high-resolution hard x-ray photoelectron spectroscopy: application of valence band and core-level...

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Nuclear Instruments and Methods in Physics Research A 547 (2005) 98–112

www.elsevier.com/locate/nima

High-resolution hard X-ray photoelectron spectroscopy:Application of valence band and core-level

spectroscopy to materials science

Keisuke Kobayashi

SPring-8/JASRI, Kouto 1-1-1, Mikazuki-cho, Sayo-gun, Hyogo 679-5198, Japan

Available online 9 June 2005

Abstract

Hard X-ray photoelectron spectroscopy using high-flux X-rays from the third-generation synchrotron radiation

sources has shown to be a valuable technique for the study of the bulk electronic and chemical states of materials. This

paper gives an overview of the feasibility tests and the development of new applications of this technology carried out

through collaboration between research institutes, SPring-8/JASRI, SPring-8/RIKEN, and HiSOR since the beginning

of measurements at BL29XU of SPring-8. An energy range from 6 to 10 keV was tested, demonstrating high resolution

and sufficient throughput for practical use. The results clearly demonstrate the wide applicability of this new method to

many areas of materials science and technology.

r 2005 Elsevier B.V. All rights reserved.

PACS: 71; 73; 81; 85

1. Introduction

Nanoscience and nanotechnology continues tomake rapid progress. Researche on the electronicand chemical states of nanomaterial is among themost important issues in both basic and appliedfield. Awareness of this situation encouraged us tosubmit a proposal for a new beam line already in2001 [1], which would have included a hard X-rayphotoelectron spectroscopy (HX-PES) station at

e front matter r 2005 Elsevier B.V. All rights reserve

ma.2005.05.016

ss: [email protected].

around 5 keV. The proposal was approved in 2002;however, it could not be carried out till now due tolack of funds.In parallel we started examining the feasibility

of HX-PES in collaboration with the IshikawaX-ray interference optics group at SPring-8/RIKEN, and conducted the first test experimentin June 2002. Since then we have continued R&Don HX-PES by improving the optics and analyzersystem and seeking to widen its potential applica-tions. The platform for these experiments has beenthe collaboration between SPring-8/JASRI (JapanSynchrotron Radiation Research Institute),

d.

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K. Kobayashi / Nuclear Instruments and Methods in Physics Research A 547 (2005) 98–112 99

SPring-8/RIKEN, and HiSOR (HiroshimaSynchrotron Radiation Center, Hiroshima Uni-versity). We also introduced user groups fromoutside of SPring-8 who are conducting researcheson advanced materials in order to further developthe application of HX-PES. Some of theresults have been presented at the HAXPESworkshop in 2003.

This paper, together with the three papers byTakata et al. [2], Chainani et al. [3], and Shimada[4], presents the activities of the collaborationmentioned above including the results obtainedsince the 2003 HAXPES workshop. For a com-plete overview of all the activities in this field, Ialso refer to the other contributions in theseproceedings. In the following section, I would liketo give a more detailed explanation of themotivation for our research and a perspective onusing HX-PES in materials science research anddevelopment. Next, I present some fundamentalresults to show the potential of the high-resolutionHX-PES, and typical results of various investiga-tions in applied physics. Finally, I mention plansfor the future development of the HX-PES.

2. Material science and photoelectron spectroscopy

Si-based Large Scale Integration (LSI) technol-ogy is still evolving according to the InternationalTechnical Roadmap of Semiconductors [5]. Thisresults in the persistent downsizing of complemen-tary-MOS (C-MOS) transistors. Downsizingcauses a decrease in gate capacitance, whichdetermines the drivability of a C-MOS transistor.To compensate for the decrease, it is necessary todecrease the thickness of the gate insulators. WhenSiO2 is used, the thickness is close to 1 nm, which isonly about four atomic layers. This is almost thematerial limit of SiO2 as an insulator. Worldwideefforts are focusing on making a breakthrough inthis area; new materials, the so-called high-kdielectrics, are being used in the Si-LSI fabricationprocess. In addition to high-k materials variouskinds of materials considered exotic in earlier Sitechnology, such as low-k insulators, nitrides oftransition metals for diffusion barriers to Cu, andCu wiring, are being introduced to LSI. Shallow

doping in the range of 5–10 nm is required for thedrain and source areas. Further intensive study ofdifferent nanolayers and interfaces is required.Nanolayers and nanostructures play essential

roles in many kinds of devices. One can mentionhere compound semiconductor devices such aslaser diodes (LDs), light-emitting diodes (LEDs),and high-mobility electron transistors (HEMTs).The main research interest in compound semicon-ductor materials has been shifting from III–V andII–VI materials to nitrides and oxides since thesuccess of the GaN blue LEDs. Memory devicessuch as hard discs, magneto-optical discs, andphase-change memory for digital versatile discs(DVD) are also typical examples. Organic electroluminescent (EL) display devices are anotherexample. Various organic materials are beinginvestigated for electron injection, hole injection,and electron–hole recombination layers. Little isknown about the electronic and chemical struc-tures at the p–i–n interfaces, between organiclayers and metal or transparent electrodes. As fortransparent electrodes, a search for materials otherthan ITO (indium tin oxide) is necessary due to thescarcity of In resources. Another challenge is thesearch for the materials for transparent thin-filmtransistors, which are required to achieve flexibledisplays [6]. Recently, much effort has beendedicated to researching and developing spinelectronics, including diluted magnetic semicon-ductors and strongly correlated electronic materi-als. Nanostructured materials such as fullerenes,carbon nanotubes, clathrates, and metal clusters,continue to be intensively studied.Most of the materials mentioned above are in

the form of nanostructures, whether they areprepared by artificial fabrication processes orself-organization processes. In situ preparation ofthese nanomaterials is quite difficult in most cases,and consequently conventional photoelectronspectroscopy (PES) with its high surface sensitivityhas severe problems. We need bulk-sensitive PEScapable of penetrating the structures below theirsurface to enable us to investigate their electronicand chemical states, and understand how stablethey are. This will contribute to the developmentof new devices and provide sound guidance in thesearch for new materials.

