Ion-bombardment-induced reduction in vacancies and itsenhanced effect on conductivity and reflectivity in hafnium nitridefilms

Zhiqing Gu1 • Jiafu Wang1 • Chaoquan Hu1 • Xiaobo Zhang1 • Jianchen Dang1 •

Sam Zhang2 • Jing Gao1 • Xiaoyi Wang3 • Hong Chen4 • Weitao Zheng1

Received: 8 December 2015 / Accepted: 24 July 2016

� Springer-Verlag Berlin Heidelberg 2016

Abstract Although the role of ion bombardment on elec-

trical conductivity and optical reflectivity of transition

metal nitrides films was reported previously, the results

were controversial and the mechanism was not yet well

explored. Here, we show that proper ion bombardment,

induced by applying the negative bias voltage (Vb), sig-

nificantly improves the electrical conductivity and optical

reflectivity in rocksalt hafnium nitride films regardless of

level of stoichiometry (i.e., in both near-stoichiometric

HfN1.04 and over-stoichiometric HfN1.17 films). The

observed improvement arises from the increase in the

concentration of free electrons and the relaxation time as a

result of reduction in nitrogen and hafnium vacancies in the

films. Furthermore, HfN1.17 films have always much lower

electrical conductivity and infrared reflectance than

HfN1.04 films for a given Vb, owing to more hafnium

vacancies because of larger composition deviation from

HfN exact stoichiometry (N:Hf = 1:1). These new insights

are supported by good agreement between experimental

results and theoretical calculations.

1 Introduction

The Group-IVB, VB and VIB transition metal nitride

(TMNx, TM = Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W) are well

known as a class of fascinating and technologically

important materials in the fields of electrochemical power

sources, microelectronic devices, cutting- and machining-

tool industries [1–3]. They have attracted much attention

because of their excellent electrical conductivity [4], high

infrared reflectance [5], large free-carrier concentrations

[6] coupled with other properties such as diffusion resis-

tance [7], thermal stability [8], high hardness [9, 10],

abrasion [11] and corrosion resistance [12]. This makes

them competitive for use as electrodes in fuel cell and Li-

ion microbatteries [13], alternative plasmonic materials [6],

solar control coatings on windows [14], decorative coatings

[15], gate electrode for MOS technology [16], conductive

diffusion barrier layers [17, 18], thin film resistors [19]

highly reflecting back-contacts in solar cells [20] and light-

emitting-diode (LED) devices [21]. In these applications,

the electrical and optical properties of the films are of

utmost importance, as they determine the ultimate perfor-

mances of the devices. It is very crucial, therefore, to

understand how to control them.

Studies have revealed [22, 23] that ion bombardment via

applying negative bias voltage (Vb) during film deposition

Electronic supplementary material The online version of thisarticle (doi:10.1007/s00339-016-0308-0) contains supplementarymaterial, which is available to authorized users.

& Chaoquan Hu

cqhu@jlu.edu.cn

& Weitao Zheng

wtzheng@jlu.edu.cn

1 School of Materials Science and Engineering, Key

Laboratory of Mobile Materials, MOE, and State Key

Laboratory of Superhard Materials, Jilin University,

Changchun 130012, China

2 School of Mechanical and Aerospace Engineering, Nanyang

Technological University, 50 Nanyang Avenue,

Singapore 639798, Singapore

3 Key Laboratory of Optical System Advanced Manufacturing

Technology, Changchun Institute of Optics, Fine Mechanics

and Physics, Chinese Academy of Sciences,

Changchun 130033, China

4 Department of Control Science and Engineering, Jilin

University, Changchun 130025, China

123

Appl. Phys. A (2016) 122:776

DOI 10.1007/s00339-016-0308-0

is an effective way to alter the electrical and optical

properties in TMNx films. Although extensive researches

on the influence of Vb on electrical conductivity and optical

reflectivity have been carried out and numerous results

reported [22–24], the following three questions still remain

unanswered: (1) what is the fundamental role that the Vb

plays in electrical conductivity and optical reflectivity? The

reported results are controversial. Some argue that the Vb

causes increase in the electrical conductivity and optical

reflectivity [24–26], while others claim the opposite

[22, 27, 28]. Clearly, further studies are needed to clear the

mist and elucidate the effect, especially in d-HfNx films as

no systematic investigations on these films are conducted.

(2) What governs the Vb-induced evolution in electrical

conductivity and optical reflectivity? Even though the

phenomenon has been reported previously where electrical

conductivity and optical reflectivity in TMNx films depend

strongly on Vb, the mechanism is unknown. (3) How will

the effect of Vb change with change of chemical stoi-

chiometry in the film? It is known that non-stoichiometric

TMNx films were usually obtained experimentally because

point defects exist when x deviates from unity [29, 30]. It

has been a conjuncture that Vb has different influence when

the stoichiometry is different, and that, in turn, affects

electrical and optical properties differently. To date, how-

ever, no research on this has been attempted yet.

