Structure and phase separation of Ag–Cu alloy thin films

www.actamat-journals.com

Acta Materialia 55 (2007) 1617–1628

Structure and phase separation of Ag–Cu alloy thin films

Hao Chen, Jian-Min Zuo *

Department of Materials Science and Engineering and Frederick Seitz Materials Research Laboratory,

University of Illinois at Urbana-Champaign, IL 61801, USA

Received 30 September 2006; received in revised form 18 October 2006; accepted 20 October 2006Available online 27 December 2006

Abstract

We report an in situ electron diffraction study and microscopy observation of phase separation in Ag–Cu solution alloy thin films ofdifferent compositions and film thicknesses. The results show that the as-deposited films consist of nanocrystalline Ag- and Cu-richphases of a few nanometers in size. Upon annealing, two stages of thin-film transitions are observed. In the first stage, the growth ofCu crystallites is responsible for phase separation in Ag50Cu50 or Ag-rich alloy. For the Cu-rich samples, the phase separation processis much slower. The second stage transition involves thin film dewetting. A combination of Z-contrast scanning transmission electronmicroscopy imaging and electron diffraction reveal grain growth and phase separation between Ag and Cu in both stages. Atomic forcemicroscopy results show the surface morphology of thin films only changes significantly at the late stage of phase separation when thethin film dewets.� 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Cu–alloys; Thin-films; Spinodal decomposition; Scanning transmission electron microscopy; Electron diffraction

1. Introduction

The eutectic Ag–Cu system is a glass-forming binaryalloy system with a positive heat of mixing in both solidand liquid states (the heat of mixing is 3 kJ mol�1 for thecomposition of Ag60Cu40 according to [1]). The other bin-ary alloy systems include Co–Cu, Cu–Zr, Cu–Fe, Cu–Ta,Cu–V and Cu–W [2]. Ag and Cu have a large miscibilitygap and a relative atomic size difference of 12%. The alloysystem has been studied extensively for its amorphous andnanocrystalline structure. At equilibrium, the maximumsolubility of Cu in Ag is 8.27 at.% Cu at the eutectic tem-perature of 779.1 �C, and Ag in Cu is 7.47 at.% [3].

Solid solution, amorphous or nanocrystalline Ag–Cualloys can be prepared by several different methods, includ-ing splat quenching [4,5], vacuum evaporation [6–9], ionbeam mixing [10,11], solid-state reaction in metal multilay-ers [12] and mechanical alloying [2,13–16]. Previous studies

1359-6454/$30.00 � 2006 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2006.10.036

* Corresponding author.E-mail address: [email protected] (J.-M. Zuo).

have shown that compositional homogeneous solid solu-tions are formed at a very high quench rate, achieved byrapid cooling to liquid nitrogen temperature. The structureof as-formed Ag–Cu alloys was characterized by X-ray dif-fraction [17] and transmission electron microscopy [8]. ForAg–Cu, the solid solution phase shows a single set of broadfcc peaks, which sit in between the corresponding bulk Agand Cu peaks. Wagner etc. performed radial distributionfunction (RDF) analysis of Ag–52 at.% Cu alloy film byX-ray diffraction and compared it with that of the liquid,and they concluded that the alloy film was nanocrystallinewith a face-centered cubic (fcc) structure and an averageparticle size of less than 16 A [7]. More recently, Bellon’sgroup has employed atomic probe field ion microscopy tostudy mechanically alloyed Ag–Cu powder at different tem-peratures and demonstrated composition homogeneity[15].

The solid solution of Ag and Cu alloy, or other binaryalloy systems with a positive heat of mixing in general, isunstable since there is no barrier against decomposition.At higher cooling temperatures or upon reheating, the Ag

rights reserved.

mailto:[email protected]

1618 H. Chen, J.-M. Zuo / Acta Materialia 55 (2007) 1617–1628

and Cu phases separate, or evolve, into a crystalline struc-ture from the initial amorphous or nanocrystalline state [8].Mader and Nowick conducted resistivity measurementswith vacuum evaporated alloy films, and compared theresults with bulk data for the pure solid metals in both solidand liquid states [8,9]. The results from resistivity measure-ments of the thin films during annealing show two stages.In the first stage, the amorphous phase transforms intothe crystalline solid solution over a narrow range oftemperatures. In the second stage, an equilibrium two-phase structure is formed. It was suggested that the firststage of transition is due to the nucleation and growth pro-cess, while spinodal decomposition is regarded as the pri-mary mechanism for phase separation. The subsequentappearance of separated Ag and Cu fcc peaks in diffractionpatterns also indicate phase separation at elevated temper-atures [4].

The phase separation is unlike any other phase transfor-mation and involves significant rearrangement of theatomic structure through short- or long-range diffusion.Structural changes during phase separation are thus inevi-table. While in situ analysis has been employed for study ofphase separation in bulk metallic glasses [18], experimentalstudies of metastable binary alloy thin films are surpris-ingly few.

Here we report an in situ electron diffraction andmicroscopy study of phase separation in Ag–Cu alloy thinfilms and give a quantitative description of the structuralevolution during phase separation.

