Group III nitrides of sphalerite structure
阅读说明：本技术 闪锌矿结构iii族氮化物 (Group III nitrides of sphalerite structure ) 是由 大卫·约翰·沃利斯 马丁·弗伦特鲁普 门诺·约翰尼斯·卡珀斯 苏曼拉塔·莎欧塔 于 2018-03-29 设计创作，主要内容包括：公开了一种制造包括(001)取向的闪锌矿结构III族氮化物层的半导体结构的方法,例如GaN。该层形成在硅衬底上的3C-SiC层上。形成成核层,进行再结晶,然后通过MOVPE在750-1000℃范围内的T3温度下形成厚度至少为0.5μm的闪锌矿结构Ⅲ族氮化物层。还公开了相应的包括闪锌矿结构III族氮化物层的半导体结构,当通过XRD表征时,该层的大部分或全部由闪锌矿结构III族氮化物形成,优先于纤锌矿结构III族氮化物。(A method of fabricating a semiconductor structure, such as GaN, including a (001) oriented zincblende structure group III nitride layer is disclosed. The layer is formed on a 3C-SiC layer on a silicon substrate. Forming a nucleation layer, performing recrystallization, and then forming a zinc blende-structured group III nitride layer with a thickness of at least 0.5 μm by MOVPE at a temperature T3 in the range of 750-1000 ℃. Also disclosed are corresponding semiconductor structures comprising a sphalerite structure group III nitride layer, most or all of which is formed from sphalerite structure group III nitride in preference to wurtzite structure group III nitride when characterized by XRD.)
1. a method of fabricating a semiconductor structure comprising a substantially (001) oriented zincblende structure group III nitride layer, the method comprising the steps of:
providing a silicon substrate;
providing a 3C-SiC layer on the silicon substrate;
growing a group III nitride nucleation layer;
performing a nucleation layer recrystallization step; and
the zinc blende structure group III nitride layer is deposited and grown by MOVPE at a temperature T3 in the range of 750-1000 ℃ to a thickness of at least 0.3 μm.
2. The method of claim 1, wherein the group III nitride nucleation layer is a sphalerite structure group III nitride nucleation layer.
3. The method as set forth in claim 1 or 2 wherein the 3C-SiC layer is subjected to a nitridation step at a temperature T1 in the range of 800-1100 ℃ prior to growing the group III nitride nucleation layer.
4. The method of any one of claims 1 to 3, wherein the group III-nitride nucleation layer is grown at a T2 temperature in a range of 500 ℃ to 700 ℃.
5. The method of any of claims 1-4, wherein the group III-nitride nucleation layer is grown to a thickness greater than 3nm and not greater than 100 nm.
6. The method of any one of claims 1 to 5, wherein the group III-nitride nucleation layer is grown to a thickness in the range of 10-50 nm.
7. The method of any one of claims 1 to 6, wherein, in the step of depositing and growing the zincblende structure group III nitride layer, a reactor pressure is not greater than 500 Torr.
8. The method of any one of claims 1 to 7, wherein, in the step of depositing and growing the zincblende structure group III nitride layer, a reactor pressure is not greater than 300 Torr.
9. The method of any one of claims 1 to 3, wherein, in the step of depositing and growing the zincblende structure group III nitride layer, the ratio of V to III is not greater than 300.
10. The method of any one of claims 1 to 9, wherein the ratio of V to III is no greater than 150.
11. The process as claimed in any one of claims 1 to 10, wherein the temperature T3 is in the range of 800-920 ℃.
12. The process as claimed in any one of claims 1 to 11, wherein the temperature T3 is in the range 820-890 ℃.
13. The method of any of claims 1-12, wherein the group III nitride layer is In-basedxAlyGa1-x-yAnd N, wherein x is more than or equal to 0 and less than or equal to 1, and y is more than or equal to 0 and less than or equal to 1.
14. The method of any one of claims 1 to 13, wherein the silicon substrate has a diameter of at least 100 mm.
15. A semiconductor structure comprising a sphalerite structure group III nitride layer, wherein:
the group III nitride layer has a thickness of at least 0.3 μm; and is
The group III nitride layer is a single crystal sphalerite structured group III nitride having an intensity I due to a wurtzite structured group III nitride 10-11 reflection when the group III nitride layer is subjected to XRD characterization10-11And intensity I due to sphalerite structure group III nitride 002 reflection002The following relationship is satisfied:
16. the semiconductor structure of claim 15, wherein the intensity I due to wurtzite structure group III nitride 10-11 reflection10-11And intensity I due to sphalerite structure group III nitride 002 reflection002The following relationship is satisfied:
17. the semiconductor structure of claim 16, wherein the intensity I due to wurtzite structure group III nitride 10-11 reflection10-11And intensity I due to sphalerite structure group III nitride 002 reflection002The following relationship is satisfied:
18. the semiconductor structure of any one of claims 15 to 17, wherein the intensity I attributed to wurtzite structure group III nitride 10-11 reflection10-11And intensity I due to sphalerite structure group III nitride 002 reflection002Determined by two-dimensional reciprocal space mapping to form a measured reciprocal space map comprising expected reflections of sphalerite structure group III nitride 002 and wurtzite structure group III nitride 10-11.
19. The semiconductor structure of claim 18, wherein the intensity I attributed to wurtzite structure group III nitride 10-11 reflection at a position in reciprocal space10-11Is caused by stacking faults formed on the 111 planes of the zincblende structure group III nitride, as evidenced by the elongated stripes in the measured reciprocal space diagram between the reflection of the zincblende structure group III nitride 002 and the reflection of the expected wurtzite structure group III nitride 10-11.
20. A semiconductor structure comprising a sphalerite structure group III nitride layer, wherein:
the group III nitride layer has a thickness of at least 0.3 μm; and is
The group III nitride layer is a single crystal sphalerite structure group III nitride, as for the group III nitrideWhen the nitride layer is subjected to XRD characterization, the volume V of the group III nitride with the sphalerite structurezbVolume V of group III nitride of wurtzite structurewzSatisfies the following relationship:
wherein VwzIs based on a wurtzite structure group III nitride 1-103 reflection evaluation, VzbIs based on the zinc blende structure group III nitride 113 reflection evaluation, the evaluation basis is:
wherein:
Vuczbis the volume of the zincblende structure group III nitride unit cell,
Vucwzis the volume of the wurtzite structure group III nitride unit cell,
F113is the structural amplitude of the sphalerite structure group III nitride 113 reflection,
F1-13is the structural amplitude of the wurtzite structure group III nitride 1-103 reflections,
2θ113is the 2 theta angle of reflection of the sphalerite structure group III nitride 113,
2θ1-13is the 2 theta angle reflected by the wurtzite structure group III nitrides 1-103,
I113is the integrated intensity of the sphalerite structure group III nitride 113 reflection,
I1-13is the integrated intensity of the wurtzite structure group III nitride 1-103 reflections.
21. The semiconductor structure of claim 20, wherein the sphalerite structure group III nitride volume VzbVolume V of group III nitride of wurtzite structurewzSatisfies the following relationship:
22. the semiconductor structure of any one of claims 15 to 21, wherein the zincblende structure group III nitride layer is substantially (001) oriented.
23. The semiconductor structure of any one of claims 15 to 22, wherein the zinc blende structure group III nitride layer has a reflective layer interposed between the zinc blende structure group III nitride layer and a substrate.
24. The semiconductor structure of any one of claims 15 to 23, wherein the zincblende structure group III nitride layer has a dimension of at least 1mm in two directions orthogonal to each other and orthogonal to a thickness direction.
25. A semiconductor device incorporating the semiconductor structure of any one of claims 15 to 24, wherein the semiconductor device is selected from the group consisting of a Light Emitting Diode (LED), a laser, a diode, a transistor, a sensor.
Technical Field
The present invention relates to the formation of zinc blende-structured group III nitride layers, such as GaN, AlGaN, InGaN, InAlN, and more generally InxAlyGa1-x-yAnd N is added. Characterization of these layers and methods of forming them are disclosed herein. These materials have particular, but not necessarily exclusive, application in the field of semiconductor structures and devices, for example, light emitting applications such as LEDs, lasers, and other devices such as transistors, diodes, sensors, and the like.
Background
Disclosure of Invention
In order to obtain single-phase epitaxial films with reasonable crystal quality, the growth process must be supported by robust structural characterization techniques.
The inventors devised the present invention in order to solve at least one of the above problems. Preferably, the present invention reduces, ameliorates, avoids or overcomes at least one of the above problems.
Accordingly, in a first aspect, the present invention provides a method of fabricating a semiconductor structure comprising a substantially (001) oriented zincblende structure group III nitride layer, the method comprising the steps of:
providing a silicon substrate;
providing a 3C-SiC layer on the silicon substrate;
growing a group III nitride nucleation layer;
performing a nucleation layer recrystallization step; and
the zinc blende structure group III nitride layer is deposited and grown by MOVPE at a temperature T3 in the range of 750-1000 ℃ to a thickness of at least 0.3 μm.
The inventors have found that these steps provide a sphalerite-structured group III nitride layer with improved crystalline quality, particularly with respect to reducing the formation of wurtzite-structured group III nitride inclusions.
In the present disclosure, some numerical ranges are expressed as open-ended ranges having upper or lower limits or closed-ended ranges having upper and lower limits. It is expressly noted that preferred ranges are disclosed herein, which are combinations of upper and/or lower limits of different ranges of the same parameter.
