Quantum dots (QDs) epitaxially grown on Si are promising for monolithic integration of light sources on a Si photonics platform. Unlike quantum well (QW) lasers on Si, 1.3 μm InAs QD lasers on Si show similar threshold current to those grown on GaAs owing to their better dislocation tolerance. To date, research on dislocation-tolerant QDs has exclusively focused on materials emitting at telecom wavelengths. In this work, we report visible InP QDs on Si with photoluminescence (PL) intensity similar to their counterparts grown on GaAs despite high threading dislocation density (TDD). In contrast, visible InGaP QWs grown on Si with the same TDD value show 9× degradation in PL intensity compared to QWs grown on GaAs. The dislocation tolerance of InP QDs arises from their high density relative to TDD and the lateral carrier confinement that they provide. InP QDs on Si with bright PL are promising for low-cost light emitters and integrated photonics applications requiring monolithic red-light sources.

The silicon photonics platform possesses well-developed processes for fabrication of modulators, detectors, and both Si and SiN optical waveguides,1,2 but monolithic integration of light sources remains challenging.3 The 4% lattice mismatch between GaAs and Si leads to a typical threading dislocation density (TDD) of 106–108 cm−2,4 causing strong non-radiative recombination in GaAs epitaxially grown on Si. The last several years witnessed tremendous progress in the performance of 1.3 μm InAs QD light emitters monolithically integrated on Si as light sources for silicon photonics. Despite the high TDD, InAs quantum dot-in-a-well (QDWELL)5,6 active regions on GaAs/Si show high luminescence efficiency comparable to growth on GaAs due to the lateral carrier confinement of the QDs;7,8 once captured by a QD, injected carriers are no longer free to diffuse toward dislocations. QDs also prevent carriers from diffusing to free surfaces, making them promising for micro-scale light emitting diodes (micro-LEDs) and lasers.9,10 In a typical QDWELL, the InAs QD density (QDD) is ∼5 × 1010 cm−2, and thus, even when TDD is 108 cm−2, QDs can out-compete the dislocations for capture of carriers and enable efficient radiative recombination.11 In contrast, InGaAs QWs grown on GaAs/Si with such high TDD show 10× reduction in luminescence intensity compared to QWs grown on bulk GaAs. InGaAs QW lasers also suffer from an ∼60× increase in threshold current density (Jth) when grown on Si, as compared to the ∼2× increase for InAs QD lasers on Si.12,13

Despite the fact that III-V QDs can be designed to emit over wavelengths spanning from 0.6–6 μm,14–16 research on dislocation-tolerant QD emitters has exclusively focused on telecom wavelengths of 1.24–1.55 μm.17,18 Phosphide-based QDs integrated on Si emitting in the visible and near infrared (NIR) range could have a wide variety of applications including integrated optogenetics,19,20 biophotonic sensing,21 low-cost monolithically integrated micro-LEDs,22,23 and quantum optics.24,25 InP QD lasers on GaAs with a high QDD of 5 × 1010 cm−2 operate with a low Jth of 190 A/cm2, a high output power of >150 mW, and a high characteristic temperature of >69 K.26–28 The emission wavelength of InP QDs on GaAs is tunable from 640 nm to 750 nm and can be extended to >800 nm by alloying with As to form InAsP QDs.29,30 Despite impressive development in the growth of InP QD active regions on GaAs, there are no reports demonstrating dislocation-tolerant QD emitters in the visible-NIR regime on Si substrates at room temperature.

In this work, we show that InP QDs with visible-NIR emission on GaAs/Si with TDD = 3.3 × 107 cm−2 exhibit minimal degradation in photoluminescence (PL) intensity compared to samples grown on GaAs. In contrast, In0.48Ga0.52P (hereafter InGaP) QWs grown on GaAs/Si with similar TDD show a 9× reduction in PL intensity compared to QWs grown on GaAs, because carriers captured in QWs are free to diffuse laterally to dislocations. The high dislocation tolerance of InP QDs compared to InGaP QWs results from the four orders of magnitude higher QDD compared to TDD and the impeded carrier diffusivity due to the lateral confinement of the QDs.

