**3. Structural analysis**

loss through non-radiative recombination, PL intensities and quenching temperatures can be increased by acting to raise the radiative recombination rate. In a very recent study, Jurczak et al. demonstrated a 10-fold enhancement of InAs NWs PL emission using an InP core-shell layer that passivates the surface states to reduce the rate of non-radiative recombination [35]. This research direction is attracting many researchers today in order to develop advanced optoelectronic devices and nanoscale photonic applications [13]. Progress in this direction will provide further insight into the optical emission and energy band-gap properties, hence improving the use of these materials, especially for infrared detectors and emitters. This chapter discusses the concept of developing novel InAsSb/InAs multi-quantum wells (MQWs) NWs on Si (111) substrate structures within InAs nanowires, as significant step towards viable nano- and quantum emitters in the extended IR wavelength range. We review the growth process for these structures, the crystal structural characterisation. Finally, we discuss their

optical properties along with developing a band structure for the NWs of this material.

The growth of high-quality nanowires should be achieved by avoiding the common growth process that employs a foreign catalyst such as gold to nucleate the wires. Au is well known to introduce deep level traps in the material band gap as contaminations [36, 37], which then limits the performance of the devices functionality on Si. Therefore, using lithographically

an accurate control over positions and diameters of the grown NWs that determines where the growth occurs and allows homogeneous arrays by controlling the nucleation position

In this work, the selective area molecular beam epitaxy (SA-MBE) technique has been used as a first step to grow the InAsSb/InAs MQWs NWs on Si (111) wafers and was achieved with nano-hole patterns produced by EBL [38]. The substrate is masked with a patterned dielectric

As and In adatoms start to form critical nuclei of certain sizes on the surface inside the pat-

The MQW nanowire growth was initiated by impinging As flux followed by exposure of the sample to the In flux 20 s later. After the growth of an initial pure InAs section for 1 h, the

of 180 and 27 s, respectively. This was expected to form 10 repeats of 25 nm InAs and 8 nm

, giving a total active region thickness of 330 nm as shown in **Figure 1(a)**. Finally, the wires were finished with an InAs cap, grown for 10 min. All growths began with a pure InAs section and were finished with an InAs cap. For further comparison, a planar bulk InAs sample was grown as a reference under the same growth conditions. The MQW NWs were grown in arrays of 50-nm diameter holes patterned using e-beam lithography in a 330-nm pitch square array defined within a number of 200 × 200 μm areas, on each silicon

without any catalysts by the selective area epitaxy technique.

MQW active region was grown as 10 repeats of InAs/InAs1−xSb<sup>x</sup>

template is an additional benefit besides avoiding catalysts which enables

layer with a thickness below 100 nm. The NW growth starts when the

MQWs with growth durations

**2. The growth**

58 Nanowires - Synthesis, Properties and Applications

predefined SiO<sup>2</sup>

layer, normally a SiO<sup>2</sup>

terned holes.

InAs1−xSb<sup>x</sup>

The lattice structure of the InAs/InAsSb MQW NWs has been investigated using Scanning Transmission Electron Microscopy (STEM). A close inspection of high-resolution STEM images shows both the angled nature of the well and a continued WZ structure within the InAsSb well. **Figure 2** shows a full diameter image of the wire together with a zoomed-in view, within which the lattice structure can be discerned. From the zoomed-in view, it is clear that the structure of the InAsSb well does not change to ZB, as it would do in a pure InAsSb NW. However, in the case of the QWs, the short growth durations are insufficient to allow a flat top to form and so the WZ phase is maintained. Hence, the Sb fraction in the wells is currently best estimated at 6–7%, based on comparable bulk wires.

The nanowires had a regular hexagonal cross section with {10¯ 10} sidewalls, faceted tips and a twinned WZ crystal structure with stacking faults, which is a characteristic of InAs NWs [39, 40]. In common with other researchers, it has been found that the addition of Sb to form bulk InAsSb NWs forces a rotation in the layer stacking, leading to a predominantly ZB structure with a flat top. Bulk InAsSb NWs grown under these conditions contained 6–7% Sb, in good agreement with the earlier work that reports saturation at this concentration [33]. However, when the InAsSb growth is limited to nanoscale quantum wells, energy-dispersive X-ray (EDXS) mapping revealed preferential incorporation of the Sb on specific crystal planes. This results in the formation of novel quantum wells having faceted, flat-topped, conical shape, with open or partially closed flat tops, as shown in **Figure 3(a)**. It is noted that in other NW studies, the authors have considered nanoscale axial heterostructures as both quantum dots (QDs) [41, 42] and QWs [43–45]. The QD model is understandably favoured for lower-wire diameters; however, in light of the larger 100-nm diameter and the weak lateral conferment of the higher effective mass holes, the heterostructures described here are considered to be QWs.

