Shape Engineering of InP Nanostructures by Selective Area Epitaxy

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Shape Engineering of InP Nanostructures by Selective Area Epitaxy Naiyin Wang,† Xiaoming Yuan,*,‡ Xu Zhang,*,†,§ Qian Gao,†,⊥ Bijun Zhao,† Li Li,∥ Mark Lockrey,∥,# Hark Hoe Tan,† Chennupati Jagadish,† and Philippe Caroff†,@

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†

Department of Electronic Materials Engineering, Research School of Physics and Engineering, The Australian National University, Canberra, ACT 2601, Australia ‡ Hunan Key Laboratory for Supermicrostructure and Ultrafast Process, School of Physics and Electronics, Central South University, 932 South Lushan Road, Changsha, Hunan 410083, P. R. China § National Center for International Joint Research of Electronic Materials and Systems, Henan Key Laboratory of Laser and Opto-electric Information Technology, School of Information Engineering, Zhengzhou University, Zhengzhou, Henan 450052, P. R. China ∥ Australian National Fabrication Facility ACT Node, Research School of Physics and Engineering, The Australian National University, Canberra, ACT 2601, Australia S Supporting Information *

ABSTRACT: Greater demand for III−V nanostructures with more sophisticated geometries other than nanowires is expected because of the recent intensive investigation of nanowire networks that show great potential in all-optical logic gates, nanoelectronics, and quantum computing. Here, we demonstrate highly uniform arrays of InP nanostructures with tunable shapes, such as membrane-, prism-, and ring-like shapes, which can be simultaneously grown by selective area epitaxy. Our in-depth investigation of shape evolution confirms that the shape is essentially determined by pattern confinement and the minimization of total surface energy. After growth optimization, all of the different InP nanostructures grown under the same growth conditions show perfect wurtzite structure regardless of the geometry and strong and homogeneous photon emission. This work expands the research field in terms of producing nanostructures with the desired shapes beyond the limits of nanowires to satisfy various requirements for nanoelectronics, optoelectronics, and quantum device applications. KEYWORDS: InP, selective area epitaxy (SAE), MOVPE, nanomembranes, nanorings, nanofaceting, surface energy

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higher-order nanostructures [two-dimensional (2D) and threedimensional (3D) shapes] will be required to fulfill the needs of these advanced applications.5,11−15 A top-down etching approach is an attractive route for fabricating nanostructures of different geometries16 but requires rather complex technological processing and further processing steps to remove the unavoidable surface damage.17−19 Metal-seeded epitaxy has also been used to grow membrane-like nanostructures by manipulating the catalyst shape and position, the introduction of structural defects (lateral twins and mirror twins), and/or the growth conditions.20−26 However, this method has a very limited

xtensive study of group III−V semiconductors at the nanometer scale over the past two decades has revolutionized nanoelectronics and nanophotonics.1,2 The nanowire structure, representing one of the most important building blocks of nanoscience and nanotechnology, has been demonstrated to be important for both fundamental research3 and device applications.4 Still, this elongated “onedimensional” (1D) nanostructure faces many challenges in practical device applications in terms of synthesis, assembly, and fabrication processes.5 Utilizing nanowires to form complex networks is drawing a great deal of attention as a way to reach improved device functionality. Despite their promising applications in all-optical logic components, nanoelectronics, and quantum devices,6−11 there are greater demands on the flexibility of the geometry, uniformity, and structural and optical qualities of nanostructures grown from the bottom up. Furthermore, there are more expectations that © 2019 American Chemical Society

Received: April 17, 2019 Accepted: June 10, 2019 Published: June 10, 2019 7261

DOI: 10.1021/acsnano.9b02985 ACS Nano 2019, 13, 7261−7269

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Cite This: ACS Nano 2019, 13, 7261−7269