2. Experiments

All Ag–Cu thin films were deposited in a load-lockedhigh-vacuum stainless-steel DC magnetron sputter deposi-tion system, with a base pressure of 5 · 10�8 Torr (7 ·10�6 Pa). The targets, a 99.995% pure Cu disk and a99.995% pure Ag disk, were co-sputtered in pure Ar(99.9995%) discharge maintained at 5 mTorr (0.67 Pa)using a capacitance manometer and an automatic mass-flow controller. Sputtering was performed at differentpower of Ag and Cu targets, (typically 15 W, with 328 Vand 38 mA for the Ag target and 29 W with 413 V and63 mA for the Cu target in Ag50Cu50 samples) in order tocontrol the growth rate of individual Ag or Cu targets.The film thickness was controlled by deposition time.Ag–Cu thin film samples were supported by an amorphousholey carbon film at room temperature. Sample thicknessand composition were calibrated by RBS measurementand X-ray reflectivity measurement of co-sputtered Ag–Cu thin film on a clean Si(1 11) substrate.

The Ag–Cu films were characterized by transmissionelectron microscopy (TEM), scanning transmission elec-tron microscopy (STEM), atomic force microscopy(AFM) and energy dispersed X-ray (EDX) spectroscopy.The TEM structure characterization was carried out inJEOL 2010F field-emission (scanning) transmission elec-tron microscope at 200 kV. We used the Gatan Model

901 heating stage for in situ observation of structural andphase evolution during heating. High-angle annular dark-field (HAADF) STEM images were recorded in JEOL2010F in STEM mode employing the approximate Z2

(where Z is the atomic number) dependence of theHAADF-STEM image intensities for detecting phase sepa-ration in Ag–Cu alloys. (The ratio of Z2 for Ag, Cu(ZAg = 47, ZCu = 29) is 2.63.)

Quantitative structure analysis was carried out usingelectron in the nanobeam diffraction (NBD) mode in theJEOL 2010F electron microscope with a submicron probesize, which allows us to record diffraction patterns usinga highly parallel illumination for high-resolution electrondiffraction. To characterize the atomic structure, werecorded electron diffraction patterns from the samples atdifferent annealing temperatures. The diffraction intensitieswere recorded on image plates (IPs) at NBD mode using alarge condenser aperture, which gives an illuminated areaof �600 nm in diameter. This relatively large area was usedto average over the sample. The diffraction intensity pro-files were obtained from the recorded diffraction patternsby first converting raw IP date from the logarithmic scaleto the linear scale, then integrating the radial intensity ofthe diffraction pattern. The amorphous carbon backgroundfrom the supported film was subtracted off using the dif-fraction intensity profiles recorded from amorphous car-bon film alone. The diffraction intensity was used toinvestigate the structure and for radial distribution func-tion analysis of the Ag–Cu films.

3. Results

This section describes the experimental investigation ofstructure evolution and phase separation for Ag/Cu sam-ples with three nominal film thicknesses (28, 57 and 84 A)as measured by using a quartz film thickness monitor andcompositions (Ag60Cu40, Ag50Cu50 and Ag33Cu67).

3.1. Ag50Cu50 of different film thicknesses

Three types of samples were prepared by co-sputteringusing the conditions specified previously, with differentthickness of 28, 57 and 84 A, and the Ag/Cu atomic ratiois 1:1. These samples were heated in situ from room tem-perature then maintained at several different temperaturesfor approximately 30 min and characterized by STEMimaging and electron diffraction.

3.1.1. Phase separation and morphology evolution of samples

in different annealing temperatures

Fig. 1 shows the morphology evolution of the Ag–Cualloy films from almost continuous thin films to varioussizes of clusters as the temperature increases. From theimages of as-deposited samples (Fig. 1, first row), the filmmorphology changes from non-continuous to continuousas the film thickness increases. The as-deposited films allshow almost uniform contrast in HAADF-STEM images.

Fig. 1. High angle annual dark field STEM images of Ag50Cu50 thin films of three different thicknesses annealed to different temperatures. The horizontaldirection is for sample thicknesses, and the vertical direction is for temperatures. The first row shows STEM images of as-deposited samples of 28, 57 and84 A nominal thickness. The 28-A-thick film is discontinuous, with a fractal morphology. As the thickness increases, the film morphology becomes morecontinuous. The second row shows STEM images of samples annealed to 200 �C. Dewetting and significant image contrast changes are seen in the thinnerfilms. The alloy thin films are fully dewetted for all three samples at 260 �C (third row), and the sizes of the final clusters are larger in thicker sample.Separation of the Ag and Cu phases into large clusters is clearly observed in all three thickness samples.

H. Chen, J.-M. Zuo / Acta Materialia 55 (2007) 1617–1628 1619

For samples annealed to 200 �C or above (Fig. 1, secondrow), the almost continuous thin films evolve into clustersof Ag, Cu and Ag–Cu. At 200 �C, phase-separated areasare visible in the thinner sample. When the samples areannealed to 260 �C (Fig. 1, third row), the alloy thin filmsare nearly all dewetted into discrete islands. The shape ofthe clusters becomes irregular as samples get thicker, andthe sizes of the final clusters are also larger in thicker sam-ple. Phase separation in large clusters is clearly observed inall different thickness samples.

3.1.2. Structure of Ag50Cu50 films at different annealing

temperatures

Fig. 2 shows the background subtracted electron diffrac-tion intensity profiles using the methods specified in Section2. Fig. 2(c)–(e) are the diffraction intensity profiles ofAg50Cu50 of three different thicknesses after annealing atfive different annealing temperatures (100, 150, 200, 230and 270 �C) for 0 < sin h/k < 1 A�1. For the as-depositedsample at room temperature, only a single peak atsin h/k = 0.213 A�1 is visible with a small shoulder at

sin h/k = 0.238 A�1, which indicates a good mixing of Agand Cu atoms in the alloy thin films. The peak position iscloser to that of Ag (sin h/k = 0.212 A�1) than to Cu(sin h/k = 0.240 A�1). The small shoulder is closer to theCu peak position. The atomic scattering factor of Ag is38.15 (sin h/k = 0.20 A�1) and Cu is 21.69 (sin h/k = 0.25 A�1). As sample is annealed to higher tempera-tures, the intensity of Cu peak increases and form two dis-tinct peak at T = 200 �C with peak position at sin h/k = 0.212 and 0.240 A�1. The pronounced Cu shoulder at100 �C is an indication of the early stage of phase separationat this temperature. When the sample is annealed at 150 �C,phase separation accelerates, and becomes well separated at200 �C. At this temperature, the Cu peak intensities arecomparable with those of the Ag peaks.

