Engineering, 2013, 5, 887-901
Published Online November 2013 (
Open Access ENG
Functional Acrylic Polymer as Corrosion Inhibitor of
Carbon Steel in Autoclaved Air-Foamed Sodium
Silicate-Activated Calcium Aluminate/Class
F Fly Ash Cement
Toshifumi Sugama, Tatiana Pyatina
Sustainable Energy Technologies Department, Brookhaven National Laboratory, Upton, USA
Received September 4, 2013; revised October 4, 2013; accepted October 11, 2013
Copyright © 2013 Toshifumi Sugama, Tatiana Pyatina. This is an open access article distributed under the Creative Commons Attri-
bution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly
The study focused on investigating the effectiveness of functional acrylic polymer (AP) in improving the ability of air-
foamed sodium silicate-activated calcium aluminate/Class F fly ash cement (slurry density of 1.3 g/cm3) to mitigate
the corrosion of carbon steel (CS) after exposure to hydrothermal environment at 200˚C or 300˚C. Hydrothermally-initi-
ated interactions between the AP and cement generated the formation of Ca-, Al-, or Na-complexed carboxylate deriva-
tives that improved the AP’s hydrothermal stability. A porous microstructure comprising numerous defect-free, evenly
distributed, discrete voids formed in the presence of this hydrothermally stable AP, resulting in the increase in compres-
sive strength of cement. The foamed cement with advanced properties conferred by AP greatly protected the CS against
brine-caused corrosion. Four major factors governed this protection by AP-incorporated foamed cements: 1) Reducing
the extents of infiltration and transportation of corrosive electrolytes through the cement layer deposited on the under-
lying CS surface; 2) Inhibiting the cathodic reactions at the corrosion site of CS; 3) Extending the coverage of CS by the
cement; and 4) Improving the adherence of the cement to CS surface.
Keywords: Calcium Aluminate Foamed Cement; Thermal Shock Resistance; Corrosion Protection; Carbon Steel
1. Introduction
The major thrust in assembling and constructing En-
hanced Geothermal Systems (EGSs) lies in creating a
hydrothermal reservoir in a hot dry rock stratum 200˚C,
located at ~ 3 - 10 km below the ground surface. In this
operation, water at a low temperature of ~25˚C is
pumped down the injection well to initiate the opening of
existing fractures. Multi-injection wells are required to
create a desirable network of permeable fractures. After
that, a production well is installed within the fracture’s
During the water injection, considerable attention must
be paid to a significant differential between the tempera-
tures at the bottom of the well and of the water from the
injection well. This differential leads to a sudden tem-
perature drop of ~180˚C at the cement sheath sur-
rounding the down-hole casing. Such a large thermal
gradient due to the cooling effect of the injection water
can give an undesirable thermal stress to the cement
sheath, causing its potential failure that will entail a cata-
strophic blowout of the well. To mitigate such tempera-
ture differential-caused stresses, the cement placed in
EGS is required to possess adequate resistances to ther-
mal cycle fatigue and thermal shock.
To deal with this issue and develop thermal shock-re-
sistant cements, our previous study investigated the ef-
fectiveness of sodium silicate-activated Class F fly ash in
improving such resistance of refractory calcium alumi-
nate cement (CAC) [1]. We conducted a multiple air
heating (500˚C)-water (25˚C) cooling quenching cycle
tests for evaluating thermal shock resistance of 200˚C-
autoclaved CAC/Class F fly ash/sodium silicate blend
cements. The phase composition of the autoclaved CAC/
Class F fly ash blend cements comprised four crystal-
line hydration products, boehmite, katoite, hydrogros-
sular, and hydroxysodalite, which were responsible for
strengthening cement. Among them, the hydroxysodalite
was transformed into nano-scale crystalline carbonated
sodalite during this cycle test. This phase transition not
only played a pivotal role in densifying cementitious
structure and in sustaining the original compressive
strength developed after autoclaving, but also it improved
the CAC’s resistance to thermal shock.
Another critical issue in formulating well casing ce-
ments is the density of the cement slurry. Cement slurry
of a typical density, 1.8 - 2.0 g/cm3, may create undesir-
able fractures zones in a weak unconsolidated rock for-
mation due to the high hydrostatic pressure needed for its
circulation and cause an issue of lost circulation. Thus, it
is vital to use cement slurry with the minimum density
possible for lowering this hydrostatic pressure. However,
our study of the properties of fine air bubble-incorpo-
rated foamed cement slurry strongly suggested that the
shortcoming of this hydrothermally cured foamed cement
was its low protection of carbon steel (CS) casings against
corrosion, compared with that of non-foamed cement [2].
As is well documented [3-8], when the surfaces of CS
come in contact with alkaline pore solution of pH > 12 in
Ordinary Portland Cement (OPC), they form a passive Fe
oxide film that protects the CS against corrosion. How-
ever, the stability of this passivation layer depends pri-
marily on the surrounding environment and exposure
conditions. In geothermal environments, there are two
major chemical factors governing the destruction of the
corrosion-preventing passive film and initiation of casing
corrosion: One is the prevalence of carbonation-related
carbonate ions; and, the other is the attack of chloride
ions on the passive film in which the risk of corrosion of
CS commonly is equated with the chloride content, usu-
ally taking into account the chloride/hydroxyl [(Cl)/
(OH)] ratio. For the former factor, carbon dioxide (CO2)
can penetrate into the pores in OPC, and then react with
water to form carbonic aid, H2CO3, thereby lowering the
pH of the pore solution in OPC layers adjacent to the
CS’s surfaces. Thus, such depassivation promotes the
localized corrosion of CS, and then, the corrosion prod-
ucts formed on CS’s surface engender a volumetric ex-
pansion of CS, generating a stress cracking and the spal-
lation of OPC sheath at the interfacial boundary regions
between the OPC and CS-based well casing. Further, the
failure of OPC sheath at the initial stage of CS’s corro-
sion not only accelerates its corrosion rate, but also
shortens the service lifespan of casings; such impairment
of the integrity of well structure entails the need for
costly repairs and restorative operations including re-
drilling to reconstruct the damaged well, and occasion-
ally necessitates the assemblage of new wells. To avoid
such catastrophic failure of well structure, very expen-
sive metal components incurring a high capital invest-
ment, for example, titanium alloy, stainless steel, and
inconel components frequently are employed in fabricat-
ing corrosion-preventing well casings. Thus, the ideal
approach to reducing capital investments along with the
operational- and maintenance-costs in EGS wells is using
inexpensive CS-based casings covered with corrosion-
inhibiting well cementing materials. In this concept, the
cementing materials play a pivotal role in extending the
lifecycles of the CS-made casings and in retaining the
well’s integrity, so eventually lowering the costs of elec-
tricity generated from geothermal power plants.
Presently, there are two prevalent methods to mini-
mize and alleviate the corrosion of CS [9,10]: The first
one is to apply coatings to CS’s surface as cathodic pro-
tection; the other is to use anodic corrosion inhibitors.
Our previous work on the former method [11] centered
on evaluating the ability of a hybrid coating system con-
sisting of a water-borne acrylic emulsion as the matrix
and CAC as the hydraulic binder to mitigate the corro-
sion of CS in CO
2-laden geothermal environments at
250˚C. We identified two major factors that significantly
contributed to inhibiting the corrosion of CS in such hy-
drothermal environment: One was the hydrothermally
stable products formed by the interactions between
acrylic polymer and cement; the other was the develop-
ment of a dense microstructure by the combination of
well-formed calcite and boehmite crystals with poly-
mer-cement reaction products.