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ACS Nano ability to control the nanostructure geometry and position, with poor yield and uniformity. Instead, selective area epitaxy (SAE) is widely applied to produce nanostructure arrays with high uniformity and potential for scalability,12,13,27−31 which is a promising route for addressing these issues.32,33 For instance, twin-free GaAs nanosheets have been epitaxially grown, which demonstrates the advantage of a sheet-like shape in material synthesis and the different growth mechanism from their nanowire counterparts.12 Homogeneous emission of triangular-shaped GaAs nanomembrane arrays at the submicrometer and submillimeter scale has also been reported.13 GaAs nanomembranes have been used as a platform to form nanowire networks due to their exceptional quality and the ability to form Y-shaped structures in situ.11 However, vertical GaAs nanomembranes can be formed only along the ⟨112̅⟩ orientations and typically evolve into a triangular shape, which may limit their controllability and usefulness.12,13,34−38 In addition, V-shaped InAs nanomembranes,39 InAs nanofins,40 and GaN nanosheets/nanowalls41−43 have also been reported. Unfortunately, a systematic investigation of the manipulation of the desired shapes and understanding the transformation mechanism of nanostructures is still lacking. In this work, we demonstrate highly uniform arrays of wurtzite (WZ) InP nanostructures with representative membrane-, prism-, and ring-like geometries, grown by catalyst-free SAE. Via an in-depth investigation of shape evolution with pattern dimension and time, we explore the growth mechanism of these nanoshapes with the help of a thermodynamic model. This also provides the underlying theoretical concepts for designing and achieving nanostructures of other shapes. By changing the pattern on the mask combined with the optimization of the growth conditions, we demonstrate exquisite control of the shape, uniformity, and scalability of nanostructures to satisfy the various requirements in practical applications. All InP nanostructures grown under optimal conditions show perfect WZ crystal structure regardless of their geometries. Moreover, cathodoluminescence (CL) results show strong and highly homogeneous emission from both ensemble and single elements of these nanostructures, indicating their excellent optical properties. This work shows the possibilities of obtaining other functional nanostructures using a bottom-up pathway and producing shapeengineered devices. Prior to growth, a standard preparation process for patterned substrates was used.31 Briefly, ∼30 nm thick SiOx was deposited on (111)A InP substrates as a mask, followed by electron beam lithography and reactive ion etching to fabricate different patterns (nanoslots and nanorings) on the mask layer. As shown by the schematic in Figure 1a, the directions of nanoslots were designed to be along ⟨101̅⟩, ⟨112̅⟩, and 15° off (i.e., the angle between nanoslot direction and ⟨101̅⟩ is 15°) directions. A large range of nanoslot lengths and nanoring diameters were patterned to investigate the shape evolution of these nanostructures and understand the formation mechanism. After trim etching,31,44 the patterned substrates were immediately loaded into the metalorganic chemical vapor deposition (MOCVD) reactor. Despite using two different systems, a close coupled showerhead (CCS) reactor and a horizontal flow reactor, for epitaxial growth (see Methods) the same nanostructures could be grown, demonstrating the high reproducibility of our work.

Figure 1. InP nanostructures grown from the patterned nanoslots with different directions. (a) Schematics of the designed nanoslots and grown InP nanostructures. The nanoslots are along the ⟨101̅⟩, 15° off (with respect to ⟨101̅⟩), and ⟨112̅⟩ directions of the (111)A InP substrate. The 30° tilted scanning electron microscope (SEM) images of InP nanostructure arrays with the nanoslots along (b) ⟨101̅⟩, (c) 15° off, and (d) ⟨112̅⟩ directions, respectively. All nanostructures were simultaneously grown under the optimal growth conditions (see Methods). The {101̅0} and {112̅0} side facets are representatively highlighted by orange and aqua colors, respectively.

RESULTS AND DISCUSSION Figure 1a schematically illustrates the SAE InP nanostructures grown from the patterned nanoslots of different orientations. As will be shown and discussed below, the nanostructures are all of the WZ crystal structure, and hence, we refer to them using the corresponding four-index scheme. Mainly bounded by the {101̅0} and/or {112̅0} side facets, they all grow vertically along the [0001] direction. Depending on the direction of the nanoslots, the cross sections of nanostructures may not fully inherit the geometries of the underlying patterns in spite of their confinement in the lateral directions. After optimization of the growth conditions (see Supporting Information section 1), the morphologies of the nanostructures grown from the ⟨101̅⟩, 15° off, and ⟨112̅⟩ oriented nanoslots are shown in panels b−d of Figure 1, respectively. For the ⟨101̅⟩ and ⟨112̅⟩ oriented nanoslots, the facets confined by the pattern correspond to the {101̅0} and {112̅0} planes, respectively. They both are the dominant side facets in panels b and d of Figure 1, respectively, leading to the formation of nanomembranes. We observe that in-plane ⟨101̅⟩ oriented nanomembranes are thinner than in-plane ⟨112̅⟩ oriented nanomembranes. Because thinner membranes are taller, surface diffusion is playing a role in this difference in geometry, but as we will show below, the difference in surface energy and radial growth of the {101̅0} and {112̅0} facets also plays a substantial role. When the nanoslot does not correspond to any low-index direction, for instance the 15° off nanoslot, the sidewalls of nanostructure predominantly consist of both {101̅0} and {112̅0} facets (see Figure 1c). Instead of the nanomembrane, the nanostructure is grown in the shape of a parallelogram prism. This coexistence of {101̅0} and {112̅0} facets in WZ InP would increase the likelihood of engineering the shapes of the nanostructures. More interestingly, the highly uniform arrays of InP nanostructures with 7262


