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Luminescence properties of Ce-activated YAG optical ceramic scintillator materials
JOURNAL OF
LUMINESCENCE EISEVIER
Journal of
Luminescence 75 (1997) 193-203
Luminescence properties of Ce-activated YAG optical ceramic scintillator materials E. Zycha>b>*,C. Brecher a, A.J. Wojtowicz a,c, H. Lingertat a a Department of Chemistry, Boston University, 590 Commonwealth Avenue, Boston, MA 02215, USA b Faculty of Chemistry, Wroclaw University, Wroclaw, Poland c Institute of Physics, N. Copernicus University, Torun, Poland
Received 17 December 1996; received in revised form 17 February
1997; accepted
10 March
1997
Abstract
Scintillation and luminescence characteristics of a highly dense transparent YAG : Ce-ceramic are reported and compared to those of a single crystal. When excited with gamma-rays both types of materials display the same dominant ~85 ns decay, but the ceramic also shows a new rapid component of ~20 ns that is completely absent in the single crystal. The scintillation output from the ceramic reaches about 50% of that from the single crystal, having been diminished by unusual loss processes caused by the deformed lattice at the grain boundary interfaces. A phenomenological model involving distortion of the band structure is proposed to explain the results of kinetic measurements. Keywords:
Scintillator; Ceramic; YAG : Ce; Energy transfer; Microstructure
1. Introduction For optically isotropic materials, powder processing technology offers an attractive alternative to single crystal growth for the fabrication of large transparent optical components. Capable of producing material with optical quality high enough for both passive (window) [l] and active (laser) [2,3] applications, this technology has obvious potential for scintillation detectors as well. Radiation detection, however, is a far more demanding application, and high transparency and efficient photoluminescence are no guarantors of good scintillation performance. Low-density (low stopping power) excludes YAG : Ce from most practical applications. However, it is a very good, cheap host for preliminary study and * Corresponding author. Tel.: +I 617 353-2024; e-mail: zych@ bu.edu. 0022-2313/97/$17.00 0 1997 Elsevier PIISOO22-2313(97)00103-8
Science B.V. All rights reserved
to get some understanding of the properties associated with the ceramic medium. Since the scintillation behavior of single crystals has already been thoroughly characterized [4] we have an excellent opportunity to compare the behavior of single crystal and ceramic materials. The experience gained from experiments with YAG: Ce may also be very useful while investigating other materials (such as the denser lutetium analog) in the future.
2. Materials and experiments Two different transparent ceramic samples of Cedoped YAG were synthesized, at a molar concentration of 0.5% with respect to yttrium. These specimens were prepared by hot-pressing of an appropriate yttrium aluminum garnet powder at 1700°C for 3 h (sample #l ) or 10 h (sample #2), under a pressure of
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340atm. After sintering, both samples were heated in air at 145O’C for 4 h. To use as comparison, two single crystal specimens, grown by the Czochralski method, were also prepared (#4 and #5), at the same 0.5 mol% Ce concentration as their ceramic counterparts. In order to minimize any possible chemical discrepancies, the garnet powder for ceramic sample #l was prepared by directly grinding in a mortar a piece of previously grown single crystal, from the same boule from which the Ce-doped single crystal specimens were cut. For ceramic sample #2, however, the YAG: Ce powder was prepared from scratch, by firing an appropriate mixture of YzOs, Al203 and Ce(N03)j .6H20 at 1700°C for 4 h in an Ar/O. l%Hz atmosphere on a molybdenum foil. Before heating, the powder mixture was wet-milled for 20 h with zirconia balls in an ethanol slurry. After conversion to the garnet phase, the material was ground in a mortar and the fine powder formed into a pellet and put in a graphite furnace for hotpressing. Boron nitride powder was used to separate the sample from the graphite punches and sleeve. Of the two single-crystal specimens (#4 and #5), the latter was fired in the graphite furnace in conditions similar to those applied for ceramic hot-pressing and then heated in air at 1450°C for 4 h. This procedure subjected single crystal specimen #5 to essentially the same thermal history as that experienced by the ceramic samples. The other crystal was investigated “as grown”. The luminescent emission of the various specimens was excited either optically or by gamma irradiation, and measured under both steady-state and pulsed conditions. For optical excitation, a Xe-lamp was used in the visible and near-UV and a McPherson #632 deuterium lamp in the VUV region, and the spectra were recorded with a McPherson 0.35m monochromator using a single-photon counting technique. Gamma excitation was accomplished with a 13’Cs source and the resultant emission spectra recorded with a simple 0.25 m Jarrell-Ash monochromator. Light output measurements were performed by comparing the photopeak position with that of bismuth germanate (BGO), the commonly used standard. All specimens were cut to similar dimensions (about 1.0 x 0.7 x 0.03 cm) to minimize the possibility of geometric artifacts, and were polished before any spectroscopic measurements
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Fig. 1. Photograph of the three YAG : Ce specimens upon which most of the measurements were performed: upper left, single crystal #5; middle, large-grain ceramic #l; lower right, small-grain ceramic #2.
were made. A photograph is shown in Fig. 1.
of the relevant
specimens
3. Results The highly reducing conditions present in the graphite furnace during the hot pressing caused substantial discoloration of the material, imparting a dark green to deep reddish color. This discoloration, presumably associated with oxygen vacancies, occurred in both ceramic and single crystal specimens, but was much more profound in the former. Moreover, the finer the powder used for hot-pressing (i.e., the higher the surface/volume ratio), the greater the discoloration of the material. In all cases (both ceramic and single crystal), however, firing in air at 1450°C for 4 h was sufficient to restore the greenish yellow color typical of YAG : Ce. The colors of the “as grown” single crystal and the ceramic fired in air were indistinguishable by eye, with virtually identical absorption spectra in the visible. As shown in Fig. 2, ceramic sample #l (prepared from powder made by crushing a single crystal) exhibited a much greater average grain size (about 40 urn) than sample #2 (,-lo pm). This is not surprising in view of the way the starting YAG: Ce powders were prepared in the respective cases. In the case of sample #I, the particles in the starting powder were much larger from the very beginning,
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foreign phase at the grain boundaries, as well as some residual micropores, both of which are readily observable in high magnification micrographs (see Fig. 2). These defects are far more common in the specimen with the smaller grains (occupying an estimated 0.5% of the sample volume), demonstrating that our processing technique is still far from optimal. The presence of residual scattering centers in optical ceramics is a rather common problem. Their quantity can be markedly reduced, however, through the use of appropriate sintering aids added to the material during the processing. This, combined with careful optimization of the processing technique for each specific material usually makes it possible to achieve a high degree of optical quality [3]. Nevertheless, at this stage of our investigation we deliberately avoided such sintering aids, so as not to introduce any foreign ingredients which might conceivably interfere with the energy transfer processes, diminishing the scintillation performance of the ceramic and distorting comparison with the single crystal. 3.1. Scintillution spectra und kinetics
Fig. 2. Micrographs of YAG: Ce ceramic specimens #l (top) and #2 (bottom). Note the smaller average grain size in the latter, and the apparent presence of voids and foreign phase at the grain boundaries.
necessarily resulting in a larger grain size in the final product. What was surprising, however, was that sintering to full density could be accomplished in only 3 h for specimen #l, as against the 10 h needed for #2, whose starting powder was much finer. This observation is important in terms of fabrication, suggesting that the hotpressing technique may not require very fine (submicron) starting materials for full densification. If confirmed this would provide significant advantages in both time and energy over conventional processes. Although both ceramic specimens were close to theoretical density, the optical quality of #l was clearly superior. The transmission loss of this specimen due to scattering was only about 10 cm-’ as against 40 cm-’ for #2. The relatively high value for the latter is probably associated with the presence of some unidentified
Gamma-excited luminescence from YAG : Ce, for both ceramic specimens (#l and #2) and single crystal (#4), are presented in Fig. 3. The main features of the spectra are the well-known Ce emissions located in the 480-640nm spectral range and peaking at about 525 nm. These emissions, when normalized, have virtually identical shapes in crystal and ceramic, although the peak positions for the latter are shifted slightly but consistently towards the red. This shift appears to scale with grain size, reaching some 6-7 nm in the case of the smaller-grain sample #2. A secondary feature, associated with a host lattice defect [S, 61, is also observed, at around 3 10 nm; this has some bearing on the luminescence kinetics, as will be seen. The spectroscopic behavior of the thermally treated single crystal specimen (#5) was essentially identical to that of the as-grown sample (#4). The spectra of the optically excited emission were indistinguishable from those measured under y-excitation. As seen in Fig. 3, however, the actual emission intensities differ quite substantially. The most intense Ce-emission is observed from the single crystal specimen, being about twice as great as that from the ceramic sample # 1 (having the larger grains) and at least
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-
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(a)
300
400
Wavelength
500
600
(nm)
Fig. 3. Luminescence spectra of YAG:Ce under y-excitation: (a) single crystal; (b) large-grain ceramic, #l; (c) small-grain ceramic, #2. When normalized to the same peak intensity, the latter two spectra show a small but systematic red shift.
