Figure 5.7: Schematic representation of the NiO thin film morphology deposited on silicon according to the oxygen partial pressure in the deposition chamber. The deeper the red, the more defects. At low oxygen concentration the thin film is considered to be homogeneous. When the oxygen concentration is increased the defects accumulate primarily at the grain boundaries before entering the whole crystal structure. As detailed in Appendix A.2, cationic inter-diffusion occurs at the interface which leads to the formation of a SiOx layer on top of a nickel rich region.
the samples prepared with 17% and 20% of oxygen concentration indicates that the grain boundaries would primarily support electrical conductivity of the RT-NiO thin films.
A schematic representation of the distribution of the chemical species responsible for the oxygen hole in the NiO thin film has been represented in Figure 5.7 according to the oxygen content in the deposition chamber. With regards to the STEM-EELS measurements, the oxygen hole can be considered as absent if the thin film is prepared at low oxygen concentration but gradually appears at grain boundaries before being present in the whole structure at higher oxygen concentration. The threshold when the oxygen holes appear in the bulk of the grain can be triggered when grain boundaries are saturated with oxygen holes.
in nickel oxide. The left-hand side of equation (5.1) are electronic defects (the dopants) which can be introduced in nickel oxide under oxygen-rich condition while the right-hand side represents the charge compensating species of the dopants.
Regarding the left-hand side (the dopants), it is generally accepted that p-type conductivity in NiO originates from nickel vacancies VN i00 [83, 164, 165].
Thus, Lany et al. have shown that in NiO crystals, under oxygen-rich conditions, nickel vacancies VN i00 and not oxygen interstitials O00i, are the defects leading to nonstoichiometry [82]. However, in most cases, the NiO thin films contain grain boundaries and to the best of our knowledge, the possibility of stabilizing oxygen interstitials at grain boundaries in a nickel oxide thin film has not yet been discussed in literature. In addition, it has been reported that the grain boundaries differ from bulk NiO as this is a structure where, in comparison to bulk NiO, the defect formation energy might be lower [38]
and the mobility of nickel vacancies is larger [180]. So a non-stoichiometry associated to oxygen interstitial at the grain boundaries has to be considered in this chapter.
Regarding the right-hand side of equation (5.1), associated to the compensating species, it is expected that the formation of one compound to be driven by thermodynamic rules. This would promote charge compensation mechanisms having the lowest formation energy. For instance, the formation of oxygen vacancies, VO••, in oxygen-rich conditions is very unlikely. Also for room temperature (RT) NiO thin films, it has been proposed that delocalized holes (free holes) would be interacting with oxygen atoms and localized holes might be found on both nickel and oxygen atoms [83]. These results available in literature would suggest that dopant in NiO (left hand side of equation (5.1)) can be compensated by free (delocalized) holes (h•), by localized holes on oxygen (O•O) and nickel (N i•N i) atoms.
5.4.2 Charge compensation in NiO
As a charge transfer material [75, 77, 79], the mobile charges in the valence band of a pure crystallographic NiO structure are transferred continuously from O 2p6 to Ni 3d8 bands during electrical displacement. Therefore, the charges have to overcome a potential barrier and this can be done with the help of thermal energy in the form of a phonon [97, 98]. The addition of the energy of the phonon to the charges (hole) produce what is called a polaron.
The conductivity (σ) is the product of the contribution of the number of charge carriers (holes) and their mobility:
σ(T) =e p, µp (5.2)
with p the hole density, e the elementary charge and µp the hole mobility.
According to Fermi statistics, the concentrationpof holes in the valence band
is:
p=Nv exp
−EF
kBT
(5.3) withNvthe effective density of states in the valence band,EFthe Fermi energy, T the temperature and kB the Boltzmann constant. Temperature activation of the mobility µp for a polaron hopping transportation in NiO follows an Arrhenius law [97, 98] following:
µ(T) =µ0
T exp
−EA
kBT
(5.4) Where µ0 is a temperature-independent pre-factor and the activation energy EA is basically the required energy for the charge carrier to overcome the barrier between two energy wells [97, 99]. It should be mentioned that as the Fermi energy EF and the activation energy EA can be both determined with temperature dependent conductivity experiments, these parameters can be easily inter-mixed.
