Impact of Surface Point Defects on Electronic Properties and p-Type

Article pubs.acs.org/JPCC

Impact of Surface Point Defects on Electronic Properties and p‑Type Doping of GaAs Nanowires Haibo Shu,*,†,§ Xiaodong Yang,† Pei Liang,† Dan Cao,‡ and Xiaoshuang Chen§ †

College of Optical and Electronic Technology and ‡College of Science, China Jiliang University, 310018 Hangzhou, China National Laboratory for Infrared Physics, Shanghai Institute of Technical Physics, Chinese Academy of Science, 200083 Shanghai, China

§

S Supporting Information *

ABSTRACT: Gallium arsenide (GaAs) nanowires possess a great potential as a fundamental building block for the next-generation electronic and optoelectronic devices, but their applications are limited by the p-type doping. Improving the p-type doping efficiency of GaAs nanowires depends on understanding the doping limits and developing effective methods to reactive p-type dopants. Here the stability of various surface point defects and their role in electronic structure and p-type doping of GaAs nanowires are studied by using first-principles calculation within density functional theory. Our results demonstrate that As antisite (AsGa) is a highly stable surface point defect under As-rich condition irrespective of bare and H-passivated GaAs nanowires. The formed AsGa defects bring deep donor-type levels into the band gap, which is responsible for the deactivation of p-type dopants in GaAs nanowires. To suppress the impact of AsGa defects on the p-type doping of GaAs nanowires, we propose two feasible methods by reducing As chemical potential (or As partial pressure) and passivating with appropriate surface species (e.g., NO2), respectively. This work provides a new insight into the origin of p-type doping limits and the guidance for the realization of highly efficient p-type doping in III−V semiconductor nanowires.

1. INTRODUCTION III−V semiconductor nanowires (NWs) have received great attention due to their importance for both fundamental physics and technology fields in the development of nanoscale electronic, photonic, and sensing devices.1−3 Among these nanowires, GaAs NWs are one of the most studied targets because of high carrier mobility, large absorption coefficient, and moderate direct bandgap (∼1.43 eV in bulk)4 as well as the compatibility with the Si technology,5,6 which make them as the potential building blocks in nanoelectronics and optoelectronics, such as light-emitting diodes (LEDs),7 field-effect transistors (FETs),8 single-photon sources,9 solar cells,10 and compact terahertz sources and detectors.11,12 The realization of the wide range of applications, on the one hand, depends on the controlled growth of NWs, including of growth direction, size, crystal phase, and morphology. A large number of studies have been carried out in the past decade and demonstrated that vapor−liquid−solid (VLS) methods can induce GaAs NWs with controlled diameter and crystal phase.13−15 For example, apart from the bulk phase of GaAs, the VLS-grown GaAs NWs exhibit both zinc-blende (ZB) and wurtzite (WZ) structures.16,17 The crystal phase of GaAs NWs can be well controlled by adjusting the V/III ratio and growth temperature.18−20 On the other hand, the manipulation of carrier concentration and conductivity type is also a key factor for the © XXXX American Chemical Society

realization of NW devices. Doping is regarded as a fundamental method to provide free carriers, but the doping efficiency of III−V semiconductor NWs, such as GaAs, InAs, and InSb,21−23 is far lower than expected, especially for the p-type doping. To improve p-type doping efficiency of III−V semiconductor NWs, extensive effort has been devoted to exploring possible mechanisms. Three potential factors are responsible for the ptype doping limits. The first one is high ionization energy of dopants. The ionization energy of NWs was found to increase with the reduction of size due to the quantum confinement effect,24 resulting in lower free carrier concentration. Such a situation can be improved by introducing surface passivation layer on NW surface.25,26 For instance, high charge carrier mobility and concentration in GaAs NWs can be achieved by surface coating with optimized AlGaAs shells.25 The second factor is the so-called self-purification effect.27,28 The impurity atoms prefer to be segregated to surfaces/interfaces of semiconductor nanomaterials in the doping process, in particular for in-situ doping. In order to overcome the doping limit, Li et al.29 developed a remote doping technology that was implemented by introducing p-type impurity atoms at the shell Received: July 29, 2016 Revised: September 8, 2016

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DOI: 10.1021/acs.jpcc.6b07624 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C

