Optical Characterization of Semiconductors || Case Studies: Photoluminescence Characterization

5 Case Studies: Photoluminescence Characterization

5.1 General introduction to case studies

In Chapters 1 through 4 I have provided the motivation, and the scientific and technical background, for the optical characterization of semiconductors. In the next three chapters, I illustrate these principles by case studies, actual characterization applications taken from the literature. This chapter deals with photoluminescence methods; Chapter 6, with Raman analysis; and Chapter 7, with infrared characterization. Each chapter is organized similarly. For each of the three optical methods, I list the main semiconductor properties it measures, such as band behavior, or lattice modes. Under each property, I work my way through specific applications. Most of my case studies involve silicon, G a A s , or A l x G a x _ x A s , the most widely used elemental, binary, and ternary materials respectively. But when it is valuable to illustrate an important capability, or to show diversity, some cases draw on other binary and ternary group I I I - V semiconductors, and more exotic systems yet, such as group I I -V I ternaries. Hence there are applications to InSb, CdTe , ZnSe, I n ^ ^ G a ^ A s , Inx-^Ga^PÂsx-^, Hg i -^Cd^Te , Cdx-^Mn/Te , Z n ^ ^ S e , and other materials. There is a general progression in the applications, although I do not always strictly adhere to it, from bulk samples and films to the more complex quantum wells and superlattices.

The discussion of each application is meant to remind the reader of the principles underlying it; to show its usefulness in meeting a specific problem or class of problems in the analysis of semiconductors; and to point out important details of the optical methods, especially when they are unusual or novel. It is not possible, of course, to give every detail of the instrumentation and techniques; but these can be found in the original references, chosen partly for their appropriate descriptions.

M y intention is not to provide anything like complete coverage of the optical characterization of semiconductors. The field and its literature are too vast to allow this even if I chose to do so. Also there are so many specialized needs that it would be impossible to illustrate them all in this book. It is more fruitful, I believe, to show a number of carefully selected examples. These are meant to present principles and techniques which can serve in many different settings. It is more important to illustrate what can be done than to show every possible example. This approach may help readers determine which methods can serve for their own purposes, and to generate ideas for new applications.

5.2 Introduction to photoluminescence case studies

Photoluminescence ( P L ) is one of the most useful optical methods for the semiconductor industry, with its powerful and sensitive ability to find impurities and defects in

62 Case studies: photoluminescence characterization

silicon and group I I I - V element semiconductors, which affect materials quality and device performance. A given impurity produces a set of characteristic spectral features. This fingerprint identifies the impurity type, and often several different impurities can be seen in a single P L spectrum. In some cases P L goes beyond bare identification, to measure impurity concentrations. In another use, the half widths of PL peaks are an indication of sample quality and crystallinity, although such analysis has not yet become highly quantitative. Finally, P L is sensitive to stress, and can measure its magnitude and direction.

Photoluminescence can also determine semiconductor band gaps. This is important for A l x G a ] _ x A s and other ternary alloys whose gap varies with the compositional parameter x, yet must be accurately known for applications. When the relation between gap energy and χ is known, the P L measurement of gap can be inverted to determine x. From this, a two-dimensional map of alloy composition can be obtained as the exciting laser beam is scanned across the face of a sample, a useful tool to determine inhomogeneity. Among the optical characterization methods I treat, P L is probably the best developed to carry out such spatial scanning, with commercial equipment available.

There are several good overviews of P L and its use in characterization. Goldberg (1966) gives a general introduction to luminescent processes. The reviews by Lightowlers (1990) and Schroder (1990) focus on P L characterization of semiconductors. Thewalt et al. (1990) review P L spectroscopy using Fourier transform methods. Conzelman (1987) gives in concise form some of the difficulties in using P L as a quantitative tool. Domanevskii et al. (1988) present a theoretical analysis of radiative recombination in GaAs doped with shallow donors and acceptors.

5.3 Band emission

5.3.1 Band gaps and offsets

Wilson (1989) has surveyed the P L analysis of buried interfaces in heterostructures, which are inherently difficult to probe by direct physical or electrical contact. She points out the fundamental importance of the energy alignment, or offset of the band levels, between the materials making a heterostructure. Theory cannot yet precisely predict these offsets, yet they are necessary to calculate superlattice band structures. Moreover, stress at the interface, the orientation of the growth surface, and the formation of interface dipoles all affect the band alignments. The band alignment can be found from P L measurements, but in ways which have been model dependent. Wilson and coworkers, however, have noted that in A l ^ G a ^ ^ A s / G a A s systems, the recombination radiation between conduction electrons and valence holes in adjacent layers is a direct measure of the band alignment. Figure 5.1 shows the conduction and valence band offsets in A l ^ G a ^ ^ A s versus the fraction of aluminum, as derived from such P L data.

The usefulness of P L to measure band gaps is illustrated in the work of Wagner (1984). He made a complete P L and photoluminescence excitation ( P L E ) analysis of n- and p-silicon doped from 10 1 7 to 4 x 10 2 0 cm~ 3 with phosphorus, boron, or arsenic. The optical measurement of band gap involves complications. A n absorption measurement determines the gap between the top of the valence band and the Fermi

Band emission 63

level (see Fig. 5.2). When the doping is heavy, the Fermi energy is significant and the absorption band gap—called the 'optical gap' EGi—is greater than the conduction band gap EG2 measured from the valence band to the bottom of the conduction band. In conventional P L , the emitted photon must arise from the states between the bottom of the conduction band and the Fermi level; hence P L spectra represent a range between EGX and EG2, as shown in the figure.

Another complication is that the band gap decreases under heavy doping, because of the interaction of the electrons; hence EGl is reduced from its value in a pure sample. This is important information for device modeling, since the carrier density may be high enough to affect the band. One motivation for Wagner's work was to explain observed discrepancies among data from absorption, P L , and transport measurements. Another was to provide a consistent set of accurate data to compare with theoretical calculations for the reduced gap.

For his P L spectra, Wagner used conventional methods with Ar 4 " or K r + lasers. For the P L E work, a dye laser pumped a tunable color-center laser which produced 100-250mW over the range 1.01-1.13μπα. His data showed that absorption and P L


Fig. 5.2 Schematic diagram of the band structure near the gap of a heavily doped n-type semiconductor, showing the optical gap EGi, the reduced gap EG2, and the Fermi energy EF. (After Wagner (1984).)

spectra agreed for the reduced gap, and that this gap depends on carrier concentration as nm. Wagner (1985) continued to investigate band-gap shrinkage in highly doped silicon. He extended to 300 Κ the earlier measurements at 10 K , to obtain results more useful for device characterization, and obtained plots of the reduced band gap versus carrier density at 20 and 300 Κ for both n- and p-type samples.

A n extension of such fundamental work is important in a related, more complex system, the Si i-^Ge^ alloy. This can be made in disordered form over the whole composition range χ = 0 to 1. The alloy band gap spans the range 0.66-1.12 e V , and its electrons are more mobile than in silicon. These considerations make Six-^Ge^ an interesting material for devices. Weber and Alonso (1989) carried out fundamental P L measurements in bulk Six-^Ge* made by a zone-leveling method, and Six-^Ge* layers grown by liquid-phase epitaxy ( L P E ) and vapor phase epitaxy ( V P E ) . The bulk samples were cut from poly crystalline ingots with large grains. Those nominally undoped were p-type with hole concentrations near 10 1 5 cm~ 3 . Others, doped with arsenic, had electron concentrations up to 3 x l 0 1 7 c m ~ 3 . High-resolution spectra were obtained at sample temperatures between 1.6 Κ and room temperature. The

Band emission 65

Photon energy (eV) 1.00 0.95 0.90 0.85 0.80 0.75 0 70 Π I I 1 1 1 f -

1200 1300 1400 1500 1600 1700 1800

Wavelength (nm)

Fig. 5.3 Photoluminescence spectra of epitaxial Sii_ x Ge x films on silicon, at 2K. The main peak in each trace is a bound exciton transition, very near the gap value, whose energy decreases as χ increases. The free exciton (FE) lines in some traces also lie near the gap. T O and L A phonon-related lines appear. (After Hansson et al. (1990).)

bulk samples gave P L spectra which varied smoothly with χ and were of higher quality than those from the L P E and V P E specimens. No-phonon gap transitions were observed, as well as those involving T A and T O phonons to conserve momentum, required because of the indirect gaps in silicon and germanium. These gave, for the


gap energies in electron-volts versus χ at low temperatures,

E x = 1.155 - 0Λ3χ + 0.206x2

(

£ L = 2.010 — 1.270* ^ }

at the X and L points, which are identified in Fig. 3.2. Such data are extremely useful in characterization and device applications. Later, Hansson et al. (1990) examined LPE-grown S i ^ G e * on silicon substrates. They reported the narrowest lines yet seen for this alloy, and confirmed that the Ε versus χ relation derived by Weber and Alonso (1989) correctly determined sample compositions. The variation with composition is clear in Fig. 5.3.

5.3.2 Stress

Stress is an important issue in the heterostructures essential to semiconductor devices. Even if two semiconducting materials have the same crystalline structure, they generally have different lattice constants and thermal expansion coefficients. This makes it difficult to grow epitaxial layers or complex layered structures of high quality.

One issue of immediate interest is the heteroepitaxy of GaAs on silicon; it has received much attention, for instance in the recent proceedings edited by Choi et al. (1988). Since silicon is the basis for electronic technology, devices using GaAs or A l ^ G a ^ ^ A s are expected to be built on silicon and to interface with it. In fact, silicon could act as a cheap substrate for G a A s and A l ^ G a ^ ^ A s films and microstructures. A long-term possibility is to combine photonic devices based on group I I I - V element materials with integrated circuits based on well-developed silicon technology. But the mismatch in properties produces one difficulty in the growth of GaAs on silicon to achieve electronic compatibility, the accumulation of stress. The lattice constant of GaAs is 4% larger than that of silicon, which produces compressive stress in the GaAs . The thermal expansion coefficients of the two materials also differ enough (<*GaAs = 6.8 x 1 0 " 6 o C _ 1 , a S i = 2.6 χ 1 0 " 6 o C _ 1 ) to produce a tensile strain of 0.1%, opposite to the strain due to lattice mismatch (Bugajski et al., 1988). Epitaxial films of GaAs on silicon (GaAs/Si) show disorder and poor crystalline quality, with a high dislocation density at the interface or in the film. In photonic devices, such interface strain can form electrically active defects which quench optical emission. Photoluminescence is one of the most powerful means to characterize the quality of the GaAs layer and the GaAs-silicon interface. ( I discuss Raman scattering, also extremely useful for this application, in Chapter 6 ) .