Right-hand side of equation (5.1)
In-situ XPS measurements of the samples prepared in chamber #2 suggest that the Fermi energy is not lowered with increasing oxygen concentration during preparation of the samples (Figure 5.3). Moreover, looking at the in-situ electrical measurements realized on the samples prepared in chamber
#2 (Figure 5.2), the conductivity increases by three orders of magnitude when the oxygen concentration increases from 5 % to 15 %.
Assuming a constant hole mobility, according to equation 5.2, the increase of the conductivity would originate from an increase of the charge carrier density pin the valence band. It would imply that:
EF,15%−EF,5%=−kBTln p15%
p5%
(5.5) Therefore, an increase of the holes in the valence band by three orders of magnitude should be accompanied by a decrease of the Fermi energy by about 0.2 eV. This is not what is observed experimentally. In Figure 5.3 the Fermi energy is rather constant for the three samples. In RT-NiO thin films, the charge carrier concentration is therefore expected to be independent to the defects concentration. Eventually, the comparison of the in-situ XPS measurements and the in-situ electrical measurement would discard that the charge compensation is realized by holes (h•). The remaining possibilities are the formation of a positive charge on oxygen which lead to the formation of peroxo species(OO•) or on nickel which would form Ni3+ (N i•N i). Thus, to explain the increase of the conductivity, the mobility of the charge carrier has to increase with increasing oxygen concentration or that the compensating species (O•O and N i•N i) can be also considered as a charge carrier having a certain mobility.
The XPS data of the O 1s region (Figure 5.3) display a shoulder (O 1s(Def.)) which has been associated in literature to peroxo species (OO•) [166]. Also the O 1s(Def.) intensity increases with increasing oxygen concentration in the chamber. It might correlate with the appearance of the oxygen holes peak at 529 eV in the O-K edge spectra obtained by EELS, especially at high oxygen concentration. Therefore, the oxygen holes could be related to positive charge on the oxygen (OO•) arising from the charge compensation of the dopants.
However, the presence of Ni3+ (N i•N i) in the thin film cannot be completely discarded as the Ni 2p region provides a more shallow valley with increasing oxygen concentration (Figure 5.3). Indeed, materials having formally nickel in the +III state, as for NiOOH and Li doped NiO, the Ni 2p spectra provides a more prominent Ni 2p(Sat.) peak in comparison to the Ni 2p(Main) peak [173, 174]. Thus, the disappearance of the valley between the Ni 2p(Sat.) peak and the Ni 2p(Main) peak with increasing oxygen concentration, may be related to the presence of positive charges on nickel (N i•N i). It should be mentioned that the Ni3+ species might be controversial as its energy of formation could be relatively high in a pure NiO crystal structure [70].
However, we assume that positive charge compensation over nickel might be possible at or nearby grain boundaries.
Left-hand side of the equation (5.1)
STEM-EELS measurements highlight that the charge compensating species accumulate primarily at the grain boundaries before entering in the bulk of the grains with increasing oxygen concentration (Figure 5.6 and Figure 5.7).
Assuming that the charge compensating species are formed in the proximity of the dopants, it implies that dopants would accumulate also at the grain boundaries. This assumption makes it difficult to determine what the majority dopant in RT-NiO thin films. Indeed, as detailed above, the stabilization of oxygen interstitials at the grain boundaries cannot be discarded. Therefore, no conclusion can be reached about the true nature of the dopant in RT-nickel oxide at this stage and it is assumed that both nickel vacancies (VN i00 ) and oxygen interstitial (O00i) can dope a RT-NiO thin film.
5.4.3 Doping mechanism in RT-NiO thin films
Finally, as detailed above and in line with the Chapter 4, the dopant might be compensated by positive charge on oxygen O•O and to a lesser extent by a positive charge on nickelN i•N ibut not by holes (h•). Also, the identification of the dopant in the RT-NiO thin film is not obvious and we assume that nickel vacancies can coexist as much as oxygen interstitial at the grain boundaries.
