VSEBINA – CONTENTS
IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES
A failure criterion for single-crystal superalloys during thermocyclic loading
Merilo za prelom monokristala superzlitine pri termocikli~ni obremenitvi
L. Getsov, A. Semenov, A. Staroselsky . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3
Accelerated creep testing of new creep resisting weld metals
Preizkusi pospe{enega lezenja zvarov novega jekla, odpornega proti lezenju
S. T. Mandziej, A. Výrostková, M. [olar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13
Zveza med analiznimi rezultati – karbonatna bomba in termi~na analiza
Connection between analysis results – carbonate bomb and thermal analysis
@. Poga~nik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27
STROKOVNI ^LANKI – PROFESSIONAL ARTICLES
Developing and testing a new type-8K mould for tool-steel ingot casting
Razvoj in preizkus nove kokile vrste 8K za ulivanje ingotov iz orodnega jekla
M. Balcar, L. Sochor, R. @elezný, P. Fila, L. Martínek, L. Kraus, D. Ke{ner, J. Ba`an . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
An AES investigation of brushed AISI 304 stainless steel after corrosion testing
AES-preiskave krta~enega nerjavnega jekla AISI 304 po korozijskem preskusu
M. Torkar, D. Mandrino, M. Lamut . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39
Using a FIB to prepare Al(OH)3 samples for the TEM
Uporaba FIB za pripravo vzorcev Al(OH)3 za TEM
I. Nikolic, V. Radmilovic, T. Z. Sholklapper, D. Blecic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45
ISSN 1580-2949
UDK 669+666+678+53 MTAEC9, 42(1)1–47(2008)
MATER.
TEHNOL.
LETNIK
VOLUME 42
[TEV.
NO. 1
STR.
P. 1–47
LJUBLJANA
SLOVENIJA
JAN.-FEB.
2008

L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
A FAILURE CRITERION FOR SINGLE-CRYSTAL
SUPERALLOYS DURING THERMOCYCLIC LOADING
MERILO ZA PRELOM MONOKRISTALA SUPERZLITINE PRI
TERMOCIKLI^NI OBREMENITVI
Leonid Getsov1, Artem Semenov1, Alexander Staroselsky2
1St. Petersburg State Polytechnical University, Russia, St. Petersburg, Polytechnicheskaja 29
2Pratt and Whitney, MS 165-16, East Hartford, CT, 06108 USA
guetsov online.ru
Prejem rokopisa – received: 2007-09-12; sprejem za objavo – accepted for publication: 2007-12-18
The accumulation of deformation affects the lifetime of a monocristal submitted to low cyclic mechanical and thermal loading.
An analytical method was developed considering the damages due to cyclic and axial plastic and creep deformation. The
method was checked with tesst in vacuum on specimens of monocrystals with different space orientation, also specimens with
stress concentrators. The distribution of stressses was FEM modelled. The fracture depends on the space orientation of the
specimen and of the loading parameters. A new experimental-computational method using deformation-fracture criteria is
suggested for the evaluation of the crystal life time
Keywords: single crystals,cyclic and axial loading, thermal fatigue, failure criteria
Kopi~enje deformacije vpliva na trajnostno dobo kovine, ki prena{a malocikli~no mehansko in termi~no obremenitev. Razvita
je bila analiti~na metoda, ki se{teva po{kodbe zaradi cikli~nih in enoosnih plasti~nih deformacij in deformacije z lezenjem.
Metoda je preverjena s preizkusi v vakuumu na vzorcih monokristalov z razli~no prostorsko orientacijo. Preizkusi so bili
izvr{eni tudi z vzorci z obliko, ki je povzro~ala lokalno koncentracijo napetosti. Porazdelitev napetosti je bila modelirana po
metodi kon~nih elementov. Prelom je odvisen od prostorske orientacije in od parametrov obremenitve. Predlagana so nova
deformacijsko-prelomna merila za oceno trajnostne dobe monokristala na podlagi eksperimentalno-ra~unske analize.
Klju~ne besede: monokristali, cikli~na in aksialna obremenitev, termi~na utrujenost, pogoji za prelom
1 INTRODUCTION
The deformation criterion
D1( pli) + D2( ci) + D3(pli) + D4(ci) =1 (1)
for the fracture of metal in conditions of low-cycle and
thermal-cycle loading was proposed in 1. The quanti-
tative result of the lifetime test on the basis of this
criterion depends on the choice of the deformation
parameters.
The proposed deformation criterion is based on the
summation of the damage caused by cyclic plastic
deformations D
C n
1 i
pl k= ∑1
1
( )∆ε , cyclic creep defor-
mations D
C n
2 i
pl k= ∑1
2
( )∆ε , unilaterally accumulated
plastic deformations D
C n
3
i
pl
r
= ∑1
3
ε
ε
and unilaterally
accumulated creep deformations D
C n
4
i
pl
r
= ∑1
4
ε
ε
.
In the case of interference of the above-mentioned
types of damage the relevant corrections are made by
considering the limiting characteristics of the material,
C1, C2, r, and cr. This criterion was tested experi-
mentally under conditions of uniaxial and composite
tension for various samples, and the GTE details made of
isotropic heat-resistant steels and alloys that fracture are
caused by low-cycle or thermal fatigue (see Figure 1
and  2–4 and others).
Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12 3
UDK 539.42:669.018 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(1)3(2008)
Figure 1: The correlation between the experimental and calculated
values of the lifetime: I-D1+D2, II-D3+D4, III-D1+D2+D4, IV-D1+D3
+D4, V-D1+D2+D3 +D4
Slika 1: Korelacija med eksperimentalnimi in izra~unanimi vred-
nostmi za trajnostno dobo I-D1+D2, II-D3+D4, III-D1+D2+D4, IV-D1
+D3+D4, V-D1+D2+D3+D4
The quantitative result of the details’ lifetime test on
the basis of this criterion depends on the choice and
precision of its deformation parameters. As for the
isotropic materials, normally the role of such parameters
is played by the intensities of the total amplitudes of the
deformations
∆ ∆ ∆ ∆ ∆ ∆ ∆ε ε ε ε ε ε εi
pl
1
pl
2
pl
2
pl
3
pl
3
pl
1
pl= − + − + −2
9
2 2( ) ( ) ([ ])2 (2a)
[ ]∆ ∆ ∆ ∆ ∆ ∆ ∆ε ε ε ε ε ε εic 1c 2c 2c 3c 3c 1c= − + − + −29
2 2 2( ) ( ) ( ) (2b)
and the intensities of the accumulated deformations,
[ ]ε ε ε ε ε ε εipl 1pl 2pl 2pl 3pl 3pl 1pl= − + − + −29
2 2 2( ) ( ) ) (3a)
[ ]ε ε ε ε ε ε εic 1c 2c 2c 3c 3c 1c= − + − + −29
2 2 2( ) ( ) ( ) (3b)
In this research we studied the possibility of
generalizing the above-mentioned criterion for the case
of single-crystal alloys. This generalization was based on
the following relationships for fatigue-crack initiation
and growth in the references 5–12.
According to the experimental data, the increase of
temperature 5 and the reduction in the of frequency 6
change the mechanism of fracture in single crystals,
which means that the transition from crack growth along
crystallographic (octahedral) planes to growth in other
directions does not depend on the crystal orientation, but
on the conditions of loading (according to Mode I). The
non-crystallographic growth is observed predominantly
on the boundary of the ’ phases.
One of the possible explanations 20 for such a
transition is the initiation of additional damage owing to
the influence of the environment. Oxygen-related brittle
behavior in the vicinity of the crack tip takes place
according to the diffusion mechanism, and is sensitive to
the temperature, the time (frequency) and the concen-
tration of oxygen. A comparison of the results of the
experiments made in air 5 and in vacuum 7 has shown
that the transition temperature is considerably higher in
vacuum. When the frequency is decreased, the time of
the brittle behavior within the cycle increases, which
promotes the further penetration of oxygen, and this
initiates the transition from the crystallographic mode of
crack growth to the non-crystallographic mode I. Also,
the increase in the temperature accelerates the diffusion
process, promoting the transition from a crystallographic
to a non-crystallographic mechanism of crack growth.
As shown in 8, the threshold value for non-crystallo-
graphic fractures, ∆KthI, is lower than the value of
∆Kth(111) for cracks growing along a crystallographic
plane, and when ∆Keq becomes lower than ∆Kth(111) the
crack cannot continue growing according to the
crystallographic mechanism and goes over to growth
according to mode I. Thus, the experiments show that
the transition from the crystallographic to the non-cry-
stallographic stage is controlled by the temperature T,
the frequency f, and the total amplitude of SIF, ∆K.
For the set of above-mentioned parameters it is
possible to create maps of mechanisms for the growth of
fatigue cracks (analogous to the maps for static loading
in 9, showing the boundaries between the crystallo-
graphic and non-crystallographic areas of crack growth).
As an example of such a map, the reader is referred to
the map in 10 (see Figure 2), obtained on the basis of an
examination of thermally activated slip processes in the
vicinity of the crack tip.
The transition from the crystallographic mode of
growth at high temperatures is also promoted by the
decrease in the anisotropy of the elastic and strength
characteristics of crystals. So, at high temperatures
(980 °C) the growth rate of the fatigue cracks does not,
in practice, depend 11 on the orientation of single crystals
under the conditions of a uniaxial and composite
deformation mode. This probably explains the absence
of anisotropy in high-temperature multi-cycle fatigue.
It is important to note that the arborescent structure
also influences the change of the orientation of the
fatigue-crack growth and it could provoke the transition
from one crystallographic plane to another. The
equalization of the crack-growth rates for single crystals
with orientations  111 and  001 can be explained 11,12
by the fact that as ∆K grows, the fatigue-crack trajectory
diverts from the plane (111) perpendicular to the loading
axis and evolves into the plane (001), where
inter-dendrite areas are concentrated, i.e., to the same
plane in which the fatigue cracks grow in samples with
the orientation  001.
As a two-dimensional model, interpreting simply the
fatigue crack’s growth in the area surrounding the
cooling duct, the problem of crack growth in a plate with
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
4 Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12
Figure 2: Map of fatigue-crack growth mechanisms for monocrystals
<001>/<010> PWA 1484 10
Slika 2: Shema mehanizma rasti utrujenostne razpoke za monokristale
<001>/<010> PWA 1484 10
a circular hole can be considered. The plate is in
non-uniform temperature conditions with the minimum
value of the temperature on the edge of the hole: the
greater is the distance from the hole, the higher are the
temperature and the SIF. The crack-growth process on
the plane of parameters corresponding to the map of
mechanisms of fatigue-crack growth in a single crystal is
shown schematically in Figure 5 with the line ABC.
When the boundary between the crystallographic and
non-crystallographic stages is crossed at the point B, the
crack changes its growth direction. With the change in
the mechanism and the direction of the crack growth, the
speed of growth also changes.
Depending on the load level and the frequency, the
sample orientation, the geometrical parameters of the
plate with a hole and the material properties, different
schemes of crack growth are possible. Among them are
both the modes without a direction change (in one plane,
crystallographic or non-crystallographic) and modes
with one or multiple direction changes.
The extrapolation of the research results described in
12 for the conditions of thermal-cycle loading and, as a
final goal, the development of a new fracture criterion,
have become the object of the present research. At the
same time, we are not discussing here the problems of
the choice of models for the visco-elastic plasticity for
the calculation of parameter values of criterion (1).
2 TEST METHODS
The thermal fatigue tests were performed for samples
made of a single-crystal high-temperature alloy with five
different orientations, including <001>, <011> and
<111>. During the tests, the samples in sand-glass form
were rigidly fixed in vacuum 13.
A comprehensive procedure is developed in NPO
CKTI for the definition of the thermal fatigue resistance
of various materials and coatings, applying a special
appliance that allows us to clamp flat the sand-glass
shaped test pieces and to ensure their cyclic heating with
a conducting current (Figure 3). The heating takes place
according to a specified program (Figure 4), maintained
automatically during the testing. The appliance is fixed
in a vacuum chamber. The ultimate cycle temperature,
defined by the rate of oxide-film formation on the
sample surface, increases with the vacuum. The sample
material’s behavior at various surface points on the
sample is observed using a microscope with a 250×
magnification. During the test the following parameters
are recorded: the characteristic properties of the
deformation relief defining the mechanism of the
accumulation of thermal fatigue damage; the number of
cycles to the first microcrack formation in various
elements of the metal and the coating; the growth rate of
the incipient cracks; the number of cycles to sample
failure and the accumulated deformations in the ruptured
zone.
The range of the conditionally elastic stresses   and
the range of the total deformation   in the specimen’s
working part in the cycle were calculated from thee
equations:
∆σ α α ϕ= −( )max maxE T E Tst st21 1 2
∆σ α α ϕ= −( )max max1 2T T (4)
ϕ = −1 ∆ ∆k l/
where Est is the static elastic modulus; l is the free
travel of the test points during heating from Tmi n to
Tmax; and  k is the measured value of the displacement
of the control microhardness marks, applied to the
sample surface along its working-part edges during the
cycle.
In this calculation the values  k and  were used
after being averaged on the basis of the hypotheses of
their linear summability.
Some 9-mm-wide plates of a single-crystal alloy with
different orientations were prepared for the tests (see
Table 1).
Table 1: Sample orientation
Tabela 1: Orientacija preizku{ancev
Number of the
sample series Orientation
Deviation from
the exact axial
orientation in
degrees
Azimuthal
orientation of
the crystallo-
graphic planes
in degrees
1 111 5.64 8.26
2 011 4.51 11.27
3 011 8.33 14.43
4 011 9.67 7.86
5 001 5.47 41.97
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12 5
Figure 3: Shape of the specimen (a) and a sketch of the specimen with
a coating and with a central hole (b)
Slika 3: Oblika preizku{anca (a) in shema preizku{ancev s prekritjem
in z osrednjo izvrtino (b)
Some tests were made on samples with stress
concentrators in the form of holes with a diameter of 0.5
mm. The tests were conducted under different maximum
temperatures in the cycle Tmax (850–1050 °C) with and
without 2–5 min of holding at Tmax. The positions of the
slip lines on the sample’s surface were calculated and
confirmed experimentally, as was the crack orientation
up to the moment of the sample’s rupture. The values of
the irreversible deformations (characterizing the
ratcheting) taking place in the central part of the samples
were also determined.
For the interpretation of the test results we: a)
analyzed the character of the fracture from the four sides
of the sample (see Figure 5) and b) carried out
finite-element calculations for the change of the
sample’s deformation mode in order to find the
orientations of the rupture surfaces with the help of
methods for deformable solid-body mechanics and to
compare the results with crystallographic predictions and
experimental data.
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
6 Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12
Figure 4: Experimentally determined distribution of the temperature
along the length of the sample (à) and an example of the change of the
maximum temperature in the working part of the sample during the
cycle (b)
Slika 4: Eksperimentalno dolo~ena porazdelitev temperature po
dol`ini preizku{anca (a) in primer spremembe najvi{je temperature v
delovnem delu preizku{anca med ciklom (b)
O – Crystallographic fracture, O – Non-crystallographic fracture.
Figure 5: Interpretation of the fracture character of the samples 3-2
and 3-5
Slika 5: Interpretacija zna~ilnosti preloma preizku{ancev 3-2 in 3-5
3 EXPERIMENTAL RESULTS AND THEIR
ANALYSIS
The results of the performed experiments were
compared to the results of the crystallographic analysis
and the calculation of the samples’ deformation mode.
The stress analysis of the samples was performed with
the help of the PANTOCRATOR finite-element program
14 (see http://www.pantocrator.narod.ru) based on the
following assumptions:
1) in the tension phase, the cleavage cracks can
originate and grow, and their growth direction is
determined by the orientation of the platform of the
maximum principal value of the stress tensor,
2) in the compression phase, the shear cracks can
originate and grow, and their growth direction is
determined by the orientation of the platforms of the
maximum tangent stresses.
From now on the concept of non-crystallographic
fracture mode is used in the phenomenological context,
as a fracture (observed at the macro-level) not coinciding
with any of the crystallographic slip planes.
As an example, let us consider the place of origin of
a shear crack on the basis of the hypothesis of maximum
tangential stresses for a sample of the 3rd group tested in
the mode 150–900 °C. We found that there are two
equal, centrally symmetric maximums max = 856 MPa
with a bias located at a distance of 0.64 mm from the
center (see Figure 6), and the stresses localized in the
lateral zones (with a maximum on the center line) (see
Figure 6b).
For the same sample, we identified the tensile-crack
initiation point using the hypothesis of maximum
principal stresses. We found that there are two equal,
centrally symmetric maximums 1 = 1559 MPa at a
distance of 0.45 mm from the centre (see Figure 7);
there is also a localization of stresses in the lateral zones
(with a maximum at the edge) (see Figure 7b).
An analysis of the calculation results showed that the
distribution of stresses x and the locations of the
maximum stress zones are very different for different
groups of samples.
An example of the experimental data analysis is
shown in Table 2, containing the results for sample 2-1
tested under the mode 150–900 °C.
For this sample the crystallographic and the
finite-element results are similar and show only an
insignificant difference compared to the experiment (a
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12 7
Figure 7: Distribution of fields of maximum principal stress (sample
3-0, 3rd cycle)
Slika 7: Porazdelitev polj najve~jih glavnih napetosti (preizku{anec
3-0, 3. cikel)
Figure 6: Distribution of the fields of maximum tangential stress
(sample 3-0, 3rd cycle)
Slika 6: Porazdelitev polj najve~jih tangencialnih napetosti (preizku-
{anec 3-0, 3. cikel)
maximum of 4°). Symbolically, we designate fractures
of this type as "Mode 90 / 45".
