- Email: [email protected]
Carbon delta doping in chemical beam epitaxy using CBr4
Journal
ELSEVIER
of Crystal
Growth
175/176 (1997) 377-382
Carbon delta doping in chemical beam epitaxy using CBr, T.B. Joycea,*, T.J. Bullough”,
T. Farrella, B.R. Davidsonb,
D.E. Sykes”, A. Chew”
a Department of Materials Science and Engineering, The University of Liverpool, Liverpool L69 3BX, UK b Interdisciplinary Research Centre for Semiconductor Materials, The Blackett Laboratory, Imperial College of Science, Technology and Medicine, Prince Consort Road, London SW7 ZBZ, UK ‘Institute of Surface Science and Technology, Loughborough University, Loughborough, Leicestershire LEI I 3TU, UK
Abstract
We describe the growth of carbon s-doped GaAs by chemical beam epitaxy (CBE) using CBr4 as the dopant during growth interrupts. Characterisation of the S-doped samples using secondary ion mass spectrometry (SIMS) at a series of impact energies showed that the C was confined to a planar sheet less than 1 nm thick. This is in agreement with HRXRD measurements on C &doping superlattices. The results of SIMS profiling of layers with a range of interrupt times indicate that there is an initial surface coverage of C from CBr4 of approximately 9 x 10” cmm2 which increases relatively slowly during an extended interrupt.
1. Introduction There is an increasing requirement for highly doped GaAs and Al,Ga, _,As in devices such as heterojunction bipolar transistors (HBTs) and distributed Bragg reflector (DBR) stacks. Carbon is well established as a p-type dopant in GaAs for HBT base layers because of its low diffusivity combined with the high active doping level [ 11. There is considerable interest in carbon as a dopant in GaAs/AlAs or GaAs/AlGaAs DBRs for vertical cavity surface emitting lasers (VCSELs) [2]. The p-type dopants used in DBRs are typically Be in molecular beam epitaxy (MBE) and Zn in metal-
*Corresponding [email protected].
author.
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organic vapour phase epitaxy (MOVPE). Neither are ideal dopants for DBRs because of the need for periodic doping, with a high doping level at the heterojunctions to smooth out the valence band discontinuities and a lower doping level where the optical field is high to minimise free carrier absorption [3]. Be doping in AlAs is difficult, particularly at the high levels required to achieve significant smoothing of the valence band discontinuities. The level of active dopants is uncertain and at the normal MBE growth temperature of around 58O”C, significant degradation of the interface may occur. This degradation can be avoided by compromising on the Bragg structure, using Be-doped GaAs/ Al o.ssGa0.67 As grown at about 48O”C, but this requires the growth of a thicker DBR (typically 30 periods compared with 18 for GaAs/AlAs) with poorer thermal properties.
1997 Elsevier Science B.V. All rights reserved
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ofCyysta1
A more attractive solution is to focus on a different dopant. Carbon acts as a p-type dopant in GaAs and AlAs and unlike Be it can be introduced at high levels and localised close to the AlAs/GaAs interface [l]. C is far more stable than Be or Zn during high-temperature processing, having a diffusion constant some three orders of magnitude lower [4]. C doping by MBE has proved difficult in the past, solid C sources tending to introduce 0 contamination or graphitic C, whereas CBE is an ideal technique for C doping using gaseous precursors [5,6]. Using CBr, as the C source GaAs and AlAs can be doped at up to 5 x 10” cmm3 under normal growth conditions with no adverse effects on material quality (Be doping in AlAs is limited to 6 x 1017 cmp3 at this growth temperature). Carbon-doped GaAs was found to be free of impurities such as 0 (seen when a C-filament doping source was used in MBE) or H-CA, compensation (seen in MOVPE material) [7]. At very high doping levels ([Cl > 10zo cm -3) there is a reduction in the GaAs growth rate due to etching of the GaAs by Br species arising from the decomposition of CBr, [S]. Carbon doping in Al(G is similar to that in GaAs, whether using CBr, [S] or Ccl4 [9] there is little change in incorporation with Al fraction in AlGaAs. CBE also offers excellent control of layer composition, in particular the direct flux control available using vapour sources simplifies the growth of graded interfaces. The use of delta-doping at the interfaces of DBRs, where a high concentration of the dopant is confined to a planar sheet in the high bandgap material, leads to a low series resistance [lo]. In this paper we describe studies of carbon &doping in GaAs which demonstrate the localisation that can be achieved using CBE. True s-doping requires a spread less than the spatial extent of the ground state wave function in the V-shaped potential well generated by the ionised impurities [Ill]. For p-type material the wave function spread is = 30 A [ 123 and so confinement of the acceptors to a sheet with a thickness of = 10 A or less is required for true a-doping to be obtained. We have demonstrated that carbon S-doping in GaAs grown by CBE can meet these conditions [13] and present further results in this paper.
