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Materials-related reliability aspects of III–V optical devices
Materials Science and Engineering, B20 (1993) 9-20
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Materials-related reliability aspects of III-V optical devices O. Ueda Fujitsu Laboratories Ltd., 10-1 Morinosato-Wakamiya, Atsugi 243-01 (Japan)
Abstract This paper describes the current understanding of materials issues in III-V compound semiconductor heterostructures,
and the degradation of optical devices, lasers and light-emitting diodes fabricated from these materials. Generation of defects and thermal instability are among these issues for these systems. Defects introduced during crystal growth are classified into two types: interface defects and bulk defects. Interface-type defects are stacking faults, V-shaped dislocations, dislocation clusters, microtwins, inclusions and misfit dislocations. The bulk defect group includes precipitates and dislocation loops. Defects in the substrate also can be propagated into the epilayer. Structural imperfections owing to thermal instability are also found. They are quasi-periodic modulated structures, owing to spinodal decomposition of the crystal either at the liquid-solid interface or growth surface, and owing to atomic ordering, which also occurs on the growth surface through the migration and reconstruction of the deposited atoms. In each phenomenon, the nature and generation mechanism of defects and structures are clarified. Three major degradation modes--rapid degradation, gradual degradation and catastrophic failure--are discussed. Rapid degradation occurs via either recombinationenhanced dislocation climb or glide. The ease with which these phenomena occur in different heterostructures are presented. Based on the results, the dominant parameters involved in the phenomena are discussed. Gradual degradation takes place presumably owing to recombination-enhanced point defect reactions in GaA1As/GaAs-based optical devices. This mode is also enhanced by the internal stress owing to lattice mismatch. However, we do not observe such degradation in InGaAsP/InP-based optical devices. Catastrophic failure is found to be due to catastrophic optical damage at a mirror or at a defect in GaA1As/GaAs double-heterostrncture (DH) lasers but not in InGaAsP/InP DH lasers. In each degradation mode, the effect of the defects and/or structures on the degradation is discussed and methods for the elimination of degradation are proposed.
1. Introduction Since the discovery of semiconductor lasers in the early 1970s [1], a variety of optical devices have been developed from III-V alloy semiconductors. Current fields of application include fiber optical communication systems, digital audio systems and optical printers. In the development of these devices, high reliability has been one of the most important goals. Degradation is a major hindrance to reliability. Moreover, it has been clarified that defects in crystals very often cause device degradation. In this paper, we describe materials issues in III-V alloy semiconductors, i.e. the generation of defects during growth and phenomena induced by the thermal instability of the material, and our current understanding of degradation in III-V semiconductor lasers and light-emitting diodes (LEDs) in terms of three major degradation modes--rapid degradation, gradual degradation and catastrophic failure. For each mode, we discuss the role of defects in the degradation and the elimination of degradation. 0921-5107/93/$6.00
2. Materials issues in III-V alloy semiconductor thin films Materials issues in the growth of III-V alloy semiconductor thin films are focused on. These include the generation of various kinds of defects during growth, the generation of modulated structures owing to spinodal decomposition and atomic ordering associated with growth kinetics on the surface.
2.1. Defect generation Defects introduced during crystal growth are classified into two types, namely interface defects and bulk defects. Defects of the interface type are stacking faults, V-shaped dislocations, microtwins, inclusions and misfit dislocations.
2.1.1. Stacking faults Stacking faults are generated from the heterointerface, extending on the four equivalent {111} planes, forming a stacking fault tetrahedron (see Fig. l(a)). Figure 2(a) illustrates a typical optical micrograph of © 1993 - Elsevier Sequoia. All rights reserved
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Materials-related refiability o1711- V optical devices
(b)
U (c)
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Fig. 1. Interface defects resultingfrom contamination,thermal damage of the substrate or local segregation. the surface of a GaAsSb crystal grown at 590 °C. A high density of surface defects are observed. They are hillock-like defects lying along the (110) direction and are very similar to the oval defects in molecular beam epitaxy (MBE) Ga(A1)As crystals [2]. The defect structure observed by transmission electron microscopy (TEM) at these surface defects is shown in Fig. 2(b) [3]. Stacking faults lying on the four equivalent {111} planes are seen. Since the thin foil includes only the surface region of the epilayer, one can expect that the surface defects correspond to stacking fault tetrahedra. The stacking faults often multiplied, as denoted by X in Fig. 2(b). In addition, threading dislocations and small stacking faults are found inside the tetrahedra. By considering the length of the stacking faults at the surface, their basal {111} planes and the epilayer thickness, it is deduced that these defects originated at the heterointerface. The density of these defects can be reduced by lowering the growth temperature and cleaning the substrate with an appropriate As overpressure. Therefore, the dominant factors affecting the defect generation could be the evaporation of P atoms from the substrate and/or the existence of some oxide or contamination on the substrate surface. 2.1.2. V-shaped dislocations Pair pits are often observed on the surface of InGaP crystals etched with HF + HNO 3. In the TEM image of the area including pair pits, small dislocation segments are observed, which correspond to the etch pits. From the shape of the dislocations, this dislocation pair is
assumed to be of V-shaped type (see Fig. l(b)). We also have found preferential vapor etching at the V-shaped dislocations during growth [4].
Z 1.3. Microtwins Microtwins are occasionally observed in epilayers grown at higher growth temperatures. They also originate from the heterointerface, extending to the surface area with {111} boundaries (see Fig. l(c)). Some of them terminate in midlayer, with {211} boundaries. A high resolution TEM image of a GaAsSb crystal grown at 590 °C from a region close to the heterointerface is shown in Fig. 3. Brighter regions along the interface are probably due to the existence of some oxides. In this micrograph, microtwins (MTs) are generated in two regions, i.e. I and II. In both regions, amorphous-like areas (A) and shallow, pit-shaped patterns (SP) are observed at the interface just below the amorphous-like area. Microtwins are generated only in the region adjacent to the areas. Based on these results, the following mechanism is suggested for the generation of microtwins: (1) evaporation of P atoms from the substrate surface (formation of In-rich droplets on the surface); (2) generation of amorphous areas as a consequence of an In-rich melt; (3) growth of crystals with different lattice constants around the amorphous areas; (4) generation of microtwins by the relaxation of the stress resulting from the lattice mismatch between the normal GaAsSb areas and those formed in step (3).
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Thus, the elimination of microtwins can be achieved by methods similar to those described for stacking faults.
2.1.4. Inclusions A typical TEM image of inclusion defects in a liquid phase epitaxy(LPE)-grown InGaAsP layer on a (001) GaAs substrate is shown in Fig. 4 [5]. The inclusions are lying in the (100) direction and, from stereoscopic observation, they are found to originate from the interface between the InGaAsP active layer and the InGaP cladding layer. These defects are thought to be formed by the development of some inclusions from the interface in the direction (01 n).