Owing to the absence of pure modes of crystallo-
graphic or non-crystallographic fracture and, in some
cases, to the ambiguities of the choice between the
modes caused by the neighboring prognoses of the
crystallographic and the finite-element analysis, in order
to map the fracture mechanisms we use the classification
of modes on the basis of the orientation of the middle
fracture "plane" instead of the concepts of crystallo-
graphic or non-crystallographic fracture modes the four
modes are considered:
• Mode 90 / 45
• Mode 90 / 90
crystallographic or SSS
• Mode 45 / 45
• Mode 45 / 90
crystallographic
In Table 3 the results of the analyses of the fracture
modes for the samples from group 3 are shown. The
results show that the sample’s rupture occurred in the
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
8 Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12
Table 2: Comparison of the experimental data with the results of the
crystallographic and finite-element analyses of specimen 2-1
Tabela 2: Primerjava eksperimentalnih podatkov in rezultatov
kristalografske analize in analize po metodi kon~nih elementov za
preizku{anec 2-1
Expe-
riment crystallography
FEA
cleavage
mode
shear
mode
A 84 / 49 84 / 51 (1-11) 90 / 90 88 / 45
B 90 / 47 84 / 51 (1-11)
C 86 / 49 84 / 51 (1-11)
D 86 / 49 84 / 51 (1-11)
Table 3: Dominant fracture modes of the samples of the 3rd group ( -duration of cycle)
Tabela 3: Prevladujo~i na~ini preloma preizku{anca iz 3. skupine ( -trajanje cikla)
Non-FE, non-forecasting with the
finite-element analysis
crystallographic, the non-crystallographic and the mixed
modes. The conditions for the occurrence of the crack
type (dominant type (mode) of fracture) are generalized
with the help of maps of fracture mechanisms plotted in
the coordinates TMAX –  T (see, for example, Figure 8).
With the help of such maps of fracture mechanisms it is
possible to separate the loading conditions.
Fracture criterion. The above-mentioned facts
make it evident that, depending on the fracture mode of
single-crystal alloys, in each concrete case it is necessary
to choose the corresponding deformation-fracture
criteria:
• criterion 1:
D1( pl) + D2( c) + D3(pl) + D4(c) = 1 (4)
• criterion 2
D1( pl1) + D2( c1) + D3(pl1) + D4(c1) = 1 (5)
Namely, for crystallographically oriented fracture
modes, the following criterion parameters must be
chosen: shear deformations and their total amplitudes pl
and  pl (c,  c), and for non-crystallographic fracture
modes, the tensile deformations and their total ampli-
tudes pl1,  pl1, c1, and  c1.
Let us take as a basis the fact that the performed
thermal fatigue tests are described, on the one hand, by
the total number of cycles and, on the other hand, by the
level of the ratcheting deformation.
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12 9
Table 4: Comparison of the ratcheting deformation max with the thermal fatigue and plasticity (short-time plasticity r and creep-rupture pla-
sticity cr) at high temperature
Tabela 4: Primerjava prelomnih deformacij  max s termi~no utrujenostjo in plasti~nostjo (kratkotrajna plasti~nost  r in plasti~nost pri prelomu z
lezenjem  cr) pri visoki temperaturi
Orientation Tmax/  Time ofstay, min  max/%
 max1/
 max2
N N2/N1
 r
%
at Tmax
 cr
%
at Tmax
D3 D4
001 900 0 9.0 0.75 560 0.17 27 12.8–16.4 0.1 0
1000 0 12 95 19 8–23 0.25 0
111 900 0 10.724.3*
823
50 19,5 14–24
0.22
0.5 0
2 21 140 0 0.35–0.6
5 18.3 16 0 0.3–0.52
1000 2 23.7 194 21,5 7.2–23.8 0 0.4–1.3
011(2) 850 0 9.3* 2952
900 0 15.3 100
1000 0 20 0.67 472 0.67
5 29.7 317
0 10.6* 0.77 187 0.33
2 13.7* 62
011(3) 900 0 10.3 951
950 0 4.3 0.4 2535 0.18
950 0 10.7 450
1000 0 7.3 0.6 1220 0.05
1000 0 12 63
1050 0 14.7 356
011(4) 900 0 10 0.470.55 308
0.06
0.08
900 2 21.3 17
900 5 18 26
900 0 9.7* 25
950 0 9.7 0.94 626 0.20
950 2 10.3 128
*-specimen with concentrator
∆
T/
°C
Tmax/°C
Figure 8: Map of the fracture mechanisms for the samples of the 3rd
series
Slika 8: Shema mehanizma preloma za preizku{ance iz 3. serije
The data in Table 4 show that a) a decrease of the
ratcheting deformation results in a sample lifetime
increase, b) the number of cycles to fracture in tests with
holding at TMAX is considerably lower than the numbers
of cycles in the case of tests without such holding. The
research of Sizova R.N.15 has shown that, in conditions
of compression, the time to fracture is 2.5–3.0 times
longer than in the case of tension. Using the deformation
criterion of destruction, it is equivalent to the updating of
the ratio D3 = max/r and D4 = max/cr, with the help of
the correction factor 0.4. We then have:
D3 = 0.4 max/r (6)
D4 = 0.4 max/cr (7)
where r and cr are the corresponding deformations in
the case of short-time fracture and fracture in the creep
test at TMAX under tension.
Let us compare the deformations max for two tests,
max1 and max2. We can see that the relation (max1/max2 –
N2/N1) varies from 0.12 up to 1.00. This means that the
ratcheting deformation value cannot characterize the
fracture conditions and only helps to identify the trend.
Here are max = max (1, 2), where 1 and 2 are the
deformations measured in two directions of the sample’s
cross-section.
To apply the proposed deformation criterion (1) to
the results of the performed tests it is necessary: a) to
calculate appropriately (using creep parameters and
deformation curves in conditions of short-time tension)
the width of the hysteresis curve, with reference to the
known temperature-change graph, b) to have the values
of the short-time and long-time plasticity at TMAX. A
comparison of the single-crystal alloy plasticity values r
obtained during the simple tension tests and the r in
long-term tests prove that, like for polycrystalline mate-
rials, in creep conditions the plasticity of the material is
less than the short-time plasticity.
Having the experimental values of ratcheting
deformations and supposing that the exponents in the
expressions for D1 and D2 are equal to 2.0 and 1.25 (like
in the case of polycrystalline materials), the calculation
of the number of cycles until fracture turns out to be
dependent on the two material parameters and the maxi-
mum temperature in the cycle. As a first approximation,
these material parameters can be evaluated from the
values r and cr (C1 = (0.5r)2, C2 = (0.75 cr)1,25).
Let us consider some results of the calculation of the
interruption mechanism damage. The calculation of the
damage D1 and D2 (see Table 5) was carried out with
reference to the number of cycles before the occurrence
of the main crack, Nm. The size of the range of the
non-elastic deformation  pl and  c was determined as
the difference of the range of full deformations,  , and
the sizes 2-1/E. Thus, it was accepted that the cyclic
limit of the elasticity is equal to the fatigue strength at
the maximum temperature of the cycle. For the calcu-
lation the D1 and D2 values of plasticity were accepted
under a simple tension (r) and at the long rupture time
(cr) given in Table 6. The values of the damage D3 and
D4 are calculated by means of Equations (6) and (7) and
the experimentally obtained values of ratcheting and
plasticity during compression. The data in Table 5
confirm the satisfactory description of the conditions of
the origin of the main cracks found during thermocyclic
loading of a single-crystal alloy.
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
10 Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12
Table 6: Input data
Tabela 6: Vhodni podatki
Temperature/
°C Elongation ( r)/% Long time elongation ( cr)/% Fatigue limit -1/MPa
001 111 001 111 001 011 approximately 111
900 27 19.5 8.6–16.4 14.1–24 350 360 430
1000 19 21.5 8–23 7.2–23.8 – – –
1100 22 21.5 – – 170 170 170
Table 5: Calculation of damage by means of (5)
Tabela 5: Izra~un po{kodb z uporabo (5)
Orien-
tation
The
minimum
temperature
of a cycle,
°C
The
maximum
tempera
ture of a
cycle, °C
Time of
cycle,
sec
Number of
cycles before
formation of
the main crack,
Nm
 
%
2-1/E
 el
%
 pl,c
%
D1 D2 D3 D4 Di
111
150 900
72 190 0.73 860/2.4·105 0.36 0.37 0.273 0 0.22 0 0.59
247 50 0.69 860/2.4·105 0.36 0.33 0 0.34 0 0.35 1.27
378 12 0.73 860/2.4·105 0.36 0.37 0 0.09 0 0.52 0.61
500 1000 149 80 0.64 648/2.3·105 0.28 0.36 0 0.60 0 0.40 1.00
001
150 900 72 500 0.86 700/0.998·105 0.70 0.16 0.013 0 0.10 0 0.11
250 1000 48 40 0.855 500/0.95·105 0.52 0.33 0.18 0 0.25 0 0.43
500 1000 28 1400 0.655 500/0.95·105 0.52 0.13 1.13 0 0.25 0 1.38
Shear-strain computation. The criterion (4) re-
quires a shear-strain ( pl,  c, pl, and c) computation.
The shear strain nl at the plane with the normal vector n
in the slip direction l is defined on the basis of the strain
tensor  by the relation
γ ε εnl nl= = ⋅ ⋅2 2n l (8)
or, by using a presentation with components
γ ε ε ε εnl nl= = + + +2 2(n l n l n lx xx x x xy y x xz z
+ + + +n l n l n ly xy x y yy y y yz zε ε ε (9)
+ + +n l n l n lz xz x z yz y z zz zε ε ε )
The presented equations are valid for all cases,  pl,
 c, pl and c, with a special choice of the strain tensor
for each case.
The components of the normal vector n and the slip
direction l in the global coordinate system are related to
the components in the crystallographic coordinate
system by the following relations
n A n A n A n
n A n A n
x
y
= + +
= +
< > < > < >
< > <
11 100 21 010 31 001
12 100 22 010 32 001
13 100 23 010 33 001
> < >
< > < > < >
+
= + +
A n
n A n A n A nz
(10)
and
l A l A l A l
l A l A l
x
y
= + +
= +
< > < > < >
< > <
11 100 21 010 31 001
12 100 22 010 32 001
13 100 23 010 33 001
> < >
< > < > < >
+
= + +
A l
l A l A l A lz
(11)
where Aij is the matrix of crystallographic orientations,
which is different for each specimen.
The results of the computations of nl for all the
specimen series are given in Table 7. The diagonal
components of the strain tensor are obtained from the
experimental data and the non-diagonal components are
taken from the finite-element simulations.
The results of the calculation of damage with the use
of criterion (4) are shown in Table 8.
For the definition of the damage size the following
relations were used:
D
C n
1 i
pl k= ∑1
3
( )∆γ ; D
C n
2 i
pl k= ∑1
4
( )∆γ ;
D3 max r= 0 4. /γ γ ; D4 max cr= 0 4. /γ γ (12)
where the values of the constants C3, C4,  , and cr
were defined using the ratio
C3 = 2.25  1; C4 = 1.66  2; r = 1.5 r; cr = 1.5 r (13)
For the orientation  111, the sum of the damage
agrees with the criterion (4) at the moment of formation
of the main crack, and appears to be much more than 1.0,
while calculations that use this criterion (5) are in good
agreement with the experiment. Thus, criterion (4) gives,
in this case, an underestimated number of cycles. At the
same time the use of this criterion for two samples with
the orientation  001 gives the best agreement with the
experiment in comparison to criterion (5), and for one
specimen the agreement is worse.
The use of diagrams of maps of destruction (see
Figure 8) allows us to predict the criterion of destruction
during the thermal cyclic loading of objects made from
single-crystal alloys.
The use of the forgoing modification of the defor-
mation fracture criterion will make it possible to solve
reliably the problems of the prediction of the conditions
of initiation of the thermal-fatigue cracks in the turbine
blades made of single-crystal alloys.
4 CONCLUSIONS
The performed experimental investigation and the
analysis led to the following conclusions:
1. The character of fracture of single-crystal alloys
under thermal-cycle loading depends on the cry-
stallographic orientation of the material and on the
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12 11
Table 8: Calculation of damages by means of (4)
Tabela 8: Izra~un po{kodb z uporabo (4)
Orien-
tation
The
minimum
tempera ture
of a cycle,
°C
The
maximum
temperature
of a cycle,
°C
Time of
cycle, s
Number of
cycles before
formation of
the main
crack, Nm
pl,c
%
D1(pl) D2(c) D3(pl) D4(c) Di
111
150 900
72 190 1.57 2.19 0 0.62 0 2.81
247 50 1.40 0 3.6 0 0.99 4.59
378 12 1.57 0 0.16 0 1.47 1.63
500 1000 149 80 1.53 0 11.7 0 1.13 12.83
001
150 900 72 500 0. 345 0.027 0 0.144 0 0.171
250 1000 48 40 0.71 0.37 0 0.36 0 0.73
500 1000 28 1400 0.28 2.34 0 0.36 0 2.70
Table 7: Relations between the shear strain γnl and the axial strain  x
Tabela 7: Razmerje med stri`no deformacijo nl in axialno defor-
macijo  x
Specimen series Specimenorientation nl/ x
1. 111 4.25
2. 011 4.50
3. 011 4.09
4. 011 2.80
5. 001 2.16
temperature and time parameters of the loading
mode.
2. The difference of the fracture mechanisms for each
case of crystallographic orientation of the material
can be represented with the help of the proposed
fracture maps using the coordinates TMAX – T.
3. New definitions for the deformation fracture criteria
of single-crystal alloys in conditions of thermal-cycle
loading have been proposed; they make it possible to
evaluate the lifetime of blades made of single-crystal
alloys using a computational-experimental approach.
In summary, the use of the deformation criterion for
the calculation of fracture conditions for single-crystal
alloys submitted to thermal-cycle loading requires, on
the one hand, a knowledge of the creep and plasticity
characteristics of the material for different crystallo-
graphic orientations in the working temperature range,
and, on the other hand, the appropriate choice of the
thermo-visco-elasto-placticity model.
5 REFERENCES
1 L. B. Getsov: Problems of creation of the "universal" theory of
destruction of materials. Journal of machinery manufacture under
reliability problem. (2001) 5, 49–55
2 K. M. Kononov, L. B. Getsov: Failure criteria of materials under
cyclic loading. Strength of materials, N2, 1984
3 P. A. Pavlov, L. B. Getsov: Deformation criteria for material rupture
estimation in complex stress-strein conditions under cyclic loadings.
Strength of materials, N3, 1989
4 L. B. Getsov, A. A. Nigin, M. G. Kabelevskiy: Use of finite-element
method for numerical evaluation of thermal cycling endurance of
disks. Strength of materials, N4, 1979, 17
5 Cunningharn S. E., DeLuca D. P., Haake F. K.: Crack growth and
life prediction in single crysta1 nickel Superalloys. Vol. I,
WL-TR-94-4089, 1994
6 Leverant G. A., Gell M.: The influence of temperature and cyclic
frequency on the fatigue fracture of cube oriented nickel-base
superalloy single crystal. Metall. Trans. A. 6 (1975), 367–371
7 Telesman J., Ghosn L. J.: Fatigue Crack Growth Behavior of a PWA
1484 Single Crystal Superalloy at Elevated Temperatures. ASME
Paper 95-GT-452, 1995
8 Chan K. S., Feiger J., Lee Y.-D., John R., Hudak S. J.: Fatigue Crack
Growth Thresholds of Deflected Mixed-Mode Cracks in PWA1484;
Journal of Engineering Materials and Technology.127 (2005), 2–7
9 Frost H. J., Ashby M. F.: Deformation- mechanism maps. 1982
10 Report P&W. Dayton, 2005
11 Shalin R. E., Svetlov I. L., Kachanov E. B. et al. Single crystals of
nickel base superalloys. Machinery, 1997, 333
12 Dulnev R. A., Svetlov I. L., Bychkov N. G. et al. Orientation depen-
dence of thermal fatigue of nickel alloy single crystal; Strength
problems (1988) 11, 3–9. (In Russian)
13 A. I. Rybnikov, L. B. Getsov: New technique and results of thermal
fatigue tests of superalloys and coatings; Proceedings of the sixth
International congress on thermal stresses. Vienna, Austria, (2005) 1,
305–309
14 Semenov A. S. PANTOCRATOR – Finite-element program spe-
cialized on the solution of non-linear problems of solid body
mechanics; Proc. V-th Int. Conf. "Scientific and engineering
problems of predicting the reliability and service life of structures
and methods of their solution", St-Petersburg, 2003, 466-480
15 Thermal strength of machine details. The theory. Experimental
researches. Calculation. Editor Birger I. A. and Shorr B. F. Mecha-
nical engineering, 1975 (In Russian)
L. GETSOV ET AL.: A FAILURE CRITERION FOR SINGLE-CRYSTAL SUPERALLOYS ...
12 Materiali in tehnologije / Materials and technology 42 (2008) 1, 3–12
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
ACCELERATED CREEP TESTING OF NEW CREEP
RESISTING WELD METALS
PREIZKUSI POSPE[ENEGA LEZENJA ZVAROV NOVEGA JEKLA,
ODPORNEGA PROTI LEZENJU
Stan T. Mandziej,1 Anna Výrostková,2 Mojca [olar3
1Advanced Materials Analysis, P.O.Box 3751, 7500 DT Enschede, Nederland
2Institute of Materials Research SAS, Watsonova 47, 043 53 Kosice, Slovakia
3Elektrode Jesenice, Cesta @elezarjev 8, 4270 Jesenice, Slovenia
100422.3355 compuserve.com
Prejem rokopisa – received: 2007-09-04; sprejem za objavo – accepted for publication: 2007-10-11
New creep-resisting weld metals of P91 grade were manufactured by MMA process producing multi-bead multi-layer test
coupons, from which specimens were taken for accelerated creep testing on Gleeble physical simulator. The recently developed
accelerated creep testing (ACT) procedure on Gleeble allows transforming in a relatively short time (less than 100 hours) the
microstructure of creep-resisting materials near to the thermodynamic equilibrium state, resembling that of multi-years
application at creep condition. Such advanced transformation of microstructure in the investigated weld metals after ACT was
confirmed with microscopy and microanalytical study. The results also appeared to be in agreement with Thermocalc
calculations. Using optimum combination of chemical compositions of the electrodes with welding procedure and post-weld
heat-treatment, advantageous mechanical properties were achieved as well as results of ACT indicating potentially long creep
life in exploitation conditions. In discussion of the ACT results, comparisons with results of conventional creep testing have
been made. The ACT procedure appears to be useful in fast screening of newly developed creep resisting materials.
