larson acta mat review paper 04

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Overview No. 138

Information storage materials: nanoscale characterisation bythree-dimensional atom probe analysis

D.J. Larson a,*, A.K. Petford-Long b, Y.Q Ma b, A. Cerezo b

a Seagate Technology LLC, Recording Heads, One Disc Drive, NRE304, Bloomington, MN 55435, USAb Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK

Received 15 December 2003; received in revised form 11 March 2004; accepted 13 March 2004

Available online 22 April 2004

Abstract

The development of nanoscale magnetic materials for applications in information storage systems relies heavily on the ability to

engineer the properties of the layered structures from which such materials are fabricated. These properties are strongly dependent

on the nature of the interfaces between the individual nanoscale magnetic layers, so knowledge of the interface chemistry is crucial.

In this paper, we discuss the application of three-dimensional atom probe analysis to the characterisation of layered magnetic

materials, including details of specimen preparation techniques required for this type of analysis. Recent results are presented on the

characterisation of interfaces in Co/Cu or CoFe/Cu multilayers, which form part of the read sensor in magnetic recording heads, and

Co/Pd multilayers, which are being considered for use as perpendicular recording media.

2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Atom probe; Layered structures; Magnetic thin films

1. Introduction

In recent years there has been increasing interest in

thin-film layered structures because of the novel prop-

erties that they exhibit. Multilayer film (MLF) structures

are formed by alternate deposition of two or more dif-

ferent elements or compounds, usually with layers that

are 15 nm in thickness. Of particular interest are layered

structures formed from one or more transition-metal

magnetic materials because they have applications in the

field of data recording. Examples include high density

longitudinal or perpendicular magnetic recording media,

and sensors based on the giant magnetoresistance

(GMR) phenomenon such as those used for read heads

[1,2]. For a recent review of modern magnetic materials

in data storage see [3].

Layered structures based on Co, Ni, Fe, Cu and

their alloys form sub-structures of devices such as the

spin-valve [4], which exhibit GMR [2]. Changes in

microstructure (e.g. interfacial roughness, chemical

intermixing, etc.) are believed to influence the magni-

tude of the GMR effect and there has been much

research devoted to the dependence of the GMR ratio

on the specific nature of the interfaces in layered

structures. However, inadequate structural character-

ization has led to contradictory results arising from

different research groups even for the same materials

system (see, for example, [57]). This arises in part

because of the difficulty in distinguishing between

chemical mixing and roughness for most of the tech-

niques used to characterise the films. Interface

roughness may be defined as a spatial variation of

the interface dividing two layers and intermixing as

the chemical mixing of the elements that comprise the

two layers, at or near the interface. These two mi-

crostructural features are likely to have very different

effects on the properties, therefore establishing the

correlation between structure and properties requires

the ability to distinguish between roughness and

intermixing.

When characterising layered structures, the very

small layer thicknesses makes it extremely difficult to

* Corresponding author. Tel.: +1-952-402-7140; fax: +1-952-402-

7734.

E-mail address: [email protected] (D.J. Larson).

1359-6454/$30.00 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

doi:10.1016/j.actamat.2004.03.015

Acta Materialia 52 (2004) 28472862

www.actamat-journals.com
http://mail%20to:%[email protected]/http://mail%20to:%[email protected]/


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determine the nature of each individualinterface unless a

technique is used that allows independent interface

analysis with high spatial resolution. Low-angle X-ray

diffraction (X-ray reflectivity) has been widely used to

study layered thin film materials (see for example [8]).

However, the X-ray reflectivity profiles that are obtained

are summed data that include contributions from all theinterfaces in the structure. It can therefore prove ex-

tremely difficult to obtain a unique fit to the data for

multilayer films if the interfaces between different layers

do not contain the same degree of roughness and/or

interdiffusion through the multilayer stack. Neutron

reflectivity has also been applied to the study of inter-

faces in epitaxial superlattice structures and to poly-

crystalline sputtered films but this technique is not as

good at determining structural information as X-ray

reflectivity, and is therefore more usually applied using

polarised neutrons to determine magnetic structure (see

[9] for a recent review).

One of the most commonly used techniques that can

be used to analyse each interface independently is

transmission electron microscopy (TEM) of cross-sec-

tion samples, which can be used to study both the mi-

crostructure and compositional profile across the layers

[10]. However, a limitation of TEM is that it is a pro-

jection technique and integration through the specimen

thickness can make it difficult to distinguish interfacial

mixing and fine-scale roughness except under very spe-

cific conditions. An example is given by Hytch of sep-

arating these two effects in Co/Cu multilayer films using

detailed image analysis of through-focal series of Fresnel

fringe images [11].The technique of three-dimensional atom probe

(3DAP) analysis [1215] is unique in having the capa-

bility to characterise internal interfaces, grain bound-

aries and precipitates with sub-nanometre resolution in

all three dimensions. With this technique it is therefore

possible to measure the extent of interdiffusion or in-

terface segregation at the atomic scale, and separate

these effects from nanometre-scale topological features.