Based upon this information, the CAC/Class F fly ash/
sodium silicate blend cements, formulated as the thermal
shock-resistant cement (TSRC), would be required to
possess two indispensable properties: First, the slurry
density must be lowered; and, second, the ability of hy-
drothermally cured TSRC to protect the CS casing
against corrosion must be assured. Hence, we formulated
air bubble-foamed TSRC for preparing low density ce-
ment slurry at room temperature, and then the foamed
TSRC was modified with acrylic-based polymer emul-
sion as corrosion-inhibiting additive to lessen the corro-
sion of CS under a hydrothermal environment at 200˚C
and 300˚C. In characterizing the polymer-modified
foamed TSRC, our study focused on eight major object-
tives: 1) Characterizing the hydrolysis-hydration reac-
tions of the cement slurries at 85˚C and assessing the
total energy evolved during these reactions; 2) identify-
ing the reaction product formed by interactions between
cement and acrylic-based polymer (AP) at 200˚C and
300˚C, and assessing its thermal stability; 3) determining
the compressive strength of 200˚C- and 300˚C-auto-
claved cements; 4) identifying the crystalline phases and
their composition formed in the cement after autoclaving
at 200˚C and 300˚C; 5) exploring the microstructure de-
veloped within the cements; 6) measuring the conducti-
vity of corrosive ions through the cement layers depos-
ited on the CS’s surfaces by the AC electrochemical im-
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pedance spectroscopy; 7) evaluating the ability of cement
coating to reduce the corrosion rate of CS using the DC
electrochemical potentiodynamic polarization; and 8)
assessing the extent of adherence of the cements to the
underlying CS.
Integrating all the data obtained from the objectives
described above would clarify the potential of the AP-
modified foamed TSRC as corrosion-alleviating well
cement for CS casings in EGS environment at 200˚C and
2. Experimental Procedures
2.1. Materials
Class F fly ash was obtained from Boral Material Tech-
nologies, Inc., and its chemical composition detected by
micro energy-dispersive X-ray spectrometer (EDX) was
as follows; 50.4% SiO2, 34.8% Al2O3, 7.1% Fe2O3, 3.1%
K2O, 2.7% CaO, 1.6% TiO2, and 0.4% SO3. A sodium
silicate granular powder under the trade name “Metso
Beads 2048,” supplied by the PQ Corporation was used
as the alkali activator of Class F fly ash. Its chemical
composition was 50.5 wt% Na2O and 46.6 wt% SiO2.
Secar #80, supplied by Kerneos Inc. was used as refract-
tory calcium aluminate cement (CAC). The X-ray pow-
der diffraction (XRD) data showed that the crystalline
compounds of Class F fly ash had three major phases,
quartz (SiO2), mullite (3Al2O3·2SiO2), and hematite
(Fe2O3), while CAC encompassed three crystalline phas-
es, corundum (
-Al2O3), calcium monoaluminate
(CaO·Al2O3, CA), and calcium dialuminate (CaO·2Al2O3,
The AISI 1008 cold rolled steel test panel according to
ASTM D 609C was used as the carbon steel (CS) sub-
strate, supplied by ACT Test Panels, LLC. An alkaline
cleaner #4429, from American Chemical Products, was
employed to remove surface contaminants from it. This
cleaner was diluted with deionized water to prepare 5
wt% cleaning solution.
Acrylic emulsion under the trade name “HYCAR 26-
0688,” supplied by Lubrizol, was evaluated as the corro-
sion-inhibiting additive for cement. This water-disperse-
ble acrylic emulsion consisted of 49.5 wt% solid acrylic
polymer (AP) and 41.5 wt% aqueous medium, and its pH
was 2.32.
Halliburton supplied the cocamidopropyl dimethyl-
amine oxide-based foaming agent (FA) under the trade
name “ZoneSealant 2000.”
The formula of dry-blend cement employed in this
study consisted of 60 wt% CAC and 40 wt% Class F fly
ash. The sodium silicate at 6.2% by total weight of the
blended cement was added to 60/40 CAC/Class F fly ash
ratio to prepare the one dry-mix cement component. The
amounts of AP used to modify the cement were 0.5, 1.0,
and 2.0% by total weight of this dry mixture. The FA
was added at 1.0% by total weight of water. The follow-
ing sequence was employed to make the foamed AP-
containing cement slurry: First, the proper amount of
water-miscible FA was blended in water, and then the
proper amount of acrylic emulsion was incorporated into
it; second, this water-based solution was added to the dry
cement component to prepare a pumpable slurry; the
water/(cement + acrylic emulsion) [W/(C + AE)] ratios
for 0, 0.5, 1.0, and 2.0 wt% AP-modified cements were
0.44, 0.37, 0.28, and 0.15, receptivity; and, finally, the
AP-containing foamed cement slurry was mixed thor-
oughly in a shear-blender for 30 sec. This shear mixing
made it possible to prepare aerated slurry containing a
vast number of air bubbles, leading to the transformation
of the original stiff slurry into a smooth creamy one.
The surfaces of the 65 mm × 65 mm CS test panels
were coated with the foamed AP-containing cements in
the following sequence: First, the “as-received” test pan-
els were immersed in a 5 wt% alkali-cleaning solution at
40˚C for 10 min; second, alkali-cleaned panels’ surfaces
were rinsed with a tap water at 25˚C, and dried for 24
hours in air at room temperature; third, the panels were
dipped into a soaking bath of cement slurries at room
temperature, and withdrawn slowly; fourth, the cement
slurry-covered panels were left for 3 days at room tem-
perature, allowing the slurry to convert into a solid layer;
and, finally, the cement-coated panels were autoclaved
for 24 hours at 200˚C and 300˚C before conducting the
electrochemical corrosion tests. The thickness of the de-
posited cement layers on the CS’s surfaces ranged from
~1.0 to ~1.4 mm.
2.2. Measurements
The density, g/cm3, of the foamed cement slurries was
determined using the fluid density container. Then, the
slurries were poured in cylindrical molds (20 mm diam.
and 40 mm long) and left for 3 days at room temperature
before compressive strength measurements. Thereafter,
the hardened foamed cements were removed from the
mold and autoclaved at 200˚C and 300˚C for 24 hours,
then left for 24 hours at room temperature.
TAM Air Isothermal Microcalorimeter was adapted to
investigate the hydrolysis-hydration reactions in cement
slurries and to determine the cumulative heat flow
evolved during these reactions at 85˚C. The changes in
compressive strength of the 200˚C- and 300˚C-24 h-
autoclaved cements as a function of AP’s content were
obtained using Instron Model 5967. The reaction prod-
ucts between the AP and cement at hydrothermal tem-
peratures of 85˚C, 200˚C, and 300˚C were identified us-
ing Fourier Transform Infrared Spectroscopy (FT-IR),
and then the thermal stability of identified reaction prod-
ucts was surveyed using Thermo Gravimetric Analyzer
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(TGA) at the heating rate of 20˚C/min in a N2 flow. The
composition of crystalline phases formed in 200˚C- and
300˚C-autoclaved foamed cements was identified by X-
ray diffraction (XRD). Alterations of the microstructure
developed in foamed cements at 200˚C by AP were ex-
plored with the High-Resolution Scanning Electron Mi-
croscopy (HR-SEM). The adherence behavior of AP-
modified and non-modified foamed cements to CS’s sur-
face was surveyed using
EDX. DC electrochemical
testing for corrosion of the underlying CS was performed
with the EG&G Princeton Applied Research Model
326-1 Corrosion Measurement System. For this assess-
ment, the cement-coated CS specimen was mounted in a
holder, and then inserted into an EG&G Model K47
electrochemical cell containing a 1.0 M sodium chloride
electrolyte solution. The test was conducted under an
aerated condition at 25˚C, on an exposed surface area of
1.0 cm2. The polarization curves were measured at a scan
rate of 0.5 mVs1 in the corrosion potential range from
0.05 to 0.75 V. AC electrochemical impedance spec-
troscopy (EIS) was used to evaluate the ability of the
cement layers to protect the CS from corrosion. For this,
the coated CS specimens with a surface area of 13 cm2
were mounted in a holder, and then inserted into an elec-
trochemical cell containing a 1.0 M sodium chloride
electrolyte at 25˚C; single-sine technology with an input
AC voltage of 10 mV (rms) was employed over a fre-
quency range of 105 to 103 Hz. To estimate the protec-
tive performance of the cements, the pore resistance, Rp,
(ohm-cm2) was determined from the plateau in Bode-plot
scans occurring in low frequency regions.
