GaP

Article pubs.acs.org/crystal

Cite This: Cryst. Growth Des. XXXX, XXX, XXX−XXX

Effective Suppression of Antiphase Domains in GaP(N)/GaP Heterostructures on Si(001) Alexey D. Bolshakov,† Vladimir V. Fedorov,‡ Olga Yu. Koval,‡ Georgiy A. Sapunov,‡ Maxim S. Sobolev,‡ Evgeniy V. Pirogov,‡ Demid A. Kirilenko,†,§ Alexey M. Mozharov,‡ and Ivan S. Mukhin†,‡ †

ITMO University, Kronverkskij 49, 197101, St. Petersburg, Russia St. Petersburg Academic University, Khlopina 8/3, 194021, St. Petersburg, Russia § Ioffe Institute, Politekhnicheskaya 29, 194021, St. Petersburg, Russia

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‡

ABSTRACT: III−V planar semiconductor heterostructures based on GaPN alloy with a nitrogen concentration up to 2.12% were grown on Si(001) by plasma assisted molecular beam epitaxy. Dependence of nitrogen incorporation on the growth conditions and its effect on the crystal structure were investigated via analysis of X-ray diffraction, transmission electron microscopy, and Raman spectroscopy data. Continuous redshift and a substantial increase in intensity of the photoluminescence emission spectra were observed upon increase of nitrogen content. The effect of antiphase disorder in GaP buffer on the GaPN epilayer properties was studied. It was found that antiphase boundaries, protruding from the GaP/Si to the GaPN/GaP heterointerface, change their orientation and self-annihilate in the dilute nitride layer even with a low (0.5%) nitrogen content.

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INTRODUCTION

incorporation and to reach high crystalline quality, an optimization of the growth parameters should be carried out. One of the main problems during the heteroepitaxial growth of III−V alloys on Si apart from the crystal lattice misfit and thermal expansion mismatch14 is the formation of heterovalent interface,15 while nonwetting conditions leads to the Volmer− Weber three-dimensional (3D) islands formation at the initial growth stage.16 Growth of such domains will be out of phase in terms of lattice occupation with group-III and -V elements, and their coalescence will lead to the formation of antiphase boundaries (APBs) considered as nonradiative recombination centers.17 Moreover, crystal polarity and the sign of the second-order optical susceptibility tensor χ(2) in antiphase domains (APDs) will change.18,19 Thus, spatially controlled nucleation of the APDs in III−V films allows steering of its second-order nonlinear optical response as it was shown in ref 20 for GaP films. However, it is inevitable to fully avoid formation of APBrelated defects or microtwining due to the low growth temperatures usually chosen for the nucleation layers.21,22 For the best reported GaP/Si(001) structure, the thickness of a defect region can be decreased down to 50 nm.23,24 This value closely approaches the critical thickness of pseudomorphic GaP on Si (∼90 nm),25 increasing probability of threading and

One of the bottlenecks limiting development of cost-effective optoelectronic solutions today is the lack of technology for integration of direct bandgap III−V materials on Si platform.1 Many approaches are still being studied including growth on metamorphic graded III−V buffers2 and synthesis of selfassembled III−V nanostructures allowing fast relaxation of mechanical stress during growth of lattice mismatched materials.3,4 Despite several advances in the field, still no scalable production technologies exist today. Gallium phosphide is a III−V semiconductor having an indirect bandgap of 2.26 eV and small lattice mismatch (0.37%) with silicon.5 Dilution of GaP with nitrogen leads to a significant decrease of the lattice parameter6,7 and affects the band structure leading to sufficient bowing and occurrence of a direct bandgap at a concentration of 0.6%.8−10 Importantly, GaN0.02P0.98 solution is lattice-matched to Si,11 making GaP alloys promising materials for realization of metamorphic heterostructures grown on Si. It is well-known that epitaxial growth with molecular beam epitaxy (MBE) and metal−organic chemical vapor deposition (MOCVD) is a nonequilibrium process, allowing synthesis of metastable crystal phases. An example of such material is GaPN alloy, whose synthesis is difficult due to low miscibility of GaP and GaN and competition of N and P due to a big difference in their covalent radia.12 To date, the highest homogeneity of these alloys can be reached via use of MOCVD.13 In order to overcome problems of group V atoms © XXXX American Chemical Society