With the above three questions in mind, we prepared

near-stoichiometric (x & 1.04) and over-stoichiometric

(x & 1.17) rocksalt hafnium nitride films (d-HfNx) at dif-

ferent Vb and studied the effect of ion bombardment on

structure, electrical conductivity and optical reflectivity of

the films by experimental measurements in combination of

the first-principle calculations. The experimental and the-

oretical results agree well, proving that: (1) in both HfN1.04

films and HfN1.17 films, proper Vb promotes significant

increase in both electrical conductivity and infrared

reflectance; too low or high Vb results in negative effect;

(2) the changes in electrical conductivity and infrared

reflectance with Vb arise from the variation in vacancies

concentration, because the formation of vacancies causes

reduction in the concentration of free electrons and electron

relaxation time; (3) although Vb has similar effect on both

HfN1.04 and HfN1.17 films, the electrical conductivity and

infrared reflectance of HfN1.04 film are always higher than

that of the HfN1.17 film at a given Vb. The mechanisms

behind these phenomena are discussed in detail.

2 Experimental and computational details

d-HfNx films were simultaneously deposited on both

single-crystal Si (001) and optical glass substrates by

radio frequency (RF) reactive sputtering. A pure Hf

target was used in the mixed discharge gases of Ar and

N2. The distance between the target and substrate holder

was fixed at 55 mm, and the chamber was evacuated by

a turbomolecular pump to 4 9 10-4 Pa prior to depo-

sition. Before being introduced into the deposition

chamber, Si (001) and glass substrates were cleaned

ultrasonically in acetone, alcohol and distilled water

ether, consecutively. The flow rate of Ar and N2 was

controlled by independent mass flow controllers (MT50-

2 J). During the deposition, the argon flow rate, RF

power, work pressure and substrate temperature were

kept at 80 sccm, 150 W, 1.0 Pa and 200 �C, respec-

tively. The nitrogen flow rate was kept at 3.6 and

6.0 sccm, respectively. Different negative bias voltages

(Vb) ranging from -6 V (floating) to -200 V were

applied to the substrate during deposition.

X-ray diffraction (XRD) measurements were taken in

grazing-incidence (GIXRD) modes by a Bruker D8tools

X-ray diffractometer using Cu Ka as the incident radi-

ation. The stoichiometry x and valence band spectra of

the films were determined by X-ray photoelectron

spectroscopy (XPS) measurements (VG ESCALAB

MKII), in which a monochromatized Al Ka (1486.6 eV)

X-ray source was used, and Ar? cleaning procedure

lasting 180 s was applied to all samples prior to XPS

quantitative analysis to remove the adventitious carbon

and absorbed oxygen from the sample surface. The

microstructure and surface morphology of the films were

characterized using a high-resolution transmission elec-

tron microscope (HR-TEM, JEOL TEM-2010), field-

emission scanning electron microscope (FE-SEM, JEOL

JSM-6007F) and atomic force microscopy (AFM,

VEECO). Raman measurements were taken using a

785-nm line Ar? laser excitation with a laser current of

0.5 A and an accumulation time of 30 s under the fre-

quency range of 50–1200 cm-1. The dc conductivity r0was obtained by four-point probe electrical conductivity

measurements. The reflectivity spectrum in the range of

300–2200 nm was obtained by a PerkinElmer Lambda

900 Ultraviolet–Visible–near-infrared (UV–Vis–NIR)

spectrometer. The substrate has no influence on the

optical reflectivity measurements due to the high thick-

ness of films. The complex dielectric function, concen-

tration n and relaxation time s of free electrons were

obtained by fitting the reflectance spectra using a Drude–

Lorentz model [5]. The details of Drude–Lorentz fitting

of optical reflectivity spectra were reported elsewhere

[5]. The thicknesses of the films were determined by

using a Dektak3 surface profile measuring system.

Density functional theory (DFT) calculations were per-

formed to explore the effect of N and Hf vacancies on

the electronic structure of d-HfNx. The details of com-

putation methods were reported elsewhere [5].

776 Page 2 of 10 Z. Gu et al.

123

3 Results and discussion

3.1 Vb-induced changes in vacancy concentrations

Figure 1a shows a typical HRTEM lattice image of HfNx

films, wherein well-crystallized nanograins with a size of

*16 nm are uniformly distributed on the film surface and

a highly ordered crystal lattice is observed from the

enlarged image of nanograin (Fig. 1b). The measured

interplanar distances match well with the (111) and (200)

plane spacings of perfect rocksalt hafnium nitride, indi-

cating that the films have the rocksalt structure (d-HfNx).

