Intrinsic defects in non-irradiated silicon carbide crystals

A comprehensive study of the intrinsic defects in sublimation-grown SiC crystals, depending on the growth conditions and thermal annealing is carried out. Complexes of the intrinsic defects including carbon vacancy (V C ) and impurities atoms are found in the Si-rich SiC crystals grown by physical vapor transport at low temperatures below 2200 °C. Similar defects are also observed in the SiC crystals irradiated with high-energy particles. Intrinsic defects in grown SiC crystals are characterized by high thermal stability, which is associated with the presence of active metastable clusters. Experimental evidence for the presence of the active clusters in the wide temperature range (up to 2600 °C) is presented. It is shown that intrinsic defects can be also introduced in the SiC crystal by high-temperature diffusion from the p-type epitaxial layer. Paramagnetic defects in SiC are considered a material platform for sensing, quantum photonics, and information processing at ambient conditions.


Introduction
Silicon carbide (SiC), together with diamond, gallium nitride, and aluminum nitride, is a typical representative of the wide band-gap semiconductor materials, characterized by high bonding energy.The undoubted advantages of SiC are its exceptionally high thermal stability and chemical resistivity.
SiC -based semiconductor devices can operate at elevated temperatures, up to 1000-1300 °C, which is 500-800 °C higher than in the case of Si or GaAs.Valuable qualities of SiC are high breakdown voltages (an order higher than for silicon), high electron saturation rate, and good thermal conductivity [1].The availability of a large number of SiC polytypes differ greatly in their intrinsic properties signifi cantly expanding the potential capabilities of SiC.This combination of the unique properties of SiC makes this material very attractive for creating particularly stable high-temperature, high-frequency, and power devices, as well as optoelectronic systems designed to work in extreme conditions, such as space, aviation, and mining.Signifi cant progress in silicon carbide semiconductor electronics is facilitated by the fact that to date, the growth of large (up to 6 inches in diameter) SiC bulk crystals used for the manufacture of devices in optoelectronics and power electronics has been realized.
The parameters of the devices are signifi cantly affected by intrinsic defects introduced into the crystal during growth, doping, or irradiation with high-energy particles.To date, a lot of theoretical and experimental material on the research of intrinsic defects in SiC has accumulated [2][3][4][5][6] with a signifi cant focus on the spin and optical properties of these defects with the purpose of using them for the development of quantum technologies [2][3][4].The nature of many centers introduced in SiC by irradiation with high-energy particles (electrons, neutrons.and ions) is established.The most important defect centers are silicon and carbon vacancies, antisite, and various complexes, including both native point defects and impurities [4].Indeed, deep-level defects such as silicon vacancies, carbon vacancies, and their structures are nowadays considered

Experimental
In this paper, we study the SiC crystals grown by the sublimation sandwich method (SSM) [7] which is one of the modifi cations of PVT growth.The basic principles of the SSM method are described in Ref [8].The important advantages of the SSM are very wide temperature (1700-2600 °C) and pressure (10 -2 -10 4 Torr) growth ranges.High pure and heavy doped bulk crystals and epitaxial layers of different SiC polytypes may be grown by this method.Growth of Si-rich and C-rich SiC crystals is also possible [8].
Various groups of SiC samples differing in the preparation conditions are investigated in this work.The fi rst group comprised SiC crystals of different polytypes (4H, 6H, 15R, and others) grown in graphite containers at high temperatures (Tg > 2400-2600 °C) (HT SiC).The second group includes samples of 6H-SiC grown at low temperatures of 1700-2100 °C (LT SiC).Some samples of this group are grown from a source enriched with silicon or in tantalum crucibles (Si-rich LT SiC).
The SiC samples of the third group are grown in an excess of carbon.We fi nd [8] that C-rich SiC crystals can also be grown by introducing group 4 impurities (Ge, Sn, Pb) into the growth zone.Impurities Sn, and Pb have a very low solubility in SiC (below 10 15 cm -3 ), making it possible to grow high pure SiC crystals enriched with carbon.The properties of the crystals are given in Table 1.
The majority of the samples were of the n-type conductivity due to doping with Nitrogen impurities.Donor concentration is calculated from the Hall effect and electron paramagnetic resonance (EPR) measurements.The concentration of acceptor impurities is found by the secondary ion mass spectrometry (SIMS) and neutron-activation analysis (NAA).High-doped p-type SiC (Al) samples (Na = 10 20 -10 21 cm -3 ) were also used.
Both n-and p-type SiC crystals were grown by the sublimation sandwich method.The dislocation density in the samples was found by X-Ray topography.X-Ray anomalous transmission method was used.The threading dislocation density was found by chemical etching in KOH melt at 500 °C.Nitrogen, Boron Aluminum, Boron 3×10 16 -10 20 (1.5-50) × 10 16  10 20 -10 21 (2-4) ×10 16   LT SiC Nitrogen, Boron Aluminum (1.2-200) ×10 16  (1-60) ×10 16  (6-10) ×10 15   C-rich SiC 2300-2400 6H N-type 10 3  Nitrogen, Boron (6-40) ×10 16  (2-35) ×10 16   Note: HT SiC: SiC crystals grown at High Temperature; LT SiC: SiC crystals grown at Low Temperature Citation: Mokhov  technique is described in detail in Ref [8].In the purest SiC samples, a characteristic luminescence with the D 1 spectrum occurs, which is observed usually in the crystals irradiated with high-energy particles, and is caused by an intrinsic defect [9].The study of optical absorption in LT SiC reveals a number of features not related to impurities.These features are similar to those found in irradiated samples [10] but are not observed in n-type HT 6H-SiC samples grown at temperatures as high, as 2500-2600 °C.

