recent r&d trends in inorganic single‐crystal scintillator...

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Recent R&D Trends in Inorganic Single-Crystal Scintillator Materials for Radiation Detection

Martin Nikl * and Akira Yoshikawa

Dr. M. Nikl Institute of Physics, AS CR Cukrovarnicka 10, 16253 Prague , Czech Republic E-mail: [email protected] Prof. A. Yoshikawa Institute for Materials Research Tohoku University 2–1–1 Katahira , Aoba-ku , Sendai 980–8577 , Japan

DOI: 10.1002/adom.201400571

the scintillation mechanism and a defi ni-tion of its effi ciency criteria had already been developed in the 1970s [ 5 ] and later further refi ned. [ 6 ] The scintillation mecha-nism can be divided into three consecu-tive stages: conversion, transport, and luminescence. In the conversion stage, depending on the energy of the entering photon, its initial multistep interaction with the scintillator lattice occurs pre-dominantly through i) the photoelectric effect (below approximately a few hun-dred keV, depending on the density and effective atomic number of the material), ii) Compton scattering (from a few hundred keV to 8000 keV), and iii) pair production (above the latter limit). Hot electrons and deep holes that are created in this process are gradually thermalized in the conduc-tion and valence band edges, respectively.

Thermalization of carriers within the conduction and valence bands is sometimes considered a separate stage in the scintil-lation process due to its importance in the study of nonpro-portionality issues in the scintillation mechanism. [ 7 ] The entire conversion stage lasts typically few ps (see the literature for a more detailed description. [ 7–9 ] In the transport stage, the sepa-rated electrons and holes have to reach the emission centers, i.e., migrate through the host material in the conduction and valence bands, respectively: they can be repeatedly trapped or even nonradiatively recombined at trapping levels arising from lattice defects. A considerable delay in charge-carrier delivery to the luminescent centers can be introduced due to such trapping processes which are responsible for slow components in scintil-lation decay and afterglow, typically occuring on microsecond and millisecond time scales, respectively. This stage is the least

In this review, the major achievements and research and development (R&D) trends from the last decade in the fi eld of single crystal scintillator materials are described. Two material families are included, namely, those of halide and oxide compounds. In most cases, the host crystals are doped with Ce 3+ , Pr 3+ or Eu 2+ rare earth ions. Their spin- and parity-allowed 5d–4f transitions enable a rapid scintillation response, on the order of tens to hundreds of nanoseconds. Technological recipes, extended characterization by means of optical and magnetic spectroscopies, and theoretical studies are described. The latter provide further support to experimental results and provide a better understanding of the host electronic band structure, energy levels of spe-cifi c defects, and the emission centers themselves. Applications in medical imaging and dosimetry, security measures, high-energy physics and the high-tech industry, in which X(γ)-rays or particle beams are used and monitored, are recognized as the main driving factor for R&D activities in this fi eld.

1. Introduction

A scintillator works as a spectral and energy transformer: it converts a high-energy photon from the X- or gamma-ray range into a bunch of UV–vis ones, i.e., to a fl ash of light. Alterna-tively, accelerated charged particles (electrons, protons, or heavier ions) or even neutrons can be detected through their energy deposit in interactions with a scintillator host, which is again converted into a fl ash of light. Scintillation is a common property of matter in general, and is observed in gas, liquid, or solid. Among solids, two big families, namely, organic and inorganic scintillators, have been developed. Inorganic scintil-lators are probably the most used in practical applications now-adays and intense research and development (R&D) activities have continuously taken place over last few decades [ 1,2 ] with big market potential and prospects for the coming years. [ 3 ] In prac-tice, a scintillator detector consists of two parts: the scintillating material itself, and a photodetector which converts the UV–vis photons into an electrical signal, I(t) [ 4 ] (see Figure 1 ).

Dielectric or semiconductor wide bandgap materials are employed for such a task. A phenomenological description of

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1 photon n photons

(keV - GeV) (2 - 4 eV)

scintillator

Opticalcoupling

Photo-detector

I(t)

scintillator

Figure 1. Workings of a scintillator material and set-up of a scintillation detector.

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http://doi.wiley.com/10.1002/adom.201400571

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introduce energy levels into the forbidden gap and strongly modify/degrade scintillation performance. These phenomena are strongly dependent upon the manufacturing technology. [ 10 ] During the fi nal luminescence stage, the trapping and radia-tive recombination of the electron and hole at the luminescent center give rise to the desired luminescence light. Excitons or other intermediate excited states can be formed and their energy transferred to the luminescent centers as well.

In the case of scintillators, X(γ)-ray photon counting con-sists of accumulating the generated light arriving soon after the initial conversion stage is accomplished. This is because the scintillator works as a high-energy photon counter, usually using the so-called pulse-height measurement technique in the detection electronics. [ 11 ] Signifi cantly delayed light, such as that due to the retrapping processes mentioned above, cannot be technically exploited in the counting mode. The most impor-tant characteristics of scintillation materials are the following:

i. Scintillation effi ciency ii. Light yield (LY) iii. Linearity of light response with the incident X(γ)-ray photon

energy – energy resolution and nonproportionality iv. X(γ)-ray stopping power v. Scintillation response in time vi. Spectral matching between the scintillator and the

photodetector vii. Chemical and mechanical stability viii. Radiation resistance ix. Price

The overall scintillation effi ciency of X(γ)-ray-to-light conver-sion is determined both by intrinsic and extrinsic material char-acteristics. The number of UV–vis photons, N ph , produced in the scintillation conversion per energy E of incoming X(γ)-ray photon can be expressed as: [ 5,6 ]

N

E

ESQph

gβ= ×

(1)

where E g represents the forbidden gap of the material, S and Q are the quantum effi ciencies of the transport and luminescence stages, respectively, and β is a phenomenological parameter which is typically found to be between 2 and 3 for most mate-rials. The relative effi ciency can then be obtained as:

η =

E N

Evis ph

(2)

where E vis is the energy of the generated UV–vis photons. The most effi cient material among the phosphors and scintil-lators today is ZnS:Ag, with η ≈ 0.2. A simple, practical way to estimate the relative scintillation effi ciency of a scintillator is to measure its radioluminescence spectrum under X-ray or γ-ray excitation in the steady-state mode, and to compare it on an absolute scale with a standard scintillator sample measured under the same experimental conditions. [ 4 ] Though some errors on the order of a few tens of percents might arise due to the different nonproportionality characteristics of materials under

comparison in the energy range of several tens of keV, such a comparison is simple to make and meaningful.

The light yield (LY) of a scintillator is always less than what is given by Equation ( 1) , since it represents only a fraction of the visible photons generated, namely, those arriving at the photodetector within a certain time gate defi ned by the detec-tion electronics (the so-called shaping time) after high-energy photon absorption. The values of the shaping time are usually set between 100 ns and 10 µs. A more detailed description of other scintillator characteristics can be found in the literature, [ 4 ] and the measurement techniques themselves are well described in the book of Knoll. [ 11 ]

Research on scintillation materials started with the discovery of X-rays in November 1895 by W. C. Roentgen. [ 12 ] For X-ray registration, a simple photographic fi lm was found to be rather ineffi cient. That is why the search for materials able to convert this new invisible radiation into visible light started immedi-ately, in order to use sensitive photographic fi lm-based detec-tors effi ciently. CaWO 4 powder was employed for this purpose just a few months later in early 1896 and, together with the ZnS-based powders introduced later, these powder phosphor materials became widely used for the detection of X-rays in such combined phosphor and fi lm detectors. [ 13 ]

The history of bulk single-crystal scintillators begins in the late 1940s with the introduction of NaI:Tl and CsI:Tl

Martin Nikl graduated in 1981 from the Faculty of Nuclear Science and Physical Engineering, Czech Technical University, and obtained his PhD in 1986 at the Institute of Physics, CAS. His research interests include lumines-cence and scintillation mecha-nisms in wide bandgap solids, energy-transfer processes, and the role of material defects in

them. He currently serves as the department head at the Institute of Physics, AS CR.

Akira Yoshikawa received his PhD from the University of Tokyo, Japan (1999). He gained experience in crystal chemistry, crystallography, and complex oxide and halide single crystal growth from the melt. His research target is to develop novel single crystal materials for scintillators, lasers, nonlinear optics, and piezoelectric

applications. He is a Professor at the Institute for Mate-rials Research and New Industry Creation Hatchery Center, Tohoku University.



