A comparison between ion implantation into sapphire and polycrystalline alumina

Nuclear Instruments North-Holland

and Methods

in Physics

Research

B59/60

1129

(1991) 1129-1141

Section IX. Insulators, ceramics, and polymers

A comparison between ion implantation and polycrystalline alumina

into sapphire

N. Moncoffre Insritut de Physique NuclPaire de Lyon, IN2P3-CNRS F-69122 Villeurbanne Cedes. France

et UniuersitP Claude Bernard. 43 Bd du II Nouemhre 1918.

The effect of ion implantation and of post-annealing treatments on the surface modifications of sapphire has been widely investigated. In particular, relationships between the microscopic state and the macroscopic properties of implanted sapphire have been established. High-dose ion implantation has been carried out in polycrystalline alumina to a smaller extent. Surface mechanical properties are modified due to surface segregation and precipitation of new phases induced by post-implantation heat treatments. These results are reviewed and a comparison between implantation and annealing effects on single-crystalline and polycrystalline materials is conducted. There have been a lot of reports on surface amorphization of sapphire which depends on the particular species implanted and on the implantation temperature. Observed results for polyctystalline alumina do not differ in this respect for the cases studied. The main differences that will be discussed in this article are attributed to grain boundaries. Charged particle bombardment notably disturbs charge equilibria between grain boundaries and grain bulk. Moreover, it has been shown that during post-implantation heat treatments the orientation relationships of newly created phases are strongly dependent on the initial sapphire orientation. These phase orientations are difficult to put in evidence for polycrystals for which each grain represents a single crystal itself. A comparison of the :nechanical property changes obtained for both types of alumina structures is also discussed.

1. Introduction

The mechanical strength of ceramic materials is. more than for other materials, very sensitive to surface conditions. Most failures under applied loads (fracture) occur from pre-existent surface flows or cracks. Surface treatments like ion implantation and subsequent annealing can noticeably strengthen ceramic properties. Ion implantation whose principle is to introduce a controlled amount of foreign elements in the near-surface region of a material can modify the surface composition and structure (creation of defects, injection of impurities, generation of a compressive stress layer) and consequently its mechanical resistance. without affecting the bulk properties. Up to now a great deal of work has been done on ion implantation of single-crystalline alumina. The earliest studies were concerned with changes in optical properties [l]. Then microstructural changes induced by ion implantation and post annealing treatments have been extensively studied [2-51. Also, their influence on the mechanical property modification has been of large interest in these last years [4.6-81. The point is that only very few investigations have been reported on ion-implanted polycrystalline alumina [9911]. One major reason can be that the primary phenomena resulting from ion beam processes (collision 0168-583X/91/$03.50

?‘ 1991 - Elsevier Science Publishers

cascades) are identical in both, single and polycrystals. Their characterization is, however, easier in single crystals and the understanding of the basic phenomena is much simplified in very pure materials. The aim of this article is to compare the effects of ion implantation into these two types of a-Al,O, and to draw the reasons why they could be different, based on the literature and on the work that has been performed in Lyon. Similarities and differences that are observed in the microstructural state as well as in the mechanical behavior of the two implanted alumina structures will be discussed. The reason why differences between single and polycrystals can be observed is essentially due to the existence of grain boundaries that characterize polycrystals. Impurities can segregate in these grain boundaries and charging effects can occur, related to the different energies of formation of cationic and anionic vacancies. These boundaries also correspond to preferential defect zones. Some consequences can appear during implantation at low fluences (channeling effects, nature of defects depending on orientation, ) and after implantation, because of epitaxial growth and of segregation at grain boundaries during annealing. Differences in charging effects during implantation can also occur due to a possible electric conductivity difference between single and polycrystals.

B.V. (North-Holland)

IX. INSULATORS/CERAMICS/POLYMERS

1130

N. Moncof/re

/ Ion implaniation

We will try to discuss these various points but in the first part of this article. we will summarize some generalities on ion-implanted sapphire.

2. General remarks The ion implantation of many species, either metallic (Cr, Ti. Fe. Ni. Cu, Zr. .. ) or gaseous (N, Ar, ... ), has often been carried out in oriented single-crystalline alumina. In a very synthetic and interesting paper (121. White et al. refer to the works performed in alumina: implanted species, implantation conditions (dose, energy. temperature) and annealing conditions (temperature, environment). In general, ceramics (including alumina) are brittle materials. They do not display noticeable plastic deformation (compared to metals) and have a bad resistance to thermal and mechanical shocks. In spite of their high hardness and their refractory aspect, their fracture toughness is rather unsatisfactory. The purpose of ion implantation is thus: (1) to create new microstructures, (2) to produce radiation damage (defect strengthening), (3) to create new phases using higher implantation doses (second-phase precipitation hardening). The residual stresses, the dislocation microstructure and the amorphization induced by implantation can inhibit the crack propagation responsible for the material brittleness and modify the tribological and mechanical properties of the surface. In most cases the first two points are reached by implanting relatively low doses in order to avoid the amorphization which degrades the mechanical properties. These low-dose implantation effects have been largely studied in single-crystalline alumina. Indeed. ion implantation into alumina induces the creation of numerous defects which are essentially charged point defects, since in an insulator the introduction of a defect or an impurity must provide for electrical neutrality [5]. Impurities cannot be added without creating charge compensating defects. A way to quantify these defects is to estimate the number of displacements per atom (dpa). It depends in particular on the ion nuclear energy loss, on the implanted fluence and on the displacement energy of the matrix. For a given ion dose, the lowest is the substrate temperature during implantation, the highest is the number of remaining defects. For instance. at 77 K, the amorphization of single-crystal alumina is reached at 3-4 dpa, whereas at 300 K a 300 dpa radiation damage is necessary. In the low-temperature range, the dynamic recovery is suppressed and the disorder is rapidly accumulated. By contrast, at room temperature and above. the dynamic annealing prevents the accumulation of defects and higher values of deposited energy are required to reach amorphization.