Growth 175/176 (1997) 377-382
2. Experimental
procedure
All structures were grown by chemical beam epitaxy (CBE) in a VG V80H system configured for all vapour sources [14] with pressure control of flux through a fine orifice without the need for a carrier gas. Arsine which was thermally cracked in an injector cell was used as the group V source and triethylgallium (TEGa) was used as the group III precursor. The fluxes were held constant for all of these samples giving a V : III ratio of 10 : 1 (an arsine BEP of 2 x 10m4 mbar and a TEGa BEP of 2 x lo-’ mbar) and a growth rate of 1.1 urn/h at substrate temperatures in the range 50&6OOC. Carbon dopant was introduced using CBr,; C incorporation shows little dependence on substrate temperature or V : III ratio [15]. Carbon h-doping was carried out using a growth interrupt; the highest practicable CBr, flux, which gave a BEP of 2 x 10m6 mbar, was introduced during an interruption in the TEGa flux. This CBr4 flux would give a doping level of 3 x 102’ cm-3 if introduced during growth under these conditions. Secondary ion mass spectrometry (SIMS) profiling was carried out using a Cameca ims 3f and the data quantified using implanted reference materials. Crater depths were measured using optical interference microscopy: layer thicknesses were cross-checked by cross-sectional transmission electron microscopy (TEM) and HRXRD. HRXRD measurements were carried out using a Philips high-resolution X-ray diffractometer with a 4-reflection Ge220 monochromator using Cu Kcli radiation. The 400 Bragg reflection was measured along a (1 1 0) azimuth for all samples. Simulations using dynamical diffraction theory were performed assuming Vegard’s law to be valid, giving a linear variation of strain with composition, and also assuming that all the carbon was present as isolated atoms occupying As lattice sites [13, 161. The lattice parameter of “GaC” was calculated to be 4.688 A (resulting in 17% strain) using tabulated covalent radii of the individual atoms.
3. Results and discussion A nine period &doping was used for SIMS analysis
superlattice (Table 1) of delta layer thickness
T.B. Joyce et al. /Journal
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of Crystal Growth 175/I 76 (1997) 377-382
Table 1 Structure of the 9-period F-doped sample, M601 SIMS [Cl, (lOI atoms cm-‘)
50 nm undoped GaAs cap 6-C layer 1 (2 s interrupt) 100 nm undoped GaAs 6-C layer 2 (5 s interrupt) 100 nm undoped GaAs 6-C layer 3 (2 s interrupt) 100 nm undoped GaAs 6-C layer 4 (4 s interrupt) 100 nm undoped GaAs 6-C layer 5 (2 s interrupt) 100 nm undoped GaAs 6-C layer 6 (3 s interrupt) 100 nm undoped GaAs 6-C layer 7 (2 s interrupt) 100 nm undoped GaAs 6-C layer 8 (1 s interrupt) 100 nm undoped GaAs 6-C layer 9 (2 s interrupt) 250 nm undoped GaAs buffer C interface contamination N’ GaAs substrate
Cs d,ata
0 data
1.40
1.23 (0.103)
1.41
1.33 (0.142)
.30
1.13 (0.127)
.50
1.275 (0.113)
.32
1.14 (0.126)
.49
1.24 (0.039)
1.28
1.16 (0.118)
1.14
1.16 (0.126)
1.29
1.20 (0.120)
21
21 (by definition)
Note: 0 data is averaged over a series of measurements, the standard deviation for each set of data is given in parentheses. These data were quantified against the C signal for the interface layer which was assumed to be constant for all 0 and Cs profiles.
