Structure, Luminescence, and Application of a Robust CarbidonitrideBlue Phosphor (Al1−xSixCxN1−x:Eu

2+) for Near UV-LED Driven SolidState LightingLe Wang,† Xiaojun Wang,‡ Takashi Takeda,‡ Naoto Hirosaki,‡ Yi-Ting Tsai,§ Ru-Shi Liu,§,∥

and Rong-Jun Xie*,‡,⊥

†College of Optical and Electronic Technology, China Jiliang University, Hangzhou, Zhejiang 310018, China‡Sialon Group, Sialon Unit, Environment and Energy Materials Division, National Institute for Materials Science (NIMS), Namiki1-1, Tsukuba, Ibaraki 305-0044, Japan§Department of Chemistry, National Taiwan University, Taipei 106, Taiwan∥Department of Mechanical Engineering and Graduate Institute of Manufacturing Technology, National Taipei University ofTechnology, Taipei 106, Taiwan⊥College of Materials, Xiamen University, No. 422 Siming-nan Road, Xiamen, Fujian 361005, China

*S Supporting Information

ABSTRACT: As an extension of nitride luminescent materials,carbidonitride phosphors are also attracting great attention due totheir superior thermal stability. This paper reports a blue-emittingcarbidonitride phosphor Al1−xSixCxN1−x:Eu

2+ suitable for nearultraviolet (UV) light emitting diodes (LEDs), which is formulatedby introducing SiC into AlN:Eu2+. With the introduction of carbon(silicon), the lattice abnormally shrinks along both a- and c-axes atlow x values (x ≤ 0.08), due to the formation of a dense interlayerfor accommodating the luminescence center Eu2+. Both of theRaman spectra and solid state NMR spectroscopy show that bothSi and C are dissolved in the AlN lattice. A single blue emissionband (λem = 472−477 nm) is observed for compositions of x >0.05 by cathodoluminescence measurements. Under the 365 nmexcitation, the maximum luminescence is attained for the composition of x = 0.06 that has an external quantum efficiency of 61%and absorption efficiency of 74.4%, which is about 11−15% higher than the corresponding carbon-free nitride sample. Thethermal quenching of Al1−xSixCxN1−x:Eu

2+ reduces with increasing C (SiC) content, and the sample of x = 0.06 shows a smallloss of ∼4.0% in quantum efficiency even at 200 °C. Using this phosphor in a near UV-driven white LED, a superhigh colorrendering index of Ra = 95.3 and R9 = 72 as well as a color temperature of 3533 K are achieved.

I. INTRODUCTION

GaN-based solid state lighting or white light-emitting diodes(wLEDs) are described as the fourth-generation lighting sourcefollowing gas lamps, incandescent bulbs, and fluorescent lights,which promise high efficiency, energy saving, quick response,long operating life, compactness, and low negative environ-mental impact.1−3 In solid-state lighting, phosphors play keyroles in converting the light from GaN-based chips intoappropriate visible emissions, and thus determine the luminousefficiency, color rendering properties, color temperature, andlifetime of wLED devices.4 Therefore, phosphors with superiorluminescent properties are continuously pursued to achievehigh color rendition, efficient, and reliabile wLEDs. Althoughblue-LED driven wLEDs, i.e., blue LED + YAG:Ce3+ or blueLED + green/red phosphors, have high efficiency and are moresafe, wLEDs pumped by near UV-LEDs (i.e., near UV-LED +blue/green/red phosphors) are superior in color rendering

properties and light uniformity. For example, Lee et al. reporteda color rendering index of 94 by combining blue BaMgA-l10O17:Eu

2+, green (Ba,Sr)2SiO4:Eu, and orange BaLa2Si2S8:Euwith a near UV-LED chip (405 nm).5 A higher color renderingindex of Ra = 96 was achieved by using blue KMg4(PO4)3:Eu,green (Ba,Sr)2SiO4:Eu, and red CaAlSiN3:Eu.

6 Daicho et al.reported a glareless white LED with a color rendering index of87, using a violet chip and blue (Ca,Sr)5(PO4)3Cl:Eu, green β-sialon:Eu, and red CaAlSiN3:Eu.

7

As near UV light is shorter in wavelength or stronger inenergy than the blue light, phosphors for near UV-LED chipsare thus required to be more thermally stable and robust. Tomeet this end, the host crystal of phosphors should contain the

Received: November 10, 2015Revised: November 24, 2015Published: November 25, 2015

Article

pubs.acs.org/cm

© 2015 American Chemical Society 8457 DOI: 10.1021/acs.chemmater.5b04384Chem. Mater. 2015, 27, 8457−8466

high hardness/rigidity building blocks, such as [SiN4], [SiC4],or [Si(C, N)4].