into sapphire und olumrnn

RANDOM ,_~~--_-~~~~~-_--__~~~~~~ (amorphous)

.;iI y/EL_ CRYSTAL

LATTICE

DAMAGE

(dpa)

Fig. 1. Schematic representation of disorder in the cation lattice as a function of ion dose and temperature (from ref. [g]).

Fig. 1 (from ref. [S]) illustrates this effect. It represents the disorder in the cation lattice as a function of ion dose and temperature. It can be added that the disorder induced by implantation is a function of the single-crystal orientation relative to the implantation direction [S] and this disorder in sapphire has been shown to increase faster for an orientation of the c-axis parallel to that direction. Fig. 2 (from refs. [5.12,13]) represents the damage state (from the xAI value of the alumina sublattice in alumina) induced by ion implantation of different species as a function of the dpa number. When amorphization is reached xA, is equal to 1 whereas in a perfect single crystal it is equal to 0. At 300 K. for a low value of displacements ( < 10 dpa) the damage effect increases linearly with dose. Above 10 dpa, a stationary regime is reached (with xA, - 0.68) corresponding to a dynamical equilibrium between the number of defects created and the damage recovery during implantation. For high doses ( > 100 dpa) the chemical surface composition differs significantly from that of alumina. For instance a 6 x 10” Cr cm-’ implantation at 30 keV in alumina (- 600 dpa) leads to a surface composition where one Cr atom corresponds to one Al,O, pattern

I

.

“’ f%

K

300

Ku

1,

. cu

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AGE3

TI

vMn

Zr

A

Cr

0

Sn

o Fe

.

Nb

*

.

.

0 0

20

40

60

80

100

120

DPA

Fig. 2. Damage state (from the xA, value of the aluminum sublattice in alumina) induced by ion implantation of different species (from refs. [5,12.13]).

N. Moncoffre / Ion implantation into sapphire and alumina

[14]. Amorphization is obtained in this case. Although between 10 and 100 dpa, the stationary regime is almost independent of the ion, there are some exceptions (Zr, Sn) for which amorphization is reached for lower dpa values. At room temperature chemical effects must be added to the amorphization enhancement. At 220 keV Zr and Nb. which have almost the same atomic masses, do not lead to amorphization in the same implantation conditions (identical deposited energies) [15]. Only Zr induces amorphization at 2 x lOI6 ions cm-2 whereas 4 X lOI Nb cm-z does not amorphize sapphire. At 150 keV and 300 K a 5 x lOI Nb cmm2 implantation is found to produce amorphization [16]. This suggests another effect independent of radiation damage. The authors propose a “chemical effect” to explain these results. In the case of Sn [17] and Zr [lo] implantation, a possible role of the 4 + oxidation state on the amorphization enhancement has been proposed. Our purpose when implanting high doses of metallic elements into polycrystalline alumina (around 10” ions cm-‘) followed by different air annealings is to produce new phases, to favour the formation of precipitates. Indeed, as it was for instance suggested by Naramoto et al. [18] who studied Ti and Zr implantations into sapphire. surface hardening can be the result of precipitation hardening induced by thermal annealings. Cochran and coworkers [19] have already implanted Zr at relatively high dose (1.5 X 10” ions cmm2) into single crystals in order to produce insoluble ZrO, precipitates after annealing in an oxygen atmosphere. They intended to improve the surface mechanical properties of the material. They measured hardness and fracture toughness on samples heated at 1100, 1220 and 1440°C and did not observe any difference between implanted and unimplanted surfaces. They suggested that if subsurface precipitates could be grown instead of surface one, a surface transformation toughening should result. In previous papers, we have also characterized polycrystalline alumina implanted with metallic ions (Zr, Fe. Cu) and have investigated the evolution of the implanted surface (microstructure, formed phases) after air annealings [lO,ll]. These results will be reviewed in the second part of this article. The characterization of as-implanted and subsequently annealed samples is carried out using many complementary techniques. In single crystals Rutherford backscattering spectroscopy in channeling position (RBS-C) allows to determine the distribution profile of the implanted element and also the damage distribution and its evolution versus the post thermal treatments. This possibility to detect amorphous regions using RBS is of course not offered when analyzed polycrystals with a classical RBS setup. Transmission electron microscopy (TEM). scanning electron microscopy (SEM),

1131

X-ray photoelectron spectroscopy (XPS), secondary ion mass spectroscopy (SIMS), conversion electron Mossbauer spectroscopy (CEMS) are the techniques frequently used to perform a fine characterization of the implanted layers (determination of the surface topography, new phases. chemical and oxidation states), independently of the crystallographic structure (single or polycrystals).

