Cao Z. (Ed.) Thin Film Growth: Physics, materials science and applications

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198 Thin ﬁlm growth

can say with condence that the Cu phases emerge from the decomposed

stoichiometrical Cu

N, revealed by XRD, only at a temperature of ~350°C

by annealing under the protection of nitrogen at a pressure of 0.9 Pa.

8.3 Nitrogen re-emission

The thermal instability of copper nitrides constitutes essential difculties

for their synthesis and structural characterization. By reactive magnetron

sputtering deposition of copper nitride lms, the effect of nitrogen re-emission

on the lm morphology may be directly observable, from which the effect

on the growth process and on the microscopic structure of deposits can be

inferred. The investigation of this topic is also helpful in understanding the

failure mechanism of copper nitrides in usage, which may cause serious

instability problems in practical applications.

To demonstrate the effect of nitrogen re-emission, copper nitride thin lms

were grown with mixed nitrogen (of 60% to 80% in ow rate) and argon as

working gas, maintained at a pressure of 1.1–1.3 Pa. The power supply was

set at 150 W. Under the given conditions, the deposits are substoichiometric

with copper contents exclusively below 77.6%. Under SEM, the lms grown

for 15 minutes, which are roughly 0.3 mm thick, display a morphology with

isolated round blisters (Fig. 8.10). These blisters are typically about 10

microns in dimension. The coverage by these protruding features is about

20%. In lms deposited for 30 minutes, the protruding features turned into

Intensity (arb. units)

20 30 40 50 60

2q (deg)

N(001)

Cu(111)

Cu(002)

N(002)

420°C

350°C

300°C

220°C

8.9 XRD patterns of the Cu

N ﬁlms after annealing at temperatures

varying from 220°C to 420°C.

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199Thin ﬁlm growth for thermally unstable noble-metal nitrides

a ramied rosette-like structure. Such copper nitride rosettes are distributed

homogeneously across the sample surface, and show a striking size and

shape uniformity, with a lateral dimension around 23 microns (Fig. 8.10).

By close inspection of the individual rosettes, it is observed that the rays

bifurcate at an average angle of ~74°, slightly larger than the wedge angle

of 71.53° for a tetrahedron. In some well-developed ones, vefold symmetry

can be conrmed at the center of the rosettes, though not as strict as in the

geometrical sense.

At rst glance, this morphology reminds us of island formation as frequently

encountered in growing crystalline lms, which can be generally modeled with

(a)

20 µm

(b)

100 µm

8.10 Effect of nitrogen re-emission on the morphology of copper

nitride ﬁlms. A 0.3 mm thick ﬁlm exhibits distinct round blisters (a),

while in a ﬁlm grew to ~0.6 mm thick, rosettes with bifurcated rays

were observed (b). The rosettes are generally quintilobed, often

ﬁvefold symmetry can be identiﬁed at the centers.

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200 Thin ﬁlm growth

an atomistic view. But we found that this is not the case. Both the smaller

round caps and the larger rosettes are hollow, as conrmed by the presence

of some cracked entities. When probing through the crack of a broken rosette,

the energy-dispersive X-ray spectrometry detected a dominantly strong signal

only for silicon and a weak signal for copper – the nitrogen is completely

absent. This is to say that the protruding features are due to a complete local

detachment of the lm; and in that process some tiny amount of copper is

left behind on the Si substrate (Ji et al. 2005, 2006a). Nevertheless, such

protruding structures cannot have developed from a previously well-deposited

coating due to the accumulated stress. This point can be claried by briey

examining the necessary area expansion of the material for the formation

of such hollow features. As a conservative estimation, consider a round

cap 20 mm in lateral dimension and 1 mm high. This corresponds to an area

expansion of ~5.0% with regard to the underlying circular base. Such a large

area expansion coefcient is prohibitive for rigid materials such as the ionic

copper nitride. Moreover, such a morphology is formed in the growth stage,

post-growth degradation of an originally at lm due to nitrogen re-emission

leaves behind a at, but Cu-rich surface. The aging of lms with a compact

morphology in the ambient will not invoke any rosette structure.