In a 3DAP, pulsed field evaporation is used to re-

move individual atoms from the surface of a specimen,

which is in the form of a needle with end radius $50 nm.A position-sensitive detector is used for measurement of

both flight time (to identify atomic species) and impact

position (which enables the original position of each

atom on the specimen surface to be recorded). The po-

sition of an atom in the depth direction is calculated

from the sequence of ion detection. The 3DAP has a

depth resolution of a single atomic layer and sub-

nanometre lateral resolution. An intrinsic requirement

of the 3DAP microscopy technique is the production of

a very high electric field (2050 V/nm) at the surface of

the specimen in order to achieve field evaporation. The

field is produced by applying a high voltage (typically

$10 kV) to a needle-shaped specimen with an apex ra-

dius of curvature of 10100 nm. For most materials,

these sharp needles can be produced by electrochemi-

cally polishing fine wires, whiskers or blanks cut from

bulk material [1216]. However, preparing a specimen

from certain sample geometries can be quite compli-

cated. Examples include certain multi-phase alloys,

semiconductors, ceramics and, in particular, thin filmmaterials. Over the years, various methods have been

applied to 3DAP specimen preparation in order to

overcome these problems [13,14].

This paper gives an overview of the results of recent

3DAP studies, both from our group and from other

groups, of magnetic layered film structures (e.g. Ni/Cu,

Co/Cu, CoFe/Cu, Co/Pd, etc.), concentrating on mate-

rials with potential applications in information storage

systems. These materials science studies have been made

possible by recent progress in the application of high

resolution focused ion beam systems to the preparation

of 3DAP specimens from complex layered systems, and

this is also reviewed briefly.

2. Development of 3DAP specimens from thin film

samples

In order to investigate thin layered structures in the

atom probe, these features must be present in the apex

region of the specimen. The preparation of 3DAP

specimens from thin film structures is difficult, but is a

crucial step in obtaining 3DAP data that provide the

highest possible spatial resolution across the feature of

interest, in this case the interfaces in the layered struc-ture. The specimen preparation process is difficult be-

cause the total sample thickness (region of interest

containing the thin layers) is often only of the order of

20100 nm. There are three main methods that have

been used to prepare 3DAP specimens from thin film

structures and these are outlined below.

A number of papers have been produced detailing the

results of 3DAP studies carried out on thin film mate-

rials in which the films of interest were deposited directly

onto pre-evaporated needle-shaped specimens [1730].

Fig. 1(a) shows a schematic of the geometry of a speci-

men that has been prepared in this way, but note that

the curvature of the tip relative to the layer thickness is

not necessarily to scale (the tip radius is of the order of

50 nm and the layer thickness is typically a few nano-

metres). The small volume from which the 3DAP data

are taken (of the order of 20 nm in diameter) means that

within the volume analysed the substrate approximates

to a flat surface, and the structure of the layers is often

assumed to be reasonably similar to that of the same

layers deposited on a planar substrate. However the

microstructure is unlikely to be exactly the same, in part

because the substrate material (in this case a metallic tip)

does not correspond well to the substrates used in ap-

2848 D.J. Larson et al.: Overview No. 138 / Acta Materialia 52 (2004) 28472862


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plications (most commonly an amorphous underlayer

followed by an oriented seed layer). This fact, combined

with the curvature of the end of the tip, will mean that

the stress configuration in the film is likely to be different

from that in a standard layered film. In addition, in

order to make a direct correlation of the structural and

composition results obtained by 3DAP with magnetic or

transport measurements, 3DAP analysis must be per-

formed on layered films deposited on flat substrates

because it is not possible to make standard magnetic and

transport measurements on the films deposited on tips.

It is also often desirable to investigate the microstructure

of films that have been deposited in standard industrial

deposition systems, rather than in special deposition

chambers designed specifically to enable the study of

films deposited onto pre-evaporated needles, as has also

been pointed out by Lang and Schmitz [30]. In spite of

these detriments, this technique has been widely usedand does have certain advantages.

A preparation method that addresses some of the

above issues has been reported by Hasegawa et al. for

specimens containing a single layer alloy film [31]. In this

method the film to be analysed is deposited on a sub-

strate, such as a Si wafer, covered in photoresist. The

film is then lithographically patterned to form small

samples with a shape similar to that shown in Fig. 1(b).

Following patterning the photoresist is dissolved, leav-

ing the samples that are comprised of the layer(s) of

interest. An individual sample must then be picked up

and attached to the end of a sharpened metallic tip using

conducting epoxy. Final shaping of the post section of

the sample, which will form the specimen to be analysed,

is carried out using a pulsed micropolishing technique.

This method results in a specimen with the film(s) in its

plane and has been used to investigate Co- and Fe-based

magnetic films [3234].

The Hasegawa technique does result in the fabrica-

tion of a specimen from a layered film deposited on a flat

substrate using a standard deposition system, but it

requires the film material(s) to be deposited onto a

non-standard seed layer and to be amenable to electro-

chemical polishing. Problems arise when the different

layers in the film require very different polishing solu-tions, as is the case for Cu and Co, two commonly used

materials in information storage applications. In such

cases, another method of specimen sharpening, such as

ion beam milling, must be used.

In principle, specimen preparation using ion milling

[3538] has several advantages over electrochemical

techniques [12,39]. These advantages include the elimi-

nation or reduction of contamination and the reduction

of preferential etching problems in multi-phase materi-

als such as layered films. The application of ion milling

to prepare 3DAP specimens from thin film(s) or near-

surface regions has not been widely investigated. Liddle

et al. [40], Larson et al. [41] and Kvist et al. [42] used a

broad ion beam to mill without the capability to image

the specimen during preparation. In-situ imaging and

ion milling of atom probe specimens using a gallium

liquid metal ion source was first reported by Waugh

et al. [43] using an instrument with a focused ion spot

size of$50 nm. Further work on ion milling of atomprobe specimens was performed by Alexander et al. [44]

using a Gatan model 645 precision ion mill with a

minimum spot size of approximately 1500 nm. Although

the above investigations show the feasibility of concur-

rent ion milling and secondary electron imaging, the full

Fig. 1. (a) Schematic diagram of the method of thin film deposition

onto curved sample surface, (b) schematic diagram of the modified

lithographic specimen preparation method described by Hasegawa

et al. [31] and Larson et al. [47] and (c) schematic diagram of the

specimen geometry that results from silicon post method used by the

current authors [50].

