Surface Science 501 (2002) 155–167
www.elsevier.com/locate/susc

Scanning tunneling microscopy and spectroscopy study
on the submonolayer growth of Mn on Fe(0 0 1)
M.M.J. Bischoﬀ, T. Yamada, A.J. Quinn 1, H. van Kempen *Research Institute for Materials, University of Nijmegen, Toernooiveld 1, 6525 ED Nijmegen, Netherlands
Received 12 October 2001; accepted for publication 7 December 2001

Abstract
The submonolayer growth of Mn on Fe(0 0 1) in the temperature range between 50 and 200 °C is studied by scanningtunneling microscopy and scanning tunneling spectroscopy. For growth temperatures above 100 °C, atomically resolved STM images with chemical contrast clearly reveal the incorporation of Mn atoms in the Fe(0 0 1) substrate. The
fraction of place exchanged Mn atoms is observed to increase with growth temperature. Although on islands a cð2  2Þ
structure forms locally which is attributed to an ordered MnFe surface alloy, long range order could not be obtained forthe growth temperatures and coverages studied. Spectroscopy results are presented for clean Fe(0 0 1), pure Mn adislands, single incorporated Mn atoms and cð2  2Þ-ordered MnFe areas. In contrast to embedded Cr atoms in theFe(0 0 1) surface [Phys. Rev. Lett. 76 (1996) 4175], isolated embedded Mn atoms do not lead to double peak structures
in dI=dV curves. Nevertheless, on the ordered cð2  2Þ MnFe structure and the pure Mn monoatomic islands surface
states are detected as peaks in the dI=dV and ðdI=dV Þ=ðI=V Þ curves. However, due to the strong inﬂuence of a tipdependent background on these peaks, the corresponding surface state energies cannot be found directly from thosecurves. The real surface state energies were recovered by normalizing the measured dI=dV curves by ﬁtted quadratic
backgrounds. Thus, surface state energies of about þ0.35 and þ0.25 eV are estimated for pure Mn islands and alloyed
MnFe areas, respectively. Ó 2002 Elsevier Science B.V. All rights reserved.
Keywords: Epitaxy; Iron; Manganese; Scanning tunneling microscopy; Scanning tunneling spectroscopies; Surface electronic
phenomena (work function, surface potential, surface states, etc.); Surface structure, morphology, roughness, and topography;
Whiskers

1. Introduction*

Corresponding author. Tel.: +31-24-3653499; fax: +31-243652190.
E-mail address: hvk@sci.kun.nl (H. van Kempen).
1
Present address: National Microelectronics Research Centre, University College, Cork, Ireland.

During the last decade a large number of experimental and theoretical studies have been dedicated to elucidating the interrelation between
structure and magnetic/electronic properties oftransition metals. Molecular-beam-epitaxial techniques allow growth of low-dimensional andstrained metal structures often with novel magnetic and electronic properties. A classic example

0039-6028/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved.
PII: S 0 0 3 9 - 6 0 2 8 ( 0 1 ) 0 1 9 4 6 - X

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M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

of a system where structure, magnetism, andelectronic eﬀects are interwoven in a complex
fashion is Mn/Fe(0 0 1) [1–22].
For Mn ﬁlms thicker than 3–5 monolayers(ML), both theoretical and experimental studies
conﬁrm the layered antiferromagnetism of the Mn
layers (see e.g. [1,2] and references therein). The
Mn ﬁlm grows in a distorted bcc structure with thein-plane lattice constant equal to the Fe(0 0 1) bulk
lattice constant of 0.287 nm but with the outof-plane lattice constant expanded to 0.323 nm,
i.e. body-centered-tetragonal [3,4]. Layer-by-layer
growth is followed by a transition to threedimensional growth depending on growth temperature and substrate quality [1,2,5,6]. Between 10
and 17 ML a cð4  4Þ superstructure was reported
by Andrieu et al. [6]. Finally, thick Mn ﬁlms
(above 17 ML in [6]) relax in the Mna structure
which is the stable bulk phase of Mn at room
temperature (its cubic unit cell includes 58 atoms).
In the ultrathin limit, however, conﬂicting results have been reported in both theoretical andexperimental studies. Andrieu et al. showed by Xray magnetic circular dichroism experiments thatin the submonolayer regime the Mn atoms couple
ferromagnetically to each other and couple ferromagnetically to the substrate [6,7]. They also
showed quite convincingly that oxygen contamination changes the coupling to the substrate toantiferromagnetic [8] which could explain the
conﬂicting experimental results of other groups [9–11]. Furthermore, extended X-ray absorption
spectroscopy (EXAFS) showed that the ﬁrst
monolayer grows with an interlayer distance of
0.148 nm [6]. However, these experiments focussed
on room temperature grown Mn ﬁlms.
Recent ab initio calculations showed that thecð2  2Þ antiferrimagnetic conﬁguration and thepð2  2Þ ferromagnetic conﬁguration (i.e. three
Mn atoms ferromagnetically coupled and one antiferromagnetically) are nearly degenerate [12–16].
Calculations for a bilayer surface-ordered MnFe
alloy with cð2  2Þ arrangement showed that thismixed conﬁguration is as stable as the groundstates of the pure Mn conﬁgurations [16]. Calculations for the dilute limit of 3d atoms on theFe(0 0 1) surface by Nonas et al. showed that Mn
atoms (and in fact all 3d transition metals exceptCo) have a tendency to place exchange with Fe
substrate atoms [17]. For Cr impurities these results were in perfect agreement with the scanningtunneling microscopy (STM) and spectroscopy(STS) study of Davies et al. [18,19] and even theexperimentally observed repulsion of surface
nearest and next-nearest Cr neighbors could be
reproduced in the calculations. In contrast to Cr
impurities, for Mn only a strong repulsion on
surface nearest neighbors was predicted whose
origin was found to be frustration: the surface
nearest-neighbor Mn atoms favor antiferromagnetic coupling with each other but this is hinderedby the strong antiferromagnetic coupling with thesubstrate Fe atoms. The embedded 3d atoms in theFe(0 0 1) surface were theoretically found to giverise to characteristic double peak structures in thevacuum local density of states (LDOS) around the
Fermi level [20] and therefore chemical identiﬁcation as was shown for Cr on Fe(0 0 1) [18] mightalso be possible for these elements. However,compared to the Fe(0 0 1)-derived surface statepeak which is always present, the additional 3dimpurity peak is predicted to be strongest for theelements at the beginning of the 3d-series (i.e. V
and Cr) and weakest for Mn [20].
The interfacial alloying of the Mn/Fe(0 0 1)
system was experimentally veriﬁed by extremely
surface sensitive proton-induced Auger electron
spectroscopy measurements: Igel et al. found thatfor growth around 300 °C only 40% of the deposited monolayer Mn atoms could still be foundon top of the surface layer implying that 60% of
the deposited Mn atoms underwent place exchanges with the substrate Fe atoms [21]. Increased intermixing for a two monolayer Mn ﬁlm
on Fe(0 0 1) was also reported by Walker andHopster after annealing above 150 °C [22].
In this paper, we report on an STM and STS
study on the growth of submonolayer Mn ﬁlms on
Fe(0 0 1) whiskers in the temperature range between 50 and 200 °C. Atomically and chemicallyresolved STM imaging allow us to study the intermixing behavior of Mn atoms in the Fe(0 0 1)
surface as a function of growth temperature.
Motivated by the chemical identiﬁcation achieved
by STS for the Cr/Fe(0 0 1) system [18,19], STS
measurements will be presented for all structures

