce the irradiated area on the sample surface, the resolution of the measurement can be further improved.

[0251]

[0252] Figures 21(a) and 21(b) show the measured curvature of a 4” 3C-SiC / Si template, where the 002 SiC reflection was used. The curvature of the wafer along the measurement direction can be easily determined graphically by linear interpolation of the measured incident beam angles. A positive (negative) slope corresponds to a concave (convex) shape of the wafer. Figure 82 and Figure 83 Figures 22(a) and 22(b) give an example of a convex bending of -51.5 km Figure 82 and Figure 83 (R = -19.4 m, respectively). However, it should be noted that this technique measures the curvature of the substrate plane. If the substrate already contains a high density of dislocations or a grain structure, the plane of the substrate may already be bent before layer growth, so the measured curvature may not accurately reflect the residual stress in the wafer. Therefore, there may also be a difference between the bending measured by X-ray and the bending measured by optical techniques. We found that for the high-quality templates used in these studies, the difference between the curvature measured by X-ray diffraction and the curvature measured by optical techniques is small and can be neglected. -1 Figures 23(a) and 23(b)

[0253] In - Depth Study

[0254] Based on the experimental work reported above, the influence of reaction pressure (and other parameters and conditions) on the growth of cubic zinc blende GaN films was further investigated.

[0255] In general, cubic zinc blende GaN thin films were grown by metalorganic vapor phase epitaxy (MOVPE) on 3C-SiC / Si(001) templates and characterized using Nomarski optical microscopy and X-ray diffraction. Specifically, at a reaction pressure of 100 Torr, the surface morphology and material quality were evaluated as a function of the thickness of the low-temperature nucleation layer (3 - 44 nm in these experiments), the epitaxial growth temperature (850 to 910 °C in these experiments), and the V / III ratio (15 to 1200 in these experiments). The main difference from the earlier results reported above is that the reaction pressure was decreased from 300 Torr to 100 Torr in this case. Under these particular conditions, a window of particularly suitable MOVPE growth conditions was determined to be: temperature between 850 and 890 °C, V / III ratio between 38 and 150, resulting in relatively smooth zinc blende GaN films with wurtzite impurity content less than 1%.

[0256] The influence of decreasing the epitaxial layer reaction pressure from 300 Torr to 100 Torr on the zinc blende phase purity is shown in Figure 111 and Figure 112 These figures show that for the 100 Torr data set, the wurtzite fraction is equally low or lower under similar temperature and V / III ratio conditions. Thus, under the reduced reaction pressure conditions, the MOVPE growth window was significantly broadened to obtain good zinc blende GaN material quality.

[0257] All samples were grown in a Thomas Swan 6×2” close-coupled showerhead MOVPE reactor on 3-SiC / Si templates. The SiC templates consisted of a ~3 μm thick 3C-SiC layer on a Si(001) substrate with a thickness between 0.75 mm and 1 mm and an orientation difference of 4° towards

[110] . For GaN growth, trimethylgallium (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 liters per minute (slm). The growth process included a high-temperature thermal annealing of the substrate, followed by a low-temperature nucleation layer deposition, and finally a proper growth of the epitaxial layer at high temperature. The temperature was the temperature recorded by the Laytec EpiTT in-situ optical monitoring system, calibrated for an Al / Si eutectic wafer. The thermal annealing step of the template was carried out at 960 °C in a mixture of hydrogen and 3 slm ammonia. For sample groups A and B (see below), the GaN nucleation layer (NL) was grown to a thickness of 44 nm at 600 °C, 500 Torr and a V / III ratio of 720. The growth pressure of the epitaxial layer was kept at 100 Torr, while the thickness of the epitaxial layer was kept constant at 300 nm. Two sample groups were prepared, where the variables were the epitaxial layer growth temperature (sample group A) and the V / III ratio (sample group B). For sample group A, with a constant V / III ratio of 76 in the gas phase, the epitaxial layer growth temperature was varied between 850 and 910 °C. For sample group B, during the growth of the epitaxial layer at 880 °C, the V / III ratio was varied between 15 and 1200 by changing the ammonia flow rate with a constant TMG flow rate of 145 μmol / min. All GaN epitaxial layers were doped with Si using silane (50 ppm SiH4 in H2) to a nominal concentration of 10 18 cm -3 intermediate nominal concentration. Finally, a third sample group (sample group C) consisted of a 300 nm thick GaN epitaxial layer grown at 875 °C, 100 Torr and a V / III condition of 76 on a low-temperature nucleation layer with a thickness varying between 3 and 44 nm.