Figure 1c displays a typical SEM cross-sectional image of

d-HfNx films with a thickness of *1.0 lm on Si (001)

substrate. There are no obvious pinholes or surface cracks

in the film. Figure 1d exhibits a typical AFM image of d-HfNx films with an area of 2.0 lm 9 2.0 lm, in which a

very flat and smooth surface with nanoscale island-like

features is observed. The HRTEM, SEM and AFM mea-

surements are in good agreement, proving that the present

d-HfNx films are smooth [root-mean-square (RMS)

roughness below 3.34 nm], compact and well-crystallized,

desirable for the following investigation on their structure,

electrical and optical properties.

Figure 2a, b plots the GIXRD spectra for near-stoi-

chiometric HfN1.04 and over-stoichiometric HfN1.17 films

deposited at different substrate bias voltages (Vb), wherein

the five diffraction peaks (2h & 33�, 39�, 57�, 68� and 72�)attributed to the (111), (200), (220), (311) and (222) in d-

HfNx phase appear simultaneously, indicating that two sets

of d-HfNx samples remain crystallized in the rocksalt

structure as the absolute value of Vb increases from 6

(floating) to 200 V. The results agree well with the

HRTEM measurements. The lattice parameters (a) ob-

tained from GIXRD spectra for d-HfNx films deposited at

different Vb are shown in Fig. 2c, d. interestingly enough,

as the absolute value of Vb increases from 6 (floating) to

200 V, the a for both sets of samples first increases and

then decreases and reaches maximum values at

Vb = -80 V (x & 1.04) and Vb = -120 V (x & 1.17),

respectively. The lattice constant a for the two sets of

samples exhibits the same evolution with Vb, pointing to

the changes in a may have the same microscopic origins. In

our previous investigation [31, 32], we proved that the

introduction of N and Hf vacancies in the d-HfNx films

caused a decrease in a due to lattice contraction. Hence, the

observed changes in a (Fig. 2c, d) may indicate that as Vb

is applied to the substrates the concentrations of vacancies

in HfN1.04 and HfN1.17 films first reduce and then increase,

and arrive at minimum values at Vb = -80 V and

-120 V, respectively. This is further confirmed by Raman

and dielectric function spectra of these samples.

Figure 3a, b plots the Raman spectra for HfN1.04 and

HfN1.17 films deposited at different Vb. The first-order

acoustic (A) and optical (O) peaks at *110 and

*500 cm-1 occur simultaneously in all spectra, indicating

that all of the films have the rocksalt phase containing N

and Hf vacancies [31, 33]. As the absolute value of Vb

Fig. 1 A typical HRTEM

lattice image of d-HfNx films

(a) and enlarged image (b) ofthe detailed atomic arrangement

of the white square framed

region in a, as well as typicalSEM cross-sectional (c) andAFM (d) images of the films

Ion-bombardment-induced reduction in vacancies and its enhanced effect on conductivity and… Page 3 of 10 776

123

increases from 6 (floating) to 200 V, both intensities of

A and O peaks first decrease and then increase (Fig. 3c, d),

and at the same time their positions shift to lower fre-

quencies firstly and then to higher ones (Fig. 3e, f). The

intensities and position of A and O peaks for HfN1.04 and

HfN1.17 samples reach minimum values at Vb = -80 V

and -120 V, respectively. These results consistently prove

that as the absolute value of Vb increases the concentrations

of N and Hf vacancies in HfN1.04 and HfN1.17 films first

decrease and then increase, and they reach a minimum

value at Vb = -80 V in the x & 1.04 film and at

Vb = -120 V in the x & 1.17 film, respectively [31, 33].

This agrees well with results from the GIXRD spectra.