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scintillators, [ 14,15 ] which continue to be used in a number of applications. Since that time, a number of other material sys-tems have been reported (see reference [ 16 ] for an historical overview). NaI:Tl and CsI:Tl halide scintillators and the fi rst oxide-based CdWO 4 scintillator, [ 17 ] together with Bi 4 Ge 3 O 12 (BGO) introduced later, [ 18 ] became widespread materials and are often used as “standard samples” to evaluate the perfor-mance of new materials under study. Within the last two decades, there has been considerable activity in this fi eld. It was mainly triggered by the needs of high-energy physics in the 1990s to fi nd new scintillators for the Large Hadron Col-lider calorimetric detectors. This led to the optimized PbWO 4 scintillator. [ 19,20 ] Challenging concepts for further scintilla-tion material R&D have recently been introduced in high-energy physics by so-called dual read-out calorimetry [ 21 ] and the inorganic single-crystal fi ber concept, which paves the way to homogeneous or high sampling-fraction calorimeter designs with high-granularity 3D imaging as well as a par-ticle identifi cation capability. [ 22 ] The advanced imaging and dosimetric applications in medicine [ 23 ] established scintilla-tion ceramics [ 24 ] in computed tomography (CT) and Ce-doped orthosilicates in positron emission tomography (PET). [ 25,26 ] Various high-tech industrial applications use aluminium per-ovskite and garnet scintillators. [ 19,27 ] Most recently, security-related applications have called for materials of special com-positions for neutron detection. [ 28,29 ] The absolute majority of new single-crystal scintillators, reported in this period, have been based on Ce 3+ -doped and Pr 3+ -doped materials, due to the fast decay time (typically 10–60 ns) and high quantum effi ciency of the 5d–4f radiative transitions of these centers at room temperature (RT), and in some cases even well above that. [ 30 ] For a recent broad survey of materials see, for example, reference. [ 31 ] More recently, the Eu 2+ center and its 5d–4f tran-sition also has been used, especially in halide scintillators, as will be described in Section 2.3. This trend appeared due to the need for security-related techniques where high light yield and high energy resolution scintillators are needed, while a somewhat lower speed of scintillation response (units of microseconds) can be tolerated.

Due to the practical importance and relatively long history of this fi eld, there already exists a large amount of published information on the topic. Let us mention the superb survey of luminescent materials written by Blasse and Grabmaier, [ 32 ] the already-mentioned overview of methodology in use in radia-tion measurement authored by Knoll, [ 11 ] monographs devoted to scintillator materials from Rodnyi [ 1 ] and Lecoq et al. [ 2 ] and the phosphor handbook of Shionoya and Yen. [ 33 ] Numerous featured review papers on phosphor and scintillator mate-rials and their applications exist in the scientifi c literature as well. [ 4,24,34–37 ]

In this paper, we address novel material systems as well as strategies for their development and optimization that have appeared in the last decade. Apart from the materials them-selves, we will also mention the effort to understand in detail the physical fundamentals and bottlenecks for certain param-eters, e.g., scintillator nonproportionality. Further development of computational methods for the theoretical description and eventual forecast of promising material systems is demon-strated and discussed.

2. Halide Scintillators

Binary halide scintillators, namely, Tl-doped NaI and CsI single crystals, have already been mentioned above [ 14 , 15 ] as the very fi rst bulk material systems exploited in this fi eld. Alkali halides were also used for many years as model systems, and extended knowledge and a detailed understanding were gained, for example, with respect to their point defects. [ 38 ] Together with point defect-related traps, the mechanisms of charge or exciton self-trapping [ 39 ] are areas of signifi cance in scintillator physics. Interestingly, even after many years of practical exploitation, a novel interpretation of the luminescence mechanism in Tl-doped CsI was proposed and generally accepted in 1990s [ 40 ] and efforts have been expended for decades to remove degrading phenomena, as shown below.

2.1. Co-doped CsI:Tl

The CsI:Tl single-crystal scintillator was mentioned above as a widespread material in practical applications. Its drawback is a high afterglow, which disables its use in, for example, CT imaging. Co-doping with various aliovalent ions has been used in the effort to improve this parameter. It is worth men-tioning the Eu 2+ and Sm 2+ co-dopants. [ 41–43 ] The atomistic mechanism of Sm 2+ and Eu 2+ functioning is the following: in the former case, a nonradiative recombination channel is introduced that decreases the delayed radiative recombination of trapped carriers. In the latter case, deep electron traps are created, scavenging electrons from the shallow traps associ-ated with thallium. In both cases, however, the LY is seriously deteriorated. Bi co-doping was reported as an effi cient tool to decrease CsI:Tl afterglow, which became less than 0.1% after 10 ms without a strong decrease in the radioluminescence effi -ciency under low-energy X-ray (less than 30 keV) excitation. [ 44 ] However, a later study [ 45 ] found that the light yield and energy resolution of Bi-co-doped CsI:Tl crystals became much worse even for the lowest Bi concentration of about 0.005 mol% (in the melt), which makes its practical exploitation problematic. Most recently, Yb 2+ co-doping was introduced and the opti-mized CsI:Tl,Yb crystals exhibited an ultra-high light yield value of 90 000 ± 6000 photons MeV −1 , an energy resolution of 7.9%@511 keV, and suppressed afterglow down to 0.035% at 80 ms (see Figure 2 ). Simultaneous improvement of the after-glow level, light yield, and energy resolution in the Yb-co-doped CsI:Tl scintillator is considered to be a breakthrough in the optimization of this scintillator, and paves the way for its appli-cation in X-ray-based fast-imaging applications. [ 46 ]

2.2. Rare Earth Halides

Fast and effi cient scintillators were discovered within the family of Ce-doped rare-earth (RE) halides (see the review [ 47 ] ). LaX 3 :Ce single crystals, where X = Cl and Br, fi rst reported in 1999 [ 48 ] and 2000, [ 49 ] respectively, showed the highest light yield and best energy resolution. Potentially even higher performances might be achieved with LuI 3 :Ce, [ 50 ] but the high content of slower components in scintillation decay and the very high



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price of the raw materials have hampered its further develop-ment. Their characteristics are listed in Tables 1 and 2 , where the scintillation decay time, LY, and energy resolution values are given for a specifi c Ce concentration.

Single crystals of Ce-doped RE halides are usually obtained by the vertical Bridgman growth technique, using binary RE halides as starting materials. The grown crystals are highly hygroscopic. They must be sealed in appropriate ampoules for handling and characterization. The Bridgman growth of large-size LaX 3 :Ce crystals with very good optical and scintillation characteristics has been described in the literature. [ 51–53 ]

In Figure 3 , the X-ray excited emission spectra of LaCl 3 :Ce 3+ at different Ce 3+ concentrations are shown. [ 54 ] Apart from a Ce 3+ doublet peaking at 337 and 358 nm, a shoulder attrib-uted to self-trapped exciton (STE) emission is observed at about 400 nm. Its excitonic nature was established by an optically detected magnetic-resonance investigation. [ 55 ] The Ce 3+ scintillation decay includes a fast component due to the prompt trapping of holes and electrons by cerium ions, which is dominated by the Ce 3+ 5d–4f emission lifetime. The slower component, on a microsecond timescale, is interpreted

in terms of thermally assisted energy transfer from STE to cerium ions. [ 56 ]

The best scintillation performance in terms of LY, decay time, and energy resolution has been achieved for LaBr 3 :Ce. [ 49 ] The scintillation decay of LaBr 3 :Ce is dominated by a fast and approximately single-exponential component with a decay time of about 15 ns that approaches that of the Ce 3+ 5d–4f transition. Thus, effi cient direct energy localization at Ce 3+ ions occurs and the energy transfer from STE in the scintillation mechanism, if any, does not introduce any additional delay. A higher number of electron–hole pairs is generated in the LaBr 3 host compared to LaCl 3 , due to a smaller bandgap (see Table 1 ) which, together with the above-mentioned fast charge-carrier transfer to Ce 3+ centers, give rise to a very high LY of about 70 000 ph MeV −1 .

For completeness, we also mention studies dealing with Pr-doped RE halides, [ 57,58 ] namely, LuCl 3 :Pr, LuBr 3 :Pr, and LaBr 3 :Pr. None of these materials shows a fast scintillation response, since the 5d state of Pr 3+ is not populated or is nonra-diatively exhausted in the scintillation mechanism. In LaBr3:Pr, the reason is that the charge transfer from the valence band to Pr 3+ occurs at energies below the energy of the 5d–4f emission, while in LuX 3 :Pr (X = Cl, Br), overlap of the STE luminescence with the Pr 3+ 4f2–4f2 interconfi gurational transitions seems to be the most probable scenario.

A drawback of LaX 3 :Ce scintillators for low count rate appli-cations (e.g., in radioastronomy) is their intrinsic radioactivity due to the presence of the isotope 138 La, which has a natural abundance of 0.09% and half-life τ 1/2 ≈ 10 11 years. [ 59 ] That was why the R&D studies also included the binary CeX 3 (X = Cl,Br) single crystals, where the intrinsic radioactivity is at least one order of magnitude lower. They were prepared using techno-logical methods very similar to those used for LaX 3 :Ce (see for example ref. [ 60 ] . Like the latter crystal family, [ 61 ] also in the case of CeX 3 , the growth of anion-mixed crystals has recently been reported. [ 62 ]

Especially for the case of CeBr 3 , characteristics have recently been reported in the literature by several groups. [ 63–66 ] The reported emission maximum due to the 5d 1 –4f transition of Ce 3+ is at about 370 nm, and the related decay time of 17 ns is very close to that of LaBr 3 :Ce. The scintillation light yield and



Figure 2. Afterglow profi les of CsI:Tl + ,Yb 2+ group B crystals after X-ray pulse excitation. Reproduced with permission. [ 46 ] Published by the RSC.