The small Cu shoulder is more pronounced for thickersamples, which cannot be explained by either multiple scat-tering or dynamic diffraction. Thus, some phase separationmust occur during sputter deposition.

The average grain sizes can be estimated from the half-width of full-maximum (HWFM) of lower-order Ag and

Fig. 2. Diffraction intensity profiles of Ag50Cu50 thin film samples of different thicknesses. (a) An example of recorded electron diffraction pattern of as-deposited 28 A Ag50Cu50 thin film sample. (b) The radial intensity profile of (a) obtained by averaging over the radius and subtracting off the central peakusing a reference diffraction pattern recorded from the carbon film. The radial intensity profiles are then plotted up to sin h/k = 0.3 A�1 for the (c) 28 (d) 57and (e) 84-A-thick films at different annealing temperatures. The Cu(111) appears as a shoulder in the intensity profiles for as-deposited samples, whichbecomes more prominent in the thicker films. As temperature increases, the small Cu(111) shoulders in the 57 and 84 A thickness samples grow intoseparate peaks in all three samples. At 200 �C, the intensity of the Cu peaks is comparable with that of the Ag peaks.


Cu peaks (the width of higher-order diffraction peaks isaffected by strain). The width of the Ag(11 1) andCu(11 1) peaks from the diffraction patterns of 57-A-thick

Ag50Cu50 film at different annealing temperatures is mea-sured by fitting the peak profile with a Gaussian functionðf ðxÞ ¼ 1

rffiffiffiffi2pp e�ðx�lÞ2=2r2Þ. The parameter r in Gaussian


function is approximately related to the average grain sizeof Ag and Cu by

L ¼ 0:94 � kBð2hÞ cos h

� 0:94 � k4ffiffiffiffiffiffiffiffiffiffiffi2 ln 2p

� rð1Þ

where k is the wavelength of the electrons and B(2h) is theFWHM of the diffraction peak in 2h [19].

Fig. 3(a) shows the evolution of Ag and Cu grain sizes asa function of annealing temperatures. For the 28-A-thickAg50Cu50 film annealed to 100 �C, the size of the Ag-richand Cu-rich particles is 29 and 18 A respectively. The Aggrain size evolves in two stages: there is little change inthe average Ag grain sizes up to 230 �C (about 3.2 nm),whereupon it increases to 4.4 nm after annealing to300 �C. For Cu, there is a continuous increase in theaverage grain size from 1.8 nm at 100 �C to 3.4 nm at300 �C.

The (111) d-spacing of Ag and Cu phases can be mea-sured from the positions of Ag(1 11) and Cu(11 1) peaksusing the same fitting procedure with Gaussian functionsðf ðxÞ ¼ 1

rffiffiffiffi2pp e�ðx�lÞ2=2r2Þ. The results are shown in Fig. 3(b).

100 150 200 250 3001.5

2.0

2.5

3.0

3.5

4.0

4.5

5.0

Cry

stal

lite

Siz

e (n

m)

Temperature (°C)

AgCu

100 150 200 250 3002.05

2.10

2.15

2.20

2.25

2.30

2.35

Temperature (°C)

Cel

l Con

stan

t (Å

)

Ag Grain CellAg Bulk CellCu Grain CellCu Bulk Cell

a

b

Fig. 3. Temperature-dependent grain size and lattice constant for the28-A-thick Ag50Cu50 film. (a) Change of the average grain size withannealing temperature. (b) Change of the average Ag and Cu latticeconstants with annealing temperature.

To analyze the atomic structure of the co-sputtered thinfilms, we performed radial distribution function (RDF)analysis using the recorded diffraction patterns. The princi-ple of this method is similar to that of X-rays, with theexception of experimental procedures and the differencethat electrons are scattered by strong Coulomb potentials[20,21]. The multiple scattering effects in amorphous ornanocrystalline thin films are much weaker than that ofcrystals. Because of this, we neglected multiple scatteringfor the ultrathin films studied here. The radial distributionfunction gives the distance correlation between a pair ofatoms. Following Warren [19], the RDF is calculated by

rGðrÞ ¼ 4pr2½qðrÞ � q0�

¼ 2rp

Z 1

0

SIðSÞNf 2� 1

� �sin rS dS ð2Þ

where q(r) is the reduced density, which corresponds to thenumber of atoms per unit volume at a distance r from thereference atom, and 4pr2q(r)dr is the number of atomscontained in a spherical shell of radius r and thicknessdr. Other quantities involved in Eq. (1) are q0, which isthe average density of atoms in the sample S = 4p sin h/k,I(S), the recorded electron diffraction intensity minus theinelastic background, and N, the normalization factor. Æfæis the averaged atomic scattering factor, with hf i ¼P

iciKif . It is assumed that atomic scattering is propor-tional to the atomic number. f is determined from

hf 2i ¼X

i

cif 2i ¼

Xi

ciK2i f 2

and

fi ¼X

i

ci

Z 1

0

4pr2V iðrÞsin Sr

Srdr

where V is the atomic potential and c the concentration.The RDF analysis can provide a general description ofthe degrees of order, the distribution of correlation dis-tances and the average nearest-neighbor coordinationnumbers. Orientation ordering or local geometrical pack-ing of atoms cannot be retrieved from the one-dimensionalRDF analysis, and we need to resort to other techniques,like fluctuation microscopy [22]. The advantage of electrondiffraction for RDF analysis is that it can be applied to thinfilms and small volumes because of the large electron elasticcross sections. For the same reason, electron diffractionmust be limited to very thin films to avoid multiplescattering.

