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Ion implantation, ion beam mixing, and annealing studies of metals in Al2O3, SiC and Si3N4
Nuclear Instruments and Methods in Physics Research Bl (1984) 167-175 North-Holland, Amsterdam
167
Section I. Crystalline oxides ION IMPLANTATION, ION BEAM MIXING, AND ANNEALING Al,Os, SIC AND Si,N, * B.R. APPLETON, C.J. McHARGUE,
STUDIES
OF METALS
IN
H. NARAMOTO + C.W. WHITE, O.W. HOLLAND, ++ G. FARLOW, J. NARAYAN and J.M. WILLIAMS
Solid State Division, Oak Ridge National Laboratoty,
Oak Ridge, TN 37830, USA
Ion scattering/channeling, TEM, EPR and optical microscopy are utilized to determine the structural modifications of ion implanted and annealed Al,O,, SIC, and Si,N,, and to correlate these modifications to surface mechanical property measurements. Ion beam mixing is also studied for inducing increased adherence of metal films on these insulators.
1. Intmluction Ion implantation doping, ion beam mixing, and pulsed-laser annealing are being increasingly utilized to alter the near-surface properties of materials. The attractiveness of these methods arises because they are nonequilibrium processing techniques which can lead to new materials properties [l-5]. Previously these techniques have been used primarily for altering the surface properties of semiconductor and metal or alloy systems [l-3]. Some of the contributing phenomena responsible for the success of these techniques include (in the case of ions) radiation damage, cascade mixing, impurity introduction, radiation enhanced diffusion and compositional and/or structural modifications; and (in the case of pulsed laser annealing) liquid or solid diffusion (mixing), extremely rapid quenching, rapid crystallization, quenched-in phase changes, etc. When one recognizes that insulating materials potentially offer a much broader range of responses to all these phenomena than either metals or semiconducting materials it is surprising that insulators have not been more widely investigated. Insulators can have a wide variety of starting properties. The composition of many insulating materials is complex. Bonding structures can range from ionic to covalent to near-metallic; and the atomic structure can range from crystalline to glassy. Thus, these materials are particularly sensitive to the chemical and electrical
* Research sponsored by the Division of Materials Sciences, U.S. Department of Energy under contract W-7405-eng-26 with Union Carbide Corporation. + JAERI, Tokai, Japan. ++ Metals and Ceramics Division, ORNL. 0168-583X/84/$03.00 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
B.V.
character of impurities or dopants. The alterations possible in insulating materials are correpondingly complex. Ion damage can occur as a result of ionizing as well as displacement collisions; diffusion can have electronic and chemical driving forces; and phase changes can be induced by damage or thermal processing. Previous investigations of insulating and ceramic materials have emphasized the use of ion bombardment or ion implantation doping on the optical properties [4,5], induced phase changes [5,6], electrical and/or chemical properties [5.7], and ion induced disorder [S-15]. Some work has been reported on laser processing of Sic [16]. This paper reports on a continuing study of the surface alterations of insulators and ceramics achievable with ion implantation doping and thermal annealing, and ion-beam mixing with the emphasis on surface mechanical properties [17-221. Related work on pulsed-laser processing of insulators will be presented in a later paper. The intent of this work is to characterize the near surface microstructural alterations induced by ion beam processing, and to correlate these with companion mechanical property measurements. The analysis techniques used include ion scattering, ion channeling, transmission electron microscopy (TEM), electron paramagnetic resonance (EPR) and optical microscopy; the associated mechanical property studies were microhardness and fracture toughness measurements. In section 2 the effects of Cr, Zr, and Ti implantation into Al,O, followed by scheduled thermal annealing are presented. These results demonstrate the utility of correlated structural and mechanical prdperty measurements, and the versatile nature of surface modifications of ceramics. Section 3 presents similar results for Sic. In section 4, results for ion-beam mixing are presented. This paper is more in the nature of an overview and the reader is referred to previous publications for specific experimental details [17-221. I. CRYSTALLINE OXIDES
B.R. Appleton et al. / Metals .in AI,O,,
168
2. Ion implantation and thermal annealing of Al,O,
damage free samples. Half of each sample was retained as virgin reference material and the remaining half was implanted with Cr, Zr, or Ti at room temperature with the ion beam incident at 7’ from the crystal normal at fluences ranging from 10r5-10” cme2 and energies from 100-300 keV. Thermal anneals were performed in
Single crystals of Al,O, of high purity (< 100 ppm total) and low dislocation density (103-lo4 cm-‘) were cut to within f2’ of (0001) and (l?lO), and polished and annealed at 12OO’C in air for 120 h to produce
0 DEPTH (pm) 0.4
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In Q-AI~O~
Fig. 1. Thermal annealing behavior of Cr (300 keV, 1 x 101’/cm2) implanted cr-Al,O,. as-implanted and after various annealing steps to 1500°C.
Ion scattering and channeling results for virgin,
B.R. Appleton et al. / Metals in AI@3, Sic ond Si,N,
either air or N, at temperatures from 800°C to 1600°C for 1 h. 2, I. Structural aiterations and mechanical property measurements oj Cr implanted AI,O,
The history of the ion damage and its recovery, the concentration profiles of the implanted dopant, and the specific lattice locations of defects and impurities for Cr implanted AlsO, were determined from detailed ion scattering/channeling measurements [21]. Representative results are shown in fig. 1 for a high dose implant (1 X lOI Cr/cm2) after a series of anneals in air. Analysis of such ion scattering/channeling data revealed the folowing information: (1) Contrary to previous reports [6], even at these high implantation concentrations, Al,O, is not turned amorphous (fig. la, confirmed by TEM). In Al,O, the damage saturates because of dynamic damage recovery during the room temperature implants. (2) Damage recovering begins selectively in the Al sublattice at a temperature - 8OO*C (figs. la and lb) with no apparent change in the 0 sublattice or Cr distribution, and is complete at temperatures -z 12OO’C. (3) Damage recovery begins in the 0 sublattice at - 1000°C but little change occurs in the Cr substitutionality. (4) Significant incorporation of Cr into substitutional lattice sites begins from 120@-1300°C (fig. le and detailed angular scans).