Figure 1. Atomic structure of (a) zinc-blende and (b) wurtzite GaAs nanowires and (c) intrinsic defect models in these nanowires. The Ga and As atoms are colored blue and green, respectively. The capital letters A−F denote different intrinsic defects. The solid circles in the zinc-blende nanowire denote the corner regions; here Ga and As atoms have lower coordinate number (CN = 2).

approximation (GGA) of Perdew−Burke−Ernzerhof (PBE).37,38 In order to describe accurately the interaction between the adsorbed NO2 molecule and nanowire surface, the PBE functional with van der Waals correction (DFT-D2) has been used.39 The kinetic energy cutoff for the plane-wave expansion is set to 500 eV. The Monkhorst−Pack k-point of 1 × 1 × 6 is found to provide sufficient accuracy for the Brillouinzone integration of GaAs nanowires. For the geometry optimization, the convergence criteria of electronic energy and forces acting on each atom are 10−3 eV and 10−2 eV/Å, respectively. GaAs NWs are created on the basis of optimized bulk structures and oriented along the [111] direction for the ZB structure (see Figure 1a) and [0001] direction for the WZ structure (see Figure 1b). The ZB and WZ NWs are faceted by six {110} and {11̅00} facets, respectively, indicating a hexagonal cross section. Here the biggest diameter of ZB and WZ GaAs NWs arrives at 2.09 and 2.52 nm, respectively. The optimized lattice constants of GaAs bulk are a = 5.71 Å for the cubic ZB structure and a = 4.05 Å and c = 6.68 Å for the hexagonal WZ structure, which agree well with experimental values.40 We considered both bare and hydrogenated GaAs NWs in the present calculations. For hydrogenated NWs, surface Ga and As atoms are passivated by artificial hydrogen atoms with fractional charges of 1.25 and 0.75 e, respectively.41 For bare NWs, the surface dangling bonds of low coordinate-number atoms (CN = 2) at the corner of ZB NWs have been removed (see Figure 1a), since the corner effect on electronic structure is weak at the large-size NWs. This way can ensure that intrinsic GaAs NWs keep a semiconductor character (see Figure S1 in Supporting Information for the details). To avoid the spurious interaction between neighboring NWs, a vacuum region of more than 15 Å is introduced along the normal direction of NW axis. In order to investigate the role of surface defects in electronic properties and p-type doping of GaAs NWs, six potential intrinsic point defects have been considered (see Figure 1c), including of As antisite (AsGa), Ga antisite (GaAs), As vacancy (VAs), Ga vacancy (VGa), As interstitial atom (Asi), and Ga interstitial atom (Gai). The stability of various intrinsic point defects is evaluated by calculating their formation energies (Ef) as follows

layer of nanowires during the growth. The last factor is the Fermi-level pinning effect that is regarded as a key factor for causing electron accumulation layer (EAI) on NW surfaces,30 resulting in p-type dopant deactivation. It is generally accepted that the Fermi-level pinning in conduction band of III−V semiconductor NWs is caused by plenty of surface states.30,31 However, the origin of surface states and the mechanism for the p-type doping inhibited by surface states in III−V semiconductor NWs are unclear. The surface electronic states of III−V semiconductor NWs can be ascribed by two aspects: one is surface dangling bonds induced by unpassivated surface atoms, and the other is intrinsic point defects on NW surfaces. The surface dangling bonds can be eliminated by surface passivation and reconstruction.31−33 However, synthesized InAs and GaAs NWs by metal−organic vapor phase epitaxy (MOVPE) with high-concentration hydrogen also presented an n-type character, suggesting that the surface dangling bonds are not the sole factor to induce the formation of EAI on the surface of III−V semiconductor NWs.34 Therefore, it is necessary to understand the role of surface point defects in electronic properties of III− V semiconductor NWs. In this work, we present a systematic theory investigation on defect stability and the impact of surface point defects on electronic structure and p-type doping of GaAs NWs using density-functional theory (DFT) calculations. Our results demonstrate that As antisite (AsGa) is a stable surface point defect under As-rich conditions regardless of bare and Hpassivated GaAs NWs. The As antisite defect induces deep defect levels in band gap that become a trap and recombination center for hole carriers, resulting in the deactivation of p-type dopants. Hence, we propose two feasible ways to improve ptype doping of GaAs NWs. The present study provides not only a deep understanding for insight into the role of intrinsic point defects in electronic properties of GaAs NWs but also a feasible route to improve p-type doping of GaAs and other III− V semiconductor NWs.