M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

observed. For the ﬁrst time, the inﬂuence of a tip
dependent background on the apparent peak energies in dI=dV and ðdI=dV Þ=ðI=V Þ curves is discussed in the context of real experiments.

2. Experimental
The experiments were performed under ultrahigh vacuum conditions (base pressure less than
5  1111 mbar). Single crystal Fe(0 0 1) whiskers
with typical sizes of 7  1  1 mm3 grown at
Simon Fraser University (Burnaby, Canada) wereused as substrate material [23]. The whiskers werecleaned by repetitive cycles of Ar ion sputtering(750 eV) around 750 °C. In the ﬁnal cleaning step,
long post-anneals after sputtering were avoided.
The concentration of impurities (mainly oxygen)
was typically below 0.01 ML as estimated by Auger
electron spectroscopy and atomically resolved STM
images. Furthermore, this treatment is suﬃcient toheal the sputter damage and terraces with widths
up to 400 nm are generally observed. Mn was
sublimated from a Knudsen cell at a rate of 0.2
ML/min as measured by a quartz crystal microbalance. In this work, 1 ML is deﬁned as a complete overlayer with Fe(0 0 1) bulk lattice constant(1:214  1015 atoms cm2 ). By proper shielding of
this cell with a liquid nitrogen cooled baﬄe andcomprehensive degassing, the pressure was maintained below 3  1010 mbar during growth. The
sample was annealed using a resistive heater which
limited the maximum growth temperature toaround 200 °C. Temperature measurements were
performed by a thermocouple which was not in
direct contact with the sample. Therefore, the estimated error in the temperature measurements is
relatively high: 30 °C at 200 °C.
After growth the sample was transported as
quickly as possible to the room temperature STM(Omicron UHV STM-1) in the analysis chamber.
In this small chamber the background CO pressureis estimated to be 4  1012 mbar from residual gas
analysis by mass spectrometry. This low CO partial pressure is quite important considering the fastdissociation of CO on Mn ﬁlms as discussed byAndrieu et al. [8]. Electrochemically etched W
tips were used which were cleaned and sharpened

157

in situ by Ne self-sputtering [24]. Field emission
spectroscopy conﬁrmed the sharpness of the tips
(radii less than 10 nm). STS measurements were
performed by recording IðV Þ curves under open
feedback conditions. To minimize the inﬂuence ofthe drift (0.3 nm/min shortly after sample preparation), these IðV Þ curves were recorded as
quickly as possible (typically 52 ms for a 100 point
curve). The drift is mainly caused by contraction(after heating) of the whisker along its long axis
relative to the Ta sample holder. The relatively
large distance (5 mm) between tip position andthe point where the whisker is ﬁxed to the holder is
not favorable in this perspective. To be able tocompare topographic features with spectroscopic
ones, IðV Þ curves were measured simultaneouslywith conventional constant current images but at a
reduced resolution of typically 75 points  75 lines
to avoid strong distortions due to drift. The IðV Þ
curves were numerically diﬀerentiated using a 5point window to obtain dI=dV .

