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Recent developments in ion implantation in silicon
Materials Science and EngiHeering, B4 (1989) 87 94
87
Recent Developments in Ion Implantation in Silicon J. A. PALS and S. D. BROTHERTON
l'hilips Research Laboratories, Redhil[, Surrey, R i l l 5HA (U.K.) A. H, VAN OMMEN, J. POLITIEK and H.J. L1GTHART
l~hilil)s Research Lahoratories, PO Box 80 O(X), NL-5600JA Eindhoven (The Netherhmds) (Received May 31, 1989:
Abstract
The influence o1"high dose implantations in silicon on defects is discussed by means of two exampies. The first is an investigation of the leakage currents caused by defects in veO, shallow p-n fimctions, obtained by low energy boron implantation into pre-amorphized n-type material. Preamorphizalion is used to prevent channelling effects in order to be able to make ve~ shallow junctions. Interstitial defects are shown to play an important role. The second example is a discussion of the defects occurring during vet3' high dose implantation used for ion beam ~ynthesis of buried oxide layers in silicon. The physical picture of this process is still rather unclear, but it is believed that again interstitial defects play an important role. 1. Introduction Ion implantation in silicon is now a wellestablished technique for introducing, in a controlled way, donors and acceptors in silicon slices. The requirement for shallow junctions in submicron very-large-scale integrated (VLSI)circuits has aroused interest in low energy implantations. Channelling effects, especially for the lighter implanted atoms such as boron, may cause the junction depths to be larger than predicted from the ion implantation range tables. An effective technique for suppressing these channelling effects is the amorphization of the surface region of the crystal by implantation of a high dose of an inert ion. Solid state phase epitaxial regrowth at temperatures as low as 550°C is then used for recrystallization and activation of the implanted dopant atoms without any appreciable redistribution. However, this procedure may result in diodes with poor electrical characteristics owing to the introduction of defects l 1,2]. 0921-5107/89/$3.50
Another development in which high dose implantations are being used is the ion beam synthesis of buried layers [3]. The best known example is the formation of a buried oxide layer in silicon by a very high dose oxygen implantation. There is much interest in such structures with a thin silicon layer on top of an insulating layer, for instance for making very high speed complementary metal-oxide-semiconductor (CMOS) circuits. The implantation is done at elevated temperatures to prevent the remaining top silicon layer from amorphizing. Again the damage created by the high dose implantation is a very important factor in determining the quality of the resulting Si-SiO~ structure. It is the purpose of this paper to present a number of experimental results on the creation of damage by high dose implantations followed by annealing steps used for amorphization or ion beam synthesis of buried layers. 2. Pre-amorphized and annealed layers
2.1. Sample preparation The samples used for the investigation of the effects of pre-amorphization were obtained by implanting n-type Si(100)-oriented substrates with Si +, Ge + or Sn + ions. The implantation conditions and the thickness of the resulting amorphous layers, measured with Rutherford backscattering (RBS) or cross-sectional transmission electron microscopy (XTEM), are summarized in Table 1. The Si + implantations have been made with different energies to vary the position of the a - c boundary with respect to the metallurgical p +-n junction. This junction was made after pre-amorphization by implanting a dose of 5 x 10 H~BF, + ions at 25 keV. The as-implanted junction depth determined by secondary ion mass © Elsevier Sequoia/Printed in The Netherhmds
S~
'•\
TABLE I
/on
Dose ',cm
Si-
Ge~ Sn +
2x 2× 2× 2× 1x ~8 x 5x
Lnergy :',
(keV)
10 I" 11)~5 10 I~ 10 ~" 1() I~ it) ~4 1014
30 50 7(1 90 200 180 65
Amorphous layer thickness(A)
t0-2
640,720* 1(16(I 1400 1740, 176()* 2000 1440, 1420"
10 -3
?
x o a Q
30 50 70 90
o\% % \\ \\ %
10-4
*TEM measurements; all others RBS. u
spectrometry (SIMS) profiling was approximately 750 A, i.e. smaller than the a - c boundary depth for all pre-amorphizations with the exception of the 30 keV Si + implantation. Furnace annealing at temperatures between 600 and 1000 °C was used to regrow the amorphous layers and to activate the implanted boron. Finally the samples were annealed at 450 °C in wet N 2 to remove fast Si-SiO 2 interface states and metallized with cold aluminium. The diodes were circular with field plates to control surface generation currents. The SIMS profiles of the boron after annealing showed no anomalous diffusion in the Si + preamorphized layers made with energies of 50 keV or above, as will be discussed later.
2.2. Electrical characterization of the diodes Leakage current measurements of the diodes were made and the results of the Si+-amorphized diodes are shown in Fig. 1. Three regions can be distinguished. Region I, for anneal temperatures of 600 and 700 °C, shows a very high reverse-bias leakage current density with the current rising strongly with reverse bias. Similar results were obtained for the Ge*- and Sn +-implanted diodes [4, 5]. Annealing at or above 800 °C resulted in lower leakage current densities, saturating at a level of about 10-5 A cm-2 for the diodes with the higher implantation energies (region II) or becoming much lower for the diodes with very shallow pre-amorphized layers (region IIt). Room temperature C - V profiles obtained from diodes with a 10 l~ cm 3 substrate doping level and amorphized by the three different implants and annealed at 600, 700 or 800 °C are shown in Fig. 2. In each case the lower regrowth temperatures of 600 and 700 °C resulted in the presence of a high concentration of excess donors, which were removed by the 800 °C anneal. The precise inter-
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I
1
i
i
10-7
1°-~l
680
.7.0.0. . . . . ~;o Anneal temperature
9OO (~C)
Fig. 1. Reverse-bias leakage current densities measured as a function of the anneal temperature in Si+-pre-amorphized p +-n diodes.
pretation of the profiles has to be done with care because of the deep level character of the excess donors [6]. We will now discuss in more detail the defects present in the diodes in regions I, II and llI in Fig. 1.
