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Applied Surface Science 253 (2006) 2244–2250

Surface segregation in HAYNES 230 alloy

D. Pop *, K. Wolski

Ecole Nationale Superieure des Mines de Saint-Etienne, Centre SMS, 158, cours Fauriel, 42023 Saint-Etienne Cedex 02, France

Received 6 February 2006; received in revised form 12 April 2006; accepted 12 April 2006

Available online 23 May 2006

Abstract

The surface segregation in the Ni-based alloy HAYNES 230 was studied by Auger Electron Spectroscopy and X-ray Photoelectron

Spectroscopy between 400 and 1100 8C. The qualitative variations of the surface contents of S, P, W, Mo, N, Si, and Mn were determined

as a function of annealing temperature and time. It was found that at 925 8C the maximum coverage of sulphur at the alloy surface is in the range

0.06–0.15 monolayers. Chromium evaporation from the HAYNES 230 surface under UHV conditions is clearly evidenced for annealing at

1100 8C.

# 2006 Elsevier B.V. All rights reserved.

PACS: 68.35Dv; 81.05Bx; 82.80Pv

Keywords: HTR; High temperature reactor; Ni-based alloys; Inconel 617; HAYNES 230; AES; XPS; High-temperature corrosion; Surface segregation; Site

competition

1. Introduction

Nowadays it is well known that the physical processes at

surfaces have technological importance for multiple applica-

tions. In the case of materials used at high temperatures the

segregation processes become significant and consequently the

material surface can be enriched in certain elements or

impurities. This change in surface composition can negatively

affect a series of mechanical and chemical properties, in

particular the adhesion of oxide scales.

However, the comprehensive studies of surface segregation in

superalloys are relatively scarce. HAYNES 230 is a Ni-based

alloy with potential for use in the fabrication of heat exchangers

in the next generation of nuclear reactors—Very High

Temperature Reactors (VHTR), that are envisaged to work at

temperatures above 900 8C. Such a material must be resistant to

the corrosion processes that occur in the impure helium

environment of the nuclear reactor and which influence

especially its mechanical properties. The main process that

has to be taken into account in such an environment is the

formation of the protective oxide scale in order to control either

carburization (which would result in low temperature brittleness)

* Corresponding author. Tel.: +33 4 77 42 02 15; fax: +33 4 77 42 01 57.

E-mail address: diana.pop2@gmail.com (D. Pop).

0169-4332/$ – see front matter # 2006 Elsevier B.V. All rights reserved.

doi:10.1016/j.apsusc.2006.04.026

or decarburization (which strongly affects creep properties). The

adhesion of the oxide film is controlled by its chemical

composition, the morphology of the oxide/metal interface, and

the type of segregation species [1].

The aim of this work was to determine the type of segregating

atoms in the absence of any oxidation-driven phenomena. Under

these conditions, obtained in ultra-high vacuum (UHV)

environment, the only driving force results from the free

enthalpy of segregation. The same driving force will act under

VHTR oxidizing environment, presumably leading to similar

segregation phenomena at the metal/oxide interface.

Knowledge of the segregation characteristics of HAYNES

230 becomes thus important for assessing the potential of this

material for nuclear applications. We have used Auger Electron

Spectroscopy (AES) and X-ray Photoelectron Spectroscopy

(XPS) under UHV of about 5 � 10�10 mbar to investigate the

changes in the surface composition of this alloy at temperatures

between 400 and 1100 8C. These studies in ultra-high vacuum

represent a preliminary step for the understanding of the

processes that can lead to material degradation in the aggressive

environment of the nuclear reactor.