The puzzle of the enormously expanded area of the rosettes with regard to

the underlying base becomes immediately resolved with the aid of scanning

electron micrographs at an enlarged magnication. We see that the deposits

with rosette structures are composed of distinct crystallites, around 45 nm

in dimension, just like a compact lm, but the SEM images taken directly

on a rosette reveals the peculiarity in the manner of crystallite stacking. The

surface of the rosette structures displays ragged steps and terraces, and it is

thinner than the at portion of the deposits. The typical width of a terrace

is about 100 nm (Fig. 8.11). Such a morphology has never been reported,

to the best knowledge of the author. We are inspired to speculate that the

crystallites in a rosette have experienced a rearrangement process. The area

expansion in due course of relief formation is sustained through the gliding

of nanocrystals, which is facilitated when the nanocrystals are clothed with

the amorphized Cu-terminated {111}-planes; and the ongoing fast growth

prevents the lm from cracking.

In order to fully understand the formation mechanism for such protruding

features, we take a close look at the morphological prole of the rosettes.

They have the shape of a starsh with rays, but these rays generally do not

meet at the same point in the central disk. Occasionally, a symmetrically

developed rosette could be found, in which a perfect pentagram is discernible

at the center, as shown in Fig. 8.12(a). The rays show further bifurcations,

and the offshoots extend generally at an angle of about 74°, a little larger

than 72° as required by vefold rotational symmetry, which in turn is a little

larger than the wedge angle of a tetrahedron (a = 70.53°). This cannot be

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201Thin ﬁlm growth for thermally unstable noble-metal nitrides

a coincidence, bearing in mind that the tetrahedral crystallites do have the

tendency to juxtapose at an angle of about 72° with a [110] orientation in

order to ll the space, giving rise to a loosely vefold symmetry. Fivefold

symmetry evolving from this construction mechanism has been conrmed

in nanostructures of cubic metals such as Au, Ag and Cu.

As we know, the {111}-planes of the cubic Cu

N lattice comprise

exclusively either Cu atoms or N atoms. When nitrogen re-emits from a Cu

crystallite, it leaves behind crystallites with a soft clothing of amorphous

Cu layer in a habit, not necessarily tetrahedral, but with facets joining at the

wedge angle (a ª 70.53°). The clothing of Cu

N crystallites by Cu {111}-

planes also builds a particular electrical conductance path, as veried from

the percolation transition in the variation of its electrical resistivity versus

the Cu content in the lm. When driven by a sufciently large stress, here

initiated by the re-emitted nitrogen gas from beneath, the nanocrystals will

undergo a gliding motion against each other along the {111}-planes, and

conglomeration of a large quantity of such nanocrystals with facets joining

at a ª 70.35° results in vefold symmetry at larger scales. Displayed in Fig.

8.12(b) is a piece of Cu

N crystal in a specimen prepared with mechanical

milling for transmission electron microscopic investigation – it is pentagonal

with an edge length of 660 nm.

Now the picture becomes clear. An ‘orogenic movement’ model can be

formulated to illustrate the formation of relief rosettes in a growing Cu

thin lm (Fig. 8.5). Copper nitride keeps decomposing while lm growth

proceeds. The re-emitted nitrogen atoms recombine into molecules and

together with entrapped nitrogen they agglomerate to form gas bubbles at the

200 nm

8.11 Scanning electron micrographs at a large magniﬁcation of the

rosette. Typical crystallite size is 45 nm. The dashed lines are plotted

to guide the eye.

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202 Thin ﬁlm growth

lm–substrate interface by virtue of a large lateral mobility, since solubility

of nitrogen in both Si and Cu

N is low. Adhesion to the local substrate is

undermined where a gas bubble forms, and nally detachment occurs (Fig.

8.5(b)). By a delicate balance between expansion via nanocrystal glide and

fast growth, the total area and the lateral dimension of the delaminated

(a)

25 µm

(b)

300 nm

8.12 Fivefoldness in the microstructure of a copper nitride deposit.

(a) A rosette displaying perfect ﬁvefold symmetry at the center; (b)

a pentagonal assembly of the Cu

N nanocrystals under transmission

electron microscope.