Received: March 1, 2019 Revised: May 8, 2019 Published: June 18, 2019 A

DOI: 10.1021/acs.cgd.9b00266 Cryst. Growth Des. XXXX, XXX, XXX−XXX

Crystal Growth & Design

Article

Figure 1. Schematic representation of the synthesized heterostructures. The arrows with the corresponding comments designate change of GaPN layer growth conditions from sample to sample. equipped with a valved phosphorus cracker source (Tcracker = 900 °C) and Riber RF-plasma nitrogen source (13.56 MHz). Reflection highenergy electron diffraction (RHEED) images were taken in situ during sample growth to control its crystal structure and surface morphology. To monitor the group-III and -V element flux, beam equivalent pressure (BEP) was measured with the conventional ion gauge. BEP of Ga flux was kept at 2.2 × 10−7 Torr, which corresponds to a 0.58 μm/h growth rate of planar GaP. It was established that the stoichiometric P/Ga BEP ratio in our MBE setup was equal to 6 at a growth temperature (Tgrowth) of 640 °C. Thus, to keep the P-rich condition during GaP growth, its flux was set to 3.5 × 10−6 Torr, i.e., V/III ratio of 16. Prior to the load into the MBE chamber, Si(001) wafers were cleaned using a modified Shiraki method.28 The substrate temperature was monitored with a thermocouple, calibrated via direct comparison with an optical pyrometer. Anisotropy of the adatom diffusion on the Si(001) with its preferential direction along the dimer rows highly affects the nucleation and growth rate of GaP antiphase domains.29 To obtain single domain GaP films, vicinal Si(001) wafers with 2−6° miscut in the direction are conventionally used.30 This morphology promotes formation of Si surface steps with an even number of monolayers (bilayer thick steps) during substrate annealing or during the homoepitaxial growth of the Si buffer layer.31,32 In contrast to the above-mentioned approach, to intentionally induce antiphase disorder in GaP buffer layers, we used vicinal Si(001) substrates with a 4° miscut in the direction, which corresponds to the dimer rows orientation aligned by ±45° relative to the surface step.33 The chosen substrate misorientation inhibits formation of bilayer surface steps during the native oxide removal at 950 °C.34 Thickness of the Si(001) monolayer (ML) corresponds to half of III−V material ML; thus, in the case of uniform heterointerface formation, areas nucleated on the neighboring surface steps on Si(001) will be out of phase.35 As it was shown for the MOVPE grown GaP on Si(001),36 and as it will be shown further in this work, antiphase disordered GaP layers with

misfit dislocation generation. Thus, the possibility of thin defect-free GaP buffer fabrication remains questionable. Recently, it was shown that APBs formed in the GaP/Si heterostructure self-annihilate in the AlGaP overlayer.24 TEM studies demonstrated that the APBs change their orientation and kinks from {110} to {111} planes in the AlGaP layer. In contrast to the APB lying in the vertical {110} planes protruding the layer, the latter higher-indexed planes completely self-annihilate at their intersection. So it becomes possible to tailor the defect formation not only with the growth conditions but also with the layer composition. Compared to GaP, AlP has a slightly larger indirect band gap (2.45 eV) and lower electron mobility.26 Also, AlxGa1−xP ternary alloys can be prone to defect formation, due to high chemical reactivity of Al, and as a result, electron transport can be negatively affected.27 Thus, the study of different composition IIIphosphide buffer layers is very promising. As was reported previously, the appearance of mismatch dislocations is suppressed with N incorporation into GaP on Si substrate due to the higher rigidity of the synthesized alloy.12 Thereby, GaP(N) solution represents a perspective platform for the synthesis of low-defect direct bandgap heterostructures integrated on Si. This report is aimed at the study of the APBs behavior in direct bandgap GaPN layers grown on GaP/ Si(001) virtual substrates. Influence of the dilute nitride layer composition on its structure, optical properties, and morphology is investigated.

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EXPERIMENTAL SECTION

Dilute nitride phosphide heterostructures were grown using a solid source Veeco GEN-III plasma-assisted MBE (PA-MBE) machine B



Article

DRON-8 X-ray diffractometer. The XRD spectra were processed using a FullProf software package. To study effect of N incorporation on crystalline quality of the GaPN ternary alloys, their structural homogeneity and optical properties were studied with Raman microspectroscopy and photoluminescence (PL) spectroscopy with Horiba LabRam HR800 in backscattering geometry with the use of an optical microscope for excitation and collection. The measurements were carried out at room temperature with laser excitation energies below and above the GaPN band gap using a 532 nm diode-pumped solid-state (DPSS) and 785 nm external cavity diode (ECDL) lasers. Heterostructures surface morphology was studied with scanning electron (SEM) (Zeiss SUPRA 25−30−63) and atomic-force (AFM) microscopy (Bruker Bioscope Catalyst SPM). High-resolution transmission electron microscopy (HRTEM) studies were carried out using a Jeol JEM-2100F microscope, with an operating voltage of 200 kV. The samples for HRTEM studies were prepared by a conventional technique involving mechanical polishing with subsequent ion milling by Ar+ at the last stage.

