Local structure of S impurities in implanted GaAs

408

LOCAL STRUCTURE

Nuclear Instruments and Methods in Physics Research B19/20 (1987) 408-412 North-Holland, Amsterdam

OF S IMPURITIES IN IMPLANTED

GaAs

F. S E T T E , S.J. P E A R T O N , J.M. P O A T E a n d J.E. R O W E A T& T Bell Laboratories, Murray Hill, New Jersey 07974, USA

The local structure of S implanted at high doses (1015-1016 cm -2, 100 keV) into GaAs and annealed at 900°C has been determined by extended X-Ray absorption fine structure by monitoring the S (K,,) fluorescence yield. The S first neighbor shell shows a significant static broadening compared to the S second and third shells. This indicates two S configurations of approximately equal population: (1) substitutional S on an As site and (2) a complex formed by S on an As site and an As vacancy on the second shell with a S-first neighbor distance relaxation of 0.14 _+0.04 .~. The two-site configuration explains the well-known disparity between implanted S concentration and net electrical activity, a phenomenon which occurs at high concentrations for all donor species implanted in GaAs.

1. Introduction

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An enduring problem in GaAs of both technological and fundamental interest is the inability to fuUy activate high dose n-type implants [1-4]. The electrical activity after high temperature annealing is generally limited to less than 3-5 × 1018 cm -3, whereas implanted acceptor ions such as Zn, Be, Cd and Mg show hole densities approaching 10 20 cm -3 after similar annealing (fig. 1). This difference between n- and p-type dopants is particularly surprising in view of recent solubility measurements of implanted S [5], Te and Sn [6,7] in GaAs, which show that the soluble fraction is at least an order of magnitude higher than the net electrical activity. In other words, substitutionality is a necessary but not sufficient condition for electrical activity of donors in implanted GaAs [6]. This phenomenon is also evident in bulk-doped crystals and manifests itself in the use of ill-controlled, alloyed ohmic contacts to n-type material where one cannot achieve the requisite doping level needed ( - 10 20 c m - 3 ) for direct metal contacts relying on tunneling conduction. Various models of the local structure for defect centers have been proposed, but so far, the geometrical structure of the impurity atom has not been measured directly. We report here bond-distances and numbers of atoms in the atomic shells surrounding a prototypical donor, S in GaAs, obtained by measuring the extended X-ray absorption fine structure (EXAFS) signal from the l s core level of the S impurity atom [8]. The EXAFS spectra were measured by detecting the impurity K a soft X-ray fluorescence yield (FY). Our studies cover the concentration range between 2 X 1019 and 2 × 10 20 S a t . / c m 3.

0168-583X/87/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

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2. Experimental The experiments were performed using synchrotron radiation at the Stanford Synchrotron Radiation Laboratory. The fluorescence detector, used to monitor the S ( K a ) FY, was a custom-made energy dispersive proportional counter with a 125/~m Be window and an

409

F. Sette et aL / Local structure of S impurities in GaAs

energy resolution of = 800 eV fwhm [9]. The samples were prepared by implanting 100 keV 32S ions into undoped GaAs (100) at doses in the range 5 × 1012-1016 cm -2. Annealing was performed in a flowing AsH 2 ambient at temperatures in the 700-900 ° C range. The depth profiles of S were measured by secondary ion mass spectrometry (SIMS) and the electrical activity of the S was determined by Hall effect and electrochemical capacitance-voltage ( C - V ) profiling. In accordance with previous results, [10] we found a saturation electron density = 5 × 1018 cm 3, which is at least an order of magnitude less than the implanted S concentration. Remnant damage was studied by ion channeling; samples showed a high degree of crystallinity after annealing, with backscattering yields ( = 4%) close to those of virgin GaAs.

