The structure of ion implanted ceramics

Nuclear Instruments North-Holland

and Methods

THE STRUCTURE C.J. McHARGUE,

in Physics

Research

846 (1990) 79988

OF ION IMPLANTED P.S. SKLAD

CERAMICS

79

*

and C.W. WHITE

OuXRidge National Laboratory, Oak Ridge. Tn USA 3783/-6/1X

The structure of ion implanted ceramics may be crystalline with large concentrations of point defects, point defect clusters. and dislocations, or it may be amorphous. The details of the implanted microstructure depend upon the implantation parameters including ion species. fluence. and substrate temperature. For a given set of implantation parameters. the as-implanted microstructure depends upon the type of chemical bonding present in the ceramic. A second level of structure is the distribution of the implanted ions between substitutional and interstitial lattice sites and among various residual charge states. The amorphous state may contain different short-range order for different implanted ion species. Recent results for these various effects are reviewed.

1. Introduction Although

the

2. Sapphire (AI,O,) process

ently

a nonequilibrium

target

material

does

of ion process,

reflect

implantation the

its chemical

is inher-

response

of

and

structural

the

nature. The interaction of the energetic ion beam with a ceramic differs from that with a metal in several significant ways. Ceramics (insulators) generally are compounds with two or more chemical species that are distributed over at least two sublattices, usually in an ordered manner. Atoms (or ions) of one type that are displaced during elastic collisions are not likely to come to rest at the lattice sites of the other type. The various species have different atomic masses and the energy to displace them from their lattice sites may be very different. These factors influence the manner by which the energy dissipated by elastic collisions is partitioned between the species [1,2] and thus the production of defects. The type of chemical bonding influences the types of defects that exist after the cascade cool-down period. In insulators, the introduction of a defect or an impurity ion must provide for electrical neutrality; consequently. charged point defects or impurity-point defect complexes are often produced. This paper will review the structural features of implanted insulators. In order to contain the material to a manageable size, only two materials will be discussed: sapphire (Al >O,), a ceramic that has largely ionic bonding; and silicon carbide (Sic). with covalent bonding.

* Research sponsored in part by the Division of Materials Sciences, US Department of Energy. under contract DEAC05-840R21400 with the Martin Marietta Energy Systems. Inc. 0168-583X/90/$03.50 (North-Holland)

” Elsevier Science Publishers

B.V

2. I. Structural cormiderattons Sapphire belongs to the space group R3C (rhombohedral). The O’- anions are approximately in hexagonal close-packing and the Al’+ cations occupy 2/3 of the octahedral sites in an ordered arrangement [3]. While it is convenient to describe the structure on the basis of this close-packing, there are significant deviations from the ideal structure, arising from the electrostatic interactions between the ions. In terms of the rhombohedral unit cell, the rhombohedral angle is 53”47’ for perfect packing. whereas, the observed angle for cr-Al,O, is 55 “17’. As a consequence of the deviation from perfect close-packing, the Al ions that would lay on basal planes in the ideal case become noncoplanar in the real structure although the oxide planes remain approximately flat. Each A13+ ion has three O’- ions at 0.1858 nm and three at 0.1971 nm, instead of six at equal distances arranged on the vertices of an octahedron. Although many oxides exhibit deviations from stoichiometry (as much as several percent in some cases), any departure from stoichiometry in Ai,O, is so small that it cannot be detected [4]. Consequently. it is possible to substitute 3 + impurities for Al?‘. but impurities cannot be added without creating charge-compensating defects. The vacant octahedral sites are structural vacancies and are ordered to maintain the electrostatic forces in the crystal. There is ample evidence for charged point defects in sapphire. much of which comes from optical absorption measurements [5]. The presence of the F-centre (an oxygen vacancy containing two electrons) and the F+centre (an oxygen vacancy containing one electron) has been definitively established in both doped crystals and irradiated crystals. 11.CRYSTALLINE