2. COMPUTATIONAL DETAILS All DFT calculations are performed using the projector augmented wave (PAW) method35 as implemented in the Vienna ab initio Simulation Package (VASP).36 The electronic exchange-correlation energy is treated by generalized-gradient B


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The Journal of Physical Chemistry C Ef = E D − E T + Δni ∑ μi i

whether the surface defects are located at ZB or WZ NWs (see Figure 2a,c). For other point defects (VGa, VAs, Gai, and Asi), they presented relatively higher formation energies, implying that these defects are difficult to be formed on NW surfaces during the growth. When these intrinsic point defects are moved onto the subsurface of NWs, all point defects indicate increasing formation energies (see Figure 2b,d). This means that intrinsic point defects are easily formed on NW surface. To clearly understand the dependence of defect depth on GaAs NWs, we take the AsGa as an example and put the defect onto different sites of NWs with the defect depth from the NW surface to the core (see Figure S2). The formation energy of AsGa gradually increases with increasing defect depth, suggesting that the intrinsic point defects prefer to form at the surface region of NWs. The result is similar to the segregation effect of impurities in III−V semiconductor NWs reported from previous studies.27,28 Nevertheless, the point defects at the subsurface of NWs exhibit similar stability. Namely, AsGa and GaAs are the most stable intrinsic defects under As-rich and Ga-rich conditions, respectively. Since the growth of GaAs NWs is largely maintained at high As/Ga flux ratio (i.e., As-rich condition; see Table S1 in the Supporting Information), thus AsGa is a stable defect in synthesized GaAs NWs. Such a result can be supported by recent experiment42 in which a large concentration of AsGa defects has been identified by highresolution scan tunneling microscopy (STM). To gain insight into the role of surface intrinsic defects in electronic properties of GaAs NWs, band structure and chargedensity isosurface (CDI) distributions of band-edge states in both ZB and WZ NWs with an AsGa defect have been calculated, as shown in Figure 3. It is a well-known fact that the

(1)

where ED and ET are total energies of GaAs NWs with and without surface defects, respectively. μi is the chemical potential of atomic species i (i = Ga and As), and Δni is the difference of the number of Ga and As atoms between GaAs NWs with and without defects. In equilibrium conditions, μGa + μAs = μGaAs(bulk); thus, μGa can be transformed into a function of μAs. The allowable range of μAs is μAs(bulk) + ΔHf < μAs < μAs(bulk), where the upper (lower) limit corresponds to As-rich (Ga-rich) condition; μAs(bulk) and ΔHf are the energy of As atom in bulk and the heat of formation of ZB GaAs, respectively. The computed ΔHf is 0.750 eV, in good agreement with the experimental value (0.736 eV).40 To provide a comparison, both ZB and WZ thin-film models (i.e., GaAs(110) and GaAs(1100)) have also been considered. The GaAs(110) and GaAs(1100) thin films are created using the slab geometry with eight GaAs layers, and thin-film surfaces are passivated by H atoms. For the study of intrinsic defects and p-type doping in GaAs thin films, (2 × 3) and (3 × 2) supercells are used for (110) and (1100) thin films, respectively. The surfaces with in-plane periodicity are separated by an ∼15 Å vacuum layer to prohibit the interactions of neighboring surface slabs, and the Brillouin zone is sampled by the k-mesh grid of 4 × 4 × 1 within the Monkhorst−Pack scheme.

3. RESULTS AND DISCUSSION We first examine the stability of various intrinsic point defects in bare GaAs NWs. Figure 2 shows that formation energies of various defects on ZB and WZ NWs as a function of As chemical potential difference (Δμ, Δμ = μAs − μAs(bulk)). For the intrinsic point defects on NW surfaces, it can be found that the As antisite (AsGa) is the most stable surface defects under the As-rich condition and the Ga antisite (GaAs) becomes the most stable one under the Ga-rich environment no matter

Figure 3. Band structure and CDI distribution of band-edge states of bare GaAs NWs with an AsGa defect. The AsGa defect is situated at (a) the surface of ZB NW, (b) the subsurface of ZB NW, (c) the surface of WZ NW, and (d) the subsurface of WZ NW. The red and blue circles in band structures denote the position of CBM and VBM, respectively. The defect levels (DLs) in band structures are colored red, and horizontal dashed lines denote the position of the Fermi level.