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Figure 2. Evolution of the shape of SAE InP nanostructures from nanowires to nanomembranes with an increase in nanoslot length along the (a−c) ⟨101̅⟩ and (d−f) ⟨112̅⟩ directions. Panels a, c, d, and f are top views, and panels b and e are 15° tilted views. High-magnification SEM images from the highlighted areas in panels a and d are shown in panels c and f, respectively. The orange and aqua colored regions and lines indicate {101̅0} and {112̅0} side facets, respectively. The growth took place under the optimal growth conditions (see Methods) for 7 min.

f). This also applies to the case of nanoslots along any crystal direction. With the 15° off direction, for example, the experimental results are shown in Figure S6 (see Figure S7 for other directions). The shape evolution of nanostructures grown from 15° off oriented nanoslots during epitaxy is divided into four stages and schematically illustrated in Figure 3a. At the earlier stage after nucleation and filling up of the opening, shape I is formed (see Figure S11 for the SEM image), followed by the transformation toward shapes II−IV (see Figure 3b). Specifically, due to the confinement effect of nanoslots, the high-index facets appear at the initial stage I, but they are not thermodynamically stable and tend to be replaced by a combination of the metastable {112̅0} facets and the stable {101̅0} facets, leading to the formation of shape II. The evolution continues with the gradual replacement of the {112̅0} facets by the {101̅0} facets (stage III) and eventually reaches stage IV with the equilibrium prism-like shape when the {112̅0} facets completely disappear. A thermodynamic model is proposed to understand the driving force of the shape evolution depicted in Figure 3a. Considering crystal symmetry, only nanoslot directions of [101̅], [112̅], and θ off are discussed, where 0° < θ < 30° is the angle between the nanoslot direction and [101̅]. When θ = 30°, it is aligned parallel to the [112̅] direction. By assuming that the nanostructures with shapes I−IV have the same volume, we compared the total surface energy (G) of shapes I and II to that of shape IV (see Supporting Information section 3 for the details):

different shapes can be simultaneously achieved in the same growth run by just adjusting the nanoslot directions. The concept of shape engineering expands the research field from the traditional nanowires to many types of possible geometries that can play a significant role in applications, such as metamaterials and meta-optics that demand accurate control of the homogeneity and geometry of nanostructures. Hence, an understanding of the evolution of these InP nanostructures is a vital step for achieving these applications. The shape of the SAE InP nanostructures is determined by the growth conditions, duration, and designed pattern, including the geometry, dimension, and orientation. By controlling the length of the ⟨101̅⟩ and ⟨112̅⟩ nanoslots, we observed the evolution of the shape from nanowires to nanomembranes (see Figure 2). All nanostructures grown from the ⟨101̅⟩ nanoslots are mainly bounded by {101̅0} side facets, independent of the nanoslot length (Figure 2a−c). In comparison, only InP nanostructures grown from the very short ⟨112̅⟩ nanoslots mainly consist of {101̅0} side facets, resulting in a prism-like shape. As the nanoslots become longer, they start to confine lateral growth and {112̅0} facets start to appear and become more dominant (Figure 2d−f). The shape of a crystal usually reflects the stability of its facets.45,46 These observations suggest that although both {112̅0} and {101̅0} facets are formed during the SAE of InP, {101̅0} facets are more stable under the current growth conditions. Also, the shape evolution determined by facet stability can be observed either by extending the growth time (Figure S5) or by decreasing the nanoslot length (Figure 2d− 7263


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correlation as shown in Figure 2d and Figure S6d. For the sake of simplicity, we assume the growth of nanostructures is dominated by diffusion of indium adatoms from the substrate/ SiOx mask with a diffusion length of λs. The feeding area is either directly proportional to dL (when λs > d) or λsL (when λs < d), where d is the distance between two nearby nanoslots in the perpendicular direction. Accordingly, the rate of growth of the nanostructures is equal to λsLv or dLv = constant, where v (nanometers per second) is the deposition rate. tc could be expressed as tc =