OXIDES

,’ CERAMICS

2.2. The structure of rmptarlted sapphire - cn.vtalline In this section, the structure will be described with respect to stoichiometry, the kind of damage or defects. the amount of disorder, and the disposition of the implanted species. Crystals implanted along different crystallographic directions exhibit different responses to the ion beam. Unless otherwise specified. the discussion will focus upon crystals implanted with c-axis ([OOOl]) approximately parallel to the ion beam. The fluence. implantation energy. and mass of the ion species can be normalized in terms of the damage energy at a particular position, for example, at the peak damage. The values of damage energy reported herein were calculated using E-DEP-1 (Version P5.00) [6]. Rutherford backscattering measurements indicate that the .rtoichiome/n of Al,O, does not change during implantation, i.e.. there is no indication of preferential sputtering [7-lo]. Calculations of the number of displaced ions remaining in the crystal using the values of

nonchanneling yields lead to the conclusion that the numbers of residual displaced aluminum and oxygen ions also occur in the stoichiometric ratio for implantation of Pb [X.9]. Xe and Pt [lo]. and Al [ll]. Thus. regardless of how the defects are produced, those that survive the postcascade cooling period. maintain the stoichiometry of the crystal. The t,‘pe ofdumqe remaining after ion implantation has been studied by optical absorption measurements and TEM observations. Optical absorption measurements confirm the presence of F-centres and F+-centres [111. A TEM micrograph and SAD pattern for AI,O, implanted with 2 x 10lh Cr/cm’ (280 keV) at room temperature are given in fig. 1. The peak deposited damage energy was 1.6 keV/atom. This micrograph contains a high density of “black spots” typical of point defect clusters. The distortions caused by the residual stresses prevented an analysis of the nature of the defect clusters in this system. In view of the results of Pell and

Fig. 1. Back-thinned transmission electron micrograph showing “black-spot” damage characteristic of point defect clusters implanted with 2 X 10” Crjcm’ (280 keV). The insert contains the selected area electron diffraction pattern.

in Al ,O,

PEPlK DAM*GE

ENERGY ,ke”,ofom,

Fig, 2. Disorder in the Al sublattice as determined hackscattering-ion

channeling

deposited

energy

damage

work dislocations [13]. The loops were observed to he on (0001) and {lOiO} habit planes. The Burgers vector analysis gave h = 1/3(Olil) for many of the disiocations. A second set of aligned dislocations were observed with their line vectors nearly perpendicular to [OOOl]. The type of damage introduced by ion implantation is qualitatively similar to that produced by low temperature fast neutron irradiation. However, voids, characteristic of high temperature neutron irradiation [14] have not been observed in ion implanted Al,O,. The mrwunt or degree of disorder introduced by implantation is generally determined from ion-channeling measurements. Since the ratio (x) of the backscattered yield from the crystal with a crystallographic direction aligned with the analyzing ion beam to the yield from a sample oriented “randomly” with respect to the ion beam indicates the disorder in a given sublattice, the value of x is a convenient measure of the residual. The disorder in the Al sublattice is shown as a function of peak damage energy in fig. 2 for implantations of Cr. Fe. Ti. and Nb at room temperature [15]. The disorder increases with damage energy (fluence) to

spectra

for AlzO, ion species.

as a function

implanted

from

of peak

with various

Stathopoulos [12] on electron irradiated sapphire, one possibility is that these are stoichiometric, interstitial dislocation loops. The nlicrostructure of AI,O, implanted to a peak damage energy of 3.1. keV/atom with a simultaneous beam of aluminum and oxygen at room temperature consists of dense arrays of dislocation loops and net-

(4

52Cr (150

keV)

IMPLANTED

(0) Al SUBLATTICE 0

-0.1

1600 ,

1200

-

400

-

I

I

IN a-Al203

I

moo ,

(6) OXYGEN SUBLATTICE I

I

I

I

0.3

0.2

Oi

%(i50 keV) ON a-A1203 AT 77OK. (a) AI SUBLATTICE (b) OXYGEN SUBLATTICE

(b)

AT 779K

I

I

I

I

I

1

___-

‘:

I

Y

I

I

I

1

I

0.1 DEPTH (pm)

I

I

I

1

Qoo

-

800

-

400

-

-

ALIGNED,

---

ALIGNED*

--

<@IO>

----

ALIGNED.