3. Results and Discussion
3.1. Density of Slurries
Figure 1 shows the changes in slurry density as a func-
tion of AP content. Without any AP, the density of the
AP content, wt%
0.0 0.5 1.0 1.5 2.0
Slurry density, g/cm
0% FA
1% FA
Figure 1. Changes in slurry density for FA-foamed and
non-foamed cement slurries as a function of AP content.
non-foamed cement slurry denoted as 0% FA was 1.86
g/cm3. This density drastically dropped by 31% to 1.29
g/cm3, as 1% FA was added to it, demonstrating that this
FA is very effective in and suitable for reducing the den-
sity of this specific blend cement slurry. When AP addi-
tive was incorporated into non-foamed and foamed slur-
ries, their densities gradually declined with an increasing
content of AP; adding 2 wt% AP reduced the density to
1.82 g/cm3 for non-foamed and to 1.2 g/cm3 for foamed
slurry, suggesting that AP has air-entraining properties in
slurry during mixing.
3.2. Hydrothermal Stability of AP in Cement
One inevitable question that must be asked is the suscep-
tibility of AP to the hydrothermal degradation in cement
slurries at 200˚C and 300˚C. To respond to this question,
we prepared samples composed of 50 wt% acrylic emul-
sion and 50 wt% CAC/Class F fly ash/sodium silicate
blend cement at room temperature, and then autoclaved
them for 24 hours at 85˚C, 200˚C, and 300˚C, to conduct
FT-IR and TGA analyses. The “as-received” acrylic
emulsion was used as the control for FT-IR analysis.
Figure 2 shows the FT-IR spectra over frequency
range from 1850 to 1250 cm1, for the control, and APs
formed in 85˚C-, 200˚C -, and 300˚C-autoclaved cements.
The absorbance spectrum of 85˚C-autoclaved sample
encompassed four prominent bands: At 1730 cm1 attrib-
uted to the acrylic acid, -COOH, and alkyl ester, -COOR
(R, ethyl or butyl ester); at 1552 and 1405 cm1 corre-
sponding to the symmetric and anti-symmetric stretching,
respectively, of the carboxylate anion, –COO in –COO
M (M, metallic cations) linkage structure; and, at 1452
cm1 associating with CO3
2 in carbonated compounds.
In comparison, no bands related to carboxylate anion
were found in the control. The spectrum of 200˚C-ex-
17501850 1650 1550 1450 13501250
Wavenumber cm
1452 1405
Absorbance ratio of
1552 cm
/ 1730 cm
= 0.0
1552 cm
/ 1730 cm
= 0.2
1552 cm
/ 1730 cm
= 2.3
1552 cm
/ 1730 cm
= 5.0
Figure 2. Comparison of FT-IR spectral features and absor-
bance ratios of 1552 cm1 to 1730 cm1 bands for control AP
and AP autoclaved in cement slurry at 85˚C, 200˚C and
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posed sample displayed two features different from the
sample exposed to 85˚C: One was a striking decay in the
intensity of band at 1730 cm1; the other was the marked
growth of the absorbance bands at 1552 and 1405 cm1.
This fact strongly suggested that raising the hydro-
thermal temperature promoted the transformation of
acrylic acid and alkyl ester within the molecular structure
of AP into the carboxylate anion derivative. This trans-
formation was further enhanced as the temperature rose
to 300˚C, reflecting a near disappearance of the band at
1730 cm1. To support this information, we compared the
absorbance ratios of 1552 cm1 to 1730 cm1 for the con-
trol, and 85˚C, 200˚C, and 300˚C samples (Figure 2).
For the control, it was very difficult to determine the ab-
sorbance at 1552 cm1. The 85˚C sample had an absorb-
ance ratio of 0.2, strongly verifying that the reactions
between AP and cement already had occurred at such a
relatively lower temperature. As expected, the value of
this ratio increased with an increasing autoclave tem-
perature; at 300˚C, the ratio of 5.0 was twenty five fold
greater than that at 85˚C. These results clearly validated
that the acrylic acid or alkyl ester carboxylate anions
transformation occurred progressively at elevated hydro-
thermal temperatures.
One important criterion of a high-temperature corro-
sion inhibitor is its stability at 200˚C. Hence, our atten-
tion next focused on surveying the thermal stability of
AP derivatives obtained in high-temperature hydrother-
mal reactions. The same samples as those used in the
previous FT-IR analysis were dried for 24 hours at 85˚C
to eliminate any free moisture and tested by TGA. Fig-
ure 3 depicts the TGA curves between 25˚C and nearly
500˚C for the 85˚C-, 200˚C-, and 300˚C-autoclaved sam-
ples. The major thermal decomposition of AP formed in
85˚C-autoclaved samples occurred around 300˚C. De-
composition temperature of AP in 200˚C-autoclaved
sample increased to ~349˚C, which was ~49˚C higher
than for 85˚C sample.
The decomposition temperature further increased to
100 200 300 400 500
Temperature, °C
Weight, %
Figure 3. Thermal decomposition curves of AP autoclaved
in cement slurry at 85˚C, 200˚C or 300˚C.
~438˚C for AP from 300˚C-autoclaved sample.
Combined results of FT-IR and TGA studies high-
lighted that the hydrothermal stability of AP in auto-
claved cement was improved by incorporating more
Metal-complexed carboxylate structures at 300˚C, com-
pared to AP autoclaved in cement at 200˚C. Thus the AP
in cement withstood the 300˚C hydrothermal environ-
3.3. Hydration of AP-Modified Foamed Cement
To investigate the changes in cement hydration with ad-
dition of FA, AP or both we measured heat flow evolved
at the isothermal temperature of 85˚C during cement hy-
dration as a function of elapsed time.
Figure 4 depicts relations between the normalized heat
flow energy and elapsed time for non-foamed and
foamed cement slurries without AP. For all the samples,
the initial heat release peak denoted as the peak No. 1,
corresponds to the particle wetting, dissolution of sodium
silicate activator, beginning of the hydrolysis of some
CAC and Class F fly ash by the alkali activation of dis-
solved sodium silicate. This peak was observed shortly
after placing an ampoule of the sample in a calorimeter
with the maximum heat flow (MHF) been reached within
50 min. Subsequently, two other heat peaks marked as
No. 2 and No. 3 were traced for both the non-foamed and
1 wt% FA-foamed slurries. This cement system consisted
of two cement-forming reactants, CAC and Class F fly
ash, our preliminary study (not shown) demonstrated that
the hydration reactions of CAC were much faster than
those of Class F fly ash. Thus, in the non-foamed cement
slurry, the No. 2 peak emerged at elapsed time of 11 h 32
min was more likely associated with CAC hydration.
It had MHF of 0.83 mW/g. The Class F fly ash-related
No. 3 peak generated a 0.56 mW/g MHF at the elapsed
time of 47 h 13 min. When non-foamed slurry was modi-
Heat flow, mW/g
01224 36 48
Time, hour
Non-foamed cement
1 wt% FA-foamed cement
Figure 4. Microcalorimetric curves for non-foamed and 1
wt% FA-foamed cement slurries at 85˚C.
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fied with 1 wt% FA, its heat flow-time curve differed in
three aspects from that of non-foamed one: First, the time
to reach No. 2 and No. 3 peaks became shorter; second,
the MHF value of the No. 2 peak increased pronouncedly
to 2.52 mW/g; and, third, the MHF value at No. 3 peak
had declined by ~45% to 0.31 mW/g. These results sug-
gested that the foaming agent promoted the hydration
reactions of CAC, so enhancing CAC-related MHF en-
Table 1 lists the elapsed times at the onset and end of
No.2 and No. 3 reactions along with the total energies
evolved in these reactions for non-foamed and foamed-
cement slurries at 85˚C. The total energies were com-
puted from the enclosed areas of the heat flow-elapsed
time curves with the baseline extending between peaks’
starts and ends. The No. 2 reaction, attributed to CAC’s
hydration, clearly verified that it was accelerated by
adding the 1 wt% FA. The onset of this reaction was
shortened to 1 h 59 min by FA, compared to 4 h 6 min
for the FA-free slurry. The computed heat flow energy,
J/g, verified the increase of this energy by FA; in fact, a
40.4 J/g generated with 1 wt% FA was tantamount to
11.3% higher than for the non-foamed slurry. Such an
effect of FA on CAC reactions accelerated the reactivity
of Class F fly ash, thereby shortening the time to reach
the onset of the No.3 reaction to 25 h 6 min for 1 wt%
FA from 33 h 53 min required for non-foamed slurry.
However, as for the Class F fly ash-related heat release,
the contribution of FA to the increase in this energy was
minimal; in fact, the value of 11.8 J/g for non-foamed
slurry declined 1.8 times when 1 wt% FA was added to it.
