Method for producing group iii nitride semiconductor

Abstract

To provide a method for producing a Group III nitride semiconductor that is capable of forming a Group III nitride semiconductor having a good crystal quality.SOLUTION: A method for producing a Group III nitride semiconductor includes: an ammonia treatment step of supplying a gas containing ammonia to a surface of a substrate 10 containing sapphire; a thermal cleaning step of subjecting the surface of the substrate 10 to heat treatment in a hydrogen-dominated atmosphere, after the ammonia treatment step; a nitriding step of nitriding the surface of the substrate 10 by supplying a gas containing ammonia to the surface of the substrate, after the thermal cleaning step; and a crystal nucleus layer forming step of forming a crystal nucleus layer 11 on the substrate by generating nuclei 11A of GaN, AlGaN or AlN on the substrate, after the nitriding step.SELECTED DRAWING: Figure 1

Description

[Technical field]

[0001] The present invention relates to a method for producing a Group III nitride semiconductor. [Background technology]

[0002] The ultraviolet wavelength of solid-state light-emitting devices using III-nitride semiconductors corresponds to a wavelength band of approximately 210 to 400 nm. In particular, UVC (wavelength 100 to 280 nm) is known to be effective for sterilization and disinfection, and there is a growing demand for III-nitride semiconductor LEDs that emit ultraviolet light with an emission wavelength corresponding to UVC. Ultraviolet LEDs are configured by forming an AlN layer on a sapphire substrate, and then laminating an n-layer, light-emitting layer, and p-layer made of AlGaN on the AlN layer.

[0003] In order to form a high-quality AlN layer, it is necessary to grow AlN after nitriding the surface of the sapphire substrate. Non-Patent Document 1 describes a method of supplying ammonia to the sapphire substrate to nitride the sapphire substrate, and then forming AlN on the sapphire substrate. It also describes that the nitriding process is performed while increasing the temperature from a low temperature to 1200°C. [Prior art documents] [Non-patent literature]

[0004] [Non-Patent Document 1] Mitsuru Funato, Mami Shibaoka and Yoichi Kawakami, "Heteroepitaxy mechanisms of AlN on nitridated c- and a-plane sappire", JOURNAL OF APPLIED PHYSICS 121,085304(2017) Summary of the Invention [Problem to be solved by the invention]

[0005] However, in the method of Non-Patent Document 1, the nitride film (AlN) formed by the nitriding process is a crystal having a mixture of +c and -c planes, and there are problems such as a deterioration in the quality of the crystal formed on the AlN, and difficulty in controlling the occurrence of cracks and surface roughness.

[0006] The present invention has been made in view of the above background, and has an object to provide a method for manufacturing a Group III nitride semiconductor that is capable of forming a Group III nitride semiconductor having good crystal quality. [Means for solving the problem]

[0007] One aspect of the present invention is an ammonia treatment step of supplying ammonia-containing gas to a surface of a sapphire substrate; a thermal cleaning step of heat-treating the surface of the substrate in a hydrogen-dominated atmosphere after the ammonia treatment step; a nitriding process step of supplying an ammonia-containing gas to a surface of the substrate after the thermal cleaning process, and nitriding the surface of the substrate; a crystal nucleus layer forming step of generating nuclei of GaN, AlGaN or AlN on the substrate after the nitriding treatment step to form a crystal nucleus layer; The present invention relates to a method for producing a Group III nitride semiconductor having the above structure. Effect of the Invention

[0008] In the above embodiment, the sapphire substrate is treated with ammonia before the thermal cleaning, and the surface of the sapphire substrate is nitrided after the thermal cleaning, thereby improving the crystal quality of the Group III nitride semiconductor formed on the sapphire substrate. [Brief description of the drawings]

[0009] [Figure 1] 1A to 1C are diagrams illustrating a manufacturing process of a Group III nitride semiconductor according to the first embodiment. [Diagram 2] 2A to 2C are schematic diagrams showing layer configurations at each growth stage in the first embodiment. [Diagram 3]Graph showing changes in temperature and pressure over time in the first embodiment. [Figure 4] Graph showing the relationship between the temperature range of ammonia treatment and the half-width of X-ray diffraction lines. [Diagram 5] 1 is a graph showing the relationship between the ammonia treatment time and the half-width of the diffraction line of X-ray diffraction. [Figure 6] 1 is a graph showing the relationship between the starting temperature of ammonia treatment and the half-width of X-ray diffraction lines. [Figure 7] 1 is a graph showing the relationship between the hydrogen flow rate in thermal cleaning and the half-width of the diffraction line in X-ray diffraction. [Figure 8] 1 is a graph showing the relationship between nitriding treatment time and the half-width of the diffraction line of X-ray diffraction. [Figure 9] 5A to 5C are schematic diagrams showing layer configurations at each growth stage in the second embodiment. [Figure 10] Graph showing changes in temperature and pressure over time in the second embodiment. [Figure 11] A graph comparing the half-widths of the (002) diffraction lines. [Figure 12] A graph comparing the half-widths of the (102) diffraction lines. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0010] The method for manufacturing a Group III nitride semiconductor includes an ammonia treatment step of supplying a gas containing ammonia to a surface of a substrate made of sapphire, a thermal cleaning step of heat-treating the surface of the substrate in a hydrogen-dominated atmosphere after the ammonia treatment step, a nitridation treatment step of supplying a gas containing ammonia to the surface of the substrate to nitride the surface of the substrate after the thermal cleaning step, and a crystal nucleus layer formation step of generating nuclei of GaN, AlGaN or AlN on the substrate to form a crystal nucleus layer.

[0011] In the method for producing a Group III nitride semiconductor, the ammonia treatment step may be carried out at a temperature of 200° C. or more and 900° C. or less, thereby making it possible to further improve the crystal quality of the crystal nucleus layer.

[0012] In the method for producing a Group III nitride semiconductor, the ammonia treatment step may be carried out for a period of time ranging from 1 second to 60 seconds, thereby making it possible to further improve the crystal quality of the crystal nucleus layer.

[0013] In the method for producing a Group III nitride semiconductor, the ammonia treatment step may be performed with an ammonia partial pressure of 0.001 atm or more and 0.1 atm or less and a gas linear velocity of 0.01 m / min or more and 300 m / min or less, thereby further improving the crystal quality of the crystal nucleus layer.

[0014] In the method for producing a Group III nitride semiconductor, the thermal cleaning step may be carried out at a temperature higher than 900° C. This can further improve the crystal quality of the crystal nucleus layer.