When the epitaxial layer has a lattice constant larger than the substrate, as is true for GaAs on silicon, it suffers biaxial compressive stress in two perpendicular directions in the plane of the sample face. Such stress can be decomposed into a tensile hydrostatic pressure component, and a compressive uniaxial stress component perpendicular to the GaAs/Si interface. The hydrostatic component changes the lattice constant. This alters the band gap of G a A s , which shifts the P L peak at the gap. The uniaxial component of the stress reduces the crystalline symmetry of the GaAs film. This splits apart the degenerate valence bands at k = 0 (see Fig. 3.3), producing a new light-hole ( L H ) valence band and a second heavy-hole ( H H ) band. (These are also denoted the rrij = 3/2 and 1/2 bands, in reference to their spin properties.) Each valence band participates in carrier excitation and recombination. Hence uniaxial

Band emission 67

stress splits singlet P L peaks into doublets. These shifts and splits are what make P L such an effective probe of stress.

In one recent application, Bugajski et al. (1988) examined MBE-grown GaAs/Si. The 2-3 μπι thick GaAs films were formed atop a G a A s buffer, which was grown on standard 2in. (100) silicon wafers slightly misoriented toward ( O i l ) . The films were intentionally doped with silicon at concentrations over 5 x 1 0 1 7 c m - 3 , or were unintentionally doped with concentrations of 1 0 1 4 - 1 0 1 5 c m - 3 . Some were annealed. Bugajski et al. found substantial differences in P L spectra at 5 Κ from annealed and unannealed samples, and from heavily and lightly doped ones. They derived strain values from the measured energies by referring to the induced hydrostatic and uniaxial components mentioned above. Under strain, the gaps ELH and EHH associated respectively with the conduction band to light hole transition, and the conduction band to heavy hole transition, become:

- L H -2a (C i i - Cu) , , ( d i + 2 C 1 2 ) Cu C

2a(Cn ~ CM) _ b(Cu + 2C 1 2 )

Cu " c

(5.2)

where EG is the unstrained gap, the Qj are elastic coefficients, a is the hydrostatic deformation potential, b is a shear deformation potential, and ε is the strain. Since the elastic constants and deformation potentials in G a A s are known, the energies measured from the P L data yield strain values. The authors found ε = 1.83 x 10~ 3 to 1.89 x 10~ 3 for as-grown samples, and ε = 2.5 x 1 0 - 3 for annealed samples. They explained the increase in the annealed samples as coming from the strain due to the difference in expansion coefficient, expressed by:

ε = ( « G a A s - <*si) Δ Γ (5.3)

where Δ Γ is the difference between the growth temperature, and the temperature at which P L is measured. This gives ε = 3.6 x 1 0 " 3 o C _ 1 for as-grown samples, with higher values for annealed films. The authors explain the quantitative discrepancy between the calculation and the data as due to the use of a bulk value of a G a A s for a film.

Freundlich et al. (1988) combined P L and P L E spectroscopy, Raman, and X-ray methods to examine GaAs grown by metal organic vapor phase epitaxy ( M O V P E ) on (100) silicon substrates. Each method gave a quantitative measurement of stress which contributed to a complete picture. They observed P L excitonic features related to the heavy-hole and light-hole bands. From the theory for the band shifts under (100) biaxial stress, and known values of the elastic constants, they deduce the energy shifts with stress X:

- 4.3 m e V

dX - 4.3

kbar

- 9.7 m e V

dX - 9.7

kbar

(5.4)

From these relations and the positions of the P L peaks, they found a biaxial tensile stress of 1.8 ± 0.3 kbar for 2 μπι thick unannealed GaAs/Si at room temperature. This


value applied near the epilayer surface, decreasing as the interface was approached. When the sample was cooled to 1.8K, the value of 1.8 kbar increased to 3.1 kbar, indicating increased thermal stress. After annealing at 850°C, the value at 1.8 Κ increased to 3.7 kbar. In their Raman experiments, Freundlich et al. measured a down-shift of 0.7 ± 0.1 c m - 1 in the L O mode frequency, which they also ascribed to stress in the epilayer. I give details in Chapter 6.

Nemanich et al. (1988) also used both Raman and P L methods to examine strain in molecular beam epitaxy ( M B E ) grown GaAs/Si. Their material was grown with a GaAs buffer layer, deposited at a lower temperature than is best for GaAs growth. The shutter in the M B E chamber was slowly moved over the sample during growth, producing a film of graded thickness. This made it possible to relate Raman and P L data to film thickness. Both showed that tensile strain existed at all thicknesses. The P L data were taken at several locations on the wedged surface. They displayed band-to-acceptor and exciton transitions at 1.475 and 1.495 e V , which were shifted by 27 meV, corresponding to a stress of approximately —2.5 kbar. Comparison with their Raman data, which showed a strain-related shift in the L O frequency, led to the conclusion that the observed P L signals came from the overlayer region, not from the buffer or intermediate layers. This indicates that these layers may contain high densities of defects, acting as centers for nonradiative recombination. The combination of techniques also permitted these authors to draw conclusions about the origins of the stress.

Zemon et al. (1988) carried out P L and resonant P L E spectroscopy in the excitonic region, for GaAs films 3.5-7.5 μπι thick grown on silicon by organometallic vapor phase epitaxy ( O M V P E ) . These samples had carrier densities below 10 1 5 cm~ 3 , and low concentrations of shallow acceptors, as seen in the P L spectra. A dye laser made it possible to excite the system where it absorbed most; for instance, at 832.94nm, the wavelength of peak absorption for the H H exciton. The authors observed new features they ascribed to free excitons and to donors. Although this study did not explicitly evaluate strain, the authors concluded that the sample quality was excellent, because the spectral features displayed narrow line widths of 1.5-3 meV.

Lee et al. (1988) laid groundwork to characterize actual device configurations when they examined patterned epitaxial GaAs films on silicon substrates. T o motivate their work, they noted that the processing temperatures for silicon exceed what GaAs can tolerate; hence any hybrid technology would require that the silicon portion be fabricated before the GaAs portion. This would mean, for instance, that silicon devices are made first in certain portions of the substrate, and covered by an insulating layer such as S i 3 N 4 or S i 0 2 . This would be followed by the growth of epitaxial GaAs , with fabrication of GaAs devices as the last step.

These authors considered two ways to make patterned GaAs/Si films. Both used n +

silicon (100) substrates tilted 3.5° towards the (011) axis. In the first method, selective-area M B E , 80 nm of S i 3 N 4 was deposited on the substrate. Then conventional photolithography, followed by reactive ion etching of the S i 3 N 4 , was used to form stripes of exposed silicon 2 mm long and 10-100 μπι wide, along the (011) direction. GaAs was deposited by M B E over this patterned section and over an adjoining unpatterned reference area. T w o GaAs films, 1.5 and 3 μπι thick, were made. Microscopic examination showed no apparent difference between the G a A s on the patterned and unpatterned regions.

The second method used post-growth etching. After M B E growth of a 3 μπι thick

Band emission 69

12 350 12 230 12 110 11 990 11 870 11 750

Frequency (cm - 1 )

Fig. 5 .4 Photoluminescence from GaAs films 1.5 and 3 μηι thick, made on silicon by selective-area MBE. Stress in the films splits the single peak seen from a standard GaAs sample into two peaks at lower energies. Theoretical curves calculated for exciton behavior under stress are also shown. (After Lee et al (1988).)

G a A s film, conventional lithography was used to pattern the film. The pattern included squares 4 μπι on a side, and stripes 4-100 μπι wide by 1mm long. The film was etched for 8 min in 1: 8: 8 H2SO4/H2O2/H2O to form the stripes and squares. A control region was left unetched.

Figure 5.4 shows P L spectra from the reference area in the selective-area M B E samples, compared to the spectrum from a standard G a A s sample. The M B E spectra show two peaks, at longer wavelengths and greater halfwidths than the single peak in unstressed G a A s . The dual peaks arise from the split valence band in G a A s . Their shifted wavelengths, relative to the unstressed peak, give stresses of 3.2 ± 0.4 and 3.7 ± 0 . 5 kbar for the 1.5 and 3 μηι films respectively. The figure also shows calculated spectra based on a stress analysis of the exciton behavior, which agree with the data. The authors found that the P L spectra from the 10 μπι and the 100 μπι stripes were similar to those from the reference area, but those from the thinner stripes showed greater linewidths in peak Β (9.5 versus 7.9 m e V ) , perhaps representing non-uniform stress in the narrower structure.

Figure 5.5 displays P L spectra from stripes made by post-growth etching. They show distinct differences compared to a GaAs /GaAs sample, and as a function of stripe width, with the splitting seen in Fig. 5.4 again apparent in some cases. These peaks are broader than those from similar stripes that had undergone rapid thermal annealing ( R T A ) , showing that R T A improves the crystallinity. The shift in the P L peaks were also examined versus stripe width. These data gave a tensile stress of about —2.6 kbar in the stripes. Polarization analysis of the P L data showed that the


8097

Γ

Wavelength (A) 8217 8340

} 4 μηι

I . χ4μΓΠ Jfk square

CO ω

8097 8298 8511, Wavelength (A)

12 350 12 170 11 990

Wavenumber (cm"'

11 810

Fig. 5.5 Photoluminescence from GaAs stripes of varying widths (indicated), and a 4 μπι x 4 μπι GaAs square, made by post-growth etching on silicon. Compared with a spectrum from GaAs/GaAs, the GaAs/Si data show stress-induced peak broadening, shifting, and splitting. (After Lee et al (1988).)

stress in narrow stripes is uniaxial along the stripes. These results show the power of P L spectroscopy to determine both magnitude and direction of stress.

Wilson (1989) also has examined strain in GaAs/Si. She points out that when one semiconductor is grown on another, if their lattice constants are not too different, a thin epitaxial layer often grows commensurately with the substrate; the lattice misfit is completely compensated by strain. But when the film exceeds a certain critical thick-

Band emission 71

ness, the energetics are such that it is preferable to form misfit dislocations to relieve some or all of the strain. Even if a thick layer grows so that it does not match the substrate, with more or less its normal relaxed lattice constant, strain may still enter because of thermal effects. A s the film-substrate system cools from the relatively high growth temperature, any difference in the thermal contraction coefficients of the two produces strain, as expressed by equation (5 .3) .