2.1.5. Misfit dislocations
b
•........
o.s, vm
Fig. 2. (a) Optical micrograph of defects revealed on the surface of an MBE-grown GaAsSb crystal on a (001) InP substrate. (b) (001) plan view, bright-field TEM image of a stacking fault tetrahedron corresponding to the surface defect shown in Fig. 2(a).
If there is a difference in the between the lattice constant of the epilayer and that of the substrate, then misfit dislocations are generated when the film thickness exceeds the critical thickness tc, which depends on the structure and elastic properties of the materials and on the type of dislocations. There are several possible mechanisms for the generation of misfit dislocations, as shown in Fig. 5. First, threading dislocations propagated from the substrate may glide out (Fig. 5(a)) or bend out (Fig. 5(b)) to the edge of the wafer. It is also expected that two dislocations with the same Burgers vectors in the substrate terminate into one misfit dislocation at the heterointerface, forming a half-loop (righthand figure of Fig. 5(c)). If the dislocation density in the substrate is low, another mechanism is expected, as follows: (i) nucleation of microhalf-loops at the surface; (ii) expansion of loops down to the interface; (iii) gliding out of the threading segments to the edge (left-hand figure of Fig. 5(c)) [6]. Since the first three kinds of defects described above are eliminated by cleaning of the substrate surface prior to growth and/or the protection of the substrate against evaporation of substrate atoms, they are expected to be generated by contamination and/or
M, aAsSb
~P
20 nm Fig. 3. Low magnification (110) cross-sectional high resolution TEM image of the GaAsSb-(00 l)InP interface.
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thermal damage to the substrate. The inclusions are presumably caused by local segregation of solute atoms at the interface. In highly mismatched heteroepitaxy, microtwins (see Fig. 5(d)) and/or stacking faults are very often generated, particularly during the initial stage of growth [7]. They propagate along the four equivalent {111 }planes. However, they often annihilate in midlayer, i.e. terminating in {211 }-type boundaries for microtwins and coalescence of two stacking faults on different planes. 2.1.6. Bulk defects Bulk defects are classified into two types, i.e. defects propagated from the substrate and those generated by local segregation of dopant atoms or native point defects. Since the defects propagated from the substrate are simply eliminated by using substrates with low defect densities, we only consider the defects of the
Fig. 4. Plan view TEM image of inclusions in LPE-grown InGaAsP on a (001 ) GaAs substrate.
--/-- _/__-_
.
.
.
.
a)
(b)
(c)
(d)
.
Fig. 5. Interface defects resulting from lattice mismatch.
other type. Typical features of these defects are shown schematically in Fig. 6. In most cases, they are generated in heavily doped crystals, as shown in Figs. 6(a)-6(c). When supersaturated dopants are present in the crystal, they tend to segregate locally during growth or the cooling process after growth, forming precipitates (see Fig. 6(a)). Figure 7 illustrates a plan view, high resolution TEM image of one of the uniformly distributed precipitates in heavily doped InP with Fe grown by metallo-organic vapor phase epitaxy (MOVPE) [8]. The precipitates are found to be spherical FeP particles nearly lattice matched to the matrix crystal. When there is a large lattice mismatch between the precipitates and the matrix, dislocation loops are often punched out from the precipitates by the relaxation of stress (see Fig. 6(b)). Such precipitates are often observed in heavily doped LPE-grown InGaAsP crystals on InP substrates [9]. In the formation of the precipitates, excess native point defects can be introduced. They may condense at some nucleation centers to form dislocation loops (either faulted or unfaulted), as shown in Fig. 6(c). Interstitial-type dislocation loops are often observed in heavily doped LPE-grown GaA1As crystals with Ge [10]. However, dislocation loops are present even in undoped crystals. We have found both vacancy-type and interstitial-type loops in undoped LPE-InGaP crystals [11]. Since modulated structures (see the next section) are generated in the crystal, these loops are presumably caused by the condensation of excess point defects induced as byproducts of modulated structure generation.
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Fig. 6. Bulk defects resulting from local segregation of dopant atoms or native point defects.
nm Fig. 7. FeP precipitate in MOVPE-grown InP heavily doped with Fe.
2.2. Structural imperfections owing to thermal instability of the crystals 2.2.1. Modulated structures Modulated structures are generated in alloy semiconductors such as InGaAs, InGaP, InGaAsP and GaAsSb, whose compositions are inside the "spinodal region" [12]. They are observed as quasi-periodic diffraction contrast in bright- or dark-field images under two-beam or multibeam conditions. Figure 8 shows a plan view, bright-field TEM image of modulated structures in LPE-grown InGaAsP on (001) GaAs (2pe=730 nm) [13]. They develop in the two equivalent directions of (100) and (010) and are colum-
nar shaped in the (001) growth direction. In spinodally decomposed alloys, which have asymmetry in their elastic coefficients, modulated waves extend so as to minimize the strain contribution to the free energy in the solid solution, i.e. the modulated structures are induced in the direction [hkl], which minimizes the elastic coefficient Yhkl. From calculations of the values Yl00, YII0, and Ylll for GaAs, GaP, InAs and InP, it has been found that in all cases, Y100< Yi l0 < Ylll [13]. This can explain well the preferential direction of elongation of the structures. It is also assumed that they are very stable once they are formed, propagating in the (001) growth direction. Furthermore, it is clarified that both structures with long (50-200 nm) and short (5-20 nm) periodicities are formed when the composition of the crystal is inside the spinodal region (Fig. 8(a)), whereas only the structures with short periodicities are formed when the composition is outside the spinodal region (Fig. 8(b)). Thus, one can expect that the structures with long periodicities are formed during growth by spinodal decomposition and that the short periodicity structures form during the cooling process after growth. Moreover, the compositional fluctuations along the structure measured by energy-dispersive X-ray analysis are in the range 2%-3%, which is far smaller than that expected. This may be due to the fact that, in semiconductors, atomic diffusion in the bulk or on the surface is very slow compared with the case of metal alloys. The generation of modulated structures does not depend on the growth method. The decompostion process may be enhanced by atomic diffusion at the liquid-solid interface for LPE growth and at the growth surface for MOVPE, VPE and MBE.
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Materials-related reliability of Ill- V optical devices Z2.2. Ordered structures
Fig. 8. Modulated structures in LPE-grown InGaAsP. (a) 2pL= 730 nm; (b)),PL= 800 nm. From the results described above, it is suggested that modulated structures can be eliminated by the growth of the crystal at temperatures where the composition of the crystal is outside the spinodal region and subsequent rapid cooling or quenching of the crystal.