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K. Kobayashi / Nuclear Instruments and Methods in Physics Research A 547 (2005) 98–112100

Thus, we are convinced that PES with a probingdepth of greater than ca. 5–10 nm is an essentialtool for nanotechnology researches.

Fig. 1. Au Fermi edge spectrum measured at excitation energy

of 9.92 keV with pass energy of 100 eV. Measurement tempera-

ture was 30K. Total resolution reduced from curve fitting is

93.1meV, in which bandwidth of X-ray and analyzer resolution

is estimated to be 90meV and 23.82 eV, respectively. A channel

cut post monochromater of Si 555 reflection was inserted in the

beam path between the vertical focusing mirror and sample.

3. High-resolution HX-PES using BL29XU and

BL47XU at SPring-8

Pioneering experiments to test the feasibility ofHX-PES were carried out by Lindau et al. using abending magnet source in 1974 [7]; however, theweakness of the signal only enabled observation ofAu 4f core spectra, showing that the techniquecould not be applied to valence band spectroscopy.Raising the photon energy has the drawback ofcausing a severe decrease in signal intensity due toa rapid drop in the photoionization cross-sections.This is one of the main reasons that HX-PESexperiments have not been widely used except forcore-level studies and resonance Auger spectro-scopy by a group using a wiggler light source atHASYLAB [8,9]. For valence band spectroscopy,we need a resolution of better than few hundredsmeV, ideally better than 100meV or even less.

SPring-8 is the world’s largest third generationsynchrotron radiation facility, and provides high-flux X-rays with its in-vacuum planar undulators[10,11]. In discussions with the SPring-8/RIKENX-ray optics group, it was suggested that X-rayswith a bandwidth of around 50meV at around6 keV were obtainable at a flux of 2� 1011/s on thesample surface using RIKEN’s X-ray undulatorbeam line, BL29XU [12]. Thus, we modified theScienta SES2002 analyzer system for the analysisof 6-keV photoelectrons in collaboration withGammdata Scienta AB, to carry out test experi-ments. The first experiment was done in June 2002at BL29XU by a team of collaborators fromSPring-8/JASRI, SPring-8/RIKEN, and HiSOR.We succeeded in measuring an Au Fermi edge at aresolution of 0.24 eV with adequate through-put,even though the instrumentation was not opti-mized [13,14]. We have continued to makeimprovements to the analyzer and optics untilnow [2,15]. Coincidentally, several groups havealso been developing PES using high-flux hardX-rays from the third generation synchrotronradiation sources during the time. Their activities

have recently been reported in these proceedingsand also elsewhere [16,17].Fig. 1 shows the state-of-the-art performance of

our 10 keV-R4000 analyzer, which was developedin collaboration with Gammdata Scienta AB. Wehave demonstrated a total resolution of 84meV at8 keV and 93meV at 10 keV, with a pass energy of100 eV. Analysis of the curve fit gives an estimatedanalyzer resolution of 23.8meV. In 2004, westarted HX-PES measurements at BL47XU,whose characteristics were almost the same asthose of BL29XU, using the same R4000 analyzer.To test the applicability of HX-PES for ob-

servation of the valence band, the sample used wasan MBE-grown GaAs crystal without any surfacetreatment prior to the measurements [18]. Thethick solid curve in Fig. 2 is a GaAs valence bandspectrum at 40K. The thin solid curve and thedotted curve are the energy distribution curves(EDC) of the valence band photoelectrons at 40Kand room temperature, respectively, which wereobtained by subtracting the background from theexperimental spectrum. The sharpness of the banddenoted by 2 (predominantly a Ga 4s–As 4p mixedband) is more prominent at a lower temperature,providing evidence that phonon scattering

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16 14 12 10 8 6 4 2 0Binding Energy (eV)

40 K

r.t.1

2

3

Inte

nsity

(a.

u.)

Fig. 2. Bold solid curve shows experimental raw valence band

spectrum of MBE grown GaAs at photon energy of 5.95 keV at

40K. The as-grown sample was measured without any surface

treatment. Bold dotted curve is photoelectron energy distribu-

tion curve (EDC), which was obtained by subtracting the

background (shown by thin dashed curve) from the raw

spectrum. Thin solid curve shows EDC of same sample at

room temperature.


decreases at low temperature. Band 3, which ispredominantly As 4s-like, exhibits significantbroadening in comparison to the local densityapproximation (LDA) calculation. The indepen-dence of broadening on temperature suggests thatthe main contribution is lifetime broadeningcaused by radiative and Auger recombination ofthe hole with electrons in band 1.

The thin solid curves in Figs. 3 (a) and (b) arethe least-squares-fitted EDCs for GaAs and GaN,respectively. The calculated EDCs are defined asEDC(E) ¼

Pa(i)DOSi(E), where DOSi(E) is the

density of the ith partial state calculated by LDAexplicitly including the shallow 3d core states. Thea(i) are fitting parameters. Details of the LDAcalculation and fitting procedure are explainedelsewhere [18]. The peak positions of bands 1 and 2coincide well. Band 3 of the experimental EDC,however, shifts by about 0.3 eV toward a lowerbinding energy in the case of GaAs. The Ga 3d,which shows atom-like spin–orbit splitting of0.45 eV, appears at a binding energy of about3.5 eV deeper than that calculated. This differencein energy position is well known to be due to thedefect in the LDA calculation in predictingbinding energies of states with localized naturelike shallow core states [19]. The curve fitting forGaN was again successful as shown in Fig. 3(b).