16
-
-
layer
0
layer 2 linear fit - layers
---o--. layer
0
1
2
3
4
1
0
5
1 8 2
3 (linear fit)
6
7
8
Ion Impact Energy I keV Fig. 1. Full width at half-maximum (FWHM) of SIMS profiles of three carbon delta layers as a function of ion impact energy. FWHM increases with energy and with layer depth, layers 1-3 were at depths of 50, 150 and 250 nm respectively.
380
T.B. Jovcr et al. ~Journal qf Cystal Growth 175/l 76 (I 997) 3 77-382
and of area1 carbon concentration [Cl,. SIMS resolution of a delta layer is limited by ion beam induced atomic mixing effects, which increase with ion beam energy and with profiled depth. Fig. 1 shows the influence of ion impact energy on the
measured FWHM of the three topmost &layers of this sample (at depths of 50, 150 and 250 nm, respectively). The sample was sputter profiled using O-, 0, and 0, ions at primary beam energies of 10.5 and 12.5 kV with an extraction voltage of
1.6 1
I Sample M601
,.:
1.5
0
0 _:
Cl
_:’
_:’
0
1.4 1
_:.
_./
,:’ __..
o....’
.:’
intermpt I set E
1
10
15
20
25
30
35
interrupt I set
Fig. 2. Area1 carbon concentration [Cl, determined by SIMS for delta layers grown using CBr, introduced during various times. Data obtained profiling with Cs+, O-, 0; and 0, ions as described in the text. Sample M601 was a 9-period structure in Table I), M668 was a lo-period sample with an extended range of interrupt times.
interrupt (detailed
T.B. Joyce et al. /Journal
qf Cytal
4.5 kV, to give ion impact energies in the range 2-8 keV. The deeper layers have a broader FWHM due to atomic mixing and surface roughening effects, particularly layer 3, while those measured at lower energy are sharper. The first layer has a FWHM of 5 nm at the lowest energy with a linear increase at higher energies. Lower ion impact energies are not easily achieved. Extrapolating the data set for layers 1 and 2 to zero impact energy suggests that the actual FWHM is less than 1 nm, which is in agreement with the results of HRXRD [13]. Only the data from layers close to the sample surface can be analysed in this way. The data for layers 3 and beyond show little dependence on impact energy as other SIMS induced profile distortions begin to dominate the measured delta layer widths. Table 1 shows details of the nine period &doping superlattice and the results of SIMS analysis of area1 carbon concentration [Cl,. To achieve a compromise between depth resolution and sensitivity to carbon the first analysis was performed at a slow sputter rate using Cs+ primary ion bombardment and negative secondary ion detection, under these conditions the SIMS carbon background is about 5 x 10” cm- 3. These data were quantified by calibration against ion-implanted standards and the value for [CIA derived by integrating the counts under the peak for each deltadoped layer. Broadening of the SIMS peak with depth gives a lower peak value but does not effect the value of [CIA. The depth and the integrated signal for the interface contamination layer were then used to calibrate a second series of measurements which were taken using O-, 0, and 0, ions as described above. The average values for these measurements are given; note that sputtering with oxygen gives improved depth resolution but lower sensitivity to carbon. Every alternate S-layer was a reference layer with a 2 s interrupt so the variation of SIMS sensitivity and resolution with sputtered depth could be monitored; however, [CIA for the first layer is artificially high because of surface C contamination. No systematic variation in [CIA was found with layer depth or with primary ion species and energy. As shown in Fig. 2a there is an approximately linear dependence of [CIA on interrupt time over the range studied but with a large offset term for