Recently, ordered structures (or natural superlattices), in which two kinds of element atoms in column III or V lattice periodically arrange, have been observed in various II1-V alloy semiconductors [14]. In most cases, the ordered structures are generated locally in the crystal and the degree of ordering is strongly affected by the growth method and growth conditions. It is known that there are three types of substitutionally ordered structures in f.c.c, lattice of AB-type binary alloy systems. They are the types of CuAu-I (L10), CuPt (Lll) and chalcopyrite (Ell), whose atomic arrangements and corresponding reciprocal lattices are shown schematically in Fig. 9. The results for ordered structures in III-V alloy semi conductors obtained until now can be summarized as follows: (a) the most common structure in crystals grown on a (001) substrate is CuPt type (see Fig. 10) and other structures are only occasionally observed; (b) ordered structures are generated very often in crystals grown by MOVPE, VPE and MBE, whereas no ordering takes place in LPE-grown crystals (except for one report for ordered structures in InGaAs by Nakayama et al. [14], which has not been confirmed by others); (c) regarding CuPt-type structures, one can observe only two variants in the (110) cross-section (atomic steps on the growth surface are believed to play an important role in the generation of the ordered structures, since one of the two variants is preferentially enhanced when substrates tilted toward the (110) directions are used); (d) the ordering is not perfect and the ordered regions are plate-like microdomains lying on planes nearly parallel to the growth surface; (e) defects, such as antiphase boundaries, are often generated in the ordered regions; (f) the degree of ordering depends on the growth temperature, V/III partial pressure ratio (in the case of MOVPE and MBE) and rotation velocity of the substrate; (g) only strongly ordered InGaP crystals grown by MOVPE exhibit abnormal bandgap energies (up to 50 meV lower than the normal value). From these results, it is now understood that the ordered structures are generated by the migration and reconstruction of the deposited atoms on the growth surface, and are not formed under thermal equilibrium conditions. Thus, one can control the introduction of these structures, i.e. fabrication of nearly perfectly ordered structures or complete elimination of them, by choosing appropriate growth conditions.
3. Device degradation In this section, degradation phenomena induced by the presence of defects are presented in terms of three
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CuPt
A
Chalcopyrite
A
(a) v
w
A
A
(b) v
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Fig. 9. Three types of ordered structure in f.c.c, lattice: (a) unit cells; (b) reciprocal lattice.
Fig. 10. High resolution TEM image of CuPt-type ordered structure in MOVPE-grown InGaP.
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major modes, namely rapid degradation, gradual degradation and catastrophic failure. On the basis of these results, methods for the elimination of device degradation are also discussed (see ref. 15 and references therein).
3.1. Rapid degradation Characteristic phenomena in rapid degradation are a sudden decrease in output power and the formation of non-radiative regions, i.e. "dark-line defects" (DLDs) or "dark-spot defects" (DSDs) in the active region. The half-life of such devices is less than 100 h at room temperature. This degradation is explained by recombination-enhanced dislocation climb (REDC) or recombination-enhanced dislocation glide (REDG). In this paper, only REDC is considered. Figure 11 shows a TEM image of dislocation dipoles associated with (100) DLDs in a rapidly degraded GaA1As/GaAs DH laser. The dipoles are of interstitial type with Burgers vectors of the type (a/2X011) inclined at 45 ° to the junction plane. The dipoles developed from a threading dislocation which propagated from the substrate during growth. Similar results are obtained in rapidly degraded GaAIAs/ GaAs DH LEDs. Dislocation dipoles also can grow from dislocation clusters originating at the heterointerface and from dislocation loops generated by diffusion. However, in InGaAsP/InP DH LEDs, rapid degradation owing to REDC is not observed. Although dark defects were observed in the fight-emitting region after 100 h of operation at room temperature, they were all shown to have been induced before operation, i.e. during growth or fabrication. Surprisingly, InGaAsP/ InGaP DH lasers, whose active region is also a quaternary layer lattice matched to GaAs, also degrade rapidly in the presence of dislocations, dislocation clusters and dislocation loops. As described above, the ease with which rapid degradation by REDC occurs depends on the material. To explain this effect, two models are most widely accepted, although many have been proposed. The first model is the "extrinsic defect model" shown in Fig. 12(a) [16]. Here, only one type of interstitial atom, e.g. Ga, is required for climb motion. In the second model, i.e. the "intrinsic defect model" (see Fig. 12(b))[17], the emission of two types of vacancies, e.g. Ga and As vacancies in the case of GaAs, is needed. Based on these models, one can suggest several candidates responsible for the REDC. First, since the non-radiative recombination event is thought to be caused at a defect by an energy transformation from a localized multiple-phonon mode excited by the absorbed light to a lattice vibrational mode, the REDC may be dominated by the bandgap energy. This can explain well the REDC in GaAs- and GaP-related materials, and the
Fig. 11. TEM image of dislocation dipoles associated with (100) DLDs in a rapidly degraded GaAIAs/GaAs DH laser.
Threading dislocation oe •
Threading dislocation
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Fig. 12. Two proposed models for dislocation climb in degraded optical devices: (a) extrinsic defect model; (b) intrinsic defect model.
difficulty in REDC in the InGaAsP system, though not for InP. The second candidate is a deep energy level and a non-radiative recombination rate associated with defects such as dangling bonds or native point defects. This Is supported by the fact that deep levels are generated in Ga(A1)As and InGaAsP on GaAs, where REDC occurs easily, and that they are not generated in InGaAsP on InP, where REDC does not occur easily. Moreover, since point defects must be absorbed to dislocations or emitted from the dislocation core in REDC, the energies required for their generation and migration are also the key parameters for REDC. More extensive experimental and theoretical studies are required to resolve this issue fully. However, from the results described above, it is clear that the generation of various defects, such as dislocations, dislocation clusters and dislocation loops, during growth must be suppressed to eliminate the rapid degradation.
3.2. Gradual degradation Typical characteristics of gradual degradation are a slow decrease in output power, the uniform darkening of the active region or the formation of DSDs, and an increase in deep levels in GaAIAs/GaAs optical devices. Gradual degradation of GaAIAs DH LEDs has been extensively studied. After operation for 7100 h at 169 °C and at 556 A cm-2, the optical power is uniformly decreased but no
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Materials-related reliability of lll- V optical devices
dark defects are found in the active region (see Fig. 13). TEM observations reveal that numerous microloops and point-like defects (possibly point defect clusters) are generated throughout the active region (see also Fig. 13). The loops are found to be interstitial-type Frank loops. Since they are not present before operation, one can assume that they result from the generation, migration and condensation of point defects via non-radiative recombination. Two deep levels A and B are present in the active region before operation. They are very commonly observed hole traps in LPE-grown GaAs. During operation, both deep levels undergo a two-step increase. Therefore, defect complexes associated with the deep levels could be the original source of the point defect reaction. Gradual degradation might proceed as follows: (i) non-radiative recombination occurs at some defects, which causes a point defect reaction and fresh point defect generation; (ii) the new defects also can act as non-radiative recombination centers (thus we have positive feedback for these two processes); (iii) the point defects generated migrate and condense at some nucleation centers; (iv) defect clusters and/or microloops are formed as byproducts. Gradual degradation can be intensified by stress, causing rather rapid degradation. Figure 14 illustrates the strain-induced degradation in a GaA1As/GaAs channel stripe planar (CSP) laser in the short wavelength range. Figure 14(a) shows a photoluminescence (PL) image of the active region after only 100 h of operation at a constant output power of 5 mW at 50 °C. Many dark lines and/or dark bands are
(a)
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observed along both edges of the stripe. A TEM image of defects corresponding to the dark defects is shown in Fig. 14(b). Point defect clusters and small dislocation loops are observed along both edges of the stripe. From calculations of the two-dimensional stress distribution in the active layer by the finite element method, it is found that stress concentrations are present on both edges of the stripe. In these regions, the stress parallel to the active layer is highly compressive and that perpendicular to the active layer is tensile. The cubic dilation of the active layer also was calculated from these results and we found that it becomes more compressive with a decrease in the lasing wavelength. From these results, it is concluded that, in GaA1As/GaAs visible lasers (when the lasing wavelength is shorter than 780 nm), and particularly in CSP-type lasers, the diode degraded unexpectedly rapidly, owing to strong enhancement of the gradual degradation effect by the internal stress. In InGaAsP/InP optical devices, since gradual degradation was absent, a recombination-enhanced point defect reaction cannot occur as easily. Consequently, one can conclude that, to eliminate the gradual degradation effect, tight control of the stoichiometry (to reduce the deep level population), careful lattice matching in the heterostructure and avoidance of stress introduction during fabrication are required. 3.3.