8 For instance, nitride phosphors, i.e., α-sialon:Eu, β-sialon:Eu, and CaAlSiN3:Eu, usually have smallthermal quenching or high thermal stability.9,10 By accom-modating carbon into the [SiN4] or [AlN4] tetrahedron, anovel family of luminescent materials named carbidonitridephosphors is then produced. In carbidonitride tetrahedral units,the center Si or Al atoms are bonded to C and/or N atoms atfour corners. Several carbidonitride luminescent materials havebeen reported as LED phosphors, including Y2Si4N6C:Ce

3+,Sr1−yY1+ySi4N7−yCy:Ce

3+, Sr2Si5N8−[(4x/3)+z]CxO3z/2:Eu2+,

Sr7Al12−x−ySix+yNx−yCy:Eu2+, Y2(NC2)3:Ce

3+, and α-SrNCN:Eu2+.8,11−14 Tian addressed that carbidonitride phos-phors were more thermally stable than their carbon-freecounterparts.8

AlN:Eu,Si was reported to an interesting blue-emittingphosphor for both field emission display (FED) and wLEDapplications, due to its high color purity, high efficiency, andhigh thermal stability.15,16 Both Si3N4 and SiC can be used asthe Si source, but great differences in crystal chemistry andphysical properties between Si3N4 and AlN make the solubilityof Si3N4 in AlN to be quite limited.17 This indicates that a smallvariation in the Si3N4 amount would yield great changes inluminescence efficiency. On the other hand, a single solidsolution between AlN and SiC (2H) can be formed in a widerange due to their equivalent crystal structure.18,19 This optiontherefore not only solves the problem arising from the use ofSi3N4, but also develop a novel carbidonitride phosphor, i.e.,Al1−xSixCxN1−x:Eu

2+.Although extensive investigations have been devoted to the

synthesis, microstructure, and mechanical properties of AlN-SiC solid solutions, there are no reports on their photo-luminescence and applications in electronic devices.20−23 Inaddition, with the introduction of carbon into AlN:Eu,Si, thestructure is definitely altered, which could affect its photo-luminescence spectra, luminescence efficiency and thermalstability. In this work, a detailed survey is therefore tackled tounderstand the structure−property relationship by investigatingthe phase purity, local structure, luminescence, high temper-ature quantum efficiency, and thermal quenching of thecarbidonitride phosphor. The suitability of this phosphor isalso verified by combining it with a near UV chip and green/redphosphors, and a wLED with a high color rendering index of95.3 can be attained.

II. EXPERIMENTAL SECTIONPhosphor Powder Preparation. Powder samples with chemical

compositions Al1−xSixCxN1−x:Euy were prepared by using raw materialsof β-SiC (Ultrafine, Ibiden Co. Ltd., Ogaki, Japan), AlN (Tokuyama,Type F, Shunan-shi, Japan), and Eu2O3 (Shin-Etsu Chemical, Tokyo,Japan). The SiC content was varied from x = 0 to 0.4, and the Euconcentration was in the range of y = 0.0005−0.006. The powderswere mixed in a mortar by hand and packed in boron nitride crucibles.The powder mixture was fired in a gas-pressure sintering furnace witha graphite heater at 1950 °C for 2 h under 1.0 MPa N2 atmosphere.After firing, the phosphor powders were pulverized by hand usingsilicon nitride mortar and pestle. The average particle size of thephosphor, determined by the scanning electron microscope (SEM)observation, is about 10 μm.Structural Analysis. Microstructure observations and energy-

dispersed X-ray spectroscopy (EDS) measurements of the phosphorsamples were carried out at room temperature using a high-resolutionfield emission scanning electron microscope (Hitachi, S4800). Ramanspectra were performed using Ar+ laser excitation (514.5 nm) with asource power of 50 mW and resolution of 1 cm−1 (model NR-1800,

Jasco Co., Tokyo, Japan). The solid state nuclear magnetic resonance(NMR) spectra were recorded on a 14.1-T wide-bore Bruker AvanceIII spectrometer. A 4 mm (magic-angle-spinning) MAS probehead wasused for 27Al and 29Si MAS NMR spectra with the sample spinning at13.5 kHz. The Larmor frequencies for 27Al and 29Si are 156.4 and119.2 MHz, respectively. The excitation pulses were set as 1.0 μs (theπ/6-pulse) for 27Al and 2.5 μs (the π/4-pulse) for 29Si. The recycledelays were 1 and 60 s for 27Al and 29Si, respectively.

Photoluminescence and Quantum Efficiency. Photolumines-cence spectra were measured at room temperature using a fluorescentspectrophotometer (F-4500, Hitachi Ltd., Japan) with a 200 W Xe-lamp as an excitation source. The emission spectrum was corrected forthe spectral response of a monochromator and Hamamatsu R928Pphotomultiplier tube by a light diffuser and tungsten lamp (Noma, 10V, 4A). The excitation spectrum was also corrected for the spectraldistribution of the xenon lamp intensity by measuring Rhodamine-B asreference. External (η0) and internal (ηi) quantum efficiencies (QEs)were calculated by using the following equations:24

∫∫

∫∫

ηλ λ λ

λ λ λη

λ λ λ

λ λ λ λ= =

−

P

E

P

E R

( ) d

( ) d,

( ) d

{ ( ) ( )} di0(1)

where E(λ)/hν, R(λ)/hν, and P(λ)/hν are the number of photons inthe spectrum of excitation, reflectance, and emission, respectively. Theluminescence spectra for QEs measurement were recorded by anintensified multichannel spectrometer (MCPD-7000, Otsuka Elec-tronics, Japan). The reflection spectrum of Spectralon diffusive whitestandards was used for calibration. Fluorescence lifetime data werecollected by using a time-correlated single-photon counting fluor-ometer (TemPro, Horiba Jobin-Yvon) equipped with a NanoLED-370nm with the pulse duration full width at half-maximum of ∼1 ns.

