Damage accumulation and amorphization in GaAs-AlGaAs structures

Nuclear Instruments and Methods in Physics Research B74 (1993) 80-83 North-Holland

Beam Interactions with Materials&Atoms

Damage accumulation and amorphization in GaAs-AlGaAs structures J.S. Williams ‘, C. Jagadish, A. Clark. G. Li and C.A. Larsen Department of Electronic Materials Engineering, Research School of Physical Sciences and Engineering, The Australian National University, GPO Box. 4, Canberra, ACT 2601 Australia

The build up of keV to MeV ion damage in MOCVD grown GaA+AIGaAs structures has been studied by ion channeling. Increasing dynamic annealing and resistance to damage is observed in thick AlGaAs layers of increasing Al content for implantation at liquid nitrogen temperatures. For GaAs-AlGaAs multilayers, the proximity of adjacent AlGaAs layers is found to have a strong retarding effect on amorphization in GaAs, presumably as a result of defects migrating into the GaAs layers from AlGaAs. The rate of energy deposition has also a strong influence on damage accumulation in multilayers. A knowledge of such processes is crucial in determining optimum conditions for ion beam isolation of device structures.

(2 3000 A) ti,Ga,_,As

1. Introduction

Ion implantation of III-V multilayer structures has recently found technological interest as a means of locally disordering superlattice structures [l] and electrically and optically isolating devices [2]. For such applications, it is important to have a knowledge of the damage accumulation processes in different III-V materials, particularly in cases where dynamic annealing during implantation [3] can result in considerable resistance to damage for some layers. For example, it has been shown that the separate layers of GaAs-AlAs heterostructures damage very differently under ion irradiation at liquid nitrogen temperature [4,5], whereby GaAs is rendered amorphous but AlAs layers remain crystalline. Furthermore, GaAs material in layers adjacent to AlAs was found to resist damage, presumably as a result of mobile defects migrating across the interface [5]. In this study, we have examined the build up of damage in thick AlGaAs layers and GaAs-AlGaAs multilayers of different composition under irradiation with keV and MeV Si ions at different dose rates and substrate temperatures. In this paper, we concentrate on implantation at constant dose rate into substrates held at liquid nitrogen temperature.

2. Experimental GaAs-AlGaAs structures were purpose-grown for this study on the ANU MOCVD reactor. Both thick 1 Also, Microelectronics and Materials Technology Centre, RMIT, Melbourne, 3000, Australia. 0168-583X/93/$06.00

layers on bulk GaAs and GaAs-Al,Ga,_,As n@tilayers (10 periods) of average thickness 300-450 A were grown. For the thick Al,Ga,_,As layers, x = 0, 0.3, 0.5, 0.7, 0.9 and 1 values were grown, and for the multilayers, x values were 0.3, 0.7, and 1. The surfaces ff the Al,Ga,_,As layers were terminated with N 30 A of GaAs to avoid oxidation of Al-rich layers on exposure to air. The samples were irradiated with 90 keV 28Si- ions or 1 or 2 MeV “Si+ ions from the ANU 1.7 MV tandem ion implanter. A range of doses was used at a dose rate of - 1 PA cmP2 with the samples held at liquid nitrogen temperatures. The samples were analysed by ion channeling using 2 MeV He+ ions with selected samples analysed with TEM.

3. Results and discussion 3.1. Damage

build up in bulk layers

Figs. 1 and 2 show typical Rutherford backscattering/channeling (RBS-C) spectra illustrating the build up of damage with dose for 90 keV Si- ions for three Al,Ga,_,As compositions. Fig. 1 shows the expected behaviour for GaAs, where disorder increases with dose until a buried amorphous layer is formed [3] at a dose of 4 x 1013 cmP2. For higher doses, the amorphous layer thickens to form a layer continuous to the surface. TEM was employed to confirm that an amorphous layer formed when the ion channeling signal reached the random level. Fig. 2 shows similar RBS-C spectra for damage build up in Al,,Ga,,As. In this case, the alloy is resistant to damage, compared with GaAs, and a buried amorphous layer forms (channeled

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J.S. Williams et al. / Damage in GaAs-AlGaAs structures

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10

0.0 0 350

300

Dose

Channel

Fig. 1. RBS-C spectra at glancing angle geometry (detector 10” to surface) illustrating damage build up in GaAs for 90 keV Si- ions at liquid nitrogen temperature and at a dose rate of - 1 PA cmm2. Virgin aligned spectrum (A ),Random (o), 8x 101’ cnm2 (+), 1 X 1Ol3 cmm2 (0), 2X 1013cmP2 (a), 4X1013cm~2(~),1X1014cm~2(~),and5X1014cm-2(r).

spectrum coincident with random level) at a dose between 8 x 1013 cm-’ and 1.4 x 1014cm-‘. Thus, the threshold dose for amorphization of this alloy is more than a factor of 2 higher than that for GaAs in fig. 1. This is despite the fact that TRIM calculations [6] indicate only 20% lower peak nuclear energy deposition density (vacancy production) for Al,,Ga,,,As compared with GaAs. This suggests, consistent with other recent studies [7], that AlGaAs layers

Energy 50

1

.o

*b.Pi&

1.2 I I 2500

(MeV)

1.4 I Depth

1.6 1 (A)

1500

500

1.8

I

2.0

I

h

40 -

I 0 350 360 Channel

Fig. 2. RBS-C spectra of Al,,,Ga,,5As as per fig. 1. 4 X 1Ol3 cmw2 (+), 8 x 1013 cmm2 (O), 1 X 1014 cue2 (a), and 3 X 1Ol4 cme2 (m).