3. Differences between single crystal and polycrystal 3.1. Before implantation

It is well known that in ceramic materials, the grain boundaries have a large influence on many properties which are mechanical (fracture strength, toughness, plastic deformation and high temperature creep) and also electrical, and processes such as sintering, grain growth and the nucleation of new phases [20,21]. Grain boundaries are characterized by the presence of intrinsic and extrinsic dislocations. faceted configurations and indices. Their properties (energy, structure, chemical reactivity, mechanical, electrical, magnetic, intergranular diffusion, boundary migration) are affected by segregation phenomena. It has been shown [22] that second phase accumulation at grain boundaries is very common. Solutes often segregate and several origins for this segregation are proposed: - The strain energies: many results are in good agreement with the model proposed by McLean [23]. In this model, the segregation energy in Al,O, is proportional to the misfits squared (cz) where

r, being

the radius of the isolated solute ion and r,, the radius of the host site in the lattice. This explains that Ca (r, = 0.9 A) segregates along grain boundaries in sintered alumina whereas Mg (r, = 0.65 A) or Si (r, = 0.41 A) do not segregate. - The electrostatic potential or more generally the electrochemical potential [24]: the ionic nature of most ceramic oxides leads to the formation of an electrostatic potential on grain boundaries which depends on defect structure, impurity concentration and temperature. The grain boundaries can assume an excess charge [23] which can be compensated by a spacecharge cloud of opposite sign at vacancies extending 20 to 100 A from the boundary. This space-charge effect is expected to contribute to a driving force for impurity segregation. Moreover, grain boundaries are known to act as sinks for lattice imperfections (vacancies and interstitials) [20]. Diffusion is often more rapid along these particular imperfections. IX. INSULATORS/CERAMICS/POLYMERS

For most oxides with a close-packed oxygen sublattice, oxygen diffusion at grain boundaries appears to be rapid relative to diffusion within the grains. Oxygen segregation at the grain boundaries strongly depends on temperature and on oxygen pressure [22]. J-7.

During

3.2. I.

imp/uniution

Charging effects

Alpha alumina is one of the most strongly bonded compounds that exists. Energies of formation of point defects are thus large as well as the energy of electronic disorder. The energy of formation of electrons and holes by intrinsic disorder is about 5 eV per electron or hole and the concentration of defects formed by disorder at 1600” C is well below 1 ppm. Since impurity levels of the purest Al,O, available are much better than 1 ppm impurity, it is obvious that the electric properties up to 1600 o C are impurity dominated. With Fe as an impurity. the conductivity is mainly ionic at low oxygen pressure (PO,) [22] and electronic at high P,,.. In insulators under ion bombardment thk types of created defects are numerous and hence the conductivIty modes vary. However. charge effects occur and induce intense electric fields. It is likely that these intense electric fields are responsible, together with compressive stresses. for the defect zone observed beyond the implanted ion distribution [25]. Furthermore it is known [26] that under electron irradiation the charge effects are less on very clean and pure single-crystalline alumina than on poiycrystalline alumina. This result is probably due to a better conductivity in single crystals which limits the electric field value induced by charge bombardment. If one assumes that grain boundaries act

F1g.

SEM micrographs of alumina,

iron implanted

as resistive barriers (depending on the nature of the segregated impurity) each grain receiving positive ions is getting charged and a diffusion induced by the electric field and assisted by the production of defects under ion bombardment could repel ions toward grain boundaries. The difference between single and polycrystals could arise from a partial segregation of the implanted impurity towards grain boundaries during the ion implantation. We have tried to check this hypothesis in the case of iron implantation for which precipitates are clearly revealed along grain boundaries after 800 o C annealings. Fig. 3b shows the iron precipitates formed after annealing in air at 1200°C in polycrystalline alumina. These precipitates are not observed in single-crystalline alumina (fig. 3a). The problem was to determine whether the precipitation occurred during air annealing or segregation had occurred during implantation and was only revealed by these annealings. For that purpose we have performed a SIMS-LAB spectrometer SIMS analysis using equipped with a MIG300 source providing a 0.15 km lateral recolution on an iron implanted polycrystalline alumina sample. We expected to demonstrate a preferential Fe concentration at regular distances corresponding to grain boundaries. The surface atomic “Fe distribution is displayed in fig. 4. It shows that there is almost no preferential iron accumulation. Only a slight track which roughly corresponds to a factor 2 in the difference of iron concentration is observed and can be attributed to grain boundaries. This experiment tends to prove that the iron diffusion that occurs during ion implantation is weak. The diffusion is essentially activated by the air annealings.

and annealed at 1200°C (b) polycrystalline alumina.

for one hour

in cur: (a) vnple

cry

N. Moncoffre / Ion implantation into sapphire and alumina

1133

Fig. 4. Surface atomic distribution of 56Fe in iron-implanted polycrystalline alumina (10” Fe cm-‘. 110 keV).