[0258] Using a Cu-Ka1 light source equipped XRD phase analysis was performed on a PANalytical Empyrean diffractometer with a 2-bounce hybrid monochromator, a 1 / 4° slit, an Eulerian cradle, and a PIXcel solid-state area detector. Reciprocal space maps (RSMs) were measured around the 113zb-GaN and 1-103wz-GaN reflections parallel and perpendicular to the substrate's cleavage direction. The intensity distribution along the SF streak between the 113 and 1-103 reflections was extracted from the RSM and then fitted to up to three Pseudo-Voigt functions: the zinc blende and wurtzite phases, and a third ill-defined defect phase possibly related to stacking faults. The integrated intensity of the fitted profiles was used to quantify the wurtzite fraction of the GaN epilayer. This work did not quantify the residual peak intensity attributed to stacking faults. The 002 peak broadening was measured with a Philips X'Pert diffractometer equipped with an asymmetric four-crystal Bartels monochromator 5×5mm 2 A cross-slit collimator, an Eulerian cradle, and a gas proportional detector, with no other secondary optical elements. The intensity distribution of the open detector ω-scan was fitted to a Pseudo-Voigt function.

[0259] Sample group A – temperature series

[0260] The first sample group to be discussed (sample group A) consisted of six samples in which the growth temperature of the epilayer varied between 850 and 910 °C in intervals of 10 to 15 °C. In this series of experiments, the V / III ratio was kept constant at 76, which represents an intermediate value within the range of values explored in sample group B (see below). As Figures 88 - 93 shown in the Nomarski optical micrographs, the surface morphology of the samples grown at temperatures below 895 °C had elongated features or stripes. The stripes were aligned along the [1-10] direction, i.e., perpendicular to the substrate cleavage that represents in-plane anisotropy. As the temperature increased, the length of the elongated features decreased to a few micrometers. At 895 °C and higher temperatures, the surface became more granular and rougher with increasing temperature.

[0261] XRD analysis of sample group A showed that the wurtzite fraction measured perpendicular to the cleavage gradually increased with increasing growth temperature but remained below 1%, as Figure 94 shown. Measurements obtained parallel to the cleavage showed no wurtzite inclusions along this direction and are therefore not shown in the figure.

[0262] Sample group B – V / III ratio series

[0263] The second sample group (Sample Group B) consists of eight samples, where the outer layer growth temperature is kept constant at 875 °C (i.e., the median value of Group A), while the V / III ratio varies between 15 and 1200, and the coefficient between each value is approximately 2. Figures 95 - 102 The Nomarski optical micrographs in Figures 95 - 102 show the surface morphology changing from granular to striated and then to rough as the V / III ratio increases.

[0264] The XRD phase analysis results of Sample Group B indicate that when measured perpendicular to the cleavage, the wurtzite inclusions slightly increase from 0% at V / III = 15 to 1% at V / III = 300, as Figure 103 shown. At the highest V / III values of 600 and 1200, the wurtzite fraction increases more rapidly to 3% and finally reaches 11%. When measured parallel to the cleavage, no significant wurtzite phase was found in all samples and is thus omitted from the figure.

[0265] Sample Group C – NL Thickness Series

[0266] The third sample group (Sample Group C) consists of five samples, where the nucleation layer thickness increases from 3 nm to 44 nm, with an increase of 2 at each step, while the growth conditions of the epitaxial layer are a constant temperature of 875 °C, a pressure of 100 Torr, and a V / III ratio of 76. Figures 104 - 108 shows the surface morphology changes with the NL thickness. Among these five samples, the only sample that seems unusual is the one with the thinnest (3 nm) grown NL, which shows large pits in the epitaxial layer. Otherwise, the surface stripes seem to become rougher as the NL thickness increases.

[0267] The XRD phase analysis of Sample Group C (see Figure 109 ) indicates that the sample with an NL thickness of 3 nm has the largest wurtzite fraction, approximately 3%, while the wurtzite phase in other samples is less than 1%. The integrated intensity of the 002 rocking curve of Sample Group C is as Figure 110 shown. The changing peak intensity indicates that the material quality improves as the NL thickness increases, and the peak intensity reaches saturation when the NL thickness is 22 nm or greater.

[0268] Discussion of Sample Groups A, B, and C

[0269] The structural data of Sample Group A (changing temperature at a constant V / III ratio of 76) indicate that growth temperatures below 895 °C produce relatively smooth film surfaces (see Figures 88 - 93 ), and the wurtzite fraction is less than 1% (see Figure 94)。At higher growth temperatures, the surface is slightly degraded and the wz-GaN fraction approaches 1%. Therefore, at a V / III of 76 and a reaction pressure of 100 Torr, the preferred growth temperature is below about 890 °C.