2. Experimental

The nominal composition of the HAYNES 230 alloy studied

in this work is given in Table 1 and compared to the

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–2250 2245

Table 1

Composition of the studied HAYNES 230 alloy compared to the composition of Inconel 617 alloy

C Ni Fe Cr Mn Ti Cu Si Al Mo Co W S

Haynes 230 (wt%) 0.10 Bulk 1.41 21.76 0.50 0.10 0.01 0.39 0.30 1.30 0.18 14.21 0.001

IN617 (wt%) 0.06 54.19 1.13 22.3 0.06 0.35 0.1 0.11 1.26 8.7 11.7 / 0.001

Fig. 1. Qualitative variation of the sulphur surface content as a function of

temperature after 30, 60, and 105 min annealing.

composition of Inconel 617 alloy, used in the past in a series of

High Temperature Reactors (HTR) operating in the 800 8Ctemperature range.

The chemical analysis of minor elements performed in our

laboratory by combustion and infrared absorption (system

LECA CS-444LS) confirmed the given percentages for C and S.

Analysis based on the sample fusion under flowing helium gas

stream and measurement by a thermal conductivity cell (system

LECA TC-436) indicated also the existence of 0.06% N in the

alloy.

The samples consisted of discs with a diameter of 10 mm

and a thickness of approximately 0.3 mm. The front faces of

these discs were mechanically polished down to 0.25 mm.

The experiments were performed in an ultrahigh vacuum

chamber (base pressure �5 � 10�10 mbar) equipped with a

cylindrical electron energy analyzer (Cameca, MAC III), an

X-ray source (Cameca CX 700), an electron gun (VG

LEG200), and an ion sputter gun (Cameca, model ACI). A

resistive heating system placed at the rear side of the sample

allowed to obtain temperatures up to 1100 8C. The

temperature measurements were realized with a thermo-

couple spot-welded on the front face of the sample. The

Auger spectra were obtained using a 5 keV electron beam

with a 10 mm spot diameter. These spectra were recorded in

the direct acquisition mode with a resolution of 1 eV and a

step of 0.25 eV. Complementary XPS analysis were per-

formed using the Mg Ka radiation (1253.6 eV) with a beam

diameter of ca. 1 cm, the data being measured in similar

conditions to the Auger ones. Auger electron spectroscopy

was the main technique used for the quantitative and

qualitative analysis due to the relatively small spot of the

incident electron beam, which allowed the study of the

regions homogeneous in terms of temperature.

Two samples were analyzed. First of them was studied

subsequent to annealing at 400, 500, 600, 700, 780, 875, 905,

925, 960, 990, 1035, and 1100 8C. The second sample was

annealed for more than 30 min only at 910 and 925 8C. Prior to

analysis at a certain temperature, the sample surface was

cleaned at room temperature by sputtering with Ar+ ions of 3–

4 keV energy and reference spectra were recorded. Afterwards

the samples were annealed at the desired temperature for a

given time, cooled down to approximately 100 8C, and the data

were collected. For a certain temperature, several cycles of

annealing and spectra recording were performed, which

allowed to study the surface segregation as a function of

annealing time. The time for reaching the annealing

temperature was maintained in the order of several minutes,

in order to limit the effects on surface stoichiometry. For the

same reason, after annealing the heating source was switched

off immediately.

3. Results and discussions

To investigate the surface segregation of the different alloy

elements and impurities, the intensities of the corresponding

Auger lines were analyzed. In order to emphasize the changes

in the surface concentration with increasing annealing

temperature and time, and at the same time to compensate

for possible differences in the intensity of the primary electron

beam (which would lead to a change in the intensity of all

Auger lines in a spectrum), the signals of interest were

calculated relative to the intensity of the Ni 844 eVAuger peak.

Nickel is considered to not segregate or evaporate when the

sample is thermally treated. However, it should be mentioned

that these ratios are influenced by the decrease of the Ni signal

at the surface, which is likely to occur when impurities or one of

the alloy components segregate significantly.

The modifications in surface composition found after an

annealing cycle include also the changes that occur during the

time needed to reach the desired temperature and the changes

that occur during sample cooling. Provided that these time

intervals are relatively short, the associated modifications in the

surface stoichiometry (which cannot be quantified in the

present experiments) are of a much smaller importance than

those occurring during annealing.