vertically aligned APBs protruding to the surface can be obtained by this approach. Observation of the (2 × 2) RHEED reconstruction pattern after Si(001) substrate annealing confirms a double-domain (2 × 1) + (1 × 2) silicon surface structure with both possible dimer rows orientation.37 To initiate uniform formation of P−Si bonds along the interface, the substrates were preexposed under phosphorus flux for 60 s, prior to the nucleation of GaP. To prevent three-dimensional island formation, two-staged growth with the use of a seeding layer prepared with the migration enhanced epitaxy (MEE) technique based on the alternate exposure to the group-III and group-V elements was used.38,39 During MEE, the substrate temperature was set to 480 °C, and the quantity of MEE cycles was in the range of 20−150 depending on the sample number. The Ga exposure period during the single MEE cycle was set to the growth time of 1 ML GaP. After formation of the seeding layer, the growth temperature was increased to 680 °C for growth of GaP buffer in a conventional MBE regime. During GaPN layer formation, the N plasma was controlled via monitoring of its integral optical emission intensity (I) with silicon photodiode. Reflected power from the RF nitrogen plasma source was kept at a minimum (below 1 W) by an automatic impedance matching network. A schematic images of the grown structures are presented in Figure 1. To study formation of the APBs and their effect on morphology of the GaP buffer layer surface, we synthesized sample S1 representing only the GaP/Si(001) structure without the dilute nitride layer prepared with the two-stage (MEE + MBE) technique. Next, samples S2−S6 represent GaP/GaPN/GaP/Si heterostructures. It has been shown earlier that to obtain nitrogen incorporation a moderate decrease of the Tgrowth is required (50−100 °C).40,41 At the same time, higher growth temperatures promote uniform distribution of N atoms and suppress nonradiative recombination.13 To study the effect of the growth temperature on nitrogen incorporation in GaPN epilayers, the substrate temperature was reduced in comparison to the GaP MBE growth for samples S2 (Tgrowth = 640 °C) and S3 (Tgrowth = 600 °C). GaPN layers of samples S2 and S3 were grown at the same P/Ga BEP ratio as the underlying GaP buffer layer with the following operational parameters of RF plasma source: N flow rate and forward RF power was set to 0.5 sccm and 550 W correspondingly, providing an optical emission brightness of I = 3.3 au. Further optimization of dilute nitride layer growth conditions was performed via synthesis of heterostructures with micrometer-thick GaPN layers (samples S4−S6) which allows increasing of the accuracy of the X-ray diffraction (XRD) studies. It is known that growth of dilute nitride layers with the use of an RF plasma source can be prone to defect formation due to the ionized species in the active nitrogen flux.12,42 To study the aforementioned effect, samples S4 and S5 were grown at a reduced plasma brightness (I = 2.1 au) with a nitrogen flow rate of 0.2 sccm and forward RF power of 450 W and an additionally decreased growth temperature of 560 °C (S5) and 540 °C (S4). It was reported that due to competitive nature of N and P incorporation into the group-V sublattice, growth of dilute nitrides can require a decrease of the P/Ga flux ratio compared to pure GaP layers.12 The interplay between P and N fluxes on the nitrogen incorporation was studied in sample S6 grown with a P/Ga BEP ratio reduced down to 5.5 while keeping the same operational parameters of the RF plasma source as for samples S2−S3 (I = 3.3 au) and Tgrowth at 600 °C - same as for S3. Thus, we investigate formation of the GaPN layer depending both on growth temperature, activated nitrogen, and phosphorus flux values. To prevent GaPN exposure on air, these layers were later overgrown with a thin GaP layer at Tgrowth = 680 °C. An increase of the growth temperature was carried out in order to improve the crystalline quality of GaPN typically observed after annealing of this material. After the growth stopped, the structures were cooled down to 400 °C under the P flux to avoid thermal decomposition. Structural characterization was performed with X-ray diffractometry via measurement of symmetrical Θ−2Θ scans under Cu K-alpha radiation (Kα1 = 1.54439 Å and Kα2 = 1.54056 Å) wavelength with a

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RESULTS AND DISCUSSION Morphology. Typical surface morphology of the GaP buffer layer (sample S1) and sample S2 containing GaPN layer is presented in Figure 2a,b. One can note the presence of the