Catastrophic failure
This phenomenon takes place accidentally by a current surge, i.e. at high power density in lasers. The degradation occurs predominantly at the mirror
(b) D
(c)
(d) 150
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Fig. 13. An electroluminescence (EL) image of gradually degraded GaA1As DH LED and TEM images of regions (a)-(d) in the EL image.
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Materials-related reliability O/VII- V optical devices
v~ Ih i
2.4
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....
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Fig. 14. Strain-induced degradation of GaAIAs/GaAs CSP-type laser (2 = 770 nm): (a) PL image of degraded region; (b) TEM image of defect structure.
surface by "catastrophic optical damage" (COD). The COD also can be caused by strong optical excitation. In PL images of catastrophically degraded GaAIAs/GaAs DH lasers, one or more (110) DLDs generated at the mirror surface are observed. From TEM studies, these are seen to correspond to arrays of dislocation networks or multiple dislocation loops connected by dark knots (see Fig. 15). The proposed mechanism for COD is as follows: (i) when the output power density reaches a critical value, strong optical absorption occurs at the mirror surface, leading to local heating of the crystal; (ii) the bandgap energy shrinks in this region, causing further optical absorption; (iii) these processes give rise to very rapid thermal runaway and, finally, the melting of the crystal; (iv) subsequent rapid cooling generates residual defects in the degraded area. In the case of pulse operation, these processes repeat and the molten region propagates from the mirror to the inner crystal. In InGaAsP/InGaP D H lasers lattice matched to GaAs and emitting in the short wavelength range, we also have found COD during operation at high output power density. However, in InGaAsP/InP DH lasers, there have been no reports of COD events for any operating conditions. Thus, one can assume that the COD level, i.e. the critical output power density for causing COD, in this laser is very high. Since COD
occurs initially owing to optical absorption, the difference in the COD level in different materials is explained by different surface recombination velocities. Although COD occurs easily in the absence of defects, we have observed occasionally another COD event associated with a grown-in defect. In this mode, {110) DLDs are also observed. However, they do not originate from the mirror but are generated inside the active region. They often appear as short (110) DLDs. From T E M observation, it has been found that defects corresponding to the (1 i 0 ) D L D s are similar to those generated in degraded diodes by normal COD but develop from an inclusion which is generated during growth (see Figs. l(d) and 2). The COD level in this case is 0.86-1.4 MW cm-2, which is lower than that for the normal COD mode (1.6-3.1 MW cm-2). These results lead us to the conclusion that the protection of mirrors with dielectric films (SiO2 etc.) and the reduction of inclusions and/or precipitates during growth are essential for eliminating catastrophic failure in lasers. 3.4. Effect of modulated and ordered structures As described previously, a modulated structure is associated with quasi-periodic compositional variation, presumably owing to spinodal decomposition of the
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is lower and the PL peak energies have normal values. The ordered structures themselves are monolayer superlattices, e.g. of (InP/GaP) on the (1 11) plane in the CuPt structure. Thus, provided that the dislocations themselves do not act as non-radiative recombination centers, one can expect that ordered structures do protect dislocations from REDC, suppressing rapid degradation. However, such structures may involve lattice strain and it is also possible that stress is associated with the boundary between ordered and non-ordered regions and with antiphase boundaries which are often observed in the ordered region. These strains affect gradual degradation, i.e. by generation of point defects through ~recombinationenhanced local disordering". The validity of these predictions will be determined in the future by the study of degradation in InGa(A1)P visible lasers.
4. Conclusions
pO
bj Fig. 15. EL image of {110) DLDs in a catastrophically degraded GaA1As DH laser. (b) TEM image of dislocation network corresponding to the ~1 10) DLDs.
system. However, the compositional fluctuation measured by energy-dispersive X-ray analysis is 2%-3% and the domains are columnar shaped and quite stable. Thus, they do not affect strongly the optical and electrical device properties. Furthermore, in optical devices, the periodicity and amplitude of modulation do not change following degradation by any of the modes discussed earlier. However, dislocation loops formed as byproducts of modulated structure generation can cause rapid degradation, and point defects associated with composition modulation can lead to gradual degradation in some materials. However, ordered structures only affect the PL peak energies in strongly ordered InGa(AI)P crystals grown by MOVPE. In other alloy semiconductors, such as InGaAs, InAIAs and GaAsSb, the degree of ordering
The current status of the evaluation of defects in various III-V alloy semiconductors by TEM and the understanding of degradation modes in optical devices have been reviewed. The detailed nature of grown-in defects, and modulated and ordered structures are described. The roles of these defects as well as point defects and their complexes in the degradation of optical devices are presented in terms of three major degradation modes. Based on these results, techniques for eliminating the degradation problem are discussed.
Acknowledgments The author would like to express his gratitude to H. Ishikawa and O. Otsuki for their encouragement throughout this work. He also thanks H. Imai, S. Komiya, S. Yamakoshi, S. Isozumi, K. Kondo and other colleagues in Fujitsu Laboratories and Fujitsu Limited for valuable discussion.
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13 O. Ueda, S. Komiya and S. lsozumi, Jpn. J. Appl. Phys., 23 (1984) L241. 14 T. S. Kuan, T. F. Kuech, W. I. Wang and E. L. Wilkie, Phys. Rev. Lett., 54 (1985) 208. H. Nakayama and H. Fujita, Inst. Phys. Conf. Ser., 79(1986) 289. O. Ueda, M. Takikawa, J. Komeno and I. Umebu, dpn. J. Appl. Phys., 26(1987) L1824. 15 O. Ueda, J. Electrochem. Soc., 135(1988) llC. 16 P. M. Petroff and L. C. Kimerling, J. AppI. Phys., 29 (1976) 461. 17 S. O'Hara, P. W. Hutschinson and P. S. Dobson, Appl. Phys. Lett., 30(1977) 368.