Figure 2. Formation energies (Ef) of various surface and subsurface point defects in bare GaAs NWs as a function of As chemical potential differnce (ΔμAs). The intrinsic point defects are situated on (a) the surface of NW with the ZB structure, (b) the subsurface of NW with the ZB structure, (c) the surface of NW with the WZ structure, and (d) the subsurface of NW with the WZ structure. C


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difference (Δμ). We find that intrinsic point defects in hydrogenated NWs exhibit the similar stability as bare NWs: AsGa is the most stable defect under the As-rich conditions, and GaAs is stable under the Ga-rich environment irrespective of the crystal phase of NWs (see Figure 4a−d). In light of high As/Ga flux ratio (see Table S1) during the growth, AsGa thus should be a stable defect in GaAs NWs. Figure 5 indicates the band structure and the CDI distribution of band-edge states of hydrogenated GaAs NWs with an AsGa defect at the NW surface and the subsurface. Unlike the case of bare NWs, the AsGa defect at the surface of hydrogenated NWs exhibits a typical n-type character (see Figure 5a,c). This originates from that two extra electrons introduced by the AsGa defect cannot be redistributed on NW surface due to the hydrogen passivation. Accordingly, the electron polarization induces the formation of defect levels in the band structure of hydrogenated NWs. The CDI distribution of defect level indicates that electronic states of DL are completely localized around the defect. Furthermore, moving the AsGa defect onto the subsurface of NWs does not change its donor-type characteristic (see Figure 5b,d). With the increase of AsGa concentration in NWs, more defect states will be introduced into the band gap, thereby causing the Fermi level pinning of GaAs NWs in the conduction band. On the other hand, we also find that the AsGa defect has a relatively smaller effect on the VBM-state distribution of NWs, but it strongly affects the distribution of CBM state. As shown in Figure 5, the CDI of CBM state in all hydrogenated NWs indicates an obvious localization due to the AsGa defect. The result implies that electron transport properties of GaAs NWs should be strongly affected by the surface defects (i.e., AsGa). Now we want to know the impact of surface point defects on p-type doping in GaAs NWs. For this purpose, a typical p-type dopant (i.e., Mg) is introduced into hydrogenated GaAs NWs with a surface AsGa defect. Here we consider Mg as the p-type dopant of GaAs nanowires since the Mg element has been widely used in the p-type doping of III−V semiconductors, such as GaAs, GaN, and InN.45,46 The Mg dopant at different sites of NWs has been studied (see Figure 6a,c). Figure 6 indicates doping models and band structures of Mg-doped hydrogenated GaAs NWs with an AsGa defect. We find that both ZB and WZ GaAs NWs present a semiconductor character when a surface AsGa defect and an Mg dopant simultaneously appear in NWs. This arises from that the acceptor level induced by the Mg dopant is filled by the electrons of donor level caused by the AsGa defect, resulting in that the donor level moves onto the higher energy and the acceptor level moves onto the lower energy (see Figure 6b,d). In other words, an AsGa defect can lead to the deactivation of a p-type dopant in GaAs NWs. With the increase of distance between the Mg dopant and the AsGa defect (see Figure 6), the donor level of both ZB and WZ NWs moves to the lower energy and their acceptor level moves to the higher energy. The result means that the coupling between the Mg dopant and the AsGa defect becomes weak with the increase of the distance. In principle, the effect of Mg dopant and AsGa defect on electronic structure of NWs will be independent of each other when their distance is enough large. However, there are a large number of AsGa defects on the NW surface,42 and thus the intrinsic point defects must be a key factor for limiting the p-type doping of GaAs NWs. It needs to be emphasized that the size of nanowires in the present calculations is smaller than 3 nm, but the smallest