The experimental procedures that we used for RDF anal-ysis are as follows. The electron diffraction patterns arerecorded digitally using methods described in Section 2.We use a camera length of 600 mm to include s = sin h/kup to 1.27 for rotation averaging. The recorded electron dif-fraction has a strong central peak, which includes a smallangle inelastic scattering and the profile of the electronbeam. This central peak was modeled by a pseudo-Voigtfunction and subtracted off. The normalization coefficient,


N, was found by fitting to large angle scattering using theatomic scattering factors. Finally, the RDF was by directintegration of Eq. (2).

Fig. 4 shows a comparison of the RDF of the 28-A-thick Ag50Cu50 film at different annealing temperaturesand the calculated RDF results from single-crystal Agand Cu particles of 4.5 and 4.0 nm in diameter respectively.The RDF of the as-deposited Ag50Cu50 sample is in goodagreement with that obtained by Wagner et al. usingX-ray diffraction [7] except the Cu peak is slightly morepronounced in our case. From the RDF plot of theAg50Cu50 sample (Fig. 4(a) and (b)), we can see two sepa-rated nearest-neighbor peaks. The second peak corre-sponds to Ag–Ag distance, which shifts from 2.80 to2.88 A from the as-deposited sample to the annealed at200 �C. The distance of 2.88 A fits with the first nearest-neighbor peak of Ag crystal (2.875 A) very well (Fig. 4(c)

0 5 10 15 20 25 30 35 40-20

0

20

40

60

80

200°C

150°C

rG(r

)

AsDeposit

100°C

0 1 2 3 4

0

5000

10000

15000

20000

25000

30000

D

Distance (Å)

Inte

nsi

ty (

a.u

.)

ab

c

Fig. 4. Radial distribution function in the form of rG(r) obtained from the exp(a and b) Radial distribution function of the Ag50Cu50 28 A sample at differentalong the y-axis in (b) for plotting. (c) Calculated RDF of Ag and Cu single-

and (d)). At 200 �C, the first peak is situated at Cu–Cuatomic distance of 2.56 A, which fits into the position ofthe first nearest-neighbor peak of Cu particles (2.55 A) verywell (Fig. 4(c) and (d)). Interestingly, for the as-depositedsample, the Cu–Cu peak is at 2.45 A and is lower thanthe bulk Cu–Cu distance. A shorter distance is expectedfor Ag, with doping of the smaller Cu atom, but not forCu, for which we expected a longer distance due to Agdoping. The reduced distance suggests that Cu is undercompressive strain. The fourth and fifth peaks inFig. 4(a) show splitting (the Cu peak growing from theCu shoulder in the Ag peak) and shifting during theannealing process, and this phase separation is due to Cudiffusing out of the Ag matrix. A detailed list of the posi-tion of first six peaks in these RDF results is shown inTable 1. Peaks 1, 3 and 4 belong to Cu, while peaks 2, 5and 6 belong to Ag. The ratio between peaks 5 and 2 is

0 1 2 3 4 5 6 7 8 9 10-20

0

20

40

60

80

6

5

43

2rG

(r)

Distance ( )

1

200°C

150°C

AsDeposit

100°C

5 6 7 8 9 10

istance (Å)

Å

Ag Single Crystal Cu Single Crystal

erimental diffraction patterns and comparison with bulk atomic distances.annealing temperatures plotted on different scales. The RDF is offset by 20crystal particles.

Table 1Position of the peaks in the RDFs of the Ag50Cu50 28 A sample at different annealing temperatures

Position of peaks in radial distribution function (A)

1 2 3 4 5 6

As-deposited 2.44 2.80 3.60 –a 4.88 5.60100 �C –a 2.84 3.76 4.56 4.96 5.76150 �C 2.56 2.84 3.72 4.52 5.00 5.80200 �C 2.56 2.88 3.72 4.44 5.04 5.80

a These peaks appear as a shoulder, whose position is difficult to measure.


1.74, which is close toffiffiffi3p

, and the ratio between peaks 6and 2 is 2. The ratio of

ffiffiffi3p

comes from the third nearestneighbor, while 2 from the fourth. Both are in agreementwith an fcc cell. In the liquid, the ratio between peaks 5and 2 is 1.86 for close packed liquid [7]. The ratios are tem-perature dependent. The ratio between peaks 4 and 1 variesfrom 1.77 at 150 �C to 1.73 at 200 �C, while the ratiobetween peaks 5 and 2 first increases to 1.76 with annealing

Fig. 5. STEM images of samples with three different Ag/Cu ratios and annedifferent Ag/Cu ratios and the vertical direction is for three different temperdifferent Ag/Cu atomic ratios. The morphologies of as-deposited samples shoislands for the Cu-rich film. The contrast differences among samples with diffeSTEM images of samples with different Ag/Cu atomic ratios annealed tomorphology compared with the other two samples of different compositions. Ibrighter contrast appears in most areas. The third row shows STEM images ofilms are all dewetted into alloy clusters of nearly the same size, and there arCu-rich phases.

temperature and then drops slightly to 1.75 at 200 �C. Thesame trend occurs for the ratio between peaks 6 and 2.Thus, the early stage of phase separation appears to beassociated with an initial increase in strain.