D.J. Larson et al.: Overview No. 138 / Acta Materialia 52 (2004) 28472862 2849


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capability of the new generation of commercial high

resolution (


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There are, of course, disadvantages associated with

ion milling methods. For example, milling may intro-

duce implantation and defects into the near-surface

region of the specimen but these can usually be re-

moved by field evaporation before atom probe analysis

proceeds. The FIB techniques described above lead to

a region of the specimen that has been partially dam-aged by the Ga ion beam [48]. An example of the

problems associated with the use of a focused ion beam

for specimen fabrication using the modified Hasegawa

technique is shown in Fig. 5. No layers were visible in

the field ion image of a specimen containing a (Co2 nm/

Cu2 nm)100 multilayer film milled using the FIB at

10 keV [45]. Subsequent atom probe analysis showed

the gallium content vs analysis depth to range from

$6.5 at.% Ga initially to $0.5% Ga at a depth of$15nm, as seen in Fig. 5(a). For comparison, the results of

analysis of a specimen fabricated using 30 keV gallium

ions are also shown. The gallium implantation level

was still 1020% even up to a depth of $20 nm intothe specimen. Fig. 5(b) shows atom maps of the Ga

and Cu taken from the 3DAP analysis of the specimen

milled at 10 kV and whose Ga content is shown in

Fig. 5(a). The layered structure appears after $20 nmof field evaporation, as shown by 3DAP (note that

Fig. 5(b) does not include the first 10 nm of the profile

shown in Fig. 5(a)).

The FIB-damaged region can be removed and mon-

itored during 3DAP analysis. Alternatively, the dam-

aged region can be removed by DC field evaporation

while field ion imaging, or by low-energy ion sputtering

using field emission in the presence of Ne gas [53]. Note

that only analyses that are relatively free of Ga im-

plantation (less than $12% Ga) should be consideredas representative of the original thin film microstructure.

3. Studies of layered thin films

Over the last decade or so, the three-dimensional

atom probe technique has been applied to various lay-

ered systems of interest for information storage mate-

rials. These include Ni/Cu [21], Co/Cu [27,46], Ni/Al

[29], CoFe/Cu [5456], NiFe/Cu [57], CoFe/NiFe [58],

Fe/Cr [59], Co/Pd [60], as well as various magnetic ma-

terial-based single film structures. The results of some of

these studies are discussed in more detail below.

3.1. Co/Cu and CoFe/Cu multilayer films

The modified Hasegawa technique was used by

the current authors to prepare specimens from an

Fe/(Cu2 nm/Co2 nm)100 multilayer film. An Fe seed layer

was used in these samples to impart a strong (1 1 1)

Fig. 3. Silicon post (a) prior to pattern placement at low ion current (inset at high ion current), (b) the evolution of the end of the post as milling

proceeds using the$1 lm inner radius pattern, (c) the post after milling has reduced the radius of the unmilled region of the specimen to$0.4 lm and(d) the post after milling at low current has reduced the radius of the unmilled region of the specimen to less than $0.1 lm.



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crystallographic texture to the MLF. The as-deposited

films showed a magnetoresistance ratio of$5% at roomtemperature, and the magnetic data suggested that a

significant fraction of the films were coupled ferromag-

netically (contrary to the expectation of antiferromag-

netic coupling for this copper interlayer thickness). The

coercivity was $60 Oe, which is relatively high for filmsof this type. A field-ion image of the multilayer,

Fig. 6(a), shows the brightly-imaging cobalt layers: in

some regions the layers are relatively straight and in

other regions the layers are non-parallel, with a wavi-

ness amplitude of$2 nm and period of$20 nm [45].

The specimen preparation method used in this workresults in layers that run parallel to the long axis of the

specimen and the layers are thus viewed in cross-section

in the field ion image. The waviness of the layers was

also evident in the 3DAP data obtained from the same

film, as seen in Figs. 6(b) (map of Cu atoms only) and

Fig. 6(c), which shows the Co composition profile. The

fact that the layers are wavy, as shown in Fig. 6, suggests

that the ferromagnetic coupling is a result of Neel or-

ange peel type magnetostatic coupling between adja-

cent cobalt layers. In addition, in some areas adjacent

cobalt layers appear to be in contact, which will further

increase the ferromagnetic coupling. The relatively high

coercivity is expected to be a result of the poor layer

planarity, which leads to a high number of domain wall

pinning sites.

Co/Cu layered structures have also been studied by

3DAP using the method of film deposition onto a pre-

formed W needle [27]. In this work, a layered structure

of NiFe25nm/Cu20nm/Co10nm was ion beam sputtered

onto a W needle and analysed in the as-deposited and

annealed (0.5 h at 350 C) states. Schleiwies at al. [27]

observed an interfacial solid solution zone approxi-

mately 1.5 nm in width at the interfaces in these samples

(as-deposited), as well as segregation of Fe through the

Fig. 4. (a) Side view of the specimen after the $0.4 lm inner radiuspattern stage of milling. The metal film(s) are visible as a bright region

on the end of the specimen. (b) Low and (c) high magnification images

of the final specimen shape. The radius is less than 100 nm and the

region of interest for 3DAP analysis is positioned at the apex of the

needle.