2.2.1. Region I deep level defects To characterize further the defects present m the amorphized and regrown layers, deep level transient spectroscopy (DLTS, measurements have been made. A typical set of curves is shown in Fig. 3 for samples amorphized with Si +. Ge T and Sn + and annealed at 700°C. Four defect levels labelled A-D were found in the Si*implanted sample. These levels were also found in the Ge*-implanted sample, but in the Sn-implanted sample only A and D were found unambiguously, although level C could be found in samples annealed at 600 °C. The similarity of the results suggest that the basic defects are identical for all three amorphizing species. Spatial profiles of the deep level defects have been obtained by measuring the DLTS curves as a function of reverse bias and a result is given in
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Fig. 2. Room temperature C - I / p r o f i l e s obtained from p+ n diodes pre-amorphized with (a) 90 keV Si +, (b) Ge+ and (c)
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150 200 Temperature
I
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250 (K)
I
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Fig. 3. DUI'S c u r v e s obtained from p * - n d i o d e s a m o r p h i z e d w i t h (a) 70 kcV Si + , (b) G e + and (c)Sn
pre-
Sn + .
Fig. 4 for an Si+-pre-amorphized sample. Also indicated are the location of the junction and the a - c boundary dislocation loops. The deep level defects are evidently lying deeper than the ct-c boundary and in fact their concentration profile is very close to the implanted silicon profile obtained with Lindhard, Scharff and Schiott (LSS) range tables. Similar results were again obtained with the Ge +- and Sn~-amorphized samples [4, 5J, both in terms of the location of the traps beyond the a - c boundary and the closeness of their profiles to the implanted impurity profiles. These results suggest strongly that the defects are fundamental damage-related centres and may be connected to the excess silicon interstitial in the region beyond the a - c boundary. The thermal activation energies of the different centres A-D were found to be about 0.30, 0.46, 0.49 and 0.65 eV respectively; the values were dependent on the applied reverse bias on the
diodes and indicated field-enhanced emission 16J. From these trap activation energies it is concluded that the near-midgap centre D could be responsible for the large reverse-bias leakage currents measured in these structures. This was confirmed by examination of the dependence of the leakage current density on the concentration of the D-trap after annealing the samples for different periods of time to progressively remove that centre. A good correlation between trap density and leakage current was found [7].
2. 2.2. Region ll--extended defects Following an anneal at 800°C or above, the deep level centres discussed above were removed and the leakage current densities were reduced to approximately 10 -5 A cm-z for the diodes with the ct-c boundary beyond the p-n junction depth Xi .
XTEM micrographs of 90 keV Si+-amor phized and regrown material [i] have revealed a
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Fig. 4, Deep level d o n o r profiles obtained from a 90 keV Si + -pre-amorphized diode annealed at 700 °C.
band of dislocation loops, 50(/_+ 200 A wide, immediately beneath the original amorphous layer. These loops were observed to grow in size and decrease in number density with increasing anneal temperatures. With increasing anneal temperature the total silicon atom density in these loops initially increased and finally saturated at approximately 3 x 1014 cm -2 for temperatures above 800 °C. This loop growth is interpreted to be caused by the flow of excess silicon interstitials from beyond the a - c boundary. Monte Carlo calculations of Thornton and coworkers [8, 9] have predicted a silicon interstitial concentration caused by both the directly implanted and the recoil-implanted atoms of about 1015 c m -2, which is rather close to the experimentally found density in the loops. These a - c boundary dislocation loops were demonstrated to be the cause of the leakage currents in region II of Fig. 1. A number of diodes were annealed at 800 or 900°C for different times. After this the boron and fluorine profiles were measured with SIMS. From the boron pro-
file the junction depth xj could be determined and the fluorine profile acted as a marker of the damaged region in the silicon because of the aggregation of the fluorine on the disk)cation loops. In Fig. 5 we have plotted the resulting leakage current density as a function of the relative position of the junction depth compared to the far edge of the dislocation loop region. All results lie on a common curve with the leakage current being reduced by four orders of magnitude when the dislocations are contained in thep + region.
2.2.3. Region Ill--enhanced diffusion The results from the preceding section and the leakage current shown in region IH of Fig. 1 would suggest that shallow low leakage junctions may be formed by simply placing the a - c boundary shallower than the subsequently formed p ~ -n junction, as is the case for the 30 keV Si +-preamorphized diodes. This indeed results in a low leakage current but the necessary 800 °C anneal step leads in this case to an increased junction depth as shown in Fig. 6. In contrast, the profiles for the 90 keV Si*-pre-amorphized substrates. where the as-implanted boron profile was remote from the a - c boundary, showed no sign of a significant enhancement of the junction depth. The appropriate model for these effects we believe is that, because boron diffuses with an appreciable interstitialcy component, excess silicon interstitials beyond the a - c boundary are responsible for the enhancement of the diffusion in the 30 keV Si ÷ -pre-amorphized samples. The
91
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.