3. Results and discussion
3.1. Submonolayer growth at low temperature: pure
Mn islands
An STM image of the Fe(0 0 1) surface after
deposition of 0.2 ML (quartz crystal calibration)
Mn at a substrate temperature of 50 °C is shown in
Fig. 1a. The island coverage is estimated to be 0.17
ML from histogram analysis. However, the ragged
step edge in Fig. 1a shows that small Mn islands
nucleated at the step edge. These islands can easily
be missed in a histogram analysis and therefore,
coverages were determined on large terraces(>200  200 nm2 ) where the inﬂuence of Fe(0 0 1)
steps can be neglected. The histogram analysis also
shows that the height of the Mn islands is slightly
higher (0:156  0:005 nm) than the Fe(0 0 1) stepheight (0:144  0:005 nm). To avoid electronic effects, step heights were measured at large negative
bias voltages where the STS results did not reveal
any features (see below). Furthermore, apparent
barrier height measurements were performed
which showed that the eﬀect of diﬀerent local
work functions on the apparent step heights can be

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Fig. 1. (a) Constant current STM image of 0.2 ML Mn deposited on Fe(0 0 1) at 50 °C (tunneling current I ¼ 0:5 nA, sample voltage
Vs ¼ 0:5 V, scan size is 30  30 nm2 ) showing monolayer high islands (0.156 nm high). The white circle marks a small Mn island
which nucleated at the step edge. A dI=dV curve is obtained at each pixel (75  75) of this image. (b) Averaged dI=dV curves representative of four diﬀerent features in (a). Curve 1 is measured at the Fe(0 0 1) terrace, curve 2 at the islands, curve 3 at the steps, andcurve 4 at the oxygen impurities. Numbers refer to numbers in (a), (c) and (d). (c) and (d) show dI=dV maps at Vs ¼ þ0:16 and þ0.99
V, respectively.

excluded [25]. Therefore, these step heights are
believed to reﬂect the real interlayer spacings.
A dI=dV curve was measured at every pixel ofFig. 1a. Four diﬀerent types of dI=dV curves canbe distinguished and averages of these curves(typically obtained from 10 single curves) are
shown in Fig. 1b. Curve 1 is the average of thedI=dV curves measured on the Fe(0 0 1) terrace. Itshows a strong peak at þ0.16 V which is attributed
to a localized minority spin band surface state ofdz2 symmetry. This Fe(0 0 1) surface state was ﬁrst
detected in tunneling spectroscopy by Stroscio
et al. [26] and slightly later by Biedermann et al.
[27]. Compared to the earlier work for which thepeak in dI=dV was reported at þ0.17 V [26], thepeak is detected at a slightly lower energy in themeasurement shown in Fig. 1b. Furthermore,
the dI=dV feature is characterized by an almost
Gaussian line shape in contrast to a weak peak on
an exponential background reported in [26,27]. In
the present study, it is found that the surface statepeak can be observed between þ0.13 and þ0.20 V
depending on the voltage dependence of the
background. The peak is most pronounced for low
background conditions, although in this case theapparent peak energy is lower than þ0.17 V.
Curve 2 represents the average of the dI=dV curves
measured on the Mn islands. A clear peak is re-M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

solved at þ0.45 V which gives strong evidence forthe diﬀerent chemical nature of the islands. The
average of the dI=dV curves measured at theFe(0 0 1) step edges shows that the surface state iscompletely quenched at monoatomic steps (curve3). The dI=dV curves measured at island edges(not shown) show the same behavior. The observation that the high energy dI=dV backgroundmeasured at the Fe(0 0 1) steps resembles the
background observed on the islands better thanthe background observed on the terrace seems to
imply that the Fe(0 0 1) step edges are completely
wetted by Mn.
Fig. 1c and 1d shows maps of dI=dV at thevoltage of the Fe(0 0 1) surface state (þ0.16 V) andat a voltage which ﬁlters out the Mn peaks (i.e.
þ0.99 V) most clearly. The diﬀerent electronic
structure of the islands, steps and terraces is obvious from Fig. 1c. Line proﬁles over the steps (notshown) show that the surface state is quenched ona length scale of order 0.7 nm. This length scale iscomparable to the broadening of the monoatomic
steps in the constant current image of Fig. 1awhich was measured at the same setpoint and withthe same tip conﬁguration. The spherical depressions in the dI=dV map of Fig. 1c are related tooxygen impurities on the Fe(0 0 1) surface. Theirconcentration is lower than 0.01 ML even somehours after deposition which is in contrast to theobservations of Andrieu et al. who reported anoxidation rate of 0.03 ML/h for submonolayer Mnﬁlms on Fe(0 0 1) [8]. The averaged dI=dV curve
representative of these impurities is also shown in
Fig. 1b (curve 4) and reveals that at a single oxygen impurity the surface state dI=dV peak shifts
0.03 V towards higher energy and has a 30%
lower amplitude. The width of the oxygen-induced
depressions in the dI=dV map is 0.7 nm (fullwidth at half maximum). The same order of peakshift and the same range of inﬂuence was recently
reported for carbon impurities on V(0 0 1) [28].
This range is related to the lower resolution at therelatively high tunneling resistances used during
STS [29] and must therefore be considered as an
upper limit for the range of inﬂuence of the steps
and the impurities on the electronic structure.
The dI=dV map at Vs ¼ þ0:99 V (Fig. 1d)shows the homogeneity of the Mn islands and the159

Fig. 2. Atomically resolved STM image obtained on Fe(0 0 1)
covered with 0.2 ML Mn at 50 °C (I ¼ 3 nA, Vs ¼ 5 mV,
7:7  4:1 nm2 ). A stepped grey scale has been used: the black–white range corresponds to 0.01 nm on the terrace (marked 1)
and the island surface (marked 2). The black–white island
borders are artifacts of the grey scale. The image has been averaged over 3  3 data points. A very weak pð1  1Þ lattice canbe resolved of which the corrugation is less than 2 pm. The
black dots indicate the positions of the atoms. The close-packed
directions are also shown.