Keywords: P91 welds, accelerated creep tests, microstructure, Gleeble
Ve~varkovni zvari jekla P91, odporni proti lezenju so bili pripravljeni po MMA-postopku in iz njih so bili izdelani ve~varkovni
vzorci za preizkuse na simulatorju Gleeble. Pred kratkim razvit preizkus pospe{enega lezenja (ACT) na napravi Gleeble
omogo~a, da se pretvori v relativno kratkem ~asu (manj od 100 h) mikrostruktura jekla, odpornega proti lezenju, v stanje blizu
termodinami~nega ravnote`ja, ki je podobno kot po mnogih letih obremenitve z lezenjem. Tako pretvorbo mikrostrukture po
ACT smo potrdili z mikroskopsko in mikroanalitsko preiskavo. Zdi se, da so rezultati v skladu z izra~uni na podlagi podatkov
Themocalc. Z uporabo optimalne kombinacije kemi~ne sestave elektrod, postopka varjenja in toplotne obdelavo po varjenju so
bile dose`ene dobre mehanske lastnosti in rezultati ACT, ki nakazujejo na potencialno dolgotrajno stabilnost pri obremenitvi z
lezenjem. V razpravi primerjamo rezultate ACT z rezultati konvencionalnih preizkusov lezenja. Videti je, da je ACT-postopek
koristen za hitro oceno na novo razvitih jekel, ki so odporna proti lezenju.
Klju~ne besede: zvari P91, poskusi pospe{enega lezenja, mikrostruktura, simulator Gleeble
1 INTRODUCTION
Design of the power plant components and
estimation of power plants lifetime are based on
long-term creep data, which are generally available for
the manufactured plate and pipe materials but seldom
available for the welds on the components of the power
plants. This last is mainly due to the large variety of
factors appearing in the fabrication procedure, which
affect the creep strength and creep life of the welds. To
predict the exploitation behaviour of the welds under
creep, especially of the repair welds, accelerated creep
testing (ACT) procedures can be used. Such procedures
of accelerated tests often apply small strains and/or
constant slow deformation rates to speed-up transfor-
mation of microstructure towards that of metal alloys
exploited for many years at creep conditions 6. In
general, they aim to determine in a relatively short time
the remaining strength or ductility which the creep
resisting materials would have after long exposure to
stresses at the elevated temperatures characteristic of
their exploitation.
An accelerated creep test recently developed by the
Advanced Materials Analysis, Enschede, Netherlands, in
collaboration with Dynamic Systems Inc, Poestenkill
NY, USA 4, is a low-cycle thermal-mechanical fatigue
procedure executed on Gleeble physical simulator, based
on the actual knowledge of the micro- and sub-structural
changes causing decay of properties during the creep. To
simulate the situation of the material during creep, the
following demands had to be achieved in the ACT:
– Total deformation at fracture being like at real creep
– just a few percent only.
– Depletion of material’s matrix in alloying elements
similar to that of long-term crept steels and the
carbide phases at onset of cracks being not different.
– Basic temperature and applied strains in the ACT
preventing odd transformations like dissolution of
carbides and/or formation of far from equilibrium
phases.
These aims have been reached by considering that
certain dislocations can effectively dissolve and then
transport interstitial elements like C and N, and that the
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 13
UDK 539.4:621.791.05:620.18 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(1)13(2008)
annihilation of these dislocations may control the
precipitation of carbides 3. Thus accumulation of strains,
generating appropriate dislocation configurations and
increasing level of stored energy in the steel, was used to
speed-up the transformation of microstructure. An
additional factor used in the development of the ACT,
known for a long time as capable of accelerating
recrystallization of matrix and coagulation of carbides 1,
is an electric resistance heating with controlled thermal
gradients.
An initial version of this ACT procedure 4 was imple-
mented in the "SmartWeld" EU R&D project and on
purpose modified in the course of it 5. What follows in
this article is the presentation of selected ACT results
and their comparison with the conventional creep-
rupture tests. Further, the reliability of the ACT is
confirmed by the microstructure evolution i.e. the
precipitation processes and sequence of changes leading
to failure, that reproduce the situations occurring in real
creep. Presented here are data obtained for P91 type
weld metals, manufactured in the company Elektrode
Jesenice from Slovenia.
2 THE ACCELERATED CREEP TEST ON
GLEEBLE – ACT
To carry out the simulative accelerated creep tests,
Gleeble 3500 physical simulator was used. In this
simulator in a rod-like sample made of the
electrically-conductive alloy the balance between
electric resistance heating of the sample and the heat
flow by thermal conductivity from this sample towards
the cold copper jaws allows forming a uniformly heated
zone in the middle-span of the sample (see Figure 1). To
define this zone better, on samples for the ACT a gauge
portion of reduced diameter was made. Samples for the
ACT were taken across multi-bead, multi-layer P91
MMA welded joints, as shown in Figure 2. The samples
were in as-welded (marked 5s and 10s) and in post-weld
heat-treated condition (marked 6s and 8s). Their
chemical compositions are given in Table 1.
Most of the ACT samples were tested till fracture
occurred, however some of them especially those for
fine fractographic observations and microanalytical
investigations were stopped before fracturing. As the
tests for different materials were run at different
temperatures, the need occurred to compare the ACT
results obtained at various temperatures, to the duration
of the test and its temperature had to be included in the
following parameter:
PACT = (7 + lg t) · T/100
where: t = time of test in ks, and T = temperature in K.
Then, the creep strength factor in ACT has been
calculated as
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
14 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Table 1: Chemical compositions of P91 weld metals
Tabela 1: Kemi~na sestava zvarov P91
Material Element, w/%
C Mn Si Cr Ni Mo Nb V Ti W N
5s & 8s 0.06 0.59 0.36 9.49 0.94 0.98 0.005 0.21 0.000 0.21
6s 0.06 0.60 0.37 9.12 0.90 0.97 0.005 0.20 0.000 0.20
10s 0.09 0.84 0.40 8.70 0.40 0.93 0.061 0.21 0.008 0.01 +
Figure 2: Cross-weld samples used for ACT on all-weld-metal and on
weld’s HAZ in the "SmartWeld" EU project
Slika 2: Preizku{anci z zvarom uporabljeni pri ACT, ki obsegajo cel
zvar in toplotno cono pri projektu EU "SmartWeld"
Figure 1: Example of a uniformly heated zone formed in Gleeble at
the mid-span of the rod-like specimen (at temperature much higher
than used in the ACT)
Slika 1: Primer enakomerno ogrete cone na napravi Gleeble na polo-
vici dol`ine pali~astega preizku{anca (pri temperaturi, ki je mnogo
nad tisto, uporabljeno pri preizkusih Gleeble)
FACT = PACT · RS /100
where RS is the average stress of all ACT cycles
measured during relaxation on tension.
Selected examples of strain-time and stress-time
graphs from ACT are in Figures 3 to 6; in the "soft", e.g.
over-tempered material like sample 6s, the zero-stress
line on the strain-time curve resembles well a normal
creep graph with its three characteristic stages, see
Figure 3, while for the "hard" materials like as-welded
sample 5s and 10s, a straight continuous elongation in
each cycle was observed from beginning till the end of
test, Figure 5.
Selected examples of the ACT results on P91 weld
metals are given in Table 2 below. They show in general
higher creep strength of non heat-treated weld metals,
with tendency of increasing time to fracture by PWHT
due to changing of failure micromechanism (compare 5s
with 6s). They also show positive results of optimizing
heat treatment (6s vs. 8s) as well as of the improved
chemical composition (5s vs. 10s).
Table 2: Examples of ACT results
Tabela 2: Primeri ACT-rezultatov
Sample
number
Material
& state
ACT
temp.
°C
ACT
time to
fracture
ks
Tensile
relax
stress
RS/MPa
Creep
strength
factor
FACT/
MPa
5s P91-AW 600 28.3 325 240
6s P91-HT 620 45.2 177 137
8s P91-HT 600 26.3 318 233
10s P91-AW 600 82.6 336 262
In parallel, short-term creep-rupture tests (STCT)
were run on these materials and the results of both tests
compared. Selected results of the STCT are given in
Table 3 and compared with those of the ACT. The creep
strength factor – FCS – given in this table was calculated
from the STCT results like:
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 15
Figure 6: Typical stress-time graph from ACT on a "hard" weld metal
– sample 5s
Slika 6: Zna~ilna odvisnost napetost-~as za ACT trdega materiala
zvara – vzorec 5s
Figure 3: Typical strain-time graph from ACT on a "soft" weld metal
– sample 6s
Slika 3: Zna~ilna odvisnost deformacija-~as za ACT mehkega mate-
riala zvara – vzorec 6s
Figure 5: Typical strain-time graph from ACT on a "hard" weld metal
– sample 5s
Slika 5: Zna~ilna odvisnost deformacija-~as za ACT trdega materiala
zvara – vzorec 5s
Figure 4: Typical stress-time graph from ACT on a "soft" weld metal
– sample 6s
Slika 4: Zna~ilna odvisnost napetost-~as za ACT mehkega materiala
zvara – vzorec 6s
FCS = PLM · RL/100
Where RL is the initial stress of the constant-load
short-term creep experiment and PLM is the Larson-
Miller parameter: PLM = (20 + lg t) · T/100, with T =
testing temp. in K and t = time to rupture in h.
3 INITIAL MICROSTRUCTURES
3.1 As welded samples 5s and 10s
The main component of the microstructure is acicular
low-carbon martensite with high density of tangled
dislocations (Figures 7–10). In this martensite numerous
spheroidal inclusions of sub-micron size are present. In
sample 5s fine cementite and very fine MX particles
interact with dislocations, Figure 8. In sample 10s, in the
inter-bead heat-affected zones along grain boundaries of
former austenite, chains of middle-sized M23C6 carbides
appear, Figure 9. In the columnar grain regions of very
high dislocation density in sample 10s, (Figure 10),
hexagonal M2X particles as well as MX particles were
identified. This martensitic microstructure should be
considered as auto-tempered or tempered due to welding
cycles, more in sample 10s than in 5s.
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
16 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Table 3: Comparison of results: ACT vs. STCT
Tabela 3: Primerjava rezultatov ACT in STCT
Sample No. Material & state STCT results:
 temp/stress/time FCS/MPa FACT/MPa FCS/FACT
5s P91-AW 620/130/381 262 240 1.09
6s P91-HT 620/130/27 248 137 1.82
8s P91-HT 620/130/31 249 233 1.07
Figure 10: High density of dislocations in martensite of as-welded
sample 10s
Slika 10: Velika gostota dislokacij v martenzitu v vzorcu zvara 10s
Figure 8: Cementite precipitated on dislocation tangles in as-welded
sample 5s
Slika 8: Izlo~ki cementita na vozli{~ih dislokacij v vzorcu zvara5 s
Figure 9: Carbides on former austenite grain boundaries in as-welded
sample 10s
Slika 9: Karbidi na avstenitnih mejah v vzorcu zvara 10s
Figure 7: Acicular low-carbon martensite in as-welded sample 5s
Slika 7: Acikularni malooglji~ni martenzit v vzorcu zvara 5s
3.2 Post-weld heat-treated samples 6s and 8s
The post-weld heat-treated samples 6s and 8s had a
tempered martensite microstructure with numerous
subgrains, often retaining orientation of the former
martensite laths (Figures 11–13). In both 6s and 8s
samples, the former austenite grain boundaries as well as
the post-martensitic lath boundaries were marked with
coagulated M23C6 carbides. A very high dislocation den-
sity appeared in grains and subgrains of the inter-bead
heat affected zones in sample 8s where also spheroidal
carbides were present (Figure 14), while the overall
microstructure of the sample 6s was much better
recovered i.e. the dislocation density in it was visibly
lower (Figure 12).
4 DEVELOPMENT OF MICROSTRUCTURE
DURING ACT
Most of the accelerated creep tests were carried out
up to appearance of fracture, in order to gain the date
representative to the creep life of the materials.
Nevertheless, for the purpose of this study on the
sequence of precipitation processes and changes in the
weld metal matrix, several tests were interrupted and
microstructure of the partly transformed samples
examined.
4.1 ACT of sample 5s
The accelerated creep test of the as-welded sample
results in transformation of the matrix and stimulates
precipitation of carbides. Semi-recovered subgrains sized
about 1–2 µm form in the matrix before the sample
reaches 50 % of its life time in ACT (Figure 15). By this
stage also the coalescence of subgrains begins, causing
oriented chains of carbides to appear within the grains
(Figure 16). By 80 % of the ACT duration the
recrystallising ferrite grains reach 5 µm or more due to
the coalescence of finer subgrains (Figure 17). At this
stage, well-pronounced agglomerates of coagulated
carbides can be observed along the former austenite
grain boundaries, in particular near to the "triple points"
(Figure 18). These carbides were mainly of M23C6 type,
although some M6C carbides could be identified as well.
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 17
Figure 14: Spheroidal carbides in ferrite in the sample 8s
Slika 14: Sferoidirani karbidi v feritu v vzorcu 8s
Figure 12: Chain of carbides along grain / subgrain boundary in
sample 6s
Slika 12: Niz karbidov na meji zrno-podzrno v vzorcu 6s
Figure 13: Fine ferrite grains/subgrains with high dislocation density
in sample 8s
Slika 13: Majhna zrna/podzrna ferita v vzorcu 8s
Figure 11: Arrays of ferrite subgrains and elongated grains in
tempered sample 6s
Slika 11: Mre`a podrzrn ferita in podolgovata zrna v popu{~enem
vzorcu 6s
The finest precipitates dominating within the
recovered post-martensitic subgrains, were positively
identified as of cubic MX type, while some larger and
elongated appeared to be of hexagonal M2X type. On
thin foil specimens, taken from near to fracture after
completion of the ACT and examined in TEM, still some
of these fine precipitates could be seen within the
recrystallised ferrite grains.
4.2 ACT of samples 6s and 8s
In the over-tempered sample 6s, having already a
well-recovered ferrite matrix with low dislocation
density before the ACT, the ACT resulted in a fast
recrystallisation of the matrix by coalescence of
subgrains and coagulation of M23C6 carbides (Figures 19
and 20). The fine precipitates of MX carbo-nitrides
disappeared almost entirely and the fine acicular M2X
phase also could not be found. Some larger precipitates
appearing after completion of the ACT were identified as
Laves phase.
The sample 8s having an initially less transformed
matrix than 6s, with a high dislocation density, around
the half of its life in the ACT showed substantial
recovery of the dislocation substructure and intensive
precipitation of mainly MX phase on the dislocations
(Figures 21 and 22). This situation persisted till 80 % of
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
18 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Figure 18: Agglomerate of carbides in the sample 5s at 80 % of life in
ACT
Slika 18: Aglomerat karbidov v vzorcu 5s pri 80 % trajnostne dobe v
ACT
Figure 15: Formation of subgrains in sample 5s at half-life in ACT
Slika 15: Nastanek podrzrn pri polovici trajnostne dobe vzorca 5s v
ACT
Figure 19: Fine recrystallised ferrite grains after ACT in sample 6s
Slika 19: Drobna rekristalizirana zrna ferita v vzorcu 6s po ACT
Figure 17: Incipient coalescence of subgrains in sample 5s at 80 % of
life in ACT
Slika 17: Za~etek koalaescence podzrn v vzorcu 5s pri 80 %
trajnostne dobe v ACT
Figure 16: Chain of carbides inside growing subgrain in sample 5s
Slika 16: Nizi karbidov v rasto~ih podzrnih v vzrocu 5s
the ACT duration, when visible coarsening of fine
precipitates within subgrains occurred and more
coagulated carbides of M23C6 and M6C appeared on
subgrain boundaries, as well as coalescence of subgrains
progressed. Further on, till fracturing of the ACT
sample, the fine precipitates within ferrite grains still
existed (Figure 23), interacting with dislocations as well
as pinning ferrite grain boundaries (Figure 24).
Agglomerates of larger precipitates consisted mainly of
M23C6 and M6C, and minor amounts of Laves phase
particles were associated with them.
4.3 ACT of sample 10s
The as-welded initial microstructure of sample 10s
contained already the M23C6 carbides along prior
austenite boundaries and MX plus M2X phases within
martensite laths before ACT. In the first half of the ACT
duration the M23C6 carbides grew while the fine MX and
M2X prevented the dislocation recovery within subgrains
(Figures 25 and 26). Around 80 % of the ACT life,
mixed substructure dominated consisting of small
recrystallised grains co-existed with non-recovered and
well-recovered subgrains. By the end of ACT, the
recrystallised ferrite grains (Figure 27), still co-existed
with the subgrains, in which fine precipitates interacted
with the medium-density configurations of recovered
dislocations (Figure 28).
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 19
Figure 21: Subgrains with medium–density dislocations in sample 8s
at half-life in ACT
Slika 21: Podzrno s srednjo gostoto dislokacij v vzorcu 8s po polovici
ACT trajnostne dobe
Figure 20: Spheroidised carbides after ACT in sample 6s
Slika 20: Sferoidizirani karbidi v vzorcu 6s po ACT
Figure 24: Agglomerate of platelike carbides at grain boundary after
ACT; sample 5s
Slika 24: Aglomerat plo{~atih karbidov na kristalni meji po ACT
Figure 22: MX precipitates on dislocations in sample 8s
Slika 22: Izlo~ki MX na dislokacijah v vzorcu 8s
Figure 23: Recrystallised small grain with a few fine precipitates after
ACT; sample 8s
Slika 23: Majhno rekristalizirano zrno s majhnimi izlo~ki po ACT,
vzorec 8s
5 DISCUSSION
The results of this work show the behaviour of some
P91 grade weld metals during accelerated creep tests.
