Boron-doping of cubic SiC for intermediate band solar cells: a scanning transmission electron microscopy study
Patricia Almeida Carvalho, Annett Thørgesen, Quanbao Ma, Daniel Nielsen Wright, Spyros Diplas, Augustinas Galeckas, Alexander Azarov, Valdas Jokubavicius, Jianwu Sun, Mikael Syväjärvi, Bengt Gunnar Svensson, Ole Martin Løvvik
SSciPost Physics Submission
Boron-doping of cubic SiC for intermediate band solar cells:a scanning transmission electron microscopy study
P. A. Carvalho , A. Thøgersen , Q. Ma , D. N. Wright , S. Diplas , A. Galeckas , A.Azarov , V. Jokubavicius , J. Sun , M. Syv¨aj¨arvi , B. G. Svensson , O. M Løvvik SINTEF Materials Physics, Oslo, Norway University of Lisbon, Instituto Superior Tecnico, Lisbon, Portugal University of Oslo, Department of Physics, Oslo, Norway SINTEF Instrumentation, Oslo, Norway University of Oslo, Department of Chemistry, Oslo, Norway Link¨oping University, Department of Physics, Chemistry and Biology, Link¨oping, Sweden* [email protected], [email protected]
Abstract
Boron (B) has the potential for generating an intermediate band in cubic sili-con carbide (3C-SiC), turning this material into a highly efficient absorber forsingle-junction solar cells. The formation of a delocalized band demands highconcentration of the foreign element, but the precipitation behavior of B in the3C polymorph of SiC is not well known. Here, probe-corrected scanning trans-mission electron microscopy and secondary-ion mass spectrometry are used toinvestigate precipitation mechanisms in B-implanted 3C-SiC as a function of tem-perature. Point-defect clustering was detected after annealing at 1273 K whilestacking faults, B-rich precipitates and dislocation networks developed in the1573 - 1773 K range. The precipitates adopted the rhombohedral B C struc-ture and trapped B up to 1773 K. Above this temperature, higher solubilityreduced precipitation and free B diffused out of the implantation layer. Dopantconcentrations of 10 at . cm − were achieved at 1873 K. Contents a r X i v : . [ c ond - m a t . m t r l - s c i ] D ec ciPost Physics Submission An intermediate band (IB) in the energy band gap of a semiconductor allows photons withlower energy than the band gap to excite electrons from the valence to the conduction band,increasing the photocurrent generated [1]. Due to the potential for enhanced efficiency inenergy conversion, intermediate-band solar cells are promising candidates for the next gen-eration of photovoltaic devices. Theoretical efficiencies of 63% have been estimated for IBsolar cells under concentrated sunlight [1], a value considerably higher than the maximumefficiency of 40% expected for conventional single p-n junctions [2].Realization of IB solar cells has faced challenges in finding a semiconductor/dopant systemwith band gap in the 1.9 - 2.5 eV range and suitably positioned intermediate band. Cubicsilicon carbide (3C-SiC) has a nearly ideal band gap of 2.36 eV at room temperature, combinedwith useful electronic properties [3], and boron (B) has been proposed to form a deep acceptorlevel 0.7 eV above the valence edge ( E v ) of SiC [4]. Thus, B-doped 3C-SiC is a promisingabsorber system for highly efficient photovoltaic devices.The hydrogen-like model with the values of permittivity and effective hole mass repor-ted for 3C-SiC [5] estimates 10 - 10 at . cm − as the minimum concentration of shallowacceptors for their energy levels to merge into impurity bands. However, the two acceptorlevels associated with B in SiC are rather deep (with positions at ∼ E v +0 . ∼ E v +0 . High-quality 3C-SiC single crystals with an area of ∼ × µ m and thickness > µ m weregrown on 4H-SiC substrates off-oriented 4 ◦ from [0001] by sublimation epitaxy, resulting inan exposed surface close to (111) (details on the process can be found elsewhere [10]).Due to the processing conditions, nitrogen (shallow acceptor) can be present in concentrationsup to 10 cm − turning the as-grown material into an n-type semiconductor [12]. Hall-effectmeasurements were used to attest the n-type character of the 3C-SiC single crystals anddetermine their carrier density ( ∼ cm − ) and resistivity ( ∼ ± . cm) [12].Implantation with B + ions was carried out at elevated temperatures (673 and 773 K)2 ciPost Physics Submission along a direction close to (111) using multiple energies (100 to 575 keV) with a total doseof 4 - 13 × atoms.cm -2 to form box-like concentration profiles of about 1, 2 and 3 at.