Intense photoluminescence between 1.3 and 1.8 μm from strained Si1−x Ge x alloysJ.P. Noël, N. L. Rowell, D. C. Houghton, and D. D. Perovic Citation: Applied Physics Letters 57, 1037 (1990); doi: 10.1063/1.103558 View online: http://dx.doi.org/10.1063/1.103558 View Table of Contents: http://scitation.aip.org/content/aip/journal/apl/57/10?ver=pdfcov Published by the AIP Publishing Articles you may be interested in 1.3 μm photoluminescence from InGaAs quantum dots on GaAs Appl. Phys. Lett. 67, 3795 (1995); 10.1063/1.115386 Broadband (8–14 μm), normal incidence, pseudomorphic Ge x Si1−x /Si strainedlayer infrared photodetectoroperating between 20 and 77 K Appl. Phys. Lett. 61, 1122 (1992); 10.1063/1.107688 Roomtemperature 1.3 μm electroluminescence from strained Si1−x Ge x /Si quantum wells Appl. Phys. Lett. 60, 3177 (1992); 10.1063/1.106734 High photoconductive gain in Ge x Si1−x /Si strainedlayer superlattice detectors operating at λ=1.3 μm Appl. Phys. Lett. 49, 155 (1986); 10.1063/1.97209 Ge x Si1−x strainedlayer superlattice waveguide photodetectors operating near 1.3 μm Appl. Phys. Lett. 48, 963 (1986); 10.1063/1.96624

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Intense photoluminescence between 1.3 and 1.8 JLm from strained Si1_xGex alloys

J.-P. Noel, N. l. Rowell, Do C. Houghton, and D. D. Perovic Division o/Physics, National Research Council, Ottawa KlA OR6, Canada

(Received 12 April 1990; accepted for publication 20 June 1990)

Intense photoluminescence (PL) from strained, epitaxial Si l _xGex alloys grown by molecular beam epitaxy is reported with measured internal quantum efficiencies up to 31 % from random alloy layers, single buried strained layers, and multiple quantum wens. Samples deposited at _400°C exhibited low PL intensity, whereas annealing at -600°C enhanced the intensity by as much as two orders of magnitude. This anneal treatment was found to be optimal for removal of grown-in defect complexes without creating a significant density of misfit dislocations. PL peak energies at 4.2 K varied from 620 to 990 meV for Ge fractions from 0.53 to 0.06, respectively. Efficient PL was due to exciton accumulation in the strained 5i J ~Gex layers of single and multiple quantum wells, where the band gap was locally reduced. Optical transitions associated with the PL occurred without phonon assistance.

There have been several recent reports of photolumi­nescence (PL) from SimGen atomic layer superlattices (ALSs) grown on partially relaxed Sil xGex alloy buffer layers, each result with a broad peak near 800 meV. I

3

These broad PL features have been attributed to a theoret­ically predicted quasi-direct band gap formed by Brillouin zone folding. 4 In this model, the band structure in SimGen ALSs is modified by the periodicity (m + n), the Ge/Si ratio (n/m), and the strain partitioning between Ge and 8i layers controlled by the in-plane lattice parameter of the buffer layer. All the experimental evidence to date, 1-3

however, is from material with extremely high threading dislocation densities, ~ 108 to 1010 em - 2. In addition, these samples were grown at low temperatures (350 to 470°C), where a high concentration of "grown in" defect com­plexes has been shown to limit PL efficiency. These recent PL studies of AsS and B6 doped Si grown by molecular beam epitaxy (MBE) have revealed that strong, bulk-like luminescence (due to annihilation of excitons bound to impurity atoms in concentrations _1017 cm- 3

) can be ob­tained from thin (;:S 5 Itm) epitaxial layers if the point­defect density is reduced by using growth temperatures (or post-growth anneals) above .~ 600 cC. In this letter we re­port a study of annealed, low threading dislocation density ( < 105 cm- 2

) Si]~. xGex random alloys in various geome­tries and show that zone folding effects are unnecessary for efficient PL.