They allow comparing the standard / lean as-welded
material 5s with the optimized one 10s, as well as
evaluate the influence of heat treatment on the creep
strength on the example of samples 6s and 8s.
Characteristic of the optimized weld metal 10s in its
initial state were tortuous prior austenite grain
boundaries in the inter-bead heat affected zones (Figure
29). This element of the microstructure was retained
during the ACT and after completion of the test the
precipitated carbides formed a pattern along the
meandering grain boundaries (Figure 30). This might
have been one of major factors responsible for the
enhanced creep strength of the sample 10s in ACT and
for its long creep life. In the case of the post-weld
heat-treated weld metals 6s and 8s, evidently the heat
treatment reduces the elevated temperature strength and
for the samples tested by ACT in the simulated creep
conditions this meants reduction of the creep strength
and creep life. The accelerated creep test evidently
causes further transformation of the microstructure and
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
20 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Figure 28: Recovered dislocations after ACT interacting with fine
carbides; sample 10s
Slika 28: Dislokacije po popravi, ki reagirajo s drobnimi karbidi po
ACT v vzorcu 10s
Figure 26: Precipitated coagulated carbides in sample 10s
Slika 26: Izlo~eni koagulirani karbidi v vzorcu 10 s
Figure 25: Subgrains with high dislocation density at half-life in
ACT; sample 10s
Slika 25: Podzrno z visoko gostoto dislokacij v vzorcu 10s pri
polovici ACT trajnostne dobe
Figure 27: Small recrystallised ferrite grain with precipitates after
ACT in sample 10s
Slika 27: Majhno rekristalizirano zrno z izlo~ki po ACT vzorca 10s
Figure 29: Tortuous prior austenite grain boundaries in as-welded
sample 10s
Slika 29: Vijugasta meja avstenitnega zrna v zvaru vzorca 10 s
enhances the precipitation. The question to be answered
is whether all the ACT samples fail at the adequate
transformation of the microstructure, similar to that of
the real creep and are comparable with the results of
other creep tests.
The results of EDS x-ray microanalysis are
summarized in Table 4 below, together with Thermocalc
data. Here (L) indicates the Laves phase present only in
a part of the temperature range, while (Z) probably the
z-phase not confirmed by diffraction. Then (x) refers to
precipitates found on carbon extraction replica taken
from ACT fracture while  x the phases identified by
SAD in thin foil specimens in TEM.
This table shows quite good correlation of ACT
results with Thermocalc, except for the M7C3 presence
on the replicas taken from the ACT fracture of the 8s
sample. However, this non-equilibrium carbide may
form when vast amount of dislocations annihilates, like
during the recrystallisation after cold drawing of the
10Cr-2Si-Mo steel 2. In fact, the ACT procedure uses a
similar phenomenon of generating dislocations, allowing
them to "inhale" interstitials and delivering the
interstitials to the sites where dislocations annihilate.
Another question is due to the difference shown in
Table 3 between ACT and STCT results on sample 6s,
considered as over-tempered. In this sample from the
beginning till the end of the ACT the finely-dispersed
MX phase was almost absent. The MX is considered as
an important creep strengthener, stable up to above 700
°C, interacting with dislocations and grain boundaries in
steels exposed to creep 7. Thus, in a 1.5Cr-1Mo-0.25V
grade steel after exposure to creep for 18 years the MX
is still present in most of ferrite grains (Figure 31). In
the STCT tests of this research, the fine MX phase
particles were only found in the as-welded material 5s, in
minor amount of subgrains on thin foils taken from near
to grip portions of the samples. Their traces could also
be found in gauge portion of 8s STCT sample (Figure
32), but entirely not in the neck portion of 5s STCT
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 21
Figure 31: Fine MX precipitates in crept for 18 years
1.5Cr-1Mo-0.25V steel
Slika 31: Drobni MX-izlo~ki v jeklu z 1.5Cr-1Mo-0.25V, ki je bilo 18
let izpostavljeno lezenju
Figure 30: Meandering pattern of carbides precipitated after ACT in
sample 10s
Slika 30: Meandraste oblike karbidov, ki so se izlo~ili po ACT v
vzorcu 10s
Table 4: Phase compositions of selected P91 weld metals (EDS microanalysis from carbon extraction replicas)
Tabela 4: Fazna sestava izbranih P91-zvarov (EDS-analiza na ogljikovih ekstrakcijskih replikah)
Sample &
treatment
Accelerated creep test
conditions
Phase composition / carbides
M3C M23C6 M7C3 M2X M6C MX Laves etc.
5s – AW
MMA, P91 Initial state  x x x
STCT – 620 °C/130/381h x x x
ACT – 600 °C/26.6 ks x x x
8s – HT
MMA, P91
STCT – 620 °C
/130/31h x x
ACT – 600 °C/24.5 ks x (x) x x L
5s & 8s Thermocalc phase contentat 500–620 °C X X X (L)
10s – AW
MMA, P91
Initial state x x x
ACT – 600 °C/82.6 ks x  x x (Z)
10s Thermocalc phase contentat 500–620 °C X X X
sample, where refinement of ferrite grains dominated
certainly resulting from localized substantial strain. This
grain refinement in the neck portion of the STCT
samples was also characteristic of samples 6s and 8s.
As in this research the long-term creep tests were not
carried out, a comparison is made here with the P91
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
22 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Table 5: Chemical compositions of reference P91 weld
Tabela 5: Kemi~na sestava referen~nega zvara P91
Sample
Element, w/%
C Mn Si Cr Ni Mo V Ti Nb N
KA 1344/3/ weld 0.10 0.42 0.28 8.01 0.14 0.92 0.21 0.00 0.060 0.045
Table 6: Phase compositions of reference P91 weld
Tabela 6: Fazna sestava referen~nega zvara P91
Sample Creep conditions
Phase composition / carbides
M3C M23C6 M7C3 M2X M6C MX Laves
KA1344/3 weld Exploited  600 °C/
170 bar/9 years x x L
KA1344/3 Thermocalc phase
content at 550−640 °C X X L
Figure 35: Recrystallised ferrite grains with coagulated M23C6
carbides; sample 1344/3
Slika 35: Rekristalizirano zrno ferita s koaguliranimi karbidi, vzorec
1344/3
Figure 33: Different size carbides in crept for 9 years P91 weld metal;
sample 1344/3
Slika 33: Karbidi razli~ne velikosti v zvaru jekla P91, ki je bil izpo-
stavljen lezenju 9 let, vzorec 1344/3
Figure 32: Subgrains and few fine carbides in gauge portion of 8s
sample after STCT
Slika 32: Podzrna in drobni karbidi v merilni dol`ini vzorca 8s po
STCT
Figure 34: Elongated grains and subgrains of ferrite in crept sample
1344/3
Slika 34: Podolgovata zrna in podzrna ferita v vzorcu 1344/3, ki je
bilo izpostavljen lezenju
grade weld metal of a power plant component, marked
KA 1344/3, which failed after exposure to creep at  600
°C/170 bar/9 years. The chemical composition of this
reference weld is given in the Table 5, while the results
of the phase identification are summarized in the
Table 6.
Microstructure of this reference weld after 9 years
exposure to creep, consisted of post-martensitic / ferritic
matrix with precipitated phases of various sizes, some
very large (Figure 33). A closer look into this micro-
structure by means of TEM on thin foil specimens,
revealed the presence of numerous fine grains /
subgrains retaining the post-martensitic orientation
(Figure 34), as well as recrystallised equiaxial ferrite
grains (Figure 35), with usually random distribution of
carbides.
In numerous fine grains and subgrains some very fine
MX precipitates were found interacting with dislocations
and low-angle boundaries (Figure 36). On former
austenite grain boundaries and in the intersecting regions
of former martensite laths, colonies of large carbides
appeared, consisting of mainly M23C6 carbides and with
a few MX precipitates (Figure 37). Other medium-size
MX precipitates were found in the separate colonies,
interacting with ferrite grain boundaries (Figure 38).
The large M23C6 carbides often consisted of several
crystallites separated by stacking faults and low angle
boundaries (Figures 39–40). In several cases large M6C
carbides were identified, usually growing on the M23C6
carbides (Figures 41 and 42). Numerous of the large
(>500 nm) M23C6 carbides possessed a substructure of
"acicular type", which could be identified as a sandwich
of stacking faults and/or very-thin (<10 nm) plate-like
crystals (Figures 43 and 44). SAD patterns from such
substructure exhibit double / multiple diffraction effects
as well as reveal details of relative rotations between the
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 23
Figure 37: Agglomerates of M23C6 carbides in recrystallised matrix
of sample 1344/3
Slika 37: Agmomerati M23C6-karbidov v rekristalizirani matici v
vzorcu 1344/3
Figure 38: Medium size MX and M23C6 precipitates in sample 1344/3
Slika 38: Srednjeveliki izlo~ki MX in M23C6 v vzorcu 1344/3
Figure 36: Fine MX precipitates interacting with dislocations in
sample 1344/3
Slika 36: Drobni MX-izlo~ki, ki reagirajo z dislokacijami v vzorcu
1344/3
Figure 39: Large M23C6 carbide in recrystallised ferrite; sample
1344/3
Slika 39: Velik izlo~ek karbida M23C6 v rekristsliranem feritu, vzorec
1344/3
platelets. In some cases Laves phase Fe2Mo was
identified as appearing in the sandwich structure.
In general, in the reference material 1344/3 the
microstructure was inhomogeneous, which is typical for
the weld metal. Accordingly, during creep it transformed
non-uniformly, showing large precipitates next to
completely recrystallised grains and some fine particles,
mainly MX, interacting with dislocation and low angle
boundaries in the regions where post-martensitic grains
were retained.
As to the finely dispersed MX phase, which should
support the creep strength, it was rather seldomly
observed; more often the coagulated MX carbo-nitrides
were present. It was recently suggested that the
vanishing of the MX phase and decay of the creep
strength might be due to formation of the Z-phase 8.
Although the observed features and chemical compo-
sition of the 1344/3 weld metal might have suggested
that the Z-phase should appear in it, this phase was not
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
24 Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25
Figure 42: Diffraction from carbides of Figure 41, with double-
diffraction effects
Slika 42: Difrakcija karbidov na sliki 41, dvojna difrakcijska slika
Figure 44: Diffraction from Figure 43, showing double-diffraction
effects
Slika 44: Difrakcija s slike 43 z dvojnim difrakcijskim odsevom
Figure 41: Large thin plate-like M6C carbide overlapping with M23C6
carbide
Slika 41: Velik plo{~at izlo~ek karbida M6C, ki se prekriva s kar-
bidom M23C6
Figure 40: Dark field image of Figure 39: visible substructure /
defects in the carbide
Slika 40: Karbid s slike 39 v temnem polju. Vidne so podstruktura in
napake v karbidu
Figure 43: M23C6 carbides with stacking faults and fine plate-like
substructure
Slika 43: Karbidi M23C6 z napakami zloga in fino plo{~ato podstruk-
turo
found in TEM/EDS investigation neither on carbon
extraction replicas nor on thin foils.
From the above graphs (Figure 45) it is evident that
when the Z-phase is present then the MX carbo-nitride
disappears. Therefore the main question to be solved in
further research is what caused the substantial decrease
of the MX amount in the weld 1344/3, which had to be
decommissioned prematurely due to the decrease of its
high temperature strength. The lack of the Z-phase there
might be explained by the failure occurring much before
the phase equilibrium is reached. But this also confirms
that the Z-phase was not here involved in the dissolution
of the MX and decay of the properties. On the contrary,
the Z-phase (or similar compound) appeared after ACT
in the optimized P91 test weld metal 10s. However in
this 10s material up to the failure of the ACT sample a
large amount of the MX phase was still present.
6 CONCLUSIONS
1. The accelerated creep test on Gleeble (ACT)
transforms microstructure of creep resisting weld metal
in less than 50 hours similarly to that calculated by
Thermocalc for phase equilibrium conditions.
2. The ACT gives results comparable with short-term
creep tests (STCT), nevertheless microstructural
investigations are helpful to explain differences if these
occur.
7 REFERENCES
1 Kidin I. N. Fizicheskie osnovy elektrotermicheskei obrabotki
metallov i splavov; Izd. Metallurgia, Moscow 1969 (in Russian)
2 Mandziej S., Marciniak J. TEM investigations on in-situ transfor-
mation of M23C6 carbides; Proc XIth Int Congr on Electron
Microscopy, Kyoto 1986, 1271–72
3 Mandziej S. T. Low-energy dislocations and ductility of ferritic
steels; Materials Science & Engineering A, 164 (1993), 275–280
4 Mandziej S. T., Výrostková A. Accelerated creep testing (ACT)
procedure for weld metals using Gleeble systems, Gleeble Appli-
cation Note, DSI, Poestenkill NY, USA, 2001
5 Mandziej S. T., Výrostková A. Accelerated creep testing of P91 weld
metals, IIW Doc II-1461–02.
6 Pepe J. J., Gonyea D. C. Constant displacement rate testing at
elevated temperatures in: Fossil Power Plant Rehabilitation, ASM
International, Proc Int Conf, Cincinnati OH, Feb.1989, 39–45
7 Woodhead J. H., Quarrell A. G. Role of carbides in low-alloy creep
resisting steels; Journal of Iron and Steel Institute, (1965) June,
605–620
8 Danielsen H. K. and Hald J. Behaviour of Z-phase in 9–12 % Cr
steels, Energy Materials, 1 (2006) 1, 49–57
S.T. MANDZIEJ ET AL.: ACCELERATED CREEP TESTING OF NEW CREEP RESISTING WELD METALS
Materiali in tehnologije / Materials and technology 42 (2008) 1, 13–25 25
Figure 45: Thermocalc graphs for the reference P91 weld metal, for
equilibrium state, without and with Z-phase (ZP); symbols: 23 =
M23C6, LP = Laves phase
Slika 45: Thermocalc diagrami za referen~ni var P91 za ravnote`no
stanje brez in z Z-fazo; ozna~be: 23 = M23C6, LP = Lavesova faza

@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
ZVEZA MED ANALIZNIMI REZULTATI –
KARBONATNA BOMBA IN TERMI^NA ANALIZA
CONNECTION BETWEEN ANALYSIS RESULTS – CARBONATE
BOMB AND THERMAL ANALYSIS
@eljko Poga~nik
Salonit Anhovo, d. d., Vojkova 1, 5210 Deskle, Slovenija
zeljko.pogacnik salonit.si
Prejem rokopisa – received: 2007-09-27; sprejem za objavo – accepted for publication: 2007-10-24
Pri spremljanju reaktivnosti karbonatnih kamnin je treba upo{tevati njihove geolo{ke danosti, kot sta struktura in tekstura.
Obstaja ve~ na~inov dolo~anja masnega dele`a karbonata v kamnini; cementna industrija in proizvodnja apna uporabljata
najve~krat TGA-DTA in XRF analizo (zadnja ne poda absolutne vrednosti), medtem pa naftna industrija prisega na t. i. hitro
analizo, t. j. uporabo karbonatne bombe (Müller and Gastner, 1971).
V prispevku `elim predstaviti primerljivost med analiznima postopkoma, ki temeljita na kislinskem razklopu ter procesu
termi~nega razpada razli~nih sedimentnih kamnin. Analizi nista pomembni le za dolo~itev masnega dele`a karbonatov,
omogo~ata tudi dopolnjeno sledljivost porabe toplotne energije pri proizvodnji OPC-klinkerja. V ta namen so bili izbrani vzorci
fli{nih kamnin, ki nastopajo na podro~ju Anhovega. Korelacija rezultatov med TGA-DTA, XRF in XRD analizo ter kislinskim
razklopom karbonatne bombe je pokazala, da zna{a za dolo~itev Mg karbonatov s termi~no analizo in analizo karbonatne
bombe, faktor korelacije 0, 97.
Klju~ne besede: analizne tehnike, karbonatna bomba, termi~na analiza, mineralogija
During the following of carbonate rocks reactivity the geological characteristics like texture and structure must be consider.
There are several analyses method in the cement and lime industry, to determinate an amount of CaCO3 during the thermal
decomposition of rocks. The most uses the TGA-DTA and XRF and analysis (where the last one cannot give an absolute value
of carbonate) in the meantime the petroleum industry swear to carbonate bomb analysis (Müller and Gastner, 1971).
In this work I wish to present the comparison results between two analytical methods; acid decomposition and process of
thermal decomposition of carbonate minerals. These two analyses are not important only for carbonate determination, during
OPC clinker and lime production connected to dissociation of mineral raw material. At the same time they make possible
control the amount of consumed heat energy in the process per unit of product. For this purpose We analysed some samples of
flyschoid rocks from Anhovo area. The determination results of Mg carbonates with TGA-DTA, XRF, XRD analysis and results
of acid decomposition of carbonates by carbonate bomb technique demonstrate a correlation factor of 0,97.
Key words: analaysis procedure, carbonate bomb, thermal analysis, mineralogy
1 UVOD
Za izdelavo portlandskega cementa OPC se uporablja
me{anica karbonatnih in silikatnih kamnin v ustreznem
masnem razmerju. Pri izbiri pravilnega razmerja ima
klju~en pomen oksidna sestava posameznih kamnin ali v
kompleksnej{em sistemu, kot v primeru cementarne
Salonit Anhovo, d. d., litolo{kih komponent (slika 1 in
tabela 1).