[0015] In the method for producing a Group III nitride semiconductor, the nitriding step may be carried out at a temperature equal to or lower than the temperature in the thermal cleaning step for 50 seconds or longer, thereby further improving the crystal quality of the crystal nucleus layer.

[0016] The method for producing a Group III nitride semiconductor may further include a low-temperature three-dimensional growth layer formation step of growing AlGaN or AlN from the nuclei at a temperature lower than that of the crystal nucleus layer formation step after the crystal nucleus layer formation step, and merging crystals from adjacent nuclei to form a low-temperature three-dimensional growth layer, and a high-temperature three-dimensional growth layer formation step of growing AlGaN or AlN from the low-temperature three-dimensional growth layer at a temperature higher than that of the low-temperature three-dimensional growth layer formation step but equal to or lower than that of the crystal nucleus layer formation step after the low-temperature three-dimensional growth layer formation step, to form a high-temperature three-dimensional growth layer. AlGaN or AlN with good crystal quality can be formed.

[0017] (Embodiment 1) 1. Method for manufacturing Group III nitride semiconductors A description will be given of a method for manufacturing a Group III nitride semiconductor in embodiment 1. The Group III nitride semiconductor is formed by MOCVD, and the source gases used are, for example, ammonia as the nitrogen source gas, TMGa (trimethylgallium) or TEGa (triethylgallium) as the Ga source gas, and TMAl (trimethylaluminum) as the Al source gas, and hydrogen or nitrogen as the carrier gas.

[0018] In the following description of the manufacturing method, reference will be made to Figures 1 to 3 as appropriate. Figure 1 is a flow chart showing an initial step (step before the formation of nuclei 11A) in the manufacturing process of a Group III nitride semiconductor. Figure 2 is a diagram showing a schematic diagram of the layer configuration at each growth stage up to the formation of two-dimensionally grown layer 13. Figure 3 is a graph showing the time changes in temperature and pressure, and the supply timing of various gases.

[0019] 1-1. Ammonia treatment process First, a substrate 10 made of sapphire with the c-plane as the principal surface is prepared. The principal surface may be the a-plane. In addition, the substrate may have an off angle of 0.1 to 2 degrees in the m-axis or a-axis direction.

[0020] Next, the substrate 10 is placed in an MOCVD apparatus, and the apparatus is evacuated and purged with hydrogen (or nitrogen) to reduce the pressure in the MOCVD apparatus (period T11 in FIG. 3). FIG. 3 shows a case where the pressure is 4 kPa.

[0021] Thereafter, an ammonia treatment is performed on the surface of the substrate 10 (step S1 in FIG. 1, period T12 in FIG. 3). The ammonia treatment is a process in which ammonia is added to hydrogen to form a mixed gas, which is supplied into the MOCVD apparatus, and the substrate 10 is heated to place the surface of the substrate 10 in an ammonia atmosphere. FIG. 3 shows an example of the ammonia treatment performed while increasing the temperature from 200° C. to 700° C. at a pressure of 4 kPa.

[0022] The temperature (substrate temperature) at the start of the supply of ammonia may be equal to or higher than room temperature (for example, 20° C.) and lower than 900° C. In the first embodiment, the ammonia treatment is performed while increasing the temperature, and the maximum temperature is preferably 900° C. or lower.

[0023] In the first embodiment, the ammonia treatment is performed while increasing the temperature, but the temperature may be constant as long as it is between room temperature and 900° C. If the ammonia treatment is performed at a constant temperature, it is more preferable to perform it at a temperature of 20 to 900° C. If the ammonia treatment is performed while increasing the temperature, the starting temperature is preferably 20 to 900° C. and the ending temperature is preferably 200 to 900° C.

[0024] The time for the ammonia treatment is preferably 1 to 60 seconds, which can further improve the crystal quality of the crystal nucleus layer 11. The time is more preferably 10 to 50 seconds.

[0025] Since ammonia is corrosive, it is believed that the ammonia treatment causes some modification to the surface of the substrate 10. In addition, since the treatment is performed at a low temperature of 900° C. or less, the decomposition efficiency of ammonia is low, and the surface of the substrate 10 may not be sufficiently nitrided.

[0026] In the ammonia treatment, the ammonia partial pressure is preferably 0.001 to 0.1 atm, and the gas flow rate is preferably 0.01 to 300 m / min. Within these ranges, the crystal quality of the crystal nucleus layer 11 formed on the substrate 10 can be further improved. The ammonia partial pressure is more preferably 0.01 to 0.05 atm, and the flow rate is more preferably 0.1 to 150 m / min.

[0027] The pressure of the ammonia treatment may be normal pressure, but reduced pressure is preferred, for example, 1 to 90 kPa, preferably 1 to 50 kPa, and more preferably 1 to 20 kPa, and may be the same pressure as that of the thermal cleaning in the next step.

[0028] Furthermore, although a mixed gas of hydrogen and ammonia is used in the first embodiment, an inert gas such as nitrogen may be mixed in place of or in addition to hydrogen.

[0029] 1-2. Thermal cleaning process Next, the surface of the substrate 10 is thermally cleaned (step S2 in FIG. 1, period T13 in FIG. 3). Thermal cleaning is performed by stopping the supply of ammonia and supplying only hydrogen while continuing to heat the substrate 10. This thermal cleaning removes impurities attached to the surface of the substrate 10, and flattens the surface of the substrate 10 (see FIG. 2(a)). It is believed that this thermal cleaning causes impurities such as oxides and Si to separate from the surface of the substrate 10, and the Al2O3 surface of the sapphire becomes an Al-rich surface.

[0030] Thermal cleaning is preferably performed at a temperature higher than 900°C. For example, it is performed at a temperature of 1150 to 1300°C for 1 second to 15 minutes. Thermal cleaning may be performed while increasing the temperature or at a constant temperature. Fig. 3 shows an example of thermal cleaning performed while increasing the temperature from 700°C to 1190°C at a pressure of 4 kPa.

[0031] The temperature of the thermal cleaning is preferably higher than that of the crystal nucleus layer forming process described later. The temperature of the thermal cleaning is more preferably in the range of 1170 to 1250°C, and even more preferably 1180 to 1220°C. The atmosphere may be dominated by hydrogen, for example, a mixed gas containing 80% or more by volume of hydrogen. The mixed gas may be, for example, a mixed gas of hydrogen and nitrogen.

[0032] The pressure may be normal pressure, but reduced pressure is preferred, for example, 1 to 90 kPa, preferably 1 to 50 kPa, and more preferably 1 to 20 kPa.