169
(5) As annealing temperatures are increased to 1500°C, Cr becomes highly (> 95%) substitutional (fig. Id). For lower initial implantation doses more complete damage recovery occurs but dechanneling from Cr-induced lattice strains always causes yields which exceed the virgin-crystal yields. (6) Detailed angular scans 1211show that the Cr goes substitutional in the Al sublattice only. (7) Furthermore, electron paramagnetic resonance measurements [21] show that most if not all the substitutional Cr is in the 3 + valence state. (8) Redistribution of the Cr (in depth) begins to occur for anneals of - 16OO”C,presumably by substitutional diffusion. The corresponding effects on surface mechanical properties caused by these structural and compositional alterations have been reported earlier (17,181, where it was found that Cr impl~tation into Al,O, increased the surface hardness by 30-40% relative to the unimplanted region and increased the fracture toughness by 15%. The effects of thermal annealing in air on the relative hardness of an Also, single crystal implanted with 2 x lOI6 Cr/c& is shown in fig. 2. As annealing is carried to higher temperatures the relative hardness decreases from 1.35 to 1.1 in the range 800-1000°C and then remains constant at 1.1 to temperatures of 1600°C where it decreases to 1.0. When correlated with the observed structural modifications these combined results suggest that annealing in air makes oxygen available for damage recovery through intersitial oxygen
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1400
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Fig. 2. Surface microhardness of AI,& annealing temperatures for 1 h in air.
TEMPERATURE
1000
(“C)
implanted with Cr, Zr, and Ti relative to unimplanted Ai&& as a function of thermal
I. CRYSTALLINE OXIDES
170
B.R. Appleton
et
al. / Metals in AI,O,, SIC and Si,N,
diffusion. The displaced Al starts reordering - 8OO’C and is completely reordered - 1200°C. The oxygen sublattice annealing begins at temperatures > 1000°C and the first signs of Cr substitutionality appear. The fraction of implanted Cr acquiring Cr3+ substitutional sites continues to increase, with little change in the initial depth distribution, up to 1500°C. It appears from the hardness measurements that as the Al and 0 damage recovers, Al,O,-Cr,O, solid solutions are formed beginning at 1000°C and persist to - 15OO’C [21]. This conclusion is supported by the fact that the relative hardness is constant at - 1.1 in this temperature range in agreement with measurements on AI,O,-Cr,O, solid solutions grown from the melt [23,24]. Detailed angular scans taken after annealing at an intermediate stage (13OO’C) utilizing ion channeling suggest that interstitial oxygen is trapped by interstitial Cr before complete recovery occurs [21]. This is an important observation since the interstitial oxygen defect in Al,O, has been very difficult to study by other methods because of its high mobility. Above 15OO’C the Cr depth distribution changes, apparently from normal substitutional diffusion into the bulk Al,O,, and the relative hardness drops to 1.0. 2.2. Implantation of Zr in AI,O, and thermal annealing in air The importance of the chemical compatibility of the implanted species on the induced structural allterations and surface mechanical properties can be illustrated by contrasting- the results for Zr implanted in AlzO, with those for Cr [17,18,21]. Analysis for Zr, similar to those already discussed for Cr, provided the following evahrations. 1) The AlaO, single crystals could not be made amorphous by room temperature implantation. 2) Damage recovery began in the Al sublattice at 800°C as with Cr. 3) However, damage recovery did not occur in the 0 sublattice until - 1300°C. 4) Even after annealing to temperatures > 1600°C substantial disorder remained in the Al,O, lattice. 5) Zr showed no substitutional tendencies for annealing to 16OO’C in air. 6) Very little redistribution in depth occurred for annealing temperatures to 16OO’C. A reexamination of fig. 2 shows that these structural differences are mirrored in the relative hardness measurements for Al,O, implanted with 2 X 1016 Zr/cm2 and annealed in air. The Zr implant induces increases in hardness and fracture toughness which remain high even to very high temperatures. This is a significant observation for applications where it would be desirable to use hardened Al,O, as a ceramic bearing or wear material operating at high temperatures. An analysis
consistent with these correlated measurements is that the implanted Zr forms precipitates on annealing, that these precipitates account for the residual damage in the Al,O, and the lack of Zr substitutional&y, and that the hardness is due to precipitation hardening. 2.3. Comparison of thermal annealing in air and H2 of AI,O, implanted with Ti The potential versatility offered by ion-implantation doping of ceramic materials that was discussed in the introduction and illustrated by the results for Cr and Zr implants in Al,O, is further emphasized by recent Ti implantation studies. Previous preliminary investigations showed a complex anisotropic redistribution behavior for Ti implanted in Al,O, and annealed in air [21]. This effect is summarized by the contrasting behaviors of (0001) (fig. 3) and (1210) (fig. 4) Al,O, single crystals implanted with Ti and thermally annealed under identical conditions. Analysis of the ion scattering and channeling results in fig. 3 shows that annealing the (0001) crystal to - 1300°C results in a double-peaked distribution of the Ti with and accumulation at the surface, and the Ti is - 60% substitutional along (0001). Further annealing to 1500°C results in nearly complete damage recovery in the Al,O, lattice, a substantial redistribution of the Ti, but a reduction in the amount of substitutional Ti. These results suggest, and optical microscopy confimed, that needle-like precipitates (probably TiO,) formed at the surface in the range 1300-1500°c. In contrast, results from the (1210) measurements (fig. 4) showed no evidence for redistribution of Ti toward the surface, or for Ti substitutionality after annealing in air to 13OO’C. Annealing to 15OO’C showed increased damage recovery and some diffusion of Ti but no precipitates were observed at the surface. These results suggest that there is a preferential redistribution along the (0001) direction in Al,O,. To investigate the effects associated with the annealing environment, (0001) Al,O, samples were implanted with 4 x 1016 Ti/cm2 (150 keV) and annealed at 13OO’C and 1400°C in H, gas. The ion scattering/channeling results are shown in fig. 5 and should be contrasted with the (0001) spectra of fig. 3 for annealing in air. As with the anneals in air, annealing to 13OO’C in H, caused substantial damage recovery in the Al,O, lattice, a redistribution of the Ti toward the surface, and incorporation of Ti into substitutional lattice sites. There are subtle differences, however, since there is no accumulation of Ti at the surface for the Hz anneals and the fraction of substitutional Ti drops uniformly form 45% at the surface to zero at the depth of the original implant. It is likely that O,, present in the anneals of fig. 3, stabilized the Ti reaching the surface by formation of TiO, and thus led to the observed accumulation.