Preferably, the 3C-SiC layer is subjected to a nitridation step at a temperature T1 in the range of 800-. This step is advantageous to ensure that there is sufficient available N for the subsequent group III nitride formation. The temperature T1 used outside this range reduces PL NBE peak intensity and broadens the emission FWHM.
We now consider the deposition and growth conditions of the group III nitride nucleation layer. Preferably, the group III nitride nucleation layer is grown at a temperature T2 in the range of 500 ℃ to 700 ℃. More preferably, the temperature T2 is in the range of 550-650 ℃. The growth rate may be at least 0.1 nm/s. The growth rate may be as high as 1 nm/s. The thickness of the nucleation layer may be at least 3nm, but more preferably greater than 3 nm. More preferably, NL may have a thickness of at least 10 nm. The thickness of the Nucleation Layer (NL) can be up to 100 nm. Preferably, the thickness of NL can be up to 50 nm. More preferably, NL may be as thick as 40 nm. Preferably, the temperature T2 is selected to be about 40-60 ℃ 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.
After the nucleation layer is grown, a nucleation layer recrystallization step is performed. In this step, the temperature is preferably raised at a rate of 0.1 to 10 deg.C/sec. More preferably, the temperature is increased at a rate of 0.5-5 deg.C/sec. The method is suitable for achieving satisfactory recrystallization of the nucleation layer, allowing subsequent high quality epitaxial layer deposition.
The group III nitride nucleation layer is preferably a sphalerite structure group III nitride nucleation layer.
In the step of depositing and growing the zincblende structure 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 no greater than 300 torr. More preferably, the reactor pressure is no greater than 100 torr.
In the step of depositing and growing the zincblende structure group III nitride layer on the recrystallized nucleation layer, the ratio of V to III is preferably in the range of 10 to 300. More preferably, the ratio of V to III is in the range of 20-150. 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/sec. For example, a growth rate of about 0.5 nm/sec is suitable. Careful selection of the ratio of V to III within the preferred ranges improves surface morphology, sphalerite phase purity and XRD rocking curve peak width.
In the step of depositing and growing the zincblende structure group III nitride layer on the recrystallized nucleation layer, the temperature T3 is preferably in the range of 800-. More preferably, the temperature T3 is at least 810 ℃, more preferably at least 820 ℃, more preferably at least 830 ℃. The temperature T3 is preferably at most 910 ℃, at most 900 ℃ or at most 890 ℃. A particularly suitable range for T3 is 845-880 ℃. Careful selection of temperature T3 within the preferred range improves Nomarski image surface morphology, XRD rocking curve peak width and PL. For example, samples grown in the 860-880 ℃ range exhibited a relatively smooth surface and the corresponding NBE PL peak was the strongest, although also the widest of the data shown here. At higher growth temperatures, the surface becomes rough and the PL NBE peak becomes clearly narrower, but the intensity of the yellow band increases.
The change in surface roughness with growth temperature was confirmed by AFM. The phase purity as measured by X-ray diffraction shows that the amount of wurtzite inclusions can be greatly reduced when T3 is 900 ℃ or less. XRD analysis showed an increase in the reflection contribution due to the wurtzite lattice above T3 ℃, thus indicating the incorporation of hexagonal inclusions into the cubic zincblende matrix.
The inventors have found in this work that the preferred conditions for temperature T3 and the III-V ratio can be broadened by conducting epitaxial layer growth at a relatively low pressure. In an exemplary set of conditions for growing zincblende GaN epitaxial layers by MOVPE at a constant pressure of 100 torr, T3 may be in the range of 850 to 890 ℃, with V/III ratios ranging from 38 to 150, resulting in relatively smooth films with wurtzite contamination of less than 1%. The preferred thickness of NL is in the range of 10-50nm, e.g. about 22 nm.
Preferably, the group III nitride layer is In-basedxAlyGa1-x-yAnd N, wherein x is more than or equal to 0 and less than or equal to 1, and y is more than or equal to 0 and less than or equal to 1.
The silicon substrate has a diameter of at least 100 mm. Different substrate diameters are possible. Notably, the growth processes described herein can be readily extended to substrates of any suitable size, for example at least 150mm, at least 200mm, or at least 300 mm.
In a second aspect, the present invention provides a semiconductor structure comprising a zincblende structure group III nitride layer, wherein:
the group III nitride layer has a thickness of at least 0.5 μm; and is
The group III nitride layer is a single crystal sphalerite structured group III nitride having an intensity I due to a wurtzite structure group III nitride 10-11 reflection when the group III nitride layer is XRD characterized10-11And intensity I due to sphalerite structure group III nitride 002 reflection002The following relationship is satisfied:
the features of the first aspect of the invention may be combined with the features of the second aspect of the invention, individually or in any combination, unless the context requires otherwise.
Preferably, at least one of the following relationships applies:
more preferably, the intensity I attributed to the reflection of the wurtzite structure group III nitride 10-1110-11And intensity I due to sphalerite structure group III nitride 002 reflection002The following relationship is satisfied:
intensity I attributed to wurtzite structure group III nitride 10-11 reflection10-11And intensity I due to sphalerite structure group III nitride 002 reflection002Can be determined by two-dimensional reciprocal space mapping to form a measured reciprocal space map comprising the expected reflections of sphalerite structure group III nitride 002 and wurtzite structure group III nitride 10-11. Reciprocal space diagrams are well known to those skilled in the art and are effective in capturing large amounts of data to indicate the presence of crystalline phases in a film.
Intensity I attributed to wurtzite structure group III nitride 10-11 reflection at a position in reciprocal space10-11Possibly due to stacking faults formed on the 111 planes of the zincblende structure group III nitride, as evidenced by the elongated stripes in the measured reciprocal space diagram between the reflection of the zincblende structure group III nitride 002 and the reflection of the expected wurtzite structure group III nitride 10-11. In this way, there may be measurable reflected X-ray intensities at defined locations in the reciprocal space, but this does not necessarily imply the presence of wurtzite structure group III nitride inclusions. Instead, the X-ray intensity may be provided by reflection of stacking faults. Hexagonal stacking faults have less impact on the properties of the group III nitride layer than wurtzite-structure inclusions.
In a third aspect, the present invention provides a zincblende structure group III nitride layer, wherein:
the group III nitride layer has a thickness of at least 0.5 μm; and is
The group III nitride layer is a single crystal zincblende structure group III nitride that, when the group III nitride layer is XRD characterized, results in a zincblende structure group III nitride objectProduct VzbVolume V of group III nitride of wurtzite structurewzSatisfies the following relationship:
wherein VwzIs based on a wurtzite structure group III nitride 1-103 reflection evaluation, VzbIs based on the zinc blende structure group III nitride 113 reflection evaluation, the evaluation basis is:
wherein:
Vuczbis the volume of the zincblende structure group III nitride unit cell,
Vucwzis the volume of the wurtzite structure group III nitride unit cell,
F113is the structural amplitude of the sphalerite structure group III nitride 113 reflection,
F1-13is the structural amplitude of the wurtzite structure group III nitride 1-103 reflections,
2θ113is the 2 theta angle of reflection of the sphalerite structure group III nitride 113,
2θ1-13is the 2 theta angle reflected by the wurtzite structure group III nitrides 1-103,
I113is the integrated intensity of the sphalerite structure group III nitride 113 reflection,
I1-13is the integrated intensity of the wurtzite structure group III nitride 1-103 reflections.
Features of the first aspect of the invention and/or features of the second aspect of the invention may be combined with features of the third aspect of the invention, individually or in any combination, unless the context requires otherwise.
Preferably, at least one of the following relationships applies:
preferably, the zincblende structure group III nitride layer of the semiconductor structure of the second and/or third aspect is substantially (001) oriented.
The thickness of the zincblende structure group III nitride layer of the semiconductor structure of the second and/or third aspect may be at least 0.3 μm.
In any of the first, second or third aspects, the thickness of the zincblende structure 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.
The zincblende structure group III nitride layer may have a reflective layer interposed between the zincblende structure group III nitride layer and the substrate. This is useful for structures used as devices, especially in light emitting devices.
The zincblende structure group III nitride layer may have a dimension of at least 1mm in two directions orthogonal to each other and to the thickness direction. More preferably, these dimensions may be at least 2mm, at least 3mm, at least 4mm or at least 5 mm.
However, it should be understood that for some useful devices, the size of the zincblende structure group III nitride layer may be smaller than the above-mentioned sizes. For example, cutting structures for a particular device.
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: light Emitting Diodes (LEDs), lasers, diodes, transistors, sensors.
Other optional features of the invention are set out below.
Drawings
Embodiments of the invention will now be described, by way of example, with reference to the accompanying drawings, in which:
FIG. 1 schematically illustrates the optical path and the different goniometer movements in X-ray diffraction characterization.
Fig. 2 schematically illustrates a measurement geometry mapping a texture in reciprocal space and its projection on a two-dimensional map.
Fig. 3, 6, 9 and 12 provide Nomarski microscope images showing the effect of nitridation temperature on the morphology of the zincblende GaN film surface.
Fig. 4, 5, 7, 8, 10, 11, 13, 14, 16 and 17 show the variation of Photoluminescence (PL) optical characteristics with nitridation temperature.
FIG. 18 shows the growth rate of the Nucleation Layer (NL) as a function of growth temperature and V/III ratio.