[0033] The flow rate of hydrogen gas or hydrogen and nitrogen mixed gas on the substrate 10 may be 5 to 500 m / min (meters / minute), preferably 5 to 300 m / min, and more preferably 5 to 150 m / min. The flow rate in the mass flow controller is preferably 1 to 30 slm. Cracks in the two-dimensional growth layer 13 can be reduced.

[0034] 1-3. Nitriding process Next, a nitriding process is performed on the surface of the substrate 10 (step S3 in FIG. 1, period T14 in FIG. 3). The nitriding process is a process in which the Al-rich surface of the substrate 10 is nitrided by lowering the temperature than in thermal cleaning and supplying ammonia in addition to hydrogen. This nitriding process forms an AlN layer (not shown) of one to several monolayers, which serves as the starting point for nucleation, on the surface of the substrate 10. FIG. 3 shows an example of the nitriding process performed at a pressure of 4 kPa and a temperature of 1155° C.

[0035] The temperature of the nitriding treatment is, for example, 1100 to 1220°C, and may be the same as that of the subsequent crystal nucleus layer forming process. The pressure in the nitriding treatment is the same as that in the ammonia treatment or thermal cleaning. The ammonia partial pressure and the gas flow rate may be the same as those in the ammonia treatment. The time for the nitriding treatment is preferably 50 seconds or more. This can reduce cracks in the two-dimensional growth layer 13. The time is more preferably 70 to 300 seconds.

[0036] 1-4. Crystal nucleus layer formation process Next, ammonia and TMAl are supplied at the same temperature as in the previous nitriding process or at a lower temperature than in the previous nitriding process, and AlN nuclei 11A are generated on the substrate 10 to form a crystal nucleus layer 11 (Fig. 2(b), period T15 in Fig. 3). Fig. 3 shows an example where the temperature is 1155°C. The pressure is the same as in the previous process. Conventionally, the nuclei 11A are formed at about 1000°C, but in embodiment 1, the temperature is set higher than this to enhance the surface migration of the source atoms and form the nuclei 11A larger than in the past. This makes it possible to achieve high quality nuclei 11A made of AlN.

[0037] The lattice mismatch between the nuclei 11A and the substrate 10 made of sapphire causes compressive stress in the crystal nucleus layer 11. However, because the lattice mismatch between the nuclei 11A and the substrate 10 is large, the stress is relieved within a few nm, and the stress at the interface between the substrate 10 and the nuclei 11A is small.

[0038] It is preferable that the crystal nucleus layer 11 is grown to a thickness of 1 to 100 nm, so that the nuclei 11A can be made sufficiently large. It is also preferable that the size and density of the nuclei 11A are within the following ranges.

[0039] The density of nucleus 11A is 3×10 11 / cm -2 or less. The size of core 11A (average diameter in plan view) is preferably 20 to 50 nm. Here, the diameter is the diameter when core 11A is converted into a circle of the same area. By setting the size of core 11A within this range, tensile stress can be sufficiently reduced. Furthermore, the variation in size of core 11A (the difference between the maximum diameter and the average diameter, and the difference between the average diameter and the minimum diameter) is preferably 10 nm or less.

[0040] It is important that the nuclei 11A are large. When the crystal nuclei combine to form a flat film, tensile stress is generated on the surface. Therefore, if the nuclei size is small and the density is high, the stress generated on the surface when they combine to form a flat film is also large. Therefore, it is necessary to reduce the nuclei density. To achieve this, it is effective to increase the nuclei size.

[0041] A low nucleus density has another advantage. Threading dislocations are formed at the interface where the nuclei 11A join together. A low nucleus density also reduces the number of interfaces where the nuclei 11A join together. Therefore, by reducing the nucleus density, it is possible to reduce the formation of threading dislocations, and a high-quality crystal film can be formed.

[0042] The preferred growth temperature for the crystal nucleus layer 11 is 1100 to 1200°C. By setting the temperature within this range, the nuclei 11A can be made sufficiently large. This is because the surface migration of the raw material atoms can be enhanced. More preferably, it is 1125 to 1190°C, and even more preferably, it is 1150 to 1180°C. It may be the same temperature as that of the previous nitriding process. By making the nuclei 11A larger, the nucleus density is reduced. Therefore, the interface between the nuclei 11A is also reduced. Tensile stress may occur in the film when the nuclei 11A combine and become flat, but by reducing the nucleus density as described above, the stress that occurs after the nuclei 11A combine can be reduced.

[0043] The growth rate is preferably 5 to 100 nm / min. By increasing the growth rate, the diameter of the nuclei 11A can be increased. The growth rate is more preferably 10 to 80 nm / min, and further preferably 20 to 70 nm / min.

[0044] The V / III ratio is preferably 5 to 500. By setting the V / III ratio in this range, the growth rate can be controlled within the above range. It is more preferably 5 to 400, and further preferably 5 to 300.

[0045] It should be noted that although there is a possibility that -c-plane AlN is formed on the nitrided substrate 10, it can be made into a +c-plane by using growth conditions such as a low V / III ratio and a high growth rate as the growth conditions for the nuclei 11A. Usually, a crystal layer with a +c-plane is preferable, and even if +c and -c are mixed in the crystal nuclei layer 11, the proportion of -c is small, and +c becomes dominant during the growth process, and finally, a crystal layer having a surface only with a +c-plane is sufficient. Another possibility is that, at 1100°C or higher, AlON may be formed on the surface of AlN. It is known that AlON reverses polarity, and it is considered that even if -c-plane AlN is formed on the nitrided substrate 10, the polarity is reversed by AlON, resulting in a +c-plane.

[0046] 1-5. Low temperature three-dimensional growth layer formation process When the crystal nucleus layer 11 grows to a predetermined thickness, the supply of TMAl is stopped and the temperature is cooled to a lower temperature than that in the formation process of the crystal nucleus layer 11 (period T16 in FIG. 3). Then, when the temperature is cooled to the growth temperature of the low-temperature three-dimensional growth layer 12A, the supply of TMAl is resumed. As a result, crystals are three-dimensionally grown from the nuclei 11A of the crystal nucleus layer 11 as starting points, and the crystals grown from adjacent nuclei 11A are combined to form the low-temperature three-dimensional growth layer 12A (FIG. 2(c), period T17 in FIG. 3). FIG. 3 shows the case of a temperature of 1000° C. as an example. The pressure is the same as in the previous process. By combining the nuclei 11A, threading dislocations can be reduced, and a high-quality low-temperature three-dimensional growth layer 12A can be formed.