The Ga 3d peak shows a broadened shape inFig. 3(b). This is due to the hybridization betweenGa 3d and N 2s [18]. This characteristic feature ofthe Ga 3d band caused by hybridization with N 2sis clearly reproduced in the results of the LDAcalculation, although the calculated structuresappear at a lower binding energy. A simple Ga3d–N 2s hybridization model leads to the follow-ing interpretation. In the tetrahedral crystal field,the Ga 3d degeneracy is lifted and splits into t2 ande states. The small Ga 3d–N 2s energy separationgives rise to hybridization between Ga 3d and N2s. The hybridization causes splitting of the t2 stateinto the bonding (t2g) and antibonding (t2u) states,while the nonbonding state eg remains nearlyatomic, showing spin–orbit splitting. The bondingt2g state is mainly Ga 3d-like, whereas theantibonding t2u state is N 2s dominated. Thebonding t2g state is recognized at 18.47 eV,whereas traces of the nonbonding and antibondingstates are recognizable in the first derivative ofthe experimental EDC, as shown in the insert inFig. 3(b).The photoionization cross-section sðiÞ of the ith

subshell is expressed as sðiÞ ¼ AR

aðiÞDOSðiÞdE.We define the relative photoionization cross-sections by s � ðiÞ ¼ sðiÞ=s (Ga 3d). The s � ðiÞ

values are summarized in Table 1 along with therelative atomic photoionization cross-sections alsonormalized at the atomic Ga 3d cross-section [20].The relative cross-sections of the valence orbitalsin both GaAs and GaN are significantly smallerthan those of the corresponding atomic subshells[20]. In the case of GaN, the relative Ga 4s cross-section is about one order of magnitude largerthan that of Ga 4p and about three times largerthan that of N 2s. The N 2p contribution isnegligible. It is noted that the values in GaN arecloser to the atomic values than those in GaAs.This may be related to less hybridization betweencations and anions in GaN than in GaAs. Tounderstand this solid-state effect, we need to carryout a set of experimental studies of compoundswith systematic variation of cations and anions.It is well known that LDA calculations always

give smaller binding energies for states than thosedetermined by PES experiments [19]. Empirically,it is noticed that this discrepancy becomes smaller

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Table 1

Relative subshell photoionization cross-sections

Cation s Cation p Anion s Anion p

GaAs 0.348 (1.18) 0.282 1.115 (2.4) 0.782

GaN 0.748 (1.18) 0.083 0.285 (0.3) 0.232 (0.032)

Relative photoionization cross-sections of valence band partial

states in GaAs and GaN at 6 keV, normalized to those of cation

3d state. The cross-sections for these two compounds are larger

than the atomic values, which were obtained using interpolated

values of atomic calculations in the photon energy region from

21 to 8 keV (shown in parentheses).

Binding Energy (eV) 20 15 10 5 0

Binding Energy (eV)

(a) GaAs

1

2

-19 -18 -17 -16 -15

(b) GaN

a

nb

b

b nbb

0.5 eV

Inte

nsity

(a.

u.)

Inte

nsity

(a.

u.)

1.3 cV

20 15 10 5 0

Fig. 3. (a) Comparison between experimental (bold) and calculated (thin) EDCs for GaAs. Curve fitting was done by least-mean-

squares method with linear combination of partial DOS contributions. (b) Experimental (bold) and calculated (thin) EDCs for GaN.

Bonding, nonbonding, and antibonding features formed due to hybridization between Ga 3d and N 2s are distinguishable in enlarged

experimental spectrum shown in the inset and are denoted by b, nb, and ab, respectively. Note that spin–orbit splitting of nonbonding

band is recognizable in the first derivative of the experimental spectrum shown in the insert even though it is broadened due to lifetime

effects.


as the bandwidth of the target state is wider, asseen in Fig. 3 (a) and (b). To clarify this, weattempted systematic study of the valence band

and shallow core states by extending the HX-PESmeasurements to ZnO and ZnTe. Fig. 4 plots thebinding energy differences (DE) as a function ofcalculated bandwidth (W ) of each state. It is veryinteresting that DE decreases linearly with increas-ing W until W reachs W c ¼ 0:63 eV and becomeszero when W exceeds W c. The above resultsprovide an empirical method for the correction ofthe binding energy discrepancy between LDAcalculations and experimental observations at leastfor III–V and II–VI compound semiconductors.This will help the quantitative analysis of valenceband and core-level spectra of mixed crystals likeGa1�xInxN1�yAsy. The hybridization between Ga3d, In 4d and N 2s is thought to modify strongly

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4

3

2

1

03.02.52.01.51.00.50.0

Band Width W (eV)

Wc=0.63 eV

∆E0=3.96 eV

GaAs

GaN Ga 3d

ZnTe

ZnO Zn3d

ZnO O2s

GaN

ZnTe 3 ZnO

GaAs GaN 2 ZnO 1 GaAs

ZnTe

GaN GaAs ZnTe

Bin

ding

Ene

rgy

Diff

eren

ce ∆

E (

eV)

Fig. 4. Plot of binding energy differences between experimental

and calculated DE versus calculated widths W is shown for

valence band and shallow core states in GaAs, GaN, ZnTe,

ZnSe, and ZnO. Here W is determined as a full-width at half-

maximum for each band in the calculated EDCs. Linear

decrease of DE with W is common for all four compounds. On

the y-axis, the intercept gives a value of DE at zero width limit,

DE0 ¼ 3:96 eV. The chemical shift is not related to this effect

because no clear ionicity dependence is recognized in the linear

DE � W relation. Lambrecht et al. [19] found that Ga 3d

appears at 3.7 eV higher binding-energy side compared to the

position predicted by LDA. When the core hole stays at the

same atom during the photoemission process, the difference,

which consists of electron correlation and relaxation energy

terms, is approximately calculated by the so-called Dscfapproach as 4.6 eV for Ga 3d in GaN [19], in good agreement

with our results. The critical width W c of 0.57 eV, where DE

diminishes, indicates a measure of wave function extent of

0.24 nm.

-8 -6 -4 -2 0 2Initial-State Energy [ eV relative to bulk Si 1s ]

Pho

toel

ectr

on In

tens

ity [a

.u.] tOX = 1.32 nm

Si(100)

θ=20°Si 1s

hν = 5950 eV

Bulk-SiSiO2X15

Si1+Si2+Si3+

Fig. 5. Upper part of figure shows Si 1s spectrum arising from

1.3-nm-thick silicon oxide formed on Si(1 0 0) at 900 1C

measured at photon energy of 5.95 keV with take-off angle of

701. As shown in bottom part of the figure on enlarged scale,

the Si 1s spectrum shows the spectra arising from three

intermediate oxidation states of Si.


the spectral shape of the shallow core levels as wellas valence band in these kinds of materials. Thusthe correction is expected to be valuable foridentification of the structures in the PES spectra.