Wilson claims that it is this thermal difference which is often observed in GaAs/Si grown by metallo-organic chemical vapor deposition ( M O C V D ) . The biaxial stress due to cooldown from growth temperatures of 650-700 °C, to her P L measurement temperature of 6 K , is about 3.1 kbar. This would produce a light hole-heavy hole split in the exciton peak, with shifts of —4.7 and —10.3 m e V / k b a r - 1 respectively, giving a ratio of 0.46 for the two shifts. She notes, however, that these are not necessarily the splits observed in P L spectra, which may reflect inhomogeneous strain relaxation, perhaps related to crack formation. P L E data give the true split doublets, she comments, and her analysis of P L E results for the split peaks gives the predicted ratio of 0.46.

Nishino (1989) has developed a highly sensitive P L method to characterize interface stress in heterostructures, which he applies to I n j ^ G a ^ A s ^ P ^ / G a A s , Alj .Gai_j .As/ G a A s , and ZnS^Set-^/GaAs. He points out the importance of interface stress in device performance, and notes that X-ray diffraction, Raman scattering, and photo-elasticity have been the main methods of characterization. His technique relies on the fact that chromium impurity atoms in G a A s have well-known P L lines. The presence of uniaxial stress reduces the site symmetry of the impurity atoms, which shifts and splits the chromium related P L peaks. Since GaAsiCr is a high-resistivity material, it provides a good substrate for other materials, which can also monitor interface stress.

Figure 5.6 shows the P L spectrum of a GaAs:Cr wafer at 4.2 K , with extremely sharp lines near 0.839 e V . Several different models for their origin have been proposed. Nishino suggests that the lines come from a chromium atom on a gallium site, and an arsenic vacancy ( V A s ) at the nearest neighbor. Regardless of the origin of the features, however, the dependence of one of them, the peak at 0.8396 e V , has been

CO c CD

U . I I I I I I

0.76 0.78 0.80 0.82 0.84 0.86 Photon energy (eV)

Fig. 5.6 Photoluminescence spectrum of GaAs:Cr at 4.2 K, showing the extremely narrow, stress-sensitive lines near 0.839 c m - 1 from a chromium complex. (After Nishino (1989).)

http://Alj.Gai_j.As/


measured versus uniaxial stress in different directions. The results shown in Fig. 5.7 make it possible to determine the direction of stress as well as its size.

Nishino examined epitaxial layers of I n 1 _ x G a A : A s > ; P 1 _ > „ A l ^ G a ^ A s , and ZnS^Sex-* grown on GaAs.Cr substrates. The layers were several micrometers thick. The 514.5 nm radiation from an A r + laser penetrated to the top surface of the substrate, returning P L spectra from within about 1 μιη of the interface. A l l measurements were made with the samples immersed in liquid helium at 4.2 K . The P L signal was detected by a germanium p - i - n photodiode cooled to liquid nitrogen temperature, and analyzed by a grating monochromator.

Figure 5.8 shows spectra for I n ^ G a - r A s ^ ^ films (x = 0.51 to 0.53; y = 0.03) grown on (100) or (111) GaAs:Cr. The films, typically 3 μπι thick, were made by liquid phase epitaxy. For the (100) sample, the chromium-related peak moves relative to that from a GaAs:Cr wafer as a function of the lattice mismatch Δα/α. The shift reflects biaxial stress. The shift of 0.04 m e V seen at Δα/a = 0 comes from lattice mismatch caused by the difference in thermal expansion coefficients between Ini-^GaÂs^Pi-^ and GaAs . For the (111) sample, the main line splits, becoming three distinct peaks at Δα/α = 0.28%. This is due to compressive [111] uniaxial stress, as shown in Fig. 5.7. From these shifts and splittings, Nishini finds the size of the interface stress, which is plotted versus lattice mismatch in Fig. 5.9. For comparison,

Band emission 73

_ l I I ι I I 1 1 1 ' — 0.838 0.839 0.840 0.841 0.838 0.839 0.840 0.841

Photon energy (eV)

Fig. 5.8 Photoluminescence from Inx-^GaÂSyP^/GaAsiCr heterostructures, with χ = 0.51 to 0.53, and y = 0.03. (a) (100) orientation; (b) (111) orientation. Each trace is labeled with the relative lattice mismatch Δα/α, which changes as a function of x. The spectrum of a GaAs:Cr wafer is shown for comparison. The 0.8396eV chromium line moves as a function of mismatch, and splits in the right panel. (After Nishino (1989).)

Lattice mismatch Δ

Fig. 5.9 Interface stress for In^Ga^PÂs^ /GaAsiCr heterostructures at 4.2 Κ from PL data like that in Fig. 5.8, compared with the stress calculated from lattice mismatch, which is twice as large. See text. (After Nishino (1989).)


the calculated stress due solely from lattice mismatch at 4.3 Κ is also shown. The measured value is half as large. This decrease Nishino interprets as due to the relaxation of stress after the formation of misfit dislocations. Nishino reports similar results and analyses for the two other heterostructures he studied, A l x G a ! _ x A s / G a A s : C r and Z n S ^ S e ^ / G a A s i C r . H e concludes that his method can detect interface stress as small as 0.5 MPa, which makes it more sensitive than other techniques.

5.4 Impurity emission

Impurity analysis is probably the most important use of P L spectroscopy, with a long history. I have already shown, in Fig. 3.10, an early spectrum which clearly displays how the addition of arsenic to silicon dramatically changes the P L spectrum. Nakayama (1980) started early to develop means to obtain values for dopant concentrations and compensation ratios from P L spectra. Since these early efforts, an enormous amount of work has been reported in P L studies of silicon.

Lightowlers (1990) shows how specific impurities in silicon clearly appear in P L spectra, and how P L data can measure their concentrations. The data in Figs 5.10 to 5.13 were taken with the samples immersed in liquid helium at 4 .2K. They were excited by the 514 nm line from an A r + laser with a power of about 250 m W . Each spectrum was obtained in 10 min using Fourier transform photoluminescence (FT-

FE(TO)

I ι ι . ι I 1090 1110 1130 1150

Photon energy (meV)

Fig. 5 .10 Photoluminescence from high-resistivity (>20kH-cm) near-intrinsic silicon. Fingerprint features from impurity boron, phosphorus, aluminum, and arsenic are marked. Their concentrations, in units of 10 1 2cm~ 3, are 1.36, 1.69, 0.61, and 0.14 respectively. Free-exciton lines (FE) are visible. TO, LO and T A phonons participate as indicated. NP (no phonon) peaks come from bound excitons. Calibrations are discussed in the text and shown in Fig. 5.14. (After Lightowlers (1990).)

Impurity emission 75

C CD

Έ . _ ω c ο ο n

ω ο c CD ο

Ρ(ΤΟ)α,

<*3

UAXJ

No-phonon Ρ(ΝΡ)

χθ.5

s \ J

α 1

ct^LO) 1146 1148 1150 1152

FE(TO)

/ FE(LO) 28(Γ,) P(TA).

JL

Ρ ( Ν Ρ ) α ι

a , a 3 a 2

1090 1110 1130

Photon energy (meV)

1150

Fig. 5.11 Photoluminescence spectrum of silicon doped with 3 x 10 1 4cm" 3 phosphorus. At this impurity concentration, excitons become bound to the impurities. Lines marked a and β arise from bound single and multiple excitons. (After Lightowlers (1990).)

ο -_ ΩΟ)

Β3ι

B 1(TO)

FE(LO)

No-phonon x10 P(NP) t t l

As(NP) I

B 1(NP)

Β 3 Β 2 I

1146 1148 1150 1152

B(TA)

Β 3 Β 2 β 1 τ _ / τ λ ^ No-phonon ι FE(TA))

JL

I 1 1 1 ι I 1090 1110 1130 1150

Photon energy (meV)

Fig. 5 .12 Photoluminescence spectrum of silicon doped with 1.3 x 10 1 3 cm - 3 boron. Structure due to phosphorus at 1.8 x 10 1 2 cm - 3 , and arsenic at 3 x 1 0 u c m - 3 is also apparent. (After Lightowlers (1990).)


1090 1110

Photon energy (meV)

1130 1150

Fig. 5 .13 Photoluminescence spectrum of silicon doped with 2.7 x 101 4cm 3 aluminum, with evidence of contamination from boron near the TO phonon-assisted peaks. (After Lightowlers (1990).)

P L ) . Figure 5.10 is the spectrum of silicon with resistivity >20kO-cm, that is, near-intrinsic material. In such pure materials, the optically generated excitons remain free, and decay to give free exciton luminescence. A s silicon has an indirect band gap, these transitions require a phonon to conserve momentum. These free exciton peaks are strong in Fig. 5.10, where each is labelled FE with the appropriate accompanying phonon ( T O , T A , and L O ) . Also appearing are sharp no-phonon ( N P ) lines coming from extremely low concentrations of boron, aluminum, phosphorus, and arsenic, of the order of 10 1 2 cm~ 3 . I will discuss how these concentrations are derived.

A t higher impurity concentrations of 1 0 1 4 - 1 0 1 5 c m - 3 or more, the excitons no longer remain free but are captured, so that bound exciton lines now appear in the spectrum, at the photon energies given by equation (3.14) if a phonon is involved, or by equation (3.13) if a phonon is not. Figures 5.11 to 5.13 show spectra from single and multiple bound exciton complexes. In Fig. 5.11, the sample is doped with 3 x 1 0 1 4 c m - 3 phosphorus, with negligible densities of boron, aluminum, and arsenic. Phonon-assisted lines appear, as do the lines marked α and β from bound single and multiple excitons. The sample in Fig. 5.12 is doped with 1.3 x 1 0 1 3 c m - 3 boron atoms. In addition to features from these, other structures indicate the presence of phosphorus at 1.8 x 1 0 1 2 c m - 3 , and arsenic at 3 x 1 0 n c m ~ 3 . Figure 5.13 comes from a sample doped with aluminum at 2.7 x 10 1 4 cm~ 3 . Contamination from boron is evident in the region near the TO-phonon-assisted aluminum peak marked A l ( T O ) .