The imaginary parts of the dielectric function for the

HfN1.04 and HfN1.17 films deposited at different Vb are

plotted in Fig. 4a–c, f–h, respectively. Each curve is

composed of four absorption bands, the most dominant

arises from the contribution of interband transition

absorption related to free electrons, and another three

absorption bands centered at *4.50,*3.64 and*1.00 eV

originate from the contribution of interband transition

absorption corresponding to bound electrons. To

investigate the relationship between the three absorption

bands and vacancy defects, the total and partial density of

states (TDOS and PDOS) for defect-free HfN, N-vacancy-

containing Hf8N7 and Hf-vacancy-containing Hf7N8 is

calculated and shown in Fig. 4d, i. Based on these calcu-

lated DOS, the observed absorption band centered at

*4.50 eV that appeared in all d-HfNx films is attributed to

the intrinsic interband transition from the N2p electrons to

the Fermi level. The absorption band at *3.64 eV is

assigned to the interband transition from defect states of Hf

vacancies to the unoccupied states located around the

Fermi level (calculated energy separation is *3.85 eV),

while the absorption band at *1.00 is attributed to the

interband transition from defect states of N vacancies to the

unoccupied states located around the Fermi level (calcu-

lated energy separation is *1.10 eV). The simultaneous

appearance of N- and Hf-vacancy-induced absorption

bands means that both N and Hf vacancies exist in d-HfNx

films. Furthermore, as the absolute value of Vb increases

from 6 (floating) to 200 V, both relative integrated inten-

sity of N- and Hf-vacancy-induced absorption bands first

decrease and then increase, and reach minimum values at

Fig. 2 GIXRD spectra for

HfN1.04 and HfN1.17 films

prepared at different Vb (a, b),as well as lattice parameters

calculated from GIXRD spectra

for two sets of samples (c, d)

776 Page 4 of 10 Z. Gu et al.

123

Vb = -80 V (x & 1.04) or Vb = -120 V (x & 1.17),

respectively (Fig. 4e, j). These results are in good agree-

ment with the above GIXRD and Raman analysis,

demonstrating that N and Hf vacancies are formed in the

films deposited at different Vb. As the absolute value of Vb

increases gradually, the total concentration of N and Hf

vacancies first decreases and then increases.

In summary, the consistence between GIXRD, Raman

and dielectric function spectra proves that too low or too

high Vb causes a significant increase in concentration of

vacancies, while the proper Vb is energetically most

favorable to reduction in concentration of vacancies. In

HfN1.04 and HfN1.17 films, concentration of N and Hf

vacancies reaches a minimum value at Vb = -80 and

-120 V, respectively.

3.2 Vb-induced evolution in electrical conductivity

and its mechanism

Figure 5a, b plots the electrical conductivity (r0) for d-HfNx films prepared at different Vb, in which as the abso-

lute value of Vb increases from 6 (floating) to 200 V, the r0

for both sets of samples first increases and then decreases,

and reaches maximum values at Vb = -80 V (x & 1.04)

and Vb = -120 V (x & 1.17), respectively. These results

indicate that the proper Vb applied to the substrate during

deposition can promote a significant increase in r0.According to the expression of electrical conductivity

r0 = ne2s/me [22], r0 is determined by the electron

relaxation time s and the free electron concentration n. As

the absolute value of Vb increases from 6 (floating) to

200 V, both n and s first increase and then decrease, and

arrive at maximum values at -80 V (x & 1.04) and

-120 V (x & 1.17), which are shown in Fig. 5c, d. The

same Vb dependence of n, s and r0 indicates that the

observed change in r0 attributes to the evolution in n and s.Why does n experience a first increase and sequent

decrease with Vb (Fig. 5c, d)? The typical XPS valence

band spectra of the d-HfN1.04 films prepared at Vb = -6 V

(floating) and -80 V are plotted in Fig. 5e. There are two

peaks in the spectra, a relatively weak peak located near the

Fermi level and another strong peak centered at 5.0 eV.

According to calculated DOS (Fig. 4d, i), these two peaks

arise from Hf 5d states and hybridized states of N 2p and

Fig. 3 Raman spectra for HfN1.04 and HfN1.17 films prepared at different Vb (a, b), respectively, as well as the integrated intensity (c, d) andpositions (e, f) of first-order acoustic (A) and optical (O) peaks for the same samples

Ion-bombardment-induced reduction in vacancies and its enhanced effect on conductivity and… Page 5 of 10 776

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Hf 5d, respectively. The occurrence of pronounced Hf

5d peaks near the Fermi level indicate that all d-HfN1.04

films are metal like, which is in good agreement with the

calculated DOS (Fig. 4d, i) and measured high r0 (Fig. 5a).It can be clearly seen from Fig. 5e that the film containing

a higher concentration of vacancies prepared at Vb = -6 V

has a much less DOS near the Fermi level than that with a

lower concentration of vacancies grown at Vb = -80 V.