4. Applications to materials science and technology

4.1. Si-LSI-related applications

The two papers in these proceedings by Chai-nani et al. [3] and Shimada [4] describe the

application of HX-PES to solid-state physics.Here, I have chosen to rather focus on topics inmaterials science to show the wide applicability ofthis method.Fig. 5 shows a 5950-eV-photon-excited Si 1s

spectrum and its enlarged spectrum measured at aphotoelectron take-off angle of 201 for a 1.32-nm-thick silicon oxide formed on a Si(1 0 0) surface at900 1C. The Si 1s spectrum arising from Si2+ at theSiO2/Si interface can be clearly observed [21]. Thecurve-fitting procedure allows separating the Si1+

and Si3+ components as shown by the dottedcurves. The advantage of HX-PES for detectingsub-oxide peaks at the interface compared withSX-PES is that the de-convolution procedure ofsuperposing spin–orbit components is unneces-sary.The quest for high-k CMOS gate dielectrics is

the most urgent issue in the development of futureSi-ULSI devices. Because the annealing process isinevitably needed to activate implanted impurities,thermal stability of the interface structure betweenthe high-k and Si substrate is the key to LSIfabrication. Among various high-k dielectricshaving a required thickness of around 3–5 nm,HfO2 is considered to be one of the mostpromising candidates. Fig. 6 shows the Si 1sspectra of a HfO2/SiO2/Si(1 0 0) structure before

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1.0

0.8

0.6

0.4

0.2

0.01846 1844 1842 1840 1838 1836 1834

Binding Energy (eV)

HfO2/SiO2/Si(100)as deposition

1.32 nmSiO2/Si(100)

hν=5.95 keVTOA=30 deg

Si-Hf bond

HfO2/SiO2/Si(100) RTA

substrate Si1s

0.8 nm SiO2

4 nm HfO2

Si(100)

Nor

mal

ized

Inte

nsity

Fig. 6. Si 1s spectra for 1.32 nm SiO2/Si(1 0 0), 4 nm HfO2/

0.8 nm SiO2/Si(1 0 0), and annealed 4 nm HfO2/0.8 nm SiO2/

Si(1 0 0) samples measured at 5950 eV photon energy and 301

take-off angle from surface. Structure of Hf-deposited sample is

shown in upper part of the figure.

Photoelectron Take-off Angle [deg.]

30 60 900

40

80

0

40

80

120

0

(a)

(b)

[Hf]/[O]

[Si]/[O]

[Hf]/[O]

[Si]/[O]

Nor

mal

ized

Inte

nsity

(a.

u.)

Fig. 7. Take-off angle dependences of [Hf]/[O] and [Si]/[O],

shown by dashed curves, were calculated using the composi-

tional depth profile shown in Fig. 8 (a) without considering

Si–Hf bonds. Here, [Hf], [Si] and [O] denote Hf 4d spectral

intensity, Si 1s spectral intensity arising from Si substrate, and

Si 1s spectral intensity arising from oxidized Si, respectively.

Fig. 7(a) was obtained for the as-deposited sample, while

Fig. 7(b) was obtained for the sample after RTA. In calculating

these curves the effective electron escape depth leff described in

the text was used. Solid curves in Fig. 7(b) are the calculated

angle dependence of [Hf]/[O] and [Si]/[O] taking in to account

the Hf–Si bonds, of which distribution is given in Fig. 8(b). The

results indicate that interface reaction in the RTA-treated

sample is serious to form the detectable amount of Si–Hf bonds

in the interface region. In the as-deposited sample, a kind of Hf

silicate was formed at the interface; however, Si–Hf bonds were

not formed.


and after rapid thermal annealing of 5s at 1000 1Cwith a takeoff angle of 301 measured from surfacenormal [13]. The spectra are normalized to thesubstrate peaks. The sample was prepared byatomic layer deposition (ALD) of 4-nm-thickHfO2 layers on a 1 nm-thick chemical oxide(SiO2) layer. The SiO2 layer is thought to act asa barrier to interface reactions and interdiffusion.The Si 1s spectra of a sample with 1.32 nm SiO2 onSi(1 0 0) are shown after rapid thermal annealing(RTA) for 5 s at 1000 1C at a take-off angle ofy ¼ 301 for a reference. The acquisition time foreach spectrum was approximately 10min. Afterdeposition of the HfO2 film, the Si 1s peak for theintermediate layer occurred at about 0.6 eV lowerbinding energy of the SiO2 peak. This peak wasattributed to Hf silicate, indicating that formationof Hf silicate had already taken place during theHfO2 ALD. Annealing the sample at 1000 1C indry nitrogen gas for 5 s (abbreviated in thefollowing as RTA) enhanced the intensity of theHf silicate peak. This result suggests that silicateformation is related to diffusion of Si atoms in thedeposited layers from the Si substrate. Theformation of Hf–Si bonds was distinguishable asa result of the annealing by a detectable increase in

the spectral intensity appearing on the lowbinding-energy side of the substrate peak (asshown by the shaded area in Fig. 6).Fig. 7 shows the Si 1s spectral intensity, arising

from the silicon substrate ([Si]), and Hf 4d spectralintensity ([Hf]) as a function of the photoelectrontake-off angle [21]. Both the intensities are normal-ized by the Si 1s intensity arising from the oxidizedsilicon ([O]). The dashed curves in Fig. 7(a) and (b)were calculated using the depth profiles shown inFig. 8(a) without considering the Si–Hf bonds. Incalculating these dashed curves, it was assumedthat Hf forms HfO2 and the rest of the oxygenatoms form SiO2. The effect of a change in

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0 1 2 3 4 5 6 7

Si

Si-Hf

Distance from Oxide Surface [nm]

HfO2

0

0

0.5

0.5

1

1

SiO2

Si (a)

(b) HfO2

SiO2Com

posi

tion

Fig. 8. (a) Depth profile of constituent atoms in HfO2/chemical

oxide/Si(1 0 0) structure after RTA determined by applying

maximum entropy concept to angle-resolved Hf 3d and Si 1s

spectra. (b) Depth profile of constituent atoms in HfO2/

chemical oxide/Si(1 0 0) structure after RTA with (solid curves)

and without ( dushed curves) taking into account the Si–Hf

bonding explicitly. Convergence of the analyses was improved

by refering to the atomic composition profiles obtained by the

high resolution Rutherford backscattering measurements on

the same samples by K. Kimura et al. at Kyoto University.