The Si] _ xGex epitaxial layers studied in this work were grown by MBE using techniques which have been described previ.ously.7 Growth temperatures were main­tained low in the range 400 ± 25°C to avoid Ge islanding and strain relaxation. R Post-growth annealing was carried out either in vacuo or by rapid thermal annealing in a dynamic N2 ambient, Transmission electron microscopy (TEM) was performed on conventionally prepared cross­sectional specimens at 300 keY in a Philips EM 430. Table I summarizes structural parameters and PL data for the range of single and multiple quantum wells (SQW and MQW) and uncapped Si1 _ xGex alloy layers (EPI).

Details of the PL technique are given elsewhere.9 In these experiments, the PL was excited normal to the sam­ple surface with 51405 nm light over an area 3 mm in diameter (-! W cm- 2) and collected in the reverse direc­tion. For measurements at 4.2 K, the samples were im­mersed in liquid helium and above 4.2 K they were in helium gas. To determine the luminescence internal quan­tum efficiency, we have developed an absolute calibration method for the PL apparatus based on a source of known radiance, i.e., a blackbody simulator. lO In relating mea­sured intensity to the total intensity in a sample due to the laser excitation, it was necessary to account for interface effects from reflection and refraction of the emitted radia­tion at the free surface and for solid angle effects, since only a fraction of the isotropicaUy emitted PL was conected. Quantum efficiencies were verified by comparison with emission at the same wavelength from a direct gap material (100% efficiency), in this case in Ino.53Gao.47As layer on InP.

Figure 1 shows PI. spectra for sample B (a SQW, see Table 1) after various anneal treatments. The sharpest fea­tures in trace (a) from as-grown sample B were due to self annihilation of excitons bound to defect complexes 11 in the grown layers (lines N4, C, II' and 13 ) and free and bound excitons in the substrate (Fem,pto. and P~p). As the an­neal temperature (Tan) was increased above 500°C, the broad peak at ~950 meV peak (hereafter called the SLE or strain-localized exciton peak) grew in intensity such that at 575'C its maximum strength was ~ 16 times its as-grown strength. The measured internal conversion effi­ciency (11) for power integrated across the SLE peak was ,- 3%, corresponding to a quantum efficiency per photon of - 7% (1]Aouri Ai,,)' When the anneal temperature was increased above 625°C, the SLE peak disappeared abruptly in a fa<;hion closely correlated with the rapid (exponential) increase in density of misfit dislocations at the Si1 _ xGex/Si interfaces. 12

T'.vo broad features were observed in Fig. 1(a): the incipient SLE peak at 950 meV and another peak at 855

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TABLE I. Structural parameters (dimensions t in nm) of the Si I _ xGe/Si heterostructures and summary of PI. (energies in meV) features. QE is the internal quantum efficiency.

Sample Type x tSIOe lSi

A EPI 0.06 600 B SQW 0.12 130 300 C MQW 0.15 11.1 21.8 D MQW 0.17 6.2 10.3 E EPI 0.18 200 F MQW 0.25 7.3 10.9 G SQW 0.25 55 200 H SQW 0.28 50 100 I EPI(rlx) 0.40 120 J MQW 0.47 5.0 K MQW 0.49 5.5 L MQW 0.53 5.0