3.1. Sulphur segregation

Segregation of sulphur can result in the spallation of the

surface region of Ni-based alloys, even if the bulk concentration

of this element is only several ppm [2]. For the materials used in

the particular environment of a nuclear reactor this can lead to

an accelerated alteration of mechanical properties due to the

enhancement of carburization effects [1].

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–22502246

In HAYNES 230 sulphur segregation is observed already for

annealing the alloy at 230 8C, which was the sample

temperature when the chamber was heated for obtaining the

ultrahigh vacuum. Fig. 1 shows the variation of the sulphur

content at the surface as a function of annealing temperature,

after three annealing times. The dependence was obtained by

calculating the ratio between the intensity of the S 146 eV

Auger line and the intensity of the Ni 844 eV signal. The data

indicate that the sulphur amount at the alloy surface generally

raises with temperature, suffering a dramatic increase for

annealing above 700 8C.

3.2. Maximum thickness of the segregated sulphur layer at

925 8C

Fig. 2 presents the variation in the amount of sulphur at the

alloy surface as a function of annealing time at 925 8C. The

error bars for the ratios between S and Ni Auger signals at

925 8C are maximum 10%. One can note that up to 1000 min

the sulphur content varies strongly with time, while for longer

annealing this variation is very slow. For 2700 min a saturation

level is reached.

To estimate the amount of sulphur at the HAYNES 230

surface that corresponds to this saturation regime (parameter

that will be further denoted as ‘e’), a bicomponent standard

sample—FeS was also analyzed and the quantification was

made using the method presented in [3,4] and references

therein. Hence, the intensities of S and Fe Auger signals in the

spectra of HAYNES 230 alloy (after 2735 min annealing) and

in the spectra of the standard material FeS were measured.

Considering that subsequent to annealing at 925 8C, a

uniform layer of sulphur is covering the alloy surface, and that

by Auger Electron Spectroscopy we probe both this sulphur

layer and a part of the HAYNES 230 alloy underneath the

segregant layer, the intensities of the Auger signals for S and Fe

can be written as

IS ¼ KaSRalloyS N0

Sl0Sð1� k0

SÞ cos u

IFe ¼ KaFeRalloyFe Nalloy

Fe lalloyFe kFe cos u

: (1)

Fig. 2. Variation of the sulphur surface content with annealing time at 925 8C.

In the above formulae the subscripts are relative to the elements

and the superscripts relative to the matrix (‘0’ for pure material

and ‘alloy’ for HAYNES 230). The parameter K depends on the

instrumental factors and analysis conditions, while a is the

product of ionization cross-section, probability of Auger elec-

tron emission, and transmission efficiency of the analyzer. R

represents a backscattering term and was calculated using the

relation

R ¼ 1þ ð4:35� 3:93Z0:1Þ ��

EP

EL

��0:25

þ 4:85Z0:1 � 5:45;

(2)

where Z is the average atomic number of the matrix (33.6 for

the alloy), EP the primary electron energy (5 keV), and EL is the

binding energy of the core level involved in the Auger transition

(162.5 eV for S and 707 eV for Fe). The values obtained in this

way were RalloyS ¼ 1:92 and Ralloy

Fe ¼ 1:68. The parameter N

designates the number of atoms per unit volume. Using the

densities of S and HAYNES 230 alloy we have determined that

N0S ¼ 3:68� 1028 atoms=m3 and Nalloy

Fe ¼ 1� 1027 atoms=m3.