For Cu and Zr implantations into polycrystals such a preferential precipitate accumulation is not observed after annealing under identical conditions [10,27]. 3.2.2. Orientation effects It is known that the residual disorder depends on the orientation of the ion beam relative to the crystallographic axes and that crystals implanted along different crystallographic directions exhibit different responses to the ion beam [5]. For instance it has been shown that in Cr-implanted sapphire the disorder increased faster for the c-axis orientation compared to the a-axis. In the first case a 2 X 10’s Cr cm- ’ fluence is necessary to create a subsurface amorphous layer at 77 K whereas 2.4 x lOI Cr cm ’ is required in the second case. In polycrystals where all the grain orientations are presented to the beam. a nonuniform effect is observed. 3.3. Post-implantation

annealing effects

3.3.1. Epitaxial regrowth Ohkubo et al. [28] have shown that the migration mechanism of various implanted species (Cr, Mn. Ni, Xe) into sapphire during post-implantation thermal annealing is strongly related to the formation of oriented precipitates. In the case of Mn implantation the epitaxial relation found between surface precipitates and the substrates was (111) MnAl,O,

I] (0001) Al,O,,

(103) Mn,O,

I] (1120) Al,O,.

Polycrystalline or randomly oriented precipitates of Mn,O, were observed on (1702) Al,O,. The surface precipitate microstructure and morphology that develop during post implantation annealings are very sensitive to the surface on which they form. For instance, when titanium is implanted into (0001) oriented cu-Al,O, [18] needle-like precipitates are formed during annealing between 1100 and 1300 o C. These precipitates are shown to be parallel to the (0001) surface. For crystals with (ljl0) surfaces, the precipitate morphology is different. Moreover, Burnett and Page [29] have identified Al,TiO, precipitates on crystals having the (1012) surface. As a consequence, various morphologies can be obtained in single crystals after annealing, by choosing the crystallographic substrate orientation.

3.3.2. Thermodynamical evolution In constrast with the post-implantation thermal treatments performed in single crystals, where the aim is to favour a damage recovery and to follow the implanted impurity evolution versus annealing temperature (chemical evolution as well as occupancy), the thermal treatments in various atmospheres of polycrystals implanted at high doses produce new phases. In our laboratory, we have studied in detail ion implantation into sintered polycrystalline alumina of three metallic species (Fe, Cu, Zr) which are supposed to form oxides of different thermodynamical stability. The study of Fe implantation is particularly interesting since it has been performed in both single and IX. INSULATORS/CERAMICS/POLYMERS

1134

N. Moncofjre Fe DEPTH DISTRIBUTIONS

/

RBS

Fe DEPTH DISTRIBUTIONS

/

RBS

/ Ion implantation info sapphire and alumina

DEPTH

( nm )

DEPTH

( nm )

Fig. 5. Iron concentration profiles in implanted polycrystalline alumina (110 keV. 10” ions cm-*), annealed for one hour in

air at various temperatures.

polycrystals [11,25,30]. The role and the effect of grain boundaries is clearly demonstrated in the iron-implanted polycrystal as it has been already mentioned. In that work several complementary techniques have been used to follow both the distribution profiles of implanted iron and its chemical states as a function of annealing temperature. The implantation conditions were 10” Fe cm-* at 110 keV at room temperature. Fig. 5 displays the distribution profiles deduced from RBS experiments as a function of different annealing temperatures (from 600 to 1600 o C) during one hour in air. It can be noted that above 1000°C the profiles change significantly and iron migrates toward the surface. There is a great interest to compare these spectra to scanning electron micrographs. It allows to associate the formation of surface precipitates along grain boundaries to the iron migration. These precipitates appear slightly at 800 o C, grow up to 1200 o C (fig. 3b) and disappear above 1400 o C. At that latter temperature, only dispersed precipitates remain on the alumina grains. Conversion electron Miissbauer spectroscopy (CEMS). which only probes the implanted zone, was