meV (of unknown origin) which anneals out at _500°C In Fig. 1 (b), the substrate features differ from those of Fig. l(a), probably due to a slightly higher PL temperature. Nonetheless, the SLE peak was relatively unaffected by such temperature differences. The defect complex lines of Fig. 1 (a) have essentially disappeared in Fig. 1 (b). 5,6 The most important change between Figs. 1(a) and 1 (b) is in the SLE peak, which has grown eightfold. Figure 1(e), similar to Fig. 1 (b), shows the strongest SLE peak of this series, - 2 X that in Fig. 1 (b). Further optimization of annealing conditions can maximize the SLE strength, as evidenced by a 600°C, 100 s anneal (results not shown) which gave SLE peak ~ 2 X that of Fig. 1 (c). An Arrhen­ius analysis of the SLE peak intensity versus Tan yielded (for Tan < 575 OC) an activation energy Ea of 352 meV and an effective growth temperature of 420°C. The relatively low activation energy suggests that the SLE peak growth on annealing resulted from the dissociation of defect com­plexes rather than from self diffusion, the latter with Ea-5 eV. As shown in Fig. led), the SLE peak decreases slightly on annealing at 625°C. However, at higher anneal temperatures the SLE peak collapses such that it is unob­servably small. at 700 0c. This behavior correlates wen with the onset of misfit dislocations8 at Si 1 _ xGexlSi interfaces and coincides with the growth of dislocation bands (Dl - D4) in the partly relaxed heterostructure. The en­ergies of D bands in Fig. 1 (e) are 807, 873, 928, and 986 meV, which agree well with energies for un strained

1--'-- -.~'~~;::~;-~~~~I---r.-~ '-----Y-"1 (el 800 0(:

D2. FET()+Or Nl D! ,,0 Dol D4 U

Jb,-~~~JL:1 ;i I ~I (,lm'e '---~

XI /~ I I (b)',"O," ~ ~

I (,) ~ <,ow, - .-FF;TO I

I'.:llergy (meV)

FIG. 1. PL spectra at 10 K (corrected for measurement response) for (a) sample B as-grown, (b) sample B annealed in ultrahigh vacuum for 900 s at 550 T, (c) at 575 "C, (d) at 625°C, and (e) at 800 T.

1038 Appl. Phys. Lett., Vol. 57, NO.1 0, 3 September 1990

20.0 21.0 20.0

N PI, peak PL width QE%

990 150 0.4 953 87 7

20 930 100 0.2 40 910 96 5

I 885 81 1.5 40 820 84 31

1 820 85 0.7 810 SO 0.1 810 100 0.4

15 695 80 5 15 685 115 0.8 15 625 100 0.6

SiO.88GeO.12.13 Figure 2 shows three intense PL spectra from annealed

samples B (SQW), F (MQW), and J (MQW) with Ge concentrations in the range 0,12-0.47 (see Table [). At energies above 1050 meV, only substrate features were ob­served. The features that dominate below 1050 me V are the SLE peaks, which represent internal quantum efficiencies of 7, 31, and 5% for samples B, F. and J, respectively. All SLE peak widths are similar and the peak position shifts to lower energy with increasing Ge content (discussed later with Fig. 3). Shapes of the SLE features in Fig. 2 all show the same degree of asymmetry in the low-energy sides of the peaks, and the most intense of these (sample F) dearly shows a smaller peak convoluted with the main peak at ~60 meV higher energy. The energy difference is close to the TO phonon energy in Si (58 meV), suggesting that intense SLE peaks we have observed were due optical tran­sitions with no phonons (NP) participating. Photolumi­nescence excitation spectroscopy (PLE) of sample B gave a spectrum whose low-energy edge was only ~ 10 meV higher than the onset of the SLE peak by PL, thus con­firming that no phonons assisted in the process. 14 PL "backscattered" from the top (100) surface and a cleaved (110) edge of sample F had the same spectral shape and was randomly polarized (± 1 %) in both cases.

Figure 3 plots PL peak energies versus average Ge fraction x within the alloy region for samples in Table 1. Values of x for epitaxial aHoy layers and SQWs were ob-

FIG. 2. PL spectra at 4.2 K after rapid thermal annealing of (a) sample D, (b) sample F, and ec) sample J.

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0.7 "'''l.q ..