ten times greater than that from ceramic #2. Since all the samples were alike in shape and size, giving nearly identical volumes exposed to gamma excitation, and other conditions of measurements were the same, the differences in the Ce-emission intensities from these three samples provide a reasonable estimate of the differences in the scintillation efficiency in these materials. This pattern is confirmed by precise light output measurements against a BGO standard. For sample #l , with the larger grains, the light yield reaches about 90% of that of BGO, while for sample #2 it does not exceed 20%. The light output from the single crystal is substantially greater, with the “as-grown” specimen (#4) showing an output of about 205% of BGO, somewhat better than the 185% value yielded by the specimen (#5) that had been post-heated. These values show that the post-growth heating reduced the scintillation efficiency, probably by creating some defects that interfere with the scintillation process. Nevertheless, it is this specimen to which the efficiency of the ceramics should be compared, because of the similar thermal treatment. This in turn indicates that the best ceramic specimen has a light output about half that of the single crystal. Considering how far from optimum our ceramics actually were, this performance is promising indeed.
0
200
400
600
600
Time (ns) Fig. 4. Decay traces of y-excited scintillation from YAG: Ce: (a) single crystal; (b) large-grain ceramic, #l; (c) small-grain ceramic, #2. Note the anomalous fast component which is present only in the ceramic.
Seeking some insight into the reasons for the reduced scintillation efficiency in our ceramic specimens, we measured the decay kinetics of the y-excited emission. The experimentally measured decay traces (typical examples of which are illustrated in Fig. 4) were subjected to a least-squares analytical fit as a sum of not more than three exponential terms, whose resultant parameters are listed in Table 1. This procedure produced three characteristic times: a fast component in the vicinity of 20 ns; an intermediate component in the 80 ns range; and a slow component of about 400500 ns. The relative proportions of these components differed markedly as a function of the specimen microstructure. In all cases the predominant term was that of the intermediate time, and the single crystal traces never showed any evidence of the fast component (i.e., less than 1% of the total). The other two characteristic times, however, appear to apply (within a rather narrow range of variation) to all the specimens investigated, single crystal and ceramic alike. Our values are also in reasonable accord with those reported in the literature [7]. The most unusual aspect of the kinetic measurements is the fast component, which appears only in the ceramic, and only under y-excitation. Its presence was discovered only recently [8], and the magnitude of its contribution shows a strong systematic
E. Zych et al. 1 Journal of Luminescence Table 1 Parameters Excitation
460 340 280 230
nm nm nm nm
180nm
derived from analytical
fit to decay traces as sum of exponential
Single Crystal (avg.)
Ceramic
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terms Ceramic
(#l)
(#2)
Decay time
Magnitude
Decay time
Magnitude
Decay time
Magnitude
(ns)
(%)
(ns)
(%)
(ns)
(%)
67 68 66 73
100 100 100 100
65 67 67 73
100 100 100 100
62 60 60 62 300 63 300 16 76 470
100 100 100 93(>98)
85 317
96(<99) 4(>1)
Gamma 88 405
71(92) 290)
76 300 21 81 440
97(>99) 3(<1) 5(22) 65 (72) 30(6)
“Magnitude” gives the relative amount of light contained in the component having the corresponding decay time-integrated and instantaneous initial (parenthesized) values, Decay time precision is on the order of 10%
7(<2) 96(>99) 4(<1) 8(37) 60(58) 32(5) time, in terms of both
connection with the microstructure of the specimen. Totally absent from the single crystal, this fast component is clearly evident in the large-grained ceramic (avg. grain size =40 urn), comprising over 20% of the initial intensity. In the smaller-grained (~10 pm) ceramic, where the grain boundary region occupies a larger fraction of the volume, the relative contribution of this fast component increases markedly, almost doubling in magnitude. 3.2. Optical excitation spectra and kinetics To explore further the excitation processes of the Ce-ions in these highly dense polycrystalline materials, we recorded the optical excitation spectra of the 525 nm Ce luminescence. As is readily seen in Fig. 5, these spectra show five prominent features, each of which (not surprisingly) corresponds to a similar feature in the absorption spectrum. Exciting at the peak wavelengths of these features, we also recorded decay traces of the emission, whose results are summarized in Table 1. Identifying each feature by the wavelength of its peak, let us examine their significance in turn: 3.2.1. 460 nm This is the most intense of the excitation peaks, at least a factor of three stronger than any of the others, and undoubtedly originates from the lowest of the five 5d components [5]. Its intensity appears to be independent of the microstructure of the specimen,
0 300 400 Wavelength (nm)
500
Fig. 5. Optical excitation spectra of the 525 nm emission of YAG:Ce: (a) single crystal, #4; (b) large-grain ceramic, #I; (c) small-grain ceramic, #2. Note the systematic red shift of the ceramic excitation peaks.
being virtually identical in all Ce-containing single crystal and ceramic specimens. This fact tells us that the quantum efficiency of Ce emission in the ceramics is not substantially different from that in the single crystal, not much less than unity [9]. The spectral structure, however, does vary with grain size, showing a more gradual slope on the long-wavelength side and a shift of the peak towards the red by as much
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as 15 nm. The corresponding absorption peak shifts much less in the ceramic, but displays a more prominent pedestal-like tail which is probably the source of the enhanced excitation on the long-wavelength side. Decay-time measurements show a single-exponential dependence with a characteristic time of about 65 ns. If anything, the observed red shift probably understates the influence of the environment. Even in the small-grained ceramic, the majority of the Ce ions find themselves deep within the individual grains, in an environment largely unperturbed relative to the single crystal. The influence of the interface becomes progressively stronger for ions situated closer to the edge of the grain, but these make up only a small fraction of the total. Thus, the energy shift for the ions closest to the interface may be significantly greater than the observed shift in the peak position, but their contribution is submerged within the totally inhomogeneously broadened band. These energy shifts presumably correspond to those ions responsible for the anomalous fast decay of the y-excited emission already mentioned, which will be discussed in more detail in the next section. 3.2.2. 335 nm The intensity of the 335 nm peak (as of the 460 nm peak, just discussed) was essentially independent of the identity of the specimen, differing by less than 15%. This feature must also originate from the Ce levels, in this case the one immediately above the lowest, and has been so identified in the literature [S]. Its behavior is consistent with that of the 460nm peak, with the same decay behavior and intensities essentially independent of microstructure. These ceramic peaks also show a red shift relative to the monocrystal, of about the same magnitude in energy as the 460 nm peaks. 3.2.3. 280 nm Unlike the previous two, this feature does not originate in the 5d manifold of Ce, but rather in a defect from which the energy is transferred nonradiatively to the Ce-ion [5]. This feature occurs in both single crystal and ceramics, but disappears after heating in a reducing atmosphere. The decay traces from pulsed excitation at this wavelength are indistinguishable from those in the previous two cases, showing that the energy transfer is quite fast. The intensity of
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this excitation shows a mild systematic dependence on microstructure, being somewhat weaker in the smaller-grained ceramic. 3.2.4. 230 nm The 230 nm feature, like those at 460 and 335 nm, appears to originate from direct Ce-ion excitation [5, lo]. Its asymmetry suggests greater complexity than in the previously discussed cases, and Tomiki et al. [l l] claim that it in fact consists of three individual overlapping components, corresponding to the highest three of the five Sd-levels of the Ce3+ ion. Unlike the other two 5d components already discussed, however, the 230nm peak shows a strong systematic dependence on microstructure, being most intense in the single crystal and weakest in the smallgrained ceramic. The difference in behavior from the other components originating from the 5d manifold would have cast serious doubt upon the assignment, had we not also performed absorption measurements on the specimens. These revealed the presence of a broad microstructure-dependent parasitic absorption below 300nm (Fig. 6) underlying the Ce-associated features. By blocking light that would otherwise reach the Ce ion, this absorptive background could readily account for major diversion of excitation energy to nonradiative paths. The origin of this absorption is unclear, but it almost certainly is associated with a lattice defect. It is not present in as-grown single crystal, but can be generated by subjecting the specimen to a sequential reduction/oxidation heat treatment that simulates the processing regimen for the ceramics. The latter had been post-fired in air to bleach out the strong visible absorption introduced by hot-pressing under highly reducing conditions, which would otherwise block the Ce emission. While this did indeed remove the absorption in the visible, it replaced it with one in the short-wavelength UV that had not previously been present. Unlike the previous case, the responsible defect shows no ability to transfer to the Ce ion and, hence, is more likely to introduce a loss mechanism. The fact that the absorption is greatest in the small-grained ceramic suggests that the defect is not randomly distributed throughout the total specimen volume, but rather is associated specifically with material in the immediate vicinity of the grain boundaries.