Fig. 19 to fig. 27: FIGS. 19, 22 and 25 show the changes in surface morphology as a function of growth temperature (845, 860 and 880 ℃) observed from Nomarski microscope images. FIGS. 20, 21, 23, 24, 26 and 27 show the change of PL optical properties with growth temperature (845, 860 and 880 ℃).
Fig. 28 shows the surface roughness (open circle symbols) and the hexagonal-cubic ratio (solid circle symbols) as a function of growth temperature observed from an Atomic Force Microscope (AFM).
Figure 29 shows XRD 002 and 004 rocking curves FWHM as a function of growth temperature.
Fig. 30 to fig. 41: fig. 30, 33, 36 and 39 show the surface morphology of Nomarski images as a function of V-III ratio (23 to 152 range) at a constant growth temperature of 880 ℃. FIGS. 31, 32, 34, 35, 37, 38, 40 and 41 show the variation of PL optical properties with V-III ratio (23 to 152 range) at a constant growth temperature of 880 ℃.
Fig. 42-fig. 50: fig. 42, 45 and 48 show the surface morphology of Nomarski images as a function of growth pressure (100, 300 and 500 torr). Fig. 43, 44, 46, 47, 49 and 50 show the PL optical properties as a function of growth pressure (100, 300 and 500 torr).
Fig. 51-fig. 56: fig. 51 and 54 show the surface morphology of Nomarski images as a function of layer thickness grown under conditions of T860 ℃, P300 torr and V/III 80 (example layer thicknesses of 600 and 750 nm). Fig. 52, 53, 55 and 56 show the PL optical properties as a function of layer thickness (example layer thicknesses of 600 and 750 nm).
Fig. 57-62 show the surface morphology of Nomarski images as a function of silane flow rate for 500nm thick films grown under conditions of T860 ℃, P300 torr, and V/III 80.
FIGS. 63-68 show the PL optical properties as a function of silane flow rate.
Fig. 69 shows the Si concentration in the GaN film measured by SIMS based on the silane flow rate.
Fig. 70 shows the morphology of a typical green QW-structured Nomarski image, taken at the center of a single mesa, according to an embodiment of the present invention.
Fig. 71 and 72 show PL optical characteristics of the green QW structure of fig. 70 at high and low laser powers, respectively.
Fig. 73-75 show polar diagrams of different wurtzite and zincblende reflections of GaN grown on a (001) oriented 3C-SiC/Si template.
Fig. 76 shows the crystal arrangement of wurtzite and zincblende GaN phases.
Fig. 77 and 78 show two-dimensional reciprocal space plots (RSM) of zincblende GaN samples grown under non-optimal conditions that promote the formation of wurtzite inclusions (fig. 77) and under improved conditions that produce nearly 100% pure zincblende GaN (fig. 78), with stacking fault Striations (SF), Detector Striations (DS), Crystal Truncation Rods (CTR), and Bragg Rings (BR).
Fig. 79 shows XRD peak width plots of optimized sphalerite GaN samples in conventional Williamson-Hall plots with Lorentzian (n ═ 1) and gaussian shaped peaks (n ═ 2) fitted.
FIG. 80 shows the extrapolated peak width β in reciprocal spacehkl·|Ghkl| and polar angle c and scattering vector magnitude | G estimated by a series of tilt-symmetric ω -scanshklThe relationship of | is given. The circles show the measured reflections used for the extrapolation.
Figure 81 shows that the XRD ω -line width (FWHM) decreases with increasing film thickness for oriented zincblende GaN grown on related substrates 3C-SiC, GaAs, MgO and Si.
FIGS. 82 and 83 show XRD wafer bending analysis examples of a 4' 3C-SiC/Si template, showing a convex curvature of-51.5 km-1。
FIG. 84 shows PL emission wavelength and PL integrated intensity versus QW thickness.
Fig. 85 shows a process flow diagram for forming a GaN layer according to an embodiment of the invention.
FIG. 86 shows a schematic cross-sectional view of forming a GaN layer on SiC/Si.
Fig. 87 shows a schematic cross-sectional view of a GaN layer transferred onto another substrate for use as a semiconductor device (e.g., LED).
Fig. 88-93 show Nomarski optical microscopy images showing the surface of GaN epitaxial layers grown at a constant V/III of 76 and at a pressure of 100 torr over a temperature range of 850 to 910 ℃. The temperature used is marked on each image.
Figure 94 shows wurtzite fraction as a function of temperature for a GaN epitaxial layer obtained from XRD grown at a constant V/III of 76 and a pressure of 100 torr.
Fig. 95-102 show Nomarski optical microscope images showing the surface of GaN epitaxial layers grown at a temperature of 875 ℃ at V/III of 15-1200 and at a pressure of 100 torr. The V/III ratio used is marked on each image.
Fig. 103 shows wurtzite fraction as a function of V/III ratio of a GaN epitaxial layer obtained from XRD, wherein the epitaxial layer was grown at a constant temperature of 875 ℃ and a pressure of 100 torr.
Fig. 104-108 show Nomarski optical microscope images showing GaN epitaxial layers grown at a temperature of 875 ℃ at V/III of 76 and at a pressure of 100 torr as a function of GaN Nucleation Layer (NL) thickness. The thickness of the nucleation layer used is marked on each image.
Fig. 109 shows wurtzite fraction of GaN epitaxial layers from XRD as a function of GaN Nucleation Layer (NL) thickness, where the epitaxial layers were grown at a temperature of 875 ℃, V/III of 76, and pressure of 100 torr.
Fig. 110 shows the XRD 002 rocking curve integrated intensity versus GaN Nucleation Layer (NL) thickness for GaN epitaxial layers grown at 875 ℃, 100 torr, and 76V/III.
Fig. 111 and fig. 112: wurtzite fraction in zincblende GaN epitaxial layers 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). The temperature dependent samples in FIG. 111 were grown at a constant V/III of 76. The V/III dependent samples in FIG. 112 were grown at 875 deg.C (100 torr) and 880 deg.C (300 torr).
Detailed description of the preferred embodiments and other optional features of the invention
Solving the green gap problem is a key challenge for future development of LED-based lighting systems. One promising approach to achieve higher LED efficiencies in the green spectral region is to grow group III nitrides in a cubic crystalline sphalerite phase. However, the metastable and crystal growth processes of zincblende GaN typically result in phase mixing with the wurtzite phase, high mosaicity, high density of extension and point defects, and strain, all of which can compromise the performance of the light emitting device. X-ray diffraction (XRD) is the primary characterization technique for analyzing these device-related structural characteristics because it is very inexpensive and the feedback is rapid compared to other techniques. In this disclosure, we propose various XRD techniques to identify phase purity in mainly zincblende GaN films to analyze their mosaicity, strain state and wafer curvature. Different techniques were used and described on samples grown on 4 "SiC/Si wafers by MOVPE or MOCVD (metal organic vapor deposition).
X-ray diffraction (XRD) is a method suitable for the above purpose because it is non-destructive, well-established and can quickly provide detailed information about the structural characteristics of crystalline materials. Herein, we illustrate how XRD techniques can be used to identify texture, phase purity, crystal orientation, and quantify the mosaicity of cubic zincblende films. We used the proposed technique to characterize the properties of epitaxial GaN films grown on low-cost, large-area (001) cubic 3C-SiC/Si templates to provide fast feedback for further growth optimization.
Crystallographic properties of group III nitrides
III-nitride materials AlN, GaN, InN and alloy thereof, AlxGa1-xN、InyGa1-yN、InxAlyGa1-x-yN (0. ltoreq. x.ltoreq.1, 0. ltoreq. y.ltoreq.1) can crystallize in wurtzite, sphalerite and halite structures, the first two of which are the most common phases in epitaxial films [ Ambacher (1998); hanada (2009)]. The hexagonal wurtzite phase and the cubic sphalerite phase of the GaN-based semiconductor are two different polymorphs of the same material. In both structures, the bonds between the metal ions and the nitrogen ions are tetrahedrally coordinated, and the inter-ionic distances within the closely packed planes are approximately the same. The main difference between these two structures is that the stacking order of the faces is different, in wurtzite structure it is … AaBbAaBbAaBb … for the (0001) plane; whereas in the zincblende structure, for the (111) plane, it is … aabbcacaabbcc …, where Aa, Bb and Cc represent different metal N bilayers. The distance between crystal faces in the wurtzite structure is
In the zincblende structure is
[ Cullity (1978) ]. Where 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, and the formation energy of the two phases can be similar [ Yeh (1992) ], so that a certain fraction of the two polymorphs form during the growth of the material with defects.
It should be noted that as is well known, cubic structures can be described by three Miller index symbols (h, k, l), while hexagonal structures can be described by Bravais-Miller index symbols (h, k, i, l). However, in the Bravais-Miller index notation, i ═ h + k, so that the crystal planes of the hexagonal structure and the X-ray reflection can also be described completely by three Miller index notations (h, k, l).
Description of the experiments and basis for X-ray diffraction
X-ray diffraction (XRD) is one of the most common methods of characterizing crystalline samples. The method is based on the measurement of X-ray reflections, the pattern of which represents a fourier transform image of the crystal structure in reciprocal space. Diffraction angle of hkl reflection (2 θ) and distance d of (hkl) planehklAre related to each other by bragg's law:
2dhklsin θ ═ λ (formula 3).
For the X-ray characterization of cubic zincblende-type GaN films, we used two different standard laboratory diffractometers and Cu-Ka1Light source
The reciprocal space diagram was measured using a PANalytical Empyrean diffractometer equipped with a 2-bounce hybrid monochromator, 1/4 ° slit, euler ring and PIXcel solid state area detector. This configuration ensures high strength and allows rapid and accurate measurement of large maps in reciprocal space.