·~·(-····I D. o 0.1 02 0.3 0.5 06

Ge FRACTION, x

FIG. 3. PL peak energy vs x for the present samples [ce) uncapped epitaxial layer, (IE ) single quantum well, (0) multiple quantum well], Ref. 1 (0), and Ref. 2 (h)· Data points shifted toward Eg (ulI§trained) are from partly relaxed samples with threading dislocation densities ~ J08 cm- 2•

tained with ±O.02 uncertainty by x-ray diffraction (XRD) using the (004) reflection. Kinematical simulations were used to fit XRD superlattice envelopes from MQWs to obtain x with ±O.03 uncertainty. It is emphasized that our samples did not have an atomic layer superlattice (ALS) structure, and all Si l _xGex layers were fully strained (ex­cept the alloy control sample I). Data from partly relaxed ALSs of other workers are given for comparison. Also shown in Fig. 3 are the variations with x of unstrained!5 and strained l6 band-gap energies Eg• Our data from fully strained structures of three different geometries (EPI, SQW, and MQW) agree well with the dashed line, corre­sponding to the function E = (Eg strained - !l.E), where the offset fl.E is-120 meV. Data from samples that have partly relaxed are observed to be shifted upward toward unstrained values of Elf

The trend for our PL data clearly shows that the ob­served SLE peak is related to the strained anoy band-gap energy. The precise origin ofthe peak, however, remains to be determined. Shapes and widths of SLE peaks (Fig. 2) do not resemble those of free excitons observed in pure Si, pure Ge, II or unstrained Si I _ xGex bulk alloysY1 The broad SLE peak and its shape suggest the possibility that annihilation of :lectron-hole droplets or other high-density exciton phases 11 might be responsible for the observed in­tense PL. Measurements of SLE peak intensity as a func­tion of laser power (down to 1 m W em' .2) did not reveal threshold behavior characteristic of electron-hole droplet luminescence. However, the possibility that high-density exciton phases were created cannot yet be discounted since there was considerable strain-induced migration of free ex­citons (electron-hole binding radius ~5 nm) from the la­ser excitation volume into the thin Si 1 __ xGex regions, akin to a two-dimensional plasma state. 17 Annihilation of exci­tons bound to unintentional impurities (e.g., C and 0, which are omnipresent in Si MBE films at concentrations _1017 cm-- 3) or dislocations is another possibility, but binding energies ~ 120 meV are unlikely, considering that typical exciton binding energies to impurity atoms are - 5 meV in Si and 1 meV in Ge.

Irrespective of the exact mechanism, the presence of a satellite peak to lower energy (-60 meV) and the PLE

1039 AppL Phys. Lett., Vol. 57, No.1 0, 3 September 1990

result suggest that the SLE peak does not involve momentum-conserving phonons. Furthermore, alloy scat­tering of carriers in Si I _ xGex appears to favor strong no­phonon components in aU opiticaI transitions, as evidenced by the large FEN? peaks observed in unstrained Si 1- xGe .. alloys. 13 Thus, optical transitions responsible for the intense SLE peak in the present work with strained alloys are "direct" (b.k = 0), even though the band struc­ture remains indirect (no zone folding). The lower inten­sity of SLE peaks from uncapped epitaxial alloys was due to growth-related compositional modulations that were weak « 5% with - 20 nm period observed in cross­sectional TEM) compared to intentional modulations in SQWs and MQWs, thus providing a smaller driving force for exciton diffusion to lower band-gap (strained) regions.

In summary, intense PL peaks from Sit _ xGex single and multiple quantum wens have been observed with en­ergies the shift consistently and predictably with the Ge fraction x. Strain relaxation in Si1 xGex layers on anneal­ing at high temperatures was accompanied by a decrease in PL intensity and an abrupt increase in the density of misfit dislocations. Intense PL in nonrelaxed structures resulted from exciton diffusion toward and accumulation within the strained, lower band-gap Si I _ xGex regions during photo­excitation at low temperature, possibly to densities suffi­ciently high for creation of two-dimensional electron-hole plasma states. Since the dimensions of single and multiple quantum wens we have studied were ~ 5 nm, zone folding or other theoretical treatments of recent interest in the literature could not have played a role in these structures. The potential for efficient light emission (PL and electrolu­minescence) in the wavelength range 1.3-1,5 ttm from Si I _ xGex alloys at room temperature remains alluring. Low-temperature internal quantum efficiencies of ~ 31 % and the persistence of PL to high temperature (~80 K) holds great promise for optoelectronic applications.

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