E. Zych et al. /Journal
0
of’ Luminescence
-.‘....‘.r..‘....‘....‘200
250
300
Wavelength
350 (nm)
Fig. 6. UV absorption of YAG:Ce: (a) large-grain unbleached; (b) same (#I ), bleached; (c) small-grain bleached. Note the progressively stronger underlying absorption in post-heated samples.
400
ceramic #I, ceramic #2. background
3.2.5. 180-185 nm This may well be the most interesting feature of the optical excitation spectra. It appears to be associated not with localized excitation of the Ce activator ions but rather with the direct elevation of an electron from the valence to the conduction band (although not necessarily fully dissociated from its corresponding hole). The absorption associated with this excitation peak is far stronger than that associated with any of the other excitations already discussed, and, unlike the 230 nm feature, is far too intense for the background to distort its behavior as a function of microstructure. Here, we see a rough correspondence between the intensity of the excitation peaks and the scintillation light yield as measured under y-excitation, with a significantly greater discrepancy between small- and large-grained ceramic than between the latter and the single crystal. The absorption, however, is so strong that hardly any excitation can penetrate deeper than a few tenths of a millimeter into the bulk of the specimen. Indeed the apparent excitation peak itself may be only a surfaceassociated artifact, since the true absorption shows no corresponding decrease at shorter wavelengths. Because of its delocalization, the 180 nm excitation represents the closest simulation that we can achieve optically of what takes place under the ionizing effects of y-irradiation. The approximately 7 eV value
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for the band-to-band energy is in excellent agreement with the literature data [l 11. This energy, however, is just barely enough to raise a valence electron to the conduction band, with very little left over to assure liberation from its hole counterpart. Hence, the optical energy absorbed at this wavelength will go mostly to the generation of mobile excitons, in stark contrast to the free electrons and holes that are the preponderant product of y-excitation. Indeed, the high efficiency with which the Ce3+ ion can trap holes effectively suppresses any tendency for the free carriers to reassociate into excitons. Clearly, the dominant transfer mechanism under band-to-band optical irradiation must be quite different from the free-carrier mediated transfer that prevails under y-excitation [ 12, 131, with profound consequences on the kinetics of the process. Thus, it is hardly surprising to find that band-toband optical excitation produces decay patterns that differ markedly from those obtained by any other means. Unlike the traces from other optical excitation. these depart significantly from single-exponential behavior for all specimens, with slow components having characteristic times of at least 300 ns. These long decay-tails arise when some of the excitation energy is diverted into long-lived traps of the host lattice, from which a delayed transfer to the Ce ion then takes place [5, 61. A prime example of this is provided by the long-lived lattice defect emission at ~3 10 nm, whose overlap with the nearby Ce absorption provides an efficient transfer path. As seen in Table I, the slow component is present in both the monocrystal and ceramic, at about the same proportion. Further details of the lattice emissions will be presented in a separate paper [ 141. In contrast, the dominant component of the decay shows a strong dependence on microstructure, with characteristic times ranging from the expected radiative value of 63 ns in the small-grained ceramic to as high as 85 ns in the single crystal. As the excitation process begins to change from a localized optical mechanism to one that is delocalized and carrier-mediated, the pattern of the decay kinetics also begins to resemble more closely what we have already found under y-excitation. Nevertheless, the similarity of the band-to-band traces to those excited by y-irradiation is only skin deep. The latter shows a substantially higher (and somewhat slower) contribution of the long component, and. as mentioned
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earlier, are the only ones to reveal the presence of a yet faster component in the ceramic, superimposed on the normal emission. The reasons for these various differences are among the issues we must address.
4. Discussion The first issue we wish to consider is the shortfall in scintillation light output exhibited by the ceramic specimens. From photopeak measurements, we know that the efficiency of conversion of absorbed gammas into visible photons drops sharply as the grain size gets smaller, to about 50% and lo%, respectively, relative to the YAG monocrystal. The scintillation spectrum tells us that virtually all of the light output comes from the Ce activator, in its characteristic d-f emission between 480 and 640 nm, with no significant contribution from lattice defects or other intrinsic emission. These two points suggest that the high concentration of abnormal lattice in the vicinity of the grain boundaries may introduce loss mechanisms into the ceramic that would not otherwise be present in the single crystal. These losses may assume at least two different forms: a diversion of the excitation energy to nonradiative paths before it reaches the emitting level of the activator ion; or a partial draining of the emitting level of the already excited activator by some sort of nonradiative quenching. Let us consider these in turn. First, we shall examine the energy transfer process. It is now generally accepted that the primary mechanism of energy transfer is the sequential trapping of mobile holes and electrons by Ce-ions [ 12, 131, ultimately leading to the desired cerium emission. However, the carriers can also be trapped by other defects that may be present in the material, which can lead to other luminescence or (more likely) to nonradiative decay of the excitation. While always present in a single crystal, such defects must be far more numerous in a ceramic, which by its very nature is built up of densely packed but separate grains. Since the grain boundaries are necessarily regions where there is a high degree of lattice distortion and an abundance of point defects, it is reasonable to expect a much higher density of sites where the mobile carriers generated by the absorbed gamma (or band-gap energy photons) could be trapped and eventually decay nonradiatively than would be the case in the bulk crystal. This in turn
would decrease the available population of the mobile electrons and holes, which are needed to convey the excitation to the Ce activators [ 131, necessarily reducing the excited Ce population and hence the light output. Such an explanation is simple and straightforward, and would be perfectly acceptable but for one critical complication: the rapid emission component observed under y-excitation alone. This rapid emission component presents the greatest challenge in the interpretation of the results. Since the kinetics of energy transfer provide no means to accelerate the rate of decay of the emitting level, ’ the presence of a rapid component can arise from only two causes: an increase in the oscillator strength of the emitting transition, or the introduction of an alternative (nonradiative) path by which the excited level can relax to ground. But the first should be accompanied by a substantial increase in the intensity of the corresponding resonance absorption, which is not observed and would not in any event represent a loss process. While the second will, of course, introduce a loss, optical measurements offer no evidence of any significant quenching from the emitting level of the Ce activator in any of the ceramic specimens. We would appear to have reached an impasse. Seeking clues as to the source of this fast component, we subjected the ceramic specimens to a variety of additional tests. We made time-resolved measurements of the y-excited spectra, with gated detection centered at various delays after the absorption of the primary gamma, but found no discernible difference in spectral structure. We explored the full range of wavelengths accessible to optical excitation, without finding any indication of the fast component so readily generated by gammas. Nor could we find any evidence of a rise time, one of the tell-tale signs of delayed energy transfer. How can we reconcile the presence in the ceramic of a y-excited component four times faster than in the single crystal, in a proportion having an inverse relation to light yield, with the total absence ’ Note, however, that a retardation of the emission can readily be attributed to kinetic factors, specifically a slow step in the transfer process. Such a retardation is in fact observed in the YAG:Ce single crystal, where the y-excited decay time is at least 30% longer than the radiative value measured optically. Such kinetic considerations can have profound implications far beyond this particular material, and will be discussed in detail in a separate paper.
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of any evidence of a corresponding component under optical excitation at any wavelength? Nevertheless, we do feel that we can offer a viable phenomenological model to explain our observations. At the core of this model is the recognition that the positions of the d-levels within the bandgap can vary widely according to the chemical properties of the host material. Indeed, it is not at all unusual to find some or even all of the d-levels falling within the conduction band itself. When this happens, optical excitation into such levels can cause the electron to leave the activator entirely, generating measurable photoconductivity. Such an autoionization process can compete effectively with the normally rapid nonradiative relaxation to the lowest level of the manifold, diverting energy away from that erstwhile emitting state. Indeed, since the radiative transition is even slower, the lowest 5d state must be substantially below the conduction band for any d-f emission to take place at all. It is the failure to satisfy this condition that is considered to be the reason why no d-f emission is found in various Ce-activated rare earth sesquioxides and oxysulfides [ 151. YAG : Ce, with its exceptionally high cubic and tetragonal distortion components of the crystal field [9] barely satisfies this condition, with its lowest d-level at only about 10 000 cm-’ below the conduction band [ 161. Consequently, not much perturbation of the crystal field at the Ce-site and/or the band structure of the host material itself is needed to upset this balance and kill the luminescence altogether. Both these effects are likely to occur at the grain boundaries of the YAG-ceramic. While the former may or may not alter the energies of the split dlevels (and in fact no evidence for this is found), they should not materially affect the kinetics. The latter, on the other hand, can fundamentally alter both the energy transfer and emission processes. In particular, the discontinuity in lattice periodicity, and the consequent profusion of broken chemical bonds at the interface, should produce in their immediate vicinity additional states, both filled and empty, at normally forbidden energies within the bandgap (see Fig. 7). This will specifically provide an abundance of “receiving states” for free carriers (particularly electrons) at energies not otherwise accessible to them, and could readily submerge within the extended bands the previously unobstructed (and much more localized) states
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g CONDUCTION s BAND
-.