Nevertheless, the total energy evolved from No. 2 and
No. 3 reactions rose with FA.
Next, we obtained information on heat flow-elapsed
time relations for the foamed slurries with 0.5, 1.0, and
2.0 wt% AP at 85˚C.
Table 2 summarizes the elapsed times at the onset and
end of each No. 2 and No. 3 reactions along with the
total energy evolved from these two reactions for 0, 0.5,
1.0, and 2.0 wt% AP-modified foamed slurries. As is
evident from these data, addition of AP in the concentra-
tion range of 0.5 - 2.0 wt% pronouncedly retarded the
CAC reaction. The onset time of this reaction for un-
modified foamed slurry was delayed by as much as 3 and
5 hours with 1 and 2 wt% of AP, respectively. In contrast,
AP was not as effective in delaying the onset time of
Class F fly ash hydration as it was with CAC. On the
other hand, the total energy evolved from the No. 2 reac-
tion was strikingly augmented with an increasing AP
content, from 33.9 J/g for 0 wt% to 55.8 J/g for 2 wt%,
for Class F fly ash-associated No. 3 reaction. Hence, AP
preferentially retarded the hydration reaction of CAC,
rather than that of Class F fly ash. Such augmentation of
CAC reaction energy by adding AP could be due to the
energy generated by the acid-base interactions between
the AP and CAC to form M-complexed carboxylates (M-
aluminum, calcium, alkali metals) as the reaction prod-
3.4. Compressive Strength
Figure 5 plots the compressive strength of the 200˚C-
and 300˚C-autoclaved foamed cements modified with 0,
0.5, 1.0, and 2.0 wt% AP. The value of compressive
strength depended on two factors, the content of AP and
the hydrothermal temperature. For the latter, cement
autoclaved at 300˚C was more effective in developing
the compressive strength, than cement autoclaved at
For the foamed cement without AP, autoclaving at
Table 1. Calorimetric test results for non-foamed, and 1 wt% FA-foamed cement slurries at 85˚C.
No. 2 reaction No. 3 reaction
FA, wt%
Onset time, h: min End time, h: min Evolved heat, J/gOnset time, h: minEnd time, h: minEvolved heat, J/g
Total evolved heat,
0 4:06 34:28 24.5 33:53 60:41 11.8 36.3
1 1:59 22:19 33.9 25:06 48:02 6.5 40.4
Table 2. Calorimetric test results for No. 2 and 3 reactions with 0, 0.5, 1.0, and 2.0 wt% AP-modified foamed slurries.
No. 2 reaction No. 3 reaction
AP, wt%
Onset time, h: min End time, h: min Evolved heat, J/gOnset time, h: minEnd time, h: minEvolved heat, J/g
Total evolved
heat, J/g
0.0 1:59 22:19 33.9 25:06 48:02 6.5 40.4
0.5 2:20 22:49 42.5 25:15 47:10 7.1 49.6
1.0 3:59 21:30 47.2 25:46 48:21 8.8 56.0
2.0 5:33 21:51 55.8 26:03 45:35 10.8 66.6
AP content, wt%
Compressive strength, MPa
After 200°C autoclave
After 300°C autoclave
Figure 5. Changes in compressive strength of the 200˚C-
and 300˚C-autoclaved foamed cements as a function of AP
200˚C resulted in a compressive strength of 2.79 MPa.
This strength rose by 1.5-fold to 4.13 MPa after auto-
claving at 300˚C, so confirming that sodium silicate-
activated CAC/Class F fly ash blended cement system
withstood the hydrothermal temperatures up to 300˚C.
Importantly, the rise in compressive strength at 300˚C
reflected the densification of formed cement; in fact, the
bulk density of 300˚C-autoclaved cements was ~15%
higher than that one autoclaved at 200˚C.
The presence of AP improved the compressive
strength at both 200˚C and 300˚C; this strength tended to
rise progressively with increasing AP content. At 200˚C,
adding only 0.5 wt% AP resulted in the compressive
strength increase by 19% to 3.45 MPa. Further increasing
AP content to 2.0 wt% strengthened it more; reaching a
value of 6.68 MPa, equivalent to 2.4-fold improvement
compared with that of AP-free cement. A similar trend
was observed from 300˚C-autoclaved foamed cement;
adding 2.0 wt% AP developed a 7.54 MPa compressive
strength, corresponding to 1.4-fold improvements above
that of 0.5 wt% AP. Thus, the AP additive greatly con-
tributed to improving the cement’s compressive strength
in a hydrothermal environment at 200˚C and 300˚C.
3.5. Phase identification and Microstructure
One intriguing question of the incorporation of AP into
the autoclaved foamed cement was whether it would
change the composition of the crystalline phases as the
hydrothermal reaction products responsible for strength-
ening it. Accordingly, XRD was used to identify the
phases assembled in the foamed cements containing 0
and 2 wt% AP after autoclaving at 200˚C and 300˚C.
Figure 6 depicts the XRD traces, 2θ degree, ranging
from 3 to 51, for the powder samples of 200˚C-auto-
10 2040 5030
Intensity (arb. Unit)
S: Hydroxysodalite
B: Boehmite
K: Katoite
Q: Quartz
G: Hydrograssular
C: Corundum
Figure 6. XRD patterns for 200˚C -autclaved foamed ce-
ments containing 0 (bottom) and 2% AP (top).
claved foamed cements with and without AP. For AP-
free foamed cement, the XRD pattern (bottom) showed
that the cement was composed of the four crystalline
hydration reaction products, hydroxysodalite
[Na4Al3Si3O12(OH)], boehmite (γ-AlOOH), Si-free ka-
toite [Ca5Al2(OH)12], and intermediate hydrogrossular
[Ca3Al2Si2O8(OH)4] phases, and two non-reacted crystal-
line products, quartz (SiO2) and corundum (α-Al2O3).
The quartz originated from the Class F fly ash, while
corundum came from CAC. The boehmite and Si-free
katoite phases were categorized as the hydration reaction
products of CAC. Since the dissolution of sodium meta-
silicate activator in water generated two major hydrolys-
ate reactants, sodium hydroxide, NaOH, and metasilicic
acid, H2SiO3 (and its sodium salt), the hydroxysodalite in
the family of zeolite, was formed by the hydrothermal
reactions between sodium hydroxide and mullite
(3Al2O3·2SiO2) in Class F fly ash. On the other hand, the
hydrothermal reactions between CAC and quartz in Class
F fly ash led to the formation of hydrogrossular phases.
Both hydrogrossular and Si-free katoite phases are in the
hydrogarnet family. After modifying this cement with 2
wt% AP, the XRD patterns (top) closely resembled that
of AP-free one, suggesting that the AP had no significant
effect upon the final phase composition assembled in the
200˚C-autoclaved foamed cement. Thus, relating this
finding to the result of compressive strength at 200˚C,
the improvement of this strength engendered by adding
AP depended on the content of AP, but was independent
of the phase composition.
Figure 7 illustrates the XRD tracings of 300˚C-auto-
claved foamed cements with and without AP. Without
AP, there were two major differences in this XRD pat-
tern (bottom) from that of the cement made at 200˚C.
First, the intensity of d-spacing lines related to quartz
considerably decayed, implying that most of it in Class F
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10 20 30 40 50
Intensity (arb. Unit)
S: Hydroxysodalite
B: Boehmite
Z: Na-P type zeolite
K: Katoite
Q: Quartz
G: Hydrograssular
C: Corundum
Figure 7. XRD patterns for 300˚C-autclaved foamed ce-
ments containing 0 (bottom) and 2% AP (top).
fly ash hydrothermally reacted with the other cement-
forming reactants. Second, intensive d-spacing lines were
observed from hydroxysodalite, katoite, and hydrogros-
sular hydration products, emphasizing that these products
were well crystallized at 300˚C and became the major
crystalline phases in conjunction with the corundum.
We interpreted these results as follows: Since the ka-
toite phase had no silicate, seemingly the hydration reac-
tion of CAC at 300˚C created a well-formed katoite. Also,
this high temperature promoted the extent of hydrother-
mal reactions between sodium hydroxide and mullite to
form a well-crystallized hydroxysodalite. The principal
reason for the depletion of quartz was the formation of
more hydrogrossular because of an extensive reaction of
CAC with quartz in Class F fly ash.