1852 1850 1848 1846 1844 1842 1840 1838 1836

590 C

620 C

Si 1s

Oxi

diza

tion

adva

nces

Oxidized Si4+

Nor

mal

ized

Int

ensi

ty (

arb.

uni

ts)

Binding Energy (eV)

Fig. 9. Typical Si 1s spectra for poly-Si(620 1C)/HfAlO/SiON I/

I and annealing (thick solid line) and poly-Si(590 1C)/HfAlO/

SiON I/I and annealing (thin solid line).


composition along the depth direction of theeffective electron escape depth was also taken intoaccount. As shown in Fig. 8 (a) and (b), theexperimental data for the sample prior to RTAwas explained without considering the Si–Hfbonds, while the experimental data for the sampleafter RTA was explained by taking into accountthe distribution of the Si–Hf bonds as shown inFig. 10 (b). It was concluded that as a result ofRTA, the amount of SiO2 formed between theHfO2 and the silicon substrate increased; in otherwords, the amount of Hf silicate increased. Inaddition, RTA resulted in the formation of Si–Hfbonds near the silicon substrate; these are deducedto create a metallic interface as pointed out byXiong et al. [22].

Another important problem in the high-k gatestructures is the stability of the interfaces betweenthe gate electrodes. Among the Hf-based materialsunder consideration, research interest has focusedon the use of HfSiO(N) and HfAlO for poly-Sigate CMOSs because of their good transistor

performances and thermal stability [23]. Onemajor issue in the application of Hf silicatematerials to advanced poly-Si gate CMOS circuitsis the high threshold voltage (Vth) of p-channelmetal-oxide–semiconductor field effect transistors.Several mechanisms have been proposed for thisVth problem. The following factors are consideredto obstruct the electrical properties: fixed chargesin the high-k material, interface states, interactionbetween the high-k and the Si substrate/SiO2/gateelectrode, crystallization of the high-k material,phase separation, dopant penetration into thehigh-k film or Si substrate, and the formation ofbonds such as Hf–B or Hf–Si [24,25]. To clarifythe origin of the physical and chemical propertiesrelating to this Vth problem, it was consideredimportant to investigate the interface structure of apoly-Si/high-k insulator system with a depth-sensitive probe [26]. A set of poly-Si/Hf-basedhigh-k insulator structures were prepared usingvarious process conditions. Poly-Si gate electrodeswere thinned to 5 nm by dry etching. Fig. 9 showsthe Si 1s spectra of poly-Si(590 1C)/HfAlO/SiON(thick solid line) and poly-Si (620 1C)/HfAlO/SiON (thin solid line). Both samples were ionimplanted and annealed. The temperatures forhigh-k deposition are shown in parentheses. Thesespectra were quite surface sensitive, since the take-off angle from the sample surface was 81. The

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0 10 20 30 40 50 60 70 80 90

0

2

4

6

8

10

12

14

Inte

grat

ed In

tens

ity

Take Off angle (degree)

Take off angle dependence of [Si+4]/[Si]

I

II

Fig. 10. Take-off-angle dependence of Si 1s spectra. Integrated

intensities for Si4+ component are plotted as functions of take-

off-angle, where integrated intensities for Si4+ component are

normalized to Si substrate peak for each take-off-angle. Key to

symbols: ’, poly-Si(590 1C)/HfAlO/SiON I/I and annealing;

K, poly-Si(590 1C)/HfAlO/SiON as-deposited; &, poly-

Si(590 1C)/HfSiON/SiO2 I/I and annealing; J, poly-

Si(590 1C)/HfSiON/SiO2 as-deposited; m, poly-Si(620 1C)/HfA-

lO/SiON I/I and annealing; ., poly-Si(620 1C)/HfAlO/SiON

as-deposited; U, poly-Si(620 1C)/HfSiON/SiO2 I/I and anneal-

ing; U, poly-Si(620 1C)/HfSiON/SiO2 as-deposited.


binding energy of the peak at 1844.7 eV corre-sponded to fully oxidized Si4+. Fig. 9 indicatesthat oxidation of the poly-Si samples deposited at620 1C was more advanced than for samplesdeposited at 590 1C. Fig. 10 shows the dependenceof the Si 1s spectra on the take-off angle, where theintegrated intensities for the Si4+ component arenormalized to the Si substrate peak for varioustake-off angles. The take-off angle dependencecurves formed two groups and a clear correlationwith the Vth shift was observed. That is, thesamples in group I showed less Vth shift whilethose in group II showed a large shift. This resultcan be explained under the following assumptions.When oxygen diffuses into the poly-Si duringdeposition, a stable interface is formed, thus Vth isalso stable. In contrast, oxygen does not diffuseinto poly-Si during deposition, it diffuses after theboron implantation and annealing process, sig-nificantly affecting the Vth. This study is importantbecause it showed for the first time that HX-PES isvaluable for characterizing the Si-LSI process.

4.2. Wide gap diluted magnetic semiconductors

Theory based on hole-mediated ferromagnetismin the frame of the Zener model has allowedreliable estimates of the Curie temperature (T c) fordiluted magnetic semiconductors (DMSs) likeGa1�xMnxAs [27,28]. This theory basically as-sumes that the ferromagnetism is due to interac-tions between the local moments of the transitionmetal atoms mediated by itinerant holes in thematerial.Recently, very stable room temperature (RT)

ferromagnetism of Cr-doped GaN was predictedby the first principle calculation [27] and confirmedexperimentally [28–30]. Since doped transitionmetals introduce deep levels in wide band gapsemiconductors, the applicability of the hole-mediated ferromagnetism model to wide bandgap GaN-based DMS is questioned. To elucidatethis point, Kim et al. attempted to conduct HX-PES studies of Ga1�xCrxN at various Cr concen-trations [31]. The valence band HX-PES spectra ofundoped GaN (open circles) and Ga0.899Cr0.101N(filled circles) are shown in Fig. 11(a). The solidline shows the difference spectrum. Cr dopingclearly introduces new electronic levels in the bandgap (A) and causes changes in the valence bandstructure (B and C). As shown in Fig. 11 (b),intensities of the in-gap states proportionallyincreased with increase in the Cr concentration inthe range from x ¼ 0 to 0.101. Therefore, the in-gap states are closely related to the Cr 3d orbitals.To identify the origin of the new electronic states,the electronic band structure for Cr-doped GaNwas calculated based on LDA by generating asupercell that contained the objects of interest [31].For GaN doped with Cr, one of the 16 sites of theGa atoms was replaced by a Cr site in the supercellmodel. Since the Cr atoms in Ga sites aretetrahedral bonded with four nitrogen atoms, 3de(dxy, dyz, and dzx orbital) and 3dg (dx2