Evidence of the presence of native defects in as-grown crystals
Positron annihilation spectroscopy (PAS): Girka, et al. [11] studied vacancy defects in as-grown LT-and HT-SiC crystals by PAS.The positron lifetime measurements in various SiC samples are shown in Figure 2.
The longest positron lifetimes,  B = 167 ps, are observed in Si-rich LT samples (Figure 2).Since the PAS method is sensitive mainly to vacancy defects, we can assume that positrons are trapped at the vacancy defects or their complexes with impurity atoms.The average positron lifetime () in n-type HT 6H-SiC samples (Figure 2) is identical within the experimental error limits, irrespective of the donor impurity concentration.The corresponding spectra are well described by a single decay exponential with a time constant  B = (157±3) ns, where  B is the positron lifetime in the defect-free matrix.
In heavily p-type Al-doped samples, the positron lifetimes are higher than in the n-type, and  increases with increasing concentration of the acceptor impurity.The lowest values of  are obtained in the С-rich SiC (Figure 2).It means that C-rich and HT 6H-SiC samples contain the minimum concentration of the vacancy-type centers.
The ionization energy of native defects in SiC is generally high (E i > 100 meV) [12].At the typical measurement temperature (T = 300 K), therefore, the fraction of neutral vacancies in an n-type sample is not expected to be less than 30-50%.Addition-ally, the data show that increasing the concentration of electrically active impurities (nitrogen or aluminum) to the maximum solubility levels (10 20 -10 21 cm -3 ) also does not induce a signifi cant increase in the positron lifetime.Thus, the positrons in crystals of the HT SiC are annihilated primarily in  a defect-free matrix with a characteristic time constant  B = 157 ps.To estimate the density of vacancies, we use the Brandt trapping model [13].From this model, we obtain the vacancy density (CV) in the LT Si-rich crystals about (3 -5)×10 17 cm -3 .
High-temperature diffusion: Additional information about the high-temperature state of defects is obtained using a comparative study of boron diffusion in the SiC samples grown under different conditions.The distribution of boron impurities in SiC samples grown under various conditions is shown in Figure 3.
It is seen that the maximum diffusion mobility of boron atoms was observed in the LT-samples grown in excess Si vapor.
On the contrary, the minimum diffusion depth was noted in the HT and C-rich 6H-SiC samples.The pre-annealing of LT SiC samples at temperatures as high as 2450 °C led to a decrease in the diffusion depth in these samples (curve 3).An analysis of the data in Figure 3 shows that the diffusion coeffi cient of boron in LT SiC and HT samples differs by more than an order of magnitude.The bulk branch of the diffusion profi le at low diffusion temperatures (T dif ) is described by the standard erfc function, and at high T dif it becomes stepwise.The fastdiffusing state is known to be the (B-V C ) pair [14].Therefore, these experimental data confi rm the increased concentration of V C in LT SiC.Defects introduced by diffusion are metastable and are partially annealed at high temperatures A similar procedure is used for a comparative estimation of the vacancy density in the different SiC polytypes [15].These results are also presented in Attempts to detect vacancy centers by PAS mean that crystals grown in a relative excess of carbon or in samples with high doping levels by donor impurities have been fruitless.However, the latter method allowed us to establish that silicon vacancies (V Si ) carry a negative charge [16] and are typifi ed by a comparatively long positron lifetime ( B = 260 ps).Consequently, these vacancies should be very effi ciently detected by the PAS.However, V Si -related defects are not detected in the as-grown SiC, which obviously indicates a very low V Si density in such crystals.

Complexes of native defects in as-grown SiC crystals
D 1 -center: Intrinsic defects in as-grown crystals are also detected when studying the luminescent properties of LT SiC samples.The most interesting feature is the characteristic luminescence with the D 1 spectrum (D 1 PL), which was previously detected in the crystals irradiated with high-energy particles.Firstly, Vodakov, et al [17] showed that (D 1 PL) can be observed not only in the irradiated samples but also directly in the as-grown crystals.D 1 luminescence was detected in pure LT Si-rich samples grown at a high growth rate (Vg > 1.2 mm/h) [17, Figure 1].
An increase in the concentration of donor or acceptor impurities leads to a signifi cant decrease in the intensity of the D 1 PL.This indicates that the D 1 spectrum is caused by recombination at the center, which includes only its intrinsic states and does not include Al, B, and N impurities.
On the contrary, in the HT 6H-SiC samples, only impurity  (2,3,4).Boron diffusion distribution was captured using the method of track autoradiography (TA).It is evident that the diffusion rate is higher in the LT (Si-rich) sample, signifi cantly higher than in the HT samples.At low temperatures of diffusion, the bulk branch of boron distribution is described by the standard erfc-function.At temperatures above 2400°C, the distribution takes on a step-like character, which is associated with the partial decomposition of rapidly diffusing associates (B-Vc).In order to verify this conclusion, a series of experiments in which intrinsic defects were introduced into pure SiC crystals by diffusion from a highly doped p-type SiC(Al) epitaxial layer [17,18] were conducted.It is assumed that the p-type SiC contains a higher concentration of carbon vacancies V C , compared with the n-type SiC [5,6].The intensity of the D 1 PL of these samples via annealing temperature is shown in Figure 6, demonstrating that the D 1 PL appears only after annealing.With increasing radiation doses (f) D 1 PL occurs at higher T a : (820, 1620 K respectively for F = 10 17 , 10 19 cm -2 ) (Figure 6).It is seen, that luminescence intensity and annealing temperature of the D 1 PL is higher in the LT than in the HT samples (Figure 6).It is noted that annealing temperature of D 1 PL in highly neutron-irradiated LT SiC samples is about of 2000 °C that much lower than in the same as-grown crystals.