As described in compressive strength testing, the
300˚C-autoclaved cements had a higher compressive
strength than those autoclaved at 200˚C. Thus, good for-
mation at 300˚C of these three crystalline hydrate phases,
hydroxysodalite, katoite, and hydrograssular, seems to be
responsible for improving compressive strength.
Of particular interest was the XRD pattern (top) of 2%
AP-modified cement, compared with that of unmodified
one. The pattern of the former was characterized by hav-
ing three distinctive features; 1) the incorporation of a
new crystalline hydrate phase attributed to Na-P type
zeolite into the cement, 2) the presence of intensive d-
spacing related to non-reacted quartz, and 3) the expres-
sion of a relatively weak line intensity of hydroxyso-
dalite-, katoite-, and hydrograssular-associated d-spacing.
The last two results suggested that AP not only inhibited
the reactions of quartz in Class F fly ash with CAC to
form hydrogrossular, but also restrained the extent of
crystallization of hydroxysodalite and katoite, while an
additional zeolite phase belonging to Na-P type was in-
troduced in 300˚C-autoclaved cement. Similar to the re-
lationship between AP content and compressive strength
at 200˚C, the improvement of compressive strength at
300˚C depended primarily on the AP content, but not on
the crystalline phases and their composition.
Our attention next shifted to exploring the microstruc-
ture developed in the foamed cement with AP. Figure 8
shows the SEM images for the fractured surfaces of
foamed cements without and with 2 wt% AP after auto-
claving at 200˚C. Without AP, the image of the foamed
cement made by incorporating numerous air bubbles into
the slurry revealed a typical honeycomb-like porous
structure encompassing craters of ~300 to ~ 50 m. Most
of these craters had a partially defective structure, which
can be interpreted as the formation of continuous voids.
After modifying the cement with 2 wt% AP, the image
was characterized by much smaller-size creators. Addi-
tionally, the defects in individual creators were minimal,
implying that AP aided in forming defect-free small cra-
ters as discrete voids. A similar microstructure was ap-
parent from the fractured surfaces of 300˚C-autoclaved
foamed cements with and without AP (not shown).
Unlike the continuous voids-containing porous struc-
ture in the foamed cement, we can assume that the struc-
ture assembled by discrete voids would reduce the rates
of permeability and transportation of corrosive electro-
lytes through the cement, assuring its ability to protect
carbon steel (CS) against corrosion.
3.6. Corrosion Mitigation
To verify the potential of AP to mitigate the corrosion of
CS, we conducted two electrochemical corrosion studies.
The first study centered on investigating the resistance
of AP to the infiltration and transportation of a corrosive
electrolyte through the foamed cement layer covering the
underlying CS surface; and, the second one was to sur-
vey the effectiveness of AP in preventing the corrosion
of CS at its interface with the foamed cement. The latter
Figure 8. SEM images for fractured surfaces of foamed
cements without AP (top) and 2 wt% AP-modified foamed
cement (bottom).
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study also involved assessing the extent of coverage of
CS surface by cement and the degree of cathodic corro-
sion protection along with the corrosion rates of CS. An-
other important factor in such protection was the adher-
ence of coating to CS [12].
3.6.1. EIS Test
One major parameter governing the mitigation of corro-
sion by the cements is their conductivity of corrosive
electrolytes; the extent of uptake of electrolytes by the
cements plays a pivotal role in inhibiting or accelerating
the corrosion of the underlying CS casing. We deter-
mined the extent of conductivity and transportation of
ionic electrolytes through the cement layer to the under-
lying CS surfaces using EIS. The samples for EIS testing
were prepared in the following manner: First, alkaline-
cleaned CS coupons were dipped in a soaking bath con-
taining AP-modified or unmodified foamed cement slur-
ry at room temperature; second, after withdrawing them,
the slurry-covered coupons were left for 24 hours at
room temperature, allowing the slurry layer to convert
into a solid layer; third, the cured cement layer-coated
CS was autoclaved for 24 hours at 200˚C and 300˚C; and,
finally, the autoclaved cement-covered CS coupon was
left at room temperature to cool before EIS testing. Af-
terward, the coated CS coupon was mounted in a holder,
and then inserted into a flat electrochemical cell. The
coated coupon with a surface area of 13 cm2 was exposed
to aerated 1.0 M sodium chloride electrolyte at 25˚C for
10 min before the EIS test.
Figure 9 compares the Bode-plot features [the abso-
lute value of impedance |Z| (ohm-cm2) vs. frequency
(Hz)] of the coupons coated with the non-foamed and
foamed cements without AP. Particular attention in the
overall EIS curves was paid to the pore resistance, Rp,
which can be determined from the peak in the Bode-plot
Frequency, H
IZI, ohm-cm
0% FA
1% FA
Figure 9. AC electrochemical impedance curves for non-
foamed and foamed cements without AP after autoclaving
at 200˚C.
occurring at a low frequency between 101 and 102 Hz.
For the non-foamed cement coating, the Rp value was
76.5 ohm-cm2. This value reduced by nearly 53% to a
35.5 ohm-cm2, when this coating was foamed. Since the
Rp value reflects the extent of ionic conductivity gener-
ated by the NaCl electrolyte passing through the coating
layer, this reduction represented an increase in the uptake
of electrolytes by the coating. In other words, the foamed
coating displayed poorer resistance to the infiltration and
transportation of electrolyte through its layer than did the
non-foamed coating.
Figure 10 depicts the changes in pore resistance, Rp,
of 200˚C- and 300˚C-autoclaved foam cements as a func-
tion of AP content. For the foamed cements without AP,
the value of Rp at 200˚C rose with an increasing auto-
clave temperature to 300˚C, from 35.3 to 76.5 ohm-cm2.
This fact strongly demonstrated that the coating’s effi-
cacy as corrosion-preventing barrier layer formed in an
autoclave at 200˚C was improved when autoclaved at
300˚C. As discussed under the compressive strength
testing, the bulk density of 300˚C-autoclaved cement was
higher than that of 200˚C-cement. Thus, we assumed that
a densified structure of cement at 300˚C was one of the
factors affecting the abatement of the infiltration and
transportation rates of corrosive electrolytes through it.
Incorporating AP into the 200˚C- and 300˚C-auto-
claved foamed cement coatings increased the Rp value;
this increase depended on the amount of AP. For instance,
at 200˚C, the Rp value of AP-free coating rose by 14% to
40.1 ohm-cm2 after adding 0.5 wt% AP. Adding 2 wt%
AP markedly enhanced its value to 60.5 ohm-cm2, cor-
responding to ~70% higher resistance than that of AP-
free coating. In other words, AP additive decreased the
uptake of corrosive electrolytes by coating layer and
suppressed the infiltration and transportation of electro-
lytes through it. Relating these findings to the results
from our study of microstructure development, there
AP content, wt%
Pore resistance, Rp, ohm-cm
After 200°C autoclave
After 300°C autoclave
Figure 10. Changes in pore resistance of 200˚C- and 300˚C-
autoclaved foamed cements as a function of AP content.
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were two possible mechanisms for decreasing the rate of
electrolyte infiltration through the foamed cement by AP:
One was creation of porous structure composed of de-
fect-free discrete voids; the other was related to the in-
situ formation of an electrolyte-impervious AP film in
the foamed cement. Such contributions by AP played an
essential role in improving the corrosion mitigation of
CS by the foamed cement.
For the 300˚C-autoclaved coatings, the Rp-AP content
relation was similar to that of the coatings at 200˚C;
namely, Rp value rose with an increasing AP content.
However, the rate of this increase brought about by add-
ing AP was far greater than at 200˚C. A 397.3 ohm-cm2
detected with 2 wt% AP was 5.2-fold higher than that of
the AP-free one at the same autoclave temperature.