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Fig. 6, SIMS profiles of boron measured in 30 and 90 kcV Si +-pre-amorphizcd substrates before and after annealing a! 8OO °C.
presence of these interstitials is beyond doubt since they have been invoked to explain the growth of the a - c boundary dislocation loops as discussed in Section 2.2.2, they have been predicted by Monte Carlo calculations [9] and they have been inferred from X-ray diffraction lattice parameter measurements [10]. From the growth of the a - c boundary dislocation loops it was concluded that the presence of the excess silicon interstitials was transitory and indeed the enhancement of the boron diffusion coefficient was more easily observed at lower times and temperatures such as 800 °C for 15 min than after 1 h at 900 °C. This also explains why no anomalies were seen when the junction was diffused through an initially distant a - c boundary. By the time the boron had diffused to this boundary the local excess silicon interstitial population had then been depleted by the formation of the dislocation loops. 3. Ion beam synthesis Ion beam synthesis is a relatively new method of making buried layers in crystalline materials by
Fig. 7. X T E M micrograph of a SIMOX structure together with oxygen and silicon SIMS signals.
means of a very high dose implantation of a species with the peak of the implanted profile far below the surface. In order to anneal the implantation damage during implantation, and so avoid amorphization, the implantation is performed at elevated temperatures. The best-known example is the formation of a buried oxide layer in silicon (SIMOX, acronym for separation by implanted oxygen), first demonstrated by lzumi et al. [11]. Later, buried SisN ~ [12], SiC [131 and crystalline silicide [141 layers were also studied. We will discuss here a number of aspects of damage created by implantation in relation to the final quality of SIMOX structures after annealing.
3.1. Microstructure of as-implanted ,~TMOX layers An XTEM micrograph of the structure obtained by 300 keV implantation of 2.5 x 1() ~ O + ions cm 2 at a relatively low flux of 1.5 #A cm 2 and in the channelling direction of the Si(100) substrate at a temperature of 570°C is shown in Fig. 7. A monocrystalline silicon layer 510 nm thick is separated from the substrate by an SiO 2 layer 350 nm thick. In the top silicon region the oxygen concentration is seen to increase with depth to the stoichiometric concen-
02 tration in SiO> The total oxygen concentration is far above the solubility limit and precipitation occurs in the form of oxide precipitates. Three regions can be distinguished. Region I is a perfect crystalline silicon layer 80 nm thick with a low oxygen concentration caused by out-diffusion during implantation (2(D0t) ~'2= 75 nm). Region II, 280 nm thick, shows a laminar contrast caused by SiO 2 precipitates about 2 nm in diameter in the crystalline silicon. These precipitates are ordered in a simple cubic lattice with a periodicity of 50 nm along the {100} directions of the silicon substrate [15]. The thickness of this region ii is limited by the maximum concentration of oxygen that can be accommodated in the superlattice structure (1.6x1021 O cm -3) [16]. When this concentration is exceeded, the microstructure of region III results, with much larger precipitates (10 nm) and a very high density (about 109 disk> cations cm-2) of defects. It is most important to note that this high defect concentration is confined to the region close to the buried oxide interface, whereas for samples with non-channelling implantation conditions these levels of defect densities are present in the whole as-implanted silicon top layer. After annealing these samples
tk)r 1 h at 1300 °C, all the precipitates in the former region I1 have been dissolved, resulting in the growth of the larger precipitates in region Iit: These growing precipitates pin the dislocations and prevent them from extending to the surface, and a prolonged anneal of 8 h at 1300 °C completely removes the precipitates and results in a top silicon layer with a dislocatkm density below 105 cm -' !171, whereas under normal non-channelling implantation conditions a value of 10 '~ cm ~' is obtained. The microstructure of the material below the buried oxide is shown in Fig. 8. In contrast to the upper SiO~ interface, which is very sharp, the lower interface is rough and appears to be built up flom an assembly of precipitates. At greater depth there is a dark band caused by a high density of{ 1 13} defects. Individual defects can be distinguished at a somewhat greater depth and a high resolution micrograph of these plate-like defects with a thickness of about 1 nm and an average diameter of 25 nm is shown in Fig. O. Bourret [18] has proposed that these defects are associated with a local phase transformation to hexagonal silicon ()wing to precipitation of silicon interstitials.
Fig. 8. Cross-sectional micrograph of the lower interface ot the as-implanted structure in Fig. 7.
Fig. 9. High resolution electron microscopy image of I 13 defects along the [111)]zone axis.
93
3.2. Relation between microstructure and implantation conditions Variations in the implantation conditions can give rise to large differences in the defect density in the resulting silicon on insulator (SOI) layer after annealing, as indicated in the preceding section. Two predominant processes that occur during the implantation are (1) Frenkel defect generation by the collision cascades of the oxygen ions
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Fig. IlL R a m a n spectra of the as-implanted samples with a 5 x 1() j 7 0 cm 2 implantation at 200 kcV.