Fig. 5. (a) Composition profile in Co/Cu multilayer film specimen

showing the degree of Ga incorporation caused by FIB milling. (b)

3DAP atom maps for Ga and for Co from final $20 nm depth ofspecimen milled at 10 kV showing onset of appearance of the Co/Cu

layers following the Ga implantation damage. (Reproduced from [47]

with permission from Institute of Physics Publishing.)



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Cu layer to the Co/Cu interface after annealing,

although no physical mechanism for this effect was

proposed.

Fig. 7(a) shows atom maps of Ni, Co and Cu from a

3DAP dataset of a repeated (Ni20Fe5 nm/Co10Fe4 nm/

Cu3 nm/Co10Fe4 nm) multilayer film fabricated from a

sample deposited onto Si posts and prepared by the

annular FIB milling method as described above [50]. As

can be seen in the figure, for samples prepared in this

geometry, the interfaces run across the apex of the

specimen (normal to the axis) and thus the 3DAP

technique can achieve atomic resolution across the in-

terfaces, enabling accurate measurement of their

roughness and intermixing. This is clearly seen in

Fig. 7(b), which shows the (1 1 1) atomic planes across a

NiFe/CoFe bilayer, indicating the high degree of crys-

tallographic texture in the sample. In addition, micro-

structural results from specimens prepared in this

geometry can be compared directly with devices grown

for magnetic and transport measurements.

Fig. 8(a) shows an atom map indicating the atomic

positions of the Co (blue), Fe (yellow) and Cu (red)

atoms in a CoFe/Cu/CoFe trilayer (a subsection of the

multilayer structure described in the previous para-graph). The atom map shows qualitatively that there is

more intermixing at the CoFe-on-Cu interface than

there is at the Cu-on-CoFe interface, resulting in the

CoFe-on-Cu interface being wider than the Cu-on-CoFe

interface. However the strength of the 3DAP technique

lies in the fact that in addition to being able to show

such effects in a qualitative manner, the extent of the

interdiffusion can be measured quantitatively by taking

composition profiles through small sections of the data

perpendicular to the local interface plane. Fig. 8(b)

shows a composition profile across the same trilayer

region. The direction of growth is from left to right in

the profile. Measurement of the interface widths from

the composition profile (using 1090% of the Cu con-

centration) gave values of 1.08 (0.18) nm for the CoFe-on-Cu interface and 0.4 (0.14) nm for the Cu-on-CoFeinterface [61].

3.2. Comparison of 3DAP data with modelling

Molecular dynamics (MD) simulation of the growth

of the same materials system [61], Fig. 9, supports the

3DAP results presented above. As can be seen, the

match between the 3DAP data and the simulation is

extremely good, with the MD simulation giving a widthfor the CoFe-on-Cu interface of$1.44 nm and for theCu-on-CoFe interface of $0.33 nm. Further 3DAPquantitative analyses of this type gave values of 0.82

(0.10) and 0.47 (0.15) nm for the CoFe-on-Cu andCu-on-CoFe interfaces, respectively [54,58]. The differ-

ence in interface width is believed to be the result of the

lower surface free energy of Cu compared to that of Co

or Fe [62]. The molecular dynamics modelling suggests a

difference in the exchange probability during deposition

of Cu with Co and Fe at the two interfaces, with the

exchange probability at the CoFe-on-Cu interface being

higher, resulting in more mixing at the interface. Of

course, this will result from a combination of driving

forces (due to the lower surface energy for Cu) and the

lower activation barrier for exchange (due to the lower

melting temperature of Cu), and it is difficult to distin-

guish the two effects in the model. Note that the model

only takes into account rearrangements resulting from

atomic impacts, and does not cover the much longer

time scale over which thermally activated surface and

bulk diffusion may occur and contribute to intermixing.

However, since the agreement between model and ex-

periment is so good, such effects are not likely to con-

tribute greatly to the observed intermixing, at least at

Fig. 6. (a) Field ion microscope image showing Co (bright) and Cu

(dark) layers, (b) 3DAP atom map of the Co/Cu multilayer showing

degree of curvature of the Co and Cu layers (Cu atoms only for clarity)

and (c) selected-region Co composition profile across the layers.



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the deposition rates (%0.1 nms1) and substrate tem-peratures (300330 K) used here. The good agreement

between the simulated and experimental data shows that

3DAP analysis is really the first type of experiment that

can be used to check this model at the correct length

scale. A favourable comparison between model and

experiment also served to validate the MD simulation

parameters used for these materials and has encouraged

both research groups to proceed with further compari-

sons, as described below.

The data in Figs. 8 and 9 show a layered film in which

the top surface of the CoFe layer is relatively flat.

However growth of multilayer films with a strong [1 1 1]

crystallographic texture often results in the top surface

of the layers being rough, or wavy, with a period of a

few nanometres. In many multilayer systems this peri-

odicity is transferred through to subsequently deposited

layers, resulting in correlated roughness that usually

becomes worse as more layers are deposited [8]. An

example of this is shown in Fig. 10. The roughness at the

top of the NiFe layer has been transferred up to the next

CoFe layer, as can be seen from the isoconcentration

surface (surface of constant concentration of one or

more elements) shown in Fig. 10(a), plotted at a com-

position of$45 at% Co. A very interesting feature is the

fact that the upper surface of the Cu layer is not con-formal with the upper surface of the underlying CoFe

the deposition of the Cu has acted to smooth out the

correlated roughness [54]. This smoothing effect had

been suggested by Eckl et al. [7] on the basis of in-situ

resistivity measurements but was not confirmed by a

microstructural study.