**4. 4-K micro-PL measurements**

(1.2 × 10<sup>4</sup>

InAs NWs at (1.2 × 10<sup>4</sup>

The 4-K PL emission from the MQW was studied using temperature-dependent microphotoluminescence spectroscopy. Introducing the InAsSb MQW significantly changes the

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**Figure 4.** Micro-PL spectra obtained at 4 K. The emission spectra measured from (a) InAsSb/InAs MQW NWs at high power

absorption from water vapour [49] in this spectral region, as shown in (a). Figure obtained with permission from authors [38].

W cm−2). The dotted lines represent Gaussian fits to the spectra, which also reveal the atmospheric

W cm−2) and (d)

W cm−2), (b) InAsSb/InAs MQW NWs at low power (80 W cm−2), (c) InAsSb NWs at (1.2 × 10<sup>4</sup>

**Figure 2.** High-resolution TEM analysis of the QW in NW. Cross-sectional TEM images, looking across a section of wire including a QW. The lattice structure is visible, and the crystalline phase (WZ) is evidently continuous across the quantum well. Figure obtained with permission from authors [38].

**Figure 3.** (a) STEM image showing the distribution of In, As and Sb obtained from 2D EDXS mapping and the resulting unusual facetted conical shape of the InAsSb MQW and (b) the unit cell of the WZ crystal structure showing the possible growth planes for the InAsSb MQW facets. Figure obtained with permission from authors [38].

## **4. 4-K micro-PL measurements**

The nanowires had a regular hexagonal cross section with {10¯

60 Nanowires - Synthesis, Properties and Applications

twinned WZ crystal structure with stacking faults, which is a characteristic of InAs NWs [39, 40]. In common with other researchers, it has been found that the addition of Sb to form bulk InAsSb NWs forces a rotation in the layer stacking, leading to a predominantly ZB structure with a flat top. Bulk InAsSb NWs grown under these conditions contained 6–7% Sb, in good agreement with the earlier work that reports saturation at this concentration [33]. However, when the InAsSb growth is limited to nanoscale quantum wells, energy-dispersive X-ray (EDXS) mapping revealed preferential incorporation of the Sb on specific crystal planes. This results in the formation of novel quantum wells having faceted, flat-topped, conical shape, with open or partially closed flat tops, as shown in **Figure 3(a)**. It is noted that in other NW studies, the authors have considered nanoscale axial heterostructures as both quantum dots (QDs) [41, 42] and QWs [43–45]. The QD model is understandably favoured for lower-wire diameters; however, in light of the larger 100-nm diameter and the weak lateral conferment of the higher effective mass holes, the heterostructures described here are considered to be QWs.

**Figure 2.** High-resolution TEM analysis of the QW in NW. Cross-sectional TEM images, looking across a section of wire including a QW. The lattice structure is visible, and the crystalline phase (WZ) is evidently continuous across the

**Figure 3.** (a) STEM image showing the distribution of In, As and Sb obtained from 2D EDXS mapping and the resulting unusual facetted conical shape of the InAsSb MQW and (b) the unit cell of the WZ crystal structure showing the possible

growth planes for the InAsSb MQW facets. Figure obtained with permission from authors [38].

quantum well. Figure obtained with permission from authors [38].

10} sidewalls, faceted tips and a

The 4-K PL emission from the MQW was studied using temperature-dependent microphotoluminescence spectroscopy. Introducing the InAsSb MQW significantly changes the

**Figure 4.** Micro-PL spectra obtained at 4 K. The emission spectra measured from (a) InAsSb/InAs MQW NWs at high power (1.2 × 10<sup>4</sup> W cm−2), (b) InAsSb/InAs MQW NWs at low power (80 W cm−2), (c) InAsSb NWs at (1.2 × 10<sup>4</sup> W cm−2) and (d) InAs NWs at (1.2 × 10<sup>4</sup> W cm−2). The dotted lines represent Gaussian fits to the spectra, which also reveal the atmospheric absorption from water vapour [49] in this spectral region, as shown in (a). Figure obtained with permission from authors [38].

PL emission characteristics of the NWs in a number of ways. A comparison of the spectra measured at 4 K from the InAsSb MQW NWs, bulk alloy InAsSb NWs and InAs NW samples is shown in **Figure 4**.