[0047] Here, since the growth temperature of the low-temperature three-dimensional growth layer 12A is set to a lower temperature than the growth temperature of the nuclei 11A, the stress acting on the layers up to the low-temperature three-dimensional growth layer 12A becomes compressive stress due to the difference in linear expansion coefficient with the substrate 10. Generally, cracks occur in the growth layer when tensile stress is applied. However, compressive strain occurs in the layers up to the low-temperature three-dimensional growth layer 12A due to the difference in linear expansion coefficient, so cracks can be suppressed.

[0048] Furthermore, by lowering the temperature, the low-temperature three-dimensional growth layer 12A can be grown slowly, and the coalescence of the low-temperature three-dimensional growth layer 12A growing from the nuclei 11A can also proceed slowly. Tensile stress may occur when the three-dimensional nuclei 11A coalesce to become a two-dimensional flat film, but since the low-temperature three-dimensional growth layer 12A is grown slowly as described above, even if the low-temperature three-dimensional growth layer 12A on the nuclei 11A coalesces to become a continuous film, the three-dimensional surface shape is maintained due to the low-temperature growth, and cracks due to tensile stress occurring in the film can also be suppressed.

[0049] The low-temperature three-dimensional growth layer 12A is preferably grown thicker than the crystal nucleus layer 11, for example, until the thickness reaches 200 to 500 nm. By sufficiently merging the nucleus 11A as a starting point so that the nucleus shape changes to a film shape, the threading dislocation density can be reduced. Furthermore, the thicker the low-temperature three-dimensional growth layer 12A is, the more uneven the three-dimensional surface can be. This makes it possible to reduce the threading dislocations that propagate to the growth layer surface by bending the threading dislocations due to the lateral growth promoted in the next and subsequent steps. The thickness is more preferably 250 to 400 nm, and even more preferably 250 to 350 nm.

[0050] The growth temperature of the low-temperature three-dimensionally grown layer 12A is preferably 900 to 1100° C. By setting the temperature at 900° C. or higher, impurities are less likely to enter the crystal, and light absorption can be reduced. The growth temperature is more preferably 950 to 1050° C., and further preferably 975 to 1025° C.

[0051] The growth rate of the low-temperature three-dimensional growth layer 12A is set to be slower than that of the crystal nucleus layer 11. By slowing down the growth rate and slowly coalescing starting from the nuclei 11A, dislocations can be efficiently reduced. In addition, the formation of tensile stress that occurs when the nuclei 11A coalesce to form a continuous film can be alleviated. The growth rate is preferably 2 to 20 nm / min, more preferably 2 to 15 nm / min, and even more preferably 2 to 10 nm / min.

[0052] The V / III ratio of the low-temperature three-dimensional growth layer 12A is set to be larger than that of the crystal nucleus layer 11, and is preferably set to 500 to 2000. The growth rate can be controlled within the above range, more preferably 750 to 1750, and further preferably 1000 to 1500.

[0053] 1-6. High temperature three-dimensional growth layer formation process When the low-temperature three-dimensional growth layer 12A, which has grown from the nuclei 11A as starting points, has fully coalesced and grown to a predetermined thickness, the growth temperature is set higher than in the formation process of the low-temperature three-dimensional growth layer 12A but lower than the growth temperature in the formation process of the crystal nucleus layer 11, and the high-temperature three-dimensional growth layer 12B is formed by further growing on the low-temperature three-dimensional growth layer 12A (FIG. 2(d)). The pressure is the same as in the previous process.

[0054] The high-temperature three-dimensional growth layer 12B is a layer grown to smoothly transition from the three-dimensional growth of the low-temperature three-dimensional growth layer 12A to the two-dimensional growth of the two-dimensional growth layer 13. The high-temperature three-dimensional growth layer 12B is a three-dimensional growth in which the lateral growth is faster than that of the low-temperature three-dimensional growth layer 12A, and the growth temperature of the high-temperature three-dimensional growth layer 12B is higher than that of the low-temperature three-dimensional growth layer 12A, so that the lateral growth is faster than that of the low-temperature three-dimensional growth layer 12A, and the crystals can be further integrated. As a result, the threading dislocations propagating toward the surface can be reduced by bending the threading dislocations laterally, and a crystal film with a low threading dislocation density can be obtained.

[0055] Here, the temperature at which the high-temperature three-dimensional growth layer 12B is formed is higher than the temperature at which the low-temperature three-dimensional growth layer 12A is formed, and tensile stress is generated in the crystal layer up to the high-temperature three-dimensional growth layer 12B during growth of the low-temperature three-dimensional growth layer 12A due to the difference in linear expansion coefficient with the substrate 10. For this reason, it is preferable to grow the high-temperature three-dimensional growth layer 12B not as a completely flat film, but as a three-dimensional surface partially having pits or facets, thereby easing the tensile stress generated on the surface.

[0056] The high-temperature three-dimensionally grown layer 12B is preferably grown thicker than the low-temperature three-dimensionally grown layer 12A, for example, to a thickness of 750 to 2000 nm. In this range, crystals can be further combined to further reduce the threading dislocation density. A more preferred thickness is 1000 to 1750 nm, and even more preferred is 1250 to 1500 nm.

[0057] The growth temperature of the high-temperature three-dimensional growth layer 12B is preferably 1050 to 1200° C. By setting the temperature in this range, the growth mode can be efficiently converted from three-dimensional growth to two-dimensional growth. The temperature is more preferably 1075 to 1175° C., and further preferably 1100 to 1150° C.

[0058] Moreover, the growth rate of the high-temperature three-dimensional growth layer 12B is faster than that of the low-temperature three-dimensional growth layer 12A and is lower than that of the crystal nucleus layer 11, and is preferably 5 to 50 nm / min. By setting the growth rate in this range, the crystals are slowly combined, and the change in the tensile stress generated on the crystal layer surface is made gentle, thereby suppressing the occurrence of cracks when the crystals are combined. The growth rate is more preferably 10 to 40 nm / min, and even more preferably 15 to 30 nm / min.

[0059] The V / III ratio of the high-temperature three-dimensionally grown layer 12B is smaller than that of the low-temperature three-dimensionally grown layer 12A and is equal to or larger than that of the crystal nucleus layer 11, and is preferably 100 to 1000. The growth rate can be controlled within the above range. It is more preferably 150 to 700, and further preferably 200 to 500.

[0060] 1-7. Two-dimensional growth layer formation process When the high-temperature three-dimensional growth layer 12B has grown to a predetermined thickness, TMGa is supplied in addition to ammonia and TMAl at the same temperature as or higher than that in the formation process of the high-temperature three-dimensional growth layer 12B to form a two-dimensional growth layer 13 made of AlN or AlGaN containing Ga (see FIG. 2(e)). The pressure is the same as in the previous process.