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Materials-related reliability aspects of III-V optical devices O. Ueda Fujitsu Laboratories Ltd., 10-1 Morinosato-Wakamiya, Atsugi 243-01 (Japan)
Abstract This paper describes the current understanding of materials issues in III-V compound semiconductor heterostructures,
and the degradation of optical devices, lasers and light-emitting diodes fabricated from these materials. Generation of defects and thermal instability are among these issues for these systems. Defects introduced during crystal growth are classified into two types: interface defects and bulk defects. Interface-type defects are stacking faults, V-shaped dislocations, dislocation clusters, microtwins, inclusions and misfit dislocations. The bulk defect group includes precipitates and dislocation loops. Defects in the substrate also can be propagated into the epilayer. Structural imperfections owing to thermal instability are also found. They are quasi-periodic modulated structures, owing to spinodal decomposition of the crystal either at the liquid-solid interface or growth surface, and owing to atomic ordering, which also occurs on the growth surface through the migration and reconstruction of the deposited atoms. In each phenomenon, the nature and generation mechanism of defects and structures are clarified. Three major degradation modes--rapid degradation, gradual degradation and catastrophic failure--are discussed. Rapid degradation occurs via either recombinationenhanced dislocation climb or glide. The ease with which these phenomena occur in different heterostructures are presented. Based on the results, the dominant parameters involved in the phenomena are discussed. Gradual degradation takes place presumably owing to recombination-enhanced point defect reactions in GaA1As/GaAs-based optical devices. This mode is also enhanced by the internal stress owing to lattice mismatch. However, we do not observe such degradation in InGaAsP/InP-based optical devices. Catastrophic failure is found to be due to catastrophic optical damage at a mirror or at a defect in GaA1As/GaAs double-heterostrncture (DH) lasers but not in InGaAsP/InP DH lasers. In each degradation mode, the effect of the defects and/or structures on the degradation is discussed and methods for the elimination of degradation are proposed.
1. Introduction Since the discovery of semiconductor lasers in the early 1970s [1], a variety of optical devices have been developed from III-V alloy semiconductors. Current fields of application include fiber optical communication systems, digital audio systems and optical printers. In the development of these devices, high reliability has been one of the most important goals. Degradation is a major hindrance to reliability. Moreover, it has been clarified that defects in crystals very often cause device degradation. In this paper, we describe materials issues in III-V alloy semiconductors, i.e. the generation of defects during growth and phenomena induced by the thermal instability of the material, and our current understanding of degradation in III-V semiconductor lasers and light-emitting diodes (LEDs) in terms of three major degradation modes--rapid degradation, gradual degradation and catastrophic failure. For each mode, we discuss the role of defects in the degradation and the elimination of degradation. 0921-5107/93/$6.00
2. Materials issues in III-V alloy semiconductor thin films Materials issues in the growth of III-V alloy semiconductor thin films are focused on. These include the generation of various kinds of defects during growth, the generation of modulated structures owing to spinodal decomposition and atomic ordering associated with growth kinetics on the surface.
2.1. Defect generation Defects introduced during crystal growth are classified into two types, namely interface defects and bulk defects. Defects of the interface type are stacking faults, V-shaped dislocations, microtwins, inclusions and misfit dislocations.
2.1.1. Stacking faults Stacking faults are generated from the heterointerface, extending on the four equivalent {111} planes, forming a stacking fault tetrahedron (see Fig. l(a)). Figure 2(a) illustrates a typical optical micrograph of © 1993 - Elsevier Sequoia. All rights reserved
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(a)
Materials-related refiability o1711- V optical devices
(b)
U (c)
i~iiiii~
(d)
Fig. 1. Interface defects resultingfrom contamination,thermal damage of the substrate or local segregation. the surface of a GaAsSb crystal grown at 590 °C. A high density of surface defects are observed. They are hillock-like defects lying along the (110) direction and are very similar to the oval defects in molecular beam epitaxy (MBE) Ga(A1)As crystals [2]. The defect structure observed by transmission electron microscopy (TEM) at these surface defects is shown in Fig. 2(b) [3]. Stacking faults lying on the four equivalent {111} planes are seen. Since the thin foil includes only the surface region of the epilayer, one can expect that the surface defects correspond to stacking fault tetrahedra. The stacking faults often multiplied, as denoted by X in Fig. 2(b). In addition, threading dislocations and small stacking faults are found inside the tetrahedra. By considering the length of the stacking faults at the surface, their basal {111} planes and the epilayer thickness, it is deduced that these defects originated at the heterointerface. The density of these defects can be reduced by lowering the growth temperature and cleaning the substrate with an appropriate As overpressure. Therefore, the dominant factors affecting the defect generation could be the evaporation of P atoms from the substrate and/or the existence of some oxide or contamination on the substrate surface. 2.1.2. V-shaped dislocations Pair pits are often observed on the surface of InGaP crystals etched with HF + HNO 3. In the TEM image of the area including pair pits, small dislocation segments are observed, which correspond to the etch pits. From the shape of the dislocations, this dislocation pair is
assumed to be of V-shaped type (see Fig. l(b)). We also have found preferential vapor etching at the V-shaped dislocations during growth [4].
Z 1.3. Microtwins Microtwins are occasionally observed in epilayers grown at higher growth temperatures. They also originate from the heterointerface, extending to the surface area with {111} boundaries (see Fig. l(c)). Some of them terminate in midlayer, with {211} boundaries. A high resolution TEM image of a GaAsSb crystal grown at 590 °C from a region close to the heterointerface is shown in Fig. 3. Brighter regions along the interface are probably due to the existence of some oxides. In this micrograph, microtwins (MTs) are generated in two regions, i.e. I and II. In both regions, amorphous-like areas (A) and shallow, pit-shaped patterns (SP) are observed at the interface just below the amorphous-like area. Microtwins are generated only in the region adjacent to the areas. Based on these results, the following mechanism is suggested for the generation of microtwins: (1) evaporation of P atoms from the substrate surface (formation of In-rich droplets on the surface); (2) generation of amorphous areas as a consequence of an In-rich melt; (3) growth of crystals with different lattice constants around the amorphous areas; (4) generation of microtwins by the relaxation of the stress resulting from the lattice mismatch between the normal GaAsSb areas and those formed in step (3).
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Materials-related reliability of lll- V optic al devices
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Thus, the elimination of microtwins can be achieved by methods similar to those described for stacking faults.
2.1.4. Inclusions A typical TEM image of inclusion defects in a liquid phase epitaxy(LPE)-grown InGaAsP layer on a (001) GaAs substrate is shown in Fig. 4 [5]. The inclusions are lying in the (100) direction and, from stereoscopic observation, they are found to originate from the interface between the InGaAsP active layer and the InGaP cladding layer. These defects are thought to be formed by the development of some inclusions from the interface in the direction (01 n).