ÅÄÅ γ{10 1̅ 0} ÅÅ ÅÇÅ cos θ

γ{1120̅ }

+ h − sin 2θ − ÑÉÑÅÄÅ γ h + sin(θ{10−1π̅ 0}/ 6) ÑÑÑÅÅÅ2 − (sin 2θ − ÑÖÅÇ

sin(π / 6 − θ)

+

0 < h ≤ (sin 2θ − G II = G IV

γ{1120̅ }

+

γ{10 1̅ 0} sin θ

(

3 /2)L

L

1

)γÉ

ÑÑ , ÑÑ 3 / 2)L Ñ ÖÑ 3 /2

HI

(1)

γ{10 1̅ 0}

sin(π / 6 − θ) sin θ γ{10 1̅ 0} γ{10 1̅ 0} cos θ

+

sin(θ − π / 6)

(3)

where ρ = H/L is the ratio of nanostructure height and length and and f (θ ) = sin(2θ + π /3) − 3 /2 f (θ ) = [sin(2θ + π /6) − 1/2] 3 /3 (0° < θ ≤ 30°) are the geometric factor of the top facet area of shapes II and IV, respectively. In particular, at θ = 30° (i.e., the [112̅] direction), shape I with high-index facets does not exist and shapes II and IV correspond to the membrane- and prism-like shapes, respectively. The results of the calculation are plotted in panels c and d of Figure 3. Figure 3c indicates that, at the fixed nanoslot length, the time needed to form shape II increases with nanoslot direction away from [101̅] and [112̅] directions and reaches a maximum at 15° off direction. On the other hand, the transformation time for achieving the final shape IV monotonously increases when the nanoslot direction gradually rotates from [101̅] to [112̅]. Figure 3d shows that three representative directions are picked out in the calculations. At a fixed nanoslot direction, the shorter the nanoslot, the sooner the shape transformation into the intermediate shape II and final shape IV could be completed (see the dotted and solid curves in Figure 3d, respectively). Thus, at a fixed growth time and a fixed nanoslot direction, the nanostructures with different nanoslot length are in the different stages of transformation, which is responsible for the observed shape evolution of the ⟨112̅⟩ nanostructures (Figure 2d,e). The results of calculation are consistent with the experimental observations. The analysis presented above shows the possibility of WZ InP in forming nanostructures with more complex shapes. The ring-like shape, which is desirable for optical signal processing and lasing applications,19,47,48 is designed and grown by SAE, as shown in Figure 4 (see Figures S12 and S13 for more details). InP nanostructures cannot fully inherit the geometry of the underlying patterns due to the unlikely formation of curved facets but present a multifaceted ring-like shape consisting of {112̅ 0} and/or {101̅ 0} side facets. The confinement effect of the patterned nanoring on the formation of these two facets is equivalent, which makes the investigation of the actual morphology of ring-like nanostructures more interesting in terms of understanding the growth behavior. Under optimal growth conditions (higher growth temperature), the area ratio of {112̅0}/{101̅0} side facets become larger but are always smaller than 1 with an increasing diameter (Figure 4h), leading to the transformation of the nanostructures from a hexagon-like to a dodecagon-like shape (Figure 4a−d). The ring size is essentially a snapshot in time of the growth evolution that is occurring. As the rings size becomes smaller, growth around the rings evolves more and as a result the {101̅0} facets are more dominant. Conversely, for largerdiameter rings, growth is less evolved and hence both {101̅0}

Figure 3. (a) Schematic of shape evolution during epitaxial growth. The inset shows the ratios of the total surface energy between shapes I and IV and shapes II and IV as a function of nanoslot direction. (b) Representative top view SEM images of shapes II− IV. The growth took place under the optimal growth conditions (see Methods) for 7 min. The relative transformation time of nanostructures as a function of nanoslot (c) direction and (d) length.