[0061] If Ga is supplied during AlN growth, the Ga atoms, which have a higher surface migration than Al atoms, promote lateral growth in the growth mode of AlN, resulting in two-dimensional growth. Therefore, the crystal can be flattened by the two-dimensional growth layer 13. In addition, the supply amount of Ga is increased continuously and stepwise over time. This allows the lateral growth rate to be increased slowly, easing the change in tensile stress that occurs when the crystals coalesce, thereby suppressing the occurrence of cracks in the two-dimensional growth layer 13.

[0062] The amount of Ga supplied is controlled by the molar ratio of TMGa to TMAl. The molar ratio of TMGa to TMAl is preferably 0.05 to 0.5, more preferably 0.08 to 0.4, and further preferably 0.1 to 0.3.

[0063] Although the ratio of TMGa is very high, Ga is not actually incorporated into AlN at the above ratio. This is because Ga atoms evaporate more easily than Al atoms on the AlN surface, so the amount of Ga atoms actually incorporated into the AlN crystal is only a few percent even when the molar ratio of the source gas is 0.3. The higher the growth temperature, the more Ga atoms evaporate preferentially, so even fewer Ga atoms are incorporated into the AlN crystal. Although it depends on the growth temperature and other conditions, the amount of Al grown within the above molar ratio range is 1-x Ga x The Ga solid phase ratio of N is about x = 0.01 to 0.1%.

[0064] By applying a continuous and stepwise gradient within the above molar ratio range, the Ga supply rate is increased over time. The two-dimensional growth layer 13 changes from AlN to a mixed crystal of AlN and GaN (Al 1-x Ga x N). Here, the Ga composition is preferably 0.01 to 0.1.

[0065] The two-dimensionally grown layer 13 is preferably grown to a thickness equal to or greater than that of the high-temperature three-dimensionally grown layer 12B, for example, to a thickness of 750 to 2000 nm. Within this range, the surface of the two-dimensionally grown layer 13 can be sufficiently flattened. For example, the surface roughness RMS can be set to 0.5 to 5 nm. A more preferred thickness is 1000 to 1750 nm, and even more preferred is 1250 to 1500 nm.

[0066] The growth temperature of the two-dimensionally grown layer 13 is preferably 1100 to 1200° C. If the temperature is within this range, the two-dimensionally grown layer 13 can be sufficiently flattened. The temperature is more preferably 1120 to 1190° C., and further preferably 1140 to 1180° C. The temperature may be the same as that in the formation process of the crystal nucleus layer 11 and the formation process of the high-temperature three-dimensionally grown layer 12B.

[0067] The growth rate of the two-dimensional growth layer 13 is preferably 5 to 50 nm / min, which is higher than the growth rate of the high-temperature three-dimensional growth layer 12B and lower than the growth rate of the crystal nucleus layer 11. If the growth rate is within this range, the growth rate is sufficiently slow, and the change in tensile stress due to the flattening of the crystal can be made gentle, and the occurrence of cracks can be suppressed. It is more preferably 10 to 40 nm / min, and even more preferably 15 to 30 nm / min.

[0068] The V / III ratio of the two-dimensional growth layer 13 is set to be equal to or lower than the V / III ratio of the high-temperature three-dimensional growth layer 12B and equal to or higher than the V / III ratio of the crystal nucleus layer 11, and is preferably set to 50 to 500. The growth rate can be controlled within the above range. It is more preferably set to 100 to 400, and further preferably set to 150 to 300.

[0069] Thereafter, a semiconductor layer may be continuously grown on the two-dimensional growth layer 13 to fabricate an element such as a light-emitting device, or the fabrication may be terminated at the stage where the two-dimensional growth layer 13 is grown, and the substrate may be used as a template substrate. Also, the substrate 10 may be peeled off by laser lift-off (LLO) and used as a template.

[0070] In the first embodiment, the period from the formation of the nuclei 11A to the formation of the two-dimensionally grown layer 13 is controlled by the growth temperature, but it may be controlled by the growth rate. The growth rate can be controlled by, for example, the growth temperature and the V / III ratio.

[0071] As described above, the method for producing a Group III nitride semiconductor according to the first embodiment can improve the crystal quality of the crystal nucleus layer 11. The reason for the improvement in crystal quality is not clear, but can be inferred as follows. When the ammonia treatment is performed before the thermal cleaning, some modification is considered to occur on the surface of the substrate 10 because the ammonia is corrosive. Therefore, it is considered that the effect of removing impurities, oxides, or particles from the surface of the substrate 10 during the thermal cleaning is improved. As a result of the further reduction of impurities and the like on the surface of the substrate 10, good quality AlN can be formed on the surface of the substrate 10 during the nitriding treatment after the thermal cleaning, and the crystal quality of the crystal nucleus layer 11 formed on the AlN is also improved.

[0072] 2. Experimental Results Next, various experimental results relating to the first embodiment will be described.

[0073] Experiment 1 After the substrate 10 was subjected to an ammonia treatment, thermal cleaning, and nitriding treatment in this order, a crystal nucleation layer 11, a low-temperature three-dimensional growth layer 12A, a high-temperature three-dimensional growth layer 12B, and a two-dimensional growth layer 13 were formed on the substrate 10 in this order from the substrate 10 side to prepare a sample. Three types of samples were prepared by changing the temperature range in the ammonia treatment. For comparison, a similar sample was also prepared except that the ammonia treatment was not performed.

[0074] The temperature and pressure were as shown in the graph of Figure 3. The crystal nucleus layer 11 had a growth rate of 30 nm / min and a thickness of 50 nm, the low-temperature three-dimensional growth layer 12A had a growth rate of 4 nm / min and a thickness of 250 nm, the high-temperature three-dimensional growth layer 12B had a growth rate of 18 nm / min and a thickness of 1300 nm, and the two-dimensional growth layer 13 had a growth rate of 18 nm / min and a thickness of 1400 nm. The molar ratio of TMGa to TMAl in the two-dimensional growth layer 13 was continuously changed from 13% to 30%.

[0075] The rocking curve measurement was carried out in X-ray diffraction of the two-dimensionally grown layer 13 of the prepared sample, and the full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line was obtained. Figure 4 is a graph showing the relationship between the temperature range of the ammonia treatment and the full width at half maximum. The temperature range of the ammonia treatment was three times: when the temperature was raised from 200°C to 400°C, when the temperature was raised from 200°C to 900°C, and when the temperature was raised from 200°C to 1190°C.