The size of the particles can be measured by taking thedistance at which oscillation decreases to the noise levelin the RDF spectrum. Using this, we find the size for theas-deposited film to be about 30 A, which corresponds to

aled to two different temperatures. The horizontal direction is for threeatures. The first row shows STEM images of as-deposited samples, withw highly percolated islands for the Ag-rich films and fractal, individual

rent Ag/Cu ratios are visible but not very obvious. The second row shows200 �C. The Ag-rich sample shows a fractal, percolated but less densen the Cu-rich sample, most areas appear dark, and in the Ag-rich sample,f samples with different Ag/Cu atomic ratios annealed to 260 �C. The thine clear contrast differences in these images corresponding to the Ag- and

0.15 0.20 0.25 0.30

0

200

400

600

800

1000

Inte

nsity

(a.

u.)

sinθ/λ (1/Å)

Cu(111)

Ag(111) As-Deposit100°C150°C200°C

0.15 0.20 0.25 0.30

0

500

1000

1500

2000

2500

Inte

nsity

(a.

u.)

sinθ/λ (1/Å)

Cu(111)

Ag(111) As-Deposit100°C

150°C200°C

0.15 0.20 0.25 0.30-200

0

200

400

600

800

1000

sinθ/λ (1/Å)

Inte

nsity

(a.

u.)

Cu(111) As-Deposit100°C150°C200°C

a b

c

Fig. 6. Diffraction intensity profiles of 28-A-thick samples with different Ag/Cu ratios under different annealing temperatures. (a) Ag/Cu = 3/2. (b) Ag/Cu = 1/1. (c) Ag/Cu = 1/2. Initially, only Ag peaks are visible in the Ag50Cu50 and Ag-rich samples, while only a broad peak is visible in the Cu-richsample. As temperature increases, in the Ag50Cu50 and Ag-rich samples, Cu peaks begin to appear and become more pronounced. Only two small Agshoulders can be seen in the Cu-rich sample.


the grain size of Ag measured from the width of the firstdiffraction peak. Upon annealing, the strength of firstand fourth peaks increases, which is a strong indicationof the growth of Cu-rich phase. Peak 2 decreases, whichis associated with the decrease of the nearest neighbors atthat distance. These findings can be taken as an indicationthat Cu separated out from the solutions. Substantialgrowth in the overall crystal size is observed at 200 �C inthe RDF. Below that temperature, the overall grain sizechanges little, although the peak shape analysis indicatesthe growth of Cu particles.

3.2. AgxCu1�x thin films of different Ag/Cu ratios

To see the effects of composition on phase separation,we investigate here Ag–Cu alloy films of three differentcomposition ratios at a constant thickness. These sampleswere prepared by co-sputtering by controlling the sputter-ing rate of Ag and Cu using the conditions specified inSection 2. Films 28 A thick were prepared with the Ag/Cu atomic ratios of 1:1 (from the previous section), 1:2

(Cu-rich) and 3:2 (Ag-rich). The samples were investigatedusing the same techniques as in the previous section.

3.2.1. Phase separation and morphology evolution of

AgxCu1�x thin film, x = 0.33, 0.50 and 0.60

Fig. 5 examines phase separation and morphology evo-lution of AgxCu1�x thin films of three different composi-tions. With increasing temperature, the morphology ofthe alloy films changes from nearly continuous thin filmsto clusters of various sizes. The morphologies of as-depos-ited samples show highly percolated islands for the Ag-richfilms and fractal, discrete islands, for the Cu-rich film(Fig. 5, first row). The contrast differences among as-depos-ited samples with different Ag/Cu ratio are visible. Whenannealed to 200 �C (Fig. 5, second row), the Ag-rich sam-ple shows a fractal, percolated but less dense morphologycompared with the other two samples of different composi-tions. Also, we can see some contrast difference amongthree samples. In the Cu-rich sample, most areas appeardark, and in the Ag-rich sample, the image contrast isreversed with most areas appearing bright. After these

0 1 2 3 4 5 6 7 8 9 10-20

0

20

40

60

80

rG(r

)

200°C

150°C

AsDeposit

100°C

0 1 2 3 4 5 6 7 8 9 10-20

0

20

40

60

80

rG(r

)

200°C

150°C

AsDeposit

100°C

Distance (Å)

Distance (Å)

a

b

Fig. 7. Radial distribution function of (a) the Cu-rich sample and (b) theAg-rich sample as a function of annealing temperature. The first peakevolves into two separated peaks in the Ag-rich sample (b) but not in theCu-rich case (a). (b) is same as Fig. 4b.


samples were annealed to 260 �C (Fig. 5, third row), thethin films are all dewetted into clusters of different sizes,and there are obvious contrast differences between differentclusters and within the same cluster. The brighter contrastis associated with the Ag-rich phases and the much darkercontrast comes from the Cu-rich phases, as shown by EDXanalysis.

3.2.2. Structure evolution of samples in different annealing

temperaturesThe film atomic structure was investigated by electron

diffraction. All electron diffraction patterns were recordedusing the same conditions as specified in Section 2, andthe intensity profiles were obtained by the same methodsas explained before.

Fig. 6 shows the diffraction intensity profiles of samplesof different compositions at different annealing tempera-

tures. Initially, for as-deposited samples at room tempera-ture, it is clear that while only Ag peaks are visible in theAg50Cu50 sample and Ag-rich sample, only a single solu-tion peak is visible in the Cu-rich sample.