GI = G IV

ρL2f (θ ) ρL2f (θ ) or tc = λs v dv

(2)

where γ{101̅0}, γ{112̅0}, and γHI [γ{101̅0} < γ{112̅0} < γHI] are the surface energies of {101̅0}, {112̅0}, and high-index side facets, respectively, and L and h are the nanoslot length and the halfwidth of shape I, respectively. Through the calculation, we found GI/GIV > GII/GIV > 1, independent of nanoslot length and orientation (see the inset of Figure 3a). This suggests that epitaxial stages I and II are energetically unfavorable, and therefore, shape I induced initially by the pattern confinement effect evolves to shape II and eventually to the final shape IV, via another intermediate shape III, due to minimization of the total surface energy. As a special case, the transformation of the [112̅] nanostructures (θ = 30°) from the membrane-like shape to the prism-like shape complies with this model (see Supporting Information section 3 for the details). The dynamic shape evolution process provides a “road map” for tuning the shape of InP nanostructures. For more precise shape engineering, the relation between the time (tc) required to reach the critical shapes II and IV and the pattern design, including the orientation and dimension, needs to be established because we have already observed this strong 7264


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Figure 4. (a−g) Top view SEM images of ring-like nanostructures with different growth temperatures and diameters. The horizontal row of SEM images (a−d) represents the ring-like nanostructures with different diameters grown under optimal conditions. The vertical column of SEM images (b and e−g) shows the ring-like nanostructures with a 4 μm diameter formed at different temperatures. For each sample, epitaxial growth was carried out for 3 min. (h and i) Relation of the area ratio of {112̅0}/{101̅0} side facets to the diameter and growth temperature, respectively.

in determining the facet surface energy.50,51 The nanostructure shapes are affected by the growth temperature, particularly at lower temperatures where there are some unwanted effects such as irregular shape, rough surface, and non-uniformity (see Figure 4g and Figures S2, S3, and S13). The SAE InP nanostructures with different shapes show good morphology under optimal growth conditions. Still, there is a very stringent requirement on their crystalline quality for device applications because it was shown that relevant figures of merit such as quantum efficiency are directly affected by the presence of crystalline defects.31 The in-depth structural investigation in Figure 5 indicates that these nanoshapes have perfect WZ crystal structure. For the ⟨112̅⟩ oriented nanomembranes (Figure 5a), the dominant side facets are perpendicular to the ⟨112̅0⟩ zone axis, which is the ideal incident direction of electrons for the planar defect investigation. The selective area diffraction pattern (SADP) confirms the WZ phase of the nanomembranes (Figure 5b). An atomically resolved high-angle annular dark-field (HAADF) image taken at the interface (Figure 5c) shows that a twin plane is formed above the zinc blende (ZB) substrate. Except for this single twin, the subsequent atomic layers follow the WZ stacking sequence without observation of any misaligned layer along the [0001] direction until the end of the membrane, evidenced by the atomically resolved HAADF images taken from the middle and top sections (Figure 5d,e). Structural analysis of ⟨101̅ ⟩ nanomembranes is more challenging because the transmission electron microscope (TEM) lamella needs to be rotated 30° to the ⟨112̅0⟩ zone axis, which weakens the quality of TEM images. Still, careful

and {112̅0} facets are dominant. Thus, the metastable {112̅0} facets will gradually evolve to {101̅0} facets with further growth, eventually leading to the formation of the hexagon-like equilibrium crystal shape (Figure S14). These results are consistent with the earlier analysis of nanostructures grown from nanoslots that found that the {101̅0} facet is more stable than the {112̅0} facet. The role of growth temperature in facet stability and shape evolution was investigated by analyzing the ring-like nanostructures with a 4 μm diameter (Figure 4b,e−g). The area ratio of {112̅0}/{101̅0} side facets monotonously decreases from 4.7 to 0.53 as the temperature increases from 593 to 680 °C (Figure 4i). This shows that while {112̅0} facets are dominant at the low growth temperature, gradually the {101̅0} facets dominate with an increase temperature, suggesting the relative facet stability and nanostructure shapes can be tuned by controlling the growth conditions. This also shows that although our modeling presented above shows that the shape evolution is predominantly driven by the growth thermodynamics, growth kinetics may also play a role. The phenomenon of the {112̅0} facet being more stable at the low temperature is also observed during the growth of the nanomembranes (see Figure S15). Generally, the stability of the crystal facets is inversely proportional to their surface energy (γ); thus, our results indicate that {101̅0} facets have a lower surface energy at higher temperatures, while the reverse is true for {112̅0} facets. More precisely, the effect of temperature on the surface energy is probably caused by the effect of temperature on the percentage decomposition of the PH3 precursor,49 and therefore, the phosphorus chemical potential plays a key role 7265