--

RANDOM,

(61

VlRGiN REGDN 0.24 I( lO’%m

ALIGNED, OBO x ~0’%,2 2 4 x !O’6/,m2

IMPLANTED

REGION

rc/’

0 0.10

0

0.10 DEPTH

020

030

(mtcronsi

channeling spectra from the Al and 0 sublattices for AI,O, implanted with Cr (150 keV) to various fluences at 77 K. The scattering yield is plotted as a function of equivalent depth in each sublattice. Implanted with (a) the c-axis and (b) the u-axis approximately parallel to the ion beam.

Fig. 3. Backscattering-ion

II. CRYSTALLINE

OXIDES

,,’ CERAMICS

a value of x - 0.67 at DE = 0.8 keV/atom and remains at this value for damage energy densities as high as 9 keV/atom. Dynamic recovery processes apparently prevent further defect accumulation. Very high fluences (6 x 10” ions/cm’) of 300 keV chromium implanted at room temperature do produce an amorphous structure in crystals having (1072) surfaces normal to the ion beam [16]. These conditions for amorphization correspond to a peak damage energy and a chromium condeposition of = 50 keV/atom centration of 50% (cation). As might be expected, the higher the substrate temperature during implantation, the lower the residual disorder or damage [17]. Implantation of Cr. Ti, or MO at = 650-670 K produces about the same amount of damage in the Al sublattice at the depth of peak disorder but the RBS spectra show evridence for considerable recovery in the immediate surface region. Bull reports similar results for titanium and yttrium implantation into samples with {IOiZ} planes normal to the ion beam [1X]. The amount of residual disorder for a given deposited damage energy is strongly dependent upon the orientation of the ion beam relative to the crystallographic axes [l&19]. Fig. 3 contains the RBS spectra for specimens of sapphire implanted at 77 K with various fluences of chromium with the direction of the beam = 7O from (A) the C-axis and (B) the u-axis. The disorder increased faster for the c-axis orientation and the channeling measurements indicate the presence of a subsurface amorphous layer after a fluence of 2 X 10li Cr/cm’ and which extended to the surface after 3 X 10” Cr/cm’ [19]. The damage accumulation pattern and the fluence necessary for amorphization were quite different for the o-axis orientation. fig. 3b. The material is considered amorphous at a given depth when the backscattered yield for the aligned implanted sample reaches that of the implanted sample oriented “randomly” (i.e., x = 1). The channeling results for the Al sublattice indicate relatively uniform disorder extending from the surface to the mean range but no “random” scattering for fluences as high as 8 X 10” Cr/cm’. The fully amorphous surface layer was present only after a fluence of 2.4 X 10lh Cr/cm’. approximately an order of magnitude greater than for the c-axis crystals. Until recently, most of the information about the frrtul dispositron of the implanted ions was limited to RBS-ion channeling measurements of lattice location. Canera et al. [9] used axial and planar scans to determine the location of lead ions implanted into Al,O,. The lead ions occupied sites along the (0001) rows but displaced 0.05 nm from the oxygen planes. Such precise location of an impurity species is limited to low fluences or to annealed samples in which the disorder is low. An indication of the degree of “substitutionality” in the sublattices may be obtained for a given crystallo-