D-center (deep boron) -(EPR, DLTS, Phl measurements):
Boron is widely used as an acceptor dopant in SiC and creates two levels in the band gap, corresponding to two different centers.The shallower center with an ionization energy E i = 0.35-0.39eV, is usually observed in electrical measurements in crystals doped with boron during HT-growth.This center is the isolated boron atom substituting a silicon in the SiC lattice.
According to our EPR measurements [2], the second center with E i = 0.7 eV is (B Si -V C ) -pair, also known as D-center.This center attracts considerable interest, primarily because it is responsible for room-temperature luminescence (RT PL) with the peak in the yellow part of the spectrum in the case of the 6H-SiC (E m ~ 2.05 -2.15 eV).
Vodakov, et al. [20] found that the intensity of this RT PL strongly depends on the growth temperature.The highest PL  intensity was found in LT 6H-SiC samples (Figure 7), around one order of magnitude higher than in the HT SiC.It is possible to increase the intensity of the luminescence in HT 6H-SiC samples by boron diffusion, provided that the concentration of the donor nitrogen impurity N D > 1×10 18 cm -3 (Figure 7).
The intensity of the room temperature (RT) PL depends on the boron concentration.Nevertheless, regardless of the doping method, a higher RT PL intensity is achieved in LT SiC.
The difference in RT PL intensity is especially signifi cant when the boron impurity content is low (Figure 8).
Both boron centers can be observed simultaneously by capacitance methods [21,22].These centers were classifi ed as В and D-centers, respectively.Ballandovich, et al. [21] studied DLTS spectra in the HT and LT SiC 6H samples doped with boron.DLTS curves measured on HT and LT 6H-SiC(B) samples are presented in Figure 9.In all cases, high-temperature annealing leads to a decrease in the luminescence intensity.However, in LT 6H-SiC samples, the RT PL effi ciency decreases more smoothly.As a result, the luminescence is kept up to higher annealing temperatures.The   A similar effect of the Si:C ratio on the concentration of deep and shallow boron-containing acceptor centers was observed in 4H-SiC epitaxial layers grown by the CVD method [23].As shown in the study [23], the concentration of shallow acceptor centers associated with boron in layers grown at a low Si:C ratio was two orders of magnitude higher than the concentration of D centers.In these layers, only low-temperature luminescence   As seen, when plotted in the (1/C 2 -V) coordinates, the C-V characteristic of the control structure is linear, which is typical of p + -n junctions with a uniformly doped base.The slope of this characteristic is determined by the concentration of donors in the n-substrate, N d = 1.8×10 18 cm -3 .Diffusion of boron into the starting sample leads to a many-fold decrease in the capacitance of the p + -n structure, with the capacitance depending only slightly on the reverse bias applied to the sample (Figure 13, curve 2).
It is known from the general theory that, in the case of a measuring bridge with a parallel equivalent circuit, a decrease in the equivalent capacitance of a sample may be due to an increase in the series resistance in the measuring circuit.However, we observed no frequency dispersion of capacitance, inherent in a diode with a high-resistivity base.The above specifi c features are characteristic of the behavior of a p-i-n structure, which in turn points to the presence in the vicinity of the p + -n junction of an extended, strongly compensated region.
DLTS measurements show that the appearance of this region is due to the introduction of D-centers by diffusion into the n-type substrate of the starting sample.Subsequent hightemperature annealing of the sample leads to the dissociation of the D-centers, with the capacitance of the structure again increasing (Figure 13, curve 2).This means that the annealing of D-centers causes an increase in the concentration of uncompensated donors in the diode base.This result is in agreement with the conclusion that part of the boron-related acceptor centers passes into an electrically inactive state in the course of annealing.The C-V characteristic of an annealed sample corresponds to the C-V characteristic of a p-n junction with a linear distribution of the impurity in the base, which is typical of diffusion structures.

Clusters of native defects:
Experimental data obtained by various methods confi rm the presence of intrinsic point defects in LT SiC crystals, primarily of the vacancy type.However, the reasons for the increased stability of these defects during thermal annealing remain unclear.It is assumed that this is due to the presence of active clusters.For LT 6H-SiC specimens grown in conditions of relative excess of silicon, faceted inclusions of unknown composition were observed by TEM [24].This type of defect is absent in HT 6H-SiC crystals.High magnifi cation permits visualization of scaly inclusions of the 2nd phase located in the center of such areas.
The results of the study show that LT 6H-SiC samples contain an increased concentration of intrinsic defects, which signifi cantly affect the properties of the material.These defects are characterized by increased thermal stability and can be maintained at high temperatures during heat treatment.Additional information on clustering processes in SiC was obtained by studying the annealing of radiation defects [24][25][26][27][28][29][30][31].
Girka, et al. researched the annealing of radiation defects in the SiC crystals, irradiated by fast electrons [25,26], reactor neutrons [27], and heavy ions [28] by positron annihilation spectroscopy (PAS).Figure 13 represents the dependence of the positron lifetime т of electron and neutron-irradiated samples on annealing temperature.
In general, one can distinguish up to fi ve stages of positronsensitive defect annealing: two low-temperature stages that are characterized by temperature interval ΔT а = 100 -600 °C and defect migration activation energy E a ≈ 1 eV; the stage o f "negative" defect annealing (ΔT a = 1000-1200 °C, Е a = 1.9 eV); two high-temperature stages of defect annealing (ΔT a = 1400-1500°C and E a = 2.2 eV, Т a = 1700-1900 °C and E a = 2.8 eV, respectively).
The most interesting stage is the "negative" annealing stage.This stage is observed only in highly doped nitrogen  16 to 1×10 12 cm -3 (Figure 15).Considering the high thermal stability of these clusters, it could be assumed that they include antisites or dislocated matrix atoms.
Sitnikova, et al. [31] studied neutron-irradiated SiC crystals by transmission electron microscopy.SiC samples irradiated with high doses (F=10 20 -10 21 cm -2 ) of reactor neutrons and annealed at temperatures of 1800-2600 °C were used.The TEM method made it possible to observe separately the stages of annealing both vacancy and interstitial defects in the samples [31].
At low doses of irradiation of the SiC crystals (F ≈ 10 20 cm - 2 ) after subsequent annealing at Ta = 1800 °C, clusters with sizes no more than 40 nm are generated.Analysis of the image contrast of the defects suggests that these defects are small dislocation loops lying in prismatic planes.The concentration of these defects is 3×10 14 cm -3 .
The concentration of dislocation loops in the samples irradiated with high doses of neutrons F ≈ 10 21 cm -2 increases by about an order of magnitude.The average size of the loops is 70-90 nm, but individual defects with sizes of about 150 nm are also observed.The analysis performed allows us to conclude about the vacancy character of the observed dislocation loops.This conclusion agrees with the results obtained by the PAS method [25], according to which the maximum size of vacancy clusters in irradiated SiC crystals is observed at an annealing temperature of Ta ≈ 1800 °C.Annealing temperature increase, above 2000-2300 ° C, results in defects disappearance is also observed by TEM.
Defects with black-and-white contrast in the samples irradiated with a dose of F ≈ 10 21 cm -2 can also be observed after annealing at a temperature (T a ) of about 2300 °C.Analysis of these defects images shows that these are clusters of interstitial atoms.These defects are fairly evenly spread in the samples, and their concentration is (2-6)×10 12 cm -3 .It should be noted that vacancy and interstitial dislocation loops lie in planes perpendicular to each other.