Table 1. Optical and scintillation characteristics of the most effi cient modern halide scintillators. The number in parenthesis in the scintillation decay time column represents the percentage of the total intensity governed by the indicated decay time, while slower components characterize the decay of the remaining intensity. Times in parenthesis in the LY column represent the shaping times over which the LY was evaluated (if not provided, typical values are within 1–4 µs). (n.r.: not reported)

Crystal Density [g cm −3 ]

Bandgap [eV]

Ce 3+ (Eu 2+ ) 5d–4f emission [nm]

Ce 3+ (Eu 2+ ) 4f–5d absorption [nm]

Ce (Eu) conc. [mol%]

Scintillation decay time [ns]

LY [10 3 Ph MeV −1 ]

Energy res. [%] @662 keV

LaCl 3 :Ce 3.86 7 337,358 243,250,263,274,281 10 24 (60%) 50 3.1

LaBr 3 :Ce 5.03 5.6 355,390 260,270,284,299,308 5 16 70 2.6

LuI 3 :Ce 5.6 n.r. 475,520 ≈300,390,419 0.5

2

<50 (50%) 42 (0.5 µs) 51 (10 µs)

58 (0.5 µs) 71 (10 µs)

4.7

CeBr 3 5.18 n.r. 370,390 n.r. 100 17 60 4.1

CeBr 3 :Sr 5.18 n.r. 370,390 n.r. 99.5 17 55 3

LaBr 3 :Ce,Sr 5.03 5.6 355,390 260,270,284, 299,308 5 18 (78%) 82–2500 (22%) 77 2.0

SrI 2 :Eu 4.6 5.5 435 n.r. 5 600–1600 80–120 2.6–3.7

CsBa 2 I 5 :Eu 4.8 n.r. 435 n.r. 7 48 (1%) 383 (6%)

1500 (68%) 9900 (25%)

80–97 3.8


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energy resolution under comparable experimental conditions are about 10% and 25% lower compared to those in LaBr 3 :Ce, respectively (see Table 1 ). However, due to its dramatically reduced intrinsic radioactivity, CeBr 3 gamma-ray detection sensitivity is about an order of magnitude better than that of LaBr 3 :5%Ce at energies within 1–3 MeV. [ 66 ] CeBr 3 also offers an excellent time resolution of 93 ps at 511 keV for 1 cm 3 crystals, which is 2–3 times better that that of LaBr 3 :Ce [ 67 ] or LSO:Ce (see Section 3.3) under comparable conditions. This feature points to the great potential of this material for time-of-fl ight-based applications, including PET.

Strengthening the crystal lattice of brittle lanthanide halides appeared to be needed in order to increase the ruggedness of these scintillators for application. For this purpose, aliovalent doping by selected divalent (Ca 2+ , Sr 2+ , Ba 2+ , Zn 2+ , Cd 2+ , and Pb 2+ ) and tetravalent (Zr 4+ and Hf 4+ ) cations was attempted at a level of 500–1000 ppm in CeBr 3 . [ 68 ] Only Pb 2+ appeared to be detrimental to the scintillation properties of CeBr 3 , and all the doped crystals showed a noticeable increase in the frac-ture toughness compared to undoped CeBr 3 . [ 69 ] Another study,

focussing on the energy-resolution enhancement in monova-lent and divalent ion-doped CeBr 3 , found that Sr dopant clearly improves the energy resolution down to at least 3%@662 keV [ 70 ] due to the diminished nonproportionality (see Table 1 ). In fact, the same positive effect was found for Sr-co-doped LaBr 3 :Ce (see Figure 4 ). This material established the world record in energy resolution for single-crystal inorganic scintillators at 2%@662 keV [ 71 ] (see Table 1 ). Several proposals were consid-ered to explain these experimental observations. [ 72 ] Contrib-uting factors were considered to be a reduction of the nonra-diative recombination rate and/or an increase of the so-called escape rate of the carriers from the quenching region, and/or an increase in the trapping rate of Ce 3+ . The mechanism for this improvement was further studied by electronic structure cal-culations [ 73 ] with the conclusion that Sr-co-doping in LaBr 3 :Ce gives rise to shallow Sr La –V Br electron traps which effectively reduce the free-electron density in the fi rst instants (2–10 ps) of the scintillation process. Nonradiative Auger recombination has been determined to be the dominant quenching process in this time period of the scintillation process, [ 74 ] showing a

Table 2. Optical and scintillation characteristics of selected oxide-based scintillators. Dopant concentration shows a typical value for the crystal. Scin-tillation decay time of the dominant component is provided. Shaping times of 1–4 microseconds were used in LY measurements. Intervals of values refl ect those reported in the literature and/or measured in the author’s laboratories using the high-quality single crystals currently available.

Crystal Density [g cm −3 ]

Bandgap [eV]

Ce 3+ (Pr 3+ ) 5d 1 –4f emission [nm]

Ce 3+ (Pr 3+ ) 4f–5d 1 absorption/exc [nm]

Ce(Pr) conc. [mol%]

Scintillation decay time [ns]

LY [10 3 Ph MeV −1 ]

Energy res. [%] @662 keV

YAG :Ce 4.56 7.5 550 458 0.2 90–100 28–30 6–7

LuAG :Ce 6.67 7.8 525 448 0.15 55–65 24–26 6–7

GGAG:Ce 6.2 6.5–7.0 540 440–450 0.3 90–170 50–58 4.2–5.2

LuAG:Pr 6.67 7.8 308 284 0.1 20–22 18–20 4.6–5

LuYAG:Pr 6.2–6.5 7.7 310 286 0.1 20–22 27–33 4.4–6

YAP:Ce 5.35 8.2 365 303 0.2 19–25 22–25 4.5–5.5

YAP:Pr 5.35 8.2 247 215 0.1 8–10 6–12 11–13

LYSO:Ce,Ca 7.2 7.2 400 357 0.1 30–35 30–32 8–9

(Gd,La)PS:Ce 5.4–5.7 6.6–6.8 365–370 338 0.3 45–50 36–41 5–6



Figure 3. X-ray excited emission in LaCl 3 :Ce as a function of temperature and concentration. Unless otherwise indicated, the spectra are separated by 50 K temperature intervals. Reproduced with permission. [ 54 ] Copyright 2005, Wiley-VCH.

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cubic dependence on electron and hole densities. Given the reduction of free-electron density mentioned above, nonradia-tive quenching will be dramatically reduced and, consequently, the nonproportionality reduced in the low-energy region, as observed in the experiments. [ 70–72 ]

2.3. Eu-Doped Halides

Divalent and trivalent europium are the most common lumi-nescent ions in practical use. Usage of Eu 2+ in the last decade was further enhanced due to the intense R&D and applications in the fi eld of persistent luminescence materials [ 75,76 ] and phos-phors for solid-state lighting. [ 77 ] Despite the fact that the life-time of the 5d–4f luminescence transition in Eu 2+ is at least one order of magnitude longer than that of the Ce 3+ center, R&D for Eu 2+ -doped scintillators became a hot topic in the last decade. The reason is upcoming security-related applications, [ 28,29 ] where high light yield and excellent energy resolution are crit-ical parameters, while a scintillation response on the order of a few microseconds is still acceptable. Due to the 2+ charge state of the europium center, other host matrices can be considered and exploited, rather than those suited for doping with Ce 3+ or Pr 3+ emission centers. In fact, one of the currently most studied systems, recently rediscovered, are single crystals of SrI 2 :Eu, [ 78 ] even if the material was patented almost 50 years ago. [ 79 ] The reason for its commercial failure in the 1960s was its high price and the availability of smaller crystals compared to NaI:Tl, which satisfi ed the absolute majority of applications in those times.

Single crystals of SrI 2 :Eu are usually grown from anhydrous iodides by the vertical Bridgman technique in evacuated silica ampoules. [ 80 ] The growth of 2 inch-diameter crystals has been reported recently. [ 81 ] The use of a modifi ed micro-pulling-down technique was also reported, achieving 1 inch-diameter crys-tals grown in a graphite crucible [ 82 ] (see Figure 5 ). The emis-sion peak of the Eu 2+ center is positioned at 435 nm. Very high light yield values are reported, between 80 000–120 000

photons MeV −1 with energy resolutions of 2.6–3.7%@662 keV (see Table 1 ). The reported decay time values vary in the range 0.6–1.6 µs, depending on the Eu concentration, sample size, and measurement geometry due to a small Stokes shift and resulting reabsorption. [ 83,84 ] Reabsorption is one of the dis-advantages of this system, together with its low tolerance to moisture, which is worse than that of NaI:Tl. Experimental data regarding the STE emission (and its temperature dependence) peaking at 3.4 eV in undoped SrI 2 were also published. [ 85 ] They may help in further understanding the energy-transfer process towards doped emission centers in SrI 2 .