All samples were grown by solid-source molecular beam epitaxy (MBE). We grew GaAs on GaP/Si (001) templates31 available commercially from NaAsPIII-V GmbH using a 4.3 μm thick GaAsyP1-y step-graded buffer with a grading rate of 1%/μm; high-resolution x-ray diffraction (HRXRD) showed that the GaAs cap layer was ∼100% relaxed at room temperature. GaAs/Si substrates were cleaved into several pieces and co-loaded with pieces of bulk GaAs for growth of QD and QW PL structures. We grew the PL active regions using a typical MBE growth window for indium-containing phosphide layers32 with a substrate temperature of 480 °C, a V/III ratio of 10–30, and growth rates of ∼0.2–0.5 μm/h. The QD samples had an additional layer of surface dots grown on top for atomic force microscopy (AFM), while we capped the QW samples with a thin 5 nm layer of GaAs to prevent oxidation of Al-containing layers. All phosphide layers were calibrated to be lattice-matched to GaAs through a combination of HRXRD and PL for quaternary compounds and through HRXRD for ternary compounds. The PL structures on both GaAs and GaAs/Si underwent post-growth rapid thermal annealing (RTA) in an AG 610 system at 700 °C–1000 °C for time between 1 s and 5 min to improve PL intensity, as is typical for MBE-grown phosphides.33 For QD samples, the surface InP and underlying AlGaInP layer were etched prior to RTA, exposing the GaAs capping layer (Fig. 1). For RTA experiments, the samples were placed on a Si wafer and covered with a GaAs wafer to prevent As desorption from the epitaxial GaAs capping layer. The RTA conditions were optimized using integrated PL intensity of the samples, with maximum enhancement in PL intensity for InP QDs and InGaP QWs observed at 750 °C for 5 min and 1000 °C for 1 s, respectively. The RTA conditions employed here improved the PL intensity of both QDs and QWs without any significant shift in the emission wavelength.

FIG. 1.

PL structure for InP QDs co-grown on GaAs and GaAs/Si. The surface InP QDs and 50 nm AlGaInP layers were etched away prior to RTA, and the 30 nm GaAs capping layer was removed prior to PL experiments.

FIG. 1.

PL structure for InP QDs co-grown on GaAs and GaAs/Si. The surface InP QDs and 50 nm AlGaInP layers were etched away prior to RTA, and the 30 nm GaAs capping layer was removed prior to PL experiments.

Close modal

We used a JEOL 2010F to conduct bright-field cross-sectional transmission electron microscopy (BF-XTEM) and an aberration-corrected JEOL 2200FS to conduct high-angle annular dark field (HAADF) scanning TEM (STEM) imaging, both at an accelerating voltage of 200 kV. Cathodoluminescence (CL) mapping of InP QD and InGaP QW samples was performed in a JEOL 7000F analytical scanning electron microscope at an accelerating voltage of 3 kV using a Gatan Mono CL system. Both CL and TEM studies were conducted on as-grown samples before RTA. After etching the GaAs capping layer and InP surface QDs, we performed steady state PL at room-temperature using a 532 nm diode-pumped solid-state laser with an incident power density of 5 W/cm2 and an Ocean Optics spectrometer.

Figure 1 shows a schematic of QD PL structures grown on GaAs and GaAs/Si substrates consisting of three monolayer (ML) InP QDs capped with an 8 nm 1.9 eV InGaP QW and surrounded by 200 nm-thick, 2.1 eV (Al0.33Ga0.67)0.52In0.48P (AlGaInP hereafter) barriers. We also grew 2.3 eV In0.49Al0.51P carrier blocking layers (InAlP CBLs) above and below the active region to reduce surface recombination losses. All layers are lattice-matched to GaAs except the InP QDs, which have an ∼3.7% compressive mismatch that drives self-assembly via the Stranski–Krastanov mechanism.34,35 In addition, we grew InGaP QW PL structures on both substrates with the same layer structure except without the InP QDs. Both the InP QD and InGaP QW PL structures possess a type-I band alignment.36,37 The calculated conduction and valence band offsets between InP and AlGaInP are ΔEc = 0.46 eV and ΔEv = 0.23 eV, respectively; for the InGaP and AlGaInP, the offsets are reduced to ΔEc = 0.13 eV and ΔEv = 0.06 eV, consistent with previously reported values.38 The ground state emission from the InP QDs blueshifts by 0.4 eV from the bulk bandgap energy of InP due to both compressive strain and quantum confinement effects.39,40

The TEM analysis of InP QD PL structures grown on GaAs and GaAs/Si shows high-density InP QDs capped with a smooth InGaP QW without nucleation of misfit dislocations around the active region. The g = ⟨220⟩ BF-XTEM image in Fig. 2(a) shows the InP QD PL structure grown on GaAs/Si, along with the top of the GaAsyP1-y step-graded buffer used to grow GaAs on GaP/Si. Figures 2(b) and 2(c) show the active region of the PL structures with coherently strained InP QDs exhibiting mottled, dark strain contrast capped by an InGaP QW and surrounded on both sides by AlGaInP barriers; the strain field of the InP QDs appears similar on both GaAs/Si [Fig. 2(b)] and GaAs [Fig. 2(c)]. We observe no misfit dislocations around the active region, an issue commonly observed with InAs DWELLs grown on Si.41–43 The BF-XTEM images show a high buried QDD of 1 × 1011 cm−2 measured over multiple images on both substrates. Figure 2(d) shows an atomically resolved HAADF-STEM image of InP QDs on GaAs/Si revealing structural details that are difficult to discern in BF-XTEM. The smooth interface between the InGaP QW and upper AlGaInP barrier confirms that a planar surface morphology recovered after the QD growth, consistent with in situ reflection high-energy electron diffraction observations. The mean height and diameter of the lens-shaped InP QDs nucleated on AlGaInP and buried by InGaP are 2.4 nm and 20 nm, respectively, similar to InAs QDs capped with (In)GaAs.44,45

FIG. 2.