Nekatere komponente v tehnolo{kem procesu lahko
nadomestimo tudi s sekundarnimi mineralnimi suro-
vinami 2, s ~imer zni`amo porabo toplotne energije ter
emisije toplogrednih plinov (TPG). Portland 3 navaja, da
je cementna industrija odgovorna za masni dele` od 8 do
12 % emisij TPG na letnem nivoju, upo{tevajo~ izgo-
Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32 27
UDK 55:666.942 ISSN 1580-2949
Izvirni znanstveni ~lanek/Original scientific article MTAEC9, 42(1)27(2008)
Tabela 1: Razlaga poimenovanja litolo{kih komponent 1, za katere velja, da delci med seboj niso bili povezani med procesom sedimentacije
(AD, AS in AL – notranja nomenklatura komponent mineralne surovine)
Table 1: Deffinition of the lithological component 1. AD, AS and AL are internal nomenclature of lithological units
MUDSTONE Sestavljajo karbonatni muljevec in dele` zrn, ki ne presega 10 mas.% premera med 0,03 mm in
2 mm • Mud supported matrix, with les than 10 % of grains which diameter ranged between
0,03 mm and 2 mm
AD
WACKESTONE Sestavljajo prete`no karbonatni muljevec kot osnova ter 10 mas.% do 50 mas.% zrn (zrna
"lebdijo" v muljasti osnovi) • Matrix consists mainly of mud (particles of clay and fine silt) with
more than 10 % of grains larger than 2 mm.
AD
PACKSTONE Ima `e zrnato podporo oz. zrna se med seboj dotikajo. Karbonatni muljevec nastopa le v
medzrnskem prostoru • Grain supported with mud particles of clay and fine silt size. Grains are
larger than 2 mm.
AS
GRAINSTONE Ne vsebuje karbonatnega muljevca, nastopajo le zrna v medprostoru, iz katerega se izlo~a
karbonatni cement • Grain supported without lime mud. AL
revanje vseh fosilnih goriv, kjer je poglavitna emisija
CO2. Letno zna{a ta dele` pribli`no 1 Gt od vseh 13 Gt
zavedenih TPG-emisij.
Namen te raziskave je potrditi domnevo, da je ve~ji
del MgO vezan v nekarbonatnih mineralih. V termo-
dinami~nem pogledu je najzanesljivej{i a zamuden
vpogled pridobljen s termi~nimi metodami. Termi~ni
razpad karbonatnega minerala, ki sestavlja sedimentno
kamnino je odvisen od njegove oblike in elementne
sestave, ki sta odsev njegovega nastanka. Stopnja
dekarbonatizacije karbonatnih sedimentnih kamnin , je
definirana kot  = 1 – (CO2(CaCO3))/(CO2(celotni)) v tem-
peraturnem intervalu od 400 °C do 1000 °C 4. Funkcija
dekarbonatizacije podaja reakcijo v trdi fazi, ki je
pomembna pri {tudiji procesa kalcinacije v toplotnem
izmenjevalcu 5. Raziskave mnogih avtorjev, ki jih
Bucchi navaja v svojem prispevku 4, ka`ejo linearno
~asovno odvisnost med dekarbonatizacijo zrna in njego-
vim premerom. V cementni industriji se za napoved
masnega dele`a karbonata uporablja rentgenska
fluorescen~na analiza – XRF, ki pa koli~ino analita
podaja na osnovi prera~unanega dele`a CaCO3, v kar je
v{teta tudi sistemati~na analitska napaka, vezana na
koli~ino CaO in MgO, ki nista vezana na karbonat.
Na~in dolo~itve dele`a karbonata v betonih je na
podoben na~in uporabila tudi Géraldine s sodelavci 6.
Slednja je komplementarno zdru`ila termogravimetri~no
(TGA) in klasi~no kemijsko analizo.
2 IZHODI[^A, METODE DELA IN IZBOR
VZORCEV
Rezultati predhodnih raziskav 2 ka`ejo, da je dele` p.
CaO (prostega apna) odvisen predvsem od genetskih
lastnosti kamnin. Visokokarbonatne litolo{ke kompo-
nente so kamnine tipa packstone – grainstone toplotno
manj potratne (v pogledu procesa dekarbonatizacije),
bolj reaktivne pa so kamnine strukturnega tipa mudstone
– wackestone 7. Geneza teh je vezana na globokomorske
gravitacijske procese (slika 1).
Za dolo~itev dejanskega dele`a Ca in drugih
Mg-karbonatov, s tem pa emisije TPG, so bile izvedene
rentgenska difrakcijska-XRD in flourescen~na-XRF-
analiza, termi~na – TGA-DTA- in primerjalna analiza s
@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
28 Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32
Slika 1: Idealiziran presek litolo{kih komponent ciklotema v kamnolomu Rode` s pripadajo~o geologijo (panoramska fotografija) in zgodovino
nastanka
Figure 1: An idealised presentation of lithological component from Rode` quarry (upper panoramic photo) with its corresponding sedimento-
logical process and lithology of cyclothem
karbonatno bombo – KB. Statisti~no opredeljivi vzorci
mineralne surovine 2 so bili zmleti in homogenizirani v
laboratorijskem mlinu Herzog, katerega kapaciteta zna{a
100 g merjenca. S takim na~inom priprave vzorca je bil
zagotovljen minimalni vpliv velikosti delcev 8,9,10 na
analitski postopek. Vzorce smo izpostavili intervalu
mletja do 30 s; interval je bil opredeljen na osnovi
masnega dele`a karbonata.
Snemalni pogoji na rentgenski cevi in{trumenta ARL
8480 za XRF-analizo so bili postavljeni na napetost 30
kV in tok 80 mA, snemalni ~as je bil 40 s. Difraktogrami
merjencev so bili posneti na Kemijskem in{titutu in
Institutu "Jo`ef Stefan" v Ljubljani (difraktometer
Siemens D-5000). Uporabljena je bila refleksijska teh-
nika in CuKα `arek v obmo~ju od 2° do 74°, po koraku
2θ 0,036°, in integracijskim ~asom 2 × 5 s, z diver-
gen~no in protisipno re`o, nastavljeno na 20 mm
osvetlitve vzorca, ter sprejemno re`o 0,6 mm 11.
Difraktogrami klinkerjev so bili posneti z refleksijsko
tehniko in CuKα-radiacijo v obmo~ju 2θ od 10° do 65°,
s korakom 2θ 0,03°, integracijskim ~asom 2 × 3 s, z
divergen~no in protisipno re`o, nastavljeno na 20 mm
osvetlitve vzorca, ter sprejemno re`o 0,6 mm 12.
Merjenci so bili analizirani z ra~unalni{kim programom
Crystallographica Search Match 11.
Kislinski razklop vzorcev je bil izveden na osnovi
protokola analize s "karbonatno bombo-KB", ki sta ga
objavila Müller in Gastner 13. Pri termi~ni preiskavi smo
uporabili analizator STD 2960 Simultaneus TGA-DTA.
Meritve smo opravljali z ustaljenim protokolom pri
laboratorijski temperaturi 1200 °C in s temperaturnim
korakom segrevanja 10 °C/min. Ta omogo~a spremljavo
temperaturnega razpada mineralnih struktur in s tem
dolo~itev temperaturnih parametrov dekarbonatizacije na
osnovi literaturnih podatkov 8,14,10. Za kalibracijo in{tru-
mentov smo uporabili standard kalcita 1.
Reaktivnost merjenca je bila dolo~ena z meritvijo
koli~in, ki so neposredno ali posredno vezane na
strukturo mineralov 4. V tehnolo{kem procesu je tempe-
ratura pri~etka (onset temperatura) reakcije najpomemb-
nej{i faktor, saj pove, pri kateri temperaturi se pri~ne
masa in zgradba minerala spreminjati zaradi dovedene
toplote. Maksimalna temperatura – Tmax ali temperatura
pri kateri je proces dosegel svoj vrh, je bila dolo~ena s
prvim odvodom gravimetri~ne krivulje, temperatura
pri~etka reakcije, kot jo to poimenuje Yen s sodelavci 15
in jo `e Ramachandran naka`e v svojem delu 14, pa na
osnovi drugega odvoda.
3 REZULTATI
V tabeli 2 so navedeni masni dele`i oksidov in
mineralov, pridobljeni s posameznimi preiskovalnimi
metodami.
Primerjavo kvantitativne emisije CO2 glede na posa-
mezno preiskovalno metodo ponazarja slika 2.
Iz rezultatov, ki jih ponazarjati tabeli 2 in 3 lahko
izra~unamo dele` MgO, ki ni vezan v kamnini v obliki
karbonata.
@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32 29
Tabela 2: Rezultati uporabljenih preiskovalnih metod (XRF, XRD, TG-DTA analize in karbonatne bombe K.B.)
Table 2: Results of XRF, XRD, TG – DTA and "carbonate bomb" analysis
XRF XRD K. B. DTA-TG
(w/%)
MERJENEC
SAMPLE
CaO MgO CO22 CaCO3
Kalcit
Calcite
CO**2
Enstatit
Enstatite
Muskovit
Muscovite
Klorit
Chlorite
CaCO3 CO2 CaCO3
AD 01V01 42,97 1,36 33,76 76,70 85,50 37,62 3,27 1,21 3,06 81,09 32,84 74,64
AS 01V01 46,04 1,26 36,17 82,18 88,12 38,77 3,09 0,69 2,28 84,58 36,59 83,16
AL 01V04 47,92 1,10 37,65 85,54 90,86 39,98 1,72 0,42 1,31 84,58 36,88 83,82
AS 02V03 47,44 1,11 37,26 84,67 90,44 39,79 2,15 0,62 1,32 85,57 36,48 82,91
AL 02V03 47,90 1,09 37,63 85,49 89,97 39,59 2,47 0,27 1,45 89,05 38,02 86,41
AS 04V03 46,51 1,16 36,54 83,02 89,78 39,50 2,22 0,45 1,81 83,58 36,25 82,39
AS 05V04 45,68 1,23 35,88 81,54 88,74 39,05 2,04 0,83 1,86 81,59 35,97 81,75
AD 05V04 44,27 1,30 34,78 79,02 87,03 38,29 2,54 0,96 2,84 80,60 33,51 76,16
AS 08V03 46,03 1,18 36,16 82,16 89,54 39,40 1,62 0,81 1,6 81,59 35,03 79,61
AL 09V02 50,95 0,91 40,03 90,95 95,10 41,84 1,07 0,21 0,56 91,54 39,67 90,16
AD 09V03 44,04 1,23 34,60 78,61 87,37 38,44 1,66 1,25 2,6 79,60 33,52 76,18
AS 10V03 47,28 1,11 37,15 84,40 90,40 39,78 1,78 0,7 1,49 85,57 36,66 83,32
AL 12V03 50,68 0,93 39,82 90,47 94,70 41,67 1,05 0,19 0,61 90,05 39,37 89,48
AD 13V02 45,06 1,21 35,40 80,43 88,07 38,75 2,03 0,93 2,03 81,09 34,43 78,25
1 CARLO ERBA REAGENTI, calcium carbonate ACS-ISO-for analysis, CaCO3 minimum essay, 99,5 % – komplexometric.
2 ** Prera~unan na dele` CaCO3 (Calculated on w/% of CaCO3).
4 DISKUSIJA
Ob dolo~itvi dele`a karbonata s termi~no analizo je
treba upo{tevati obliko in polo`aj endotermnega vrha 8,
ki je odvisen od stati~nega pritiska CO2 na povr{ini
vzorca med procesom dekarbonatizacije. Razlika v
procesu dekarbonatizacije se poka`e v tipu lokalne
atmosfere, kjer poteka reakcija: ~e poteka proces v zraku
– laboratorijski atmosferi, je temperatura razpada ni`ja
kot pa v atmosferi CO2. Polo`aj endotermnega vrha je
izpostavljen tudi nekaterim eksotermnim reakcijam
kovinskih oksidov, ki nastanejo ob razpadu karbonata.
Na obliko vrha imajo vpliv tudi soli ali Fe v magnezitu
(nastanek breunnerita ((Mg,Fe) CO3) z atomskim raz-
merjem Mg : Fe od 90 : 10 do 70 : 30 8), ki navidezno
povi{a maksimalno temperaturo, ki jo vrh krivulje
dose`e.
Iz pridobljenih rezultatov ugotavljamo, da CaO in
MgO nista v celoti vezana v karbonatne minerale.
TGA-analiza standarda kalcita ka`e na 0,469-odstotno
nenatan~nost analize in kvantitativne dolo~itve emisije
CO2. Ob primerjavi natan~nosti TGA-analize merjencev
z XRF-analizo je v zadnji opaziti pribitek CaO, vezanega
na karbonat (kljub upo{tevani natan~nosti metode pri
dolo~itvi CaO), kar ponazarja slika 3.
Razliko v odmiku rezultatov XRF–TGA analize gre
pripisati uporabi standarda pri dolo~itvi CaO z XRF-
metodo. Prav tako ima tudi TGA-metoda svoje pomanj-
kljivosti pri dolo~itvi karakteristi~ne temperature
razpada karbonata. Treba je namre~ upo{tevati prekri-
vanje endotermnih vrhov pri analizi polimineralnih
merjencev razli~ne kristalini~nosti, lastnosti, ki so
vezane na nastanek karbonatnih mineralov sedimentnih
kamnin (vpliv oblike osnovne celice na temperaturo
razpada; tabela 5).
Tako sta dele`a CaO in MgO (tabela 4) dolo~ena z
XRF-metodo lahko vezana tudi na nekatere minerale
glin (tabela 6).
Dele` MgO, ki ni vezan na enstatit in klorit (tabela
4), bi lahko bil vezan v karbonatih le:
1. ~e bi bil s TGA- in XRD-analizo dolo~en dolomit
CaMg(CO3)2; kot monomineral vsebuje masni dele`
MgO do 21,86 %; dolomit je pokazatelj diagenetskih
@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
30 Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32
Tabela 3: Oksidna sestava mineralov, ki nastopajo v vezivu sedimentne kamnine-merjenca (http://webmineral.com/), dolo~ena z XRD-analizo
Table 3: Total oxide composition (http://webmineral.com/) of minerals (detected by XRD analysis)
Mineral Emperi~na formulaEmpirical formula
Oksidna sestava (w/%)
Total oxide composition
CaO Al2O3 SiO2 Fe2O3 K2O Na2O MgO (OH)n/H2O
Klorit
Chlorite Na0,5Al4Mg2Si7AlO18(OH)12·5(H2O) 26,18 43,19 1,59 8,28 20,35
Enstatit
Enstatite Mg2Si2O6 59,85 40,15
Muskovit
Muscovite KAl3Si3O10(OH)1,8F0,2 38,36 45,21 11,81 4,07
0,00
5,00
10,00
15,00
20,00
25,00
30,00
35,00
40,00
45,00
A
D
-0
1V
01
A
S-
01
V
01
A
L-
01
V0
4
A
S-
02
V
03
A
L-
02
V0
3
A
S
-0
4V
04
A
S
-0
5V
04
A
D
-0
5V
04
A
S-
08
V
03
A
L-
09
V0
2
A
D
-0
9V
03
A
S-
10
V
03
A
L-
12
V0
3
A
D
-1
3V
02
w
/
%
CO2 (XRF) CO2 (TGA-DTA) CO2 (KB) CO2 (XRD)
Slika 2: Masni dele` CO2 glede na rezultate preiskovalnih metod
(TGA-DTA in KB izmerjena, XRF in XRD prera~unana na masni
dele` CaO oz. CaCO3)
Figure 2: Amount of CO2, concording the results of mineralogical
(XRD), chemical (XRF and "carbonate bomb") and thermical
(TG-DTA) analysis
Tabela 4: Masni dele` MgO, ki ni vezan na karbonat
Table 4: Amaount of MgO oxide which is not a part of carbonate
minerals
MERJENEC
SAMPLE
(w/%)
XRF XRD ∆3
CaO MgO4 CO2 EnstatitEnstatite
Klorit
Chlorite MgO
AD 01V01 42,97 1,36 33,76 3,27 3,06 -0,20
AS 01V01 46,04 1,26 36,17 3,09 2,28 -0,17
AL 01V04 47,92 1,10 37,65 1,72 1,31 0,30
AS 02V03 47,44 1,11 37,26 2,15 1,32 0,14
AL 02V03 47,90 1,09 37,63 2,47 1,45 -0,03
AS 04V03 46,51 1,16 36,54 2,22 1,81 0,12
AS 05V04 45,68 1,23 35,88 2,04 1,86 0,25
AD 05V04 44,27 1,30 34,78 2,54 2,84 0,05
AS 08V03 46,03 1,18 36,16 1,62 1,6 0,40
AL 09V02 50,95 0,91 40,03 1,07 0,56 0,43
AD 09V03 44,04 1,23 34,60 1,66 2,6 0,34
AS 10V03 47,28 1,11 37,15 1,78 1,49 0,28
AL 12V03 50,68 0,93 39,82 1,05 0,61 0,45
AD 13V02 45,06 1,21 35,40 2,03 2,03 0,23
3 Dele` prostega oksida, ki bi lahko bil vezan v obliki karbonata (Atributed to CaCO3).
4 Napaka meritve dele`a MgO z XRF-analizo zna{a 0,09 w/% (The error of MgO detection by XRF analysis is 0,09 w/%).
procesov, ki na obmo~ju Anhovega ob nastanku
kamnin niso bili prisotni, velja pa omeniti, da lahko
nastopa kot predstavnik starej{ih erodiranih in
prelo`enih kamnin v sedimentacijsko okolje, kar bi v
primeru merjencev pomenilo kvantitativno od 0,22 %
do 2,08 % minerala; z opravljeno XRD-analizo, s
katero lahko lo~imo minerale z masnim dele`em nad
1,0 %, ni bila dokazana prisotnost dolomita, prav
tako ne za merjenec z 2,08 %;
2. ~e bi bil Mg vezan v obliko aragonita, ki je pri
normalnih razmerah nestabilna-metastabilna mine-
ralna faza in ga med procesom diageneze najve~krat
zamenja kalcit. V kamninah pelagi~ne sedimentcije,
ki so pokazatelj globalnih sprememb, so nosilci
karbonata mikro- in kriptokristalne strukture 16,17 oz.