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Figure 5. Crystalline quality of InP nanostructures with different shapes grown under the optimal conditions. (a) 30° tilted SEM image. (b) SADP, atomically resolved HAADF images taken along the ⟨112̅0⟩ zone axis from the (c) bottom, (d) middle, and (e) top sections of the ⟨112̅⟩ nanomembrane. (f) 30° tilted SEM image. (g) Low-magnification TEM image and high-resolution HAADF images taken along the ⟨112̅0⟩ zone axis from the (h) bottom, (i) middle, and (j) top regions of the prism-like nanostructure with 15° off nanoslot. (k) 30° tilted SEM image and (l) TEM image taken along ⟨112̅0⟩ zone axis from the bottom of the InP ring-like nanostructure. (m and n) High-resolution HAADF images from the highlighted area marked by the orange and green boxes in panel l, respectively. (o) ⟨112̅0⟩ zone axis bright-field TEM image indicating that stacking faults can be found only next to the interface and do not penetrate the whole layer. The dotted red lines in panels a, f, and k indicate the direction and position where the TEM lamellas were prepared by a focused ion beam (FIB). The dotted white lines and arrows indicate the defects in panels c, h, and n.

TEM examinations confirm that the ⟨101̅⟩ nanomembrane shows a pure WZ structure without twins or stacking faults (see Figure S16). In addition, the prism-like nanostructure grown from the 15° off nanoslot (Figure 5f−j) and the ringlike nanostructure (Figure 5k−o) also have a perfect WZ structure except at the bottom section. A twin defect close to the interface is found in the prism-like nanostructure, but the WZ phase starts from the first layer, showing an atomically sharp interface for the phase transformation (Figure 5h). As for the ring-like nanostructures, the initial stage of epitaxy (i.e., nucleation and filling up of opening) seems to be more complicated because nine layers of the ZB segment with the rotational twinning planes are formed inside of the pattern (Figure 5l,n). In addition to the geometric factor of the patterns, the minor fluctuation caused by the unstable growth conditions at the beginning of epitaxy might be responsible for the formation of these defects. After the initial growth stage, the crystal stacking sequence of all examined nanostructures becomes stable, growing as the pure WZ phase. This is not unlike the case for SAE InP nanowires31 where better structural and morphological properties are realized at the higher growth temperatures. More interestingly, we noticed that the planar defects at the bottom section of the ring-like nanostructure disappear

somewhere instead of penetrating the whole layer (Figure 5o).52 Because the dimension of the patterned openings can be larger than the adatom diffusion length,31 multiple nuclei are most likely formed in the initial stages, which subsequently expand and merge with each other to form a single crystal. The merging of nuclei could interrupt the spread of the planar defects. In addition to the excellent crystalline structure, superior optical properties as the prerequisite for photonic applications are achieved. Figure 6 shows the SEM and panchromatic CL images of highly ordered arrays of InP nanostructures with membrane-, prism-, and ring-like shapes. Strong and uniform luminescence from each element is observed for all of the nanostructure arrays regardless of their geometries, which is also confirmed by the intensity profiles along the dashed lines (see insets of Figure 6d−f). More surprisingly, the InP nanomembranes show uniform emission even when their length is up to 50 μm (Figure S17). The room-temperature photoluminescence spectra have a sharp emission peak at 871 nm for all of the different nanoshapes, consistent with the bandgap of the WZ crystal structure (Figure S18). Decreasing the growth temperature deteriorates the emission intensity and uniformity (see Figure S19), indicating the poor crystal quality of InP nanostructures grown at the lower temperature, 7266