(a) Low- and (b) high-magnification XTEM of the InP QD active region grown on GaAs/Si. (c) XTEM of InP QDs grown on GaAs showing a nearly identical morphology [same scale as 2(b)]. (a)–(c) are taken with g = ⟨220⟩ two-beam conditions showing compressive strain fields around individual QDs. (d) False-colored high-resolution HAADF-STEM image of the InP QD active region showing 2.4 nm tall and 20 nm wide InP QDs. The smooth and planar interface between the InGaP QW and AlGaInP barrier shows the recovery of planarity after capping InP QDs.

FIG. 2.

(a) Low- and (b) high-magnification XTEM of the InP QD active region grown on GaAs/Si. (c) XTEM of InP QDs grown on GaAs showing a nearly identical morphology [same scale as 2(b)]. (a)–(c) are taken with g = ⟨220⟩ two-beam conditions showing compressive strain fields around individual QDs. (d) False-colored high-resolution HAADF-STEM image of the InP QD active region showing 2.4 nm tall and 20 nm wide InP QDs. The smooth and planar interface between the InGaP QW and AlGaInP barrier shows the recovery of planarity after capping InP QDs.

Close modal

Consistent with the high QDD seen in TEM, AFM shows a high surface QDD of 1.3 × 1011 cm−2 with a bimodal height distribution on both GaAs/Si [Fig. 3(a)] and GaAs [Fig. 3(b)]. The InP QDD we report is higher than in previous reports46,47 and results from nucleation on an Al-rich surface35 at a relatively low growth temperature (480 °C for MBE vs 650 °C for MOVPE).48 The 20–30 nm lateral size of QDs on GaAs is slightly larger than that of QDs on GaAs/Si due to the slight variation in the surface temperature on the two substrates. The 4–7 nm height of the surface QDs shown in Fig. 3 is 2–3× higher than that of the buried QDs observed in HAADF-STEM [Fig. 2(d)] due to mass transport upon capping, as seen with capped InAs QDs.49,50Figure 3 also shows a bimodal height distribution with distinct height peaks at 4 nm and 7 nm for InP/AlGaInP surface QDs on both GaAs and GaAs/Si, indicating the need for further growth optimization for a homogeneous QD morphology.46,47 The size of InP QDs grown here is similar to InAs QDs grown on GaAs and Si, while their density is significantly higher than the 5 × 1010 cm−2 typical for InAs QDs.51,52 A high QDD observed for InP QDs is essential for efficient luminescence and defect-tolerance of visible QD-based emitters on Si.

FIG. 3.

AFM images of 3 ML surface InP QDs grown on (a) GaAs/Si and (b) GaAs with a similar QDD of 1.3 × 1011 cm−2.

FIG. 3.

AFM images of 3 ML surface InP QDs grown on (a) GaAs/Si and (b) GaAs with a similar QDD of 1.3 × 1011 cm−2.

Close modal

The panchromatic CL map of InP QDs grown on GaAs/Si in Fig. 4(a) shows a TDD of 3.3 × 107 cm−2, four orders of magnitude lower than the QDD. We verified a similar TDD using electron channeling contrast imaging (not shown). Figure 4(b) shows the CL map of InP QDs grown on GaAs showing no dark spots, as expected; the GaAs wafers were specified at a TDD of 0.5–1.0 × 104 cm−2. Consistent with BF-XTEM images [Figs. 2(a)–2(c)], planar-view CL shows no misfit dislocations in the active region grown on GaAs/Si over a measured area >150 μm2. The lack of misfit dislocations in Fig. 4(a) may result from the ability to perfectly lattice-match the InGaP QW to GaAs, which is not possible in the InAs/InGaAs QDWELL system. Furthermore, the lattice mismatch of 3.7% between InP and GaAs is much smaller than the lattice mismatch of 7% between InAs and GaAs, reducing the driving force for formation of misfit dislocations at the interface between the AlGaInP barrier and the InP wetting layer/QDs. Unlike previous reports on InAs QDWELLs on Si,42 no glide of dislocations was observed in either CL or electron channeling contrast imaging.