Figure 2. SEM images (plane-view) of the III-phosphide/Si heterostructure surface: (a) sample S1 - double staged GaP buffer layer, (b) sample S2 with GaPN layer grown on the double staged GaP buffer.

dark curves with an irregular shape on the both obtained SEM images. As was discussed in the experimental part, use of vicinal Si(001) substrates with a 4° miscut in the direction inevitably leads to the nucleation of the APDs during the heteroepitaxial growth of III−V material. It is known that APBs tend to minimize their surface area due to the interface energy cost, which leads to the formation of trenches bordering the APDs on the sample surface.14 So we assume that the above-mentioned SEM contrast features observed in Figure 2a,b correspond to the APBs in the phosphide layer.43 Analysis of the images in Figure 2 allows us to distinguish GaP domains with the predominant lattice orientation occupying a large surface area (especially pronounced for sample S2, visible as a brighter area in Figure 2b. Interestingly, despite a lower growth temperature, sample S2 (Figure 2b) demonstrates an at least five times larger mean lateral size of the domains compared to the GaP buffer layer (Figure 2a). Such a distinction in morphology can be caused by either the difference in the growth mechanisms or difference in the APB interface energy values for GaP and GaPN layers, which affect the propagation of the APBs. More detailed analysis of the APBs formation and propagation was performed using TEM microscopy and presented in the next sections. C



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Figure 3. RHEED pattern images taken along the Si azimuthal direction after deposition of each heterostructure layer of sample S6: (a) MEE LT lower GaP buffer (spotty RHEED pattern, low crystal coherency), (b) MBE HT upper GaP buffer, smoother surface, (2 × 2) RHEED pattern (antiphase regions), (c) GaPN broadened dashed reflexes (surface roughening), (d) GaP capping, (4 × 2) superstructure (surface smoothening, single domain GaP growth).

Figure 4. (a−f) Experimental Θ−2Θ XRD scans of the samples across the Si(004) symmetric Bragg reflection and theoretical fits approximated with the pseudo-Voigt function corresponding to Si, GaP, and GaPN lattices.

appears (1/4 fractional-order reflexes are marked with arrows in Figure 3d) acting as evidence of a surface smoothening and single domain GaP growth. XRD Study. Figure 4 shows Θ−2Θ XRD scans of the samples taken along the Si(004) symmetric Bragg reflection and results of the fitting using Pseudo-Voigt function. As can be seen in Figure 4, no thickness oscillations were observed, indicating a high interface roughness.45 Beside the narrowest and brightest Si(004) reflex, one can find the broadened diffraction peaks corresponding to the epitaxial layers. As can be seen from the data for sample S1 with a single GaP epilayer, the broadened peak appears at a lower diffraction angle compared to Si(004) and can be attributed to (004) Bragg reflection of the pseudomorphic GaP layer.46 Position of the GaP peak for most of the samples appears at a lower angle compared to relaxed GaP (004) (Θ = 68.853 deg.), which indicates an increase of the interplanar distance of the GaP lattice in the [001] direction induced by in-plane compressive strain. However, the observed tetragonal lattice distortion strain is less than for the case of strained GaP layers pseudomorphically grown on Si(001). Since the buffer layer

Additional information on both structure and morphology was obtained from the RHEED patterns taken after the growth of each consecutive heterostructure layer. In Figure 3 one can find RHEED pattern images taken along Si direction during sample S6 growth. A spotty RHEED pattern observed in Figure 3a is a characteristic feature for the MEE-grown GaP epilayer and indicates low crystal coherency during the low temperature growth.21 A GaP layer grown on the MEE-seed at high temperature demonstrates the presence of the streaky RHEED pattern and appearance of 1/2 fractional-order streaks (marked with arrows in Figure 3b along azimuthal directions). A streaky RHEED pattern indicates smoothening of the sample surface, while a double-domain (2 × 2) RHEED pattern indicates the presence of antiphase regions with both possible orientations of the GaP lattice. During growth of the GaPN layer, streaks on the RHEED pattern change to the broadened dashed-shaped Bragg reflexes corresponding to surface roughening more likely due to surface faceting44 and transition to three-dimensional growth (Figure 3c). The reflexes again become narrower and streaky during the capping GaP layer growth, and importantly the (4 × 2) superstructure D



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Table 1. Relation between the Growth Parameters and Calculated N Fraction in GaPN Layer lattice parameter a, Å sample

plasma brightness I(a.u.)