Ce-levels
s VALENCE ==zzz=z s BAND _~..____ ____
A
_.
Fig. 7. Schematic diagram of the effect of the grain boundaries upon the band structure of YAG:Ce ceramic. Note the local decrease in the bandgap, with the conduction band dropping to or below the energy of the Ce emitting level.
for the Ce optical transition, thereby largely quenching the emission. The effect must certainly depend on the distance from the interface, which we are not in a position to quantify. We propose that the fast (~20 ns) component in the y-excited decay of Ce-ions in the YAG ceramic is due to a quenching process by means of autoionization or tunneling of electrons from the excited ions to “receiving” states residing at the grain boundaries of the host. These receiving states can be viewed as continuations of levels within the conduction band characteristic of the crystal lattice, but moved to lower energies by the lattice discontinuities at the interface. Conceptually, these are the counterparts of the surface states of the original crystalline granules, but now folded into the bulk of the ceramic. The rate of the tunneling is enough to produce a rapid, partially quenched but still readily detectable ~20 ns component in the decays of the ceramic excited by gamma irradiation. Indeed, we find that the contribution of the fast component clearly decreases with increasing grain size. As the number of Ce ions situated in this layer and hence subject to the
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quenching process constitute only a small fraction of the total number of Ce ions in the ceramic, their contribution under optical excitation should be very low, and indeed is not detectable in our measurements. How then can such a small number of ions give rise to such an easily detectable fast component under y-excitation? This we attribute to the fact that such excitation in the ceramic generates a very different distribution of excited Ce-ions than does optical excitation, or even y-excitation in the single crystal. This is inherent in the nature of the cascade that generates the carriers that convey excitation to the Ce activators. This process begins with the absorption of a gamma, generating a relatively small number of highly energetic carriers. These, on collision with lattice electrons, give up some of their excess kinetic energy in the creation of yet more free carriers, with the cascade continuing until the remaining excess kinetic energy per carrier is too low to create more. Thus, by the time these carriers reach thermal equilibrium and become available for recombination, momentum conservation has given rise to a veritable wave front of carriers, moving outward and forward a substantial distance from the point of initial gamma absorption. In the single crystal, this distance is essentially irrelevant; the lattice is uniform and ordered, and each activator is essentially identical. Under optical excitation, the distance is zero, and hence also irrelevant. In the ceramic, however, the lattice discontinuities at the grain boundaries create a network of trenches athwart the path of the hot carrier wave front, in which, as we have already seen, the conduction band is lowered (and the valence band raised) relative to the values in the bulk crystal. Thus, if the average distance a hot carrier travels before thermalization is comparable to the grain size, it can fall into the trench and lose enough kinetic energy to be unable to escape. Then, for the remainder of its mobile life, it will be constrained to travel within this narrow trench defined by the grain boundary, available to excite only a Ce ion in the immediate vicinity of the interfacial region. But, as we have also already seen, these are precisely the ions most subject to nonradiative quenching by interaction with the distorted conduction band. In this manner we see how hot carriers generated under y-excitation can become channeled to the interfacial regions between the grains, substantially altering the distribution of the excitation density in the grain itself by increas-
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ing the relative probability that the Ce emitting centers in such regions become excited. This does not happen under optical excitation, which either targets Ceions directly or, under band-to-band excitation, favors an excitonic mechanism of energy transfer which is less prone to the effect we have just described. Such a “focusing” of the gamma-induced excitation into the interfacial regions could well overcome their limited population of emitting centers, too low to be detected optically. And any such “focusing” must also necessarily increase the probability that the mobile carriers themselves will be trapped at non-emitting defect sites (such as, perhaps, those responsible for the VUV background absorption), preventing their energy from reaching the activator ion at all. The decay traces and light yield measurements under y-excitation even provide us with tools to estimate the relative contributions of the two loss mechanisms, carrier attrition and activator quenching. Table 1 gives us the decay times for the anomalous fast (-20 ns) component in the ceramic specimens, as extracted from a three-term exponential fit. If we assume that the normal (unquenched) decay time would have been the same ~58.5ns as is found in the single crystal, we can calculate the effective quantum efficiency for this component as the ratio of the two decay times. Then, since we know the magnitude of the contribution that this component makes to the total light, we can derive the amount of light we would have had if there had been no quenching, and hence the effective quantum efficiency of the total Ce emission in the ceramic, relative to that in the single crystal. Finally, the factor by which this ratio differs from the light yield of the ceramic (also relative to the single crystal) gives us the fraction of carriers that have been lost. These results are summarized in Table 2. While we should not take these values as definitive, it is interesting to note that it is the carrier attrition that is the more serious of the two loss mechanisms, being both larger than the activator quenching and more sensitive to the grain size. At this point we can offer no positive evidence for this model. It is only phenomenological and must await further study, both mathematical and experimental, to develop it into a consistent and mature quantitative theory. Until then it will remain purely conjectural, supported only by the inapplicability of all others that we have been able to conceive. Yet it is
E. Zych et al. /Journal Table 2 Estimated losses in gamma-excited ceramic specimens”
emission
from
qf’ Luminescence
YAG:Ce
Avg. grain si7e
Carrier attrition (%)
Activator quenching
40 pm IOpm
39 84
18 32
7.5 II9971
193-203
203
CA62330-03. We also express our great appreciation to Prof. A. Lempicki of Boston University, for invaluable advice and assistance during the course of the work.
(%)
Reference [I] W.H. Rhodes, J. Am. Ceram. Sot. 64 (1981) 13; G.C. Wei.
“Relative
to the single crystal.
axiomatic that if all other explanations are ruled out by the evidence, then the only one remaining, no matter how speculative, must be accepted [17].
[2] [3] [4]
5. Conclusions [5]
We have shown that transparent ceramic specimens of YAG: Ce can be reproducibly prepared by means of a hot-pressing technique. The scintillation light output from these specimens is significantly lower than in the single crystal, showing an inverse relationship with grain size. The interfacial regions between the grains are the sites of two spatially localized loss processes: an attrition of the free carrier population through nonradiative recombination at defect sites; and a quenching of the emitting level of the Ce ion through its interaction with the deformed band structure at the grain boundaries. Both of these effects appear to be enhanced by an unusual transport process, wherein the free carriers tend to be channeled preferentially into the interfacial region. There is no doubt that the quality of the material can be significantly improved by optimizing the fabrication procedure, but the extent to which the actual scintillator performance can be improved remains to be seen.
[6] [7]
[8]
[9] [IO] [II]
[I21 [13]
[14]
[15]
Acknowledgements The authors gratefUlly acknowledge the support of the U.S. Department of Health and Human Services, Public Health Service (NIH), under Grant 1 R01
[16] [17]
M.R. Pascucci, E.A. Trickett, C. Brecher. W.H. Rhodes, SPIE Proc. 968 (1988) 5; C. Brecher, G.C. Wei, W.H. Rhodes, J. Am. Ceram. Sot. 73 (1990) 1473. C. Greskovich, J.P. Chemoch. J. Appl. Phys. 44 ( 1973 ) 4599; 45 (1974) 4495. A. Ikesue, T. Kinoshita, J. Am. Ceram. Sot. 78 (1995) 1033. see e.g. A. Lempicki, A.J. Wojtowicz. C. Brecher, in: S.R. Rotman (Ed.), Wide-Gap Luminescent Materials: Theory and Applications, Kluwer Academic Publishers. Norwell. MA, 1996. CM. Wong, S.R. Rotman. C. Warde. Appl. Phys. Lett. 44 (1984) 1038. D.J. Robbins, B. Cockayne. B. Lent, C.N. Duckworth, J.L. Glasper. Phys. Rev. B 9 (1979) 1254. M. Moszynski, T. Ludziejewski. D. Wolski, W. Klamra. N.O. Norlin, Nucl. Instr. and Meth. A 345 (1994) 461. E. Zych, C. Brecher, H. Lingertat, Scintillation properties of YAG:Ce optical ceramics, in: ICDIM, Proc. 13th lnt. Conf. on Defects in Insulating Materials, Wake Forest IJniversity. Winston-Salem, NC, July, 1996, in press. D.J. Robbins, J. Electrochem. Sot. 126 (1979) 1550. S.R. Rotman. C. Warde, J. Appl. Phys. 5X (1985) 522. T. Tomiki, H. Akamine, M. Gishiken, Y. Kinjoh, M. Miyazato, T. Miyazato, N. Toyokawa. M. Hiraoka. N. Hirata, Y. Ganaha, P. Futemma. J. Phys. Sot. Japan 60 (1991) 2437. D.J. Robbins, B. Cockayne, J.L. Glasper. B. Lent, J. Electrochem. Sot. 126 (1979) 1213. A.J. Wojtowicz, Scintillation mechanism: The significance of variable valence and electron-lattice coupling in R.E.activated scintillators, in: P. Dorenbos, C.W.E. van Eijk (Eds.). Stint 95, Proc. Int. Conf. on Inorganic Scintillators and their Applications, Delft University Press, 1996. p. 95. E. Zych, C. Brecher, H. Lingertat, Host-associated luminescence from YAG optical ceramics under gamma and optical excitation, to be published. G. Blasse. W. Schipper, J.J. Hamelink, Inorg. Chim. Acta 189 (1991) 77. D.S. Hamilton, S.K. Gayen, G.J. Pogatshnik, R.D. Ghen, W.J. Miniscalco, Phys. Rev. B 39 (1989) 8807. A.C. Doyle. The sign of the four. Lippincott Monthly 45 (1890) 172.