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Upon annealing to 1400°C in H, (WO”C less than the comparable spectrum in fig. 3) the Ti diffused to much greater depths, and has developed a profile consistent with normal diffusion. The optical micrastructures of
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the H,-anneakd samples developed with annealing as follows: 1) After 12OPC Eine precipitates formed. 2) After 13OCPCthe fine precipitates grew and new
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and annealed in air. I. CRYSTALLINE
OXIDES
B.R. Appleton et al. / Metals in AI,O,, SC and Si,N,
172 0 DEPTH
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Fig. 5. Behavior of (0001) oriented a-Al,O,
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needle-like precipitates surrounded by regions devoid of fine precipitates formed. 3) After 14OO’C both fine and needle-like precipitates disappeared and were replaced with a few precipitates with a different optical contrast. 4) After 1600°C only a few smal precipitates remained. These observations and the spectra of fig. 5 suggest that small precipitates form initally (- 12OO’C) but ripen and give way to larger precipitates at higher anneals (- 1400°C). The fact that these precipitates dissolve easily at higher temperatures, in contrast to those formed in air anneals, is probably due to the lack of stabilizing oxygen. These are, however, speculations, and experiments are presently in progress to identify the phase formation and stability of these systems. Referring again to fig. 2, implantation of (0001) Al,O, to 1 x lOI Ti/cm’ produces an increase in hardness to > 1.25. Annealing in air at temperatures of - 800-lOOO”C, causes a decrease, but then the hardness increases again to - 1.25 at 12OO’C and remains at this level to > 1500°C. The initial increase in hardness may be due to small TiO, precipitates formed and stabilized in the damaged Al,O, lattice. The hardness begins to decrease at the annealing temperatures where the Al (- 800°C) and 0 (- 1000°C) sublattices begin to anneal [21]. At greater temperatures (13OO’C) the Ti moves
and annealed in H, gas.
toward the surface, becomes partially substitutional, and accumulates with the probable formation of TiO, precipitates again and the hardness increases. The higher temperature annealing (- lSOO°C) causes redistribution and decreased substitutionality probably due to the ripening of the TiO, precipitates. Measurements are presently in progress to evaluate the effects of H, annealing on the hardness of Ti implanted Al,O,.
3. Ion implantation doping, damage and thermal annealing studies in SIC Investigations of the changes in surface mechanical properties due to ion implantation doping were extended to Sic because it is a refractory ceramic which figures prominently in a number of energy options of interest to the Department of Energy, and also because, like silicon, it has a covalent bond structure and should have different structural modification properties than say, Al,O, [19,20]. Initial studies showed that implantation of 2 X 1016 Cr/cm’ (280 keV) into SIC single crystals increased the fracture toughness by - 101, a result similar to that observed for Al,O,, but caused a decrease of 25% in surface hardness, which is opposite to the change for Al,O, [19]. Furthermore, scratch wear tests showed a decrease in the ratio of tangential to
B.R. Appleton et al. / Metals in AI,O,, N DOSE (cm-21 10’4 rlTrll[
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402’
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AREA kV/cm2)
Fig. 6. Step height increases of Sic implanted with N+ and Cr+ as a function of total damage energy per unit area (see
text).
normal forces (> 25%) for the implanted surface and the scratch was much smoother in appearance. Surface profilometry measurements revealed that ion implantation of Cr caused a noticable step height between the unimplanted and implanted regions, corresponding to a - 20% swelling for implanted single-crystalline Sic samples and - 30% for polycrystalline samples [19]. Changes in hardness similar to those mentioned above have been reported for N+ implantations in Sic but swelling measurements were not made [25]. Detailed investigations were performed to identify
5ooo
8
SC and Si,N,
173
the structural alterations responsible for the swelling and surface mechanical property changes. These studies utilized ion scattering/channeling techniques and transmission electron microscopy and revealed a close correlation between ion induced damage and the large swelling [19,20]. Scheduled implantations of both Cr and Ni were made up to doses of 3 X 10n’ cmm2 and 8 x 1016 cmm2 respectively and the induced damage was monitored at incremental doses in situ by ion channeling. Ion energies of 62 keV N and 260 keV Cr were chosen to match the damage-energy versus depth as closely as possible for the two ions. Correlation of damage measurements and calculations with the observed swelling data are summarized in fig. 6 where step height is plotted versus total damage energy per implanted area. This latter quantity is the product of the calculated damage energy per ion and the fluence [20]. The N and Cr doses corresponding to these damage energies are shown by the scales at the top of fig. 6. Plotting the results in this manner seems to unify the swelling data for Cr and N implantations and suggests a direct relationship between the total damage and the swelling. There are, however, several interesting observations which require further study [20]. For example, the measured depths over which the Sic was turned amorphous were found to be considerably greater for Cr at a given damage energy than that for N even though the damage energy versus depth was carefully matched for the two ions. We believe this results from knock-on C atoms which have greater ranges when initiated by 260 keV Cr than 62 keV N. It should be reiterated,
2.0MEV HE -> AGO, / CR I I I
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*o,
j
c 3500:
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0 .
0 k
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.
Fig. 7. 2.0 MeV He+ scattering from Cr on Al,O, 5 X lOI Ar+/cm2 (filled circles).
heated in vacuum to 400°C (open circles) and ion bombarded
at 4UO’C with
I. CRYSTALLINE
OXIDES
174
B.R. Appleton et al. / Metals in AI,03,
however, that when normalized to the total damage energy the induced swelling for these two ions is the same as fig. 6 shows. The surface mechanical property changes in Sic appear to be due to the amorphization of the surface which makes the surface much more forgiving (increased fracture toughness) than the normal brittle properties of Sic. The mechanical properties are roughly the same for Cr and N implantations. Experiments are in progress to determine the importance, if any, of the chemical nature of the implanted (damaging) species, the importance of implantation temperature and depth, and effects associated with the starting structure of the Sic.