Growth 175/176 (1997) 377 3X2
381
zero interrupt. This would imply that there is an initial surface coverage of C from CBr, which increases relatively slowly during an extended interrupt. A second sample was grown with a range of interrupt times from 0.1 to 30 s. The results of SIMS profiling using 10 kV 0; ions are summarised in Fig. 2b. There is clearly a linear dependence on exposure time, with no evidence of saturation at higher coverage. This linear dependence appears to extend to times less than 1 s but the flux response is probably limited by pressure transients in the gas line rather than by valve switching. For both samples the value of [Cl, for zero interrupt time is close to the calculated level of 8.5 x 10” cmm2 per monolayer for GaAs homogeneously doped at 3 x 10” cmm3, the expected level for these TEGa and CBr4 fluxes. SIMS values of [Cl, are in reasonable agreement with those measured by other techniques. SIMS profiling using Csf gave values between 1.3 and 2.0 x 1013 cme2 for a series of samples with a 2-3 s interrupt. As reported previously [13], HRXRD gave a value for [C,JA of 1.5 x 1013 cmA2 for a 50 period doping superlattice. IR absorption spectra showed a [CAJA concentration of 1.1 x 1013 cmm2 per Z-layer with no detectable hydrogen passivation, in agreement with Hall and ECV measurements of area1 acceptor concentration of 1.1 x lOI cmm2 per &layer. The C F-doping superlattices grown for HRXRD analysis comprised 50 carbon &layers separated by 50 nm GaAs spacers. Comparison of measured and simulated HRXRD data for such a structure (sample M533, with C doping during 2 s growth interrupts) gave an excellent fit for a simulation which was obtained using a rectangular doping profile with an area1 carbon concentration [CA,], of 1.5 x lOi3 cme2 and a thickness rd of 5 A [13, 161. Decreasing td to one monolayer (2.8 A) while [CA,]* was kept constant gave an equally good fit while increasing t, to 10 A with the same value of [C,,], gave a slightly worse fit. Replacing the abrupt doping profile in the simulation with a triangular profile (linear spreading on both sides) showed that the spread of the &layers was less than 10 A. To improve the fit a random period fluctuation of 2% was included in this simulation which broadened the satellite lines. Greater precision in
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ofCrystal Growth 1751176 (1997) 377-382
determining the &layer thickness would be possible only if more orders of satellites could be detected since it is the high-order satellites that result from abrupt interfaces [ 171.
4. Conclusions We have demonstrated carbon delta-doping in GaAs grown by CBE. Area1 C concentrations of up to 3.5 x lOi cmm2 per delta layer were obtained using CBr, introduced during a growth interrupt. The results of SIMS profiling indicate that there is a rapid initial surface coverage of C from CBr4 and a relatively slow increase during an extended interrupt. The delta layers were confined to a planar sheet less than 1 nm thick as measured by HRXRD and SIMS.
Acknowledgements The authors would like to thank L. Hart and R.C. Newman for helpful discussions and to acknowledge the EPSRC for their financial support.
References [I] T.H. Chiu, J.E. Cunningham, and S.N.G. Chu, J. Crystal
J.A. Ditzenberger, W.J. Jan Growth 111 (1991) 274.
