1. Introduction

Photon detectors for thermal imaging in the long-wave infrared (LWIR) atmospheric transmission window (8-14 µm) are currently dominated by epitaxially grown structures of the small-bandgap bulk alloy HgCdTe and GaAs based quantum-well infrared photodetectors (QWIPs), both of which necessitate costly cryogenic cooling solutions for optimal operation [1]. The possibility of improved on-wafer uniformity and elevated detector operation temperatures, which would enable increased performance and reduced system size and cost, has given rise to various alternative approaches; most notably intraband based quantum dot infrared photodetectors (QDIPs) [2–6] and strained layer superlattices (T2SLs) [7–10]. Recently, an InSb quantum dot (QD) based structure grown on a GaSb substrate was demonstrated as a means of extending the detection wavelength of InAsSb photoconductors with a so-called nBn barrier design [11]. The detection mechanism is based on transitions from bound hole states in the QDs to continuum states in the surrounding bulk material. This configuration is expected to benefit from long carrier lifetime due to the type-II nature of the involved transitions, suppressed Shockley-Read-Hall generation due to the relatively large bandgap matrix material (as compared to the detection wavelength) and reduced Auger generation due to the high strain.

In this article we study InSb and InGaSb QDs grown on InAs (001) substrates as a means to create a band structure with type-II optical transitions in the LWIR region that could allow for a new class of thermal imaging devices with superior dark current characteristics. The heterostructure discussed here has a very similar strain situation and has a matrix material conduction band energy level close to that of the material combination presented in Ref. 11 but is significantly simpler to grow and characterize. This is due to the binary, nominally antimony free matrix material and a full epitaxial structure comprising solely one heterojunction type. Previous studies on this heterostructure include material grown with metal-organic vapor-phase epitaxy (MOVPE) [12], molecular-beam epitaxy (MBE) [13] and liquid-phase epitaxy (LPE) [14]. Photoluminescence (PL) has been demonstrated with peaks located in the range 3.3-4.6 µm [15], which is slightly shorter than the 5.5 µm reported for the InSb QD/InAsSb/GaSb system [11].

Here, we report on the growth, fabrication and optical properties of InSb and InGaSb QD layer structures and investigate the possibilities of photon detection in the LWIR range. A mesa etched photovoltaic device was fabricated and photoresponse was measured up to 6 µm at 80K. Photoluminescence was furthermore demonstrated in a series of single-QD layer test structures with peak emission wavelengths extending up to 8.5 µm at 77 K. In all, the current study demonstrates a potential for realizing an interband QD-based photon detector for the LWIR spectral range.

2. Experimental details

Samples for PL measurements were grown on undoped InAs (001) substrates at 470-530 °C by MOVPE using an Aixtron 200/4 system with hydrogen as carrier gas. The reactor pressure was 100 mbar with a total gas flow of 15 standard liters per minute. Trimethylindium (TMIn), triethylgallium (TEGa), trimethylstibine (TMSb), diethylzinc (DEZn), silane (SiH4) and arsine (AsH₃) were used as precursors. The growth rate for InAs was 0.39 nm/s, calibrated using x-ray diffractometry (XRD) while the InSb growth rate was estimated to 0.06 nm/s from growth of relaxed InSb bulk material at 490 °C. InSb thin films were grown with thicknesses ranging between 3 and 14 monolayers (MLs), which exceed the critical thickness of 1.7 ML for InSb QD formation by the Stranski-Krastanov growth mechanism in the material system [16]. In addition, InGaSb thin films were grown with nominal thicknesses up to 14 ML to evaluate reduced strain conditions. During the growth the V/III molar input flux ratio was 150 for InAs and was varied between 0.8 and 1.6 for the growth of the InSb and InGaSb QD layers. The PL test structures consist of a single QD layer grown on an undoped 200 nm thick InAs buffer layer and are capped with 50-80 nm InAs. The p-i-n diode structure consists of a 500 nm thick zinc-doped p-type bottom-contact layer doped to 2x10¹⁸ cm⁻³, a 800 nm undoped buffer layer, an undoped active region consisting of ten InSb QD layers grown at 470 °C, each covered with 80 nm InAs and a 300 nm n-type top-contact layer which is silicon-doped to 4x10¹⁷ cm⁻³. The QD layers in the p-i-n diode have a nominal thickness of 8 ML and were grown at a V/III input flow ratio of 0.8.