2.1.5. Misfit dislocations
b
•........
o.s, vm
Fig. 2. (a) Optical micrograph of defects revealed on the surface of an MBE-grown GaAsSb crystal on a (001) InP substrate. (b) (001) plan view, bright-field TEM image of a stacking fault tetrahedron corresponding to the surface defect shown in Fig. 2(a).
If there is a difference in the between the lattice constant of the epilayer and that of the substrate, then misfit dislocations are generated when the film thickness exceeds the critical thickness tc, which depends on the structure and elastic properties of the materials and on the type of dislocations. There are several possible mechanisms for the generation of misfit dislocations, as shown in Fig. 5. First, threading dislocations propagated from the substrate may glide out (Fig. 5(a)) or bend out (Fig. 5(b)) to the edge of the wafer. It is also expected that two dislocations with the same Burgers vectors in the substrate terminate into one misfit dislocation at the heterointerface, forming a half-loop (righthand figure of Fig. 5(c)). If the dislocation density in the substrate is low, another mechanism is expected, as follows: (i) nucleation of microhalf-loops at the surface; (ii) expansion of loops down to the interface; (iii) gliding out of the threading segments to the edge (left-hand figure of Fig. 5(c)) [6]. Since the first three kinds of defects described above are eliminated by cleaning of the substrate surface prior to growth and/or the protection of the substrate against evaporation of substrate atoms, they are expected to be generated by contamination and/or
M, aAsSb
~P
20 nm Fig. 3. Low magnification (110) cross-sectional high resolution TEM image of the GaAsSb-(00 l)InP interface.
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Materials-related reliabili~ of lll- V optical devices
thermal damage to the substrate. The inclusions are presumably caused by local segregation of solute atoms at the interface. In highly mismatched heteroepitaxy, microtwins (see Fig. 5(d)) and/or stacking faults are very often generated, particularly during the initial stage of growth [7]. They propagate along the four equivalent {111 }planes. However, they often annihilate in midlayer, i.e. terminating in {211 }-type boundaries for microtwins and coalescence of two stacking faults on different planes. 2.1.6. Bulk defects Bulk defects are classified into two types, i.e. defects propagated from the substrate and those generated by local segregation of dopant atoms or native point defects. Since the defects propagated from the substrate are simply eliminated by using substrates with low defect densities, we only consider the defects of the
Fig. 4. Plan view TEM image of inclusions in LPE-grown InGaAsP on a (001 ) GaAs substrate.
--/-- _/__-_
.
.
.
.
a)
(b)
(c)
(d)
.
Fig. 5. Interface defects resulting from lattice mismatch.
other type. Typical features of these defects are shown schematically in Fig. 6. In most cases, they are generated in heavily doped crystals, as shown in Figs. 6(a)-6(c). When supersaturated dopants are present in the crystal, they tend to segregate locally during growth or the cooling process after growth, forming precipitates (see Fig. 6(a)). Figure 7 illustrates a plan view, high resolution TEM image of one of the uniformly distributed precipitates in heavily doped InP with Fe grown by metallo-organic vapor phase epitaxy (MOVPE) [8]. The precipitates are found to be spherical FeP particles nearly lattice matched to the matrix crystal. When there is a large lattice mismatch between the precipitates and the matrix, dislocation loops are often punched out from the precipitates by the relaxation of stress (see Fig. 6(b)). Such precipitates are often observed in heavily doped LPE-grown InGaAsP crystals on InP substrates [9]. In the formation of the precipitates, excess native point defects can be introduced. They may condense at some nucleation centers to form dislocation loops (either faulted or unfaulted), as shown in Fig. 6(c). Interstitial-type dislocation loops are often observed in heavily doped LPE-grown GaA1As crystals with Ge [10]. However, dislocation loops are present even in undoped crystals. We have found both vacancy-type and interstitial-type loops in undoped LPE-InGaP crystals [11]. Since modulated structures (see the next section) are generated in the crystal, these loops are presumably caused by the condensation of excess point defects induced as byproducts of modulated structure generation.
O. Ueda
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(b)
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Materials-related refiabifity of ill- V optical devices
/
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@
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Fig. 6. Bulk defects resulting from local segregation of dopant atoms or native point defects.
nm Fig. 7. FeP precipitate in MOVPE-grown InP heavily doped with Fe.
2.2. Structural imperfections owing to thermal instability of the crystals 2.2.1. Modulated structures Modulated structures are generated in alloy semiconductors such as InGaAs, InGaP, InGaAsP and GaAsSb, whose compositions are inside the "spinodal region" [12]. They are observed as quasi-periodic diffraction contrast in bright- or dark-field images under two-beam or multibeam conditions. Figure 8 shows a plan view, bright-field TEM image of modulated structures in LPE-grown InGaAsP on (001) GaAs (2pe=730 nm) [13]. They develop in the two equivalent directions of (100) and (010) and are colum-
nar shaped in the (001) growth direction. In spinodally decomposed alloys, which have asymmetry in their elastic coefficients, modulated waves extend so as to minimize the strain contribution to the free energy in the solid solution, i.e. the modulated structures are induced in the direction [hkl], which minimizes the elastic coefficient Yhkl. From calculations of the values Yl00, YII0, and Ylll for GaAs, GaP, InAs and InP, it has been found that in all cases, Y100< Yi l0 < Ylll [13]. This can explain well the preferential direction of elongation of the structures. It is also assumed that they are very stable once they are formed, propagating in the (001) growth direction. Furthermore, it is clarified that both structures with long (50-200 nm) and short (5-20 nm) periodicities are formed when the composition of the crystal is inside the spinodal region (Fig. 8(a)), whereas only the structures with short periodicities are formed when the composition is outside the spinodal region (Fig. 8(b)). Thus, one can expect that the structures with long periodicities are formed during growth by spinodal decomposition and that the short periodicity structures form during the cooling process after growth. Moreover, the compositional fluctuations along the structure measured by energy-dispersive X-ray analysis are in the range 2%-3%, which is far smaller than that expected. This may be due to the fact that, in semiconductors, atomic diffusion in the bulk or on the surface is very slow compared with the case of metal alloys. The generation of modulated structures does not depend on the growth method. The decompostion process may be enhanced by atomic diffusion at the liquid-solid interface for LPE growth and at the growth surface for MOVPE, VPE and MBE.
14
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Materials-related reliability of Ill- V optical devices Z2.2. Ordered structures
Fig. 8. Modulated structures in LPE-grown InGaAsP. (a) 2pL= 730 nm; (b)),PL= 800 nm. From the results described above, it is suggested that modulated structures can be eliminated by the growth of the crystal at temperatures where the composition of the crystal is outside the spinodal region and subsequent rapid cooling or quenching of the crystal.