2Si ~+ 2 0 i =(SiO~) ~+ Sii 4o N
46 A ~
in which the superscripts indicate substitutional(s) or interstitial(i) lattice positions. The first process is the dominant source of point defects since each oxygen ion typically generates 1 ( ) 2 10 -~ Frenkel pairs. The internal oxidation reaction gives rise to the generation of silicon interstitials because the volume has to be conserved. The dynamic equilibrium distribution of the point defects during implantation will depend on the implantation rate, the mobility of the defects determined by the substrate temperature, and the existence of sinks for the defects, i.e. the surface or extended defects such as precipitates or dislocations. The reaction rate for the internal oxidation will depend on the point defect concentration. A supersaturation of silicon interstitials will shift the reaction equilibrium such that the oxidation becomes much smaller and may even become zero. The roughness of the lower SiO~-Si interface as observed in Fig. 8 is likely to be caused by a supersaturation of interstitials. As soon as the oxide becomes a continuous layer, the excess interstitials caused by the oxidation reaction can no longer escape to the surface and the oxidation stops at the lower interface. In order to investigate further the effect of implantation conditions on the resulting microstructure, we have made 5 x l 0 1 7 O + cm 2 implantations at 200 keV into Sill00) substrates at 550 and 600 °C in either a random implantation direction or the {100/channelling direction. Implantation in the channelling direction will strongly reduce the generation rate of Frenkel pairs. After these implantations we have measured the stress conditions in the structure by
Raman measurements. Figure 10 shows the resulting Raman spectra. For the implants at 600°C the spectra are only slightly broadened compared with the spectrum for a virgin silicon sample. The 550 °C implantations resulted in a shift corresponding to a tensile strain in the top silicon layer, the strain in the randomly implanted sample being higher than in the sample with channelling implantation conditions. After an anneal at 1350°C for 4 h, the implanted oxygen coalesces into an almost continuous SiO2 layer with a thickness of about 100 nm. For the 550 °C randomly implanted sample a dislocation density of 3 x 1 0 7 c m 2 was found after this anneal, whereas for the 600°C randomly implanted sample the density was below the TEM detection limit of about 10 s cm 2 This illustrates that there is a relation between the point defect densities during implantation, the generated internal strain and the resulting precipitation of the silicon oxide. As a conclusion we may remark that a relatively good quality SOI structure formed by high dose oxygen implantation can be obtained by optimization of the implantation conditions. The build-up of strain during the implantation leads to the nucleation of the excess point defects into extended defects which ew)lve into dislocations during annealing. It has been shown that all these effects are related to the implantation conditions, although it has to be said that there is not yet a clear physical picture of what exactly is going on. Better physical understanding of the relations between the complex observed phenomena may eventually lead to further improvements in the quality of layers obtained with this promising technology of ion beam synthesis.
94
References 1 s. D. Brothertom J. R Gowers, N, D. Young, J. B. Clegg and J. R. Ayres, J. Appl. Phy.~'.,60 (1986) 3567. 2 F. Cembali, M. Servidori, E. Landi and S. Sohni, l'hys. Stares Solidi A, 94 ( 1986 ) 315. 3 P.L.F. Hemment, MRS 5)'rap. l'roc., 53 (1986) 207. 4 J. R. Ayres, S. D. Brotherton, J. B. Clegg and A. Gill, J. Appl. Phys., 62(1987)3628. 5 J. R. Ayres, S. D. Brotherton, J. B. Clegg, A. Gill and J. P. Gowers, Semicond. &'i. l~,chnol. 6 S. D. Brotherton, J. R. Ayres, J. B. Clegg and J. P. Gowers, Eh'ctron. Mater.. 18 (1989) 173. 7 S. D. Brotherton, J. R. Ayres, B. J. Goldsmith and A. Gill, SP1E Conf Proc., 797(1987) 26. 8 J. Thornton, K. C. Paus, R. P. Webb, J. H. Watson and G. R. Brooker, J. Phys. D: Appl. Phys,, 21 (1988) 334. 9 .1. Thornton, R. R Webb, J. H. Watson and K. C. Paus, Semicond. Sci. Technol.. ,~'(1988) 281.
10 K. Yamada, M. Kashiwagi and K. Taniguchi, Jpn. ,I. AI)pL Phys., 22-1 (1983) 157. 11 K. Izumi, M. Doken and 1-t. Aziyoshi, Electron. l, eu,, t4 1978) 593. 12 K. J. Reeson, Nut/. lttslrtmt. '~lel/mds ]4, 19-20 1987 269. 13 K. J. Reeson, P. L. F Hemment, J. Smemenos, J R. Davis and G. K. Celler. Inst. I'h)~s. CoJtf Set,, 87 ( 1987 J 427. 14 A. Ic;. White, K. T. Short, R. C. Dynes, J. P. Garno and J. M. Gibson, Appl. l'hys, l,ett.. :,0 (1987) 95. 15 A. H. van Ommen, B. H. Koek and M. R A, Viegers. ,tplft. Phys. Lett., 49(1986) 1062, 16 A. fl. van Ommen. MRSS),mp, Proc., 107!1988) 43. 17 A. It. van Ommen, H. J. lAgthart, ,1. Politik and M. P. A, Viegers, MRS 5~vmp. Proc., {Z~'~1987) I 19. 18 A. Bourret, hist. Phys. (ot{f25er, $7(1987) 39. 19 A. It. van O m m e n and M. P. A. Viegers, Appt. Sur]2 Sei, 30(1987) 383.
