**1. Introduction**

488 Solar Cells – New Aspects and Solutions

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**5. References** 

With only 3 % of N and 9 % of In, InGaAsN with a band gap of 1.04 eV was obtained and could be lattice matched to GaAs and Ge. This dilute nitride semiconductor has been selected as a promising candidate for high efficiency multijunction tandem solar cells (Geisz and Friedman, 2002). However, the diffusion length of minority carriers and the mobility are still lower than of that in GaAs or InGaAs and showed a considerable degradation with increasing the N concentration. These electrical properties are insufficient to insure the current matching in the multijunction solar cell structure AlInGaP/GaAs/InGaAsN/Ge (Friedman et al., 1998). An obvious reason of such degradation is the high density of Nrelated lattice defects that can be formed during growth to compensate for the tensile strain caused by the small atomic size of N compared with that of arsenic (As) and to the large miscibility of the gap between GaAs and GaN. These defect centers are expected to act as active recombination and/or scattering centers in the forbidden gap of the alloy (Zhang & Wei, 2001). However, no experimental evidence has yet been reported. On the other hand, the conductivity of undoped *p*-type InGaAsN or GaAsN and their high background doping (Friedman et al., 1998; Kurtz et al., 1999; Moto et al., 2000; Krispin et al., 2000) prevent the design of wide depletion region single junction solar cell and the fabrication of intrinsic layer to overcome the short minority carrier lifetime. This serious problem was expected in the first stage to the density of unintentional carbon in the film (Friedman et al., 1998; Kurtz et al., 1999; Moto et al., 2000). However, the carrier density in some InGaAsN semiconductors was found to be higher than that of carbon (Kurtz et al., 2002). Furthermore, the high density of hydrogen (up to 1020 cm−3) and the strong interaction between N and H in InGaAsN to form N-H related complex were confirmed to be the main cause of high background doping in InGaAsN films (Li et al., 1999; Janotti et al., 2002, 2003; Kurtz et al., 2001, 2003; Nishimura et al., 2007). In addition, N-H complex was found theoretically to bind strongly to gallium vacancies (VGa) to form N-H-VGa with a formation energy of 2 eV less than that of isolated VGa (Janotti et al., 2003). These predictions were supported experimentally using positron annihilation spectroscopy results (Toivonen et al., 2003).

Investigation of Lattice Defects in GaAsN

**2.1 Fundamental concept of DLTS 2.1.1 Capacitance transient** 

Grown by Chemical Beam Epitaxy Using Deep Level Transient Spectroscopy 491

current (TSC) (Leonard & Grossweiner, 1958; Bube, 1960), the admittance spectroscopy (Losee, 1974), the increase or decay curves of photoconductivity (Rose, 1951; Devore, 1959), the optically stimulated conductivity (Lambe, 1955; Bube, 1956), and the analysis of spacecharge-limited currents as function of applied voltage (Smith & Rose, 1955; Rose, 1955; Lampert, 1956). The basic concept of using the change of capacitance under bias conditions by the filling and emptying of deep levels was already anticipated fifty years ago (Williams, 1966). The thermally stimulated capacitance (TSCAP), which gives the temperature dependence of junction capacitance (Sah et al., 1978; Sah & Walker, 1973), was used. By dressing the properties of all these techniques, D. V. Lang found that they lacked the sensitivity, the speed, the depth range of recorded trap, and the spectroscopic nature to make them practical for doing spectroscopy on non-radiative centers. For that, in 1974, he proposed DLTS as a characterization method of lattice defects that can overcome the disadvantages of the other methods (Lang, 1974). DLTS is based on the analysis of the change of capacitance due to a change in bias condition at different temperatures. It can be applied to Schottky contacts and *p-n* junctions. DLTS has advantages over TSC due to its better immunity to noise and surface channel leakage currents. It can distinguish between majority and minority carrier traps, unlike TSC, and has a strong advantage over admittance spectroscopy, which is limited to majority-carrier traps. Comparing with TSCAP, DLTS has much greater range of observable trap depths and improved sensitivity. Despite the success of optical techniques such as photoluminescence to characterize superficial levels, they are rarely used in the study of non-radiative deep levels. Furthermore, such experiences must be done in the infrared domain. However, sensors are less sensitive than in the visible domain. Thus, we need a technique, which can separate between minority and majority traps and evaluate their concentrations, their energies, and their capture cross sections.