4. Ion beam mixing Although the defects and phase changes induced in insulators by ion bombardment have been studied for some time (see for example, refs. 4-6 and references therein), the materials interactions which can be induced by ion bombardment of this film deposits on insulators have not been investigated extensively to date. These processing techniques (so-called ion beam mixing) have been very successful in metal/semiconductor and metal/metal systems [2,3], and although there are some materials limitations to their use with some insulators, these same techniques should find wide spread utilization in the future. In this section a few specific applications will be mentioned with emphasis on some of the problems that can be encountered. The success encountered in the thermal annealing studies of Cr implanted Al,O, discussed earlier suggested that ion beam mixing of a Cr thin film on Al,O, might be an effective means of introducing Cr into the surface. Ion scattering results are shown in fig. 7 for an Al,O, single crystal with a - 400 A Cr film evaporated on the surface after bombardment with - 5 x 1016 Ar/cm2 at energies of 150 keV in a vacuum of 10-6-10-7 Torr. The sample was heated to 4OOOC during ion bombardment in an attempt to enhance diffusion of the Cr in the presence of the interstitial and vacancy defects created by the Ar damage. Half of the sample was masked from the beam during Ar+ bombardment (open circles) while the other half was not (filled circles). The entire sample was heated at 4OO’C for the duration of the implantation and analysis (-30m). Interpretation of the results from the sample in fig. 7 is somewhat complicated because of a systematic variation in the thickness of the sputter deposited Cr film across the virgin and implanted areas. From measurements before mixing, however, it was determined that the initial film thickness in the mixed area was slightly thinner than the unmixed film in fig. 7. Allowing for
Sic and Si,N,
this, the results in fig. 7 indicate the following. Simple heating of the sample (open circles) caused the formation of a surface oxide as the peak at - 1.05 MeV on top of scattering from Al in the Al,O, shows. Ion bombardment of the heated sample (filled circles) may have produced a small amount of mixing at the Al,O,/Cr interface, but the dominant occurrence appears to be enhanced oxidation of the film. This is shown both by the steps on the Cr portion of the spectrum and the increase in the 0 peak. We have seen an even larger oxidation effect for Al bombarded at 300°C with Bi (4 x lOI cme2, 240 keV) where several thousand angstroms of Al,O, was formed at the surface. The exact origins of these effects are under investigation. These mixed samples showed an increased resistance to removal of surface material by abrasive polishing. Similar effects were observed for MO coated samples that were subjected to ion beam mixing. Such effects are probably due to the mixing at the interface or possibly increased surface hardness. Judging from the thermal annealing results of ion implanted Cr in Al,O,, which showed that significant diffusion occurred only after temperatures of - 16OO”C, it is likely that temperatures considerably greater than 4OO’C during implantation will be required for enhanced transport of Cr. However, prevention of the apparent enhanced oxidation of Cr will probably require mixing in ultrahigh vacuum environments. Ion mixing results for metal films on SIC appear to be much more promising. Initial backscattering results indicate that Ni mixes to a considerable
Fig. 8. Cross section transmission electron micrograph of 1000 A Ni on 1000 A Si,N, on a Si substrate after mixing with 200 keV 1 X 1015 Ar+/cm’.
B.R. Appleton et al. / Metals in AI,O,,
degree but these observations need to be verified by cross-sectional electron microscopy since backscattering alone is not a reliable indicator. Both Zr and Cr were successfully introduced into Al,O, by bombarding - 160 A thick evaporated films on Al,O, with high doses (2 X 1017 cme2) of 1 MeV Fe+ [22]. The mixing mechanism in this case appeared to be recoil implantation. It is also interesting to note that the recoil implanted Zr as well as the Cr were reported as substitutional in Al,O, even in the absence of thermal annealing [22]. This is in marked contrast to the results for implanted Zr which showed no substitutionality even after annealing to 1600°C where most damage to the Al,O, had been removed. Mechanical property measurements also showed increased hardness of most samples in this study [22]. An example of the usefulness of cross-sectional transmission electron microscopy for understanding ion beam mixing can be seen from fig. 8 [26]. Samples consisting of 1000 A of Ni on a 1000 A film of Si3N, on a silicon substrate, were bombarded with 1 X 1015 Ar/cm2 at energies of 200 keV to induce mixing between the Ni and Si,N,. Ion scattering analysbs suggested that interface mixing occurred over a 100 A thick region at the interface but film nonuniformities rendered these conclusions suspect. The TEM analysis in fig. 8 clearly shows the presence of a mixed phase 75 A thick in the interfacial region. In the micrograph, the substrate Si,N, exhibits light contrast and Ni film shows dark contrast. The micro-micro-diffraction analysis indicated the mixed phase to be nickel-rich silicide. In general ion scattering is an excellent indicator of ion mixing, however, cross-section electron microscopy coupled with micro-analysis techinques provides complementary information to achieve a complete understanding of these complicated phenomena.
References
PI Treatise on materials
PI
science and technology, vol. 18 Ion implantation, ed., J.K. Hirvonen (Academic Press, New York, 1980). Laser and electron beam interactions wih solids, eds., B.R. Appleton and G.K. Celler (North-Holland, New York, 1982).
SC and Si,N,
175
(31 Metastable materials formation by ion implantation, eds., ST. Picraux and W.J. Choyke (North-Holland, New York, 1982). [4] P.D. Townsend and S. Valetle, ibid. ref. 1, ch. 11, p. 446. [5] Nucl. Instr. and Meth. 182/183 (1981) 727. (61 R. Kelly, Nucl. Instr. and Meth. 182/183 (1981) 351. [7] R.R. Hart, H.T. Dunlap and O.J. Marsh, in Ion implantation in semiconductors, eds., I. Ruge and J. Graul (Springer, Berlin, 1971) p. 134. [8] G.W. Arnold, G.B. Krefft and C.B. Norris, Appl. Phys. Lett. 25 (1974) 540. [9] G.B. Krefft, W. Bee&old and E.P. EerNisse, IEEE Trans. Nucl. Sci. NS-22 (1975) 2247. [lo] T.F. Luera, J.A. Borders and G.W. Arnold, in Ion implantation in semiconductors, eds., F. Chemov, J.A. Borders and D.K. Brice (Plenum, New York, 1976). [ll] B.D. Evans, H.D. Hendricks, F.D. Bazzarre and J.M. Bunch, ibid., p. 265 [12] H.M. Naguib, J.F. Singleton, W.A. Grant and G. Carter, J. Mater. Sci. 8 (1973) 1633. [13] A.V. Drigo, S. LoRosso, P. Mazzoldi, P.D. Goode and N.E.W. Hartley, Rad. Eff. 33 (1977) 161. [14] A. Camera, A.V. Drigo and P. Mazzoldi, Rad. Eff. 49 (1980) 29. [15] A. Turos, H.J. Mat&e and P. Rabette, Phys. Stat. Sol. A 64 (1981) 565. [16] V.V. Makarov, T. Tuomi, K. Naukkarinen, M. Loumajarvi and M. Riihonen, Appl. Phys. Lett. 35 (1975) 922. [17] C.J. McHargue, M.B. Lewis, B.R. Appleton, H. Naramoto, C.W. White and J.M. Williams, in Int. Conf. on Science of hard materials, Jackson, NY (Plenum, New York; in press). [18] C.J. McHargue, H. Naramoto, B.R. Appleton, C.W. White and J.M. Williams, ref. 3, p. 147. [19] C.J. McHargue and J.M. Williams, ref. 3, p. 303. [20] J.M. Williams, C.J. McHargue and B.R. Appleton, Proc. Conf. on Ion beam modification of materials Grenoble, France (1982); Nucl. Instr. and Meth. 209/210 (1983) 317. (211 H. Naramoto, C.W. White, J.M. Williams, C.J. McHargue, O.W. Holland, M.M. Abraham and B.R. Appleton, J. Appl. Phys. 54 (1983) 683. [22] M.B. Lewis and C.J. McHargue, ref. 3, p. 85. (23) R.C. Bradt, J. Am. Ceram. Sot. 50 (1967) 54. [24] B.B. Ghate, W.C. Smith, C.H. Kim, D.P.H. Hasselman and G.E. Kane, Ceram. Bull. 54 (1975) 210. (251 S.G. Roberts and T.F. Page, in Ion implantation into metals, eds., V. Ashworth, W.A. Grant and R.P.M. Proctor (Pergamon, New York, 1982) (in press). [26] J. Narayan, O.W. Holland, B.R. Appleton, P. Becker and G. Lucovsky, work in progress. [27] B.C. Larson, C.W. White, T.S. Noggle, J.F. Barhorst and D.M. Mills, Appl. Phys. Lett. 42 (1983) 282.