%B (used henceforth to designate the concentration in the material). The profiles extendedabout 600 nm in depth either directly below the free surface or buried with a start at ∼ . × s. Combinations of different B concentration level andannealing temperature/time have been investigated to infer precipitation and defect formationtrends within a large window of processing parameters. Prior to annealing, the samples wereprotected by a pyrolized resist film (carbon cap) after native oxide etching. The pyrolysis wasperformed in forming gas at 1173 K for 6 - 9 × s and the carbon cap was removed bydry thermal oxidation at 1173 K for 2 - 3 × s before the measurements. The crystallinityof the as-grown, as-implanted and annealed samples was attested by grazing incidence X-raydiffraction and Rutherford backscattered spectroscopy [13].Absence of extended structural defects in the sublimation-grown 3C-SiC single crystalswas confirmed by conventional transmission electron microscopy (TEM) prior to the B im-plantation. The observations were performed close to (cid:104) (cid:105) zone axes for easy detection ofthe typical stacking faults on { } planes. The implanted and annealed samples were inves-tigated by TEM and by annular bright-field (ABF), low-angle annular dark field (ADF) andhigh-angle annular dark field (HAADF) STEM. The microscopy work was performed with aDCOR Cs probe-corrected FEI Titan G2 60-300 instrument with 0.08 nm of nominal spa-tial resolution. Chemical information was obtained by X-ray energy dispersive spectroscopy(EDS) with a Bruker SuperX EDS system, comprising four silicon drift detectors, and byelectron energy loss spectroscopy (EELS) with a GIF Quantum 965 EELS Spectrometer.TEM sample preparation was performed by focused ion beam with Ga + ions accelerated at30 kV using a JEOL JIB 4500 multibeam system. Lattice images were indexed using fastFourier transforms (FFT) and strain was evaluated by geometric phase analysis (GPA) usingthe FRWRtools plugin [15] implemented in Digital Micrograph (Gatan Inc). The crystallo-graphic orientations between precipitate variants and the 3C-SiC matrix were investigatedusing crystallographic data retrieved from the literature [16]. The Carine Crystallography 3.1package [17] was employed to associate stereographic projections to the respective lattices.Phase diagrams of the B-C-Si system [18–20] have been used to interpret the microstructuralconfigurations.Boron concentration across sample depth was measured by SIMS using a Cameca IMS 7fmicroanalyzer with a primary sputtering beam of 10 keV O − ions rasterized over 150 × µ m with lateral resolution of 1 µ m and detection of B + secondary ions. Absolute values of Bconcentration were obtained via calibration with ion-implanted reference samples. The depthof the sputtered crater was measured using a Dektak 8 stylus profilometer and a constanterosion rate was assumed for conversion of sputter time to depth. Cross-sectional STEM images obtained directly after implantation and after annealing at 1273K are shown in figure 1 together with the corresponding concentration profiles obtained bySIMS. At low collection angles, the implanted regions exhibited homogenous but distinctive3 ciPost Physics Submission contrast compared to the underlying material (figure 1(a)). This is expected to result fromdisplaced Si atoms rather than from the introduction of the weakly-scattering B atoms [21] (orfrom displacing the also weakly-scattering C atoms). In silicon, scattering induced by latticedisplacement around a single substitutional B atom peaks at 40 mrad and several atomssuperimposed along the atomic columns can add up to significant scattering [22]. Boron hasbeen proposed to preferentially substitute Si in 4H-SiC [23]. However, boron-nitrogen donor-acceptor transitions (DAT) indicate that the behavior is more complicated, with B presumablyalso substituting C [24, 25], and possibly forming a complex in which a B atom in a Si site ispaired with a carbon vacancy (B Si -V C ), the so called deep boron-related D-center at ∼ E v +0 . { } planes (seestraight lines in figure 3 (a) and (b)). The absence of these defects in the samples annealed atlower temperatures (see Figures 1 (c) and 2 (a)) implies that additional thermally activatedlattice rearrangements occurred at 1673 K. A mottled contrast compatible with the presenceof clusters or precipitates was barely discernible in BF TEM (see arrows in figure 3 (b)), whileno significant changes were detected in the B concentration profiles (figure 3 (c)).4 ciPost Physics Submission Figure 1: Annular dark-field STEM images of a sample implanted with 1 at.