**4.1.1 DLTS spectra and properties of a N-related electron trap**

The DLTS spectrum of Fig. 2(a) shows an electron trap (*E*2) at 0.69 eV below the CBM of GaAs. After rapid thermal annealing at 720C for 2 min, *E*2 disappears completely whereas a new electron trap (*E*3) appears at 0.34 eV below the CBM. From the Arrhenius plots of Fig. 2(d), the capture cross sections of *E*2 and *E*3 are calculated to be E2 = 8.1 × 10-15 cm2 and E3 = 7.5 × 10-18 cm2, respectively. Based on previous results about native defects in *n*-type GaAs, *E*2 and *E*3 are independent of *N* and considered to be identical to *EL*2 and *EL*3, respectively (Reddy et al., 1996). In order to focus only on *N*-related lattice defects, these two energy levels will be excluded from the DLTS spectra of *N*containing *n*-type GaAsN. The addition

Fig. 2. DLTS spectra of (a) *N* free as grown and annealed GaAs, (b) as grown n-type GaAs0.998N0.002, (c) annealed n-type GaAs0.998N0.002, and (d) Arrhenius plots of DLTS spectra.

Investigation of Lattice Defects in GaAsN

size of N compared with that of As.

**4.1.2 Nature of the electron trap E1** 

**4.1.2.1 Origin of reverse bias current in GaAsN** 

conditions, the reverse bias current *I*d(*T*) can be expressed by

*d*

in the depletion region.

where *I*

stable electron trap.

Grown by Chemical Beam Epitaxy Using Deep Level Transient Spectroscopy 497

rapidly. This behavior is explained by the large value of E1, compared with that of *E*2, *E*3, and other native defects in GaAs. The adjusted densities of *E*1 (*N*E1) in as grown and annealed samples are plotted in Fig. 3(d). *N*E1 increases considerably with increasing [*N*] in the film and persists to post thermal annealing. This indicates that *E*1 is a *N*-related and a

The defect center *E*1 was not observed previously in *N* free GaAs grown by CBE despite the existence of *N* species in the chemical composition of the As source. This can be explained through three possible scenarios. First, the absence of tensile strain in GaAs prevents the formation of *E*1. Second, the *N* atom in the atomic structure of *E*1 comes from the *N* source. Finally, the *N* atom comes from the *N* and *As* compound sources and in presence of tensile strain *E*1 can be formed. The tensile strain was reported in most theoretical and experimental studies. As given in Fig. 3(e) and using the ICTS, this idea is supported by the uniform distribution of *N*E1 in the bulk of GaAsN. This indicates that *E*1 is formed during growth to compensate for the tensile strain in the GaAsN films caused by the small atomic

Furthermore, the properties of *E*1 are identical to that of the famous electron traps reported by Johnston et Kurtz (Johnston and Kurtz, 2006) and Krispin et al. (Krispin et al., 2003) in MOCVD and MBE grown *n*type GaAsN, respectively. As illustrated in Fig. 3(d), the densities of these traps are approximately similar to *N*E1 despite the large difference in the density of residual impurities between the three growth methods. Therefore, the atomic structure of *E*1 may be free from impurities. Furthermore, by carrying out DLTS measurements for minority carriers in undoped *p*-type film, *E*1 was also observed. This

Two methods are used to verify whether *E*1 is a recombination center or not. The first method is indirect, in which the activation energy of deep levels is correlated with that of the reverse bias current in the depletion region of *n*-type GaAsN Schottky junction and *n*+- GaAs/*p*-GaAsN heterojunction. These two device structures are selected because the current is due mainly to electrons. The second method is direct, in which DC-DLTS is used to show the behavior of the electron traps in simultaneous injection of majority and minority carriers

The temperature dependence of the reverse bias current in the depletion region of *n*-type GaAsN Schottky junction and *n*+-GaAs/*p*-GaAsN is shown in Fig. 4(a) for reverse bias voltages of 0.5 and -0.5 V, respectively. At lower temperature, the dark current changes slowly in the two structures, then fellows an Arrhenius type behavior. As shown in Fig. 4(b), the same result is obtained by applying reverse bias voltages of 1 and -1 V. Under these

> *<sup>E</sup> IT I kT*

energy of the reverse bias current, the Boltzmann constant, and the temperature, respectively. The I-V characteristics deviate in the two samples from the thermionic emission. This is

, *E*, *k*, and *T* denote the limit of the high-temperature current, the thermal activation

( ) exp( ) (14)

indicates clearly that *E*1 is independent of doping atoms (see Fig. 3(f)).

of a small atomic fraction of *N* to GaAs leads to the record of a new electron trap (*E*1), at an average activation energy 0.3 eV below the CBM of GaAsN. The DLTS spectra of as grown and annealed *n*-type GaAs0.998N0.002 are given in Figs. 2. (b) and (c), respectively. The activation energies (*E*E1) and the capture cross sections (E1) of *E*1 for *N* varying GaAsN samples are given Fig. 3 (a) and (b), respectively. The fluctuation of *E*E1 from one sample to another can be explained by the effect of PooleFrenkel emission, where the thermal emission from *E*1 is affected by the electric field (Johnston and Kurtz, 2006). As illustrated in Fig. 3(c), with increasing the filling pulse duration, the DLTS peak height of *E*1 saturates

Fig. 3. Nitrogen dependence of (a) thermal activation energy, (b) capture cross section, and (d) adjusted density of *E*1 in as grown and annealed GaAsN samples. The large capture cross section is confirmed with (c) the filling pulse width dependence of the DLTS peak height of *E*1. (e) Density profiling of *E*1 in the bulk of GaAsN films, and (f) DLTS spectrum of undoped *p*-type GaAsN grown by CBE.

rapidly. This behavior is explained by the large value of E1, compared with that of *E*2, *E*3, and other native defects in GaAs. The adjusted densities of *E*1 (*N*E1) in as grown and annealed samples are plotted in Fig. 3(d). *N*E1 increases considerably with increasing [*N*] in the film and persists to post thermal annealing. This indicates that *E*1 is a *N*-related and a stable electron trap.

The defect center *E*1 was not observed previously in *N* free GaAs grown by CBE despite the existence of *N* species in the chemical composition of the As source. This can be explained through three possible scenarios. First, the absence of tensile strain in GaAs prevents the formation of *E*1. Second, the *N* atom in the atomic structure of *E*1 comes from the *N* source. Finally, the *N* atom comes from the *N* and *As* compound sources and in presence of tensile strain *E*1 can be formed. The tensile strain was reported in most theoretical and experimental studies. As given in Fig. 3(e) and using the ICTS, this idea is supported by the uniform distribution of *N*E1 in the bulk of GaAsN. This indicates that *E*1 is formed during growth to compensate for the tensile strain in the GaAsN films caused by the small atomic size of N compared with that of As.

Furthermore, the properties of *E*1 are identical to that of the famous electron traps reported by Johnston et Kurtz (Johnston and Kurtz, 2006) and Krispin et al. (Krispin et al., 2003) in MOCVD and MBE grown *n*type GaAsN, respectively. As illustrated in Fig. 3(d), the densities of these traps are approximately similar to *N*E1 despite the large difference in the density of residual impurities between the three growth methods. Therefore, the atomic structure of *E*1 may be free from impurities. Furthermore, by carrying out DLTS measurements for minority carriers in undoped *p*-type film, *E*1 was also observed. This indicates clearly that *E*1 is independent of doping atoms (see Fig. 3(f)).