[2] M. Sugimoto, I. Ogaru, H. Saito, A. Yasuda, K. Kurihara. H. Kosaka and T. Numai, J. Crystal Growth 127 (1993) 1. [3] M. Sugimoto, H. Kosaka, K. Kurihara, 1. Ogaru. T. Numai and K. Kasahara, Electron. Lett. 28 (1993) 385. [4] J. Nagle, R.J. Malik and D. Gershoni, J. Crystal Growth 111 (1991) 264. [S] T.B. Joyce, S.P. Westwater, P.J. Goodhew and R.E. Pritchard, J. Crystal Growth 164 (1996) 371. [6] CR. Abernathy, Mater. Sci. Eng. R14 (1995) 2033254. [7] M.J. Ashwin, B.R. Davidson, K. Woodhouse, R.C. Newman, T.J. Bullough, T.B. Joyce, R. Nicklin and R.R. Bradley, Semicond. Sci. Technol. 8 (1993) 625. [S] T.B. Joyce, T.J. Bullough and T. Farrell, Appl. Phys. Lett. 65 (1994) 2193. [9] C.R. Abernathy, S.J. Pearton, F. Ren, W.S. Hobson and P.W. Wisk, J. Vat. Sci. Technol. A 12 (1994) 1186. [lo] M.G. Peters, B.J. Thibeault, D.B. Young, A.C. Gossard and L.A. Coldren, J. Vat. Sci. Technol. B 12 (1994) 3075. [ 1 l] E.F. Schubert, in: Doping in III-V Semiconductors, Cambridge Studies in Semiconductor Physics and Microengineering, Vol. 1 (Cambridge University Press, Cambridge, 1993) ch. 1 I. [12] F.A. Reboredo and CR. Proetto, Phys. Rev. B 47 (1993) 4655. [13] B.R. Davidson, L. Hart, R.C. Newman, T.B. Joyce and T.J. Bullough, J. Crystal Growth 164 (1996) 383. 1141 T.B. Joyce, T.J. Bullough, P. Kightley, Y.R. Xing and P.J. Goodhew, J. Crystal Growth 120 (1993) 206. [15] T.J. de Lyon, N.I. Buchan, P.D. Kirchner, J.M. Woodall, G.J. Scilla and F. Cardone, Appl. Phys. Lett. 58 (1991) 517. 1161 L. Hart, M.R. Fahy, R.C. Newman and P. Fewster, Appl. Phys. Lett. 62 (1993) 2218. 1171 L. Hart. P.F. Fewster, M.J. Ashwin, M.R. Fahy and R.C. Newman, J. Phys. D: Appl. Phys. 28 (1995) Al 54.
ELSEVIER
of Crystal
Growth
175/176 (1997) 377-382
Carbon delta doping in chemical beam epitaxy using CBr, T.B. Joycea,*, T.J. Bullough”,
T. Farrella, B.R. Davidsonb,
D.E. Sykes”, A. Chew”
a Department of Materials Science and Engineering, The University of Liverpool, Liverpool L69 3BX, UK b Interdisciplinary Research Centre for Semiconductor Materials, The Blackett Laboratory, Imperial College of Science, Technology and Medicine, Prince Consort Road, London SW7 ZBZ, UK ‘Institute of Surface Science and Technology, Loughborough University, Loughborough, Leicestershire LEI I 3TU, UK
Abstract
We describe the growth of carbon s-doped GaAs by chemical beam epitaxy (CBE) using CBr4 as the dopant during growth interrupts. Characterisation of the S-doped samples using secondary ion mass spectrometry (SIMS) at a series of impact energies showed that the C was confined to a planar sheet less than 1 nm thick. This is in agreement with HRXRD measurements on C &doping superlattices. The results of SIMS profiling of layers with a range of interrupt times indicate that there is an initial surface coverage of C from CBr4 of approximately 9 x 10” cmm2 which increases relatively slowly during an extended interrupt.
1. Introduction There is an increasing requirement for highly doped GaAs and Al,Ga, _,As in devices such as heterojunction bipolar transistors (HBTs) and distributed Bragg reflector (DBR) stacks. Carbon is well established as a p-type dopant in GaAs for HBT base layers because of its low diffusivity combined with the high active doping level [ 11. There is considerable interest in carbon as a dopant in GaAs/AlAs or GaAs/AlGaAs DBRs for vertical cavity surface emitting lasers (VCSELs) [2]. The p-type dopants used in DBRs are typically Be in molecular beam epitaxy (MBE) and Zn in metal-
*Corresponding [email protected].
author.
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0022-0248/97/$17.00 Copyright PII SOO22-0248(96)00957-8
+ 44 151 794 4675; e-mail:
0
organic vapour phase epitaxy (MOVPE). Neither are ideal dopants for DBRs because of the need for periodic doping, with a high doping level at the heterojunctions to smooth out the valence band discontinuities and a lower doping level where the optical field is high to minimise free carrier absorption [3]. Be doping in AlAs is difficult, particularly at the high levels required to achieve significant smoothing of the valence band discontinuities. The level of active dopants is uncertain and at the normal MBE growth temperature of around 58O”C, significant degradation of the interface may occur. This degradation can be avoided by compromising on the Bragg structure, using Be-doped GaAs/ Al o.ssGa0.67 As grown at about 48O”C, but this requires the growth of a thicker DBR (typically 30 periods compared with 18 for GaAs/AlAs) with poorer thermal properties.