The question then arises as to how the Cu grows in

order to smooth out the roughness. Eckl et al. pro-

posed that the initial Cu layer was deposited confor-

mally on the underlying Co, followed by nucleation of

islands of Cu in the troughs at the grain boundaries.

While the 3DAP data in Fig. 10 show the Cu

smoothing effect, the actual growth process cannot be

understood just from the 3DAP data as the technique

does not allow us to follow the process as a function

of time. However, this can be achieved using the same

MD simulation technique as discussed above [61].

Fig. 11 shows a series of images obtained during a

MD simulation of the growth of a Cu layer on a

rough CoFe layer [63]. Fig. 11(a) shows that the initial

monolayers of Cu are indeed deposited conformally

on the CoFe surface, as proposed by Eckl et al. The

MD simulations indicate that as the Cu layer gets

thicker, facets with (11 1) crystallographic planes

Fig. 7. (a) 3DAP atom maps of Ni, Co and Cu atoms in the multilayer (volume is $20 nm$20 nm$35 nm), (b) selected-region compositionprofile across two of the layers shown in (a).



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Fig. 8. (a) 3DAP atom map of Co (blue), Fe (yellow) and Cu (red)

atoms in a CoFe/Cu/CoFe trilayer and (b) a composition profile across

the same region (growth direction is left to right in the profile). (Re-

produced from [61] with permission from Elsevier Publishing.)

Fig. 9. (a) Molecular dynamics simulation showing an atom map of Co

(blue), Fe (yellow) and Cu (red) atoms in a a CoFe/Cu/CoFe trilayer

and (b) a composition profile across the same region (MD-simulated

growth direction is left to right in the profile). (Reproduced from [61]

with permission from Elsevier Publishing.)

Fig. 10. (a) Isoconcentration surface at $45% Co for a NiFe/CoFe/Cu/CoFe multilayer, and (b) and (c) 3DAP atom maps of Co (blue), Fe

(yellow) and Cu (red) atoms in the same region. (b) Shows the first few

atoms of the Cu layer and (c) shows the structure up to the point at

which deposition of the Cu ends.

Fig. 11. Molecular dynamics simulation [63] of the growth of a thin Cu

layer onto a rough CoFe layer. (a) The first few monolayers of Cu

grown conformally on the CoFe and (b) and (c) further Cu growth

results in a smooth film surface.



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begin to form at the peaks and troughs of the un-

dulations. The (1 1 1) surfaces are low-energy surfaces

and to maximise their area, Cu atoms lying on the

sides of the troughs diffuse across the growing surface

and fill in the troughs as seen in Figs. 11(b) and (c).

The final result is that the upper surface of the Cu

layer is much flatter than the CoFe surface on whichit grew. Both this MD simulation and the 3DAP data

(Fig. 10(b)) show that deposition of Cu onto a rough

CoFe layer results in relatively little intermixing with

the CoFe. When the next CoFe layer is deposited, the

enhanced interdiffusion at the CoFe-on-Cu interface

serves to remove any remaining correlated roughness

at the upper surface of the Cu layer, as shown in

Fig. 10(c). The combination of the 3DAP data and

the MD simulation can thus be used to provide a full

picture of the way in which the growth of the Cu

layer smoothes out the CoFe roughness and at the

same time yields an intermixed upper Cu surface.

3.3. Effect of oxygen during growth on thin film micro-

structure

One factor that can significantly affect the GMR

properties of thin films structures is the presence of

impurities, such as oxygen, in the sputtering chamber

during deposition. Different proposals have been put

forward to explain this effect, such as a smoothing of

the interfaces during growth via a reduction of surface

free energy [64], a reduction in pinhole density across

the Cu spacer layers [65] and a reduction in grain size

and interface roughness caused by partial oxidation ofthe interfaces [66]. However, to date, relatively little

nanoscale characterisation has been carried out to

support these suggestions. There are therefore several

issues to be clarified with respect to oxygen-doped

multilayer growth such as the amount and position

of residual oxygen trapped in the layered structures

and the exact effects of oxygen on the nature of the

interfaces.

The 3DAP has been used to investigate the thin

film microstructure resulting from the addition of an

oxygen surfactant during growth of the Cu layers in

{Si//Ni alloy5 nm/Co10Fe3 nm/(Cu1:8 nm/Co10Fe3 nm)5/

cap50nm} films [55]. In this structure, the Ni-based seed

layer leads to a h1 1 1i crystallographic orientation inthe films, and this can be seen in the field ion image of

the multilayer in Fig. 12 which shows strong crystal-

lographic orientation along the specimen axis, together

with with the bright (CoFe) and dark (Cu) contrast

regions which indicate the layers in the structure. The

sputter gas pressure during deposition of the Cu and

CoFe layers was of the order of several mTorr, and

deposition of the Cu layers was carried out either in Ar

or in an Ar + oxygen gas mixture with the total pres-

sure being kept constant.