PBE functional underestimates the band gap of semiconductor nanowires, but previous studies have indicated that electronicstructure characteristic of nanowires cannot be affected by the use of the PBE functional.43,44 We find that both ZB and WZ NWs indicate intrinsic semiconductor character with a direct band gap (see Figure 3a,c) when the AsGa defect is located at the NW surfaces. The CDI distribution indicates that valenceband maximum (VBM) and conduction-band minimum (CBM) are typical extended states and mainly from the contribution of s states of Ga and As atoms and p states of As atoms, respectively, which agrees with the case of the bare NW without the AsGa defect. This originates from that an AsGa defect brings two additional electrons that are redistributed at an As−As dimer on the NW surface, forming a closed shell of electronic configuration.43 Therefore, the formation of AsGa on NW surfaces has little influence on electronic structure of bare GaAs NWs. When the AsGa defect lies at the subsurface of NWs, band structures of both ZB and WZ NWs include deep donor-type defect levels (DLs) in band gap (see Figure 3b,d). It is seen from the CDI distribution that electronic states of DLs are localized around the AsGa defect. The deep DLs originate from the polarization of two additional electrons induced by the AsGa defect. However, we find that there are two DLs in the WZ NW and a DL in the ZB NW. It is because one of two DLs (or the higher-energy DL) in the ZB NW is higher than the position of CBM. Overall, the change of defect sites from the surface to the subsurface will cause the electronic-structure transition of bare GaAs NWs from an intrinsic semiconductor to an n-type one. Since hydrogen abundantly exists in currently popular growth techniques for III−V semiconductor NWs, such as MOVPE, chemical beam epitaxy (CBE), and gas source molecular beam epitaxy (MBE), it is thus necessary to understand the stability and properties of intrinsic point defects in hydrogenated NWs. Figure 4 shows formation energies of various point defects at the surface and the subsurface of hydrogenated GaAs NWs as a function of As chemical potential

Figure 4. Defect formation energies (Ef) of various surface and subsurface defects on hydrogenated GaAs NWs as a function of ΔμAs. The intrinsic point defects are located at (a) the surface of a ZB NW, (b) the subsurface of ZB NW, (c) the surface of WZ NW, and (d) the subsurface of WZ NW. D


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Figure 5. Band structure and CDI distribution of band-edge states of hydrogenated GaAs NWs with an AsGa defect. The AsGa defect is situated at (a) the surface of ZB NW, (b) the subsurface of ZB NW, (c) the surface of WZ NW, and (d) the subsurface of WZ NW. The red and blue circles in band structures denote the position of CBM and VBM, respectively. The defect levels (DLs) in band structures are colored red, and horizontal dashed lines denote the position of the Fermi level.

Figure 6. (a) Doping models and the corresponding band structures of Mg-doped hydrogenated GaAs NWs with an AsGa defect. (a) Doping configurations in ZB NW, (b) the band structures of ZB NW with the Mg dopant in different sites, (c) the doping configurations in WZ NW, and (d) the band structures of WZ NW with the Mg dopant in different sites. The numbers 1−8 denote different doping sites of Mg dopant in NWs. The acceptor and donor levels in band structures are colored green and red, respectively. The horizontal dashed lines denote the position of the Fermi level.

effect. We find that both GaAs nanowires and thin films with an AsGa defect indicate a typical n-type character (see Figures S3 and S4). Moreover, these GaAs nanowires and thin films with the AsGa−Mg complex exhibit a semiconductor character (see Figures S5 and S6). The above results suggest that the formation of AsGa defects on NW surface will introduce donortype defect levels that can trap hole carriers of p-type dopants, resulting in the deactivation of p-type dopants. To reactive the p-type dopants of GaAs NWs, the effect of AsGa defects on electronic structure of NWs must be suppressed. There are two feasible methods to achieve this

diameter of experimentally reported GaAs nanowires is beyond 10 nm.47 Hence, the defect/dopant concentration of nanowires may be overestimated. However, the investigation of larger size nanowires is obviously beyond the scope of first-principles calculation. To prove the fundamental effect of AsGa defect on electronic structure and p-type doping of GaAs nanowires, electronic properties of smaller sized GaAs nanowires and GaAs thin films with an AsGa defect and the AsGa-Mg complex have been studied, respectively. The investigation of GaAs thin films is because two-dimensional thin films can be analogized to the large-size nanowires due to the considerable quantum size E