(err?)

Fig. 3. Normalized disorder vs. dose for various Al,Ga,_,As compositions. x = 0 (O), x = 0.3 (III), x = 0.5 (A), X = 0.7 (0) x=0.9(*)andx=l.O(m).

of increasing Al contents are more resistant to damage for keV implantation at liquid nitrogen temperature. Indeed, Al,Ga,_,As layers with x 2 0.9 could not be amorphized for 90 keV Si irradiation to doses up to 5 x 1015 cm-’ at liquid nitrogen temperature. In fig. 3, we show the normalized disorder as a function of dose for 90 kev Si ions for each of the compositions studied. The disorder level (1.0) corresponds to the point where the channeled spectrum is coincident with the random level at the depth of the peak in nuclear energy deposition. As the concentration of Al in the AlGaAs layer increases, so does the resistance to damage for liquid nitrogen temperature implantation. Indeed, up to doses of 5 X 1015 cm-‘, the Al,,Ga,,,As and the AlAs layers did not amorphize, but rather exhibited a higher dechanneling level in RBS-C spectra, compared with a virgin spectrum, which is consistent with accumulation of defect complexes and extended defects in crystalline material [8]. Furthermore, the nature of damage build up is very different in AlGaAs compared with GaAs. For example, in GaAs, the disorder builds up in a manner consistent with the additive accumulation of damage from individual ion tracks, leading ultimately to amorphization. Alternatively, for Al,.,Ga,.,As layers, the initial damage builds up very slowly with dose such that, by 1.4 X 1014 cm-‘, little disorder is detected by RBS-C, despite the fact that every atom has been displaced several times. However, at 3 X 1014 cm-’ a thick amorphous layer is produced. We suggest that, when the level of accumulated defect clusters and extended defects reaches a critical concentration, the highly defective layer either “collapses” to an amorphous phase upon subsequent bombardment or preferIV. MATERIALS SCIENCE

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entially nucleates an amorphous phase via further defect accumulation. This behaviour will be illustrated further in a subsequent publication [9]. Before moving to the multilayer data, we comment on the trends observed with changing implantation parameters. Very strong dose rate effects were observed with AlGaAs layers, whereby decreasing the dose rate decreased the measured damage for the same dose. This is consistent with previous GaAs data at room temperature [lo], whereby slower damage production rates favour dynamic annealing and reduced damage. Indeed, raising the substrate temperature also enhanced dynamic annealing, making it more difficult to amorphize AlGaAs. At 100°C for example, it was not possible to amorphize Al,,Ga,,As, even at doses exceeding 5 X 10”’ cm-‘. Furthermore, increasing the ion energy, increased the threshold dose for amorphization, consistent with previous observations [7] and with the lower average energy deposition densities per ion. These data will be presented in detail elsewhere 191.

structures Energy

1.2 ,

50 40 -

d

0 250

1.4 I

(MeV) 1.8 I

1.6 /

2.0 /

AlGaAs AIGaAs AL P*.3AS ’ G;As 4 G;As J G;As

/

300

dl_.........,....i 400

350

450

Channel

Fig. 4. RBS-C spectra for an A&Ga,,,As-GaAs

multilayer under conditions of fig. 1. 6X 1013 cm-’ (+), 2 x 1014 crne2 (0) and 1 x 101’ cm-’ ( n ).

3.2. Damage in multilayem In general, the presence of adjacent AlGaAs layers had a substantial effect on the build up of damage in GaAs for liquid nitrogen temperature implants. This effect was more apparent for MeV impl~tation but was also observed in the 90 keV case, as indicated in fig. 4. Fig. 4 shows RBS-C spectra depicting damage build up with dose for Ala,,Ga,,As-GaAs multilayers. The peak in the nuclear energy deposition is about 2/3 of the depth into the first AlGaAs layer. At a dose of 6 x 1Ol3 cm-‘, about Z/3 of the first GaAs layer and