[0076] As shown in FIG. 4, the half-width of the (102) diffraction line was smaller in all temperature ranges than when no ammonia treatment was performed. In addition, when the temperature was raised from 200°C to 1190°C, the half-width was slightly larger than in other temperature ranges. As for the (002) diffraction line, the half-width became larger as the temperature range became wider and the maximum temperature value became higher. As a result, it was found that a temperature range of 200°C to 400°C or 200°C to 900°C is suitable for improving the crystal quality of the two-dimensionally grown layer 13. In addition, it was found from this result that ammonia treatment at 900°C or lower is preferable.

[0077] Experiment 2 The ammonia treatment time was varied to prepare samples in the same manner as in Experiment 1. The ammonia treatment time was the time it took to raise the temperature from 200°C, and the heating rate was set to 3 to 5°C / s. The full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line were then obtained in the same manner as in Experiment 1.

[0078] Figure 5 is a graph showing the relationship between the treatment time of the ammonia treatment and the half-width. As shown in Figure 5, for the (102) diffraction line, the half-width significantly decreased with increasing treatment time up to about 60 seconds, and when the treatment time exceeded 60 seconds, the half-width gradually increased with increasing treatment time. Furthermore, for the (002) diffraction line, the half-width increased with increasing treatment time. From this result, it was found that the treatment time of the ammonia treatment is preferably 1 to 60 seconds in order to improve the crystal quality.

[0079] Experiment 3 The temperature at the start of the ammonia treatment was varied in various ways to prepare samples in the same manner as in Experiment 1. The temperature rise time for the ammonia treatment was 20 seconds, and the temperature rise rate was the same as in Experiment 2. The full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line were then obtained in the same manner as in Experiment 1.

[0080] Fig. 6 is a graph showing the relationship between the temperature at the start of the ammonia treatment and the half-width. As shown in Fig. 5, for the (102) diffraction line, the half-width decreased with increasing temperature up to about 700°C, and the half-width increased with increasing temperature beyond that point. For the (002) diffraction line, the half-width increased gently with increasing temperature up to about 700°C, and the half-width increased significantly with increasing temperature beyond that point. As a result, it was found that the temperature at the start of the ammonia treatment is preferably 100 to 700°C in order to improve the crystal quality.

[0081] Experiment 4 The flow rate of hydrogen gas in the thermal cleaning was varied to prepare samples in the same manner as in Experiment 1. The full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line were then obtained in the same manner as in Experiment 1.

[0082] FIG. 7 is a graph showing the relationship between the flow rate of hydrogen gas and the half-width during thermal cleaning. As shown in FIG. 7, for the (102) diffraction line, the half-width increased with an increase in the flow rate. Also, for the (002) diffraction line, the half-width decreased with an increase in the flow rate. Also, the number of cracks in the two-dimensionally grown layer 13 decreased as the flow rate decreased. As a result, it was found that in order to improve the crystal quality, a hydrogen flow rate of 1 to 30 slm during thermal cleaning is preferable.

[0083] Experiment 5 The nitriding time was varied to prepare samples in the same manner as in Experiment 1. The full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line were then measured in the same manner as in Experiment 1.

[0084] Fig. 8 is a graph showing the relationship between the processing time of the nitriding process and the half-width. As shown in Fig. 8, for the (102) diffraction line, the half-width decreased with increasing processing time up to about 120 seconds, and after that, the half-width increased with increasing processing time. For the (002) diffraction line, the half-width gradually increased with increasing processing time. Furthermore, when the processing time was 50 seconds or more, the number of cracks in the two-dimensional growth layer 13 decreased as the processing time was increased, and when the processing time was 150 seconds or more, no cracks occurred. As a result, it was found that a processing time of 70 seconds or more is preferable for improving crystal quality.

[0085] (Modification of the first embodiment) As the nucleus 11A in the first embodiment, AlGaN or GaN may be used instead of AlN. If GaN or AlGaN is used, the difference in lattice constant with the sapphire substrate 10 becomes larger, and the strain relaxation at the interface between the substrate 10 and the nucleus 11A becomes larger. In addition, when the nucleus 11A is made of AlGaN or GaN, the lattice constant of the low-temperature three-dimensional growth layer 12A formed thereon has a higher Al composition than the nucleus 11A, so that the low-temperature three-dimensional growth layer 12A grown on the nucleus 11A receives tensile strain. This is because the growing crystal is formed in conformity with the lattice constant of the base. Therefore, it is possible to relax the compressive strain generated after the low-temperature three-dimensional growth layer 12A due to the thermal stress of the substrate and the crystal layer generated when the growth is completed and the temperature becomes room temperature.

[0086] In addition, in the strain relaxation at the interface between the substrate 10 and the nuclei 11A, the nuclei 11A are formed discretely, not in a film shape, so that the occurrence of cracks is suppressed. In addition, the low-temperature three-dimensional growth layer 12A formed on the crystal nuclei layer 11 receives tensile stress from the crystal nuclei layer 11 due to the difference in lattice constant, but the occurrence of cracks due to the tensile stress is suppressed because of the three-dimensional growth. As a result of the above, the threading dislocation density of the two-dimensional growth layer 13 can be reduced.

[0087] Furthermore, when the nuclei 11A are made of GaN or AlGaN, the surfactant effect of Ga allows the nuclei 11A to be made larger, and the nucleus density to be reduced. As described above, a low nucleus density allows the tensile stress generated on the surface when the nuclei combine to form a flat film to be reduced. In addition, the number of meeting surfaces where the nuclei 11A combine is reduced, which reduces the formation of threading dislocations, and allows the formation of a high-quality crystal film.

[0088] The core 11A may be GaN doped with Al. The core 11A is preferably AlGaN rather than GaN, and the Al composition is preferably 60% or more and 98% or less. This allows the strain to be further relaxed and the core 11A to be larger. More preferably, it is 60% or more and 98% or less.

[0089] When the nucleus 11A is GaN or AlGaN containing Ga, the substrate 10 can be easily separated by laser lift-off. In other words, since the band gap energy is smaller than that of AlN, it is easy to make the crystal nucleus layer 11 absorb laser light. In addition, since Ga has a very low melting point compared to Al, the crystal nucleus layer 11 becomes liquid when decomposed by irradiation with laser light. Therefore, when the crystal nucleus layer 11 contains Ga, the substrate 10 can be easily separated. Thus, when the nucleus 11A is GaN or AlGaN containing Ga, the crystal nucleus layer 11 is effective as a sacrificial layer for laser lift-off.

[0090] When the cores 11A are made of GaN or AlGaN, the size of the cores 11A is preferably 20 to 100 nm. By setting the size of the cores 11A in this range, the tensile stress can be sufficiently reduced. In addition, the variation in size of the cores 11A (the difference between the maximum diameter and the average diameter, and the difference between the average diameter and the minimum diameter) is preferably 20 nm or less.