% B at 773 K.Same region observed with (a) 48 - 200 and (b) 99 - 200 mrad collection angles. The samplesurface corresponds to the images top. (c) Annular dark-field STEM image of the same layerafter annealing at 1273 K for 3600 s obtained with 22 - 98 mrad collection angles. Convergenceangle: 21 mrad. (c) SIMS concentration profiles measured from the samples in (a,b) and (c).Different types of contrast were observed after heat treatment at even higher tempera-tures, as shown in figure 4. Extended structural defects, such as dislocations and stackingfaults, separating mosaics with slightly different crystallographic orientations, were presentin ABF and ADF images after annealing at 1773 K (figure 4 (a,d) and (b,e), respectively),whereas low-mass precipitates were the dominating feature in HAADF images (figure 4 (c,f)).Annealing at 1873 K resulted in lower defect density but large precipitates pinned curveddislocations (figure 4 (g) to (i)). The lower precipitate density and larger precipitate sizeat 1873 K (2 at. % B) compared to 1773 K (3 at. % B) points to lower nucleation rate,as expected for lower undercooling and/or lower supersaturation (lower initial concentrationand higher solubility). Ostwald ripening after precipitation may also have contributed to theincreased particle size [28]. 5 ciPost Physics Submission
Figure 2: (a) and (b) ABF and HAADF STEM images of 3C-SiC implanted with 1 at.% Bat 773 K and annealed at 1273 K. The images were obtained simultaneously from the sameregion with a convergence angle of 31 mrad and collection angles of respectively, 11 - 21 mradand 99 - 200 mrad. (c) and (d) Strain maps obtained with the 111 and 002 reflections from theinsets in the (a) and (b) raw images. The x and y directions are indicated in (c) and yellowarrows point to apparent compressive/tensile strain fields characteristic of edge dislocations. (e) and (f) Bragg-filtered images of (a) and (b) insets.6 ciPost Physics Submission
Figure 3: (a) and (b) Bright-field TEM images at different magnifications of a 3 at.% Blayer implanted at 673 K and annealed at 1673 K for 3600 s. The arrows in (b) indicateclusters/precipitates and the straight lines are stacking faults parallel to { } planes. (c)SIMS concentration profiles measured in the as-implanted and annealed states.The microstructural changes occurring during the heat treatments are summarized infigure 5 (a) using ADF images (where both extended structural defects and low-mass pre-cipitates present distinctive contrast) for different concentration and annealing conditions.The observations suggest that precipitation was controlled by diffusion at the lower annealingtemperatures and by driving force (supersaturation and undercooling) at the higher temper-atures, with highest rate at around 1773 K (see also Figure 4 (f)). Quantitative evaluation ofprecipitation parameters was complicated by the diffusive fluxes out of the implanted layer,which promoted precipitate dissolution. Nevertheless, a qualitative description is schematizedin (figure 5 (b)).Complete elimination of defects was not accomplished in the conditions studied, sinceLomer-Cottrell locks [29] and precipitates stabilized by extended defects were detected afterannealing at 2073 K for 1 . × s (see arrows in figure 5 (a)). Structural defects, suchas precipitates, stacking faults and dislocations, are inherently associated with electronictransitions and can contribute to the optical and electrical activity of the material, poten-tially masking/mimicking IB behavior in absorption/emission spectra.7 ciPost Physics Submission Figure 4: (a), (b) and (c) ABF, ADF and HAADF STEM images obtained simultaneouslywith a convergence angle of 22 mrad from a layer implanted with 3 at.% B at 673 K andannealed at 1773 K. (d) to (f) Magnified details of (a) to (c). (g), (h) and (i) ABF, ADF andHAADF STEM images obtained simultaneously with a convergence angle of 31 mrad froma sample implanted with 2 at.% B at 673 K and annealed at 1873 K. The arrows indicatestacking faults in the SiC matrix ending at a precipitate. Collection angles for ABF: 11 - 21mrad, ADF: 22 - 98 and HAADF: 99 - 200 mrad.