used to determine the chemical iron environment. The CEMS spectra obtained on polycrystals and single crystals as-implanted and annealed at various temperatures, are displayed in fig. 6. The registered CEMS spectra are the same in both as-implanted matrices. They indicate the presence of a single line characteristic of small a-Fe precipitates. As in the case of Fe implantation at 160 keV and lOi ions/cm* in sapphire. the surface is shown to remain crystalline by TEM investigations. The iron precipitates have an average size of 2.5 nm. Two paramagnetic doublets characteristic of the Fe,, state are ascribed to the FeAl,O, spine1 compound and to a spine1 structure associated with oxygen vacancies since this second doublet is consistent with a higher electronic density around iron atoms. Two other doublets are assigned to the Fe,, state with high spin configurations [30]. This state corresponds to the distribution of iron in a distorded FeO, octahedron and is metastable (it disappears from the earliest stage of annealing around 400 o C). At 700-800 o C. the Fe precipitates have grown and remain superparamagnetic whatever the matrix structure. The FeAl,O, spine1 phase is still observed. However, the Fe,,, oxidation state has different hyperfine parameters in single and polycrystals. In polycrystals the doublet is characteristic of the mixed oxide (Fe,_,Al,),O, and the typical sextet of Fe,O, slightly begins to appear as shown in fig. 6b. In single crystal only one doublet attributed to the Fe(Fe,-_,Al.),O, spine1 is detected. Indeed TEM experiments [25] have shown that after 1OOO“C annealing, c1- and y-Fe,O, precipitates are preferentially located along grain boundaries whereas inside the grains the spine1 and the mixed oxide phases are identified. In sapphire, the spine1 compound and a small amount of hematite are identified. SEM micrographs of fig. 3 clearly reveal this contrast in the precipitate distribution between single and polycrystalline alumina. At 1200 o C Fe,O, precipitates are preferentially identified in the polycrystal whereas in the single crystal only the spine1 phase is observed. At 1400” C in the polycrystal a 30 at.% iron loss is observed on RBS spectra. It could be correlated with the Fe,O, phase vanishing. The process of this vanishing could either have a mechanical origin (as already observed for ZrO, nodules in the case of Zr implantation [lo]) or could be an evaporation. At 1400 o C CEMS analysis suggest that the remaining iron may be incorporated into the alumina lattice in the (Fe,Al,-,),O, solid solution or in the form of small particles of diamond cubic Fe,O, containing aluminum. In addition grazing angle X-ray diffraction (GXRD) experiments detect the Fe,O, structure under its high pressure form and considering the thermodynamic phase diagrams the presence of the spine1 phase is very likely [27]. From the RBS results of

N. Moncoffre

1135

/ Ion rmplantation into sapphtre and alumina

j,

As implanted

\

I_ 3.

3, 1.

a‘1

1 _ 800°C

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Fig. 6. CEMS spectra

-6

of iron-implanted

-3

0

alumina

3

“EL0cIT;8(m;nG.S7; 6 9

-2

(lOI ions cme2, 110 keV) and annealed matrix, (b) polycrystalline matrix.

McHargue et al. [31]. it seems that the iron distribution remains stable for higher annealing temperature (1500° C) in single crystals compared to polycrystals. This difference could be explained by a grain boundary effect which favours the Fe,O, precipitate growth and then the iron loss in this chemical state. Fig. 7 displays four TEM micrographs of polycrystalline and single-crystalline alumina implanted with IO” Zr cmP2 and annealed at 1400” C for one hour. The observed precipitates described as T-ZrO, are homogeneously distributed on both matrix surfaces [lo]. Similar pictures are observed with Cu [27]. These different behaviours between Fe and Cu or Zr can be attributed to the fact that in iron-implanted polycrystals, a noticeable amount of iron is free to move during an-

8

for one hour in air. (a) Single-crystalline

nealings (as Fe 3* in solid solution in corundum or as Fe’) and is trapped at the grain boundaries which constitute barriers. The presence of oxygen there can favour the precipitate growth. By contrast, Zr is embedded in very stable ZrO, precipitates [lo]. Moreover, the fact that ZrO, and Al,O, are two refractory oxides completely insoluble in each other hinders the zirconium migration towards boundaries. Fig. 8 represents TEM micrographs of Cu-implanted polycrystalline alumina. In the implanted region, at the maximum Cu concentration, alumina remains crystalline although its structure is strongly disturbed. However, in localized regions (in front of and behind the maximum). the a-A1,03 crystalline structure is perfectly maintained. In addition, a dislocation zone containing IX. ~NSV~ATORS,‘CERAM~CS/F’OLYMERS

1136

N. Moncoffre / Ion implantation into sapphire and alumina

Fig. 7. SEM micrographs of Zr-implanted alumina (10”

ions cm-“). single crystalline (a. c) or polycrystalline (b. d) and annealed at 1400°C or 1600°C.

essentially linear defects is observed in the back of the implanted zone. GXRD spectra associated with XPS spectra indicate that Cu is only in its metallic state. After annealing at 800°C, CuO precipitates are identified by GXRD and SEM and at 1200” C Cu is no longer detected [27]. Moreover, TEM analysis performed on CL-implanted samples annealed at 1000 o C has allowed to observe kinds of Cu pockets located in triple nodes (fig. 9) [32]. These pockets, which also contain impurity elements like Ca and Y, are partly crystallized. One can assume that they were first amorphous and that the inside migration of Cu made them grow and crystallize. The Cu crystallization in these particular grain boundaries corresponds to a degradation of the alumina surface and probably of its mechanical properties.