I. CRYSTALLINE
OXIDES
167
Section I. Crystalline oxides ION IMPLANTATION, ION BEAM MIXING, AND ANNEALING Al,Os, SIC AND Si,N, * B.R. APPLETON, C.J. McHARGUE,
STUDIES
OF METALS
IN
H. NARAMOTO + C.W. WHITE, O.W. HOLLAND, ++ G. FARLOW, J. NARAYAN and J.M. WILLIAMS
Solid State Division, Oak Ridge National Laboratoty,
Oak Ridge, TN 37830, USA
Ion scattering/channeling, TEM, EPR and optical microscopy are utilized to determine the structural modifications of ion implanted and annealed Al,O,, SIC, and Si,N,, and to correlate these modifications to surface mechanical property measurements. Ion beam mixing is also studied for inducing increased adherence of metal films on these insulators.
1. Intmluction Ion implantation doping, ion beam mixing, and pulsed-laser annealing are being increasingly utilized to alter the near-surface properties of materials. The attractiveness of these methods arises because they are nonequilibrium processing techniques which can lead to new materials properties [l-5]. Previously these techniques have been used primarily for altering the surface properties of semiconductor and metal or alloy systems [l-3]. Some of the contributing phenomena responsible for the success of these techniques include (in the case of ions) radiation damage, cascade mixing, impurity introduction, radiation enhanced diffusion and compositional and/or structural modifications; and (in the case of pulsed laser annealing) liquid or solid diffusion (mixing), extremely rapid quenching, rapid crystallization, quenched-in phase changes, etc. When one recognizes that insulating materials potentially offer a much broader range of responses to all these phenomena than either metals or semiconducting materials it is surprising that insulators have not been more widely investigated. Insulators can have a wide variety of starting properties. The composition of many insulating materials is complex. Bonding structures can range from ionic to covalent to near-metallic; and the atomic structure can range from crystalline to glassy. Thus, these materials are particularly sensitive to the chemical and electrical
* Research sponsored by the Division of Materials Sciences, U.S. Department of Energy under contract W-7405-eng-26 with Union Carbide Corporation. + JAERI, Tokai, Japan. ++ Metals and Ceramics Division, ORNL. 0168-583X/84/$03.00 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
B.V.
character of impurities or dopants. The alterations possible in insulating materials are correpondingly complex. Ion damage can occur as a result of ionizing as well as displacement collisions; diffusion can have electronic and chemical driving forces; and phase changes can be induced by damage or thermal processing. Previous investigations of insulating and ceramic materials have emphasized the use of ion bombardment or ion implantation doping on the optical properties [4,5], induced phase changes [5,6], electrical and/or chemical properties [5.7], and ion induced disorder [S-15]. Some work has been reported on laser processing of Sic [16]. This paper reports on a continuing study of the surface alterations of insulators and ceramics achievable with ion implantation doping and thermal annealing, and ion-beam mixing with the emphasis on surface mechanical properties [17-221. Related work on pulsed-laser processing of insulators will be presented in a later paper. The intent of this work is to characterize the near surface microstructural alterations induced by ion beam processing, and to correlate these with companion mechanical property measurements. The analysis techniques used include ion scattering, ion channeling, transmission electron microscopy (TEM), electron paramagnetic resonance (EPR) and optical microscopy; the associated mechanical property studies were microhardness and fracture toughness measurements. In section 2 the effects of Cr, Zr, and Ti implantation into Al,O, followed by scheduled thermal annealing are presented. These results demonstrate the utility of correlated structural and mechanical prdperty measurements, and the versatile nature of surface modifications of ceramics. Section 3 presents similar results for Sic. In section 4, results for ion-beam mixing are presented. This paper is more in the nature of an overview and the reader is referred to previous publications for specific experimental details [17-221. I. CRYSTALLINE OXIDES
B.R. Appleton et al. / Metals .in AI,O,,
168
2. Ion implantation and thermal annealing of Al,O,
damage free samples. Half of each sample was retained as virgin reference material and the remaining half was implanted with Cr, Zr, or Ti at room temperature with the ion beam incident at 7’ from the crystal normal at fluences ranging from 10r5-10” cme2 and energies from 100-300 keV. Thermal anneals were performed in
Single crystals of Al,O, of high purity (< 100 ppm total) and low dislocation density (103-lo4 cm-‘) were cut to within f2’ of (0001) and (l?lO), and polished and annealed at 12OO’C in air for 120 h to produce
0 DEPTH (pm) 0.4
AI DEPTH (urn)
0.2
00.4
Cr DEPTH (pm)
0.2
0
0.2
0 DEPTH (pm)
0
0.4
h 1 ’ 1 ’ IV ’ 1 ’ IV1 1 1 1
I-’
-I
II
SIC and Si,N,
Cr DEPTH (pm)
DE:TH (pm)
0.2
I,
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I (VI
0
( 1 ,
103
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102
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I 0.4
I
I 0.6
I
I 0.8
I
I 1.0
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I 1.2
I
I 1.4
I
I 1.6 ENERGY
0.4
0.6
0.0
1.0
1.2
1.4
1.6
(MeV)
52Cr (300 keV, 1X10i7/crn2)
In Q-AI~O~
Fig. 1. Thermal annealing behavior of Cr (300 keV, 1 x 101’/cm2) implanted cr-Al,O,. as-implanted and after various annealing steps to 1500°C.