In the search for iodide host crystals suitable for Eu 2+ doping, several ternary compounds have also been recently reported in the literature. Eu-doped CsBa 2 I 5 and Cs(Sr , Ba)I 3 in the form of single crystals have been grown using the vertical Bridgman method, [ 86–90 ] where the CsBa 2 I 5 shows better tolerance to moisture compared to SrI 2 :Eu. The Eu 2+ emission peaks in the blue spectral region, similar to that of the SrI 2 host. For 2%Eu:CsBa 2 I 5 and a 1.0 × 1.0 × 1.0 cm 3 sample dimension, a very good energy resolution of 3.9% at 662 keV, a high light output of about 80 000 photons MeV −1 , and a principal decay time of about 900 ns were reported [ 89 ] (see Table 1 ). The pres-ence of a monovalent cation site in these materials provides a good opportunity to dope with ns 2 ions, [ 91 ] namely, In + and Tl + , where the former especially provides a new doping strategy in halide scintillators, rarely considered before. In the CsBa 2 I 5 host, encouraging values of light yield within 33 000–40 000 photons MeV −1 and energy resolution 7.1%@662 keV were reported for these dopants. [ 92 ]

2.4. Theoretical Studies

In synergy with the experimental studies mentioned above, more and more effort has been devoted to theoretical calcula-tions and modelling of the electronic band structure of the host materials, energy-level positioning, properties, and interactions of the emission centers with both the host and particular defect-related energy levels in the host forbidden gap, which gives rise to charge-carrier traps affecting the conversion and transfer stages of the scintillation mechanism.

Electronic band structure calculations have recently been published for LaBr 3 . [ 93 ] Unlike standard density functional theory (DFT) calculations that cannot correctly reproduce some



Figure 5. a) Photograph of SrI 2 :Eu crystal grown by a modifi ed micro-pulling-down method. [ 82 ] b) SrI 2 :Eu crystal sealed in aluminium can with a quartz window.

Figure 4. The X-ray response curves for CeBr 3 and Sr 2+ co-doped CeBr 3 and LaBr 3 :Ce scintillators compared with that of a standard LaBr 3 :5%Ce scintillator. Reproduced with permission. [ 71 ] Copyright 2013, AIP Pub-lishing LLC.


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experimental quantities such as the bandgap, the currently used quasiparticle self-consistent method correctly reproduces this value as well as the level ordering and spacing of the 4f and 5d states. Thus, it provides an excellent starting point for inves-tigating the electronic structure of excited states, charge self-trapping, and activator energy-level positioning in LaBr 3 and related materials. Also for CeX 3 (X = Cl, Br), the structural, elec-tronic, and optical properties have been studied with the help of DFT calculations within the generalized gradient approxima-tion. [ 94 ] The calculated lattice parameters are in good agreement with the experimental data. Their optical properties, including dielectric function, absorption spectra, refractive index, extinc-tion coeffi cient, and refl ectivity, are calculated as well. In SrI 2 , improved DFT calculations [ 95 ] revealed that, despite the struc-tural anisotropy (orthorhombic crystal structure, space group Pbca), the optical constants of SrI 2 are almost isotropic over a wide energy range. This indicates the possibility of preparing such a material and also other binary and ternary halides with such a property [ 96 ] in the ceramic form, which might be advan-tageous, especially for some imaging applications.

As for the emission centers, the systematic study of chemical trends of electronic and optical properties of ns 2 ions in halides has already been mentioned. [ 91 ] DFT calculations provide fur-ther insight into the luminescence mechanism of ns 2 ion-doped alkali halides. The activator-halogen hybridization strength and the ionicity of the host material strongly affect the positions of the activator levels relative to the valence and conduction band edges and provide strategies for exploring different combina-tions of host materials and activators for desired luminescence mechanisms. Similar DFT calculations were performed for Eu 2+ centers in various hosts [ 97 ] with the aim of distinguishing prospective scintillator hosts for Eu 2+ doping, such as the pro-posed CsBa 2 Br 5 . This approach was designed as a systematic high-throughput method to aid in the discovery of new bright scintillating materials by prioritization and down-selection of the large number of potential new materials. In the case of the Ce 3+ center, the information about the 5d level positioning in a number of halide compounds was summarized and used to calculate the centroid shift and the crystal-fi eld splitting of the 5d confi guration. [ 98 ] The observed trends provide insights into the relationship between the spectroscopic properties of Ce 3+ 5d levels and the crystalline environment.

For the case of defects in halides, we mention the detailed fi rst principle calculations of electronic structure, lattice relax-ation, and formation energies of iodine vacancy defects in SrI 2 for the one-electron, two-electron, and ionized charge states. [ 99 ] They predict certain characteristics, such as absorption transi-tions, confi guration coordinate curves, vibrational lineshape, and thermal trap depth of the luminescence centers. The above-mentioned study of Sr 2+ -doped LaBr 3 [ 73 ] proposes the Sr La –V Br shallow electron trap with its suggested specifi c role in the fi rst picoseconds of the scintillation mechanism.

Probably the most intense effort in theoretical calculations and modelling in the last decade in the fi eld of inorganic scin-tillators is devoted to the study of the nonproportionality of scintillators: its causes, and ways to diminish this unwanted feature degrading the energy resolution. Given the fact that the total energy deposit of each entering high-energy photon or particle (of the same energy) is defi ned by a unique sequence

of elementary interactions with the scintillator material, the resulting number of generated electron–hole pairs is practically never exactly the same, nor is the amount of generated scintil-lation photons; this is the reason for the energy resolution deg-radation. For the history, defi nition of the problems, and basic description, see reference [ 100 ] . Studies and models focus on the electron and hole thermalization processes within the conduc-tion and valence bands, respectively, [ 4,101–103 ] and their related distribution in space and time. The interaction among closely spaced elementary excitations and free carriers in the fi rst pico-seconds of the scintillation process has been recognized as the dominant mechanism of nonradiative losses (2 nd and 3 rd order nonlinear quenching processes), which are the key issue in scintillator nonproportionality. Qualitative effects of high or low values of longitudinal optical phonon frequency ω LO , thermal-ized effective diffusion coeffi cient D eff , and hot-electron velocity v g on host-dependent trends of LY and the shape of the electron energy response are illustrated schematically in Figure 6 . The three material parameters mentioned determine branching at successive decision points. LY and electron energy response curves are represented along the bottom of this fi gure as results of branching on values of the physical parameters. [ 4 ]

3. Oxide Scintillators

The fi rst single-crystal oxide scintillator, CdWO 4 , was reported at the beginning of the 1950s, [ 17 ] at around the same time as the CsI:Tl halide mentioned above. This material, as well as all the material systems mentioned below, are prepared by the Czochralski method, the principle of which is sketched in Figure 7 . The Czochralski method is one of the very few melt-growth techniques for single crystals frequently used in industry due to a favourable combination of quality, dimen-sions, and cost of the crystals produced. This method is one of the oldest and most developed. An adequate understanding of the physical phenomena observed during the solidifi cation process has been achieved which, in turn, enabled its practical expansion into industrial, large-scale production. It allows the controllable formation of single-crystalline cylindrical ingots of various inorganic scintillation materials A review has recently been published summarising the results of Czochralski growth of a number of scintillation materials. [ 104 ]

Bulk single crystals of all the material systems described below in this section are grown from the melt using the Czochralski method. Due to their high melting temperature (above 1700 °C), an iridium crucible must be used. The cru-cible is heated by an inductive radio-frequency (RF) system. The crystal diameter is controlled using an automatic diameter control system that is operated using a signal received from the weight sensor. The typical growth atmosphere is Ar or N 2 , in order to prevent possible oxidation of the iridium crucible. Only in the case of gallium-containing garnets is a few per-cent of O 2 admixed in order to prevent the decomposition of β-Ga 2 O 3 . A typical growth rate is around 0.1–1 mm per hour.

In the melt growth of crystals using the Czochralski or Bridgman methods, the dopant is distributed inhomogene-ously along the crystal growth axis due to a segregation phe-nomenon. Generally, different concentrations of the dopant




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occur in the melt and crystal grown, and their ratio changes with increasing crystal growth time as the volume of the melt in the crucible diminishes. For instance, in the case of a Ce dopant in a Gd 3 Ga 3 Al 2 O 12 (GGAG) host, the segregation coeffi cient (K(Ce 3+ ) = 0.36) is considerably greater than that in the Y 3 Al 5 O 12 (YAG) host (K(Ce 3+ ) = 0.082). [ 105 ] Also, the Ce content in the YAG crystals produced by the Czochralski method was 8.4 times less compared to the Ce-content in the melt. [ 106 ] This corresponds to a K(Ce 3+ ) of about 0.1. These observations demonstrate that GGAG crystals accept the Ce dopant much more easily than those of YAG. That is a very reasonable conclusion considering the large size of the Ce 3+ cation and the greater lattice parameter of GGAG compared to that of YAG.

All the material systems described below in this section belong to the class of fast scintillators where the radiative tran-sition is based on the 5d–4f emission of a Ce 3+ or Pr 3+ ion. The 5d–4f transition in these ions is completely allowed, and typical photoluminescence lifetimes are within 20–60 ns and 8–20 ns, respectively. Consequently, the scintillation response is domi-nated by similar decay-time values. Due to carrier trapping in the transfer stage, slower components are present as well.