Table 1 Fraction of Implanted sites [I I]

ions occupying auhstltutlonal

lattice

Implanted lo”:,

loflrc radius

ValerIce

Fraction substitutional “I

Cla c‘u Mn Fe

0.62 0.69 0.x0 0.76 0.64 0.78 0.69 0.62 0.68 0.70 0.6X 0.80 0.41 0.74

3 2 2 2 3 2 3 6 6 5 4 4 4 2

< 0.05 < 0.05 0.45 0.50

Ni Cr Mo W Nh TI Zr Si Zn

0.50 0.55 0.60 0.70 0.70 0.90 -0 -0 -0

,I1 The fraction substitutional is ohtalned from the minimum yields of RBS spectra taken with the c,-axis parallel to the ton beam.

graphic direction by examining the backscattered yield from the impurity in an aligned and a “randomly” oriented crystal. Room temperature implantation of Cr. Ti. Fe, W. Ga. Mn. Ni. Cu and Nb fail to give a simple correlation between the fraction of implanted ions that have substituted for Al (FS) and ionic size. valence, or electronegativity of the implanted species. table 1 [II]. Both Cr and Ga commonly have a valence of 3+ and have ionic radii that are 38% (Cr) and 24%’ (Ga) larger than A13*. yet the channeling results indicate FS values of 0.55 for Cr and 0.04 for Ga. The transition metals Cr. Fe, Ni and Mn exhibit FS values of 0.45 to 0.55; W. Nb, and Ti have values of 0.7 to 0.9: while those for Cu and Ga are less than 0.05. There is no detectable substitution of Zr, Si. or Zn. The Mossbauer effect provides a unique probe for investigating the structural aspects of solids on the atomic scale. Measurements of the hyperfine interactions provide information on symmetry, ordering. and chemical bonding in the immediate environment of the probing nucleus. The structure of the matrix and the resultant charges on iron ions were determined for “Fe implanted into Al?O, at 300 K [20]. The matrix was crystalline after implantation. Four components were identified in the spectra obtained by conversion electron Mbssbauer spectroscopy (CEMS). These components were assigned to a ferric ion (Fe?‘), two forms of ferrous iron (Fe:+ and Fe:,+). and to metallic iron clusters (Fe”). The Fef ’ component has Miissbauer parameters similar to those of the Fee0 bond in wtistite, whereas, Fe:,+ component represents a more ionic state and is

83

Fig. 4. Transmission electron micrograph of as-implanted microstructure in AllO, implanted with 1 x 10” Fe/cm’ at room temperature. showing - 2 nm diameter a-Fe precipitates. Inset is a selected area diffraction pattern for the same region,

with bonding in FeAl,O,. The TEM observations did not detect any second phase at the fluence for which all the iron was distributed between these states,

leading to the conclusion that the regions corresponding to these bonds are small and probably do not represent second-phase precipitates.

consistent

ION FLUENCE

Id 0 I

0

2 /

(X lO’=~ons ~cr81

4 I

6 I

40

20

CONCENTRATION

Fig. 5. Relative

amounts

of iron in various

8 I

states

10 I

1

30

Fe/Al 1%)

plotted as functions of fluence temperature and (b) 77 K.

FLUENCE

(ions/cm2)

for AlzO,

implanted

II. CRYSTALLINE

with “Fe

OXIDES

at: (a) room

,’ CERAMICS

Ferric ion, Fe’+, has parameters that are comparable to iron substitutionally located in the sapphire lattice, i.e.. (Al,_,Fe,),O,. A single line in the spectrum was attributed to small metallic precipitates that behave supermagnetically at room temperature but show magnetic splitting at 4 K. The TEM photographs of Sklad et al. [21] contain images of particles having a size of = 2 nm, fig. 4. Selected area electron diffraction (SAD) patterns suggest that the precipitate is a-Fe, and the isomer shift indicate that the small particles are under very high hydrostatic pressure. The relative amount of Fe in each state is given as a function of fluence in fig. Sa. At the lowest fluence studied, all the iron is distributed between the two 2 + states. A small amount of the 3 + state is present at 2 x lO’(’ Fe/cm2, and its relative amount increases to about 20% at the higher fluences. The Fe” state appears at the higher fluences and constitutes almost 508 of the iron at 10” Fe/cm’. Bourdillon et al. [22] used X-ray absorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) to study the charge states and atomic environments of Ti+ ions implanted into single crystal Al,O, ({ lOi2) orientation). This group concluded that the initial titanium implant exists in both Ti2+ and Ti”’ charge states and is located in a range of sites in the radiation damaged matrix. 2.3. The structure of implunted sapphire - amorphous From the earliest studies of ion implantation into Al 20, there have been conflicting reports of amorphization. An understanding of the phenomenon is beginning