FIG. 4.

CL images of InP QDs grown on (a) GaAs/Si showing TDD = 3.3× 107 cm−2 with dark spots correlating with threading dislocations and (b) GaAs showing no dislocations over an area >150 μm2.

FIG. 4.

CL images of InP QDs grown on (a) GaAs/Si showing TDD = 3.3× 107 cm−2 with dark spots correlating with threading dislocations and (b) GaAs showing no dislocations over an area >150 μm2.

Close modal

InP QDs grown on dislocated GaAs/Si emit in the visible-NIR regime with PL intensity comparable to their counterparts grown on nearly dislocation-free GaAs substrates. We subjected all InGaP QW and InP QD PL samples to rapid thermal annealing (RTA) at temperatures of 700–1000 °C to remove point defects and increase the emission intensity.33Figure 5 shows that InGaP QWs on GaAs emit at 649 nm (1.91 eV), closely matching the expected ground state emission wavelength with a full width at half maximum (FWHM) of 24 meV and an additional shoulder peak at 621 nm corresponding to the first excited state. InP QDs on GaAs emit at 713 nm (1.74 eV) with a FWHM of 65 meV, similar to previous reports.47,53 The comparatively higher FWHM of InP QDs can be attributed to inhomogeneous broadening due to the distribution in QD size.27 We also observe an additional peak at 680 nm due to the bimodal QD size distribution observed in the AFM scans. The PL spectra of QD and QW samples grown on GaAs/Si are slightly redshifted from samples grown on GaAs due to the tensile strain arising from the thermal mismatch between the III and V layers and Si.54,55

FIG. 5.

Room temperature PL spectra of InGaP QW and InP QDs grown on GaAs (dashed) and GaAs/Si (solid). The InGaP QW (blue) grown on GaAs/Si shows 9× degradation of PL intensity compared to the QW grown on GaAs, whereas InP QDs (black) show nearly identical PL intensity on both GaAs and GaAs/Si (inset: visible emission observed from InP QDs grown on GaAs/Si at pump power 5 W/cm2).

FIG. 5.

Room temperature PL spectra of InGaP QW and InP QDs grown on GaAs (dashed) and GaAs/Si (solid). The InGaP QW (blue) grown on GaAs/Si shows 9× degradation of PL intensity compared to the QW grown on GaAs, whereas InP QDs (black) show nearly identical PL intensity on both GaAs and GaAs/Si (inset: visible emission observed from InP QDs grown on GaAs/Si at pump power 5 W/cm2).

Close modal

InGaP QWs grown on GaAs/Si show a 9× reduction in PL intensity compared to QWs grown on GaAs due to strong non-radiative recombination at threading dislocations. The integrated emission intensity of InP QDs on Si was ∼8× higher than that of InGaP QWs on Si, demonstrating the viability of dislocation-tolerant, visible, phosphide-based light emitters on Si. The inset of Fig. 5 shows the intense, short-wavelength tail emitted by the InP QDs grown on Si, which is visible to the naked eye. InP QDs are dislocation-tolerant due to the large disparity between QDD and TDD, which in turn leads to efficient carrier capture to the QDs. A reduced diffusion length of carriers due to lateral carrier confinement11,56 also makes InP QD active regions insensitive to dislocations compared to InGaP QW structures, where carriers freely diffuse to dislocations. Finally, differences in the energy level and capture cross section for dislocation-related traps in InP and InGaP could also partially account for the observed discrepancies in dislocation tolerance.

In conclusion, we demonstrated dislocation-tolerant InP QDs with comparable PL intensity on both GaAs and GaAs/Si. In contrast, InGaP QWs grown on GaAs/Si showed 9× PL degradation compared to QWs grown on GaAs and 8× lower intensity compared to InP QDs grown on Si. The dislocation tolerance of InP QDs arises from the orders of magnitude higher QDD compared to the TDD, as well as impeded lateral diffusivity of carriers in the QD layer, similar to InAs QDs. However, unlike InAs QDs, the InP QD-based active regions described here do not exhibit any misfit dislocations, potentially due to the ability to lattice-match the InGaP QW used to cap the QDs to GaAs. The lack of misfit dislocations around the active region may prove beneficial for improving the performance and reliability of InP QD based emitters,42,43 as the climb of misfit dislocations in InAs QD lasers on Si during device operation significantly increases non-radiative recombination and operating current.43 The high density QDs demonstrated here are tolerant to the presence of threading dislocations, opening the possibility of realizing efficient emitters on Si over a wider wavelength range than was previously possible.

R.D.H. was supported by a National Aeronautics and Space Administration (NASA) Space Technology Research Fellowship under Grant No. 80NSSC18K1171.

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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