87
Recent Developments in Ion Implantation in Silicon J. A. PALS and S. D. BROTHERTON
l'hilips Research Laboratories, Redhil[, Surrey, R i l l 5HA (U.K.) A. H, VAN OMMEN, J. POLITIEK and H.J. L1GTHART
l~hilil)s Research Lahoratories, PO Box 80 O(X), NL-5600JA Eindhoven (The Netherhmds) (Received May 31, 1989:
Abstract
The influence o1"high dose implantations in silicon on defects is discussed by means of two exampies. The first is an investigation of the leakage currents caused by defects in veO, shallow p-n fimctions, obtained by low energy boron implantation into pre-amorphized n-type material. Preamorphizalion is used to prevent channelling effects in order to be able to make ve~ shallow junctions. Interstitial defects are shown to play an important role. The second example is a discussion of the defects occurring during vet3' high dose implantation used for ion beam ~ynthesis of buried oxide layers in silicon. The physical picture of this process is still rather unclear, but it is believed that again interstitial defects play an important role. 1. Introduction Ion implantation in silicon is now a wellestablished technique for introducing, in a controlled way, donors and acceptors in silicon slices. The requirement for shallow junctions in submicron very-large-scale integrated (VLSI)circuits has aroused interest in low energy implantations. Channelling effects, especially for the lighter implanted atoms such as boron, may cause the junction depths to be larger than predicted from the ion implantation range tables. An effective technique for suppressing these channelling effects is the amorphization of the surface region of the crystal by implantation of a high dose of an inert ion. Solid state phase epitaxial regrowth at temperatures as low as 550°C is then used for recrystallization and activation of the implanted dopant atoms without any appreciable redistribution. However, this procedure may result in diodes with poor electrical characteristics owing to the introduction of defects l 1,2]. 0921-5107/89/$3.50
Another development in which high dose implantations are being used is the ion beam synthesis of buried layers [3]. The best known example is the formation of a buried oxide layer in silicon by a very high dose oxygen implantation. There is much interest in such structures with a thin silicon layer on top of an insulating layer, for instance for making very high speed complementary metal-oxide-semiconductor (CMOS) circuits. The implantation is done at elevated temperatures to prevent the remaining top silicon layer from amorphizing. Again the damage created by the high dose implantation is a very important factor in determining the quality of the resulting Si-SiO~ structure. It is the purpose of this paper to present a number of experimental results on the creation of damage by high dose implantations followed by annealing steps used for amorphization or ion beam synthesis of buried layers. 2. Pre-amorphized and annealed layers
2.1. Sample preparation The samples used for the investigation of the effects of pre-amorphization were obtained by implanting n-type Si(100)-oriented substrates with Si +, Ge + or Sn + ions. The implantation conditions and the thickness of the resulting amorphous layers, measured with Rutherford backscattering (RBS) or cross-sectional transmission electron microscopy (XTEM), are summarized in Table 1. The Si + implantations have been made with different energies to vary the position of the a - c boundary with respect to the metallurgical p +-n junction. This junction was made after pre-amorphization by implanting a dose of 5 x 10 H~BF, + ions at 25 keV. The as-implanted junction depth determined by secondary ion mass © Elsevier Sequoia/Printed in The Netherhmds
S~
'•\
TABLE I
/on
Dose ',cm
Si-
Ge~ Sn +
2x 2× 2× 2× 1x ~8 x 5x
Lnergy :',
(keV)
10 I" 11)~5 10 I~ 10 ~" 1() I~ it) ~4 1014
30 50 7(1 90 200 180 65
Amorphous layer thickness(A)
t0-2
640,720* 1(16(I 1400 1740, 176()* 2000 1440, 1420"
10 -3
?
x o a Q
30 50 70 90
o\% % \\ \\ %
10-4
*TEM measurements; all others RBS. u
spectrometry (SIMS) profiling was approximately 750 A, i.e. smaller than the a - c boundary depth for all pre-amorphizations with the exception of the 30 keV Si + implantation. Furnace annealing at temperatures between 600 and 1000 °C was used to regrow the amorphous layers and to activate the implanted boron. Finally the samples were annealed at 450 °C in wet N 2 to remove fast Si-SiO 2 interface states and metallized with cold aluminium. The diodes were circular with field plates to control surface generation currents. The SIMS profiles of the boron after annealing showed no anomalous diffusion in the Si + preamorphized layers made with energies of 50 keV or above, as will be discussed later.
2.2. Electrical characterization of the diodes Leakage current measurements of the diodes were made and the results of the Si+-amorphized diodes are shown in Fig. 1. Three regions can be distinguished. Region I, for anneal temperatures of 600 and 700 °C, shows a very high reverse-bias leakage current density with the current rising strongly with reverse bias. Similar results were obtained for the Ge*- and Sn +-implanted diodes [4, 5]. Annealing at or above 800 °C resulted in lower leakage current densities, saturating at a level of about 10-5 A cm-2 for the diodes with the higher implantation energies (region II) or becoming much lower for the diodes with very shallow pre-amorphized layers (region IIt). Room temperature C - V profiles obtained from diodes with a 10 l~ cm 3 substrate doping level and amorphized by the three different implants and annealed at 600, 700 or 800 °C are shown in Fig. 2. In each case the lower regrowth temperatures of 600 and 700 °C resulted in the presence of a high concentration of excess donors, which were removed by the 800 °C anneal. The precise inter-
10-6 i
I
1
i
i
10-7
1°-~l
680
.7.0.0. . . . . ~;o Anneal temperature
9OO (~C)
Fig. 1. Reverse-bias leakage current densities measured as a function of the anneal temperature in Si+-pre-amorphized p +-n diodes.
pretation of the profiles has to be done with care because of the deep level character of the excess donors [6]. We will now discuss in more detail the defects present in the diodes in regions I, II and llI in Fig. 1.