Cathodoluminescence (CL). CL measurements were conductedusing an ultrahigh vacuum SEM with Gemini electron gun (Omicron,Bavaria, Germany) equipped with a CL system.25 The diameter of theelectron beam is in the order of 10 nm. The phosphor powders wereput on a conductive carbon tape and then covered with a copper grid.The specimens were irradiated for 1 h with the electron beam of 5 kVand 1000 pA, with the aim to stabilize the luminescence intensity. TheCL spectra were measured at room temperature.

Thermal Stability Evaluations. The temperature-dependentluminescence was measured by an intensified multichannel spec-trometer. The phosphor powders were heated up to 250 °C in 25 °Cintervals at a heating rate of 100 °C/min and held at each temperaturefor 5 min. The temperature-dependent quantum efficiency wasevaluated by using a QE-1100 phosphor quantum yield spectropho-tometer (Otsuka Electronics, Japan).

Fabrication of White LEDs. Warm wLEDs were prepared bycombining a near UV LED chip (λem = 380 nm) with blueAl1−xSixCxN1−x:Eu

2+ and commercial green (β-sialon:Eu2+ andLu3Al5O12:Ce

3+) and red [(Sr,Ca)AlSiN3:Eu2+] phosphors. The optical

properties of the wLED were measured by using an integrating spherespectroradiometer system (LHS-1000, Everfine Co., Hangzhou,China).

III. RESULTS AND DISCUSSIONLattice Parameters and Structure Description. Both of

AlN and SiC have the hexagonal crystal structure with the spacegroup of P63mc.26 In addition, the bond length of Al−N andSi−C in the AlN and SiC is 1.867 and 1.895 A, respectively.This enables them to form an interesting solid solution in awide range of compositions. It was reported that the solubilityof SiC in AlN could be up to ∼76 mol %.18,19 As seen in Figure1, XRD patterns of Al1−xSixCxN1−x:Eu

2+ with different SiCcontents indicate a pure phase of AlN when x < 0.4, evidencingthe dissolution of SiC in AlN and the formation of solidsolutions between them. Above x = 0.4, a minor phase β-SiC isseen. This implies that the solubility of SiC in AlN is less than x= 0.4 (40 mol %), and the difference in solubility from the

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literature is ascribed to the Eu2+-doping that results in a Sr-sialon polytype-like local structure between AlN4 tetrahedralnetworks.27

A close look at the shift of the (100) and (002) diffractionpeaks with increasing the SiC content reveals an interestingphenomenon, as shown in Figures 1b and c. It is seen that the(100) peak shifts monotonically toward higher angles with thedissolution of SiC, indicating a lattice shrinkage along the a-axis(i.e., along the ab plane). On the other hand, the (002)diffraction peak first shifts to higher angles and then towardlower angles when x ≥ 0.08. It suggests a lattice shrinkage alongthe c-axis at first, and then a lattice expansion finally. Asillustrated in Figure 2, the measured lattice parameters validatesuch observations. Above x = 0.08, the lattice parameters (a andc) of the AlN-SiC solid solution obey the Vegard’s law, and thevariations of these parameters are in a good agreement with theliterature.28,29 In addition, for samples with x = 0.3 and 0.4, thediffraction peaks at 33.0−33.5° and 35.6−36.2° are split intotwo, with the low angle assigned to SiC and the high angle toAlN, respectively. This implies that the solubility of SiC inEu2+-doped AlN is smaller than 30 mol %.The abnormal change in lattice constants at x ≤ 0.08

suggests a different mechanism for the accommodation of SiCin Eu2+-doped AlN. As proposed by Takeda, the Si (Si3N4)atoms in AlN:Eu2+ occupy two sites: one is to substitute Al

atoms by forming a Sr-sialon polytypoid (SrSi9Al19ON31) likestructure between wurtzite AlN4 blocks (i.e., the Eu-containinglayer due to the similar ionic size of Sr2+ and Eu2+) and theother is located at a condensed layer for the wurtzite polarityinversion in between two Eu layers.27 This unique localstructure can also be anticipated in the current work by usingSiC as the Si source, as schematically shown in Figure 3. The

difference to the use of Si3N4 is that the C atoms substitute theN ones simultaneously. Therefore, the mechanism for theabnormal change in lattice parameters can be interpreted byforming a Si condensed layer that compresses the c-axis and aEu-containing layer that promotes the shrinkage along the abplane. Above the critical point of x = 0.08, these uniquestructures are totally formed, and SiC starts to dissolve in thoseAlNs that do not participate in both the condensed layer andthe Eu layer, forming normal Al1−xSixCxN1−x solid solutions.