3.3.3. Mechanical

properties

3.3.3.1. A comparison between single and po!vcrystals. Two reasons why differences in mechanical properties can be observed between single and polycrystals are essentially due to the presence of grain boundaries which are brittle zones in polycrystals and to the single crystal anisotropy. Indeed the mechanical properties in the a-axis are different from those in the c-axis. For instance it has been shown by O’Hern et al. [33] that the relative hardness increase is slightly higher on c-axis implanted specimens compared to a-axis specimens. Moreover, indentation on the basal plane shows a lower critical load to induce plastic deformation. Hardness. residual compressive stress and fracture toughness have

1137

N. Moncoffre / Ion implantation into sapphire and alumina

0

0.1 pm

a

’

0

I

OJum

b

(3

A’203

IMPLANTED AREA

47nm 85 nh

I ; /

/ DEFECT ZONE

25nm

I I I

20 nm

‘37nm

Fig. 8. Cross section of Cu implanted polycrystalline alumina (110 keV, 10” ions cm ‘). (a) Bright-field image of the implanted area and of the defect zone. (b) Evidence of the sublayers constitutive of the implanted layer (upper and lower parts). (c) Dark-field image of the same region using a diffraction vector of the alumina lattice.

‘3

!i

i

Al203

GRAIN

i

Fig. 9. TEM micrograph of Cu-implanted polycrystalline alumina after a 1000 o C anneal for one hour in air. It displays Cu pockets located in triple nodes.

been measured on a-Al z03 single crystals Cr-implanted with c- and u-axis orientations. These properties were found to increase to a higher extent for the c-axis specimens. The compressive strength of virgin crystalline sapphire was also measured. The bend strength of sapphire was found to be a function of orientation of stress relative to the crystallographic axis of the specimen [34]. Perpendicular to the c-axis, it was found to be 3120 kg/cm2. Parallel to the a-axis, it amounted to 8100 kg/cm?. This is a strong effect of anisotropy on the strength property. It was checked that the average random strength was ~(8100) + f(3120) that is 4750 kg/cm2. The elastic constants of corundum as well were found to be strongly dependent on orientation. However, in sintered alumina, mechanical properties which depend on grain sizes and on the sintering quality have not shown substantial differences with average properties of single-crystalline corundum [34]. Furthermore. it is logical that as the ion implantation dose is increased the crystal anisotropy is decreased. Therefore for low implantation dose mechanical properties remain different depending on the implantation orientation. As an example, it has been shown by Krefft et al. [35] that the absolute stress values display strong crystallographic orientation dependence. For up to 1Ou’ ions cm-’ Ar or H ion implantation, similar stress values are observed along [1120] and [OliO] IX. INSULATORS/CERAMICS/POLYMERS

N. Moncoffre / Ion implantation into sapphwe and alumma

1138

hardness regimes: the first one where hardness increases monotonically with increasing dose; the next one, above a critical dose, where amorphization has started to occur thus softening the surface; the last one where the hardness decreases below the value of the original surface. Many studies [3,4,7] agree that the introduction of defects hinders the dislocation motion and that radiation-induced defects are partly responsible for much of the hardness increase (radiation hardening, part I). The amorphous transformation softens the material surface and as a consequence the hardness decreases (parts II and III). The hardness increase is less for elevated temperature implantation reflecting the lower damage. Very recently the effects of high-dose implantation 17 tons cm -?, 110 keV at room temperature) have (10 been tested (elastic modulus and hardness measurements) on polycrystalline alumina implanted with Cu. Zr or Fe [37]. One annealing temperature was also tested for each ion. The results are displayed in fig. 10. The vertical displacement is continuously measured as a function of the applied load. Thus elastic modulus and relative hardness are directly deduced from the load

into Al,O,, whereas implantation along [OOOl] results in lower stress values. These experiments again show, as previously mentioned, that defect formation and corresponding lattice expansion occur more rapidly along the c-axis. The resulting integrated stress values obtained for polycrystalline alumina lie between those for implantation parallel and perpendicular to the c-axis. This is the result of random orientation of crystallites in the ceramic. For high implantation doses (> 10” ions cmm2) differences between poly- and single crystals are attenuated and progressively disappear since the material tends to become amorphous. 3.3.3.2. Example: hardness and elastic modulus. Using a nanoindentation tester, the elastic modulus and hardness can be measured on very thin depths corresponding to the implanted layers. It has been shown [36] that the elastic modulus in unimplanted single-crystalline alumina was 539 GPa and that this value decreases down to 175 GPa when the surface was made amorphous by stoichiometric implant of Al and 0 at 77 K. Burnett and Page have proposed a surface hardness evolution as a function of dose [3]. They distinguish three different

Zr 0.8.

I 0.4.

I

I

I

ITI I I cu

Fe

f

*I f

1

0.0

J_

4-

0

0

20

40

60

80

4

20

40

60

80

0

20

40

60

80

100

Plastic depth (nm) performed on as-implanted (with Zr. Cu or Fe) and annealed polycrystalline alumina samples. (0) Virgin alumina. (a) as-implanted (10” ions cm _ ‘. 110 keV). (m) Annealed (at 1000°C for Zr. 800 o C for Cu. 1200 o C for Fe). Fig.