At 150 �C, a small Cu(11 1) shoulder and peak begin toshow up in both the Ag50Cu50 sample and the Ag-rich sam-ple, while one small Ag shoulder can be seen in the Cu-richsample. In each case, the small, but visible, separated peaksindicate the early stage of phase separation in both theAg50Cu50 and Ag-rich samples. However, the effect of Agseparation from Cu in the Cu-rich sample is much smallercompared with Cu separation in the Ag50Cu50 and Ag-richsamples.

The RDF functions of the Cu-rich and Ag-rich 28 Asamples (Fig. 7) are obtained following the same procedureused in Section 2. From the RDF plot of the Cu-rich sam-ple (Fig. 7(a)), we can see that the first nearest-neighborshows no obvious shift or splitting during the annealingprocess, and the peak is situated at a Cu–Cu atomic dis-tance of 2.56 A. In the Ag-rich sample (Fig. 7(b)), the ini-tial first nearest-neighbor peak corresponds to an Ag–Agatomic distance of 2.88 A, and later this peak is trans-formed into two separate peaks similar to Ag50Cu50, whichindicates separate Ag- and Cu-rich phases. These resultsindicate that for the Ag-rich samples, the Cu crystallitesgrow in almost the same way as in the Ag50Cu50 samples;while in the Cu-rich samples, Ag atoms diffusing out of theCu matrix are barely noticeable. The difference must bedue to the difference in the diffusion rates of Ag and Cuatoms, since the order of diffusion rates is Cu in Cu < Agin Cu < Ag in Ag. The double peaks in the as-depositedCu-rich sample at 150 �C do not show up in 100 and200 �C samples, which could result from the RDF analysisprocedure. The Cu-rich samples have a large backgroundin the recorded diffraction pattern, which makes back-ground subtraction difficult.

3.3. Surface morphology study of samples at differenttemperatures

TEM characterization techniques only give 2-D projec-tion information about the thin films. It is difficult to seesurface morphology directly from TEM or STEM images.The AFM technique meanwhile is an excellent techniquefor studying film morphology. Digital Instruments Dimen-sion 3000 AFM was utilized to investigate film morphologyevolution in the z-direction during the sample annealingprocess. The AFM experiment was performed for differentAg50Cu50 84 A thin-film samples, which have previouslybeen annealed at different annealing temperatures.

AFM images (Fig. 8) show that the Ag/Cu films are con-tinuous flat thin film after deposition, and that subsequentheat treatments roughen the film surface. As the annealingtemperature increases, the surface roughness also increases(Fig. 8). Fig. 9 plots the surface roughness as a function oftemperature. The results show that there is no significantchange in film morphology below 200 �C, although we

Fig. 8. AFM surface images of 84-A-thick Ag50Cu50 samples annealed to different temperatures. The scale bars at the right side of the images indicate theheight scales of the images in that row.

2

4

6

8

10

12

14

16

Sur. Rms (Rq)

Sur

face

Rou

ghne

ss (

nm2 )

Annealing Temperatures (°C)

270230200150AsDeposit 100

Early Stage of Phase Separation (ED)

Cluster Forming (STEM)

Fig. 9. Surface roughness of different samples (with different annealingtemperatures) increases with the annealing temperature increases. The sumof squares of all height differences between every point and a certainreference (RMS) is plotted here for each temperature.


cannot exclude atomic scale events on the surfaces. Thechange in surface roughness is small up to �150 �C, where-upon it starts to increase. A roughening transition occurssomewhere between 230 and 270 �C, where surface rough-ness increases dramatically and dewetting and formation ofindividual cluster occur. Thus, the Ag–Cu alloy thin filmsbegin dewetting and cluster forming at temperature some-where between 200 and 230 �C, and complete the processabove 230 �C. Electron diffraction shows that the earlystage of phase separation occurs around 100 �C, at whichstage no obvious surface morphology change is observed.

4. Discussion

The in situ and ex situ experimental observations ofmorphology and structure of AgxCu1�x thin films bymicroscopy and electron diffraction show two stages oftransformation. The first stage is associated with the for-mation of two terminal phases (Ag- and Cu-rich) from theAg/Cu solid solution phase of as-deposited films. Thetransformation involves no significant morphologicalchange in the thin films. The second stage of transitionis associated with thin film dewetting and formation oflarge (�101 nm) Ag and Cu grains. STEM observationduring in situ annealing shows that the transition fromthe first to the second stage of transformation occursaround 200 �C, depending on the film thickness. Thin-filmdewetting is completed at �260 �C for all three thicknessstudied (28, 57 and 84 A). Surface roughness measured byAFM shows that the 84-A-thick Ag50Cu50 alloy thin filmsbegin dewetting and cluster forming at temperature some-where between 200 and 230 �C, with the film being fullydewetted above 230 �C. Electron diffraction shows thatthe first stage of phase separation occurs around 100 �C,at which no obvious surface morphology change isobserved by AFM.

The structure of the as-deposited thin films depends onthe film thickness. A double peak develops in the 84-A-thick film for the first diffraction peak, while only a small,but noticeable shoulder is seen in the 28-A-thick film. Thegrain size estimated from the width of the first diffractionpeak shows a large difference between thicker and thinnerfilms. The structure of the 28-A-thick film can be consid-ered as close to that of a solid solution. Thus, phase segre-gation can develop during thin film deposition. The films

(a) (b)

r

Fig. 10. (a) As-deposited flat and continuous alloy thin film. (b) Film dewets into a cluster with a semi-spherical shape of radius r on the substrate surface.


studied here were prepared by co-sputtering onto carbonfilm substrate at room temperature. The relatively poorthermal conductivity of the carbon film (0.20 W m�1 K�1

at room temperature [23]) and the high substrate tempera-ture probably contribute to the phase segregation. TheRDF spectrum of 28-A-thick Ag50Cu50 film is consistentwith the earlier published results and the nanocrystallinestructure suggested by Wagner et al. [7].