presence of Mn islands connected to the Fe(0 0 1)
step. The island homogeneity also follows from a
careful check of the single dI=dV curves obtainedon the islands.
An atomically resolved STM image of thissample is shown in Fig. 2. An undisturbed pð1  1Þlattice with an extremely low corrugation (below
2 pm) was always observed. In contrast to
growth at higher temperatures, for growth at 50 °C
no indications for Mn incorporation in theFe(0 0 1) substrate layer were found. The closepacked directions of the Fe(0 0 1) lattice are indicated in this ﬁgure. Note that they do not make aperfect angle of 90° which is attributed to a thermal drift of typically 0.3 nm/min shortly after
deposition. In this paper, the STM images are notcorrected for this drift.
In conclusion, from both the STM and STS
results of the growth at 50 °C it is evident that nointermixing occurs at this temperature.
3.2. Submonolayer growth at elevated temperatures: surface alloying
3.2.1. Chemically resolved STM
An STM image of the Fe(0 0 1) surface after
deposition of 0.37 ML Mn at a substrate temperature of 150 °C is shown in Fig. 3. The terraces and the islands (0:156  0:005 nm high)now look completely diﬀerent compared to thegrowth at 50 °C. At the low tunneling resistance

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M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

Fig. 3. STM image (I ¼ 1:4 nA, Vs ¼ 5 mV, 14:9  14:9 nm2 )of 0.37 ML Mn deposited on Fe(0 0 1) at 150 °C. A stepped grey
scale has been used: the black–white range corresponds to 0.1
nm on the terrace (marked T) and the island surfaces (marked
I). Chemical contrast is achieved: asymmetric depressions on
the terrace are attributed to embedded Mn atoms and protrusions on the islands to incorporated Fe atoms. The protrusions
on the islands are locally ordered in a cð2  2Þ structure. The
area indicated by the white arrow does not show chemical
contrast. The close-packed directions correspond to those indicated in Fig. 2.

used (i.e. Vs ¼ 5 mV, I ¼ 1:4 nA), it is generally
not possible to keep a stable tip conﬁguration
while scanning over the islands. Consequently, tipchanges are frequently observed in Fig. 3. No
lateral shift occurs in the image after a tip change
and the perpendicular shift is typically less than 0.1
nm. This indicates that only the tip apex atom is
aﬀected. Nevertheless, with particular tip conditions chemical contrast is achieved in the larger
part of the image (e.g. no chemical contrast isobserved in the area marked by the white arrow).
The terrace is covered with asymmetric depressions which are 40 pm deep, while on the islands 20
pm high protrusions are observed. The strong
dependence of the appearance of these features on
tip conditions and the rather low stability of the tip
conﬁgurations leading to these features makes itlikely that the same mechanism is responsible forthe chemical contrast as was reported by Schmid
et al. in their studies on PtNi alloyed surfaces [30–32]. In these studies, a tip–sample interaction mediated by an unknown adsorbate on the tip was
believed to lead to chemical contrast. The adsorbate (e.g. sulphur or oxygen impurity atoms)
forms a precursor of a chemical bond, similar to ahydrogen bridge, more eﬀectively with one type of
surface atom than the other.
The depressions on the terrace are ascribed tosingle embedded Mn atoms and the protrusions on
the islands to Fe atoms which have been released
from the terrace after place exchange events. For
this particular amount of Mn deposited at 150 °C(i.e. 0.37 ML), the concentration of Mn atoms in
the terrace is 5%, while the concentration of Fe
atoms in the islands is estimated to be 21%. 0.05
ML Mn of the 0.37 ML deposited, i.e. 14% of the
Mn deposited, has therefore been place exchanged
at this growth temperature. As can clearly be seenfrom this ﬁgure, the protrusions on the islands
order locally in a cð2  2Þ structure which indicates the absence of a surface next-nearest neighbor repulsion as was reported for Cr/Fe(0 0 1)
[18,19] and which is in agreement with theoreticalpredictions [17].
From STM images like Fig. 3 the concentrationof Mn atoms embedded in the Fe(0 0 1) terrace andFe atoms incorporated in the islands can be estimated as a function of growth temperature withreasonable accuracy. The results are shown in
Table 1. The onset substrate deposition temperature for surface alloy formation is 100 °C, while
submonolayer growth at 200 °C leads to islands
which include large numbers of intermixed Fe
atoms. At this growth temperature the step heightof the island is found to be reduced to 0:148  0:005
nm which is considered to be a consequence of thehigh Fe concentration in the islands. From theexp
concentration of Fe atoms in the islands (CFeðiÞ
),the concentration of Mn atoms in the terracecal(CMnðtÞ
) can be estimated under the assumptionthat the Mn atoms stay in the surface layer: i.e.
expcalCMnðtÞ
¼ CFeðiÞ
 coverage. The calculated concentrations are shown in the table and the good
agreement with the observed Mn concentrations
exp
on the terrace (CMnðtÞ
) shows that the Mn atoms do
indeed stay in the surface layer for the growthM.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