Magnetoresistance values measured on films depos-

ited on unpatterned samples (using a standard four-

point probe measurement) were found to be $1% and$7% for the Ar and Ar + oxygen Cu deposition condi-tions, respectively, as shown in Fig. 13. Fig. 14 shows

composition profiles across the CoFe/Cu layers for de-

positions (a) without and (b) with oxygen present for

deposition of the Cu layers [55]. The most obvious dif-

ference is the degree of intermixing in the Cu layers,

which is considerably reduced for the Ar + oxygen de-

position, resulting in a higher Cu content at the centre of

the layers. As expected from the surface free energy

argument discussed above, the CoFe-on-Cu interface

shows more intermixing than the Cu-on-CoFe interface,

and the effect of the oxygen has been to reduce the width

of both interfaces (see Table 1). See [55] for further de-

tails of the interfacial width measurements in these

structures.

Fig. 12. Field ion image of four repeats of the CoFe/Cu bilayer

showing the expected [1 1 1] crystallographic texture.

0

1

2

3

4

5

6

7

-1500 -1000 -500 0 500 1000 1500

H (Oe)

dR/R(%)

O-doped as-grownO-doped,annealedNo-O, as-grownNo-O,annealed

Fig. 13. GMR curves for as-deposited and annealed CoFe/Cu multi-

layer films for which the Cu layers were deposited either in Ar or in

Ar + oxygen.



11/16

As can be seen from Table 1, the Ar + oxygen depo-

sition has resulted in a decrease in the width of the

CoFe-on-Cu interface by a factor of approximately two

relative to the deposition without oxygen. This clearlyshows that the presence of oxygen suppresses mixing

during growth. Egelhoff et al. [67] have suggested that

this occurs because the oxygen favours the formation of

CoO (or FeO) bonds rather than CuO bonds (the

heats of formation of the oxides are: CoO)58 kcal/mol and CuO)38 kcal/mol [68]). The lowest energystate thus has Co at the surface rather than Cu, leading

to less interfacial mixing. This argument would tend to

suggest that the width of the Cu-on-CoFe interface

might increase in the presence of oxygen. However the

higher cohesive energy of Co compared to that of Cu

may restrict the atomic exchange [61] at the Cu-on-CoFe

interface, so that no increase in intermixing is observed.

Indeed the intermixing is found to be reduced slightly,

which suggests that the presence of oxygen on the

growing surface reduces the number of atomic ex-

changes at the surface. This will also contribute to

the reduction in interface width at the CoFe-on-Cu

interface.

Another interesting observation is that there is anincrease of only approximately a factor of 2 in the re-

sidual oxygen level (average of $260 appm oxygen forthe no-oxygen condition compared to an average of

$530 appm oxygen) for the Ar + oxygen condition [55].The oxygen present in the films grown without added

oxygen is most likely a result of the dissociation of water

vapour in the deposition chamber and of oxygen in-

corporated into the sputter deposition targets. The rea-

son that the incorporated oxygen content increases only

by a factor of two, despite a large difference in the ox-

ygen pressure in the chamber, may be the very low

solubility of O in Cu and Co [69]. Further evidence for

the low incorporation of oxygen for the Ar + oxygen

deposition can be seen in the elemental map shown in

Fig. 15. The oxygen atoms are not positioned prefer-

entially at the layer interfaces but are distributed rela-

tively randomly through the layers, and there is no

evidence for the formation of an oxide phase in these

samples, as has been suggested in the literature [66].

However, there is some evidence for preferential reten-

tion of oxygen near grain boundaries, as previously

suggested [64].

A further effect of the presence of oxygen during

sputtering is a decrease in the conformal roughness of

the layers, as observed previously by Miura et al. [66],which would also be expected to increase the GMR

ratio. An estimate of the conformal interfacial rough-

ness was made from TEM cross-section images by

measuring the maximum peak-to-peak amplitude of the

spatial variations in each layer. The mean amplitude

values (averaged across all five bilayers) are 1.89 (0.08)nm and 1.62 (0.08) nm for the Ar and Ar + oxygendeposition conditions, respectively [55]. The presence of

oxygen is believed to reduce long-range surface diffusion

[70] and thus decrease grain boundary grooving this

would drive the system away from equilibrium and thus

lead to decreased grooving, as observed in the TEM

data.

Fig. 14. Composition profiles across CoFe/Cu multilayers for deposi-

tions of the Cu layers in (a) Ar and (b) Ar + oxygen. The growth di-

rection is from right to left in both profiles. (Reproduced from [55]

with permission from The American Physical Society.)

Table 1

Interface widths for CoFe/Cu MLFs for which the Cu layers were deposited either in Ar or in an Ar + oxygen mixture, for the as-deposited films and

for films annealed at 300 C for 30 min

Cu deposition condition Interface widths (nm) Ave. Cumax Conc. (%)

CoFe-on-Cu Cu-on-CoFe

Ar as-grown 1.21 (0.05) 0.49 (0.03) 65.113.3A r + O2 as-grown 0.55 (0.03) 0.31 (0.02) 77.111.72Ar (300 C/0.5 h) 1.16 (0.09) 0.61 (0.02) 72.17.6A r + O2 (300 C/0.5 h) 3.01 (0.12) 2.38 (0.12) 55.67.4



12/16

3.4. Annealing studies

Diffusivity in Ni/Cu layered films (as-deposited and

annealed) has been studied by Keilonat et al. [21]. Ac-

tivation energies for interdiffusion, as a function of Cuconcentration, were measured and compared well with

previous literature data. Interfacial mixing in NiFe/Cu

multilayers was reported using 3DAP by Larson et al.