The first transformation is associated with the develop-ment of two phases, which corresponds to the spinodaldecomposition observed previously [8,9,13]. Electron dif-fraction and RDF analysis results show the first stage ofphase separation in Ag50Cu50 and the Ag-rich sample butnot in the Cu-rich sample. The phase separation processis much slower in the Cu-rich samples than in the othertwo samples with a higher Ag concentration. Decomposi-tion is thermally activated in this case. Mader and Nowich[9] interpreted the activation energy as a barrier for inter-diffusion. Thus, the decomposition rate, which is propor-tional to the amount of alloyed phase x, can be relatedto the diffusion constant [13]:

dxdt¼ �K0DðT Þx

where K0 is a constant and D(T) is the temperature-depen-dent diffusion constant. The observation of thermal decom-position at the low temperature of 373–473 K (100–200 �C)is consistent with the results of Klassen et al. [13] obtainedfor ball-milled Ag50Cu50 alloy, in which they found thatball milling at 473 K led to fully decomposed Ag and Cualloys. In the work of Klassen et al., decomposition ofAg–Cu films at low temperatures was attributed to the low-er activation energy for diffusion due to vacancy-mediateddiffusion. In the case of thin films, the surface may play thesame role as vacancies in the bulk.

From the energetic point of view, the phase separationprocess can be regarded as a competition between the heatof mixing (DHmix) in the solution phase and the reductionin entropy and increase of interface plus strain energies inthe separated phases. During phase separation, the initiallymixed Ag and Cu atoms gradually separate from eachother to form the Ag-rich and Cu-rich phases. This processreduces DHmix, while increases interface and strain ener-gies. The change of heat of mixing can be viewed as thechange in cohesive energy from different Ag and Cu atomsin the solution phase to the cohesive energy among thesame atoms in the mixed solid solution phase, as DHmix =EAg–Cu � EAg–Ag � ECu–Cu and DGmix = DHmix � TDSmix

is positive for the three Ag/Cu alloy compositions studied

here. The change of interface energy can be written asDc = cinterface · r2 and the strain energy can be regardedas part of the interface energy because of the incoherentinterface from the large lattice mismatch. The criticalcluster size r* is r� ¼ 2cinterface

DG . When cinterface is much largerthan DGmix, the barrier for nucleation approaches zeroand spinodal decomposition occurs by composition fluctu-ation and uphill diffusion. The approximate value of r* isclose to 4.2 nm when we use DHmix = 3 kJ mol�1 = 3.44 ·108 J m�3 [1], DSmix = 3.07 J mol�1 K�1 = 3.52 · 105 Jm�3 K�1 [24] and cinterface = 0.5 J m�2 [25]. From ourexperimental observation, a Cu-rich phase of �2 nm andan Ag-rich phase of �3 nm exist in as-deposited films.

The second stage of transformation involves film dew-etting. The driving force comes from the interface energybetween metal thin films and the support and their surfaceenergies. Large grains in the order of 101 nm are formed inthis process driven which is associated with significantamount of interdiffusion. In Ag/Cu alloy, the surface/inter-face energies include that of substrate surface, cluster sur-face, the substrate-cluster interface and Ag and Cuinterphase boundaries. If we simplify the cluster shape asa semisphere (Fig. 10) with radius r, the surface area ofas-deposited thin film as S, and the Ag–Cu interface areaas s ¼ 1

2pr2, we can write the change of surface/interface

energy during film dewetting as DG ¼ cS=F � ðpr2 � SÞþcF � ð2pr2 � SÞ þ cs � ðS � pr2Þ þ cAg–Cu � N � pr2=2, inwhich cS/F, cF, cs and cAg–Cu are the interface energy of sub-strate–film, film–vacuum, substrate–vacuum and Ag–Cuinterphase boundary. The factor of N comes from theassumption that the dewetted cluster size is N times largerthan the cluster size before dewetting. The term of inter-phase boundary must play a role here since film dewettingis observed after phase separation at relatively higher tem-peratures (>200 �C).

The phase separation is composition-dependent. Whileit is observed in Ag50Cu50 and Ag-rich samples, it is barelynoticeable in Cu-rich samples. This phase separation pro-cess difference, reflected from both STEM images and elec-tron diffractions, must be due to the difference in thediffusion rates of Ag and Cu atoms, since the order of dif-fusion rates is Cu in Cu < Ag in Cu < Ag in Ag.

The RDF and grain size measurement show that initialphase separation is dominated by the growth of Cu crystal-lites, which are very small (less than 1 nm in size) in as-deposited thin films. The development of a clear, separatedpeak as a function of annealing temperature indicates anincreasing number of Cu–Cu and Ag–Ag pairings. Thisand the increase in the peak intensity at the third Cu–Cu


neighbor distance are associated with the growth of Cucrystallites.

Electron diffraction proves to be much more sensitive tophase separation than imaging. The phase separation thatwe observed in the STEM images at higher temperatures(P200 �C) is mainly due to the growth of a few Ag crystal-lites while the average size remains about the same below230 �C.