to emerge from current studies of metal ion implantation to various fluences and at various substrate temperatures, and with the use of several techniques to characterize the resulting structures. It now appears that displacements alone will produce the amorphous state if the temperature is low enough to suppress dynamic recovery and that “chemical effects” are important for many implanted species. Implantation at low substrate temperatures ( < 100 K) will produce the amorphous state by damage UCcumulation alone at relatively low values of deposited damage energy [19]. Fig. 2 shows that the random value of backscattered yield for the Al sublattice is reached at 0.3 keV/atom for 260 keV Cr implantation at 77 K (c-axis orientation). Transmission electron microscopy and selected area diffraction patterns confirm the amorphous nature of the implanted zone. That the amorphization is due to damage accumulation and not to impurity effects was confirmed by stoichiometric implants of aluminum and oxygen [19]. In many instances, amorphization by room temperature implantation appears to be due to as yet unidentified “chemical effects”. Implanting “‘Zr and “Nb to the same fluence at similar ion beam energies give different amounts of disorder [23]. Since these elements differ by only three atomic mass units, the details of the cascade formation should be similar: hence, the differences in the disorder must originate in some chemical effect during the cool-down period. Fig. 6 shows the RBS spectra for samples implanted with 4 x 10lh Zr/cm* (7.1 keV/atom) and 4 X lO’(’ Nb/cm* (6.9 keV/atom). The zirconium produced a subsurface amorphous layer extending from about 40 to 100 nm from the surface. As the fluence was increased,

I

!

I

-

I

I



ALIQNED.

I

1

T

VIRGIN - - -



ALIGNED,

IMPLANTED ~RANDDM. IMPLANTED

0.5

0.6

0 7

0.8

0.9

1.0

ENERGY

Fig. 6. Backscattering-ion

channeling

spectra

11

1.2

1.3

1.4

1.5

1.6

(MeV)

from AI,O, implanted at 300 K with (a) 4 X 10 ” “‘Zr/cm* ” Nb/cm’ (220 keV).

(175 keV) and (b) 4

x

10”

x5

0.40

o.i5 DISTANCE

0.20

0.25

QiO

0.45

0.20

DISTANCE

tnm)

Fig. 7. Partial radial distribution functions determined from extended energy loss

this layer broadened and reached the surface. The RBS spectra for the Nb-implanted crystal indicate only a damaged surface region. The fraction of Nb ions that reside in substitutional Al sites (as viewed along the c-axis) is 0.7. The spectra for Zr-implanted crystals show no substitutionality at any fluence studied. The existence of a critical concentration (of Zr) for amorphization is indicated by an analytical electron microscopy study [24]. Measurements of the Zr-concentration profile were made in a series of cross-sectioned specimens which had been implanted to give buried amorphous layers of different widths. The amorphous region was centered on the peak of the zirconium depth profile and its boundaries corresponded to a Zr/Al atomic fraction of - 0.065. The position of the boundaries exhibited no correlation with the rate of damage energy deposition, i.e., defect production. There are still differences reported by different laboratories on the amorphization of sapphire at room temperature. A part of these differences may arise from rate effects, either variation in ion beam currents or damage energy deposition rates. Since there are chemical or impurity effects in the amorphization. contamination of the samples in the ion beam chamber may also affect the amorphization. Information on the structure of the amorphous phases has been obtained and is reported in detail in two papers in these proceedings [25.26]. The radial distribution function, which gives the Al-O nearest-neighbor distances, can be obtained from an analysis of the fine structure of the electron energy loss spectra [27]. Fig. 7 shows the radial distribution functions for cu-Al>O,, y-Al,O,, amorphous Al,O, produced by a stoichiometric implantation of aluminum and oxygen at 77 K. and