LUMINESCENCE EISEVIER
Journal of
Luminescence 75 (1997) 193-203
Luminescence properties of Ce-activated YAG optical ceramic scintillator materials E. Zycha>b>*,C. Brecher a, A.J. Wojtowicz a,c, H. Lingertat a a Department of Chemistry, Boston University, 590 Commonwealth Avenue, Boston, MA 02215, USA b Faculty of Chemistry, Wroclaw University, Wroclaw, Poland c Institute of Physics, N. Copernicus University, Torun, Poland
Received 17 December 1996; received in revised form 17 February
1997; accepted
10 March
1997
Abstract
Scintillation and luminescence characteristics of a highly dense transparent YAG : Ce-ceramic are reported and compared to those of a single crystal. When excited with gamma-rays both types of materials display the same dominant ~85 ns decay, but the ceramic also shows a new rapid component of ~20 ns that is completely absent in the single crystal. The scintillation output from the ceramic reaches about 50% of that from the single crystal, having been diminished by unusual loss processes caused by the deformed lattice at the grain boundary interfaces. A phenomenological model involving distortion of the band structure is proposed to explain the results of kinetic measurements. Keywords:
Scintillator; Ceramic; YAG : Ce; Energy transfer; Microstructure
1. Introduction For optically isotropic materials, powder processing technology offers an attractive alternative to single crystal growth for the fabrication of large transparent optical components. Capable of producing material with optical quality high enough for both passive (window) [l] and active (laser) [2,3] applications, this technology has obvious potential for scintillation detectors as well. Radiation detection, however, is a far more demanding application, and high transparency and efficient photoluminescence are no guarantors of good scintillation performance. Low-density (low stopping power) excludes YAG : Ce from most practical applications. However, it is a very good, cheap host for preliminary study and * Corresponding author. Tel.: +I 617 353-2024; e-mail: zych@ bu.edu. 0022-2313/97/$17.00 0 1997 Elsevier PIISOO22-2313(97)00103-8
Science B.V. All rights reserved
to get some understanding of the properties associated with the ceramic medium. Since the scintillation behavior of single crystals has already been thoroughly characterized [4] we have an excellent opportunity to compare the behavior of single crystal and ceramic materials. The experience gained from experiments with YAG: Ce may also be very useful while investigating other materials (such as the denser lutetium analog) in the future.
2. Materials and experiments Two different transparent ceramic samples of Cedoped YAG were synthesized, at a molar concentration of 0.5% with respect to yttrium. These specimens were prepared by hot-pressing of an appropriate yttrium aluminum garnet powder at 1700°C for 3 h (sample #l ) or 10 h (sample #2), under a pressure of
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340atm. After sintering, both samples were heated in air at 145O’C for 4 h. To use as comparison, two single crystal specimens, grown by the Czochralski method, were also prepared (#4 and #5), at the same 0.5 mol% Ce concentration as their ceramic counterparts. In order to minimize any possible chemical discrepancies, the garnet powder for ceramic sample #l was prepared by directly grinding in a mortar a piece of previously grown single crystal, from the same boule from which the Ce-doped single crystal specimens were cut. For ceramic sample #2, however, the YAG: Ce powder was prepared from scratch, by firing an appropriate mixture of YzOs, Al203 and Ce(N03)j .6H20 at 1700°C for 4 h in an Ar/O. l%Hz atmosphere on a molybdenum foil. Before heating, the powder mixture was wet-milled for 20 h with zirconia balls in an ethanol slurry. After conversion to the garnet phase, the material was ground in a mortar and the fine powder formed into a pellet and put in a graphite furnace for hotpressing. Boron nitride powder was used to separate the sample from the graphite punches and sleeve. Of the two single-crystal specimens (#4 and #5), the latter was fired in the graphite furnace in conditions similar to those applied for ceramic hot-pressing and then heated in air at 1450°C for 4 h. This procedure subjected single crystal specimen #5 to essentially the same thermal history as that experienced by the ceramic samples. The other crystal was investigated “as grown”. The luminescent emission of the various specimens was excited either optically or by gamma irradiation, and measured under both steady-state and pulsed conditions. For optical excitation, a Xe-lamp was used in the visible and near-UV and a McPherson #632 deuterium lamp in the VUV region, and the spectra were recorded with a McPherson 0.35m monochromator using a single-photon counting technique. Gamma excitation was accomplished with a 13’Cs source and the resultant emission spectra recorded with a simple 0.25 m Jarrell-Ash monochromator. Light output measurements were performed by comparing the photopeak position with that of bismuth germanate (BGO), the commonly used standard. All specimens were cut to similar dimensions (about 1.0 x 0.7 x 0.03 cm) to minimize the possibility of geometric artifacts, and were polished before any spectroscopic measurements
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Fig. 1. Photograph of the three YAG : Ce specimens upon which most of the measurements were performed: upper left, single crystal #5; middle, large-grain ceramic #l; lower right, small-grain ceramic #2.
were made. A photograph is shown in Fig. 1.
of the relevant
specimens
3. Results The highly reducing conditions present in the graphite furnace during the hot pressing caused substantial discoloration of the material, imparting a dark green to deep reddish color. This discoloration, presumably associated with oxygen vacancies, occurred in both ceramic and single crystal specimens, but was much more profound in the former. Moreover, the finer the powder used for hot-pressing (i.e., the higher the surface/volume ratio), the greater the discoloration of the material. In all cases (both ceramic and single crystal), however, firing in air at 1450°C for 4 h was sufficient to restore the greenish yellow color typical of YAG : Ce. The colors of the “as grown” single crystal and the ceramic fired in air were indistinguishable by eye, with virtually identical absorption spectra in the visible. As shown in Fig. 2, ceramic sample #l (prepared from powder made by crushing a single crystal) exhibited a much greater average grain size (about 40 urn) than sample #2 (,-lo pm). This is not surprising in view of the way the starting YAG: Ce powders were prepared in the respective cases. In the case of sample #I, the particles in the starting powder were much larger from the very beginning,
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195
foreign phase at the grain boundaries, as well as some residual micropores, both of which are readily observable in high magnification micrographs (see Fig. 2). These defects are far more common in the specimen with the smaller grains (occupying an estimated 0.5% of the sample volume), demonstrating that our processing technique is still far from optimal. The presence of residual scattering centers in optical ceramics is a rather common problem. Their quantity can be markedly reduced, however, through the use of appropriate sintering aids added to the material during the processing. This, combined with careful optimization of the processing technique for each specific material usually makes it possible to achieve a high degree of optical quality [3]. Nevertheless, at this stage of our investigation we deliberately avoided such sintering aids, so as not to introduce any foreign ingredients which might conceivably interfere with the energy transfer processes, diminishing the scintillation performance of the ceramic and distorting comparison with the single crystal. 3.1. Scintillution spectra und kinetics
Fig. 2. Micrographs of YAG: Ce ceramic specimens #l (top) and #2 (bottom). Note the smaller average grain size in the latter, and the apparent presence of voids and foreign phase at the grain boundaries.