�y2 and dz2

orbital) states are separated into nonbonding (e),bonding (tb), and antibonding states (ta) in thetetrahedral crystal field. Two sharp spin-up bandsarose in the calculation results, each of whichcorresponded to the e and ta orbitals. The tb wasfound to merge with the valence band andthe spin-down band overlapped the conduction

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-8 -6 -4 -2 0 2 4

0

EF

Diffrence spectrum

Ga1-xCrxN

A

B

CPho

toem

issi

on In

tens

ity (

arb.

unit)

x=0

x=0.101

Binding Energy (eV)

Pho

toem

issi

on In

tens

ity (

arb.

uni

ts)

-1 0 1 2 3 4

Energy (eV)

undoped GaN

Cr=1.3%

Cr=3.0%

Cr=6.3%

Cr=10.1%

(a) (b)

Fig. 11. (a) EDCs for undoped GaN (open circles) and Ga0.899Cr0.101N (filled circles) at photon energy of 5.95 keV. These EDCs were

obtained by subtracting trivial background from experimental spectra. Solid line in bottom panel shows difference spectrum. (b) Cr

concentration dependence of Ga1�xCrxN valence band spectrum in band-gap region.


band. Because of Ef positioned at ta, the spin-upand -down states are separated. Thus the materialis spin polarized.

The EDCs are known to consist of Ga 4s (65%)and N 2p (24%) components according to least-squares fitting of the experimental spectrum with aLDA calculation [18]. A detailed comparison ofthe Ga 4s and Cr 3d photoionization cross-sections [19,31] led to the conclusion that the bandgap states originated mainly from a Ga 4s orbitalthrough hybridization with Cr 3d.

The open circles, solid gray circles, and solidblack circles in Fig. 12 (a) show the N 1s core-levelspectra of undoped GaN, Ga0.937Cr0.063N, andGa0.899Cr0.101N, respectively. The difference spec-tra between the undoped GaN and Ga1�xCrxN(x ¼ 0:063 and 0.101) are shown by gray and black

solid curves in the lower part of the figure. Thesedifference spectra indicate that the original peak(394.5 eV) decreased and a new chemical-shiftedcomponent BE region (393.8 eV) increased with Crdoping in the low-energy tail of the original peak.The rate of the decrease and increase wasproportional to the Cr concentration. Fig. 12 (b)shows the Ga 2p3/2 core-level spectra with thedifference spectra of the same set of samples. Theintensity of the Ga 2p3/2 core-level spectrumdecreases linearly with increasing Cr, evidencingthat Cr atoms substitute for Ga atoms in the GaNmatrix. The full-width at half-maximum (FWHM)increased with Cr content (inset in Fig. 7(b)). Theline shape of the difference spectra was asym-metric. This indicates the chemical effects of Cr onGa. The electronic states in the GaN band gap

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-398 -396 -394 -392 -390

0.0

Binding Energy (eV)

Nor

mal

ized

pho

toem

issi

on In

tens

ity (

arb.

uni

ts)

x=0.063 -x=0

x=0.101 -x=0

N 1s

x=0

x=0.063

x=0.101

-1118 -1116 -1114 -1112 -1110

0

x=0.063 - x=0

x=0.101 -x=0

Ga 2p3/2 Ga1-xCrxN

x=0

x=0.063

x=0.101

Nor

mal

ized

pho

toem

issi

on In

tens

ity (

arb.

units

)

Binding Energy (eV)

Ga1-xCrxN

0.92

0.90

0.88

0.86

0.84

FW

HM

0.00 0.05 0.10Cr contents (x)

(a) (b)

Fig. 12. Core-level spectra of (a) N 1s , and (b) Ga 2p3/2 of undoped GaN (open circles) Ga1�xCrxN (x ¼ 0:063 (solid gray circles), and

0.101(solid black circles)). Right inset in Fig. 12(b) shows the FWHM variation in Ga 2p3/2 spectra with Cr concentration. Difference

stpectra between GaAs and Ga1�xCrxN spectra are also shown in the bottom of (a) and (b).


region induced by Cr doping in a GaN matrix isfor the first time observed by HX-PES at 6 keV.These results, together with the results for N 1sand Ga 2p3/2, provide evidence for the Cr 3d–N2p–Ga 4s hybridization, which may play asignificant role in the magnetic properties ofGa1�xCrxN. Similar results for Ga1�xMnxN werealso reported by Takata et al. [14].

4.3. Strongly correlated electron materials

Tanaka and his co-workers found that thetensile strain from the substrate stabilized thedouble exchange ferromagnetism in lightly doped(La1�xBax)MnO3 thin films [32,33] and that ultra-thin films displayed (RT) ferromagnetism incomparison to a ferromagnetic Curie temperature(Tc) of 260K in unstrained systems [34]. Thisfeature was not observed in other manganites, and

is also important in relation to applicationspertaining to spintronic devices such as a ferro-magnetic (La,Ba)MnO3/Pb(Zr,Ti)O3 field effecttransistor [35] and a ferromagnetic (La,Ba)MnO3/Sr(Ti,Nb)O3 p-n diode [36] operating at roomtemperature. To elucidate the origin of thisenhanced ferromagnetism, information on theelectronic structures of the films is essential, andhelps us to understand the lattice structure andelectric/magnetic properties.The standard PES technique is very surface

sensitive even if soft X-rays (SX) with relativelyhigh energy (hn�1 keV) are used. Thus, thistechnique sometimes produces contradictory re-sults in relation to bulk physical properties.In fact, Taguchi et al. [37] and Horiba et al.[38] reported that low binding-energy satellitepeaks were clearly observed in (V,Cr)2O3, high-Tc cuprate, and (La1�xSrx)MnO3 thin films

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644 642 640 638Binding Energy (eV)

Nor

mal

ized

Pho

toel

ectr

on In

tens

ity

hv = 5.95 eVt = 20 nm

∆

∆∗U

UHB

LHB

O 2p

EF

Mn 2p3/2

28K100K

200K220K

320K

240K260K280K300K

160K180K

Fig. 13. Temperature dependence of Mn 2p3/2 spectra for

(La0.85Ba0.15)MnO3 film with 20 nm thickness. Inset shows

schematic energy diagram of valence band.