2.2.1. Region I deep level defects To characterize further the defects present m the amorphized and regrown layers, deep level transient spectroscopy (DLTS, measurements have been made. A typical set of curves is shown in Fig. 3 for samples amorphized with Si +. Ge T and Sn + and annealed at 700°C. Four defect levels labelled A-D were found in the Si*implanted sample. These levels were also found in the Ge*-implanted sample, but in the Sn-implanted sample only A and D were found unambiguously, although level C could be found in samples annealed at 600 °C. The similarity of the results suggest that the basic defects are identical for all three amorphizing species. Spatial profiles of the deep level defects have been obtained by measuring the DLTS curves as a function of reverse bias and a result is given in
89 12-
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">~ 0e- 0 . 8 D
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700~C 'E u
D T= 7 0 0
u~
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800 °c
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Ge* ,6OO °C
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)",
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ml th
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I
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Fig. 2. Room temperature C - I / p r o f i l e s obtained from p+ n diodes pre-amorphized with (a) 90 keV Si +, (b) Ge+ and (c)
I
}
J
[
150 200 Temperature
I
~
250 (K)
I
J
300
Fig. 3. DUI'S c u r v e s obtained from p * - n d i o d e s a m o r p h i z e d w i t h (a) 70 kcV Si + , (b) G e + and (c)Sn
pre-
Sn + .
Fig. 4 for an Si+-pre-amorphized sample. Also indicated are the location of the junction and the a - c boundary dislocation loops. The deep level defects are evidently lying deeper than the ct-c boundary and in fact their concentration profile is very close to the implanted silicon profile obtained with Lindhard, Scharff and Schiott (LSS) range tables. Similar results were again obtained with the Ge +- and Sn~-amorphized samples [4, 5J, both in terms of the location of the traps beyond the a - c boundary and the closeness of their profiles to the implanted impurity profiles. These results suggest strongly that the defects are fundamental damage-related centres and may be connected to the excess silicon interstitial in the region beyond the a - c boundary. The thermal activation energies of the different centres A-D were found to be about 0.30, 0.46, 0.49 and 0.65 eV respectively; the values were dependent on the applied reverse bias on the
diodes and indicated field-enhanced emission 16J. From these trap activation energies it is concluded that the near-midgap centre D could be responsible for the large reverse-bias leakage currents measured in these structures. This was confirmed by examination of the dependence of the leakage current density on the concentration of the D-trap after annealing the samples for different periods of time to progressively remove that centre. A good correlation between trap density and leakage current was found [7].
2. 2.2. Region ll--extended defects Following an anneal at 800°C or above, the deep level centres discussed above were removed and the leakage current densities were reduced to approximately 10 -5 A cm-z for the diodes with the ct-c boundary beyond the p-n junction depth Xi .
XTEM micrographs of 90 keV Si+-amor phized and regrown material [i] have revealed a
90 1 0 ~7
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zh
Edge of pre-omorphised region
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Symbol
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Fig. 5. Reverse-bias leakage current density as a function of the relative position of the junction depth x i to the far edge _%~ of the band of dislocation loops.
o
A
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o
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t
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Fig. 4, Deep level d o n o r profiles obtained from a 90 keV Si + -pre-amorphized diode annealed at 700 °C.
band of dislocation loops, 50(/_+ 200 A wide, immediately beneath the original amorphous layer. These loops were observed to grow in size and decrease in number density with increasing anneal temperatures. With increasing anneal temperature the total silicon atom density in these loops initially increased and finally saturated at approximately 3 x 1014 cm -2 for temperatures above 800 °C. This loop growth is interpreted to be caused by the flow of excess silicon interstitials from beyond the a - c boundary. Monte Carlo calculations of Thornton and coworkers [8, 9] have predicted a silicon interstitial concentration caused by both the directly implanted and the recoil-implanted atoms of about 1015 c m -2, which is rather close to the experimentally found density in the loops. These a - c boundary dislocation loops were demonstrated to be the cause of the leakage currents in region II of Fig. 1. A number of diodes were annealed at 800 or 900°C for different times. After this the boron and fluorine profiles were measured with SIMS. From the boron pro-
file the junction depth xj could be determined and the fluorine profile acted as a marker of the damaged region in the silicon because of the aggregation of the fluorine on the disk)cation loops. In Fig. 5 we have plotted the resulting leakage current density as a function of the relative position of the junction depth compared to the far edge of the dislocation loop region. All results lie on a common curve with the leakage current being reduced by four orders of magnitude when the dislocations are contained in thep + region.
2.2.3. Region Ill--enhanced diffusion The results from the preceding section and the leakage current shown in region IH of Fig. 1 would suggest that shallow low leakage junctions may be formed by simply placing the a - c boundary shallower than the subsequently formed p ~ -n junction, as is the case for the 30 keV Si +-preamorphized diodes. This indeed results in a low leakage current but the necessary 800 °C anneal step leads in this case to an increased junction depth as shown in Fig. 6. In contrast, the profiles for the 90 keV Si*-pre-amorphized substrates. where the as-implanted boron profile was remote from the a - c boundary, showed no sign of a significant enhancement of the junction depth. The appropriate model for these effects we believe is that, because boron diffuses with an appreciable interstitialcy component, excess silicon interstitials beyond the a - c boundary are responsible for the enhancement of the diffusion in the 30 keV Si ÷ -pre-amorphized samples. The
91
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L
o
o
30keY Si+I800°C - - - - 90keV Si ÷]
\\
.
