We now consider how the intrinsic native defect concentrations in Li 3 OCl respond to aliovalent doping. Figure shows, again, the Fermi-energy dependent defect formation energies and equilibrium defect concentrations in as-synthesised Li 3 OCl, but now compares these "stoichiometric" defect concentrations with those predicted for an effective supervalent dopant concentration of r[M r ] = 10 18 cm -3 (see eq. ( )) . Supervalent doping under Li-poor conditions principally increases the [X q ] / cm -3 concentration of lithium vacancies, which are the lowest formation energy negatively-charged defect species in the parent undoped material-the system therefore behaves qualitatively as predicted by simple lithium-defect charge-compensation models. Supervalent doping under Li-rich conditions principally increases the concentration of the O Cl antisite, which is now the lowest formation energy negatively-charged defect species. Under Lirich conditions, therefore, supervalent doping produces a smaller increase in lithium vacancy concentration than predicted by simple models that assume direct chargecompensation by lithium defects. The equivalent analysis for the case of subvalent doping is shown in fig. . Under both Li-rich and Li-poor synthesis conditions, subvalent doping gives only a modest increase in the concentration of lithium interstitials, and the principal effect is instead to increase the concentration of V Cl lithium vacancies. We also observe a significant decrease in the lithium vacancy concentration. Although lithium interstitials are expected to be more mobile than lithium vacancies, this analysis suggests the potential for a regime where subvalent doping causes a net decrease in ionic conductivity due to the decrease in the dominant charge-carrying lithium-defect species, i.e., lithium vacancies. We return to this point in more detail in section III D where we quantify how ionic conductivi-ties are predicted to vary as a function of supervalent or subvalent dopant concentrations.