10. Relative

hardness

and

elastic

modulus

measurements

N. Moncoffre / Ion implantation into sapphrre und alumina

displacement curves. The loading-unloading cycles are repeated several times for various depths. From these curves we first observe that ion implantation always lowers the hardness value compared to the virgin samples. It corresponds to the amorphization or highly damaged effect previously described for high doses I softened layer). A post-implantation anneal at 800’ C <,n Cu-implanted alumina lowers the hardness even further. For Zr at 1000’ C, the hardness increases very slightly. In the case of Fe, the 1200’ C anneal allows to recover nearly the initial hardness value. However. it can be noticed that the values are very spread out for iron measurements. It could be related to the fact that iron implantation leads to the formation of various phases (spine!, mixed oxide, hematite) which probably display different hardnesses. In contrast Zr and Cu only give rise to one oxide (ZrO, or CuO). The Fe value dispersion is also observed in the relative elastic modulus curves. As is typically the case, the effect of implantation is revealed more markedly at shallow depths. The values approach the unimplanted value with increasing depth due to the increased volume sampled as the indentation proceeds. In single crystals. it has been shown repeatedly that hardness and modulus decrease when the alumina becomes amorphous. In summary, high-dose implantations do not produce improved hardness properties of polycrystalline alumina.

3.3.3.3. Example: t-es&a/ srress. The creation of defects by ion implantation favors a volume increase in the implanted region. By increasing the residual compressive stress, the applied stress necessary to place the stressed surface into tension is increased. The probability of propagating a pre-existing flaw is thus reduced 1121. The residual stresses can be measured using the indentation technique or the cantilever beam method. The extent of lattice damage introduced in sapphire strongly depends on the implantation temperature. Hioki et al. [7] have measured the integrated compressive stress S as a function of the Ni dose implanted into sapphire (fig. lib). They showed that at 100 K. S increases with dose to reach a maximum value for a dose of 1.5 X 10” Ni cm-‘. Above this dose. a stress relief occurs, attributed to the amorphization phenomenon. In contrast at 300 K the S increase is weaker and the broad maximum is reached at 5 X IO” Ni cmm2. From these experiments the authors deduce the volume expansion AV/V for the two implantation temperatures (fig. lla). AV/V is higher at 100 K (higher production of lattice defects) than at 300 K. The occurrence of amorphization involves a large volume expansion. It corresponds to a relief of the compressive stresses accumulated during ion implantation [7].

1139

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3.3.3.4. Example: flexural strength. Hioki et al. have studied the relative flexural strength evolution as a function of dose in single-crystal alumina implanted with various ions (Ar, N. Ni) [6] (fig. 11~). They show that ion implantation can increase the flexural strength IX. INSULATORS/CERAMICS/POLYMERS

1140

N. Moncoffre / Ion inlpluntation into sapphire and ulurninu

up to 60%. All the species studied are effective but the heavier ions are more effective at lower doses. The observed strengthening is related to the radiation damage. For N implantation at high doses, the blister formation induces a relief of the lateral surface compressive stress and reduces the strengthening effect. For Ni implantations at 10” ions cme2. after the specimens were annealed in air for one hour up to 1300” C. the flexural strength improvement was maintained. The strengthening mechanism is attributed in this case to the new compound formation (NiAl,O,) in addition to the radiation damage. These authors have also performed flexural strength measurements on polycrystalline alumina and found lower values compared to single crystals. Eskildsen et al. [9] have measured the relative flexural strength in polycrystalline alumina implanted with Ti, Ar and N. They found for high Ti and N implantation doses, improvements in the strength of 6 and 12% respectively. For Ti. a maximum increase of 14% was obtained at 10” ions cm-.‘. They also observed that the wear rate of Al,O, balls N- and Ti-implanted decreased after a maximum between 5 X 10lh and 10” ions cme2. Hioki et al. [38] suggested that strengthening improvement was due to the surface residual compressive stress arising from the volume expansion.

4. Conclusion In comparison with the improvements that have been observed in metals, ion implantation of ceramics has not led to noticeable improvements up to now. For this reason, many studies have still to be performed. Grain boundary segregation occurring when implanting some ion species like iron, does not seem to produce differences in the mechanical properties of single and polycrystals. Moreover, while the influence of annealing has been largely studied, phenomena that occur during implantation (the role of temperature for instance) and their influence on charge effects remain to be studied.

Acknowledgements I wish to thank sincerely J. Tousset, C. Donnet, G. Marest and C. Esnouf for very helpful and fruitful discussions. I am also very grateful to M.E. O’Hern and L. Romana who performed the nanoindentation tests.

References [l] G.W. Arnold. G.B. Kreft and C.B. Norris. Lett. 25 (1974) 540.

Appl. Phys.