[0091] Moreover, when the nucleus 11A is made of GaN or AlGaN, the thickness of the crystal nucleus layer 11 is preferably 5 to 100 nm. Since the nucleus 11A is larger than when the nucleus 11A is made of AlN, the crystal nucleus layer 11 is also thicker. As a result of the larger nucleus 11A, the quality of the crystal layer formed after the nucleus 11A can be further improved. A more preferable thickness of the crystal nucleus layer 11 is 5 to 50 nm. When the nucleus is AlGaN and the Ga composition is high, the crystals become more mobile by annealing, and it becomes easier to enlarge the nucleus by solid phase growth. The size of the nucleus immediately before growing the three-dimensional growth layer can be increased by annealing, and the density can also be reduced.

[0092] In addition, when the core 11A is made of GaN or AlGaN, the density of the core 11A is 3×10 11 / cm -2 The nucleus 11A is larger than that in the case of AlN, so the density of the nucleus 11A is also smaller than that in the case of AlN. A more preferable nucleus density is 1.5×10 11 / cm -2 Less than 1×10, more preferably11 / cm -2 In addition, when the nuclei 11A are GaN or AlGaN, the nuclei 11A can be enlarged by annealing. That is, the nuclei density can be increased to 1×10 11 / cm -2 It can be as follows:

[0093] In the first embodiment, the low-temperature three-dimensional growth layer 12A is made of AlN, but it may be made of AlGaN instead of AlN. However, from the viewpoint of strain relaxation and crystallinity, it is preferable that the difference between the Al composition of the low-temperature three-dimensional growth layer 12A and the Al composition of the crystal nucleus layer 11 is 40% or less.

[0094] In the first embodiment, the high-temperature three-dimensional growth layer 12B is AlN, but may be AlGaN instead of AlN. However, from the viewpoint of strain relaxation and crystallinity, it is preferable that the difference in Al composition between the high-temperature three-dimensional growth layer 12B and the low-temperature three-dimensional growth layer 12A is 30% or less.

[0095] In addition, in the first embodiment, the two-dimensional growth layer 13 is AlN, but it may be AlGaN instead of AlN. However, from the viewpoint of strain relaxation and crystallinity, it is preferable that the difference between the Al composition of the two-dimensional growth layer 13 and the Al composition of the high-temperature three-dimensional growth layer 12B is 20% or less.

[0096] In addition, the crystal layers below the two-dimensional growth layer 13 must have an Al composition that does not absorb light if the device formed on the two-dimensional growth layer 13 is a light-emitting device. If the device to be formed on the two-dimensional growth layer 13 is an ultraviolet light-emitting LED, the Al composition of the two-dimensional growth layer 13 is preferably the same as or greater than the Al composition of the n-type layer of the LED.

[0097] For the above reasons, most of the device structure is often formed with an Al composition lower than that of the two-dimensional growth layer 13. In order to relieve the strain in the device structure, it is preferable to change the lattice mismatch difference stepwise. It is preferable that the Al composition of the layers subsequent to the three-dimensional growth layer 12, excluding the crystal nucleus layer 11, is reduced stepwise. In this way, the strain on the device structure can be relieved.

[0098] The lower the Al composition of the nucleus 11A, the more the lattice constant changes from AlN to that of GaN. That is, the larger the in-plane lattice constant. Therefore, the three-dimensional growth layer 12 formed on the nucleus 11A and having a larger Al composition than the Al composition of the nucleus 11A receives tensile strain from the nucleus 11A. The lower the Al composition of the nucleus 11A, the larger the in-plane lattice constant of the three-dimensional growth layer 12 tends to be. The three-dimensional growth layer 12 receives tensile strain during growth, increasing the possibility of crack formation, so there is an optimal value for the Al composition difference between the nucleus 11A and the three-dimensional growth layer 12, as described above. The tensile strain formed after the three-dimensional growth layer 12 during growth becomes compressive strain due to the difference in thermal expansion coefficient with the substrate 10 made of sapphire when growth is completed and the temperature is raised to room temperature. However, as in the case of growth, the larger the lattice constant of the nucleus 11A, the more the compressive strain generated after the three-dimensional growth layer 12 at room temperature is alleviated.

[0099] Furthermore, when the nucleus 11A is made of GaN or AlGaN, the thickness of the two-dimensionally grown layer 13 is preferably 0.5 to 5 μm. The surface of the two-dimensionally grown layer 13 can be sufficiently flattened. For example, the surface roughness RMS can be set to 0.5 to 2 nm. The thickness of the two-dimensionally grown layer 13 is more preferably 1 to 3 μm.

[0100] Even when the nucleus 11A is made of GaN or AlGaN, the threading dislocation density of the two-dimensionally grown layer 13 is 5×10 11 cm -2 or less, and good quality crystals can be obtained. Furthermore, the full width at half maximum (FWHM) of the (002) diffraction line in rocking curve measurement in X-ray diffraction of the two-dimensionally grown layer 13 can be, for example, 100 to 300 arcsec, and the full width at half maximum of the (102) diffraction line can be, for example, 300 to 600 arcsec. This is because the full width at half maximum can be sufficiently reduced by reducing the threading dislocation density.

[0101] When the nuclei 11A are made of GaN or AlGaN, the growth conditions such as the growth temperature and growth rate of each layer can be in the same range as in the first embodiment.

[0102] (Embodiment 2) A method for manufacturing a Group III nitride semiconductor in embodiment 2 will be described. In the description of the manufacturing method, Figs. 9 and 10 will be referred to as appropriate. Fig. 9 is a diagram showing a schematic configuration of each layer at each growth stage up to the formation of a two-dimensional growth layer 23. Fig. 10 is a graph showing the change in temperature over time and the supply timing of various gases.

[0103] In the method for producing a Group III nitride semiconductor in the second embodiment, first, the surface of the substrate 10 is treated with ammonia (period T21 in FIG. 10), and then thermally cleaned (FIG. 9(a), period T22 in FIG. 10). The various conditions for the ammonia treatment and thermal cleaning are the same as those in the first embodiment.

[0104] Next, the material is cooled to a temperature for performing the next nitriding process (period T23 in FIG. 10), and then the nitriding process is performed (period T24 in FIG. 10). The temperature is, for example, 350 to 650°C. FIG. 10 shows a case where the temperature is 390°C as an example. The nitriding process may be performed during cooling. The nitriding process may be performed under the same conditions as the nitriding process in the first embodiment, except for the temperature.