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In fig. 2 we show the S (K-edge) photoabsorption spectra and the EXAFS signal in k-space of two GaAs samples which were sampled with S to doses of i × 1015 and 1 × 1016 at./cm 2 and annealed at 900°C. The mean concentrations (~ = 2 X 1019 at./cm 3 and ~ = 2 × 10 20 at./cm 3 respectively) were calculated assuming Gaussian-like implantation distributions in agreement with SIMS measurements. Measurements of the S K L L Auger electrons, with escape depths of = 30 A showed no enrichment of S at the surface [11]. The S atoms were distributed over a mean depth of = 500 A from the surface. On the 2 × 1019 at./cm 3 sample, the S K-edge signal-to-background ratio (STB) is 220%. EXAFS experiments can be performed with STB ratios of 10% [9,11]. The S K~ absorption length in GaAs is - 1 gm. The detection sensitivity limit of soft X-ray FY EXAFS can therefore be calculated to be (10%/220%) × (0.5 g m / 1 /zm) × 2 × 1019 at./cm 3, i.e., 5 × 1017 a t . / c m 3 for S impurities in GaAs. Simile arguments apply to other impurities, such as Si, P, C1 and Ar, with K-edges in the soft X-ray region. In fig. 3 we show the Fourier transform IF(r)l of the data reported in fig. 2. Three peaks arise from the first three atomic coordination shells around S which are analyzed using experimental phase-shifts and backscattering amplitudes [8] from GaS, GeS2, and CuGaS 2. Inconsistencies in bond length due to differences between Cu, Ga and Ge phase shifts and back-scattering amplitudes were less than + 0.01 ,h,. The peaks in fig. 3 arise from Ga and As neighbors and not from other S atoms as revealed by the k-space maximum of the backscattering amplitude. Each bond distance and number of neighbors were obtained by fitting the backtransformed data of the shell with the amplitude and phase obtained from the first-neighbor shell of each

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410

F. Sette et a L / Local structure of S impurities in GaAs

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the first shell E X A F S parameters show some anomalies with respect to those of the outer shells. The n u m b e r of first-shell neighbors is 2.6 instead of 4 as expected for substitutional S atoms and the first-shell Debye-Waller factor is physically too large when compared with that of the second and third shells. Another unusual feature is the width of the first-shell peak in the IF(r)[ transform function which has fwhm 0.13 A wider than the second and third-shell width. This point is illustrated in fig. 4 for the data of the 2 × 1020 a t . / c m 3 sample. A first shell wider than the outer shells means that there

Fig. 5. S first shell filtered data (dots) with one and two configuration fits for the ~ = 2 × 102o at/cm 3 S doped GaAs sample. The first were obtained using GaS first-shell model amplitude and phase shift: (a) one-configuration fit with R = 2.33 ,~ and 2.6 neighbors (dashed line) and the two-configuration fit with R t = 2.31 A, R 2 = 2.43 ~, and 4 neighbors (solid line); (b) the residual differences between the first shell filtered data and the one-configuration fit (dashed line) or the twoconfiguration fit (solid line). Note the change in vertical scale from (a) to (b).

are interactions affecting the local structural configuration of the atoms closer to the S but not the position of the S with respect to the rest of the lattice. We checked this idea by comparing the width of the first, second and third-shell Fourier transformed EXAFS peaks in the IF(r)[ of G a and As in GaAs [13]. In this system second and third-shell peak widths are 0.55 ,~ fwhm, consistent with our results but the first-shell width is 0.45 ,~ fwhm, which is 0.23 A narrower than our case. This effect c a n n o t be explained by a single substitutional site for the S atoms but it can be explained by two different substitutional configurations which are characterized by different first-neighbor relaxation and

411

F. Sette et al. / Local structure of S impurities in GaAs

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by relatively unperturbed second and third shells. The two S site configurations are: (1) --- 50% of S atoms on As sites with 4 G a first neighbors at 2.43 + 0.04 ,~ (S^s) and (2) ---50% of S atoms on As sites with an As vacancy on the second shell and 4 Ga first neighbors at 2.31 + 0.04 ,~ (SA~,VAs). The first-shell filtered experimental data and the least-squares best fit with one and two configurations (dashed and solid line, respectively) are reported in fig. 5. The fit with two configurations decreased the statistical X 2 and residual values by a factor of 30. The two configurations have equal concentration to within _+20%, each with 4 nearest-neighbor atoms and bond distances of 2.31 and 2.43 ,~. The uncertain, ty on these distances is given conservatively as _+0.04 A. Another conclusive aspect of this model is that the D e b y e - W a l l e r factors of the two configurations are the same as the S second and third shells and the GaS standard to within 10%, whereas the single configuration fit requires an anomalously large and unphysical D e b y e - W a l l e r factor. The two-site configuration is shown in the lattice diagram of fig. 6, emphasizing the relaxation associated with the missing As atom.