3.1. Garnet Compounds

Single crystals of Y 3 Al 5 O 12 (YAG) had already been grown in the 1960s [ 107 ] and the 5d–4f photoluminescence decay kinetics of the Ce 3+ and Pr 3+ centers in a single crystal YAG host [ 108 ]

was reported soon after revealing the absence of nonradiative thermal quenching up to about 550 K and 250 K, respectively. The potential of Ce 3+ -doped YAG single crystals for fast scin-tillators was reported several years later. [ 109 ] The fi rst compre-hensive description of YAG:Ce scintillator characteristics was reported by Moszynski et al., [ 110 ] who included this material among the high fi gure-of-merit oxide scintillators. Isostructural Lu 3 Al 5 O 12 (LuAG) has a higher density and effective atomic number Z eff (6.67 g cm −3 , Z eff = 63) than YAG (4.56 g cm −3 , Z eff = 32.0), which is critically important in the case of hard X- and γ-ray detection. To obtain a fast scintillator, the Ce and Pr-doped LuAG grown by the Czochralski method became of systematic interest in the years 2000 [ 111 ] and 2006, [ 112 ] respectively. Recently, high temperatures for the onset of photoluminescence thermal quenching for Ce 3+ [ 113 ] and Pr 3+ [ 114 ] centers in LuAG host were reported, of about 790 K and 680 K, respectively. The onset tem-perature is defi ned as the one at which the decay time obtained from the single exponential approximation has dropped to half the value at the plateau before the initial drop occurs. So high a resistence of an emission center to thermal quenching points to the possible use of the material in high-temperature appli-cations. Early in the research of these modern, highly effi cient scintillators, the problem of trapping electrons in shallow traps in the transfer stage was recognized. [ 115 ] Intense slow compo-nents in the scintillation response due to the delayed radiative recombination at emission centers were reported. [ 116,117 ] The R&D of these garnet scintillators has recently been reviewed [ 118 ] and, as a result of a decade of intense research, a new ultra-effi -cient single-crystal family of so-called multicomponent garnets



Figure 6. The suggested schematic hierarchy of parameter dependence is that ω LO determines electron thermalization time relative to the time of nonlinear quenching and, consequently, the order of quenching kinetics. This also determines if thermalized or nonthermalized diffusion coeffi cients are applicable. At the second level, the value of thermalized D eff in oxides, semiconductors, and fl uorides determines electron response roll-off, while for the heavier halides, the value of hot-electron v g ( D hot ≈ v g 1/2 ) scales k 1 for defect trapping and thus determines the hump size and LY. Reproduced with permission. [ 7 ] Copyright 2012, Wiley-VCH.

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doped with cerium has been discovered [ 119,120 ] (see Figure 8 ). The balanced admixture of Gd and Ga cations in aluminium garnet effi ciently decreases the trapping effects mentioned and prevents ionization-induced quenching of the Ce 3+ excited state around room temperature. As a result, the light yield of these materials was increased almost three times compared to LuAG:Ce. [ 121,122 ] Currently, these are the most effi cient bulk single-crystal oxide scintillators. In the latest optimized material compositions, their LY approaches 60 000 photons MeV −1 , [ 123 ] which is reported as the fundamental limit of these garnet

scintillators. [ 124 ] For further details, see the review paper mentioned. [ 118 ]

The decreased bottom edge of the conduc-tion band in these compounds is mainly due to a gallium admixture. [ 125,126 ] However, the onset of thermal quenching is considerably lowered, which limits their usage to room temperature applications. This problem has been studied in detail and the ionization of the Ce 3+ excited state was determined to be its main cause. [ 127–130 ] Interestingly, another strategy preserving the high-temperature sta-bility of cerium emission centers has recently been formulated to approach the problem of electron trapping in the transfer stage of the scintillation mechanism in garnet scintilla-tors. This has an evident positive impact on the light yield, speed of scintillation response, and afterglow as well. The modifi cation of the chemical composition of the garnet scintil-lators described above results in the immer-sion of shallow electron traps in the bottom edge of the conduction band. This causes diminished charge trapping. An alternative strategy consists of the creation of an addi-tional fast radiative recombination pathway, which would effi ciently compete in electron trapping from the conduction band with the shallow electron traps in the YAG and LuAG hosts that were mentioned. Such a pathway is realized by the stabilization of a tetravalent Ce 4+ center in the garnet lattice by divalent

rare earth ion co-doping [ 113,131–134 ] and/or by air annealing. [ 135 ] The positive role of Ce 4+ in the scintillation mechanism in orthosilicates has recently been reported in the literature (see Section 3.3 below) and the same mechanism apparently also functions in garnets.

In Figure 9 step 1, in the fi rst picoseconds of the scintillation mechanism, the Ce 4+ center can effi ciently compete with any electron traps for an immediate capture of electrons from the conduction band. The stable Ce 3+ center is much less effective in this competition, since it fi rst needs to capture the hole from the valence band in step 1.

In step 2, the Ce 4+ transforms into an excited Ce 3+ center and emits the desired scintillation photon, i.e., it contributes to the fastest part of scintillation response. In the same step the Ce 3+ center, temporarily converted into a Ce 4+ center, an electron is captured from the conduction band and becomes excited.

In step 3, the return to the initial state (beginning of the cycle) is accomplished by hole capture from the valence band (Ce 4+ in the right part) and by emission of the scintillation photon (Ce 3+ in the left part). It is worth mentioning that, in the last step of the Ce 4+ scintillation mechanism (right part), the hole capture from the valence band must always be nonradia-tive, i.e., not contributing to afterglow.

Very recently, similar studies were done for the Gd 3 Al 2 Ga 3 O 12 :Ce (GAGG:Ce) multicomponent garnet [ 133,136,137 ] mentioned above. Given its high light yield, Me 2+ (Me = Ca, Mg) co-doping does not further increase the light yield value,



Figure 7. Phases of a typical Czochralski process including A) approach of the seed to the overheated melt surface, B) immersion of the seed in the melt and their thermal equilibration, C) pulling of the seed in the upward direction with a continuous increase of crystal diameter and shoulder formation, D) steady state, as the constant diameter crystal is pulled, E) ending the growth with continuous diameter decrease, and F) separation of the crystal from the melt and its subsequent cooling to room temperature. Reproduced with permission. [ 104 ] Published by the Polish Physical Society.

Figure 8. Photo of a Gd 3 Ga 3 Al 2 O 12 :Ce single crystal, diameter 5 cm, length 15 cm. Small optical elements cut from the crystal are shown in the front. Reproduced with permission. [ 245 ] Copyright 2014, the Furukawa Co.


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but the scintillation response apparently becomes faster, as seen in Figure 10 , which might be of importance for applica-tions in PET medical imaging with a time-of-fl ight option. [ 26 ] It is interesting to note that Ca 2+ co-doping decreases the light yield much more than Mg 2+ co-doping. The reason for that effect has not yet been explained.

Another possibility for increasing the light yield of alu-minium garnet scintillators was the creation of a solid solution of YAG and LuAG hosts. It was shown that, with an increase of temperature up to about 150 °C, the light yield in LuAG:Pr or LuAlO 3 :Ce (LuAP:Ce) increases by a factor of up to 1.8, while no such increase is obtained for YAG:Ce or YAP:Ce. [ 138,139 ] This is consistent with the fact that thermoluminescence characteristics show increasing trap depth when lutetium is admixed with YAG or YAP hosts in the case of cerium doping and similar production technologies. [ 116,140–143 ] At the same time, the amount of slow-decay components in the scintilla-

tion response strongly increases. [ 117,144 ] This might be related to the fact that the bandgap in LuAG is somewhat larger than that in YAG, which is apparent in their excitation spectra of intrinsic emission. [ 145 ] The same tendency might also be sup-posed to occur in a perovskite compound considering the reported refl ection spectra in the VUV spectral region. [ 146,147 ] In Ce-doped and Y-admixed LuAG, an increase in the light yield of about 10% compared to LuAG:Ce was reported. [ 148 ] The recent systematic study of light yield and other characteristics in Pr-doped (Lu x Y 1-x ) 3 Al 5 O 12 showed an enormous increase of light yield up, to 33 000 photons MeV −1 for a 1 mm-thick sample of (Lu 0.75 Y 0.25 ) 3 Al 5 O 12 :Pr. This increase is ascribed to the effect of trap depth tuning, so as to maximize the light yield just around room temperature. [ 149 ] An additional positive effect that might contribute to the light yield increase would be due to the solid solution host, an effect that has been observed in several mate-rial systems. [ 150 ] Finally, due to the fact that all these Ce- or Pr-doped garnet crystals are grown in an atmosphere which con-tains little (1–2%) or no oxygen, the post-growth annealing in air at elevated temperatures can also contribute to a light yield increase due to the reduction of the oxygen-vacancy concentra-tion and related deep electron traps. [ 135,151 ]