Since proper alignment is an important factor in accurately assessing bragg reflection and lattice properties, the main beam is calibrated for the goniometer 2 θ angle prior to each measurement. The sample is then moved into the main beam path (z-move) until half the intensity is blocked, as best practice suggested by Fewster and Andrew (1995).
The measurements are mainly directed to zincblende-type GaN thin films grown by Metal Organic Vapor Phase Epitaxy (MOVPE) on a 4' (001) cubic 3C-SiC template deposited on a Si substrate, as will be discussed in detail below. To relieve strain in large area templates, a 3 μm to 8 μm thick SiC layer was etched with a polycrystalline square grid to form a millimeter-scale length mesa structure. The present invention is not limited to the use of substrates of this size, but it should be noted that embodiments of the present invention allow the use of large area substrates, which allows growth and device processing to be performed efficiently.
A suitable SiC/Si substrate is disclosed in US 2016/0247967.
The SiC layer is polished by Chemical Mechanical Polishing (CMP) to reduce the surface roughness (measured by AFM) from about 5-10nm to less than about 1 nm.
Film preparation
Growth of zincblende GaN and InGaN using MOVPE is described below. A so-called two-step growth method is used which includes performing substrate cleaning and a nitridation step at high temperature (800 ℃ < T1<1100 ℃), growing a thin GaN nucleation layer at low temperature (500 ℃ < T2<700 ℃), a nucleation layer recrystallization step, and depositing a GaN layer at high temperature (750 ℃ < T3<1000 ℃).
The temperatures referenced herein are emissivity-corrected high temperature measurements measured using an in-situ monitoring tool EpiTT provided by LayTec AG. This value is calibrated against a Si/Al eutectic wafer.
Fig. 85 shows a process flow diagram for forming a GaN layer according to an embodiment of the invention.
FIG. 86 shows a schematic cross-sectional view of forming a GaN layer and an active device on SiC/Si. 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. Active device layer 108 is then formed.
Fig. 87 shows a schematic cross-sectional view of the GaN layer formed in fig. 86 transferred onto another substrate for use as a semiconductor device (e.g., LED). A reflective p-contact 110 is first formed on the active device layer 108. At this time, a buried n-contact layer (not shown) may also be formed, if necessary. The device structure is then bonded to a silicon handle 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 may be left in place as shown and used as part of any light extraction structures or they may be removed leaving only the active device layer 108 bonded to the silicon handle wafer 114 for further device processing.
The thin film is characterized by the well-known Photoluminescence (PL) technique, in which the thin film is illuminated above the bandgap of the GaN material with a laser (266nm 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 band gap of cubic GaN is about 3.2eV and that of hexagonal GaN is about 3.4 eV.
Substrate nitridation:
the substrate was 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. The use of temperatures below or above 960 ℃ will reduce the PL NBE peak intensity and broaden the emission FWHM as shown in fig. 4, fig. 5, fig. 7, fig. 8, fig. 10, fig. 11, fig. 13, fig. 14, fig. 16, fig. 17.
Nitridation of the surface may partially remove oxide from the SiC surface and fill the surface with N atoms when Ga is available, in preparation for forming GaN.
In an alternative embodiment, the oxide may be removed completely thermally, but this requires a treatment temperature of about 1400 ℃.
Nucleation layer deposition：
About 40nm of GaN was deposited using TMG at a flow rate of 93. mu. mol/min, ammonia at a flow rate of 0.15slm, at a reactor pressure of 500 torr and a temperature of 575 deg.C. The growth rate was about 0.3 nm/s. The growth temperature is selected to be about 40-60 ℃ 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. FIG. 18 shows the growth rate of the Nucleation Layer (NL) as a function of growth temperature and V/III ratio.
Due to the low temperature, the nucleation layer formation step is believed to form relatively small GaN nuclei.
It should be noted that a buffer layer may be provided between the SiC layer and the GaN layer, if desired, in order to manage the thermal expansion mismatch between these layers. The buffer layer may be made of AlN or AlxGa1-xN layers (in which the composition is stepwise or graded continuously) or from AlN/AlxGa1-xAnd N layers.
Recrystallization：
The temperature was raised to the epitaxial layer growth temperature at a rate of 1 deg.c/sec at a pressure of 100 torr and an ammonia flow of 0.5slm, and a hydrogen flow of 20 slm. The residence time was 30 seconds. Alternatively, a higher ammonia flow rate may be used, which helps prevent or reduce roughening of the nucleation layer.
The recrystallization step is intended to improve the crystal quality of the nucleation layer by smoothing the layer surface and preventing the formation of surface facets that may lead to wurtzite-like inclusions. In the recrystallization step, N (from NH) is provided on the surface of the GaN nucleation layer3) In order to reduce or avoid decomposition of the GaN nucleation layer at high temperatures.
Epitaxial layer growth：
GaN deposition was carried out using TMG at a flow rate of 140. mu. mol/min, ammonia at a flow rate of 0.25slm, at a reactor pressure of 300 torr and a temperature of 860 ℃ at a growth rate of about 0.5 nm/s. The observed surface morphology and PL changes of Nomarski images (see fig. 19-27) indicate that the optimal growth temperature is about 860 ℃ when the growth temperature is varied between 845 to 880 ℃ at a fixed V/III ratio; samples grown at 860 ℃ showed a relatively smooth surface with the strongest but also widest peak of NBE PL. 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 temperatures cited are based on in-situ real-time measurements using an emissivity-corrected pyrometer. The pyrometers may be calibrated with reference to an Al/Si eutectic temperature or a calibration light source such as the AbsoluT system provided by LayTec AG. Those skilled in the art of MOVPE will be aware of the calibration procedures for such emissivity corrected pyrometer systems.
The change of surface roughness with growth temperature was confirmed by AFM (see fig. 28), and the phase purity was determined from the hexagonal-cubic XRD peak ratio, indicating that no significant hexagonal inclusions were present in the cubic film grown at a temperature of 860 c or less. XRD analysis at growth temperatures above 860 ℃, under the growth conditions reported herein, showed an increase in the contribution of reflections attributed to the wurtzite lattice, thus indicating the incorporation of hexagonal inclusions into the cubic zincblende matrix.
The PL NBE peak narrows with increasing growth temperature, which may correlate with the same trend shown by the FWHM values of the 002 and 004XRD rocking curves (see fig. 29).
When the ratio of ammonia input flow to Ga precursor TMG input flow (V/III ratio, see fig. 30-41) was varied, changes in surface morphology and PL spectra were observed, indicating that higher ammonia, and an increase in PL NBE peak intensity and its FWHM improved surface morphology, while decreasing the intensity of YB.
When the growth pressure was changed from 100 torr to 300 torr to 500 torr (fig. 42-50), the observed changes in surface morphology and PL were similar in effect (from a surface morphology and PL perspective) to a corresponding increase in growth temperature of approximately 15 to 30 ℃.
When the layer thickness was increased, the surface morphology and PL spectra showed that the surface was roughened, but the PL NBE peak intensity increased and the FWHM decreased. Examples of 600 and 750nm thick layers are shown in fig. 51-56.
N-type doping:
when incorporated into the GaN lattice, the Si acts as an electron donor, rendering the GaN n-type. We used the Si precursor silane (SiH4) diluted to 50ppm in hydrogen. Other n-type dopant sources are also possible, such as disilane (Si)2H6) Germane (GeH)4) Digermane (Ge)2H6) Or an oxygen-containing precursor. N-type conductivity was observed for Si-doped layers, and high quality ohmic metal contacts were also demonstrated on Si-doped GaN. For silane flow rates up to 70sccm, the surface topography of the layer as well as the PL NBE peak and yellow emission band were almost unaffected (see FIGS. 57-61). At higher flow rates, dishing can occur (see fig. 62). The optical properties of Si-doped GaN are shown in fig. 63-68, and show the silane flow rate.
Silane input flow rate is linearly proportional to Si concentration, measured in bulk using SIMS (see fig. 69).
InGaN/GaN quantum well:
InGaN deposition was carried out using TMG at a flow rate of 8.2. mu. mol/min, TMI at a flow rate of 9.7. mu. mol/min, ammonia at a flow rate of 446mmol/min at a reactor pressure of 300 torr and a temperature of 700 to 800 ℃ at a growth rate of about 0.8 nm/min. For GaN barrier growth, the same conditions were used except for TMI flow.
For a nominal Quantum Well (QW) width of 2nm, the InGaN growth temperature to achieve a peak of 450nmPL is very similar to that used for a standard wurtzite c-plane structure with the same well width. However, increasing the QW width to 10nm, the emission wavelength can be extended to 540nm at the same growth temperature.
Fig. 70 shows an optical micrograph of a typical green QW structure, while fig. 71 and 72 show PL spectra at high and low laser excitation power, respectively. The strong, broad InGaN peak at 521nm shifts to 540nm with increasing excitation density. The structure includes 5 quantum wells, formed at 300 torr with a nominal thickness of 10 nm. It can be observed from fig. 84 that for constant laser excitation power, the peak emission wavelength increases approximately linearly as the well width increases up to a thickness of 8nm when saturation begins. The PL peak intensity shows a smaller drop with increasing well width, but due to the green gap it is much smaller than expected in wurtzite GaN Q-wells.