Recently, ordered structures (or natural superlattices), in which two kinds of element atoms in column III or V lattice periodically arrange, have been observed in various II1-V alloy semiconductors [14]. In most cases, the ordered structures are generated locally in the crystal and the degree of ordering is strongly affected by the growth method and growth conditions. It is known that there are three types of substitutionally ordered structures in f.c.c, lattice of AB-type binary alloy systems. They are the types of CuAu-I (L10), CuPt (Lll) and chalcopyrite (Ell), whose atomic arrangements and corresponding reciprocal lattices are shown schematically in Fig. 9. The results for ordered structures in III-V alloy semi conductors obtained until now can be summarized as follows: (a) the most common structure in crystals grown on a (001) substrate is CuPt type (see Fig. 10) and other structures are only occasionally observed; (b) ordered structures are generated very often in crystals grown by MOVPE, VPE and MBE, whereas no ordering takes place in LPE-grown crystals (except for one report for ordered structures in InGaAs by Nakayama et al. [14], which has not been confirmed by others); (c) regarding CuPt-type structures, one can observe only two variants in the (110) cross-section (atomic steps on the growth surface are believed to play an important role in the generation of the ordered structures, since one of the two variants is preferentially enhanced when substrates tilted toward the (110) directions are used); (d) the ordering is not perfect and the ordered regions are plate-like microdomains lying on planes nearly parallel to the growth surface; (e) defects, such as antiphase boundaries, are often generated in the ordered regions; (f) the degree of ordering depends on the growth temperature, V/III partial pressure ratio (in the case of MOVPE and MBE) and rotation velocity of the substrate; (g) only strongly ordered InGaP crystals grown by MOVPE exhibit abnormal bandgap energies (up to 50 meV lower than the normal value). From these results, it is now understood that the ordered structures are generated by the migration and reconstruction of the deposited atoms on the growth surface, and are not formed under thermal equilibrium conditions. Thus, one can control the introduction of these structures, i.e. fabrication of nearly perfectly ordered structures or complete elimination of them, by choosing appropriate growth conditions.
3. Device degradation In this section, degradation phenomena induced by the presence of defects are presented in terms of three
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Materials-related reliabilio, of lll- V optical devices
CuPt
A
Chalcopyrite
A
(a) v
w
A
A
(b) v
w
Fig. 9. Three types of ordered structure in f.c.c, lattice: (a) unit cells; (b) reciprocal lattice.
Fig. 10. High resolution TEM image of CuPt-type ordered structure in MOVPE-grown InGaP.
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Materials-related reliabilio, ~(HI- Voptical devices
major modes, namely rapid degradation, gradual degradation and catastrophic failure. On the basis of these results, methods for the elimination of device degradation are also discussed (see ref. 15 and references therein).
3.1. Rapid degradation Characteristic phenomena in rapid degradation are a sudden decrease in output power and the formation of non-radiative regions, i.e. "dark-line defects" (DLDs) or "dark-spot defects" (DSDs) in the active region. The half-life of such devices is less than 100 h at room temperature. This degradation is explained by recombination-enhanced dislocation climb (REDC) or recombination-enhanced dislocation glide (REDG). In this paper, only REDC is considered. Figure 11 shows a TEM image of dislocation dipoles associated with (100) DLDs in a rapidly degraded GaA1As/GaAs DH laser. The dipoles are of interstitial type with Burgers vectors of the type (a/2X011) inclined at 45 ° to the junction plane. The dipoles developed from a threading dislocation which propagated from the substrate during growth. Similar results are obtained in rapidly degraded GaAIAs/ GaAs DH LEDs. Dislocation dipoles also can grow from dislocation clusters originating at the heterointerface and from dislocation loops generated by diffusion. However, in InGaAsP/InP DH LEDs, rapid degradation owing to REDC is not observed. Although dark defects were observed in the fight-emitting region after 100 h of operation at room temperature, they were all shown to have been induced before operation, i.e. during growth or fabrication. Surprisingly, InGaAsP/ InGaP DH lasers, whose active region is also a quaternary layer lattice matched to GaAs, also degrade rapidly in the presence of dislocations, dislocation clusters and dislocation loops. As described above, the ease with which rapid degradation by REDC occurs depends on the material. To explain this effect, two models are most widely accepted, although many have been proposed. The first model is the "extrinsic defect model" shown in Fig. 12(a) [16]. Here, only one type of interstitial atom, e.g. Ga, is required for climb motion. In the second model, i.e. the "intrinsic defect model" (see Fig. 12(b))[17], the emission of two types of vacancies, e.g. Ga and As vacancies in the case of GaAs, is needed. Based on these models, one can suggest several candidates responsible for the REDC. First, since the non-radiative recombination event is thought to be caused at a defect by an energy transformation from a localized multiple-phonon mode excited by the absorbed light to a lattice vibrational mode, the REDC may be dominated by the bandgap energy. This can explain well the REDC in GaAs- and GaP-related materials, and the
Fig. 11. TEM image of dislocation dipoles associated with (100) DLDs in a rapidly degraded GaAIAs/GaAs DH laser.
Threading dislocation oe •
Threading dislocation
/D
o
, ~ Do•
•
~
(a)
~
\Interstitial
(b)
.
o
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,~ ~ Growth
~ .~ '°% ~.~ Growth % Vacancy
Fig. 12. Two proposed models for dislocation climb in degraded optical devices: (a) extrinsic defect model; (b) intrinsic defect model.
difficulty in REDC in the InGaAsP system, though not for InP. The second candidate is a deep energy level and a non-radiative recombination rate associated with defects such as dangling bonds or native point defects. This Is supported by the fact that deep levels are generated in Ga(A1)As and InGaAsP on GaAs, where REDC occurs easily, and that they are not generated in InGaAsP on InP, where REDC does not occur easily. Moreover, since point defects must be absorbed to dislocations or emitted from the dislocation core in REDC, the energies required for their generation and migration are also the key parameters for REDC. More extensive experimental and theoretical studies are required to resolve this issue fully. However, from the results described above, it is clear that the generation of various defects, such as dislocations, dislocation clusters and dislocation loops, during growth must be suppressed to eliminate the rapid degradation.