To fully understand DLTS, it is worth to have a basic knowledge of capacitance transients arising from the SCR of Schottky contacts or *p+-n*/*n+-p* asymmetric junctions. If a pulse voltage is applied to one of these device structures that is originally reverse-biased, the SCR width decreases and the trap centers are filled with carriers (majority or minority depending on the structure). When the junction is returned to reverse bias condition, the traps that remains occupied with carriers are emptied by thermal emission and results in a transient decay. The capacitance transients provide information about these defect centers. Here, we restrain our description to a *p+-n* junction where the p-side is more much heavily doped

The causes of change in capacitance depend on the nature of applied voltage. In case of reverse biased voltage, the junction capacitance, due to the change in SCR width, is dominant. However, when the applied voltage is forward biased, the diffusion capacitance, due to the contribution of minority carrier density, is dominant. The basic equation

> *D T T Rb D D eN N t N t Ct A <sup>C</sup> VV N N*

where A is the contact area, *V*b is the built-in potential, 0 is the permittivity of the semiconductor material, and e is the elementary charge of an electron. *C*0, *N*T, *N*D, and

<sup>0</sup> ( ) 1 exp 1 exp 2( ) 2 <sup>2</sup> (1)

than the *n*-side, which gives the SCR almost in the low doped side.

governing the capacitance transient in the *p+-n* junction is expressed by

0

On the other hand, similar electrical properties were obtained in InGaAsN grown by metalorganic chemical vapor deposition (MOCVD) and molecular beam epitaxy (MBE) despite the large difference in the density of residual impurities, which excludes them as a main cause of low mobilities and short minority carrier lifetimes. For that, lattice defects, essentially related to the N atom, were expected to be the main reason of such degradation. Several theoretical and experimental studies have investigated carrier traps in InGaAsN films. Theoretically, using the first principles pseudo-potential method in local density approximation, four N-related defects were proposed: (AsGa-NAs)nn, (VGa-NAs)nn, (N-N)As, and (N-As)As ( Zhang & Wei, 2001). While the two first structures were supposed to have lower formation probabilities, the two split interstitials (N-N)As, and (N-As)As were suggested to compensate the tensile strain in the film and to create two electron traps at around 0.42 and 0.66 eV below the conduction band minimum (CBM) of InGaAsN with a band gap of 1.04 eV, respectively (Zhang & Wei, 2001). Experimentally, the ion beam analysis provided a quantitative evidence of existence of N-related interstitial defects in GaAsN (Spruytte et al., 2001; Ahlgren et al., 2002; Jock, 2009). Furthermore, several carrier traps were observed in GaAsN and InGaAsN using deep level transient spectroscopy (DLTS). A deep level (E2/H1), acting as both an electron and a hole trap at 0.36 eV below the CBM, was observed (Krispin et al., 2001). Other electron traps in GaAsN grown by MBE were recorded: A2 at 0.29 eV and B1 at 0.27 eV below the CBM of the alloy (Krispin et al., 2003). In addition, a well known electron trap at 0.2 0.3 eV and 0.3 0.4 eV below the CBM of *p*-type and *n*-type GaAsN grown by MOCVD were observed, respectively (Johnston et al., 2006). Although the importance of these results as a basic knowledge about lattice defects in GaAsN and InGaAsN, no recombination center was yet experimentally proved and characterized. Furthermore, the main cause of high background doping in p-type films was not completely revealed.

Chemical beam epitaxy (CBE) has been deployed (Yamaguchi et al., 1994; Lee et al., 2005) to grow (In)GaAsN in order to overcome the disadvantages of MOCVD and MBE. It combines the use of metal-organic gas sources and the beam nature of MBE. (In)GaAsN films were grown under low pressure and low temperature to reduce the density of residual impurities and to avoid the compositional fluctuation of N, respectively. Furthermore, a chemical N compound source was used to avoid the damage of N species from N2 plasma source in MBE. Although we obtained high quality GaAsN films gown by CBE, the diffusion length of minority carriers is still short (Bouzazi et al., 2010). This indicates that the electrical properties of GaAsN and InGaAsN films are independent of growth method and the problem may be caused by the lattice defects caused by N. Therefore, it is necessary to investigate these defects and their impact on the electrical properties of the film. For that, this chapter summarizes our recent results concerning lattice defects in GaAsN grown by CBE. Three defect centers were newly obtained and characterized. The first one is an active non-radiative *N*-related recombination center which expected to be the main cause of short minority carrier lifetime. The second lattice defect is a N-related acceptor like-state which greatly contributes in the background doping of p-type films. The last one is a shallow radiative recombination center acceptor-like state.