Summary information on detected intrinsic defects in SiC
crystals is given in Table 2.  of these defects does not exceed 10 16 cm -3 .The same intrinsic defects were found in Si-rich crystals grown by hightemperature chemical vapor deposition (HTCVD) [33], named P6/P7, Z1/2 and EH6/7 [6].All these centers are very stable and anneal out at temperatures as high as 2000 °C.
We believe that the main intrinsic defects in LT SiC samples are carbon vacancies (V C ), which are part of the D 1 and D centers.Indeed, the conclusion that the D 1 -center includes V C is confi rmed by EPR measurements [34].PAS data show that vacancy defects in SiC are thermally stable and persist up to We associate the main difference between LT and HT 6H-SiC samples with the presence of clusters in LT SiC (or Si-rich crystals).During annealing, these clusters are in an active state, changing in size, up to extremely high temperatures of 2500÷2600 °C.Active clusters are the sources of non-equilibrium point defects [10].Such a mechanism for the formation of intrinsic point defects was described in theoretical works and confi rmed by experimental studies of the defects' behavior in crystals containing radiation defects [34].For example, it is known that Si vacancies are stable up to 750 0 C, and C vacancies -up to 1150 °C.Nevertheless, they are actually found in SiC crystals at temperatures as high as 1400-1600 °C [34].
The highest cluster activity is observed at a relatively low growth temperature (below 2000 °C), when small clusters are predominate.The behavior of a system with non-equilibrium intrinsic defects is described within the framework of a wellknown effect -Oswald ripening [35].However, in contrast to classical semiconductor materials (for example, silicon) and metals, clusters in SiC remain in an active form at very high temperatures at which crystal growth occurs.Nonequilibrium defects are stable up to temperatures of the order of 2500 -2600 °C, which is related to a low rate of relaxation processes in the bulk of the crystal.Thus, the relatively high concentration of electrically and optically active defects in SiC is a consequence of the presence of metastable active clusters.
On the contrary, in HT and C-rich crystals, the clusters are obviously not sources of electrically active centers.According to theoretical calculations, the formation energies of carbon ) and silicon (Е v Si ) vacancies in cubic position are 4.56 eV and 7.16 eV, respectively [32].However, these values differ signifi cantly from those obtained from the analysis of hightemperature diffusion processes in SiC.It is known that the activation energy of self-diffusion of C and Si atoms in n-type SiC lies in the range between 8.18 -8.2 eV [36,37] Since the migration energy of vacancies does not exceed 1.9-2.1 eV [10 ], we fi nd that E v is 6.2 eV, which is close to the calculated value The values of the vacancy formation energies obtained from the analysis of the impurities diffusion are even less consistent.
According to experimental data, the energies activation of diffusion for donor impurities of nitrogen and phosphorus are (7.6-9.3)eV [38] and 11.6 eV [39], respectively.These values are signifi cantly higher than the diffusion activation energies for acceptor impurities (B, Al, Ga) which are about 5.1-6.0 eV [8].
Since donor impurities (N and P) migrate over carbon vacancies whereas the acceptor impurities over silicon vacancies [8] we can conclude that the formation energy of the carbon vacancies signifi cantly exceeds the formation energy of the silicon vacancies.Indeed, according to calculations, the formation energies of carbon and silicon vacancies are 7-9 eV and 4-5 eV, respectively.These results signifi cantly contradict the calculated data [32].
The low formation energy of silicon vacancies can be explained by the fact that the diffusion-active states of acceptor impurities are (VA) pairs that are stable at high temperatures.
In the case of boron diffusion, this conclusion is confi rmed.
This assumption is supported by the fact that deep centers, which are complexes of impurity atoms with vacancies, are observed not only in SiC (B) but also in SiC samples doped with Al and Ga.Such samples of p-type conductivity contain signifi cant concentrations of carbon vacancies, which  At the same time, we believe that the higher values of the formation energy of carbon vacancies obtained from the analysis of high-temperature diffusion data are more reliable [40].The high formation energies of equilibrium (thermal) point defects explain the fact that in a wide temperature range, the main mechanism of vacancy generation in LT 6H-SiC is cluster dissociation.