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203Thin ﬁlm growth for thermally unstable noble-metal nitrides

part grow in pace without cracking. As can be expected, at reduced growth

rates a compact morphology of the lm ensues. Gliding motion of the

crystallites into ragged steps and terraces – similar to the movement of

rocks in mountain building, i.e. the orogenic movement – provides the path

for a complete strain relaxation; this explains the negligible residual stress

in the sample despite the presence of the relief features, as concluded from

the transmission electron microscopic and X-ray diffraction data about the

lattice constant (Fig. 8.1). The possibility to glide along the Cu-terminated

{111}-planes for the pyramidal Cu

N crystallites is prerequisite for the

formation of the relief morphology; this also explains its absence in the lms

of nearly stoichiometrical or heavily Cu-rich samples. As to the quintilobed

structure of the rosettes that ramify at a branching-off angle of ~74°, they

are the consequence of the preferred packing and rearrangement driven by

an anisotropic stress of about two million {111}-faceted Cu

N nanocrystals

(Ji et al. 2007). Unlike in the nanorods of cubic metals where the core is

vefold twin crystals textured along the [110]-axis, the vefold symmetry

here is quite ‘soft’ in the sense that it arises rather as an integral effect at

large scales.

8.4 Doping of Cu

N by co-sputtering

The structure of Cu

N is, in itself, rather interesting. It is an open structure

that can, theoretically, accommodate one more metal ion in each unit cell.

The incorporation of an excess metal atom can signicantly alter the band

structure of the material that continuous semiconducting-to-(semi)metallic

transition can be expected, through which various interesting phenomena

can be measured in the ternary compounds (Lovett 1977). Here, through

the co-sputtering of Cu and Pd targets, we demonstrate the deposition of

ternary Cu

NPd

thin lms, where a vanishing temperature coefcient of

resistivity (TCR) in the range from 240 K down to 5 K could be measured

in Cu

NPd

0.238

(Ji et al. 2006b). Such a phenomenon has not been found in

any other single solids.

The inclusion of another metal atom in each unit cell of the Cu

N lattice

results in compounds like Cu

N or the ternary Cu

NPd (Hahn and Weber 1996).

Remarkably, Cu

NPd is a semimetal due to the intersection of the energy

bands at the Fermi level. Consequently, a semiconducting-to-semimetallic

transition is anticipated in the off-stoichiometric Cu

NPd

when ‘x’ increases

continuously from zero to unity.

We prepared thin lms of Cu

N and Cu

NPd

(0 < x < 0.350) on the Si

(001) wafers, a low substrate temperature (~100°C) in combination with

a small power supply (<150 W) and a high working pressure (>0.7 Pa) is

applied. The interposition of metal atoms at the cell centers of the cubic Cu

lattice severely deteriorates the covalent bonding in the resulting materials

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204 Thin ﬁlm growth

such that the compounds Cu

N and Cu

NPd only exist in theory. In fact,

the largest reported value of ‘x’ never exceeds 0.989 in Cu

NPd

whereby

the compounds exist in ne powders (Jacobs and Zachwieja 1991).

found that those Cu

NPd

thin lm samples with x > 0.5 show a very poor

crystallinity; therefore we restrict our discussion to the samples with 0.0 ≤

x ≤ 0.350 which are well crystallized as conrmed by both X-ray diffraction

pattern (Fig. 8.13) and transmission electron micrograph (Fig. 8.14). For

samples containing only a tiny amount of Pd, e.g., Cu

NPd

0.048

, the X-ray

diffraction pattern remains unchanged, like that of the host Cu

N lattice

showing only the (001) and (002) reections. When the Pd concentration

is large enough (x ≥ 0.175), the (001) reection was replaced by the strong

(111) reection which deteriorates both in intensity and in spectral prole.

We can say that the Cu

NPd

thin lms here concerned with x < 0.35 are

crystalline. The addition of Pd atoms results in a slightly enlarged lattice

constant which increases from a = 0.383 nm for the pure Cu

N to a = 0.385

nm for Cu

NPd

0.175

(Fig. 8.14). The Cu

NPd

samples obviously have a better

electrical conductivity which in turn results in an improved image quality for

the transmission electron micrograph (cf. Fig. 8.7). The fact that the Cu

NPd

is metallic at sufciently large ‘x’ values, while Cu

N is semiconducting,

may suggest that the interposition of Pd atoms alter the energy bands near

the Fermi surface.

