Himwas et al. Nanoscale Research Letters 2011, 6:496 http://www.nanoscalereslett.com/content/6/1/496 not entirely unexpected. The GaAs emission comes lar- gely from the buffer and the substrate and thus reab- sorbed by the narrower gap CHP. But the CHP is compressively strained. Upon annealing, strains in zinc blende crystals with similar misfits relax via misfit and threading dislocations (TDs) [35]. MDs are confined in the growth plane, i.e., at the heterointerface, whereas TDs penetrate the layer. Improvement in GaAs layers is thus marred by the degradation of the CHP layer which explains why the deterioration of the InGaAs CHP sig- nal is accompanied by the appearance of the weak 1.5- eV GaAs peak. For thin InGaAs sandwiched between GaAs, however, misfit strain is small and, upon anneal- ing, interdiffusion causes a small blueshift in PL with no crystalline degradation [36]. This is not the case in our controlled CHP sample where misfit is large but neces- sary to induce the interfacial dislocation network that enables the formation of orthogonally aligned QDs. The rapid degradation of the optical quality of QDs on CHP upon high-temperature in situ annealing thus can- not result from the degradation of QDs since the con- trolled QDs subject to the same annealing conditions remain optically active despite the longest annealing times as seen in Figure 3a; it must therefore result from the degradation of the CHP itself as the PL from the controlled CHP shown in Figure 3b. The thermal budget for the overlayers on QDs on CHP is thus lower than that those on conventional QDs and must be well below 700°C. Alternatively, improvement sought from post- growth annealing may be carried out ex situ at a lower temperature and, consequently, with a quantitatively and qualitatively different improvement. For ex situ anneal- ing studies in the next section, sample D is chosen over C because of its well-resolved QD peaks and the richer PL characteristics which act as sensitive probes for material’s integrity. Ex situ annealed After 350°C ex situ annealing in a forming gas for between 30 and 120 min, the overall quality of the QDs on CHP (sample D) improves as shown in Figure 4a. The improvements are twofold. First, the QDs and WL emissions have overall increased intensities as shown in Figure 4b. Second, the 1.42-eV WL emission has reduced FWHM as shown in Figure 4c. This is the 1.42- eV WL, not the 1.34-eV WL whose changes upon annealing cannot be resolved as it is too close in energy to the 1.377-eV CHP peak. The improvement is not related to material crystallinity as the temperature is too low to have any effect. Instead, it is related to the abun- dance of hydrogen and the supplied thermal energy that is sufficiently high to dissociate the hydrogen atoms/ molecules, driving them through the structure, neutra- lizing dislocations and dangling bonds (MDs), and Page 5 of 7 ) (a) . .a u ex-situ anneal (e 120 min c n 24 meV e c s ine60 (b) (c) lum30 o to0 h P 1.2 [110] WL In In te te QDs sn rga [110] it(ya ted [1-10] ..) u [1-10] WL 65FW 60H 55M 50(m 45e 40)V 1.3 1.4 Energy (eV) 1.5 0 30 60 90 120 Annealing Time (min) Figure 4 PL spectra of sample D, and changes in integrated intensity and FWHM. (a) PL spectra of sample D [QDs on CHP] subject to 350°C ex situ hydrogen annealing for 0, 30, 60, and 120 min (offset for clarity). Changes in (b) integrated intensity and (c) FWHM of the main emission peaks as a function of annealing time. making available more free carriers. Low-temperature hydrogen annealing is a standard Si process that effec- tively neutralizes interface-trapped charges [37] since hydrogen can diffuse several microns into Si even at room temperature [38]. Figure 4b shows the integrated intensities of the WL and the QDs as a function of annealing time. The 1.42- eV WL intensity increases immediately and significantly during the first hour of annealing after which no further improvement can be made. This emission arises from the radiative recombination of carriers photoexcited in the WL itself and those captured into the WL from the overlying GaAs capping layer. The improvement is due to the fact that the GaAs capping layer is grown at a relatively low temperature of 500°C on a lattice-mis- matched layer which result in non-radiative defects and strain. In addition, the free GaAs surface is full of sur- face states which nullify any photoexcited surface car- riers. Annealing in hydrogen makes available plentiful hydrogen atoms which subsequently permeate the epi- layers and bond to dangling bonds and crystalline defects. This has two important consequences. First, it frees up carriers in the GaAs capping layer which then trickle down to the WL, increasing the intensity. Sec- ond, it relieves some strains caused by defects which induce lattice distortion. Changes in interfacial strain would result in changes in band offsets which then affect the eigenenergies of carriers confined by one or two of such interfaces. A closer inspection of Figure 4a reveals that the increased WL intensity indeed occurs together with a 24-meV blueshift, consistent with values reported by Ryu et al. who achieved similar degrees of blueshift at much higher annealing temperatures of 900° C and above [36]. Our results indicate that strains may play a much greater role than cation interdiffusion in