161

Table 1
Amount of Mn deposited (coverage) as determined from histogram analysis on large terraces, concentration of Fe atoms in the islands
exp
expCFeðiÞ
, and concentration of embedded Mn atoms in the Fe(0 0 1) terrace CMnðtÞ
as a function of growth temperature Tgr . The concal
centration of embedded Mn atoms in the terrace CMnðtÞ is calculated under the assumption that all place exchanged Mn atoms areexp
cal
incorporated into the ﬁrst substrate layer (CMnðtÞ
¼ coverage  CFeðiÞ
). The fraction of deposited Mn atoms which exchange place with
exp
substrate Fe atoms (Cex ¼ CMnðtÞ /coverage) is also shown
Tgr (°C)

Coverage [ML]expCFeðiÞ
(at.%)

expCMnðtÞ
(at.%)calCMnðtÞ
(at.%)Cex (at.%)

50  10100  20150  30175  30200  30

0:17  0:04
0:38  0:04
0:37  0:04
0:51  0:04
0:36  0:04–
18  221  217  560  5–42
52
12  220  4

0
71
81
93
22  30
11  614  624  456  13temperatures studied. For Cr deposited on
Fe(0 0 1) it was reported that for growth at 300 °C
and coverages between 0.2 and 1 ML only 0.1
ML Cr atoms were found on the surface while for
growth at lower temperatures the degree of intermixing was reduced [18].
For STM experiments performed at room
temperature, the atomic processes leading to thesurface alloy are too rapid to be imaged directly.
The low density of steps on Fe(0 0 1) whisker surfaces and the large calculated energy gain for placeexchange (0.65 eV for Mn compared to 0.57 eV for
Cr [17]) makes it very likely that Mn atoms areincorporated in the Fe(0 0 1) surface without
moving to steps ﬁrst. The incorporation may involve either a direct place exchange mechanism[33] or annihilation of vacancies which were originally formed at kinks or steps. In the latter case,the diﬀusion of the vacancies from the kink or stepto the terrace must be faster than the diﬀusion of
the Mn adatoms to the steps. If the adatoms diffuse faster, the vacancies are annihilated by Mnadatoms close to step edges and consequently thehighest concentration of incorporated Mn atoms
will be found here. This was shown to be the casefor e.g. Mn on Cu(0 0 1) [34], Pb on Cu(1 1 1) [35]and Pb on Cu(0 0 1) [36]. In the present study, thedistribution of incorporated Mn atoms is homogeneous on the terrace and there is no evidence forhigher Mn concentrations around islands andsubstrate step edges even for growth at 100 °C.
However, the observation that pure Mn islands
form at 50 °C while the onset for Mn incorporation is 100 °C seems to suggest that diﬀusion
dominates over place exchange for growth tem-

peratures below 100 °C. Therefore, it cannot be
excluded that Mn adatoms diﬀuse to islands ﬁrst
and place exchange with substrate atoms at stepedges and kink sites [34]. Subsequent step ﬂuctuations and vacancy-mediated diﬀusion may lead toa more homogeneous distribution of the Fe atoms
in the Mn islands and of the embedded Mn atoms
in the Fe(0 0 1) substrate [34].
From Table 1 the fraction of place exchanged
Mn atoms (Cex ) can be estimated by dividing theobserved number of Mn atoms embedded in theexpterrace (CMnðtÞ
) by the amount which had been
deposited (coverage). This fraction is shown as afunction of temperature both in the last row of thetable and in Fig. 4. It clearly shows that forthe coverage regime studied, the fraction of placeexchanged Mn atoms increases as a function oftemperature. However, the place exchanged frac-Fig. 4. Concentration of deposited Mn atoms which have placeexchanged with substrate Fe atoms (Cex ) as a function of
growth temperature.

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M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

tion cannot be ﬁtted by a simple activation law (i.e.
Cex  expðEa =kT Þ, with Ea an activation energy)
showing the complexity of the processes involved.
It might be wondered if for growth at even higher
temperatures the fraction of place exchanged Mn
atoms increases to 100% and if Mn atoms eventually do dissolve into the bulk. However, oursetup does not allow growth at higher temperatures and post-annealing is believed to lead todiﬀerent results.
3.2.2. STS on alloyed surface layer
In order to study the electronic structure of theMn atoms embedded in the terrace and the Featoms incorporated in the islands, STS measurements were performed. As shown in Fig. 5 thedI=dV curves measured on the terrace protrusions(curve 2), i.e. incorporated Mn atoms, do not diﬀermuch from those obtained on the clean Fe(0 0 1)
terrace (curve 1). Both curves, which are averagesof typically 10 single dI=dV curves, show a peak at

Fig. 5. STS measurement on isolated Mn atoms in the Fe(0 0 1)
terrace (I ¼ 0:5 nA and Vs ¼ 0:5 V). Curves 1 and 2 are representative of dI=dV curves measured on the clean terrace and
Mn atoms, respectively. The insets show the corresponding
topographic image (top left: scan size is 4:8  4:4 nm2 , with thistip condition the Mn atoms are visible as 5 pm high protrusions) and the dI=dV map at Vs ¼ 0:17 V (bottom right).