@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32 31
Tabela 5: Delitev karbonatnih mineralov glede na obliko osnovne celice in karakteristi~no temperaturo razpada 8
Table 5: Distribution of carbonates according their crystallographic form and characteristic temperature 8
SKUPINA KALCITA
CACLITIC GROUP
SKUPINA DOLOMITA
DOLOMITIC GROUP
SKUPINA ARAGONITA
ARAGONITIC GROUP
Oblika osnovne celice heksagonalna
Space group
R3c
Oblika osnovne celice heksagonalna
Space group
R3
Oblika osnovne celiceortorombi~na
Space group
PmCn
Mineral
emp. ena~ba
Empirical
formula
Tkarakteristi~na
T/°C
mineral
emp. ena~ba
Empirical
formula
Tkarakteristi~na
T/°C mineral
emp. ena~ba
Empirical
formula
Tkarakteristi~na
T/°C
Kalcit
Calcite
CaCO3 878,65
Dolomit
Dolomite
CaMg(CO3)2 745
Aragonit
Aragonite
CaCO3 920
Magnezit
Magnesite
MgCO3 670
Ankerit
Ankerite
Ca(Mg,Fe)
(CO3)2 730
Viterit
Witherite
BaCO3 830
Rodohrozit
Rhodochrosite
MnCO3 610
Kutnohorit
Kutnohorite
CaMn(CO3)2
Stroncianit
Strontianite
SrCO3 930
Tabela 6: Minerali glin, na katere je vezan dele` CaO (http://webmineral.com/)
Table 6: Amount of CaO and MgO within some other minerals (http://webmineral.com/) which can be found in carbonate rocks
Mineral Emperi~na formulaEmpirical formula
Oksidna sestava (w/%)
Total oxide composition (w/%)
CaO Al2O3 SiO2 Fe2O3 K2O Na2O MgO (OH)n/H2O
Odinit
Odinite
Fe3+0,7Mg0,7Al0,5Fe2+0,3Ti0,1Mn2+0,1
Si1,8Al0,2O5(OH)4
11,87 35,98 18,596 9,397 11,99
Montmorillonit
Montmorillonite
Na0,2Ca0,1Al2Si4O10(OH)2(H2O)10 1,02 18,57 43,77 1,13 36,09
Volkonskoit
Volkonskoite
Ca0,1Mg0,1Cr1,2Mg0,8Fe3+0,3Si3,5Al0,5O10
(OH)2·3,6(H2O)
1,18 5,36 44,21 5,04 7,63 17,42
Swinefordite8 LiCa0,5Na0,1Al1,5Mg0,5Si3O10(OH)1,5F0,5·
2(H2O)
7,06 19,25 45,38 0,78 5,07 12,47
Hidrobiotit
Hydrobiotite
Mg2,3Fe3+0,6K0,3Ca0,1Si2,8Al1,2O10(OH)1,8F0,2·
3(H2O)
1,21 13,17 36,21 10,31 3,04 19,95 15,12
Illite K0,6(H3O)0,4Al1,3Mg0,3Fe2+0,1Si3,5O10(OH)2·
(H2O)
17,02 54,01 1,85# 7,26 3,11 12,03
Vermikulit
Vermiculite
Mg1,8Fe2+0,9Al4,3SiO10(OH)2·4(H2O) 43,48 11,92 12,82# 14,39 17,87
Bramalit
Brammallite
Na0,66(H3O)0,33AlMg0,6Fe2+0,1Si3AlO10
(OH)2·(H2O) 26,02 46,00 1,83
# 5,22 6,17 11,47
0,00
0,10
0,20
0,30
0,40
0,50
0,60
0,70
0,80
0,90
1,00
A
D
-0
1V
01
A
S-
01
V
01
A
L-
01
V
04
A
S-
02
V
03
A
L-
02
V0
3
A
S-
04
V
04
A
S-
05
V
04
A
D
-0
5V
04
A
S-
08
V
03
A
L-
09
V0
2
A
D
-0
9V
03
A
S-
10
V
03
A
L-
1
2V
03
A
D
-1
3V
02
XRFsp.meja XRFzg.meja K.B. TGA XRF
Slika 3: TGA-dolo~itev CaCO3 v primerjavi z normiranimi rezultati
(na ordinatni osi) XRF-analize
Figure 3: Results of TGA analysis compared with normalised XRF
results
5 Povpre~na temperatura izra~unana na osnovi Tkarakteristi~na, ki jo podaja Mackenzie
8, (Averaged temperature after Mackenzie 8).
6 7,17 FeO
7 2,36 MnO
8 LiO2 = 3,76, F = 2,39
fosilni planktonski organizmi in mikroorganizmi.
Berner 18 navaja, da je najdba aragonita v globoko-
morskih sedimentih lahko posledica dolgega prenosa
materiala iz plitvovodnega okolja ali pa posledica
sedimentacije skeletov odmrlih planktonskih orga-
nizmov. Avtor prav tako navaja, da je karbonat z
visoko vsebnostjo Mg izpostavljen raztapljanju
zaradi vpliva pCO2, raztopljenega v morju, kar
potrjuje ob dodatnem vplivu kemizma morja (masno
razmerje Mg/Ca) tudi Tucker 19. V splo{nem velja, da
so vse karbonatne kamnine ne glede na njihovo
prvotno mineralno sestavo, med diagenezo izpostav-
ljene pretvorbi v kalcit z nizkim dele`em Mg 19.
Aragonita v heterogenih sedimentnih kamninah, kot
nastopajo v okolici Anhovega in so vezane na podo-
ben nastanek, ni bilo mogo~e ugotoviti, saj je njegov
prvi endotermni vrh pri 447 °C, ki ponazarja dio-
morfni prehod aragonita v kalcit 20. Ta temperatura
sovpada z obmo~jem endotermnih reakcij mineralov
glin. Kljub temu obstaja teoreti~na mo`nost pre-
tvorbe kalcita v aragonit, vendar le pri tlakih 2,0 GPa
in v temperaturnem obmo~ju 600 °C 21, kar pa v
primeru geolo{kih danosti Anhovega ne gre pri~ako-
vati.
Termi~ni razpad karbonatnega minerala, ki sestavlja
karbonatno sedimentno kamnino, je odvisen od njegove
oblike in elementne sestave, ki so odsev njegovega
nastanka (tabela 5). Iz rezultatov preiskav je razvidno,
da poda termi~na analiza najbli`je realne vrednosti
karbonatnega dele`a v surovini, saj sledi popolni emisiji
CO2 iz karbonata, ne glede na genezo nastanka kamnine.
Vendar pa korelacijski faktor (tabela 7) potrjuje, da
obstaja sprejemljiva soodvisnost tudi z analizo
karbonatne bombe in XRD-analizo.
5 SKLEP
V prispevku je dokazana primerljivost med
TGA-DTA-analizo in analizo s t. i., karbonatno bombo
po Müller-Gastnerjevi metodi 13. Iz dobljenih rezultatov
lahko potrdimo, da masni dele` MgO (do 0,27 % od
prvotno ocenjenega na 1,13 %) ni popolnoma vezan na
karbonatne minerale.
Ob sklepu je treba dodati, da detajlna mineralo{ka
analiza veziva karbonatnih kamnin {e ni bila izvedena. V
tem pogledu bodo v prihodnje izvedene {tudije, ki bodo
ovrednotile masni dele` CaO, ki ni vezan na karbonate v
osnovni mineralni surovini za proizvodnjo OPC
klinkerja.
6 LITERATURA
1 Wright, V. P. A revised classification of limestones, Sed. Geol,. 76
(1992), 177–186
2 Poga~nik, @. Reaktivnost karbonatnih kamnin in sekundarnih
surovin-indikator okoljevarstva in trajnostnega razvoja v cementni
industriji; Doctor's degree, Ljubljana: 2006, 124 str.
3 Portland, B. Geopolymer – the future? International Cement review,
(2007), 79–82
4 Bucchi, R. Influence of the nature and preparation of raw materials
on the reactivity of raw mix, V 7th International Congress on the
Chemistry of Cement – Volume 1 Principal Reports, Pariz, 1980,
1–43
5 Fidaros, D. K., Baxevanou, C. A., Dritselis, C. D., Vlachos, N. S.
Numerical modelling of flow and transport processes in a calciner
for cement production, Powder Technology, 171 (2007) 2, 81–95
6 Géraldine, V., Mickaël, T., Gérard, P. Measurement methods of
carbonation profiles in concrete: Thermogravimetry, chemical
analysis and gammadensimetry, Cement and Concrete Research, 37
(2007), 1182–1192
7 Dunham, R. J. Classifiction of carbnate Rocks According To
Depositional Texture. V Classification of Carbonate Rocks – a
simposium, (1962), 108–122
8 Mackenzie, R. C. The differential thermal investigation of clays,
London: Mineralogical society (clay minerals group), 1957, 456 str.
9 Oates, J. A. H., Lime And Limestones Chemistry And Technology,
Production And Uses, Weinheim: Wiley-VCH, 1998, 455 str.
10 Taylor, H. F. W. Cement Chemistry, 2nd ed, Oxford : Thomas
Telford, 1997, 439 str.
11 Meden, A. Poro~ilo o kvantitativni fazni analizi `lindre; speciali-
sti~na naloga, Ljubljana, 2005, 5 str.
12 Meden, A. Poro~ilo o kvantitativni fazni analizi surovin dostavljenih
v marcu 2004; specialisti~na naloga, Ljubljana, 2004, 29 str.
13 Müller, G., Gastner, M. The "Karbonat-Bombe", a simple device for
the determination of the carbonate content in sediments, soils, and
other materials: N, Jb, Mineral, Mh, 10 (1971), 466–469
14 Ramachandran, V. S., Phil, D. Applications of differential thermal
analysis in Cement Chemistry, Chemical pulishing company, 1969,
308 str.
1 Yen, F. S., Lo, H. S., Wen, H. L., Yang, R. J. θ – to α – phase trans-
formation subsystem induced by α – Al2O3 – seeding in boehmite –
derived nano – sized alumina powders, Journal of Crystal Growth,
249 (2003), 283–293
15 Marschner, H. Relationship between carbonate grain and
non-carbonate content in carbonate sedimentary rocks, In: Recent
developments in carbonate sedimentology in Central Europe,
Spinger –Verlag, Editors: Müller, G., Frideman, G. M., New York,
1968, 55
16 Sprung, S. Effect of the Burning Process on the formation and
propreties of the clinker, V 7th International Congress on the Che-
mistry of Cement – Principal Reports, Pariz, 1 (1980), 44–60
17 Berner, R. A., Honjo, S. Pelagic Sedimentation of Aragonite: Its
Geochemical Signifiance, Sience, 11 (1981), 940–942
18 Tucker, M. E. Sedimentary petrology, 3rd ed, Blackwell Publishing
company, 2005, 262 str.
19 Faust, G. T. Differentiation of Aragonite from Calcite by Differential
Thermal Analysis, Science, 110 (1949), 402–403
20 Hacker, B. R., Kirby, S. H., Bohlen, S. R. Time and Metamorphic
Petrology: Calcite to Aragonite Experiments, Science, 258 (1992),
110–112
@. POGA^NIK: ZVEZA MED ANALIZNIMI REZULTATI – KARBONATNA BOMBA IN TERMI^NA ANALIZA
32 Materiali in tehnologije / Materials and technology 42 (2008) 1, 27–32
Tabela 7: Korelacija med rezultati, podanih s posamezno preisko-
valno metodo
Table 7: Results correlation between analysis research
K, B, TGA-DTA XRD XRF
K, B, 1
TGA-DTA 0,9796655 1
XRD 0,9385986 0,9393562 1
XRF 0,8411661 0,8372388 0,8152043 1
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
DEVELOPING AND TESTING A NEW TYPE-8K MOULD
FOR TOOL-STEEL INGOT CASTING
RAZVOJ IN PREIZKUS NOVE KOKILE VRSTE 8K ZA ULIVANJE
INGOTOV IZ ORODNEGA JEKLA
Martin Balcar1, Libor Sochor1, Rudolf @elezný1, Pavel Fila1, Ludvík Martínek1,
Libor Kraus2, Du{an Ke{ner2, Ji í Ba`an3
1@AS, a. s., @ár nad Sázavou, Czech Republic
2COMTES FHT, s .r .o. Plze, Czech Republic
3V[B TU – Ostrava, Czech Republic
martin.balcar zdas.cz
Prejem rokopisa – received: 2007-09-26; sprejem za objavo – accepted for publication: 2007-10-15
The traditional production of machinery and tools at @AS, a. s., and the increasing requirements placed on the final product,
particularly those related to the internal quality of forgings, requires new technological measures for the ingot’s casting and
solidification. The development and use of the new mould type 8K9.2 for tool-steel ingot casting has reduced the share of
forgings with unacceptable quality. A substantial improvement has been achieved, especially in terms of the occurrence of
defects, such as cavities and cracks in the ingot body.
Keywords: tool steel, mould, ingot casting, solidification, forging, ultrasonic testing
Tradicionalna proizvodnja strojev in orodij v podjetju @DAS, a. s., in nara{~ajo~e zahteve pri lastnostih kon~nega proizvoda, ki
so povezane z notranjo kakovostjo izkovkov, so zahtevale nove tehnolo{ke ukrepe pri litju in strjevanju ingotov. Razvoj in
uporaba nove kokile vrste 8K9.2 za litje jeklenih ingotov sta zmanj{ala dele` ingotov nesprejemljive kakovosti. Pomembno
izbolj{anje je bilo dose`eno pri pogostosti napak, razpok in votlin v notranjosti ingota.
Klju~ne besede: orodno jeklo, kokila, ulivanje ingotov, strjevanje, kovanje, ultrazvo~na preiskava
1 INTRODUCTION
The determination of the causes of the internal
defects in tool-steel ingots and forgings leads to the
implementation of optimisation measures in production
technology 1,2. The requirements for the internal quality
of forgings from large blocks, and the conditions for
forming in the company @AS, a. s., represented by the
limit of the pressing force, has focused work on the
processes of ingot casting and solidification related to
ingot-mould shape changes. An optimised form of the
ingot 8K9.2 with a mass of 8.9 t, and with a reduced
portion of axial defects defined in particular by the
Niyama criterion 3 was proposed after a numerical
simulation of the casting process and solidification of the
forging ingot 8K8.4 with a mass of 7.6 t by applying
MAGMA software.
With the use of the new shaped mould 8K9.2 for the
casting of tool steel W.Nr. 1.2344 (X40CrMoV51) in
ingots with a mass of 8.9 t, the internal quality of the
ingots was improved. Ingots with a low share of the poor
integrity in the axial part are the basic pre-requisite for
the high-quality production of large tool-steel forgings.
2 DEFECTS IN INGOTS MADE OF TOOL
STEEL
In the company @AS, a. s. the production of tools is
carried out by processing the 8K ingots with masses of
1000 kg to 11700 kg on CKV 630, CKV 1250 and
CKV 1800 presses. In 2006 the forming force for the
CKV 1800 press was increased from 18 MN to the
current value of 22.5 MN.
The traditional assortment of tool steels consists of,
e.g., the steel grades 90MnCrV8, X37CrMoV51,
X40CrMoV51, 55NiCrMoV7, X210Cr12 and special
steels for the rolls in rolling mills: 8CrMoV,
8Cr3MoSiV, 8CrMoSiV.
When producing forgings from tool steel, the
problems are connected with the low plastic properties
related to the high content of carbon in combination with
the content of the alloying elements, chromium,
molybdenum and vanadium. Sufficient forging through
the ingot and the elimination of internal defects,
particularly in the zone of the end of solidification, in the
axial part of the ingot, requires a high degree of forming
and a sufficient deformation volume.
In the case of large forgings with a diameter that
make it impossible to achieve the necessary deformation
in the axial part of the ingot by forging, the required
internal quality is obtained with an improvement of the
ingot’s internal quality.
3 VERIFICATION OF THE INGOT 8K9.2
The design of the new mould shape is based on an
analysis of numerous modifications to the mould
Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38 33
UDK 669.14.018.252:620.17 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 42(1)33(2008)
geometry ensuing from the existing polygonal ingot of
the type 8K8.4 with a mass of 7600 kg.
The main changes to the geometry were made from
the viewpoint of slenderness and the bevel of the ingot
body. The final shape was, afterwards, adapted and
verified with modelling of the ingot 8K9.2 with a mass
of 8850 kg. A comparison of the basic parameters of the
ingot 8K8.4 and the design of the new shape of the ingot
8K9.2 are given in Table 1.
Table 1: Basic parameters of the ingot 8K8.4 and the design of the
ingot 8K9.2
Tabela 1: Osnovni parametri ingota 8k8.4 in na~rt ingota 8K9.2
Type
Mass and volume
S
slen-
derness
V
taper
ingot head body H
D
°
kg kg % kg %
8K8.4 7 600 800 11% 6 800 89% 2.1 4.8
8K9.2 8 850 1 200 14% 7 650 86% 1.3 11.0
The steel X40CrMoV51 was chosen for the
verification of the quality of the ingots and the forgings.
Two heats were cast into two ingots, 8K8.4 and 8K9.2,
each. One ingot from the heat was always submitted for
analyses of the chemical composition from the viewpoint
of segregations and structural heterogeneities, while the
second ingot was processed by forming on the CKV
1800 press to produce a bar with a diameter of 350 mm
with the maximum use of material, i.e., without the
technological waste associated with the technology of
open-die forging.
3.1 Chemical heterogeneity of the ingot
For a comparison of the chemical heterogeneity of
the ingots made from the steel grade X40CrMoV51 the
chemical composition of the ingots of the types 8K8.4
and 8K9.2 was determined on a cross-section in the
longitudinal direction according to the layout shown in
Figure 1. The results of these analyses were compared to
the results of heat analyses, and on the basis of the
differences, positive and negative values of the absolute
differences of the individual elements were determined.
The diagrams in Figure 2 and Figure 3 show the
positions of the greatest differences in the carbon
concentration, i.e., in the axis along the height of the
ingot of types 8K8.4 or 8K9.2. Figure 4 shows the
changes in the concentration of sulphur in the axis along
the height of the ingot of type 8K8.4 and then along the
height of the ingot of type 8K9.2.