1.2

1.0

0.8

0.6

0.4

0.2

0.0

Nor

mal

ized

Sat

ellit

e P

eak

Inte

nsity

0 50 100 150 200 250 300 350Temperature (K)

(V *)2

Satellite Intensity, Is(T)

M (T)

M (T)2

Fig. 14. Summarized dependence of normalized satellite

intensities of Mn 2p3/2 spectra (ISðTÞ), normalized hybridiza-

tion parameter (V*)2, magnetization (MðTÞ, and MðTÞ2) on

temperature. ISðTÞ and ðV�Þ2 are normalized at 28K (as 1.0)

and 300K (as 0).


displaying a metallic state in the core-level spectra(V 2p, Cu 2p, Mn 2p) using bulk-sensitive HX-PES at 6 keV, but only traces of the satellite arerecognizable when using SX-PES at 1.5 keV. Basedupon a cluster model calculation, they concludedthat the satellite peaks had originated from core–hole screening through evolution of the density ofthe state at bulk Fermi level (DðEf Þ).

Bulk-sensitive HX-PES was used to investigate(La,Ba)MnO3 thin films of varying film thicknessthat exhibited strain-enhanced ferromagnetism[39]. Fig. 13 shows the dependence of the Mn2p2/3 spectra on temperature for the t ¼ 20 nmfilm, the ferromagnetism of which was stronglyenhanced. When the temperature was increasedfrom 28 to 320K, the intensity of the satellite peaksystematically decreased, and almost disappearedat 300K, which corresponded to a T c of 299K forthis film. The closed squares in Fig. 14 show thedependence of the integrated satellite intensityI sðTÞ on temperature obtained by subtractionfrom the spectrum at 300K. The I sðTÞ wasnormalized to have a value of unity at the lowesttemperature of 28K. The intensity systematicallyincreased with the developing ferromagnetic orderas the temperature decreased.

Cluster calculations, which were successful for(V,Cr)2O3, high-T c cuprate, and (La1�xSrx)MnO3

[37,38], were also used to estimate the developmentof DðEf Þ (the C state denotes a coherent metallicband at Ef, as shown in the inset in Fig. 13). Theexperimental spectra were fitted by changing onlytwo parameters. One is D�, which is the chargetransfer energy between the Mn 3d and the Cstates. The other is V*, which represents thehybridization between the central Mn 3d orbitalsand the C states. Details of the fitting procedureare described elsewhere [38,39]. The (V*)2 isproportional to the DðEf Þ by an analogy with theKondo coupling between the f and conductionband states [40]. As shown in Fig. 14, (V*)2 varies(closed circle) in parallel with I sðTÞ. From thisfact, it is concluded that the integral intensity ofthe well-screened peak, I sðTÞ, probes the density ofstate at the Fermi level. In the same figure,temperature dependence of MðTÞ and fMðTÞg2

are also shown. Here M(T) stands for themagnetization of the same sample measured bySQUID magnetometer. The experimental I sðTÞ vs.T curve agrees well with the fMðTÞg2 vs. T curve.Small discrepancy remains between I sðTÞ andfMðTÞg2. In order to see whether this residue isdue to the insufficient bulk sensitivity of HX-PESusing 6 keV X-rays for excitation, a HX-PESexperiment using higher photon energy up to10 keV was carried out. No photon energy

ARTICLE IN PRESS

Inte

nsity

(arb

. um

its)

1.5 1.0 0.5 -0.5 -1.00.0

Binding Energy (eV)

Au plate7.6nm5.1nm3.0nm2.2nm1.7nm

Fig. 15. Size dependence of valence band spectra at near Fermi

level of Au nanoclusters. Au clusters with sizes larger than 5 nm

show metal-like Fermi edges, whereas a decrease in density of

states just below Fermi level is clear in smaller clusters.

Measurements were done at 5.95 keV photon energy.


dependence was found in evolution of the spectrallineshape and the intensity as functions oftemperature. These results for the Mn 2p corelevel could provide a measure of metallicity; i.e.,the DðEf Þ, in the metal–insulator transition, andinformation on the electronic structure corre-sponding to double exchange ferromagnetism inthe bulk part of the thin film of strongly correlatedelectronic materials.

4.4. Nanoclusters

As an example of the application of HX-PES tonanofunctional materials, Ikenaga et al. studiedthe cluster size dependence of the valence bandand Au 4f core spectra of polymer-stabilizedAu nanoclusters [41]. Their experiments weremotivated by the discovery of Yamamoto et al.of ferromagnetic spin polarization in polymer-stabilized Au particles by X-ray magnetic circulardichroism (XMCD) at the Au L2,3 edges [42].These Au nanoclusters showed that the saturationmagnetization begins to increase at around 4–5 nmand exhibits the maximum at around 3 nm atlow temperatures [43]. Corresponding to thismagnetic behavior, changes in the Au 4f coreand valence band spectra were observed. Themost interesting fact is the change in the valenceband at the Fermi level, which was observed usingX-rays of 6 keV, as shown in Fig. 15. Aunanoclusters with size greater than 5 nm showbulk-like Fermi edges, whereas the clusters withsmaller sizes loose density of state at the Fermilevel. Shift and broadening of the Au 4f spectra,and changes in the Au 5d bands were alsoobserved in clusters with sizes of less than 5 nm.The fact that the probing depth is sufficientlylarger than the cluster size proves this behavior isof bulk rather than of surface nature. Thisexperimental result suggests that HX-PES can beapplied to the investigation of ‘‘bulk’’ electronicsas well as chemical properties in various nanoclus-ters and particles.