Texture analysis
Polymorph identification
The phases and orientations present in the GaN film can be identified by XRD texture analysis, where different selection rules are used for reflections occurring in the zincblende and wurtzite phases. For some angles of diffraction 2 θ, the reflections of the two phases are superimposed, but other angles of diffraction are particularly suitable, with only one of the two phases being visible at a time. For example, for a 2 θ of about 34.5 °, both the wurtzite 0002 reflection and the zincblende 111 are inappropriate, as are the wurtzite 11-20 and zincblende 220 reflections (2 θ of about 57.8 °). Herres et al (1999) propose the use of cubic 200(2 theta about 40.0) and hexagonal 10-12(2 theta about 48.1) reflections, respectively, for predominantly (111)zbAnd (0001)wzOriented films and demonstrated that reasonable results can be obtained. However, for a cubic film that is predominantly (001) oriented, it should be more inclined to use different sphalerite reflections, since the 100 reflections from the crystal planes are typically very weak and add surface scattering effects, making it difficult to determine the in-plane relationship of the film. There are several other suitable reflection combinations that we can use for texture analysis, e.g. 113zbReflection (2 theta about 69.0 deg.) and 1-103wzReflections (2 θ is about 63.4 °) because they are well separated in reciprocal space and the characteristic diffraction patterns are relatively easy to interpret.
Weave the picture
For XRD texture analysis, the angular distribution of selected reflections in reciprocal space was measured by plotting the surface of a hemisphere with the radius given by the specific bragg condition (see fig. 2). To do this, the sample is rotated around its surface normal (φ scan) and gradually tilted toward the beam path plane (χ steps) after each scan. The measured intensity is shown as a polar projection (radius χ) or a volume 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 edges of the figure (χ ═ 90 °) represent the directions in the surface plane.
In order to determine the phase purity, principal orientation and crystallographic relationship between different textures of nitride films, at least two texture patterns, one for each phase, must be measured.
FIG. 73. Fig. 74 and 75 show the weave patterns collected at 2 θ 34.5 °, 2 θ 68.9 °, and 2 θ 63.4 ° for GaN epitaxial layers grown on a 3C-SiC/Si template under non-optimized conditions, respectively. In fig. 73, four strong reflections at χ of about 57 ° can be clearly seen. From the four-fold symmetry, it can be seen that these reflections are likely to represent 111 of the sphalerite phasezbAnd (4) reflecting. However, this result alone cannot demonstrate the absence of wurtzite phase. The reflection measured in FIG. 73 may represent 0002 even though no cubic phase is present at allwzPeaks derived from four different twins of hexagonal wurtzite composition grown on the {111} plane of 3C-SiC. In a more likely case, a mixture of the two phases has grown, which will help measure the reflection from the two phases.
For further examination, we measured zincblende 113zbReflexes (FIG. 74) and wurtzite 1-103wzThe distribution of reflections (fig. 75), which do not overlap with the reflections of the other phase. 113zbThe reflection pattern (FIG. 74) shows four-fold symmetry, with three reflections (2 at approximately 72 ° χ and 1 at approximately 25 °) surrounding a common 111zbThe poles are arranged. This clearly confirms that the cubic film has a main orientation of (001), corresponding to the SiC/Si template orientation, which is indicated by the 004Si weak reflection at the center of the pole figure. In FIG. 75, {1-103} of wurtzite phasewzThe facets cause many reflections that form a distorted hexagonal pattern around their common central 0001 pole (not visible). The results show that the hexagonal wurtzite phase is also present in the GaN film, but it is a minority phase, with much weaker reflection intensity compared to the cubic sphalerite reflection.
The phase purity of the GaN mixture can be estimated by integrating the reflected intensity of the phases and determining the ratio of these values. This may be sufficient to provide rapid feedback for optimizing crystal growth activity, as suggested by Herres et al (1999). However, in order to quantify the volume fractions of the zincblende and wurtzite GaN phases more accurately, additional corrections need to be made taking into account the different scattering efficiencies of the two structures and their crystallographic planes. Since the sphalerite and wurtzite phases have the same absorption coefficient, only the different structure factors F are consideredhklCell, cellVolume VUCAnd a geometric correction called the Lorentz Polarization (LP) factor. Neglecting smaller correction factors (e.g. absorption correction and temperature correction), the integrated intensity of the single reflection hkl is related to the material volume VphaseIn certain proportions, given by:
wherein the Lorentz polarization factor is
Amplitude of structure is
The psi factor depends on the diffraction angle (psi ═ sin) of the powder sample-1(theta) and a constant for single crystal crystals [ Reynolds (1986)]. Coordinate xj,yj,zjIs the position of each atom in the unit atom and is listed in table 1. Atomic scattering factor f of Ga and NjIn proportion to the number of electrons per atom, there is also a complex dependence on diffraction angle and wavelength. Relevant details are described in the literature: [ Cullity (1956); international Tables for Crystallography, Volume C (2004); waasmaier and Kirfel (1995)]And an online database [ Cromer-Mann coefficient: http:// www.ruppweb.org/Xray
Htm and http:// www.ruppweb.org/new _ comp/scattering _ factors, htm and DABAX library (ESRF) http:// tx.technique.ac.il/. katrin/f0_ CromerMann.txt]. With these corrections, {113 }in FIGS. 74 and 75zbAnd {1-103}wzThe reflections revealed that the volume fraction of zincblende GaN in this sample was about 69 vol%. The detection accuracy at which we collect data is a few volume percent and is mainly affected by large variations in the integrated intensity of the reflections in different crystal directions.
Thus, when XRD is performed on a group III nitride layerIn characterization, the volume V of group III nitride of sphalerite structurezbVolume V of group III nitride of wurtzite structurewzThe relative volume ratio of (A) to (B) is:
wherein VwzAre evaluated based on wurtzite structure group III nitride 1-103 (i.e., 1-13 in conventional Miller index notation), while VzbAre evaluated based on the sphalerite structure group III nitride 113 reflection.
Rearranging the already given equations:
wherein:
Vuczbis the volume of the zincblende structure group III nitride unit cell,
Vucwzis the volume of the wurtzite structure group III nitride unit cell,
F113is the structural amplitude of the sphalerite structure group III nitride 113 reflection,
F1-13is the structural amplitude of the wurtzite structure group III nitride 1-103 reflections,
I113is the integrated intensity of the sphalerite structure group III nitride 113 reflection,
I1-13is the integrated intensity of the wurtzite structure group III nitride 1-103 reflections.
The crystallographic relationship between the sphalerite and wurtzite phases can be obtained by combining the pole figures in FIGS. 63-65, which is (111)zb||(0001)wzWherein two of the four unequal {111} planes in the zincblende gallium nitride are [ 11-2%]zb||[-1010]wzAnd [ -110 ]]zb||[1-210]wz. This unit cell arrangement is shown in fig. 76, and this is not surprising since the closely packed faces of each structure are parallel to each other and differ only in stacking order. Other families observed similar alignment of the two phases by XRD [ Qu et al (2001); tsuchiya et al(1998)]And measured by transmission electron microscopy [ Trampertet al (1997)]And it has also been found that the sphalerite and wurtzite phases are arranged in alternating zb-wz-lamellae [ Wu et al (1997)]. As we point out in the example of FIG. 75, not necessarily in the four unequivalent {111}zb(0001) wurtzite GaN was formed on each of the faces with equal probability.
Table 1: ideal atomic positions in wurtzite and zincblende unit cells
The sample was grown under optimal conditions to maximize the zincblende phase, which reflected negligible X-ray intensity at the expected location, i.e. only slightly above the background noise level. In these cases, it is likely that the signal is not due to diffraction of hexagonal wurtzite inclusions in the epitaxial layer, but rather originates from diffuse scattering on surface defects (e.g., stacking faults). This can be illustrated by measuring a two-dimensional reciprocal space map, as described below.
In addition to wurtzite inclusions, the zincblende GaN film may also contain twinned zincblende regions. Similar to stacking faults, these are introduced by stacking errors of the individual (111) faces, but contrary to stacking faults, the zincblende substrates continue in a different stacking sequence … aaccbbaaccbb. Twinning [1-10] of zincblende relative to the surrounding GaN matrix]The axis is inclined by about 70.4 DEG, and therefore, the relationship between the twin crystal and the matrix is (111)twin||(115)matrix[Tsuchiya et al(1998)]. This zinc blende twin, and possibly wurtzite-like twin materials with similar relationships, may cause a weak 111 reflection at χ of about 15 ° (circled in fig. 73), and at χ of about 83 ° (outside the range in fig. 73). Their volume fraction is in the low percentage range.
Reciprocal space map for texture analysis
Suitable two-dimensional Reciprocal Space Maps (RSM) of sphalerite and wurtzite phase reflections, combined with omega-2 theta scans and the gradual change in omega angle after each scan, can be used to analyze the phase of GaN samplesPurity, and several other structural characteristics. Suitable reflections include 002ZBAnd 10-11WZAs shown in fig. 77 and 78, the two different samples have and do not have hexagonal inclusions, respectively. In both reciprocal space plots, the 002 reflections of high intensity zincblende GaN and 3C-SiC are clearly visible. Along the edge<111>The low intensity fringes reflected across 002 are caused by the diffuse scattering caused by the 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 the GaN reflection to shift slightly from the ideal position. Another feature of 3C-SiC reflection across a 2 theta arc is the detector fringe (DS), which is caused by the instrument function of the diffractometer. With 002 perpendicular to the surfacezbThe stripes intersected by GaN reflections are so-called (X-ray) crystal cutoff bars (CTRs), the shape of which is influenced by the surface structure, consistent with atomic force microscopy observations. The partially visible bragg ring (2 θ about 35.6 °) in RSM originates from polycrystalline SiC deposited on a 3C-SiC/Si template etched gate, independent of both the SiC mesa region and the GaN epitaxial layers. Since the GaN epitaxial layers are much thinner than the SiC template, GaN grown on etched gates produces much weaker bragg-like rings and is generally not visible. The wurtzite GaN film had wurtzite-like inclusions (FIG. 67), and two additional 10-11 of wurtzite GaN phases appearedwzReflection, which is not present in the samples in which these inclusions are not present (fig. 78). Since the stacking fault striations overlap with wurtzite phase reflections, these striations are easily misinterpreted as signals of small amounts of hexagonal inclusions in the weave pattern.