This is in good agreement with the trend of measured

n (Fig. 5a), indicating that increased vacancies cause a

significant decrease in n. In order to understand the reason

that n of HfNx films decreases with increasing vacancies,

the difference between the spectra of two samples at

Vb = -6 V (floating) and Vb = -80 V is obtained and

shown in Fig. 5e. The original electrons of 0–2 eV near EF

in the film prepared at Vb = -80 V are transferred to the

new hybridized states around 2–5 eV in the film deposited

at Vb = -6 V, which is consistent with the calculated DOS

(Fig. 4i), indicating the formation of vacancies causes free

electrons near Fermi level being partly localized, and cre-

ating new Hfd- and Np-hybridized states. The similar

results of XPS valence band spectra are also found in the

HfN1.17 films (Fig. 5f). These valence band spectra of

HfN1.04 and HfN1.17 films are in good agreement with

theoretical calculation, proving that the first increase and

sequent decrease in n with Vb arises from a first reduction

and then increase in concentration of vacancies because

formation of vacancies causes free electrons near Fermi

level being partly localized.

s is the average time between two scattering events of the

free electrons [34], and in polycrystalline materials it is

directly related to point defects and grain boundaries, which

is determined by the formula 1/s & 1/sd ?1/sg [22], wheresd and sg are the relaxation times corresponding to point

defects and grain boundary scattering, respectively. In order

to investigate whether s is determined by point defects or

grain boundaries, the electron mean free path (le) is calcu-

lated by le = svF [32], where vF is the velocity at the Fermi

surface defined by vF & (�h/m*)(3p2N)1/3 according to the

free electron model [32]. Figure 5g displays le for HfN1.04

films deposited at different Vb, wherein the average grain

sizes are also given for comparison. It can be clearly seen that

le (0.3–1.2 nm) is much smaller than the grain size

(8–37 nm).Moreover, the trend of le is totally different from

variation of the grain size with Vb. These results clearly

indicate s and le in d-HfNx films are not significantly

Fig. 4 Imaginary parts of the dielectric function for HfN1.04 and

HfN1.17 films prepared at different Vb (a–c), (f–h). Density of states

(DOS) and partial density of states (PDOS) for d-HfNx containing N

vacancies (Hf8N7) and Hf vacancies (Hf7N8) as well as the perfect

HfN structure obtained by first-principles calculations (d, i). The ratioof integrated intensity of N vacancy (VN) and Hf vacancy (VHf)

interband absorption bands to that of intrinsic interband absorption

bands for HfN1.04 and HfN1.17 films prepared at different Vb (e, j)

776 Page 6 of 10 Z. Gu et al.

123

influenced by grain boundary scattering. Considering the

contribution of grain boundary scattering to s is rather small,

we believe that the electron scattering processes that involve

point defects control s. This conclusion can be confirmed by

the strong dependence between the concentration of vacan-

cies and le whenever in HfN1.04 or HfN1.17 films. Hence, the

observed increase and then decrease in s and le arise from a

reduction and sequent increase in concentration of vacancies

because in polycrystalline d-HfNx films the main electron

scattering sites are vacancies.

As such, the HfN1.04 film deposited at Vb = -80 V and

HfN1.17 film grown at Vb = -120 V exhibit the highest r0because they contain the lowest concentrationofvacancies and

hence thehighestn and longest s. Furthermore, it canbeclearly

seen from Fig. 5a, b that r0 (1 9 104–30 9 104 Sm-1) of

HfN1.17 films are always lower than that (5 9 104–

55 9 104 Sm-1) of HfN1.04 films for a given Vb, which is

attributed to more Hf vacancies formed in HfN1.17 films

because of larger composition deviation from N:Hf = 1:1.

3.3 Vb-induced evolution in optical reflectivity

and its mechanism

Figure 6a, b displays the reflectance spectra of HfN1.04 and

HfN1.17 films deposited at different Vb. From Fig. 6c, d, as

the absolute value of Vb increases from 6 (floating) to

200 V, the cutoff wavelengths at the minimum reflectance

(kRmin) for both sets of samples first decrease and then

increase, and reach minimum values at Vb = -80 V

(x & 1.04) and Vb = -120 V (x & 1.17). This is in good

agreement with the trend of n (Fig. 5c, d), because a high

n leads to a high plasma energy (Ep) according to the

formula Ep � n1/2, and thus a small kRmin [5]. The average

reflectance in the infrared range of 1000–2200 nm (RIR)

first increases and then decreases with Vb, and arrives

maximum values at Vb = -80 V (x & 1.04) and

Vb = -120 V (x & 1.17), respectively (Fig. 6c, d). The

regular variation in the cutoff wavelengths and reflectance

also leads to the gradual changes in color and brightness of

d-HfNx films, which are clearly displayed in inset of

Fig. 6a, b. The changes in RIR with Vb indicate that the

proper Vb applied to the substrate during deposition can

promote a significant increase in RIR.