Despite the low fluorescence yield of B (EDS) and the overlapping of the B-K edge with thestrong Si-L2,3 edge (EELS), the local spectroscopy techniques demonstrated that the lowmass precipitates were rich in B (red tint in figure 6). The boride precipitates tended toadopt platelet morphologies but their atomic structure was not easily resolved due to theintrinsically weaker B scattering and the strong contribution of the embedding 3C-SiC matrix(see figure 7 (a) to (c)). 8 ciPost Physics Submission
Stacking faults running in the matrix often terminated at precipitates (see arrows infigures 4 and 7), in some instances in association with Lomer-Cottrell locks (figure 7 (d)).Since stacking faults were absent both in as-implanted samples and after annealing at 1273 K(figures 1 and 2), these defects were generated during precipitate growth, probably through astress relaxation mechanism. Therefore, precipitation has additional deleterious implicationson the overall structural quality of the B-doped 3C-SiC crystals.Large precipitates presented facets parallel to { } − SiC and { } when imagedalong (cid:104) (cid:105) and exhibited high density of planar defects, as illustrated in figure 8 for twooverlapping platelets with different crystallographic orientation. The symmetry and latticespacing are consistent with the rhombohedral B C cell [16]. Variants of the (0001) B C // { } and (cid:104) (cid:105) B C // (cid:104) (cid:105) orientation relation have been found, as depicted infigure 9. Analysis of the additional spots and streaks present in the Fourier transforms of thehigh-resolution images revealed that the planar defects lied on first-order pyramidal planes,i.e., { } B C [30–32].Boron carbide is a well-known semiconductor with electronic properties dominated byhopping-type carrier transport [33]. This is relevant in the context of the present investigationbecause the optical behavior of B-doped 3C-SiC may be masked by the presence of B C precipitates acting as embedded quantum dots with electronic transitions in spectral rangessimilar to those expected for an IB in 3C-SiC. Luminescence peaks exhibited by boron carbidehave been attributed to localized gap states and transitions between such states and the energybands [34, 35]. Specific characteristics of the boride determined from optical absorption,photoluminescence and charge transport data are: band gap of 2.09 eV, several disorder-induced intermediate gap states extending 1.2 eV above E v , excitonic level at 1.56 eV above E v , electron trap level around 0.27 eV below the bottom of the conduction band, and aconductivity of 20 (Ωcm) − for the B C stoichiometry at room temperature [34, 36]. Thetypical random distribution of twins parallel to { } B C [30–32] is likely to contribute tothe complex electronic configuration of this compound and to play a significant role in chargecarrier recombination. Changes in the concentration profiles compatible with long-range diffusion were only detectedupon annealing at temperatures higher than 1773 K (see figure 10). At lower temperatures,the B C precipitates trapped the B atoms and the low concentration of solute available inthe 3C-SiC matrix prevented any significant long-range diffusion, i.e., low solubility led tostable concentration profiles, suggesting an apparent low diffusivity.Since the equilibrium solid solubility as a function of temperature, C B3C-SiC ( T ), is givenby the concentration in thermodynamic equilibrium with the B C phase, the concentrationof free B immediately below the precipitation layer can be assumed ≤ C B3C-SiC ( T ). Thus, theconsistent absence of precipitates below ∼ C B3C-SiC (1873K) ≥ at . cm − (see arrow in figure 10 (a)).This value is comparable to the solubility reported for 6H- and 4H-SiC at 1873 K [37, 38],with 10 at . cm − proposed for these polytypes at the temperature of the 3C-SiC + B C ↔ L eutectic reaction ∼ ciPost Physics Submission Figure 5: (a) ADF images of B-implanted 3C-SiC samples (same magnification) for different Bconcentration level and annealing temperature/time, as indicated for each image. The arrowsin the image of the sample annealed for 4 h at 2073 K point to Lomer-Cottrell locks (leftand center) and to a locked defect configuration involving a large precipitate and a planardefect (right). (b) Qualitative illustration of the microstructural evolution in the implantedlayer with annealing temperature. The precipitate density was estimated from the initialconcentration and approximate precipitate size assuming spherical shape.10 ciPost Physics Submission