Figure 6 and Figure 7 show the deviations of the
molybdenum concentration from the heat analysis in the
axis along the height of the ingot type 8K8.4 or 8K9.2.
From the absolute deviations of the elements
determined from the samples of steel taken during the
casting of ingots it is obvious that with the solidification
of steel in the mould 8K9.2 a more distinct segregation
of carbon, sulphur and molybdenum occurs along the
ingot axis in comparison to the mould 8K8.4.
In Figure 8 it is possible to compare the differences
in the segregation of elements at the surface, at a
distance of ½ of the radius from the surface (middle
radius), and in the axis of the ingot 8K9.2. The higher
content of carbon at the ingot surface is probably related
to the used casting powder, which contains the mass
fraction of C up to 21 %. At the place marked as ½ of the
radius (middle radius) and also in the ingot axis, at the
foot part the carbon is being negatively segregated and at
the top part a positive segregation occurs, whereas the
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
34 Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38
Figure 2: 8K8.4 – deviations of carbon concentration from the heat
analysis (0.38 wt. % C) along the ingot axis
Slika 2: 8K8.4 – odmik pri vsebnosti ogljika od analize taline v
masnih dele`ih (w(C) = 0.38 %) vzdol` osi ingota
Figure 1: Layout of samples for analyses of the chemical composition
Slika 1: Polo`aj izreza vzorcev za kemijsko analizo
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38 35
Figure 3: 8K9.2 – deviations of carbon concentration from the heat
analysis (0.38 wt. % C) along the ingot axis
Slika 3: 8K9.2 – odmik pri vsebnosti ogljika od analize taline v
masnih dele`ih (w(C) = 0.38 %) vzdol` osi ingota
Figure 6: 8K8.4 – deviations of the molybdenum concentration from
the heat analysis (1.25 wt. % Mo) along the ingot axis
Slika 6: 8K8.4 – odmik pri vsebnosti molibdena od analize taline v
masnih dele`ih (w(Mo) = 1.25 %) vzdol` osi ingota
Figure 4: 8K8.4 – deviations of sulphur concentration from the heat
analysis (0.0060 wt. % S) along the ingot axis
Slika 4: 8K8.4 – odmik pri vsebnosti `vepla od analize taline v
masnih dele`ih (w(S) = 0.006 %) vzdol` osi ingota
Figure 5: 8K9.2 – deviations of sulphur concentration from the heat
analysis (0.0035 wt. % S) along the ingot axis
Slika 5: 8K9.2 – odmik pri vsebnosti `vepla od analize taline v
masnih dele`ih (w(S) = 0.0035 %) vzdol` osi ingota
axial part of the ingot is characterised by greater
differences in the concentrations.
It is possible to verify the results of the numerical
simulation using the MAGMA software with a com-
parison of the curve in Figure 8 with the curve in Figure
9. It is evident that in real conditions more distinct
differences occur from the viewpoint of the absolute
change of the carbon concentration. The numerical
simulation also does not take into account the possibility
of carburisation of the melt by carbon pick up from
external sources. The contribution of the numerical
simulation is apparent in the area of the determination of
the distribution of positive and negative segregations on
the ingot’s cross-section.
3.2 Structural defects of ingots
Real forgings were investigated for a practical
verification of the spread of internal defects. The
following ingots were cast from the steel X40CrMoV51:
• ingot 8K8.4 – bottom casting
• ingot 8K8.4 – top casting
• ingot 8K9.2 – bottom casting
After cutting the ingots and completing a capillary
test by wet testing, the pictures that were taken were
changed to a graphical transformation with the aim to
accentuate the ingot defects. The obtained appearance of
an internal part of the ingot in cross-section is shown in
Figure 10 to Figure 12.
The pictures confirm the influence of the casting
technology and the shape of the mould on the spread of
internal defects. A more detailed investigation of the size
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
36 Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38
Figure 9: 8K9.2 – MAGMA deviations of carbon concentrations
along the height of the ingot body
Slika 9: 8K9.2 – MAGMA odmik pri vsebnosti ogljika po vi{ini
ingota
Figure 7: 8K9.2 – deviations of the molybdenum concentration from
the heat analysis (1.25 wt. % Mo) along the ingot axis
Slika 7: 8K9.2 odmik pri vsebnosti molibdena od analize taline v
masnih dele`ih (w(Mo) = 1.25 %) vzdol` osi ingota
Figure 10: 8K8.4 bottom cast ingot – ingot defects
Slika 10: 8K8.4 ingot ulit od spodaj – napake v ingotu
Figure 8: 8K9.2 – deviations of carbon concentration from the heat
analysis (0.38 weight % C) along the height of the ingot body
Slika 8: 8K9.2 – odmik pri vsebnosti ogljika od analize taline v
masnih dele`ih (w(C) = 0.38 %) po vi{ini ingota
of individual defects shows that the ingot 8K9.2 contains
cracks that are 1 mm wide and with a length up to
10 mm, while the ingot 8K8.4 contains along the ingot
axis – regardless of the casting technology – continuous
cracks and voids with widths above 5 mm and lengths up
to several centimetres, in some cases.
3.3 Results of ultrasonically testing the forging
The main part of the evaluation of the influence of
changes of the mould shape on the production of
forgings made from tool steels were tests of the internal
quality of forgings with ultrasonics in accordance with
the standard SEP 1921. The verification of the quality of
the experimental forging from the ingot 8K9.2 of steel
X40CrMoV51 was performed according to the degrees
A/a, B/b, C/c, D/d, E/e, with the level D/d or possibly
E/e, considered as acceptable for the given steel.
On the basis of measurements and from the
comparison of the results of the statistical evaluation of
the previous production of forgings from the steel grade
X40CrMoV51, summarised in Table 2, it is evident that
a change in the mould shape brought a significant
increase in the utilisation of the ingot, as well as an
enhancement of the forging’s internal quality.
Table 2: Utilisation of ingots of steel X40CrMoV51 – ultrasonic
testing according to the SEP 1921
Tabela 2: Uporaba ingotov iz jekla x40CrMoV51 – ultrazvo~na kon-
trola po SEP 1921
Type
Mass
kg
Guaranteed
US
Ingot utilisation
kg %
8K8.4 7600
B/b 5950 78%
C/c 5450 72%
D/d 4600 61%
E/e 0 0%
8K9.2 8850 E/e 7300 82%
The test bar forged from the ingot 8K9.2 satisfied the
requirements of all the degrees of ultrasonic tests and the
length of the bar corresponded to 82 % of the utilisation
of the ingot’s volume. The remaining material in the
forging could not be used due to cracks in the ingot’s
head and footing crop end due to material flow and the
necessity of aligning the forging’s front face with
machine cutting.
The optimistic conclusions based on an investigation
of one forging were confirmed by the results obtained on
real forgings of bars and blocks from tool steels with a
mass up to 7 t and a bar diameter up to 600 mm, or the
height of a block up to 500 mm.
4 CONCLUSIONS
The verification of the results of a numerical
simulation with examinations on real ingots confirms the
influence of changes of the mould geometry on the
segregation in the ingot. The degree of chemical
heterogeneity of the ingot also depends a great deal on
the type of cast steel.
The change of the casting technology did not
decrease the quantity of defects, such as cracks and
voids, along the axial part of the ingot to an acceptable
level. Important changes in the distribution and shape of
the internal defects of the ingot were noted in ingots cast
into the new shapes of the moulds of the type 8K9.2.
On the basis of the results of the investigation it was
possible to obtain a substantial enhancement of the
internal quality of the forgings with a simultaneous
reduction in the share of unsatisfactory forgings that did
not pass the ultrasonic testing, as a consequence of the
change in geometry of the forging ingot of the type
8K9.2 with a mass of 8.9 t.
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38 37
Figure 12: 8K9.2 bottom cast ingot – ingot defects
Slika 12: 8k9.2 ingot ulit od spodaj – napake v ingotu
Figure 11: 8K8.4 top cast ingot – ingot defects
Slika 11: 8K8.4 ingot ulit od zgora – napake v ingotu
The investigations were performed within the
TANDEM program of the FT_TA/061 project. The
project was funded partially with the financial support
of the Ministry of Industry and Trade of the Czech
Republic.
5 LITERATURE
1 Martínek, L., Balcar, M., Novák, J., Sochor, L.: Analysis of internal
non-integrities of the ingot made of tool steel, 6th International
metallurgical symposium, Rájecké Teplice, 2003
2 Martínek, L., Balcar, M., Novák, J., Sochor, L.: Internal structure of
the ingot made of tool steel, 20th conference Theory and practice of
making and processing of steel, Ro`nov, Tanger, 2003
3 Carlson, K. D., Shouzhu OU, Hardin, R., Beckerman, CH.: Develop-
ment of new feeding – distance rules using casting simulation: Part I.
Methodology.
4 http://www.engineering.uiowa.edu/~becker/documents.dir/
FeedingPart1.pdf
M. BALCAR ET AL.: DEVELOPING AND TESTING A NEW TYPE-8K MOULD FOR TOOL-STEEL INGOT CASTING
38 Materiali in tehnologije / Materials and technology 42 (2008) 1, 33–38
M. TORKAR ET AL.: AN AES INVESTIGATION OF BRUSHED AISI 304 STAINLESS STEEL ...
AN AES INVESTIGATION OF BRUSHED AISI 304
STAINLESS STEEL AFTER CORROSION TESTING
AES-PREISKAVE KRTA^ENEGA NERJAVNEGA JEKLA AISI 304
PO KOROZIJSKEM PRESKUSU
Matja` Torkar, Djordje Mandrino, Martin Lamut
Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia
matjaz.torkar imt.si
Prejem rokopisa – received: 2007-07-19; sprejem za objavo – accepted for publication: 2007-11-16
During a test in a wet salt (NaCl) chamber rusty spots and elongated pits appeared on the brushed surface of a prototype casing
for a household device made from AISI 304 stainless steel. The corrosion pits only formed on the brushed surface and on the
deformed edges. Metallography revealed deformation-induced martensite and delta ferrite in the microstructure. In the wet salt
environment the brushed surface showed a lower corrosion resistance than the non-brushed surface. The corrosion was
explained as being due to damage to the protective oxide layer with brushing, insufficient re-passivation and the presence of
iron particles. The corrosion propensity of the brushed surface of the casing was increased by the presence of
deformation-induced martensite, delta ferrite and, probably, the remains of iron particles. The use of AISI 316 stainless steel,
alloyed with molybdenum, would be more appropriate for avoiding the formation of corrosion pits due to chloride ions in a wet
environment.
Key words: stainless steel, pitting corrosion, passivation, deformation-induced martensite, AES analysis
Ohi{je prototipne gospodinjske naprave je bilo izdelano iz nerjavnega jekla AISI 304. Iz dekorativnih razlogov je bila povr{ina
krta~ena. Med preskusom v vla`ni slani komori se je pojavila korozija v obliki rjastih made`ev in podolgovatih jamic vzdol`
zakrivljenega roba. Predstavljeni so rezultati metalografske preiskave in analiza korozijskih produktov. Korozijske jamice se
opazi samo na predhodno krta~eni povr{ini in po deformiranih robovih. Z metalografskimi preiskavami smo odkrili v
mikrostrukturi po deformiranih robovih prisotnost deformacijskega martenzita in delta ferita. V vla`nem slanem (NaCl) okolju
je krta~ena povr{ina pokazala slab{o korozijsko odpornost v primerjavi z neobdelano povr{ino. Glavna razloga za korozijo sta
med krta~enjem po{kodovan oksidni sloj in slaba repasivacija po krta~enju. Dodatno so lahko k pojavu korozije prispevali {e
deformacijski martenzit, delta ferit in ostanki delcev `eleza na povr{ini. Posledica slab{e korozijske odpornosti je pojav
jami~aste korozije v vla`nem, slanem okolju. Za zmanj{anje mo`nosti pojava jami~aste korozije bi bila bolj primerna izbira
nerjavnega jekla AISI 316, ki je legirano z molibdenom in s tem bolj odporno proti pojavu jami~aste korozije.
Klju~ne besede: nerjavno jeklo, jami~asta korozija, pasivacija, deformacijski martenzit, AES analiza
1 INTRODUCTION
The passive film on a stainless-steel surface consists
of a mix of iron oxide and chromium oxide. The
chromium oxide film is formed instantaneously in the
air, if the stainless steel surface is clean and dry.
However, complete passivation cannot be achieved if the
product-contact surfaces are not clean or they contain
surface defects. The polishing of the stainless steel’s
surface, therefore, improves the corrosion resistance 1.
The resistance to pitting corrosion is also increased by
alloying the stainless steel with molybdenum. The
beneficial effect of molybdenum on the corrosion
resistance of stainless steels is widely recognized 2,3,4 and
has been attributed to several factors 4, such as: the
enrichment of Cr in the oxide layer; the enrichment of
Mo in the passive film or in the alloy layer just below the
passive film; the thickening of the passive film; and the
stabilisation of the Cr oxides by the presence of
six-valence Mo 5.
The interaction between the different oxides and the
passivation/corrosion characteristics of stainless steel are
very complicated and not yet fully understood.
Different surfacing techniques are applied to stainless
steel to obtain a special visual effect on the surface. The
brushing of stainless steel’s surface is popular because it
gives an attractive look to the surface. The protection
from contamination of the steel surface requires the
brushing of the stainless steel to be with stainless-steel
brushes. In the event that iron particles on the
stainless-steel surface are not removed by cleaning or by
passivation in acid, subsequent pitting corrosion usual
appears in a wet environment 6.
An additional influence on the corrosion resistance of
the material comes from different constituents of the
microstructure, like carbides, martensite, delta ferrite and
inclusions. Also well known is the harmful influence of
deformation-induced martensite 7,8, which forms during
the cold working of stainless steel. After cold working
the stainless steel needs to be annealed at 1050 °C in
order to transform the martensite to austenite.
Furthermore, residual stresses in the material can induce
stress-corrosion cracking and accelerate the appearance
of other corrosion forms when the material is exposed to
a corrosive environment. Pitting corrosion due to
chloride ions can cause considerable damage to the steel
Materiali in tehnologije / Materials and technology 42 (2008) 1, 39–43 39
UDK 669.14.08.8:620.193:543.428.2 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 42(1)39(2008)
surface; it penetrates deep into the material and causes
local weak points or even perforations of the wall.
The process of pitting corrosion is well defined, but
frequently the reasons for its starting are not clearly
defined. The pit can form only in places where the
passive layer is interrupted. The well-known standard
work on corrosion 8 states that pits form more readily on
etched or ground surfaces than on a polished surface, and
the presence of cold working can increase the propensity
to pitting.
The corrosion resistance of stainless steel also
depends on the amount of carbon in the steel. The
precipitation of carbides depletes the neighbourhood of
chromium and accelerates the intercrystalline and pitting
corrosion. Besides martensite and the presence of iron
particles, the corrosion is also accelerated by delta ferrite
and sulphide inclusions.
When the pits start to grow, the corrosion process
diminishes the pH in the pit bottom and chloride ions
move towards the pit. This accelerates the pit growth in
terms of depth, which can lead to the perforation of the
sheet and a leak in the wall or the failure of the structure.
For these reasons, many factors should be considered for
the selection of suitable stainless steel. The investigated
stainless steel showed failures typical for pitting
corrosion.
The aim of the investigation was to analyse the
corrosion residue on the stainless steel’s surfaces and to
identify the main reasons for the corrosion.
2 EXPERIMENTAL
The investigated samples were from a corrosion-
tested AISI 304 stainless-steel casing with a previously
brushed surface. The composition of the investigated
steel is listed in Table 1.
The corroded surface of the casing was carefully
visually and macroscopically surveyed and typical
corroded samples were cut for a metallographic
investigation and for an AES analysis. The samples for
metallographic observation were prepared by the
standard procedure and etched in aqua regia. The
microstructure was investigated with an optical
microscope, the surface defects were observed with a
stereo microscope and in a scanning electron
microscope. The corrosion products were analysed by
Auger spectroscopy (AES). Before the AES analysis, the
surface was cleaned for 45 min with ion etching. For the
AES analysis an electron source of Mg Ka (1253.6 eV)
was used, and the results were compared with the
reference database of binding energies 9 accessible at the
URL: http://srdata.nist.gov/xps/.
3 RESULTS AND DISCUSSION
A careful visual and stereo-microscope examination
revealed spots of rust, white remnants identified as
sodium chloride (Figure 1) and corrosion pits on the
surface of the case. After the corrosion test a difference
was observed when comparing the brushed (Figure 2)
and non-brushed surface (Figure 3) of the case. The
corrosion pits were only observed on the brushed
surface, while the non-brushed surface was without any
corrosion damage.
The rusty spots in Figure 1 represent the remains of
dried rust and were easily removed from the non-brushed
surface. Most of the surface below the rusty spots was
without any corrosion damage. The possible sources of
rust were cut edges, holes for screws and deformed
regions where the sheet was bent. It is also important not
to neglect the remnants of other material that was in
contact with the sheet, like for instance, fragments of the
brush. It seems that a liquid, contaminated with rust, was
spread over the sheet, and after drying the rust spots
remained on the non-corroded steel surface.
SEM microscopy of the brushed surface revealed
individual corrosion pits and groups of corrosion pits
(Figure 2). The longitudinal traces of brushing are
evident on one side of the sheet, while the opposite side
M. TORKAR ET AL.: AN AES INVESTIGATION OF BRUSHED AISI 304 STAINLESS STEEL ...
40 Materiali in tehnologije / Materials and technology 42 (2008) 1, 39–43
Table 1: Chemical composition of the AISI 304 steel
Tabela 1: Kemi~na sestava jekla AISI 304
Composition
Element in mass fractions wt/%
C Si Mn Cr Ni Mo N Cu S
AISI 304 0.053 0.48 1.42 17.4 8.3 0.39 0.04 0.14 0.024
Figure 1: Remains of rust and NaCl on the surface of AISI 304
stainless steel after a corrosion test in wet salt chamber. The longitu-
dinal pit is present on the bent edge
Slika 1: Ostanki rje in NaCl na povr{ini AISI 304 nerjavnega jekla po
korozijskem preskusu v vla`ni slani komori. Na upognjenem robu se
opazi vzdol`no jami~asto izjedo
of the sheet and the non-brushed surface revealed
equiaxed grains with clearly defined grain boundaries
(Figure 3). Only the brushed surface and the bent edge
of the sheet were damaged by corrosion, and on the
curved edge longitudinal pits were observed (Figure 4).