[0270] Sample group B used a constant growth temperature of 875 °C (i.e., within the most favorable temperature range), while the V / III ratio varied between 15 and 1200, with a coefficient of approximately 2 between each value. At the low and high ends of this range (see Figures 95 - 102 ), a more granular surface is shown. As long as Nomarski micrographs can be ranked according to properties, the surface of samples grown between V / III ratios of 38 and 150 can be observed to be the flattest. Despite the different surface morphologies of the samples, the contamination of wz-GaN is still very small for samples with a V / III of 300 (see Figure 103 ). The wurtzite contamination increases rapidly when the V / III ratio exceeds 300, while the surface morphology deteriorates slightly.

[0271] It is worth noting that sample group A was grown at a V / III ratio of 76, which falls within the most favorable V / III ratio range obtained from the results of sample group B. Similarly, sample group B was grown at an epitaxial layer growth temperature of 875 °C, which falls within the most preferred temperature range obtained from the results of sample group A. Therefore, within this MOVPE growth window, the surface and material properties of the zinc blende GaN film hardly change.

[0272] The study of the low-temperature nucleation layer thickness (sample group C) shows that the preferred thickness is about 22 nm. Using a thinner GaN NL results in a decrease in material quality, while a thicker NL makes the surface morphology coarser.

[0273] In summary, considering the collective data of the surface morphologies and phase purities of sample groups A to C, the preferred MOVPE growth conditions for zinc blende GaN epitaxial layers at a constant pressure of 100 Torr are: a temperature between 850 and 890 °C, a V / III ratio between 38 and 150, within which a relatively smooth film can be obtained with wurtzite contamination less than 1%. The preferred thickness of the NL is about 22 nm.

[0274] Although the present invention has been described in connection with the above exemplary embodiments, many equivalent modifications and variations will be apparent to those skilled in the art when the present disclosure is given. Therefore, the above-described exemplary embodiments of the present invention are considered to be illustrative rather than restrictive. Various changes can be made to the described embodiments without departing from the spirit and scope of the present invention.

[0275] All references mentioned above and / or listed below are incorporated herein by reference.

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[0346] Figure 71 References in:

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Claims

1. A method of manufacturing a semiconductor structure including a (001)-oriented zinc blende structure group III nitride layer, characterized in that, The method comprises the following steps: Providing a silicon substrate; Providing a 3C-SiC layer on the silicon substrate; Subjecting the 3C-SiC layer to a nitridation step at a temperature T1 in the range of 800 - 1100 °C; Growing a wurtzite structure group III nitride nucleation layer at a rate of 0.1 - 1 nm / s to a thickness in the range of 10 - 100 nm at a temperature T2 in the range of 550 - 650 °C; Subjecting the nucleation layer to a recrystallization step at a temperature T3 in the range of 800 - 920 °C; and Depositing and growing a wurtzite structure group III nitride layer to a thickness of at least 0.3 μm at a temperature T3 in the range of 800 - 920 °C by MOVPE; wherein the group-III nitride layer is based on In x Al y Ga 1-x-y N, where 0 ≤ x ≤ 1 and 0 ≤ y ≤ 1.

2. The method according to claim 1, characterized in that, Growing the group III nitride nucleation layer to a thickness in the range of 10 - 50 nm.

3. The method according to claim 1 or 2, characterized in that, In the step of depositing and growing the wurtzite structure group III nitride layer, the reactor pressure is not greater than 500 Torr.

4. The method according to claim 1 or 2, characterized in that, In the step of depositing and growing the wurtzite structure group III nitride layer, the reactor pressure is not greater than 300 Torr.

5. The method according to claim 1 or 2, characterized in that, In the step of depositing and growing the wurtzite structure group III nitride layer, the relative flow rate V / III ratio of the group V precursor to the group III precursor is not greater than 300.

6. The method according to claim 5, wherein The V / III ratio is not greater than 150.

7. The method according to claim 1, wherein The diameter of the silicon substrate is at least 100 mm.

8. The method according to claim 1, characterized in that, The semiconductor structure is incorporated in a semiconductor device selected from the group consisting of: lasers, diodes, transistors or sensors.

9. A method of manufacturing a semiconductor structure including a (001)-oriented zinc blende structured group III nitride layer, characterized in that, The method comprises the following steps: Providing a silicon substrate; Providing a 3C-SiC layer on the silicon substrate; Subjecting the 3C-SiC layer to a nitridation step at a temperature T1 in the range of 800 - 1100 °C; Growing a GaN nucleation layer at a rate of 0.1 - 1 nm / s to a thickness in the range of 10 - 100 nm at a temperature T2 in the range of 550 - 650 °C; Subjecting the nucleation layer to a recrystallization step by heating at a rate of 0.1 - 10 °C / s to a temperature T3 in the range of 800 - 920 °C; and Depositing and growing a wurtzite structure group III nitride layer at a rate of 0.1 - 1 nm / s to a thickness of at least 0.3 μm at a temperature T3 in the range of 800 - 920 °C by MOVPE, wherein the relative flow rate V / III ratio of the group V precursor to the group III precursor is 38 - 300.

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