Despite the fact that this review is devoted to single-crystal materials, the following paragraph is devoted to garnet-based scintillation ceramics, as they have appeared as serious compet-itors to the single crystals of the garnet scintillators described above. Recently, a review paper devoted to scintillating optical ceramics was published, covering all of the most important material fi elds and related R&D activities. [ 152 ] The technology of garnet optical ceramics was developed in 1990s and was triggered by solid-state laser applications where Nd-doped YAG ceramics started to be used instead of single crystals. [ 153 ] Ceramic Ce 3+ -doped YAG for fast scintillator applications was reported for the fi rst time in the late 1990s. [ 154,155 ] In those studies, the materials were not transparent but translucent, so that the scintillation light output achieved was about 50% of that of the single-crystal counterpart. In the early 2000s, a transparent ceramic YAG:Ce scintillator was developed by the Konoshima Chemicals Co. with a light output exceeding that of a single crystal. [ 156 ] Due to the reasons mentioned above for single crystals, further R&D focused on higher density and Z eff materials, i.e., LuAG-based ceramics. Over 2005–2008, transparent ceramic Ce-doped LuAG was reported by several groups. [ 157–161 ] Similar to the ceramic YAG:Ce, the emission at 300–350 nm due to antisite defects was diminished in ceramic LuAG:Ce. Moreover, the shallow electron trap concentration responsible for the dominant thermoluminescence (TSL) peaks in the range 140–190 K in single crystals [ 116 ] became compara-tively much smaller in ceramics. Lower preparation tempera-tures and the absence of the melting of raw materials in the manufacturing process are the main causes of the absence of antisite defects in the ceramic structure. Those sites are respon-sible for the UV emission and TSL traps mentioned above. However, the absolute light-yield value of the ceramic LuAG:Ce was noticeably inferior to that of the single crystal. [ 157,160 ] In fact, TSL characteristics between 200 and 300 K and the enhanced intensity of the very slow components in scintilla-tion decay in ceramic LuAG:Ce point to the occurence of deep traps situated most probably at the surface and interface of the



Figure 10. Scintillation decay curves of the Mg co-doped Ce:GAGG crys-tals. [ 137 ] Excitation by the 137 Cs radioisotope (662 keV). Curves are verti-cally shifted for clarity.

Figure 9. Sketch of the scintillation mechanism at the stable Ce 3+ (left) and Ce 4+ (right) emission centers in an aluminium garnet host. Repro-duced with permission. [ 113 ] Copyright 2014, American Chemical Society.


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grains in ceramics, [ 160 ] which are responsible for light-yield deterioration. After these pio-neering studies, the recently synthesized Ce 0.5%-doped transparent ceramic LuAG:Ce by the Konoshima Chemicals Co. did show a superior light yield compared to that of the Czochralski-grown, high-quality, single-crystal LuAG:Ce. [ 162 ] Furthermore, the Mg 2+ -co-doping mentioned above was also suc-cessfully applied to LuAG:Ce ceramics and the light yield achieved exceeded that of a high-quality single crystal. [ 132 ] The fi rst report of a transparent ceramic LuAG:Pr was published in 2009, [ 163 ] and its light yield was approximately half that of its single-crystal counterpart, similar to the other group's results. [ 164,165 ] However, recently, the Pr 0.25 mol%-doped ceramic LuAG prepared by the Konoshima Chemicals Co. exceeded the light-yield value of a high-quality single crystal by about 20%. [ 166 ] In this ceramic sample, the antisite defect-related emission in the near UV completely disappeared in the radioluminescence spectrum, similar to the YAG:Ce and LuAG:Ce ceramics discussed above.

In the family of multicomponent garnet materials, analo-gous R&D activities concerning Gd and Ga substitution in Ce:(Gd,Y) 3 (Al,Ga) 5 O 12 (Ce:GYGAG), as described above for single crystals, were also studied in optical ceramics. [ 167 ] Trans-parent ceramic GYGAG:Ce scintillators with high scintillation performance, that is, with a light yield of 50 000 ph MeV −1 and an energy resolution of 4.5–4.8% at 662 keV, were reported. [ 168 ] Similar to single crystals of this composition, the GYGAG:Ce ceramics show a dominant decay time of about 250 ns, consid-erably longer that the Ce 3+ related one (50–60 ns [ 119,120 ] meas-ured in photoluminescence decay. This phenomenon was inter-preted as being due to energy migration over the Gd 3+ sublat-tice, followed by an energy transfer to Ce 3+ . [ 168 ] An interesting observation was also reported for Gd 3 (Al,Ga) 5 O 12 :Ce regarding the relation between the afterglow and stoichiometry of the host. [ 169 ] In particular, the afterglow substantially diminished for a small excess of Gd cation with respect to a stoichiometric composition. Formation of isoelectronic traps was discussed based on the stoichiometry of the material and the difference in the electronegativity of the dopant and host ions.

Apparently, transparent ceramic and single-crystal multicomponent garnets show similar luminescence and scintillation characteristics, and are very promising for X- and gamma-ray detectors due to their high scintillation effi ciency and, especially, for their light-yield values. However, further investigations are required to reveal the details of the scintilla-tion mechanism, the underlying energy-transfer processes, and to further optimize these complex material systems.

3.2. Aluminium Perovskite Compounds

Fast 5d–4f luminescence of Ce 3+ and Pr 3+ in YAlO 3 (YAP) have been reported by Weber [ 170 ] and Gumanskaya et al., [ 171 ] respec-tively. The reported lifetimes of about 18 ns and 8 ns, respec-tively, are among the shortest within the Ce- and Pr-doped oxide-based scintillators; a feature that makes these materials attractive. Favourable properties of YAP:Ce for scintillation

applications were described later by Takeda et al. [ 172 ] and Autrata et al. [ 173 ] In the mid 1990s, similar to what was men-tioned in Section 3.1 for garnets, several laboratories replaced yttrium with lutetium in a YAP crystal to increase the density and Z eff from 5.35 and 33.5 g cm −3 , respectively, up to 8.34 and 64.9 g cm −3 for LuAP. Due to a severe instability of the LuAP perovskite phase in the process of crystal growth from the melt (see recent studies [ 174,175 ] for further details), resulting in a high price for the crystals, mixed (Lu,Y)AlO 3 :Ce scintilla-tors were fi nally chosen for industrial-scale production. [ 176 ] These LuYAP:Ce crystals were used in the ClearPET proto-type, [ 177 ] but did not fi nd a further commercial market due to the price and relatively lower light yield compared to YAP:Ce. The reasons for the light-yield decrease in Lu-rich aluminium perovskite seem to be similar to what was seen in garnets (see Section 3.1). [ 138,141,142,144–147,178 ] However, despite several attempts, e.g., by co-doping, [ 179–181 ] no successful strategy has so far been found to substantially improve it. In addition, no successful bandgap engineering of YAP by compositional varia-tion (similar to what was done in garnets [ 119,120 ] has been found either. [ 182,183 ] R&D results for the aluminium perovskite family have also been the subject of several reviews. [ 2,19,31 ]

Given the speed of the scintillation response dominated by the 8 ns decay time, [ 184 ] additional effort was devoted to Pr-doped YAP to understand its 2–3-fold lower light yield [ 185 ] compared to Ce-doped YAP. The scintillation performance of a large set of Pr-doped YAP single crystals prepared by three dif-ferent technologies, including that of Czochralski ( Figure 11 ), was evaluated by radioluminescence spectra ( Figure 12 ), photo-electron yield, and scintillation decay measurements. [ 186 ] The intrinsic scintillation effi ciency (the integral of the radiolumi-nescence spectrum) of the best Czochralski-grown YAP:Pr reached about 150% of that of a YAP:Ce standard sample. However, the photoelectron yield of the former was only about 80% of that of the latter, suggesting the presence of enhanced delayed radiative recombination processes in YAP:Pr. The dis-crepancy between scintillation effi ciency and photoelectron yield was systematically observed in all Pr-doped samples.Thermoluminescence glow curves in the 10–350 K tempera-ture range show that, in YAP:Pr, the dominant glow curve peak occurs at a noticeably higher temperature (192 K) with respect to that found in YAP:Ce (115 K). [ 187 ] The calculated RT lifetime of the 192 K peak-related trap in YAP:Pr is about one order of magnitude longer (1 ms) than that related to the 115 K peak in YAP:Ce. It can be one of the reasons for the decrease of photo-electron yield in YAP:Pr, due to an enhanced delay in the trans-port stage of the scintillation mechanism. However, since the existence of the delayed radiative recombination processes in YAP:Pr mentioned earlier is due to a defect instead of being an



Figure 11. Photo of Czochralski-grown single crystals of YAP:Pr0.5% (left) and Lu 0.1 Y 0.9 AP:Pr1% (right).


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intrinsic property of this material, the possibility of further opti-mizing the scintillation performance of YAP:Pr remains open.