Tgrowth GaPN (°C)

P/Ga BEP ratio

S1 S2 S3 S4 S5 S6

3.3

640 600 540 560 600

16

2.1 3.3

5.5

GaP 5.470 5.480 5.478 5.466 5.448 5.467

± ± ± ± ± ±

0.001 0.001 0.001 0.001 0.001 0.001

GaPN 5.468 5.461 5.442 5.436 5.433

± ± ± ± ±

0.001 0.001 0.001 0.001 0.001

N content, % 0 0.47 0.78 1.59 1.92 2.12

± ± ± ± ±

0.01 0.01 0.01 0.01 0.01

Figure 5. Room-temperature Raman spectra of the studied heterostructures (S1, S2, and S5) with N content x = 0, 0.47, and 1.92%, obtained in backscattering geometry with (a) 532 nm and (b) 785 nm laser excitation. The spectra are normalized to the intensity of the TO-LO mode of Si.

Growth temperature of 560°C (S5) results in 1.92% of incorporated nitrogen, while a further reduction of Tgrowth down to 540°C (S4) leads to its decrease to 1.59% and additional Bragg peak broadening, demonstrating the degradation of the crystalline quality. A reduced amount of the incorporated N in the latter case can be caused by an effective increase of P adatoms supersaturation on the sample surface at this temperature. The dramatic role of the interplay between the P and N fluxes on the nitrogen incorporation is demonstrated with sample S6 with a P/Ga BEP ratio decreased just below the stoichiometric value for GaP. The highest nitrogen content of 2.12% obtained in this sample even at an elevated temperature of 600°C can be explained by an increase in number of P vacancies which can be occupied by the N adatoms during GaPN growth at an understoichiometric P/Ga BEP ratio. On the basis of the experimental data, we can make the following conclusions on the optimal growth conditions for the N incorporation. Our observations confirm the findings of earlier work on GaNAs51 and GaPN alloys:40,12 a moderate decrease of the growth temperature compared to GaP growth by 50−80°C favors effective N incorporation; however, a large temperature drop of 100°C can negatively affect both N incorporation and crystalline quality. Another important parameter is the P/Ga flux ratio affecting the interplay between the P and N incorporation. For the maximum amount of the incorporated N, it is desirable to keep the P/Ga flux ratio close to the stoichiometric value. Analysis of the Composition and Structural Disorder. In Figure 5, one can find the Raman spectra of the GaP/Si(001) buffer layer (sample S1) and heterostructures with dilute nitride epilayers with low (sample S2) and high (sample S5) nitrogen content obtained with 532 and 785 nm laser

thickness exceeds a critical value for GaP/Si of 90 nm, the lattice stress could be relaxed via formation of threading dislocations.25 We assume that GaP lattice stress is only partially relaxed due to the high density of APBs, acting as a barrier for dislocations propagation.47,48 XRD scans of the samples containing a dilute nitride layer (S2−S6) show an additional broad peak, which we attribute to the (004) Bragg peak from the GaPN layer. It is clearly seen that position of the peak varies from sample to sample due to a change in the lattice constant with the layer composition. Estimation of the angular distance between the (004) GaP and GaPN layer diffraction peaks was used for calculation of the incorporated N concentration in the GaPN layer according to Vegard’s law.49,7 When the peak positions of GaPN and Si reflex coincide corresponding to their lattice matching, the N concentration was assumed to be equal to 2%.12 Overlapping of the Bragg reflexes from GaP and GaPN layers makes it difficult to precisely distinguish their peak position. However, fitting of the obtained data indicates that the position of the GaPN Bragg peak is shifted toward larger angles compared to bulk GaP (004) and approaches Si(004) reflex. Broadening of the GaPN Bragg reflex is observed with an increase of N content and may be caused by the large number of structural defects related to the formation of N pairs or N−N split interstitials.50 A summary of the XRD analysis and calculated N concentration values are presented in Table 1. Effect of the Tgrowth on N incorporation is clearly visiblea decrease of the temperature from 640 °C down to 600°C leads to an increase of the N content from 0.47% up to 0.79% in samples S2 and S3 correspondingly. This tendency is confirmed to a certain point with the samples S4 and S5 grown at additionally reduced Tgrowth and RF plasma brightness (I = 2.1 au). E



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Figure 6. Room-temperature Raman spectra of the studied heterostructures obtained with (a) 532 nm and (b) 785 nm excitation. The spectra are normalized to the intensity of the LO1 - GaP-like mode.