To explore the reason why the Vb significantly affects

the RIR, we have fitted the reflectance spectra of d-HfNx

films by the Drude–Lorentz model, wherein a series of

spectral fittings have been performed by changing the

number of Lorentz oscillators considered in the fitting

functions. The best fitting has been obtained by using one

Drude part and three Lorentz oscillators in the fitting

Fig. 5 Electrical conductivity r0 (a, b) and concentration of free

electrons n and relaxation time s (c, d) for HfN1.04 and HfN1.17 films

prepared at different Vb. The typical valence band spectra normalized

to the intensity of the respective Hf 4f core lines (e, f). The electron

mean free path (le) and average grain size for the samples prepared at

different Vb (g, h)

Ion-bombardment-induced reduction in vacancies and its enhanced effect on conductivity and… Page 7 of 10 776

123

process. It is found that the shape and peak position of

reflectivity spectra obtained by fitting are in good agree-

ment with the corresponding experimental spectra (see

Supporting Information Figs. A1 and A2), indicating that

the reflectivity spectra for d-HfNx films can be well fitted

by the Drude–Lorentz model and the observed evolution in

RIR with Vb is totally attributed to the changes in the

Lorentz and Drude parameters.

To find out the main factors contributing to RIR of the

films deposited at different Vb, we have investigated the

role of each Drude and Lorentz parameter in changing the

reflectivity spectra by simulating the reflectivity spectra at

different parameters (see Supporting Information Figs. A3

and A4). It is clearly seen that all of Lorentz parameters

almost have no influence on RIR, indicating that the chan-

ges in the Lorentz parameters are not responsible for the

evolution in RIR. In contrast, the changes in n and s (Drudeparameters) cause RIR first rising and then falling, which is

consistent with the experimentally observed trend of RIR

with Vb (Fig. 6c–d), proving that n and s are the dominant

factors contributing to RIR. Hence, the first increase and

then decrease in RIR with Vb is attributed to the first

increase and sequent decrease in n and s.In order to understand why the RIR of d-HfNx film is

more dependent on the changes in Drude parameters (n and

s), we have compared the contribution between Drude

parameters and three Lorentz oscillators to the dielectric

function spectra in the infrared region of 1000–2200 nm,

which is plotted in Fig. 6e–h. It can be seen that Drude

parameters have a much more contribution to both real and

imaginary parts of the dielectric function than three Lor-

entz oscillators, which means that the dielectric function is

more sensitive to the variation in Drude parameters.

Therefore, the changes in n and s with Vb more easily

change the dielectric function and thus more significantly

influence RIR of d-HfNx film than three Lorentz oscillators.

This explains why there is a strong dependence between

Drude parameters (n and s) and RIR (see supporting

information Figs. A3 and A4).

In summary, the HfN1.04 film deposited at Vb = -80 V

and HfN1.17 film grown at Vb = -120 V exhibits the

highest RIR because they contain the lowest concentration

of vacancies and hence the highest n and longest s. Fur-thermore, it can be clearly seen from Fig. 6c, d that RIR

(30–70 %)of d-HfN1.17 films is always lower than RIR

(49–80 %)of d-HfN1.04 films (x & 1.04) at a given Vb,

which is due to more Hf vacancies formed in HfN1.17 films

because of larger composition deviation from N:Hf = 1:1.

4 Conclusions

1. Ion bombardment via applying the negative bias volt-

age (Vb) can cause reduction or more formation of Hf

and N vacancies in hafnium nitride films. Upon

reduction, increase in concentration of free electrons

Fig. 6 Reflectance spectra and photographs (a, b), and the average

reflectance in the infrared range of 1000–2200 nm (RIR) and the

cutoff wavelengths at the minimum reflectance (kmin) (c, d) for

HfN1.04 and HfN1.17 films prepared at different Vb. The typical real

parts e1 (e, f) and imaginary parts e2 (g, h) of the dielectric function in

the infrared range of 0.56–1.24 eV (1000–2200 nm) for the d-HfNx

samples

776 Page 8 of 10 Z. Gu et al.

123

and increase in electron relaxation time take place,

causing increase in electrical conductivity and optical

reflectance. This critical bias is -80 V for near-stoi-

chiometric HfN1.04 film and -120 V for over-stoi-

chiometric HfN1.17 film with corresponding electrical

conductivity of 5.5 9 105 Sm-1 (x & 1.04) and

3.0 9 105 Sm-1 (x & 1.17) and infrared reflectance

of 80 % (x & 1.04) and 70 % (x & 1.17). Further

increase in substrate bias results in formation of more

Hf and N vacancies, which in turn causes decrease in

concentration of free electrons and decrease in electron

relaxation time, resulting in both electrical conductiv-

ity and optical reflectance to deteriorate.

2. HfN1.17 films exhibit always much lower electrical

conductivity and infrared reflectance than HfN1.04

films for a given Vb, owing to more hafnium vacancies

because of larger composition deviation from HfN

exact stoichiometry (N:Hf = 1:1).