EPR of deep-level defects based on vacancies associated with impurity atoms
In this section, we will consider group-III acceptors with deep levels in SiC.Group-III elements (B, Al, and Ga) are the most important acceptor impurities in SiC.They may be introduced into a crystal either during the growth process, by diffusion, or by implantation.Results of EPR and ENDOR studies of group-III acceptors in SiC are summarized in the book [2]. Figure 16   Spin defects in SiC as a material platform for sensing, quantum photonics, and information processing at ambient conditions: The unique quantum properties of NV centers in diamond [42] have motivated efforts to fi nd defects with similar properties in silicon carbide, which can extend the functionality of such systems [2,[43][44][45][46].SiC is a technological material that is used in various devices.This material is now taking on a new role as a platform for using new quantum technologies.A feature of SiC is the presence of different polytypes, and the properties of atomic-scale defects are unique for each polytype; moreover, even in one polytype, the center can occupy different nonequivalent lattice positions.This allows one to select a center with optical and microwave parameters suitable for a specifi c task.Atomic-scale defects in SiC are considered a material platform for applications in photonics, spintronics, quantum information processing, and environmental sensing.The spin state of the defects can be initialized, manipulated, and read out by means of ODMR, via LAC and CR [47][48][49][50][51][52][53][54][55][56][57].The presence of high concentrations of isotopes with zero magnetic moments in silicon carbide provides good coherent properties for spin defects.As a result, coherent manipulations with spin states were performed at room temperature and even at temperatures exceeding room temperature.Two families of defects were found in silicon carbide: the fi rst with the spin S = 1 and the second with the spin S = 3/2.For both families, a unique property of optical alignment of the populations of spin sublevels was discovered, which made it possible to perform spin manipulations by changing the populations of these sublevels using resonant microwave irradiation.Such manipulations lead to giant changes in the photoluminescence of defects, which makes it possible to register magnetic resonance by optical methods.
The center with spin 1 is a neutral divacancy of silicon and carbon located in adjacent positions with a covalent molecular bond.A model is proposed for a spin 3/2 defect in the form of a negatively charged silicon vacancy in the paramagnetic state, which is perturbed by the interaction with a neutral carbon vacancy in a nonparamagnetic state located at a neighboring site along the SiC symmetry axis [54].These centers are usually designated by the corresponding zero-phonon lines (ZPLs).
For this family, the ground state and the excited state have spin S = 3/2.For a number of defects, the population inversion in the ground state is created even in zero magnetic fi elds, and at room temperature using optical pumping, which leads to stimulated microwave radiation.It is assumed that if the ODMR signal coincides with the LAC, polarization of the surrounding silicon 29 Si and 13 C carbon nuclei with long relaxation times is possible through hyperfi ne interactions.Circles indicate areas of reduction of the signals with a certain polarization of 29 Si, which are present in the signals of the ODMR in the form of satellites.We assume that optical registration of nuclear magnetic resonance on polarized nuclei 29 Si and 13 C of spin centers in SiC will fi nd a number of applications (quantum computing, NMR imaging, gyroscopes, etc).In Figure 18  A spin ensemble of defects can be prepared in a coherent superposition of states, while Rabi oscillations in SiC can be maintained at room temperature up to fractions of a millisecond.The electron spin of defects can be controlled using a low-energy microwave fi eld with a frequency of 1-300 MHz, which is compatible with NMR imaging.Vacancy defects in SiC are optically active in the near-infrared, which is preferred for potential biological applications in vivo due to its deepest tissue penetration and which is compatible with fi ber optics.The sensing concept is based on variants of the ODMR method with sensitivity up to a single spin [2,46,47].The demonstrated spin properties of defects in SiC and related nanostructures open up new possibilities for modern quantum technologies and the creation of quantum sensors for submicron probing.The effect of optically induced population inversion of spin states at RT can be used to create solid-state masers and highly sensitive radio frequency amplifi ers.
Recently, the existence of NV defects (N C -V Si ) in SiC has been demonstrated in [58,59] and can be expected to hold similar promising properties as NV defects in diamond and divacancies in SiC.In the negative charge state, they have S=1, its optical properties shifted to the NIR region (around 1200nm wavelength).
The behavior of point defects and their clusters forming with deep levels in SiC band gap together with shallow and deep acceptors properties, considered in the current manuscript is of signifi cant importance for the development of quantum technologies with defect spins.The latter is well documented in the case of the negatively charged nitrogen-vacancy defect in diamond for which the doping levels with donor and acceptor impurities, the presence of other point defects in the lattice plays a crucial role on the NV defect charge stability, optical properties and spin-coherence times [60,61].The same is true for V Si defects in SiC [62].That is results presented in the current manuscript are of signifi cant importance.
The EPR and photoluminescence spectra of erbium centers Er 3+ associated with a carbon vacancy (Er 3+ -V C ) were found in SiC [2].Er 3+ ions occupy a special place in optical communication since the energy of the transition inside the 4f-shell of this ion corresponds to a wavelength of 1.54 μm, which corresponds to the minimum absorption in fi ber-optic communication systems.Qubits based on nuclear spins in solids, as a rule, have a high degree of coherence, which allows them to be used as systems for storing and processing quantum information.The electron-nuclear spin system for Er 3+ in SiC is of particular interest, since both the electron spin and the nuclear spins of the 29 Si and 13 C isotopes have minimum values (1/2), while it is possible to change the isotopic composition of the nuclear spins of silicon and carbon, as well as changes in the isotopic composition of erbium.Due to the large electronic g-factor for the Er 3+ ion [2], there is a strong interaction with resonant microwave excitation in EPR.The electronic and nuclear spins interact.The spin of an electron can be strongly excited through optical or microwave channels to create certain quantum states.The information can then be mapped into the nuclear spin of the ligand ( 29 Si and 13 C) to implement long-lived qubit memory.