We observed that at the low excitation (0.011 W cm−2), the MQW NW emission intensity is 4.2 fold enhanced with respect to a bulk InAs reference sample, compared to a 1.6-fold enhancement from the InAs NWs (see **Figure 5**). It must be noted that this comparison assumes a direct area proportionality for the optical pumping efficiency. However, it has been shown that the efficiency of optical absorption in nanowire arrays exhibits a spectral dependence arising from mode guiding, due to the geometry of the wire and the array, such that the peak field can occur either inside or outside the wires [50]. To a first order, this effect is defined by the wire diameter, and in prior work, a very similar effect has been reported from an array of InAsSb NWs in a photodetector [11], where the peak response was obtained at 1.5 μm with an FWHM of 320 nm. Consequently, there is a non-optimal coupling with the 808-nm pump laser used in the present PL studies, and hence, further enhancement of PL emission intensities is to be expected if the pump laser wavelength is correctly matched to the NW geometry.

Under low excitation conditions, bulk ZB InAs at 4 K normally exhibit characteristic emission from bound exciton and donor-acceptor transitions around 0.417 and 0.374 eV, respec-

in state filling such that a single InAs peak is observed at 0.425 eV. In the present case, the InAs NW emission is further blue-shifted with respect to the bulk ZB reference sample, due to the WZ crystal structure of the NW, with a peak emission energy ranging from 0.469 eV under low excitation, to 0.485 eV under high excitation, see **Figure 6**. The band gap for WZ InAs is known to be higher than that of ZB InAs, and our result is consistent with earlier studies of WZ InAs NWs which reported band gaps in the range of 0.477–0.540 eV [12, 32].

**Figure 6.** Power dependence of PL emission. The dependence of the peak emission energy on the power of the pump laser incident on the sample, for InAs and InAsSb/InAs MQW NWs, showing the difference in the blue shift between the pure InAs NW, with minimal quantum confinement effects and the MQW NW, with a strong quantum confinement and

W cm−2) in our micro-PL experiments results

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**5. PL power dependence**

tively [51]. The high excitation intensity (~ 0<sup>4</sup>

charging effects. Figure obtained with permission from authors [38].

**Figure 4(a)** and **(b)** shows the PL emission from the InAsSb MQW NWs at high and low excitations, respectively. The emission from the bulk InAsSb NWs is shown in **Figure 4(c)** and are deconvoluted into peaks at 0.380 eV corresponding to 6% Sb with a dominant ZB phase in agreement with previous work [46], and a shoulder on the main peak originating from ZB InAs which appears as the dominant phase in the early stage of all the NW growths [47].

The emission from the InAs NWs peaks at 0.482 eV is shown in **Figure 4(d)**, demonstrating a dominant WZ phase [3, 33]. The PL emission of the InAsSb MQW NWs collected at high and low excitations exhibits a clear increase in peak emission energy with respect to the bulk InAsSb NWs. This is due to the strong carrier confinement within the quantum wells. In addition, the InAsSb MQW NWs also exhibit an increased emission intensity and a superior temperaturequenching behaviour compared with the bulk InAsSb NWs as expected. Most notably, the emission intensity is enhanced at all temperatures, due to the quantum confinement of electrons and holes. PL originates from type II spatially indirect recombination of electrons in the InAs layers with confined hole states in the InAsSb QWs, where the spatial separation helps in reducing non-radiative Auger recombination with a corresponding increase in radiative emission [48]. Single Gaussian fitting to the spectra reveals interference from characteristic atmospheric absorption by water vapour between 0.445 and 0.485 eV [10]. The spectra can be scaled to account for the reduced cross-sectional area of the nanowire samples, where only 7% of the surface area is covered by the NWs, assuming a 100% nucleation yield in the mask sites. Accounting for this reduced active area allows the most direct comparison of emission intensity.

**Figure 5.** Comparison of emission intensities. PL spectra, under 3.2 × 10<sup>4</sup> W cm−2 excitation, showing the relative emission intensities for MQW NWs, InAs NWs and an InAs bulk sample, scaled by active area. Figure obtained with permission from authors [38].

We observed that at the low excitation (0.011 W cm−2), the MQW NW emission intensity is 4.2 fold enhanced with respect to a bulk InAs reference sample, compared to a 1.6-fold enhancement from the InAs NWs (see **Figure 5**). It must be noted that this comparison assumes a direct area proportionality for the optical pumping efficiency. However, it has been shown that the efficiency of optical absorption in nanowire arrays exhibits a spectral dependence arising from mode guiding, due to the geometry of the wire and the array, such that the peak field can occur either inside or outside the wires [50]. To a first order, this effect is defined by the wire diameter, and in prior work, a very similar effect has been reported from an array of InAsSb NWs in a photodetector [11], where the peak response was obtained at 1.5 μm with an FWHM of 320 nm. Consequently, there is a non-optimal coupling with the 808-nm pump laser used in the present PL studies, and hence, further enhancement of PL emission intensities is to be expected if the pump laser wavelength is correctly matched to the NW geometry.