The reciprocal spatial mapping is performed much faster even at high integration times, compared to the above-described texture pattern, even with a CCD detector. This increases the signal-to-noise ratio, allowing quantification of wurtzite-type GaN inclusions, which are lower than the comparative examples. However, this method assumes a fixed epitaxial relationship and does not provide additional information about the presence of cubic twins.
Mosaic degree analysis
Due to the lack of a suitable homogeneous substrate, cubic zincblende GaN-based nitrides are typically heteroepitaxially grown on heterocubic substrates, such As GaAs [ As et al (2000); yang et al (1996); shen et al (2003); qu et al (2001); tsuchiya et al (1998) ], SiC [ Wu et al (1997); chichibu et al (2003) ], Si [ Lei et al (1991) ] and various other cubic materials (e.g., GaP [ Cheng et al (1995) ], MgO [ Compe a nGarca et al (2015) ]). Lattice mismatch between different materials can lead to high mosaicism and the formation of defects at grain boundaries of the epitaxial layers. In general, mosaicing should be avoided because it can negatively affect the physical properties of the sample, e.g., create high electrical resistance at grain boundaries [ Fujii et al (2010) ]. Therefore, the mosaicism should preferably be quantified to optimize crystal growth.
In a simplified model derived from powder diffraction methods, the thin film consists of mosaic patches (grains) whose finite dimensions and orientations differ slightly from each other. The diffusion of size, tilt and twist, as well as microstrain and compositional non-uniformities (for alloys) result in broadening of X-ray reflections in reciprocal space. Mosaic tilt causes the reflection angle normal to the surface to expand, while twist causes azimuthal spread around the surface normal. Thus, the absolute broadening Δ G of the reciprocal space for mosaic tilt and twisthklAlong with scattering vector | GhklThe magnitude of | increases linearly. The finite lateral size of the mosaic grains results in broadening parallel to the interface, inversely proportional to the average actual spatial dimension L, and to the scattering vector magnitude (Δ G)hkl2 pi/L) is irrelevant. The effects of tilt, twist and finite grain size will spin with the diffusion of the reflection hkl as measured by the skewed symmetry ω -scan as follows:
[Lee et al(2005)]. Where β denotes the integration width and the exponent n takes a value between 1 and 2, depending on the fit to Pseudo-Voigt (n ═ 1+ η -2) The contribution of Gauss η and Lorentz (1- η) (see Srikant et al (1997) appendix).
The peak broadening measured is then the combination of the mosaic broadening of the sample and the instrument function (no sample). As long as the latter is much narrower than the damascene broadening, it is possibleIt is ignored. Experimentally, a beta scan can be plotted in a modified Williamson-Hall plot (not shown) by measuring a series of omega scans of different order symmetric reflections 00ln·GnFor GnTo distinguish peak broadening effects due to lateral size and tilt. Slope of the line and the slope component (beta)tilt n) In this connection, the shift in the ordinate is related to the average grain size ((2. pi./L)n). Unfortunately, the commonly used Cu-ka radiation only gives symmetric 002 and 004 zinc blende GaN reflections, which greatly limits the accuracy, especially the accuracy of determining limited dimensions. FIG. 79 shows β in a conventional Williamson-Hall diagramhkl nFor this figure, lorentz broadening (n ═ 1) is usually employed, although the lorentz curve usually does not fit well with the measurement curve. The value of the finite size L is much larger than for a gaussian fit curve (n ═ 2), as indicated by Lee et al (2005). Since the X-ray intensity curve can be described empirically as a convolution of the gaussian function and the lorentzian function, the actual transverse finite size is within these two limit values, depending on the portion of the two curves. The gaussian and lorentzian ratios can be obtained from curve fitting in general, but they may be different for a series of different reflections. However, since the lorentz portion in such fits is typically small, mosaic sizes estimated from a pure gaussian fit can yield relatively good estimates.
The azimuthal angle spread around the surface normal due to mosaic distortion can be determined from off-axis reflections with large polar angles χ measured in a helical symmetric geometry. Ideally, one of the in-plane reflections (χ about 90 °) would be used, but these reflections typically exhibit only very low intensities and are often difficult to measure. For a (001) oriented zincblende GaN film, the 331 reflection (χ about 76.7 °) can be better used. Alternatively, the integrated widths of a series of different off-axis reflections extracted from an off-symmetric ω -scan can be extrapolated using equation (9) to determine the distortion component.
FIG. 80 shows such an extrapolation, we convert equation (9) to βhkl n·|Ghkl|nThis function (for n ═ 2) was fitted to the measured peak broadening of the optimal cubic GaN sample by using the slope and finite size values of the Williamson-Hall plot in fig. 79. The measured reflections are marked with circles and indicated by contour lines in reciprocal space betahkl·|GhklExtrapolated peak width and polar angle χ and scattering vector magnitude | G in |hklThe relationship of | is given. The contour lines represent constant peak widths. The curve at χ ═ 0 ° has been shown previously (fig. 79) and is affected only by the inclination and limited size of the mosaic blocks. As the polar angle χ increases, the broadening gradually increases, showing that the inlay twist 0.864 ° (χ ═ 90 °) is slightly higher than the tilt (χ ═ 0 °)0.755 °. This trend is not apparent since the tilt and twist are very similar, but for larger scattering vectors | GhklThis trend becomes more pronounced as the contribution of finite size to peak broadening decreases.
Density of defects
In general, the mosaic tilt and twist are believed to be related to threading dislocations formed at the grain boundaries of the thin film. Thus, by following the different mosaic tilt models discussed in the literature, XRD peak broadening is sometimes used to estimate defect density in thin films. According to these models, threading dislocation density D in well-oriented damascene filmsTDAnd betatilt/twistIn a certain proportion:
[Fewster(1989)]however, in a poorly oriented film, the grain orientation is randomly and strictly distributed between the Burgers vector and the line vector, with the threading dislocation density and beta2 tilt/twistIn a certain proportion:
[Dunn and Koch(1957)]. Where the parameter L is the average transverse finite size of the grains, and bTDIndicating that when the dislocation in zincblende GaN is complete, the Burgers vector value of the dislocation is
Compared with wurtzite-type GaN material, in which the threading dislocation line vector propagates mainly in [0001] c direction, threading dislocations in zincblende-type GaN extend in a plurality of <110> directions. Therefore, the above equation cannot distinguish edge type, mixed type, or screw type dislocations of zincblende type GaN. However, it is well known that the predominant type of threading dislocation in zincblende-like structures is the 60 ° type of total dislocation [ Blumenau et al (2000) ].
In an in-depth comparative study using XRD and Transmission Electron Microscopy (TEM) to estimate defect density in wurtzite GaN films, Metzger et al (1998) found that it matched the random distribution model (equation xy) well, even though the assumption of this model could not be met at all in oriented epitaxial films. Contrary to expectations, models of oriented damascene films show threading dislocation densities one order of magnitude lower than TEM estimated values. Lee et al (2005) concluded that similar conclusions were drawn and it is common to note that there is a large difference in dislocation density measured between TEM and XRD. In general, XRD appears to slightly overestimate threading dislocation density when the distortion component is used, and underestimates it [ Lee et al (2005) ], when broadening due to tilt is used. Furthermore, it should be noted that for very thin films, XRD also samples the tilt associated with GaN/SiC interface misfit dislocations. If the Burgers vector of dislocations is randomly oriented, the associated strain fields will tend to cancel as the film thickness increases, but if the Burgers vector is not random, the tilt will persist. The discussion shows that even for the more extensively studied measurements of wurtzite GaN, the current understanding still has some limitations and requires careful handling of XRD estimated defect densities. This is particularly true when comparing samples of different layer thicknesses.
Influence of layer thickness
In general, the intensity spread of the X-ray reflection is not constant, but decreases the full width at half maximum (FWHM) of the 002 reflection in the ω -scan as the film thickness increases, as shown in fig. 81. It is also evident that in low lattice mismatch substrates (e.g. 3C-SiC (lattice constant)Number SiC:thus, the compressibility was 3.4%) and MgO (lattice constant MgO:
Several reports in the literature (see reference in FIG. 81) show that the trend in FIG. 71 is due to the reduction of threading dislocation density with increasing film thickness As a result of threading dislocation reaction, but TEM has insufficient evidence [ As (2010); Kemper (2015); R ü sing (2016); Lischka (1997)]. Theoretical modeling predicts that threading dislocation density is inversely proportional to film thickness t [ Ayers (1995)]. Combining these models with the reflection broadening caused by mosaicing, it can be found that the intensity distribution is reduced by an extent t-1Or t-1/2Depending on whether it is an oriented film (equation (10)) or a powder sample (equation (11)). As can be seen from the dashed line in fig. 81, the experimental data did not follow the predicted trend. Instead, the observed attenuation is much weaker, following about t-1/3The dependency of (c). This can be explained by the fact that new penetrating dislocations can be generated when stacking faults react with each other, which, to our knowledge, is not taken into account in the current model. Furthermore, it should be considered that the model predicts that the threading dislocation density will decrease after a certain thickness, and XRD is an integral method that provides a weighted average of the entire layer thickness. It should also be taken into account that the width of the X-ray reflection naturally decreases with increasing number of scattering atoms and with increasing layer thickness. All this makes it difficult to compare the material quality of samples with different thicknesses.