Conclusion
Native defects in SiC crystals grown at low temperatures below 2200 °C in a Si-rich system were discovered.These defects affect on electrical and optical properties of the material and the quality of p-n-junctions.They are complexes that include a carbon vacancy.The same defects also appear in SiC crystals containing radiation defects.We note that the high thermal EN, Baranov PG, Kazarova OP (2024) Intrinsic defects in non-irradiated silicon carbide crystals.Open Journal of Chemistry 10(1): 004-019.DOI: https://dx.doi.org/10.17352/ojc.000034To identify the intrinsic defects, we used different methods including Positron Annihilation Spectroscopy (PAS), EPR, Deep-Level Transient Spectroscopy (DLTS), electrical, and optical measurements.Results obtained by these methods are in good agreement with the results collected by means of transmission electron microscopy (TEM) and X-Ray Diffraction spectroscopy (XRD).To compare the native defect densities in different type samples we studied diffusion distribution of boron impurity.The nuclear reaction 10 B(n,) 7 Li was used.This

Electrical and optical measurements:
Electrical and optical measurements revealed intrinsic defects in the pure LT-SiC 6H-SiC crystals [9,10], grown by Mokhov, et al. [9] in open tantalum containers in a vacuum at temperatures below 2000 °C.Grown crystals are characterized by the donor N d ~ (1-3)×10 16 сm -3 and electron mobility μ n ~320-350 cm 2 /(V×s).The dependence of the donor concentration in the LT SiC crystals grown in a Si-enriched system on the Nitrogen doping level in the SiC sources is presented in Figure 1.It is seen that at low impurity content, the donor concentration in SiC crystals ceases to depend on its concentration in the source.This result points to the possible presence of donor-type intrinsic defects in the SiC crystals.This conclusion is confi rmed by the fact that the concentration of paramagnetic nitrogen (main donor impurity) determined by the EPR in the high pure SiC samples is signifi cantly lower than the donor concentration obtained from electrical measurements[9].In the EPR spectra of the samples, additional lines are detected or nitrogen bands are split, which is associated with intrinsic defects[10].A decrease in the nitrogen concentration (N c D ) in LT SiC below 1×10 17 сm -3 does not lead to the increase in diffusion lengths of the minority current carriers.

Figure 1 :
Figure 1: Donor concentration in LT SiC Si-rich crystals of the 6H polytype versus donor (nitrogen) content in the SiC source for different growth directions: 1-[0001] Si, 2-[0001] C. The growth temperature (Tg) is 1850 °C.The pressure of Ar is 10 -3 Torr.The crystals were grown in the Ta containers.

Figure 4 .
It is seen, that the scatter of the vacancy density in different polytypes is much smaller than for crystals with different deviations from stoichiometry.Based on the mentioned above it can be concluded that LT SiC samples contain an elevated density of intrinsic defects.It seems logical, to assume that the defects are vacancy clusters or complexes of vacancies with impurity atoms.From our point of view, the latter possibility is less probable, because positron-sensitive defects are observed in high pure SiC samples.Additionally, a change in the impurity doping level does not increase the intensity of the long-life component.At the same time, clusters of vacancies have a very high thermal stability in SiC.

Figure 3 :
Figure 3: Diffusion profi les of the boron isotope 10 B in HT 6H-SiC (1,2-red line) and LT 6H-SiC (3,4-blue line).Sample 3 was preliminary annealed at 2450 °C.Diffusion temperature is 1900 °C (1) and 2100 °C(2,3,4).Boron diffusion distribution was captured using the method of track autoradiography (TA).It is evident that the diffusion rate is higher in the LT (Si-rich) sample, signifi cantly higher than in the HT samples.At low temperatures of diffusion, the bulk branch of boron distribution is described by the standard erfc-function.At temperatures above 2400°C, the distribution takes on a step-like character, which is associated with the partial decomposition of rapidly diffusing associates (B-Vc).

Figure 4 :
Figure 4: Density of carbon vacancies (V C ) versus the fraction (D) of the hexagonal polytype phase in HT SiC crystals (1,3); LT SiC Si-rich (2-blue point); and C-rich (3red point) is defi ned by the diffusion method.The absolute values of the vacancy densities are based on PAS data [11].C-rich SiC was grown at a temperature of 2040°C and with the introduction of tin impurities into the growth zone.It is evident that more cubic polytypes, like the LT samples, have an increased vacancy concentration.
Verenchikova, et al.[18] studied the p-n-structures obtained by epitaxial growth of p+-SiC (Al) layers on the high-purity HT 6H-SiC substrates.The shift of the electroluminescence maximum to longer wavelengths in the structures grown at relatively low temperatures (Tg < 2200 °C) was observed.The assumed ionization energy of a deep donor center was about 0.6-1.0eV [18].Besides, the reverse breakdown voltage of such p-n junctions was much lower compared to theoretical values.The authors assumed that these features are connected with the formation of clusters of intrinsic defects.According to their calculation, the defect cluster includes about 60-70 atoms.The clusters were destroyed partly during high-temperature annealing at Ta > 2500 °C.The thermal stability of the D 1 centers was also studied by Vodakov, et al. [19].Figure 5 represents the comparative dependences of the intensity of the D 1 PL at the maximum of the spectrum on the annealing temperature of SiC samples containing D 1 centers.From Figure 5 it is readily seen, that the D 1 PL is stable at temperatures as high, as 2200 °C.The luminescence decreases sharply only at Ta > 2300 °C, indicating a drop in the concentration of the D 1 -centers.It is of interest to compare the behavior of the D 1 -centers in the as-grown LT 6H-SiC and in the crystals irradiated with high-energy particles.For this purpose, LT 6H-SiC and HT 6H-SiC crystals with the same doping level, irradiated with reactor neutrons were studied [19].The radiation dose was varied in the range F ~ 10 17 -10 19 сm -2 .