The parameter l represents the attenuation length, which can be

calculated according to the relation

l ¼ 538E�2C þ 0:41

ffiffiffiffiffiffiffiffiffiffiffiffiffiffid � EC

p: (3)

In this formula EC is the Auger electron kinetic energy (146 eV

for S and 700 eV for Fe) and d is the monolayer thickness of the

matrix. Considering that d for S is the cubic root of the atomic

volume (0.295 nm) and approximating that for the alloy d is an

average of the monolayer thickness of the constituents

(accounting for the percentage contribution of each element

in the alloy – 0.228 nm) the obtained values are l0S ¼

2:72 monolayers and lalloyFe ¼ 5:18 monolayers. The para-

meters denoted k0S and kFe are the attenuation factors for S

and Fe. They can be calculated according to the formula

k ¼ expð�e=ðl cos uÞÞ, where l is the attenuation length, u

the collection angle with respect to the surface normal (408 in

the case of the present system), and e is the thickness of the

segregated layer. Thus

k0S ¼ exp

�� e

2:08

�and kFe ¼ exp

�� e

4:51

�:

Using these values, the ratio between the intensities of S and Fe

Auger signals becomes:

IS

IFe

¼ aS

aFe

� 21:96� 1� expð�e=2:08Þexpð�e=4:51Þ : (4)

The value IS/IFe, determined using the intensities of the S and Fe

Auger lines in the spectra corresponding to saturation condi-

tions at 925 8C, is 227 � 35. The unknown aS/aFe ratio includes

the transmission efficiencies of the analyzer and therefore was

determined from reference measurements with the same ana-

lyzer. In order to obtain this value the Auger spectra of the

standard material FeS were measured.

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–2250 2247

The Auger intensities of Fe and S in the spectra of

homogeneous FeS can be expressed as:

I�S ¼ K 0aSR�SN�Sl�S cos u0

I�Fe ¼ K 0aFeR�FeN�Fel�Fe cos u0

where the factor K0 is identical for both elements because the

data are obtained during the same experiment. The values for

R�S and R�Fe, determined using formula (2), are 1.71 and 1.53,

respectively, whereas the ratio N�S=N�Fe for FeS is taken to be 1.

Considering that the monolayer thickness of FeS is equal to the

cubic root of the ratio between the unit cell volume and the

number of atoms per cell (4 atoms/cell), we obtained that l�Sand l�Fe are 2.48 and 5.38 monolayers, respectively. With all

these parameters and the values determined for I�S and I�Fe from

the Auger spectra, we calculate that ðaS=aFeÞ ¼ ðI�S=I�FeÞ�ððR�FeN�Fel

�FeÞ=ðR�SN�Sl�SÞÞ ¼ ð116� 35Þ � 1:93. The larger

relative uncertainty in the I�S=I�Fe value is caused by the fact

that the surface of this sample was not perfectly flat, so that data

measured on different spots had a larger dispersion. Using this

value in Eq. (4), it is obtained that at 925 8C the maximum

thickness e of the sulphur layer is in the range 0.06–0.15

monolayers.

Sulphur segregation has been extensively studied in the past.

In the particular case of a commercial purity nickel (99.3%) the

equilibrium surface concentration of sulphur, between 700 and

800 8C, was found to be 44% [5]. The fact that the value

determined in the present measurements on HAYNES 230 is

considerably smaller than the one reported in Ref. [5] can be

due to the following factors: (i) at 925 8C, P and Si still

segregate and thus occupy some of the available surface sites,

(ii) in the frame of McLean’s theory, it is possible that at 925 8Ca desegregation process is already taking place, and (iii) the

occurrence of a certain sulphur evaporation.

3.3. P segregation

A second alloy impurity that shows strong segregation,

clearly observed for annealing temperatures higher or equal to

700 8C, is phosphorus. This element is reported to generate

embrittlement in Ni–Cr steels [6].

Fig. 3. Dependence of the phosphorous surface content with annealing time at

925 8C. The error bars represent standard deviation.

Fig. 3 shows the variation in the amount of phosphorus

present at the alloy surface as a function of annealing time at

925 8C. This dependence differs significantly from that

observed for sulphur. In the case of phosphorus segregation,

there is a rather fast accumulation of the phosphorous at the

sample surface, which is followed by a slow decay. The decay

in the surface content can be attributed to two reasons. First

would be that, with increasing the annealing time, there is more

and more sulphur that segregates and a competition between the

phosphorous and sulphur segregation occurs. A second

explanation could be related to the evaporation of phosphorous,

as claimed by other studies on steels [7]. However, given the

significant change in sulphur content at the alloy surface up to

1000 min annealing time (see Fig. 2), we interpret that the

major process responsible for the decrease in phosphorous

amount is the site competition between sulphur and phosphor-

ous.