The determination of dopant and contaminant densities from such P L spectra requires care, especially in comparing calibrations from one laboratory to another. Differences in exciting laser power, spectral resolution and response of the P L system,


and actual sample temperature—which may differ from the ambient cryogenic temperature—can cause significant deviations. T o avoid these, Colley and Lightowlers (1987) and Lightowlers (1990) use internal calibrations. T o calibrate the amount of phosphorus, for instance, they ratio the N P peak to the FE peak seen in Fig. 5.11. These authors considered whether ratios of peak height, or peak area (integrated peak intensity) were more reliable. The latter is more satisfactory because it is less dependent on spectral resolution and broadening effects. The peaks are compared at an excitation intensity which saturates the multi-exciton luminescence but does not heat the luminescent region above 4.2 K , when the sample is immersed in liquid helium boiling at atmospheric pressure.

Calibration curves for phosphorus, boron, and aluminum in silicon are given in Fig. 5.14, as impurity concentration versus the ratio of the N P peak area to the FE peak height. The relations are very nearly linear, and the range of densities 1 0 1 2 - 1 0 1 5 c m - 3

is useful for silicon. More recent work is expected to extend the lower limit, and to extend these methods to thin epitaxial films. Measurements at 20 K , where the bound excitons are ionized, extends the upper limit to about 10 1 7 cm~ 3 (Lightowlers, 1990).

In addition to broad coverage of different impurities in silicon, such as I have just

10 1 !

101*

1 0 1 3

10 1

I I M I III 1 I I I M i l l » 1 I I I I 111 1 I I I M i l l 1 1 I I I M 11 1 I I I I I III

10" 6 10" 5 10" 4 io- 3

NPIine area/FE peak height (eV)

io- ;

Fig. 5 .14 Calibration curves to convert PL intensity into concentration of boron, aluminum, or phosphorus in silicon. This plot uses an internal calibration method. The area of the no-phonon (NP) peak for the particular impurity is ratioed against the height of the FE(TO) peak in the same spectrum (Fig. 5.11). (After Lightowlers (1990).)


illustrated, P L has also been used for in-depth studies of specific important impurities such as carbon and oxygen, and associated complexes or defects. Carbon in silicon, for instance, has been examined for over a decade. Lightowlers et al. (1984) probed radiation-damaged lithium doped silicon with and without additional carbon. Without carbon, the P L spectrum consists of three no-phonon lines—called the ' Q series'— near 1.045 e V . This has been identified as coming from a complex of four lithium atoms which is trigonally distorted, suggested to substitute for a single silicon atom. In samples with added carbon, a second set of no-phonon lines—the S series—appears at 1.082 e V . A full analysis, which also required measurements of P L decay times (a topic I consider in Chapter 8) confirms the identification of the Q and S lines with exciton decay at defect complexes. Both the Q and S centers are taken as complexes with four lithium atoms, with a strong suggestion that the S system includes a nearest-neighbor carbon atom. The analysis indicates that < 2 % of the lithium atoms in the sample are involved in these complexes. Results such as this show the power of P L methods to examine intricate impurity and defect arrangements.

Wagner et al. (1984) used P L and P L E to examine the so-called 4 C line', a no-phonon P L transition observed at 0.79 e V in oxygen-rich n- and p-silicon irradiated by 2 M e V electrons or neutrons. A KC1:T1 color-center laser, tunable from 1426 to 1579 nm, provided resonant excitation at the C line with typical powers of 50 m W . Features on the C-line complex were identified as arising from local vibronic modes, not electronic states as previously surmized. Their energies of 65.5, 72.5, 138.1, and 145.3 m e V are near known vibrational energies for interstitial oxygen (S i -O—Si) , carbon in substitutional sites, or carbon-oxygen complexes. These have also been seen in infrared absorption, as I discuss in Chapter 7.

Weman et al. (1985) carried out other P L work on carbon and oxygen in Czoch-ralski-grown silicon, where they form impurity complexes. These are highly dependent on the heat treatment of the material and the devices which it comprises. In this work, the authors followed the P L spectra of p-type boron doped Czochralski-grown silicon through a series of annealing cycles, a good demonstration of the power of P L (and optical methods in general) to quickly return information to the material maker and the device designer. In an example of the interaction of different optical methods, these researchers also derived the concentrations of carbon and oxygen from infrared data, using a calibration method which is described in Chapter 7.

^Tiewalt et al. (1985) also used P L to monitor silicon during processing. Studies of transport behavior had shown that acceptor impurities in silicon could be neutralized if the material were treated by a low temperature (100 °C) plasma discharge containing atomic hydrogen. The observation had been subject to some controversy. Did the neutralization occur in plasmas containing pure atomic hydrogen, or was it necessary to also include a simultaneous or earlier exposure to atomic oxygen? In the first case, the hydrogen simply neutralizes the acceptor; in the second, the acceptor joins with neutral oxygen to form a shallow acceptor complex, which only then is neutralized by the hydrogen.

Thewalt et al. used P L to examine excitons bound to donors and acceptors in ultrahigh purity silicon made by the vacuum float zone method. Samples were implanted with the shallow acceptor boron, or the deep acceptors indium and thallium, at a dose of 1 x 10 1 1 c m - 2 at 100 keV. As-implanted material was compared with that which had been subjected to a plasma of atomic hydrogen. Plasma treatment, it was found, reduced the luminescent lines from implanted acceptors by factors corres-


ponding to reductions of 10-50 in their concentrations. T o eliminate the possibility that implantation damage simply created a nonluminescent dead zone, a set of samples was made which included both the acceptors indium and thallium, and the donor arsenic, in the same region. After plasma treatment, P L spectra showed a great reduction in the indium and thallium lines, whereas the arsenic line continued as a strong feature. Hence the data confirmed that low-temperature exposure to a hydrogen plasma neutralized boron, indium, and thallium in silicon, without affecting the donor arsenic.

Photoluminescence has also been used to analyze impurities in silicon grown by molecular beam epitaxy. Robbins et al. (1985) claimed to present the first optical evidence that carbon is a persistent impurity in MBE-grown silicon. In examining epitaxial layers, one important consideration was to differentiate P L from the layers from that arising in the substrate. This was accomplished by using p-type silicon substrates so that η-type impurities could be definitely associated with the layer. Also some layers were made thicker than 10 μιη, far exceeding the penetration depth in silicon at a wavelength of 488 nm, which is less than 1 μπι. Although this eliminated direct impact of the absorbed light on the substrate, diffused carriers and excitons could conceivably reach the substrate. T o monitor this possibility, spectra were compared for layers of different thickness. Features from the substrate became progressively weaker with increasing layer thickness.

The data, excited by 488 nm radiation from an A r + laser with the sample held at 4.2 K , showed free exciton and electron-hole drop features, indicative of good film quality. Some peaks came from bound excitons, which identified electrically active shallow impurities such as phosphorus. A transitioh at 0.97 e V implicated carbon on silicon sites as part of a radiative complex. Other impurities were identified, and their dependence on growth details was tracked. When graphite retaining rings were used to support the substrate during M B E deposition, the dominant impurity was phosphorus. With tantalum rings, only features connected to the boron acceptor appeared.

In a broad-based survey, Kaminski et al. (1987) examined excitons localized at point defects in silicon, to explore the structure and formation kinetics of the defects. The samples were grown by the floating-zone method and were either pure, or doped with phosphorus at a concentration of 2 x 1 0 l 4 c m - 3 . Defects were induced by thermal neutron irradiation, after which the samples were annealed. This extensive investigation used an A r + laser, with the samples held at 2-35 K . Measurements in α magnetic field, and under uniaxial compression, helped to analyze the data. The authors conclude that the P L approach provides a means to analyze electrically inactive centers which are not otherwise easy to examine.

A thorough study of a family of impurities comes from Conzelman (1987), who treated the 3d transition metals in silicon. These can be serious contaminants because they diffuse readily, and affect the electronic properties of silicon even at extremely low densities. Defects associated with transition metals produce deep levels within the gap and can act as highly efficient centers for recombination. Considerable work has gone into identifying these impurities, using electron paramagnetic resonance, Hall measurements, and deep level transient spectroscopy ( D L T S ) . These, along with luminescence studies, have yielded the level positions for many of the 3d impurities, but their excited states and their role in carrier recombination are not well understood.

Conzelman examined n- and p-type silicon crystals grown by float-zone refining,


with some samples also made by the Czochralski method. Titanium, vanadium, chromium, manganese, iron, cobalt, or nickel were introduced into the samples by evaporation, by mechanical contact, or by ion implantation, followed by heat treatment to promote diffusion. Photoluminescence measurements were made with excitation by an A r + or K r + laser, and with the samples held at 2-300 K .

Conzelman notes the advantages of P L over electrical measurements, since the energy resolution is excellent and different defects produce very different spectra; but since the P L intensity depends on quantum efficiency, it can be difficult to measure defect concentrations. With a high quantum efficiency, even concentrations below 1 0 1 2 c m - 3 can be readily seen, whereas for defects with important nonradiative transitions, even 1 0 1 7 c m - 3 may not be very visible. With these factors in mind, Conzelman shows in detail the process of identifying and analyzing chromium diffused into boron doped silicon, starting with a P L spectrum showing the no-phonon and phonon-assisted lines arising from chromium-boron pairs.

Thewalt et al. (1990) have explicitly demonstrated how the increased sensitivity of F T - P L enhances impurity characterization. In one example, they note that although the infrared absorption of copper in germanium has been well studied, it has not yielded evidence of an exciton bound to this triple acceptor. Thewalt et al. searched for the exciton using conventional dispersive P L , but did not observe it. Using F T - P L , however, they were able to see the excitonic peak in Ge:Cu, with a signal-to-noise ratio of 3. After sufficient averaging of data, they obtained the excellent spectrum shown in Fig. 5.15.

τ 1 Γ

J I

695 705' 715

Energy (meV)

Fig. 5 .15 Infrared photoluminescence from germanium doped with copper, obtained by FT-PL methods. Lines from free excitons, and from excitons bound to copper, are seen. Subscripts denote the phonons involved. Copper is a triple acceptor in silicon. The peaks associated with the C u N P (no-phonon) transition give an ionization energy of 43.3 meV. (After Thewalt et al. (1990).)