Raman Spectra. The Raman spectrum is very sensitive tothe cluster ordering and usually used to reveal the localordering, structural distortion, etc.30,31 Figure 4 shows theRaman spectra of Al1−xSixCxN1−x:Eu

2+ with varying SiCconcentrations. As wurtzite AlN has its all atoms to occupythe C3v sites. Six Raman-active modes are possibly present:1A1(TO), 1A1(LO), 1E1(TO), 1E1(LO), and 2E2.

32−34 The E2

Figure 1. (a) XRD patterns of Al1−xSixCxN1−x:Eu2+ with varying SiC

contents, (b) shift of the diffraction peak of (100), and (c) shift of thediffraction peak of (002) with increasing the SiC content.

Figure 2. Lattice parameters (a and c) of Al1−xSixCxN1−x:Eu2+ as a

function of the SiC content.

Figure 3. Proposed crystal structure of SiC-doped AlN:Eu2+ consistingof the Si condensed layer and the Eu-enriched player. The(Al,Si)(N,C)4 tetrahedron and Eu(C,N)12 cuboctahedraon are alsoillustrated.

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(low) mode at 247 cm−1, the A1 (TO) mode at 608 cm−1, theE2 (high) mode at 657 cm−1, and the A1 (LO) at 904 cm

−1 areclearly seen, which are in accordance with the reported values.With increasing the solubility of SiC, the Raman peaks of AlNare getting diffused as a result of the overlap with those of SiC,together with the reduction in their intensity. On the otherhand, Raman peaks at 264, 764, 798, and 964 cm−1 areobserved in the samples containing 30 mol % SiC, which areassigned to 2H-SiC.35,36 The peak at 264 is attributed to thetransverse acoustic (TA) mode, the peaks at 764 and 798 cm−1

correspond to the transverse optical (TO) mode, and the peakat 964 cm−1 is the longitudinal optical (LO) mode.Solid State NMR Spectroscopy. Solid state NMR

spectroscopy provides a description of local arrangementaround atoms. Figure 5 presents the 27Al and 29Si solid state

NMR spectra of samples containing 6 and 15 mol % SiC. For27Al, chemical shifts are largely dependent on the coordinationnumber.37,38 For Al−N coordination compounds, the chemicalshift for tetrahedrally coordinated Al appears near 108 ppm.Usually, in wurtzite AlN the 27Al exhibits a resonance around110 nm.37,38 A distinct maximum at 113−114 ppm is clearlyseen in the chemical shift of 27Al for both samples (Figure 5a),

which arises from AlN4 units and agrees well with the reporteddata.37 In addition, the resonance peak intensity decreases withincreasing the SiC content, suggesting the reduction of the core27Al amount and the formation of solid solutions. The 29SiNMR spectroscopy is useful for quantitative studies of local Sisites. As shown in Figure 5b, two obvious chemical shifts at−13.8 and −20.4 ppm are seen for both compositions, whichcan be assigned to SiC3N and 2H-SiC (SiC4), respectively.

39−41

In addition, the composition of x = 0.06 has another twoshoulders in its spectrum at −25.6 and −31.7 ppm, both ofwhich are for SiCN3.

39,40 Only the should at −25.6 ppm ispresent for the composition of x = 0.15. These results indicateAlN is partially substituted by SiC, forming Si(C,N)4tetrahedra. The significantly enhanced signal at −13.8 ppmfor the composition of x = 0.15 shows further evidence of thereplacement of N by C as the C (SiC) content increases.Furthermore, the absence of the signal at ∼ −48.7 ppm for theSiN4 units indicates that Si is either bounded to C or (C, N)but not to N solely in AlN.42

EDS Mapping. Figure 6 gives the elemental mapping ofAl1−xSixCxN1−x:Eu

2+ (x = 0.06, 0.2 mol % Eu). It evidence that

the phosphor particle is indeed composed of Al, Si, Eu, N, C,and impurity O. The O impurity is originated from Eu2O3, andthe oxidized layer on the AlN starting powder.

Indispensable Role of Si. Samples with varying SiCcontents and a fixed Eu doping level were first examined bycathodoluminescence (CL) and CL mapping. As shown inFigure 7, the concentration of SiC dissolved in AlN plays acrucial role in the CL properties (e.g., the emission color andintensity) of the carbidonitride phosphor. The SiC-free sample(x = 0) exhibits a single broad band centered at 366 nm, andthe diluted sample (x = 0.01) has two distinct bands peaking at366 and 469 nm, respectively. A single broad band is again seenin samples with x in the range of 0.06−0.40, but is centered at472−477 nm depending on the SiC concentration. This is alsoobviously seen in the CL mapping of those samples (Figure 8).The difference in CL spectra of samples with varying SiCconcentrations implies that there are several types ofluminescence centers and mechanisms.The UV emission of the SiC-free sample is quite uniform for

each AlN particle (Figure 8a), suggesting that the emission is

Figure 4. Raman spectra of Al1−xSixCxN1−x:Eu2+ with varying SiC

contents.

Figure 5. 27Al (a) and 29Si (b) NMR spectra of Al1−xSixCxN1−x:Eu2+

with x = 0.06 and 0.15.