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of 100 mbar and used highly purified H2 as a carrier gas with total flows of 10 and 15 L/min, respectively. Trimethylindium (TMIn) and phosphine (PH3) were used as precursors for In and P elements, respectively. All substrates were first annealed in ambient PH3 at a susceptor surface temperature (SST) of 658 °C in the CCS reactor or at a temperature of 750 °C in the horizontal reactor. Note that the SST in the CCS rector is measured by the in situ EpiTT, while the temperature in the horizontal reactor was measured via a thermocouple imbedded in the susceptor. After a 10 min annealing, the epitaxial growth was carried out once the reactors reached the desired conditions. Specifically, the molar fractions of PH3 for CCS and horizontal reactors were 2.80 × 10−3 and 2.67 × 10−3 while those of TMIn were set to 9.41 × 10−6 and 9.07 × 10−6, respectively. The SST of CCS reactor was changed from 593 to 680 °C, while the temperature of the horizontal reactor varied from 650 to 750 °C. The optimal growth temperatures in the CCS and horizontal reactors were 680 and 750 °C, respectively. The growth time of InP nanostructures grown by the CCS reactor was fixed at 3 min. Furthermore, two different growth times, 7 and 14 min, were implemented in the horizontal reactor. Characterization. The morphology of InP nanostructures was characterized using an FEI Verios 460 scanning electron microscope. A Gatan MonoCL4 Elite CL spectroscopic system equipped with the scanning electron microscope and a time-resolved photoluminescence (TRPL) system were used to investigate their optical properties. For structural characterization, an FEI Helios Nanolab FIB system was used to prepare the lamellae by cutting the nanostructures along different crystallographic orientations. The lamellae were examined by a JEOL JEM-ARM200f aberration-corrected scanning transmission electron microscope.

Figure 6. Strong and uniform photon emission of nanostructures. (a−c) Top view SEM and (d−f) corresponding panchromatic CL images of InP nanostructure arrays with different shapes. The overlays in panels d−f show the emission intensity profiles along the dash lines. Scale bars are 2 μm.

consistent with a poorer morphology. Therefore, a higher growth temperature is crucial for the SAE of InP nanostructures with the controllable shapes, high uniformity and crystal quality, and superior optical properties.

CONCLUSIONS We have demonstrated catalyst-free SAE of InP nanostructures with membrane-, prism-, and ring-like shapes, expanding the research field from 1D nanowires to more sophisticated 2D and 3D nanostructures. WZ InP is a good candidate for shape engineering because we found that {101̅0} and {112̅0} facets are able to be formed at the same time. The in-depth study of the shape evolution with time and dimension indicates that the shape is determined by the designed pattern (geometry, dimension, and orientation) and growth parameters (temperature, time, etc.). A proposed thermodynamic model confirms that the driving force of shape evolution is the minimization of the total surface energy during growth. We also found that the stability of {101̅0} and {112̅0} facets varies with the growth conditions. The {101̅0} ({112̅0}) facet is more stable at the higher (lower) growth temperature, which we attribute to the change in their surface energy. Under the optimal conditions at high temperatures, highly uniform arrays of InP nanostructures can be obtained with the excellent WZ crystal structure. Moreover, panchromatic CL results show the strong luminescence in the nanostructure arrays with high homogeneity for each nanostructure element in the array. The understanding of growth mechanism is crucial to the design and growth of nanostructures with the desired geometries. Our results represent the important steps toward shape engineering of nanostructures to realize advanced optoelectronic and electronic devices based on sophisticated and controllable geometries.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.9b02985. Optimization of growth conditions, time-dependent growth results of InP nanostructures, growth model, more results of ring-like nanostructures, results of nanostructures grown from nanoslots at low temperatures, and structural and optical characterization (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Xiaoming Yuan: 0000-0001-6840-6136 Present Addresses

⊥ Q.G.: BluGlass Ltd., 74 Asquith St., Silverwater, NSW 2128, Australia. # M.L.: Microstructural Analysis Unit, University of Technology Sydney, Sydney, NSW 2007, Australia. @ P.C.: Microsoft Quantum Lab Delft, Delft University of Technology, 2600 GA Delft, The Netherlands.

Author Contributions

N.W. carried out the sample preparation and growth, took part in all measurements, analyzed data, and wrote the paper. X.Y. performed the part of the TEM experiments, analyzed data, and co-wrote the paper. X.Z. carried out the modeling work. Q.G. developed the processing recipes for substrate preparation. B.Z. assisted in the SEM measurements. L.L. carried out FIB for preparing TEM samples. M.L. assisted in the CL measurements and interpretation of the results. H.H.T. supervised the project and took part in the data analysis and

METHODS Growth. In this work, two different MOCVD reactors, a CCS reactor (Aixtron 3 × 2 in.) and a horizontal flow reactor (Aixtron 200/4), were utilized for epitaxial growth. The CCS reactor was used to optimize the growth conditions in terms of temperature and V/III ratio (see Supporting Information section 1). Then, we used the horizontal reactor to test the epitaxial reproducibility and carry out the time-dependent growth. Both reactors operated at a low pressure 7267