- J I L_

1140 1145 1150

Energy (meV)

Fig. 5 .16 Fourier transform photoluminescence spectrum of silicon doped with phosphorus at the medium resolution of 0.5cm - 1 (0.06 meV), showing the a series of lines from a bound exciton (a1) and from bound multiexciton complexes (a -a5). No fine structure is visible. The same series appears in the spectrum displayed in Fig. 5.11, and at higher resolution in Fig. 5.17. (After Thewalt et al (1990).)

This application required only medium resolution. A second example was the measurement by Thewalt et al of high-resolution spectra from Si:P. These show a line ( a 1 ) due to the bound exciton ( B E ) , and other lines from bound multiexciton complexes ( B M E C ) with up to five electron-hole pairs ( α 2 - α 5 ) . (The α*-α 4 lines also appear in Fig. 5.11). Figure 5.16 shows the non-phonon P L spectrum of these five lines, measured by conventional methods, at a resolution of 0.5 c m - 1 (0.06 m e V ) . The expected fine structure in the α 2 - α 5 lines is not visible. However , when data are taken with an interferometric instrument at the ultra-high resolution of 0.02 c m - 1 (2.4 μβν ) , fine structure appears as shown in Fig. 5.17. The 5.7 μ&Υ full width at half maximum ( F W H M ) of the a 1 transition is, these authors claim, the narrowest nonresonant bound exciton line ever reported. Thewalt et al further display the value of F T - P L by presenting a high-resolution analysis of photoluminescence from G a A s in a magnetic field.

Photoluminescence has proven equally valuable in semiconductor alloys and compound materials. In addition to the fundamental gap measurements in Si i -^Ge* discussed earlier, for instance, Weber and Alonso (1989) used P L to compare the quality of samples made by different methods. Polycrystalline large-grained bulk samples, made by zone-leveling, were compared with films 0.3-20 μπι thick, grown by liquid-phase epitaxy ( L P E ) and vapor-phase epitaxy ( V P E ) . The bulk samples gave P L spectra of higher quality than either the L P E or VPE-grown specimens. The spectral peaks from the bulk material were typically 3-5 times narrower than those from the L P E material. The latter also showed weak P L features near the band gap,


! _ ω

ο ο

1143.7 1141.7

Energy (meV)

Fig. 5 .17 Fourier transform photoluminescence spectrum of phosphorus-doped silicon, at the very high resolution of 0.02cm - 1 (2.5 μεν ) , showing the al-a4 lines from Fig. 5.16. Fine structure is seen in the lines from bound multiexciton complexes. The a1 line has a FWHM of 5.7 μεΥ. See text. (After Thewalt et al. (1990).)

which came from dislocation effects. Although the V P E samples also showed narrow peaks, these were always accompanied by very strong broad structure of unknown origin. The work by Hansson et al. (1990), which I discussed earlier, also points to the quality of as-grown S i i _ x G e x material, since it reports extremely narrow P L lines for this alloy.

Skromme et al. (1985) carried out an extensive set of impurity measurements on MBE-grown GaAs doped with silicon, which also illustrated the combination of optical methods. They used P L at low temperature, infrared photothermal ionization spectroscopy (described in Chapter 7 ) , and other techniques. Their samples were high-purity GaAs wafers lightly doped with silicon (carrier density at 77 Κ was 2 x 10 1 4

to 8 x 1 0 1 4 c m ~ 3 ) . Photoluminescence data were obtained at sample temperatures of 1.7-21 K , using either liquid or gaseous helium with the samples mounted in a strain-free configuration. The excitation came at 514.5 nm from an A r + laser, and the P L signal was detected by a photomultiplier tube.

Photoluminescence spectra like that shown in Fig. 5.18 established the quality of the high-purity GaAs , because they displayed lines due to free excitons, as well as sharp bound-exciton peaks. From the F W H M of the bound exciton peaks, the authors concluded that the quality of the MBE-grown material approached that of the best

• 1150.0. _ 1146.4


Energy (eVi) 1.5150 1.5125 1.5100

(D°,X),

x10 !

(A°,X)

817 818 819 820 821 822

Wavelength (nm)

Fig. 5 .18 Photoluminescence spectrum of MBE-grown GaAs at 1.7 K, excited at 514 nm ( F L = 12mWcm~ 2 l The appearance of free exciton (FE) lines and sharp neutral donor-bound exciton lines (D ,X) indicates the quality of the material. The FWHM of the (D°,X)„ = ι peak is 0.15meV, compared with 0.11 meV in the best available LPE and VPE materials. Other lines arise from neutral acceptor-bound exciton emissions ( A ° , X ) and other donor interactions. (After Skromme et al. (1985).)

liquid phase and vapor phase ( L P E and V P E ) material. Their photothermal data, which I discuss under infrared characterization, showed that sulfur donors appeared, introduced from the arsenic source. The P L data confirmed that the dominant residual acceptor was carbon, and showed that the acceptor concentration increased with arsenic flux, indicating that at least part of the carbon also comes from the arsenic source. The authors recommend growth at the lowest possible arsenic levels. Other P L features between 1.466 and 1.482eV, and 1.504 and 1.512eV (not shown) are related to defects or complexes which were not identified, but whose concentration was very low.

Further work on GaAs came from van de Ven et al. (1986), who examined residual impurities in material grown from trimethyl gallium ( T M G ) and A s H 3 by M O C V D . Their emphasis was on correlating many growth parameters—III-V ratio, temperature, axial location of the growth area in the reactor, gas sources, materials for the substrate and the susceptor, carrier gas, crystallographic orientation of the substrate, and physical misorientation of the substrate. The P L data were combined with van der Pauw-type Hall measurements.

The P L data were taken with the 514 nm line of an A r + laser, at sample tempera-


tures of 2-300K, and excitation intensities of 1 0 _ 3 - 1 0 W c m - 2 . The data were corrected for the sensitivity of the detector, a photomultiplier with an SI response. The spectra were very rich, showing evidence of free and bound excitons, neutral and ionized shallow donors, neutral shallow acceptors, deep levels in the gap, and native defect complexes. Identification of the lines was aided by measuring spectra versus temperature, and versus excitation intensity; for instance, conduction band-acceptor ( e - A ° ) transitions dominate donor-acceptor ( D ° - A ° ) lines at higher temperatures and higher intensities. One result of these detailed studies was the observation that the main acceptor impurities were zinc, silicon, and carbon. The zinc and carbon were shown to originate from the T M G gas source, and the silicon from quartz components

Z r WC A s

Energy (eV)

Fig. 5 .19 Photoluminescence from GaAs grown by MOCVD at 700 °C, for different V/III ratios in the input gas mixture. Energy ranges displaying the most significant structure are shown. Transitions involving copper, zinc, carbon, free excitons (FX) , and excitons bound to donors and acceptors (DX, A X ) appear. (After van de Ven et al. (1986).)

Fig. 5 .20 Photoluminescence from GaAs grown by MOCVD with a V/III ratio of 10 in the input gas mixture, for different growth temperatures. Transitions like those in Fig. 5.19 appear, as well as peaks related to silicon, and to a deep center ( D X ) . (After van de Ven et al. (1986).)

Fig. 5.21 Photoluminescence from GaAs grown by MOCVD at a temperature of 700 °C and a V/III ratio of 10 in the input gas mixture, for different axial positions in the growth reactor. Transitions like those seen in Figs 5.19 and 5.20 appear. (After van de Ven et al. (1986).)



in the growth cell. The P L spectra showed clear changes as a function of V / I I I ratios (Fig. 5.19), growth temperature (Fig. 5.20), and axial position in the growth reactor (Fig. 5.21), which proved an important parameter for layer quality.

In another application to G a A s , Koteles et al. (1987) used P L to compare the purity of MBE-grown material with that of VPE-grown G a A s , considered to be the highest purity material. Their M B E samples were grown on (100) undoped semi-insulating or silicon-doped GaAs . For the P L measurements, the samples were mounted strain-free in a liquid helium cryostat and cooled by exchange gas. They were excited at powers of l O O m W c m - 2 or less from a H e N e or a dye laser.

Figure 5.22 shows P L spectra from the best MBE-grown sample, from a typical M B E sample, and from a VPE-grown sample. The V P E sample, with excellent characteristics ( μ 7 7 Κ = 210000cm 2 ( V - s ) _ 1 , Ndonor = 5 x l 0 1 3 c m ~ 3 , N a c c e p t o r = 2 x 1 0 1 3 c m - 3 ) , nevertheless produces a spectrum dominated by impurity peaks. The peaks below 1.5 e V come from free-to-bound transitions involving acceptors and donor-acceptor pairs. The large peaks from 1.51 to 1.52 e V come either from excitons bound to neutral donors ( D ° , X ) or free-to-bound transitions involving neutral donors (h, D ° ) . The free exciton feature ( X ) is insignificant by comparison, but becomes important for the typical M B E sample. However , other structure is still apparent, from excitons bound to neutral acceptors ( A ° , X ) . In the best M B E sample, the free exciton peak completely dominates all other structure. Emission related to free-to-bound acceptor transitions is several orders of magnitude weaker, and other exciton-related peaks are either not apparent or extremely weak. These identifications are clearer in the high-resolution data in Fig. 5.23. T o confirm that the single narrow peak in the best sample came from free excitons, the sample was also examined by magneto-photoluminescence.

Fouquet et al. (1989) illustrated the usefulness of near-infrared P L to examine defects in GaAs and InP. Their motivation came from the use of these materials in devices. Although GaAs and InP have been extensively studied as substrates, their properties as high-resistivity epitaxial device layers have not been much investigated. In one application, GaAs buffer layers between an active layer and a substrate prevent backgating and sidegating—the undesirable influence on a device from the potential applied to an adjacent device—in metal oxide semiconductor field-effect transistors ( M E S F E T S ) . The GaAs layer is grown at low temperature (^300 ° C ) . Such layers are reported to be optically inactive, not degrading the optical or electrical performance of active layers grown atop them. In a second application, iron-doped InP layers provide lateral confinement in high speed GalnAsP/InP lasers, to give the necessary photon densities. Several confinement schemes have been tried. High-resistivity iron-doped InP effectively prevents current from leaving the active lasing region and has given excellent performance in certain laser geometries. However, the way in which iron enters InP grown by organometallic vapor phase epitaxy ( O M V P E ) is not completely understood.