025

(nm)

fine structure

analysis (EXELFS)

amorphous Al,O, produced by implantation of iron at 77 K. The Al-0 first nearest-neighbor distance in the amorphous stoichiometric-implanted sample is that of y-AlzO,, whereas that in the amorphous Fe-implanted sample is cu-Al?O,. The CEMS results for samples implanted with iron at 77 K are given in fig. 5b and should be compared with results of similar implants at room temperature. fig. 5a. There are two Fe’+, two Fe4+, and one Fe0 components for iron in the amorphous matrix, whereas there were two Fe’+, one Fe’+. and one Fe” components for the crystalline matrix. There are significant differences in the Miissbauer parameters for the two 2 + states in the two matrices. The unusual charge state of Fe4+ is .. consistent with the iron residing in tetragonally distorted octrahedra that contain oxygen vacancies or nonbridging oxygen ions. Room temperature implantation of tin also produces an amorphous layer and it is believed that this is another example of a chemical effect [26]. The CEMS spectra show the presence of Sn(lI) and Sn(lV).

3. Silicon

carbide

3.1. Structural

considerations

Silicon is tetrahedrally coordinated as a result of its covalent bonding, and exists in several polymorphs. The polytypes have similar atomic arrangements in the plane perpendicular to the symmetry axis but differ from each other in the stacking sequence. The nearest-neighbor bonding is tetrahedral but the second-nearest-neighbors determine whether the structure is cubic or hexagonal. II. CRYSTALLINE OXIDES ,’ CERAMICS

More than 45 different polytypes of Sic have been identified from X-ray diffraction studies [28]. The phase of primary interest for structural applications is the a-phase consisting mainly of the 6H polytype, This structure can be described in terms of planar stacking of silicon and carbon layers in which one type of atom has approximately close-packed layers and the other occupies one-half of the tetragonal interstices. This hexagonal arrangement has a six layer repeat along the c-axis; (ABCACB ). An alternative description that emphasizes the tetrahedral bonding consists of “puckered” graphite-like stoichiometric layers with alternating silicon and carbon atoms around the rings. These “puckered” planes are arranged in a stacking sequence of ABCA’B’C’ , where A. B. and C are related to A’. B’. and C’, respectively. by a rotation of 180°C. Information on the intrinsic defects in a-Sic is limited. Profuse stacking faults are generally present in all the polytypes. Patrick and Choyke [29] proposed that the luminescence spectrum obtained after implantation in of 5 x 10’4 He/cm2 (150 keV) could be interpreted terms of a di-vacancy although an impurity-vacancy complex could not be ruled out. Studies on the solubility of impurities in Sic indicate that In. Mg, P, SC, and the lanthanides have solubilities of 1 ppm or less. Solubilities at elevated temperatures in the range of 1 to 5 at.% have been reported for N, Al, B [301. 3.2.Structure of ion lmplunted Sic Silicon carbide is easily amorphized by ion implantation at room temperature [23,31L34]. Compared to Al,O,, random RBS spectra are obtained at relatively low fluences for all ions studies. Fig. 8 shows the results for 62 keV nitrogen and 260 keV chromium, normalized in terms of peak damage energy. Within the accuracy of the measurements, the data for the two species fall on the same curve. The critical damage energy for amorphization is about 0.02 keV/atom. That the random scattering corresponds to an amorphous surface has been confirmed by TEM and SAD. Analysis from ion channeling yield [32] and Raman spectroscopy [35] indicate that there is no decomposition or preferential sputtering. Raman spectroscopy can be used to follow changes in lattice vibrations due to the implantation. Wright et al. [36] obtained evidence for chemical trapping of Ht and Di implanted into Sic at room temperature. The spectra contained evidence for C-H, C-D, and Si-H bonds, indicating trapping by both silicon and carbon. Fig. 9 shows the Raman spectra for a virgin a-Sic crystal and the same area after implantation with 2 x 1Ol5 Cr/cm* (280 keV) at room temperature [35]. Both RBS and TEM confirmed that the sample was