4.5. Organic materials and other materials

One of the most serious shortcomings inconventional lower-energy PES for the study of

organic materials is that VUV-SX radiation read-ily damages samples. In test HX-PES measure-ments of organic conductor molecules, beta-(ET)TCNQ, kappa-(BETS)2FeBr4, at 6 keV, andalso of several lubricant molecules on hardmagnetic storage discs at 8 keV, no spectraldegradation was found over several hours ofcontinuous measurements at room temperature.In contrast, very rapid degradation was found inall the samples during the test measurements of thesame samples performed using SPring-8 SX beamlines. This advantage of HX-PES may be attribu-table to the low absorption of hard X-rays andalso the high kinetic energy of photoelectrons,which contribute less to secondary electron gen-eration.HX-PES can be applied to many kinds of

materials that are interesting from the point ofview of both basic and applied research. We havealready tried applying HX-PES to transparentconducting films like indium-oxide-based materi-als, laser annealing of amorphous Si, plasma-doped ultra thin junctions in Si for drain andsource regions in Si-LSI, Ga2Sb5Te2 thin filmsfor digital versatile discs, mesoscopic-structuredPt, electrode layers in Li solid-state batteries,and so on.

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5. Future development

Since 2003, JASRI has opened HX-PES topublic use within the framework of a nanotech-nology support project. A fixed experimentalstation is in the process of being establishedto enable wider public use beginning from autumnof 2005.

In the technical area, there are still many tochallenge. Among these is the challenge of achiev-ing higher resolutions approaching 10meV athigher energies over 10 keV. We have to overcomethe drastic decrease in intensity, for instance, usinga more intensive undulator such as the BL19LXUat SPring-8. Another challenge is the developmentof angular resolved HX-PES. For this, we needbetter focusing of the excitation X-rays as well asbetter angular resolution in the analyzer.

We are also developing an experimental stationfor a HX-PES scanning microscope, by putting awide-acceptance objective lens in front of theanalyzer. Very recently, Daimon has successfullydesigned an objective lens with acceptance angle ofaround 901 and angle resolution of less than 11[44]. The angle divergence at the outlet of thisobjective lens matches the acceptance of the10 keV R4000 analyzer, that is, 151. The take-offangle dependence of photoelectron spectra, fromwhich one can elucidate depth profile of chemicalstates, is expected to be obtainable withoutrotation of the sample. Combination of this systemwith a focused X-ray beam and a sample scanningstage will enable us to perform three-dimensionalmapping of the chemical states in nanostructuredmaterials.

Finally, it should be mentioned that core levelswith photoionization cross-sections of10�5–10�6Mb range were detectable using 8 keVX-rays, suggesting that HX-PES is applicable fordetection of most of the subshells of almost all theelements.

6. Summary

The necessity, feasibility and applicability ofHX-PES were discussed in this article. Theadvantages of this method over conventional

PES at lower photon energies include greaterprobing depth, high resolution, high throughput,high quality of obtained spectra, and wide applic-ability to materials including hard, soft, and evenwet materials. The shortcomings of HX-PES arealso outlined. One of the most obvious disadvan-tages is that some of the important partial statesare not observable due to the rapid decrease inphotoionization cross-sections with photon en-ergy. This is an especially serious problem inattempting to observe transition metal 3d states incompounds like MnTe because the strong Te 6scontribution masks that of the Mn 3d. In spite ofthis drawback, HX-PES is likely to become anessential tool for research in solid-state physics,and materials science and technology.

Acknowledgements

All the work presented in this paper is the resultof collaborations between various groups. Theissue of X-ray optics owes thoroughly to Prof. T.Ishikawa’s SPring-8/RIKEN group. The work ofDr. K. Tamasaku, Dr. D. Miwa, and Dr. Y.Nishino is acknowledged with appreciation. Dr.M. Yabashi of SPring-8/JASRI also contributedto the design of the optics for HX-PES atBL29XU. The contribution from Prof. S. Shin’sSPring-8/RIKEN group was essential for carryingout the HX-PES test experiments. Dr. Y. Takataled the design and construction of the experimen-tal station, and the test experiments with the helpof Dr. T. Tokushima, Dr. K. Horiba, and Mr. K.Yamamoto. The contribution of Dr. E. Ikenaga ofSPring-8/JASRI was also indispensable for per-forming the various feasibility test measurementsin collaboration with different groups outside ofSPring-8. The collaboration with GammdataScienta AB was essential for development of thehigh-energy analyzer system. In particular, thecontribution of Dr. B. Wannberg and Dr. S.Sodergren are much appreciated. Recently, mea-surements have begun using BL47XU. Dr. A.Takeuchi and Dr. M. Awaji of SPring-8/JASRIare acknowledged for their help in aligning theX-ray optics at this beam line. The contributionfrom the HiSOR group was also inevitable. The

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encouragement and support of Prof. M. Tanigu-chi, and Prof. H. Namatame in carrying out thecollaborations were much appreciated. Prof. K.Shimada and Dr. M. Arita were the maincontributors in performing the test experimentswith the HiSOR group. The author is also gratefulto Prof. T. Yao, Dr. H. Makino, and Dr. J. J.Kim, of Tohoku University, and Prof. T. Yama-moto of Kouchi University of Technology, fordiluted magnetic semiconductors, Prof. K. Hattoriand Prof. H. Nohira and their co-workers forhigh-k materials, and Prof. H. Tanaka and his co-workers for (La,Ba)MnO3 research. Collaborationand discussions with Prof. C. Kunz and Dr. W.Drube were stimulating and helpful. Many re-search groups, which are not mentioned above arealso acknowledged for their involvement in theseries of test experiments at BL29XU and also atBL47XU. This work is partially supported by thenanotechnology support project of the Ministry ofEducation, Culture, Sports, Science and Technol-ogy (MEXT), and also partially by MEXTthrough a Grant-in-Aid for Scientific Research(A) (no. 15206006).

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