The experiments show that for temperatures lower than

925 8C the initial accumulation of phosphorus at the alloy

surface is slower, so that the maximum in its surface amount

is reached after longer annealing times. For higher

temperatures the initial increase in the P content is even

faster than that illustrated in Fig. 3. The variation of the

phosphorus amount at the surface with annealing temperature

is shown in Fig. 4. After 30 min annealing at 960 8C the P

surface content is increased by a factor of 9 compared to that

for annealing at 700 8C. We interpret that the point for

annealing at 1035 8C indicates a smaller P content because at

this temperature the initial accumulation at the surface

occurred in less than 30 min and consequently after 30 min

the P content is already decreasing.

3.4. Segregation of carbide forming elements (Mo and W)

and segregation of N

Carburization leads to the low-temperature embrittlement of

the Ni-based alloys. In the VHTR environment tungsten and

molybdenum are strong carbide formers [8].

Fig. 4. Variation of the phosphorous surface content as a function of tempera-

ture after 30 min annealing. Error bars represent the magnitude of the standard

deviation of the data.

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–22502248

Fig. 6. (a) Nitrogen signal in the Auger spectra of HAYNES 230 alloy after

225 min annealing at 700 8C. (b) Nitrogen signal in the X-ray photoelectron

spectra of HAYNES 230 alloy after 225 min annealing at 700 8C.

Fig. 5. Variation of the tungsten surface amount with time for annealing at 600,

700, and 780 8C.

To study the segregation behavior of tungsten, the intensity

of its Auger signal centered at 164 eV was analyzed. W starts to

segregate around 600 8C. Fig. 5 shows the variation in the

intensity of the above mentioned WAuger line relative to the Ni

844 eV Auger signal for annealing times up to 225 min at 600,

700, and 780 8C. For all these temperatures the amount of W at

the alloy surface increases with time in the given range. At

600 8C, for example, the ratio between W and Ni Auger signals

changes from�0.6 at t = 0 min to�1.6 at t = 225 min. The data

also indicate that the W segregation rate increases with

temperature, reaches a maximum for 700 8C and then

decreases. For annealing temperatures equal to or above

875 8C no more segregation of W was observed.

After the annealing at 700 8C the W found on the sample

surface could not be removed by sputtering down to the initial

level. Therefore, the value for t = 0 on the 780 8C curve is

higher than the corresponding values for 600 or 700 8C and all

the points from the 780 8C curve could be slightly over-

estimated. We interpret this as a consequence of the selective

sputtering. Due to its larger atomic weight, W is harder to

sputter than the other alloy elements or impurities.

To determine the variation of the molybdenum surface

content with annealing temperature and time, the intensity of

the Mo Auger line centered at 184 eV was analyzed. Mo

segregation was first observed for annealing at 400 8C (after

30 min annealing, signals ratio Mo/Ni is approximately 0.5).

The studies indicated that for a temperature between 400 and

780 8C and annealing time up to 225 min the Mo content at the

surface increases continuously with time. The segregation rate

raises with temperature, reaches a maximum for 700 8C (after

30 min annealing, signals ratio Mo/Ni is approximately 1.2),

and then decreases (after 30 min annealing at 780 8C signals

ratio Mo/Ni � 0.5), so that for temperatures larger than 875 8Cno molybdenum is detected in the spectra.