DAMAGE

ENERGY

iheV/alom)

Fig. 8. Disorder in the Si sublattice as determined from backscattering-ion channeling spectra as a function of peak deposited damage energq for Sic implanted with Cr and N.

amorphous. In the unimplanted sample, the strong peaks to the transverse at 784 and 959 cm-’ correspond optical and longitudinal optical vibrations of the Sic lattice. Auxiliary peaks at 768 and 796 cm-’ are characteristic of the 6H Sic polytype. As Sic is disordered by ion implantation, these peaks broaden until they vanish for the amorphous material. All vibrational modes in amorphous material can contribute to the first-order Raman scattering and this gives a spectrum with broad. diffuse peaks. In the case of Sic, the completely amorphous material shows very few spectral features.

200

400

600

800 WAVE

1000

NUMBERS

1200

1400

1600

jcm 11

Fig. 9. Raman spectra from unimplanted a-Sic and implanted with 2 x 10” Cr/cm’ (260 keV) at room temperature.

Petersen et al. report that the MGssbauer spectra for ‘Sic implanted with In, Sn, Sb, and Te (Fluences of 10” to 10’4 lons/cm2) can be decomposed into two groups of lines [37]. One set could be unambiguously associated with Sn atoms on substitutional Si sites. The second component for room temperature implantation of Sn. Sb. and In indicated the presence of several complex defects. This second component in samples implanted at elevated temperatures (550@9OO”C) appeared to be associated with tin on a Si site-carbon vacancy type of complex (Sn,,-V,). Implantation at elevated temperatures is accompanied by dynamic recovery and Sic can be implanted to rather high fluences without amorphization [35]. At 750 O, fluences of 8 X 1016 N/cm2 (62 keV) and 1 X 1016 Cr/cm’ (260 keV) produced disordered by not amorphous surfaces. These implantations correspond to peak damage energies of 1.49 and 0.6 keV/atom, reseveral times the values necessary to spectively, amorphize Sic at room temperature. Approximately 50% of the chromium occupied substitutional Si sites, as viewed along the c-axis.

4.

Summarq

The structure of ceramics after ion implantation depends upon the implantation parameters of ion fluence, ion species, and substrate temperature and direction of the beam relative to crystallographic axes, and the material parameter of chemical bonding. The stoichiometries of Al,O, and Sic are not altered by implantation, and the type of damage in specimens that remain crystalline is similar to that produced by low temperature neutron damage. Covalent-bonded Sic is amorphized at deposited damage energy densities of 0.02 keV/atom at room temperature but remain crystalline to values as high as 1.6 keV/atom for implantation at 1050 K. Sapphire can be amorphized by damage alone at temperatures below effects in addition to 100 K but requires “chemical” damage-produced disorder at higher temperatures. Information on the charge (valence) states of implanted species is beginning to emerge. The amorphous phases contain short-range order that is ion speciesdependent.

References (REI - 5) Nucl. Instr. [II D.M. Parkin, these Proceedings. and Meth. B46 (1990) 26. 121 R. Smith et al., ibid. p. 7. [31 M.L. Kronberg, Acta Met. 5 (1951) 507. 141 C. Monty. in: Defects in Solids, eds. A.V. Chadwick and M. Terenzi (Plenum, New York, 1986) p. 377.

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