Ion scattering and channeling results for virgin,
B.R. Appleton et al. / Metals in AI@3, Sic ond Si,N,
either air or N, at temperatures from 800°C to 1600°C for 1 h. 2, I. Structural aiterations and mechanical property measurements oj Cr implanted AI,O,
The history of the ion damage and its recovery, the concentration profiles of the implanted dopant, and the specific lattice locations of defects and impurities for Cr implanted AlsO, were determined from detailed ion scattering/channeling measurements [21]. Representative results are shown in fig. 1 for a high dose implant (1 X lOI Cr/cm2) after a series of anneals in air. Analysis of such ion scattering/channeling data revealed the folowing information: (1) Contrary to previous reports [6], even at these high implantation concentrations, Al,O, is not turned amorphous (fig. la, confirmed by TEM). In Al,O, the damage saturates because of dynamic damage recovery during the room temperature implants. (2) Damage recovering begins selectively in the Al sublattice at a temperature - 8OO*C (figs. la and lb) with no apparent change in the 0 sublattice or Cr distribution, and is complete at temperatures -z 12OO’C. (3) Damage recovery begins in the 0 sublattice at - 1000°C but little change occurs in the Cr substitutionality. (4) Significant incorporation of Cr into substitutional lattice sites begins from 120@-1300°C (fig. le and detailed angular scans).
169
(5) As annealing temperatures are increased to 1500°C, Cr becomes highly (> 95%) substitutional (fig. Id). For lower initial implantation doses more complete damage recovery occurs but dechanneling from Cr-induced lattice strains always causes yields which exceed the virgin-crystal yields. (6) Detailed angular scans 1211show that the Cr goes substitutional in the Al sublattice only. (7) Furthermore, electron paramagnetic resonance measurements [21] show that most if not all the substitutional Cr is in the 3 + valence state. (8) Redistribution of the Cr (in depth) begins to occur for anneals of - 16OO”C,presumably by substitutional diffusion. The corresponding effects on surface mechanical properties caused by these structural and compositional alterations have been reported earlier (17,181, where it was found that Cr impl~tation into Al,O, increased the surface hardness by 30-40% relative to the unimplanted region and increased the fracture toughness by 15%. The effects of thermal annealing in air on the relative hardness of an Also, single crystal implanted with 2 x lOI6 Cr/c& is shown in fig. 2. As annealing is carried to higher temperatures the relative hardness decreases from 1.35 to 1.1 in the range 800-1000°C and then remains constant at 1.1 to temperatures of 1600°C where it decreases to 1.0. When correlated with the observed structural modifications these combined results suggest that annealing in air makes oxygen available for damage recovery through intersitial oxygen
1.6
0 2x 1016 Cr.cme2 A2~10~~
Zr.cms2
o IX 1016 Ti - cms2
I
I
I
I
I
6cm
Boo
loo0
1200
1400
ANNEALING
Fig. 2. Surface microhardness of AI,& annealing temperatures for 1 h in air.
TEMPERATURE
1000
(“C)
implanted with Cr, Zr, and Ti relative to unimplanted Ai&& as a function of thermal
I. CRYSTALLINE OXIDES
170
B.R. Appleton
et
al. / Metals in AI,O,, SIC and Si,N,
diffusion. The displaced Al starts reordering - 8OO’C and is completely reordered - 1200°C. The oxygen sublattice annealing begins at temperatures > 1000°C and the first signs of Cr substitutionality appear. The fraction of implanted Cr acquiring Cr3+ substitutional sites continues to increase, with little change in the initial depth distribution, up to 1500°C. It appears from the hardness measurements that as the Al and 0 damage recovers, Al,O,-Cr,O, solid solutions are formed beginning at 1000°C and persist to - 15OO’C [21]. This conclusion is supported by the fact that the relative hardness is constant at - 1.1 in this temperature range in agreement with measurements on AI,O,-Cr,O, solid solutions grown from the melt [23,24]. Detailed angular scans taken after annealing at an intermediate stage (13OO’C) utilizing ion channeling suggest that interstitial oxygen is trapped by interstitial Cr before complete recovery occurs [21]. This is an important observation since the interstitial oxygen defect in Al,O, has been very difficult to study by other methods because of its high mobility. Above 15OO’C the Cr depth distribution changes, apparently from normal substitutional diffusion into the bulk Al,O,, and the relative hardness drops to 1.0. 2.2. Implantation of Zr in AI,O, and thermal annealing in air The importance of the chemical compatibility of the implanted species on the induced structural allterations and surface mechanical properties can be illustrated by contrasting- the results for Zr implanted in AlzO, with those for Cr [17,18,21]. Analysis for Zr, similar to those already discussed for Cr, provided the following evahrations. 1) The AlaO, single crystals could not be made amorphous by room temperature implantation. 2) Damage recovery began in the Al sublattice at 800°C as with Cr. 3) However, damage recovery did not occur in the 0 sublattice until - 1300°C. 4) Even after annealing to temperatures > 1600°C substantial disorder remained in the Al,O, lattice. 5) Zr showed no substitutional tendencies for annealing to 16OO’C in air. 6) Very little redistribution in depth occurred for annealing temperatures to 16OO’C. A reexamination of fig. 2 shows that these structural differences are mirrored in the relative hardness measurements for Al,O, implanted with 2 X 1016 Zr/cm2 and annealed in air. The Zr implant induces increases in hardness and fracture toughness which remain high even to very high temperatures. This is a significant observation for applications where it would be desirable to use hardened Al,O, as a ceramic bearing or wear material operating at high temperatures. An analysis
consistent with these correlated measurements is that the implanted Zr forms precipitates on annealing, that these precipitates account for the residual damage in the Al,O, and the lack of Zr substitutional&y, and that the hardness is due to precipitation hardening. 2.3. Comparison of thermal annealing in air and H2 of AI,O, implanted with Ti The potential versatility offered by ion-implantation doping of ceramic materials that was discussed in the introduction and illustrated by the results for Cr and Zr implants in Al,O, is further emphasized by recent Ti implantation studies. Previous preliminary investigations showed a complex anisotropic redistribution behavior for Ti implanted in Al,O, and annealed in air [21]. This effect is summarized by the contrasting behaviors of (0001) (fig. 3) and (1210) (fig. 4) Al,O, single crystals implanted with Ti and thermally annealed under identical conditions. Analysis of the ion scattering and channeling results in fig. 3 shows that annealing the (0001) crystal to - 1300°C results in a double-peaked distribution of the Ti with and accumulation at the surface, and the Ti is - 60% substitutional along (0001). Further annealing to 1500°C results in nearly complete damage recovery in the Al,O, lattice, a substantial redistribution of the Ti, but a reduction in the amount of substitutional Ti. These results suggest, and optical microscopy confimed, that needle-like precipitates (probably TiO,) formed at the surface in the range 1300-1500°c. In contrast, results from the (1210) measurements (fig. 4) showed no evidence for redistribution of Ti toward the surface, or for Ti substitutionality after annealing in air to 13OO’C. Annealing to 15OO’C showed increased damage recovery and some diffusion of Ti but no precipitates were observed at the surface. These results suggest that there is a preferential redistribution along the (0001) direction in Al,O,. To investigate the effects associated with the annealing environment, (0001) Al,O, samples were implanted with 4 x 1016 Ti/cm2 (150 keV) and annealed at 13OO’C and 1400°C in H, gas. The ion scattering/channeling results are shown in fig. 5 and should be contrasted with the (0001) spectra of fig. 3 for annealing in air. As with the anneals in air, annealing to 13OO’C in H, caused substantial damage recovery in the Al,O, lattice, a redistribution of the Ti toward the surface, and incorporation of Ti into substitutional lattice sites. There are subtle differences, however, since there is no accumulation of Ti at the surface for the Hz anneals and the fraction of substitutional Ti drops uniformly form 45% at the surface to zero at the depth of the original implant. It is likely that O,, present in the anneals of fig. 3, stabilized the Ti reaching the surface by formation of TiO, and thus led to the observed accumulation.