Nitrogen signal is clearly observed in both Auger and X-ray

photoelectron spectra of HAYNES 230 for temperatures

between 400 and 780 8C (see Fig. 6). Its segregation behavior

is similar to that observed for W and Mo; for a given annealing

temperature the nitrogen content at the surface increases

continuously up to 225 min annealing time, the maximum in its

segregation rate being reached for annealing at 700 8C. After

30 min annealing at 500, 600, and 700 8C the ratios between the

Auger signals of N and of Ni are �0.8, �0.9, and �1,

respectively. For larger temperatures the segregation rate

decays, so that for 875 8C and above no more N is detected in

the spectra.

The reduction in the segregation rates of W, Mo, and N for

temperatures above 700 8C is most likely caused by the fact that

sulphur strongly segregates above this temperature and the site

competition between S and these elements occurs. The general

behavior of the N segregation could as well be influenced by an

effect of cosegregation with Mo and/or W, resulting in the

formation of molybdenum nitride and tungsten nitrate,

respectively. Such processes were previously observed in

Fe–Mo [9] and Fe–W [10] alloys.

3.5. Segregation of oxide forming elements (Si, Mn, Ti, Al)

In the nuclear reactor environment, the formation of an

adherent oxide layer at the material surface can protect against

the detrimental processes like carburization or decarburization

[1]. It was found that manganese, together with chromium, very

often forms a stable oxide of the spinel type MnCr2O4 [8]. The

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–2250 2249

Fig. 7. Dependence of the silicon surface content on annealing temperature. Fig. 8. Relative intensity variation of the chromium Auger lines centered at 486

and 525 eV with annealing time at 780 8C.

Fig. 9. Relative intensity variation of the chromium Auger lines centered at 486

and 525 eV with annealing time at 1100 8C. Error bars represent the standard

deviation.

nature of the compounds formed with titanium is controversial,

possibly TiCr2O4. Silicon and aluminum form pure oxide

phases: SiO2 and Al2O3, respectively [8].

Silicon segregation was observed for annealing between

500 and 1035 8C. At a given temperature, its dependence

on the annealing time proved to be similar to that of

phosphorous, consisting of a fast accumulation of this

element at the surface, followed by a slow decay. The

variation of the silicon surface content as a function of

annealing temperature is shown in Fig. 7. The data were

obtained by analyzing the intensity of the 89 eVAuger line of

Si. They indicate a continuous increase of the surface content

with temperature, the segregation rate having a significant

growth above 900 8C.

Manganese signal was clearly detected in the Auger spectra

at 925 8C only for the second HAYNES 230 sample, that

suffered annealing at 910 8C summing about 10 h, but no other

significant previous heating. In the case of the first sample,

which suffered longer annealing processes at lower tempera-

tures, the intensity of the manganese Auger peak is much

smaller. Therefore, it is concluded that manganese suffers a

segregation processes followed by a fast evaporation. This

interpretation is supported by the fact that the vapor pressure for

manganese is very high in comparison to that of the other

components of the alloy (at 1200 K, 5.31 � 10�6 atm, in

contrast to that of 2.3 � 10�10 atm for Cr, or 1.1 � 10�11 atm

for Ni) [11]. Due to the small intensity of its Auger peak in most

of the spectra, it is difficult to estimate the temperature range

where the manganese segregation occurs as well as its time and

temperature dependence.

No aluminum segregation was detected for HAYNES 230 in

the present experiments. However, in the Auger spectra

measured up to 1000 eV, as in the present experiments, a

small Al signal is difficult to distinguish due to the fact that it

lays in the vicinity of the much stronger Ni and Cr signals.

Aluminum was not detected in the X-ray photoelectron spectra

as well, but also in this case small quantities of this element are

hardly perceivable due to the low cross-section. Therefore, even

if a slight segregation of Al occurred, this could not be

experimentally proved.

Concerning the detection of titanium, the intrinsic sensitivity

of Auger/XPS techniques is much better due to the Ti favorable

position in the spectra and to the higher cross-section. However,

the Ti signal being absent in both Auger and XPS spectra, we

concluded that no Ti segregation takes place.