The only P L work reported for these buffer layers had been near their gaps, at 0.87 and 0.98 μπι for GaAs and InP respectively. Fouquet et al. examined the extended wavelength region to 1.6 μπι, the first investigation beyond 1 μιη, to understand why the buffer layer reduces backgating. T o study GaAs , they used the 632.8 nm H e N e laser line at an intensity of 3 W c m - 2 . The samples were immersed in liquid helium at 3.9-4.1 K . The authors concluded from the P L data that the GaAs buffer material contains gallium vacancies, and surmise that this is responsible for its nonconducting


Ε ο ο

(a) VPE

Exciton region 1

x10

x1

k . τ ' 1 • — ι ι X

(b)MBE-117

— ι 1 1 r

(c) MBE 11-19

1.48

Energy (eV)

Fig. 5 .22 Photoluminescence from three nominally undoped epitaxial GaAs layers at 5 K. (a) VPE-grown sample; (b) typical MBE-grown sample on semi-insulating substrate; (c) highest-quality MBE sample. The free exciton peak is labeled X. (After Koteles et al. (1987).)

nature. Their study also examined the quality of epitaxial G a A s on silicon. They concluded that the formation of gallium vacancies contributes to the low quality of such films, and feel it probable that the vacancies come because the G a A s growth is initiated at low temperatures to prevent the formation of islands.

For the InP study, the authors used high-resistivity undoped and iron-doped InP


c Β c

Ε Ο Ο

(a) VPE

(h,D°) (D°,X)

«2 τ r

(b) . χ

ΜΒΕ-117

( Α ° , Χ )

ι 1 r Χ

(c) Α ΜΒΕ 11-19

1.513 1.517

Energy (eV)

Fig. 5 .23 High-resolution photoluminescence spectra for the same samples shown in Fig. 5.22, in the exciton region. The peak labels are discussed in the text. (After Koteles et al. (1987).)

films, grown by O M V P E on InP substrates doped with iron, zinc, and sulfur. The films were examined with the 488 nm line from an A r + laser, while held at 5-13 Κ by flowing helium gas. The authors found that iron doping produces many recombination centers in InP, which reduce the efficiency of near-band edge transitions by a factor of 10 for the InP substrates and a factor of 103 for the InP epitaxial films. They found P L lines at 1.07, 1.10, and 1.35 e V that correlated with the presence of iron. These were more useful for characterizing the iron-doped InP films than the weak P L features near the band edge, which could be confused with P L bands from the substrate. Iron-

Two-dimensional mapping 89

doped films grown on substrates doped with zinc showed strong zinc acceptor features, indicating that zinc may diffuse from the substrate into the iron-doped films.

5.5 Two-dimensional mapping

Because it uses short-wavelength visible light, P L can provide two-dimensional maps of semiconductor properties at spatial resolutions suitable for probing device-size structures. Hennessy et al. (1990) and Moore and Miner (1990) have scanned G a A s , A l x G a x _ x A s / G a A s , I n ! _ x G a x A s / G a A s , and I n ^ G a ^ A S y P ^ / G a A s with the commercial system described in Chapter 4. Measurements at a single wavelength, called 'topography', could be made in 40min for the 2 x 10 5 data points from a 100 μιη grid over a 2 in. wafer. Changes in intensity at the P L peak were ascribed to variations in composition, defect density, and layer thickness. But these authors note a problem in single-wavelength topography: apparent intensity changes may come from a shift in peak wavelength. It is better to measure a full P L spectrum at each position on the wafer. Then peak wavelength can be related to alloy composition when Egap is known versus x; and peak intensity and halfwidth, to other quality factors such as variations in layer thickness and in the number of nonradiative recombination centers. Selection of wavelength can also probe specific impurities. Hence a map produced at 830 nm corresponds to the carbon distribution in the sample. These workers display their data as false-color maps of properties across the wafer, which give an easily assimilated overview, but which I cannot readily reproduce in black and white. The maps are shown in the references Hennessy et al. (1990) and Moore and Miner (1990).

Moretti et al. (1989) used another design for a scanning P L system (described in Chapter 3) to examine G a A s / A l ^ G a ^ A s quantum wells for use in laser diodes, modulators, and other photonic devices. These authors note that uniformity of growth of a wafer is central to obtaining a high yield of usable devices. The ideal technique would be nondestructive, would examine a large area, and would measure properties important for photonic devices. Photoluminescence meets these criteria. Its direct measurement of gap energy is related to well width and barrier composition in a microstructure.

Like Moore and Miner (1990), Moretti et al. point out that single-wavelength measurements are far less valuable than a full spectrum from each point on the sample. Using an optical multichannel analyzer ( O M A ) to rapidly acquire each spectrum, they examined a single quantum well grown by M B E on a 2 in. semi-insulating GaAs wafer. The structure consisted of G a A s buffer layer, followed by a 100 nm A l o . 3 G a o . 7 A s barrier, a 5nm GaAs quantum well, and a 100 nm Alo.3Gao.7As barrier layer. Figure 5.24 shows P L from the center of the wafer, and from its edge. The shift in peak position from 816.8 to 807.6 nm represents variation in the width of the well, the χ value of the barrier layers, or both.

Moretti et al. display a map of the peak wavelength distribution over the wafer in a gray-scale display. This image showed strong radial symmetry for the peak wavelength, with the wavelength decreasing radially from the center of the wafer. Since the substrate rotated during growth, the authors ascribe the radial decrease of wavelength to the distribution of aluminum and gallium flux in the M B E chamber. One long-wavelength peninsula which broke the radial symmetry corresponded to an area on

http://Alo.3Gao.7As

http://Alo.3Gao.7As


(b) (a)

750 775 800 825 850 λ (nm)

Fig. 5 .24 Photoluminescence from the center (a) and edge (b) of a 2 in. wafer with a single quantum well. The well is 5nm of GaAs between 100 nm barriers of Alo.3Gao.7As. The peak represents the η = 1 electron-heavy hole transition (see Section 5.6). The shift in peak position from (a) to (b) , 816.8 nm to 807.6 nm, comes from changes in well width, in barrier composition, or in both. This plot illustrates the difficulty of interpreting single-wavelength PL topography and the advantage of obtaining a full spectrum at each point. (After Moretti et al. (1989).)

the back surface of the wafer which lacked the indium coating used to provide thermal contact to the heater. Where indium is missing, the wafer temperature is lower, leading to greater adhesion of gallium, a thicker quantum well, and higher gallium concentration in the barriers. This would decrease the peak wavelength, just as the map showed.

Iizuka et al. (1989) used a different kind of P L topography to characterize crystal defects in In.0jGa0.9As/GaAs quantum wells and strained-layer superlattices. The single In 0 jGa 0 . 9As wells lay between GaAs barriers and were 5-60 monolayers wide (1 monolayer = 0.28nm). The superlattices comprised 20 pairs of Ino.1Gao.9As layers 7nm thick and GaAs layers 10 nm thick. Infrared emission was excited by a K r + laser operating at 647.1 nm, with the samples held at 50 Κ in a closed cycle helium refrigerator. Emission over the sample area was detected and imaged by a commercial infrared vidicon camera.

Iizuka et al. observed dark lines in the images from the quantum wells which seemed associated with lattice mismatch; the number of lines increased with well width, and did not appear for lattice-matched A l ^ G a ^ A s / G a A s . For the superlattices, the topographic pictures revealed stripes both parallel and perpendicular to the (110) direction, but Nomarski contrast microscopy revealed no such stripes on the surface of the structure. However, in bulk In 0 . iGa 0 .9As/GaAs, similar—although finer—cross stripes were seen on the surface by P L topography. The authors concluded that these features appear in the superlattice near the substrate and are related to the lattice mismatch. T o explain their observations, they proposed a model for the propagation of dislocations in these heterostructures. In films of In.0jGa0.9As on

http://Alo.3Gao.7As

http://In.0jGa0.9As/Ga

http://Ino.1Gao.9As

http://In.0jGa0.9As

Two-dimensional mapping 91

GaAs , defects produced by lattice mismatch propagate to the film surface, but in the microstructure, defects produced at the superlattice-substrate interface propagate only until they reach an interface, where they bend and continue along the interface. Hence the defects never reach the free surface.

Chen and Lyon (1989) also measured P L over a spatial region to examine a semiconductor, in this case silicon. Their interest was in determining the diffusion length of free excitons, for fundamental reasons and because diffusion processes affect any spatial P L measurements of impurities in silicon. The measurement required a specially constructed sample. Starting with silicon doped with phosphorus to a concentration of 4 x 1 0 1 5 c m - 3 , boron was diffused in to make 200 μπι stripes separated by 15 μηι. The stripes were 10 μΐΉ deep, and contained a boron concentration of approximately 1 0 I 6 c m - 3 . A n A r + laser operating at 514.5 nm was the excitation source, with the radiation focused to a spot size of approximately 10 μηι. With the sample cooled to helium temperatures, the luminescence due to free excitons was measured as a function of beam position on the sample, perpendicular to the boron trenches.

Figure 5.25 illustrates the measurement geometry and shows a plot of P L intensity (measured at 1.1296 μηι, the location of the free exciton line) versus position. The high impurity concentration in the boron-doped regions captures excitons, reducing

19 Κ

200 400 600 Χ(μΓΠ)

800

Fig. 5 .25 Spatially resolved photoluminescence from Si:Ρ with boron added to create the geometry shown. As the beam sweeps the sample, PL intensity at the free exciton line (at 1.1296 μηι) decreases in regions with boron and increases where boron is absent. The plot yields the free exciton diffusion length. (After Chen and Lyon (1989).)


the number of free excitons compared to the areas without boron. This causes the variation of signal with position displayed in the figure, which shows that the free exciton concentration decreases exponentially as the beam traverses the interface between the boron-doped and undoped regions. This behavior was analyzed to give the free exciton diffusion length L. In a related measurement using pulsed excitation from an A r + laser, these researchers also measured the free exciton lifetime τ . From the formula D = L 2 / t , where D is the exciton diffusivity, they obtain values for D much lower than in bulk silicon, which they ascribe to influence of the shallow impurities.

5.6 Interfaces and microstructures

The exploration of stress at interfaces is only one aspect of the P L study of semiconductor microstructures, from films on substrates to quantum wells and superlattices. Photoluminescence methods also return a great deal of other important information about such structures. They have played a major role in theoretical understanding and materials development for systems based on A l x G a ! _ x A s . This work provides an excellent set of case studies to show the power of P L analysis.