[57]. Highly intermixed interfacial regions at the NiFe-

on-Cu and Cu-on-NiFe interfaces were reported al-

though the geometry of the samples was such that the

interfaces were parallel to the specimen axis (Fig. 1(b))

so that the accuracy of the measurements can only be

considered to be $1 nm.The interaction of an Al film deposited onto Ni (using

ion beam sputtering and electron beam evaporation) has

been studied by Jeske et al. [29] also using the method of

film deposition onto a preformed needle of Ni or W. It

should be noted that although Ni/Al layers are not of

interest per se in magnetic information storage systems,

the formation of the aluminium oxide tunnel barriers in

spin tunnel junction structures [3] is usually carried out

by depositing metallic Al on the underlying ferromag-

netic layer (such as a Ni, Fe or Co alloy) and then ox-

idising. Interdiffusion of the Al and the underlying layer

can therefore occur, which can lead to a decrease in the

spin polarisation in the ferromagnetic layer. In the as-

deposited state, Jeske et al. identified an initial solid

solution zone about 1 nm in width in both types of

samples. After annealing, an off-stoichiometry B2 NiAl

phase was observed to form in samples formed by

electron beam deposition of an Al layer onto a Ni nee-

dle, in which no grain boundaries were observed.However, in samples where a Ni/Al bilayer was sputter

deposited onto a W needle, NiAl3 regions were observed

to form at grain boundaries in the Al layer. The authors

ascribed this to a difference in the flux of Ni into and out

of the thin inter-reaction layer at the Al/Ni interfaces.

In films containing grain boundaries, rapid diffusion

of Ni along the grain boundaries to the NiAl3 precipi-

tates leads to depletion of nickel in the interfacial region

and thus inhibits the formation of the intermetallic

phase [29].

Annealing of the Cu2 nm/Co2 nm multilayer films de-

scribed above has been studied by the current authors.

Samples that were annealed for 1h at 360 C showed an

increase in grain size but no apparent change in the in-

terlayer mixing between Cu and Co [71]. Studies of the

CoFe/Cu films deposited in the presence of oxygen,

following annealing, tend to confirm that the presence of

oxygen drives the system further from equilibrium. This

can be seen from the changes in the GMR ratio for the

two types of film, as shown in Fig. 13. For the film de-

posited in Ar, the decrease in GMR ratio following

annealing at 300 C for 30 min was from 1% to 0.5%,

whereas for the film deposited in Ar + oxygen, the de-

crease was from 7% to 0.2%. The changes to the trans-

Fig. 15. Atom map from 3DAP data for the as-deposited CoFe/Cu

MLF deposited using Ar + oxygen. The oxygen atoms (large, spheres)

and Cu atoms (small, dark spheres) are shown. The presence of a

columnar grain boundary is indicated. (Reproduced from [55] with

permission from The American Physical Society.)

0

10

2030

40

50

60

70

80

90

100

0 1 2 3 4 5 6 7 8

Distance (nm)

Con

centration(at%)

%Co %Cu

0

10

20

30

40

50

60

70

0 5 10 15 20 25 30 35

Distance (nm)

Concentration(at%)

Co% Cu%

(a)

(b)

Fig. 16. Composition profiles (showing only Co and Cu) across CoFe/

Cu multilayer films following annealing at 300 C for 30 min (a) for

deposition of Cu layers in Ar, and (b) for deposition of Cu layers in

Ar + oxygen.



13/16

port properties can be accounted for by changes to the

microstructure that occur during annealing. Fig. 16

shows composition profiles taken from 3DAP analyses

across annealed CoFe/Cu multilayer films with the Cu

layers deposited (a) in Ar and (b) in Ar + oxygen. As can

be seen from Table 1, the increase in interfacial mixing is

relatively small for the film deposited in Ar (the Culayers are still well defined), but for the film deposited in

the presence of oxygen, the CoFe-on-Cu interface width

increases by a factor of 5, and the Cu-on-CoFe in-terface by a factor of 6.8. This is observed even thoughthe distribution of the oxygen surfactant does not ap-

pear to change after annealing, as can be seen from a

comparison of the atom maps shown in Fig. 15 (as-de-

posited) and Fig. 17 (following annealing at 300 C for

30 min). It is likely that in the layers that are grown with

oxygen, and thus are further from equilibrium, the re-

duced conformal roughness provides a driving force for

intermixing at the CuCoFe interfaces to occur during

annealing. On the other hand intermixing in the layers

grown without oxygen appears limited to that which

occurs by atomic exchange during sputter deposition

[61]. However it is not clear from our results whether the

intermixing has occurred by long-range changes in the

interface position in order to establish the equilibriumshape, or whether the non-equilibrium configuration

raises the free energy of the system sufficiently to make

the CoFe and Cu layers miscible [56].

3.5. Media materials

New information storage systems require not only the

development of new read/write heads but also of new

storage media. The media used currently in hard disk

drives consist of CoCr alloys (to which additional al-

loying elements such as Ta are frequently added) in

which the information is stored with the magnetisation

vectors in the plane of the film (longitudinal media). The

high storage density is achieved in the media by having a

polycrystalline film in which the grains are small and are

ideally magnetically decoupled from each other. Mag-

netisation studies of the films suggested that this occurs

because of compositional inhomogeneities in the CoCr

alloy layer [72] and this was confirmed by TEM studies

[73]. However the exact way in which the inhomogenei-

ties were distributed in three-dimensions had not been

well understood. The grain size in this type of material is

typically of the order of a few nanometres and this is

therefore an ideal topic to be addressed by 3DAP.