necessarily resulting in a larger grain size in the final product. What was surprising, however, was that sintering to full density could be accomplished in only 3 h for specimen #l, as against the 10 h needed for #2, whose starting powder was much finer. This observation is important in terms of fabrication, suggesting that the hotpressing technique may not require very fine (submicron) starting materials for full densification. If confirmed this would provide significant advantages in both time and energy over conventional processes. Although both ceramic specimens were close to theoretical density, the optical quality of #l was clearly superior. The transmission loss of this specimen due to scattering was only about 10 cm-’ as against 40 cm-’ for #2. The relatively high value for the latter is probably associated with the presence of some unidentified
Gamma-excited luminescence from YAG : Ce, for both ceramic specimens (#l and #2) and single crystal (#4), are presented in Fig. 3. The main features of the spectra are the well-known Ce emissions located in the 480-640nm spectral range and peaking at about 525 nm. These emissions, when normalized, have virtually identical shapes in crystal and ceramic, although the peak positions for the latter are shifted slightly but consistently towards the red. This shift appears to scale with grain size, reaching some 6-7 nm in the case of the smaller-grain sample #2. A secondary feature, associated with a host lattice defect [S, 61, is also observed, at around 3 10 nm; this has some bearing on the luminescence kinetics, as will be seen. The spectroscopic behavior of the thermally treated single crystal specimen (#5) was essentially identical to that of the as-grown sample (#4). The spectra of the optically excited emission were indistinguishable from those measured under y-excitation. As seen in Fig. 3, however, the actual emission intensities differ quite substantially. The most intense Ce-emission is observed from the single crystal specimen, being about twice as great as that from the ceramic sample # 1 (having the larger grains) and at least
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-
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(a)
300
400
Wavelength
500
600
(nm)
Fig. 3. Luminescence spectra of YAG:Ce under y-excitation: (a) single crystal; (b) large-grain ceramic, #l; (c) small-grain ceramic, #2. When normalized to the same peak intensity, the latter two spectra show a small but systematic red shift.
ten times greater than that from ceramic #2. Since all the samples were alike in shape and size, giving nearly identical volumes exposed to gamma excitation, and other conditions of measurements were the same, the differences in the Ce-emission intensities from these three samples provide a reasonable estimate of the differences in the scintillation efficiency in these materials. This pattern is confirmed by precise light output measurements against a BGO standard. For sample #l , with the larger grains, the light yield reaches about 90% of that of BGO, while for sample #2 it does not exceed 20%. The light output from the single crystal is substantially greater, with the “as-grown” specimen (#4) showing an output of about 205% of BGO, somewhat better than the 185% value yielded by the specimen (#5) that had been post-heated. These values show that the post-growth heating reduced the scintillation efficiency, probably by creating some defects that interfere with the scintillation process. Nevertheless, it is this specimen to which the efficiency of the ceramics should be compared, because of the similar thermal treatment. This in turn indicates that the best ceramic specimen has a light output about half that of the single crystal. Considering how far from optimum our ceramics actually were, this performance is promising indeed.
0
200
400
600
600
Time (ns) Fig. 4. Decay traces of y-excited scintillation from YAG: Ce: (a) single crystal; (b) large-grain ceramic, #l; (c) small-grain ceramic, #2. Note the anomalous fast component which is present only in the ceramic.
Seeking some insight into the reasons for the reduced scintillation efficiency in our ceramic specimens, we measured the decay kinetics of the y-excited emission. The experimentally measured decay traces (typical examples of which are illustrated in Fig. 4) were subjected to a least-squares analytical fit as a sum of not more than three exponential terms, whose resultant parameters are listed in Table 1. This procedure produced three characteristic times: a fast component in the vicinity of 20 ns; an intermediate component in the 80 ns range; and a slow component of about 400500 ns. The relative proportions of these components differed markedly as a function of the specimen microstructure. In all cases the predominant term was that of the intermediate time, and the single crystal traces never showed any evidence of the fast component (i.e., less than 1% of the total). The other two characteristic times, however, appear to apply (within a rather narrow range of variation) to all the specimens investigated, single crystal and ceramic alike. Our values are also in reasonable accord with those reported in the literature [7]. The most unusual aspect of the kinetic measurements is the fast component, which appears only in the ceramic, and only under y-excitation. Its presence was discovered only recently [8], and the magnitude of its contribution shows a strong systematic
E. Zych et al. 1 Journal of Luminescence Table 1 Parameters Excitation
460 340 280 230
nm nm nm nm
180nm
derived from analytical
fit to decay traces as sum of exponential
Single Crystal (avg.)
Ceramic
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terms Ceramic
(#l)
(#2)
Decay time
Magnitude
Decay time
Magnitude
Decay time
Magnitude
(ns)
(%)
(ns)
(%)
(ns)
(%)
67 68 66 73
100 100 100 100
65 67 67 73
100 100 100 100
62 60 60 62 300 63 300 16 76 470
100 100 100 93(>98)
85 317
96(<99) 4(>1)
Gamma 88 405
71(92) 290)
76 300 21 81 440
97(>99) 3(<1) 5(22) 65 (72) 30(6)
“Magnitude” gives the relative amount of light contained in the component having the corresponding decay time-integrated and instantaneous initial (parenthesized) values, Decay time precision is on the order of 10%
7(<2) 96(>99) 4(<1) 8(37) 60(58) 32(5) time, in terms of both
connection with the microstructure of the specimen. Totally absent from the single crystal, this fast component is clearly evident in the large-grained ceramic (avg. grain size =40 urn), comprising over 20% of the initial intensity. In the smaller-grained (~10 pm) ceramic, where the grain boundary region occupies a larger fraction of the volume, the relative contribution of this fast component increases markedly, almost doubling in magnitude. 3.2. Optical excitation spectra and kinetics To explore further the excitation processes of the Ce-ions in these highly dense polycrystalline materials, we recorded the optical excitation spectra of the 525 nm Ce luminescence. As is readily seen in Fig. 5, these spectra show five prominent features, each of which (not surprisingly) corresponds to a similar feature in the absorption spectrum. Exciting at the peak wavelengths of these features, we also recorded decay traces of the emission, whose results are summarized in Table 1. Identifying each feature by the wavelength of its peak, let us examine their significance in turn: 3.2.1. 460 nm This is the most intense of the excitation peaks, at least a factor of three stronger than any of the others, and undoubtedly originates from the lowest of the five 5d components [5]. Its intensity appears to be independent of the microstructure of the specimen,
0 300 400 Wavelength (nm)
500
Fig. 5. Optical excitation spectra of the 525 nm emission of YAG:Ce: (a) single crystal, #4; (b) large-grain ceramic, #I; (c) small-grain ceramic, #2. Note the systematic red shift of the ceramic excitation peaks.
being virtually identical in all Ce-containing single crystal and ceramic specimens. This fact tells us that the quantum efficiency of Ce emission in the ceramics is not substantially different from that in the single crystal, not much less than unity [9]. The spectral structure, however, does vary with grain size, showing a more gradual slope on the long-wavelength side and a shift of the peak towards the red by as much
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as 15 nm. The corresponding absorption peak shifts much less in the ceramic, but displays a more prominent pedestal-like tail which is probably the source of the enhanced excitation on the long-wavelength side. Decay-time measurements show a single-exponential dependence with a characteristic time of about 65 ns. If anything, the observed red shift probably understates the influence of the environment. Even in the small-grained ceramic, the majority of the Ce ions find themselves deep within the individual grains, in an environment largely unperturbed relative to the single crystal. The influence of the interface becomes progressively stronger for ions situated closer to the edge of the grain, but these make up only a small fraction of the total. Thus, the energy shift for the ions closest to the interface may be significantly greater than the observed shift in the peak position, but their contribution is submerged within the totally inhomogeneously broadened band. These energy shifts presumably correspond to those ions responsible for the anomalous fast decay of the y-excited emission already mentioned, which will be discussed in more detail in the next section. 3.2.2. 335 nm The intensity of the 335 nm peak (as of the 460 nm peak, just discussed) was essentially independent of the identity of the specimen, differing by less than 15%. This feature must also originate from the Ce levels, in this case the one immediately above the lowest, and has been so identified in the literature [S]. Its behavior is consistent with that of the 460nm peak, with the same decay behavior and intensities essentially independent of microstructure. These ceramic peaks also show a red shift relative to the monocrystal, of about the same magnitude in energy as the 460 nm peaks. 3.2.3. 280 nm Unlike the previous two, this feature does not originate in the 5d manifold of Ce, but rather in a defect from which the energy is transferred nonradiatively to the Ce-ion [5]. This feature occurs in both single crystal and ceramics, but disappears after heating in a reducing atmosphere. The decay traces from pulsed excitation at this wavelength are indistinguishable from those in the previous two cases, showing that the energy transfer is quite fast. The intensity of
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this excitation shows a mild systematic dependence on microstructure, being somewhat weaker in the smaller-grained ceramic. 3.2.4. 230 nm The 230 nm feature, like those at 460 and 335 nm, appears to originate from direct Ce-ion excitation [5, lo]. Its asymmetry suggests greater complexity than in the previously discussed cases, and Tomiki et al. [l l] claim that it in fact consists of three individual overlapping components, corresponding to the highest three of the five Sd-levels of the Ce3+ ion. Unlike the other two 5d components already discussed, however, the 230nm peak shows a strong systematic dependence on microstructure, being most intense in the single crystal and weakest in the smallgrained ceramic. The difference in behavior from the other components originating from the 5d manifold would have cast serious doubt upon the assignment, had we not also performed absorption measurements on the specimens. These revealed the presence of a broad microstructure-dependent parasitic absorption below 300nm (Fig. 6) underlying the Ce-associated features. By blocking light that would otherwise reach the Ce ion, this absorptive background could readily account for major diversion of excitation energy to nonradiative paths. The origin of this absorption is unclear, but it almost certainly is associated with a lattice defect. It is not present in as-grown single crystal, but can be generated by subjecting the specimen to a sequential reduction/oxidation heat treatment that simulates the processing regimen for the ceramics. The latter had been post-fired in air to bleach out the strong visible absorption introduced by hot-pressing under highly reducing conditions, which would otherwise block the Ce emission. While this did indeed remove the absorption in the visible, it replaced it with one in the short-wavelength UV that had not previously been present. Unlike the previous case, the responsible defect shows no ability to transfer to the Ce ion and, hence, is more likely to introduce a loss mechanism. The fact that the absorption is greatest in the small-grained ceramic suggests that the defect is not randomly distributed throughout the total specimen volume, but rather is associated specifically with material in the immediate vicinity of the grain boundaries.