In very early work Dingle et al. (1974), and Dingle (1975) pointed out that the one-dimensional potential well, a central problem of quantum mechanics, is a key to the behavior of thin semiconducting layers. They report on the bound states that result in potential wells formed from thin layers of G a A s set between barrier layers of A l x G a ! _ x A s . Their discussion of the fundamental physics makes clear how optical analysis gives detailed quantitative information about electrons, light and heavy holes, and excitons in thin GaAs layers; and notes how these results from quantum theory are important for technology as well.

Equation (3.6) gives the approximate energy levels for the states quantized in the ζ direction in an energy well of width Lz. In the real world of A l ^ G a ^ A s heterostruc-tures, separated heavy and light hole bands exist because there is always some uniaxial stress to split the valence bands, as I discussed in the section on stress analysis. Hence excitons can form between electrons and light holes, and electrons and heavy holes, leading to two series of excitonic peaks, each following the quantized form of equation (3.6) . The earliest experiments to display the quantum nature were carried out by optical absorption, not by P L , and illustrate the basic physics very well. Figure 3.7, for example, shows absorption spectra at 2 Κ for different values of well thickness Lz. The peaks behave as expected with thickness. The lowest peak from the well with Lz = 14 nm barely shows a hint of a second resolved feature, evidence for a second set of excitons.

Photoluminescence and P L E studies of quantum wells soon came in great abundance, and clearly showed both transitions. One of the earlier was made by Miller et al. (1980), who examined structures consisting of 30 periods of 26 nm GaAs wells and 26 nm A l 0 2 sGa 0 7 2 A s barriers, over 7-150 K . They observed both light hole ( L H ) and heavy hole ( H H ) excitonic features, as seen in Fig. 5.26. In one significant result, they found that the H H excitons sustained broadened transitions. They ascribed this to variation in the widths Lz of the GaAs wells. Since the energies of excitons confined in a narrow well depend on Lz through equation (3 .6) , a spread in width produces a spread in energy, which broadens the transitions. The broadening Miller et al.

Interfaces and microstructures 93

meV meV meV

U I 1 1 1 1 1 1 1.52 1.54 1.56 1.58 1.60 1.62 1.64

Energy (eV)

Fig. 5 .26 Photoluminescence excitation (PLE) spectrum of a G a A s - A l ^ G a ^ A s multiquan-tum well at 7 K. Excitation is at 1.519 eV, the peak of the PL spectrum. Peaks labeled Enh and Ene (n = 1, 2, 3, 4) come from the heavy hole and light hole excitons respectively. The former broaden with n. This is ascribed to variation in the well widths. See text. (After Miller et al. (1980).)

observed corresponded to spread in well width of about 2 nm. This was one indication of how useful P L data could be in characterizing microstructures. Later Miller et al. (1981) examined single and multiwell structures with Lz = 4.2 to 14.5 nm. They used P L E spectroscopy, where the luminescence at 1.6288 e V was measured versus excitation pump energy, to observe H H and L H excitons (Fig. 5.27) including evidence of their excited states. T o describe the measured energies, they used a theory for the two-dimensional H H and L H excitons that exist in narrow wells, and derived their energy dependence on well width.

The sensitivity of P L probes to extremely small effects in heterostructures was further displayed in the work of Miller et al. (1982), who examined undoped M B E -grown single and double G a A s wells. They concluded that the first well grown in a series differs from the rest. Its first interface is typically rough on the scale of 0.6-1.5 nm, and a luminescing impurity—which they could not clearly identify—appears in the first few nanometers of GaAs grown. Maki et al. (1983) made specific connections with growth properties, as they developed methods to grow A l ^ G a x - ^ A s - G a A s heterojunctions for devices. They presented P L results from single quantum wells grown by M B E on Alo.2Gao.8As buffers. Their P L linewidth of 0.7 m e V for the

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94 Case studies : photoluminescenc e characterizatio n

Ê 1 :1H

Ο Ο

sz CL

Έ

1.64 1.66 1.68 Energy (eV)

1.70 1.72 1.74

Fig. 5 .27 Photoluminescence excitation spectrum from a single GaAs quantum well, 4.2nm wide, between 0.65 μιη layers of A l 0 3 7Ga() 6 3 A s . The dominant peaks come from the η = 1 LH and HH excitons, and the shoulders marked 2S represent their excited states. These data were analyzed with the quantum theory for two-dimensional excitons. (After Miller et al. (1981).)

confined η = 1 electron-heavy hole transition was the narrowest reported at the time for any epitaxial material. The quality of the P L spectrum was found to depend on the annealing procedure used on the substrate before growth, and on the ratio of the V i l l i fluxes during growth.

Petroff et al. (1984) provided further understanding of the relation between P L spectra and interface quality in G a A s - A l ^ G a ^ A s structures. They especially noted that A l x G a ] _ x A s - G a A s and G a A s - A l ^ G a ^ A s interfaces behave differently. Interface roughness depends on the order in which one material is grown on the other. Their samples were MBE-grown GaAs single quantum wells with Lz between 10 and 17.2 nm, set between A l ^ G a ^ A s cladding layers which themselves contained GaAs quantum wells with Lz ^ 5 nm. These were excited by a tunable continuous wave dye laser at an intensity of 1 W c m - 2 , to obtain P L and P L E spectra at 6 K . Earlier work by Weisbuch et al. (1981) had shown that the widths of the P L peaks due to the confined light hole and heavy hole free excitons correlated with interface smoothness. Using this and other criteria, Petroff et al. ranked the interface quality of six different single-well samples.

The ranking showed that the more GaAs quantum wells or interfaces below the particular well under examination, the greater the luminescence efficiency. These and similar observations, combined with data from cathodoluminescence and transmission electron microscopy ( T E M ) , led the researchers to conclude that interfaces trap nonradiative centers, and that the responsible impurities or impurity complexes originate both from the substrate and A l ^ G a ^ A s layers. Although the specific trapping mechanism was not identified, one proposal was that misfit strain across the interfaces provided a gettering effect.


Woodbridge et al. (1984) used P L to find whether improvements in single quantum wells due to the layer sequence prior to the well represented gettering, or interface smoothing. They examined G a A s - A l ^ G a ^ A s wells grown in an M B E system with high-speed sample rotation to enhance uniformity. One evaluation compared a 2.5 nm GaAs well grown on 1 μιτι of A l 0 4 G a 0 6 A s followed by a 0.1 μηι A l o . 4 G a o . 6 A s cap, with a second structure with 1 μηι of A l 0 4 Ga 0 . 6As , followed by a 1 nm A l A s prelayer, followed by 15 nm of Al0.4Ga0.eAs, followed by a quantum well and cap as in the first sample. Figure 5.28 shows P L spectra from these samples with and without the A l A s prelayer (Nos 121 and 122, respectively). The large, broad asymmetric peak for N o . 122, which shifts with excitation power, comes from recombination between an η = 1 electron and an acceptor, probably carbon. In N o . 121, this impurity peak is much weaker than a new peak at 1.69 e V , from the η = 1 e lec t ron-HH free exciton. Its width of 12 m e V shows that the interface is very smooth. Previous work had suggested that prelayers influence P L spectra by smoothing interfaces. However , T E M photographs showed no significant difference in smoothness between Nos 121 and 122, suggesting a gettering effect instead. These researchers also examined multiple wells, the P L spectra of which showed excitonic features even at room temperature. This attests to the quality of the samples grown on the M B E system with a revolving stage.

Fujiwara and Ploog (1984) studied the effect of cladding on the quality of GaAs single quantum wells. They noted a possible reason that heterostructure configuration affects P L or quantum efficiency, namely the poor crystal quality of the cladding A l x G a ! _ x A s layers, leading to poor interface properties. A s an alternative, they studied single wells clad between layers consisting of G a A s - A l A s short-period super-

Fig. 5 .28 Photoluminescence from identical GaAs single quantum wells with (No. 121) and without (No. 122) an AlAs prelayer. The asymmetric peak for No. 122, which shifts with excitation power, represents recombination between an η = 1 electron and an acceptor, probably carbon. In No. 121, this impurity peak becomes much less important than the new 733 nm (1.69 eV) peak from an electron-HH free exciton. (After Woodbridge et al. (1984).)

700 Wavelength (nm)

800

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850 800 750 700 650

Wavelength, λ (nm)

Fig. 5 .29 Low-temperature photoluminescence spectra for a single GaAs quantum well clad between: (a) Alo.5Gao.5As layers; (b) short-period GaAs/AlAs superlattices (SPS). The FWHM of the line near 800 nm in each spectrum, the confined η = 1 electron-HH transition, indicates that SPS cladding gives better interfaces. (After Fujiwara and Ploog (1984).)

lattices (SPS) . Their MBE-grown wells were 6.4 or 9.2 nm thick. Figure 5.29 compares P L spectra from two 9.2 nm wells, one clad between layers of A l 0 5 G a 0 . 5 A s , and the second clad between G a A s / A l A s SPS with layer thickness 2.4 and 2.6 nm respectively and with 10 periods. The dominant line in each spectrum (at 799.9 nm (1.550 e V ) for the Al x Ga!_ x As-c lad sample, and at 798.3 nm (1.553 e V ) for the SPS-clad sample) is associated with the confined η = 1 e lec t ron-HH transition. This line displays a F W H M of 11 m e V for the SPS-clad sample, a reasonable value for a single quantum well, whereas the much larger value of 19 m e V for the Alo. 5Gao.5As-clad sample indicates inferior interface properties.

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1.57 1.58 1.59 0.2 2 20

Energy (eV) Excitation intensity (W cm - 2 )

Fig. 5 .30 (a) Photoluminescence from a 5nm thick GaAs/Alo.isGao.ssAs quantum well at different excitation intensities, (b) Dependence of the PL signal on excitation intensity for the peaks in (a). These dependences and its position indicate that the low energy peak is related to acceptor impurities. The high energy peak comes from the electron-HH exciton. (After Meynadier et al (1985).)