Energy 1.2 50 a> ,

1.4

l/3 of the second GaAs layer are amorphized, but little disorder is observed in the AlGaAs layer. Note that the back edge of the first GaAs layer is not completely amorphous. This presumably arises from the pro~mi~ of the underlying AlGaAs layer and the migration of defects into the GaAs layer, as previously observed for GaAs-AlAs multilayers [4,5,7]. At a dose of 1 X 1015 cm-‘, the first two GaAs layers appear to be amorphous and the first AlGaAs layer is close to amorphous (spectrum height not quite to the random level),

Energy

(MeV) 1.6

1.8

(MeV)

2.0

/

L

/

10

360

Channel

350

Channel

Fig. 5. RBS-C spectra for multilayers irradiated with 2 MeV Sic ions. Other conditions as per fig. 1. (a) GaAs-AI,,Ga,,As: Virgin spectrnm (+>, Random (0) 2x 1Ol4 cm-’ (A) and 1 X 101’ cm-* (0). (b) GaAs-Al,,Ga,,,As: Virgin spectrum (A 1, Random (o), 2 X lOI cm-* (0) and 1 X 1015 cn-’ ( i- 1.

J.S. Williams et al. / Damage in Gds-AlGaAs

The proximity effects of AlGaAs layers on GaAs damage build up are more evident for MeV irradiation, where the nuclear energy deposition density per ion is lower (by a factor of 8 for 2 MeV Si). This is illustrated in fig. 5. For the AlO.,Ga,,,As-GaAs multilayer (fig. 5a), a dose of 1 X 101’ cm-‘, which should have resulted in sufficient integrated nuclear energy density to amorphize the first twd GaAs layers, compared with the 90 keV data, has not amorphized the GaAs and has produced little damage in AlGaAs. The irradiation of the Al,,,,Ga,sAs multilayer is even more striking, whereby the same 1 X 1015 cm-’ dose has resulted in little damage in both the AlGaAs and GaAs layers. Clearly, the influence of mobile defects from the adjacent AlGaAs layers is having a profound influence on damage produced in the GaAs layers. We speculate that the difference between the 90 keV and 2 MeV cases arises from the large difference in energy deposition density per ion in the two cases. For 2 MeV irradiation, the less dense disordered zones favour dynamic annealing in a manner analogous to the dose rate effect. Indeed, lower dose rates for 90 keV irradiation are also observed to favour dynamic annealing in adjacent GaAs layers. These energy deposition density and dose rate differences in dynamic annealing of AlGaAs and adjacent GaAs layers may indicate that discrete mobile defects (e.g. vacancies) on the Al/Ga sublattice may be the major cause of dynamic annealing.

4. Conclusions Consistent with previous studies [7], we have shown that AlGaAs layers of increasing Al content are more resistant to ion damage, even at liquid nitrogen temperatures. Damage accumulation in GaAs layers is reduced as a result of the proximity of AlGaAs layers in heterostructures. This effect is favoured at higher beam energies and lower dose rates, where dynamic

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annealing rates of discrete defects can dominate damage production rates. At liquid nitrogen temperatures, it is speculated that amo~hization of Al,Ga,_,As layers of high x (x 2 0.5) is initiated by the build up of crystalline defect clusters (and extended defects) which “nucleate” an amorphous phase at a critical defect level.

Acknowledgement

The ARC and the Special Research Centre’s Scheme is acknowledged for partial financial support of the RMIT part of this study.

References [l] B. Tell, B.C. Johnson, J.L. Zyskind, J.M. Brown, J.W. Sulhoff, K.F. Brown-Goebeler, B.I. Miller and U. Keren, Appl. Phys. Lett. 51 (1988) 1428.

121S.J. Pearton, F, Ren, J.R. Lothian, T.R. Fullowan, A. Katz, P.W. Wisk, C.R. Abernathy, R.F. Kopf, R.G. Elliman, M.C. Ridgway, C. Jagadish and J.S. Williams, J. Appl. Phys. 71 (1992) 4949. [3] J.S. Williams, Mat. Res. Sot. Bull. 17 (1992) 47. [4] A.G. Cullis, P.W. Smith, DC. Jacobson and J.M. Poate, J. Appl. Phys. 69 (1991) 1279. [5] D.J. Eaglesham, J.M. Poate, DC. Jacobson, M. Cerollo, L.N. Pfeiffer and K West, Appl. Phys. Lett. 58 (1991) 523. [6J J.P. Biersack and L.G. Haggmark, Nucl. Instr. and Meth. 174 (1980) 237. [7] A.G. Cullis, A. Polman, P.W. Smith, D.C. Jacobson, J.M. Poate and C.R. Whitehouse, Nucl. Instr. and Meth. in press. 181J.S. Williams and R.G. Elliman, in: Ion Beam for Materials Analysis, eds. J.S. Williams and J.R. Bird (Academic Press, Sydney, 1989), chapt. 6. [9] J.S. Williams, C. Jagadish, A. Clark, G. Li and C.A. Larsen, to be published. [lOI T.E. Haynes and 0-W. Holland, Appl. Phys. Lett. 58 (3991) 62.

IV. MATERIALS SCIENCE