Figure 6. EDS elemental mapping of Al1−xSixCxN1−x:Eu2+ with x =

0.06 and 0.2 mol % Eu2+.

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originated from either AlN itself or Eu2+ ions doped in AlN.However, the Eu2+ emission can be excluded. As evidenced bythe EDS elemental mapping, without the SiC codoping Eu2+

ions are not dissolved in the AlN lattice but enriched at thegrain boundary phases (Figure S1). A similar phenomenon wasencountered in Si3N4-doped AlN:Eu2+.17 Furthermore, eventhrough Eu2+ could occupy a highly distorted Al sitecoordinated by four nitrogen in AlN, it emitted a green colorat 540 nm.43,44 In the Si3N4-doped AlN:Eu2+, besides the UVemission there was also a green emission at 550 nm.17 Thisgreen emission was interpreted by the 5d → 4f electronictransitions of Eu2+ enriched in the grain boundary phase.However, the green emission is not observed in the currentsample, perhaps due to its extremely low intensity or itsintensity changing from one sample to another. Therefore, theUV emission at 366 nm is only from the blank AlN. As theband gap emission of AlN is located at ∼200 nm (6.2 eV), the366 m emission can be attributed to the Al vacancy-oxygencomplex defects (VAl−ON).

45,46 It is more clear from theluminescence of the Eu-free Si-doped AlN.47

With the SiC dissolved in AlN, a blue emission is intensifiedand finally dominated solely. The blue emission can beattributed to the Eu2+ luminescence center in AlN, which canbe verified by both CL mapping (Figure 8) and EDS mapping(Figure S1). In the CL mapping is seen the homogeneousemission of the blue color from each phosphor particle.Further, the EDS mapping shows the uniform distribution ofEu2+ in AlN particles, totally differing from that of the SiC-freesample. It thus validates that the dissolution of SiC in AlNpromotes the accommodation of Eu2+ in the same lattice.Takeda et al. used high-resolution transmission electronmicroscope (HRTEM) and HAADF to clarify that Eu2+ is

located in a single layer with Si condensations between thewurtzite AlN4 blocks, forming a EuN12 cuboctahedron that issimilar to Sr(O,N)12 polyhedron in SrSi9Al19ON31.

27,48 In thiswork, the partial substitution of N by C occurs, and thus, Eu2+

is coordinated to 12 N/C atoms, the emission of which thusoccurs at high energies or short wavelengths due to the higherformal charge of C (i.e., large crystal-field splitting). The 472−477 nm blue emission observed in this work is comparable tothe Eu2+ emission in SrSi9Al19ON31 (λem = ∼470 nm).49,50

The SiC concentration also impacts an influence in the CLintensity of the title phosphor. As seen in Figure 7, the CLintensity is maximized at x ∼ 0.06. The luminescence isdramatically reduced as more SiC is dissolved.

Photoluminescence Properties. The diffuse reflectancespectra of the phosphor are given in Figure 9. Two absorption

bands are present in the spectral range of 250−350 and 350−450 nm, respectively. The first one is ascribed to the opticalband gap of the Al1−xSixCxN1−x hosts, and the second one is theabsorption band of Eu2+ due to the 4f7 → 4f65d transitions. It isseen that the optical absorption edge is red-shifted withincreasing SiC, implying that the band gap of the host reduces.It is roughly estimated as 5.04, 4.84, 4.78, 3.63, and 3.55 eV forx = 0.06, 0.1, 0.2, 0.3, and 0.4, respectively. This result agreeswell with the literature data.10,51 Dallaeva et al. addressed thefact that the band gap of (SiC)1−z(AlN)z reduced as the SiCcontent increased.19 Kurbanov et al. reported that the band gapof the solid solution varied from 3.3 to 5.8 eV depending on thecomposition, and a direct-band was for the SiC content smallerthan 0.4.51 In addition, the absorption is getting increased asthe SiC content increases, resulting in a gray body color.The normalized excitation spectra of Al1−xSixCxN1−x:yEu

2+ (y= 0.002) are given in Figure 10a. The introduction of SiC has alarge influence in the excitation spectra of Eu2+ in AlN,compared to the case of changing the concentration of Eu2+

(Figure S2). As the Si−C bonds substitute for Al−N ones, theexcitation spectrum redshifts in both the maximum and the leftwing. Furthermore, the shoulder at ∼350 nm is enhancedobviously, typically for the samples of x = 0.06−0.08. Theenhanced right wing of the excitation spectrum indicates thatthe Al1−xSixCxN1−x:yEu

2+ phosphor with optimized composi-tions is suitable for use with a UV-LED chip. These shifts in theexcitation spectrum are attributable to the variations of 5denergy levels of the excited state of Eu2+, caused by the changesof the band gap of AlN-SiC solid solutions. This will bediscussed later.The emission spectra of samples with varying SiC contents

are also given in Figure 10b. A broad band is observed, due to

Figure 7. Cathodoluminescence of Al1−xSixCxN1−x:Eu2+ (x = 0−0.30).