I21 N. 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. [31 P.J. Burnett and T.F. Page. Radiat. Eff. 97 (1986) 123. Relationships in [41 C.J. McHargue, in: Structure-Property Surface Modified Ceramics. eds. C.J. McHargue, R. Kossowsky and W.O. Hofer. NATO AS1 Series (1989) p. 253. [51 C.J. McHargue. P.S. Sklad and C.W. White, Nucl. lnstr. and Meth. B46 (1990) 79. and 0. [61 T. Hioki, A. Itoh. S. Noda, H. Doi, J. Kawamoto Kamigaito. Nucl. Instr. and Meth. B7/8 (1985) 521. S. Noda, H. Doi. J. [71 T. Hioki, A. Itoh. M. Ohkubo, Kawamoto and 0. Kamigaito. J. Mater. Sci. 21 (1986) 1321. Nucl. Instr. and Meth. B19/20 (1987) PI C.J. McHargue. 797. F. Vago. C.A. Strcede and G. Sorensen. [91 S.S. Eskildsen. Proc. 5th Int. Congr. on Tribology, 1989. vol. 5, eds. K. Holmberg and 1. Nieminen. p. 442. J. Tousset and G. 1101 C. Donnet, J. Jaffrezic. N. Moncoffre, Fuchs, Nucl. Instr. and Meth. B46 (1990) 89. [111 C. Donnet, H. Jaffrezic, G. Marest. N. Moncoffre and J. Tousset. Nucl. Instr. and Meth. B50 (1990) 410. P.S. Sklad. L.A. Boatner and WI C.W. White, C.J. McHargue. G.C. Farlow, Mater. SCI. Rep. 4 (l-2) (1989). and B.R. Apu31 G.C. Farlow, C.W. White, C.J. McHargue pleton, Proc. Mater. Res. Sot. Symp. 27 (1984) 395. 1141 P.J. Burnett and T.F. Page, J. Mater. Sci. 19 (1984) 845. G.C. Farlow, C.W. White. J.M. Williams. u51 C.J. McHargue. B.R. Appleton and H. Naramoto. Mater. Sci. Eng. 69 (1985) 123. P. Thevenard. B. Canut, G. Massouras. R. U61 L. Romana. Brenier and M. Brunei, Nucl. Instr. and Meth. B46 (1990) 94. P.S. Sklad, J.C. McCallum. C.W. White. [I71 C.J. McHargue. A. Perez. E. Abonneau and G. Marest. Nucl. Instr. and Meth. B46 (1990) 74. C.J. McHargue, C.W. White. J.M. Wil1181 H. Naramoto, liams, O.W. Holland, M.N. Abraham and B.R. Appleton, Nucl. Instr. and Meth. 209/210 (1983) 1159. S.G. Pope, K-0. Legg and H.F. SolnickiiI91 J.K. Cochran. Legg. Proc. Amer. Sot. Metals 1st Nat. Conf. on the Applications of Ion Plating and Implementation to Materials. Atlanta. GA, USA. 1985. Metals/Materials Technology Series. [201 W.D. Kingery. J. Am. Ceram. Sot. 57 (1974) 1. Pll S. Lartigue and L. Priester, J. Phys. (Paris) C4 (suppl. 46) (1985) 101. WI W.D. Kingery. J. Am. Ceram. Sot. 57 (1974) 74. 1231 W.C. Johnson, Met. Trans. A8 (1977) 1413. v41 S.K. Tiku and F.A. Kroger. J. Am. Ceram. Sot. 63 (1980) 183. ~251 A. Rahioui and C. Esnouf. Proc. Eur. Mater. Res. Sot. Symp.. Strasbourg. 1990, Surf. Coatings Technol.. to be published. Ml C. Jardin, Le Vide. Les Couches Minces (1988) 95. G. Marest. N. Moncoffre. J. Tousset, A. v71 C. Donnet, Rahioui, C. Esnouf and M. Brunel. these Proceedings (7th Int. Conf. on Ion Beam Modification of Materials. Knoxville, TN, USA. 1990) Nucl. Instr. and Meth. B59/60 (1991) 1205.

N. Moncoffre [28] M. Ohkubo,

[29]

[3O] [31]

[32]

/ Ion implantation

T. Hioki and J.I. Kawamoto. J. Appl. Phys. 60 (1986) 1325. P.J. Burnett and T.F. Page, in: Plastic Deformation of Ceramic Materials, eds.’ R.C. Bradt and R.E. Tressler (Plenum, New York. 1984) p. 669. C.J. McHargue, P.S. Sklad. C.W. White, J.C. McCallum, A. Perez and G. Marest. J. Mater. Res.. to be published. C.J. McHargue, G.C. Farlow, P.S. Sklad. C.W. White. A. Perez, N. Kornilios and G. Marest, Nucl. Instr. and Meth. B19/20 (1987) 813. C. Esnouf. unpublished results.

into sapphire and alumina

1141

[33] M.E. O’Hem, C.J. McHargue, C.W. White and G.C. Farlow, Nucl. Instr. and Meth. B46 (1990) 171. [34] E. Ryshkewitch. Oxide Ceramics (Academic Press. New York, 1960). [35] G.B. Krefft and E.P. Eernisse. J. Appl. Phys. 49 (1978) 2725. [36] W.C. Oliver and C.J. McHargue. Thin Solid Films 153 (1987) 185. [37] M.E. O’Hern and L. Romana. unpublished results. [38] T. Hioki, A. Itoh, S. Noda, H. Doi, J. Kawamoto and 0. Kamigaito, J. Mater. Sci. Lett. 3 (1984) 1099.

IX. INSULATORS/CERAMICS/POLYMERS