possible explanations for the appearance of an additional lowenergy shoulder (indicated as SO) located just to the right of LO1 peak: disorder-activated Raman scattering from the zoneedge phonons56 or surface optical phonons (SO).57 Since the surface-to-volume ratio for the epilayers is small in comparison with nanostructures, most likely this peak can be attributed to disorder-activated scattering. This peak becomes more prominent in the sample with high nitrogen content and can be caused by various structural imperfections such as APBs, dislocations, or layer inhomogeneity.58 The presence of N atoms in GaP matrix causes the emergence of additional Raman-active vibrational mode visible under 532 nm excitation, located between the TO1 and LO1 GaP-like modes (marked as X) (Figure 6a). The occurrence of this mode was previously discussed in ref 59 and was attributed to the alloying effect in ternary alloy semiconductors such as GaPN,60 GaPAs,61,62 and InGaAs.58 This peak can be caused by disorder activated optical phonons in the region of N atoms clustering because its intensity increases with increasing N content52 (Figure 6a). As can be seen in Figure 6, a composition dependent redshift of GaP LO1 and TO1-modes is observed with an increase of N more likely caused by the N-induced local lattice strain in GaPN epilayers. The shift of LO1 GaP-like vibrational mode as a function of N content was investigated in detail for all of the samples (see Figure 7). Frequency of this mode in unstrained bulk GaP is 403 cm−1,62, 63 while according to our experimental data (see Figure 6) this mode is more than 1 cm−1 shifted for the sample S1 with a strained GaP epilayer without N. Analysis of the experimental data obtained with both types of excitation shows that the LO1 GaP-like mode shifts linearly with N concentration as shown in Figure 7. GaP/GaPN lattice mismatch-driven frequency shift was also calculated theoretically assuming 1.13 cm−1/dx (%)64 and presented in the graph with a dotted line. Good agreement between the experimental data and the theoretical fit allows us to conclude that analysis of the LO1−GaP vibrational mode position can be used to determine the N concentration in GaP1−xNx solid solutions. PL Measurements. PL spectra of the studied heterostructures with GaPN layers obtained at 300 K demonstrate the presence of typical for GaPN alloys broad PL emission band (Figure 8).64 Investigation of the pure GaP sample without N demonstrates the absence of the PL signal,

excitation. With an exception of N-rich GaPN sample (S5), the most intense Raman scattering peak is located at a frequency of ∼520 cm−1 and corresponds to the degenerated transverselongitudinal optical (TO-LO) phonon mode in Si substrate. At high N-content, direct band gap transition occurs in GaPN,52 and thus the appearance of a broadband signal corresponding to the high energy shoulder of the GaPN PL peak is prominent for sample (S5) under the above bandgap excitation at 532 nm. In addition, the TO-LO phonon mode of the Si substrate becomes barely visible. This effect can be explained by a higher optical absorption in the GaPN layer at a 532 nm wavelength compared with 785 nm, and thus a smaller light penetration depth and higher surface sensitivity. Because of the higher penetration depth, second-order Raman scattering in Si is visible under 785 nm excitation (Figure 5b); namely, a flat plateau of the two phonon transverse optical mode (2TO) is observed at 920−1000 cm−1,53 and a low-frequency peak can be found at ∼300 cm−1, corresponding to the two-phonon transverse acoustic mode (2TA) at the edges of the Brillouin zone.54 The Raman spectra region between 350 to 420 cm−1, highlighted with a dotted line in Figure 5a,b, corresponds to the Raman scattering on the GaP1−xNx/GaP layers and requires detailed analysis. An enlarged view of this region is presented in Figure 6. A strong Raman signal located at ∼400− 405 cm−1 is observed in all of the spectra. This peak can be attributed to GaP-like LO1 phonon mode matching the selection rules in the backscattering geometry. A weak peak corresponding to the GaP-like TO1-mode forbidden in this geometry appears at 367 cm−1. Occurrence of this peak relates to violation of the Raman selection rules due to the disturbance of the backscattering geometry due to the finiteness of the spectrometer aperture or antiphase disorder in thin films. The broadening of Raman peaks can be attributed to both the size effects and the strain generated at the interfaces between the epilayers. One can notice that Raman spectra obtained using an excitation wavelength of 532 nm (Figure 6a) demonstrate the appearance of additional scattering peaks located in between the LO and TO modes which are nonvisible with 785 nm excitation (Figure 6b). As was discussed earlier, this effect can be attributed to the higher surface sensitivity and above bandgap near-resonant scattering conditions.55 There are two F



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9a,b correspondingly. As was discussed in refs 65 and 66, these parameters were chosen to provide contrast between APDs.

Figure 9. TEM images of (a) (S1) GaP layer grown on Si(001) (white arrows indicate MEE-grown GaP interface layer, (b) (S2) GaP/GaPN/GaP/Si heterostructure (note that APB propagation is suppressed in GaPN layer). Figure 7. A frequency shift of the GaP-like (LO1) mode as a function of N content in the GaPN layer. Green stars correspond to the excitation with 532 nm laser, red circles with 785 nm laser. Dotted line - calculated frequency shift caused by the mismatch-driven stress.