In corrosion-damaged regions we observed either the
aggregation of shallow pits or individual deep pits
(Figures 2 and 5).
The metallography revealed martensite in the de-
formed regions in the austenitic matrix. Also, elongated
delta ferrite (Figure 6) was present in the matrix. Both
phases were a possible source of the corrosion, when in
contact with the corrosion medium.
Because of the wet environment and the presence of
chloride ions in the salt chamber, the local pitting
corrosion process started and advanced to the middle of
the sheet, as revealed by the optical microscopy in the
cross-section (Figure 5). The corrosion process was also
accelerated by the presence of deformation-induced
martensite. To transform the martensite into austenite, it
is necessary to anneal the stainless steel after cold
deformation.
M. TORKAR ET AL.: AN AES INVESTIGATION OF BRUSHED AISI 304 STAINLESS STEEL ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 39–43 41
Figure 6: Microstructure of austenitic stainless steel AISI 304 with
elongated delta ferrite. In deformed regions the deformation-induced
martensite was also present. Etched with aqua regia.
Slika 6: Mikrostruktura avstenitnega nerjavnega jekla AISI 304 z
razpotegnjenim delta feritom. V deformiranih podro~jih je bil prisoten
tudi deformacijski martenzit. Jedkano v zlatotopki.
Figure 3: Untreated surface after the corrosion test. SEM
Slika 3: Neobdelana povr{ina po korozijskem preskusu. SEM
Figure 5: Micrograph of cross-section of the corrosion pit on the bent
edge. Local corrosion advanced deep into the material of the case wall
Slika 5: Posnetek preseka korozijske jamice na upognjenem robu.
Lokalno je korozija prodrla globoko v material stene ohi{ja
Figure 2: Corrosion pits on the brushed surface after the corrosion
test. SEM
Slika 2: Jami~asta korozija na krta~eni povr{ini po korozijskem
preskusu. SEM
Figure 4: Longitudinal corrosion pit on the bent edge. SEM
Slika 4: Vzdol`ne korozijske jamice na upognjenem robu. SEM
The AES analysis of the corroded steel surface
(Figures 7 and 9) after the wet salt chamber test
revealed that in the corrosion pits either sodium and
chlorine or only chlorine were present (Figure 8 and
10). The binding energy of the analyzed element sodium
corresponded to NaCl (Figure 8). The source of the
NaCl was the wet salt chamber. No traces of sodium or
chlorine were detected on the non-corroded surface
(Figure 10, P2). But in typical use a possible source of
chloride ions is from the environment (sprayed brine, the
salt in food, cleaners).
The investigation of the prototype case for use in
domestic and outdoor activities, e.g., a barbecue,
revealed several deficiencies. The AISI 304 stainless
steel is less appropriate for a wet environment where
chloride ions are present. The brushing of the surface,
due to aesthetic reasons, damaged the passive film on the
steel’s surface. A possible accelerator of the corrosion
was the presence of deformation-induced martensite and
delta ferrite in the microstructure of the steel, as well as
unremoved iron particles on the steel’s surface.
4 CONCLUSIONS
The corrosion test in the wet salt chamber revealed
that the brushed surface was sensitive to pitting
corrosion, but the unbrushed surface was corrosion
resistant. From this we concluded that the corrosion of
the prototype stainless-steel casing is very probably due
to the smoothing that damaged the passive layer on the
steel’s surface. The re-passivation was too weak and the
pitting corrosion appeared during the corrosion test due
to the presence of iron particles and chloride ions in the
M. TORKAR ET AL.: AN AES INVESTIGATION OF BRUSHED AISI 304 STAINLESS STEEL ...
42 Materiali in tehnologije / Materials and technology 42 (2008) 1, 39–43
Figure 10: AES spectrum in the areas A1 and A2 (Figure 7), and at
the points P1 and P2 (Figure 9) Only in the corroded areas A1 and A2
was the presence of Na and Cl detected. In the pit P1 only Cl was
detected. Out of the pit no Na or Cl was detected.
Slika 10: AES spekter s podro~ja A1, A2 (slika 7) in v to~kah P1, P2
(slika 9). Samo na korodiranih podro~jih A1 in A2 je bila odkrita
prisotnost Na in Cl, v jamici P1 pa samo prisotnost klora. Izven jamice
ni bilo Na ali Cl.
Figure 8: AES spectrum on the corroded places A1 and A2 marked in
Figure 7. Presented peak for Na KLL at 990.3 eV in non-diffe-
rentiated spectrum of AES corresponds to Na in NaCl 9
Slika 8: AES-spekter na podro~jih korozije A1 in A2 ozna~enih na
sliki 7. Predstavljen vrh za Na KLL pri 990,3 eV v nediferenciranem
spektru AES ustreza Na v NaCl 9
Figure 9: Surface with corrosion pit. At the points P1 and P2 the AES
analyses were performed. The presence of Cl was detected only in the
pit. SEM
Slika 9: Povr{ina z jami~asto korozijo. To~ki P1 in P2 sta bili analizi-
rani z AES. Samo v jamici je bil prisoten Cl
Figure 7: Surface damaged by corrosion (SEM). A1 and A2 mark the
areas analysed with AES. Before the analyses the surface was ion
etched for 45 minutes.
Slika 7: Povr{ina po{kodovana s korozijo (SEM). A1 in A2 sta
podro~ji, ki sta bili analizirani z AES. Pred analizo je bila povr{ina 45
minut ionsko jedkana.
wet environment. The appearance of pitting corrosion is
also to be expected during use.
The investigation of the prototype casing revealed
that the wrong type of stainless steel had been selected
for the proposed application.
Compared to AISI 304 the AISI 316 stainless steel
should be more suited to the application in a wet
environment where chloride ions are also present.
Alloying with 2–3 wt% of molybdenum reduces the
sensitivity of the steel to pitting corrosion.
It is also advisable to passivate the surface of the
stainless steel with diluted nitric acid to remove the iron
particles from the surface. Later, during the use of the
casing, it is important to clean thoroughly and to rinse
with clean water all surfaces that may come into contact
with chloride ions.
5 REFERENCES
1 Leonardo S. Andrade, Sandro C. Xavier, Romeu C. Rocha-Filho,
Nerilso Bocchi, Sonia R. Biaggio, Electropolishing of AISI-304
stainless steel using an oxidizing solution originally used for
electrochemical coloration, Electrochimica Acta, 50 (2005) 13,
2623–2627
2 Z. Szklarska-Smialowska, Pitting corrosion of metals, Houston,
NACE, TX, 1986, 143–157
3 A. Kocijan, ^. Donik, M. Jenko, Electrochemical and XPS studies of
the passive film formed on stainless steel in borate buffer and
chloride solutions, Corrosion Science, 49 (2007) 5, 2083–2098
4 M. F. Montemor, A. m. p. Simões, M. G. S. Ferreira, M. Da Cunha
Belo, The role of Mo in the chemical composition and semi-
conductive behaviour of oxide films formed on stainless steels,
Corrosion Science, 41 (1999) 1, 17 – 34
5 A. Frontini, I. Milo{ev, B. Pihlar, Vpliv vsebnosti molibdena na
biokompatibilnost nerjavnega jekla, Kovine zlit. tehnol., 33 (1999)
3–4, 185–188
6 R. R. Maller, Passivation of stainless steel, Trends in Food Science
& Technology, 18 (2007) 1, S112–S115
7 Torkar, M., Corrosion of AISI 316 Ti in 50% KOH due to defor-
mation induced martensite. Eng fail. anal. 13 (2006) 4, 624–628
8 M. G. Fontana, N. D. Greene, Corrosion Engineering, 1978, 2nd Edi-
tion, McGraw Hill, New York
9 Reference database of binding energies: http://srdata.nist.gov/xps/
M. TORKAR ET AL.: AN AES INVESTIGATION OF BRUSHED AISI 304 STAINLESS STEEL ...
Materiali in tehnologije / Materials and technology 42 (2008) 1, 39–43 43

I. NIKOLIC ET AL.: USING A FIB TO PREPARE Al(OH)3 SAMPLES FOR THE TEM
USING A FIB TO PREPARE Al(OH)3 SAMPLES FOR THE TEM
UPORABA FIB ZA PRIPRAVO VZORCEV Al(OH)3 ZA TEM
Irena Nikolic1, Velimir Radmilovic2, Toll Z. Sholklapper3, Dragoljub Blecic1
1 University of Montenegro, Faculty of Metallurgy and Technology, 81000 Podgorica, Montenegro
2 LBNL, NCEM, MS-72, University of California, Berkeley, CA 94720, USA
3 LBNL, MSSM, University of California, Berkeley, CA 94720, USA
irena cg.ac.yu
Prejem rokopisa – received: 2007-09-25; sprejem za objavo – accepted for publication: 2007-10-23
One of the basic applications of a focused ion beam (FIB) tool is the preparation of samples for transmission electron
microscopy (TEM) analysis. A dual-beam FIB system combines multiple experimental systems within the same chamber and
the combinations of different techniques can be accommodated by unique sample preparation, manipulation and analysis
methods. The system contains both the focused Ga+ ion beam and a field-emission scanning electron column. In this paper we
described a procedure for preparing Al(OH)3 samples for TEM analysis using a FIB.
Key words: focused ion beam, Al(OH)3, milling, deposition, sample preparation
Ena od temeljnih vrst uporabe fokusiranega ionskega curka (FIB) je priprava vzorcev za presevno elektronsko mikroskopijo.
Dualni FIB-sistem kombinira v isti komori ve~ eksperimentalnih sistemov, kombinacija razli~nih tehnik pa omogo~a enkratno
pripravo vzorcev, manipulacijo in analizo. Sistem ima oboje: fokusirani in ozki curek Ga+ in vrsti~no elektronsko kolono s
poljsko emisijo. V ~lanku je opisan postopek uporabe FIB za pripravo vzorcev Al(OH)3 za presevno elektronsko mikroskopijo.
Klju~ne besede: fokusirani ionski curek, Al(OH)3, mletje, depozicija, priprava vzorcev
1 INTRODUCTION
The two main features of the FIB tool are the
capability to remove material from the sample by
sputtering (micro-machining) and to add materials to the
sample by depositing at sub-micron dimensions. There
are a wide variety of applications for a FIB, but they can
be segmented into three main categories 1-3, i.e.,
• IC review and modification, where sputtering is used
to create a cross-section or modification to the
sample, and where cutting is desired together with
the deposition of either metallic or insulating
materials to either modify existing structures or to
create new structures.
• TEM and STEM sample preparation, where the FIB
is used to form a thin slice of material by sputtering
trenches on either side of a slice.
• Producing thin-film specimens.
In this paper we explain a step-by-step approach to
preparing samples of Al(OH)3 for an investigation with
transmission electron microscopy using an FEI Strata
Dual Beam FIB. Investigating the mechanism of
Al(OH)3 crystal growth required the use of both scanning
and transmission electron microscopy (SEM and TEM).
The data obtained with the SEM are not sufficient to
draw a reliable conclusion about the mechanism of
Al(OH)3 growth, with respect features such as twinning,
and a determination of the twinning plane and direction.
This means that TEM is required and a special way of
preparing samples for such an investigation needs to be
developed.
2 EXPERIMENT
The crystals of Al(OH)3 were obtained by a
crystallization process from a caustic soda solution. They
crystallized in the form of small particles joined together
in larger agglomerates. When such particles were
exposed to the ion beam, they started to dissipate. In
order to prevent the dissipation process we put Al(OH)3
powder in the cannon mod and then filled it up with
liquid epoxy. The vacuum, liquid epoxy filled up all the
vacancies inside the Al(OH)3 crystals. After 12 h the
epoxy was solidified and made the particle interior
stronger. After this, the cannon samples were polished in
order to remove excess epoxy and coated with gold so
they were conductive. The samples prepared in such a
way were ready for the FIB treatment.
3 DISCUSSION
The preparation of the Al(OH)3 sample for the TEM
analysis using the FIB was carried out in the National
Center for Electron Microscopy, Lawrence Berkeley
Laboratory, University of Berkeley, California. A picture
of the FIB is shown in the Figure 1.
Figure 2 shows a step-by-step approach to the prepa-
ration of a sample for transmission electron microscopy
using the selective focused ion beam (FIB) milling
technique.
After the area of interest is found, the first step is the
deposition of platinum. The platinum is deposited, either
under an ion or electron beam, in the strip with
dimensions of approximately 25 x 1 x 2) µm (length,
Materiali in tehnologije / Materials and technology 42 (2008) 1, 45–47 45
UDK 620.18:543.2:669.715 ISSN 1580-2949
Professional article/Strokovni ~lanek MTAEC9, 42(1)45(2008)
width, height). We used an ion beam with a beam current
corresponding to double the area of the deposition
(Figure 2a).
The next step is milling of the cleanup trenches on
both sides of a platinum strip (Figure 2b). Before the
milling procedure begins, the milling stair step pattern
needs to be defined in terms of shape and position. The
shape of the stair-step pattern is shown in Figure 3, 1.
The milling procedure consists of two steps. A large
beam current (20 000 pA) is used in the first step of the
bulk removal with the aim to reduce the time of milling
when the ion beam is positioned far from the platinum
strip. In the next step, after the bulk of the material has
been removed, the face of the specimen slice must be
cleaned up using a lower beam current. Again, a stair
step pattern must be defined. After the final milling the
platinum strip must retain the defined thickness.
After the final milling is finished, the sample is tilted
by 45° in order to make a cross-section of the obtained
sample slice. Under a beam current of 7000 pA a thin
slice of the specimen is cut free from the three sides
leaving a small "bridge" in the upper-left corner (Figure
2c).
In the next stage (Figure 2d), the sample stage in the
chamber (z-position) is lowered (to avoid any needle
damage) and a platinum needle is inserted into the
chamber. Then, the sample stage is raised again. Using a
manipulator the Pt needle is navigated and attached to
the free upper (right) corner. The needle is welded to the
specimen slice by depositing platinum in the same way
as in the case of the deposition of a platinum strip. For
this test a smaller pattern (approximately (1x1) µm) and
a lower beam current were used. Finally, the remaining
bridge connecting the slice sample to the sample bulk is
cut free using the 300 pA beam current. In the next step
the sample stage in the chamber must be lowered slowly
to lift out the membrane – the thin specimen slice.
I. NIKOLIC ET AL.: USING A FIB TO PREPARE Al(OH)3 SAMPLES FOR THE TEM
46 Materiali in tehnologije / Materials and technology 42 (2008) 1, 45–47
Figure 1: FEI Strata Dual Beam FIB in NCEM LBNL, University of
California, Berkeley
Slika 1: Strata dualni curek FIB pri NCEM LBNL, Univerza v Kali-
forniji, Berkeley
Figure 2: Step-by-step approach to preparing Al(OH)3 samples for TEM using a FIB
Slika 2: Stopnje priprave vzorcev za TEM z uporabo FIB
The last sequence is welding the membrane to the
half-cut copper grid. Prior to this the copper grid is
placed in the holder in the chamber and aligned
horizontally. The membrane is then placed in a
pre-milled trench located near the flat side of the grid
and welded to the grid with platinum deposition (Figure
2e). When the membrane is welded to the grid, the
platinum needle is cut free from the membrane. After
this step the specimen slice is still too thick for any TEM
analysis and the bulk of the FIB dust (the result of
milling) is present on the surface of the membrane. Thus,
a final thinning on the grid is necessary and, indeed,
vital; this is the last step in the preparation of the
Al(OH)3 sample for TEM analysis. There is, however, a
final step; it is good practice to let the ion beam scan (for
a short time) the thin section using a low beam current
(50–100 pA). The picture of the sample after the final
thinning on the grid is shown in Figure 2f. Such a thin
sample (about 50 nm) is appropriate for a TEM
investigation.
4 CONCLUSION
The application of a focused ion beam in the
preparation of Al(OH)3 crystals for TEM is shown to be
a good technique. Problems related to the dissipation of
Al(OH)3 particles under the ion beam can be
successfully avoided by using epoxy. A general problem
related to the FIB technique is the re-deposition of dust
(removed material) during the milling step. In addition,
the problem that may occur during the preparation of the
sample for the TEM analysis is related to the last steps
when the sample slice is very thin and the handling
needs to be very delicate. The procedure of cleaning the
sample slice on the grid must be carefully performed
with the choice of an appropriate beam current. The
lower beam current is recommended, as is the repetition
of the cleaning procedure several times until the desired
sample thickness is achieved. If the sample is too thin
the sample slice may break.
5 REFERENCES
1 J. Orloff, M. Utlaut, L. Swanson, High Resolution Focused Ion
Beams: FIB and Its Applications, Kluwer Academic /Plenum
Publishers, New York, 2003
2 L. A. Giannuzzi, F. A. Stevie, Introduction to focused ion-beams:
instrumentation, theory, techniques and practice, Springer, New
York, 2005
3 Z. Rahman, S. Shkula, H. Cho, S. Seal: Microscopy and Analysis, 20
(2006) 3, 11–13
I. NIKOLIC ET AL.: USING A FIB TO PREPARE Al(OH)3 SAMPLES FOR THE TEM
Materiali in tehnologije / Materials and technology 42 (2008) 1, 45–47 47
Figure 3: Shape of stair-step pattern used for the bulk removal of
material
Slika 3: Shema stopni~astega na~ina odvzema materiala