A small number of atom probe field ion microscopystudies had been made of CoCr alloy thin films, both by

deposition of the films on curved surfaces [22] and using

the Hasegawa method [32] and more recently studies

have been made using 3DAP. Hono et al. [74] used the

Hasegawa method to prepare thin films of a Co

22at%Cr alloy that had been deposited on a heated

substrate, as would be the case for the thin film media

layer on a hard disk. Their results showed that the films

consisted of a ferromagnetic phase and a paramagnetic

Cr-rich phase, both with a lamellar-type structure, which

coexist within each grain of the film. The lamellar shape

of the Co-rich ferromagnetic phase results in the particles

having a high shape anisotropy which increases their

thermal stability, and the presence of the paramagnetic

phase acts to isolate the ferromagnetic lamellae mag-

netically from each other. Further work on a similar

system was carried out by Nishikima et al. [75] on a Co

12Cr2Ta alloy thin film deposited at elevated temper-

ature. The addition of the Ta to the alloy resulted in

segregation of the Cr to the grain boundaries, leading to

Co-rich ferromagnetic grains, as shown in the Co con-

centration map in Fig. 18(a), enveloped by a Cr-rich non-

magnetic phase, seen in Fig. 18(b), but still with some

non-uniformity in the Cr composition within the grains.

Fig. 17. Atom map from 3DAP data for the CoFe/Cu MLF deposited

using Ar+ oxygen, following annealing at 300 C for 30 min. The

oxygen atoms (large, pale spheres) and Cu atoms (small, dark spheres)

are shown.



14/16

The authors proposed that the enhanced segregation

occurred because the Ta atoms are larger than both Co

and Cr and thus reduced the migration energy of the Cr

to the boundaries. The Ta was found to be uniformly

distributed through the film, as seen in Fig. 18(c), and did

not segregate to the grain boundaries with the Cr.

One recent innovation has been the use of perpen-

dicular magnetic storage in which the bits of informa-

tion are written into a magnetic layer with the

magnetization vectors pointing perpendicular to the

plane of the magnetic film [76]. This allows for a con-

siderably higher information density than can be

achieved with conventional longitudinal storage media

[77]. A number of materials are currently being consid-

ered for perpendicular storage media, one of the most

popular being a Co/Pt or Co/Pd multilayer film in which

the Co layers are very thin (typically


15/16

composition profile taken from 3DAP data for the Co/

Pd multilayer film. Four Pd layers and three Co layers

are visible plus part of the NiFe cap (which shows atomic

planes). As can be seen, the Co and Pd layers are not

clearly differentiated; the Co concentration varies from

$20% to $40% across the Co layers. The binary phasediagram for CoPd [69] shows complete solubility be-

tween the two elements and, in addition, the deposition

was carried out in a relatively low Ar pressure, leading to

high adatom energy and increased intermixing. Thus it is

not entirely surprising that a high degree of intermixing

was observed in the Co/Pd samples. However, such a

high degree of intermixing was unexpected from a con-

sideration of the magnetic data, which show a degree of

perpendicular anisotropy greater than would be expected

without inclusion of an interfacial contribution arising

from the Co/Pd interfaces.

4. Summary and conclusions

The aim of this paper has been to show that 3DAP

analysis is a technique that can be very effectively ap-

plied to the understanding of the correlation between

microstructure, composition and physical properties in

thin films with applications in information storage. The

technique is particularly appropriate to films in which

the composition is inhomogeneous on the nanometre

scale, for example because the film consists of thin lay-

ers, or because of chemical segregation at grainboundaries, and for which the properties can vary

greatly even with atomic scale differences in chemical

distribution. The validity of the technique has been

demonstrated by comparison with molecular dynamics

models of the growth mode in CoFe/Cu layered films,

which allows us to place reliance on the experimental

data that are obtained using 3DAP.

However, as for all techniques, there are outstanding

issues with using 3DAP for the analysis of thin films, of

which one is specimen preparation and another is the size

of the analysis volume. In order to make valid correla-

tions of the composition profile and microstructure with

the magnetic and transport properties, the authors be-

lieve that it is crucial to analyse films that have been de-

posited on planar substrates. As has been illustrated

above, this is now an achievable goal, and the results that

have been obtained have been invaluable in enabling

phenomena such as the effect of surfactants on GMR in

CoFe/Cu multilayers to be elucidated. The analysis vol-

ume in 3DAP remains limited, although novel instrument

designs are currently helping to address this issue too. In

any case, careful comparison with other techniques such

as TEM can ensure that the data that are obtained using

3DAP are typical of the structure as a whole.

In conclusion, we believe that atomic-scale nano-characterisation using 3DAP, possibly combined with

molecular dynamics modeling, provides a unique tool

that can be applied both to the design of nanostructured

magnetic materials for information storage applications

and to the further understanding of the fundamental

physical properties on which their operation relies.

Acknowledgements

The authors thank Drs. T.P. Nolan, P.H. Clifton,

S.P. Bozeman, H.L. Brown, E.W. Singleton, R. Grove,

R. Hipwell (Seagate Technology LLC), M. Abraham, O.Dimond, R.M. Langford and G.D.W. Smith FRS

(University of Oxford) for their various contributions to

this research. The authors especially thank Dr. X.W.

Zhou (University of Virginia) for his provision of the

MD model results on a rough substrate. We are also

grateful to the EPSRC for funding.

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