.
.
.
.
90key Si';
.
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I
I
I
I
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I J
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Fig. 6, SIMS profiles of boron measured in 30 and 90 kcV Si +-pre-amorphizcd substrates before and after annealing a! 8OO °C.
presence of these interstitials is beyond doubt since they have been invoked to explain the growth of the a - c boundary dislocation loops as discussed in Section 2.2.2, they have been predicted by Monte Carlo calculations [9] and they have been inferred from X-ray diffraction lattice parameter measurements [10]. From the growth of the a - c boundary dislocation loops it was concluded that the presence of the excess silicon interstitials was transitory and indeed the enhancement of the boron diffusion coefficient was more easily observed at lower times and temperatures such as 800 °C for 15 min than after 1 h at 900 °C. This also explains why no anomalies were seen when the junction was diffused through an initially distant a - c boundary. By the time the boron had diffused to this boundary the local excess silicon interstitial population had then been depleted by the formation of the dislocation loops. 3. Ion beam synthesis Ion beam synthesis is a relatively new method of making buried layers in crystalline materials by
Fig. 7. X T E M micrograph of a SIMOX structure together with oxygen and silicon SIMS signals.
means of a very high dose implantation of a species with the peak of the implanted profile far below the surface. In order to anneal the implantation damage during implantation, and so avoid amorphization, the implantation is performed at elevated temperatures. The best-known example is the formation of a buried oxide layer in silicon (SIMOX, acronym for separation by implanted oxygen), first demonstrated by lzumi et al. [11]. Later, buried SisN ~ [12], SiC [131 and crystalline silicide [141 layers were also studied. We will discuss here a number of aspects of damage created by implantation in relation to the final quality of SIMOX structures after annealing.
3.1. Microstructure of as-implanted ,~TMOX layers An XTEM micrograph of the structure obtained by 300 keV implantation of 2.5 x 1() ~ O + ions cm 2 at a relatively low flux of 1.5 #A cm 2 and in the channelling direction of the Si(100) substrate at a temperature of 570°C is shown in Fig. 7. A monocrystalline silicon layer 510 nm thick is separated from the substrate by an SiO 2 layer 350 nm thick. In the top silicon region the oxygen concentration is seen to increase with depth to the stoichiometric concen-
02 tration in SiO> The total oxygen concentration is far above the solubility limit and precipitation occurs in the form of oxide precipitates. Three regions can be distinguished. Region I is a perfect crystalline silicon layer 80 nm thick with a low oxygen concentration caused by out-diffusion during implantation (2(D0t) ~'2= 75 nm). Region II, 280 nm thick, shows a laminar contrast caused by SiO 2 precipitates about 2 nm in diameter in the crystalline silicon. These precipitates are ordered in a simple cubic lattice with a periodicity of 50 nm along the {100} directions of the silicon substrate [15]. The thickness of this region ii is limited by the maximum concentration of oxygen that can be accommodated in the superlattice structure (1.6x1021 O cm -3) [16]. When this concentration is exceeded, the microstructure of region III results, with much larger precipitates (10 nm) and a very high density (about 109 disk> cations cm-2) of defects. It is most important to note that this high defect concentration is confined to the region close to the buried oxide interface, whereas for samples with non-channelling implantation conditions these levels of defect densities are present in the whole as-implanted silicon top layer. After annealing these samples
tk)r 1 h at 1300 °C, all the precipitates in the former region I1 have been dissolved, resulting in the growth of the larger precipitates in region Iit: These growing precipitates pin the dislocations and prevent them from extending to the surface, and a prolonged anneal of 8 h at 1300 °C completely removes the precipitates and results in a top silicon layer with a dislocatkm density below 105 cm -' !171, whereas under normal non-channelling implantation conditions a value of 10 '~ cm ~' is obtained. The microstructure of the material below the buried oxide is shown in Fig. 8. In contrast to the upper SiO~ interface, which is very sharp, the lower interface is rough and appears to be built up flom an assembly of precipitates. At greater depth there is a dark band caused by a high density of{ 1 13} defects. Individual defects can be distinguished at a somewhat greater depth and a high resolution micrograph of these plate-like defects with a thickness of about 1 nm and an average diameter of 25 nm is shown in Fig. O. Bourret [18] has proposed that these defects are associated with a local phase transformation to hexagonal silicon ()wing to precipitation of silicon interstitials.
Fig. 8. Cross-sectional micrograph of the lower interface ot the as-implanted structure in Fig. 7.
Fig. 9. High resolution electron microscopy image of I 13 defects along the [111)]zone axis.
93
3.2. Relation between microstructure and implantation conditions Variations in the implantation conditions can give rise to large differences in the defect density in the resulting silicon on insulator (SOI) layer after annealing, as indicated in the preceding section. Two predominant processes that occur during the implantation are (1) Frenkel defect generation by the collision cascades of the oxygen ions
SP = Si~+ V
~ 2 (200 keY) ~,-,~ u .' c m - ~ / u u / - Si
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300
1
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E 200
,,//
cc 100 ~ ' " J / 51
'
/ ~' 520
530
Stokes shift (cm ~)
(2) reaction of the implanted oxygen ions to SiO~
Fig. IlL R a m a n spectra of the as-implanted samples with a 5 x 1() j 7 0 cm 2 implantation at 200 kcV.