Compared with this, at 200˚C, the resistance increase by
2 wt% AP was only 1.7-fold. This finding not only high-
lighted an excellent hydrothermal stability of AP at tem-
peratures up to 300˚C, but also demonstrated that the
hydrothermal temperature of 300˚C aided AP in display-
ing a better performance in a pronounced reduction of the
electrolyte’s uptake by coating, compared with that at
200˚C. As is described in FT-IR study, the transforma-
tion of the acrylic acid and alkyl ester within AP into a
hydrothermally stable M-complexed carboxylate struc-
ture was accelerated with an increasing hydrothermal
temperature, underscoring that incorporating more AP
would result in the formation and distribution of a large
number of complexed carboxylate structures in the
foamed cement. Thus, there were two possible reasons
for an enhanced pore resistance of 300˚C-autoclaved
foamed cement: One was the creation of a dense cement
structure; the other was that such a complexed molecular
structure served as barrier layers controlling the trans-
portation rate of electrolytes through the foamed cement.
3.6.2. DC Potenti od yn ami c Pol arization Test
For this test, we prepared the samples in the same man-
ners as those employed in the EIS test. Figure 11 illus-
trates the typical cathodic-anodic polarization curves
where the potential voltage, E, is plotted versus current
density, A/cm2, for CS coupons covered with the 200˚C-
autoclaved foamed cements unmodified and modified
with 2 wt% AP. The shape of the curve reveals the tran-
sition from cathodic polarization region at the onset of
the most negative potential to the anodic polarization
region at the end of the test at a less negative potential.
The potential at the transition point from cathodic to an-
odic is normalized as the corrosion potential, Ecorr. Com-
pared with the features of the curve from AP-free foamed
cement denoted as 0% AP, there were two noticeable
differences for the CS coated with 2 wt% AP-modified
cements: One was a shift in the Ecorr value to a less nega-
tive potential; the other was the reduction of cathodic
0 % AP
2 % AP
-0.70-0.80-0.60 -0.50-0.40 -0.30 -0.20 -0.10
Potential volts vs. SCE
Current Density, A/cm
Figure 11. DC electrochemical potentiodynamic cathodic-
anodic polarization diagrams for foamed cements unmodi-
fied and modified with 2% AP after autoclaving at 200˚C.
current density (A/cm2) between 0.6 and 0.7 V. The
first difference reflects the extent of cement coverage of
the CS surface; namely, a good coverage by a continuous
void-free coating layer at the contact zones with CS sur-
face is responsible for moving the Ecorr value to a less
negative potential. This shift of Ecorr underscored that
adding 2% AP increased the extent of coverage by the
foamed cements. For the second difference, the decline
in the cathodic current density signified that the cathodic
reaction at the corrosion site of CS was restrained, par-
ticularly the oxygen reduction reaction, 2H2O + O2 + 4e
4OH. Correspondingly, such a reaction leading to the
cathodic corrosion of CS appeared to be inhibited by
depositing the AP-modified foamed cement on the CS
surfaces, thereby highlighting the ability of AP to abate
the corrosion of CS. Thus, AP conferred two advanced
corrosion-mitigating properties on the foamed cements
for providing a better protection of CS against brine-
caused corrosion: First, it afforded a better coverage of
cement over the steel surfaces; second was its efficacy in
upgrading the ability of cement to inhibit the cathodic
corrosion reaction of CS.
Based upon the potentiodynamic polarization curve
(Figure 12), we determined the absolute corrosion rates
of CS, expressed in the conventional engineering units of
milli-inches per year (mpy). The Equation (1) proposed
by Stern and Geary [13] was used in the first step:
corra cacR
 
 (1)
In the equation the Icorr is the corrosion current density
in A/cm2; βa and βc in V/decade of current refer to the
anodic and cathodic Tafel slops, respectively, obtained
from the log I vs. E plots encompassing both anodic and
cathodic regions; and, PR is the polarization resistance
determined from the corrosion potential, Ecorr. With Icorr
computed through Equation (1), the corrosion rate (mpy)
can be obtained from the following expression:
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Corrosion rate = 0.13 Icorr (EW)/d; where EW is the
equivalent weight of the corroding CS in g; and d is the
density of the corroding CS in g/cm3. Table 3 gives the
Icorr and corrosion rate calculated for CS panels coated
with unmodified and modified with AP cement after au-
toclaving at 200 and 300˚C for 24 hours. For the
200˚C-autoclaved test panels, the corrosion rate of the
CS for AP-free coating was 175.7 mpy with accompany-
ing Icorr of 38.5 × 105 A/cm2. These values conspicu-
ously dropped to 68.9 mpy and 15.07 × 105 A/cm2 when
the cement was modified with 2% AP, clearly validating
the effectiveness of the AP in mitigating the CS’s corro-
sion and in resisting the cathodic corrosion reactions.
displayed a better performance in inhibiting the carthodic
reaction and mitigating the CS’s corrosion. A further
reduction of both the corrosion rate and Icorr value was
realized by adding AP to it; in fact, a 2% AP provided a
very low corrosion rate and Icorr value of 6.8 mpy and
1.49 × 105 A/cm2, respectively, reflecting 6.3-fold low-
ering over those of AP-free cement. This finding also
verified that AP not only withstood exposure to a hydro-
thermal environment at 300˚C, but also improved CS
protection against brine-caused corrosion by 300˚C-
autoclaved foamed cement.
Hence, AP had a potential as a high-temperature cor-
rosion-inhibiting additive responsible for mitigating the
CS’s corrosion at temperatures up to 300˚C.
A very interesting data on corrosion mitigation was
obtained from the 300˚C-autoclaved test panels; namely,
the CS’s corrosion rate of 43.1 mpy and Icorr value of
9.42 × 105 A/cm2 by AP-free foamed cement was con-
siderably lower than those of the 200˚C-autoclaved AP-
free one. This corrosion rate represented nearly a 4-fold
reduction compared with that of 200˚C; correspondingly,
the similar reduction was observed in the Icorr value,
suggesting that the foamed cement autoclaved at 300˚C
3.6.3. Adherence of Foamed Cement to CS
Since a better interfacial bond between coating and CS
was another factor governing the alleviation of the CS’s
corrosion, we evaluated the extent of adherence of the
200˚C and 300˚C-autoclaved foamed cements modified
with 0, 0.5, 1.0, and 2.0 wt% AP to CS surfaces by
EDX-elemental mapping and -oxide quantitative analy-
The samples were prepared in the following sequences:
First, the CS coupons’ surfaces were coated with
AP-modified and unmodified foamed cements; second,
the coated CS coupons were autoclaved at 200˚C or
300˚C for 24 hours; third, the autoclaved coupons were
placed on the center-loading bending apparatus to gener-
ate the tensile shear-bonding failure at the side of cement
coating of CS, leading to the delamination of cement
layer from underlying CS surface; and fourth, the CS
side separated from the cement layer was used for the
EDX mapping and the quantitative analysis of metal
oxides. Assuming the cement adhered well to CS sur-
faces and the failure of interfacial bonding occurs in ce-
ment layers, our attention centered on obtaining the fol-
lowing information: One was to visualize the distribution
of cement-related Ca, Al, and Si elements remained
Figure 12. Typical Tafel plot from a polarization experi-
Table 3. Tafel analyses of potentiodynamic polarization curves for steel panels covered with AP-modified and unmodified
foamed cements.
Temperature, ˚C AP content, wt% Ecorr (I = 0), (V)
a, (V/decade)
c, (V/decade) Icorr, (A/cm2) Corrosion rate, (mpy)*
200 0 0.4551 0.2556 0.3814 38.50 × 105 175.7
200 2.0 0.4099 0.1623 0.6835 15.07 × 105 68.9
300 0 0.3226 0.1562 0.2756 9.42 × 105 43.1
300 0.5 0.4668 0.0955 0.3224 7.03 × 105 32.1
300 1.0 0.3128 0.2046 0.4188 4.29 × 105 19.6
300 2.0 0.3414 0.1700 0.3136 1.49 × 105 6.8
mpy: milli-inches per year.
over the separated CS surfaces; the other was the changes
in the content of metal oxides, in particular, Fe2O3, CaO,
Al2O3, and SiO2, as a function of AP content. Thus, if we
observed a widespread distribution of these elements and
a great deal of these oxides, a possible interpretation was
that the cement was well adhered to CS surfaces.
Also, it should be noted that since the penetration
depth of X-ray from µEDX is ~2 µm, all data related to
elemental mapping and oxide compositions came from a
top ~2 µm thick surface layer.