[0105] Next, ammonia and TMAl are supplied at the same temperature as the nitriding process in the previous step to generate nuclei 21A of AlN on the substrate 10 to form a crystal nucleus layer 21 (FIG. 9(b), period T25 in FIG. 10). Note that the temperature may be different from that of the nitriding process as long as it is in the range of 350 to 650°C.

[0106] Next, the temperature is raised to the growth temperature of the three-dimensional growth layer 22 (period T26 in FIG. 10), and the nuclei 21A of the crystal nucleus layer 21 are annealed to promote solid-phase growth and change into a columnar shape. The temperature is raised to, for example, 1100 to 1200° C. Then, ammonia and TMGa are supplied to three-dimensionally grow GaN on the crystal nucleus layer 21, thereby forming the three-dimensional growth layer 22.

[0107] Next, the temperature is increased to a level higher than that of the three-dimensional growth layer 22, and GaN is grown two-dimensionally on the three-dimensional growth layer 22 to form the two-dimensional growth layer 23. Increasing the growth temperature promotes migration of the source material, and the two-dimensional growth can be achieved by speeding up the lateral growth. Instead of increasing the temperature, the V / III ratio may be increased to speed up the lateral growth and achieve two-dimensional growth.

[0108] In the second embodiment, GaN is grown as the three-dimensional growth layer 22 and the two-dimensional growth layer 23, but any III-nitride semiconductor may be used. However, AlGaN with a low Al composition is preferable. For example, AlGaN with an Al composition of 20% or less. GaN is particularly preferable as in the second embodiment. Unlike AlGaN or AlN with a high Al composition, the lateral growth of AlGaN or GaN with a low Al composition is easy to control by temperature or V / III ratio. This is because Ga migration is faster than Al. Therefore, even if the three-dimensional growth layer has large unevenness, it can be easily flattened by the two-dimensional growth layer.

[0109] As described above, in the method for producing a Group III nitride semiconductor in the second embodiment, the ammonia treatment is performed before the thermal cleaning and the nitridation treatment is performed after the thermal cleaning, so that the crystal quality of the crystalline nucleus layer 21 can be improved even when the crystalline nucleus layer 21 is formed at a low temperature. Therefore, the crystal quality of the Group III nitride semiconductor, particularly GaN, formed on the crystalline nucleus layer 21 can be improved.

[0110] Next, various experimental results according to the second embodiment will be described. After the substrate 10 was subjected to an ammonia treatment, thermal cleaning, and nitriding treatment in this order, a crystal nucleus layer 21, a three-dimensional growth layer 22, and a two-dimensional growth layer 23 were formed on the substrate 10 in this order from the substrate 10 side to prepare a sample (hereinafter referred to as an example). For comparison, a sample was prepared in the same manner as the example except that the ammonia treatment was not performed (hereinafter referred to as a comparative example). The temperature and pressure were as shown in the graph of FIG.

[0111] The rocking curves in X-ray diffraction of the two-dimensionally grown layer 23 of the prepared sample were measured to determine the full width at half maximum (FWHM) of the (002) diffraction line and the (102) diffraction line. Figures 11 and 12 are graphs comparing the full width at half maximum between the example and the comparative example, where Figure 11 shows the full width at half maximum of the (002) diffraction line and Figure 12 shows the full width at half maximum of the (102) diffraction line.

[0112] As shown in FIG. 11, there was no significant difference between the half-width of the (002) diffraction line in the example and the comparative example. On the other hand, the half-width of the (102) diffraction line in the example was smaller than that in the comparative example. As a result, it was found that the crystal quality can be improved even when the crystal nucleus layer 21 is formed at a low temperature by performing an ammonia treatment before thermal cleaning and a nitriding treatment after thermal cleaning. Although the effect of improving the half-width is smaller than that in the first embodiment, it was confirmed that the same effect can be obtained even when the growth temperature of the nucleus 21A is low. From the comparison between the first embodiment and the second embodiment, it is possible that the effect of the substrate surface treatment by the ammonia treatment, the thermal cleaning, and the nitriding treatment is higher when high-quality nuclei are formed at a high temperature. [Explanation of symbols]

[0113] 10: Substrate 11, 21: Crystal nucleation layer 11A, 21A: Nucleus 12A: Low temperature three-dimensional growth layer 12B: High temperature three-dimensional growth layer 22: Three-dimensional growth layer 13, 23: Two-dimensional growth layer

Claims

1. an ammonia treatment step of supplying ammonia-containing gas to a surface of a sapphire substrate; a thermal cleaning step of heat-treating the surface of the substrate in a hydrogen-dominated atmosphere after the ammonia treatment step; a nitriding process step of supplying an ammonia-containing gas to a surface of the substrate after the thermal cleaning process, and nitriding the surface of the substrate; a crystal nucleus layer forming step of generating nuclei of GaN, AlGaN or AlN on the substrate after the nitriding treatment step to form a crystal nucleus layer; The present invention relates to a method for producing a Group III nitride semiconductor having the above structure.

2. The method for producing a Group III nitride semiconductor according to claim 1 , wherein the ammonia treatment step is carried out at a temperature of 20° C. or more and 900° C. or less.

3. The method for producing a Group III nitride semiconductor according to claim 1 , wherein the ammonia treatment step is carried out for a period of time ranging from 1 second to 60 seconds.

4. 2. The method for producing a Group III nitride semiconductor according to claim 1, wherein the ammonia treatment step is performed with an ammonia partial pressure of 0.001 atm or more and 0.1 atm or less and a gas linear velocity of 0.01 m / min or more and 300 m / min or less.

5. The method for producing a Group III nitride semiconductor according to claim 1 , wherein the thermal cleaning step is carried out at a temperature higher than 900° C.

6. The method for producing a Group III nitride semiconductor according to claim 1 , wherein the nitriding step is carried out for 70 seconds or longer at a temperature equal to or lower than that of the thermal cleaning step.

7. a low-temperature three-dimensional growth layer formation step in which, after the crystal nucleus layer formation step, AlGaN or AlN is grown from the nuclei at a temperature lower than that of the crystal nucleus layer formation step, and crystals from adjacent nuclei are combined to form a low-temperature three-dimensional growth layer; The method for producing a Group III nitride semiconductor according to any one of claims 1 to 6, further comprising, after the low-temperature three-dimensional growth layer formation step, a high-temperature three-dimensional growth layer formation step of growing AlGaN or AlN from the low-temperature three-dimensional growth layer at a temperature higher than that of the low-temperature three-dimensional growth layer formation step and equal to or lower than the temperature of the crystal nucleus layer formation step to form a high-temperature three-dimensional growth layer.

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