It is also worth mentioning the temperature stability of the Pr 3+ center in YAP, with the onset of thermal quenching above 600 K. Knowledge of this property is derived from photolumi-nescence decay time measurements (see Figure 13 ). Thermal quenching of the emission center itself is a dominant factor in the onset of quenching of the scintillation effi ciency and light yield. [ 188,189 ] Applying the same criterion for the temperature of thermal quenching, as in the case of Pr-doped LuAG in Sec-tion 3.1, we obtain a temperature of 690 K, which is slightly higher than that in a LuAG host (680 K). This indicates a com-parable or slightly higher thermal stability of Pr- and Ce-doped

YAP compared to a LuAG host. Consequently, Ce- and Pr-doped Lu-admixed aluminium perovskites should show even higher thermal stability compared to a simple YAP host. [ 114 ]

Finally, an extended correlated electron paramagnetic reso-nance (EPR) and TSL study of undoped and several rare earth-doped YAP single crystals gave deep insight into charge-car-rier trapping in this material family and its infl uence on the transfer stage of the scintillation mechanism. [ 187,190 ] In addition to two already known O − hole centers, [ 191 ] four other O − centers were identifi ed; these arose from holes being captured from the valence band by oxygen anions. The centers differ in their thermal stability, which is characterized by the thermal activa-tion energy E a starting from 0.024 eV (most probably a self-trapped hole) up to more than 0.5 eV. The latter, the most stable center, can survive even at room temperature for a few days. The holes are most probably stabilized by an impurity ion at the Y site or a cation vacancy. Electronic-type trapping sites are assigned to the yttrium antisite ions, which become paramag-netic Y Al 2+ centers after trapping an electron. They are found in four structurally different confi gurations with a thermal sta-bility around or above 300 K. This enables radiative recombina-tion of thermally liberated holes with electrons that have been localized in this way. In two of the centers, the trapped electron is additionally stabilized by an oxygen vacancy. Yttrium antisite positions in a YAP lattice were directly identifi ed by 89 Y nuclear magnetic-resonance measurements. [ 190 ]

3.3. Ortho and Pyrosilicates

Scintillation characteristics of the Ce 3+ -doped rare earth oxy-orthosilicate Gd 2 SiO 5 :Ce (GSO:Ce) were reported for the fi rst time in 1983 [ 192 ] and those of Lu 2 SiO 5 :Ce (LSO:Ce) were intro-duced in 1992. [ 193 ] Later, they became well known and commer-cially successful single-crystal scintillators due to a favourable combination of high density, effective atomic number, and fast scintillation response, dominated by a decay time of sev-eral tens of nanoseconds and with the further advantages of mechanical and chemical stability. Their fundamental optical and luminescence characteristics, including Y 2 SiO 5 :Ce (YSO:Ce), were reported by Suzuki et al. [ 194 ] This study revealed two Ce 3+ emission centres, Ce1 and Ce2, embedded in two sites of the RE 3+ cation in the orthosilicate structure ( Figure 14 ), and showed emission bands at about 400–440 nm and 500 nm, respectively. LSO:Ce crystals, however, identical to the LuAG or LuAP-based crystals described above, exhibit an intrinsic back-ground signal of a few hundred Hz cm −3 due to the presence of the radioactive isotope 176 Lu, while GSO:Ce or YSO:Ce does not have this disadvantage. Consequently, GSO:Ce can be used in low-signal-count-rate applications such as hard X(gamma)-ray astronomy. [ 195 ] Thanks to the greater temperature stability of the Ce1 center in the GSO host, it is widely used in the oil-well industry and geophysical explorations up to at least 150 °C. This is despite the fact that its light yield at room temperature is less than half of that of LSO:Ce due to thermal quenching of the Ce2 site above 200 K. [ 194,196 ] On the other hand, optimized LSO:Ce and especially yttrium-admixed LYSO:Ce (introduced in year 2000 [ 197 ] show a light yield exceeding 30 000 phot MeV −1 and are used in the latest generation of scintillation detectors



Figure 13. Temperature dependence of the photoluminescence decay time of Pr 3+ 5d–4f emission in a YAP host, obtained from a single expo-nential approximation of the decay (symbols). Dashed line is to guide the eye.

Figure 12. Radioluminescence spectra of Ce- and Pr-doped YAP at RT. They are compared on an absolute scale. The doped ion, its concentra-tion, and the kind of crucible used (made from Ir or Mo) is in the legend. Excitation by X-ray, 40 kV. Reproduced with permission. [ 186 ] Copyright 2010, IEEE.


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in PET imaging. [ 20 ] Large crystals up to 8 cm in diameter and 20 cm in length have been grown (see Figure 15 ). [ 198 ] Due to the early excited state ionization of both Ce centers, [ 199,200 ] however, they cannot be used in applications above room temperature. The Y-admixture, among others, increases the temperature of the onset of this ionization process for both Ce centers, which positively infl uences the afterglow and daylight sensitivity of this Ce-doped orthosilicate scintillator. [ 201 ] Structural, optical, luminescence, and scintillation characteristics of these ortho-silicate materials have been previously reviewed. [ 31 ]

A solid solution of Ce-doped LSO and GSO hosts, so called LGSO:Ce, was discovered to be an effi cient, dense scintillator with high light yield, fast decay, and weak afterglow. [ 202–204 ] In a

systematic compositional study, the phenomenon of light-yield improvement for intermediate mixed crystal compositions was demonstrated. [ 205 ] This is analogous to the solid solutions of other materials mentioned above. [ 150 ] LGSO may possess mono-clinic P21/c or C2/c structures depending on the Lu/Gd ratio in the host. The comparative advantage of LGSO:Ce in compar-ison with LSO:Ce or GSO:Ce in terms of the overall fi gure-of-merit is most evident around room temperature. [ 196 ]

Recent studies of LSO:Ce and YSO:Ce scintillators have shown that Ca 2+ co-doping plays a positive role in their per-formance, [ 206–208 ] but a detailed explanation of the underlying physical mechanism was not proposed. In 2013, a detailed study of the effect of Me 2+ co-doping in LYSO:Ce was published in which optical and photoelectron (XANES) spectroscopy tech-niques were combined to provide evidence for the presence and explain the role of a stable Ce 4+ center in the scintillation mechanism. [ 209 ] In fact, the charge transfer (CT) absorption characteristics of the Ce 4+ center in LSO in the UV spectral region were already revealed a long time ago in LSO:Ce. [ 210 ] They were also found and discussed regarding Ce 4+ participa-tion in the scintillation mechanism in Ce-doped silica glass. [ 211 ] The absorption fi ngerprint of a stable Ce 4+ center provides a very sensitive tool to reveal its presence in oxide hosts. Me 2+ co-doping and annealing in air [ 212,213 ] stabilize the Ce 4+ center, which positively infl uences several scintillation characteris-tics: light yield, speed of scintillation response, and afterglow. This center forms a new fast radiative recombination pathway based on immediate electron capture from the conduction band, radiative de-excitation of the excited Ce 3+ center, and hole capture to return to the Ce 4+ stable initial state. This has been sketched and described for garnet scintillators in Section 3.1 and in Figure 9 . Such a mechanism works in parallel with the standard one based on a stable Ce 3+ center. Consequently, these two centers do not compete with each other. However,their rela-tive ratio must be properly tuned for material optimization. [ 113 ] In fact, very similar values of the onset of this CT absorption process occur in the orthosilicates and in the aluminium garnet structures mentioned above (see Figure 16 ). This provides strong support for such an interpretation taking into account the nature of charge transfer transitions. [ 32 ]

More than a decade ago lutetium pyrosilicate Lu 2 Si 2 O 7 (LPS) was also found to be a potentially interesting scintillator host. [ 214 ] A comparative EPR study of the Ce 3+ -doped LSO and LPS showed that the Ce ion in an LPS structure substitutes for Lu in its single crystallographic site, while in the structure of LSO it is found in both Lu crystallographic sites. [ 215 ] The light yield of LPS:Ce single crystals, which were grown from the melt, can reach a value comparable to that of LSO:Ce. The scintilla-tion decay time of the Ce 3+ is around 37 ns with no observable afterglow. [ 214,216 ] Furthermore, as for LSO:Ce [ 212,213 ] post-growth annealing in air at elevated temperatures was found to be effec-tive for increasing the scintillation effi ciency. [ 217 ] The lack of afterglow in LPS:Ce, in contrast to its observation in LSO:Ce, was correlated with the signifi cantly higher temperature maxima of the TSL glow peaks above room temperature. [ 218 ]

More recently, Gd 2 Si 2 O 7 :Ce (GPS:Ce) pyrosilicate was intro-duced, showing much higher light output and a faster scintilla-tion response than GSO:Ce. [ 219 ] GPS:Ce shows an incongruent growth from the melt [ 220 ] which disables its preparation by



Figure 14. The structure of Lu 2 SiO 5 (LSO), displaying two Lu sites and surrounding SiO 4 tetrahedra. Three oxygen ions with cut bonds form OLu 4 tetrahedra and do not participate in the Si–O bonds.

Figure 15. Crystal of LYSO:Ce, diameter 72 mm, length 155 mm, grown in the Shanghai Institute of Ceramics CAS, China. Image courtesy of D. Ding.