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ACS Nano

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writing of the paper. C.J. supervised the project and took part in the writing of the paper. P.C. designed and supervised the project as well as the initial set of growths and co-writing of the paper. All authors contributed to the discussion. Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS The Australian Research Council is acknowledged for its financial support. N.W. thanks the China Scholarship Council and the Australian National University for scholarship support. X.Y. thanks the National Natural Science Foundation of China (51702368), the Hunan Provincial Natural Science Foundation of China (2018JJ3684), and the Innovation-Driven Project of Central South University (2018CX045) for financial support. Access to the facilities is made possible through ACT node of the Australian National Fabrication Facility and the ANU node of the Australian Microscopy & Microanalysis Research Facility. The authors thank Drs. Gilberto Casillas Garcia, Felipe Kremer, Xiangyuan (Carl) Cui, Naeem Shahid, and Shagufta Naureen for their assistance and discussions. REFERENCES (1) Duan, X.; Huang, Y.; Cui, Y.; Wang, J.; Lieber, C. M. Indium Phosphide Nanowires as Building Blocks for Nanoscale Electronic and Optoelectronic Devices. Nature 2001, 409, 66−69. (2) Wang, Z.; Tian, B.; Pantouvaki, M.; Guo, W.; Absil, P.; Van Campenhout, J.; Merckling, C.; Van Thourhout, D. Room-Temperature InP Distributed Feedback Laser Array Directly Grown on Silicon. Nat. Photonics 2015, 9, 837−842. (3) Tizei, L.; Craven, A.; Zagonel, L.; Tencé, M.; Stéphan, O.; Chiaramonte, T.; Cotta, M.; Ugarte, D. Enhanced Eshelby Twist on Thin Wurtzite InP Nanowires and Measurement of Local Crystal Rotation. Phys. Rev. Lett. 2011, 107, 195503. (4) Eaton, S. W.; Fu, A.; Wong, A. B.; Ning, C.-Z.; Yang, P. Semiconductor Nanowire Lasers. Nat. Rev. Mater. 2016, 1, 16028. (5) Rogers, J. A.; Lagally, M. G.; Nuzzo, R. G. Synthesis, Assembly and Applications of Semiconductor Nanomembranes. Nature 2011, 477, 45−53. (6) Desplanque, L.; Fahed, M.; Han, X.; Chinni, V.; Troadec, D.; Chauvat, M.; Ruterana, P.; Wallart, X. Influence of Nanoscale Faceting on the Tunneling Properties of Near Broken Gap InAs/ AlGaSb Heterojunctions Grown by Selective Area Epitaxy. Nanotechnology 2014, 25, 465302. (7) Gazibegovic, S.; Car, D.; Zhang, H.; Balk, S. C.; Logan, J. A.; de Moor, M. W.; Cassidy, M. C.; Schmits, R.; Xu, D.; Wang, G.; et al. Epitaxy of Advanced Nanowire Quantum Devices. Nature 2017, 548, 434−438. (8) Yang, H.; Khayrudinov, V.; Dhaka, V.; Jiang, H.; Autere, A.; Lipsanen, H.; Sun, Z.; Jussila, H. Nanowire Network−Based Multifunctional All-Optical Logic Gates. Sci. Adv. 2018, 4, No. eaar7954. (9) Vaitiekėnas, S.; Whiticar, A.; Deng, M.-T.; Krizek, F.; Sestoft, J.; Palmstrøm, C.; Marti-Sanchez, S.; Arbiol, J.; Krogstrup, P.; Casparis, L.; Marcus, C. M. Selective-Area-Grown Semiconductor-Superconductor Hybrids: A Basis for Topological Networks. Phys. Rev. Lett. 2018, 121, 147701. (10) Krizek, F.; Sestoft, J. E.; Aseev, P.; Marti-Sanchez, S.; Vaitiekėnas, S.; Casparis, L.; Khan, S. A.; Liu, Y.; Stankevič, T.; Whiticar, A. M.; et al. Field Effect Enhancement in Buffered Quantum Nanowire Networks. Phys. Rev. Mater. 2018, 2, 093401. (11) Friedl, M.; Cerveny, K.; Weigele, P.; Tütüncüoglu, G.; MartíSánchez, S.; Huang, C.; Patlatiuk, T.; Potts, H.; Sun, Z.; Hill, M. O.; et al. Template-Assisted Scalable Nanowire Networks. Nano Lett. 2018, 18, 2666−2671. 7268


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Shape Engineering of InP Nanostructures by Selective Area Epitaxy

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