Figure 5 :
Figure 5: Dependence of the D 1 PL intensity on annealing temperature (T a ) for different SiC crystals: 1 -as-grown LT crystals (Тg = 1900 °C) and 2 -p(Al)-nstructures (T dif = 2100 °C).It is evident that at temperatures above 2200°C, the annealing of D 1 centers is observed.
The DLTS measurements show that the B-centers are present in the all-initial SiC samples.The peak near 130 К on the DLTS spectra corresponds to these centers.The boron diffusion leads to a signifi cant increase in the concentration of the В-centers.It also induces the appearance of the D-centers, with corresponding peaks on the DLTS spectra near 280 K (Figure 9 (b), (c), (d)), indicating the presence of the deep-level center).The relative concentrations of the D and В-centers are very sensitive to the growth method and to the conditions of the original samples doping.In the HT 6H-SiC samples with an initial concentration of uncompensated donors below 6×10 17 cm -3 , doped by boron via diffusion we observe only В centers (Figure 9a).At a nitrogen concentration of 8×10 17 cm -3 , a weak signal of the D-center appears (Figure 9b), which becomes dominant at a nitrogen concentration of 2×10 18 cm -3 (Figure 9c).In contrast with the HT 6H-SiC, the diffusion of boron into LT 6H-SiC with the same nitrogen concentration leads to the predominant formation of D-centers (Figure 9d).This behavior agrees well with the difference in the luminescence properties of the samples.Specifi cally, in the HT 6H-SiC samples (without D-centers), the RT PL is not observed.The crystals with (N d -N a ) > 1×10 18 cm -3 with the D-centers, exhibit a faint luminescence at room temperature.Finally, both the LT and HT samples with predominant concentrations of the D-centers, exhibit an intense room-temperature luminescence.These results confi rm the conclusion that the D-centers are involved in the high-temperature boron luminescence.In [22], it was that the thermal stability of deep boron centers in LT and HT samples differs signifi cantly.The dependence of the RT PL intensity on the annealing temperature for the HT (curve 1) and LT SiC (curve 3) samples doped with boron diffusion is shown in Figure 10.For comparison, the same dependence is shown for the initial LT-sample before boron diffusion (curve 2).

Figure 8 :
Figure 8: Dependence of the RT PL intensity on the boron concentration in different samples: LT 6H-SiC (B) (1 -blue points) and HT 6H-SiC(B) (2 -red points).Nitrogen impurity concentration is 3×10 18 cm -3 .It is seen that the emergence of boronrelated room temperature photoluminescence (RT PL) in LT SiC (B) is observed at signifi cantly lower boron concentrations compared to HT samples.

Figure 15 :
Figure 15: Dependence of the average cluster size and cluster concentration on the annealing temperature in SiC crystal irradiated by reactor neutrons (F = 10 21 сm -2 ).
high annealing temperatures.The rapidly diffusing state of the boron impurity is the D-center or (B Si V C ) complex.In this case, the heavily doped SiC of p-type conductivity may host the (B S V C ) and (C Si V C ) pairs, created by means of a hightemperature annealing.Our data indicate that native defects are not observed in carbon-rich SiC crystals.Obviously, this is due to the fact that neutral carbon antisites (C Si ) are electrically non-active centers in SiC[12].We show that deep boron (B Si V C ) centers have high thermal stability and are preserved upon annealing to very high temperatures of 2500÷2600 °C.It is important to note that the annealing temperatures of the D-centers in as-grown crystals are signifi cantly higher than those for the same centers created by irradiation.Boron diffusion data show that in the LT-samples these centers almost do not dissociate during diffusion at temperatures below 2400 °C.This is confi rmed by the fact that at T d < 2200 0 C the boron impurity concentration profi le in LT 6H-SiC is described by the standard erfc-function (Figure3, curve 4).Contrary, in the HT 6H-SiC samples the steep part of boron distribution on the "tails" of diffusion profi les is observed, which indicates a sharp decrease in the relative concentration of D -centers (Figure3, curve 2).
, p-n-junctions neutrons irradiated samples, TEM, XRD[24,30,31], C-Vmeasurements Note: LT SiC: SiC crystals grown at low temperatures; EPR and DLTS are for electron paramagnetic resonance and for Deep Level Transient Spectroscopy, respectively; TEM, XRD, and C-V are for Transmission Electron Microscopy, X-Ray Diffraction, and Capacitance-voltage method, respectively.Tg means the growth temperature of SiC crystals.Citation: Mokhov EN, Baranov PG, Kazarova OP (2024) Intrinsic defects in non-irradiated silicon carbide crystals.Open Journal of Chemistry 10(1): 004-019.DOI: https://dx.doi.org/10.17352/ojc.000034obviously contributes to a decrease in the observed diffusion activation energy.Upon diffusion annealing, such layers are sources of vacancies responsible for the appearance of defect luminescence bands (D 1 and D spectra).
shows the EPR spectra of deep B (a), deep Al (b), and deep Ga (c) in 6H-SiC crystal recorded in the X band (9.3 GHz) in various crystal orientations at 4.5 K.In Figure 16(a), the high-fi eld signal denoted as sB and marked with a horizontal bar belongs to the shallow boron, since it can be described by a spin Hamiltonian with g factors and hyperfi ne (HF) structure parameters of shallow boron in two quasi-cubic and hexagonal positions in the lattice.In the low-fi eld region at angles close to  = 0 0 two unresolved overlapping broad lines are observed; they are denoted dB (deep B) in Figure 16(a).Their intensity correlates with the characteristic light-yellow luminescence and ODMR signals of deep boron centers.Finding and determination of structure of Al acceptors with deep levels (deep Al) in SiC is of great interest concerning its application since aluminum is the main acceptor dopant in devices based on SiC.Appearing of Al acceptors with deep levels worsens the electric properties of a material; so the minimization of the content of deep aluminum in device elements is an important problem.Figure 16(b) shows the angular dependence of the EPR spectra in the 6H-SiC crystals doped with Al.The spectra consist of several groups of peaks which depend strongly on the angle between the fi eld and the crystal axis c.One of the groups of anisotropic lines belongs to the shallow Al acceptors in inequivalent crystallographic positions.This group is seen in the low-fi eld region in Figure 16(b) as two broad lines (in orientations close to B ||c).Vertical lines show the angular dependence of the ESR transitions of the shallow Al acceptors corresponding to cos.This behavior points to the fact that the shallow acceptors may be described within the effective mass approximation.Upon heating, the ESR signals of the shallow Al acceptors disappear above 5 K.The other group of anisotropic lines has angular and temperature dependences resembling those of the deep B. Thus, we will denote the new Al center as deep Al, analogous to the deep B. The signals of the shallow and deep Al acceptors can be distinguished by their angular dependences: the deep Al signals have a large intensity of up to 7 K while the shallow Al signals disappear at temperatures above 5 K.The difference in the peak positions for the shallow and deep Al is clear at angles exceeding 30.At orientations close to = 0 0 the signals of the shallow and deep Al overlap.EPR spectra of deep gallium were observed in epitaxial layers of 6H-SiC doped with Ga.The spectra, which depend strongly on the angle between the crystal axis c and the magnetic fi eld, are shown for different orientations of the