1997 Elsevier Science B.V. All rights reserved
378
T.B. Joyce et al. /Journal
ofCyysta1
A more attractive solution is to focus on a different dopant. Carbon acts as a p-type dopant in GaAs and AlAs and unlike Be it can be introduced at high levels and localised close to the AlAs/GaAs interface [l]. C is far more stable than Be or Zn during high-temperature processing, having a diffusion constant some three orders of magnitude lower [4]. C doping by MBE has proved difficult in the past, solid C sources tending to introduce 0 contamination or graphitic C, whereas CBE is an ideal technique for C doping using gaseous precursors [5,6]. Using CBr, as the C source GaAs and AlAs can be doped at up to 5 x 10” cmm3 under normal growth conditions with no adverse effects on material quality (Be doping in AlAs is limited to 6 x 1017 cmp3 at this growth temperature). Carbon-doped GaAs was found to be free of impurities such as 0 (seen when a C-filament doping source was used in MBE) or H-CA, compensation (seen in MOVPE material) [7]. At very high doping levels ([Cl > 10zo cm -3) there is a reduction in the GaAs growth rate due to etching of the GaAs by Br species arising from the decomposition of CBr, [S]. Carbon doping in Al(G is similar to that in GaAs, whether using CBr, [S] or Ccl4 [9] there is little change in incorporation with Al fraction in AlGaAs. CBE also offers excellent control of layer composition, in particular the direct flux control available using vapour sources simplifies the growth of graded interfaces. The use of delta-doping at the interfaces of DBRs, where a high concentration of the dopant is confined to a planar sheet in the high bandgap material, leads to a low series resistance [lo]. In this paper we describe studies of carbon &doping in GaAs which demonstrate the localisation that can be achieved using CBE. True s-doping requires a spread less than the spatial extent of the ground state wave function in the V-shaped potential well generated by the ionised impurities [Ill]. For p-type material the wave function spread is = 30 A [ 123 and so confinement of the acceptors to a sheet with a thickness of = 10 A or less is required for true a-doping to be obtained. We have demonstrated that carbon S-doping in GaAs grown by CBE can meet these conditions [13] and present further results in this paper.
Growth 175/176 (1997) 377-382
2. Experimental
procedure
All structures were grown by chemical beam epitaxy (CBE) in a VG V80H system configured for all vapour sources [14] with pressure control of flux through a fine orifice without the need for a carrier gas. Arsine which was thermally cracked in an injector cell was used as the group V source and triethylgallium (TEGa) was used as the group III precursor. The fluxes were held constant for all of these samples giving a V : III ratio of 10 : 1 (an arsine BEP of 2 x 10m4 mbar and a TEGa BEP of 2 x lo-’ mbar) and a growth rate of 1.1 urn/h at substrate temperatures in the range 50&6OOC. Carbon dopant was introduced using CBr,; C incorporation shows little dependence on substrate temperature or V : III ratio [15]. Carbon h-doping was carried out using a growth interrupt; the highest practicable CBr, flux, which gave a BEP of 2 x 10m6 mbar, was introduced during an interruption in the TEGa flux. This CBr4 flux would give a doping level of 3 x 102’ cm-3 if introduced during growth under these conditions. Secondary ion mass spectrometry (SIMS) profiling was carried out using a Cameca ims 3f and the data quantified using implanted reference materials. Crater depths were measured using optical interference microscopy: layer thicknesses were cross-checked by cross-sectional transmission electron microscopy (TEM) and HRXRD. HRXRD measurements were carried out using a Philips high-resolution X-ray diffractometer with a 4-reflection Ge220 monochromator using Cu Kcli radiation. The 400 Bragg reflection was measured along a (1 1 0) azimuth for all samples. Simulations using dynamical diffraction theory were performed assuming Vegard’s law to be valid, giving a linear variation of strain with composition, and also assuming that all the carbon was present as isolated atoms occupying As lattice sites [13, 161. The lattice parameter of “GaC” was calculated to be 4.688 A (resulting in 17% strain) using tabulated covalent radii of the individual atoms.