OEPOTH OEPTH ~I.‘ml Af
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Upon annealing to 1400°C in H, (WO”C less than the comparable spectrum in fig. 3) the Ti diffused to much greater depths, and has developed a profile consistent with normal diffusion. The optical micrastructures of
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the H,-anneakd samples developed with annealing as follows: 1) After 12OPC Eine precipitates formed. 2) After 13OCPCthe fine precipitates grew and new
E
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and annealed in air. I. CRYSTALLINE
OXIDES
B.R. Appleton et al. / Metals in AI,O,, SC and Si,N,
172 0 DEPTH
Al OEPTH
IPI 0.4
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I.0 ENERGY
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Fig. 5. Behavior of (0001) oriented a-Al,O,
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needle-like precipitates surrounded by regions devoid of fine precipitates formed. 3) After 14OO’C both fine and needle-like precipitates disappeared and were replaced with a few precipitates with a different optical contrast. 4) After 1600°C only a few smal precipitates remained. These observations and the spectra of fig. 5 suggest that small precipitates form initally (- 12OO’C) but ripen and give way to larger precipitates at higher anneals (- 1400°C). The fact that these precipitates dissolve easily at higher temperatures, in contrast to those formed in air anneals, is probably due to the lack of stabilizing oxygen. These are, however, speculations, and experiments are presently in progress to identify the phase formation and stability of these systems. Referring again to fig. 2, implantation of (0001) Al,O, to 1 x lOI Ti/cm’ produces an increase in hardness to > 1.25. Annealing in air at temperatures of - 800-lOOO”C, causes a decrease, but then the hardness increases again to - 1.25 at 12OO’C and remains at this level to > 1500°C. The initial increase in hardness may be due to small TiO, precipitates formed and stabilized in the damaged Al,O, lattice. The hardness begins to decrease at the annealing temperatures where the Al (- 800°C) and 0 (- 1000°C) sublattices begin to anneal [21]. At greater temperatures (13OO’C) the Ti moves
and annealed in H, gas.
toward the surface, becomes partially substitutional, and accumulates with the probable formation of TiO, precipitates again and the hardness increases. The higher temperature annealing (- lSOO°C) causes redistribution and decreased substitutionality probably due to the ripening of the TiO, precipitates. Measurements are presently in progress to evaluate the effects of H, annealing on the hardness of Ti implanted Al,O,.
3. Ion implantation doping, damage and thermal annealing studies in SIC Investigations of the changes in surface mechanical properties due to ion implantation doping were extended to Sic because it is a refractory ceramic which figures prominently in a number of energy options of interest to the Department of Energy, and also because, like silicon, it has a covalent bond structure and should have different structural modification properties than say, Al,O, [19,20]. Initial studies showed that implantation of 2 X 1016 Cr/cm’ (280 keV) into SIC single crystals increased the fracture toughness by - 101, a result similar to that observed for Al,O,, but caused a decrease of 25% in surface hardness, which is opposite to the change for Al,O, [19]. Furthermore, scratch wear tests showed a decrease in the ratio of tangential to
B.R. Appleton et al. / Metals in AI,O,, N DOSE (cm-21 10’4 rlTrll[
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TOTAL DAMAGE ENERGY/lMPLANTED
I
I Ill1
402’
1022
AREA kV/cm2)
Fig. 6. Step height increases of Sic implanted with N+ and Cr+ as a function of total damage energy per unit area (see
text).
normal forces (> 25%) for the implanted surface and the scratch was much smoother in appearance. Surface profilometry measurements revealed that ion implantation of Cr caused a noticable step height between the unimplanted and implanted regions, corresponding to a - 20% swelling for implanted single-crystalline Sic samples and - 30% for polycrystalline samples [19]. Changes in hardness similar to those mentioned above have been reported for N+ implantations in Sic but swelling measurements were not made [25]. Detailed investigations were performed to identify
5ooo
8
SC and Si,N,
173
the structural alterations responsible for the swelling and surface mechanical property changes. These studies utilized ion scattering/channeling techniques and transmission electron microscopy and revealed a close correlation between ion induced damage and the large swelling [19,20]. Scheduled implantations of both Cr and Ni were made up to doses of 3 X 10n’ cmm2 and 8 x 1016 cmm2 respectively and the induced damage was monitored at incremental doses in situ by ion channeling. Ion energies of 62 keV N and 260 keV Cr were chosen to match the damage-energy versus depth as closely as possible for the two ions. Correlation of damage measurements and calculations with the observed swelling data are summarized in fig. 6 where step height is plotted versus total damage energy per implanted area. This latter quantity is the product of the calculated damage energy per ion and the fluence [20]. The N and Cr doses corresponding to these damage energies are shown by the scales at the top of fig. 6. Plotting the results in this manner seems to unify the swelling data for Cr and N implantations and suggests a direct relationship between the total damage and the swelling. There are, however, several interesting observations which require further study [20]. For example, the measured depths over which the Sic was turned amorphous were found to be considerably greater for Cr at a given damage energy than that for N even though the damage energy versus depth was carefully matched for the two ions. We believe this results from knock-on C atoms which have greater ranges when initiated by 260 keV Cr than 62 keV N. It should be reiterated,
2.0MEV HE -> AGO, / CR I I I
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*o,
j
c 3500:
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0 .