4. D i s c u s s i o n and concl usi ons

Our E X A F S results show that the S is on the As site with or without an As vacancy on the second shell (SAs or (SA~,VAs)). A missing As atom on the second shell implies four broken bonds, one on a Ga atom of the first shell, two on two Ga atoms of the third shell and one on a G a atom on the fifth shell. Therefore the As vacancy is a weak perturbation on the eleven secondshell As atoms and on ten of the twelve third shell Ga atoms, which do not have broken bonds. However, the first-shell G a atom with a broken bond will cause a contraction of the average S - G a distance. Therefore we assign the value of 2.31 + 0.04 A to the S - G a bond of this complex. Moreover the remaining 11 bonds between the first-shell Ga and second-shell As constrain the S first shell so that the S atom is close to a substitutional site and the second and third S shells are relatively unperturbed with respect to the GaAs lattice. It could be argued that the S is equally distributed on As (SAs) (unrelaxed configuration, S - G a bond 2.43 /k) and G a (S~a) (relaxed configuration, S - A s bond III. SEMICONDUCTORS

412

F Sette et al. / Local structure of S impurities in GaAs

2.31 A) sites. However, this arrangement is ruled out by specific lattice site location measurements performed on implanted samples similar to our, using the channeling technique, which show S to be uniquely on As sites [5]. Moreover, electrical activity would not saturate in the SGa + SAs model as SGa is a double donor. In contrast, electrical activity measurements are consistent with the As vacancy model where the (SAs,VAs) complex is an acceptor [5] that will compensate the donor SAs. Therefore, with two species present in roughly equal amounts, the concentration of net electrically active sites, (i.e. uncompensated donors) should be much smaller than the total S atom concentration. Indeed there are experimental observations which show that the electrical activity changes when the VA~ concentration is altered. The electrical activity of donors decreases when GaAs crystals are grown under As-deficient conditions [10] (i.e. an increase in VA~). Moreover, when our samples are co-implanted with As to the same dose and range (i.e. a decrease in VA~) the electrical activity increases. In this work we have demonstrated that EXAFS measurements using soft X-ray FY detection are a powerful tool for the determination of the local defect environment of impurities in semiconductors. Such information is crucial for a complete description of these systems. We have shown that the structure of S in implanted GaAs can be correlated with the electrical properties of this material. A similar correlation most probably holds for the other single site donors in GaAs, namely Se and Te. In the case of the amphoteric species Si, Ge and Sn, the role of self-compensating donoracceptor pairs must also be considered. The application of EXAFS to directly probe the local dopant environment is likely to continue to shed new light on dopantdefect interactions and compensation mechanisms in GaAs.

We wish to acknowledge useful discussions and support from our colleagues at A T & T Bell Laboratories. We also wish to thank the staff at SSRL, which is supported by the Office of Basic Energy Sciences of the US Department of Energy and the Division of Materials Research of the National Science Foundation.

References [1] F.H. Eisen, Radiat. Eft. 47 (1980) 99. [2] J.P. Donnelly, Nucl. Instr. and Meth. 182/183 (1981) 553. [3] K.G. Stephens, Nucl. Instr. and Meth. 209/210 (1983) 589. [4] J.S. Williams, in Laser Annealing of Semiconductors, eds., J.M. Poate and J.W. Mayer (Academic Press, New York, 1982) pp. 383. [5] R.S. Bhattacharya, A.K. Rai, Y.K. Keo, P. Pronko and Y.S. Park, Nucl. Instr. and Meth. 209/210 (1983) 637; R.S. Bhattacharya, P. Pronko and S.C. Ling, Appl. Phys. Lett. 42 (1983) 880. [6] K.G. Orrman-Rossiter, S.T. Johnson and J.S. Williams, Nucl. Instr. and Meth. B7/8 (1985) 448. [7] S.J. Pearton, J.M. Poate, F. Sette, J.M. Gibson, D.C. Jacobson and J.S. Williams, these Proceedings (IBMhi '86) Nucl. Instr. and Meth. B19/20 (1987) 369. [8] P.A. Lee, P.H. Citrin, P. Eisenberger and B.M. Kincaid, Rev. Mod. Phys. 53 (1981) 769. [9] J. Stohr, E.B. Kollin, D.A. Fisher, J.B. Hastings, F. Zaera and F. Sette, Phys. Rev. Lett. 55 (1985) 1468. [10] C.M. Wolfe and G.E. Stillman, Appl. Phys. Lett. 27 (1975) 564. [11] P.H. Citrin, J.E. Rowe and P. Eisenberger, Phys. Rev. Lett. 28 (1983) 2299. [12] DV. Lang in: Deep Levels in Semiconductors, ed., S.T. Pantelides (Gordon and Breach, New York, 1985). [13] J.C. Mikkelsen and J.B. Boyce, Phys. Rev. B28 (1983) 7130. The same results are found for Ge: E.A. Stem, B.A. Bunker and S.M. Heald, Phys. Rev. B21 (1980) 5521.