The Eu doping level is about 0.2 mol %. The spectra were measuredunder an accelerating voltage of 15 kV and a beam current of 100 pA.

Figure 8. CL mapping and corresponding CL spectra ofAl1−xSixCxN1−x:Eu

2+ (x = 0, 0.05, and 0.10).Figure 9. Diffuse reflectance spectra of Al1−xSixCxN1−x:Eu

2+ sampleswith varying SiC and Eu concentrations.

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the dipole-allowed 4f65d → 4f7 electronic transitions of Eu2+.The full-width at half-maximum is about 56 nm for all samples,which is quite smaller than ∼90 nm usually for Eu2+ in manyhosts. It is believed that such a narrow-band emission isascribed to the highly symmetric site of Eu2+ where it forms aEu(N,C)12 cubo-octohedra polyhedron. It has also beenencountered in β-sialon:Eu2+ and Sr[Li3Al]N4:Eu

2+.52,53 Inaddition, the peak maximum of the emission band redshifts

only by several nanometers (465 → 471 nm) for samples withvaried contents of SiC but almost remains the same for sampleswith different Eu concentrations (Figure S3). The longestemission (471 nm) is observed in samples of x = 0.08.Moreover, the maximum luminescence is observed at x = 0.06when the Eu concentration is y = 0.002 (Figure 10c).Concentration quenching was investigated for the sample

with x = 0.06. As presented in Figure S4, the optimal Eu2+

concentration is found at 0.3 mol % (y = 0.003) under the UVlight irradiations. The low concentration quenching value isperhaps due to the limited solubility of Eu2+ in the local layeredstructure of Al1−xSixCxN1−x. This is quite different from theEu2+-doped SrSi3Al19ON31 that basically has no concentrationquenching due to the large distance between Eu2+ ions, but verysimilar to β-sialon:Eu2+ that also does not have a distinctcationic site for Eu2+.49,50,54

The external quantum efficiency (η0) of samples withdifferent SiC contents is shown in Figure 10d, which exhibitsa similar tendency with the luminescence intensity. The highestη0 of 61% is found for the sample of x = 0.06 when excited at365 nm. This sample has an absorption efficiency (αabs) of74.4%. These efficiencies are about 15% (η0) and 11% (αabs)higher than those previously reported for the sample usingSi3N4 as the Si source.16

The crystal-field strength (Dq) can be estimated by using theflowing equation.55

= Ze rR

Dq6

2 4

5 (2)

Where Z is the charge or valence of the anion, e is the charge ofthe electron, r is the radius of the d wave function, and R is thedistance between the central ion and its ligands. The Eu-(C,N)bond length decreases due to the lattice shrinkage at low SiCcontents (x ≤ 0.08) and increases due to the lattice expansionat high SiC contents (x > 0.08), and it is anticipated that thecrystal field strength tends to increase at first and then decreasewith increasing the SiC content. Actually, the crystal fieldsplitting is roughly calculated as 12 220, 12 837, 12 993, 11 177,and 11 160 cm−1 for x = 0.03, 0.06, 0.08, 0.15, and 0.2,respectively, which is in a good agreement with the expectedresults.Taking into account of the changes in the band gap, crystal-

field splitting, and the shape and position of the excitationspectrum with varying the SiC content, we then can speculatethat the energy levels of the Eu2+ excited states also vary withthe introduction of SiC. As illustrated in Figure 11, at lower SiCcontent (x ≤ 0.08), the lowest energy level of Eu2+ moves

Figure 10. (a) Excitation spectra, (b) emission spectra, (c) peakemission intensity, and (d) external quantum efficiency of sampleswith varying SiC contents. The excitation spectra were measured bymonitoring the 470 nm emission at room temperature, and theemission spectra were recorded by exciting the samples under 290 nm.

Figure 11. Schematics of the changes in 5d energy levels of the Eu2+

excited state. CFS, CB, and VB are crystal field splitting, conductionband, and valence band, respectively.

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downward and the highest energy level remains unchanged,resulting in an increased crystal-field splitting and enhancedright wing of the excitation spectrum. On the other hand, athigh SiC content (x > 0.08), the position of the highest energylevel of the Eu2+ excited states is lowered while that of thelowest energy level of Eu2+ does not alter, which is deducedfrom the reduction in crystal field splitting and the band gap ofthe host.Decay Time. The luminescence decay curves were

measured by monitoring the 470 nm emission, as shown inFigure 12. For samples with varying SiC and a fixed Eu2+