Thus, as already shown in Chapter 4, the charge compensation mechanism in RT-NiO thin films must be:
5 Ni×Ni+ 4 O×O+O2→N iO+ 4 (1−δ)(N i•N i+O×O)
+ 4δ(OO• +N i×N i) +VN i00 +Oi00 (5.6)
with δ a parameter comprised between 0 and 1 to take into account charge compensation discrepancy over oxygen and nickel atoms.
5.4.4 Mechanism of conductivity in RT-NiO thin films
As represented in Figure 5.7, the oxygen holes would first accumulate at the grain boundaries before appearing in the bulk of the grain with increasing oxygen concentration. As it can be assumed that p-type doping is obtained in an oxygen-rich thin films, the RT-NiO thin films must be also oxygen-rich and these electrically active oxygen holes can be related to the accumulation of an oxygen-rich species at the grain boundaries. Although this hypothesis cannot be directly supported by EELS measurements, where the determination of the chemical composition is not accurate enough to emphasize an oxygen concentration distribution of less than 10 % inside the thin films [176, 179], it has been assumed that that higher conductivity is supported at the grain boundaries by the oxygen-rich secondary phase.
As mentionned above, charge transportation in NiO is in theory realized through the polaron hopping mechanism [73]. It means that charges in NiO (holes) need to gain enough thermal energy (phonon) to overpass an energy barrier Ea (see equation 5.4) located over a nickel atom separating two energy wells situated over oxygen atoms [97, 98] (Figure 5.8). Interstingly, literature reports low electrical activation energy in the 0-0.15 eV range for RT-NiO thin films [150, 175] whereas for pure NiO, the activation energy is higher with at least 0.6 eV [50].
Moreover, the pinning of the Fermi energy for the RT-NiO thin films suggests that the creation additional positive species in the valence band with temperature would lead to the formation of peroxo O− or Ni3+ species. So for such materials as RT-NiO thin films, the concentration of charge carrier in the valence band is temperature independent which means that pin equation 5.2 is constant. Therefore, the activation energy reported in RT-NiO thin films could be related to the energy barrier that mobile charges as to overcome in the RT-NiO thin films.
As detailed previously, the interpretation of the situ electrical and in-situ XPS data suggests that the increase of the conductivity of RT NiO thin films with increasing oxygen concentration can be associated to an increase of the mobility of the charge carriers or that the compensating species provide support for conductivity.
Regarding the mobility of holes in the valence band, as the grain boundaries is supposedly an oxygen-rich material, it means that the alternative transfer from cationic to anionic site, typical of polaron hopping, could be replaced by a continuous band such as, for example, an O 2p band at the grain boundaries.
Thus, EAis reduced at the grain boundaries while it would have a higher value in the bulk of the grain. The low EAfor RT-NiO thin films would suggest that the energy barrier provided by the nickel atoms, and seen by the polarons, is
Figure 5.8: Schematic representation of the potential profile encountered by a charge in a stoichiometric NiO structure or in the case of the presence of a nickel vacancies.
The electrical activation energy EA value are taken in literature [50, 150, 175].
substantially reduced and the charges can flow from atoms to atoms without encountering substantial energy barrier (Figure 5.8).
Regarding the fact that the compensating species can also participate in the conduction process, as mentionned in the Chapter 4, as reported by optical adsorption measurements on Li doped NiO, a double feature is observed in the bandgap and can be assigned to the presence of an acceptor state∼1 eV above the valence band [86]. This acceptor state could be made of the species compensating defects for RT-NiO thin films and could be also an alternative electrical path to the charges participating in the conduction process.
Eventually, the assumption that grain boundaries support electrical path correlates also with the reported dependence of the conductivity on the thin film thickness for the RT-NiO thin films in the 100 nm – 1000 nm range, where it can be observed that the thin film conductivity decreases with thickness [87, 150]. Indeed, as seen previously, when deposited by sputtering, the RT-NiO thin films are polycrystalline and composed of elongated grains (Figure 5.4). As proposed by Deuermeier et al. for Cu2O thin films, it can be assumed that the grain boundary density is inversely proportional to the thin film thickness [181]. Therefore, similarly with NiO, in assuming that the quantity of grain boundaries decreases with thickness, as the results in this chapter supports the fact that grain boundaries are the main conduction paths, it can explain why the RT-NiO thin films conductivity decreases with thickness.