þ0.17 V, although this peak has a 25% lower intensity on the Mn atoms. The insets in this ﬁgureshow the topographic image taken at the setpointof spectroscopy (top left) and the corresponding
dI=dV map at the surface state energy þ0.17 V(bottom right). The minima in this dI=dV map arecorrelated with the protrusions in the topographic
image.
The topographic image shows that at the STS
setpoint the Mn atoms have an apparent height of
5 pm. Since the Mn atoms can be imaged either as
protrusions or depressions depending on tunnelingconditions, it seems reasonable to ascribe thisapparent height to electronic eﬀects and/or tip–
sample interaction eﬀects rather than to a realgeometric eﬀect. However, care has to be taken
since it is not trivial to distinguish between geometric and electronic eﬀects in STM images (e.g.
see [37]). It should also be realized that the atomic
radius of Mn is 9 pm larger than of Fe [38] andbuckling eﬀects cannot be excluded [12,13] (although for the antiferromagnetically coupled
substrate Mn atom the interlayer distance is always lowered in the calculations, 3 pm in [12] and0.3 pm in [13]). Nevertheless, if these eﬀects are
ignored, the apparent height of 5 pm, thus, implies
that the tip retracts 5 pm when over the Mn atom.
This retraction leads to a 10% drop in the tunneling conductance, expð2jDÞ ¼ 0:90 (with j ¼
1  1010 m1 and tip retraction D of 5 pm) [29]. The
reduction of 25% on top of the Mn atom cantherefore only be partially ascribed to a larger tip–
sample distance.
The calculations of Papanikolaou et al. show a
30% reduction in the LDOS 0.55 nm above thesurface when a surface Fe atom is replaced by a
Mn atom [20]. According to these calculations thispeak arises from the pure Fe surface: it was shownthat its relative weight (e.g. compared to the additional peak observed in the calculations for Cr
impurities) increases for larger tip–sample distances due to the larger area that is probed by thetip. This prediction seems to be in qualitative
agreement with our experimental results taking
into account the diﬀerent tip–sample distances for
dI=dV measurements on Fe(0 0 1) and on surface
substitutional Mn sites. However, the absolute tip–
sample distances in the measurements are notM.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

known which makes a quantitative comparison
impossible.
The dI=dV curve measured on the Mn atomalso shows a higher dI=dV signal below Vs ¼
0:2 V compared to the clean Fe(0 0 1) curve. Thisfeature might be related to the weak and broad
peak at 0.7 eV found in the calculated minority
band LDOS in the vacuum (at 0.55 nm above thesurface) [20]. However, the experimental feature is
too weak to base any strong claims on it. Observation of occupied sample LDOS peaks is not very
favorable due to the higher tunneling barrier experienced for tunneling from these states (e.g. thetunneling probability is reduced by a factor 0.2 forsample states 0.5 eV below the Fermi level at 1 nm
above the surface [26]). Nevertheless, this decay isalready included in the calculated vacuum LDOS.
Furthermore, unlike for the other impurities, the163

calculations for Mn predict a strong peak aroundþ1 eV in the majority band. This peak was,however, never observed in the experiments (seee.g. Fig. 5).
Results of STS measurements performed on theintermixed islands are shown in Fig. 6. dI=dV
curves were obtained on the surface area shown in
the atomically resolved STM image of Fig. 6a(same island as imaged in Fig. 3). Based upon the
dI=dV map at Vs ¼ þ0:99 V (Fig. 6b) three kindsof dI=dV spectra can be distinguished. Averages ofthese spectra (typically obtained from 100
curves) are shown in Fig. 6c. Curve 1 is measured
on the clean Fe(0 0 1) terrace and shows theFe(0 0 1) surface state as a peak at þ0.20 eV.
Comparing the topographic image (Fig. 6a), thedI=dV map (Fig. 6b) and the dI=dV curves (Fig.
6c) shows that curve 2 (with low dI=dV at þ1 V) is

Fig. 6. STS measurement on Fe(0 0 1) covered with 0.37 ML Mn deposited at 150 °C. (a) Topographic image (I ¼ 0:5 nA, Vs ¼ 5 mV,
4:8  4:4 nm2 ). At this area dI=dV is measured at setpoints I ¼ 0:5 nA, Vs ¼ 0:54 V. (b) dI=dV map at Vs ¼ þ0:99 V. (c) Three
diﬀerent types of dI=dV curves (curves are averaged) are obtained: curve 1 is measured on the Fe(0 0 1) terrace, curve 2 on the local
cð2  2Þ MnFe surface alloy and curve 3 on the pð1  1Þ Mn areas on the islands. (d) dI=dV curves of (c) normalized by I=V .

164

M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

correlated with the protrusions, and curve 3 (withhigh dI=dV at þ1 V) with the topographic dark
pð1  1Þ areas. Curve 2 shows a peak at þ0.35 eV,
while no peaks are visible in curve 3. Nevertheless,
it is clear that at positive bias voltages there is anadditional contribution to dI=dV curve 3 on top of
the exponential background (see also Fig. 8). As
discussed previously, the protrusions on the islands were identiﬁed as Fe atoms which have beenreleased from the Fe(0 0 1) terrace after place exchange events and intermixed with Mn islands.
Curve 2 is therefore representative of a cð2  2Þlike clustering of Mn and Fe atoms and curve 3 fora pure Mn adlayer on Fe(0 0 1).
The absence of a peak on the pure Mn part ofthe island in the dI=dV measurement of Fig. 6cmight seem conﬂicting with the observation of apeak at þ0.45 V in the measurement of Fig. 1b. A
diﬀerent electronic structure cannot be excluded
due to the possibility of the Mn atoms being differently coordinated, i.e. four (zero) Fe (Mn)nearest neighbors for growth at 50 °C and four orless (zero or more) Fe (Mn) nearest neighbors for
growth at 150 °C. Growth at diﬀerent temperatures may also stabilize diﬀerent magnetic unit
cells with inequivalent electronic structures [15].
Nevertheless, the eﬀect can be perfectly explainedby an intrinsic artifact of STS measurements: thepresence of a voltage dependent background.
In the dI=dV curve measured on the clean
Fe(0 0 1) surface shown in Fig. 1b almost no exponential background is visible on the high-energy
side of the peak and consequently, the peak exhibits a Gaussian-like shape. In Fig. 6c, however,the clean Fe(0 0 1) dI=dV curve has a much larger
exponential background on the high-energy side of
the peak. The background is clearly related withthe tip conﬁguration: voltage pulses (up to 10 V)
almost always lead to a change of this background. In Fig. 7 two extreme cases are shown: theleft dI=dV curve shows a Gaussian-like shape witha weak exponential background, while for theright curve this background is high on the positiveenergy side. To categorize the diﬀerent peak
shapes observed with diﬀerent tip conditions, aparameter which characterizes the amount ofbackground is deﬁned. A possible parameter,
which has the advantage that it is easily obtained