Figure 16 :
Figure 16: Electron paramagnetic resonance of Boron impurities in 6H-SiC.(a) Angular dependence of the EPR spectra at X-band (9.3 GHz) in the 6H-SiC crystal doped with 11 B. Dashed lines show the spectra of the 6H-SiC crystal doped with 10 B. (b, c) Angular dependence of the EPR spectra in the X band in the 6H-SiC:Ga (b) and 6H-SiC:Al (c) crystal.Calculated dependence for the shallow Ga and shallow Al is shown with short bars.T=4.5 K. (c, inset) A model of the deep B (Al, Ga) acceptor.A distribution of the spin density of the deep boron.B (Al, Ga) shows the position of the group-III impurity in the Si site; V C is the carbon vacancy [41].

Figure 17 (
Figure 17 (a) shows an example of the room temperature photoluminescence (PL) spectra of spin-3/2 centers in SiC (see Chapter 6 in Ref. [2]). Figure 17 (b) presents a typical experimental microwave (MW) dependence of the ODMR signal of the spin-3/2 centers in SiC, Δ is the zero-fi eld splitting (ZFS).The ODMR signal is detected under laser excitation  = 808 nm in zero and B=0.34 mT magnetic fi elds.The vertical bar indicates the ODMR contrast.A hyperfi ne structure (HFS) due to the interaction with 29 Si nuclei for twelve Si atoms in the next nearest-neighbor shell of the Si vacancy is shown by arrows.The inset shows the energy level diagram and corresponding resonance transitions.The schemes of optically induced alignment of spin level populations in a zero magnetic fi eld at room temperature for spin-3/2 centers (possible confi guration is depicted) in different SiC polytypes are shown in Figure 16 (c).

Figure 18 (
Figure 18 (a), as an example, shows signals of room temperature ODMR and level anti-crossing (LAC) of the V2 spin centers in a 4H-SiC crystal (see [2]).The signals were recorded by lock-in detection of the change in the PL in the near-IR range with the application of a constant magnetic fi eld and an oscillating low-frequency magnetic fi eld.The energy of the corresponding spin levels is shown at the top.The LAC in the absence of the MW fi eld (MW "off"), denoted as LAC1 and LAC2; for LAC1 B=D and for LAC2 B=2D, where ZFS Δ=2|D|, and two ODMR signals recorded at the MW fi eld on 40 and 45 MHz are shown.The choice of frequencies was due to the fact that when the ODMR signal approaches the LAC1 signal, the ODMR signal increases several times and approaches the LAC1 signal intensity.The ODMR line shift ΔB at frequencies of 40 and 45 MHz is shown demonstrating the principle of measuring magnetic fi elds by the ODMR method.A magnetic-fi eld position of the LAC1 signal of spin-3/2 centers in the ground state weakly depends on the orientation of the crystal, and therefore LAC1 signal is narrow enough in powder SiC.Thus, the LAC1 signal of spin-3/2 centers in nanosized samples of SiC can be used for all-optical measuring of magnetic fi elds with

Figure 17 :
Figure 17: Optically detected magnetic resonance of spin-3/2 centers in 6H-and 4H-SiC crystals.(a) An example of the room temperature photoluminescence (PL) spectra of spin-3/2 centers in SiC.(b) A typical ODMR signal of the spin-3/2 centers in SiC in zero and B=0.34 mT magnetic fi elds.A hyperfi ne structure (HFS) with twelve Si atoms of the next nearest-neighbor shell is shown.(c) The schemes of optically induced alignment of spin level populations for spin-3/2 centers and possible confi gurations of spin-3/2 centers.

Figure 18 :
Figure 18: Level anti-crossing and electron paramagnetic resonance spectroscopy of spin-3/2 centers in 4H-and 15R-SiC crystals.(a) Room-temperature detection of the change in PL of spin-3/2 centers in a 4H-SiC: the level anti-crossing (LAC) signals (MW off) and ODMR.ODMR line shift ∆B in a magnetic fi eld for two frequencies demonstrates the principle of measuring magnetic fi elds.(b) (top) The PL ZPLs in 15R-SiC of the corresponding spin-3/2 centers.(bottom) X-band direct-detected (DD) EPR spectra induced with ZFLs V2, V3, and V4 [54].(c) Optically induced ESE spectra measured in 15R-SiC at W-band, (inset) the energy diagram and the light-induced inverse population of the spin sublevels of V2 centers.
stability of intrinsic defects in as-grown crystals in comparison with similar radiation defects is explained by the presence of active clusters in crystals, which are sources of mobile point defects at elevated temperatures.Experimental evidence has been obtained of the presence of such clusters, which are retained upon annealing to temperatures as high as 2500-2600 °C.That is results presented in the current manuscript are of signifi cant importance.Quantum technologies based on coherent control of the spins of color centers in silicon carbide have developed fantastically over the past decade.The combination of supersensitive optical techniques for highresolution image detection and reliable coherent control using magnetic resonance are key components in the development of the fi rst quantum devices based on these materials.The possibility of highly accurate scaling of color centers and achieving long coherence times opens up new prospects for socalled hybrid quantum processes, in which color centers are associated with different types of qubits.There is no doubt that this area of research will develop rapidly in the near future and, as can be expected, will lead to the emergence of new quantum technologies.

Table 1 :
Growth conditions and properties of SiC crystals.

Table 2 :
Intrinsic defects in SiC crystals.