The mosaic structure of the GaP film with an average lateral grain size of 50 nm is clearly seen in Figure 9a. Contrast between the grains is only visible in dark field mode using (002) reflection, which also confirms that the grain boundaries are APBs.67,68 Antiphase domain boundaries of the GaP buffer layer in the S1 tend to lie close to the {110} planes and protrude to the sample surface. In the GaP/Si heterointerface region, one can also observe contrast marked with the white arrows in Figure 9a, which we associate with APB in the MEEgrown GaP interface layer.69,70 Thus, the chosen Si(001) substrate miscut and two-staged growth regime allow us to effectively suppress APB kinking and self-annihilation and obtain the antiphase disordered GaP buffer to study further behavior of APBs in the dilute nitride layer. One can note that a much lower density of the APBs protruded to the sample surface is observed in GaP/GaPN/ GaP/Si(001) heterostructure (S2) compared to GaP/Si (S1) (Figure 9b). The characteristic lateral grain size on the surface of the samples S1 and S2 is consistent with the SEM observation presented in Figure 2 and proves that surface morphology relates to the antiphase domain structure in the III−V layer. As can be seen, the APB density in the GaP buffer layer of S1 and S2 are comparable; however, the number of APBs decreases dramatically in the GaPN layer. To prove the presence of the APBs in the S2 buffer layer the high-resolution TEM (HRTEM) image of the GaP/Si interface region was obtained; see Figure 10a. Formation of vertical APB lying in the {110} plane is clearly seen. To visualize the APB, we carried out Fourier space filtering of the

Figure 8. Room-temperature PL spectra of the studied heterostructures S2, S3, and S6.

indicating that nitrogen incorporation increases the quantum yield of the PL, due to the indirect-to-direct band structure transition. The PL band shape depends on the N content in the GaPN epilayer, and its peak position varies from ∼630 to ∼680 nm, corresponding to a change of the bandgap related to the increase of the N content from 0.3 to 2.12%, which correlates well with the results of the XRD data analysis. The broad PL response can be caused by the low crystalline quality of the heterostructures (for example, the presence of defects). Analysis of the PL peak position demonstrates that it is red-shifted in comparison with the expected bandgap of GaPN alloy with the given N content estimated with the use of the BAC model:7 Eg(2.12%) = 1.95 eV (636 nm); Eg (0.78%) = 2.07 eV (600 nm); Eg (0.47%) = 2.11 eV (588 nm), which confirms the disorder-induced band-tail nature of the PL response.10 APB Formation and Propagation. TEM microstructural analysis of the synthesized dilute nitride heterostructures allows us to get the details of the APB formation and propagation across GaP and GaPN epilayers. Cross-section dark-field TEM images taken along the direction using g = (002) reflection of GaP/Si(001) (S1) and GaP/GaPN/ GaP/Si(001) (S2) heterostructures are presented in Figure

Figure 10. (a) HRTEM image of GaP/Si(001) interface region of sample S2, (b) Fourier space filtered HRTEM image - contrast corresponds to the periodicity of the (001) lattice planes, white arrows indicate APBs. On inset: corresponding FFT image - dashed circles indicating selected frequency components in the filtered image. G



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the relative energy of the {111} facets with respect to the {100}, which was observed earlier for the GaAsN alloys with a high nitrogen content.44 Therefore, it cannot be excluded that changes in the GaPN growth mechanism also affect the propagation of APB. We assume that the surface mobility of Ga should not be affected by N adatoms, since impingement of N flux is low compared to the GaP growth rate.73 However, it is known that a decrease in the substrate temperature which is required for the dilute nitrides growth significantly affects the adatom mobility; the surface diffusion length of Ga on the GaP (001) surface decreases by more than an order of magnitude when T decreases by 150 °C.74 In turn, the diffusion length of N adatoms is considered to be small in comparison with P and Ga.73,75 The above-mentioned factors together with a change in the ratio of the {111} and {001} planes surface energy in the GaPN layer leads to a higher density of nucleation events and promotes the transition from planar growth to the nucleation of three-dimensional islands, which facilitates the reorientation of APBs. Our findings indicate that GaPN interlayers or GaP/GaPN superlattices can be used for effective suppression of formation and propagation of APBs and threading dislocations during III−V growth on Si.