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the Czochralski method. However, heavy Ce-doping (at least 10 mol%) in a GPS host does enable its congruent growth. [ 221 ] Unfortunately, at such a high Ce concentration, the light output is signifi cantly reduced because of self-absorption and concen-tration quenching. An optimal cerium concentration in oxide hosts is usually in the range 0.1–1 at%. Congruent crystal growth of GPS:Ce is achieved by expansion of the average ionic radius in the Gd site, resulting from Ce-doping. At the same coordination number, the La 3+ ion has a very similar ionic radius to that of Ce 3+ , so that the substitution of La for Ce can also be applied to stabilize the pyrosilicate phase avoiding unwanted concentration quenching. The optical and scintilla-tion properties of (Ce 0.01 , Gd 0.90 , La 0.09 ) 2 Si 2 O 7 were reported for the fi rst time by Suzuki et al. [ 222 ] where these single crystals were grown by the fl oating zone method in an argon atmos-phere. Using a Si-avalanche photodiode detector, excellent values of light output of 41 000 ± 1000 photons MeV −1 and a full width at half maximum (FWHM) energy resolution at 662 keV of 4.4 ± 0.1% were achieved. [ 223 ] The impact of the La and Sc admixture in GPS:Ce prepared by the top-seeded solu-tion growth with SiO 2 self-fl ux was also studied regarding their structure and optical and scintillation properties. [ 224 ]

In a very recent study, [ 225 ] the absorption spectra, photolumi-nescence spectra, decays, and selected scintillation characteris-tics were studied for Ce-doped LPS, GPSLa30%, and GPSLa48% single crystals grown by the Czochralski technique (see Figure 17 ). The 4f–5d x (x = 1–5) Ce 3+ absorption bands in GPSLa30% were determined to be at 338, 320, 294, 242 and 219 nm, respectively. The 5d – 4f emission of Ce 3+ peaks were at 377 nm and 372 nm in LPS and GPSLa hosts, respectively. The onset of nanosecond decay-time shortening appears around 380 K (LPS:Ce) and 440 K (both GPSLa30%:Ce and GPSLa48%:Ce) and is due to thermally induced excited-state ionization. The Ce 3+ ionization onset, which fortunately occurrs well above RT, provides an opportunity to exploit LPS:Ce and particularly GPSLa:Ce in high-temperature applications. The evaluated scintillation effi ciency (integral of

the radioluminescence spectrum) reaches about 250%, 1210%, and 1530% of that of the BGO single crystal standard for LPS:Ce, GPSLa48%:Ce, and GPSLa30%:Ce, respectively. In the latter com-pound, the effi ciency is almost doubled with respect to that of commercial high-performance LYSO:Ce,Ca. [ 226 ] The afterglow in La-admixed gadolinium pyrosilicates is fairly low and tends to get less intense with increasing La concentration, becoming compa-rable to that of BGO. Taking into account the approximately two orders of magnitude lower intrinsic radioactivity (due to the 138 La isotope, 0.09% natural abundance, τ 1/2 ≈ 10 11 years) compared to Lu-based scintillators, the La-admixed GPS:Ce single crystals show a combination of characteristics highly favourable for med-ical imaging, the oil industry, and geophysical applications.

3.4. Theoretical Studies

For oxide scintillators, a systematic effort has also been devoted to theoretical calculations and modelling. The electronic band structure of host materials, energy-level positioning, properties, interactions of emission centers with the host, and particular defect-related energy levels in the host forbidden gap are being studied. The latter energy levels give rise to charge-carrier traps affecting the conversion and transfer stage of the scintillator mechanism. Electronic band-structure calculations have been performed for all the yttrium-based materials mentioned above, including YAG, [ 227 ] YAP, [ 228 ] and both YSO and YPS. [ 229 ] The top of the valence band is dominated by the energy levels of the oxygen anion, while the bottom of conduction band is governed by 4d orbitals of the yttrium cation. That is why in the defect studies by EPR O − hole centers are mostly found in YAP [ 190 ] and YSO. [ 230 ] In electron traps, the 4d level of the yttrium cation is frequently involved (see also the review in the literature. [ 231 ] The defect creation energy and energy-level positioning in the bandgap have been calculated, e.g., for antisite defects in alu-minium garnet and perovskites. [ 232–237 ] The astochiometry of the host [ 238 ] and aliovalent dopants [ 239 ] can strongly affect the creation of Y Al antisite defects. In silicates, the critical role of the Si-unbound oxygen site has been emphasised in defect and trap creation [ 240 ] and further refl ected in the EPR study of para-magnetic point defects in orthosilicates. [ 230 ]

Bandgap engineering has successully been applied in the technology of the garnet family for scintillator optimization,



Figure 17. Photo of a Ce-doped GPSLa30% single crystal grown by the Czochralski technique.

Figure 16. The Ce 4+ -related induced absorption in Ce-doped LYSO and LuAG single crystals. Reproduced with permission. [ 113 ] Copyright 2014, American Chemical Society.


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as already mentioned, [ 119,120 ] and has been complemented by theoretical calculations to clarify the effect on the electronic band structure of an admixture of gallium into the aluminium garnets. [ 125,126 ]

Furthermore, new insight into the 5d energy-level positioning of Ce 3+ in YAG has been obtained. [ 241 ] The Ce 3+ multisites were characterized in YAG and LuAG using fi ne infrared spectros-copy and correlated with the theoretical analysis of the 4f ground state multiplets of Ce 3+ . [ 242,243 ] The spectroscopy of the lantha-nide dopants in the RE 3 (Al 1-x Ga x ) 5 O 12 (RE = Gd, Y, Lu and x = 0, 0.2,0.4, 0.6, 0.8, 1) family of garnet compounds was reviewed, pro-viding information on the red shift, the centroid shift, the charge-transfer energies, and the host exciton creation energies. [ 244 ]

4. Conclusions

In this review paper, we have provided a selected survey of recent trends in the development of high fi gure-of-merit single-crystal scintillation materials. We have demonstrated the extended activities and success in the development of new sys-tems and optimization of existing scintillators based on halide and oxide compounds.

In the case of Ce-doped halide systems, among the new materials we focused on CeBr 3 and CeBr 3 :Sr, which are compet-itive alternatives to LaBr 3 :Ce in those cases where low intrinsic activity is critical in applications, such as radioastronomy. A successful example of material optimization was described for Sr-co-doped LaBr 3 :Ce, which constitutes a scintillator with the best energy resolution nowadays among single-crystal bulk materials: 2%@662 keV. Such a characteristic, together with an extremely high light yield exceeding 70 000 photons per MeV, makes it a very competitive candidate for both medical imaging and security-related techniques.

In the case of Eu 2+ -doped halides, the revival of SrI 2 :Eu is described and explained as being due to new applications in the security fi eld, where classical NaI:Tl cannot compete due to its considerably worse energy resolution. Because of the rela-tively high sensitivity of SrI 2 :Eu to moisture, there is an intense search for other iodide single crystals with comparable scintil-lation performance, but better tolerance to moisture. The most interesting candidate is found in the CsBa 2 I 5 :Eu single crystal.

Among the oxide-based materials, the biggest breakthrough was achieved in the discovery of multicomponent garnet scin-tillators based on the general formula (Gd,Lu,Y) 3 (Ga,Al) 5 O 12 doped with a Ce 3+ center. For selected compositions, the light yield approaches 60 000 photons MeV −1 , which is a funda-mental limit for garnet scintillators and is almost double that of LYSO:Ce,Ca. It is also the best existing example from recent R&D efforts in the scintillator fi eld to develop new materials by compositional engineering where the electronic band structure is suitably modifi ed. Also worth mentioning is the La-admixed GPS, a new competitive candidate among high-performing scin-tillators. By a compositional change, it overcomes the problem of the incongruent growth of GPS and enables the growth of larger size, high-quality, single crystals with excellent scintilla-tion performance. We also showed that, in mixed composition/solid solution hosts, the scintillation effi ciency and light yield can be enhanced relative to the limit (end) compositions.

The defect engineering strategy, on the other hand, is focused on the suppression or creation of specifi c defects which are harmful or benefi cial in the scintillation mechanism. It is most frequently realized through co-doping by aliovalent ions and/or postgrowth annealing in a tailored atmosphere. Here, the most surprising issue is defi nitely the discovery and expla-nation of the positive role of stable a Ce 4+ center in both garnets and orthosilicates. This center, in proper concentration relative to stable Ce 3+ , considerably enhances performance regarding light yield, speed of scintillation response, and afterglow. At the same time, one should note the sensitivity of materials to the nature of the co-dopant when, for example, Ca 2+ or Mg 2+ is used in garnets. On the contrary, no substantial difference in the effects achieved has been reported in orthosilicates for the same co-dopants. Thus, these concepts and strategies are highly compound-specifi c and cannot be mechanically copied from one host to another.

Experimental activities are complemented by extended theoretical studies, which provide further support and under-standing of the electronic band structure, energy-level posi-tioning of defects, and emission centers in the forbidden gap of the host. In some cases, these computational methods try to forecast suitable material compositions for further development to avoid costly and time-consuming experimental search. The latter mostly use a combinatorial approach. Success of com-putational methods for providing a guiding concept is so far rather limited.

The intensity and extent of R&D in the scintillator fi eld in the last two decades have been driven by a number of appli-cations requiring tailored materials with specifi c parameters of critical importance for a particular use. While in the 1990s, applications were mainly in high energy physics, later, med-ical imaging and fi nally security-related techniques came into play. Though less known, high-tech industry often, on a very short time scale, also calls for specifi c scintillator materials for their devices; electron microscopy and the detection of particle beams in general are two examples.

Acknowledgements This research and article preparation was supported by grants from the Czech Science foundation (no. P204/12/0805) and MEYS, Kontakt, (no. LH14266). Partial support by the JST Program, Adaptable & Seamless Technology Transfer Program through Target-driven R&D (A-STEP) and Development of Systems and Technology for Advanced Measurement and Analysis are also greatly appreciated. Thanks are due to V. Babin for providing the unpublished experimental data for Figure 13 , D. Ding for the photo in Figure 15 , E. Mihokova for language corrections and E. Jurkova for technical help in the manuscript preparation.

Received: December 2, 2014 Revised: January 24, 2015

Published online: February 25, 2015

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