Meynadier et al. (1985) summarized earlier work and contributed their own study of prelayers. T o obtain qualitative and quantitative information about impurities near the well, these researchers examined single wells grown between different arrangements of A ^ G a ^ ^ A s prelayers which varied in their number and aluminum concentration. The wells were MBE-grown at a substrate temperature of 680 °C, with all parameters and procedures chosen to yield high-quality samples. The P L spectra all showed two peaks at low excitation energies, as shown in Fig. 5.30. The higher peak is assigned to the η = 1 e lec t ron-HH exciton, as confirmed by P L E spectroscopy. A s excitation power increases, the low energy peak diminishes. This effect, related to saturation in recombination, and the location of the peak, strongly indicate that it comes from the free-electron to acceptor transition; hence it is a marker for acceptor impurities.

Figure 5.31 compares the integrated intensity of the acceptor-related peak with the total integrated intensity from single quantum wells with no prelayers, one prelayer, and three prelayers. The single quantum well, 5nm thick, is embedded in Alo.13Gao.87As. The prelayers consist of 2nm thick GaAs wells. The emission due to acceptors decreases as the number of prelayers increases, suggesting that the prelayers are impurity-trapping centers. Figure 5.32 traces the influence of different aluminum concentrations (x = 0, 0.13, and 0.32) in the A l ^ G a ^ A s prelayers. The acceptor-related P L decreases relative to the excitonic P L as the aluminum concentration decreases. The authors interpret this behavior as showing that impurity trapping is more efficient in GaAs than in A l ^ G a ^ A s layers. This supports other observations

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Fig. 5.31 Acceptor-related PL intensity relative to total PL intensity versus laser power, for the quantum well in Fig. 5.30: (1), no prelayer; (2) one prelayer; (3) three prelayers. Each prelayer is a 2nm GaAs well. The acceptor-related intensity decreases as the number of prelayers increases. (After Meynadier et al. (1985).)

1.58 1.57 1.56 1.55 Energy (meV)

-400 -200 0 ζ (A)

Fig. 5 .32 Photoluminescence spectra for the quantum well in Fig. 5.30, for A l ^ G a ^ A s prelayers with different aluminum concentrations: (1) χ = 0.32; (2) χ = 0.13; (3) χ = 0. The acceptor-related peak at 1.57-1.58 eV decreases as χ increases. (After Meynadier et al. (1985).)

that G a A s / A l A s structures provide better cladding than do A l ^ G a ^ A s layers, such as those by Fujiwara and Ploog (1984) discussed above.

The authors also drew conclusions about the spatial distribution of impurities. From the difference between the energies of the electron-acceptor peak and the free exciton peak, they found the binding energy of the acceptors, assuming they were located at


Energy (meV)

Fig. 5 .33 (a) Assumed acceptor concentration profile at the interface of a quantum well like that in Figs 5.30 to 5.32; (b): experimental ( ) and calculated ( ) spectra for the electron to acceptor PL peak for a well 10nm wide. (After Meynadier et al. (1985).)

the well interfaces. Then they calculated a P L line shape for the recombination of electrons with acceptors spread out in space, using the distribution shown in Fig. 5.33. The distribution is asymmetric around the interface because of the way impurities are incorporated. The figure shows that a P L spectrum calculated with these distributions agrees with the data for a 10 nm wide well, and other calculations (not shown) also agreed with data from 5 and 15 nm wells. This analysis showed that acceptor distributions within the well spread over 1.2-3.0nm (four to ten monolayers), and those outside the well covered 0.6-0.8nm (two to three monolayers).

In other P L work assessing G a A s - A l ^ G a ^ ^ A s wells, Miller et al. (1984a) examined quantum wells grown by M O C V D rather than M B E . A tunable dye laser gave P L and P L E spectra, at sample temperatures as low as 5 K . For single wells, the η = 1 L H and H H peaks were seen, and the η = 2 H H exciton. From the peak widths, the authors concluded that the G a A s - A l ^ G a ^ ^ A s interfaces consisted of islands 30 nm or greater in extent and slightly thicker than a monolayer (0.28nm), so they are extremely smooth. Complementary results from transmission electron microscopy indicated that the width of compositional grading across the interface was less than 0.8 nm. The P L data could not resolve such grading effects, but further P L studies of multiple quantum wells indicated no interfaces with gradations greater than three monolayers.

Wilson (1989) has also considered the effect of interface disorder on P L spectra. Disorder, she notes, includes physical roughness at the interface due to the growth process, and defects or impurities concentrated at the interface. The reduction of these is an important goal in the refinement of growth methods for new microstructure devices. One technique explored to increase surface smoothing is interrupted growth. The growth process is stopped for a short time at the top of each layer, to allow the newly deposited atoms to achieve a smooth surface. Figure 5.34 shows the manifestation of these effects in P L spectra from MBE-grown GaAs quantum wells. The lower trace shows data from wells grown without any pauses in the process. The exciton peaks are broadened, due to variations in well thicknesses on the scale of a monolayer. The upper curve comes from wells grown with a 2 min interruption at each interface. The multiple sharp peaks are associated with extended regions which are


1.0 22 11

Interrupted growth 7

Continuous growth c C

1.0

0.0 A λ Λ_ 1.5 1.6 1.7

Energy (eV) 1.8

Fig. 5 .34 Photoluminescence from two MBE-grown samples, each with four single GaAs quantum wells of different thickness. The lower trace shows that with continuous growth the exciton peaks are broadened by interface roughness of about a monolayer. In the upper trace, the peaks sharpen when growth is interrupted for 2min at each interface, indicating that the emission comes from extended areas flat to within a monolayer. Each peak is marked with the number of monolayers in the corresponding well. (After Tu et al. (1987).)

flat to within a monolayer. However , this spectrum also shows an unfavorable outcome of interrupted growth, the incorporation of additional impurities at the interface during the delay. This is seen in the shoulders on the low-energy side of several of the main peaks, which are associated with free-to-bound transitions at carbon sites. (Wi l son goes on to show how the determination of lifetimes by pulsed spectroscopy also probes surface disorder. I comment further on time-resolved work in Chapter 8.)

Recent work in A l x G a ! _ x A s / G a A s quantum wells, even after growth techniques have become well established, continues to rely on P L to define sample quality. Fujiwara et al. (1988) examined MBE-grown undoped GaAs single quantum wells 6.1 nm wide confined by 200 nm of A l 0 ^ G a o ^ A s . They observed at room temperature (see Fig. 5.35) a sharp P L peak with a F W H M of 14 meV from the η = 1 H H free exciton, and a shoulder from the η = 1 L H exciton. The main peak does not change when the sample is illuminated at different positions as shown, indicating that the layer thickness is uniform. The observation of an exciton at room temperature, its small linewidth, and its independence of the spatial location of the exciting light are all indicators of high sample quality. These authors comment that the further understanding of room-temperature behavior is important for device development.

Photoluminescence also provides a sensitive way to examine novel microstructures, and to further explore the complex issue of band offsets. Miller et al. (1984b), for instance, examined MBE-grown GaAs-Alo.3Gao.7As quantum wells with parabolic compositional well profiles. Each well consisted of 20 layers of Alo.3Gao.7As and 21 layers of GaAs , with the thickness of the former increasing quadratically with distance

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PC / , j

. Λ / 14 meV \l ' \ • a T

• b Λ , 7 mm ι Π • c _l

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(b) PL

J ^ _ (c) 1HH 1LH 1

850 800 750

Wavelength (nm)

Fig. 5 .35 Photoluminescence at 300 K ( ) from a single GaAs quantum well (Lz = 6.1 nm) between Alo.24Gao.76As barriers. The peaks come from the η = 1HH and L H free excitons. Their presence at room temperature indicates good sample quality. The similarity of curves (a), (b) , and (c) , from illumination at different points (inset), indicates uniform well width. The dashed line is photocurrent, which increases at the HH transition, representing photoexcited carriers due to excitonic resonant absorption. (After Fujiwara et al. (1988).)

from the well center while the thickness of the G a A s decreased. Such wells support simple harmonic oscillator wavefunctions, which lead to approximate confined energies:

2JLZ\ m* j

where QAEg is related to the energy-gap discontinuity between G a A s and A ^ G a ^ A s . This important parameter does not appear in the theory for square wells, which is one reason to examine parabolic systems. Figure 5.36 shows P L and P L E spectra from a sample with ten periods, each consisting of a parabolic well with Lz = 51 nm, and A l o . 3 G a o . 7 A s barriers with L B = 23.7 nm. The excitation spectrum shows many labelled transitions and calculated values from a more accurate version of equation

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1.52 1.53 1.54

I I I ι I • I • I • I • 1 ι I ι I ι I ι I ι I ι I ι I ι l ι I ι I ι ι ι I t » • I • I ι » ι l 1.52 1.56 1.60 1.64 1.68 1.72

Photon energy (eV)

Fig. 5 .36 Photoluminescence spectrum (inset) and photoluminescence excitation (PLE) spectrum at 5 Κ from parabolic G a A s - A l ^ G a ^ A s quantum wells (Lz = 51.0nm). The PL spectrum is excited at 1.6eV. The PLE spectrum was excited at the peak of the PL emission, 1.531 eV. It shows an L H excitonic transition (Eu) and HH excitonic transitions (Enf, η = 1, 2, . . . ) . Their energies, compared with a more exact version of equation (5.7), yield the band offsets between GaAs and A l ^ G a ^ A s . (After Miller et al (1984b).)

(5.5) . However, calculation and data could be made to agree only when the energy-gap discontinuity was split equally between the conduction and the valence band wells, whereas other results suggest an allocation of 85-15%. The authors did not explain the discrepancy. Nevertheless the work shows the usefulness of P L to examine unusual microstructures.

Although applications to A l ^ G a ^ A s based structures dominate, P L characterization is equally useful for other quantum well and superlattice systems. Abstreiter et al. (1989), for instance, have combined Raman and P L data to examine S i - G e superlattices. I discuss their Raman analysis in Chapter 6. Figure 5.37 shows P L spectra from these authors, for S i 6 G e 4 superlattices (which have six silicon monolayers and four germanium monolayers per period) with different strain distributions. The single large peak which shifts with strain suggests that gaps with direct transition character exist in these systems, although it is not definite that the feature is an intrinsic one.


0.70 0.75 0.80 0.85 0.90 0.95

Energy (eV)

Fig. 5 .37 Photoluminescence spectra from Si 6 Ge 4 superlattices with six monolayers of silicon and four of germanium per period, each with the indicated strain ε in the silicon layers. The peak shifts with strain are consistent with the existence of an energy gap in the alloy with a direct transition character. (After Abstreiter et al. (1989).)

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