3. This study discovers that the electrical conductivity and

optical reflectivity of hafnium nitride films can be

enhanced via ion bombardment, which is useful in

improving the ultimate performance of transition metal

nitrides for important technological applications such as

electrical back-contacts for light-emitting devices, thin

film resistors and conductive diffusion barrier layers.

Acknowledgments The authors gratefully acknowledge the financial

support from National Natural Science Foundation of China (Grant

Nos. 51572104, 51102110 and 51372095), National Major Project for

Research on Scientific Instruments of China (2012YQ240264),

Technology Development Project (2015220101000836), Program for

studying abroad of China Scholarship Council. Also, the authors are

grateful for the support from the Academic Research Fund of Sin-

gapore (Tier 1, RG187/14).

References

1. K. Dejun, G. Haoyuan, Analysis of structure and bonding

strength of AlTiN coatings by cathodic ion plating. Appl. Phys. A

119, 309–316 (2015)

2. M.M. Larijani, M. Kiani, E. Jafari-Khamse, V. Fathollahi,

Temperature dependence of the optical properties of ion-beam

sputtered ZrN films. Appl. Phys. A 117, 1179–1183 (2014)

3. R. Gupta, S. Soni, D.M. Phase, Improvement of oxidation resis-

tance of TiCN films prepared by laser alloying. Appl. Phys. A

118, 191–196 (2015)

4. P. Pedrosa, D. Machado, J. Borges, M.S. Rodrigues, E. Alves,

N.P. Barradas, N. Martin, M. Evaristo, A. Cavaleiro, C. Fonseca,

F. Vaz, Agy:TiNx thin films for dry biopotential electrodes: the

effect of composition and structural changes on the electrical and

mechanical behaviours. Appl. Phys. A 119, 169–178 (2015)

5. C. Hu, Z. Gu, J. Wang, K. Zhang, X. Zhang, M. Li, S. Zhang, X.

Fan, W. Zheng, Nature of tunable optical reflectivity of rocksalt

hafnium nitride films. J. Phys. Chem. C 118, 20511–20520 (2014)6. G.V. Naik, V.M. Shalaev, A. Boltasseva, Alternative plasmonic

materials: beyond gold and silver. Adv. Mater. 25, 3264–3294(2013)

7. G. Abadias, V.I. Ivashchenko, L. Belliard, P. Djemia, Structure,

phase stability and elastic properties in the Ti1-xZrxN thin-film

system: experimental and computational studies. Acta Mater. 60,5601–5614 (2012)

8. D. Craciun, G. Socol, G. Dorcioman, S. Niculaie, G. Bourne, J.

Zhang, E. Lambers, K. Siebein, V. Craciun, Wear resistance of

ZrC/TiN and ZrC/ZrN thin multilayer grown by pulsed laser

deposition. Appl. Phys. A 110, 717–722 (2013)

9. L. Tsetseris, N. Kalfagiannis, S. Logothetidis, S.T. Pantelides,

Role of N defects on thermally induced atomic-scale structural

changes in transition-metal nitrides. Phys. Rev. Lett. 99, 125503(2007)

10. G. Abadias, M.B. Kanoun, S. Goumri-Said, L. Koutsokeras, S.N.

Dub, P. Djemia, Electronic structure and mechanical properties of

ternary ZrTaN alloys studied by ab initio calculations and thin-

film growth experiments. Phys. Rev. B 90, 144107 (2014)

11. E. Camps, L. Escobar-Alarcon, I. Camps, S. Muhl, M. Flores,

Tribological characterization of TiCN coatings deposited by two

crossed laser ablation plasma beams. Appl. Phys. A 110, 957–961(2013)

12. K. Vasu, G.M. Gopikrishnan, M. Ghanashyam Krishna, A. Pad-

manabhan, Optical reflectance, dielectric functions and phonon-

vibrational modes of reactively sputtered Nb-substituted TiN thin

films. Appl. Phys. A 108, 993–1000 (2012)

13. S. Dong, X. Chen, X. Zhang, G. Cui, Nanostructured transition

metal nitrides for energy storage and fuel cells. Coord. Chem.

Rev. 257, 1946–1956 (2013)

14. Z. Gu, C. Hu, H. Huang, S. Zhang, X. Fan, X. Wang, W. Zheng,

Identification and thermodynamic mechanism of the phase tran-

sition in hafnium nitride films. Acta Mater. 90, 59–68 (2015)

15. L. Tsetseris, N. Kalfagiannis, S. Logothetidis, S.T. Pantelides,

Trapping and release of impurities in TiN: a first-principles study.