Figure 13 represents the individual µEDX maps of Fe
and Ca elements for CS surface separated from the
200˚C- and 300˚C-autoclaved AP-free foamed cement
layers. The µEDX analyses were conducted on the map-
ping areas of 4.0 × 3.0 mm (12 mm2) and targeted for
detecting the two specific elements, Fe belonging to un-
derlying CS and Ca attributed to cement. During the
mapping operation, X-ray intensities of 6.4 and 3.7 keV
lines for Fe Kα and Ca Kα respectively were recorded
and used for creating color-coded maps of their relative
distribution. The maximum amount, in percent, of each
tested element is given above the color-barcode. Further,
the concentration of these elements is ranked by the dif-
ference in colors; namely, the area in white color coding
indicates the highest concentration of these elements
present on the CS surfaces, the dark red color is their
secondary concentrated areas, while no presence or neg-
ligible amount of these elements is signified by the areas
in the darkest blue color. Correspondingly, the greenish
and yellowish colors reflect a moderate presence of these
elements. For the 200˚C-autoclaved foamed cement, the
elemental mapping of Fe revealed mainly two different
regions on the CS surface: One was the top- and bot-
tom-central regions representing the distributions of the
4.0 x 3.0mm
0.80mm FeKa
[% ]
4.0 x 3.0mm
0.80mm CaKa
4.0 x 3.0mm
0.80mm FeKa
4.0 x 3.0mm
0.80mm CaKa
Figure 13. EDX elemental mapping of Fe (top) and Ca
(bottom) for interfacial CS sides separated from AP-free
foamed cements after autoclaving at 200˚C (left) and 300˚C
red signal as major color code, and some green signal as
well as blue signal as minor one; the other was the rest of
this detected area, which mostly was white. The last re-
sult meant that this region was dominated by Fe element
originating from the underlying CS, implying no or poor
converge of CS surface by any other elements. In con-
trast, the former region was covered by some different
elements. Hence, although the amount of Ca was only
22.4% at maximum on the studied area, Ca mapping ex-
hibited some presence of this element over the top- and
bottom-central regions, signifying that some cement re-
mained locally on the CS surface separated from the ce-
ment layer. This Ca-presence increased to 30.4%, when
the foamed cement was autoclaved at 300˚C. In fact, as
seen in Fe map, a large area over the CS surface was
covered by other elements coexisting with Fe, compared
with that of 200˚C-autoclaved foamed cement/CS inter-
face sample. This finding demonstrated that the climbing
the temperature to 300˚C from 200˚C improved the ad-
herence of the foamed cement to CS.
Figure 14 gives the
EDX mapping results of four
elements, Fe, Ca, Al, and Si for the interfacial CS dis-
lodged from the 1.0 wt.% AP-modified foamed cement
layer after autoclaving the cement/CS sample at 300˚C.
X-ray intensity counts of Al and Si were taken from Al
Kα and Si Kα at 1.5 and 1.7 keV, respectively.
When the configuration of Fe mapping was compared
with that of Ca over the CS surfaces, the distribution
pattern of these two elements were very similar; thus, the
highly concentrated region of Fe marked as the white-
color code was directly opposite to the lowest concentra-
tion region of Ca (dark blue). In contrast, although a high
amount of 93.8% and 72.3% was detected for Al and Si,
respectively, the configuration of their mapping images
was quite distinct from those of Fe and Ca.
4.0 x 3.0mm
0.80mm FeKa0.000
4.0 x 3.0mm
0.80mm CaKa
[% ]
4.0 x 3.0mm
0.80mm AlKa0.000
4.0 x 3.0mm
0.80mm SiKa
Figure 14. Comparison of Fe, Ca, Al, and Si elemental
mapping images for interfacial CS surfaces dislodged from
the 1.0 wt% AP-modified foamed cement after autoclaving
at 300˚C.
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A possible interpretation of this finding was that Ca
preferentially reacted with CS surfaces, rather than Al
and Si. Similar results were observed from all tested pa-
nels with and without AP after autoclaving at 200˚C and
Next, we investigated the effectiveness of AP in im-
proving the adherence of foamed cement to the CS at
200˚C and 300˚C. Using
EDX, our approach to obtain-
ing this information was made by tracing the contents of
oxide compounds such as Fe2O3, CaO, Al2O3, and SiO2,
as a function of AP content, for the interfacial CS sur-
Thus, a high content of three cement-related oxides,
CaO, Al2O3, and SiO2, on CS surface, we judged as good
cement coverage of the CS, and an excellent adherence
of cement to CS.
Figure 15 reveals the concentrations of oxides on CS
for 200˚C-autoclaved foamed cements containing 0, 0.5,
1.0, 1.5, and 2.0 wt% AP. Without AP, the composition
of oxides was 92.8% Fe2O3, 1.3% CaO, 0.0% Al2O3, and,
7.1% SiO2, revealing that some Ca and Si oxides related
to cement were deposited on CS. Adding 0.5% AP to
cement increased the content of CaO and SiO2, to 2.1%
and 8.8%, respectively. A further increase in Ca and Si
oxides to 3.6% and 10.9% was detected for 2.0% AP,
while the CS-related F2O3 content declined with an in-
creasing AP content. Of particular interest in the deposi-
tion of cement-related oxides on CS was Al2O3; its depo-
sition began at 1.0 wt.% AP, beyond that, at 2.0 wt.% AP
the deposition was 13.6% Al2O3 , which was the highest
content, among these cement-related oxides. Thus, the
incorporation of more AP into foamed cement not only
assured the extensive coverage of cement over the CS, so
representing the improved adherence of cement to CS,
but also led to the formation of Al2O3-rich cement layer
adhering to the CS.
AP content, %
0.0 0.5 1.01.5 2.0
Oxide content, %
Figure 15. Changes in the content of oxides present at in-
terfacial CS surface as a function of AP content for 200˚C-
autoclaved cement/CS samples.
For the 300˚C-autoclaved cement/CS samples (Figure
16), the oxide composition of interfacial CS surfaces was
80.6% Fe2O3, 2.5% CaO, 0.0% Al2O3, and, 16.9% SiO2.
Like, in the 200˚C-autoclaved sample without AP, no
Al2O3 was detected on the CS surface. However, the
concentration of Fe2O3 was ~13% lower than that of
200˚C-autoclaved sample.
Correspondingly, the concentrations of CaO and SiO2
rose by nearly 2-fold the former oxide and 2.4-fold for
the latter one, underscoring that the extent of the adher-
ence of foamed cement to CS increased with the raising
hydrothermal temperature to 300˚C. When the foamed
cement was modified with 0.5% AP, the cement layer
adhering to CS had a relatively high Al2O3 content of
20.5% coexisting with 3.5% CaO and 18.2% SiO2. The
content of these oxides gradually rose with an increasing
AP content, contrarily, the Fe2O3 content declined. With
2.0 wt% AP, a 44.9% Fe2O3 detected corresponded to
lowering of ~22% from that of 0.5% AP, while the in-
crease of ~47%, ~10%, and ~37% was observed for all
cement-related oxides, Al2O3, CaO, and SiO2, respec-
tively, implying that although the temperature was ele-
vated to 300˚C, the AP was as effective in improving the
adherence of foamed cement as was the case at 200˚C.
Hence, the M-complexed AP compounds formed in
300˚C-autoclaved foamed cement withstood the hydro-
thermal temperature of 300˚C; they played an essential
role in enhancing bonding of foamed cement to CS, the-
reby resulting in the cohesive failure mode wherein in-
terfacial bonding failure took place in the cement layer
near the interface regions between CS and cement.
4. Conclusion
Air bubble-foamed cement (slurry density of 1.3 g/cm3)
consisting of refractory calcium aluminate cement (CAC),
Class F fly ash, sodium silicate activator, and cocamido-
AP content, %
0.0 0.51.0 1.52.0
Oxide content, %
Figure 16. Changes in the content of oxides present at in-
terfacial CS surface as a function of AP content for 300˚C-
autoclaved cement/CS samples.