Fig. 7. Peak energies observed in STS as a function of background. The background parameter a=b is deﬁned as the ratio
between the local minimum dI=dV on the low-energy side of theFe(0 0 1) surface state peak (a) and the local minimum dI=dV
on the high-energy side (b). Therefore, a low (high) backgroundon the high-energy side corresponds to high (low) a=b. The
dI=dV peak energies obtained on the clean Fe(0 0 1) terrace,
MnFe cð2  2Þ alloyed areas, and Mn pð1  1Þ areas are shownas black squares, circles and diamonds, respectively. Grey
symbols are used for the corresponding peak energies in
ðdI=dV Þ=ðI=V Þ. The data points are ﬁtted to linear functions toguide the eye.

from experimental data, is the ratio a=b between
the (local) minimum in dI=dV on the low-energy
side of the clean Fe(0 0 1) peak a and the (local)
minimum in dI=dV on the high-energy side b. aand b are indicated in Fig. 7. With this deﬁnition,
Gaussian peak shapes correspond to a=b  1,
while dI=dV curves which exhibit a high exponential background on the high-energy side correspond to a=b < 1. In Fig. 7 the dI=dV peakenergies obtained on the clean Fe(0 0 1) surface areplotted as a function of this parameter (blacksquares). The apparent peak energy varies between
þ0.13 and þ0.20 V with a clear trend that a low(high) peak energy is found if the high energy
background is absent (present).
Normalizing dI=dV by I=V does not help toomuch (grey squares in Fig. 7): now the peak energy
shifts between þ0.09 and þ0.15 V with the same

M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

trend. The normalized curves obtained on the Mn
islands grown at 150 °C are shown in Fig. 6d. Itshows that peaks at high energies which areheavily distorted by the background in dI=dV are
more pronounced in ðdI=dV Þ=ðI=V Þ although foreven higher energies the peak is disturbed by the
background in ðdI=dV Þ=ðI=V Þ as well (curve 3 inFig. 6d). Therefore, a tip conﬁguration which
suppresses this background is most desirable toresolve the high-energy peaks in dI=dV (andðdI=dV Þ=ðI=V Þ). This is the case for the measurement shown in Fig. 1b.
The peak energies observed on the MnFe alloyed island areas (circles) and the pure pð1  1ÞMn island areas (diamonds) are also shown in Fig.
7 as a function of the background parameter a=b.
Peak energies in dI=dV and ðdI=dV Þ=ðI=V Þ areindicated by black and grey symbols, respectively.
As a function of background parameter a=b,
dI=dV peak energies between þ0.20 and þ0.43 V
can be found for the alloyed areas and between
þ0.45 and þ0.50 V for the pure Mn islands (although these peaks are only observed for quite
high a=b for which the background is suppressedwhich explains why there are only two data points
for the dI=dV peaks measured on pure Mn islands
in Fig. 7). The peak energies in ðdI=dV Þ=ðI=V Þ
vary between þ0.17 and þ0.25 V for the cð2  2Þ
MnFe areas, and between þ0.29 and þ0.42 V for
the pure Mn islands. Obviously, low (high) peakenergies are found for low (high) background, andpeak energies in ðdI=dV Þ=ðI=V Þ are always lower
than those in dI=dV . The ﬁtted lines in Fig. 7 areintended as a guide to the eye.
Ukraintsev has shown that neither dI=dV nor
ðdI=dV Þ=ðI=V Þ reﬂect the real sample LDOS even
within a simpliﬁed one-dimensional WKB approach [39]. Both quantities represent the LDOS
convoluted with a tunneling probability function
which also contains the tip density of states at the
Fermi level for positive bias voltages. Ukraintsev
proposed to normalize dI=dV with its ﬁtted tunneling probability function: this deconvolution
technique is considered to lead to the best recoveryof the sample LDOS, at least within the onedimensional WKB approximation of tunneling.
However, the deconvolution technique proposedby Ukraintsev is not very practical since it requires