E. Zych et al. /Journal
0
of’ Luminescence
-.‘....‘.r..‘....‘....‘200
250
300
Wavelength
350 (nm)
Fig. 6. UV absorption of YAG:Ce: (a) large-grain unbleached; (b) same (#I ), bleached; (c) small-grain bleached. Note the progressively stronger underlying absorption in post-heated samples.
400
ceramic #I, ceramic #2. background
3.2.5. 180-185 nm This may well be the most interesting feature of the optical excitation spectra. It appears to be associated not with localized excitation of the Ce activator ions but rather with the direct elevation of an electron from the valence to the conduction band (although not necessarily fully dissociated from its corresponding hole). The absorption associated with this excitation peak is far stronger than that associated with any of the other excitations already discussed, and, unlike the 230 nm feature, is far too intense for the background to distort its behavior as a function of microstructure. Here, we see a rough correspondence between the intensity of the excitation peaks and the scintillation light yield as measured under y-excitation, with a significantly greater discrepancy between small- and large-grained ceramic than between the latter and the single crystal. The absorption, however, is so strong that hardly any excitation can penetrate deeper than a few tenths of a millimeter into the bulk of the specimen. Indeed the apparent excitation peak itself may be only a surfaceassociated artifact, since the true absorption shows no corresponding decrease at shorter wavelengths. Because of its delocalization, the 180 nm excitation represents the closest simulation that we can achieve optically of what takes place under the ionizing effects of y-irradiation. The approximately 7 eV value
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for the band-to-band energy is in excellent agreement with the literature data [l 11. This energy, however, is just barely enough to raise a valence electron to the conduction band, with very little left over to assure liberation from its hole counterpart. Hence, the optical energy absorbed at this wavelength will go mostly to the generation of mobile excitons, in stark contrast to the free electrons and holes that are the preponderant product of y-excitation. Indeed, the high efficiency with which the Ce3+ ion can trap holes effectively suppresses any tendency for the free carriers to reassociate into excitons. Clearly, the dominant transfer mechanism under band-to-band optical irradiation must be quite different from the free-carrier mediated transfer that prevails under y-excitation [ 12, 131, with profound consequences on the kinetics of the process. Thus, it is hardly surprising to find that band-toband optical excitation produces decay patterns that differ markedly from those obtained by any other means. Unlike the traces from other optical excitation. these depart significantly from single-exponential behavior for all specimens, with slow components having characteristic times of at least 300 ns. These long decay-tails arise when some of the excitation energy is diverted into long-lived traps of the host lattice, from which a delayed transfer to the Ce ion then takes place [5, 61. A prime example of this is provided by the long-lived lattice defect emission at ~3 10 nm, whose overlap with the nearby Ce absorption provides an efficient transfer path. As seen in Table I, the slow component is present in both the monocrystal and ceramic, at about the same proportion. Further details of the lattice emissions will be presented in a separate paper [ 141. In contrast, the dominant component of the decay shows a strong dependence on microstructure, with characteristic times ranging from the expected radiative value of 63 ns in the small-grained ceramic to as high as 85 ns in the single crystal. As the excitation process begins to change from a localized optical mechanism to one that is delocalized and carrier-mediated, the pattern of the decay kinetics also begins to resemble more closely what we have already found under y-excitation. Nevertheless, the similarity of the band-to-band traces to those excited by y-irradiation is only skin deep. The latter shows a substantially higher (and somewhat slower) contribution of the long component, and. as mentioned
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earlier, are the only ones to reveal the presence of a yet faster component in the ceramic, superimposed on the normal emission. The reasons for these various differences are among the issues we must address.
4. Discussion The first issue we wish to consider is the shortfall in scintillation light output exhibited by the ceramic specimens. From photopeak measurements, we know that the efficiency of conversion of absorbed gammas into visible photons drops sharply as the grain size gets smaller, to about 50% and lo%, respectively, relative to the YAG monocrystal. The scintillation spectrum tells us that virtually all of the light output comes from the Ce activator, in its characteristic d-f emission between 480 and 640 nm, with no significant contribution from lattice defects or other intrinsic emission. These two points suggest that the high concentration of abnormal lattice in the vicinity of the grain boundaries may introduce loss mechanisms into the ceramic that would not otherwise be present in the single crystal. These losses may assume at least two different forms: a diversion of the excitation energy to nonradiative paths before it reaches the emitting level of the activator ion; or a partial draining of the emitting level of the already excited activator by some sort of nonradiative quenching. Let us consider these in turn. First, we shall examine the energy transfer process. It is now generally accepted that the primary mechanism of energy transfer is the sequential trapping of mobile holes and electrons by Ce-ions [ 12, 131, ultimately leading to the desired cerium emission. However, the carriers can also be trapped by other defects that may be present in the material, which can lead to other luminescence or (more likely) to nonradiative decay of the excitation. While always present in a single crystal, such defects must be far more numerous in a ceramic, which by its very nature is built up of densely packed but separate grains. Since the grain boundaries are necessarily regions where there is a high degree of lattice distortion and an abundance of point defects, it is reasonable to expect a much higher density of sites where the mobile carriers generated by the absorbed gamma (or band-gap energy photons) could be trapped and eventually decay nonradiatively than would be the case in the bulk crystal. This in turn
would decrease the available population of the mobile electrons and holes, which are needed to convey the excitation to the Ce activators [ 131, necessarily reducing the excited Ce population and hence the light output. Such an explanation is simple and straightforward, and would be perfectly acceptable but for one critical complication: the rapid emission component observed under y-excitation alone. This rapid emission component presents the greatest challenge in the interpretation of the results. Since the kinetics of energy transfer provide no means to accelerate the rate of decay of the emitting level, ’ the presence of a rapid component can arise from only two causes: an increase in the oscillator strength of the emitting transition, or the introduction of an alternative (nonradiative) path by which the excited level can relax to ground. But the first should be accompanied by a substantial increase in the intensity of the corresponding resonance absorption, which is not observed and would not in any event represent a loss process. While the second will, of course, introduce a loss, optical measurements offer no evidence of any significant quenching from the emitting level of the Ce activator in any of the ceramic specimens. We would appear to have reached an impasse. Seeking clues as to the source of this fast component, we subjected the ceramic specimens to a variety of additional tests. We made time-resolved measurements of the y-excited spectra, with gated detection centered at various delays after the absorption of the primary gamma, but found no discernible difference in spectral structure. We explored the full range of wavelengths accessible to optical excitation, without finding any indication of the fast component so readily generated by gammas. Nor could we find any evidence of a rise time, one of the tell-tale signs of delayed energy transfer. How can we reconcile the presence in the ceramic of a y-excited component four times faster than in the single crystal, in a proportion having an inverse relation to light yield, with the total absence ’ Note, however, that a retardation of the emission can readily be attributed to kinetic factors, specifically a slow step in the transfer process. Such a retardation is in fact observed in the YAG:Ce single crystal, where the y-excited decay time is at least 30% longer than the radiative value measured optically. Such kinetic considerations can have profound implications far beyond this particular material, and will be discussed in detail in a separate paper.