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propyl dimethylamine oxide-based foaming agent, was
modified with acrylic polymer (AP) employed as a high-
temperature cathodic corrosion inhibitor of carbon steel
(CS) after exposure to the hydrothermal environment at
200˚C or 300˚C. Under these conditions, the functional
acrylic acid and alkyl ester groups within AP reacted
with the metal cations (M), such as Ca2+, Al3+, and Na+,
liberated from sodium silicate-activated CAC and Class
F fly ash, leading to the formation of M-complexed car-
boxylate group-containing AP in the cement body at
85˚C. Such in-situ transformation of these functional
groups into complexed groups progressively occurred as
the hydrothermal temperature rose. This transformation
improved the thermal stability of AP. Correspondingly,
the complexed carboxylate-rich AP withstood a hydro-
thermal temperature at 300˚C, ensuring that AP as the
corrosion inhibitor was capable of protecting the CS
against corrosion at this temperature. Addition of AP
delayed the onset of cement set while increasing the in-
tegrated heat released during cement hydration in calo-
rimetric experiments at 85˚C.
At 200˚C, AP did not cause any significant changes of
crystalline hydrate composition assembled in the AP-free
foamed cement; namely, its composition comprised four
hydrothermal hydration reaction products, hydroxyso-
dalite [Na4Al3Si3O12(OH)], intermediate hydrogrossular
[Ca3Al2Si2O8(OH)4], boehmite (γ-AlOOH), and Si-free
katoite [Ca5Al2(OH)12] phases that were responsible for
strengthening the 200˚C-autoclaved foamed cement. The
hydroxysodalite phase was formed by hydrothermal in-
teractions between the sodium silicate activator and mul-
lite phase in Class F fly ash, while the quartz in Class F
fly ash reacted with CAC to form hydrogrossular. On the
other hand, hydration of CAC engendered two other
phases, boehmite and Si-free katoite. Although, similar
crystalline phases were formed in cement with and with-
out AP, the compressive strength rose with an increasing
AP content. At the hydrothermal temperature of 300˚C,
the crystalline phases and their quantities differed from
those observed at 200˚C. For AP-free cement, three
phases, hydroxysodalite, katoite, and hydrogrossular,
were well formed and crystallized; in particular, a sub-
stantial amount of quartz in Class F fly ash hydrother-
mally reacted with CAC to form more hydrogrossular.
Consequently, these well-formed crystalline compounds
aided in improving further the cement’s compressive
strength, compared with that at 200˚C. Adding AP re-
strained the formation of these three crystalline hydrate
phases, while Na-P type zeolite was formed as additional
crystalline phase. Like the findings at 200˚C, the devel-
opment of compressive strength depended on AP content;
namely, strength rose with an increasing AP content.
The microstructure developed in the autoclaved
foamed cements was characterized by a honeycomb-like
porous structure constituted of numerous defected mi-
cro-size craters. Adding AP conferred two beneficial
alterations in this microstructure: First, the defected cra-
ters were transformed to defect-free discrete voids; and,
second, the size of craters became much smaller. Thus,
creating defect-free, small craters was one reason why
the AP-modified foamed cements developed a good
compressive strength. The other benefit from the pres-
ence of such advanced microstructure was the reduction
of infiltration and transportation of corrosive electrolytes
through the foamed cement layer, thereby mitigation of
the corrosion of underlying CS.
We believe that the AP addition to the foamed cement
significantly reduced corrosion rate of CS at the hydro-
thermal temperature of 300˚C because of the following
two key factors: the formation of 300˚C-withstanding
barrier layers constituted of complexed carboxylate-rich
AP and the improved adherence of the cement to CS
surfaces. For the latter, incorporating more AP yielded a
better cement adherence. Additionally, among the Ca-,
Si- and Al-oxides in hydraulic cement, Ca oxide prefer-
entially coupled to the Fe2O3 layer arrayed at the surface
of CS at 200˚C. The coverage of CS surface by Ca oxide
extended with an increasing temperature, resulting in
better adherence of 300˚C-autoclaved cement to CS,
compared with that of 200˚C-autoclaved one. The fol-
lowing three important factors governed the mitigation of
CS’s corrosion: 1) Minimized conductivity of corrosive
ionic electrolytes through the foamed cement layer; 2)
inhibited cathodic reactions at corrosion site of CS; and 3)
increased coverage of CS surface by a foamed-cement
layer at the interfacial boundary regions between the ce-
ment and CS.
For AP-free foamed cements, the corrosion rate, 175
milli-inch per year (mpy), of CS after autoclaving at
200˚C, reduced by 4 times when the autoclaving tem-
perature increased to 300˚C. For AP-modified foamed
cements, the corrosion rate of CS coated with AP-free
cement at 200˚C fell 2.6 times with 2 wt% AP. At 300˚C,
2 wt% AP lowered conspicuously the CS’s corrosion rate
to only 6.8 mpy from 43 mpy for AP-free one.
[1] S. Gill, T. Pyatina and T. Sugama, “Thermal Shock-Re-
sistant Cement,” Geothermal Resources Council Trans-
action, Vol. 36, 2012, pp. 445-451.
[2] T. Sugama L. E. Brothers and T. R. Van de Putte,
“Air-Foamed Calcium Aluminate Phosphate Cement for
Geothermal Wells,” Cement and Concrete Composite,
Vol. 27, No. 7-8, 2005, pp. 758-768.
[3] K. Y. Ann, H. S. Jung, H. S. Kim, S. S. Kim and H. Y.
Moon, “Effect of Calcium Nitrite-Based Corrosion In-
hibitor in Preventing Corrosion of Embedded Steel in
Open Access ENG
Open Access ENG
Concrete,” Cement and Concrete Research, Vol. 36, No.
3, 2006, pp. 530-535.
[4] M. Saremi and E. Mahallati, “A Study on Chloride-In-
duced Depassivation of Mild Steel in Simulated Concrete
Pore Solution,” Cement and Concrete Research, Vol. 32,
No. 12, 2002, pp. 1915-1921.
[5] P. Ghods, O. B. Isgor, G. A. McRae and G. P. Gu, “Elec-
trochemical Investigation of Chloride-Induced Depas-
sivation of Black Steel Rebar under Simulated Service
Conditions,” Corrosion Science, Vol. 52, 2010, pp. 1649-
[6] Y. M. Tang, Y. F. Miao, Y. Zuo, G. D. Zhang and C. L.
Wang, “Corrosion Behavior of Steel in Simulated Con-
crete Pore Solutions Treated with Calcium Silicate Hy-
drates,” Construction and Building Materials, Vol. 30,
2012, pp. 252-256.
[7] A. R. Boga and I. B. Topcu, “Influence of Fly Ash on
Corrosion Resistance and Chloride Ion Permeability of
Concrete,” Construction and Building Materials, Vol. 31,
2012, pp. 258-264.
[8] J. Hu, D. A. Koleva and K. van Breugel, “Corrosion Per-
formance of Reinforced Mortar in the Presence of Poly-
meric Nano-Aggregates: Electrochemical Behavior, Sur-
face Analysis, and Properties of the Steel/Cement Past
Interface,” Journal of Material Science, Vol. 47, 2012, pp.
[9] S. X. Wang, W. W. Lin, S. A. Ceng and J. Q. Zhang,
“Corrosion Inhibition of Reinforcing Steel by Using
Acrylic Latex,” Cement and Concrete Research, Vol. 28,
1998, pp. 649-653.
[10] R. Selvaraj, M. Selvaraj and S. V. K. Iyer, “Studies on the
Evaluation of the Performance of Organic Coatings Used
for the Prevention of Corrosion of Steel Rebars in Con-
crete Structure,” Progress in Organic Coatings, Vol. 64,
No. 4, 2009, pp. 454-459.
[11] T. Sugama, “Hydrothermally Self-Advancing Hybrid
Coatings for Mitigating Corrosion of Carbon Steel,”
Brookhaven National Laboratory, 2006, BNL-77335.
[12] P. R. Sere, A. R. Armas, C. I. Elsner and A. R. Di Sarli,
“The Surface Condition Effect on Adhesion and Corro-
sion Resistance of Carbon Steel/Chlorinated Rubber/Ar-
tificial Sea Water Systems,” Corrosion Science, Vol. 38,
1996, pp. 853-866.
[13] M. Stern and A. L. Geary, “Electrochemical Polarization I.
A Theoretical Analysis of the Shape of Polarization
Curves,” Journal of Electrochemical Society, Vol. 104,
No. 1, 1975, pp. 56-62.