3.6. Cr segregation and depletion

Chromium is the major alloying element that, as a function

of the environment of the nuclear reactor, can either form a

protective Cr2O3 oxide layer at the surface, or produce carbides

[1].

To study the chromium behavior in HAYNES 230, the

intensities of the most important Cr Auger lines (centered at

525 and 486 eV kinetic energy) were analyzed as a function of

annealing temperature and time. A chromium segregation was

observed for annealing the alloy at 700 and 780 8C. Fig. 8

illustrates this process at 780 8C, indicating the variation of the

ratio between the Cr and Ni Auger signals as a function of

annealing time.

Nickel-based alloys and stainless steels that are used at

temperatures above 900 8C present the risk of a chromium

D. Pop, K. Wolski / Applied Surface Science 253 (2006) 2244–22502250

Fig. 10. Chromium depletion profile of HAYNES 230 after repetitive anneal-

ing. The data were obtained by EDX measurements on a polished transverse

section.

depletion [12] at their surfaces. The evaporation of chromium

from stainless steels during vacuum annealing was considered

to be a process similar to that of chromium loss due to selective

oxidation [12]. For HAYNES 230, the variations in the

intensities of chromium Auger lines present a monotonous

decrease with annealing time at 1100 8C, as shown in Fig. 9.

The data indicate that an evaporation of chromium occurs at

this temperature. Within the precision limits of the present

experiment, no systematic dependencies of the chromium

signal on annealing time were evidenced when the sample was

heated at 875, 905, and 925 8C. A more detailed determination

of the onset temperature for Cr evaporation will be performed in

a further study.

To evaluate the depth of the alloy subsurface region which is

depleted in chromium, Energy Dispersive X-ray analysis

(EDX) were performed on the sample that was repeatedly

annealed between 400 and 1100 8C. The sample was cut

transversally to the surface and its two halves were analyzed

using the electron beam of a Scanning Electron Microscope

operating at 20 kV. The spot size of the electron beam was ca.

1 mm. The depletion profiles for chromium were obtained by

performing an elemental analysis in depth of the sample, at

different distances compared to the surface that was previously

studied by AES and XPS. All the spots used for measurements

were situated on a same line, transverse to the surface. The

results are shown in Fig. 10. They indicate that the annealing

generated a depth of approximately 40 mm from the sample

surface that is depleted in chromium.

The EDX analysis point to a concentration of chromium at

the surface that is approximately five times smaller than the

bulk concentration. On the other side, the AES results from

Fig. 9 indicate only a reduction by a factor of 1.5–1.7 of the

chromium signal after 90 min annealing relative to the initial

intensity at time t = 0. This discrepancy can be attributed to two

main factors: first—a possible slow chromium evaporation at a

temperature inferior to 1100 8C, at a too small rate to determine

systematic changes the spectra, and second—the segregation of

Cr at the extreme surface, as clearly evidenced after annealing

at 700 and 780 8C.

4. Conclusions

Auger spectroscopy and XPS studies of the HAYNES 230

alloy were performed to determine the changes in composition

that occur at the material surface at different temperatures

under UHV conditions. The experimental observations allow to

conclude that:

� S

everal constituent elements and impurities of this alloy (S, P,

W, Mo, N, Si, Mn) segregate to the surface.

� T

he strong segregation of sulphur above 700 8C leads to a site

competition between this element and W, Mo, N, Si, and P,

respectively.

� A

t temperatures above 875 8C only S, P, and Si still segregate,

with the amount of P and Si increasing with temperature.

� I

n the particular case of sulphur, it was found that the

maximum coverage at 925 8C is in the range 0.06–0.15

monolayers, which might influence the spallation of the

protective oxide scale.

� A

nnealing of HAYNES 230 at 1100 8C under UHV

conditions results in chromium evaporation.

Acknowledgements

The authors thank F. Rouillard and C. Cabet from CEA-

Saclay, France, for providing the alloy material and for the

helpful scientific discussions. The contribution of P. Passet to

the experimental work is also highly acknowledged.

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