E. Zych et al. /Journal
of Luminescence
of any evidence of a corresponding component under optical excitation at any wavelength? Nevertheless, we do feel that we can offer a viable phenomenological model to explain our observations. At the core of this model is the recognition that the positions of the d-levels within the bandgap can vary widely according to the chemical properties of the host material. Indeed, it is not at all unusual to find some or even all of the d-levels falling within the conduction band itself. When this happens, optical excitation into such levels can cause the electron to leave the activator entirely, generating measurable photoconductivity. Such an autoionization process can compete effectively with the normally rapid nonradiative relaxation to the lowest level of the manifold, diverting energy away from that erstwhile emitting state. Indeed, since the radiative transition is even slower, the lowest 5d state must be substantially below the conduction band for any d-f emission to take place at all. It is the failure to satisfy this condition that is considered to be the reason why no d-f emission is found in various Ce-activated rare earth sesquioxides and oxysulfides [ 151. YAG : Ce, with its exceptionally high cubic and tetragonal distortion components of the crystal field [9] barely satisfies this condition, with its lowest d-level at only about 10 000 cm-’ below the conduction band [ 161. Consequently, not much perturbation of the crystal field at the Ce-site and/or the band structure of the host material itself is needed to upset this balance and kill the luminescence altogether. Both these effects are likely to occur at the grain boundaries of the YAG-ceramic. While the former may or may not alter the energies of the split dlevels (and in fact no evidence for this is found), they should not materially affect the kinetics. The latter, on the other hand, can fundamentally alter both the energy transfer and emission processes. In particular, the discontinuity in lattice periodicity, and the consequent profusion of broken chemical bonds at the interface, should produce in their immediate vicinity additional states, both filled and empty, at normally forbidden energies within the bandgap (see Fig. 7). This will specifically provide an abundance of “receiving states” for free carriers (particularly electrons) at energies not otherwise accessible to them, and could readily submerge within the extended bands the previously unobstructed (and much more localized) states
75 (1997)
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g CONDUCTION s BAND
-.
Ce-levels
s VALENCE ==zzz=z s BAND _~..____ ____
A
_.
Fig. 7. Schematic diagram of the effect of the grain boundaries upon the band structure of YAG:Ce ceramic. Note the local decrease in the bandgap, with the conduction band dropping to or below the energy of the Ce emitting level.
for the Ce optical transition, thereby largely quenching the emission. The effect must certainly depend on the distance from the interface, which we are not in a position to quantify. We propose that the fast (~20 ns) component in the y-excited decay of Ce-ions in the YAG ceramic is due to a quenching process by means of autoionization or tunneling of electrons from the excited ions to “receiving” states residing at the grain boundaries of the host. These receiving states can be viewed as continuations of levels within the conduction band characteristic of the crystal lattice, but moved to lower energies by the lattice discontinuities at the interface. Conceptually, these are the counterparts of the surface states of the original crystalline granules, but now folded into the bulk of the ceramic. The rate of the tunneling is enough to produce a rapid, partially quenched but still readily detectable ~20 ns component in the decays of the ceramic excited by gamma irradiation. Indeed, we find that the contribution of the fast component clearly decreases with increasing grain size. As the number of Ce ions situated in this layer and hence subject to the
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quenching process constitute only a small fraction of the total number of Ce ions in the ceramic, their contribution under optical excitation should be very low, and indeed is not detectable in our measurements. How then can such a small number of ions give rise to such an easily detectable fast component under y-excitation? This we attribute to the fact that such excitation in the ceramic generates a very different distribution of excited Ce-ions than does optical excitation, or even y-excitation in the single crystal. This is inherent in the nature of the cascade that generates the carriers that convey excitation to the Ce activators. This process begins with the absorption of a gamma, generating a relatively small number of highly energetic carriers. These, on collision with lattice electrons, give up some of their excess kinetic energy in the creation of yet more free carriers, with the cascade continuing until the remaining excess kinetic energy per carrier is too low to create more. Thus, by the time these carriers reach thermal equilibrium and become available for recombination, momentum conservation has given rise to a veritable wave front of carriers, moving outward and forward a substantial distance from the point of initial gamma absorption. In the single crystal, this distance is essentially irrelevant; the lattice is uniform and ordered, and each activator is essentially identical. Under optical excitation, the distance is zero, and hence also irrelevant. In the ceramic, however, the lattice discontinuities at the grain boundaries create a network of trenches athwart the path of the hot carrier wave front, in which, as we have already seen, the conduction band is lowered (and the valence band raised) relative to the values in the bulk crystal. Thus, if the average distance a hot carrier travels before thermalization is comparable to the grain size, it can fall into the trench and lose enough kinetic energy to be unable to escape. Then, for the remainder of its mobile life, it will be constrained to travel within this narrow trench defined by the grain boundary, available to excite only a Ce ion in the immediate vicinity of the interfacial region. But, as we have also already seen, these are precisely the ions most subject to nonradiative quenching by interaction with the distorted conduction band. In this manner we see how hot carriers generated under y-excitation can become channeled to the interfacial regions between the grains, substantially altering the distribution of the excitation density in the grain itself by increas-
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ing the relative probability that the Ce emitting centers in such regions become excited. This does not happen under optical excitation, which either targets Ceions directly or, under band-to-band excitation, favors an excitonic mechanism of energy transfer which is less prone to the effect we have just described. Such a “focusing” of the gamma-induced excitation into the interfacial regions could well overcome their limited population of emitting centers, too low to be detected optically. And any such “focusing” must also necessarily increase the probability that the mobile carriers themselves will be trapped at non-emitting defect sites (such as, perhaps, those responsible for the VUV background absorption), preventing their energy from reaching the activator ion at all. The decay traces and light yield measurements under y-excitation even provide us with tools to estimate the relative contributions of the two loss mechanisms, carrier attrition and activator quenching. Table 1 gives us the decay times for the anomalous fast (-20 ns) component in the ceramic specimens, as extracted from a three-term exponential fit. If we assume that the normal (unquenched) decay time would have been the same ~58.5ns as is found in the single crystal, we can calculate the effective quantum efficiency for this component as the ratio of the two decay times. Then, since we know the magnitude of the contribution that this component makes to the total light, we can derive the amount of light we would have had if there had been no quenching, and hence the effective quantum efficiency of the total Ce emission in the ceramic, relative to that in the single crystal. Finally, the factor by which this ratio differs from the light yield of the ceramic (also relative to the single crystal) gives us the fraction of carriers that have been lost. These results are summarized in Table 2. While we should not take these values as definitive, it is interesting to note that it is the carrier attrition that is the more serious of the two loss mechanisms, being both larger than the activator quenching and more sensitive to the grain size. At this point we can offer no positive evidence for this model. It is only phenomenological and must await further study, both mathematical and experimental, to develop it into a consistent and mature quantitative theory. Until then it will remain purely conjectural, supported only by the inapplicability of all others that we have been able to conceive. Yet it is
E. Zych et al. /Journal Table 2 Estimated losses in gamma-excited ceramic specimens”
emission
from
qf’ Luminescence
YAG:Ce
Avg. grain si7e
Carrier attrition (%)
Activator quenching
40 pm IOpm
39 84
18 32
7.5 II9971
193-203
203
CA62330-03. We also express our great appreciation to Prof. A. Lempicki of Boston University, for invaluable advice and assistance during the course of the work.
(%)
Reference [I] W.H. Rhodes, J. Am. Ceram. Sot. 64 (1981) 13; G.C. Wei.
“Relative
to the single crystal.
axiomatic that if all other explanations are ruled out by the evidence, then the only one remaining, no matter how speculative, must be accepted [17].
[2] [3] [4]
5. Conclusions [5]
We have shown that transparent ceramic specimens of YAG: Ce can be reproducibly prepared by means of a hot-pressing technique. The scintillation light output from these specimens is significantly lower than in the single crystal, showing an inverse relationship with grain size. The interfacial regions between the grains are the sites of two spatially localized loss processes: an attrition of the free carrier population through nonradiative recombination at defect sites; and a quenching of the emitting level of the Ce ion through its interaction with the deformed band structure at the grain boundaries. Both of these effects appear to be enhanced by an unusual transport process, wherein the free carriers tend to be channeled preferentially into the interfacial region. There is no doubt that the quality of the material can be significantly improved by optimizing the fabrication procedure, but the extent to which the actual scintillator performance can be improved remains to be seen.
[6] [7]
[8]
[9] [IO] [II]
[I21 [13]
[14]
[15]
Acknowledgements The authors gratefUlly acknowledge the support of the U.S. Department of Health and Human Services, Public Health Service (NIH), under Grant 1 R01
[16] [17]
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