The device processing was done using standard photolithography techniques with a H₃PO₄:H₂O₂:H₂O-based wet etch to define the mesas. P and n contacts were both formed using a Ti/Pt/Au configuration. The devices were annealed at 255 °C for 5 minutes in vacuum and etched with a sodium hypochlorite solution after which the mesa sidewalls were encapsulated with polymerized photoresist.

The photoluminescence measurements were carried out at 77 K with a solid state laser emitting at 532 nm as excitation source. A Bruker V70 Fourier transform infrared (FTIR) spectrometer, fitted with a 16 µm cut-off wavelength HgCdTe detector and step-scan functionality, was used to measure the PL whereas the responsivity was measured using a dispersive IR spectrometer system with a calibrated spectral irradiance. Cross-sectional scanning tunneling microscopy (X-STM) images were recorded on in situ cleaved samples using an Omicron VT-STM at room temperature in a chamber with a base pressure lower than 1x10⁻¹⁰ mbar [17].

3. Results and discussion

The PL of samples grown at 490 °C with a V/III-ratio of 1.6 with different nominal thicknesses was studied to establish an appropriate interval of starting conditions for InSb QD formation and is presented in Fig. 1 . In all cases a peak at 3.1 µm is observed, corresponding to the InAs bandgap. The shoulder towards the low energy side is possibly related to shallow levels in InAs previously discussed by Fisher and Krier [18]. The peaks in the range 3.8-5.0 µm are related to the InSb QD layer and are here suggested to originate from the type-II transitions schematically indicated in the inset of Fig. 1. Samples with increasing QD layer thickness exhibit a redshift of the emission wavelength as expected for larger dots due to a reduction in the confinement energy. However, we here also see a reduced PL intensity with increasing layer thickness which possibly indicates defect-mediated strain relaxation or an altered QD formation mechanism resulting in a lower QD density.

 

Fig. 1 PL measured at 77 K from single InSb QD layer samples grown at a V/III ratio of 1.6. The band energies with a type-III broken gap band alignment as calculated by Pryor and Pistol [19] are included in the inset together with the indicated type-II transition.

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In Fig. 2 , the PL from InSb QD layer samples with 5 ML QD layer thickness grown at 490 °C is plotted as a function of V/III ratio. Besides the emission from recombination over the InAs bandgap at 3.1 µm, QD layer-related emission is identified with peaks in the interval 4.0-6.2 µm. The peak wavelength is notably red-shifted with decreasing V/III ratio, suggesting that the V/III ratio has profound effects on the growth dynamics of the QDs. The double peak feature of the 0.8 V/III ratio sample is possibly a result of a bimodal dot size distribution or the presence of a second bound state.

 

Fig. 2 PL measured at 77 K from single InSb QD layer samples grown at different V/III input flow ratios.

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Figure 3 shows the PL from InSb QD layer samples with a QD layer thickness of 5 ML and a V/III ratio of 0.8 grown at different temperatures. All four samples show InAs-related luminescence at 3.1 µm as well as peak emission in the wavelength range 4.3-6.2 µm interpreted to originate from the QD layer. The samples grown at the two higher temperatures are observed to exhibit a significantly shorter emission wavelength than their lower-temperature counterparts, which show similar PL intensity with a slightly reduced linewidth for the sample grown at 470 °C. This result can partially be explained by a thermally limited nucleation that produces sparse islands which grow large and/or reduced intermixing between the QD layer and the matrix material at low temperatures.

 

Fig. 3 PL measured at 77 K from single InSb QD layer samples grown at indicated temperatures in the interval 470-530 °C at a V/III ratio of 0.8 and a thickness of 5 ML.

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In Fig. 4 the PL from InSb QD-layer samples grown at 470 °C with a V/III ratio of 0.8 and varied QD layer thickness is shown. The peaks visible at 3.1 µm and in the range 5.5-6.2 µm are attributed to InAs and QD-layer PL respectively, where the latter extends up to 7.5 µm. It is observed that as the QD layer thickness is increased, the PL feature at 5.5 µm becomes less prominent whereas the peak feature at 6.2 µm increases in relative intensity. This is an indication that there are two dominant dot size distributions of which the larger is promoted when the QD-layer thickness is increased. For samples with QD layers thicker than 8 ML, not shown in Fig. 4, there is no further wavelength-shift but a reduction in PL-intensity which suggests an approximate limit for the onset of non-elastic strain-relaxation.

 

Fig. 4 PL measured at 77 K from single InSb QD layer samples grown at 470 °C with varied QD layer thickness. The dotted lines indicate the emission peaks of two possible dot size-distributions.