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The Journal of Physical Chemistry C goal: the first one is to reduce the density of AsGa in NWs, and the other is to passivate the defects. Owing to the defect concentration C ∝ exp(−Ef/kT), where Ef, k, and T are defect formation energy, Boltzmann constant, and temperature, respectively, the reduction of AsGa density can be implemented by adding its formation energy. It is seen from Figures 2 and 4 that the formation energy of AsGa increases with the decrease of As chemical potential. Therefore, the density of AsGa defects in GaAs NWs can be controlled by reducing the partial pressure of As-based molecules (e.g., As4 and AsH3) during the growth. The passivation of intrinsic defects depends on that introduced molecules or species can efficiently suppress the influence of surface point defects. Therefore, choosing appropriate passivating molecules or species is crucial for the reactivation of p-type dopants in III−V semiconductor NWs. From previous studies,32 NO2 molecules were proved as an efficient passivating species for eliminating surface states of semiconductor nanowires. Therefore, here we choose the NO2 molecule as an example to demonstrate the role of molecular passivation in the reactivation of p-type dopants in GaAs NWs. The adsorption of NO2 molecule at the AsGa site of NW surface leads to the reduction of O−N−O angle (N−O bond length) from 134° (1.19 Å) in the free molecule to ∼126° (1.24 Å) in the adsorbed one, suggesting that there is a strong interaction between NO2 molecule and NW surface. Figure 7 shows band

efficiency of Si and III−V NWs.48−50 For instance, Hang et al.50 reported that surface passivation of Cd-doped InAs NWs with organic ligands leads to the transition from n-type conductivity to the p-type one. Hence, the present study is helpful for the understanding of surface passivation effect of p-type doping in III−V semiconductor NWs.

4. CONCLUSIONS In summary, we have performed systematic DFT calculations to gain deep insight into the role of intrinsic point defects in electronic structure and p-type doping of GaAs nanowires. Our results have indicated that AsGa is a stable surface defect under a general growth condition (i.e., As-rich condition) no matter what crystal phase and surface environment of NWs are. The produced AsGa defects bring deep donor-type levels into the band gap, which is responsible for surface Fermi level pinning in conduction band and the deactivation of p-type dopants of GaAs NWs. In order to suppress the effect of surface point defects on the p-type doping of GaAs NWs, we proposed two feasible strategies by reducing As partial pressure during the growth and using the surface passivation with appropriate adsorbed species (e.g., NO2), respectively. The present results can be also applied to explain the surface Fermi level pinning effect and guide the reactivation of p-type dopants in other III− V semiconductor NWs.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.6b07624. Corner effect on electronic structure of zinc-blende GaAs NWs, part experimental data for growth parameters of GaAs NWs, the stability of AsGa at different depths of GaAs NWs, band structures of hydrogenated GaAs NWs and thin film with an AsGa defect, and band structures of hydrogenated GaAs NWs and thin film with an AsGa-Mg complex (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Ph 86-0571-86875622 (H.S.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported in part by the National Natural Science Foundation of China (Grants 11404309, 11347109, 51402275, 11334008, and 61290301) and the Fund of Shanghai Science and Technology Foundation (Grant 13JC1408800). Computational resources from the Shanghai Supercomputer Center are acknowledged.

Figure 7. Band structure and CDI distribution of dopant level in Mgdoped (a) ZB and (b) WZ GaAs NWs with the passivation of NO2 molecule at the AsGa site. The acceptor-type level is colored red, and the Fermi level is shifted to zero.

structure of Mg-doped GaAs NWs with the passivation of NO2 molecule at the AsGa site. It can be found that both ZB and WZ NWs indicate p-type character due to the adsorption of NO2 molecule. The result originates from that extra electrons of AsGa defect are compensated by adsorbed NO2 molecule, resulting in the reactivation of Mg dopant. The further CDI analysis indicates that electronic states of acceptor level are mainly localized around NO2 molecule and Mg dopant (see Figure 7a,b). The results suggest that the passivation of surface AsGa defects improves the p-type doping efficiency. In fact, such a method has been widely used to improve the p-type doping

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REFERENCES

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