concentration, a single exponential decay of luminescenceintensity is obtained for all samples except the composition of x= 0.4. The single exponential decay indicates that only onerelaxation process involved in the decay of the emission, whichis in a good agreement with the fact that there is only onecrystallographic site for Eu2+. Obviously, the decay timedecreases with increasing the SiC content, and it is 0.345,0.339, and 0.325 μs for x = 0.05, 0.1, and 0.2, respectively. Thedecay time of the title carbidonitride phosphor falls in the rangeof 0.2−2.0 μs for the radiative lifetime of Eu2+ in most hosts.The reduction in decay time is perhaps due to (i) increasedsolubility of Eu2+ in the lattice with increasing Si (SiC) thatenhances the energy migration between Eu2+ ions; (ii) reducedthe defect/trap density (e.g., VAl−ON complex pairs) thatdecreases the possibility of nonradiative transitions. On theother hand, for samples of x = 0.3 and 0.4, a double exponentialdecay is seen. The presence of two decay times (τ1 = 0.236 μs,τ2 = 0.831 μs for x = 0.3, and τ1 = 0.144 μs, τ2 = 1.099 μs for x= 0.4) implies that there are two relaxation processesresponsible for the luminescence decay. The reason for thisbiexponential decay is not clear yet. It could be ascribed toadditional defects or traps generated by the excess Si/C thatsignificantly improves the nonradiative transitions. Furtherstudies are required to clarify it.The samples with varying Eu concentrations (y = 0.0005−

0.006) and a fixed SiC content (x = 0.06) exhibit singleexponential luminescence decays (Figure S5). The decay timebasically does not change with the dopant concentration, whichvaries in the range of 0.443−0.404 μs. Usually, the radiativelifetime shortens with increasing the doping levels as a result ofnonradiative energy transfers. We believe that, the smallvariation in decay time is due to a very dilute concentrationof Eu2+ that makes the energy transfer hard to occur.

Temperature-Dependent Luminescence. Thermalstability is a key parameter that controls the reliability andefficiency of a phosphor itself and wLEDs using it. Weevaluated the thermal stability of the title phosphor bymeasuring both the temperature-dependent emission maximumintensity (i.e., thermal quenching) and quantum efficiency. Theeffect of the solubility of SiC on the thermal quenching ofcarbidonitride phosphors is given in Figure 13a, and an

enhanced thermal stability with increasing the SiC content inAl1−xSixCxN1−x:yEu

2+ is clearly seen. At 150 °C, theluminescence intensity of the sample with x = 0.02 declinesby 21% compared to the initial intensity measured at roomtemperature, whereas it reduces only by 10% for the samplewith x = 0.15. Using the Arrhenius equation IT/I0 = [1 + C·exp(−Ea/κT)]

−1 (I0 is the initial luminescence intensity, IT isthe intensity at a given temperature T, C is a constant, and κ isBoltzman’s constant), the activation energy for thermalquenching was calculated as 0.15 eV for x = 0.02 and 0.21eV for x = 0.15 (see inset in Figure 13a). The increased thermalstability is attributable to the partial substitution of Al−N bondsby more covalent Si−C ones that leads to a higher structuralrigidity, which has already been observed in carbon-dopedSr2Si5N8:Eu.

8

Further, the temperature-dependent quantum efficiency ofthe sample with x = 0.06 and y = 0.002 was measured under the365 nm excitation. As shown in Figure 13b, the quantumefficiency reduces slowly when the temperature rises up to 200°C. The internal and external quantum efficiencies only loss by4.4 and 3.7% at 200 °C respectively, indicating high thermalstability of the title phosphor. In addition, this sample has anabsorption efficiency of 66.7% and an external quantumefficiency of 54.3% at room temperature when excited at 365

Figure 12. Luminescence decays of samples with varying SiC.

Figure 13. (a) Thermal quenching of samples with varying SiCcontents (x = 0.02−0.15) and a fixed Eu concentration (y = 0.002).(inset) Activation energy for thermal quenching as a function of theSiC content. (b) Temperature-dependent quantum efficiency of thesample with x = 0.06 and y = 0.002. The sample was measured under356 nm excitation.

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nm. These data are smaller than those shown in Figure 10d,because they were obtained in different measuring systems. Thehigh thermal stability and high quantum efficiency enable thetitle carbidonitride phosphor a robust and efficient luminescentmaterial for UV-pumped wLEDs.Ultrahigh Color Rendition White LEDs. The suitability of

the title phosphor was verified by fabricating warm wLEDsusing the “RGB phosphors + UV-LED” method. Two samplesof wLEDs were prepared by blending the title blue phosphorwith different types of green and red phosphors and combiningthe phosphor blend with a UV-LED chip with a peak emissionwavelength of 380 nm. In Sample A, β-sialon:Eu2+ (λem = 532nm) and Sr-rich CaAlSiN3:Eu

2+ (λem = 610 nm) were used asthe green and red phosphors, respectively. In Sample B, a greenLu3Al5O12:Ce

3+ (λem = 545 nm) and a red Sr-lessCaAlSiN3:Eu

2+ (λem = 630 nm) were selected. All of thesegreen and red phosphors are commercially available. Forcomparison, a commercial blue phosphor, Sr5(PO4)3Cl:Eu

2+,was also mixed with the same red and green phosphors used inSample A, and combined with the same UV-LED chip toprepare a warm wLED.The optical properties and emission spectra of two wLED

samples are given in Table 1 and Figure 14. As seen, both

wLEDs show warm white light (3000−3500 K) and high colorrendering indexes (Ra > 85). By selecting appropriate greenand red phosphors, Sample B even achieves a superhigh Ra of95.3. The reference sample exhibits a color rendering index of92.1. In addition, the luminous efficacy is 7.16 and 5.05 lm/Wfor Samples A and B, respectively. It is quite similar to 8.71 lm/W of the reference sample. The low luminous efficacy isascribed to the low efficiency of the used UV-LED chip (η =0.59 lm/W).Furthermore, the R9 index is as high as 72 for Sample B,

which is three times larger than that for Sample A (see Table2). This big difference is ascribed to the longer-wavelength redphosphor used in Sample B that enhances the red spectral part.The R9 of the reference sample is 44. These results validatesthe suitability of the blue-emitting Al1−xSixCxN1−x:Eu

2+

phosphor in UV-LED pumped white LEDs.