Intensity (arb. units)

Si (002)

(111)

(001)

(002)

x = 0.349

x = 0.238

x = 0.175

x = 0.048

20 30 40 50 60

2q (deg)

8.13 X-ray diffraction patterns of Cu

NPd

with varying ‘x’ values.

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205Thin ﬁlm growth for thermally unstable noble-metal nitrides

Figure 8.15 displays the temperature dependence of the electrical resisitivty

for the nearly stoichiometric Cu

N and for the different ternary Cu

NPd

compounds. Upon cooling from room temperature, the electrical resistivity

in the intrinsic Cu

N lm rises rapidly (Fig. 8.15(a)). This observation is

consistent with the wide band-gap nature of this semiconducting material. The

insertion of Pd atoms confers Cu

NPd

an improved electrical conductivity

owing to the narrowed band-gap. Even with a tiny amount of Pd insertion

as in the samples with x = 0.048 and x = 0.071, the rapid onset of resistivity

upon cooling is shifted to below 50 K, indicating a narrowed band-gap for

the current samples (Fig. 8.15(b) and (c)). A further addition of Pd atoms

initiates the semiconducting-to-semimetallic transition such that the TCR

changes its sign. For Cu

NPd

0.175

, its electrical resistivity rises roughly linearly

with decreasing temperature in a range as wide as 270 K, and the TCR is

–0.00039 K

–1

(Fig. 8.15(d)). Taking T = 280 K as the reference point, the

resistance over such a broad temperature range changes only by ~5.0%. The

resistivity for Cu

NPd

0.349

also displays a good linearity from 50 K to 280

K, but now the TCR is 0.00117 K

–1

(Fig. 8.15(f)), indicating unambiguously

that now the compound has changed to be metallic. In between these two

cases of opposite tendencies for the variation of the electrical resistivity with

temperature, we may expect a vanishing TCR over a considerable temperature

range at some ‘x’ value. This occurs in the sample Cu

NPd

0.238

. From 240

K down to 5 K, the TCR is nearly zero (< 3.0 ¥ 10

–6

/K), and the relative

variance of the electrical resistivity in this temperature range measures as

small as 1.42 ¥ 10

–5

(Fig. 8.15(e)).

d = 0.3855 nm

2 nm

8.14 Transmission electron micrograph for the sample Cu

NPd

0.175

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206 Thin ﬁlm growth

The temperature dependence of the electrical conductivity for solids is

susceptible to many inuencing factors; no hard-and-fast rules are available.

For a semiconductor, the electrical conductivity can be formally written as

s = e(n

+ n

), where e and h denote electron and hole, while n and

m stand for the carrier density and carrier mobility, respectively. Both the

carrier mobility and the carrier density, or the number of the carriers, are

temperature dependent,

yet the temperature dependence of the electrical

conductivity is dominated by that of the latter, since the mobility changes

only in a mild way with the lattice temperature while the number of carriers

Resistivity (10

–1

Wm)

3.0

2.0

1.0

0.0

160 200 240 280

Temeprature (K)

(a)

Resistivity (Wm)

3.0

2.0

1.0

0.0

0 50 100 150 200 250 300

Temperature (K)

(b)

8.15 Temperature (in K) dependence of the electrical resistivity

measured in crystalline Cu

NPd

thin ﬁlms. (a) x = 0.0; (b) x = 0.048;

to (f), the electrical resistivity at room temperature is reduced by four

orders of magnitude.

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207Thin ﬁlm growth for thermally unstable noble-metal nitrides

Resistivity (10

–2

Wm)

1.5

1.0

0.5

0.0

0 50 100 150 200 250 300

Temperature (K)

(c)

Resistivity (10

–5

Wm)

2.08

2.04

2.00

1.96

0 50 100 150 200 250 300

Temperature (K)

(d)

Resistivity (10

–5

Wm)

1.14

1.12

1.10

1.08

0 50 100 150 200 250 300

Temperature (K)

(e)

8.15 Continued.

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