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Figure 5 shows the measured photoresponse at 80 K from a 10 QD-layer photodiode together with the PL of a corresponding test structure. The photoresponse has a large contribution from the InAs band-to-band excitation with a cut-off at approximately 3.1 µm and a small contribution around 3.2 µm possibly related to shallow levels in the InAs bandgap. In addition, the shoulder towards longer wavelengths extending up to 6 µm can be assigned to type-II transitions. The reason for the modest responsivity is presently not clear. As discussed by Ting et al. it may be associated with the trapping energy of the confined hole state but it may also be related to non-ideal contacts or simply the limited number of QD layers [20]. The dark current in the same device was measured to 1x10⁻⁶ A/cm² at 0.05 V reverse bias and 80 K.

 

Fig. 5 Responsivity measured in a 10 QD-layer device structure at zero bias and 80 K. The PL from a corresponding sample measured at 77 K is also presented.

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The QD sizes in the InSb QD/InAs material system are rather small compared to other III/V QD systems as reported by Norman et al. [12] and Ivanov et al. [21] where dot diameters were observed to be in the range 2.5-10 nm. Corresponding results are obtained in this study from X-STM measurements on a sample with 3 ML QD layer thickness, as presented in Fig. 6 . The figure indicates that 3D-islands have formed which extend around 2 nm in the growth direction, thus providing strong quantum confinement effects. A detailed STM study will be published elsewhere.

 

Fig. 6 A 35 x 35 nm² X-STM micrograph of a 10 QD-layer sample with a QD-layer thickness of 3 ML grown at 490 ̊C with a V/III ratio of 1.6. The image was acquired with a sample bias of −0.6 V.

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The reported short-wavelength emission in the InSb QD/InAs system [12,13,15], the PL results from this study and the expected band alignment where the valence band edge in the QD is located 0.30 eV above the InAs conduction band at 0 K [19] are all in favor of the interpretation that InSb QDs are grown in a size-sensitive domain where the QD bound state energy is strongly affected by confinement. This also indicates that type-II transitions in the LWIR range can be obtained from an optimization of the QD size.

A promising method to facilitate growth of larger QDs is to reduce the lattice mismatch. In this material system it is possible to reduce the strain by addition of gallium with limited impact on the valence band energy, which is a determining parameter for the type-II effective bandgap [19]. In Fig. 7 the PL from single In_0.6Ga_0.4Sb QD-layer samples with a QD-layer thickness of 12 ML grown in the temperature range 470-530 °C is shown. The peaks present in the 4.5-7.4 µm range are attributed to QD-layer PL which extends up to 9 µm for the sample grown at 530 °C. A strong influence of the growth temperature on emission wavelength is observed, with longer wavelength and potentially larger QDs in the samples grown at higher temperature, in contrast to what was observed for the temperature dependence of InSb QD-layer growth in Fig. 3. This can be explained by the expected larger equilibrium dot size for the case of reduced strain.

 

Fig. 7 PL measured at 77 K from InGaSb QD layer samples grown with a thickness of 12 ML at a V/III ratio of 0.8 at different temperatures.

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Figure 8 shows a PL spectrum from a single QD-layer sample where the Ga content has been increased to 50%. As before, the peaks observed at 3.1 µm and 8.5 µm are attributed to the InAs matrix material and QD-related type-II transitions respectively. Notably, the signal corresponding to the type-II transitions is now significantly longer in wavelength compared to the cases of lower Ga content and pure InSb QDs. We interpret this as a consequence of the reduced strain leading to larger QDs and thereby reduced confinement energy. Our current efforts are focused on the optimization of such long-wavelength QD material and the fabrication of LWIR photodiodes.

 

Fig. 8 PL measured at 77 K from a single QD-layer In_0.5Ga_0.5Sb QD sample grown with a thickness of 14 ML and a V/III ratio of 1.2 at 530 °C.

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4. Conclusion

In conclusion, a photovoltaic device based on InSb QDs has been demonstrated with photoresponse up to 6 µm at 80 K. It was furthermore shown that peak PL from InSb- and InGaSb-QD layer structures can be obtained at 6.2 µm and 8.5 µm respectively at 77 K. This result can be seen as a first step towards the realization of an interband QD-based photon detector material for the LWIR regime.

This work was supported by FLIR systems, the Swedish Defence Materiel Administration (FMV), the Knowledge Foundation (KK-stiftelsen) and the Swedish Governmental Agency for Innovation Systems (VINNOVA) through the industry excellence center IMAGIC. Additional support is acknowledged from the Swedish Foundation for Strategic Research (SSF) and the Göran Gustafsson Foundation.