3. Results and discussion A nine period &doping was used for SIMS analysis
superlattice (Table 1) of delta layer thickness
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of Crystal Growth 175/I 76 (1997) 377-382
Table 1 Structure of the 9-period F-doped sample, M601 SIMS [Cl, (lOI atoms cm-‘)
50 nm undoped GaAs cap 6-C layer 1 (2 s interrupt) 100 nm undoped GaAs 6-C layer 2 (5 s interrupt) 100 nm undoped GaAs 6-C layer 3 (2 s interrupt) 100 nm undoped GaAs 6-C layer 4 (4 s interrupt) 100 nm undoped GaAs 6-C layer 5 (2 s interrupt) 100 nm undoped GaAs 6-C layer 6 (3 s interrupt) 100 nm undoped GaAs 6-C layer 7 (2 s interrupt) 100 nm undoped GaAs 6-C layer 8 (1 s interrupt) 100 nm undoped GaAs 6-C layer 9 (2 s interrupt) 250 nm undoped GaAs buffer C interface contamination N’ GaAs substrate
Cs d,ata
0 data
1.40
1.23 (0.103)
1.41
1.33 (0.142)
.30
1.13 (0.127)
.50
1.275 (0.113)
.32
1.14 (0.126)
.49
1.24 (0.039)
1.28
1.16 (0.118)
1.14
1.16 (0.126)
1.29
1.20 (0.120)
21
21 (by definition)
Note: 0 data is averaged over a series of measurements, the standard deviation for each set of data is given in parentheses. These data were quantified against the C signal for the interface layer which was assumed to be constant for all 0 and Cs profiles.
16
-
-
layer
0
layer 2 linear fit - layers
---o--. layer
0
1
2
3
4
1
0
5
1 8 2
3 (linear fit)
6
7
8
Ion Impact Energy I keV Fig. 1. Full width at half-maximum (FWHM) of SIMS profiles of three carbon delta layers as a function of ion impact energy. FWHM increases with energy and with layer depth, layers 1-3 were at depths of 50, 150 and 250 nm respectively.
380
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and of area1 carbon concentration [Cl,. SIMS resolution of a delta layer is limited by ion beam induced atomic mixing effects, which increase with ion beam energy and with profiled depth. Fig. 1 shows the influence of ion impact energy on the
measured FWHM of the three topmost &layers of this sample (at depths of 50, 150 and 250 nm, respectively). The sample was sputter profiled using O-, 0, and 0, ions at primary beam energies of 10.5 and 12.5 kV with an extraction voltage of
1.6 1
I Sample M601
,.:
1.5
0
0 _:
Cl
_:’
_:’
0
1.4 1
_:.
_./
,:’ __..
o....’
.:’
intermpt I set E
1
10
15
20
25
30
35
interrupt I set
Fig. 2. Area1 carbon concentration [Cl, determined by SIMS for delta layers grown using CBr, introduced during various times. Data obtained profiling with Cs+, O-, 0; and 0, ions as described in the text. Sample M601 was a 9-period structure in Table I), M668 was a lo-period sample with an extended range of interrupt times.