Phys. Rev. B 78, 094111 (2008)

16. H. Van Bui, A.W. Groenland, A.A.I. Aarnink, R.A.M. Wolters, J.

Schmitz, A.Y. Kovalgin, Growth kinetics and oxidation mecha-

nism of ALD TiN thin films monitored by in situ spectroscopic

ellipsometry. J. Electrochem. Soc. 158, H214–H220 (2011)

17. A. Salamat, A.L. Hector, P. Kroll, P.F. McMillan, Nitrogen-rich

transition metal nitrides. Coord. Chem. Rev. 257, 2063–2072

(2013)

18. L. Yu, C. Stampfl, D. Marshall, T. Eshrich, V. Narayanan, J.M.

Rowell, N. Newman, A.J. Freeman, Mechanism and control of

the metal-to-insulator transition in rocksalt tantalum nitride.

Phys. Rev. B 65, 245110 (2002)

19. C. Stampfl, A.J. Freeman, Stable and metastable structures of the

multiphase tantalum nitride system. Phys. Rev. B 71, 024111(2005)

20. S. Mahieu, W.P. Leroy, K. Van Aeken, M. Wolter, J. Colaux, S.

Lucas, G. Abadias, P. Matthys, D. Depla, Sputter deposited

transition metal nitrides as back electrode for CIGS solar cells.

Sol. Energy 85, 538–544 (2011)

21. R.S. Ningthoujam, N.S. Gajbhiye, Synthesis, electron transport

properties of transition metal nitrides and applications. Prog.

Mater Sci. 70, 50–154 (2015)

22. D. Valerini, M.A. Signore, A. Rizzo, L. Tapfer, Optical function

evolution of ion-assisted ZrN films deposited by sputtering.

J. Appl. Phys. 108, 083536 (2010)

23. P. Patsalas, S. Logothetidis, Optical, electronic, and transport

properties of nanocrystalline titanium nitride thin films. J. Appl.

Phys. 90, 4725–4734 (2001)

24. T. Migita, R. Kamei, T. Tanaka, K. Kawabata, Effect of dc bias

on the compositional ratio of WNx thin films prepared by rf-dc

coupled magnetron sputtering. Appl. Surf. Sci. 169–170, 362–365(2001)

25. C. Mitterer, P.H. Mayrhofer, W. Waldhauser, E. Kelesoglu, P.

Losbichler, The influence of the ion bombardment on the optical

Ion-bombardment-induced reduction in vacancies and its enhanced effect on conductivity and… Page 9 of 10 776

123

properties of TiNx and ZrNx coatings. Surf. Coat. Technol.

108–109, 230–235 (1998)

26. M. Del Re, R. Gouttebaron, J.P. Dauchot, M. Hecq, Study of the

optical properties of AlN/ZrN/AlN low-e coating. Surf. Coat.

Technol. 180–181, 488–495 (2004)

27. W.-J. Chou, G.-P. Yu, J.-H. Huang, Bias effect of ion-plated zirco-

nium nitride film on Si(100). Thin Solid Films 405, 162–169 (2002)28. A. Rizzo, M.A. Signore, D. Valerini, D. Altamura, A. Cappello,

L. Tapfer, A study of suppression effect of oxygen contamination

by bias voltage in reactively sputtered ZrN films. Surf. Coat.

Technol. 206, 2711–2718 (2012)

29. L. Tsetseris, N. Kalfagiannis, S. Logothetidis, S.T. Pantelides,

Structure and interaction of point defects in transition-metal

nitrides. Phys. Rev. B 76, 224107 (2007)

30. C. Stampfl, A.J. Freeman, Metallic to insulating nature of TaNx:

role of Ta and N vacancies. Phys. Rev. B 67, 064108 (2003)

31. Z. Gu, C. Hu, X. Fan, L. Xu, M. Wen, Q. Meng, L. Zhao, X.

Zheng, W. Zheng, On the nature of point defect and its effect on

electronic structure of rocksalt hafnium nitride films. Acta Mater.

81, 315–325 (2014)

32. C. Hu, X. Zhang, Z. Gu, H. Huang, S. Zhang, X. Fan, W. Zhang,

Q. Wei, W. Zheng, Negative effect of vacancies on cubic sym-

metry, hardness and conductivity in hafnium nitride films. Scr.

Mater. 108, 141–146 (2015)

33. M. Stoehr, H.-S. Seo, I. Petrov, J.E. Greene, Effect of off stoi-

chiometry on Raman scattering from epitaxial and polycrystalline

HfNx (0.85 B x B 1.50) grown on MgO (001). J. Appl. Phys.

104, 033507 (2008)

34. S.D. Yoo, K.D. Kwack, Theoretical calculation of electron

mobility in HgCdTe. J. Appl. Phys. 81, 719–725 (1997)

776 Page 10 of 10 Z. Gu et al.

123