2Si ~+ 2 0 i =(SiO~) ~+ Sii 4o N
46 A ~
in which the superscripts indicate substitutional(s) or interstitial(i) lattice positions. The first process is the dominant source of point defects since each oxygen ion typically generates 1 ( ) 2 10 -~ Frenkel pairs. The internal oxidation reaction gives rise to the generation of silicon interstitials because the volume has to be conserved. The dynamic equilibrium distribution of the point defects during implantation will depend on the implantation rate, the mobility of the defects determined by the substrate temperature, and the existence of sinks for the defects, i.e. the surface or extended defects such as precipitates or dislocations. The reaction rate for the internal oxidation will depend on the point defect concentration. A supersaturation of silicon interstitials will shift the reaction equilibrium such that the oxidation becomes much smaller and may even become zero. The roughness of the lower SiO~-Si interface as observed in Fig. 8 is likely to be caused by a supersaturation of interstitials. As soon as the oxide becomes a continuous layer, the excess interstitials caused by the oxidation reaction can no longer escape to the surface and the oxidation stops at the lower interface. In order to investigate further the effect of implantation conditions on the resulting microstructure, we have made 5 x l 0 1 7 O + cm 2 implantations at 200 keV into Sill00) substrates at 550 and 600 °C in either a random implantation direction or the {100/channelling direction. Implantation in the channelling direction will strongly reduce the generation rate of Frenkel pairs. After these implantations we have measured the stress conditions in the structure by
Raman measurements. Figure 10 shows the resulting Raman spectra. For the implants at 600°C the spectra are only slightly broadened compared with the spectrum for a virgin silicon sample. The 550 °C implantations resulted in a shift corresponding to a tensile strain in the top silicon layer, the strain in the randomly implanted sample being higher than in the sample with channelling implantation conditions. After an anneal at 1350°C for 4 h, the implanted oxygen coalesces into an almost continuous SiO2 layer with a thickness of about 100 nm. For the 550 °C randomly implanted sample a dislocation density of 3 x 1 0 7 c m 2 was found after this anneal, whereas for the 600°C randomly implanted sample the density was below the TEM detection limit of about 10 s cm 2 This illustrates that there is a relation between the point defect densities during implantation, the generated internal strain and the resulting precipitation of the silicon oxide. As a conclusion we may remark that a relatively good quality SOI structure formed by high dose oxygen implantation can be obtained by optimization of the implantation conditions. The build-up of strain during the implantation leads to the nucleation of the excess point defects into extended defects which ew)lve into dislocations during annealing. It has been shown that all these effects are related to the implantation conditions, although it has to be said that there is not yet a clear physical picture of what exactly is going on. Better physical understanding of the relations between the complex observed phenomena may eventually lead to further improvements in the quality of layers obtained with this promising technology of ion beam synthesis.
94
References 1 s. D. Brothertom J. R Gowers, N, D. Young, J. B. Clegg and J. R. Ayres, J. Appl. Phy.~'.,60 (1986) 3567. 2 F. Cembali, M. Servidori, E. Landi and S. Sohni, l'hys. Stares Solidi A, 94 ( 1986 ) 315. 3 P.L.F. Hemment, MRS 5)'rap. l'roc., 53 (1986) 207. 4 J. R. Ayres, S. D. Brotherton, J. B. Clegg and A. Gill, J. Appl. Phys., 62(1987)3628. 5 J. R. Ayres, S. D. Brotherton, J. B. Clegg, A. Gill and J. P. Gowers, Semicond. &'i. l~,chnol. 6 S. D. Brotherton, J. R. Ayres, J. B. Clegg and J. P. Gowers, Eh'ctron. Mater.. 18 (1989) 173. 7 S. D. Brotherton, J. R. Ayres, B. J. Goldsmith and A. Gill, SP1E Conf Proc., 797(1987) 26. 8 J. Thornton, K. C. Paus, R. P. Webb, J. H. Watson and G. R. Brooker, J. Phys. D: Appl. Phys,, 21 (1988) 334. 9 .1. Thornton, R. R Webb, J. H. Watson and K. C. Paus, Semicond. Sci. Technol.. ,~'(1988) 281.
10 K. Yamada, M. Kashiwagi and K. Taniguchi, Jpn. ,I. AI)pL Phys., 22-1 (1983) 157. 11 K. Izumi, M. Doken and 1-t. Aziyoshi, Electron. l, eu,, t4 1978) 593. 12 K. J. Reeson, Nut/. lttslrtmt. '~lel/mds ]4, 19-20 1987 269. 13 K. J. Reeson, P. L. F Hemment, J. Smemenos, J R. Davis and G. K. Celler. Inst. I'h)~s. CoJtf Set,, 87 ( 1987 J 427. 14 A. Ic;. White, K. T. Short, R. C. Dynes, J. P. Garno and J. M. Gibson, Appl. l'hys, l,ett.. :,0 (1987) 95. 15 A. H. van Ommen, B. H. Koek and M. R A, Viegers. ,tplft. Phys. Lett., 49(1986) 1062, 16 A. fl. van Ommen. MRSS),mp, Proc., 107!1988) 43. 17 A. It. van Ommen, H. J. lAgthart, ,1. Politik and M. P. A, Viegers, MRS 5~vmp. Proc., {Z~'~1987) I 19. 18 A. Bourret, hist. Phys. (ot{f25er, $7(1987) 39. 19 A. It. van O m m e n and M. P. A. Viegers, Appt. Sur]2 Sei, 30(1987) 383.