IV. CONCLUSIONSA blue-emitting carbidonitride phosphor Al1−xSixCxN1−x:Euywas generated by forming solid solutions between AlN andSiC. The solubility of SiC in AlN:Eu was smaller than 30 mol %(x = 0.3) when the samples were fired at 1950 °C. Theintroduction of SiC first had the lattice to shrink in both a and caxes if x ≤ 0.08, and then to expand only in the c-axis if x >0.08. The abnormal lattice shrinkage (typically in the c-axis) inthe low SiC content was attributed to the formation of a denselayer that accommodates the Eu activators. Both the Ramanshift and the 29Si solid state NMR indicated that Si wasdissolved in AlN and bonded to C dominantly. Under the 365

nm excitation, the Al1−xSixCxN1−x:Euy carbidonitride phosphorshowed a broad emission band centered at 472−475 nm andhad an optimal SiC content of x = 0.06 for the maximumphotoluminescence. The absorption, internal, and externalquantum efficiencies of the sample (x = 0.06, y = 0.002) were66.7, 81.4, and 54.3% under 365 nm excitation, respectively.The introduction of C obviously enhanced the thermal stabilityof AlN:Eu, with the luminescence maximum and the externalquantum efficiency reducing by 10 and 3% at 200 °C,respectively. A high color rendering wLED with Ra = 95.3,R9 = 72 and color temperature of 3533 K was achieved bycombining the caridonitride blue phosphor with a near UV chip(λem = 380 nm) and LuAG:Ce3+ (green) and (Sr,Ca)-AlSiN3:Eu

2+ (red) phosphors. It thus suggests that theAl1−xSixCxN1−x:Euy carbidonitride phosphor would be an

Table 1. Optical Properties of High Color Rendering wLEDsUsing Al1−xSixCxN1−x:Eu

2+

samplechromatic

coordinates (x, y)

colortemperature

(K)luminous efficacy

(lm/W) Ra

A (0.4357, 0.4035) 3015 7.16 90.1B (0.4018, 0.3854) 3533 5.05 95.3ref (0.4113, 0.4057) 3491 8.71 92.1

Figure 14. Electroluminescence spectra of wLEDs using (a)Al1−xSixCxN1−x:Eu (x = 0.06) + β-sialon:Eu + (Sr,Ca)AlSiN3:Eu(Sample A); (b) Al1−xSixCxN1−x:Eu (x = 0.06) + Lu3Al5O12:Ce +(Ca,Sr)AlSiN3:Eu (Sample B); (inset) enlarged visible part of theemission spectra; and (c) Sr5(PO4)3Cl:Eu + β-sialon:Eu + (Sr,Ca)-AlSiN3:Eu (reference). (d) Photo of wLED (Sample B).

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interesting thermally robust blue phosphor for near UV-drivenwhite LEDs.

■ ASSOCIATED CONTENT*S Supporting InformationThe Supporting Information is available free of charge on theACS Publications website at DOI: 10.1021/acs.chemma-ter.5b04384.

EDS mappings of samples with and without SiC;excitation spectra of samples with varying Eu2+

concentrations; emission spectra of samples with varyingEu2+ concentrations; concentration quenching of thesample with x = 0.06; luminescence decay time ofsamples with varying Eu2+ concentrations (PDF)

■ AUTHOR INFORMATIONCorresponding Author*E-mail: Xie.Rong-Jun@nims.go.jp.NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSWe give our thanks to Dr. Benjamin Dierre from Saint-Gobainfor help with CL experiments and to Mr. Sheng Lin at SunpuOpto Semiconductor Co. Ltd. for measuring the opticalproperties of wLEDs. The financial support of the JSPSKAKENHI (No. 23560811), National Natural ScienceFoundation of China (Nos. 51272259, 61177050, 61575182,and 51572232), and Natural Science Foundation of ZhejiangProvince (No. Y16F050012) is acknowledged. Y.-T.T. and R.-S.L. are grateful for the financial support of the Ministry ofScience and Technology of Taiwan (Contract Nos. MOST104-2113-M-002-012-MY3 and 104-2119-M-002-027-MY3).

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Table 2. Color Rendering Indexes of wLEDs Using Al1−xSixCxN1−x:Eu2+

sample R1 R2 R3 R4 R5 R6 R7 R8 R9 R10 R11 R12 R13 R14 R15

A 88 93 96 88 87 94 87 88 22 82 88 85 89 97 80B 96 97 97 95 96 98 95 89 72 92 96 95 96 98 93

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