3.2. Gradual degradation Typical characteristics of gradual degradation are a slow decrease in output power, the uniform darkening of the active region or the formation of DSDs, and an increase in deep levels in GaAIAs/GaAs optical devices. Gradual degradation of GaAIAs DH LEDs has been extensively studied. After operation for 7100 h at 169 °C and at 556 A cm-2, the optical power is uniformly decreased but no
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Materials-related reliability of lll- V optical devices
dark defects are found in the active region (see Fig. 13). TEM observations reveal that numerous microloops and point-like defects (possibly point defect clusters) are generated throughout the active region (see also Fig. 13). The loops are found to be interstitial-type Frank loops. Since they are not present before operation, one can assume that they result from the generation, migration and condensation of point defects via non-radiative recombination. Two deep levels A and B are present in the active region before operation. They are very commonly observed hole traps in LPE-grown GaAs. During operation, both deep levels undergo a two-step increase. Therefore, defect complexes associated with the deep levels could be the original source of the point defect reaction. Gradual degradation might proceed as follows: (i) non-radiative recombination occurs at some defects, which causes a point defect reaction and fresh point defect generation; (ii) the new defects also can act as non-radiative recombination centers (thus we have positive feedback for these two processes); (iii) the point defects generated migrate and condense at some nucleation centers; (iv) defect clusters and/or microloops are formed as byproducts. Gradual degradation can be intensified by stress, causing rather rapid degradation. Figure 14 illustrates the strain-induced degradation in a GaA1As/GaAs channel stripe planar (CSP) laser in the short wavelength range. Figure 14(a) shows a photoluminescence (PL) image of the active region after only 100 h of operation at a constant output power of 5 mW at 50 °C. Many dark lines and/or dark bands are
(a)
17
observed along both edges of the stripe. A TEM image of defects corresponding to the dark defects is shown in Fig. 14(b). Point defect clusters and small dislocation loops are observed along both edges of the stripe. From calculations of the two-dimensional stress distribution in the active layer by the finite element method, it is found that stress concentrations are present on both edges of the stripe. In these regions, the stress parallel to the active layer is highly compressive and that perpendicular to the active layer is tensile. The cubic dilation of the active layer also was calculated from these results and we found that it becomes more compressive with a decrease in the lasing wavelength. From these results, it is concluded that, in GaA1As/GaAs visible lasers (when the lasing wavelength is shorter than 780 nm), and particularly in CSP-type lasers, the diode degraded unexpectedly rapidly, owing to strong enhancement of the gradual degradation effect by the internal stress. In InGaAsP/InP optical devices, since gradual degradation was absent, a recombination-enhanced point defect reaction cannot occur as easily. Consequently, one can conclude that, to eliminate the gradual degradation effect, tight control of the stoichiometry (to reduce the deep level population), careful lattice matching in the heterostructure and avoidance of stress introduction during fabrication are required. 3.3.
Catastrophic failure
This phenomenon takes place accidentally by a current surge, i.e. at high power density in lasers. The degradation occurs predominantly at the mirror
(b) D
(c)
(d) 150
IJm
0.5 Hm It
!
Fig. 13. An electroluminescence (EL) image of gradually degraded GaA1As DH LED and TEM images of regions (a)-(d) in the EL image.
18
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m"
Materials-related reliability O/VII- V optical devices
v~ Ih i
2.4
"5 p m
....
....
Fig. 14. Strain-induced degradation of GaAIAs/GaAs CSP-type laser (2 = 770 nm): (a) PL image of degraded region; (b) TEM image of defect structure.
surface by "catastrophic optical damage" (COD). The COD also can be caused by strong optical excitation. In PL images of catastrophically degraded GaAIAs/GaAs DH lasers, one or more (110) DLDs generated at the mirror surface are observed. From TEM studies, these are seen to correspond to arrays of dislocation networks or multiple dislocation loops connected by dark knots (see Fig. 15). The proposed mechanism for COD is as follows: (i) when the output power density reaches a critical value, strong optical absorption occurs at the mirror surface, leading to local heating of the crystal; (ii) the bandgap energy shrinks in this region, causing further optical absorption; (iii) these processes give rise to very rapid thermal runaway and, finally, the melting of the crystal; (iv) subsequent rapid cooling generates residual defects in the degraded area. In the case of pulse operation, these processes repeat and the molten region propagates from the mirror to the inner crystal. In InGaAsP/InGaP D H lasers lattice matched to GaAs and emitting in the short wavelength range, we also have found COD during operation at high output power density. However, in InGaAsP/InP DH lasers, there have been no reports of COD events for any operating conditions. Thus, one can assume that the COD level, i.e. the critical output power density for causing COD, in this laser is very high. Since COD
occurs initially owing to optical absorption, the difference in the COD level in different materials is explained by different surface recombination velocities. Although COD occurs easily in the absence of defects, we have observed occasionally another COD event associated with a grown-in defect. In this mode, {110) DLDs are also observed. However, they do not originate from the mirror but are generated inside the active region. They often appear as short (110) DLDs. From T E M observation, it has been found that defects corresponding to the (1 i 0 ) D L D s are similar to those generated in degraded diodes by normal COD but develop from an inclusion which is generated during growth (see Figs. l(d) and 2). The COD level in this case is 0.86-1.4 MW cm-2, which is lower than that for the normal COD mode (1.6-3.1 MW cm-2). These results lead us to the conclusion that the protection of mirrors with dielectric films (SiO2 etc.) and the reduction of inclusions and/or precipitates during growth are essential for eliminating catastrophic failure in lasers. 3.4. Effect of modulated and ordered structures As described previously, a modulated structure is associated with quasi-periodic compositional variation, presumably owing to spinodal decomposition of the
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is lower and the PL peak energies have normal values. The ordered structures themselves are monolayer superlattices, e.g. of (InP/GaP) on the (1 11) plane in the CuPt structure. Thus, provided that the dislocations themselves do not act as non-radiative recombination centers, one can expect that ordered structures do protect dislocations from REDC, suppressing rapid degradation. However, such structures may involve lattice strain and it is also possible that stress is associated with the boundary between ordered and non-ordered regions and with antiphase boundaries which are often observed in the ordered region. These strains affect gradual degradation, i.e. by generation of point defects through ~recombinationenhanced local disordering". The validity of these predictions will be determined in the future by the study of degradation in InGa(A1)P visible lasers.
4. Conclusions
pO
bj Fig. 15. EL image of {110) DLDs in a catastrophically degraded GaA1As DH laser. (b) TEM image of dislocation network corresponding to the ~1 10) DLDs.
system. However, the compositional fluctuation measured by energy-dispersive X-ray analysis is 2%-3% and the domains are columnar shaped and quite stable. Thus, they do not affect strongly the optical and electrical device properties. Furthermore, in optical devices, the periodicity and amplitude of modulation do not change following degradation by any of the modes discussed earlier. However, dislocation loops formed as byproducts of modulated structure generation can cause rapid degradation, and point defects associated with composition modulation can lead to gradual degradation in some materials. However, ordered structures only affect the PL peak energies in strongly ordered InGa(AI)P crystals grown by MOVPE. In other alloy semiconductors, such as InGaAs, InAIAs and GaAsSb, the degree of ordering
The current status of the evaluation of defects in various III-V alloy semiconductors by TEM and the understanding of degradation modes in optical devices have been reviewed. The detailed nature of grown-in defects, and modulated and ordered structures are described. The roles of these defects as well as point defects and their complexes in the degradation of optical devices are presented in terms of three major degradation modes. Based on these results, techniques for eliminating the degradation problem are discussed.
Acknowledgments The author would like to express his gratitude to H. Ishikawa and O. Otsuki for their encouragement throughout this work. He also thanks H. Imai, S. Komiya, S. Yamakoshi, S. Isozumi, K. Kondo and other colleagues in Fujitsu Laboratories and Fujitsu Limited for valuable discussion.
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Materials-related reliability of lll- V optical devices
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