Material parameters for strain analysis
The lattice parameters of zincblende type III nitrides have not been determined well experimentally because such films exhibit stacking disorder, undoubtedly high density of line defects and wurtzite inclusions, resulting in local strain variations and relatively broad reflections. Furthermore, most X-ray diffraction experiments on mainly zincblende GaN films have focused on phase purity analysis, rather than high resolution lattice parameter measurement.
We measured the zincblende GaN lattice parameters using a high resolution 2 theta-omega scan with 8 on-axis and off-axis reflections and a least squares fit to give a value of (4.50597 + -0.00038)
TABLE 2 lattice parameters and elastic constants of wurtzite and zincblende GaN, InN and AlN
*1:
*2:
*3:
*4Experiment of
*5That work (experiment)
*6Vurgaftman and Meyer (2003) recommendations
However, to our knowledge, the exact lattice parameters of zincblende InN and AlN, determined experimentally, are not mentioned in the literature. Therefore, in these cases it is necessary to derive from the perfect wurtzite lattice parameter awzAnd cwzThese values are derived. However, in wurtzite group III nitrides, the strong internal electric field causes the unit cell to deform from the ideal shape, its cwz/awzRatio ofIn fact, cwzIs generally less than idealA and awzIt is slightly greater than ideal and therefore, as can be seen from the values in table 2, the estimated zincblende parameters for nominally unstrained type III nitride may vary greatly. Due to awzRatio cwzLess affected by the distortion of the wurtzite unit cell, so that the parameter can be the zincblende phase azbProvides reasonable values for the natural lattice constant of (a). Alternatively, a slave cell volume may be used
Table 2 also contains the spring constant C for the zincblende group III nitrides described in Vurgaftman and Meyer (2003)11And C12And the method can be used for calculating stress and strain.
Strain of
During the growth of thin films on foreign substrates and during the growth of heterostructures with alloys of different composition, the thin films are subjected to varying stresses, which generally lead to elastic deformations of the crystal lattice. Such lattice strain has a significant impact on the physical properties and performance of semiconductor devices. Therefore, it is important to know and monitor these strains during device development. In the following sections we will discuss different sources of strain and describe how to measure the strain in zincblende GaN films.
Lattice mismatch strain
In an epitaxial thin film, when two lattice dimensions are forced to match each other, the lattice mismatch between the thin film and the underlying template produces biaxial in-plane strain. Three different states are commonly used to describe film deformation. When the crystal lattice of the thin film matches the size of the template crystal lattice at the common interface, the thin film will be fully deformed, and when the crystal lattice of the thin film is undeformed and has a natural size, the thin film will be fully relaxed. The state between the two extremes is called partial relaxation.
In the reciprocal space, the lattice-mismatched strain causes the Reciprocal Lattice Point (RLP) of the GaN thin film to be displaced from its intended position relative to the RLP of the substrate. The relative distance between the slice and the buffer peak can be measured in several separate omega-2 theta scans, or more commonly by collecting reciprocal space maps of asymmetric geometries. The latter generally better outlines the relationship between the X-ray reflections of the different layers, but for lattice mismatch strain assessment, correction of sample miscut by secondary scanning is required. Furthermore, it should be taken into account that the layer used as reference may also be affected by the substrate, which may reduce the accuracy of the method. Ideally, the substrate peak should be used as a reference, but for systems with large mismatches, there may be large separation of the reciprocal spaces.
The strain of the film in a certain direction is as follows:
wherein a is0Is a natural lattice constant, and aiAre measurement constants in the same direction. Since the material quality is generally relatively low in zincblende nitride materials, there is no accurate reference in the literature for the natural lattice parameters, as described in the previous section. For GaN, the experimentally determined values we provide in table 2 can be used. For other group III nitrides, we propose to use values derived from the wurtzite a parameter or wurtzite unit cell volume (see table 2), since the wurtzite lattice parameters are well known.
Assuming that the film is stress-free in the growth direction (generally labeled z), the strain in the growth direction of the (001) oriented film is given according to hooke's law:
wherein epsilonxAnd εyIs the strain in both in-plane directions, C11And C12Is the elastic constant of the material (see Table 2) [ Dunstan (1997)]. For isotropic plane strain (. epsilon.)x＝εy) The above formula can be further simplified. OthersThe strain relationship in direction differs from the above formula and is published in other documents, such as Dunstan (1997).
Thermal mismatch strain and growth induced strain
Small strains in epitaxial films arise from thermal mismatch between the substrate used and the epitaxial layer, or formation early in the growth. It is typically much less than the strain due to lattice mismatch, but it may be much larger than the residual mismatch strain in the partially relaxed film.
Since GaN has a larger coefficient of thermal expansion than SiC and Si [ Wahab et al (1994); la Via (2012); okada and Tokumura (1984)]Residual thermal strain causes the GaN film to develop tension at the interface with the substrate after cooling from the growth temperature. For a typical zincblende GaN growth temperature between 700 ℃ and 1000 ℃, the theoretical thermal strain is 1.1 × 10 when using a Si substrate-3To 1.6X 10-3In the meantime.
Growth-induced strain is due to coalescence of islands during nucleation on the substrate early in growth. Its amplitude is given by the minimum gap Δ between two islands and the average size of the islands in a particular in-plane direction:
[Hoffman(1976)]。
relative lattice parameter measurements as described above are not sufficient to accurately determine such small strains because resolution is typically low and the substrate itself may also be affected by strain. Instead, to analyze very small strains, a large set of high resolution 2 θ - ω scans of different reflections is required to make absolute measurements of the lattice parameters. The measured interplanar spacing d is then fitted using least squaresjMatching with the face spacing of the model crystal:
to improve the methodThe accuracy of (2) sometimes requires the use of a weighting factor WjE.g. 2 thetaj/Δ(2θj)[Roder etal(2006)]Or dj -2/Δ(dj -2)＝0.5·tan(θj)/Δ(2θj) (this work) for taking into account the measured value Δ (2 θ)j) Is not accurate.
In general, a suitable coordinate system must be chosen that better describes the geometry of the problem than the natural lattice. The following example illustrates this. Table 3 lists the 2 θ values for the different reflections, from a value at [110 ]]Measured on a zinc blende GaN film grown on a 3C-SiC/Si template obliquely cut in direction by 4 degrees. All reflected bragg angles that are tilted along the miscut direction hhl are much smaller than similarly reflected bragg angles that are tilted away from the miscut direction h-hl, indicating that the lattice sizes in these two directions are different. Thus, the natural lattice is slightly sheared in the growth plane. It can be simplified by using a new coordinate system x ', y', z '(where x' (y ') is obliquely cut parallel (perpendicular) to the sample and z' points in the growth direction). It should be noted that by transforming this coordinate, the new cell is √ 2 × √ 2 × 1 larger than the cell in the natural lattice. Using this method, a least squares fit (see above) together with bragg's law gives the unit cell size x ═ in the new coordinate system (6.39566 ± 6.7 × 10-4)
TABLE 3 reflection of the natural (hkl) and rotational (h ' k ' l ') coordinate systems for zincblende GaN samples, where 2 θ and delta (2 theta) from high resolution 2 theta omega scanning
Substrate miscut is known to cause strain relaxation in epitaxial films due to alignment threading dislocations [ Young et al (2010); chen et al (2007)]But with the zincblende GaN layer in the off-cut direction (ε)x') is less than in the vertical direction (εy') strain. Since we observe the opposite, we can rule out the relaxation mechanism of this sample. Instead, the results indicate that the strain anisotropy may be due to coalescence of small islands of different sizes in the two in-plane directions, which is observed in atomic force microscopy images.
Wafer curvature analysis
In heteroepitaxial films, stress above a certain level can be relieved by the formation of defects, or tensile surface stress can be relieved by the formation of cracks. Furthermore, the stress in the film can be reduced by bending the entire sample. This is typically the case in thick, moderately stressed epitaxial layers (e.g., template and buffer layers). Thermal strain also causes a large amount of wafer bow. This is especially a problem for large area templates with diameters up to 8 ", where even small bends can cause significant deviations in uniformity during growth and processing. Therefore, it is important to control and manage strain and wafer bow.