0 k
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.
Fig. 7. 2.0 MeV He+ scattering from Cr on Al,O, 5 X lOI Ar+/cm2 (filled circles).
heated in vacuum to 400°C (open circles) and ion bombarded
at 4UO’C with
I. CRYSTALLINE
OXIDES
174
B.R. Appleton et al. / Metals in AI,03,
however, that when normalized to the total damage energy the induced swelling for these two ions is the same as fig. 6 shows. The surface mechanical property changes in Sic appear to be due to the amorphization of the surface which makes the surface much more forgiving (increased fracture toughness) than the normal brittle properties of Sic. The mechanical properties are roughly the same for Cr and N implantations. Experiments are in progress to determine the importance, if any, of the chemical nature of the implanted (damaging) species, the importance of implantation temperature and depth, and effects associated with the starting structure of the Sic.
4. Ion beam mixing Although the defects and phase changes induced in insulators by ion bombardment have been studied for some time (see for example, refs. 4-6 and references therein), the materials interactions which can be induced by ion bombardment of this film deposits on insulators have not been investigated extensively to date. These processing techniques (so-called ion beam mixing) have been very successful in metal/semiconductor and metal/metal systems [2,3], and although there are some materials limitations to their use with some insulators, these same techniques should find wide spread utilization in the future. In this section a few specific applications will be mentioned with emphasis on some of the problems that can be encountered. The success encountered in the thermal annealing studies of Cr implanted Al,O, discussed earlier suggested that ion beam mixing of a Cr thin film on Al,O, might be an effective means of introducing Cr into the surface. Ion scattering results are shown in fig. 7 for an Al,O, single crystal with a - 400 A Cr film evaporated on the surface after bombardment with - 5 x 1016 Ar/cm2 at energies of 150 keV in a vacuum of 10-6-10-7 Torr. The sample was heated to 4OOOC during ion bombardment in an attempt to enhance diffusion of the Cr in the presence of the interstitial and vacancy defects created by the Ar damage. Half of the sample was masked from the beam during Ar+ bombardment (open circles) while the other half was not (filled circles). The entire sample was heated at 4OO’C for the duration of the implantation and analysis (-30m). Interpretation of the results from the sample in fig. 7 is somewhat complicated because of a systematic variation in the thickness of the sputter deposited Cr film across the virgin and implanted areas. From measurements before mixing, however, it was determined that the initial film thickness in the mixed area was slightly thinner than the unmixed film in fig. 7. Allowing for
Sic and Si,N,
this, the results in fig. 7 indicate the following. Simple heating of the sample (open circles) caused the formation of a surface oxide as the peak at - 1.05 MeV on top of scattering from Al in the Al,O, shows. Ion bombardment of the heated sample (filled circles) may have produced a small amount of mixing at the Al,O,/Cr interface, but the dominant occurrence appears to be enhanced oxidation of the film. This is shown both by the steps on the Cr portion of the spectrum and the increase in the 0 peak. We have seen an even larger oxidation effect for Al bombarded at 300°C with Bi (4 x lOI cme2, 240 keV) where several thousand angstroms of Al,O, was formed at the surface. The exact origins of these effects are under investigation. These mixed samples showed an increased resistance to removal of surface material by abrasive polishing. Similar effects were observed for MO coated samples that were subjected to ion beam mixing. Such effects are probably due to the mixing at the interface or possibly increased surface hardness. Judging from the thermal annealing results of ion implanted Cr in Al,O,, which showed that significant diffusion occurred only after temperatures of - 16OO”C, it is likely that temperatures considerably greater than 4OO’C during implantation will be required for enhanced transport of Cr. However, prevention of the apparent enhanced oxidation of Cr will probably require mixing in ultrahigh vacuum environments. Ion mixing results for metal films on SIC appear to be much more promising. Initial backscattering results indicate that Ni mixes to a considerable
Fig. 8. Cross section transmission electron micrograph of 1000 A Ni on 1000 A Si,N, on a Si substrate after mixing with 200 keV 1 X 1015 Ar+/cm’.
B.R. Appleton et al. / Metals in AI,O,,
degree but these observations need to be verified by cross-sectional electron microscopy since backscattering alone is not a reliable indicator. Both Zr and Cr were successfully introduced into Al,O, by bombarding - 160 A thick evaporated films on Al,O, with high doses (2 X 1017 cme2) of 1 MeV Fe+ [22]. The mixing mechanism in this case appeared to be recoil implantation. It is also interesting to note that the recoil implanted Zr as well as the Cr were reported as substitutional in Al,O, even in the absence of thermal annealing [22]. This is in marked contrast to the results for implanted Zr which showed no substitutionality even after annealing to 1600°C where most damage to the Al,O, had been removed. Mechanical property measurements also showed increased hardness of most samples in this study [22]. An example of the usefulness of cross-sectional transmission electron microscopy for understanding ion beam mixing can be seen from fig. 8 [26]. Samples consisting of 1000 A of Ni on a 1000 A film of Si3N, on a silicon substrate, were bombarded with 1 X 1015 Ar/cm2 at energies of 200 keV to induce mixing between the Ni and Si,N,. Ion scattering analysbs suggested that interface mixing occurred over a 100 A thick region at the interface but film nonuniformities rendered these conclusions suspect. The TEM analysis in fig. 8 clearly shows the presence of a mixed phase 75 A thick in the interfacial region. In the micrograph, the substrate Si,N, exhibits light contrast and Ni film shows dark contrast. The micro-micro-diffraction analysis indicated the mixed phase to be nickel-rich silicide. In general ion scattering is an excellent indicator of ion mixing, however, cross-section electron microscopy coupled with micro-analysis techinques provides complementary information to achieve a complete understanding of these complicated phenomena.
References
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I. CRYSTALLINE
OXIDES