image (see Figure 10b). Frequency components corresponding to the structure periodicity in the [001] direction were chosen, allowing us to visualize contrast related to GaP and Si(001) lattice planes. GaP planes shifted by the half period at the APB region can be seen (indicated with arrows in Figure 10b). Thus, according to the performed TEM analysis, we conclude that APBs start to self-annihilate in the GaPN layer with a low nitrogen content of 0.5%. Moreover it is turned out that the efficiency of the APB annihilation depends on the N concentration. Cross-section dark-field TEM images of GaP/ GaPN/GaP/Si heterostructure (sample S5) with N content of 1.92% is presented in Figure 11a,b. It is clearly seen that at the

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Figure 11. TEM images of GaP/GaPN/GaP/Si heterostructure with 1.92% nitrogen content (S6). (a) Cross-section of the structure, (b) close-up view on the GaPN/GaP interface region. Arrows indicate areas of threading dislocation generation in the GaPN layer.

CONCLUSIONS The study of the GaP/GaPN/GaP/Si (100) heterostructures synthesized via PA-MBE has been carried out. In situ analysis of the RHEED pattern has been used to qualitatively characterize the growing layer surface morphology. It is found that during the GaPN layer growth the RHEED pattern possessed a broadened dashed shaped Bragg reflexes corresponding to surface roughening. Analysis of the XRD spectra of the samples allowed us to estimate the lattice constant of the grown GaPN layer and consequently the N molar fraction in the alloy. Further confirmation of the layer composition was obtained via analysis of the LO1 GaP-like Raman scattering mode position which shifts linearly with N concentration. An influence of the growth parameters on incorporation of N in GaPN layers is investigated. Effect of the growth temperature on the GaPN layer composition is demonstrated. Decrease of the GaPN growth temperature from 640 °C down to 560 °C leads to an increase of the N content. A further decrease to 540 °C leads to a slight decrease of the N fraction. The obtained experimentally almost 3-fold increase of the N fraction with the 3-fold decrease of the P BEP corresponds to P−N competition during the growth. Crucial influence of the GaPN layer on the APB propagation is demonstrated. TEM investigation of the samples shows that the APBs protruding vertically through the GaP buffer layer effectively annihilate after reorientation in the GaPN layer. The latter result is very promising for direct integration of high crystalline quality GaPN direct band gap alloys on Si.

GaP/GaPN interface APBs kink and change their orientation from vertical {110} to inclined planes. APBs in the dilute nitride are lying at angles of 53−55° and 34−36° relative to the (001) indicating their {111} and {112} orientation correspondingly. As a result, APBs effectively self-annihilate at the distance comparable with the lateral size of APD. As can be seen in TEM image (Figure 11b), all APBs observed previously in the GaP layer are completely annihilated on the first 100− 200 nm of dilute nitride layer thickness. A large thickness of the GaPN layer and its lattice mismatch with GaP buffer lead to the formation of misfit dislocations. In some areas, marked with arrows in Figure 11b, one can see how threading dislocation arises at the point of APB annihilation. Detailed analysis of the S5 TEM images shows that the APB density in the GaP buffer exceeds 1010 cm−2, while in the GaPN layer mainly threading dislocations are present with a density of 108 cm−2. Dislocations density in the topmost GaP layer additionally decreases by several times, meaning that during the transition to the GaP capping layer, a significant part of the threading dislocations become blocked at the GaP/GaPN interface. We assume that APB reorientation process can be driven by the difference in formation energy between {111}, {112}, and {110} APB planes, which changes with the addition of N in the III-phosphide compound. As was shown in refs 71 and 72 {110} and {112}, APB planes are charge-neutral and have an equal number of III−III and V−V bonds; thus transition to the {112} oriented APBs can be caused by reduction of the kinetic barrier for the APB kinking in the dilute nitride layer.66 On the other hand, {111}-oriented APBs are nonstoichiometric, and depending on domain polarity consist only of III−III or V−V bonds. Thus, the demonstrated {111} reorientation is unobvious; a possible reason is the tendency for the {111} faceted growth of the dilute nitride layer due to the lowering of

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AUTHOR INFORMATION

ORCID

Alexey D. Bolshakov: 0000-0001-7223-7232 Notes

The authors declare no competing financial interest. H



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ACKNOWLEDGMENTS This work was carried out with the support of the Russian Science Foundation (Grant No. 18-72-00219). TEM characterizations were performed using equipment of the Federal Joint Research Center “Material Science and Characterization in Advanced Technology” supported by the Ministry of Education and Science of the Russian Federation (id RFMEFI62117X0018). M.S.S. and E.V.P. acknowledge support from the grant of Minobrnauki No. 16.9789.2017/BCh.

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