Can be measured by XRD on sample xjAt different positions of the light beam, angle omega of the symmetrically reflected incident light beamjThe bow of the wafer is determined. In the curved sample, the lattice planes also curve with the curvature of the wafer. Therefore, the angle of incidence needs to be corrected for different positions along the wafer diameter. Then, through ωjAnd xjTo obtain the wafer curvature and bend radius R:
[Inaba(2014)]. Due to reflection of the zincblende GaN epitaxial layerTypically relatively wide and therefore it is more appropriate to use a narrower symmetrical reflection of the underlying template. The resolution of the measurement can be further improved by using a larger range of measurement positions and using a beam mask to reduce the illuminated area of the sample surface. FIGS. 82 and 83 show the measured curvature of the 4' 3C-SiC/Si template, where the 002SiC reflections were used. The curvature of the wafer along the measurement direction can be easily determined in a graphical manner by linear interpolation of the measured incident beam angle. The positive (negative) slope corresponds to the concave (convex) shape of the wafer. FIGS. 82 and 83 show-51.5 km-1An example of convex curvature (R ═ 19.4m, respectively). It should be noted, however, that this technique measures the curvature of the plane of the substrate. If the substrate already contains a high density of dislocations or grain structures, the plane of the substrate may already be curved prior to layer growth, and thus the measured curvature may not accurately reflect the residual stress in the wafer. Thus, 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 was small and negligible.
Depth study
The effect of reaction pressure (as well as other parameters and conditions) on the growth of cubic zincblende GaN films was further investigated according to the experimental work reported above.
In summary, cubic zincblende GaN films were grown by Metal Organic Vapor Phase Epitaxy (MOVPE) on 3C-SiC/Si (001) templates and characterized using Nomarski optical microscopy and X-ray diffraction. Specifically, the surface morphology and material quality were evaluated as a function of low temperature nucleation layer thickness (3-44 nm in these experiments), epitaxial growth temperature (850-910 ℃ in these experiments), and V/III ratio (15-1200 in these experiments) at a reaction pressure of 100 torr. The reduction of the reaction pressure from 300 torr to 100 torr here is the main difference from the earlier results reported above. In this particular case, the window in which particularly suitable MOVPE growth conditions are determined is: the temperature is between 850 and 890 ℃, and the V/III ratio is between 38 and 150, so that a relatively smooth zincblende GaN film is obtained, wherein the wurtzite impurity content is less than 1%.
The effect of reducing the epitaxial layer reaction pressure from 300 torr to 100 torr on the sphalerite phase purity is shown in fig. 111 and 112. These figures show that under similar temperature and V/III ratio conditions, the wurtzite fraction is equally low or lower for the 100 torr data set. Thus, under reduced reaction pressure conditions, the MOVPE growth window is greatly broadened to obtain good sphalerite GaN material quality.
All samples were grown in a Thomas Swan 6X 2 "close coupled showerhead MOVPE reactor on a 3-SiC/Si template. The SiC template consists of an approximately 3 μm thick 3C-SiC layer on a Si (001) substrate, the thickness of which is between 0.75mm and 1mm, oriented [110 ]]The orientation difference of (2) is 4 deg.. For GaN growth, Trimethylgallium (TMG) and ammonia are used as precursors for Ga and N, respectively, while hydrogen is used as a carrier gas. The total gas flow was kept constant at 20 standard liters per minute (slm). The growth process includes high temperature thermal annealing of the substrate, followed by low temperature nucleation layer deposition, and finally appropriate growth of epitaxial layers at high temperature. The temperature is recorded by a Laytec EpiTT in-situ optical monitoring system and is calibrated aiming at the Al/Si eutectic wafer. The thermal annealing step of the template was performed at 960 ℃ in a mixture of hydrogen and 3slm ammonia. For sample sets a and B (see below), a GaN Nucleation Layer (NL) was grown to a thickness of 44nm at 600 ℃, 500 torr, and a V/III ratio of 720. The growth pressure of the epitaxial layer was maintained at 100 torr while the thickness of the epitaxial layer was kept constant at 300 nm. Two sample sets were prepared, where the variables were epitaxial layer growth temperature (sample set a) and V/III ratio (sample set B). For sample set a, the constant V/III ratio in the gas phase was 76 and the epitaxial layer growth temperature was varied between 850 and 910 ℃. For sample set B, the V/III ratio was varied between 15 and 1200 by varying the ammonia flow rate at a constant TMG flow rate of 145. mu. mol/min during the growth of the epitaxial layers at 880 ℃. Using silane (50ppm SiH4At H2In) doping all GaN epitaxial layers with Si to 1018cm-3An intermediate nominal concentration. Finally, a third sample set (sample set C) consisted of 300nm thick GaN epitaxial layers grown at a thickness between 3 and 44nm at 875 deg.C, 100 Torr and 76V/IIIOn the altered low temperature nucleation layer.
Using a light source equipped with Cu-Ka1
Sample set A-temperature series
The first sample set to be discussed (sample set a) consisted of six samples, in which the growth temperature of the epitaxial layers was varied between 850 and 910 ℃ with a variation interval of 10 to 15 ℃. In this series of experiments, the V/III ratio was kept constant at 76, representing the median value within the range of values explored in sample set B (see below). As shown in the Nomarski optical micrographs shown in fig. 88-93, the surface morphology of the samples grown at temperatures below 895 ℃ had elongated features or striations. The stripes are aligned in the [1-10] direction, i.e., in the direction perpendicular to the substrate miscut, which represents in-plane anisotropy. As the temperature increases, the length of the elongated features shortens to several microns. At 895 ℃ and above, the surface becomes more grainy and rougher with increasing temperature.
XRD analysis of sample group a showed that wurtzite fraction measured perpendicular to miscut gradually increased with increasing growth temperature, but remained below 1%, as shown in fig. 94. The measurement results obtained parallel to the beveling showed no wurtzite inclusions in this direction and are therefore not shown in the figure.
Sample set B-V/III ratio series
The second sample set (sample set B) consisted of eight samples, where the outer layer growth temperature was kept constant at 875 ℃ (i.e. the median value of set a), while the V/III ratio varied between 15 and 1200, with a factor of about 2 between each value. The Nomarski optical micrographs in fig. 95-102 show that the surface morphology changes from granular to striated to rough as the V/III ratio increases.
XRD phase analysis of sample set B showed that wurtzite inclusions increased slightly from 0% at a V/III of 15 to 1% at a V/III of 300 when measured perpendicular to the miscut, as shown in FIG. 103. At the highest V/III values of 600 and 1200, wurtzite fraction increased more rapidly to 3% and finally to 11%. When measured parallel to the beveling, no significant wurtzite phase was found in all samples and was therefore omitted from the figure.
Sample set C-NL thickness series
The third sample set (sample set C) consisted of five samples, in which the nucleation layer thickness was increased from 3nm to 44nm, with an increase of 2 per step, while the epitaxial layer was grown under conditions of 875 ℃ at a constant temperature, 100 torr at a V/III ratio of 76. Fig. 104-108 show the surface morphology change as a function of NL thickness. Of these five samples, the only seemingly unusual sample was one that grew the thinnest (3nm) NL, which showed large pits in the epitaxial layers. In addition to this, the surface striations appear to be roughened as the NL thickness increases.
XRD phase analysis of sample group C (see fig. 109) showed that the sample with NL thickness of 3nm had the largest wurtzite fraction of about 3%, while the other samples had wurtzite phase less than 1%. The integrated intensity of the 002 rocking curve for sample set C is shown in fig. 110. The varying peak intensities indicate that the material quality improves with increasing NL thickness, and that the peak intensities saturate at NL thicknesses of 22nm or greater.
Discussion sample groups A, B and C
The structural data (temperature change at constant V/III ratio of 76) for sample set a indicates that growth temperatures below 895 ℃ resulted in relatively smooth film surfaces (see fig. 88-93) with wurtzite fractions of less than 1% (see fig. 94). At higher growth temperatures, the surface was slightly degraded and the wz-GaN fraction was close to 1%. Thus, at a V/III of 76 and a reaction pressure of 100 Torr, a preferred growth temperature is less than about 890 ℃.
Sample set B used a constant growth temperature of 875 ℃ (i.e., in the most favorable temperature range), while the V/III ratio varied between 15 and 1200, with a factor of about 2 between each value. At the low and high ends of this range (see fig. 95-102), a more granular surface is shown. As long as the Nomarski micrographs can be ranked by nature, it can be observed that the surface of the sample grown between V/III ratios of 38 to 150 is flattest. Despite the differences in surface morphology of the samples, contamination of wz-GaN was still very small for the sample with a V/III of 300 (see FIG. 103). When the V/III ratio exceeds 300, wurtzite contamination rises rapidly and the surface morphology is slightly deteriorated.
It is worth noting that sample set A grew at a V/III ratio of 76, which falls within the most favorable V/III ratio range obtained from the findings of sample set B. Likewise, sample set B was grown at an epitaxial layer growth temperature of 875 ℃ which falls within the most preferred temperature range obtained from the findings of sample set a. Therefore, in this MOVPE growth window, there was little change in the surface and material properties of the zincblende GaN film.
Studies on the thickness of the low temperature nucleation layer (sample set C) showed that the preferred thickness is about 22 nm. Using a thinner GaN NL results in a reduced material quality, while a thicker NL coarsens the surface morphology.
In summary, considering the collective data of surface morphology and phase purity for sample groups a through C, the preferred MOVPE growth conditions for zincblende GaN epitaxial layers at a constant pressure of 100 torr are: at temperatures between 850 and 890 ℃ and V/III ratios between 38 and 150, relatively smooth films can be obtained within this parameter window, with wurtzite contamination of less than 1%. The preferred thickness of NL is about 22 nm.
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While the invention has been described in conjunction with the exemplary embodiments outlined above, many equivalent modifications and variations will be apparent to those skilled in the art given this disclosure. Accordingly, the exemplary embodiments of the invention set forth above are considered to be illustrative and not restrictive. Various changes may be made to the described embodiments without departing from the spirit and scope of the invention.
All references mentioned above and/or listed below are incorporated herein by reference.
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