interrupt (detailed
T.B. Joyce et al. /Journal
qf Cytal
4.5 kV, to give ion impact energies in the range 2-8 keV. The deeper layers have a broader FWHM due to atomic mixing and surface roughening effects, particularly layer 3, while those measured at lower energy are sharper. The first layer has a FWHM of 5 nm at the lowest energy with a linear increase at higher energies. Lower ion impact energies are not easily achieved. Extrapolating the data set for layers 1 and 2 to zero impact energy suggests that the actual FWHM is less than 1 nm, which is in agreement with the results of HRXRD [13]. Only the data from layers close to the sample surface can be analysed in this way. The data for layers 3 and beyond show little dependence on impact energy as other SIMS induced profile distortions begin to dominate the measured delta layer widths. Table 1 shows details of the nine period &doping superlattice and the results of SIMS analysis of area1 carbon concentration [Cl,. To achieve a compromise between depth resolution and sensitivity to carbon the first analysis was performed at a slow sputter rate using Cs+ primary ion bombardment and negative secondary ion detection, under these conditions the SIMS carbon background is about 5 x 10” cm- 3. These data were quantified by calibration against ion-implanted standards and the value for [CIA derived by integrating the counts under the peak for each deltadoped layer. Broadening of the SIMS peak with depth gives a lower peak value but does not effect the value of [CIA. The depth and the integrated signal for the interface contamination layer were then used to calibrate a second series of measurements which were taken using O-, 0, and 0, ions as described above. The average values for these measurements are given; note that sputtering with oxygen gives improved depth resolution but lower sensitivity to carbon. Every alternate S-layer was a reference layer with a 2 s interrupt so the variation of SIMS sensitivity and resolution with sputtered depth could be monitored; however, [CIA for the first layer is artificially high because of surface C contamination. No systematic variation in [CIA was found with layer depth or with primary ion species and energy. As shown in Fig. 2a there is an approximately linear dependence of [CIA on interrupt time over the range studied but with a large offset term for
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zero interrupt. This would imply that there is an initial surface coverage of C from CBr, which increases relatively slowly during an extended interrupt. A second sample was grown with a range of interrupt times from 0.1 to 30 s. The results of SIMS profiling using 10 kV 0; ions are summarised in Fig. 2b. There is clearly a linear dependence on exposure time, with no evidence of saturation at higher coverage. This linear dependence appears to extend to times less than 1 s but the flux response is probably limited by pressure transients in the gas line rather than by valve switching. For both samples the value of [Cl, for zero interrupt time is close to the calculated level of 8.5 x 10” cmm2 per monolayer for GaAs homogeneously doped at 3 x 10” cmm3, the expected level for these TEGa and CBr4 fluxes. SIMS values of [Cl, are in reasonable agreement with those measured by other techniques. SIMS profiling using Csf gave values between 1.3 and 2.0 x 1013 cme2 for a series of samples with a 2-3 s interrupt. As reported previously [13], HRXRD gave a value for [C,JA of 1.5 x 1013 cmA2 for a 50 period doping superlattice. IR absorption spectra showed a [CAJA concentration of 1.1 x 1013 cmm2 per Z-layer with no detectable hydrogen passivation, in agreement with Hall and ECV measurements of area1 acceptor concentration of 1.1 x lOI cmm2 per &layer. The C F-doping superlattices grown for HRXRD analysis comprised 50 carbon &layers separated by 50 nm GaAs spacers. Comparison of measured and simulated HRXRD data for such a structure (sample M533, with C doping during 2 s growth interrupts) gave an excellent fit for a simulation which was obtained using a rectangular doping profile with an area1 carbon concentration [CA,], of 1.5 x lOi3 cme2 and a thickness rd of 5 A [13, 161. Decreasing td to one monolayer (2.8 A) while [CA,]* was kept constant gave an equally good fit while increasing t, to 10 A with the same value of [C,,], gave a slightly worse fit. Replacing the abrupt doping profile in the simulation with a triangular profile (linear spreading on both sides) showed that the spread of the &layers was less than 10 A. To improve the fit a random period fluctuation of 2% was included in this simulation which broadened the satellite lines. Greater precision in
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determining the &layer thickness would be possible only if more orders of satellites could be detected since it is the high-order satellites that result from abrupt interfaces [ 171.
4. Conclusions We have demonstrated carbon delta-doping in GaAs grown by CBE. Area1 C concentrations of up to 3.5 x lOi cmm2 per delta layer were obtained using CBr, introduced during a growth interrupt. The results of SIMS profiling indicate that there is a rapid initial surface coverage of C from CBr4 and a relatively slow increase during an extended interrupt. The delta layers were confined to a planar sheet less than 1 nm thick as measured by HRXRD and SIMS.
Acknowledgements The authors would like to thank L. Hart and R.C. Newman for helpful discussions and to acknowledge the EPSRC for their financial support.
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