5. Conclusions

We have studied the phase separation and structure evo-lution process in co-sputtered Ag–Cu alloy thin films byin situ electron diffraction and composition-sensitivehigh-angle annual dark-filled STEM as well as ex situ filmmorphology characterization using AFM. The resultsshow two stages of transformation. The first stage is asso-ciated with decomposition from the solid solution phase totwo terminal phases, while the second stage is associatedwith film dewetting and the formation of crystal grainsof 101 nm. The first stage transition corresponds to spin-odal decomposition observed earlier in amorphous Ag–Cu alloys. The structure analysis by electron diffractionshows that the initial stage of decomposition is associatedwith the growth of Cu crystallites in Ag50Cu50 and Ag-richfilms, and the initial phase separation is also accompaniedby strain.

Acknowledgements

The work was supported by NSF DMR 0449790. Theelectron microscopy characterization was supported byDOE DEFG02-01ER45923 and DEFG02-91ER45439 car-ried out in the Center for Microanalysis of Materials, Uni-versity of Illinois, which is partially supported by the USDepartment of Energy under grant DEFG02-91-ER45439.

References

[1] Liu BX, Lai WS, Zhang Q. Irradiation induced amorphization inmetallic multilayers and calculation of glass-forming ability fromatomistic potential in the binary metal systems. Mater Sci Eng R:Reports 2000;29:1.

[2] Zghal S, Twesten R, Wu F, Bellon P. Electron microscopy nanoscalecharacterization of ball milled Cu–Ag powders. Part II: Nanocom-posites synthesized by elevated temperature milling or annealing.Acta Mater 2002;50:4711.

[3] Hayes FH, Lukas HL, Effenberg G, Petzow G. A thermodynamicoptimization of the Cu–Ag–Pb system. Z Metallk 1986;77:749.

[4] Duwez P, Willens RH, Klement W. Continuous series of metastablesolid solutions in silver–copper alloys. J Appl Phys 1960;31:1136.

[5] Sheng HW, Wilde G, Ma E. The competing crystalline andamorphous solid solutions in the Ag–Cu system. Acta Mater 2002;50:475.

[6] He JH, Sheng HW, Lin JS, Schilling PJ, Tittsworth RC, Ma E.Homogeneity of a supersaturated solid solution. Phys Rev Lett2002;89:125507.

[7] Wagner CNJ, Light TB, Halder NC, Lukens WE. Structure of avapor-quenched AgCu alloy. J Appl Phys 1968;39:3690.

[8] Mader S, Nowich AS, Widmer H. Metastable evaporated thin films ofCu–Ag and Co–Au alloys – I occurrence and morphology of phases.Acta Metall 1967;15:203.

[9] Mader S, Nowich AS. Metastable evaporated thin films of Cu–Agand Co–Au alloys – II kinetics of the transformations. Acta Metall1967;15:215.

[10] Tsaur BY, Lau SS, Mayer JW. Continuous series of metastable Ag–Cu solid-solutions formed by ion-beam mixing. Appl Phys Lett1980;36:823.

[11] Enrique R, Wu F, Bellon P. A new approach for the direct synthesisof nanocomposite thin films by ion-beam processing. Surf CoatTechnol 2002;150:1.

[12] Liu BX, Lai WS, Zhang ZJ. Solid-state crystal-to-amorphoustransition in metal–metal multilayers and its thermodynamic andatomistic modelling. Adv Phys 2001;50:367.

[13] Klassen T, Herr U, Averback RS. Ball milling of systems withpositive heat of mixing: effect of temperature in Ag–Cu. Acta Mater1997;45:2921.

[14] Zghal S, Hytch MJ, Chevalier JP, Twesten R, Wu F, Bellon P.Electron microscopy nanoscale characterization of ball-milled Cu–Agpowders. Part I: Solid solution synthesized by cryo-milling. ActaMater 2002;50:4695.

[15] Wu F, Bellon P, Melmed AJ, Lusby TA. Forced mixing andnanoscale decomposition in ball-milled Cu–Ag characterized byAPFIM. Acta Mater 2001;49:453.

[16] Ma E, Atzmon M. Phase-transformations induced by mechanicalalloying in binary-systems. Mater Chem Phys 1995;39:249.

[17] Klement K, Willens RH, Duwez P. Non-crystalline structure insolidified gold-silicon alloys. Nature 1960;187:869.

[18] Wang XL, Almer J, Liu CT, Wang YD, Zhao JK, Stoica AD, et al.In situ synchrotron study of phase transformation behaviors in bulkmetallic glass by simultaneous diffraction and small angle scattering.Phys Rev Lett 2003;91:265501.

[19] Warren BE. X-ray diffraction. New York: Dover Publications, Inc;1969.

[20] Cockayne DJH, Mckenzie DR. Electron-diffraction analysis ofpolycrystalline and amorphous thin-films. Acta Crystallogr Sect A1988;44:870.

[21] Hirotsu Y, Ishimaru M, Ohkubo T, Hanada T, Sugiyama M.Application of nanodiffraction to local atomic distribution functionanalysis of amorphous materials. J Electron Microsc 2001;50:435.

[22] Ma E. Alloys created between immiscible elements. Prog Mater Sci2005;50:413.

[23] Bullen AJ, O’Hara KE, Cahill DG, Monteiro O, von Keudell A.Thermal conductivity of amorphous carbon thin films. J Appl Phys2000;88:6317.

[24] Ozolins V, Wolverton C, Zunger A. Cu–Au, Ag–Au, Cu–Ag, andNi–Au intermetallics: first-principles study of temperature-compo-sition phase diagrams and structures. Phys Rev B 1998;57:6427.

[25] Porter DA, Easterling KE. Phase transformations in metals andalloys. New York: Van Nostrand Reinhold Company; 1981.

Structure and phase separation of Ag–Cu alloy thin films

Documents

Transcript of Structure and phase separation of Ag–Cu alloy thin films