165

dI=dV data over a large voltage range to allow foran accurate ﬁtting of the background. Such a large
voltage range is usually not available. Stroscio
et al. deduced the surface state energy of clean
Fe(0 0 1), i.e. þ0.17 V, by a ﬁt of a Gaussianfunction plus quadratic background to the dI=dV
data [26] which can be considered as a simple approximation to the Ukraintsev approach. We
veriﬁed that normalizing the dI=dV curve measured on the clean Fe(0 0 1) shown in Fig. 6 (andmany other curves) by a quadratic function shifts
the Fe(0 0 1) surface state from þ0.20 to þ0.17 V
as well. This gives a strong motivation to use thismethod for LDOS recovery also for the dI=dV
curves measured on the pure Mn pð1  1Þ island
areas and the MnFe cð2  2Þ alloyed areas.
Fig. 8 shows the results of this deconvolution.
In Fig. 8a the dI=dV curve measured on the pure
Mn island (see Fig. 1b) and two quadratic functions ﬁtted to the high energy background of thiscurve are shown. Due to the small voltage range
used (up to þ1.5 V) the ﬁts are not very accurate
and the two ﬁts shown in Fig. 8a reﬂect two extremes. However, the two normalized curves arealmost equal and both show a peak at þ0.34 V.
This gives an indication that the ﬁts to the tunneling probability do not need to be perfect to givea reasonable estimate of the surface state energy.
Also shown in Fig. 8a is a dI=dV curve measured
on pure Mn which shows a strong background inwhich no peak can be resolved (see Fig. 6c).
Normalizing this curve by its quadratic ﬁt (which
can only be a very rough estimate of the realtunneling probability function), leads to a curveshowing a peak at þ0.37 V.
In Fig. 8b two dI=dV curves measured on the
MnFe cð2  2Þ alloyed areas are shown (one curveis the curve of Fig. 6c). Normalizing these curves
by quadratic functions ﬁtted to their high energy
backgrounds recovers peaks at þ0.26 V for bothcases. Therefore, although dI=dV curves need tobe measured over large voltage ranges to allow fora more accurate ﬁtting of the background, ourresults imply that even imperfect ﬁts to quadratic
functions lead to a relatively accurate recovery of
the real surface state energies. In the present study,
we estimate the surface state energies of thepð1  1Þ Mn islands and the cð2  2Þ MnFe alloy

166

M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

Fig. 8. dI=dV curves (black drawn) measured on pure Mn pð1  1Þ island areas (a) and MnFe cð2  2Þ alloyed areas (b). In both (a)and (b) two dI=dV curves with diﬀerent backgrounds due to diﬀerent tip conﬁgurations (see text) are shown. The background on thehigh-energy side of these curves have been ﬁtted by quadratic functions T (A, B, C, D and E; black dotted curves). dI=dV curves with
diﬀerent backgrounds show the same peak energy (within 0.05 V) in ðdI=dV Þ=T (grey drawn curves). In (a) the dI=dV curve showing apeak was normalized using two separate backgrounds (A and B). The corresponding dI=dV curves are almost identical showing thatthe ﬁt of T to the background in dI=dV does not have to be extremely accurate.

to be þ0:35  0:05 V and þ0:25  0:05 V, respectively.
The diﬀerent backgrounds appearing underdiﬀerent tip conﬁgurations may be a result of theapex having either metallic (e.g. W, Fe or Mnatom) or non-metallic (e.g. O) character. Measurements of apparent barrier heights showed thatwith tips leading to small backgrounds high apparent barrier heights are found (5 eV), whilewith tips showing an extreme high background low
apparent barrier heights are found (3 eV). Bothkind of tips image monoatomic steps with thesame sharpness and tip changes mostly lead to tip
displacements of less than 0.1 nm. Both observations imply that the tip changes only involve an
exchange of the apex atom. Electronegative adsorbates like oxygen are known to increase work
functions on tungsten surfaces (see e.g. [40,41])
which implies that the absence of background in
dI=dV might be related to slightly contaminatedtip conﬁgurations. Furthermore, atomically sharp
tungsten tips are expected to have smaller work
functions than close-packed surfaces [42] whichmight explain why apparent barrier heights as low
as 3 eV are measured. Moreover, it is not unlikely that the lower conductivity of non-metallic
tips leads to smaller tip–sample distances at agiven setpoint: the tunneling probability function
within a one-dimensional WKB approximation(see e.g. [39]) shows that at smaller tip–sample
distances the voltage dependence of the tunnelingprobability function is reduced. The inﬂuence of
diﬀerent tip conﬁgurations on the appearanceof the voltage dependent background in dI=dV
curves can now be qualitatively understood.

4. Conclusion
STM measurements showed that Mn atoms
become incorporated in the Fe(0 0 1) surface forsubmonolayer growth at temperatures above 100°C. From images showing chemical contrast thefraction of Mn atoms that have undergone placeexchanges with Fe substrate atoms was deduced asa function of growth temperature. A local cð2  2Þ
structure was observed on the islands which wasinterpreted as an ordered MnFe alloy. For thegrowth temperatures studied (up to 200 °C) the
Mn atoms were found to stay in the surface layer.
STS measurements did not reveal a double peakstructure in the dI=dV spectra on the single Mn
impurities in contrast to Cr/Fe(0 0 1). The inﬂuence of a tip dependent background on the apparent surface state energies in dI=dV andðdI=dV Þ=ðI=V Þ was discussed and carefully analyzed. By normalizing the dI=dV curves by ﬁtted
quadratic backgrounds the surface state energies

M.M.J. Bischoﬀ et al. / Surface Science 501 (2002) 155–167

were recovered. Pure pð1  1Þ Mn islands were
found to show a surface state around þ0.35 V,
while the cð2  2Þ MnFe alloyed areas on the islands showed a surface state around þ0.25 V.

Acknowledgements
This work was supported by the Stichting voor
Fundamenteel Onderzoek der Materie (FOM),
which is ﬁnancially supported by the Nederlandse
Organisatie voor Wetenschappelijk Onderzoek(NWO), and the European Growth project
MAGNETUDE. T.Y. acknowledges Gakushuin
University (Tokyo, Japan) for a travel grant. We
gratefully thank B. Heinrich (Simon Fraser University, Burnaby, Canada) for supplying the Fe
whiskers.
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