Figure 6: Sample implanted with 3 at.% B at 673 K and annealed at 1673 K. (a) ADF STEMimage acquired with collection angles of 48 - 200 mrad. (b) Overlapped B and C X-ray maps.(c) Si-K, (d) C-K and (e) Si-K X-ray maps. (f) Overlapped EDS spectra of the matrix (gray )and precipitate (red). (g) ABF STEM image acquired with collection angles of 11 - 22 mrad.(h) B-K and Si-L2,3 EELS maps. (i) Overlapped EELS spectra of the matrix (blue) andprecipitate (red). Convergence angle: 22 mrad.11 ciPost Physics Submission
Figure 7: (a) to (d) ADF STEM image of a sample implanted with 3 at.% B at 673 K andannealed at 1673 K acquired with a convergence angle of 22 mrad and collection angles of 48- 200 mrad. The arrows indicate planar defects in the SiC matrix. Zone axis: [1¯10] .Figure 8: (a), (b) and (c) ABF, ADF and HAADF STEM images obtained simultaneouslywith a convergence angle of 22 mrad from a layer implanted with 2 at.% B at 673 K annealedat 1973 K. Collection angles for ABF: 11 - 22 mrad, ADF: 22 - 98 and HAADF: 99 - 200mrad. The arrows point to a stacking fault in the matrix. Zone axis: [1¯10] .In the present study, the low precipitate density and high dilutions achieved hindered anevaluation of solubility at temperatures higher than 1873 K. Precise determination of solubilityat high temperatures is challenging for nanometric sources given the relatively poor lateralresolution of SIMS and the transient nature of the B concentration profiles [39], with fastdiffusion due to steep gradients around the precipitates. In addition, precipitates stabilizedby extended defects were present in layers implanted with 1 at.% B even after annealing at2073 K for 1 . × s (see figure 5 (a)), although the B concentration measured at those depths12 ciPost Physics Submission Figure 9: ABF and HAADF STEM images obtained simultaneously with a convergence angleof 22 mrad from a sample implanted with 2 at.% B at 673 K and annealed at (a) 1973 Kand (d) 1773 K. Collection angles for ABF: 11 - 22 mrad, and HAADF: 99 - 200 mrad. (b)and (e) Fourier transforms of the images in (a) and (d), respectively. (c) and (f) Simulateddiffraction patterns mirrored perpendicularly to { } B C and multiple scattering justifythe additional spots and streaks observed in (b) and (e). Reflections of the 3C-SiC matrixare encircled for clarity. 13 ciPost Physics Submission Figure 10: Concentration profile with 2 at.% B in as-implanted state and after annealing at1873 K. (b) Cross section of the same sampleas observed by HAADF with a convergence angle of 31 mrad and collection angles of 99 -200 mrad.( ∼
200 nm) was only 10 at . cm − . Therefore, dissolution rather than diffusion may be therate-controlling mechanism defining the concentration profiles at high temperatures, i.e., the3C-SiC matrix may be rapidly depleted of B by fast diffusive fluxes in the neighborhood ofslowly dissolving B C precipitates. In this case, the concentration measured immediatelybelow the precipitation layer after relatively long heat treatments may, in fact, be significantlylower than the equilibrium solubility at the annealing temperature. Annealing after implantation led to precipitation of B C that trapped B up to 1773 K. There-fore, the low solubility induced stable concentration profiles and resulted in apparently low Bdiffusivity. The precipitates adopted the (0001) B C // { } and (cid:104) (cid:105) B C // (cid:104) (cid:105) ciPost Physics Submission orientation relation with the matrix and exhibited planar defects on { } B C . Strain gen-erated by precipitate growth induced the formation of stacking faults and dislocations in3C-SiC that contributed to lower the overall structural quality of the crystal. The presence ofextended defects and precipitates may mask the presence of an IB in 3C-SiC and total elimi-nation of these structures is challenging due to the locked configurations adopted. Correlationbetween SIMS concentration profiles and STEM observations enabled to infer that the solu-bility of B in perfect 3C-SiC crystals is ≥ at . cm − at 1873 K. Our results suggest thatfor IB formation, alternative methods for introducing high concentrations of B into 3C-SiCshould be sought. Acknowledgements
The authors acknowledge the Norwegian Center for Transmission Electron Microscopy, NORTEMNational Infrastructure (Grant agreement 197405/F50)
Funding information
The work reported here has been undertaken as part of the projectSUNSiC - Efficient exploitation of the sun with intermediate band silicon carbide funded bythe Research Council of Norway (Grant Agreement 229711/O20).
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